Microstructure, Mechanical Properties, and Corrosion Behavior of Al-4.0Cu-1.1Li-0.5Mg-xAg Alloys

Abstract

The influence of various Ag contents on the microstructure, mechanical properties, and corrosion behavior of extruded Al-4.0Cu-1.1Li-0.4Mg-xAg-0.2Mn-0.2Zr (x = 0.4 and 0.9, wt.%) alloys was investigated. The alloy with 0.9 Ag content contains higher number density of slender T1 (Al₂CuLi) precipitates along with some θ’ (Al₂Cu) phases in the matrix than the alloy with 0.4 Ag content, which is associated with a more rapid hardening response and higher mechanical properties and corrosion resistance, particularly for aging at 130 °C. When aging at high temperatures (above 160 °C), the increase of Ag content mitigates hardness loss by preventing the T1 precipitates from coarsening, and makes the alloy decorate more coarse precipitates at grain boundaries, which leads to the fracture morphology mainly occupied by intergranular fracture. Furthermore, due to the simultaneous promotion of T1 precipitates at grain boundaries and in grain interiors, the 0.9 Ag-containing Al-Cu-Li-Mg-Ag alloy has almost no improvement in corrosion resistance.

1. Introduction

Weight reduction, as one of the latest development trends in the transportation industry, can decrease fossil fuel consumption and moderate greenhouse gas emissions [1,2]. Researchers found that adding Li to Al can reduce the alloy’s density and increase its stiffness per percentage weight [3]. For each 1 wt.% of Li added to Al, the density decreases by 3%, and the elastic module increases by approximately 6% [3]. In combination with the addition of Cu, Al-Li alloys present remarkable potential for precipitation strengthening. After the development of the last decades, the third generation Al-Cu-Li alloys containing 0.8–1.2 wt.% Li provide not only weight reduction but also excellent corrosion resistance, high stiffness to weight ratio, more favorable damage tolerance, and a good combination of strength and toughness [4,5]. Although many advances have been achieved, the conflict between strength and corrosion behaviors in Al-Cu-Li alloys has limited their widespread application in the military, aerospace, and commercial sectors.

In general, regulating precipitates is considered an essential method for age-hardened Al alloys to improve their overall performance. Al-Cu-Li alloys contain a complex precipitation sequence, including T1 (Al₂CuLi), θ’ (Al₂Cu), S’ (Al₂CuMg), δ’ (Al₃Li), and σ (Al₅Cu₆Mg₂) phases [6,7,8]. In contrast to other precipitates, the T1 phases precipitate simultaneously at grain boundaries and in grain interiors, which is considered to be the primary contributor to the properties of third-generation Al-Cu-Li alloys [9]. Therefore, considerable research has been conducted to control the precipitation of the T1 phase in Al-Cu-Li alloys. Generally, the T1 phase is believed to nucleate via a stacking fault mechanism, for which pre-deformation is necessary before artificial aging to enhance T1 precipitation [10,11]. Furthermore, precipitation kinetics of T1 phases are also determined by alloying compositions. Besides the main alloying elements Cu and Li, the addition of other minor solutes, such as Mg and Zn, can also affect the precipitation behavior of the T1 phases, resulting in desired properties [12,13,14].

Among the micro-alloying elements, it is known that Ag additions could improve the precipitation strength for Al-Cu-Li alloys. Gumbmann et al. [12] suggested that Ag elements could trap vacancies to inhibit the formation of θ’ phases, contributing to a high number density of T1 phases. Generally, alloying with only Ag has limited its hardening potential [13]. According to Huang et al. [15], adding Ag to Al-Cu-Li alloys in conjunction with Mg could reduce the stacking fault energy and facilitate the formation of dislocations under pre-deformation, which subsequently benefited the precipitation of T1 phases [16] and improved the age-hardening response significantly [7,12,13,17]. The results were also in good agreement with the microstructure in which dislocations were enriched with Cu, Mg, and some Ag atoms at the initial aging stage [18]. Boukos et al. [19] found that the addition of 0.5 wt.% Ag in Al-1.3Cu-2.4Li-0.7Mg-0.1Zr alloys diminished the size of T1 precipitates. Furthermore, Ag element distributed on the T1/matrix interface seemed to relieve the coherency strain using high-angle annular dark field-scanning transmission electron microscopy (HAADF-STEM) and atom probe tomography [13,18], which was similar to the role of Ag on the Ω/matrix interface [20]. Therefore, Ag addition plays an essential role in promoting the precipitation kinetics for the T1 phase and helps to reach a higher volume fraction of T1 phases. The analysis of T1 composition in the alloy AA2198 after long-term aging at 155 °C demonstrated that Ag may contribute to the coarsening resistance of T1 phases [21]. More recently, Liu et al. [22] offered contradictory findings that Ag addition in the Al-3.2Cu-1.0Li-0.4Mg-0.3Mn-0.1Zr alloy promoted the formation of the coarse and discrete Cu-rich phases at the grain boundaries, leading to a decreased hardness at the overage condition.

As mentioned before, although previous works have demonstrated that the addition of minor Ag on Al-Cu-Li-Mg alloys could reduce the interfacial energy, increase Li solubility, and provide heterogeneous nucleation sites for T1 phases, the influence of higher Ag additions (>0.5 wt.%) for novel Al-Cu-Li alloys on the microstructures and related properties during different aging conditions is still not clear. Therefore, further studies are needed to understand the correlation of high Ag content and aging-microstructure-properties. This study aims to investigate the effect of varying Ag contents on the mechanical properties, corrosion behavior, and precipitation behavior of extruded Al-Cu-Li-Mg alloys under different aging conditions. The mechanism of Ag content in microstructure evolution and corresponding properties will be discussed.

2. Experimental Procedures

Al-Cu-Li-Mg-Ag alloys (

Table 1. The chemical composition (wt.%) of the studied alloys.

Alloys		Cu	Li	Mg	Ag	Mn	Zr	Al
0.4 Ag	Nominal	4.0	1.1	0.5	0.4	0.2	0.2	Bal.
	Actual	3.99	1.12	0.42	0.33	0.25	0.16	Bal.
0.9 Ag	Nominal	4.0	1.1	0.5	0.9	0.2	0.2	Bal.
	Actual	3.98	1.13	0.45	0.88	0.21	0.20	Bal.

) were fabricated in a vacuum resistance melting furnace (ZGG-0.025A, Shanghai, China) using commercial pure Al, pure Mg, pure Zn, pure Ag, pure Li, and master alloys of Al-50Cu, Al-10Mn, and Al-5Zr. The actual composition of the alloys (Table 1) was characterized by inductively coupled plasma atomic emission spectrometry (ICP-AES, iCAP7600, Thermo Fisher Scientific, Waltham, MA, USA). The as-cast ingots were homogenized for 24 h at 500 °C and then cooled in the air furnace (Zhonghuan, Tianjin, China). The homogenized ingots were lathed into cylinders with a diameter of 98 mm. The cylinders were subsequently hot extruded into sheets with a size of 60 (width) mm × 13 (thickness) mm, and the extrusion ratio was 9. For hot extrusion, the ingots were heated for 60 min at 450 °C. After a 60-min solution treatment at 530 °C followed by water quenching, the sheets were cold-rolled with a thickness reduction of 5% at room temperature and then aged at 130 °C, 145 °C, 160 °C, and 185 °C.

The hardness of the samples was tested using HVS-50 microhardness tester (Huayin, Laizhou, China) under a load of 3 kgf for 15 s. The hardness value of each sample was obtained based on the average value of seven measurements. The tensile tests were performed on a universal testing machine (Instron 3369, Boston, MA, USA) equipped with a contact extensometer (Epsilon, Jackson, FL, USA). The tested samples with the dimension of 6 mm in width and 2 mm in thickness were processed along the sheet extrusion direction at a displacement rate of 2 mm/min according to GB/T 228 standard. For each condition, three samples were conducted to calculate the average tensile properties, including ultimate tensile strength (UTS), yield strength (YS, offset = 0.2%), and elongation (EL).

The corrosion characteristics were investigated using intergranular corrosion (IGC) tests and electrochemical behaviors. Before corrosion tests, each sample was polished and further cleaned with ethanol. According to GB/T 7998-2005, the intergranular corrosion (IGC) test samples were immersed in the etching solution (57 g/L NaCl + 10 mL/L H₂O₂) at 35 ± 2 °C for 6 h. After immersion in the etching solution, an optical microscope (OM, Leica MEF4A/M, Wizz, Germany) was performed to observe the cross-section of the corroded samples. The dynamic polarization curves were measured by electrochemical workstation (CHI660C, Shanghai Chenhua, Shanghai, China) in 57 g/L NaCl solution at 0.1 mV/s. The surface area of each sample exposed to the solution was 0.5 cm².

Scanning electron microscope (SEM, Quanta 250, FEI, Hillsboro, OH, USA) equipped with an energy dispersive spectrometer (EDS, Oxford instrument, Oxford, UK) was used to observe the fracture morphology of the alloys under different treatments. Transmission electron microscopy (TEM, Tecnai G2 F20, FEI, Hillsboro, OH, USA) was used to characterize the microstructure of the alloys with different aging temperatures. At least three pictures were manually measured using the Image J analysis software (1.8.0, 2021, National Institutes of Health, Bethesda, MD. USA) to estimate the diameter distribution of T1 precipitates. Thin foils for TEM observation were cut from tensile samples, then mechanically ground to approximately 70 μm, and finally electro-thinned at −30 °C in a mixed solution of 75% CH₃OH and 25% HNO₃ (vol. %).

3. Results

3.1. Hardness

Figure 1 shows the evolution of the hardness as a function of aging time at different aging temperatures for the two alloys. The 0.9 Ag alloy experiences a steeper increase in hardness in the early aging stage at various temperatures. As aging proceeded, the hardness of the 0.4 Ag alloy at aging temperatures of 130 °C, 145 °C, 160 °C, and 185 °C reached its peak values at 90 h (187 HV), 70 h (193 HV), 20 h (187 HV), and 10 h (192 HV), respectively. By contrast, the 0.9 Ag alloy reached the peak aging more easily under different conditions with the hardness 198 HV (130 °C, 40 h), 197 HV (145 °C, 35 h), 198 HV (160 °C, 15 h), and 197 HV (185 °C, 4 h). The peak aging can be more easily achieved at higher temperatures for the two alloys, contributing to an acceleration in the hardening rate. Furthermore, compared with 0.4 Ag and 0.9 Ag alloys, the hardness of the 0.9 Ag alloy is a little higher than that of the 0.4 Ag alloy under all experimental conditions, indicating that the increasing of Ag content could improve the hardness of the alloys. According to Figure 1, it can be concluded that the age-hardening response is expected to be improved by increasing aging temperatures and Ag contents. Furthermore, it is interesting that the high Ag content alloy remains harder with extended aging time (in overage state), particularly when aged at 185 °C.

3.2. Tensile Properties

Figure 2 illustrates the tensile properties of the two alloys aged for 8 h and 32 h at different temperatures. The tensile strength exhibits an increasing trend with raising aging temperature for the two alloys, while elongation continues to decrease. At various temperatures, the 0.4 Ag alloy shows lower overall strength than the 0.9 Ag alloy, even as aging time increases. Compared to the low Ag content alloy, the YS in the 0.9 Ag alloy increases from 335 MPa to 468 MPa after 36 h at 130 °C, and the UTS increases from 515 MPa to 573 MPa. As the aging temperature rises to 145 °C, the YS and UTS of the 0.9 Ag alloy are 144 MPa and 107 MPa higher than that of the 0.4 Ag alloy. Further increasing the aging temperature to 160 °C, the 0.9 Ag alloy aged for 8 h has the highest UTS and YS of 691 MPa and 657 MPa, respectively, while the elongation reduces to 6.5%. Therefore, the tensile properties of the alloys can be improved by increasing the aging time at a certain temperature. However, when the aging temperature reaches 185 °C, the drop in tensile strength of the 0.4 Ag alloy is more significant than in the 0.9 Ag alloy as aging continues, which indicates that the loss in strength can be inhibited by increasing Ag content. A similar phenomenon is observed by considering the hardness curve in Figure 1.

Figure 3 shows SEM images of the fracture of the alloys at different aging temperatures. The dimples and intergranular cracks are indicated by white and black solid arrows, respectively. When the alloys are subjected to aging at 130 °C for 36 h, the fracture of the 0.4 Ag alloy (Figure 3a) exhibits numerous larger and deeper ductile trans-granular dimples covering the entire fracture surface compared with the 0.9 Ag alloy, indicating higher plasticity of the 0.4 Ag alloy. A closer look at Figure 3a reveals that many large particles are at the bottom of large dimples. As the aging temperature is raised to 160 °C, as shown in Figure 3b,e, the fracture surface still consists of large dimples, and there is also a micro dimple network formed by fine precipitates. However, the typical intergranular fracture characteristic is observed in the 0.9 Ag alloy. After further increasing the aging temperature to 185 °C, the large area of fracture surface in the 0.9 Ag alloy is occupied by intergranular fracture.

3.3. Corrosion Behaviors

Figure 4 presents IGC morphologies of the cross-section (TD-ED plane) of the two alloys aged at different temperatures for 36 h, and the corresponding maximum IGC depth is measured. The maximum IGC depth of the alloys remarkably decreases with increases in the aging temperature. The maximum IGC depths of the 0.4 Ag alloy aged at 130 °C, 160 °C, and 185 °C are 49.50 μm, 33.75 μm, and 27.60 μm, respectively. For the 0.9 Ag alloy, the maximum IGC depths under different aging temperatures are 37.50 μm (130 °C), 33.78 μm (160 °C), and 28.50 μm (185 °C), respectively. This observation suggests that the 0.9 Ag alloy sample has lower maximum IGC depth than the 0.4 Ag alloy at low aging temperature (130 °C). However, there is little difference among the corrosion degrees of the two alloys aged at higher temperatures (above 160 °C). The results show that the maximum IGC depth is affected by Ag content and the precipitation during the aging process.

To better understand the influence of Ag contents on corrosion performance, the corresponding polarization curves of the alloys aged for 36 h at different temperatures are shown in Figure 5, and the electrochemical parameters derived from the polarization curves are summarized in

Table 2. Electrochemical parameters of extruded 0.4 Ag and 0.9 Ag alloys derived from polarization curves.

Aging Condition	0.4 Ag Alloy		0.9 Ag Alloy
	Ecorr (VSCE)	Icorr (A/cm2)	Ecorr (VSCE)	Icorr (A/cm2)
130 °C/36 h	−0.601	4.139 × 10−6	−0.633	2.578 × 10−6
160 °C/36 h	−0.682	3.251 × 10−6	−0.699	1.732 × 10−6
185 °C/36 h	−0.680	2.244 × 10−6	−0.672	1.411 × 10−6

. The polarization curves all exhibit a similar variation tendency in the two alloys. When aging at 130 °C, the corrosion potential (E_corr) and corrosion current density (I_corr) of the 0.4 Ag alloy is −0.601 V and 4.139 × 10⁻⁶ A/cm², respectively. However, increasing Ag content to 0.9 wt.% results in a more negative E_corr (−0.633 V) and an associated reduction in I_corr (2.578 × 10⁻⁶ A/cm²). Therefore, the 0.9 Ag alloy exhibits better corrosion resistance under this condition. Upon increasing the aging temperature to 160 °C, the E_corr of the alloys becomes negative. Furthermore, the corrosion potential of the 0.9 Ag alloy is −0.017 V higher than that of the 0.4 Ag alloy, indicating limited difference in microstructures. The E_corr of the alloys shifts positively when aging at 185 °C, and the E_corr of the 0.4 Ag alloy is more negative than the 0.9 Ag alloy. The I_corr of the alloys also has a downward tendency with increasing aging temperatures. Obviously, the 0.9 Ag alloy exhibits lower I_corr compared to the 0.4 Ag alloy. Based on the above observations, the increase of Ag content on the corrosion properties of the alloys decreases with raising aging temperatures, which is consistent with the observations of intergranular corrosion (Figure 4).

3.4. TEM Observation

Figure 6 shows the bright-field (BF) images, high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) images, and corresponding selected area electron diffraction (SAED) patterns of the two alloys aged for 36 h at 130 °C taken along the [11] _Al direction. As shown in Figure 6a, there are not only distinct continuous diffraction streaks along the <111> _Al direction but also diffraction spots on the 1/3 <220> _Al and 2/3 <220> _Al. In addition to the before-mentioned Al matrix diffraction spots, SAED patterns contain no characteristic reflections from θ′ phases formed on the {001} _Al planes due to their low number density [23]. These prove that the T1 phase aligned in {111} _Al is the primary precipitate in this condition. It can be observed from Figure 6b that plate-like T1 phases are relatively small in dimension and non-uniformly distributed in the matrix, which is further consistent with the reflections observed in the diffraction patterns. At the same time, there are also entangled dislocations and a small amount of fine θ′ phases (Figure 6b,c). The HAADF-STEM micrograph reveals that all dislocations marked by black solid arrows exhibit higher brightness relative to the surrounding matrix (Figure 6c). Compared with the 0.4 Ag alloy, when the Ag content reached 0.9 wt.% (Figure 6d), except for the diffraction fringes and spots produced by the T1 phase in SAED, distinct diffraction spots in 1/2 <200> _Al direction and discontinuous streaking along <001> _Al directions which arose from the θ′ phase are detected [23]. The BF image also shows a dramatic change in the morphology of precipitates presented in the matrix. The number density of the uniform distributed T1 phases in the 0.9 Ag alloy increases considerably, while a relatively small number of θ′ phases exit between the intersecting T1 phases (Figure 6e,f).

The BF images and precipitate size distribution of the alloys with different Ag contents taken along the [110] _Al zone axis aged at 160 °C and 185 °C for 36 h are presented in Figure 7. After increasing aging temperatures, a visible change in SAED patterns occurs: the spots in the middle of the rhombus disappear while the intensity of T1 spots and streaks become the strong ones, indicating that T1 phases are the dominating precipitates. When aging at 160 °C, a large number of precipitates distribute in the Al matrix uniformly, and the length of T1 phases in the 0.4 Ag alloy becomes longer. While there is no diffraction characteristic of the θ′ phase in the SAED image, a small number of θ′ phases are still presented in Figure 7a. Furthermore, the diffraction spots which originated from the θ′ phase disappear in the 0.9 Ag alloy, revealing that the number density of θ′ phases reduces drastically to the extent that it cannot be detected. The diffraction spots and streaks related to the T1 phase are sharper, suggesting a significant increase in the number density of T1 phases (Figure 7b). In addition, the frequency plot in Figure 7c provides the detail that increasing Ag content raises the proportion of larger T1 plates effectively (e.g., diameter ≥150 nm). In Figure 7c,d are the microstructures of the alloys aged at 185 °C for 36 h, where both the plate-shaped T1 and θ’ phases distribute in the two alloys. Compared with the 0.4 Ag alloy, the particle size of T1 phases is more uniform and shorter than those in the 0.9 Ag alloy (Figure 7f). Additionally, these T1 precipitates become obviously coarse after aging at 185 °C, particularly in the 0.4 Ag alloy (Figure 7d), leading to decreases in the number density of T1 phases. On the contrary, partially coarse T1 phases can be seen in the 0.9 Ag alloy.

Figure 8 illustrates the morphology of grain boundaries of the two alloys treated for 36 h at different aging temperatures. As shown in Figure 8a,d, the 0.4 Ag alloy with 32 h heat treatment at 130 °C has no obvious precipitations in GBs, while the 0.9 Ag alloy has many fine precipitations decorating the GBs. After increasing the heat treatment temperature to 160 °C, the precipitates densely distribute on the GBs in the two alloys (Figure 8e,f). In contrast to the 0.4 Ag alloy, more intensive and continuous phases precipitate at GBs in the 0.9 Ag alloy. Based on a similar precipitation plane of T1 phases within the grain (Figure 8g), the high-resolution transmission electron microscopy (HRTEM) image of the grain boundary precipitate in Figure 8h and its corresponding fast Fourier transform (FFT), the cross-precipitated grain boundary precipitations are verified as T1 phases. As the aging temperature reaches 185 °C (Figure 8c,f), T1 precipitates separate from each other and coarsen into discontinuous particles. By comparison with the 0.4 Ag alloy, precipitates at the GB in the 0.9 Ag alloy are coarser and more discrete. According to Figure 8i, these coarse grain boundary phases are also T1 phases.

4. Discussion

4.1. Influence of Ag Contents on Aging Precipitation

The hardness and hardening kinetic are improved as the Ag content increases (Figure 1), indicating precipitation response differences between the two alloys aged at various temperatures. The effect of Ag content on the precipitation, such as T1 and θ′ phases, could be interpreted in terms of different nucleation mechanisms [15,17]. The nucleation and growth of the T1 phase, located on the {111} _Al plane, are between the two partial dislocations (a/6 <112>) produced by cross-slip on the adjacent {111} _Al plane [24,25]. Ag addition may reduce the stacking fault energy and improve the dislocation density [15,16], so the solutes Cu and Li in the 0.9 Ag alloy are depleted in a shorter timescale than the 0.4 Ag alloy. Furthermore, Ag addition increased the solubility of Li in the matrix, which is beneficial for reducing the size and increasing the number density of T1 phases [19]. By trapping vacancies after quenching, alloys containing traces of Ag elements also possess a high density of dislocations [26]. Hence, the increase of Ag content retards vacancy-assisted formation of the GP-zone, which is regarded as the precursor of θ′ phases [17], and promotes the precipitation of T1 phases (Figure 6 and Figure 7).

Additionally, as shown in Figure 7, there are more coarse T1 phases in grain interiors of the 0.4 Ag alloy than that of the 0.9 Ag alloy when aging at high temperatures. Gault et al. [27] found that Ag atoms in solid solution first decrease during precipitation of T1 and then increase due to release from the precipitates, revealing that Ag atoms had a catalytic effect on T1 precipitation. Further investigation has shown that the solute Ag atoms only presented at the T1/Al matrix interfaces gradually return from T1 phases to the matrix with aging prolongation [12,18]. As Al and Ag have only 0.5% differences in atomic radii, more Ag additions could remarkably reduce interfacial energy between T1 phases and the Al matrix. Thus, the T1 phase can exist in the matrix stably by reducing coarsening kinetics when aging at high temperatures. Therefore, adding Ag to Al-Cu-Li alloys leads to a lower coarsening ability of the T1 phase, particularly in the alloys with a high Ag content. Accordingly, a combined influence of the above factors results in an accelerated age-hardening response at the initial aging process and maintains hardness with aging prolongation in the 0.9 Ag alloy (Figure 1).

4.2. Effect of Ag Contents on Mechanical Properties

The enhancement in strength is mainly ascribed to precipitation strengthening for age-hardened Al alloys. Aging treated Al-Cu-Li alloys produce a high number density of shear-resistant T1 precipitates and the plate-shaped θ′ precipitates. These precipitates can effectively hinder the movement of dislocations in the slip planes, improving the mechanical properties of the alloys [28]. Based on the analysis of the results above, the microstructures of the two alloys are primarily dominated by T1 plates, which mainly affect the mechanical properties of Al-Cu-Li alloys. It is not complicated to confirm that the change in strength during different aging conditions (Figure 2) is overwhelmingly dependent on the evolution of the T1 morphology. According to Dorin et al. [21,28], the precipitation strengthening related to the contribution of the T1 phase, neglecting the influence of other phases (θ’ and S′), can be expressed as:

$$\Delta {\tau =D}^2{N}^{\frac{1}{2}}{t}^{-}{}^{\frac{3}{2}}$$

(1)

where Δτ is precipitation hardening provided by the T1 phase; t, D, and N are the average thickness, number density, and average diameter of the T1 phase, respectively. It can be deduced from Equation (1) that the increase in the thickness t of the T1 phase adversely affects the strength, while the diameter D and number density N have a positive effect on strength. Therefore, a higher strength contribution from the T1 phase can be achieved by a larger diameter and a higher volume fraction of T1 precipitates with a single unit thickness.

As confirmed by TEM observations (Figure 6 and Figure 7), it has been found that the aging temperature significantly affects the precipitation response. With increasing aging temperature from 130 °C to 160 °C, the number density of T1 phases increases. The T1 precipitates are coarsened with further increasing aging temperature to 185 °C, which suggests that the elevated temperature promotes the precipitation kinetics and nucleation efficiency of the T1 phase leading to increased strength [29]. When aging at 130 °C, the 0.9 Ag alloy has higher strength than the 0.4 Ag alloy, which can be attributed to increasing the number density of T1 precipitates. The strength almost reaches the peak level because of the fully precipitated condition at 160 °C. Although there are many T1 phases in the 0.4 Ag alloy, the homogeneously distributed T1 phase in the 0.9 Ag alloy has greater number density and longer diameter. Furthermore, the increased precipitation density also decreases the distribution distance of precipitates; the smaller the space is, the more effective pinning on dislocations is [30]. Thus, the 0.9 Ag alloy exhibits a higher strength. With the aging temperature rise to 185 °C, the coarse T1 phases at the grain boundary and in grain interior are the primary factors for the decrease in the strength of the alloys. The number density of coarse T1 phase of the 0.4 Ag alloy is higher than that of the 0.9 Ag alloy, which means that the alloy with high Ag content presents superior thermal stability (Figure 7c,d). Moreover, the discrete coarse particles would result in a low grain boundary stress tolerance limit [31,32], which may contribute to intergranular fracture behavior and reduce the mechanical properties of the 0.9 Ag alloy. From this perspective, it is also necessary to consider the morphology of grain boundaries to further enhance the mechanical properties of Al-Cu-Li alloys.

4.3. The Role of Ag Contents in the Corrosion Behavior

It is well known that the corrosion behaviors in Al-Cu-Li alloys, dominated by intergranular corrosion, are usually influenced by precipitates [33]. In the present study, with the increase in the Ag contents and aging temperatures, the number density of T1 phases in grain interiors increases, and T1 phases grow gradually along the grain boundary and transform from continued to discontinued. Therefore, the morphology of T1 phases along grain boundaries and in grain interiors can significantly affect the corrosion behavior of Al-Cu-Li alloys [34]. The micro-galvanic cell action happens at T1 precipitates as the T1 phase is more active than the surrounding solid solution Al matrix in the NaCl solution [34]. Since T1 precipitates are easier and denser to form along the grain boundary, the corrosion is prone to propagate along grain boundaries during the corrosion process [35]. However, Proton et al. [36] suggested that increasing the number density of intragranular precipitates leads to the depletion of Cu atoms in the solid solution, reducing the potential difference between the grain boundary and grain interior, resulting in the corrosion being more likely to occur within the grains. Accordingly, the corrosion resistance of the alloys can be improved by increasing the number density of intragranular T1 phases.

At low temperature (130 °C), the 0.9 Ag alloy has denser and more uniform dispersion of T1 phases within grains combining with the depletion of matrix Cu, resulting in the corrosion potential of the matrix becoming more negative and thus a significant elevation in IGC resistance (Figure 5). Due to the increase in the aging temperature (160 °C), the two alloys precipitate more T1 phases in the grain interior, further reducing the corrosion potential. Compared with the 0.9 Ag alloy, the density of precipitates within the grain is higher than that of the 0.4 Ag alloy, leading to the corrosion potential shifting to a more negative value (Figure 5 and Table 2). However, more continuous T1 phases distributed along the grain boundaries in the 0.9 Ag alloy are susceptible to the formation of corrosion channels for electron transfer [37], resulting in the corroded area beneath the surface with large corrosion depth (Figure 5). When the aging temperature reaches 185 °C, the continuous GBP tends to break. The increase of Ag content not only promotes the precipitation inside the grains but also facilitates the growth of GBPs with larger sizes. As a result of these contradictory factors, the corrosion behavior of the two alloys is less different under high-temperature aging conditions.

5. Conclusions

This study investigated the microstructure, mechanical properties, and corrosion behavior of Al-4Cu-1.1Li-0.5Mg alloys with various Ag contents, and the conclusions were as follows:

(1) When aging at low temperatures (130 °C), more Ag addition improves the precipitation of the thermosensitive T1 phases and inhibits the formation of θ’ phases, leading to a more rapid increase in hardness and a significant improvement of mechanical properties. As the aging temperature rises (above 160 °C), increasing Ag content prevents T1 phases from coarsening and promotes the dense precipitation of T1 precipitates at grain boundaries, which creates less hardness loss and facilitates intergranular fracture formation.

(2) The increase of Ag content can improve the corrosion resistance of the alloys aged at 130 °C. However, the corrosion resistance of the high Ag content alloy is almost unchanged with the rise of aging temperature, which may be attributed to decreasing precipitation differences between alloys in grain interiors and increasing the coarse T1 phases at the grain boundaries.