Effect of Y on the microstructure and corrosion resistance of the as-cast Al-Mg-Mn alloys

Abstract

Rare earth can simultaneously enhance mechanical properties and corrosion resistance of Al-Mg-Mn alloys. Among them, yttrium (Y) possesses the advantages of lower cost and high crust abundance. Effect of Y content (0.1, 0.2, 0.3 wt.%) on microstructure and corrosion resistance of as-cast Al-Mg-Mn alloys was examined. With increasing Y content, the number of Al₃Mg₂ and Al₆(Fe, Mn) is gradually decreased, and the number of Y-containing precipitates is increased. After 0.3 wt.% Y addition, Al₆(Fe, Mn) disappears and the Y-containing phases dominate. Galvanic corrosion is reduced since the potential difference between the α-Al and Y-containing phases is lower than that of the Al₆(Fe, Mn) phase. Furthermore, after Y addition, volume fraction of Al₃Mg₂ is reduced, which improves intergranular corrosion resistance.

Keywords

Al-Mg-Mn alloys Yttrium Y-containing phases Galvanic corrosion.

Introduction

Al-Mg series aluminum alloys are widely used in aircraft (such as the fuselage skins and welded structural components), automobiles (such as bumper beams and energy-absorbing boxes), shipbuilding (such as structural components of hulls, yachts, warships, and offshore platforms), and architectural structures (such as pressure vessels, mechanical structural components) due to their excellent mechanical properties and corrosion resistance.^1–6 With the growing demand for high-performance aluminum alloys in fields such as aerospace and marine engineering, the strength and corrosion resistance of the conventional Al-Mg alloys are still unsatisfactory. In particular, when the Mg content is higher than 3.5 wt.%, the continuous β phase (Al₃Mg₂) tends to form along the grain boundaries. For Al-Mg series alloys, β phase will be preferentially dissolved in corrosive environments, consequently decreasing corrosion resistance and mechanical properties, which ultimately limits the applicability of Al-Mg alloys under harsh service conditions.^7,8

In recent years, microalloying with rare earth (RE) elements has been considered an effective method to modify the microstructure and enhance the properties of aluminum alloys. Among them, Sc is one of the most commonly used and effective microalloying rare earth elements in Al-Mg alloys. The addition of Sc promotes the formation of numerous fine Al₃Sc precipitates. Al₃Sc is coherent with the matrix and can serve as nucleation cores to promote grain refinement.^9–11 The addition of Sc can significantly enhance the mechanical properties of Al-Mg alloys by grain refinement strengthening and precipitation strengthening. Algendy et al.¹² found that the strength of 5454 alloy can be enhanced by approximately 50 MPa after adding 0.1 wt.% Sc. In addition, Sc addition can promote the formation of AlMn dispersoids in 5454 alloy. Al₃Sc precipitates and AlMn dispersoids can inhibit subgrain boundary migration and hinder dislocation motion, which leads to the improvement of the yield strength from 187 MPa to 239 MPa. Meanwhile, Sc also significantly affects the corrosion behavior of Al-Mg alloy. Liang et al.¹³ demonstrated that the intergranular corrosion resistance of the 5182 alloy can be significantly enhanced by the addition of 0.1 wt.% Sc. The decrease of intergranular corrosion depth is attributed to the inhibition of recrystallization due to the formation of Al₃(Sc, Zr) phase. According to the report of Gao et al.,¹⁴ the stress corrosion resistance of Al-Mg alloys can be enhanced with the addition of 0.5 wt.%Sc and 0.1 wt.%Zr, as evidenced by the decrease of the corrosion current density from 3.7 × 10⁻⁶ to 1.3 × 10⁻⁶ A/cm^-2. The main reason is that the grain refinement by the addition of Sc promotes the formation of a passivation film.

However, the high cost of Sc limits its application in aluminum alloys. Yttrium (Y), as a rare earth element in the same group as Sc, has the advantages of lower cost and higher chemical activity. Y can also form stable intermetallic compounds with Al (e.g., Al₃Y intermetallic), which can serve as nucleation cores, promote grain refinement, and reduce the size of harmful precipitates.^15,16 In Al-Si alloys, Y can significantly reduce the secondary dendrite arm spacing (SDAS) and modify the eutectic Si from flake or needle-like to fibrous or even spheroidized, which can significantly improve tensile properties.^17–19 In Al-Mg-Si alloys, the microalloying with Y refines the grains and inhibits the growth of the β-AlFeSi and Mg_xSi phases,²⁰ which is beneficial to improve the mechanical properties and corrosion resistance. Wang et al.²¹ demonstrated that I_corr, E_corr-E_prot, E_pit-E_corr, and E_ptp (I_corr is corrosion current density; E_corr is corrosion potential；E_prot is repassivation/protection potential; E_pit is pitting potential and E_ptp is pit transition potential) in the cyclic polarization tests were significantly improved after the addition of Y. For Al-Zn-Mg alloys, Y addition promotes the formation of precipitates such as Al₈Cu₄Y and (Al, Cu)₁₁Y₃, and accelerated the dissolution of the Mg₂Si phase,^22,23 which improved the mechanical properties and corrosion resistance. Ru et al.²⁴ found that the L1₂-type Al₃Y phase formed after the addition of Y and Zr in Al-Mg alloys, which inhibited the enrichment of Mg atoms at grain boundaries and hindered the precipitation and coarsening of the β phase along the grain boundaries, thereby reducing the susceptibility of intergranular corrosion (IGC).

In summary, Y has demonstrated considerable strengthening effects in various aluminum alloy systems. However, for Al-Mg-Mn alloys, only the effect of Y on the IGC has been investigated. The effect of Y on microstructure and corrosion resistance of the as-cast Al-Mg-Mn alloys needs further investigation. In this work, corrosion resistance was evaluated using a combination of potentiodynamic polarization, salt spray, weight loss tests, X-ray photoelectron spectroscopy (XPS), and Kelvin probe force microscopy (KPFM). The role of Y in regulating the microstructure, mechanical properties and corrosion behavior of Al-Mg-Mn alloy was discussed, which provides a theoretical foundation and technical support for the performance optimization of Al-Mg-Mn alloys toward practical applications.

Materials and experimental procedure

In order to facilitate the understanding of the layout of the present work, the experimental design and technical flowchart are presented. Figure 1 summarizes the entire process, including alloy preparation, microstructure analysis (EBSD, DSC, EMPA and SEM), corrosion resistance evaluation (electrochemical tests, salt spray tests, weight loss tests, and KPFM tests), and mechanism analysis.

Figure 1.

Schematic diagram of the experimental procedure in this work.

Sample preparation

According to the work of Ru et al.,²⁴ after the addition of 0.1 wt.% Y and 0.1 wt.% Zr, the grains of the Al-Mg-Mn alloy were refined and the IGC resistance was enhanced. Additionally, Wei et al.¹⁷ designed a Y content gradient of 0.1, 0.2 and 0.3 wt.% in the A356 alloy. It was found that when the Y content was 0.2 wt.%, the dendrites were markedly refined and mechanical properties was significantly enhanced. Thus, in this work, we designed the Y content gradient as 0.1, 0.2 and 0.3 wt.%. The alloys used in this work were prepared by using Al-5.8Mg-0.7Mn alloy, Al-10Y master alloys, Al-10Mn master alloys, pure Mg, and pure Zn in a pit furnace (JSL 1400) to produce as-cast Al-Mg-Mn-xY alloys (x = 0, 0.1, 0.2, 0.3 wt.%). The detailed steps are as follows: Firstly, the Al-5.8Mg-0.7Mn alloy was placed in a graphene crucible and heated to 760 °C in the pit furnace, and then held for 5-10 minutes to ensure complete melting of the alloy. Subsequently, Al-10Y and Al-10Mn master alloys were added to the melt, and after complete melting the mixture was thoroughly stirred to ensure uniform composition distribution. Then, the first skimming was performed to remove the oxidized dross on the melt surface, after which the temperature was raised to 760 °C and held for 5 minutes. Pure Mg and pure Zn were wrapped in aluminum foil and placed at the bottom of the melt using crucible tongs, followed by thorough stirring. Then, refining agents (0.2% of the total weight) were added to the melt, and the mixture was stirred thoroughly before undergoing the second skimming. The temperature was maintained at 760 °C for 10 min to ensure complete melting and thorough mixing. Finally, the melt temperature was lowered to 710 °C, and the melt was poured into a preheated graphene mold (200 °C) to form rectangular ingots. The chemical composition of the as-cast aluminum alloy ingots was determined by Inductively Coupled Plasma Atomic Emission Spectrometry (ICP-AES), and the compositions were listed in Table 1.

Table 1.

Chemical compositions of the four Al-Mg-Mn as-cast alloys (wt.%).

Alloys	Mg	Mn	Zn	Fe	Si	Y	Al
Al-Mg-Mn	5.83	0.77	0.65	0.13	0.03	-	-
Al-Mg-Mn-0.1Y	5.90	0.75	0.57	0.12	0.03	0.09	-
Al-Mg-Mn-0.2Y	5.88	0.86	0.54	0.11	0.03	0.21	-
Al-Mg-Mn-0.3Y	5.84	0.72	0.56	0.12	0.03	0.33	-

Microstructural characterization

The effect of Y content on grain size was investigated by electron backscatter diffraction (EBSD, ZEISS Sigma 300, Germany). EBSD specimens were prepared by mechanical grinding and electropolishing, with an acceleration voltage of 20 kV, a sample tilt of 70°, and a scanning step size of 4 μm. The thermal behavior during the cooling process was evaluated using thermogravimetric-differential scanning calorimetry (TG-DSC, Netzsch STA409PC, Germany). Prior to the TG-DSC tests, the samples were dried to eliminate surface moisture. During the TG-DSC tests, cylindrical specimens (Φ3 × 2 mm) were placed in a crucible under a flowing argon atmosphere to prevent oxidation. The cooling rate during the TG-DSC tests was set as 5 °C/min, and the tested temperature ranged from 300 to 700 °C.

To determine the composition of the precipitates in the alloys, the chemical composition of the precipitated phases was determined using an electron probe microanalyzer (EPMA, JEOL JXA-iHP200F, Japan). Statistics of the number of the precipitates in the alloys were performed using a scan electron microscope (SEM, ZEISS Sigma 360, Germany).

Potentiodynamic polarization tests

For potentiodynamic polarization tests, specimens with dimensions of 10 × 10 × 3 mm were used. The specimens were ultrasonically cleaned in a solution of anhydrous ethanol and detergent to remove oil, grease, and other impurities. The specimens were then cold-mounted in epoxy resin and subjected to mechanical grinding and polishing after solidification. Potentiodynamic polarization tests were conducted using a CHI760E electrochemical workstation in a 3.5 wt.% NaCl solution at room temperature with a scan range of –0.5 V to 0.5 V (vs. SCE) and a scan rate of 0.00167 V/s. All the measurements were carried out inside a Faraday cage to minimize external electromagnetic interference. A standard three-electrode system was used, where the sample served as the working electrode, the platinum electrode as the counter electrode, and the saturated calomel electrode (SCE) as the reference electrode. Before the tests, the samples were soaked in solution for about 20 h, and then the open circuit potential (OCP) for 600 s to ensure that a relatively stable state can be reached before potentiodynamic polarization tests. The key electrochemical parameters such as corrosion current density (i_corr), corrosion potential (E_corr) and polarization resistance (R_P) were derived by fitting the data using instrument software (CHI760E). To ensure data reliability, tests were conducted on three replicate specimens for each alloy composition.

Salt fog tests

Salt spray tests were conducted in a small neutral salt spray chamber according to GB/T 10125-2021, using a 3.5 wt.% NaCl solution as the corrosive medium. The tested temperature was controlled at 35 ± 1 °C, and the spray method was continuous spray. The sample size was 10 × 10 × 3 mm, and the exposed surface was polished after being ground with 400#, 800#, 1500#, 3000#, and 4000# sandpapers, then cleaned with deionized water and anhydrous ethanol and dried. The samples were placed on the sample holder at an inclination angle of 45° to ensure uniform coverage of the exposed surface by the spray, and were periodically taken out for overall corrosion observation using an optical microscope (OM). To examine the change in corrosion progress over time, corrosion durations of 0 h, 24 h, and 48 h were set. After 48 hours of salt spray testing, the corrosion products on the sample surfaces were analyzed by X-ray photoelectron spectroscopy (XPS, Thermo Scientific K-Alpha, USA). Using a monochromatic Al Kα X-ray source, all the peaks were calibrated using C1 s peak with 284.8 eV binding energy at 100 W under vacuum condition, with an energy and steps of 50 and 0.1 eV, respectively. The Avantage software was used to fit the response peaks.

Weight loss tests

Prior to the weight loss measurements, the specimens were ultrasonically cleaned and dried to remove surface oils and particulate debris. After cleaning, the specimens were dried with a stream of warm air and stored for subsequent tests. The overall corrosion rate was determined using the weight loss method. A pretreatment, including sequential immersion in a 50 g/L NaOH solution at 80 °C for 1 min, followed by rinsing and drying, was applied to the specimens. Subsequently, the specimens were immersed in a concentrated HNO₃ solution (1.4 g/ml) for 30 s to achieve a bright surface, before a final rinse with distilled water and drying. The corrosive medium was a HNO₃ solution, the corrosion temperature was controlled at 30.0 ± 0.1 °C, and the corrosion time was 24 h. The initial weight (W₁) and final weight (W₂) of the specimens were measured using a precision electronic balance with an accuracy of 0.01 g. The total exposed surface area for corrosion was 3.2 cm².

KPFM characterization

Atomic force microscopy (AFM) was used to characterize the three-dimensional morphologies of microstructures. The Volta potential, measured by Kelvin probe force microscopy (KPFM), serve as a key indicator for evaluating the corrosion tendency of metallic materials. To investigate the influence of different precipitates on the corrosion resistance of the Al-Mg-Mn alloys, the Volta potentials of the Al₃Mg₂, Al₆(Fe, Mn), and Y-containing phase were measured using KPFM. Prior to the KPFM measurements, the specimens were ground, polished, rinsed with anhydrous ethanol, and dried. All the measurements were conducted at room temperature with a scan resolution of 256 × 256 pixels and a scan rate of 1.0 Hz.

Results and discussion

Microstructure

To further confirm the effect of Y on the grain size of the as-cast Al-Mg-Mn alloys with different Y contents, electron backscatter diffraction (EBSD) was employed. Figure 2 shows that the grain sizes of the four alloys with 0, 0.1, 0.2, and 0.3 wt.% Y addition are 222.6 μm, 243.5 μm, 236.1 μm, and 307.8 μm, respectively. Thus, within the range of 0∼0.3 wt.%, the addition of Y promotes grains coarsening for the as-cast Al-Mg-Mn alloys. The coarsening of grain size is closely related to the degree of undercooling during solidification

Figure 2.

IPF-Z orientation maps of the as-cast Al-Mg-Mn alloys with different Y contents. (a)Al-Mg-Mn; (b) Al-Mg-Mn-0.1Y; (c) Al-Mg-Mn-0.2Y; (d) Al-Mg-Mn-0.3Y.

To determine the degree of undercooling, the TG-DSC cooling tests were conducted on the alloys with different Y contents. The TG-DSC results of the alloys with different Y additions are shown in Figure 3. TG-DSC curves exhibit only one α-Al exothermic peak, with no other peaks detected, as shown in Figure 3. This indicates that the phase transformation of α-Al dominates the majority of the thermal effects during solidification. It is known that nucleation of α-Al from the liquid requires a certain degree of undercooling (ΔT), which can be expressed as

\begin{matrix} Δ T = T_{m} - T_{s} \end{matrix}

(1)

where T_m represents the theoretical solidification temperature and T_s denotes the actual solidification temperature. For Al-5.8Mg-0.7Mn alloy, T_m is approximately 640 °C. As the content of Y element increases, the actual solidification temperature of the Al-Mg-Mn alloy slightly rises. The calculated degrees of undercooling ΔT is 9.2, 8.6, 7.1, and 6.6 °C, for the as-cast Al-Mg-Mn alloys with 0, 0.1, 0.2, and 0.3 wt.% Y addition, respectively. Commonly, the increase of undercooling can enhance the nucleation rate during solidification, which is key to obtaining fine equiaxial grains. The results of DSC show that the undercooling of the alloys is gradually reduced and the grain size is coarsened with the increase of Y content, which is contrary to the commonly reported conclusion that rare earth element can refine the grains. Y, as a rare earth with high chemical activity, has a large atomic size and trends to aggregate at the solid-liquid interface, which may slow down the grain growth rate due to solute dragging effect, thereby increasing constitutional undercooling. However, the decrease of the measured values of undercooling indicates that heterogeneous nucleation behavior in the solidification process is dominant. The decrease of undercooling means that the critical nucleation work increases during solidification, resulting in difficult formation of heterogeneous nucleation cores and nucleation rate is decreased, ultimately resulting in coarsened grain size.

Figure 3.

DSC cooling curves of the as-cast Al-Mg-Mn-xY alloys (x = 0, 0.1, 0.2, 0.3 wt.%).

To further elucidate the effect of Y content on the precipitates in the as-cast Al-Mg-Mn alloys, Figure 4 presents the morphology and distribution of the precipitates in the alloys with different Y contents. Table 2 lists the corresponding chemical compositions of these precipitates. By analyzing the compositions of multiple representative precipitates (A-G), the influence of Y on the types and compositional evolution of precipitates in the alloy can be discussed. Previous studies²⁵ indicated that β phase (Al₃Mg₂) and network-like Al-Mn-Fe intermetallic compounds are commonly present in the as-cast Al-Mg-Mn alloys, with Zn enrichment often observed within the β phase. Figure 4 indicates that phase A is enriched in Al, Mg, and Zn. According to the Al-Mg-Zn ternary phase diagram, no stable ternary phase with a comparable atomic ratio is documented. Therefore, this region is likely a phase formed by Zn dissolving into an Al-Mg phase. Given that the atomic radius of Zn (1.33 Å) is close to those of Al (1.43 Å) and Mg (1.60 Å), Zn atoms can dissolve into the Al-Mg phase via substitutional solid solution. The results of EPMA further show that the sum of atomic percentage ratio of Al and Zn to that of Mg is approximately 3:2, which is consistent with the stoichiometry of the β phase and also indicates that Zn partially replaces Al in the phase.²⁶ Therefore, this region can be identified as a (Al, Zn)₃Mg₂-type phase with a grayish-white color, and usually precipitated at grain boundaries.²⁷ In contrast, phase B exhibits a skeleton-like or blocky morphology. Its atomic percentage ratio of Al to (Fe + Mn) is approximately 4.7:1, which is close to the composition of the Al₆(Fe, Mn) phase. Thus, phase B can be determined as a typical solid-solution intermetallic compound, the Al₆(Fe, Mn) phase, which presents as a bright white precipitate. During the solidification of Al-Mg-Mn alloys, metastable α-Al solid solutions with varying Mg solubility typically form, and these can subsequently transform into β-phase precipitates during heating.²⁸ Phase C is primarily composed of Al, Fe, Mn and a small amount of Mg. Research work has shown that in aluminum alloys containing Fe and Mn, phases with similar elemental compositions are identified as the Al₆(Fe, Mn) phase.²⁹ Consequently, this phase is confirmed to be Al₆(Fe, Mn) phase.

Figure 4.

SEM images of the precipitates and determined by EPMA in the as-cast Al-Mg-Mn alloys without and with Y addition. (a) - (b) Y-free alloy; (c) - (e) Y-modified alloys.

Table 2.

Composition of the precipitates in the alloys with different Y contents (at%).

Precipitates	Al	Mg	Mn	Fe	Zn	Si	Y
A	58.23	37.93	-	-	3.95	-	-
B	82.44	-	8.84	8.34	-	-	-
C	82.99	3.81	8.60	4.48	-	-	-
D	57.86	37.35	-	-	4.71	-	-
E	73.73	13.64	6.96	2.62	-	1.63	1.38
F	86.89	4.89	2.81	2.54	-	-	2.72
G	86.87	8.13	1.66	0.76	0.59	-	1.87

In alloys with difference Y contents (0.1, 0.2, and 0.3 wt.%), the composition of phase D is similar to that of phase A, which is the β phase (Al₃Mg₂ phase with dissolved Zn). In addition, phase E is enriched in Al (73.73%), Mg (13.64%), Mn (6.96%), Fe (2.62%), Si (1.63%), and Y (1.38%). Phase F is similar to phase E, but contain no Si. Due to the strong segregation tendency of Y during solidification in the alloys, it is more likely to form precipitated phases with other elements. Therefore, the aforementioned phases E, F, and G are all referred as Y-containing multi-component phases.

Furthermore, the variation in the number and size of Al₃Mg₂ phase with Y content is shown in Figure 5. Statistical analyses were performed using 3-5 images for each alloy to evaluate the number of Al₃Mg₂ phase to ensure the reliability of the data. The number of Al₃Mg₂ phase in the Figure 5 represents the median of the statistical counts. In the Y-free alloy (Figure 5(a)), the Al₃Mg₂ phase is densely distributed, with a counted number of 26 particles and exhibiting a size predominantly ranging from several tens to one hundred micrometers. After the addition of Y, both the number and size of the Al₃Mg₂ phase are decreased. As shown in Figure 5(b) to (d), the number of Al₃Mg₂ phase is reduced to 22, 13, and 9, respectively. When the addition of Y content is 0.1 wt.%, the volume fraction and size of Al₃Mg₂ phase almost does not exhibit a significant decrease. When the Y addition comes to 0.2 wt.%, the number of Al₃Mg₂ phase is markedly decreased, and its size is not significantly decreased. When Y addition is further increased to 0.3 wt.%, the number of larger Al₃Mg₂ phase diminishes significantly, and the size is significantly reduced.

Figure 5.

SEM images of Al₃Mg₂ phase distribution in the as-cast Al-Mg-Mn alloys with different Y contents. Red circle: Al₃Mg₂ phase; Blue circle: Al₆(Fe, Mn) phase; Yellow circle: Y-containing phase. (a) Al-Mg-Mn; (b) Al-Mg-Mn-0.1Y; (c) Al-Mg-Mn-0.2Y; (d) Al-Mg-Mn-0.3Y.

In addition, the number and size of Al₆(Fe, Mn) phase are also decreased with the increase of Y content. When no Y is added, the number of Al₆(Fe, Mn) phase is 27 as shown in Figure 5(a), and the size of Al₆(Fe, Mn) phase is relatively large about a few tens of microns. For the alloys with 0.1, 0.2, and 0.3 wt.% Y addition, the number of Al₆(Fe, Mn) phase are gradually decreased to 26, 18 and 7, respectively, and the size is reduced, as shown in Figure 5(b) to (d). Meanwhile, for the alloys with 0.1, 0.2, and 0.3 wt.% Y addition, the number of Y-containing phase (bright white precipitated phase) is increased to about 7, 39, and 50, respectively.

In conclusion, the above results demonstrate that the addition of Y changes the solidification path of the alloy, so that Al₆(Fe, Mn) phase and Al₃Mg₂ phase translate to the more stable Y-containing phase. Y, as a rare earth with high chemical activity and strong tendency to segregate, has a strong chemical affinity with Mg, Fe, and Si elements.^30,31 During solidification, Y is preferentially enriched in the liquid phase at the solid-liquid interface, and preferentially reacts with impurity elements such as Fe and Si to form more stable Y-containing metallic compounds more stable than Al₆(Fe, Mn) phase and Al₃Mg₂ phase.³² Therefore, the formed Y-containing phases consume the Mg, Fe, and Si elements, thereby decreasing the number of Al₆(Fe, Mn) and Al₃Mg₂ phases.

Effect of Y content on the corrosion resistance

Figure 6 presents the potentiodynamic polarization curves of as-cast Al-Mg-Mn-xY alloys (x = 0, 0.1, 0.2, and 0.3) in the 3.5 wt.% NaCl solution, with the corresponding electrochemical parameters listed in Table 3. The corrosion current density (i_corr) reflects the corrosion rate, and a higher i_corr indicates more severe corrosion. The corrosion potential (E_corr) evaluates the corrosion susceptibility, with more negative E_corr values suggesting higher corrosion susceptibility. Meanwhile, a higher polarization resistance (R_p) signifies better corrosion resistance. As observed in Figure 6, the polarization curves of the four alloys exhibit similar shapes, and none show distinct passivation plateaus. Thus, the as-cast Al-Mg-Mn alloys do not form a dense passivation film in 3.5 wt.% NaCl solution. With the increase of Y content, the i_corr is first increased from 2.633 × 10⁻⁶A·cm^-2 to 4.716 × 10⁻⁶A·cm^-2 and then is decreased to 1.597 × 10⁻⁷A·cm^-2. When the Y content is 0.1 wt.%, the increase in i_corr may be due to the increase in the number of phases as cathodes. In contrast, the sharp decrease in i_corr in the alloy with 0.3 wt.% Y indicates significant suppression of corrosion activity. Furthermore, the E_corr reaches its most negative value at 0.2 wt.% Y, indicating the highest corrosion susceptibility, while the alloy with 0.3 wt.% Y exhibits the most positive E_corr. Combined with its lowest i_corr and highest R_p, the alloy with 0.3 wt.% Y demonstrates the best overall corrosion resistance among the tested four alloys.

Figure 6.

Potentiodynamic polarization curves of the as-cast Al-Mg-Mn-xY (x = 0, 0.1, 0.2, 0.3) alloys in 3.5 wt.% NaCl solution.

Table 3.

Electrochemical parameters derived from the potentiodynamic polarization curves.

Alloys	i_corr (A·cm^-2)	E_corr (VSCE)	R_p (Ω·cm²)	Tafel slope / (mV/dec)
Alloys	i_corr (A·cm^-2)	E_corr (VSCE)	R_p (Ω·cm²)	Cathodic	Anodic
Al-Mg-Mn	2.633 × 10−6	-0.964	1.27 × 104	161.26	146.58
Al-Mg-Mn-0.1Y	4.716 × 10−6	-0.921	0.56 × 104	195.69	87.32
Al-Mg-Mn-0.2Y	2.676 × 10−6	-1.02	1.36 × 104	161.29	174.16
Al-Mg-Mn-0.3Y	1.597 × 10−7	-0.783	1.14 × 105	208.72	52.301

In addition, according to the Tafel slope analysis in Table 3, the corrosion mechanism of Al-Mg-Mn-xY alloy can be further elucidated. For Al-Mg-Mn alloys, the anodic Tafel slope is 146.6 mV/dec and the cathodic Tafel slope is 161.3 mV/dec, indicating that the kinetics of anodic dissolution and cathodic reduction reaction are relatively balanced during corrosion. When the Y content is 0.1 wt.%, the anodic Tafel slope is significantly decreased to 87.3 mV/dec, indicating that the anodic dissolution process is accelerated, consistent with the observed the increase of i_corr. Although the cathodic Tafel slope is increased to 195.7 mV/dec, implying some inhibition of the cathodic reaction, the overall corrosion rate is still dominated by the enhanced corrosion process, which may be attributed to the formation of additional galvanic cells. When the Y content reaches 0.2 wt.%, the anodic Tafel slope is increased to 174.2 mV/dec, indicating that anodic dissolution becomes more difficult, while the cathodic Tafel slope remains almost constant (161.3 mV/dec). The corrosion resistance is improved compared with Al-Mg-Mn-0.1Y alloy. Notably, for the alloy containing 0.3 wt.% Y, the anodic Tafel slope drops sharply to 52.3 mV/dec, indicating a significant acceleration of anodic dissolution kinetics. However, the cathodic Tafel slope is increased to 208.7 mV/dec simultaneously, the highest value among all the tested alloys, indicating that the cathodic reaction is strongly inhibited. Thus, despite the enhanced anodic activity, the severely inhibited cathodic reaction causes the i_corr to drop sharply to 1.597 × 10⁻⁷ A·cm^-2, the lowest value among the four alloys. In summary, the excellent corrosion resistance of the Al-Mg-Mn-0.3Y alloy is mainly attributed to the significant inhibition of cathodic reaction, rather than the formation of passivation film or the inhibition of anodic dissolution.

Figure 7 displays the evolution of corrosion morphologies of as-cast Al-Mg-Mn alloys with different Y contents under various salt spray exposure times (top to bottom: 0, 24 and 48 hours). With the increase of salt spray time, the corrosion degree of the alloys with different Y contents is aggravated, and the corrosion products are increased correspondingly.

Figure 7.

Optical micrographs of the as-cast Al-Mg-Mn alloys with different Y contents after salt spray exposure. Rows: exposure time for 0, 24, and 48 h, Columns: (a1) ∼ (a3) Al-Mg-Mn; (b1) ∼ (b3) Al-Mg-Mn-0.1Y; (c1) ∼ (c3) Al-Mg-Mn-0.2Y; (d1) ∼ (d3) Al-Mg-Mn-0.3Y.

In the Y-free alloy (Figure 7a1∼a3), until 24 h (Figure 7 a2), the pitting corrosion is formed, and the corrosion products are gradually increased and adhered to the surface of the alloy. When the salt spray time comes to 48 h (Figure 7 a3), a large number of pitting corrosion pits and corrosion products appear on the surface, and the size of some pitting pits is increased. Combined with the above polarization curves, the alloy does not form a dense passivation film, and Cl^- can cause corrosion to continue. For the Y-containing alloys, when the salt spray time is 24 h, pitting and corrosion products of varying severity are also formed. For the alloy with 0.1 wt.% Y (Figure 7 b2), the pitting corrosion is gradually intensified with the increase of exposure time, and the performance is more sever than the yttrium-free alloys, which is consistent with the i_corr in the potentiodynamic polarization curves. However, when the Y content is increased to 0.2 wt.%, after exposure to the corrosive medium for 24 h (Figure 7 c2), the number of the pitting corrosion pits on the surface is higher than that the alloy with 0.1 wt.% Y, which may be related to the most negative result of E_corr in the polarization curve. This will be discussed according to the XPS results. After 48 h exposure as shown in (Figure 7 c3), the size of the pits is increased. It is worth noting that when the Y content is further increased to 0.3 wt.%, there are only a few small pitting appear on the surface as the salt spray time increases (Figure 7 d2). When the salt spray time is increased to 48 h (Figure 7 d3), although there are corrosion products, a large area of corrosion products cannot be observed, indicating that the pitting corrosion is significantly inhibited.

In order to comprehensively evaluate the effect of Y on the corrosion behavior of as-cast Al-Mg-Mn alloys, the weight loss experiments were performed. The weight loss reflects the intergranular corrosion rate of a material in a certain corrosive medium. Figure 8 shows the weight loss results of Al-Mg-Mn alloys with different Y contents, and the detailed data are shown in Table 4. The weight loss per unit area W was calculated by the following equation

\begin{matrix} W = \frac{Δ W}{A} = \frac{W_{1} - W_{2}}{A} \end{matrix}

(2)

where

W_{1}

and

W_{2}

are the initial weight and final weight of the specimens (mg), A is the exposed corrosion area (3.2 cm²), W is the weight loss per unit area (mg/cm²), which is a direct measure of the corrosion rate. Average weight loss per unit area

\overset{W}{↼}

is calculated from three replicates and used as the final average corrosion rate for each alloy to ensure the reliability of the experiments.

Figure 8.

Weight loss curve in as-cast Al-Mg-Mn alloys with different Y contents.

Table 4.

Experimentally measured weight loss of the alloys.

Alloys	W₁/mg	W₂/mg	Mass loss W/(mg/cm²)	Average mass loss( $\bar{W}$ /(mg/cm²))
Al-Mg-Mn	729.5	704.2	7.9062	10.6979
	736.3	698.8	11.7188
	735.6	695.7	12.4688
Al-Mg-Mn-0.1Y	691.9	649.5	13.2500	12.1563
	757.9	731.5	8.2500
	679.7	631 .8	14.9688
Al-Mg-Mn-0.2Y	762.6	743.8	5.8750	6.4896
	756.3	736.9	6.0625
	757.5	733.4	7.5312
Al-Mg-Mn-0.3Y	767.8	752.8	4.6875	3.9896
	769.6	754.5	4.7187
	760.5	752.3	2.5625

As shown in Figure 8, the average weight loss per unit area $\overset{W}{↼}$ for the Al-Mg-Mn, Al-Mg-Mn-0.1Y, Al-Mg-Mn-0.2Y, and Al-Mg-Mn-0.3Y alloys are 10.6979 mg/cm², 12.1563 mg/cm², 6.4896 mg/cm², and 3.9896 mg/cm², respectively. When the Y content is 0.1 wt.%, the weight loss $\overset{W}{↼}$ has increased. However, as the Y content is increased to 0.2 wt.% and 0.3 wt.%, $\overset{W}{↼}$ is decreased significantly. In particular, the Al-Mg-Mn-0.3Y alloy exhibits the lowest $\overset{W}{↼}$ of only 3.9896 mg/cm², indicating that the addition of Y effectively reduces the corrosion rate of the Al-Mg-Mn alloy. This result is consistent with the potentiodynamic polarization tests.

It has been reported that coarse grain leads to the reduction of grain boundary density in the non-passivating environment, which is beneficial to improving corrosion resistance.^33,34 When the addition of Y content is 0.1 wt.%, the number of grain boundaries due to grain coarsening is decreased without passivation film protection, while the volume fraction of Al₃Mg₂ phase does not decrease significantly, which will increase the density of the Al₃Mg₂ phase at the grain boundary, making it easier for corrosion to occur at the grain boundaries to result in an increase in the corrosion rate between grain boundaries. When the Y content is increased to 0.2 and 0.3 wt.%, the intergranular corrosion rate is gradually decreased due to the significant decrease of the Al₃Mg₂ phase and the decrease in the number of grain boundaries due to grain coarsening.

To investigate the effect mechanism of Y addition on the corrosion behavior of the as-cast Al-Mg-Mn alloys, Figures 9 to 11 present the morphological characteristics and contact potential difference (V_CPD) of Al₃Mg₂ phase, Al₆(Fe, Mn) phase, and Y-containing phase. The contact potential difference between the sample and the tip can be directly measured using the KPFM technique.^35–37 The measured contact potential difference between the sample and the tip is defined by the following equation (3)

\begin{matrix} V_{CPD} = \frac{Φ_{tip} - Φ_{sample}}{- e} \end{matrix}

(3)

where Φ_tip and Φ_sample represent the work functions of the tip and the sample, respectively, and e denotes the elementary charge. The work function is defined as the minimum energy required for electrons to escape from the surface of a solid material. This parameter is regarded as an important indicator for evaluating the corrosion resistance of metallic materials. Generally, metallic materials with higher work functions have lower corrosion activity.^35–37 In this study, the work function and contact potential difference between the three phases and α-Al matrix were obtained by KPFM tests. Since the work function of probe (Φ_tip) is known, the work function of the sample (Φ_sample) can be calculated using Equation (3). According to Equation (3), Φ_sample is positively correlated with the measured V_CPD. Therefore, higher V_CPD values correspond to higher electrode potential and better corrosion resistance.

Figure 9.

(a) AFM morphology of Al₃Mg₂ phase; (b) KPFM image; (c) Line profile taken from the line on (b) across Al₃Mg₂ phase and the surrounding α-Al matrix.

Figure 10.

(a) AFM morphology of Al₆(Fe, Mn) phase; (b) KPFM image; (c) Line profile taken from the line on (b) across Al₆(Fe, Mn) phase and the surrounding α-Al matrix.

Figure 11.

(a) AFM morphology of Y-containing phase; (b) KPFM image; (c) Line profile taken from the line on (b) across Y-containing phase and the surrounding α-Al matrix.

Figure 9 shows that the Al₃Mg₂ phase is approximately 95.9 nm higher than α-Al matrix, and its V_CPD is lower than that of the surrounding α-Al matrix, which means its electrode potential is lower than that of α-Al matrix. Therefore, the Al₃Mg₂ phase will be preferentially dissolved as the anodic phase during galvanic corrosion. The heights of the Al₆(Fe, Mn) phase and the Y-containing phase are about 272.0 nm and 300 nm, respectively. The V_CPD of both Al₆(Fe, Mn) phase and the Y-containing phase is higher than α-Al matrix (Figures 10 and 11). The potential difference between the Al₆(Fe, Mn) phase and α-Al matrix is about 240.78 mV, which is higher than 137.02 mV of the Y-containing phase.

To characterize the corrosion products generated during the corrosion process, the specimens after 48 h of salt spray exposure were analyzed by X-ray Photoelectron Spectroscopy (XPS). During corrosion, these phases act as cathodes, leading to preferential dissolution of the surrounding α-Al matrix and the formation of corrosion products such as Al₂O₃ and Al(OH)₃ around the precipitates (Figure 12). Moreover, the corrosion products after salt spray tests are primarily loose and porous Al(OH)₃. Since no distinct passivation region can be observed in the polarization curves, while XPS analysis (Figure 12) confirms the presence of Al₂O₃ and Al(OH)₃. Thus, it can be inferred that the corrosion products are only formed around the precipitated phases and are relatively loose. The galvanic corrosion around the Al₆(Fe, Mn) phase is more severe than that around the Y-containing multi-component phases. The loose and porous corrosion products generated in these areas fail to effectively inhibit the galvanic corrosion caused by the Al₆(Fe, Mn) phase.³⁸ Thus, after the addition of Y, the galvanic corrosion can be alleviated due to the formation of the Y-containing phases with lower potential difference with α-Al matrix.

Figure 12.

XPS spectra of Al (a, c, e, g) and O (b, d, f, h) elements in the corrosion products on the Al-Mg-Mn alloys with different Y contents. (a)∼(b) Al-Mg-Mn; (c)∼(d) Al-Mg-Mn-0.1Y; (e)∼(f) Al-Mg-Mn-0.2Y; (g)∼(h) Al-Mg-Mn-0.3Y.

Since the relative ratio of Al₂O₃ and Al(OH)₃ in the corrosion products can directly reflect the density of the passivation film, the quantitative analysis of the relative content of Al₂O₃ or Al(OH)₃ by XPS can help to elucidate the effect mechanism of Y addition on the corrosion resistance of the as-cast Al-Mg-Mn alloys. The corresponding peak areas of Al₂O₃ and Al(OH)₃ in Figure 12 were fitted by Origin software. Since the relative sensitivity factor (RSF) of Al₂O₃ and Al(OH)₃ is the same, the relative contents of Al₂O₃ and Al(OH)₃ in the corrosion products can be calculated by the following equation (4)

\begin{matrix} Relative conten t_{p} (%) = \frac{{Area}_{p}}{{Area}_{{Al}_{2} O_{3}} - {Area}_{Al (OH)_{3}}} \times 100 % \end{matrix}

(4)

where p is Al₂O₃ or Al(OH)₃, and Area_p is the corresponding peak area. Table 5 shows that when the Y content is 0.1 wt.%, the relative content of Al₂O₃ does not change significantly. When the Y content increases to 0.2 wt.%, the relative content of Al₂O₃ is decreased to 4.31%, indicating that corrosion is easier to occur at this content, which is consistent with the most negative value of E_corr. Subsequently, when the Y content is further increased to 0.3 wt.%, the relative content of Al₂O₃ is enhanced to 21.75%, which is a slight increase compared with the alloy with 0.1 wt.% Y. Thus, 0.2 wt.% Y addition has an detrimental effect on the passivation film, rendering the alloy more susceptible to corrosion. While 0.1 and 0.3 wt.% Y addition have little effect on the passivation behavior of the alloy.

Table 5.

The relative content of Al₂O₃ and Al(OH)₃ in the corrosion products of Al-Mg-Mn alloys with different Y content after 48 h salt spray tests (%).

Alloys	Al₂O₃/%	Al(OH)₃/%
Al-Mg-Mn	20.18	79.82
Al-Mg-Mn-0.1Y	21.27	78.73
Al-Mg-Mn-0.2Y	4.31	95.69
Al-Mg-Mn-0.3Y	21.75	78.25

When the Y content is 0.1 wt.%, the number of Al₃Mg₂ and Al₆(Fe, Mn) phases remains almost unchanged and a small amount of Y-containing multi-component phase is formed in α-Al matrix. Thus, the total number of precipitates is an increased, which leads to more corrosion initiation sites. Since both Al₆(Fe, Mn) phase and the Y-containing phase act as cathodes (Figures 10 and 11), more galvanic couples form in the alloy. Additionally, the composition of corrosion products is almost unchanged, resulting in higher corrosion current density (i_corr) and reduced corrosion resistance.

When the Y content is increased to 0.2 wt.%, the number of Al₆(Fe, Mn) phase is decreased and the Y-containing multiple phases are increased, which significantly reduces the potential difference between the precipitated phase and α-Al matrix, which weakens the galvanic corrosion. However, as shown in Figure 12, the ratio of dense Al₂O₃ in corrosion product layer is significantly decreased, while the content of loose and porous Al(OH)₃ is increased, which is beneficial to the development of corrosion. In addition, pitting in salt spray testing is primarily caused by the reduction of the relative content of Al₂O₃ in corrosion products, which exposes the matrix. The above two changes counteract each other, and the corrosion current density (i_corr) is decreased compared with the alloy with 0.1 wt.% Y addition, making it comparable to the Y-free alloy.

When the addition of Y is further increased to 0.3 wt.%, the precipitates in the matrix are dominated by the Y-containing multiple phases (Figure 11). This significantly reduces the number of the strong cathode Al₆(Fe, Mn) phase, which greatly decreases the galvanic corrosion in the matrix. Furthermore, the Al₃Mg₂ phase, as an anodic phase (Figure 9), becomes smaller and the number is reduced, which fundamentally decreases the driving force of corrosion. The content of dense Al₂O₃ in the corrosion product is increased from 4.31% to 21.75%, resulting in the i_corr is decreased and the corrosion rate is decreased, ultimately the corrosion resistance of the alloy is enhanced.

In terms of the effect of Y addition on the corrosion behavior Al-Mg alloys, only intergranular corrosion can be found. Li et al.³⁹ reported that the addition of 0.1 wt.% Y to Al-9.2Mg-0.7Mn alloy effectively inhibited the formation of the grain boundary β-Al₃Mg₂ phase, thereby reducing the weight loss of intergranular corrosion by about 50%. Similarly, Ru et al.⁴⁰ showed that after the addition of 0.1 wt.% Y and 0.1 wt.% Zr, the grains of the Al-5.5Mg-0.3Mn alloy can be refined and AlMgYZr phase forms, inhibit the segregation of Mg at the grain boundary, and mitigate galvanic corrosion.

However, previous studies have focused on IGC, and there is limited on the pitting behavior of the Y microalloyed Al-Mg-Mn alloys. In this work, the i_corr is decreased from 2.633 × 10⁻⁶ to 1.597 × 10⁻⁷ A/cm² after adding 0.3 wt.% Y, and the corrosion resistance is improved by about 16.5-fold improvement.

Conclusions

The effect of Y addition on the microstructure, and corrosion resistance of the as-cast Al-Mg-Mn alloys was examined using potentiodynamic polarization, EPMA, EBSD, and KPFM techniques. The main findings of this work can be summarized as follows:

The grain size of as-cast Al-Mg-Mn alloys with different Y contents (0.0, 0.1, 0.2, 0.3 wt.%) gradually become coarser with the increase of Y content, which are 222.6 μm, 243.5 μm, 236.1 μm and 307.8 μm, respectively. In the Y-free alloy, the main precipitates are the β-phase (Al₃Mg₂) and Al₆(Fe, Mn) phase. Both the size and volume fraction of Al₃Mg₂ phase and the number of Al₆(Fe, Mn) precipitates are decreased with the increase of Y content. Meanwhile, various multi-component phases Y-containing form with the increase of Y content.

The corrosion current density (i_corr) in the 3.5 wt.% NaCl solution is first increased and then decreased, the corrosion potential (E_corr) shifts positively, and the polarization resistance (R_p) is increased markedly with the increase of Y content in the as-cast Al-Mg-Mn alloy. Salt spray tests further confirm this conclusion: after 48 hours of exposure, the alloy with 0.3 wt.% Y shows no significant pitting, only corrosion product deposition, demonstrating excellent pitting resistance.

Average weight loss is decreased with the increase of Y content, which is in agreement with the trend of electrochemical and salt spray tests. Since the Al₃Mg₂ phase tends to precipitate at the grain boundary, the precipitated phase at the grain boundary decreases with the increase of Y content, and the channels of intergranular corrosion decreases, enhances the resistance of IGC. Therefore, an appropriate Y addition, particularly 0.3 wt.%, can significantly enhance the corrosion resistance of the as-cast Al-Mg-Mn alloy.

Although this study confirms the beneficial effect of Y on the corrosion resistance of Al-Mg-Mn alloys, certain aspects remain to be explored deeply. Future research should design a wider range of Y contents to optimize the alloy composition, evaluate the influence of deformation processes on microstructural evolution and corrosion behavior and assess mechanical properties.

Footnotes

Acknowledgements

This work was financially supported by Guangxi Science and Technology Programs (Guike XT2503980012, and Guike AD25069019).

ORCID iD

Junwei Fu

Funding

This work was finacially supported by Guangxi Science and Technology Programs (Guike XT2503980012, and Guike AD25069019).

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship and/or publication of this article.

Author contribution(s)

Jingzeng Tang: Data curation; Investigation; Methodology; Writing – original draft.

Junwei Fu: Conceptualization; Formal analysis; Methodology; Writing – review & editing.

Chongyu Liu: Data curation; Investigation.

Yunfang Wan: Data curation; Investigation.

Yuanqing Chen: Data curation; Investigation.

Maomi Zhao: Data curation; Investigation.

Baorong Hou: Conceptualization; Supervision.

References

Pan

Zhang

, et al. Mechanical properties, corrosion behavior and microstructure evolution of zinc and scandium co-strengthened 5xxx alloy. J Rare Earths 2023; 41(11): 1819–1826.

Zhang

, et al. Analysing the degree of sensitisation in 5xxx series aluminium alloys using artificial neural networks: a tool for alloy design. Corros Sci 2019; 150(15): 268–278.

Kotov

Mochugovskiy

Mosleh

, et al. Microstructure, superplasticity, and mechanical properties of Al-Mg-Er-Zr alloys. Mater Char 2022; 186: 111825.

Haoran

Zeyu

, et al. Mechanical property and microstructure evolution of novel Al-Mg-Zn(-Cu) alloys under dynamic impact. Mater Sci Technol 2021; 37(9): 852–862.

Moradpour

Khodabakhshi

Eskandari

. Microstructure-mechanical property relationship in an Al-Mg alloy processed by constrained groove pressing-cross Route. Mater Sci Technol 2018; 34(8): 1003–1017.

Song

Dong

Zhang

, et al. Recent progress of Al-Mg alloys: forming and preparation process, microstructure manipulation and application. J Mater Res Technol 2024; 31: 3255–3286.

Jiang

Zhang

, et al. Intermetallic phase evolution of 5059 aluminum alloy during homogenization. Trans Nonferrous Met Soc China 2013; 23(12): 3553–3560.

Jain

Lim

MLC

Hudson

, et al. Spreading of intergranular corrosion on the surface of sensitized Al-4.4Mg alloys: a general finding. Corros Sci 2012; 59: 136–147.

Qiu

Yang

Chen

, et al. Sanders.Recrystallization behaviors of AlMgScZr alloys with different Al₃(Sc,Zr) dispersoid distribution. Mater Char 2025; 228: 115361.

10.

Deng

Zhang

Yang

, et al. Microstructure characteristics and mechanical properties of new aerospace Al-Mg-Mn alloys with Al₃(Sc_1-xZr_x) or Al₃(Er_1-xZr_x) nanoparticles. Mater Char 2019; 153: 79–91.

11.

Ahmed

Liu

Paul

, et al. Effects of AlMn dispersoids and Al₃(Sc,Zr) precipitates on the microstructure and ambient/elevated-temperature mechanical properties of hot-rolled AA5083 alloys. Mater Sci Eng A 2022; 855(10): 143950.

12.

Algendy

Rometsch

Chen

. Tailoring microstructure and mechanical properties of an AA5454 extruded aluminum alloy with Sc/Zr microalloying and processing conditions. Mater Sci Eng A 2025; 948: 149318.

13.

Qiu

Yang

, et al. The influence of Sc and Zr additions on microstructure and corrosion behavior of AA5182 alloy sheet. Corros Sci 2022; 199(1): 110181.

14.

Gao

Xie

Huang

, et al. Effect of trace Sc addition on microstructure, mechanical and stress corrosion cracking properties of Al-Mg alloys fabricated by wire arc additive manufacturing (WAAM). J Alloys Compd 2025; 1021(5): 179575.

15.

Mochugovskiy

Aremu

Yakovtseva

, et al. Influence of Y and Er on the microstructure and mechanical properties of the Al-Mn-Zr-based alloys processed by mechanical alloying and hot press sintering. Mater Sci Eng A 2025; 944: 148926.

16.

Ding

Zhao

Chen

, et al. Effect of rare earth Y and Al-Ti-B master alloy on the microstructure and mechanical properties of 6063 aluminum alloy. J Alloys Compd 2020; 830(25): 154685–154685.

17.

Jin

Xiong

, et al. Study on microstructure and mechanical properties of ZL107 alloy added with Yttrium. J Rare Earths 2013; 31(2): 198–203.

18.

Wei

Lei

Yan

, et al. Microstructure and mechanical properties of A356 alloy with yttrium addition processed by hot extrusion. J Rare Earths 2019; 37(6): 659–667.

19.

Zhang

Guo

Bai

, et al. Effect of rare earth Ce and Y on microstructure and viscosity of ZL114A alloy. J Alloys Compd 2024; 1004(5): 175755.

20.

Ding

Sun

Wei

, et al. Thermal deformation behavior and microstructure evolution of AA6061 aluminum alloy with trace rare-earth Y added. J Alloys Compd 2025; 1031(5): 180979.

21.

Wang

Pan

Wang

, et al. Improved heat and corrosion resistance of high electrical conductivity Al-Mg-Si alloys by multi-alloying of Ce, Sc and Y. Corros Sci 2024; 226: 111695.

22.

Peng

Yuan

Wang

, et al. Achieving high strength and plasticity in a novel Al-Zn-Mg-Cu-Y-Cr alloy via thermomechanical treatment. Mater Sci Eng A 2025; 941: 148623.

23.

Zhang

, et al. Effect of combined addition of Zr, Ti and Y on microstructure and tensile properties of an Al-Zn-Mg-Cu alloy. Mater Des 2022; 223: 111129.

24.

Wei

Liu

, et al. Effect of Y and Zr addition on the intergranular corrosion of Al-Mg alloy after sensitization treatment. Micron (Oxford, England: 1993) . 2025; 191: 103789.

25.

Hou

Zhang

Ding

, et al. Solute clustering and precipitation of Al-5.1Mg-0.15Cu-xZn alloy. Mater Sci Eng A 2019; 759(24): 465–478.

26.

Meng

Zhang

Hua

, et al. Mechanical properties, intergranular corrosion behavior and microstructure of Zn modified Al-Mg alloys. J Alloys Compd 2014; 617(25): 925–932.

27.

Jianing

Yingwei

Ziyu

, et al. The Synergistic Effect of β-Al₃Mg₂ and Al-Mn Intermetallic Particles on the Corrosion Mechanism of 5383 Al Alloy. J Alloys Compd 2025: 1048.

28.

Jie

Zou

Wang

, et al. Thermal stability of Al-Mg alloys after solidification under high pressures. J Alloys Compd 2014; 584(25): 507–513.

29.

Yang

Gao

Liu

, et al. Formation mechanism of refined Al₆(Mn, Fe) phase particles during continuous rheo-extrusion and its contribution to tensile properties in Al-Mg-Mn-Fe alloys. Mater Sci Eng A 2023; 872(8): 144952.

30.

Liu

Han

, et al. The electrochemical Co-reduction of Mg-Al-Y alloys in the LiCl-NaCl-MgCl₂-AlF₃-YCl₃ melts. Metall Mater Trans B 2015; 46(2): 644–652.

31.

Guo

Zheng

, et al. Phase transformation and property improvement of Al-0.6Mg-0.5Si alloys by addition of rare-earth Y. Sci Eng Compos Mater 2025; 32(1).

32.

Klosek

Tobola

Vemiére

, et al. Electronic structure and magnetic properties of RMnX (R = Mg, Ca, Sr, Ba, Y; X = Si, Ge) studied by KKR method. Eur Phys J B 2004; 42(2): 219–230.

33.

Handel

Nickel

Lampke

. Effect of different grain sizes and textures on the corrosion behaviour of aluminum alloy AA6082. Mater Sci Eng Technol 2011; 42(SP7): 606–611.

34.

Ralston

Birbilis

. Effect of grain size on corrosion: A review. Corros 2010; 66(7): 075005-1–075005-13.

35.

Revilla

Terryn

Graeve

. Growth kinetics and passive behavior of the native oxide film on additively manufactured AlSi10Mg versus the conventional cast alloy. Corros Sci 2022; 203(15): 110352.

36.

Zhang

Chou

Zhou

, et al. Corrosion-resistant L1₂-strengthened high-entropy alloy with high strength and large ductility. Corros Sci 2023; 225: 111593.

37.

Cheng

Wang

, et al. Origin of the pitting corrosion in the as-rolled and annealed ferritic stainless steel in 3.5wt.% NaCl solution. Electrochim Acta 2025; 520(20): 145899.

38.

Wang

Qian

Xie

, et al. Effect of corrosion product on corrosion behavior of PEO coating based on its structure and semiconducting properties. Corros Sci 2025; 246(15): 112727.

39.

Xia

Yan

, et al. Enhancing the intergranular corrosion resistance of high Mg-alloyed Al-Mg alloy by Y addition. Mater Corros 2020; 71(11): 1802–1811.

40.

Wei

Liu

, et al. Effect of Y and Zr addition on the intergranular corrosion of Al-Mg alloy after sensitization treatment. Micron 2025; 191: 103789.