Heat Treatment Analysis and Mechanical Characterization of a Recycled Gravity Die Cast EN 42000 Alloy

Abstract

Recycled aluminum–silicon alloys provide significant environmental benefits by reducing the consumption of raw materials and lowering carbon emissions. However, their industrial application is limited by the presence of iron-based intermetallic compounds and the insufficient investigation in the literature regarding their effects on mechanical behavior. This study focuses on a recycled EN 42000 alloy, comprising 95% recycled aluminum, with a focus on the effect of its elevated iron content (0.447 wt%) on aging behavior and mechanical performance. Laboratory-scale specimens were produced through gravity die casting and subjected to T6 heat treatment, consisting of solution, quenching, and artificial aging from 160 °C to 190 °C for up to 8 h. To investigate overaging, analyses were conducted at 160 °C and 170 °C for durations up to 184 h. Tensile tests were conducted on specimens aged under the most promising conditions. Based on innovative quality indices and predictive modeling, aging at 160 °C for 4.5 h was identified as the optimal condition, providing a well-balanced combination of strength and ductility (YS = 258 MPa, UTS = 313 MPa, and e% = 3.9%). Mechanical behavior was also assessed through microstructural and fractographic analyses, highlighting the capability of EN 42000 to achieve properties suitable for high-performance automotive components.

1. Introduction

Aluminum is the most extensively produced non-ferrous metal worldwide, with applications across a wide range of industries, including the energy, aerospace, and automotive sectors ([1,2]). Aluminum alloys are particularly important in the transport industry due to their favorable mechanical and physical properties, such as high strength-to-weight ratio, excellent workability, good corrosion resistance, and high thermal conductivity [1]. Casting is by far the most widely adopted manufacturing process for the production of vehicle components, and further growth in the use of casting is expected as it allows for lower energy consumption [3]. Among casting aluminum alloys, those of the Al–Si–Mg system are the most commonly used for engine blocks, cylinder heads, and other powertrain components [4], and more recently also in emerging applications for electric vehicles (EVs). One notable example is EN 42000 (AlSi7Mg), which offers an optimal balance between strength and castability, making it the preferred choice for structural components subjected to high mechanical loads. Compared to the widely used EN 42100 (AlSi7Mg0.3), EN 42000 exhibits a higher tolerance to impurities, allowing for increased concentrations of iron, manganese, and copper [5]. This characteristic enables a higher recycled content by allowing greater use of scrap derived from industrial waste and end-of-life components. As the transition toward recycled materials is becoming increasingly imperative, this feature offers significant industrial relevance.

Primary aluminum production is an energy-intensive process, accounting for approximately 1% of global anthropogenic greenhouse gas emissions and about 1% of global electricity consumption, equivalent to roughly 13 EJ annually [6]. This is largely due to the electricity-intensive Hall–Héroult electrolytic reduction process, which is often powered by fossil fuels. As a result, the aluminum industry is increasingly focusing on recycling, which preserves raw materials and cuts energy use, consuming only 2.8 kWh/kg to remelt Al scraps, compared to 45 kWh/kg for primary production [6]. As a result, the aluminum industry is increasingly focused on recycling, both for environmental and economic reasons.

However, castings produced from recycled aluminum alloys typically contain higher iron levels than those made from primary alloys, as the iron content tends to rise with the increasing amount of recycled aluminum within the alloy. No clear model or equation is currently available in the literature that directly correlates the number of remelting cycles with the progressive increase in Fe content. Nevertheless, it is generally acknowledged that Fe tends to accumulate over successive remelting cycles. In this regard, studies such as [7,8] report that the observed increase in Fe content may be partially attributed to contamination from steel crucibles during remelting. For example, [7] measured an approximately 15% increase in Fe content after the seventh remelting cycle. When impurity content reaches high levels, one of the most recommended industrial practices for cleaning molten aluminum is treatment with salt fluxes. However, this method is not specifically targeted at iron removal. To reduce Fe content more effectively, some authors have investigated physical separation techniques, such as sedimentation, although these methods tend to be energy intensive and leave residue ([9,10]). Since the complete removal of iron is difficult and expensive, its accumulation during recycling promotes the formation of Fe-based intermetallic compounds. These compounds can negatively affect key material properties by increasing the size and volume fraction of casting defects, such as porosity and interdendritic shrinkage, ultimately compromising the structural integrity ([11,12]). As a result, a detailed understanding of the impact of Fe-based intermetallics on mechanical performance is required to define the limits and potential applications of recycled aluminum alloys in the automotive field. Since the complete removal of iron remains industrially unfeasible, its detrimental effect on ductility is typically mitigated by adjusting the alloy’s chemical composition through the addition of neutralizing elements such as manganese and chromium. Manganese addition prevents the formation of the harmful acicular β-Al₅FeSi phase, instead favoring the formation of the less harmful script-like α-Al-Fe-Mn-Si phase. Notably, the β-Al₅FeSi phase forms primarily when the Mn/Fe ratio falls below 0.5 [13].

To overcome the detrimental effects of intermetallic compounds and casting defects, particularly in recycled alloys, heat treatment can be effectively employed to enhance the mechanical performance of the material. The T6 heat treatment is well known for enhancing the mechanical properties of aluminum alloys. For primary alloys, the optimal parameters have been extensively investigated in the literature, depending on the specific manufacturing process, whether conventional casting [14] or additive manufacturing [15]. However, in recycled alloys, the different chemical composition and microstructure affect the heat treatment response, requiring tailored processing parameters to optimize strength, preserve ductility, and meet environmental sustainability requirements.

For primary AlSiMg alloys, the T6 heat treatment begins with a high-temperature solution stage, typically conducted at 540 ± 5 °C, aimed at dissolving coarse intermetallic phases, particularly Mg-rich compounds, formed during solidification [16]. This is followed by rapid water quenching to obtain a supersaturated solid solution. The final stage, known as artificial aging, is carried out at temperatures between 90 °C and 260 °C, commonly around 155 °C for 3–5 h for the AlSi7Mg alloys ([17,18]). During this stage, the controlled precipitation of the precursors of the Mg₂Si phases, with the formation of coherent and semi-coherent precipitates within the α-Al matrix, significantly enhances strength and hardness [18]. Aging kinetics are temperature dependent. While higher aging temperatures accelerate peak hardening, they generally reduce the maximum achievable mechanical properties, such as yield strength (YS) and ultimate tensile strength (UTS). For instance, Pezda et al. [19] reported that UTS values around 320 MPa in AlSi7Mg alloys were obtained through a solution treatment at 540–550 °C for 30–90 min, followed by aging at 165 °C for 5–8 h. Interestingly, improvements in elongation (e%) were only observed when aging temperatures exceeded 300 °C.

Despite these findings, data on the effects of extended aging on EN 42000 alloys, especially in recycled variants, is limited. No comprehensive studies have systematically evaluated the influence of prolonged thermal exposure or overaging on their mechanical behavior.

To fill this knowledge gap, the present study evaluates the effects of T6 heat treatment on a lab-scale gravity-die-cast EN 42000 alloy with a 95% recycled Al content. The aim is to gain deeper insights into its mechanical performance and microstructural evolution.

This study provides new experimental data on the heat treatment response and mechanical behavior of a recycled EN 42000 alloy with an Fe content of 0.447 wt.%, addressing the current lack of literature on recycled aluminum alloys. The results may support the potential industrial use of recycled EN 42000 as an alternative to alloys commonly used in the automotive sector, such as primary EN 42100, with potential benefits in terms of sustainability. The T6 heat treatment parameters selected in this work were designed to be compliant with industrial and environmental constraints, including reduced solution times and moderate quenching conditions. This study also introduces novel quality indices and empirical correlations that link mechanical performance to microstructural features and processing history, with the aim of identifying optimized heat treatment parameters and supporting the design of mechanical components made from recycled aluminum alloys.

2. Materials and Methods

2.1. Materials and Casting Process

The castings employed in this study were produced by remelting recycled EN 42000 ingots with 95% recycled aluminum, supplied by Raffmetal S.p.A, in a laboratory-scale gravity die casting setup. Those secondary ingots were produced from scraps originating from post-consumer sources.

The studied alloy is characterized by a high iron content (0.447 wt.%), thus exceeding the typical threshold recommended for high-performance automotive applications (the requirement is generally below 0.15 wt.%, but, for specific applications, even values below 0.08 wt.% may be required).

The chemical composition was measured directly on the as-cast samples. The values reported in

Table 1. Sample chemical composition (wt.%.), obtained by GD-OES, of the analyzed alloy.

Alloy	Recycling Rate	Si	Fe	Cu	Mn	Mg	Ni	Zn	Ti	Al
EN 42000	95%	7.018	0.447	0.108	0.255	0.508	0.013	0.054	0.12	Bal.
		±0.101	±0.004	±0.002	±0.004	±0.009	±0.001	±0.010	±0.007

represent the average of three measurements, using Glow Discharge Optical Emission Spectroscopy (GD-OES, GDA-650, Spectruma Analytik GmbH, Hof, Bavaria, Germany). This equipment offers high analytical precision and can detect elemental concentrations with a resolution of up to 0.001 wt.%. The chemical composition in Table 1 is consistent with the limits specified in EN 1706:2020 for EN 42000 [5].

The casting process was carried out using a vacuum-assisted vertical die casting machine, specifically designed for laboratory-scale experiments (TVC10S, TOPCAST S.r.l., Arezzo, Italy). Melting was performed under a protective argon atmosphere to limit oxidation, with an external vacuum pump ensuring suction within the mold chamber. The process began by loading 1 kg of recycled EN 42000 ingots into the upper chamber. An addition of 15 g of AlTi₅B₁ grain refiner (~0.075 wt.% Ti) was introduced to refine the microstructure and reach the Ti content specified by EN 1706:2020 (0.20–0.25 wt.%) [5], considering the initial Ti content of the base alloy (0.12–0.20 wt.%). To modify the eutectic Si morphology, 0.8 g of AlSr₁₀ master alloy was added, targeting a Sr content of approximately 100–200 ppm. This moderate addition was chosen to stay within the commonly adopted Sr range for eutectic Si modification, avoiding excess levels that could promote porosity, as reported in the literature [20]. Prior to melting, an initial inerting phase using argon gas was carried out to establish a protective atmosphere that limits oxidation and contamination. Once the chamber was properly inerted, the alloy was melted at 750 °C. A second argon purge was then performed to further stabilize the environment and eliminate residual gases. The molten metal was then poured into the mold, preheated to 300 °C—which is the maximum temperature allowed by our equipment—to reduce the thermal gradients between the mold and molten metal. This aimed to minimize the risk of hot tearing and better replicate industrial casting conditions. After complete solidification, the casting was removed from the mold. As shown in Figure 1, from each casting, two semi-finished parts were extracted and later machined into tensile specimens (according to the ISO 6892-1:2019 standard procedure [21]), as well as metallographic samples for microstructural characterization and additional samples for heat treatment analysis.

2.2. Microstructural Characterization

Microstructural analyses were conducted on metallographic samples prepared according to standard procedures. Samples were mounted in conductive resin, then ground using silicon carbide (SiC) abrasive papers up to P2500 grit. Subsequent polishing was carried out with polycrystalline diamond suspensions, progressively down to a final particle size of 1 µm, following the ASTM E3-11 standard [22]. To enhance the contrast between the aluminum matrix and the eutectic microstructural constituent, etching was performed using Keller’s reagent for 10 s, following the ASTM E407-07 standard [23]. Microstructural characterization was conducted using both Optical Microscopy (Axioscope 7, Zeiss, Oberkochen, Germany) and Field Emission Gun Scanning Electron Microscopy (FEG-SEM, Mira3, Tescan, Brno, Czech Republic), equipped with an Energy-Dispersive X-ray Spectroscopy (EDS) system. The analysis focused on identifying casting defects (e.g., gas porosities and interdendritic shrinkages), as well as intermetallic compounds. All specimens were extracted from the bottom region of the castings to ensure consistency. Quantitative image analysis was performed using ImageJ software (version 1.54p) to evaluate the key microstructural parameters. The cross-sectional area fraction of defects was quantified from OM micrographs with a field of view of 2800 × 1600 µm, evaluating seven images per sample. Secondary Dendrite Arm Spacing (SDAS) was measured according to the procedure indicated in [24], from higher magnification micrographs (1400 × 800 µm), analyzing a minimum of ten dendrites per image across three images per sample. The modification level of eutectic silicon was assessed on optical micrographs with a field of view of 260 × 150 µm and classified according to the American Foundry Society guidelines [25]. In this guideline, eutectic silicon modification in aluminum–silicon alloys is classified based on the morphology of silicon particles. The AFS system categorizes modification levels from unmodified to fully modified, depending on the fineness and distribution of eutectic silicon.

2.3. Heat Treatment Analysis

T6 heat treatment was adopted in this study. Solution treatment was carried out in a laboratory furnace (LT 9/14, Nabertherm GmbH, Lilienthal, Germany) by heating the specimens to 535 °C for 4.5 h, in accordance with parameters used in the author’s previous study [26] and consistent with those reported in the literature for castings produced with a similar alloy [27]. The temperature profile was monitored using two thermocouples (accuracy ±1% of the measured value) to ensure precise and uniform thermal conditions throughout both the solution and aging stages. The solution treatment promotes the dissolution of Mg-containing phases, thereby increasing the magnesium content in the solid solution, which is essential for the subsequent precipitation of strengthening phases during artificial aging. Upon completion of the solution step, the samples were immediately quenched in water at 60 °C to obtain a supersaturated solid solution. The quenching temperature of 60 °C was selected to better align with industrial practices, where warm water quenching is used to reduce residual stresses and distortions in large castings. Aging was then performed without delay to avoid pre-aging in an electric furnace at temperatures ranging from 160 °C to 190 °C, with holding times from 0.5 to 8 h, followed by air cooling. Additionally, prolonged aging treatments were carried out to evaluate the long-term thermal stability of the alloy. In detail, aging temperatures of 160 °C and 170 °C were investigated, with durations up to 184 h. The aging parameters were selected based on the Differential Thermal Analysis (DTA) data available for AlSi7Mg alloys and aligned with industrial heat treatment practice. The optimization of the aging process was achieved by systematically varying both the aging temperature and duration, as summarized in the experimental plan shown in

Table 2. Scheme of the tested aging parameters.

Solution Temperature and Duration	Quenching Conditions	Aging Temperature [°C]	Aging Time [h]
535 °C for 4.5 h	H2O at 60 °C	160	0.5 1 1.5 2 3 4.5 5 6 8
		170
		180
		190

.

Aging curves were built using Brinell hardness as a key parameter. Hardness measurements were performed for each aging condition using a QATM durometer (QNESS 150 CS, ATM Qness GmbH, Mammelzen, Germany). Prior to testing, the sample surfaces were carefully ground to enhance measurement precision and repeatability. Brinell hardness tests (HBW 2.5/62.5) were carried out with a 2.5 mm tungsten carbide ball and a 62.5 kgf load. For each aging condition, a minimum of three indentations were carried out, and average values along with standard deviations were calculated.

2.4. Mechanical Characterization

Tensile testing was conducted on specimens subjected to the four most promising aging conditions, using a screw-driven testing machine (Giuliani Italsigma, Forlì, Italy). These conditions were selected to provide a comprehensive overview of the alloy’s mechanical response across different stages of the aging curve, including the first hardness peak (probably associated with the formation of GP zones), the peak-aged condition, and the pre-aged and overaged states. The samples were machined following standardized geometries specified in ISO 6892-1:2019 [21]. Key tensile properties, including the elastic modulus (E), yield strength (YS), ultimate tensile strength (UTS), elongation to fracture (e%), and strain hardening exponent (n), were determined according to ISO 10275:2020 [28]. To consolidate tensile strength and elongation into a single performance metric, the quality index (QI) proposed by Drouzy et al. [29] was employed, as defined by Equation (1):

$$QI=UTS+150\ast {\mathit{log}}_{10}e\%$$

(1)

To further evaluate mechanical performance, two innovative quality indices were considered: the ductility quality index (QD) and the toughness quality index (QT). QD, as formulated by Angella et al. [30], evaluates the ratio between the theoretical uniform strain at necking onset (e_u) and the experimentally measured fracture strain, i.e., elongation to rupture (e_r). It can be determined by the following equation:

$$Q_D={\displaystyle \frac{e_u}{e_r}}$$

(2)

where e_u and e_r are defined as follows:

$$e_u={\epsilon }_c·\mathit{ln}\left({\displaystyle \frac{1+{\epsilon }_c}{{\epsilon }_c}}{\displaystyle \frac{{\sigma }_v-{\sigma }_0}{{\sigma }_0}}\right)$$

(3)

$$e_r=\max \left({\epsilon }_{tp}\right)$$

(4)

$${\epsilon }_{tp}={\epsilon }_t-{\displaystyle \frac{{\sigma }_t}{E\ast 1000}}$$

(5)

$${\sigma }_V={\Theta }_0·{\epsilon }_c$$

(6)

where ε is the engineering strain [-], σ is the stress [MPa], σ_t is the true stress [MPa], ε_t is the true strain [-], ε_tp is the true plastic strain [-], and E is the elastic modulus [GPa]. To calculate the theoretical uniform strain at necking onset (e_u), σ₀ is the stress at ε = 0, while ε₀ is a parameter denoted by the Voce equation as a characteristic strain. Here, 1/ε_c and Θ₀ are, respectively, the slope and intercept of best fit of the differential form of the Voce equation [30] at high stresses in a θ = dσ_t/dε_t,p versus σ_t graph. A QD value below this unit suggests the presence of internal defects adversely affecting ductility. QT, based on the method proposed by Tiryakioglu et al. ([31,32]), compares the theoretical maximum energy absorption capacity (in the absence of porosity or inclusions) with the actual measured toughness. Toughness Ψ can be determined by Equation (7):

$$\Psi ={\int }_0^{{\epsilon }_f}\sigma d\epsilon ={\displaystyle \frac{{YS}^2}{2E}}+{\sigma }_{\infty }\left({\epsilon }_f-{\displaystyle \frac{YS}{E}}\right)+{\epsilon }_0({\sigma }_{\infty }-{\sigma }_0)\left(\mathit{exp}\left({\displaystyle \frac{{-\epsilon }_f}{{\epsilon }_0}}\right)-\mathit{exp}\left({\displaystyle \frac{-YS}{E{\epsilon }_0}}\right)\right)$$

(7)

where σ_∝ is the saturation stress (the stress at which full plasticity is reached) and ε_f is the true fracture strain. In addition, with Equations (8) and (9), it is possible to define the target toughness Ψ_c, as shown in Equation (10).

$${\sigma }_c={\displaystyle \frac{{\sigma }_{\infty }}{1+{\epsilon }_0}}$$

(8)

$$e_{lc}={\left({\displaystyle \frac{{\sigma }_{\infty }-{\sigma }_c}{{\sigma }_{\infty }-{\sigma }_0}}\right)}^{-{\epsilon }_0}-1$$

(9)

$${\Psi }_c={\displaystyle \frac{YS+\frac{{\sigma }_c}{1+e_{lc}}}{2}}{e}_{lc}$$

(10)

The toughness quality index (QT) is then calculated as follows:

$$Q_T={\displaystyle \frac{\Psi }{{\Psi }_c}}=\left({\displaystyle \frac{YS+UTS}{YS+\frac{{\sigma }_c}{1+e_{lc}}}}\right){\displaystyle \frac{{el}_f}{{el}_c}}$$

(11)

For each aging condition, a minimum of three tensile specimens were tested, and average values and standard deviations were reported to ensure statistical reliability.

2.5. Fractographic Characterization

Fractographic analyses were performed using Field Emission Gun Scanning Electron Microscopy (FEG-SEM). The fracture surfaces of tensile samples were examined using secondary electron (SE) imaging to evaluate morphology, and backscattered electron (BSE) imaging was used to reveal compositional contrast and heterogeneities. To investigate specific regions of interest, Energy-Dispersive X-ray Spectroscopy (EDS) was used for qualitative point analyses, particularly for identifying intermetallic compounds (field of view: 520 × 520 µm). The purpose of the fractographic analyses was to identify fracture mechanisms and to evaluate the influence of casting defects and intermetallic compounds on the tensile behavior. A quantitative assessment of the intermetallic phase fraction (IM%) was also performed by analyzing BSE images, where Fe-rich intermetallics can be distinguished from the α-Al matrix based on their grayscale contrast.

2.6. Correlations Between Microstructural Features and Mechanical Properties

Quantitative data from image analysis and mechanical testing were processed through descriptive statistics, including mean values and standard deviations. The main objective was to identify correlations between mechanical properties and selected microstructural indicators, as follows: percentage of intermetallic compounds on the fracture surface (IM%); true cross-sectional area at fracture (CS%), defined as 100%$-$%area of defects on the fracture surface; area fraction of eutectic silicon (A) observed in metallographic specimens; and Brinell hardness (HBW 2.5/62.5). These analyses were carried out to elucidate the influence of microstructural features on mechanical behavior, with the aim of providing meaningful insights for the development and optimization of recycled AlSi7Mg alloys.

3. Results and Discussion

3.1. Microstructure

The microstructure of the recycled EN 42000 alloy in the T6 condition is shown in Figure 2a,b. These micrographs depict a typical hypoeutectic Al–Si microstructure, consisting of a primary α-Al matrix in dendritic form surrounded by eutectic constituent.

SDAS, measured on micrographs such as that shown in Figure 2c, was 19 ± 2 µm, consistent with values typically observed on alloys produced under controlled laboratory-scale gravity die casting solidification conditions. The measured porosity, equal to 3.25 ± 0.25%, was higher than the values typically reported for recycled Al–Si alloys [11]. This elevated porosity is attributed to the significant presence of iron (0.447 wt.%) and copper (0.108 wt.%), unintentionally introduced through the recycling of post-consumer scrap. Iron promotes the formation of Fe-rich intermetallic phases, which, due to their morphology and distribution, impair melt fluidity during casting. The reduced fluidity results in incomplete mold filling and promotes the formation of interdendritic shrinkage, ultimately increasing the overall porosity, negatively affecting the alloy’s mechanical performance [12]. Furthermore, copper is known to exacerbate microporosity, favoring hot tearing. Its presence has been shown to increase the volume fraction of dispersed porosity by up to a factor of four compared to Cu-free Al–Si–Mg alloys [33]. Despite the high Fe content, no needle-like β-Al₅FeSi intermetallic phase was observed. This is attributed to the presence of manganese (0.255 wt.%), which promotes the transformation of the β-phase into less detrimental α-phases [13]. With a Mn/Fe ratio of 0.57, the alloy predominantly forms script-like Al–Si–Fe–Mn intermetallics, which are morphologically less harmful, although still exhibit sharp edges, as shown in Figure 2d.

Further insights into the microstructure were obtained through BSE FEG-SEM analyses, as shown in Figure 3, and complemented by EDS analyses for phase identification.

A representative α-Fe intermetallic phase is shown in Figure 3, accompanied by eutectic Si particles exhibiting various morphologies. The α-phase, commonly identified as α-Al₁₅(Mn,Fe)₃Si₂, is known to contribute to embrittlement and act a stress concentrator. According to the literature [12], it retains its morphology and size following T6 heat treatment. In contrast, the eutectic Si particles undergo substantial morphological evolution. In the as-cast condition, modification is incomplete (classified as Class 4, according to AFS standards), whereas, after the T6 treatment, a fully modified morphology (Class 5, according to AFS) was achieved. The refined and fibrous morphology of the eutectic silicon is the result of the thermal effects induced during the solution treatment stage. This process promotes spheroidization and fragmentation of the eutectic silicon lamellae through diffusion-driven mechanisms, leading to a more rounded and homogeneous morphology. Such microstructural modification is known to enhance both the ductility and overall mechanical performance of the alloy [34].

3.2. Heat Treatment and Aging Response

This section focuses on the aging response of the recycled EN 42000 alloy, a topic that remains poorly documented in the literature, despite its growing industrial relevance. Figure 4a presents a 3D surface plot illustrating the evolution of Brinell hardness (HBW 2.5/62.5) as a function of aging time and temperature following T6 heat treatment. The corresponding isohardness contours are displayed in Figure 4b, enabling the identification of the optimal time–temperature combinations for achieving targeted hardness levels. The aging behavior of the studied alloy is consistent with that reported by Mohamed et al. [18] for Al–Si–Mg alloys. The initial hardness peak is likely associated with the formation of Guinier–Preston (GP) zones. As aging progresses, these metastable clusters are expected to dissolve, leading to the precipitation of Mg₂Si particles in the form of coherent and semi-coherent precipitates, which are responsible for the observed peak hardness. Beyond this point, the coarsening of these nanoscale precipitates, undetectable by SEM, marks the onset of the overaging regime, accompanied by a gradual decrease in hardness.

The overall aging trend observed in the recycled EN 42000 mirrors the one reported in the literature for its primary counterpart [19]. Peak hardness is achieved more rapidly at higher temperatures. For instance, at 190 °C, maximum hardness is reached within just 2 h, followed by a steep drop due to accelerated overaging. Conversely, aging at lower temperatures (160 °C–170 °C) proceeds more gradually, yielding a slower hardness increase and a delayed peak, often beyond the 8 h observation window. Under these conditions, the alloy retains elevated hardness over extended durations, which is advantageous for ensuring a more uniform hardness distribution, even in thick-wall or geometrically complex castings. To further evaluate long-term thermal stability, extended aging experiments up to 184 h were conducted at 160 °C and 170 °C. As shown in Figure 5, aging at 170 °C resulted in a peak hardness of 111 HBW 2.5/62.5 at 8 h, followed by a gradual softening. In contrast, at 160 °C, a significantly higher peak of 118 HBW 2.5/62.5 was achieved around 38 h and remained relatively stable over time. Based on this behavior, 160 °C was selected as the optimal aging temperature for subsequent mechanical testing.

3.3. Mechanical Properties

Tensile tests were performed on specimens subjected to selected aging treatments at 160 °C, chosen for its slower aging kinetics and enhanced hardness stability. The selected aging times were as follows: 1.5 h (underaged condition); 4.5 h (representative of typical industrial practice for internal combustion engines (ICE) components produced with similar primary alloys [26]); 38 h (peak hardness); and 128 h (overaged condition). As shown in Figure 6a, the elastic modulus remained constant (69 GPa) across all tested conditions, indicating that aging does not affect the elastic behavior of the alloy. In contrast, the plastic region of the stress–strain curves displayed significant variations, with yield strength (YS), ultimate tensile strength (UTS), and elongation to fracture (e%) depending on the applied aging condition. Notably, the strain hardening exponent (n) decreased from 0.13 at 1.5 h to 0.05 at 128 h, suggesting a loss of strain-hardening capacity with prolonged aging. The mechanical response reveals a typical strength–ductility trade-off, with an increasing mechanical strength from underaged to peak-aged conditions accompanied by a progressive reduction in ductility, particularly evident in the overaged state.

As clearly shown in Figure 6b, the underaged condition (160 °C × 1.5 h) exhibited the highest elongation but limited strength. At 4.5 h, strength increases significantly while ductility decreases. Peak aging at 38 h yields maximum strength due to the formation of finely dispersed precipitates yet further reduces elongation. Upon overaging (128 h), strength slightly declines due to precipitate coarsening, and ductility is further compromised. When compared to primary EN 42000 alloys with similar Fe content [35], the recycled alloy studied in this work consistently exhibits comparable strength but lower ductility. This reduction in elongation is mainly attributed to the increased presence of Fe-rich intermetallic phases, which hinder the flow of interdendritic liquid during feeding, thereby promoting the formation of shrinkage porosity [12] and compromising mechanical behavior. Figure 7 summarizes the results of the quality indices evaluation for the alloy aged at 160 °C.

QI remains relatively stable up to 38 h, but it drops significantly at 128 h, mainly due to the marked reduction in elongation observed in the overaged state. QD shows a similar trend, decreasing noticeably at 128 h, which indicates a reduced fracture resistance. QT, conversely, remains essentially unchanged between 38 and 128 h, suggesting that the strength–ductility trade-off plateaus beyond peak aging. This decoupling between QD and QT underlines how overaging primarily affects the alloy toughness. The divergence of QD, QT, and QI at prolonged aging times emphasizes the need for comprehensive interpretations of mechanical performance indicators. Collectively, the results support the conclusion that aging at 160 °C for up to 4.5 h offers the best compromise between strength and ductility for the recycled EN 42000 alloy.

3.4. Fracture Behavior

The fracture surfaces of the recycled EN 42000 alloy showed a high density of interdendritic shrinkage cavities (Figure 8), along with α-Al₁₅(Fe,Mn)₃Si₂ intermetallic compounds. The investigated shrinkages have a maximum size of roughly 11,500 µm², while intermetallics exhibited an average size of 202 ± 34 µm², indicating relevant distribution within the microstructure.

Porosity and large intermetallics acted as preferential crack initiation sites [36], thereby explaining the lower elongation to fracture observed in this recycled alloy compared to that observed in the literature data [35]. The observation of fracture surfaces of samples subjected to different aging conditions qualitatively revealed no substantial changes in defect morphology or in the size and distribution of intermetallic phases. However, the heat treatment significantly influenced the mechanical behavior, as reflected in the evolution of hardness (Figure 5) and fracture surface morphology (Figure 9).

An increase in hardness was observed from the underaged to overaged conditions, following the typical strength–ductility trade-off. This was evident on the fracture surfaces, where, from the 160 °C × 1.5 h condition characterized by deep and well-defined dimples, the fracture surfaces gradually evolved into flatter and more fragmented features after 128 h of aging at 160 °C.

3.5. Microstructure–Tensile Property Correlations

Microstructural parameters were quantified through image analysis conducted on the fracture surfaces and metallographic specimens of the most representative specimen from each condition. These parameters were then correlated with both UTS and e%.

$$UTS=5.8\ast {\%CS}^{0.5}\ast {\%IM}^{0.2}\ast A^{-0.2}\ast {HBW}_{2.5/62.5}^{0.4}$$

(12)

$$e\%=256\ast {\%CS}^{0.927}\ast {\%IM}^{-0.01}\ast A^{-4.033}\ast {HBW}_{2.5/62.5}^{-0.01}$$

(13)

where:

• %CS = True cross-sectional area (100%—percentage of area of defects on the fracture surface);
• %IM = Area percentage of intermetallic compounds on the fracture surface;
• A = Area fraction of eutectic silicon from metallographic specimens;
• HBW 2.5/62.5 = Brinell hardness.

Equation (12) (R² = 0.91) reveals a positive coefficient for %CS, confirming that casting integrity enhances ultimate tensile strength (UTS). A similar positive correlation is observed for the fraction of hard intermetallic phases. In contrast, the size of eutectic silicon particles exhibits a negative coefficient, highlighting their detrimental effect on strength when coarsened. Equation (13) (R² = 0.97) confirms that elongation is strongly influenced by casting integrity, as reflected in the high exponent for %CS. All other microstructural indices (%IM, A, HBW 2.5/62.5) exhibit negative coefficients, indicating that both the amount of intermetallics and coarser eutectic silicon particles contribute to reduced ductility by promoting localized stress concentration. Hardness (HBW 2.5/62.5) confirms its role as a trade-off indicator, exhibiting a positive influence on UTS (Equation (12)) and a moderate effect on elongation (Equation (13)). Finally, Equation (14) (R² = 0.94) demonstrates that yield strength is predominantly influenced by Brinell hardness, as evidenced by:

$$YS=4.1204\ast {HBW}_{2.5/62.5}^{-0.01}-160.02$$

(14)

This inverse relationship reflects the complex role of precipitation hardening during T6 heat treatment, which simultaneously enhances hardness and strength, in agreement with the findings reported by Ceschini et al. [26].

4. Conclusions

This study focused on the mechanical and microstructural characterization of a recycled EN 42000 alloy T6-heat-treated with 95% recycled aluminum and 0.447 wt.% Fe, produced via gravity die casting. The aim was to identify the most suitable aging condition to extend the alloy’s applicability and promote its wider adoption towards sustainable high-performance castings. The main conclusions drawn from this investigation are as follows:

• Microstructural analysis revealed a homogeneous structure with fine dendrites (SDAS ≈ 20 µm) and well-refined eutectic silicon, resulting from controlled processing under argon, which limited oxide inclusions. However, shrinkage cavities were still observed, mainly linked to script-like α-Al₁₅(Fe,Mn)₃Si₂ intermetallics that hinder feeding during solidification and reduce castability.
• A T6 heat treatment consisting of a solution at 535 °C for 4.5 h, quenching in water at 60 °C, and aging at 160 °C provided improved thermal stability compared to aging at 170 °C. The aging curve at 160 °C revealed an initial hardness peak at 4.5 h and a peak-aged condition at 38 h.
• Aging at 160 °C for 4.5 h was identified as the optimal condition in terms of mechanical performance. Based on the evaluation of both ductility and toughness quality indices, this treatment provided a well-balanced combination of strength and ductility, with a yield strength of 258 MPa, an ultimate tensile strength of 313 MPa, and an elongation to fracture of 3.9%.
• A high Fe content promotes the formation of coarse intermetallic compounds, which act as stress concentrators and reduce ductility through the increased formation of shrinkage cavities. An inverse exponential correlation between Fe content and elongation to fracture was established through an innovative microstructure–mechanical property analysis, highlighting the combined influence of intermetallics and casting defects on fracture behavior.