Effects of Combined Cr, Mn, and Zr Additions on the Microstructure and Mechanical Properties of Al–6Cu Alloys Under Various Heat Treatment Conditions

Abstract

This study investigates the synergistic effects of Cr–Zr and Mn–Zr additions on the microstructural evolution and mechanical properties of Al–6 wt.%Cu alloys. Alloys were designed with solute concentrations positioned below, near, and above their maximum solubility limits, and were evaluated under as-cast, T4, and T6 heat treatment conditions. Mechanical testing revealed distinct behavioral trends depending on the heat treatment: the T4 heat treatment condition generally exhibited superior hardness and yield strength, whereas the T6 heat treatment condition resulted in a slight reduction in hardness but facilitated a significant recovery in tensile strength and structural stability, particularly in alloys designed near the solubility limit. To elucidate the crystallographic origins of these mechanical variations, X-ray diffraction analysis was conducted to monitor changes in lattice parameters, dislocation density, and micro-strain. The results showed that T4 heat treatment induced lattice contraction and a decrease in dislocation density, suggesting that the high strength under T4 heat treatment conditions arises from lattice distortion caused by supersaturated solute atoms. Conversely, T6 aging led to lattice relaxation approaching that of pure aluminum, yet simultaneously triggered a re-accumulation of dislocation density and micro-strain due to the coherency strain fields surrounding precipitates, which effectively impede dislocation motion. Therefore, rather than proposing a single, definitive optimization condition, this study aims to secure foundational data regarding the correlation between these microstructural descriptors and mechanical behavior, providing a guideline for balancing the strengthening contributions in transition metal-modified Al–Cu alloys.

1. Introduction

Precipitation-hardened Al–Cu alloys remain indispensable structural materials for aerospace and transportation applications, where high specific strength and mature processing routes provide an exceptional balance of performance and manufacturability [1,2]. Their strengthening originates primarily from fine, plate-shaped ${\mathsf{\theta }}^{\prime }$ precipitates that form along the canonical sequence SSSS → GP zones → ${\mathsf{\theta }}^{″}$→ ${\mathsf{\theta }}^{\prime }$ → $\mathsf{\theta }$ [3,4]. A high number density of thin ${\mathsf{\theta }}^{\prime }$ plates effectively impedes dislocation glide, thereby imparting superior strength. However, this strengthening is intrinsically unstable, as rapid diffusion at service temperatures above 250 °C accelerates ${\mathsf{\theta }}^{\prime }$ coarsening and its transformation to the equilibrium $\mathsf{\theta }$ phase with a reduced aspect ratio [5]. The result is a catastrophic loss of strength that has long defined the upper temperature limit of conventional Al–Cu alloys. Overcoming this thermal stability ceiling is therefore a critical and unresolved challenge in alloy design.

To address this limitation, two primary strategies have been explored. The first strategy involves introducing thermally stable dispersoids to suppress coarsening at elevated temperatures [6,7,8,9,10,11], while the second strategy focuses on directly stabilizing ${\mathsf{\theta }}^{\prime }$ precipitates through micro-alloying with transition metals. Among the most promising examples of this latter approach, Mn and Zr co-additions have been reported to extend strengthening up to 350 °C [12,13,14,15,16]. Their synergy arises from a unique interfacial mechanism. Mn kinetically prolongs the persistence of ${\mathsf{\theta }}^{\prime }$, enabling sluggish Zr atoms to segregate to Al/${\mathsf{\theta }}^{\prime }$ interfaces, reduce interfacial energy, and form a core–shell structure stabilization effect. Similarly, computational studies have identified Cr as another potent candidate due to its strong tendency to segregate at Al/${\mathsf{\theta }}^{\prime }$ interfaces, low cost, and ability to suppress deleterious phases [17]. Aligning with this second strategy, the present study investigates the synergistic effects of these transition metals on the microstructural evolution and mechanical properties of Al–Cu alloys.

While previous studies have successfully confirmed that Cr–Zr and Mn–Zr combinations can maintain hardness stability at elevated temperatures up to 350 °C [17], there remains a critical lack of comprehensive data regarding their overall mechanical properties at room temperature. Furthermore, the fundamental relationships between varying solute contents particularly relative to their solubility limits and the resulting mechanical performance have not been fully elucidated. This absence of systematic data hinders the optimization of alloy compositions for broader industrial applications.

Consequently, the objective of this work is to elucidate the synergistic effects of Cr–Zr and Mn–Zr combination additions on the microstructural evolution and mechanical properties of Al–6 wt.% Cu alloys relative to their solubility limits. Mechanical properties were systematically evaluated under As-cast, T4, and T6 heat treatment conditions to determine the influence of thermal history, while scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS) were employed to investigate the morphological characteristics and phase compositions of the precipitates. Furthermore, X-ray diffraction (XRD) analysis was performed to quantify crystallographic parameters, including lattice parameter, micro-strain, and dislocation density, thereby linking microstructural defects to the observed deformation behaviors. Finally, the strengthening mechanisms were discussed by examining the relationships between the micro-alloying content, microstructural descriptors, and mechanical performance. The findings of this study the interplay between solute supersaturation and precipitate morphology, providing practical guidance for alloy design and serving as foundational data for future evaluations of high-temperature mechanical properties. Finally, by analyzing how varying alloying contents and heat treatment conditions influence microstructural evolution, this study provides a guideline for alloy design via transition metal additions.

2. Materials and Methods

2.1. Materials

The investigated alloys were based on an Al–6 wt.% Cu matrix prepared using master alloys of Al–30 wt.% Cu, Al–5 wt.% Cr, Al–10 wt.% Mn, and Al–5 wt.% Zr, together with commercially pure Al (99.8 wt.%). Approximately 8 kg per batch was melted in a graphite crucible inside an electric resistance furnace. The addition levels of Cr, Mn, and Zr were precisely controlled relative to their maximum solid solubility in the Al matrix at 540 °C. The maximum solid solubility of each transition element in the Al-6 wt.%Cu alloy was calculated using JMatPro software V13.0, as illustrated in Figure 1. Based on these calculations, the alloy compositions were designed to represent three distinct solute levels, namely below the solubility limit (0.13 wt.% Cr/Mn–0.07 wt.% Zr), near the maximum solubility limit (0.25 wt.% Cr/Mn–0.13 wt.% Zr), and above the solubility limit (0.40 wt.% Cr/Mn–0.20 wt.% Zr). This strategic variation was intended to systematically evaluate the effects of solute supersaturation and excess intermetallic formation during solidification and subsequent heat treatment. The melts were held at 750 °C for 30 min, followed by degassing with high-purity Argon gas for 30 min and cast into a permanent steel mold to ensure rapid cooling to ambient temperature. The cooling rate was calculated to be 168 K/min based on temperature measurements recorded using a data logger (Graphtec GL240, Graphtec, Yokohama, Japan). The actual chemical compositions were verified by laser-induced breakdown spectroscopy (LIBS), and the results are summarized in

Table 1. Chemical compositions of Al–6Cu alloys with different micro-alloying treatments (wt.%).

Alloy	Cu	Cr	Mn	Zr	Fe	Si	Al
2CZ	6.06	0.13	-	0.08	0.15	0.06	Bal.
4CZ *	6.10	0.25	-	0.14	0.15	0.06	Bal.
6CZ	5.99	0.43	-	0.19	0.15	0.06	Bal.
2MZ	6.07	-	0.13	0.07	0.14	0.06	Bal.
4MZ *	6.02	-	0.25	0.13	0.15	0.06	Bal.
6MZ	6.05	-	0.41	0.20	0.15	0.06	Bal.

* Maximum solubility in the Al–6Cu matrix.

.

2.2. Heat Treatment Conditions

Three metallurgical conditions were investigated to systematically evaluate the effects of heat treatment on microstructural evolution and mechanical properties. First, the as-cast condition was characterized to establish a baseline. Second, T4 heat treatment was conducted by solutionizing the samples at 540 °C for 6 h, followed by immediate water quenching to room temperature. This process was designed to dissolve coarse intermetallic phases and achieve a supersaturated solid solution. T6 heat treatment consists of the T4 heat treatment followed by artificial aging at 240 °C for 6 h. The aging temperature of 240 °C, which is relatively higher than conventional aging temperatures, was specifically selected based on prior research [14]. Based on prior research [14], a relatively elevated aging temperature of 240 °C was strategically selected to promote the formation of thermally stable θ′ precipitates that exhibit superior coarsening resistance up to 350 °C. Consequently, this specific condition was adopted to establish a microstructural baseline for subsequent investigations into high-temperature mechanical properties.

2.3. Characterization of Microstructure

For microstructural characterization, the specimens were mounted and mechanically ground using silicon carbide (SiC) papers up to #2000grit. Subsequently, the samples were metallographically polished using 3 μm and 1 μm diamond suspensions, followed by a final polishing step with 0.04 μm colloidal silica to achieve a mirror-finish surface. The microstructures were observed using a field-emission scanning electron microscope (FE-SEM, Clara, Tescan, Brno, Czech Republic) located at the Eco-friendly Shipbuilding Core Research Support Center. High-resolution imaging was conducted at an acceleration voltage of 5 kV to ensure surface clarity, while energy-dispersive X-ray spectroscopy (EDS) analysis was performed at an acceleration voltage of 15 kV for precise phase composition analysis. X-ray diffraction (XRD, Smartlab, Rigaku, Tokyo, Japan) was employed to evaluate the micro-strain of the α–Al matrix using CuKα radiation (λ = 1.54 Å) generated at 40 kV and 30 mA. The diffraction patterns were recorded over a 2$\mathsf{\theta }$ range of 20° to 120° with a scanning speed of 5°/min and a step size of 0.02°.

2.4. Mechanical Properties Test

The mechanical properties of the investigated alloys were evaluated through Vickers hardness and tensile tests. Vickers hardness tests were conducted using a Vickers hardness tester (HV-110D, Mitutoyo, Tokyo, Japan). The measurements were performed on metallographically polished specimen surfaces to ensure clarity. A load of 10 kgf was applied with a dwell time of 10 s. To ensure statistical reliability, ten indentations were performed for each condition, and average values with standard deviations were reported. Tensile tests were carried out at room temperature using a universal testing machine (KDMT-158, Kyung Do Precision, Gyeonggi-do, Republic of Korea). The tensile specimens were machined according to the specific geometry shown in Figure 2. The tests were conducted at a constant crosshead speed of 1 mm/min. To guarantee reproducibility, a total of six specimens were tested for each condition.

3. Results

3.1. Microstructural Evolution

Figure 3 illustrates the microstructural evolution of the CZ combinations (2CZ, 4CZ, and 6CZ) as a function of composition and heat treatment conditions (as-cast, T4, and T6).

Table 2. The area fraction of precipitated intermetallic compounds within the Al matrix of CZ alloys (%).

Alloy	As-Cast	T4	T6
2CZ	3.45 (±0.74)	1.46 (±0.33)	1.47 (±0.45)
4CZ	3.56 (±0.56)	0.82 (±0.27)	0.94 (±0.30)
6CZ	4.08 (±0.64)	0.91 (±0.22)	2.25 (±0.15)

provides the area fraction of intermetallic compounds measured to estimate the general trend of phase evolution, and

Table 3. EDS results for each point of the CZ alloy (wt.%).

No.	Al	Cu	Fe	Phase
1~11	39–82	7–52	-	Al2Cu
12~23	45–78	11–34	3–11	Al7Cu2Fe

presents the corresponding EDS point analysis results. In the as-cast condition, the microstructure is characterized by an α–Al matrix and bright eutectic Al₂Cu phases forming a continuous network along the dendritic cell boundaries, resulting in a relatively high area fraction of the secondary phases. Furthermore, it was confirmed that the main constituent phases remained identical even as the alloying content increased from 2CZ to 6CZ, predominantly consisting of the α–Al matrix and eutectic Al₂Cu phases. Regarding the EDS analysis, regions exhibiting elevated Cu concentrations were attributed to solute segregation during solidification. Consistent with prior findings [19,20], these phases were identified as eutectic Al₂Cu. Notably, precipitates containing Cr or Zr were not detected in the observed microstructures. Upon T4 heat treatment, significant microstructural changes were confirmed as the coarse eutectic Al₂Cu network fragmented and either dissolved into the matrix or remained as fine spherical particles, which led to a substantial reduction in the measured area fraction. Conversely, while the area fractions remained similar between the T4 and T6 heat treatment conditions up to the 4CZ alloy, a distinct increase in the area fraction was observed in the 6CZ–T6 condition. This increase indicated the formation of precipitates, and it was also confirmed that the Al₇Cu₂Fe phase grew into distinct needle- and rod-like shapes.

Figure 4 illustrates the microstructural evolution of the MZ combinations (2MZ, 4MZ, and 6MZ) as a function of composition and heat treatment conditions (as-cast, T4, and T6).

Table 4. The area fraction of precipitated intermetallic compounds within the Al matrix of MZ alloys (%).

Alloy	As-Cast	T4	T6
2MZ	4.69 (±0.92)	1.23 (±0.11)	1.45 (±0.14)
4MZ	4.28 (±0.48)	0.44 (±0.10)	1.01 (±0.26)
6MZ	4.62 (±0.44)	1.26 (±0.23)	1.59 (±0.32)

provides the area fraction of intermetallic compounds measured to estimate the general trend of phase evolution, and

Table 5. EDS results for each point of the MZ alloy (wt.%).

No.	Al	Cu	Fe	Mn	Phase
1~12	40–87	5–51	-	-	Al2Cu
13~19	45–74	15–47	2–9	-	Al7Cu2Fe
20~31	46–63	21–34	4–9	1–4	Al–Cu–Fe–Mn

presents the corresponding EDS point analysis results. In the as-cast condition, the microstructure is characterized by an α–Al matrix and bright eutectic Al₂Cu phases forming a continuous network along the dendritic cell boundaries, resulting in a relatively high area fraction of the secondary phases. Furthermore, it was confirmed that the main constituent phases remained identical even as the alloying content increased from 2MZ to 6MZ, predominantly consisting of the α–Al matrix and eutectic Al₂Cu phases. Mn containing intermetallic compounds was not significantly detected in the as-cast state. Upon T4 heat treatment, significant microstructural changes occurred as the coarse eutectic Al₂Cu network fragmented. However, as the heat treatment progressed, EDS analysis confirmed the precipitation of a significant amount of Al–Cu–Fe–Mn intermetallic compounds. Since the precise crystallographic identification of these complex compounds has not yet been fully established in current research, they are collectively referred to as the ‘Al–Cu–Fe–Mn phase’ in this study [21]. In the T6 heat treatment condition, these Al–Cu–Fe–Mn phases were observed to predominantly form distinct needle- and rod-like shapes.

3.2. Mechanical Properties

Figure 5 illustrates the evolution of Vickers hardness for the investigated alloys under as-cast, T4, and T6 heat treatment conditions. A distinct dependence on thermal history was observed across all compositions. In the as-cast state, the hardness values were relatively low, generally ranging between 50 and 80 HV. However, the application of heat treatment resulted in a substantial enhancement in hardness for all alloy systems. Notably, the T4 heat treatment condition consistently yielded the highest hardness values, surpassing those observed in the T6 heat treatment condition. For instance, the 4CZ and 4MZ alloys exhibited significant increases to 119.7 HV and 108.0 HV, respectively. Furthermore, a slight upward trend in hardness was noted in the T4 heat treatment condition as the alloying content increased (from 2CZ to 6CZ and 2MZ to 6MZ), suggesting a compositional dependency driven by solid solution strengthening. In contrast, the hardness values in the T6 heat treatment condition decreased relative to the T4 heat treatment condition, although they remained significantly higher than the as-cast baseline. Despite this reduction, the 6CZ and 6MZ alloys maintained hardness values of 86.4 HV and 86.8 HV, respectively. When comparing the two alloy systems in the T6 heat treatment condition, the hardness levels were generally comparable, with the MZ series exhibiting slightly higher stability.

The variations in Ultimate Tensile Strength (UTS) for the CZ and MZ alloy series are summarized in Figure 6. Regardless of the alloying elements, the heat-treated samples demonstrated a substantial improvement in UTS compared to the as-cast state. For the CZ combinations (Figure 6a), the UTS under the T6 condition showed a progressive increase with higher alloying content. Specifically, the 6CZ alloy achieved a maximum UTS of 347.5 MPa after T6 treatment, which significantly surpassed the values obtained in the as-cast (181.9 MPa) and T4 (314.1 MPa) conditions. This indicates that for high-solute CZ alloys, artificial aging is crucial for maximizing strength. Similarly, the MZ combinations (Figure 6b) exhibited excellent strength characteristics. The 4MZ and 6MZ alloys subjected to T6 aging reached UTS values of 329.7 MPa and 343.8 MPa, respectively. It is noteworthy that while the T4 condition previously exhibited the highest micro-hardness, the T6 condition generally provided superior UTS values, particularly for compositions with higher solute contents (e.g., 4CZ, 6CZ, and 6MZ). This suggests that artificial aging effectively enhances the maximum load-bearing capacity through precipitation strengthening [22].

The variations in yield strength (YS) and elongation for the CZ and MZ alloy series are summarized in Figure 7 and Figure 8, respectively. To evaluate the structural performance, the relationship between strength and ductility was analyzed with a focus on the trade-off behavior typically observed in precipitation-hardened alloys. As shown in Figure 7, the CZ series exhibited a classic inverse correlation between strength and ductility. For the 4CZ alloy, the T4 condition resulted in the lowest yield strength (115.8 MPa) but the highest elongation (37.7%), indicating a matrix dominated by dislocation glide capability. However, upon T6 aging, the yield strength surged to 260.8 MPa while the elongation dropped sharply to 15.7%. This drastic shift demonstrates that in the CZ system, strength gains are achieved at the direct expense of ductility. In distinct contrast, the MZ series, particularly the 4MZ alloy (Figure 8), demonstrated an exceptional capability to overcome this conventional trade-off. Under the T4 condition, the 4MZ alloy showed a yield strength of 185.8 MPa and an elongation of 22.9%. Remarkably, the T6 aging treatment triggered a simultaneous enhancement in both properties with the yield strength increased to 250.5 MPa, and the elongation improved to 27.4%. Unlike the CZ series, where aging significantly reduced ductility, the 4MZ–T6 condition achieved a superior strength–ductility balance. This suggests that the microstructural evolution in the 4MZ–T6 alloy effectively manages internal stress, maintaining high deformability even at elevated strength levels.

4. Discussion

This work aimed to demonstrate the synergistic effects of Cr–Zr and Mn–Zr additions on the microstructural evolution and mechanical properties of Al–6 wt.%Cu alloys relative to their solubility limits. Microstructural analysis revealed distinct phase formation behaviors depending on the alloying elements. Specifically, Cr and Zr precipitates were not detected, which is attributed to an insufficient driving force for precipitation under the specific heat treatment conditions employed in this study [23]. In contrast, Mn was confirmed to form needle- or rod-shaped Al–Cu–Fe–Mn phases by combining with Fe. Building on these observations, mechanical assessments and microstructural observations suggested that distinct behavioral trends, specifically the variation in hardness and the strength–ductility balance, were strictly driven by phase transformations during heat treatment. To provide a fundamental explanation for these macroscopic properties, the variations in internal defect structures with respect to alloying content and heat treatment conditions were identified through XRD line-broadening analysis. Consequently, this investigation elucidates the correlation between crystallographic parameters, such as lattice parameter and dislocation density, and the resultant mechanical performance.

4.1. Correlation Between Lattice Parameter and Strengthening Mechanisms

The variation in lattice parameter serves as a critical indicator of the dominant strengthening mechanism for each heat treatment condition. Here, the accurate lattice parameter value was calculated from the lattice parameter of different ($\mathrm{hkl}$) reflections by the Nelson–Riley extrapolation function. The Nelson–Riley function was calculated through Equation (1), where $\mathsf{\theta }$ is the Bragg angle [24].

$$\mathrm{F}\left(\mathsf{\theta }\right)={\displaystyle \frac{1}{2}}\left({\displaystyle \frac{{\cos }^2\mathsf{\theta }}{\sin \mathsf{\theta }}}+{\displaystyle \frac{{\cos }^2\mathsf{\theta }}{\mathsf{\theta }}}\right)$$

(1)

The variations in lattice parameter are shown in Figure 9. In the T4 heat treatment condition, a decrease in the lattice parameter was observed compared to the as-cast state. This reduction is attributed to the dissolution of coarse eutectic phases into the aluminum matrix. Although solute atoms dissolve to form a supersaturated solid solution, the lattice relaxation effect resulting from the dissolution of large precipitates dominates, leading to a net decrease in the lattice parameter. Consequently, solid solution strengthening becomes the primary mechanism. The solute atoms increase the lattice friction stress, effectively impeding dislocation motion and resulting in the highest Vickers hardness values observed in this study [25]. Conversely, in the T6 heat treatment condition, the lattice parameter exhibited a renewed increase. This expansion is likely attributed to the formation of lattice strain fields associated with the coherent or semi-coherent interfaces of newly formed precipitates, such as ${\mathsf{\theta }}^{\prime }$ (Al₂Cu) and transition metal compounds (Al₇Cu₂Fe, Al–Cu–Mn–Fe) [26]. In this stage, precipitation strengthening (Orowan mechanism) appears to be the dominant factor, where the fine and dense dispersion of precipitates forces dislocations to bypass them, thereby significantly enhancing the yield strength [27].

4.2. Influence of Dislocation Density on Deformation Behavior

Dislocation density served as the primary quantitative metric for strengthening mechanisms, while micro-strain was analyzed as a supplementary indicator to assess the general trend of lattice distortion induced by thermal processing. Therefore, the micro-strain values were interpreted qualitatively to understand the relaxation or accumulation of internal stresses relative to the heat treatment condition, rather than being treated as absolute governing factors for mechanical properties. First, the micro-strain (σ) was calculated using the Size Strain Plot (SSP) method according to Equation (2) [28]:

$${\left({\displaystyle \frac{{\mathrm{d}}_{\mathrm{hkl}}{\mathsf{\beta }}_{\mathrm{hkl}}\cos \mathsf{\theta }}{\mathsf{\lambda }}}\right)}^2={\displaystyle \frac{\mathrm{K}}{\mathrm{D}}}\left({\displaystyle \frac{{{\mathrm{d}}_{\mathrm{hkl}}}^2{\mathsf{\beta }}_{\mathrm{hkl}}\cos \mathsf{\theta }}{\mathsf{\lambda }}}\right)+{\left({\displaystyle \frac{\mathsf{\sigma }}{2}}\right)}^2$$

(2)

where d_hkl is the interplanar spacing, β_hkl is the integral breadth of the diffraction peak, $\mathsf{\theta }$ is the Bragg angle, K is a constant, λ is the X-ray wavelength, and D is the crystallite size.

As shown in

Table 6. Micro-strain (σ) of the investigated alloys under different heat treatment conditions.

	Micro-Strain (σ) (×10−3)
Alloys	As-Cast	T4	T6
2CZ	3.73	1.18	3.27
4CZ	2.48	2.28	4.68
6CZ	4.19	2.27	3.78
2MZ	4.36	2.64	3.82
4MZ	3.88	1.69	3.88
6MZ	4.10	1.97	4.22

, the micro-strain values exhibited distinct variations depending on the heat treatment. In the T4 heat treatment condition, the micro-strain was significantly lower compared to the as-cast and T6 heat treatment conditions across all compositions. For instance, the 4CZ alloy showed a minimum micro-strain of approximately 2.28 × 10⁻³ in the T4 heat treatment condition, which is considerably lower than the values observed in the as-cast (2.48 × 10⁻³) and T6 (4.68 × 10⁻³) heat treatment conditions. Similarly, the 4MZ alloy exhibited the lowest micro-strain of 1.69 × 10⁻³ after T4 treatment. This reduction in lattice strain is attributed to the relief of internal stresses and the homogenization of the solute distribution during the high-temperature solution treatment [29].

The dislocation density results for each alloy are presented in

Table 7. Dislocation density of the investigated alloys under different heat treatment conditions.

	Dislocastion Density (×1012 m−2)
Alloys	As-Cast	T4	T6
2CZ	2.14	1.45	4.76
4CZ	1.66	0.92	6.98
6CZ	10.3	1.38	2.52
2MZ	3.84	1.89	2.72
4MZ	2.25	0.45	2.03
6MZ	3.46	0.75	2.56

. Subsequently, the dislocation density ($\mathsf{\rho }$) was determined using Equation (3), as shown below, based on the Williamson–Hall method [30].

$$∆\mathrm{K}={\displaystyle \frac{0.9}{\mathrm{D}}}+\left({\displaystyle \frac{\mathsf{\pi }\mathrm{M}{\mathrm{b}}^2}{2}}\right)·{\mathsf{\rho }}^{{\displaystyle \frac{1}{2}}}·{\mathrm{K}}^2{\stackrel{\mathrm{-}}{\mathrm{C}}}_{\mathrm{hkl}}$$

(3)

where D is the crystallite size, M is the Wilkens arrangement parameter characterizing the dislocation distribution (with a value of approximately 1–2), and b is the Burgers vector of the Al alloy (approximately 0.286 nm). $∆\mathrm{K}$ and $\mathrm{K}$ are the full width at half maximum (FWHM) of the broadened peak in reciprocal space and the diffraction vector, respectively, which are given by

$$∆\mathrm{K}={\displaystyle \frac{\cos \mathsf{\theta }[∆2\mathsf{\theta }]}{\mathsf{\lambda }}},\mathrm{K}={\displaystyle \frac{2\sin \mathsf{\theta }}{\mathsf{\lambda }}}$$

(4)

where $\mathsf{\theta }$ is the diffraction angle, $∆2\mathsf{\theta }$ is the FWHM of the diffraction peak at $\mathsf{\theta }$, and $\mathsf{\lambda }$ is the wavelength of the X-ray (0.15406 nm).

${\stackrel{\mathrm{-}}{\mathrm{C}}}_{\mathrm{hkl}}$ is the average contrast factor for each specific $\mathrm{hkl}$ plane. $\mathrm{h}$, $\mathrm{k}$, and $\mathrm{l}$ are the Miller indices of each peak. The variation in the average contrast factor for different peaks is used to account for broadening anisotropy [30]:

$${\stackrel{\mathrm{-}}{\mathrm{C}}}_{\mathrm{hkl}}={\stackrel{\mathrm{-}}{\mathrm{C}}}_{\mathrm{h}00}(1 - \mathrm{q}{\mathrm{H}}^2)$$

(5)

The value of ${\mathrm{H}}^2$ depends only on the crystal diffraction plane indices. The values of ${\stackrel{\mathrm{-}}{\mathrm{C}}}_{\mathrm{h}00}$ and $\mathrm{q}$ are influenced by the crystal structure and the material anisotropic elastic constants. For the aims of this study, a feasible estimation of dislocation density can be achieved by employing a simplified homogeneous methodology (Equation (3)) to calculate contrast factors. This is assumed to be 1 for simplicity in this study.

Consistent with the micro-strain results, the dislocation density in the T4 heat treatment condition was found to be the lowest among the tested conditions, generally remaining in the range of 0.45–1.89 (×10¹² m⁻²). Specifically, the 4CZ–T4 alloy exhibited a dislocation density of 0.92 (×10¹² m⁻²), and the 4MZ–T4 alloy showed an even lower value of 0.45 (×10¹² m⁻²). The exceptionally low dislocation density and micro-strain observed in the 4CZ–T4 condition create an environment where mobile dislocations can travel long distances without significant obstruction. This long mean free path promotes extensive plastic deformation. The fine clusters formed near the solubility limit are inferred to be shearable, which delays dislocation pile up and necking, resulting in a remarkable elongation of 37.7% despite the lower yield strength (115.8 MPa). In contrast, the T6 heat treatment condition resulted in a sharp increase in both micro-strain and dislocation density. For the 4CZ alloy, the dislocation density increased to 6.98 (×10¹² m⁻²) after T6 aging. This significant rise is can be explained by the generation of Geometrically Necessary Dislocations (GNDs) around non-shear able precipitates to accommodate the lattice misfit during precipitation [31]. While these high-density dislocations contribute to strength via work hardening and interaction hardening, they simultaneously restrict the mean free path of mobile dislocations [32], leading to a reduction in ductility compared to the T4 heat treatment condition.

4.3. Effect of Precipitate Morphology on Ductility Degradation

A critical factor influencing the ductility, particularly the premature failure observed in certain conditions (Figure 3i and Figure 4e), is the morphology of the precipitates. As revealed in the SEM analysis (Figure 3 and Figure 4), the formation of needle- or rod-like phases (Al₇Cu₂Fe, Al–Cu–Mn–Fe) significantly affects the fracture mechanism. Needle-like or high-aspect-ratio precipitates act as severe stress concentrators within the soft aluminum matrix. During tensile deformation, dislocations pile up rapidly at the interfaces of these incoherent or semi-coherent elongated particles. The local stress at the tips of these needles can easily exceed the critical fracture strength of the particle or the interface, leading to micro-void nucleation or particle cracking at low strain levels [33].

The 4MZ alloy in the T4 heat treatment condition exhibited a uniquely poor balance of mechanical properties, showing a reduced yield strength (185.8 MPa) combined with limited elongation (22.9%), unlike the highly ductile 4CZ–T4. Although the matrix was soft due to low solid solution strengthening, the microstructure contained distinct needle- like Mn-rich phases, as shown in Figure 4e. These needles are considered to facilitate early crack initiation via stress concentration, preventing the matrix from fully utilizing its plastic deformation potential. This indicates that even with a low dislocation density, the morphological factor of second-phase particles can play a dominate role in the failure process. Similarly, the 4CZ alloy after T6 aging showed a recovery in strength but a significant drop in elongation. This can be attributed to the reprecipitation and coarsening of Al₂Cu and Al₇Cu₂Fe into rod shapes (Figure 3i). The combined effect of high dislocation density (restricting slip) and morphological stress concentration (promoting void nucleation) accelerated the transition to fracture [34].

Fractographic analysis was performed to investigate the failure mechanisms. Since variations in fracture characteristics based on alloying content were indistinct, the analysis focused on the clear transitions observed across heat treatment conditions. As presented in Figure 10, distinct fracture features were observed depending on the thermal history. In the as-cast condition, the fracture surfaces exhibited brittle features, such as cleavage facets and cracks along dendrite cell boundaries, caused by stress concentration at the coarse eutectic Al₂Cu network. Upon T4 heat treatment, the fracture mode transitioned to a ductile rupture characterized by deep and uniform dimples, indicating that the dissolution of coarse intermetallics enhanced plastic deformation. In contrast, the T6 heat treatment condition revealed a mixed fracture mode consisting of both brittle and ductile features. While ductile dimples were present, they were frequently accompanied by cleavage facets and micro-voids initiated at precipitate interfaces. This observation serves as fractographic evidence aligning with the microstructural analysis: the lattice expansion induced by precipitation necessitates a significant increase in dislocation density to accommodate the lattice misfit. The presence of brittle features indicated that these accumulated dislocations restricted plastic flow, thereby confirming the correlation between the increased lattice parameter and the elevated dislocation density.

4.4. Optimization of Mechanical Performance

The 4MZ–T6 condition emerged as the optimal processing route, achieving a high yield strength of 250.5 MPa and a superior elongation of 27.4%. This balanced performance is particularly noteworthy when compared to recent reports on similar transition-metal-modified Al–Cu alloys. For instance, M. Amer et al. reported that Al–Cu–Y–Zr alloys with the addition of Cr achieved higher yield strengths (about 308–315 MPa) but suffered from significantly limited elongation (2.0–3.3%) due to the formation of complex intermetallics [33]. S.K. Kairy et al. demonstrated that while Sc and Zr additions significantly increased hardness and strength by refining ${\theta }^{\prime }$ precipitates and forming core–shell Al₃(Sc, Zr) particles [35], achieving high ductility alongside strengthening remains a challenge. In contrast, the current 4MZ–T6 alloy exhibits a distinct advantage in maintaining high plasticity (27.4%), suggesting that the specific Mn–Zr combination offers a more favorable pathway for overcoming the strength–ductility trade-off.

This performance was attributed to a synergistic effect. A high lattice parameter indicated robust precipitation strengthening by fine, dispersed particles [36]. Controlled precipitate morphology is unlike the T4 heat treatment condition. T6 heat treatment in the 4MZ alloy is thought to promote the formation of precipitates that effectively pin dislocations (high strength) while potentially maintaining a morphology or distribution that mitigates catastrophic stress concentration compared to the coarse needles found in other conditions. In these phases, dislocations are accumulated, and the increased dislocation density supported work hardening, delaying instability [37].

In contrast to the tensile improvements, the hardness values in the T6 heat treatment condition were observed to be lower than those in the T4 heat treatment condition. This decrease in hardness was attributed to the over-aging effect induced by the high aging temperature. During the T6 artificial aging process, as the aging time presumably surpassed the optimal peak point, the strengthening precipitates (e.g., ${\theta }^{\prime }$) coarsened from a fine and uniform state, thereby reducing the resistance to dislocation movement. With prolonged aging, these precipitates tended to transform into stable phases ($\theta$) diminishing the hardening effect, while grain growth at the high aging temperature reduced the Hall–Petch effect. These microstructural changes indicated that the alloy entered the over-aging regime during the T6 heat treatment [38,39,40,41]. However, to explicitly confirm the loss of coherency and the specific phase evolution associated with this over-aging phenomenon, further microstructural analysis via Transmission Electron Microscopy (TEM) is required. TEM investigation is essential to rigorously verify the crystallographic coherency between the precipitates and the matrix, as well as to quantitatively characterize their size and distribution.

5. Conclusions

In this study, the individual and synergistic effects of Cr–Zr and Mn–Zr additions on the microstructural evolution and mechanical properties of Al–6 wt.% Cu alloys were investigated under various heat treatment conditions. The notable conclusions are summarized as follows:

(1)

In the as-cast state, the microstructure was characterized by a continuous network of coarse eutectic Al₂Cu phases. The T4 heat treatment effectively fragmented and dissolved these coarse phases. During T6 artificial aging at 240 °C, the precipitation behavior varied by composition; notably, in the MZ combinations, EDS analysis confirmed the formation of Al–Cu–Fe–Mn intermetallic compounds, which were observed to grow into needle- and rod-like shapes morphologies during aging.

(2)

The micro-hardness in the T4 heat treatment condition was consistently higher than that in the T6 condition across all alloys. This trend is attributed to the relatively high aging temperature (240 °C), which likely induced an over-aging effect. This condition led to the coarsening of strengthening precipitates and a reduction in coherency strain, resulting in lower hardness compared to the T4 heat treatment state, where solid solution strengthening was maximized.

(3)

A distinct difference in tensile behavior was observed between the two alloy systems. The CZ series exhibited a typical trade-off where T6 aging increased strength but significantly reduced ductility. In contrast, the MZ series, particularly the 4MZ–T6 alloy, achieved a superior balance of mechanical properties (yield strength: 250.5 MPa; elongation: 27.4%). This indicates that the specific combination of Mn and Zr effectively enhances strength while maintaining high ductility.

(4)

XRD line-broadening analysis revealed that the T6 aging treatment induced lattice expansion, which necessitated a significant increase in dislocation density to accommodate the lattice misfit. This accumulation of defects contributed to the strengthening mechanism.

(5)

The study demonstrates that mechanical performance is dependent on the optimization of alloying elements and heat treatment conditions. The 4MZ–T6 condition emerged as the optimal processing route, demonstrating that controlling the precipitation of transition metal phases (such as Al–Cu–Fe–Mn) and managing dislocation density are critical strategies for the effective design of Al–Cu alloy systems.