**3. Results and discussion**

#### **3.1 Potentiodynamic polarization curve and weight loss at 25 o C**

Fig. 1 shows the anodic dissolution behaviour of alloys in 0.5 M H2SO4, while Table 2 summarizes relevant data. This figure reveals a well-defined passive region of 0 VSHE to 1.2 VSHE in all curves. All curves, except for the one at x = 0.25 (C-0.25), show a secondary passive region at 0.15 VSHE. This passivation is attributed mainly to the further oxidation or hydroxidation of the passive oxide film, thus altering the valence of Cr [18,19]. Fig. 1 also indicates that the secondary passive regions of C-0.50 and C-1.00 are more prominent than those of C-0 and C-0.25. This observation is due to the selective dissolution in the duplex FCC-BCC structure for C-0.50 and in the BCC-ordered BCC structure for C-1.00, as compared with C-0 and C-0.25 which are single FCC phase. In the active-passive transition region, different compositions at different secondary passivation potentials reveal different dissolution rates owing to a selective dissolution. This observation resembles that observed in duplex phase stainless steel [20,21].


Table 2. Potentiodynamic polarization curve diagram parameters of alloys C-x and SS 304 at 25oC.

Fig. 1. Potentiodynamic polarization curve diagrams for alloys C-x and SS 304 at 25oC.

compositions. Finally, samples were examined before and after 3 days immersion of 0.5 M

Fig. 1 shows the anodic dissolution behaviour of alloys in 0.5 M H2SO4, while Table 2 summarizes relevant data. This figure reveals a well-defined passive region of 0 VSHE to 1.2 VSHE in all curves. All curves, except for the one at x = 0.25 (C-0.25), show a secondary passive region at 0.15 VSHE. This passivation is attributed mainly to the further oxidation or hydroxidation of the passive oxide film, thus altering the valence of Cr [18,19]. Fig. 1 also indicates that the secondary passive regions of C-0.50 and C-1.00 are more prominent than those of C-0 and C-0.25. This observation is due to the selective dissolution in the duplex FCC-BCC structure for C-0.50 and in the BCC-ordered BCC structure for C-1.00, as compared with C-0 and C-0.25 which are single FCC phase. In the active-passive transition region, different compositions at different secondary passivation potentials reveal different dissolution rates owing to a selective dissolution. This observation resembles that observed

Ecorr (VSHE) Icorr (A/cm2) Epp (VSHE) Icrit (A/cm2) Ipass (A/cm2)

C-0 -0.081 15.8 0.002 42.8 4.5 C-0.25 -0.095 16.7 0.008 87.4 7.1 C-0.50 -0.084 13.4 0.017 117.2 6.4 C-1.00 -0.094 13.1 0.010 198.0 13.9 SS 304 -0.185 45.3 -0.071 603.0 19.1

Table 2. Potentiodynamic polarization curve diagram parameters of alloys C-x and SS 304 at

Fig. 1. Potentiodynamic polarization curve diagrams for alloys C-x and SS 304 at 25oC.

**C** 

**3.1 Potentiodynamic polarization curve and weight loss at 25 o**

H2SO4.

**3. Results and discussion** 

in duplex phase stainless steel [20,21].

Alloys C-x & SS304

25oC.

Table 2 reveals that the corrosion potential (Ecorr) and the corrosion current density (Icorr) for all of the alloys differ only slightly, and no obvious trends occur for Ecorr and Icorr vs. x variation. The above phenomenon can be attributed to the spontaneous passivation of pure Al in H2SO4 [22]. Al metal spontaneously passivates in H2SO4, explaining why its corrosion potential is ready in the passive region, i.e., this passivation explains why the polarization curve of Al does not display an apparent active-passive transition region. However, elements such as Cr and Fe exhibit a large critical current density (Icri) for passivation, explaining why Cr and Fe dissolve more than Al before the alloy reaches its passive state. Thus, the variation of Al affects the active region of the polarization curves slightly. Furthermore, in H2SO4, all Al, Co, Cr, Fe, and Ni metals show passivity. Among them, Al has a relatively high passive current density (Ipass) [22,23] because only Al oxide can easily form a porous film on the metal surface [24]. Therefore, protection by oxide layer on the alloys with higher Al content is inferior to that with lower Al content. Fig. 1 thus reveals that Ipass increases with x.

The results of potentiodynamic polarization were compared via performing 15-day-dipping and weight loss experiments. In the 15-day-dipping and weight loss experiments, the corrosion rates for C-0.50 and C-1.00 were markedly higher than those of C-0 and C-0.25 (Fig. 2). This observation differs substantially from the values of Icorr obtained from polarization experiment (Fig. 1), in which the two groups only differ slightly, despite the fact that the trend is the same. A previous study found a similar deviation in corrosion current densities obtained from weight loss test and potentiodynamic polarization method [25].

Fig. 2. Diagram showing change in corrosion rate (g m-2 h-1) in the 15-day-dipping and weight loss measurement for alloys C-x.

Fig. 1 shows potentiodynamic polarization diagrams for the AlxCoCrFeNi alloys and SS 304. The alloys have better overall general corrosion behaviour, with a larger Ecorr and smaller Icorr, Icri, and Ipass than SS 304.

Electrochemical Passive Properties of AlxCoCrFeNi

(d) C-1.00 at various temperatures.

(x = 0, 0.25, 0.50, 1.00) High-Entropy Alloys in Sulfuric Acids 139

Fig. 3. Potentiodynamic Polarization Curve diagrams for (a) C-0, (b) C-0.25, (c) C-0.50, and

Fig. 4. The Arrhenius plots for alloys C-x and SS 304 at 20oC - 65oC.

#### **3.2 Effect of temperature on potentiodynamic polarization**

Fig. 3 shows polarization diagrams of AlxCoCrFeNi alloys at various temperatures. A rising temperature decreased the Tafel slopes of anode (Table 3), increased Icorr, and increased Ecorr and Et (the transpassive potential) slightly. The corrosion rate is directly related to Icorr, according to Arrhenius equation, Icorr = A exp(−Ea/RT) [13,26], where the pre-exponential factor A is generally independent of temperature and is a constant of alloys, where R denotes the gas constant, T denotes temperature, and Ea denotes activation energy for corrosion. In the case of small experimental temperature range, Ea is assumed to be independent of T. Consequently, Ea can be obtained from ln(Icorr) vs. 1/T plot. Fig. 4 shows such plots for the alloys and SS 304, indicating that Ea increases with x. This finding suggests that the corrosion rate is more sensitive to temperature for a larger Al content than for a smaller Al content. The ln(Icorr) vs. 1/T curves intersect with each other in a range of 23oC – 27oC. Beyond this temperature range, Icorr increases with x. The situation is reversed at temperatures lower than 23oC, which is inconsistent with a situation in which Eas for all alloys increase with x from 20oC to 65oC. Hence, Ea, i.e., an intrinsic property of metal, and A, i.e., a surface property of metal, are determinative factors of Icorr. While Ea depends only on x, A depends on both x and temperature (Table 4). Therefore, although Ea increases with x, A also increases with x. Combining the effects of Ea and A explains the different corrosion behaviours of the alloys with an increasing x at temperatures exceeding 27oC and lower than 23oC. Thus, the performance of passive films, when Al is added, at higher temperatures becomes inferior to that without addition of Al. In determining Icorr, A is more important than Ea at temperatures exceeding 27oC, while Ea is more important than A at temperatures lower than 23oC.


a Anodic Tafel slope βa in mV/decade, the measured Tafel regions are with 40~50 mV of overvoltage.

b Cathodic Tafel slope βc in mV/decade, the measured Tafel regions are with 150~170 mV of



Table 4. The fits for A(x, T) and Ea(x) in Icorr(x, T) = A(x, T) exp(-Ea(x)/RT).

overvoltage.

Fig. 3 shows polarization diagrams of AlxCoCrFeNi alloys at various temperatures. A rising temperature decreased the Tafel slopes of anode (Table 3), increased Icorr, and increased Ecorr and Et (the transpassive potential) slightly. The corrosion rate is directly related to Icorr, according to Arrhenius equation, Icorr = A exp(−Ea/RT) [13,26], where the pre-exponential factor A is generally independent of temperature and is a constant of alloys, where R denotes the gas constant, T denotes temperature, and Ea denotes activation energy for corrosion. In the case of small experimental temperature range, Ea is assumed to be independent of T. Consequently, Ea can be obtained from ln(Icorr) vs. 1/T plot. Fig. 4 shows such plots for the alloys and SS 304, indicating that Ea increases with x. This finding suggests that the corrosion rate is more sensitive to temperature for a larger Al content than for a smaller Al content. The ln(Icorr) vs. 1/T curves intersect with each other in a range of 23oC – 27oC. Beyond this temperature range, Icorr increases with x. The situation is reversed at temperatures lower than 23oC, which is inconsistent with a situation in which Eas for all alloys increase with x from 20oC to 65oC. Hence, Ea, i.e., an intrinsic property of metal, and A, i.e., a surface property of metal, are determinative factors of Icorr. While Ea depends only on x, A depends on both x and temperature (Table 4). Therefore, although Ea increases with x, A also increases with x. Combining the effects of Ea and A explains the different corrosion behaviours of the alloys with an increasing x at temperatures exceeding 27oC and lower than 23oC. Thus, the performance of passive films, when Al is added, at higher temperatures becomes inferior to that without addition of Al. In determining Icorr, A is more important than Ea at temperatures exceeding 27oC, while Ea is more important than A at temperatures

20 °C 35 °C 50 °C 65 °C

C-0 158 218 128 158 134 162 89 158 C-0.25 158 178 103 167 89 168 92 149 C-0.50 94 158 113 178 138 159 89 198 C-1.00 104 148 93 173 98 242 100 220 a Anodic Tafel slope βa in mV/decade, the measured Tafel regions are with 40~50 mV of overvoltage.

Alloys A(x, T), A/cm2 Ea, kJ/mol 20 oC (293 K) 35 oC (303 K) 50 oC (323 K) 65 oC (338 K) C-0 1.16 x 10-4 1.16 x 10-4 1.16 x 10-4 1.07 x 10-4 3.96 C-0.25 1.90 x 10-4 2.03 x 10-4 1.77 x 10-4 1.90 x 10-4 5.35 C-0.50 7.17 5.64 3.70 8.41 31.24 C-1.00 1.78 x 109 7.46 x 108 1.89 x 109 1.31 x 109 78.61 SS 304 1.18 x 10-4 1.28 x 10-4 1.70 x 10-4 2.07 x 10-4 9.87

b Cathodic Tafel slope βc in mV/decade, the measured Tafel regions are with 150~170 mV of

Table 4. The fits for A(x, T) and Ea(x) in Icorr(x, T) = A(x, T) exp(-Ea(x)/RT).

Table 3. Fit data for Tafel slopes of alloys C-x in 20 oC - 65 oC.

<sup>b</sup>βa βc βa βc βa β<sup>c</sup>

**3.2 Effect of temperature on potentiodynamic polarization** 

lower than 23oC.

βaa β<sup>c</sup>

Alloys C-x

overvoltage.

Fig. 3. Potentiodynamic Polarization Curve diagrams for (a) C-0, (b) C-0.25, (c) C-0.50, and (d) C-1.00 at various temperatures.

Fig. 4. The Arrhenius plots for alloys C-x and SS 304 at 20oC - 65oC.

Electrochemical Passive Properties of AlxCoCrFeNi

(x = 0, 0.25, 0.50, 1.00) High-Entropy Alloys in Sulfuric Acids 141

(b)

(c) Fig. 5. (a). The Nyquist plots for alloys C-x at 25oC. (b). The Bode plot for C-0.25. (c). The

Bode plot for C-1.00.

#### **3.3 EIS test at 25o C**

Figs. 5 and 6 summarize the EIS results of alloys in a sulfuric solution and their schematic equivalent circuit diagrams, respectively. Table 3 lists related parameters of the equivalent circuit diagrams, where Rs, Rf, and Rct denote impedances of the sulfuric solution, oxide layer, and adsorption layer, Qf and Qad denote capacitances of constant phase element (CPE) for oxide layer and adsorption layer, respectively. Next, the oxide layer thickness is evaluated by using the Helmholtz model [27] and expressing the layer thickness of the oxide layer, d, as d = εεoS/Qf, where εo denotes the permittivity of free space (8.85 x 10-14 F/cm), ε denotes the dielectric constant of the medium, and S denotes the surface area of the electrode. Assuming that ε and S for all oxide layers of alloys are the same allows us to compare relative values of d for all samples by 1/Qf. Fig. 7 reveals that 1/Qf values are proportional to x, implying that d increases with Al content x. However, according to this figure, the impedance of oxide layer Rf decreases with x and, in Fig. 8, the impedance of the oxide layer is inversely proportional to Ipass. Restated, a thinner oxide layer implies a larger value of impedance. To explain this phenomenon, besides the thickness of oxide layer, the density of oxide layer is also considered. As mentioned in Section 3.1, Al oxide easily forms a porous film on the metal surface [24]. Therefore, it is easily understood that in addition to causing a thicker oxide layer, Al element promotes the dispersive oxide layer. Combining these two effects obviously reveals that Rf decreases with x.

Figs. 5 and 6 summarize the EIS results of alloys in a sulfuric solution and their schematic equivalent circuit diagrams, respectively. Table 3 lists related parameters of the equivalent circuit diagrams, where Rs, Rf, and Rct denote impedances of the sulfuric solution, oxide layer, and adsorption layer, Qf and Qad denote capacitances of constant phase element (CPE) for oxide layer and adsorption layer, respectively. Next, the oxide layer thickness is evaluated by using the Helmholtz model [27] and expressing the layer thickness of the oxide layer, d, as d = εεoS/Qf, where εo denotes the permittivity of free space (8.85 x 10-14 F/cm), ε denotes the dielectric constant of the medium, and S denotes the surface area of the electrode. Assuming that ε and S for all oxide layers of alloys are the same allows us to compare relative values of d for all samples by 1/Qf. Fig. 7 reveals that 1/Qf values are proportional to x, implying that d increases with Al content x. However, according to this figure, the impedance of oxide layer Rf decreases with x and, in Fig. 8, the impedance of the oxide layer is inversely proportional to Ipass. Restated, a thinner oxide layer implies a larger value of impedance. To explain this phenomenon, besides the thickness of oxide layer, the density of oxide layer is also considered. As mentioned in Section 3.1, Al oxide easily forms a porous film on the metal surface [24]. Therefore, it is easily understood that in addition to causing a thicker oxide layer, Al element promotes the dispersive oxide layer. Combining

(a)

**3.3 EIS test at 25o**

**C** 

these two effects obviously reveals that Rf decreases with x.

Fig. 5. (a). The Nyquist plots for alloys C-x at 25oC. (b). The Bode plot for C-0.25. (c). The Bode plot for C-1.00.

Electrochemical Passive Properties of AlxCoCrFeNi

(x = 0, 0.25, 0.50, 1.00) High-Entropy Alloys in Sulfuric Acids 143

Fig. 7. Impedance and relative thickness (1/Qf) of oxide layer vs. Al content x plots.

Fig. 8. Impedance and Ipass of oxide layer vs. Al content x plots.

ability (Figs. 9 (b) to (d)).

**3.4 Polarization behaviour for alloys in a chloride-containing H2SO4 solution** 

Fig. 9 shows potentiodynamic polarization curve diagrams for the alloys in 0.5 M H2SO4 solution containing various concentrations of chloride ions, as well as in simple 0.5 M H2SO4 solution as a comparison. According to Fig. 9(a), oscillation occurs in a passive region for C-0 in 0.5 M H2SO4 containing 0.5 M and 1 M of chloride. This phenomenon has been attributed to the cycling process for small pitting and re-passivation with the duration of several seconds for each cycle [34]. Oscillation in the passive region in potentiodynamic polarization curve is a metastable state [35]. This metastable state generally reflects the difficulty of pitting, i.e., alloy C-0 has good anti-pitting ability, while those containing aluminum (C-0.25, 0.50, and 1.00) with no metastable state show an inferior anti-pitting

According to Fig. 6, only C-1.00 reveals a component of inductance in the equivalent circuit (See Fig. 5(c)). In previous studies [28-31], alloys with Al and Ni readily react with (OH)− and (SO4)2− in a sulfuric solution and adsorbed on the surface of the alloys, which increases the amount of the ions in the adsorption layer. Therefore, Qad increases with x, as listed in Table 5. As x value increases to 1.00, the inductance appears in the equivalent circuit in Fig. 6(b). This effect normally occurs in the case of a severe corrosive condition [32]. Origin of the inductance can generally be influenced by some adsorbed intermediates or can be attributed to a space at the interfaces [33]. In C-1.00, a microstructure with an Al and Ni-rich phase which is seen as a reactive phase from metallograph, not only causes adsorption in these Al and Ni-rich areas in corrosion process, but also decreases the impedance in the low frequency area owing to their continuous dissolution. The fact that Rct decreases with x demonstrates a higher dissolution rate for alloys with a higher Al content.


\* Lad = 122.4 Henry

Table 5. EIS equivalent circuit parameters for alloys C-x.

Fig. 6. EIS equivalent circuits for alloys C-0, C-0.25, C-0.50, and C-1.00.

According to Fig. 6, only C-1.00 reveals a component of inductance in the equivalent circuit (See Fig. 5(c)). In previous studies [28-31], alloys with Al and Ni readily react with (OH)− and (SO4)2− in a sulfuric solution and adsorbed on the surface of the alloys, which increases the amount of the ions in the adsorption layer. Therefore, Qad increases with x, as listed in Table 5. As x value increases to 1.00, the inductance appears in the equivalent circuit in Fig. 6(b). This effect normally occurs in the case of a severe corrosive condition [32]. Origin of the inductance can generally be influenced by some adsorbed intermediates or can be attributed to a space at the interfaces [33]. In C-1.00, a microstructure with an Al and Ni-rich phase which is seen as a reactive phase from metallograph, not only causes adsorption in these Al and Ni-rich areas in corrosion process, but also decreases the impedance in the low frequency area owing to their continuous dissolution. The fact that Rct decreases with x demonstrates a higher dissolution rate for alloys with a higher Al

nf Rf

C-0 3.271 54.57 0.9094 992.2 636.7 0.7444 7691 C-0.25 3.758 56.61 0.9081 610.5 1525 0.6347 1932 C-0.50 2.994 46.55 0.9223 642.8 3221 0.6454 819.1 C-1.00 3.462 47.16 0.9614 518.1 Lad - 66.81

( cm2)

Qad (μF/cm2) nad Rct

( cm2)

content.

Alloys C-x

\* Lad = 122.4 Henry

Rs ( cm2)

Qf (μF/cm2)

Table 5. EIS equivalent circuit parameters for alloys C-x.

Fig. 6. EIS equivalent circuits for alloys C-0, C-0.25, C-0.50, and C-1.00.

Fig. 7. Impedance and relative thickness (1/Qf) of oxide layer vs. Al content x plots.

Fig. 8. Impedance and Ipass of oxide layer vs. Al content x plots.

#### **3.4 Polarization behaviour for alloys in a chloride-containing H2SO4 solution**

Fig. 9 shows potentiodynamic polarization curve diagrams for the alloys in 0.5 M H2SO4 solution containing various concentrations of chloride ions, as well as in simple 0.5 M H2SO4 solution as a comparison. According to Fig. 9(a), oscillation occurs in a passive region for C-0 in 0.5 M H2SO4 containing 0.5 M and 1 M of chloride. This phenomenon has been attributed to the cycling process for small pitting and re-passivation with the duration of several seconds for each cycle [34]. Oscillation in the passive region in potentiodynamic polarization curve is a metastable state [35]. This metastable state generally reflects the difficulty of pitting, i.e., alloy C-0 has good anti-pitting ability, while those containing aluminum (C-0.25, 0.50, and 1.00) with no metastable state show an inferior anti-pitting ability (Figs. 9 (b) to (d)).

Electrochemical Passive Properties of AlxCoCrFeNi

**3.5 Metallographic examination and EDS analysis** 

Table 6 lists the EDS composition for each phase in different alloys.

Table 6. EDS analyses (at %) for alloys C-0, C-0.25, C-0.50, and C-1.00.

(x = 0, 0.25, 0.50, 1.00) High-Entropy Alloys in Sulfuric Acids 145

and C-1.00, decrease to 0.2-0.5 VSHE at a chloride concentration of 0.25 M (Fig. 10). A higher Al concentration in the alloys implies a lower value of Epit. For C-0, deterioration of the passive layer is attributed to the evolution of oxygen. Meanwhile, for C-0.25, C-0.50, and C-1.00, the deterioration of passive layer is attributed to the pitting process. An increasing chloride ion concentration causes the chloride ions to cluster at the defect sites of the passive layer and severely attack the passive layer. Consequently, Epit shifts to a more active region.

Fig. 10. Histogram of Epit for alloys C-x and SS 304 in solution of different Cl− ion molarity (M).

Microstructures for not H2SO4-immersed alloys C-0, C-0.25, C-0.50, and C-1.00 are with single FCC, single FCC, duplex FCC-BCC, and BCC-ordered BCC phases, respectively [6].

Alloys Phases and states Al Co Cr Fe Ni C-0 Overall, not immersed 0 25.93 25.73 24.21 24.13 Overall, immersed 0 24.45 26.39 24.83 24.32 C-0.25 Overall, not immersed 6.16 23.27 23.58 23.59 23.40 Overall, immersed 6.18 23.65 24.41 23.04 22.71 C-0.50 Overall, not immersed 11.01 22.77 22.61 21.70 21.92 FCC matrix, not immersed 8.36 24.74 23.48 22.77 20.65 BCC, not immersed **13.94** 21.11 20.48 20.53 **23.94**  Overall, immersed 8.35 22.75 27.19 23.60 18.11 FCC matrix, immersed 9.96 22.57 23.16 23.59 20.72 Wall-shaped BCC, immersed **3.82** 22.50 **36.33** 25.58 **11.77**  C-1.00 Overall, not immersed 18.88 20.55 20.63 20.01 19.93 Overall, immersed 12.45 19.80 29.87 23.42 14.45 BCC, immersed 17.14 20.67 21.96 20.89 19.34 Ordered BCC, immersed **3.04** 17.34 **47.53** 27.96 **4.14** 

Fig. 9. Potentiodynamic Polarization Curve diagrams for (a) C-0, (b) C-0.25, (c) C-0.50, and (d) C-1.00 at 25 oC in chloride-containing sulfuric acid solution at various Cl− molarity (M) values.

From an adsorption viewpoint, adsorption competition always prevails on the alloy surface between chloride ions and dissolved oxygen atoms. Notably, no oxide layer forms once chloride ions adsorb on the alloy surface, in which the metal ions readily dissolve. Therefore, the adsorption of chloride ions increases the reacting current density (as indicated by a comparison of Figs. 1 and 9), subsequently increasing the rate of metal dissolution.

Rapid dissolution of alloys in chloride-containing solution is discussed next. When chloride ions are adsorbed on the interface of passive layer and a sulfuric solution, metastable ion complexes gradually form from the anions of a passive layer. These metastable ion complexes enable the anions to dissolve. Once the ion complexes that are on the passive layer/solution interface dissolve into the sulfuric solution, the inner ion complexes of the passive layer move to the passive layer/solution interface in order to correlate with the applied potential. The inability of the anions to form oxide implies the continuous formation of metastable ion complexes and dissolution of ions. Since Al easily forms [Al(SO4 )]+ with (SO4)2-, and Al(OH)SO4 with (SO4)2- and (OH)−, respectively [36], these metastable ion complexes combine with Cl− and dissolve afterwards. Therefore, pitting easily occurs on the surface of aluminum alloys. Next, the aluminiferous passive layer and non-aluminiferous passive layer are compared. Fig. 10 shows the pitting potential (Epit) of the alloys and SS 304 in different solutions. The value of Epit for C-0 is almost independent of chloride concentration. The value of Epit for C-0.25 decreases abruptly for a chloride concentration exceeding 0.50 M. This value is close to that of SS 304. The values of Epit, for C-0.25, C-0.50,

Fig. 9. Potentiodynamic Polarization Curve diagrams for (a) C-0, (b) C-0.25, (c) C-0.50, and (d) C-1.00 at 25 oC in chloride-containing sulfuric acid solution at various Cl− molarity (M) values.

From an adsorption viewpoint, adsorption competition always prevails on the alloy surface between chloride ions and dissolved oxygen atoms. Notably, no oxide layer forms once chloride ions adsorb on the alloy surface, in which the metal ions readily dissolve. Therefore, the adsorption of chloride ions increases the reacting current density (as indicated by a

Rapid dissolution of alloys in chloride-containing solution is discussed next. When chloride ions are adsorbed on the interface of passive layer and a sulfuric solution, metastable ion complexes gradually form from the anions of a passive layer. These metastable ion complexes enable the anions to dissolve. Once the ion complexes that are on the passive layer/solution interface dissolve into the sulfuric solution, the inner ion complexes of the passive layer move to the passive layer/solution interface in order to correlate with the applied potential. The inability of the anions to form oxide implies the continuous formation of metastable ion complexes and dissolution of ions. Since Al easily forms [Al(SO4 )]+ with (SO4)2-, and Al(OH)SO4 with (SO4)2- and (OH)−, respectively [36], these metastable ion complexes combine with Cl− and dissolve afterwards. Therefore, pitting easily occurs on the surface of aluminum alloys. Next, the aluminiferous passive layer and non-aluminiferous passive layer are compared. Fig. 10 shows the pitting potential (Epit) of the alloys and SS 304 in different solutions. The value of Epit for C-0 is almost independent of chloride concentration. The value of Epit for C-0.25 decreases abruptly for a chloride concentration exceeding 0.50 M. This value is close to that of SS 304. The values of Epit, for C-0.25, C-0.50,

comparison of Figs. 1 and 9), subsequently increasing the rate of metal dissolution.

and C-1.00, decrease to 0.2-0.5 VSHE at a chloride concentration of 0.25 M (Fig. 10). A higher Al concentration in the alloys implies a lower value of Epit. For C-0, deterioration of the passive layer is attributed to the evolution of oxygen. Meanwhile, for C-0.25, C-0.50, and C-1.00, the deterioration of passive layer is attributed to the pitting process. An increasing chloride ion concentration causes the chloride ions to cluster at the defect sites of the passive layer and severely attack the passive layer. Consequently, Epit shifts to a more active region.

Fig. 10. Histogram of Epit for alloys C-x and SS 304 in solution of different Cl− ion molarity (M).
