Microstructural Characteristics and Mechanical Properties of Al–5Cu–0.4Mg–0.1Zr (–0.4Ag) Alloys Processed by Continuous Cast and Conform Processes

Abstract

The Al–Cu–Mg–Ag alloys have excellent specific strength, good heat resistance and have a wide range of applications in the aerospace and automotive industries. However, industrial production of such alloys is a great challenge owing to the addition of Ag, which limits their widespread application. In this work, the industrial continuous cast and continuous extrusion (Conform) processes were employed to prepare Al–5Cu–0.4Mg–0.1Zr (–0.4Ag) alloys, and the effects of Ag addition on the microstructural characteristics and mechanical properties during processing and heat treatment were investigated. The results indicated that Ag addition significantly refined grain size, increased high-angle grain boundary fraction and grain orientation difference in as-cast Al–5Cu–0.4Mg–0.1Zr (–0.4Ag) alloys, and suppressed excessive grain coarsening during homogenizing annealing. During Conform, Ag further refined grain size, increased subgrain number and enhanced grain orientation difference in extruded alloys. For the aging heat treatment, the T6 process demonstrated superior strengthening effects compared to the T5 process. Following T6 treatment, Ag promoted efficient and uniform precipitation of the Ω (Al₂CuMgAg) phase and then significantly enhanced peak hardness (160 HV) and tensile strength (511.46 ± 2.06 MPa). Additionally, Ag accelerated second-phase dissolution throughout the entire process and produced finer, denser ductile dimples on tensile fracture surfaces to gain good strength–ductility balance.

1. Introduction

Amid the global shift in manufacturing toward high-end, lightweight, and intelligent development, fields like aerospace and automotive engineering impose stringent requirements on the performance of structural materials—including high strength, high-temperature resistance, low density, and long service life [1]. However, traditional aluminum alloy structural components face considerable challenges. In aerospace applications, when service temperatures range from 150 to 250 °C, the strengthening phases in Al-Cu-Mg alloys such as the θ’ phase and S phase undergo substantial coarsening, which leads to a rapid decline in alloy strength and hinders the maintenance of stable mechanical properties. In the automotive sector, components including battery housings and motor brackets for new energy vehicles demand lightweight materials that also ensure structural reliability under high-temperature conditions [2,3,4]. Therefore, Al–Cu–Mg–Ag heat-resistant aluminum alloys—featuring typical heat-treatable strengthening behavior, distinct compositional design, and excellent comprehensive properties—have emerged as one of the core materials to address the aforementioned challenges. Compared with conventional Al–Cu–Mg alloys, microalloying with Ag alters the primary age-precipitation strengthening phase from the metastable θ’ phase (Al₂Cu) to the thermodynamically more stable Ω phase (Al₂CuMgAg) [5]. This phase transformation refines the alloy microstructure and enhances its mechanical properties, with further performance optimization achievable via precise control of the Ag content. Increasing the Ag content from 0.41 to 0.97 wt.% increases the number of Mg-Ag co-precipitates, which accelerates Ω phase nucleation and promotes the formation of denser Ω precipitates, thereby yielding a more pronounced age-hardening response at 180 °C [6]. Concurrently, the coalescence/coarsening of S precipitates is significantly inhibited. The synergistic effect of these two precipitate phases imparts enhanced mechanical properties to the alloy. However, while Ag addition imparts excellent overall properties to Al–Cu–Mg alloys, it also brings technical challenges to industrial production. Consequently, unresolved process issues remain in current industrial manufacturing and engineering use [7]. The traditional production process for Al–Cu–Mg–Ag alloys primarily involves “gravity casting-homogenization annealing-indirect extrusion-solution treatment-aging”. Among these steps, gravity casting, owing to its relatively slow cooling rate, tends to result in microstructural inhomogeneity in castings, causing defects such as columnar grains and Ag segregation. Indirect extrusion, conducted at low speeds, may induce non-uniform grain orientation, leading to inconsistent alloy properties. This poses a challenge to simultaneously satisfying the requirements of high quality and industrial production [8,9].

To address existing issues, this study fabricated Al–5Cu–0.4Mg–0.1Zr (–0.4Ag) alloys via a combined route of continuous casting + Conform + Ag microalloying. Alloy billets were first prepared by continuous casting, whose rapid cooling property effectively inhibits Ag segregation, ensures the uniform distribution of Ag in the alloy matrix, and lays a high-quality microstructural foundation for subsequent extrusion. Subsequently, the Conform process was adopted to process the billets; with the dynamic recrystallization effect during extrusion, the grain size was significantly refined, the alloy microstructure was optimized, and the material’s mechanical properties were further improved [10,11]. Compared with the traditional “gravity casting-indirect extrusion” process, the “continuous casting + Conform” stepwise process in this study not only avoids the slow cooling defect of gravity casting but also overcomes the mechanical property fluctuation problem of low-speed indirect extrusion. It greatly simplifies the traditional production process, effectively avoids secondary damage to alloy structure and properties caused by intermediate processes, significantly enhances the stability of alloy mechanical properties, and further improves production efficiency and reduces industrial production costs [12,13,14]. This innovative integrated process achieves synergistic improvement of alloy performance and industrial applicability, filling the research gap in systematic investigation into such combined processes for heat-resistant Al–Cu–Mg–Ag alloys. This work systematically studied the effect of Ag addition on the microstructure and mechanical properties of the alloy in as-cast, extruded, and heat-treated states, aiming to clarify the Ag microalloying mechanism and its compatibility with the process, thus providing theoretical and technical support for industrial production.

2. Experimental Procedures

2.1. Alloy Preparation

Al–Cu–Mg–Zr (A1) and Al–Cu–Mg–Zr–Ag (A2) alloys were fabricated into Φ18 mm ingots by continuous casting, and their chemical compositions are listed in

Table 1. The chemical compositions and contents (wt.%) of the studied Al–Cu–Mg–Zr (–Ag) alloys.

Alloy/Element (wt.%)	Si	Fe	Cu	Mn	Mg	Zr	Ag	Al
Al-Cu-Mg-Zr (A1)	0.0550	0.0742	4.8	0.354	0.371	0.126	0.00056	Bal.
Al-Cu-Mg-Zr-Ag (A2)	0.0381	0.0765	5.12	0.359	0.408	0.103	0.370	Bal.

.

2.2. Thermomechanical Processing

The cast rods were homogenized at 500 °C for 6 h, then subjected to Conform to Φ12 mm rods followed by immediate water quenching. Conform parameters: bar preheating at 360–400 °C, extrusion at 480–490 °C (material temperature during processing). Two heat treatments were applied to the extruded alloys: (1) T5 treatment: direct artificial aging at 160–190 °C for 0–24 h; (2) T6 treatment: solution treatment at 510 °C for 1.5 h, water quenching, then artificial aging at 160–190 °C for 0–24 h.

2.3. Mechanical Testing

2.3.1. Vickers Microhardness Testing

After homogenization and aging, samples were ground and polished. Vickers microhardness was measured using a Digital Display Vickers Microhardness Tester (Model: HVS-1000, Shanghai Lidun Instrument & Meter Testing Technology Co., Ltd., Shanghai, China) with a load of 5 kgf and dwell time of 15 s; the average of five indentations per sample was recorded.

2.3.2. Room-Temperature Tensile Testing

Tensile tests were performed at room temperature using an Instron-3382 Electronic Universal Testing Machine (Instron Corporation, Norwood, MA, USA), equipped with an extensometer. The tests were carried out at 25 °C with a crosshead speed of 2 mm/min, following GB/T 228.1-2021 [15]. Circular proportional specimens (outer diameter Φ12 mm, gauge diameter Φ8 mm, gauge length L₀ = 50 mm) were prepared by lathing, with at least three specimens for each condition.

2.4. Microstructural Characterization

Microstructures and tensile fracture surfaces were observed by optical microscopy (OM, CK-300, Shanghai Caikon Optical Instrument Co., Ltd., Shanghai, China) and scanning electron microscopy (SEM, QUANTA 200, FEI Company, Hillsboro, OR, USA); specimens were ground, polished, cleaned and dried before testing without further treatment.

For electron backscatter diffraction (EBSD), specimens were electropolished after grinding and polishing to eliminate surface stress. The electrolyte was 80 vol.% anhydrous ethanol + 20 vol.% perchloric acid, cooled to −35~−25 °C with liquid nitrogen; electropolishing was performed at 25 V for 25 s (sample as anode, steel ruler as cathode). Specimens were ultrasonically cleaned in alcohol for 1 min and dried, then 2–3 micrographs were acquired from the central region of each specimen. Raw data were analyzed using OIM Analysis 6 to determine average grain sizes and standard deviations of A1 and A2 alloys.

For transmission electron microscopy (TEM), 0.5 mm × 1 cm × 1 cm slices were cut by a low-speed low-frequency wire saw, manually polished to 60–80 μm with 400–5000-grit sandpaper, then punched into Φ3 mm disks. Disks were thinned by electrolytic double-jet polishing (electrolyte: 70 vol.% methanol + 30 vol.% nitric acid, cooled to ~−30 °C) at 15 V for ~25 s. After ultrasonic cleaning in alcohol for 1 min and drying, specimens were observed by a Tecnai G2 F20 S-TWIN field-emission transmission electron microscope (FEI Company, Hillsboro, OR, USA).

3. Results and Discussions

3.1. Microstructural Characteristics of Continuous Cast and Conform Processed Al–Cu–Mg–Zr (–Ag) Alloys

Figure 1a,b show the microstructures of the as-cast A1 (Al–Cu–Mg–Zr) and A2 (Al–Cu–Mg–Zr–Ag) alloys. A2 exhibits a significantly smaller average grain size (GS_avg) (240.48 ± 32.27 μm) than A1 (398.62 ± 69.71 μm), corresponding to a reduction of 158.14 μm (39.67%). A2 also exhibits a notably higher proportion of high-angle grain boundaries (HAGBs, 62.4%) and a larger average grain orientation difference (GOD, 36.1°) compared to A1 (32.3% and 14.9°, respectively). This demonstrates that Ag effectively refines the grain size of as-cast alloys and enhances intergranular orientation differences. After homogenization, the average grain sizes of both alloys increased. A1 exhibited a grain size increment of 28.83 μm (7.23%), rising to 427.45 ± 33.36 μm, whereas A2 showed an increase of 62.73 μm (26.09%), reaching 303.21 ± 23.15 μm. Both the HAGB proportion and average GOD decreased, but A2 retained finer grains, a higher HAGB proportion (47.4%), and a larger average GOD (22.35 °C)—consistent with its as-cast grain characteristics. This indicates that Ag can partially suppress excessive grain growth during homogenization, reduce HAGB loss, and thus maintain intergranular orientation differences. Homogenization exerts the following effects on the alloy grain structure: A1 undergoes substantial grain growth, a significant reduction in HAGB proportion, and diminished orientation differences, resulting in a microstructural homogenization trend. In contrast, while A2 also experiences grain growth, Ag not only inhibits excessive grain coarsening but also preserves a higher HAGB proportion and larger orientation differences—facilitating subsequent microstructural evolution and performance enhancement during Conform. Overall, Ag refines grains and increases HAGB proportion and orientation differences in the as-cast state. After homogenization, Ag retains the superior microstructural features of fine grains and high HAGB content, alleviating the microstructural coarsening/homogenization induced by homogenization.

After Conform (Figure 2a,b), the GS_avg of A2 alloy is 124.81 ± 12.93 μm, which is 10.81 μm (7.97%) finer than that of A1 alloy (135.62 ± 19.93 μm), which confirms that Ag effectively retains its grain-refining effect during Conform. Although the proportion of fine high-angle grain boundaries (F_HAGBs) in the A1 alloy (13.0%) is higher than in the A2 alloy (9.0%), A2 exhibits a larger average orientation difference (θ_avg, 20.78°), reflecting more significant grain boundary orientation differences. This is attributed to Ag-mediated regulation of dynamic recovery/recrystallization during extrusion. After extrusion, both alloys are dominated by low-angle grain boundaries (LAGBs) and contain numerous visible subgrains, with A2 having significantly more subgrains than A1. As shown in Figure 2c,d, the two alloys exhibit distinct microstructural responses to solution treatment. The grains of A1 alloy are significantly refined, with the GS_avg reduced to 61.22 ± 12.67 μm, corresponding to a reduction of 54.86% (74.40 μm), while the F_HAGB proportion increases substantially to 23.1% and θ_avg rises to 20.95°. This indicates that A1 undergoes complete recrystallization during solution treatment, transforming from the “deformation microstructure” of the Conform state to a “recrystallized microstructure”. In contrast, A2 exhibits a much weaker grain refinement effect: its average grain size (GS_avg) decreases by 14.00 μm (11.22%), from 124.81 ± 12.93 μm to 110.81 ± 17.63 μm, the F_HAGB decreases to 7.7% (with LAGBs still dominant), and θ_avg decreases to 14.53°. This suggests that Ag influences element diffusion and grain boundary migration, thereby inhibiting recrystallization kinetics. As a result, the primary microstructural evolution of A2 during solution treatment is grain growth rather than “recrystallization refinement”, retaining more deformation features from the Conform state. The orientation difference distribution (Figure 2e) shows distinct peak shapes for A1 and A2 in the extruded state, reflecting differing post-extrusion orientation structures. After solution treatment, A1 exhibits broader orientation difference peaks with more significant quantitative changes—likely due to increased orientation diversity from recrystallization—while A2 shows more concentrated peaks with smaller variations. This aligns with grain growth-dominated orientation evolution in A2, consistent with the grain and grain boundary characteristics of A2 in the inverse pole figure. After solution treatment, fine grains and a high HAGB proportion of A1 are conducive to plastic deformation. Although A2 has coarser grains post-solution treatment, Ag compensates for grain coarsening through subsequent precipitation hardening (via formation of special strengthening phases), ultimately achieving high strength. Overall, Ag refines grains and increases orientation differences during extrusion, while suppressing recrystallization in A2 during solution treatment. This makes grain growth the dominant mechanism in the microstructural evolution of A2, leading to significant differences in the solution-treated microstructures of the two alloys.

Figure 3 shows the SEM microstructures of the Al–Cu–Mg–Zr (–Ag) alloys under different processing conditions, and

Table 2. Composition of the phases in Figure 3 by EDS (wt.%).

Position/Element (wt.%)	Al	Cu	Mg	Zr	Ag
P1	75.14	23.33	0.56	0.18	/
P2	84.6	14.7	0.50	0.10	/
P3	95.07	4.07	0.43	0.04	/
P4	70.32	28.85	0.35	0.23	/
P5	94.59	4.27	0.41	0.00	0.43
P6	68.28	31.08	0.39	0.00	0.01
P7	72.86	26.41	0.39	0.00	0.00
P8	59.92	38.81	0.24	0.28	0.00
P9	94.48	4.79	0.37	0.04	0.06

presents the chemical compositions of the corresponding phases and the average secondary phase sizes measured by EDS. The area fraction of the secondary phase continuously decreases in both A1 and A2 alloys during the whole process, which is attributed to the synergistic effect of high temperature and mechanical stress; homogenization eliminates the inhomogeneous distribution of the as-cast secondary phases and promotes the dissolution of some coarse phases, and also completely transforms the metastable Al–Cu–Mg ternary phase (Al₂CuMg) unique to A2 into the equilibrium Al₂Cu phase, with Ag atoms originally enriched at the ternary phase boundaries dissolving into the Al matrix, while Conform further accelerates the dissolution and refinement of the dominant Al₂Cu phase through mechanical fragmentation and thermal effects. Solution treatment dissolves the remaining undissolved secondary phases as much as possible, so that most secondary phases dissolve into the matrix in both alloys with only a small number of coarse particles remaining sparsely distributed, verifying the rationality of the heat treatment processes [16]. With the progress of processing, the average secondary phase size shows an overall decreasing trend in both alloys with obvious differences in reduction amplitude: A1 decreases from 1.86 μm to 0.655 μm with a total reduction of 64.8%, among which homogenization provides the largest reduction of 43.8%, and A2 decreases from 1.09 μm to 0.62 μm with a total reduction of 43.1%, while its reduction during solution treatment reaches 25.7%, much higher than the 1.8% of A1. In comparison, the average secondary phase size of A2 is significantly smaller than that of A1 in the as-cast state (1.09 μm vs. 1.86 μm), and A2 remains smaller than A1 in the solution-treated state (0.62 μm vs. 0.655 μm); meanwhile, the secondary phase area fraction of A2 (9.03%) is higher than that of A1 (7.67%) in the as-cast condition, indicating that Ag can promote the nucleation and enrichment of secondary phases in the as-cast microstructure and plays a significant refining effect mainly during solidification and final heat treatment, thereby contributing to the improvement of the comprehensive properties of the alloy.

3.2. Effect of Aging Temperature on the Precipitation Behaviors of Conform Processed Al–Cu–Mg–Zr (–Ag) Alloys

Figure 4 summarizes the age-hardening curves under different aging temperatures. The results show that the age-hardening response of A2 is significantly stronger than that of A1. The peak hardness after T6 treatment (solution treatment + aging) is significantly higher than that after T5 treatment (direct aging). When aged in the range of 170–180 °C, the alloy requires a moderate time to reach peak hardness, with a stable aging process. However, low-temperature aging (160 °C) exhibits slow age-hardening kinetics, while high-temperature aging (190 °C) is relatively sensitive to over-aging. Thus, the optimal aging temperature range is 170–180 °C. The optimal aging processes for achieving peak hardness are as follows: for T5 treatment of A1, the process is 170 °C-16 h, with a hardness of 127 HV; for T6 treatment of A1, it is 170 °C-12 h, with a hardness of 142 HV. For T5 treatment of A2, the process is 180 °C-12 h, with a hardness of 153 HV; for T6 treatment of A2, it is 180 °C-6 h, with a hardness of 160 HV.

Figure 5 illustrates the TEM subgrain morphology of Conform A1 and A2 alloys under various conditions. In the Conform state of the A1 alloy, the subgrains are relatively coarse with coarsely spaced subgrain boundaries. Additionally, the subgrain population is low, subgrains within the grains are incompletely developed, and dislocations are partially retained. After T5 heat treatment, the dislocation density in the alloy matrix decreases significantly, rendering dislocations nearly undetectable. This phenomenon arises because dislocations tend to aggregate and form subgrains, which gradually transform into subgrain boundaries. However, due to the absence of a solution treatment step in the T5 process, grain evolution can only proceed based on the as-processed state. Although precipitates formed during aging pin the subgrain boundaries, suppressing grain growth and promoting subgrain formation, fully developed subgrains are still not achieved. Following T6 heat treatment, subgrain boundary migration and coalescence occur, accompanied by subgrain growth, and well-defined subgrain boundaries are formed within the grains. This is primarily attributed to the solution treatment, which fully dissolves solute atoms, enabling precipitates to strongly pin the subgrain boundaries during the subsequent aging process. Compared with the A1 alloy, the A2 alloy exhibits finer subgrains with sharper boundaries. The main reason is that during the Conform continuous extrusion process, Ag promotes subgrain nucleation and grain refinement, thereby regulating the microstructural evolution of the alloy [17]. After T5 heat treatment, although Ag microalloying provides more nucleation sites for precipitates, accelerating their formation and enhancing subgrain pinning, the aging temperature of T5 treatment is significantly lower than the solution temperature, leading to sluggish atomic diffusion and thus minimal changes in subgrain characteristics. After T6 heat treatment, the microalloying effect of Ag is further enhanced, resulting in fully developed subgrains with distinct boundaries within the grains. This may be explained by the fact that atomic diffusion during solution treatment facilitates the “consumption” of smaller subgrains by larger ones, thereby promoting subgrain growth [18]. Furthermore, during aging, the fine and uniformly dispersed precipitates strongly pin the subgrain boundaries, enabling the complete formation of well-defined subgrain boundaries. This achieves synergistic regulation between precipitates and subgrains, markedly augmenting the microstructural strengthening potential.

As shown in Figure 6, TEM characterization along the [110]_Al crystal direction reveals that no obvious difference in second-phase particles is observed between the two alloys after Conform extrusion and before heat treatment. The matrix of both alloys is mainly a supersaturated solid solution, in which solute atoms predominantly exist in the solid solution state without extensive precipitation nucleation. The corresponding selected area electron diffraction (SAED) patterns show typical diffraction features of the Al matrix, and no additional diffraction spots from massive precipitates are present.

Table 3. Precipitate Area Analysis of Al–Cu–Mg–Zr (–Ag) Alloys under Different Aging Conditions.

Alloy	Condition	S + θ’ (nm2)	Ω (nm2)
Al-Cu-Mg-Zr (A1)	T5 (170 °C-16 h)	2711.36685	/
	T6 (510 °C-1.5 h + 170 °C-12 h)	7231.5368	/
Al-Cu-Mg-Zr-Ag (A2)	T5 (180 °C-12 h)	3253.74	8897.5744
	T6 (510 °C-1.5 h + 180 °C-6 h)	10,266.927	5696.5836

presents the results of the precipitate area analysis of Al–Cu–Mg–Zr (–Ag) alloys under different aging conditions. At the T5 state, the total area of S + θ’ phases in the A1 (Al–Cu–Mg–Zr) alloy is 2711.37 nm². For the Ag-containing A2 (Al–Cu–Mg–Zr–Ag) alloy, the total area of S +θ’ phases is 3253.74 nm², with a large amount of Ω phase (8897.57 nm²). In the A1 alloy (a1) after T5 treatment, the precipitates are mainly acicular and rod-like Al₂CuMg (S phase), accompanied by a small amount of metastable Al₂Cu (θ’ phase) [19]. The SAED pattern confirms the nucleation and growth of these two phases. The microstructure exhibits precipitation-hardening characteristics. However, due to the lack of solution treatment, the solute dissolution is incomplete, which limits the uniformity and refinement degree of the precipitates. Compared with the A1 alloy, the precipitates in the A2 alloy (b1) after T5 treatment are more abundant, including θ’ phase, S phase and Ω phase (Al₂CuMgAg phase) [20,21], with higher density and diversity. The SAED pattern verifies the cooperative precipitation of multiple phases. It indicates that Ag can accelerate the aging precipitation kinetics, provide nucleation sites for the Ω phase, and promote the formation of strengthening phases.

At the T6 state, the area of S + θ’ phases in the A1 alloy increases to 7231.54 nm², with an increment of 166.71%. This is mainly due to the sufficient decomposition of the supersaturated solid solution and the precipitation of solute atoms promoting phase growth during T6 treatment. For the A2 alloy, the area of S + θ’ phases increases to 10,266.93 nm² (increment of 215.55%), while the area of the Ω phase decreases to 5696.58 nm² (reduction of 35.98%). The latter is attributed to the refinement and fracture of the Ω phase during T6 aging, as well as the competition for solute atoms with S + θ’ phases. In terms of microstructure, the precipitates in the A1 alloy (a2) after T6 treatment are finer and more uniform. This benefits from the complete dissolution of solute atoms by solution treatment and the promotion of uniform nucleation and growth by aging treatment, thus improving the strengthening efficiency. The A2 alloy exhibits the characteristics of uniformly dispersed ultra-fine Ω phases and increased number of rod-like S phases. It confirms that the T6 process combined with Ag microalloying can obtain the optimal strengthening microstructure [22,23]. The overlapping of multiple diffraction spots in the SAED pattern further verifies that Ag and T6 process synergistically achieve multi-phase strengthening.

3.3. Mechanical Properties of Heat-Treated Al–Cu–Mg–Zr (–Ag) Alloys

Figure 7 and

Table 4. Tensile properties (ultimate tensile strength, yield strength, and elongation) of Al–Cu–Mg–Zr (–Ag) alloys under different processing and heat treatment conditions.

Alloy	Condition	UTS (MPa)	YS0.2 (MPa)	EI (%)
Al-Cu-Mg-Zr (A1)	Cast	307.77 ± 1.75	144.19 ± 3.97	18.15 ± 0.78
	HT (500 °C-6 h)	361.75 ± 6.72	199.48 ± 0.73	13.92 ± 0.54
	Conform	451.89 ± 3.03	301.38 ± 2.30	14.9 ± 0.14
	T5 (170 °C-16 h)	445.28 ± 2.43	347.75 ± 3.19	13.75 ± 0.21
	T6 (510 °C-1.5 h + 170 °C-12 h)	450.62 ± 3.36	345.77 ± 5.98	15.13 ± 0.47
Al-Cu-Mg-Zr-Ag (A2)	Cast	318.51 ± 2.70	137.5 ± 4.95	17.97 ± 0.71
	HT (500 °C-6 h)	370.83 ± 10.14	193.11 ± 4.40	13.00 ± 0.57
	Conform	480.57 ± 3.64	332.63 ± 4.76	14.54 ± 0.20
	T5 (180 °C-12 h)	494.21 ± 5.95	423.33 ± 3.77	11.25 ± 0.64
	T6 (510 °C-1.5 h + 180 °C-6 h)	511.46 ± 2.06	440.45 ± 6.29	12.79 ± 0.41

summarize the tensile properties of Al–Cu–Mg–Zr (–Ag) alloys under different processing and heat treatment conditions. The results reveal that homogenization treatment renders the as-cast structure more uniform, with uniformly distributed dispersed phases. This contributes to increased tensile and yield strengths of the alloys, but reduced elongation due to decreased coordination of structural deformation. Subsequent Conform after homogenization induces grain refinement and strain hardening via plastic deformation, thereby simultaneously optimizing the strength and elongation of the alloys. The figure also clearly shows that the strength of A2 consistently exceeds that of A1, while the elongation of A2 remains lower than that of A1; additionally, the tensile properties of the alloys after T6 treatment are consistent with their post-aging hardness, consistently surpassing those after T5 treatment. Furthermore, under the identical process of continuous extrusion combined with T6 heat treatment (500 °C solution treatment for 1 h and 180 °C artificial aging for 10 h), the conventional 2024 Al-Cu-Mg alloy exhibits an ultimate tensile strength of 497.6 MPa, a yield strength of 412 MPa and an elongation of 12.93%. In specific contrast, the A2 alloy achieves a higher tensile strength of 511.46 ± 2.06 MPa while maintaining a comparable elongation of 12.79 ± 0.41%, showing a significant strength improvement with almost no loss of ductility [24].

As shown in Figure 8, all tensile fracture surfaces exhibit numerous ductile dimples, which reflects the aggregation of micropores (merged into a single term, material science convention) during plastic deformation. Consequently, the fractures are predominantly ductile. Analysis of the size, depth, and distribution of these dimples enables further assessment of the effects of alloy composition (Ag addition) and aging treatments (T5/T6). The A1 in the Conform state exhibits deeper dimples with more uniform sizes, indicating the highest plasticity among the tested alloys. After T5 aging, the number of ductile dimples decreases while their depth increases, maintaining a certain degree of toughness alongside enhanced strength. T6 aging produces finer and denser ductile dimples, as solution treatment combined with aging results in smaller, more uniform strengthening phases, balancing strength and plasticity. The A2 in the Conform state exhibits deep dimples, though less uniformly distributed than those of A1, leading to higher strength but inferior plasticity. T5 aging of A2 yields finer, denser dimples, attributed to Ag promoting efficient precipitation and refinement of Ω phases, which facilitates the formation of more micropores. T6 aging of A2 produces uniformly fine dimples with denser precipitation of strengthening phases, fully demonstrating the strengthening potential of Ag. Comparative analysis shows that under identical processing conditions, A2 alloys exhibit finer and denser dimples. This indicates that Ag not only enhances strength but also improves the strength–ductility balance by refining strengthening phases, resulting in superior fracture toughness.

3.4. Discussions

Direct solution treatment plus aging of continuous cast Al-Cu-Mg alloys frequently causes microstructural segregation, residual stresses, coarse second phases, and grain size inhomogeneity, resulting in unstable mechanical properties. The combined process of homogenization, Conform forming, and Ag microalloying effectively overcomes these casting defects. Homogenization was performed at 500 °C for 6 h, which promotes diffusion of Cu and Mg atoms from grain boundaries into the matrix, reduces compositional gradients, and eliminates interdendritic segregation. Meanwhile, coarse Al₂Cu phases dissolve into the matrix to form a supersaturated solid solution. In the A1 alloy, the strengthening effect of the θ’ phase is limited due to its low nucleation density. Ag tends to form Mg-Ag co-clusters, segregates on specific crystal planes, and exerts chemical pinning on dislocations to restrain grain boundary migration. This lowers the nucleation energy barrier of the Ω phase and increases its density during post-extrusion aging. Compared with θ’, the Ω phase exhibits higher atomic packing density and stronger coherence with the matrix, leading to superior dislocation blocking ability. The tensile strength increases from 307.77 ± 1.75 MPa to 318.51 ± 2.70 MPa, verifying that Mg-Ag co-clusters can regulate the precipitation pathway of strengthening phases during homogenization [11,25,26]. During casting, the low lattice mismatch between Ag and the Al matrix promotes the preferential precipitation of Ag nanoparticles, which act as effective heterogeneous nucleation sites. The addition of Ag also broadens the undercooled zone at the solidification front, further restraining dendritic arm growth and favoring grain refinement and equiaxed grain formation [27]. As the core deformation step in the synergistic process, Conform achieves grain refinement and precipitate optimization through large strain, high strain rate and dynamic thermal effects. The high temperature and deformation during extrusion create a deformation–thermal coupling environment that favors dynamic recrystallization (DRX). In the early stage of Conform, deformation occurs mainly via dislocation slip, forming massive subgrains. With increasing strain, subgrain boundaries transform into high-angle grain boundaries (HAGBs) through dislocation climb and coalescence, greatly raising the DRX nucleation rate [28]. Most importantly, most undissolved second phases after homogenization—such as Al₂Cu and Al₆(MnFe) phases—are intensively fragmented by shear stresses during extrusion. This eliminates the discontinuity caused by coarse second phases in the matrix while increasing the specific surface area of these phases. Consequently, it provides more heterogeneous nucleation sites for subsequent aging, thereby achieving grain refinement [29]. More crucially, the deformation–heat coupling environment enhances the regulation of Ag-Mg clusters in the Ω phase. High-angle grain boundaries formed during extrusion provide extra nucleation sites for Mg-Ag bilayers at the Ω/matrix interface, inhibiting Cu diffusion and Ω coarsening at elevated temperatures to improve its thermal stability. Extrusion-induced dislocation networks also act as preferential sites for the dynamic precipitation of the Ω phase, which starts to precipitate during extrusion, greatly accelerating strengthening efficiency and enabling precise control of precipitate size [30,31]. The synergistic process of homogenization, Conform forming and Ag microalloying tailors the microstructure of continuous cast Al–Cu–Mg alloys via composition homogenization, dynamic recrystallization, strain-induced precipitation and interface control. Combined grain refinement, precipitation and interface strengthening contribute to the final alloy with high strength and excellent stability [32].

Direct solution treatment and aging of continuous cast Al–Cu–Mg alloys easily lead to microstructural segregation, residual stresses, coarse second phases and grain inhomogeneity, resulting in unstable mechanical properties. The synergistic process of homogenization, Conform forming and Ag microalloying effectively solves these casting-related problems. Homogenization at 510 °C for 6 h promotes the diffusion of Cu and Mg atoms from grain boundaries into grain interiors, reduces the compositional gradient, eliminates interdendritic segregation, and dissolves coarse Al₂Cu phases into the matrix to form a supersaturated solid solution. Meanwhile, trace elements in the alloy form a dual pinning system: Mn precipitates in situ as Al₆(MnFe) dispersed phases or intermetallic compounds, which exert a Zener pinning effect to restrain grain boundary migration and grain growth, thus improving microstructural uniformity [33,34]. Meanwhile, trace elements in the alloy form a dual pinning system: Mn precipitates in situ as Al₆(MnFe) dispersed phases or intermetallic compounds, which exert a Zener pinning effect to restrain grain boundary migration and grain growth, thus improving microstructural uniformity [6]. For the A2 alloy, Ag microalloying exerts distinct effects on precipitation behavior under T5 and T6 treatments. During T5 treatment, Ag atoms aggregate at defects to form nanoclusters in the early stage, providing more nucleation sites for GP zones, raising the nucleation rate and inhibiting precipitate-free zones; with aging, GP zones transform into needle-like S phases and coherent Ω phases, which exhibit much higher strengthening efficiency than the S phase due to stronger matrix coherency [35,36]. During T6 treatment, solution treatment and rapid quenching dissolve Cu, Mg, and Ag into a uniform supersaturated solid solution. Compared with T5, Ag more significantly improves the nucleation efficiency of GP zones in T6 through defect segregation and reduced formation energy. This promotes the formation of hexagonal Ω phases. Meanwhile, GP zones gradually transform into needle-like S precipitates and rod-shaped θ’ precipitates during aging [8,9,10,11,12,13,14,16,17,18,19,20,21,22,23,24,25,26,27,28,29,30,31,32,33,34,35,36,37]. The Ω phase stabilized by Ag-Mg interfacial segregation has strong coherency with the matrix and significant lattice distortion, and its strengthening efficiency is superior to that of S phases and θ’ phases. Ag inhibits the coarsening of precipitates by pinning dislocations and hindering solute diffusion, ensuring the uniform distribution of their sizes. The above mechanisms synergistically endow the T6-state alloy with high strength and excellent thermal stability [32]. The strengthening effect of Ag is fully exerted during T6 aging of the A2 alloy: Ag completely dissolved after solution treatment forms far more Ag-rich nanoclusters in the early stage of aging than in the T5 process, increasing the nucleation density of GP zones by more than 50% compared with Ag-free alloys, achieving uniform distribution, and eliminating the problem of local sparse precipitates [38,39]. With the progress of aging, GP zones transform into S phases and refined Ω phases, constructing a three-dimensional strengthening network.

4. Conclusions

In summary, Al–5Cu–0.4Mg–0.1Zr (–0.4Ag) alloys were successfully fabricated via industrial continuous casting and Conform processes. The effects of Ag addition on microstructure evolution and mechanical properties during processing and heat treatment were systematically investigated. The main conclusions are as follows:

(a)

Ag addition significantly refined the as-cast grains, increased the fraction of high-angle grain boundaries and grain orientation difference, and suppressed excessive grain coarsening during homogenizing annealing.

(b)

During Conform extrusion, Ag further refined the microstructure and increased the subgrain density. The T6 heat treatment exhibited a better strengthening effect than T5. Under the optimized T6 condition, the Ag-containing alloy achieved a tensile strength of 511.46 ± 2.06 MPa and an elongation of 12.79 ± 0.41%.

(c)

After T6 treatment, Ag promoted the sufficient and uniform precipitation of the Ω (Al₂CuMgAg) phase, accelerated the dissolution of second phases, and induced finer and denser ductile dimples, thus contributing to the outstanding strength–ductility synergy.