**4. Al–Co Alloys**

The Al–Co alloys composed of SCIPs were investigated by Lekatou et al. [79–82] The authors prepared a series of novel Al–Co alloys with 3.3–10.3 at.% Co by arc-melting. The microstructures obtained were ranging from fully eutectic to hypereutectic microstructures with primary precipitation of structurally complex Al9Co2. Relatively uniform and directional microstructures were obtained (Figure 11, [81]). The fraction of directionally solidified Al9Co2 was increasing with increasing Co concentration. Microstructures of the materials before and after corrosion are compared in Figure 11. The alloys displayed a similar corrosion behavior in 3.5 wt.% NaCl. The corrosion attack resulted in a preferential dissolution of Al solid solution (ss).

**Figure 11.** Microstructure of the rapidly solidified Al–Co alloys before and after corrosion in aqueous NaCl, adapted from reference [81].

A potentiodynamic polarization behavior of the Al–Co alloys in aqueous NaCl has been studied in refs. [81,82]. The Al96.7Co3.3 alloy showed a slightly superior corrosion performance compared to the remainder of the alloys as it had a relatively low corrosion current. Nevertheless, all Al–Co had a substantially higher corrosion resistance compared to Al. The anodic dissolution behavior was found to consist of four stages [81,82]. The individual stages were assigned to the following processes:


At higher Co concentrations, a fragmentation of Al9Co2 in the corroded alloys occurred. The fragments piling in the gaps resulting from Al(ss) dissolution retarded the corrosion attack of the electrolyte. Al9Co2 had a higher electrochemical potential compared to Al(ss).

The corrosion behavior of Al–Co complex metallic alloys with 24–29 at.% Co was studied by Palcut et al. [83,84]. The following relative nobility of Al–Co CMAs has been found:

$$\text{Al(ss)} < \text{Al}\_3\text{Co}\_2 < \text{Al}\_{13}\text{Co}\_4 < \text{Al}\_5\text{Co}\_2 < \beta(\text{AlCo}) \tag{10}$$

The nobility of IMCs increases with increasing Co concentration. The volume fractions of the phases and physical contacts between them play an important role in the alloy corrosion behavior. Results indicate that a galvanic mechanism is involved. Moreover, it should be mentioned that Al–Co IMCs are brittle [85]. Therefore, a piling of noble but brittle particles, such as β(AlCo), in pores resulting from massive dissolution of surrounding less-noble phases may significantly influence the alloy stability [82]. The structural defects in the alloy may act as rapid diffusion paths leading to a significant material degradation over time. The galvanic coupling of noble IMCs with more active phases may be critical to the alloy corrosion stability in halide-containing environments.

The parallel occurrence of SCIPs with similar chemical compositions has a positive effect on the corrosion susceptibility of the alloy [84,85]. The Al74Co24 alloy was composed of three phases with close chemical compositions (Z–Al3Co, Al5Co2, and Al13Co4, [84]). The Al74Co24 alloy had a higher corrosion potential compared to the remaining alloys which is an indicator of a superior corrosion resistance. The inspection of the alloy after corrosion testing revealed a relatively uniform phase dissolution [84]. The potential differences between constituent phases were probably small enough to initiate galvanic corrosion. The alloy corrosion could only be initiated at high electrode potentials. A polarization at high potentials resulted into a massive degradation of this alloy.

To further investigate the corrosion susceptibility of individual SCIPs with close chemical composition, an annealing of the Al74Co26 alloy at 1000 ◦C for 330 h has been carried out [86]. The annealing resulted in equilibrium microstructure of the alloy composed of Z–Al3Co and Al5Co2. The Z–Al3Co phase in the as-annealed Al74Co26 alloy was significantly less attacked. Although the bulk of this phase comprises less aluminum, it appears to be nobler and less susceptible to pitting corrosion compared to Al5Co2. The reason for this behavior could stem in a different structure of the phase surface. The Al5Co2 surface is terminated in puckered layers [87]. The surface of Z–Al3Co, on the other hand, is more densely populated compared to Al5Co2 [88]. Therefore, Z–Al3Co was less prone to corrosion attack.

The corrosion behavior of the as-annealed Al74Co26 alloy was investigated in neutral (NaCl, 0.6 mol dm−3), alkaline (NaOH, 10−<sup>2</sup> mol dm−3) and acidic electrolytes (HCl, 10−<sup>2</sup> mol dm<sup>−</sup>3) by cyclic potentiodynamic polarization [86]. The potentiodynamic curves are shown in Figure 12. Anodic parts of the curves measured in HCl and NaCl solutions displayed a passive region which was followed by an abrupt current density increase. When the polarization scan was reversed, a positive hysteresis was found. These features indicate pitting corrosion. The polarization behavior in NaOH, on the contrary, corresponds to uniform alloy corrosion.

**Figure 12.** Potentiodynamic cyclic polarization curves of near-equilibrium Al74Co26 alloy in different electrolytes, re-plotted from reference [86].

The forward curves were evaluated by Tafel extrapolation and corrosion currents and corrosion potentials were obtained [86]. The lowest corrosion potential and highest corrosion current were found for NaOH. The highest corrosion potential and lowest corrosion current, on the other hand, were observed for the HCl solution. This behavior is in accordance with equilibrium E–pH diagram of Al (Figure 7).

#### **5. Al–Cr Alloys**

The Al–Cr alloys are expected to demonstrate a good corrosion resistance due to high concentrations of Al and Cr [89]. Both are passivating elements producing protective scales. The corrosion resistance of an Al70Cr20Fe10 alloy was studied by Li et al. [90]. The authors used commercial gas-atomized Al70Cr20Fe10 powders that were consolidated by spark plasma sintering. The phases present in the sintered Al–Cr–Fe pellets were the following: an icosahedral phase (i–Al–Cr–Fe), decagonal phase (d–Al–Cr–Fe) and crystalline Al8(Cr,Fe)5 and Al9(Cr,Fe)4 phases. Authors measured an open circuit potential (OCP) of the alloy in 3.5 wt. % NaCl and found that the OCP was nobler compared to Al. The OCP of the alloy was stable over time, indicating that an equilibrium has been rapidly established on the alloy surface. The Al70Cr20Fe10 alloy had a nobler corrosion potential and hence a lower susceptibility to corrosion compared to Al. It passivated in saline solution spontaneously due to significant amount of Cr. The alloy had a higher corrosion rate compared to pure Al [90]. Nevertheless, the corrosion rate was close to that of 316 stainless steel and smaller than AISI 440C stainless steel or AISI H13 tool steel.

The passivation behavior of Al–Cr–Fe alloys was studied by Ott et al. [91] The authors used a flow microcapillary plasma mass spectrometry. The schematic of the experimental set up is shown in Figure 13 [91]. In the experiment, a tiny microcapillary was positioned on the alloy surface and continuously filled with the desired solution. The flow injection was operated in loops by switching the valve. The loop volume was continuously transferred to the inductively coupled plasma mass spectrometer (ICP MC) for element analysis. The microcapillary was refilled with fresh electrolyte from the reservoir. The circulation was ensured by a peristaltic pump. A microscope was included to control precise positioning of the capillary on the alloy surface. The technique provided time-resolved information about transient electrochemical processes and element-specific dissolution at the metal–electrolyte interface.

**Figure 13.** Schematic of the microcapillary flow ICP MS setup, re-drawn from reference [91].

The authors prepared and studied a polycrystalline γ-phase Al64.2Cr27.2Fe8.1 alloy (composition given in at.%, [91]). The corrosion behavior was studied in two acidic solutions: H2SO4 (pH 0) and HCl (pH 2). In sulfuric acid, very low element dissolution rates were found. Neither Fe nor Al is stable at low pH [69]. Therefore, Cr is an essential element in the passive film stability. It helps to stabilize the Al cations within the passive film, as evidenced by a low Al release over 2 h. A possibly mixed oxy-hydroxide of Al and Cr was suggested to have been formed on the alloy surface (Figure 14, [91]). The dynamic passivation mechanism is related to the fact that the cation dissolution occurring at the oxyhydroxide–solution interface (<sup>2</sup> ) is compensated by additional film growth at the metal–oxyhydroxide interface (<sup>1</sup> ). Longer air-aging was found to be beneficial for stabilizing the passive film.

**Figure 14.** Schematic of the passive film evolution on γ-phase Al64.2Cr27.2Fe8.1 alloy, re-drawn from reference [91].

In chloride-containing hydrochloric acid, ten times higher Al dissolution rates were found at the OCP, suggesting a decreasing stability of the spontaneously formed passive film [91]. The thickness of the dissolved passive film was much higher compared to H2SO4 (Table 3). But even in HCl, a potentiostatic polarization at 0.18VSCE slowed down the dissolution processes at the oxyhydroxide–solution interface by a factor of 6. The electrochemical polarization at low passive potentials induces electrical field generated oxide film modification, thereby increasing the chemical stability at the oxyhydroxide– solution interface. In the high potential passive region, a localized attack was initiated with subsequent active metal dissolution.



<sup>a</sup> Volts versus saturated calomel electrode, <sup>b</sup> Open circuit potential.

The passivation behavior of Al–Cr–Fe complex metallic alloys in NaCl+HCl mixtures was further investigated by Beni et al. [92] The authors prepared three alloys: polycrystalline single phase Al79.5Cr12.5Fe8.0 (composed of orthorhombic phase), Al64.2Cr27.2Fe8.1 (single phase, composed of cubic γ phase) and single crystalline orthorhombic Al79.0Cr15.0Fe6.0. The corrosion behavior of the different alloys could be explained considering the passivating role of Cr combined with Fe oxyhydroxide precipitation. The anticipated reaction mechanism is presented in Figure 15 [92]. The Al79.5Cr12.5Fe8.0 alloy was found to undergo an active dissolution in the electrolyte, as proven by the high element concentrations in solution measured by ICP MS. The chromium concentration (12.5 at.%) was small but sufficient to stabilize the initially air-formed oxyhydroxide for 22 days, as evidenced by the constant low pH of the solution and low dissolution compared to Al. The concentration of Cr was, however, too low to provide a long-term protection. A thick and non-protective layer has been formed on the surface. With increasing Cr concentration, a protective layer on the alloys started to form. The Cr concentration of 15.0 at.% was sufficient to stabilize the passive film up to 78 days. A complete and long-lasting protective scale was finally achieved at 27.2 at.% Cr [92,93].

**Figure 15.** Schematic of passivation/dissolution processes on pure Al and Al–Cr–Fe alloys in aqueous HCl + NaCl mixture with initial pH 2, re-drawn from reference [92].

#### **6. Al–Noble-Metal Alloys**

Massiani et al. [35] investigated the corrosion behavior of crystalline and quasicrystalline phases in the Al–Cu–Fe(–Cr) alloys by potentiodynamic polarization in strongly acidic and alkaline solutions. They found that the corrosion resistance was determined by the alloy chemical composition. The complex crystal structure had only a minor influence. Rüdiger and Köster [94] found that the corrosion behavior of quasicrystals and their approximants in the Al–Cu–Fe alloy system could be explained based on the electrochemical behavior of the component elements. The surface of the icosahedral Al63Cu25Fe12 was covered by a thick non-protective layer composed of Cu2O, Al(OH)3 and metallic Cu. The scale chemical composition was comparable to crystalline Al7Cu2Fe. The complex crystal structure thus did not have a substantial influence on the corrosion resistance [94]. Furthermore, the authors observed a formation of porous Cu layer in i–Al63Cu25Fe12 phase at pH 0.

While Rüdiger and Köster studied single phase quasicrystalline alloys, Huttunen et al. investigated Al–Cu–Fe alloys composed of several different phases (Table 4, [95]). The corrosion behavior was determined by anodic polarization. The microstructural features and phase constitution of the alloys before and after the polarization were studied by scanning electron microscopy and X-ray diffraction.


**Table 4.** Chemical composition of SCIPs in Al–Cu–Fe alloys (in at.%) studied in reference [95].

The study was focused on four different Al–Cu–Fe alloys: Al67.5Cu20Fe12.5, Al65Cu20Fe15, Al62.5Cu25Fe12.5 and Al60Cu27.5Fe12.5 [95]. The authors found that the presence of structurally complex phases did not improve the alloys corrosion resistance [95,96]. The chemical composition of the phases, however, was of great importance. The corrosion potentials of Al–Cu–Fe alloys with Cu-rich phases were nobler and had lower corrosion rates compared to Cu-lean alloys [95]. Relative amounts of the phases and their electrical contacts were also significant factors influencing the overall corrosion behavior. Phases with high Cu concentration remained virtually unaffected by corrosion. The phases with low Cu atomic fractions were susceptible to corrosion attack. This behavior could be explained by the higher electrode potential of Cu compared to Al and Fe (Table 2). The corrosion was found to occur by galvanic mechanism near phase boundaries. The corrosion behavior of Al–Cu–Fe alloys was studied in alkaline, neutral, and acidic solutions. In alkaline and neutral electrolytes, an oxidation of Al and Cr occurred on the surface of the alloys. The oxidation was accompanied with Cu deposition on the alloy surface. The Cu deposition interfered with passive layer formation and introduced pores into the oxide film. The icosahedral ψ–Al65Cu20Fe15 was the only phase capable of forming a stable passive layer on the surface [95].

The Al–noble-metal alloys are interesting materials from electrochemical point of view. The alloys are prone to selective dissolution of less noble elements (leaching) because of markedly different electrode potentials of the constituent metals [97]. The less noble elements tend to dissolve in the electrolyte, leaving behind their vacant positions. As a result of leaching, a porous de-alloyed structure forms on the alloy surface [97]. The leaching can be either uniform or localized. Examples of leaching include preferential dissolution of Zn from brass (de-zincification) or Fe removal from gray cast iron (a so-called graphitic corrosion) [98]. Other examples include de-aluminification, de-nickelification and de-cobaltification [99].

Mishra et al. studied a chemical leaching of Al–Cu–Co decagonal quasicrystals [100]. The authors prepared two alloys with Al65Cu15Co20 and Al65Cu20Co15 chemical compositions and studied their corrosion behavior in aqueous NaOH (10 mol L−1). The alloys were immersed in the alkaline solution at room temperature for 8 h. Most Al atoms were removed (Figure 16, [100]). A nearly uniformly distributed metallic Cu, Co and Co3O4 nanoparticles were found on the alloy surface after leaching. The crystallite size was calculated from the XRD reflections' broadening and further confirmed by TEM [100]. The nanostructure formation of the leached layer was controlled by Al dissolution rate during leaching. The dispersed Cu and Co nanoparticles were stable in the leached layer and Cu agglomeration was suppressed.

**Figure 16.** Chemical composition (in at.%) of as-cast and as-leached Al–Cu–Co alloys, plotted from data in reference [100].

Porous nanostructures composed of noble metals are important catalysts. The formation of Cu-rich porous nanostructure from decagonal Al65CoxCu35−<sup>x</sup> alloys (x = 12.5, 15, 17.5 at.%) was studied by Kalai Vani et al. [101] A selective dissolution of Al and Co was achieved by combined immersion of the alloys in both NaOH (5 mol L−1) and Na2CO3 (0.5 mol L−1) electrolytes. A high specific surface of 30 m<sup>2</sup> g−<sup>1</sup> of the porous Cu structure was achieved.

The electrochemical de-alloying of binary Al–noble metal alloys was also studied [102]. It has been shown that nano-porous Pd, Ag and Au with various structures can be produced through electrochemical leaching of the Al–based alloys in NaCl aqueous solution. Galvanic interactions between coexisting phases dominate during corrosion of double phase alloys. The level of de-alloying depends on the critical de–alloying potential [103], diffusion of the noble element and reactivity of the noble element and chloride anion. The porosity evolution is a dynamic process. It is not a simple excavation of the less noble phase from two phase material. The formation of the porous nanostructure involves selective leaching of Al and is accompanied with coarsening of the noble element due to surface diffusion.

The corrosion behavior of Al–Pd alloys composed of SCIPs was studied in references [104,105]. The open circuit potentials are given in Figure 17. The OCPs decrease in the following order:

$$\text{Al}\_{67}\text{Pd}\_{33\prime}\text{Al}\_{72}\text{Pd}\_{28}\text{ (group I)} > \text{Al}\_{77}\text{Pd}\_{23\prime}\text{ Al}\_{88}\text{Pd}\_{12}\text{ (group II)}\tag{11}$$

**Figure 17.** Open circuit potentials of Al–Pd alloys in aqueous NaCl, re-plotted from data in reference [105].

The OCPs of the alloys decrease with decreasing Pd concentration. This observation is in accordance with expectations since Al is electrochemically more active compared to Pd (Table 2). The corrosion resistance of both as-annealed and as-solidified alloys was comparable. A large difference, however, between OCP and Ecorr has been found for group I alloys (Al67Pd33 and Al72Pd28). The OCPs of these alloys were comparable to their pitting potentials obtained by potentiodynamic polarization. The Al67Pd33 and Al72Pd28 alloys were probably in a pitting corrosion stage during the OCP measurement. This suggestion was manifested by oscillations of OCP resulting from a possible pitting behavior (Figure 17). The anodic dissolution of the alloy at pits requires a generation of cathodic current from the surrounding surface. The electric current bursts are transient and cause a temporary decrease in the OCP value. The pitting corrosion sites are usually very small. However, the current densities during transient bursts inside the pits can be up to 1 A/m2 [106]. The significant corrosion rates of the alloy are due to aggressive environments developed inside the pits. Although the pits are small, they may affect the electrochemical response of much larger surface areas. Therefore, the differences in current densities on separated anodic and cathodic sites are reflected in potential oscillations (so-called electrochemical noise associated with localized corrosion).

Interactions between phases with different chemical composition play a significant role in alloy corrosion. The Al67Pd33 and Al72Pd28 alloys were found to be composed of structurally complex ε<sup>n</sup> (Al3Pd) and δ(Al3Pd2). The electrochemical nobility of Al–Pd phases in aqueous NaCl (0.6 mol L<sup>−</sup>1) increases in the following order

$$\delta(\text{Al}) < \varepsilon\_n(\text{Al}\_3\text{Pd}) < \delta(\text{Al}\_3\text{Pd}\_2) \tag{12}$$

The δ phase has a higher concentration of Pd. It serves as a cathode, and thereby further accelerates the anodic dissolution of the surrounding ε<sup>n</sup> phase. The corrosion mechanism of Al–Pd alloys in aqueous NaCl involves a rapid passivation stage on the alloy surface [105,106]. However, once a breakdown potential is reached during potentiodynamic polarization, the passive layer becomes unstable and susceptible to local attack by chloride anions. Consequently, chloro–aluminum complex cations are formed and released into the solution. The local disruption of the passive layer reveals a naked alloy surface which becomes more susceptible to further corrosion attack.

The microstructures of as-annealed and as-solidified Al72Pd28 and Al67Pd33 alloys had similar features after corrosion testing. In the alloys a high number of inter-penetrating channels have been found [105,106]. Pits were also observed in the inter-connection between the channels. The formation of channels was driven by pitting and de–alloying. The pits were probably initiation sites of the channels. A preferential de-alloying of Al (de-aluminification) has also been observed. The preferential leaching of Al led to initiation of microcracks. During rapid solidification residual stresses have been accumulated in the alloys. The stresses were released during leaching, resulting in continuous tunnels inter-penetrating the surfaces of de-alloyed materials. A similar corrosion behavior was also found for the Al–Pd–Co alloys (Figure 18, [107]). The de-alloying of Al was more pronounced in the as-solidified alloys. This is probably a consequence of their higher defect concentrations compared to as-annealed alloys. The de–alloying behavior was significantly reduced in as-annealed alloys [105].

**Figure 18.** Microstructure of as-corroded Al–Pd–Co alloys: backscatter scanning electron microscopy images of Al70Pd25Co5 (**a**) and Al74Pd12Co14 (**b**) and confocal laser scanning images of Al70Pd25Co5 (**c**) and Al74Pd12Co14 (**d**). Reproduced from reference [107].
