In this study, the electrochemical properties of synthesized intermetallic compounds (IMCs) similar to those commonly found in AA2024-T3 were studied in neutral 0.1 M NaCl at 10, 30, 50, and 70°C using the electrochemical microcell method. Results on the synthesized IMCs were combined with and supported by analysis of free corrosion experiments performed on AA2024-T3. Results reveal that, in general, corrosion rates of the IMCs increase with temperature and pitting potentials of most IMCs show a slight decrease with temperature. Increased dealloying kinetics of S-phase with increasing temperature was evident in both open-circuit potential transients and post-exposure corrosion morphology characteristics, leading to increased Cu surface areas and a higher likelihood of the propagation of localized corrosion into the matrix. The corrosion potential of synthesized Al-Cu-Mn-Fe type particles was found to be cathodic to the matrix at low temperatures, but anodic to the matrix at high temperatures. This high-temperature behavior led to dissolution of Al-Cu-Mn-Fe type particles, not commonly found under ambient conditions. After dealloying, these particles became cathodic to the matrix, leading to trenching similar to that found under low-temperature conditions. The limiting current density on most synthesized IMCs, as a result of the oxygen reduction reaction, was maximized at around 50°C, consistent with calculations using a simplified Cottrell equation and based on the temperature dependence of oxygen solubility and oxygen diffusion.

INTRODUCTION

Copper and magnesium are added to aluminum to improve its mechanical properties such as strength and toughness through precipitation strengthening and thermochemical processing.1-2  However, these alloying additions react with other alloying elements (e.g., Mn) and impurity elements (e.g., Fe) to form constituent intermetallic compounds (IMCs) such as Al2Cu, Al2CuMg (S-phase), and a range of Al-Cu-Mn-Fe type particles during solidification processing.1,3  These constituent particles are large, 0.1 μm and greater in diameter. They form as a result of dendritic solidification and are distinct from precipitation particles and dispersoids that form as a result of heat treatment. Constituent IMC particles often possess different corrosion potentials from the surrounding matrix phase, leading to localized galvanic effects and localized corrosion.4-12  IMC particles whose corrosion potential is higher than that of the matrix remain relatively noble.13-18  These particles catalyze oxygen reduction in aerated aqueous environments and lead to the dissolution of the matrix immediately adjacent to the particle, generating what is often referred to as trenching corrosion.5,19-23  IMC particles such as S-phase, whose corrosion potential is initially lower than that of the matrix phase, are relatively active. These particles tend to be selectively dissolved and undergo corrosion of the IMC, leaving insoluble remnants with different electrochemical properties than the original particle at the particle location.8,24-28 

All aspects of localized corrosion in aluminum alloys are related to secondary phase IMCs, whose electrochemical properties are strongly dependent on environmental conditions. The corrosion potentials of most IMC particles decrease as [Cl] increases. As pH affects oxide film stability and the composition of a protective oxide film is related to the elements contained in a given IMC, the effect of pH on corrosion and the pitting potentials of IMC particles depend on the composition in a complex manner.5,19-20,29  To have a full understanding of localized corrosion, it is necessary to have a comprehensive electrochemical characterization of IMC particles.

The micro-capillary electrochemical cell (microcell) allows electrochemical measurements to be performed on a small area scale using a glass capillary tube drawn to a small opening whose diameter ranges from a few micrometers to a few millimeters.30-32  The microcell approach has been used to characterize the electrochemical properties of small, specially synthesized IMC crystals in aqueous solutions over a range of different pH and NaCl concentrations.5,20  This information can then be used to develop mechanistic interpretations of localized corrosion damage accumulation.

Exposure temperature is also an important environmental factor that affects localized aqueous corrosion. For example, in alkaline chloride solutions, the anodic and cathodic kinetics of Cu and Fe increase with increasing temperature under mixed reaction control.33  In acid solutions (0.1 M NaCl + 0.1 M HCl), corrosion rates of pure metals (Cu, Fe, Zn, and Ni) increase exponentially with temperature in accordance with the Arrhenius equation.33  For corrosion processes that involve oxygen reduction, temperature can exert a secondary influence through its effect on equilibrium oxygen solubility and the dissolved oxygen diffusion coefficient.34-35  Furthermore, the solubility of hydrated corrosion products is temperature dependent, which can affect the corrosion process.36 

The electrochemical characteristics of IMC particles exposed to dilute chloride solutions at temperatures other than room temperature have not been widely reported. In this work, the electrochemical properties of synthesized IMCs at temperatures ranging from 10°C to 70°C were studied using the microcell method, and the data gathered were used to interpret corrosion morphology of actual AA2024-T3 (UNS A92024(1)) samples after exposure in a NaCl solution at the temperature of interest.

EXPERIMENTAL PROCEDURES

Materials and Sample Preparation

A 0.1 M NaCl solution was prepared by dissolving reagent grade NaCl in deionized (DI) water (18 MΩ·cm). Synthesized IMCs of interest in the study included Al2CuMg and Al2Cu, Al7Cu2Fe, Al20Cu2Mn3, and Al-4%Cu. Al7Cu2Fe and Al20Cu2Mn3 are proxy phases used to represent a range of Al-Cu-Fe-Mn type particles, while Al-4%Cu is an analog for the matrix phase of AA2024-T3. These IMCs were specially synthesized and heat treated to achieve a grain diameter size of several hundred micrometers to facilitate electrochemical characterization by the microcell method.20  All samples were polished from 600 grit to 1200 grit using SiC paper lubricated by ethyl alcohol, polished using 1 μm diamond paste lubricated by light weight oil, degreased ultrasonically in ethanol, and dried using compressed air.

Electrochemical measurements were performed immediately after polishing. A specially designed Cu stage, which had a hole to accommodate the samples, was used to control the temperature of the samples, as shown in Figure 1. The Cu stage was precooled or preheated to the desired temperature with flowing water from a circulator, and the temperature was measured using a contact thermometer with an uncertainty of ±1°C. The samples were also cut thin (~5 mm) so that they reached the target temperature quickly and uniformly.

FIGURE 1.

Microcell setup with cross section of Cu heating stage.

FIGURE 1.

Microcell setup with cross section of Cu heating stage.

Potentiodynamic Polarization Measurements

Anodic and cathodic polarization curves on the synthesized IMCs were collected in 0.1 M NaCl solution using the microcell at 10, 30, 50, and 70°C. Details of the microcell setup have been reported previously and a schematic is presented in Figure 1.5,29  Measurements were made using a standard three-electrode setup, with a platinum wire as the counter electrode and saturated calomel electrode (SCE) as the reference electrode. The measurement was controlled and data were recorded using a Gamry Reference 600 potentiostat. The opening of the capillary was about 100 μm, and a layer of silicon rubber was used to dress the capillary tip to make a leak-free seal with the sample surface. Anodic and cathodic polarization curves were initiated, after an approximately 60 s open-circuit potential (OCP) measurement, at a scan rate of 0.01 V/s. At least six replicate polarization curves were collected, and the average values of the corrosion potential, pitting potential, corrosion rate, and limiting current density were used to characterize each IMC’s electrochemical response.

Corrosion Morphology Characterization

Free exposure corrosion tests were performed on actual AA2024-T3 samples in a circulating water-jacketed cell. The NaCl solution was precooled or preheated to the designed temperature, and then the polished samples were placed at the bottom of the cell. After a 1 h exposure, samples were taken out, rinsed with DI water, dried with compressed dry air, and then stored in a desiccator. The corrosion morphology was characterized using a scanning electron microscope (SEM) coupled with energy dispersive x-ray spectroscopy (EDS) enabling the identification of the IMCs. Some particles were cross sectioned for characterization of corrosion damage in the near-surface region using focused ion beam analysis with a 30 kV Ga ion beam.

RESULTS

Potentiodynamic Responses of Synthesized Intermetallic Compounds and Al-4%Cu as a Function of Temperature

The anodic polarization curves for Al-4%Cu, Al2Cu, Al7Cu2Fe, and Al20Cu2Mn3 presented here were chosen from the replicate curves based on the proximity of their corrosion potential or pitting potential to the average value of replicate measurements (Figure 2). Figures 3 and 4 show the corrosion potential and pitting potential distributions and the average values, respectively, for each of the tested IMCs and for the AA2024-T3 matrix. Measurements made on the AA2024-T3 matrix were achieved by placing the microcell probe over an area on AA2024-T3 at which no IMC particles were visible through the microcell’s optical microscope at a magnification of 50×. Corrosion rates were estimated by extrapolating the linear portion of the anodic polarization curves to the intersection of corrosion potential, and these data are summarized in Figure 5. Determined from the cathodic polarization curves at −1.00 VSCE (Figure 6), the limiting cathodic current density (Figure 7) was used to compare the ability of each IMC to support oxygen reduction and evaluate the effect of temperature on oxygen reduction kinetics.

FIGURE 2.

Anodic polarization curves of IMCs and Al-4%Cu in 0.1 M NaCl at 10, 30, 50, and 70°C.

FIGURE 2.

Anodic polarization curves of IMCs and Al-4%Cu in 0.1 M NaCl at 10, 30, 50, and 70°C.

FIGURE 3.

Corrosion potential of IMCs in 0.1 M NaCl at 10, 30, 50, and 70°C (the corrosion potential of S-phase was obtained from the open-circuit potential curves at 80 s).

FIGURE 3.

Corrosion potential of IMCs in 0.1 M NaCl at 10, 30, 50, and 70°C (the corrosion potential of S-phase was obtained from the open-circuit potential curves at 80 s).

FIGURE 4.

Pitting potential of IMCs in 0.1 M NaCl at 10, 30, 50, and 70°C.

FIGURE 4.

Pitting potential of IMCs in 0.1 M NaCl at 10, 30, 50, and 70°C.

FIGURE 5.

Corrosion rate (lg icorr) of IMCs in 0.1 M NaCl as function of 1/temperature.

FIGURE 5.

Corrosion rate (lg icorr) of IMCs in 0.1 M NaCl as function of 1/temperature.

FIGURE 6.

Cathodic polarization curves of IMCs and Al-4%Cu in 0.1 M NaCl at 10, 30, 50, and 70°C.

FIGURE 6.

Cathodic polarization curves of IMCs and Al-4%Cu in 0.1 M NaCl at 10, 30, 50, and 70°C.

FIGURE 7.

Cathodic current density as a function of temperature of each sample (−1 VSCE value of cathodic polarization scan in aerated 0.1 M NaCl).

FIGURE 7.

Cathodic current density as a function of temperature of each sample (−1 VSCE value of cathodic polarization scan in aerated 0.1 M NaCl).

Al-4%Cu

This solid solution was spontaneously passive at all temperatures tested, and it was not prone to corrosion. The corrosion potential varied in a range of −0.800 VSCE to −0.950 VSCE over the temperatures tested. Up to 70°C there was little variation in the pitting potential, which fell in the range of −0.350 VSCE to −0.450 VSCE. At 70°C, another breakdown potential was also observed at around −0.75 VSCE. The tendency of this phase to repassivate was strong and repassivation potentials were all greater than −0.70 VSCE (Figure 2).

The corrosion rate of Al-4%Cu did not exhibit strong temperature dependence, ranging from a few tenths of one μA/cm2 at 10°C to 3 μA/cm2 at 70°C. Cathodic kinetics increased with increasing temperature, as shown in Figure 6.

Al2Cu

The corrosion potential of this phase decreased from about −0.450 VSCE to −0.600 VSCE over the temperature range 10°C to 70°C (Figure 3), but remained higher than that of the AA2024-T3 matrix at all temperatures. The pitting potential decreased from about −0.300 VSCE to −0.550 VSCE over the temperature range 30°C to 70°C (Figure 4). It is notable that no obvious pitting potential was observed at 10°C, suggesting that Al2Cu cannot spontaneously passivate at 10°C. Although a small passive window (difference between corrosion potential and pitting potential) was observed at all temperatures, the smallest occurred at 70°C and was only about 80 mV. The corrosion rate of Al2Cu increased from about 2 μA/cm2 at 10°C to nearly 100 μA/cm2 at 70°C.

The Al2Cu phase did not show a strong tendency to repassivate under potentiodynamic polarization. In fact, under all temperatures, repassivation corrosion potentials on the reverse scan during anodic polarization were all more negative than corrosion potentials measured on forward scans (Figure 2[b]). Repassivation potentials did, however, increase with increasing temperature.

Cathodic kinetics increased modestly along the full cathodic polarization curve as temperatures increased from 10°C to 50°C and in the diffusion-controlled region of the curve over the entire temperature range (Figure 6). The mass transport limited current density for the oxygen reduction reaction (ORR) was in the range of 20 μA/cm2 to 40 μA/cm2 for all temperatures.

Al7Cu2Fe

Similar to the Al-4%Cu phase, passivity of the Al7Cu2Fe phase was robust at all temperatures as judged by the large difference in pitting potential and corrosion potential at each temperature (Figure 2). This phase showed a very strong tendency to repassivate and in many instances repassivated shortly after the scan direction was reversed in the anodic polarization curve. At all temperatures tested, the repassivation potential was well above the corrosion potential and in some experiments was observed to be above the observed pitting potential. At 50°C and 70°C, repassivation on the reverse polarization scans occurred near or above the pitting potential.

The corrosion potential was not strongly temperature dependent, varying from −0.79 VSCE to −0.70 VSCE (Figure 3), while the pitting potential dropped by about 50 mV as temperature increased from 10°C to 70°C (Figure 4). Generally, the passive current density increased with temperature, while the corrosion rate increased exponentially with temperature. The cathodic kinetics increased from 10°C to 30°C, and then slightly decreased at 50°C and 70°C (Figure 7).

Al20Cu2Mn3

Al20Cu2Mn3 also proved to be spontaneously passive over the temperature range investigated in this study (Figure 2[d]). This phase showed a strong tendency to repassivate, and repassivation was somewhat promoted by increasing temperature. As with Al7Cu2Fe, the repassivation corrosion potential was always observed to be more positive than the corrosion potential.

The corrosion potential was modestly dependent on temperature, though no trend was apparent across the four temperatures. It increased between 10°C and 30°C, but then decreased by about 200 mV between 30°C and 70°C, through the range of −0.700 VSCE to −0.900 VSCE. The pitting potential decreased from about −0.350 VSCE to −0.450 VSCE from 10°C to 50°C, but increased somewhat at 70°C. Passive current densities increased from about 1 μA/cm2 to 10 μA/cm2 with increasing temperature up to 50°C, but no obvious change was observed from 50°C to 70°C. As with all phases tested (except Al-4%Cu), an increase of corrosion rate was observed with increasing temperature (Figure 5). Figure 6(e) reveals that cathodic kinetics were slowest at 70°C, followed by those at 10°C, while those at 30°C and 50°C were similar.

Al2CuMg (S-Phase)

S-phase is reactive as a result of the presence of Mg, making its corrosion process complex. S-phase corrodes initially by selective dissolution of Mg and Al. This process enriches the surface of S-phase with Cu, which ennobles the phase and renders it a net cathode with time. Dealloyed S-phase can support oxygen reduction at an appreciable rate.24-25,27  The dealloying process is electrochemically complicated and the OCP of the phase oscillates between noble and active values as the process occurs.37  This complicated behavior was observed in 0.1 M NaCl even at temperatures as low as 10°C, as shown in Figure 8. As the OCP delay increased from 13 s to 430 s, the corrosion potential increased from −0.933 VSCE to −0.580 VSCE. A further increase of the OCP delay to 2,000 s did not change the corrosion potential, but no passive response or breakdown potential was observed up to applied potentials of 0.3 VSCE. This response suggests ennoblement of the phase by dealloying to produce a nonprotective, possibly porous surface enriched in Cu.

FIGURE 8.

The anodic polarization curves of S-phase after increasing delay time in 0.1 M NaCl at 10°C.

FIGURE 8.

The anodic polarization curves of S-phase after increasing delay time in 0.1 M NaCl at 10°C.

These results show that the effect of temperature on the dealloying kinetics and polarity reversal between S-phase and the matrix cannot be evaluated fully from anodic polarization curves alone. To address this, OCP transients were used to characterize the effect of temperature on the corrosion properties of S-phase (Figure 9). At 10°C, the OCP was fairly stable at about −910 mVSCE, and there was no obvious variation during the 2,000 s measurement. At 30°C, a potential jump up to −680 mVSCE after 670 s at a stable OCP (around −910 mVSCE) was observed, but the elevated potential only lasted for about 70 s before returning to −910 mVSCE. This low potential remained for an additional ~200 s, and was then followed by continuous oscillation. At 50°C, the OCP demonstrated great fluctuation immediately upon exposure to 0.1 M NaCl, which lasted for the entire duration of the measurement, albeit with some stabilization at −970 mVSCE after ~1,350 s. The most positive potential observed in the 50°C reading was quite high at −370 mVSCE. During the first ~100 s of the 70°C reading, oscillations similar to those observed at 50°C were detected, but after ~100 s, the OCP gradually increased with little variation to −240 mVSCE, a value close to the corrosion potential of pure Cu at 70°C. For reference, OCP transients acquired at each of the four temperatures on the AA2024-T3 alloy are shown in Figure 10.

FIGURE 9.

OCP changes of S-phase as a function of time in 0.1 M NaCl at different temperatures.

FIGURE 9.

OCP changes of S-phase as a function of time in 0.1 M NaCl at different temperatures.

FIGURE 10.

OCP transient of AA2024-T3 alloys at different temperature in 0.1 M NaCl.

FIGURE 10.

OCP transient of AA2024-T3 alloys at different temperature in 0.1 M NaCl.

Localized Corrosion Morphology of Secondary Phase Intermetallic Compounds Found in AA2024-T3 at Different Temperatures

The localized corrosion morphology of AA2024-T3 after a 1-h exposure in 0.1 M NaCl at 10, 30, 50, and 70°C reveals different extents of corrosion associated with two important secondary phase IMCs: S-phase particles and Al-Cu-Mn-Fe type particles (Figure 11). At 10°C, no trenching corrosion was observed around the Al-Cu-Mn-Fe type particles, but S-phase showed dealloying corrosion and obvious deep trenching with a width of ~300 nm (Figure 11[b]). A magnified view (Figure 12[a]) shows a nanoscale porous structure on the dealloyed S-phase. As the exposure temperature increased to 30°C, a similar dealloying and trenching morphology was observed on S-phase particles, albeit with larger trenching width and pore sizes (Figures 11[d] and 12[b]). It is noteworthy that dealloying corrosion was limited to the surface only at 10°C and 30°C. For Al-Cu-Mn-Fe type particles exposed at 30°C, a slight trench formed around the particles that was partially covered by corrosion products (Figure 10[c]). The trenching that occurred around Al-Cu-Mn-Fe type particles became larger at 50°C, as shown in Figure 11(e). However, a detailed view of the surface reveals porous and loose corrosion products on the surface. The cross section of S-phase particles exposed at 50°C, which is covered by porous corrosion products (Figure 11[f]), clearly shows severe dealloying corrosion penetrating half of the particle’s depth (Figure 12[c]). At 70°C, many small pits (less than 5 μm) formed in S-phase, attributable to dissolution of Mg and Al (Figure 11[h]), whereas with Al-Cu-Mn-Fe type particles some large portions completely dissolved and formed larger pits (Figure 11[g]). EDS examination of S-phase remnants exposed at 70°C showed an almost complete absence of Mg, further supporting extensive dealloying. Fast dealloying kinetics at high temperature is also supported by the cross section of an S-phase remnant (Figure 12[d]), which shows that dealloying penetrated through the entire particle and large holes were formed after only 1 h of exposure. Al-Cu-Mn-Fe type particles exposed at 70°C displayed severe corrosion along with a significant drop in the Al concentration. The significant dissolution of Al in these particles leads to a fragmentary morphology, which was observed in back scattered electron images (not shown).

FIGURE 11.

The localized corrosion morphology of secondary phase IMCs after 1 h exposure in the 0.1 M NaCl.

FIGURE 11.

The localized corrosion morphology of secondary phase IMCs after 1 h exposure in the 0.1 M NaCl.

FIGURE 12.

Magnified views of S-phase after 1 h exposure in 0.1 M NaCl at each temperature. (a) and (b) Surface views. (c) and (d) Cross-sectional views.

FIGURE 12.

Magnified views of S-phase after 1 h exposure in 0.1 M NaCl at each temperature. (a) and (b) Surface views. (c) and (d) Cross-sectional views.

It should be noted with regards to microcell experiments versus free corrosion testing that the dealloying of S-phase can be accelerated in the free corrosion experiments by galvanic coupling with the noble surrounding matrix. This effect can lead to faster dissolution rates and more rapid formation of detached Cu-rich nanoparticles by dealloying than in the case of electrochemical testing of a single S-phase particle with the microcell. Lastly, most of the secondary phase IMCs were at least somewhat covered by nonconductive corrosion products, as demonstrated by the bright areas in the SEM images arising from charge accumulation.

DISCUSSION

General Effects of Temperature on Intermetallic Compounds and Their Galvanic Interactions

The data in this study show that temperature exerts a strong effect on the electrochemical properties of IMCs. For example, a comparison of the corrosion potential values for synthesized Al20Cu2Mn3 and Al7Cu2Fe at the temperatures tested shows that their corrosion potential is higher than that of the matrix at 10°C and 30°C but lower at 50°C and 70°C. As a result, their galvanic relationship with the matrix changes from cathodic at low temperatures to anodic at high temperatures. This transition indicates that Al-Cu-Mn-Fe type particles would be expected to undergo dissolution upon exposure at 50°C and above, leading to more dissolution of the IMC. This is a strong example of the importance of considering temperature when discussing the effects of the corrosion of aluminum alloys. It is well established that the corrosion of aluminum alloys is attributable to the galvanic interactions between IMCs and the alloy matrix. Two types of localized corrosion morphologies are associated with this type corrosion.5,7,11,24,38-39  First, there is corrosion that occurs primarily in the matrix in areas surrounding IMCs. Second, dissolution of an IMC arising from dissolution of active phases within the IMC can result in deep pits in the alloy in the place of the original IMC. This second type of pit can contain remnants of the original IMC. In the first instance, the trenching corrosion is believed to arise from IMCs that are noble to the surrounding matrix. In the second instance, the portions of the IMC that are active relative to the matrix dissolve and the remnants that are left behind are typically noble to the matrix. In the case of noble IMCs (initial or after dealloying), trenching often occurs around the IMC, attributable to dissolution of the Al matrix surrounding a noble IMC resulting from the increased alkalinity generated by the ORR on the noble IMC.39  The role of Al-Cu-Mn-Fe type particles, revealed in this work, indicate that the morphology of trenching corrosion induced by galvanic interactions between Al-Cu-Mn-Fe type particles and the matrix would be dependent on temperature. At low temperatures (commonly studied), these particles would be associated with trenching in the matrix areas surrounding the IMC, with little effect on the IMC. As the Al-Cu-Mn-Fe particles are assumed to be cathodic throughout their exposure at 10°C and 30°C, the increase in trenching observed at 30°C is supported by the increased cathodic current observed on the similar synthesized particles at that temperature (Figure 6). However, at higher temperatures, more aggressive corrosion associated with the loss of at least part of the particle would be expected in early stages, likely leading to different long-term results. EDS analysis on Al-Cu-Mn-Fe particles in AA2024-T3 exposed at 50°C revealed that the Al wt% dropped by 28% to 85%, while the Cu wt% increased by a factor of 1.3 to 4.3. The expected combination of dissolution of the IMC (occurring during the initial anodic phase) and trenching (occurring after the IMC changes from anodic to cathodic) can be observed in Figure 11(h) (Al-Cu-Mn-Fe particle in AA2024-T3 exposed at 70°C). This is in contrast to observations, in this and previous work, which revealed no sign of attack on Al-Cu-Mn-Fe particles in AA2024-T3 when exposed in aqueous NaCl at room temperature.7  The dissolution of initially anodic Al-Cu-Mn-Fe particles at high temperatures and their subsequent change to cathodic behavior is further supported by the high repassivation potentials observed on the synthesized particles (Figures 1[c] and [d]).

In general, the pitting potential of most IMCs slightly decreased with increasing temperature, as shown in Figure 4. The decreasing trend of pitting potential with temperature implies that the passive film on IMCs becomes less stable and can break down at high temperatures. However, it should be noted that the fast corrosion rate during the 60 s OCP measurement at 70°C could lead to surface composition changes (as revealed in the EDS data). This may result in an increased pitting potential. Although the erratic behavior of S-phase precludes the measurement of a temperature-dependent pitting potential, the difference in the OCP measurement made at 70°C versus the other temperatures (Figure 9) combined with the cross-section morphology of Figure 12 strongly suggests that the dealloying of the active S-phase occurs much more quickly at 70°C than at lower temperatures. As is typically assumed (at ambient temperatures) regarding the galvanic interactions of S-phase with the matrix, the dissolution of the active portions of the phase switch these IMCs from anodes to cathodes and thereby results in the loss of volume resulting from the partial loss of the IMCs and, subsequently, to the loss of matrix volume when the IMC becomes noble to the matrix. Although (unlike Al-Cu-Mn-Fe particles) S-phase particles dealloy at all temperatures, the process is clearly much faster at higher temperatures, possibly leading to even more complexities in the role of S-phase in corrosion of aluminum alloys. As previously reported, dissolution of S-phase particles can occur without propagating into the Al matrix, but propagation of corrosion into the matrix will occur if the local cathodes can support high reaction rates.7  The concentration of Cu was as high as 62.1 wt% in S-phase particles exposed at 70°C. Furthermore, the exceptionally fast rate at which S-phase dealloys at higher temperatures, supported by Figures 10 and 11, would quickly provide an increase in Cu surface area, thereby providing conditions under which localized corrosion would be more likely to propagate into the matrix and trenching would be more severe.

Effect of Temperature on the Cathodic Response of Intermetallic Compounds

In aqueous aerated solutions, the cathodic reaction is generally the ORR, which is primarily a diffusion-controlled process.40  The diffusion-limited reaction leads to a limiting current density, which is closely related to the oxygen diffusion coefficient (D), oxygen solubility (C), and diffusion thickness (δ), and can be expressed by the following equation:

formula

where F is the Faraday constant and n is the number of equivalents per mole of O2 (4 eq/mol).

The temperature dependence δ is not well characterized. Under hydrodynamic conditions such as a rotating electrode, δ increases only slightly by about 10% as temperature increases from 20°C to 80°C.41  The temperature dependence of both D and C is well documented. Oxygen solubility decreases with increasing temperature, while diffusion increases with increasing temperature.34-35  Li, et al., calculated the limiting current density of the ORR as a function of temperature based on Equation (1) and determined a peak limiting current density at around 55°C.42  In this work, most IMCs showed a maximum cathodic current density at 50°C (Figure 7), which is consistent with the above calculation. The cathodic kinetics on Al2Cu was much faster than on the other IMCs at all temperatures, which is attributed to a relatively high amount of Cu and its ability to support the ORR.

Also of note with regards to cathodic activity on IMCs is the high corrosion rate of the matrix at high temperatures (Figure 5). Clearly, the high cathodic behavior observed on IMCs at high temperatures could drive the high anodic behavior observed on the matrix at high temperatures.

CONCLUSIONS

Electrochemical measurements performed with a microcell provided information on the effects of temperature on synthesized IMCs similar to those found in AA2024-T3. The microcell results were combined with results obtained involving actual IMCs in AA2024-T3 samples subjected to free corrosion experiments.

  • In general, the corrosion rate of synthesized IMCs increased with increasing temperature.

  • Temperature generally decreased the pitting potential of most synthesized IMCs, indicating IMCs are more susceptible to corrosion at high temperature.

  • Temperature affects an IMC’s corrosion potential and, therefore, its role in localized galvanic corrosion. Corrosion potentials measured on synthesized Al-Cu-Mn-Fe type particles indicate that similar particles in AA2024-T3 would be expected to behave as cathodes at 10°C and 30°C. At 50°C and 70°C, however, these particles would be expected to dealloy as they would be anodic to the matrix. After dealloying at the higher temperatures, these particles would be expected to switch to cathodic behavior. This high-temperature behavior leads to a combination of corrosion of the Al-Cu-Mn-Fe type particles (not found at lower temperatures) followed by trenching similar to, but more severe, than that found at lower temperatures. The fragmentary morphology displayed by Al-Cu-Mn-Fe type particles exposed at high temperature (70°C) is a result of selective dissolution of Al.

  • At high temperatures, S-phase particles undergo faster, more severe dealloying, leading to a faster transformation of anodic behavior to cathodic behavior and a higher Cu surface area, thereby making propagation of localized corrosion into the matrix more likely.

  • Most IMCs show maximum limiting current density for the ORR at 50°C, attributable to the effect of temperature on oxygen solubility and diffusivity, and diffusion thickness, which is consistent with the calculated theoretical limiting current density.

(1)

UNS numbers are listed in Metals and Alloys in the Unified Numbering System, published by the Society of Automotive Engineers (SAE International) and cosponsored by ASTM International.

Trade name.

ACKNOWLEDGMENTS

This research is sponsored by the U.S. Air Force Academy under agreement number FA7000-12-2-0015. The U.S. Government is authorized to reproduce and distribute reprints for Governmental purposes notwithstanding any copyright notation thereon. The views and conclusion contained herein are those of the authors and should not be interpreted as necessarily represented official policies or endorsement, either expressed or implied, of the U.S. Air Force Academy or the U.S. Government.

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