Influence of Zn and Cr Additions on the Microstructure and Mechanical Properties of Al–Mg–Si–Zr–Cu Multicomponent HPDC Alloys

Abstract

This study develops novel multicomponent Al–Mg–Si–Zr–Cu-based alloys for high-pressure die casting (HPDC) with improved mechanical properties and thermal stability. Four compositions were designed through Zn and Cr additions, supported by thermodynamic modelling. XRD and SEM/EDS analysed phase formation and microstructure, while density, electrical conductivity, hardness, and mechanical behaviour under tensile and compressive loading at room temperature and 200 °C were evaluated. Hardness increased from 166 to 214 HV3 with Zn and Cr due to the formation of complex intermetallic phases. The Al–Mg–Si–Zr–Cu alloy showed the best balance of strength and ductility under tensile loading, whereas Zn and Cr additions reduced tensile performance. In compression, Zn significantly improved strength, reaching the ultimate compressive strength of 697 MPa. Compared with the reference AlSi9Cu3 alloy, the new alloys achieved up to 30% higher yield strength, 13% higher ultimate tensile strength, and improved thermal stability. Among the studied compositions, Al72Mg10Si5Zr3Cu10 showed the best overall performance, while Al67Mg10Si5Zr3Cu10Zn10 was optimal for compression-dominated applications.

1. Introduction

The increasing demand for lightweight structural materials with enhanced mechanical performance and thermal stability has driven the development of advanced alloy design strategies. Among these, multicomponent alloy concepts have emerged as a promising approach, inspired by the compositional complexity paradigm originally introduced in high-entropy alloys (HEAS). Rather than relying on a single dominant strengthening mechanism, these alloy design strategies exploit the synergistic interaction of multiple alloying elements to tailor phase formation, microstructural stability, and mechanical performance [1]. Recent studies have also extended these concepts to aluminium casting alloys and high-pressure die casting (HPDC) systems, highlighting the growing interest in applying compositional complexity principles to lightweight alloys processed by industrial routes [2]. Unlike conventional HEAs, aluminium multicomponent alloys must simultaneously preserve low density, good castability, and compatibility with established high-productivity manufacturing technologies. Consequently, alloy development is not only governed by configurational complexity but also by the careful selection of alloying additions capable of promoting thermally stable strengthening phases while maintaining satisfactory processing characteristics. Recent advances in multicomponent aluminium alloys have demonstrated that this approach offers considerable potential for improving the mechanical performance and thermal stability of cast aluminium alloys without sacrificing their industrial applicability [3].

More recently, integrated computational alloy design approaches combining CALPHAD thermodynamic modelling, machine learning, and experimental validation have become increasingly important for accelerating the development of advanced aluminium casting alloys and guiding alloy optimisation before experimental fabrication [4,5].

Besides computational alloy design, tailoring the microstructure has become another key strategy for improving the overall performance of aluminium alloys. For example, hybrid reinforcement introduced by friction stir processing has been reported to simultaneously enhance strength, ductility, and wear resistance in Al5052 alloys. Although the processing route differs substantially from HPDC, these findings further demonstrate that careful control of microstructural features is essential for achieving a balanced combination of mechanical and tribological properties in aluminium-based materials [6].

In this context, HPDC is characterised by rapid solidification under applied pressure, leading to pronounced microstructural refinement [7]. Under these conditions, the formation of coarse equilibrium phases is reduced, while metastable and non-equilibrium intermetallics are favoured, often resulting in enhanced mechanical performance [8]. Cooling rates on the order of 10² °C/s significantly reduce grain size and/or secondary dendrite arm spacing and modify solute redistribution during solidification. Consequently, the resulting microstructures deviate markedly from equilibrium predictions, making phase identification and analysis of solidification pathways essential for understanding structure–property relationships [9]. Similarly, studies on other cast alloy systems have demonstrated that alloy composition and cooling rate act synergistically in determining phase constitution and microstructural evolution during solidification, highlighting the importance of jointly considering both parameters when designing advanced casting alloys [10]. Moreover, the increasing service temperatures of automotive components, particularly in electrified vehicles, have further highlighted the limitations of conventional Al–Si–Cu HPDC alloys and the growing need for thermally stable strengthening phases capable of retaining mechanical performance at elevated temperatures [11].

From a compositional perspective, Al–Mg alloys exhibit excellent ductility and fracture resistance in the as-cast condition, although their limited castability restricts their industrial application. In contrast, Al–Si-based alloys dominate HPDC applications due to their superior fluidity and ability to produce complex geometries, with alloys such as AlSi9Cu3 being widely used. However, these alloys typically suffer a significant reduction in hardness at temperatures above 100–150 °C, mainly due to precipitate coarsening and solute redistribution within the aluminium matrix [12,13].

A potential strategy is to combine the advantages of both systems. The addition of magnesium (Mg) promotes the formation of Mg₂Si (at the expense of free silicon), modifying both the morphology and distribution of second-phase particles. In particular, Mg additions of approximately 3 wt.% can be sufficient to suppress the formation of the acicular primary Si phases, favouring the formation of Mg₂Si particles [14] instead. The characteristics of these particles strongly depend on the alloy composition, especially on maintaining an appropriate Mg/Si ratio close to the stoichiometric requirement for Mg₂Si formation. When Mg and Si contents are sufficiently high, the alloy composition may shift towards the hypereutectic region of the Al–Mg₂Si pseudo-binary system, leading to the formation of primary Mg₂Si prior to α-Al solidification [15]. These particles are particularly beneficial for high-temperature performance due to their high melting temperature, low density, high hardness, and relatively high elastic modulus. Furthermore, compared to acicular silicon, Mg₂Si provides a more favourable morphology, reducing stress concentration and enhancing matrix strengthening.

Copper (Cu) addition promotes the formation of θ-Al₂Cu and S-Al₂CuMg strengthening phases [15], whose stability is strongly influenced by the Mg content. Given the limited solubility of Cu in aluminium (around 4 wt.%), compositions close to this threshold are typically preferred to increase solid-solution and precipitation strengthening, while avoiding the formation of coarse intermetallics. However, in highly alloyed systems, lower Cu contents may be more effective [16]. For instance, in AlSi10Mg alloys, the addition of 2 wt.% Cu has been reported to increase pore size [17], whereas in Al–Cu–Mg systems, increasing the Cu content from 4 to 8 wt.% led to improvements of approximately 30% in UTS at room temperature (RT), and up to 62% at 300 °C, attributed to the formation of a higher density of fine strengthening phases [18]. Similarly, in Al–Mg alloys, the increased Cu content up to 6 wt.% has been associated with grain refinement, as well as simultaneous improvements in strength and ductility [19].

Zinc (Zn) is another promising alloying element. It helps modify the solidification process by reducing dendritic structures, promoting a more equiaxed morphology, and increasing the fraction of eutectic phases [20]. In Al–Mg–Cu alloys, Zn stimulates the precipitation of coherent T-Mg₃₂(Al,Zn)₄₉ phases, while suppressing S-Al₂CuMg and β-Al₃Mg₂ phases and promoting the dispersion of θ-Al₂Cu [21]. The combined addition of Zn and Cu has shown particularly promising results. In Al–Mg alloys, the addition of approximately 6 wt.% Cu and 6 wt.% Zn led to the formation of MgZn₂ and Al₂CuMg phases, resulting in improved mechanical performance [22]. This final phase has been reported to improve the thermal stability of high-entropy alloys or multicomponent Al–Mg–Si–Cu alloys [23]. MgZn₂ phases provide significant strengthening, particularly in the form of metastable η′ phases due to their high coherency and resistance to dislocation motion; however, their thermal stability is limited, as η′ tends to transform into the equilibrium η phase at elevated temperatures [24].

The Zn/Mg ratio plays a critical role in determining phase formation. At low Mg contents or low Zn/Mg ratios, strengthening is primarily associated with η-MgZn₂ formation [22]. In contrast, higher Mg contents favour the formation of T-type phases, such as T-Mg₃₂(Al,Zn)₄₉. Additional phases, including T-Al₂Mg₃Zn₃, T-Al₂Mg₃Zn₃Cu_3-x, or η-type Mg(Zn,Cu,Al)₂, may also precipitate depending on the composition [25,26]. These more complex Zn-containing phases generally exhibit enhanced thermal stability compared to metastable η′ precipitates, due to their compositional complexity and reduced diffusion kinetics. Higher Zn contents combined with lower Mg and Cu levels promote MgZn-rich phases instead of S-Al₂CuMg [27]. Since η-MgZn₂ dissolves more readily during heat treatments, coarse Al₂CuMg particles may act as crack initiation sites [28]. Furthermore, Cu contents above approximately 2.5 wt.% tend to favour the S-phase formation rather than dissolution in the matrix or precipitating as in T-type phases [25].

Microalloying additions such as zirconium (Zr) and chromium (Cr) further influence microstructural stability. Zr promotes the formation of thermally stable Al₃Zr dispersoids within the α-Al matrix, enhancing recrystallisation resistance and high-temperature stability [29]. However, due to its very low solubility in aluminium, high processing temperatures may be required, limiting its widespread application. In addition, Zr can form intermetallics such as Al₃Zr or (Al,Si)₃(Zr,Ti) during early solidification stages, whose morphology strongly affects mechanical properties. Typically, additions are limited to around 0.2 wt.% to avoid undesirable interactions with other elements [30]. Cr, on the other hand, contributes to matrix strengthening and improves structural stability at elevated temperatures. It can also refine the microstructure and increase tolerance to Fe impurities compared to Al–Si–Mn systems [31]. However, Cr is rarely used in Al–Si alloys due to its tendency to form coarse intermetallic sludge particles in interdendritic regions, which can deteriorate mechanical properties [32,33].

Table 1. Representative recent studies on multicomponent and HEA-inspired aluminium casting alloys.

Alloy System	HEA-Inspired	Process	Main Objective	Main Limitation	Ref.
AlMgSiCu	Yes	HPDC	Strength	No Zn/Cr/Zr	[15]
AlCuMgSiZn	Yes	HPDC	Hardness	No Cr/Zr	[34]
AlSi + HEA	Yes	HPDC	Strength	No Mg, Cu, Zn, Cr or Zr	[35]
AlMgFe	No	HPDC	Ductility	Conventional alloy	[36]
AlSiMgCu	No	Gravity sand	Strength	Conventional alloy	[37]
AlZnMgCu	No	Gravity sand	Strength	Conventional alloy	[38]
AlTiVZrNb	Yes	Gravity die	Hardness	No HPDC; No AlMgSi/Zn/Cr	[39]
AlZnCuNiSiMg	Yes	Gravity die	High entropy	No HPDC; No AlMgSi/Zr/Cr	[40]
AlSiNiFeMn	Yes	Gravity die	Strength and thermal stability	No HPDC; No AlMgSi/Zr/Zn/Cr	[41]
This work	Yes	HPDC	Strength and thermal stability	Cu + Zn + Cr + Zr in HPDC	-

summarises the latest representative multicomponent and HEA-inspired aluminium casting alloys reported in the literature, together with their processing route, main objective, and the remaining research gap addressed in each study. As can be seen, although several studies have explored compositionally complex aluminium alloys, most of them have been produced by gravity sand or die casting. In contrast, HPDC studies remain predominantly focused on conventional aluminium alloys or on systems containing only a limited number of alloying additions. Consequently, the development of compositionally complex aluminium alloys specifically designed for HPDC applications is still at an early stage. To the best of the authors’ knowledge, no previous study has systematically investigated the combined addition of Cu, Zn, Cr and Zr in Al–Mg–Si alloys processed by HPDC while simultaneously evaluating the influence of each alloying element on phase evolution, microstructure, and mechanical behaviour. Therefore, the present work aims to bridge this gap by combining a compositionally complex alloy design with HPDC processing and comprehensive thermodynamic, microstructural, and mechanical characterisation.

2. Materials and Methods

2.1. Alloy Design

The alloy design strategy adopted in this study is based on a progressive modification of a conventional Al–Mg–Si system base alloy, selected due to its widespread use in the automotive industry and its excellent castability. The base alloy was first modified by the addition of Zr and Cu and subsequently tailored through the incorporation of Zn and Cr, as shown in Figure 1, leading to the development of multicomponent alloys. This design approach aims to combine multiple strengthening mechanisms while preserving the fundamental solidification and precipitation behaviour of the base alloy.

Each alloying element was selected according to its specific role: Cu and Zn were introduced to enhance precipitation hardening; Cr to promote the formation of thermally stable intermetallic phases, contribute to microstructural refinement, and improve thermal stability; and Zr to provide dispersoid strengthening and further enhance thermal stability. The combined effect of these elements is expected to generate thermally stable obstacles to dislocation motion. To systematically evaluate the influence of Cr, its content was progressively increased up to 3 wt.%, not to identify the optimum alloy composition, but to investigate the upper limit of Cr addition in this multicomponent HPDC system and to determine the transition from beneficial strengthening towards excessive primary intermetallic formation and microstructural embrittlement. Based on this design strategy, four experimental alloy compositions were defined for subsequent thermodynamic evaluation and experimental validation.

Based on this alloy design strategy, four experimental multicomponent alloys, labelled H1–H4, were designed and subsequently manufactured. The target nominal chemical compositions of these alloys are summarised in

Table 2. Nominal chemical composition of alloys.

Alloy	Ref.	Al	Mg	Si	Zr	Cu	Zn	Cr
AlMgSiZrCu	H1	72	10	5	3	10	-	-
AlMgSiZrCuZn	H2	67	10	5	3	10	5	-
AlMgSiZrCuZnCr1	H3	66	10	5	3	10	5	1
AlMgSiZrCuZnCr3	H4	64	10	5	3	10	5	3

. These nominal compositions were initially used as input for the CALPHAD calculations during the alloy design stage. After alloy production, the actual chemical compositions were determined experimentally and employed for a second set of thermodynamic simulations, which are presented and discussed in Section 3.

2.2. CALPHAD Methodology

To support the alloy design, thermodynamic calculations were performed using FactSage software version 8.3 (2023) coupled with the FTlite database (2023). During the alloy design stage, equilibrium and Scheil solidification calculations based on the targeted nominal compositions were carried out to optimise the elemental ratios and predict phase formation. Particular attention was given to avoiding the formation of brittle primary silicon phases and to promoting thermally stable strengthening phases that could improve the mechanical performance of the alloys. Following alloy production, the actual chemical compositions were determined, and slight deviations from the targeted nominal compositions are inherent to alloy manufacturing; a second set of equilibrium and Scheil calculations were performed using the experimentally measured compositions. These updated simulations provided a more accurate representation of the solidification behaviour of the manufactured alloys and are therefore the results presented and discussed throughout Section 3.

2.3. Materials and Casting Procedure

The raw materials were melted and chemically homogenised in a 500 kg electric furnace (EBC model, Dugopa, Fuenlabrada, Spain). The initial metal charge consisted of AlSi9Cu3, AM60, recycled AZ91, and Zamak 5, which served as base materials. Alloying elements were added through master alloys and specific additives, including zinc supplied as spherical pellets, metallurgical-grade silicon, Al–Mn (Mn80-Al20), Al–Cu (Cu80-Al20), Al-Cr (Cr80-20Al), and Al-Zr (80Zr-Al20) briquettes.

Zr and Cr were added at the initial stage of the melting process due to their high melting temperatures, followed by the addition of Si to promote its dissolution. Cu- and Mn-containing master alloys were subsequently introduced shortly before casting to minimise oxidation and element loss.

Once the target chemical composition for each alloy was achieved, the melt temperature was stabilised at 700 °C. The molten alloys were then injected into a steel die preheated to 300 °C using a 950 t HPDC machine (PT-650 model, Pretansa, Tarragona, Spain). After solidification, the castings were extracted from the die and immediately quenched in water at 50 °C to reduce internal stress. An overview of the HPDC setup and representative cast components produced under these conditions is presented in Figure 2.

The chemical compositions of the alloys were analysed both prior to and after casting using inductively coupled plasma–optical emission spectrometry (ICP-OES) with a SPECTROX spectrometer (Spectro Analytical Instruments GmbH, Kleve, Germany) to verify compositional consistency during processing. In addition, local compositional analysis was performed by scanning electron microscopy coupled with energy-dispersive spectroscopy (SEM–EDS) at a magnification of 1000×. Prior to casting, no additional compositional modifications were introduced, as the aim of this work was to assess the as-cast performance of the alloys under conditions representative of industrial processing. To ensure melt cleanliness, a fluxing treatment was first carried out to facilitate oxide removal, followed by careful skimming of the melt surface. Subsequently, a degassing step using nitrogen (N₂) was performed to minimise dissolved gases and enhance melting quality.

In addition to the production of different specimens by HPDC (tensile, wear samples), the alloys were also cast in commercial sand moulds for thermal analysis, to validate the number and relative fraction of the solidification phases predicted by thermodynamic calculations, as well as to identify their characteristic formation temperatures.

Cooling curves were derived from temperature–time data collected during solidification using a high-speed data acquisition system (FieldLogger v1.5x, Novus Automation, Canoas, Brazil) connected to a laptop, operating at a sampling rate of 10 Hz. To ensure reproducibility, each thermal analysis measurement was repeated at least three times. The experiments were carried out in a custom-designed metallic mould, yielding an average cooling rate of approximately 0.5 °C/s. It should be noted that the thermal analysis (TA) was performed under relatively slow cooling conditions, which are considerably lower than those experienced during the HPDC process. Consequently, the TA results were intended as a complementary tool to experimentally identify the main solidification reactions and compare the relative transformation temperatures among the investigated alloys. Similarly, the equilibrium thermodynamic calculations performed using FactSage represent low-cooling or near-equilibrium conditions, under which a larger number of phases are predicted to form than under the rapid solidification conditions characteristic of HPDC. Therefore, both TA and equilibrium calculations provide a useful reference for assessing the maximum phase formation potential of the investigated alloys, while the Scheil simulations and the experimental observations better represent the non-equilibrium solidification behaviour occurring during HPDC.

2.4. Microstructural Characterisation

2.4.1. X-Ray Diffraction Analysis

Phase identification was conducted by X-ray diffraction (XRD) using a Philips X’Pert Pro MPD PW3040/60 diffractometer (Malvern Panalytical Ltd., Malvern, UK) with Cu Kα radiation operating at 40 kV and 40 mA (1.6 kW). Diffraction patterns were acquired over a 2θ range of 10–90°, with a step size of 0.02° and a counting time of 2 s per step. The crystalline phases were identified by comparison with reference patterns from the PDF-2 database of the International Centre for Diffraction Data (ICDD). The analyses were performed on samples extracted from the gauge section of the tensile specimens.

2.4.2. Optical and Scanning Electron Microscopy

Samples for metallographic analysis were prepared following standard preparation procedures, including selection of the extraction area, cutting, mounting, grinding, and polishing. Samples were extracted from the neck region of the tensile specimens to obtain representative microstructural information using a SiC disc cutter, model Mecatome T 255/300 (Presi, Grenoble, France). Subsequently, the samples were hot-mounted in resin using a MultiFast mounting press, model CitoPress-1 (Struers, Madrid, Spain). Finally, the samples were ground using SiC papers of different grit sizes (#240, #400, #600, #800, #1200, and #2500) and polished with a cloth and a 3 μm diamond suspension using a Mecatech 334 polishing machine (Presi, Grenoble, France).

The microstructure of each alloy was examined using an optical microscope (Leica DMI5000 M, Leica, Wetzlar, Germany) and a scanning electron microscope (Quanta 450, FEI, Hillsboro, OR, USA) equipped with energy-dispersive X-ray spectroscopy (EDS) for compositional analysis.

ImageJ software (version 1.54t) was used to perform the quantitative image analysis of the α-Al grain structure and the primary intermetallic particles. The average grain size was determined by analysing more than 50 α-Al grains from optical micrographs acquired at 1.000× magnification. The primary intermetallic particles were quantitatively analysed using micrographs acquired at 200× magnification. Their area fraction, equivalent diameter, maximum Feret diameter, and circularity were determined by considering only particles with an area greater than 50 μm², thereby excluding fine secondary particles and ensuring a representative characterisation of the primary intermetallic population.

2.5. Physical Properties

The density of the alloys was determined using the Archimedean method with a high-precision balance (BC Memory, Orma, Milan, Italy) with an accuracy of 0.01 mg.

Volumetric porosity was estimated by comparing the experimental density with the theoretical density calculated from the actual alloy composition. At least three samples were analysed for each condition to ensure reproducibility. In addition, both qualitative and quantitative analyses of the pore area fraction were carried out on optical micrographs acquired at 200× magnification using ImageJ software.

Electrical conductivity (EC) measurements were carried out using a portable Autosigma 3000 device (SigmaTest GmbH, Aachen, Germany). Up to five measurements were taken at different locations to ensure representativeness and reproducibility. Electrical conductivity is closely related to the amount of solute in solid solution and can be used as an indirect indicator of precipitation state, often exhibiting an inverse relationship with mechanical strength.

2.6. Mechanical Properties

2.6.1. Hardness

Vickers hardness tests were performed at RT according to UNE-EN ISO 6507-1 [42] using a Shimadzu FV-700 hardness tester (Shimadzu Corporation, Kyoto, Japan) with an applied load of 3 kgf. Measurements were taken from the surface towards the interior of the metallographic samples. A clear gradient was observed, with higher hardness values noted near the surface than in the core, which is consistent with the typical behaviour of HPDC components. This phenomenon is commonly associated with the formation of a refined surface layer (“skin”), characterised by very fine primary aluminium dendrites that enhance the mechanical response of the material [43]. The reported hardness values represent the average of five indentations per sample.

2.6.2. Tensile Properties

Tensile tests at RT were conducted in accordance with UNE-EN ISO 6892-1 [44] using a universal testing machine (Instron 5500R6025, Norwood, MA, USA) with a load capacity between 1 and 100 kN. Temperature tensile tests at 200 °C were carried out following UNE-EN ISO 6892-2 [45]. Before testing, specimens were heated from RT to 200 °C over a period of 30 min and subsequently held at this temperature for 15 min by using a controlled heating chamber (Instron 3119-007, Norwood, MA, USA).

The tensile specimens had a cylindrical geometry with a reduced gauge section. The overall length was 95 ± 0.2 mm, with a gauge length of 33 ± 0.2 mm and a gauge diameter of 6.325 ± 0.2 mm. The grip section was designed with an enlarged diameter of 11.85 ± 0.5 mm to facilitate gripping. Smooth transitions between the gauge section and the ends were incorporated to avoid stress concentration and promote uniform deformation within the gauge region. At least three specimens per alloy and condition were tested to ensure reproducibility.

2.6.3. Compressive Properties

Uniaxial compression testing was conducted at RT following the ASTM 9 standard [46]. Test specimens with a cylindrical geometry (12 mm in diameter and 20 mm in length) were prepared by sectioning the ends of previously tested tensile samples. The experiments were conducted using the same universal testing machine employed for tensile testing (Instron 5500R6025, Norwood, MA, USA). A minimum of three tests per condition was performed to ensure consistent results.

2.6.4. Fracture Surface Analysis

The fracture surfaces obtained after tensile and compression testing were analysed by scanning electron microscopy (SEM) at magnifications up to ×2500. Energy-dispersive X-ray spectroscopy (EDS) was also used to identify elemental distributions and characterise fracture mechanisms.

3. Results and Discussion

3.1. Chemical Composition and Casting Quality

Table 3. Elemental compositions of experimental alloys.

Targeted Alloy	Ref.	Method	Al	Mg	Si	Cu	Zn	Cr	Zr	Others
Al72Mg10Si5Zr3Cu10	H1	OES	76.3	9.9	3.6	7.6	0.6	-	1.9-	Balance
		SEM	72.8	12.2	5.2	7.9	1.35	-	-	Balance
Al67Mg10Si5Zr3Cu10Zn10	H2	OES	65.2	9.0	3.0	10.6	11.5	-	1.9	Balance
		SEM	73.2	12.3	5.8	6.8	1.3	-	-	Balance
Al66Mg10Si5Zr3Cu10Zn10Cr1	H3	OES	70.9	7.6	2.7	8.3	9.3	0.6	1.9	Balance
		SEM	63.7	10.7	4.5	7.6	8.9	0.8	-	Balance
Al64Mg10Si5Zr3Cu10Zn10Cr3	H4	OES	68.9	7.6	2.7	8.3	9.3	2.6	1.9	Balance
		SEM	65.8	11.5	4.6	7.8	9.3	0.6	3.1	Balance

summarises the final chemical composition of the investigated alloys as determined by OES and SEM-EDS techniques. The OES results represent the bulk chemical composition of the alloys and generally showed good agreement with the nominal compositions used during alloy fabrication. Owing to its larger sampling volume and higher quantitative accuracy for bulk chemical analysis, OES provides a more representative and reliable measurement of the overall alloying composition. In contrast, the compositions obtained from SEM-EDS showed noticeable variations. These differences arise from the localised nature of SEM-EDS analysis, which is highly sensitive to microstructural heterogeneities such as the Al matrix, eutectic regions, intermetallic particles, and elemental segregation. Therefore, SEM-EDS data should be considered as local compositional information corresponding to specific microstructural constituents rather than as a measure of the overall alloy composition. Accordingly, SEM–EDS was used to identify and characterise the chemical composition of individual phases, while OES was considered the reference technique for determining the bulk alloy composition.

The measurements of Cu, Cr, and Zr contents showed good agreement with the nominal compositions, with only minor deviations that can be attributed to limited dissolution, element losses during melting, and slight compositional inhomogeneities. During processing, Zr and Cr exhibited relatively high stability in the aluminium melt; however, their limited solubility may lead to incomplete dissolution or the formation of intermetallic phases. Although Cu is generally stable under the melting conditions employed, the slightly lower measured Cu content may be associated with incomplete homogenization during alloy preparation. For the Zn-containing alloys, the initial target composition considered approximately 5 wt.% Zn. However, an excess amount of Zn was intentionally added during alloy fabrication to compensate for the potential Zn losses associated with its relatively high vapour pressure during melting. As a result, the final bulk Zn contents determined by OES in H2, H3, and H4, were close to 10 wt.%, and these measured compositions were therefore used as the reference throughout the present study.

3.2. Solidification Behaviour and Thermal Analysis

Figure 3 shows a representative Scheil solidification for the H4 alloy, calculated using its experimentally measured chemical composition; however, the complete set of predicted phases, based on their experimentally measured chemical compositions, is summarised in

Table 4. Summary of FactSage-calculated phases under equilibrium and Scheil conditions.

Alloy	Equilibrium Phases (Maximum wt.%)	Scheil Phases (Maximum wt.%)
H1	ZrSi (1107—528; 2.48), Mg2Si (594; 9.83), FCC (574; 73.84), Al3Zr (542—520/481; 3.52/3.58), Al7Cu (525—447; 5.47), Al2CuMg (510; 16.96), Mg32(Al,Zn)49 (212; 1.00)	ZrSi (1107; 2.48), Mg2Si (594; 7.87), FCC (574; 74.24), Al2CuMg (509; 14.63), MgZn2 (467; 0.44)
H2	ZrSi (1125—489; 2.48), Mg2Si (581; 8.17), FCC (545; 63.83), Al7Cu (498—230; 5.80), MgZn2 (480; 12.49), Al2CuMg (480; 10.62), Mg2Zn11 (479—307; 12.07), AlCuZn (307; 9.42), Al3Zr (230; 3.59)	ZrSi (1125; 2.48), Mg2Si (581; 6.20), FCC (545; 64.32), Al2CuMg (483; 6.19), MgZn2 (481; 11.54), Mg2Zn11 (480; 9.25)
H3	ZrSi (1103—486; 2.48), Al11Cr2 (730—647; 1.78), Al7Cr (648—514; 2.69), FCC (558; 69.14), Mg2Si (557; 7.36), MgZn2Al (520—351; 3.56); Al7Cu (496—230; 5.77); Mg2Zn11 (481—362; 11.36), Al2CuMg (480; 12.54), MgZn2 (480; 10.20), Al13Cr4Si4 (362—160; 1.93), AlCuZn (286; 9.52), ZrAl (230; 3.59)	ZrSi (1103; 2.48), Al11Cr2 (730; 1.77), Al7Cr (648; 0.56); FCC (558; 69.35), Mg2Si (557; 5.41), Al2CuMg (483; 4.96), MgZn2 (481; 8.34), Mg2Zn11 (480; 7.13)
H4	ZrSi (1132—440; 2.48), Al11Cr2 (812—629; 9.66), Al7Cr (633—482/463—385; 11.98/3.92), Mg2Si (574; 7.36), FCC (552; 58.85), MgZn2Al (521; 16.29), Mg2Zn11 (482—362; 16.56), Al7Cu (482—250; 5.73), Al2CuMg (387—150; 16.57), Al13Cr4Si4 (385—6.90), MgZn2 (362; 11.17), ZrAl (275; 3.59), AlCuZn (715; 14.15)	ZrSi (1132; 2.48), Al11Cr2 (812; 9.66), Al7Cr (633; 0.37), Mg2Si (571; 5.45), FCC (553; 61.19), Al2CuMg (483; 4.96), MgZn2 (482; 9.07), Mg2Zn11 (480; 6.83)

.

For all alloys, solidification was dominated by the formation of the FCC α-Al matrix, which remained the major constituent, accounting for approximately 59–74 wt.% depending on the alloy composition and solidification conditions. The solidification sequence was initiated by the precipitation of primary Zr-containing phases (ZrSi), followed by Mg₂Si and, in the Cr-containing alloys, Cr-rich intermetallics such as Al₁₁Cr₂ and Al₇Cr. Although the fundamental Al–Mg–Si solidification framework was preserved throughout the alloy series, the progressive addition of Zn and Cr substantially modified the secondary phase constitution, particularly during the final stages of solidification.

Thermodynamic calculations predicted that the first solidification reactions occurred between approximately 1100 and 1130 °C with the precipitation of ZrSi, well before the nucleation of the α-Al matrix. The aluminium-rich FCC phase subsequently formed between approximately 574 and 545 °C, while the remaining Cu-, Zn- and Cr-containing intermetallics precipitated over a broad temperature range that extended to lower temperatures. This wide solidification interval reflects the increasing alloying complexity and the progressive enrichment of the remaining liquid during solidification.

As expected, equilibrium calculations predicted a larger number of stable intermetallic phases than Scheil simulations. Under Scheil conditions, the predicted phase assemblage became considerably simplified owing to the limited diffusion in the solid state, favouring segregation-driven phase selection and suppressing the formation of several low-temperature equilibrium phases. This behaviour is fully consistent with the rapid solidification conditions characteristic of high-pressure die casting.

For the H1 (Al–Mg–Si–Zr-Cu) alloy, both equilibrium and Scheil predictions showed a relatively simple solidification sequence, dominated by Mg₂Si and α-Al, together with the early precipitation of ZrSi. Under equilibrium conditions, additional Al₃Zr particles were predicted to form during the later stages of solidification, while Cu mainly precipitated as Al₇Cu and Al₂CuMg. Under Scheil conditions, the phase assemblage became simpler, with Al₂CuMg remaining the only significant Cu-containing intermetallic, whereas Al₇Cu and Al₃Zr were suppressed. Only a minor fraction of MgZn₂ was predicted under Scheil conditions, which is consistent with the low residual Zn content measured in this alloy. Overall, the small differences between equilibrium and Scheil calculations indicate limited microsegregation and confirm that the solidification behaviour of H1 is primarily governed by the Al–Mg–Si system, with minor contributions from Cu- and Zr-containing intermetallics.

The introduction of Zn in H2 (Al–Mg–Si–Zr-Cu-Zn) substantially increased the complexity of the system. In addition to ZrSi, Mg₂Si, and the FCC matrix, equilibrium calculations predicted the formation of several Zn-containing intermetallic phases, including MgZn₂, Mg₂Zn₁₁, and AlCuZn, together with Cu-rich phases such as Al₇Cu and Al₂CuMg. Under Scheil conditions, the phase assemblage was simplified, with the MgZn₂ and Mg₂Zn₁₁ fractions (19%) becoming the dominant Zn-bearing phases, whereas AlCuZn and Al₇Cu were suppressed. This behaviour reflects the suppression of some low-temperature equilibrium phases under rapid solidification while maintaining the preferential segregation of Mg and Zn in the remaining interdendritic liquid.

For the H3 (Al–Mg–Si–Zr-Zn–Cu–Cr) alloy, the addition of 1 wt.% Cr modified the solidification behaviour by promoting the precipitation of primary Cr-rich intermetallics, mainly Al₁₁Cr₂ and Al₇Cr, which formed before or during the early stages of α-Al solidification owing to the extremely low solubility of Cr in aluminium.

Equilibrium calculations also predicted the formation of additional Zn-containing phases, including MgZn₂, Mg₂Zn₁₁, MgZn₂Al, and AlCuZn, together with Cu-rich intermetallics. Under Scheil conditions, the predicted phase assemblage became less complex, with Al₁₁Cr₂ remaining the main Cr-rich phase, while the fraction of Al₇Cr was considerably reduced and several low-temperature equilibrium phases were suppressed. Simultaneously, MgZn₂ and Mg₂Zn₁₁ became the dominant Zn-containing intermetallics, reflecting the progressive enrichment of Zn and Mg in the remaining liquid during non-equilibrium solidification.

For the H4 (Al–Mg–Si–Zr–Cu–Zn–Cr, 3 wt.% Cr) alloy, the higher Cr content produced the greatest modification of the solidification sequence. Equilibrium calculations predicted a marked increase in the fraction of primary Cr-rich intermetallics, particularly Al₁₁Cr₂ and Al₇Cr, together with Zn-containing phases such as MgZn₂, Mg₂Zn₁₁, MgZn₂Al, and AlCuZn, as well as Cu-rich phases, including Al₂CuMg and Al₇Cu. Under Scheil conditions, the phase assemblage was considerably simplified, with Al₁₁Cr₂ remaining the dominant primary Cr-rich phase, whereas Al₇Cr was predicted to form only in trace amounts (<1 wt.%) during the initial stages of solidification. The suppression of several equilibrium phases together with the preferential formation of MgZn₂ and Mg₂Zn₁₁ illustrates the strong influence of rapid HPDC solidification on phase selection. Furthermore, the substantial increase in the predicted fraction of Cr-rich intermetallics is consistent with the larger and more numerous primary Cr-containing particles observed experimentally in this alloy.

Overall, the comparison between equilibrium and Scheil calculations clearly demonstrated the strong influence of rapid solidification on phase selection. Whereas equilibrium calculations predicted a larger variety of stable intermetallic landscapes, the Scheil simulations indicated that non-equilibrium solidification favours a reduced phase assemblage dominated by the FCC α-Al matrix together with Mg₂Si, Cr-rich intermetallics, and Mg–Zn phases. This behaviour is consistent with the limited solid-state diffusion expected under high-pressure die casting conditions and provides a solid thermodynamic framework for interpreting the evolution of the microstructure as a function of alloy composition.

Cooling Curves and Phase Formation Temperatures

Figure 4 illustrates a representative cooling curve and the associated phase transformations for H4, and

Table 5. TA characteristic phases and solidification temperatures for the investigated alloys.

Solidification Parameters	H1	H2	H3	H4
Ta nucleation liquidus (Ta Prel. Nuc.)	628.7	615.6	630.2	628.7
Ta nucleation liquidus (Ta Nuc. Liq.) FCC-Al	597.8	581	596.5	596.1
Ta maximum on the liquidus (Ta Max. Liq.) FCC-Al	576.9	557.3	556.5	555.4
Ta nucleation eutectic (Ta Nuc. Eut.)	573.9	551.7	552.0	551.2
Ta maximum on the eutectic (Ta Max. Eut)	572.4	548.7	549.0	549.6
Ta nucleation post-eutectic1 (Ta Nuc. PEut.1)	502.1	476.1	474.8	482.7
Ta maximum on the post-eutectic1 (Ta Max.PEut.1)	505.1	474.5	474.3	474.1
Ta nucleation post-eutectic2 (Ta Nuc. PEut.2)	-	469.4	468.8	469.4
Ta maximum on the post-eutectic2 (Ta Max.PEut.2)	-	466.9	467.2	467.8
Ta solidus	487.4	446.3	447.0	447.8

summarises the solidification temperatures obtained for the corresponding phases in the four experimental multicomponent alloys. Thermal analysis results revealed systematic changes in the solidification behaviour of the studied alloys as a function of alloying additions.

Alloy H1 showed the simplest solidification path, with a relatively narrow solidification interval and a final post-eutectic event around 502–505 °C, consistent with the formation of Cu-rich intermetallics such as Al₂Cu and Al₂CuMg. In addition, eutectic reactions occurred relatively close to the liquidus temperature, indicating a limited solute redistribution during solidification.

The addition of Zn in H2 shifted both the eutectic and post-eutectic reactions to lower temperatures and significantly decreased the solidus temperature, indicating enhanced solute enrichment in the interdendritic liquid and a broader solidification interval. The appearance of additional post-eutectic reactions, which were absent or less pronounced in H1, suggested the formation of Zn-rich intermetallic phases during the final stages of solidification, which is in agreement with the calculated appearances of Mg₂Zn₁₁, AlCuZn, and MgZn₂.

For H3 alloy, the higher preliminary nucleation temperature was consistent with the early formation of Cr-rich intermetallic compounds predicted by Thermo-Calc. However, the subsequent liquidus and eutectic events were expected to remain close to those of H2-H4, indicating that Cr mainly affected the early stages of solidification, while the later stages were still governed by the segregation of Zn-, Mg-, and Cu-rich phases in the residual liquid.

In H4 alloy, the higher Cr content further modified the solidification path by promoting the early formation of Cr-rich intermetallic phases. The low solidus temperature, comparable to that of H2 and H3, indicated a strong tendency toward microsegregation. Unlike H1, the post-eutectic reactions occurred over a broader temperature range and without clear recalescence, suggesting a more gradual, diffusion-controlled formation of intermetallic phases. This behaviour is consistent with strong solute segregation and with the diffusion-controlled phase formation in the residual liquid.

Overall, the progressive decrease in solidus temperature from H1 to H4 indicated an increasing tendency toward microsegregation as the alloying complexity increased. These thermal analysis results support the thermodynamic predictions and provide experimental validation of the increasingly complex solidification sequences expected in the multicomponent alloys processed by HPDC.

3.3. Microstructural Characterisation

3.3.1. Phase Identification by XRD

Figure 5 presents the XRD patterns obtained for the four investigated multicomponent alloys. The experimentally identified phases were compared with the equilibrium and Scheil thermodynamic predictions obtained using the experimentally measured chemical compositions.

The base alloy (Al–Mg–Si–Zr–Cu) exhibited an aluminium matrix with a face-centred cubic (FCC) crystal structure (Fm$\overline{3}$m), together with the Mg₂Si and Al₂CuMg phases.

In addition, the presence of primary Al₃Zr was confirmed, which is in agreement with the thermodynamic prediction of Zr-containing intermetallic formation during the later stages of equilibrium solidification, although the specific phase constitution differed slightly between the equilibrium and Scheil conditions. No Zn-containing phases were detected by XRD, which agrees with the thermodynamic calculations predicting only a negligible MgZn₂ fraction (0.44 wt.%) under Scheil conditions, resulting from the low residual Zn content (0.6 wt.%) measured in this alloy. This limited amount is likely below the detection limit of laboratory XRD.

With the addition of Zn in the H2 alloy (Al–Mg–Si-Zr-Cu–Zn), the same major phases were detected. However, no Zn-containing intermetallics were identified by XRD despite the prediction of MgZn₂ and Mg₂Zn₁₁ formation. This discrepancy may arise from the rapid solidification conditions inherent to HPDC, which can modify the precipitation sequence and suppress the formation of thermodynamically stable Zn-rich intermetallics. Furthermore, the simultaneous precipitation of the S-Al₂CuMg phase consumes part of the available Mg, locally altering the Zn/Mg ratio and affecting the nucleation and growth of Mg–Zn phases. Consequently, although thermodynamic calculations predict the stability of MgZn₂ and Mg₂Zn₁₁, their actual volume fraction or crystallite size after solidification may remain below the detection limit of laboratory XRD. Previous studies have shown that the stability of Zn-rich intermetallics is strongly influenced by the local Zn/Mg ratio and by competition with other Mg-containing phases [47].

In alloys H3 and H4, after the addition of Cr (1 wt.% and 3 wt.%, respectively), similar phases were identified, including the aluminium matrix, Mg₂Si, and Al₃Zr. In contrast to H2, the Cu-rich intermetallic changed from Al₂CuMg to Al₂Cu, while the Zn-containing phase MgZn₂ was clearly detected by XRD. The replacement of Al₂CuMg by Al₂Cu was expected to increase the amount of Mg available to react with Zn, thereby favouring the formation of MgZn₂ in the Cr-containing alloys. These observations are in good agreement with both the equilibrium and Scheil calculations, which predicted MgZn₂ as one of the major Zn-bearing intermetallics in H3 and H4.

Although the thermodynamic calculations also predicted the formation of Mg₂Zn₁₁, this phase was not experimentally identified by XRD. This difference is attributed to the limitations of CALPHAD-based Scheil simulations, which neglect nucleation barriers, solid-state diffusion, and possible phase transformations during post-solidification cooling. Under the high cooling rates characteristic of HPDC, phase selection is strongly influenced by kinetic effects, which may favour the preferential formation or retention of MgZn₂ while suppressing minor equilibrium phases such as Mg₂Zn₁₁. Similar CALPHAD experiment differences have previously been reported for rapidly solidified Al–Mg–Zn and Al–Zn–Mg–Cu alloys [48,49].

It is noticeable that, in the 2θ ranges of 60–70° and around 78°, the intensity of the α-Al diffraction peaks decreased in H3 and H4 after Cr addition. This reduction suggests a lower matrix volume fraction, consistent with the increased precipitation of secondary phases predicted by the thermodynamic calculations. Additionally, at around 40°, the Mg₂Si diffraction peak in the Cr-containing alloys slightly shifted toward higher 2θ values, indicating a reduction in interplanar spacing [50]. Rather than indicating direct Cr incorporation into the Mg₂Si lattice, this behaviour is more likely associated with lattice distortion and microstrain arising from the increased alloying complexity and the precipitation of Cr-rich intermetallic particles.

Overall, the XRD results were in good agreement with the thermodynamic predictions regarding the major phases present in the investigated alloys. In particular, the experimentally observed evolution of the Zn-containing phases and the Cr-rich intermetallics closely followed the trends predicted by both the equilibrium and Scheil calculations. The remaining differences were mainly associated with minor Zn-rich intermetallics, highlighting the influence of rapid non-equilibrium solidification on phase selection during HPDC.

3.3.2. Optical and SEM + EDS Microstructures

Optical micrographs of the H1–H4 alloys are presented in Figure 6 and Figure 7, showing the characteristic microstructural features developed under HPDC conditions. In Figure 6, a clear distinction between the surface layer and the interior region was observed in all the alloys. The surface layer exhibited a more refined microstructure, along with a reduced presence of coarse intermetallic phases compared to the interior. This refined surface layer has been reported to exhibit improved mechanical properties.

Figure 7 shows the microstructure in the interior of the castings, where the cooling rate is lower. All alloys exhibited a predominantly α-Al matrix with a relatively fine and homogeneous microstructure. Due to the high cooling rates inherent to HPDC, the formation of well-defined dendritic structures was largely suppressed, resulting in a compact morphology with a partially globular appearance. This microstructural gradient is consistent across all compositions and reflects the strong thermal gradients typical of HPDC processing.

Dark, polygonal particles were homogeneously distributed throughout the matrix and were attributed to the primary Mg₂Si phase, based on their morphology and optical contrast. Their presence is consistent with thermodynamic predictions and their formation at relatively high temperatures during solidification.

Also, from H1 to H4, an increase in the number density and refinement of particles was observed, suggesting enhanced nucleation and/or restricted growth associated with the addition of Zn and Cr.

In the base alloy (H1), Mg₂Si particles appeared relatively coarse and sparsely distributed. With the addition of Zn (H2), a slight increase in the amount of fine secondary phases was observed, particularly within the interdendritic regions, indicating enhanced solute segregation during the final stages of solidification. The addition of Cr in H3 and H4 led to a more pronounced refinement of the microstructure, with a higher density of finer particles and a more homogeneous distribution of intermetallic phases throughout the matrix. In H4 (3 wt.% Cr), this effect becomes more evident, and some elongated or plate-like features, likely associated with Cr-rich intermetallics, were distinguished.

In addition to the primary Mg₂Si particles, finer intermetallic constituents formed during the final stages of solidification were observed within the interdendritic or interglobular regions. These phases appeared as light grey features and were associated with Cu- and Zn-rich compounds, predicted by thermodynamic calculations. Their distribution was more pronounced in the alloys containing Zn (H2–H4), indicating increased segregation and solute enrichment in the remaining liquid during the last stages of solidification.

No significant differences in the overall α-Al matrix morphology were observed among the alloys. Minor casting defects were occasionally detected; however, their fractions remained below 5%, and their sizes were typically below 20 µm, which is under the critical threshold typically reported for HPDC alloys. No clear correlation between porosity and alloy composition was observed [51,52].

It is also worth noting that Zr-containing phases were not clearly visible, likely due to their fine dispersion within the matrix. Nevertheless, the presence of Al₃Zr was confirmed by XRD, in agreement with the thermodynamic predictions of Zr-containing intermetallic formation. In the case of Cr-containing alloys (H3 and H4), these intermetallic phases were not easily distinguishable at lower Cr contents and only became more apparent in H4 (3 wt.% Cr), where darker grey particles could be observed. This trend is consistent with the Scheil calculations, which predicted only minor fractions of Cr-rich intermetallics in H3, while considerably higher fractions were predicted in H4.

SEM micrographs of the H1–H4 alloys at a higher magnification are shown in Figure 8, revealing the distribution and morphology of the main intermetallic phases. In H1, the interdendritic regions were mainly composed of Al₂CuMg, together with primary Mg₂Si particles and the associated eutectic structure. Upon Zn addition (H2), although no Zn-containing phase could be unequivocally identified by XRD, SEM observations revealed the presence of Zn-rich interdendritic particles, which, together with the thermodynamic predictions, are consistent with the formation of MgZn₂. The absence of distinct MgZn₂ diffraction peaks may be attributed to the limited amount of this phase observed by SEM, which may be below the XRD detection limit, as well as to peak overlap with other intermetallic phases. A significant microstructural change was observed after Cr addition (H3 and H4), where the Cu-rich phase changed from Al₂CuMg to Al₂Cu. Simultaneously, MgZn₂ became clearly detectable by XRD, while SEM observations revealed a more pronounced presence of Zn-rich interdendritic constituents. This transformation is consistent with the reduced consumption of Mg by the Cu-rich phase, increasing the amount of Mg available to combine with Zn and favouring the formation of MgZn₂. Furthermore, Cr-rich intermetallic particles, mainly identified as Al₇Cr, became progressively more abundant with an increasing Cr content, particularly in H4, in agreement with the thermodynamic predictions, which predicted the precipitation of Cr-rich intermetallics during the early stages of solidification. The presence of Al₃Zr particles was also confirmed in the alloys with higher alloying additions.

3.3.3. Microstructural Evolution and Quantitative Image Analysis

To further clarify the elemental distribution, Figure 9 presents the corresponding EDS elemental maps of the H4 alloy. The results indicated that Zn was mainly concentrated within the interdendritic regions, where it was spatially overlapped with Cu and Al. In several areas, Zn also co-localised with Mg, confirming that Zn segregates into the final interdendritic liquid during solidification and contributes to the formation of a Mg–Zn-rich intermetallic. Although EDS mapping represents chemically complex multiphase regions rather than individual stoichiometric compounds, the enrichment of Zn within the interdendritic regions is consistent with the XRD results and the thermodynamic predictions, which support the formation of MgZn₂ as the predominant Zn-bearing phase.

Also, to provide a quantitative description of the microstructural evolution, the average α-Al grain size, together with the main characteristics of the primary intermetallic particles, were measured by image analysis and are summarised in

Table 6. Quantitative microstructural parameters of the alloys investigated.

Alloy	Grain Size (µm)	Primary Intermetallic Area Fraction (%)	Primary Intermetallic Equivalent Diameter (µm)	Primary Intermetallic Maximum Feret Diameter (µm)	Primary Intermetallic Circularity
H1	8.2	16.65	14.48	66.91	0.214
H2	6.2	16.14	15.18	93.40	0.119
H3	4.7	20.61	12.84	107.80	0.220
H4	4.7	26.02	18.04	125.40	0.132

. The results confirmed a progressive grain refinement from H1 (8.2 µm) to H3 (4.7 µm), while no further refinement was observed for H4, indicating a saturation of the refinement effect at higher Cr contents. At the same time, the area fraction of the primary intermetallic particles increased progressively with alloying complexity, reaching its maximum value in H4. Likewise, the maximum Feret diameter increased from approximately 67 µm in H1 to 125 µm in H4, indicating a progressive coarsening of the primary intermetallic particles with increasing Zn and Cr additions. Although the equivalent particle diameter remained relatively constant for H1–H3, a marked increase was observed in H4, confirming the formation of coarser primary intermetallic phases.

3.4. Mechanical Properties

3.4.1. Hardness Distribution

Figure 10 and

Table 7. Hardness values of investigated multicomponent alloys.

Alloy	H1	H2	H3	H4	AlSi9Cu3
HV3	166 ± 19	177 ± 9	211 ± 12	214 ± 35	103

summarise the Vickers hardness (HV3) values of the investigated alloys, compared with those of an AlSi9Cu3 alloy produced under the same processing conditions. All experimental alloys exhibited significantly higher hardness than AlSi9Cu3 [14], indicating the strong strengthening effect of the multicomponent alloy design.

Hardness evolution can be directly related to the quantitative microstructural parameters presented in Table 6. From H1 to H3, the average α-Al grain size decreased from 8.2 to 4.7 μm, while the primary intermetallic area fraction increased from approximately 16% to 21%. According to the Hall–Petch relationship, grain refinement contributes to increasing hardness by restricting dislocation motion, whereas the higher fraction of hard primary intermetallic particles provides additional reinforcement. These two mechanisms explain the progressive increase in hardness from 166 HV3 in H1 to 211 HV3 in H3.

Among the investigated compositions, the H1 (Al-Mg-Si-Zr-Cu) base alloy exhibited the lowest hardness. Although it contains approximately 10 wt.% Cu, which is known to contribute to strengthening, the overall hardness remained limited compared to the other alloys. The presence of Zr contributed to strengthening through grain refinement and the formation of thermally stable particles; however, its effect was less pronounced than that of Zn or Cr additions. Previous studies on Al–Zn–Cu systems have shown that relatively high Cu contents are often required to significantly enhance hardness [53], yet the values obtained in the present work are comparable or higher due to the combined alloying strategy.

The addition of Zn to the base alloy led to a clear increase in hardness. This improvement can be attributed to the combined effect of solid-solution strengthening and the formation of Zn-containing strengthening phases such as η-MgZn₂ or complex Mg(Zn,Cu,Al)₂ intermetallics. This behaviour is consistent with previous studies on Al–Cu–Zn alloy systems, where Zn promotes both precipitation and solid-solution strengthening mechanisms [54]. Quantitative image analysis further supported this interpretation. Although the primary intermetallic area fraction remained nearly unchanged (16.6% in H1 and 16.1% in H2), the maximum Feret diameter increased from approximately 67 to 93 μm, indicating the formation of larger primary intermetallic particles. At the same time, the average α-Al grain size decreased from 8.2 to 6.2 μm, providing an additional Hall–Petch strengthening contribution. The combined effect of grain refinement and Zn-containing strengthening phases explained the increase in hardness from 166 to 177 HV3.

Increasing the Cr content from 1 to 3 wt.% (H4) did not produce a proportional increase in hardness despite the larger fraction of primary intermetallics (26%). Image analysis revealed that the equivalent particle diameter increased from 12.8 to 18.0 μm, while the maximum Feret diameter increased from approximately 108 to 125 μm. These results indicate a transition from a finer and more uniformly distributed intermetallic population towards coarser primary particles. Consequently, the strengthening associated with the higher intermetallic fraction was partially offset by particle coarsening, leading to a saturation of hardness [55].

As a result, hardness increased progressively from H1 to H4, from 166 HV3 to 214 HV3, with the largest increment occurring between the Cr-free and Cr-containing alloys (approximately 20%). However, the difference in hardness between H3 and H4 remains relatively small (≈1%) despite the higher Cr content in H4, indicating a saturation of the strengthening effect [55]. This behaviour can be attributed to the formation of coarser Cr-containing intermetallic particles and increased microsegregation during solidification at higher Cr levels, leading to a more heterogeneous microstructure. Consequently, the additional Cr did not result in a proportional increase in hardness, as also reflected by the higher scatter observed in H4.

Similar trends have been reported in other HPDC multicomponent Al–Mg–Si-based alloys, where the solubility of Cr is limited and its addition is typically restricted to low levels to control the formation of intermetallic phases [55]. In contrast, the present alloys achieved significantly higher hardness values, even when compared to conventional high-temperature Al–Si–Cu–Mg alloys [56].

3.4.2. Tensile Properties at RT and 200 °C

Table 8. Tensile test values of the investigated alloys tested at RT and 200 °C.

	RT			200 °C
Alloy	YS (MPa)	UTS (MPa)	E (%)	YS (MPa)	UTS (MPa)	E (%)
H1	257 ± 37	294 ± 15.6	0.6 ± 0.1	236 ± 12	272 ± 7.6	0.6 ± 0.02
H2	222 ± 1.7	274 ± 16.9	0.4 ± 0.1	204 ± 35.2	249 ± 16.9	0.5 ± 0.1
H3	238 ± 7.4	243 ± 12.9	0.3 ± 0.01	224 ± 21	239 ± 20.5	0.4 ± 0.1
H4	218 ± 7.4	227 ± 17.7	0.3 ± 0.06	190 ± 26	205 ± 36.7	0.3 ± 0.06

presents the tensile engineering properties: yield strength (YS), ultimate tensile strength (UTS), and elongation (E), of the investigated multicomponent aluminium at RT and 200 °C, while Figure 11 illustrates the evolution of these properties as a function of the combined addition of alloying elements.

At RT, the addition of 5 wt.% Zn (H2) to the base alloy H1 (AlMgSiZrCu) resulted in a clear deterioration of mechanical properties, with reductions of approximately 13% in YS, 7% in UTS, and 20% in elongation. This behaviour can be attributed to the influence of Zn on phase evolution. In particular, the formation of Mg–Zn and Mg–Zn–Cu intermetallic phases may consume part of the available Mg, leading to a redistribution of Mg from Mg₂Si, which is the primary strengthening phase in Al–Mg–Si alloys [57], towards Zn-rich intermetallic phases and increasing microstructural heterogeneity. As a result, the overall strengthening effect was diminished. In addition, Zn tends to segregate in interdendritic regions during solidification, forming coarse Zn-rich particles that promote chemical heterogeneity and localised stress concentration sites, thereby degrading mechanical performance [58].

In this context, the Zn/Mg ratio plays a critical role in controlling solidification behaviour. Previous studies have shown that variations in this ratio significantly affect phase formation and defect susceptibility in Al–Mg–Zn systems [59]. Therefore, the increased Zn content in H2 likely modified the local solidification path and phase distribution, contributing to the observed reduction in tensile properties. This behaviour is consistent with earlier works reporting that the formation of complex Mg–Al–Zn–Cu intermetallic phases can be detrimental to mechanical performance [60].

The addition of Cr further modified the mechanical response of the alloys. For an addition of 1 wt.% Cr (H3), the yield strength increased by approximately 7%, while UTS decreased by about 11% and elongation by 25%. The increase in YS can be attributed to the formation of hard Al–Cr intermetallic particles (e.g., Al₇Cr), which act as obstacles to dislocation motion and enhance resistance to plastic deformation. However, these particles tend to be relatively coarse and heterogeneously distributed, limiting their strengthening efficiency under tensile loading. As the Cr content increased to 3 wt.% (H4), further coarsening, agglomeration, and possible precipitation at grain boundaries occurred. As a result, their effectiveness in strengthening obstacles is reduced, leading to a decrease in the yield strength and a further deterioration of ductility. The formation of these coarse primary Cr-rich intermetallic particles was consistent with the well-known tendency of Cr to form sludge-like intermetallic compounds once its solubility limit in α-Al is exceeded. Consequently, these particles become less effective as strengthening features and instead act as preferential stress concentrators during tensile loading. Therefore, H4 should be regarded as an upper-limit composition intended to investigate the effects of excessive Cr additions rather than as a candidate industrial alloy.

The quantitative microstructural analysis presented in Table 6 provides further insight into the observed tensile behaviour. The evolution of ductility cannot be explained solely by grain refinement, since H3 and H4 exhibited similar average grain sizes (4.7 μm) yet different intermetallic characteristics and mechanical responses. Instead, the size and morphology of the primary intermetallic particles played a dominant role. From H1 to H4, the maximum Feret diameter increased progressively from approximately 67 to 125 μm, while the equivalent particle diameter reached its highest value in H4 (18.0 μm). These coarse primary intermetallic particles acted as preferential stress concentrators, promoting earlier crack initiation and limiting the aluminium matrix’s ability to accommodate plastic deformation. Although the average circularity did not exhibit a monotonic trend, the lower values measured for H2 and H4 indicated a more irregular particle morphology, which is expected to further increase local stress concentrations. These observations are in excellent agreement with the fractographic analysis, where crack initiation was predominantly associated with fractured intermetallic particles rather than with porosity.

Although some studies have reported improvements in ductility with small additions of Cr in Al–Zn–Mg–Cu alloy [61], the behaviour observed in the present work differed significantly. This discrepancy can be attributed to the higher Cr contents used (1–3 wt.% compared to ~0.3 wt.% in the literature), as well as to the increased compositional complexity of the investigated alloys (Al–Mg–Si–Zr–Cu–Zn–Cr). In such multicomponent systems, Cr promotes the formation of a higher volume fraction of intermetallic phases, including Al–Cr and Al–Cr–Fe–Mn particles, which become coarser and more heterogeneously distributed under HPDC conditions.

Furthermore, the combined presence of Zn and Cr enhanced solute segregation and the formation of complex interdendritic constituents, resulting in an increased microstructural heterogeneity. Consequently, both strength and ductility may deteriorate despite the addition of Cr, particularly at higher concentrations, where the beneficial effects of grain refinement and particle strengthening were outweighed by the detrimental influence of coarse intermetallics and segregation [62].

When compared with the reference AlSi9Cu3 alloy [14], the investigated multicomponent alloys exhibited improvements of up to 30% in YS and 13% in UTS, while maintaining similar elongation values. Similarly, when compared with Al–Zn-based alloys containing similar Cu contents, the maximum tensile strength of the present alloys remained competitive or superior [53]. In addition, compared to other multicomponent Al–Mg–Si-based alloys alloyed with low levels of Cr, Sc, or Zr, the YS values obtained in this study were higher, although the UTS was slightly lower [55]. In addition, compared to Al–Mg–Si–Cu-based alloys with lower Cu contents, both YS and UTS of the investigated alloys were significantly improved, although in the case of Cr-containing alloys, the values were comparable.

A comparison between tensile properties and hardness revealed that they do not correlate directly. While hardness increased due to the higher volume fraction of intermetallic phases and precipitates, which acted as reinforcement phases and enhanced resistance to localised plastic deformation, these phases were often coarse, brittle, and heterogeneously distributed. As a result, they negatively affected the overall mechanical behaviour under tensile loading by acting as stress concentrators and promoting crack initiation [58].

At 200 °C, a general reduction in YS and UTS was observed for all alloys, while elongation tended to increase slightly, except for sample H4. The H1 alloy showed a decrease of approximately 8% in YS and 7% in UTS, with a slight increase in elongation. The Zn-containing alloy (H2) exhibited reductions of 8% in YS and 9% in UTS. The alloy with 1 wt.% Cr (H3) showed a moderate decrease of 6% in YS and 2% in UTS, indicating improved thermal stability. In contrast, the alloy with 3 wt.% Cr (H4) exhibited more pronounced reductions of 13% in YS and 9% in UTS, while elongation remained nearly constant.

Compared to the reference AlSi9Cu3 alloy, which exhibits a significant degradation in mechanical performance at an elevated temperature (with reductions of up to 40% in UTS), the multicomponent alloys demonstrated superior thermal stability, retaining a higher fraction of their mechanical properties at 200 °C.

Similar improvements have also been reported when compared with other multicomponent HPDC alloys, where both RT and elevated temperature tensile properties were enhanced while maintaining comparable elongation values [34].

Although the present study is limited to 200 °C, previous studies on Al–Mg₂Si alloys have shown that Mg₂Si and complex intermetallic phases exhibited good thermal stability. However, at temperatures above approximately 250–300 °C, these phases may coarsen, leading to a gradual loss of strengthening. Therefore, the conclusions regarding thermal stability were limited to the investigated temperature range (up to 200 °C).

Overall, the Cu-containing alloy (H1) exhibited the best combination of mechanical properties at both RT and 200 °C, whereas the alloy with 1 wt.% Cr (H3) showed the highest thermal stability, with the smallest reduction in mechanical properties at elevated temperatures. Nevertheless, the reductions observed in all alloys were relatively limited, indicating generally good thermal stability.

3.4.3. Compressive Behaviour

Table 9. Compressive test values of the investigated alloy at RT.

Alloy	YS (MPa)	UCS (MPa)	D (%)
H1	275 ± 17.2	592 ± 11	6.3 ± 1.01
H2	433 ± 31.5	697 ± 19.1	4.6 ± 0.9
H3	425 ± 24.2	696 ± 19.7	3.9 ± 0.4
H4	386 ± 6.8	690 ± 11.4	3.8 ± 0.4

presents the compressive engineering properties, yield strength (YS), ultimate compressive strength (UCT), and deformation (S) of the multicomponent aluminium at RT.

The results showed a progressive increase in maximum compressive stress from H1 to H4, accompanied by a reduction in the total strain. The H1 alloy exhibited the lowest strength (592 MPa) but the highest deformability (6.3%), consistent with its more homogeneous macroscopic deformation and delayed strain localization.

The addition of Zn in the H2 alloy led to a significant increase in compressive strength (57% YS and 18% UCT) but also promoted earlier strain localization and macroscopic cracking under tensile loading. This behaviour could be attributed to the combined effect of solid-solution strengthening and the formation of hard Zn-rich interdendritic intermetallic phases. Under compressive loading, these phases could act as rigid load-bearing constituents, transferring part of the applied load from the softer α-Al matrix. At the same time, the surrounding matrix constrains the particles and suppresses crack opening, limiting crack propagation and increasing resistance to plastic deformation. In contrast, under tensile loading, the elastic and plastic mismatch between the hard interdendritic particles and the ductile α-Al matrix generates local stress concentrations at the particle/matrix interface. These regions become preferential sites for interfacial decohesion or particle fracture, facilitating crack initiation and propagation. Therefore, although the Zn-rich intermetallics contribute positively to compressive strength through a load-transfer mechanism, their coarse interdendritic morphology acts as a stress concentrator under tensile loading, reducing ductility and tensile strength. Previous studies have reported that finely dispersed MgZn₂ precipitates improve mechanical performance by promoting a more homogeneous stress distribution and reducing stress concentration [58]. In contrast, the relatively coarse interdendritic Zn-rich phases observed in the present alloys favour strain localization and premature fracture.

With the incorporation of 1 wt.% Cr in H3, the YS and UCT values remained relatively stable, while the total strain decreased by approximately 15%. This behaviour suggests that the addition of Cr did not provide a significant additional increase in compressive strength but altered the deformation behaviour of the alloy. The Cr-rich intermetallic particles likely acted as obstacles to dislocation motion, while their brittle nature promoted local stress concentration, thereby reducing the material’s ability to accommodate uniform plastic deformation. Although these particles may contribute to strengthening, their effectiveness strongly depends on their size, morphology, and distribution. In addition, Cr promoted grain refinement and modified the phase distribution, leading to a more constrained deformation behaviour and, consequently, lower ductility.

Finally, in H4 (3 wt.% Cr), the UCT remained comparable, while a slight decrease (≈9%) in YS was noted. This behaviour can be attributed to the increased formation of Cr-rich intermetallic phases with a coarser size and more heterogeneous distribution. As the Cr content increased, these particles tended to agglomerate and may also segregate at grain boundaries [51], reducing their effectiveness as strengthening obstacles. At the same time, their coarse and brittle nature promoted slightly higher stress concentration and facilitated crack initiation, thereby limiting both YS and S. This indicates that the strengthening effect of Cr tended to saturate and may even become detrimental at higher concentrations, where the negative impact of microstructural heterogeneity outweighs the potential benefits of particle strengthening. This suggests that Cr did not contribute to further strengthening but promotes the formation of Al–Cr intermetallic particles, which, in the present alloys, tended to be relatively coarse and heterogeneously distributed. Additionally, the introduction of Cr modified the microstructure through grain refinement and changes in phase distribution. Together with the presence of Cr-rich intermetallic particles, these microstructural changes reduced the alloy’s ability to accommodate plastic deformation, resulting in the lowest compressive strain among the investigated alloys.

H1 provided the best overall mechanical performance due to its superior compressive properties, together with the highest ductility and the most balanced compressive response. In contrast, H2 exhibited the highest UCT, making it a promising option for compression-dominated applications, although this is at the expense of reduced tensile performance and ductility.

Compared to the reference AlSi9Cu3 alloy, the H1 alloy exhibited a significant improvement in mechanical performance, with an increase of approximately 100% in YS and 129% in UCT, although this was accompanied by an 80% reduction in S [14]. The substantial increase in strength could be attributed to the higher alloying content and the formation of multiple strengthening phases, such as ZrSi, Mg₂Si, and Al₂CuMg, which contributed through both precipitation and solid-solution strengthening mechanisms. However, this enhanced strengthening was accompanied by reduced deformability, which can be attributed to the higher volume fraction of hard intermetallic phases that restricted plastic deformation and promoted local stress concentration.

Compared with recently developed multicomponent HPDC alloys [34], the compressive strength values of the present alloys were comparable in terms of maximum load-bearing capacity. However, those alloys generally exhibited higher deformation values, indicating a better balance between strength and ductility. Furthermore, when compared with Cantor-type alloys, the present alloys demonstrate a superior strength-to-density ratio, highlighting their potential for lightweight structural applications where high specific strength is required [63,64].

Overall, the mechanical behaviour is governed by the interplay between phase distribution, intermetallic morphology, and microstructural heterogeneity, which affect tensile and compressive responses differently.

3.5. Fracture Behaviour

3.5.1. Fracture Surfaces Under Tensile Loading

Figure 12 presents the fractography of the fracture surfaces of the investigated alloys after RT-tensile testing. In agreement with the mechanical results, the fracture mode progressively evolved from a mixed ductile–brittle mechanism in H1 towards an increasingly brittle fracture in H4. The fractographic observations also indicated that crack initiation was predominantly associated with brittle intermetallic particles rather than with casting defects or porosity. In addition, no evidence of critical porosity defects or large shrinkage cavities acting as fracture initiation sites was observed, which is consistent with the image-based porosity characterisation reported in the following section.

Alloy H1 showed a relatively more homogeneous fracture surface, characterised by the presence of shallow dimples and microvoid coalescence, indicating a limited but more uniform plastic deformation before failure. The absence of large cleavage regions and critical casting defects indicates that fracture was mainly controlled by microvoid nucleation and coalescence within the aluminium matrix, explaining the comparatively higher ductility of this alloy.

In contrast, the addition of Zn in H2 led to a more heterogeneous fracture surface, with flatter regions and a higher fraction of cleavage facets. Although some dimples were still observed, their distribution was less uniform, indicating reduced local plastic deformation. The transition towards flatter fracture regions and cleavage facets indicated that crack initiation increasingly occurred at brittle Zn-containing intermetallic particles, reducing the ability of the matrix to accommodate plastic deformation. This interpretation is further supported by quantitative image analysis, which revealed a lower average circularity of the primary intermetallic particles (0.119 in H2 compared with 0.214 in H1), indicating a more irregular particle morphology, which is expected to increase the severity of local stress concentrations at the particle/matrix interface. This explains the pronounced decrease in elongation.

With the incorporation of 1 wt.% Cr in H3, the fracture surface showed a further increase in brittle features, with more pronounced cleavage planes and a reduced presence of dimples. In addition, a higher number of transgranular cracks was observed, indicating that crack propagation preferentially occurred through the grains rather than along grain boundaries. These observations were consistent with a transition towards a more brittle fracture mode, in agreement with the reduced elongation (0.3%). The presence of Cr-containing intermetallic particles likely acted as local stress concentrators and further restricted the material’s ability to accommodate uniform plastic deformation.

Finally, H4 exhibited the most brittle fracture behaviour. Large cleavage facets, long transgranular cracks, and fractured coarse Cr-rich intermetallic particles were frequently observed. These features indicate that the coarse intermetallic particles acted as preferential crack initiation sites, promoting rapid crack propagation with very limited plastic deformation. Quantitative image analysis also showed the highest intermetallic area fraction together with the largest equivalent particle size and maximum Feret diameter. In addition, the relatively low circularity indicated a more irregular particle morphology, which further increased local stress concentration and facilitated crack initiation. Although H4 did not exhibit the highest pore area fraction, it presented the lowest elongation, further confirming that the deterioration in tensile behaviour was primarily controlled by the evolution of the intermetallic network rather than by porosity.

3.5.2. Fracture Surfaces Under Compressive Loading

Figure 13 presents the fractography of the fracture surfaces of the investigated alloys after RT-compressive testing. All alloys exhibited predominantly ductile fracture features, although the extent of plastic deformation decreased progressively from H1 to H4.

Alloy H1 exhibited a relatively homogeneous fracture surface characterised by the presence of slip lines and microvoid coalescence, indicating a ductile fracture mechanism. The surface appeared relatively smooth at the microscale, with evidence of plastic flow. This behaviour was consistent with its lower UTC and highest deformability, suggesting that the material could accommodate significant plastic deformation before failure.

In contrast, the addition of Zn in H2 resulted in a more heterogeneous fracture surface. Although dimples were still observed, their distribution was less uniform, and regions of cleavage-like facets became more apparent. This was consistent with the significant increase in compressive strength and the reduction in deformation, suggesting that the formation of Zn-rich phases enhanced strength but limited the capacity for uniform plastic deformation.

With the addition of 1 wt.% Cr in H3, the fracture surface exhibited a mixed-mode behaviour, with both ductile and brittle features. Microvoids and dimples were still present, but more pronounced cleavage facets and fragmented regions could be observed, indicating increased strain localization. The presence of Cr-containing intermetallic particles likely contributed to the nucleation of damage and limited plastic deformation. This was consistent with the high compressive strength and further reduced deformability.

Finally, H4 showed the most localised deformation behaviour, with fracture surfaces characterised by larger cleavage regions and reduced dimple density. The presence of cracks and more pronounced brittle features indicated a quasi-brittle fracture mechanism under compression. Despite maintaining a high compressive strength, the H4 alloy exhibited the lowest deformability, suggesting that the increased Cr content promoted the formation of coarse intermetallic particles and enhanced microstructural heterogeneity, leading to early damage initiation and limited plastic flow.

3.6. Physical Properties

Density and Electrical Conductivity

Table 10. Physical properties of the investigated multicomponent alloys.

Alloy	Theoretical Density (g/cm3)	Real Density (g/cm3)	Porosity (%)	%IACS
H1	2.77	2.72 ± 0.01	1.8%	17.6 ± 1.39
H2	2.86	2.90 ± 0.01	1.6%	15.0 ± 0.87
H3	2.88	2.89 ± 0.04	2.8%	10.8 ± 1.32
H4	2.92	2.92 ± 0.03	3.0%	9.2 ± 1.36

summarises the theoretical and measured values of density, porosity, and relative electrical conductivity (%IACS) for the investigated alloys (H1–H4). To further evaluate the possible influence of casting defects on mechanical behaviour, a qualitative and quantitative image analysis of porosity was performed (

Table 11. Image-based porosity characterisation of the investigated alloys.

Alloy	Pore Area Fraction (%)	Equivalent Pore Diameter (µm)	Maximum Pore Size (µm)	Mean Circularity
H1	0.86	4.00	24.73	0.795
H2	2.20	4.48	21.74	0.813
H3	1.22	5.41	33.00	0.739
H4	0.46	3.30	21.70	0.796

). While the porosity values reported in Table 10 correspond to the volumetric porosity estimated from the Archimedes method, the parameters reported in Table 11 provide two-dimensional information regarding pore morphology and distribution.

In particular, the gradual increase in density from H1 to H4 was attributed to the incorporation of heavier alloying elements such as Zn and Cr. At the same time, the presence of Mg₂Si, which exhibited relatively low density, contributed to maintaining the overall density within a narrow range. It should be noted that the measured density values may be influenced by the presence of internal porosity. As shown in Table 10, the estimated porosity values for the investigated alloys ranged from approximately 1–3%, which are consistent with typical values reported for HPDC alloys [52]. A slight increase in volumetric porosity from H3 to H4 with increasing Cr content was observed, which may be attributed to the formation of complex intermetallic phases and the resulting increase in microstructural heterogeneity. However, the pronounced reduction in tensile strength and elongation observed between H1 and H2 cannot be explained by porosity, as both alloys exhibited very similar levels. Likewise, no direct correlation was found between the pore characteristics determined by image analysis and the mechanical behaviour. Although H4 exhibited the lowest pore area fraction determined by image analysis, it also showed the lowest elongation. It should be noted that the Archimedes method was used to estimate the overall volumetric porosity of the specimens, whereas image analysis was employed to characterise pore morphology, including pore size and shape. Since the two techniques evaluate different aspects of the pore structure, slight differences between the measured values are expected and should be regarded as complementary rather than directly comparable [65,66].

As shown in Table 11, all alloys exhibited relatively small pores, with equivalent pore diameters below 6 μm and maximum pore sizes below 33 μm.

Furthermore, the pores presented high circularity values (0.74–0.80), indicating that gas porosity predominated over irregular shrinkage porosity. Previous studies on HPDC aluminium alloys have shown that fine, nearly spherical gas pores have only a limited influence on tensile behaviour, whereas large irregular shrinkage pores, typically 300–500 μm in size, act as preferential crack initiation sites and significantly reduce both ultimate tensile strength and ductility [51,67]. The maximum pore sizes measured in the present work are therefore approximately one order of magnitude smaller than the critical values reported in the literature. These observations are further supported by the fractographic analysis, which revealed that crack initiation occurred predominantly at brittle intermetallic particles rather than at pores. Consequently, the marked reduction in tensile properties is primarily attributed to the increasing fraction of coarse intermetallic phases and the associated microstructural heterogeneity.

In terms of electrical conductivity, a progressive decrease was observed from H1 to H4, with values dropping from 17.6%IACS to 9.2%IACS. This reduction is consistent with the increasing alloying complexity and the addition of Zn and Cr, which enhance electron scattering [68]. Alloying elements in solid solution are known to strongly reduce electrical conductivity, whereas their effect is less pronounced when they are present as secondary phases. Therefore, the decrease in conductivity can be attributed to the combined effect of solute atoms and the formation of complex intermetallic phases [69].

From H1 to H2, the addition of Zn resulted in a reduction in conductivity, which could be associated with both solid-solution strengthening and the formation of Zn-rich intermetallic phases (e.g., MgZn₂ or Mg(Zn,Cu,Al)₂), which exhibit lower conductivity. Additionally, increased segregation contributed to enhanced electron scattering.

For H3 and H4, the further decrease in conductivity was mainly related to the addition of Cr. Cr has a strong effect on electron scattering, both in solid solution and through the formation of Cr-containing intermetallic phases. Furthermore, the increased microstructural heterogeneity and the presence of fine and coarse intermetallic particles contributed to reducing effective electron transport. As a result, the conductivity of these alloys was significantly lower than that of H1, although it remained within the expected range for multicomponent aluminium alloys. Also, in comparison with AlSi9Cu3, the values were comparable [14].

4. Conclusions

Four multicomponent Al-based alloys (H1–H4) were developed within the Al–Mg–Si–Zr-Cu–Zn–Cr system and evaluated for HPDC applications.

• Thermodynamic predictions using FACTSAGE, supported by thermal analysis, provided a reliable framework for understanding phase formation and solidification behaviour. These predictions were validated through experimental characterisation using XRD and SEM/EDS.
• Microstructural and phase analyses confirmed that solidification was dominated by an α-Al matrix and Mg₂Si phases, with the formation of Zr-, Cu-, Zn-, and Cr-containing intermetallics increasing from H1 to H4.
• The alloys exhibited similar densities (~2.77–2.92 g/cm³) and typical HPDC porosity levels (≈1–3%). Electrical conductivity decreased progressively due to an increasing alloying content and enhanced electron scattering.
• Hardness increased from 166 to 214 HV3 with the addition of Zn and Cr. However, the strengthening effect of Cr tended to saturate due to the formation of coarse intermetallic particles.
• The Al-Mg-Si-Zr-Cu system showed the best tensile performance at both RT and 200 °C, while Al-Mg-Si-Zr-Cu-Zn and Al-Mg-Si-Zr-Cu-Zn-Cr reduced tensile properties due to increased microstructural heterogeneity and brittle phase formation.
• Compressive strength was significantly enhanced by Zn addition, reaching ~700 MPa in H2, while ductility decreased with the addition of Cr alloying, maintaining a high compressive strength but promoting strain localization.
• Overall, Al-Mg-Si-Zr-Cu provided the best balance of mechanical properties, whereas Al-Mg-Si-Zr-Cu-Zn was the most suitable for compression-dominated applications. Higher Cr contents increased hardness but reduced ductility and provided limited additional strengthening.

These results demonstrate the potential of Al–Mg–Si–Zr-Cu–Zn-based multicomponent alloys as promising alternatives to conventional HPDC alloys.

Nevertheless, their high yield strength, hardness, and thermal stability make them promising candidates for HPDC components where stiffness, compressive strength, wear resistance, and dimensional stability are prioritised over tensile ductility.

Future work will focus on optimising the Al–Mg–Si–Zr-Cu–Zn alloy system and developing tailored heat treatments to further enhance its mechanical and tribological performance for advanced HPDC applications.