Effects of Copper Content and Thermo-Mechanical Treatment on Microstructure and Mechanical Properties of AlMgSi(Cu) Alloys

Abstract

This study presents the results of research on the influence of different contents of copper in aluminium alloys based on the 6xxx series on mechanical and structural properties. The investigation started with the alloying, and casting four billet variants with different copper content—0.8% Cu; 2B—1% Cu; 3A—1.2% Cu; and 3B—1.4% Cu. The prepared materials were homogenised and extruded on a 500T horizontal press with two different process temperatures and ram speeds ranging from 1 mm/s to 9 mm/s. After heat treating to the T6 and T5 tempers, their mechanical properties were tested. On this basis, the two most promising alloys 2A and 3B were selected and subjected to further tests. After extrusion and heat treatment of the profiles (to F, T1, T2, T5, and T5510), their mechanical properties were determined to select the preferred process parameters. Finally, a structural test based on crystallographic orientation based on the EBSD technique and TEM observations allowed for the characterisation of grain size, dispersoids, and phase analysis. Bright-field (BF) analysis allowed us to compare the deformed areas for T1, T5, and T5510 temperatures. The results showed significant growth in the mechanical properties of all the subjected alloys, and the best properties were shown for a Cu content of 1.4% with a tensile strength of 460 MPa and an elongation of 16% (T5510 tempering). The structural test showed an average grain size of 18 µm to 23 µm and solid solution decomposition differences for different heat-treating parameters.

1. Introduction

Aluminium alloys are increasingly applied in the automotive industry, replacing steel where it is technically feasible. This is primarily due to the significant reduction in vehicle weight when components are made from aluminium alloys. Recent studies on the conversion of vehicle parts, such as doors, body pillars, and other structural components, to lightweight aluminium alloys indicate weight savings of about 35–47% [1]. Due to the high mechanical properties of aluminium alloys with increased alloying element content, their use in the automotive industry has the potential for further growth, surpassing current application levels.

The 6xxx series aluminium alloys are in high demand in the industry, accounting for as much as 90% of extruded aluminium products [2]. In the automotive industry, 6xxx series aluminium alloys are commonly used due to their high mechanical properties, good corrosion resistance, and high weldability. For more demanding applications where strength is critical, 2xxx series aluminium alloys with high copper content are used. Although Al-Cu alloys exhibit high strength, their corrosion resistance is significantly lower compared to the 6xxx series alloys. Numerous studies show that copper as an alloying element negatively impacts the corrosion resistance of aluminium, especially in chloride-containing environments.

In addition, research indicates that an increase in copper content increases the susceptibility of Al-Cu alloys to pitting corrosion [3]. In recent years, intensive research has been conducted to develop new alloys on the border between the 2xxx and 6xxx series. The main objective of these studies is to improve the mechanical properties of the 6xxx series alloys by adding copper while maintaining their favourable properties such as corrosion resistance and weldability.

The properties of metal alloys are determined not only by material parameters but also by the processing conditions. The extensive state of knowledge of the thermo-mechanical treatment of aluminium alloys, and its influence on phase transformations, provides a substantial base of knowledge that determines further research directions, especially for non-standard alloys. Thermo-mechanical treatment, combining plastic deformation with heat treatment, allows for controlled modification of the alloy’s mechanical properties by influencing the precipitation processes taking place in the material. A key aspect of this technology is the ability to induce phase transitions under conditions of high defect density in the crystal lattice, which accelerates the ageing process and enhances the alloy’s strengthening efficiency [4,5,6].

The application of thermo-mechanical processing in aluminium alloys, especially in the automotive industry, enables the development of materials with significantly improved mechanical properties, resulting in wider applications. An example is the EN AW-6060 alloy, in which low-temperature thermo-mechanical treatment leads to an increase in strength of 25–30% compared to traditional precipitation hardening methods [5]. Due to these properties, aluminium alloys are promising materials for advanced engineering applications, including aerospace, automotive, and construction industries, where high strength and durability are required.

One of the 6xxx series alloys used for extruded profiles is EN AW-6061, which contains basic alloying elements such as magnesium (Mg), silicon (Si), and copper (Cu). These elements significantly influence its mechanical properties by forming a reinforcing phase under the appropriate conditions. The copper content in EN AW-6061 typically ranges from 0.15% to 0.4% by weight [7]. The addition of copper is not typical of 6xxx series aluminium alloys. Its presence contributes to improved strength properties by increasing the density and stability of precipitates in the alloy microstructure [2,8]. Published studies indicate that in the ageing treatment, the EN AW-6061 alloy achieves a significant increase in yield strength and tensile strength. This is due to the easier precipitation of MgSi particles (so-called β′′), which effectively impede the dislocation movement in the microstructure. According to available research results, optimum mechanical properties for the EN AW-6061 alloy with a copper content of around 0.3% by weight are obtained after artificial ageing at about 200 °C for 2 h. This results in maximum material strength and hardness. Under these conditions, the yield strength is approximately 275 MPa, and the tensile strength reaches 321 MPa [2].

Among the 6xxx series aluminium alloys, the highest mechanical properties are obtained for alloys with a high content of Si and Mg additives, such as EN AW-6082 or EN AW-6182. These alloys, depending on the production method and heat treatment, reach a yield strength of ~290 MPa and a tensile strength of ~380 MPa [3]. If we would like to refer to the aluminium alloys of the 2xxx series (for example, EN AW-2014 and EN AW-2024), the Cu content in these alloys usually ranges from 3.5% to 5%. Single alloys with a Cu content of 1.8% are used, but they contain other additives such as Fe and Ni [7]. The 2xxx series alloys, depending on the content of alloy additions and the temper of heat treatment, reach yield strengths of 300 MPa and above and tensile strengths of 420 MPa and more [4,9].

This article analyses the impact of thermo-mechanical treatment on the properties of 6xxx series aluminium alloys with copper additions (Al-Mg-Si-Cu). These alloys exhibit a complex microstructure influenced by both equilibrium and metastable phases. The primary alloying elements, Mg and Si, form strongly reinforcing Mg₂Si precipitates, denoted as β. A typical precipitation sequence in these alloys is shown below [10]:

Al SSS → GP zones → β′′ → β′,U1,U2,B′ → β (Mg₂Si)

Once a supersaturated solid solution is obtained, the atoms initially form small clusters, known as Guinier–Preston zones. The next stage after the GP zones is the formation of the β′′ phases, followed by the possible appearance of four metastable phases, depending on the chemical composition, which are the immediate precursors of the stable Mg₂Si phase [11].

In recent years, it has been shown that adding Cu to Al-Mg-Si alloys can enhance the age-hardening effect of the alloy [12,13,14,15,16]. Sunde et al. [14] found that relatively small changes in Cu levels and the Si:Mg ratios significantly impacted the dissolution of the solid solution, thus influencing structural evolution and thermal stability. The addition of copper accelerates the phase transition from β-Mg₂Si to Q′ [17,18]. The formation of the Q phase, which is stable in the presence of the elements Al, Mg, Si, and Cu, is crucial for improving the mechanical properties of the material [19,20]. Q (α = 10.39 Å, c = 4.02 Å, ƴ = 120°) is an equilibrium phase and can be grown as macroscopic crystals outside the Al matrix, and its structure has been determined by X-ray diffraction [18]. Q′ grows as coherent needles inside the Al matrix with needle directions along <100> Al. This phase has a slightly different cell that is more coherent with the matrix (α = 10.32 Å, c = 4.05 Å, ƴ = 120°) in comparison to the Q phase [21].

Studies [22,23] have demonstrated that Cu additions in the range of 0.3–0.5% do not significantly affect precipitation processes in alloys containing magnesium and silicon. It is only when a certain threshold value of Cu addition is exceeded that new Q-type precipitates form. The researchers note that increasing the copper-to-silicon and magnesium ratio leads to a greater number of Q-type precipitates while reducing the number of β-type precipitates. Additionally, the researchers emphasise that Q-type precipitates are more stable and may cause the disappearance of β precipitates in the final phase of the process due to an insufficient amount of constituents for their formation [21]. Intermediate states may occur. Further increases in copper content can lead to the disappearance of β′ and β′′ phases, which are replaced by the θ′ and θ′′ phase characteristics of 2xxx series aluminium alloys. Numerous studies [10,19] have shown that introducing Cu into Al-Mg-Si alloys expands the phase fields and leads to the formation of new equilibrium phases, such as the θ phase (CuAl₂), which, in some cases, may precede the precipitation of the θ′ phase. This affects not only the mechanical properties of the alloys but also the phase transformation sequence during ageing, which is critical for process optimisation.

The optimisation of technological processes together with the design of new materials allows high properties to be achieved. The well-known and well-applied thermo-mechanical treatment leads to an additional effect where second-phase precipitates are released along dislocations and deformation grain boundaries. A transmission electron microscopy (TEM) study revealed that local dislocation-rich regions favour the nucleation and growth of irregularly shaped deposits along defects. Furthermore, homogeneous nucleation of needle-like phases occurs in regions free of deformation-induced defects, but the solid solution is locally deficient in Mg and Si due to previous defect-induced diffusion [24]. Since the diffusion coefficient of copper at 160 °C is an order of magnitude lower [25] than that of Mg and Si, Cu atoms will form phases more slowly as artificial ageing continues [24].

The main problem posed by the copper addition to Al-Mg-Si-type aluminium alloys is a significant increase in their susceptibility to intergranular corrosion (IGC). This is particularly evident in the case where copper segregation occurs at grain boundaries. As demonstrated by Larsen et al. [26], the presence of a nearly continuous copper-rich layer along grain boundaries, surrounded by a copper- and silicon-depleted zone, leads to microgalvanic coupling between the noble copper layer and the adjacent impurity-free active zone. This susceptibility decreases after artificial ageing of the Al-Mg-Si-Cu alloys, which is associated with the disappearance of the copper-rich layer due to the precipitation of the Q phase. The precipitation of the strengthening Q phase inside the grains further reduces the driving force responsible for IGC, thereby lowering the material’s corrosivity. Research shows that optimising the ageing process of these alloys is crucial not only for achieving high mechanical properties but also for ensuring adequate corrosion resistance [27].

The effect of chromium addition on the crystal structure of Al-Si alloys is well known [28]. Chromium acts as a recrystallisation inhibitor in aluminium alloys, and during heat treatment, by forming fine α-Al₁₂Cr₃Si phases, it significantly improves the mechanical properties of the alloys [29,30]. In recent years, much attention has been paid to eliminating the deleterious effects of impurities, such as iron (Fe), which negatively affect the microstructure and mechanical properties of Al alloys. One method of addressing this problem is microalloying [31,32,33,34]. It has been shown that chromium addition significantly reduces the negative effects of iron by transforming the Fe-containing β/α phases into finer chromium-rich dispersoids which are fully coherent with the surrounding matrix [35]. Additionally, these uniformly distributed high-density dispersoids effectively inhibit dislocation movement, leading to improved mechanical properties through dispersion enhancement [33,35].

The aim of the research presented here was to characterise and design the best variants in terms of mechanical properties of the 6xxx series aluminium alloys—AlMgSi with copper additions in the range of 0.8–1.4%. The research work was divided into three stages. In the first stage of the research, optimum extrusion process parameters were determined to ensure the desired level of on-line supersaturation, which is crucial for the proper heat treatment of these alloys. The extrusion susceptibility of the materials studied was also investigated. The second part of the research focused on optimising the heat treatment of F, T5, and T6 alloys by selecting the chemical composition of the alloys based on mechanical properties. The third part of the research was aimed at analysing the influence of the thermo-mechanical treatment of selected alloys on their structure and mechanical properties.

2. Materials and Methods

The chemical compositions of the AlMgSi(Cu) alloy variants for research were designed by the authors on the basis of the presently used 6xxx series aluminium alloys. They do not have exact equivalents in the standards. Four different variants of AlMgSi(Cu) alloys with different copper content were selected: 2A—0.8% Cu; 2B—1% Cu; 3A—1.2% Cu; 3B—1.4% Cu. The designations of the alloy variants were defined by the authors in another project (part of the results was also presented in paper [4]) and the authors decided to continue the designation scheme. All variants had copper contents different from or higher than the normative 6xxx series alloys. The main alloying elements, such as Mg, Si, and Mn, were at identical levels for all alloys. The difference between the assumed and obtained chemical compositions for the main alloying elements, Mg, Si, and Cu, were within the range of ±0.05%. Alloys of variants 2A and 2B had a chromium content of 0.2% and variants 3A and 3B had a chromium content of 0.4% level. The chemical composition of the tested alloys is shown in

Table 1. Chemical composition of casted 4” billets intended for investigation.

Chemical Composition wt. %
Alloy	Si	Fe	Cu	Mn	Mg	Cr	Ti	Zr
AlMgSi(Cu)-2A	1.20	0.04	0.81	0.61	0.79	0.23	0.02	0.15
AlMgSi(Cu)-2B	1.22	0.04	0.98	0.60	0.79	0.22	0.02	0.15
AlMgSi(Cu)-3A	1.21	0.06	1.22	0.62	0.80	0.41	0.02	0.15
AlMgSi(Cu)-3B	1.22	0.05	1.41	0.62	0.80	0.38	0.02	-

. All materials (alloys) for the study in the form of profiles were produced on experimental, semi-industrial semi-continuous casting and extrusion lines installed at the Łukasiewicz Research Network—Institute of Non-Ferrous Metals in Skawina. As the feedstock for alloying, aluminium ingots were used, and the melting process was conducted in a Monometer crucible resistance furnace with a crucible capacity of 300 kg. To the molten pure aluminium at 730–750 °C, the other metallic components were added sequentially: Si, Cu, Mn, Cr, and Mg. Elements Si, Cu, and Mg were added in the metallic form; however, Cr and Mn were added as master alloys AlCr75 and AlMn75. Once all alloying additives were melted, barbotage refining was carried out using a UR0 200 device to remove non-metallic and gaseous impurities. In the next step, a grain refiner (AlTi5B) was added. After the alloy preparation procedure was completed, a sample was taken from the alloy, for chemical composition testing. The chemical composition study was performed on an ARL 4460 spectrometer. The prepared alloys were cast in the form of 100 mm (4 inches) diameter ingots by semi-continuous vertical casting using two Hot-Top crystallisers with continuous lubrication. Casted ingots were approximately 4 m long, free of defects (such as tears or cracks), and had a smooth surface (Figure 1A). After casting, all billets were homogenised (two-stage 530 °C/4 h + 545 °C/4 h and cooling 500 °C/h) in order to obtain a more even chemical composition inside grain structures, and in the grain boundary zones. The homogenisation process can also result in the dissolving of large precipitates in the solution. The structure of the billets after homogenisation was correct; the average grain size was within the range of 120–160 μm. The final step of billet preparation was machining to 96 mm in diameter (to fit into the press container and remove the billet skin that may contain some defects and impurities). The extrusion process was conducted on a semi-industrial line on a ZAMET 5MN press equipped with a 100 mm (4 inches) diameter container (Figure 1B). Tools and ingots were heated in an electric chamber furnace CT 250 EKSO. The dies were heated to 480 °C and the ingots to 490–530 °C. A constant press container temperature of 490 °C was used for the extrusion tests. The press was equipped with an 11 m runway with an “on-line” cooling system. The extruded profiles were cooled by a water wave. The test materials were extruded in the form of profiles of 30 × 9 mm (extrusion ratio λ = 30) and 80 × 5 mm (extrusion ratio λ = 10), as shown in Figure 1C.

The aim of this stage of the research was to determine the optimum parameters for the extrusion process. This was necessary for further analysis of various types of thermo-mechanical treatment on the properties of the extruded profiles. In the first stage of the research, the extrusion of flat profiles was carried out. The extrusion was performed on a 5 MN (500 t) horizontal hydraulic press (ZAMET BUDOWA MASZYN S.A., 83 Zagórska Str., 42-680 Tarnowskie Góry, Poland) in direct mode with a recording of the extrusion force, ram speed, and profile temperature. For this purpose, 30 × 9 mm and 80 × 5 mm flat dies were used. The profile geometries were selected to obtain different plastic deformation factors, expressed in λ (extrusion ratio): λ = 10 for the 80 × 5 mm profile, and λ = 30 for the 30 × 9 mm profile. All 4 variants of the alloy were extruded. The materials for the extrusion process were billets Ø100 × 300 mm heated to 2 different temperatures: 490 °C and 530 °C for the 30 × 9 mm profile, and 490 °C for the 80 × 5 mm profile. This was intended to achieve diversity of material preheating. Both the billets and the tooling were heated using the electrical furnace CT 250 EKSO.

A static tensile test was performed in accordance with the requirements of the standard PN-EN ISO 6892-1:2020-05 on an Instron 5582—max load of 100 kN. In general, each result was based on three test repetitions, and the average value and standard deviation were calculated. Strain deformation was measured with an extremely accurate video extensometer (Instron, Norwood, MA, USA), with a crosshead first speed of 0.37 mm/min to determine yield strength and a crosshead second speed of 3 mm/min until the end of the test. The gauge length was 20 mm. For hardness tests, a Brinnel tester was used—Dia Testor 2c model no. 4074 produced by Otto Wolpert-Werke Ludwigshafen a.R.h.

A Velocity plus camera integrated with an Inspect F50 scanning electron microscope was used to measure the crystallographic orientation by EBSD. The EBSD analysis was performed using Apex and EDAX TSL OIM Analysis 8.0, as well as the ICCD PDF 2 database format. The samples were polished using standard metallographic techniques.

They were then subjected to ion milling using a Leica EM RES101 device. Tests were performed at an accelerating voltage of 20 kV and a working distance of WD = 15 mm. Thin films for TEM observations were made of the samples by grinding and electropolishing both sides on a TENUPOL device in Struers A2 reagent. Observations were carried out on a Tecnai G20 X-twin transmission electron microscope (TEM) with an FEI HAADF and EDS attachment. Phase analyses were performed using ICDD PDF 5+ 2024.

3. Results and Discussion

3.1. Determination of Optimal Supersaturation Parameters for the Extrusion Process

An extrusion test was carried out on all four materials with two variants of temperatures, and a wide set of ram speeds was tested to obtain proper process parameters. Due to the simplified temperature condition of the process, the tolling was preheated to the same level as the material for the process. Based on the authors’ previous experience, profiles were extruded at different deformation rates depending on the billet and tolling temperatures:

-

a 30 × 9 mm profile was extruded at a slider speed of 3, 6, and 9 mm/s at 490 °C, and 3, 5, and 7 mm/s at 530 °C

-

an 80 × 5 mm profile was extruded at a slider speed of 5 mm/s at 530 °C

The piston speed was selected individually based on the first extrusion trials when the maximum extrusion speed was tested. During the trials, each alloy variant and process temperature was extruded with increasing ram speed until hot cracking occurred. As the higher temperature (530 °C) for the 30 × 9 mm profile resulted in lower material strength, the crack appeared at a lower ram speed (around 8 mm/s); therefore, for the mentioned temperature, lower ram speeds were assumed. According to the 80 × 5 mm profile, lower ram speed values were selected due to its thin geometry resulting in hot cracking on the edges with relatively low extrusion speeds.

Online readings of process parameters, such as ram speed, profile temperature at the die output, extrusion force, and the preheating temperature of the billets, die, and container temperatures, allowed for a detailed analysis of the process. Most importantly, it was possible to characterise the influence of the aforementioned parameters on the final structural and mechanical properties of the material. The optimum parameters for a proper extrusion process were also determined. Figure 2, Figure 3 and Figure 4 present the results of the mechanical property tests of each variant after heat treatment to a tempering temperature of T5 (on-line cooling and ageing for 7 h at 175 °C).

Alloy 2A is characterised by a yield strength ranging from 359 MPa when extruded at a speed of 3 mm/s and a process temperature of 530 °C to 392 MPa when extruded at a speed of 7 mm/s and the same process temperature of 530 °C, which represents the lowest level among the tested alloys across the entire range of applied process parameters. It is worth noting the good ductility of the material, which falls within the range of 12.7% to 14.4%. In many cases, the elongation of alloy 2A is higher than that of alloy 2B.

Compared to alloy 2A, alloy 2B is characterised by higher yield strength (392–404 MPa) and tensile strength (437–465 MPa), depending on the process parameters. At an extrusion temperature of 490 °C, alloy 2B is surpassed only by alloy 3B. At higher temperatures, alloy 2B exhibits lower strength properties than alloys 3A and 3B. However, in comparison to alloy 2A, alloy 2B was more difficult to extrude, had a poorer surface quality, and showed more frequent occurrences of hot cracking.

Alloy 3A achieved the highest yield strength of 426 MPa and demonstrated a tensile strength similar to alloy 3B, amounting to 475 MPa (both results were obtained during extrusion at a speed of 7 mm/s with a heating temperature of 530 °C).

As for alloy 3B, high-strength properties were obtained, with a yield strength of 426 MPa during extrusion at a speed of 7 mm/s, a heating temperature of 530 °C, and a tensile strength of 475 MPa under the same process conditions. At the same time, very good ductility was observed, as indicated by an elongation of 15%, achieved during extrusion at a speed of 7 mm/s and a heating temperature of 530 °C.

The analysis of the mechanical properties of aluminium alloys at different extrusion speeds and billet temperatures (490 °C and 530 °C) indicates that the process parameters significantly affect tensile strength (R_m), yield strength (R_p0.2), and elongation (A). The tensile strength (R_m) increases with extrusion speed, particularly for alloys extruded at 530 °C, where the maximum values reach 475 MPa at a speed of 7 mm/s. At a temperature of 490 °C, the values are lower and more uniform.

The yield strength (R_p0.2) also increases with extrusion speed and at higher temperatures, reaching a maximum of 426 MPa at a speed of 7 mm/s and a temperature of 530 °C. The elongation (A) shows greater variability, especially for alloy 3B, where a decrease in ductility is observed at higher speeds, reaching 8.1% at a speed of 9 mm/s.

The optimum process parameters are extrusion speeds of 5–7 mm/s and a temperature of 530 °C, providing high strength and stable results, especially for alloy 3B, which obtains a favourable combination of strength and ductility under these conditions.

The selected temperature of 530 °C showed a significant increase in the strength properties of the extruded profiles along with an increase in the extrusion speed. This phenomenon allows us to assume that the extruded rods have a temperature high enough to achieve a supersaturated solution in the structure. The second parameter determining the achievement of high mechanical properties is the initial time of the cooling process in the water wave, which results from the extrusion speed (ram speed). For a temperature of 530 °C, the minimum speed should be 5 mm/s. At lower extrusion speeds or when the material was heated up to 490 °C, the results of the mechanical properties showed that the alloys were not fully saturated (too low of a material temperature at the start of cooling/supersaturation). For the lower temperature of 490 °C, even at the highest ram speed, the extruded profiles showed much lower properties. The decrease in properties, in this case, was also related to the fact that the heat of plastic deformation generated at high speed caused local overheating of the material and the formation of microcracks (the tested alloys are sensitive to the ram speed and susceptible to hot cracking in the extrusion process).

3.2. Research into Alloys in Basic Heat Treatment Tempers

Profiles of 80 × 5 mm (all variants) extruded at a temperature of 530 °C with 5 mm/s ram speed were subjected to mechanical property tests in order to select heat treatment temper and analyse their influence on the final properties. For this purpose, samples of each extrusion variant were heat-treated to T6 and T5 temper according to parameters shown in

Table 2. Heat treatment parameters of the 80 × 5 mm extruded profile.

Temper	Solutionising	Artificial Ageing
F	Open air cooling on the press runway	-
T5	On-line cooling with water wave (cooling speed > 100 °C/s)	175 °C/7 h
T6	525 °C/2 h/water quenching	175 °C/7 h

. Heat-treating parameters were carefully selected from internal Łukasiewicz Research Network studies that are based on previous research conducted on high copper content alloys (2xxx and 6xxx series). The mentioned research was carried out on the laboratory scale and on both industrial casting and extrusion lines. To compare the properties of the untreated material, samples of the extruded profiles were also tested as the reference material. The results of a mechanical test of the 80 × 5 mm profile are presented in Figure 5.

In general, both heat treatment and increased copper addition resulted in an increase in the strength properties of the alloys studied. Similar results have been obtained in other studies on this subject. The study by Eskin and Kharakterova found that the addition of 0.5–0.6% copper to Al-Mg₂Si alloys (type 6009) significantly increased the tensile strength, reaching up to 445 MPa with natural ageing and 10% deformation, although this improvement was at the expense of reduced ductility, and in the case when artificial ageing was applied after stretching, elongation fell to as low as 7% [36].

For alloy 2A, based on the obtained results, a significant increase in R_p0.2 and R_m values can be observed between the F and T5 conditions, indicating an effective improvement in mechanical properties after heat treatment. In the T6 condition, this alloy reaches a yield strength of 389 MPa and a tensile strength of 455 MPa, making it comparable to alloy 3B.

Alloy 2B had the lowest yield strength among the tested alloys in the T6 condition—370 MPa—and its value for the T5 condition (368 MPa) is also one of the lowest among the tested alloys. It can similarly be classified in terms of tensile strength, with R_m values of 431 MPa for the T5 condition and 442 MPa for the T6 condition, which are among the lowest in the examined set of alloys. The ductility result, expressed as elongation percentage, at 13.5% for the T5 condition, ranks alloy 2B similarly to other alloys. However, the value for the T6 condition, which is 12.3%, was the lowest in the studied group.

Alloy 3A exhibited high-strength properties, with a yield strength of 390 MPa in the T5 condition and 376 MPa in the T6 condition. Its tensile strength was also high, at 453 MPa for the T5 condition and 455 MPa for the T6 condition. However, regarding ductility expressed as elongation percentage (A), this alloy showed noticeably lower values than the other alloys, especially in the T5 condition, where elongation was at 12.8%. In the T6 condition, elongation was also low, at 12.5%.

Alloy 3B achieves the highest tensile strength in the T6 condition, reaching 473 MPa. This indicates that it is the strongest alloy compared to the other alloys tested. Its yield strength in the T6 condition is 397 MPa, which also places it at the top. When analysing the graphs, it is noticeable that alloy 3B shows a rising trend for both R_m and R_p0.2 across all three heat treatment conditions (F, T5, T6), with only small differences between the T5 and T6 conditions. This trend suggests that furnace heat treatment did not improve the supersaturation of the alloy and did not result in a significant further increase in strength compared to treatment at extrusion temperatures.

In terms of ductility, expressed as percentage elongation (A), alloys 2A and 3B show similar values in the T6 condition, with 13.1% and 15.4%, respectively. Although heat treatment significantly improves the strength of these alloys, it leads to a decrease in ductility, which is a typical trend. Alloys 2A and 3B, while maintaining a relatively low reduction in elongation, still exhibit a favourable balance between strength and ductility.

Based on the analysis of the graphs, alloys 2A and 3B can be considered optimal in terms of mechanical strength. Their properties make them very promising in applications requiring high tensile strength and yield strength while maintaining acceptable ductility.

3.3. Structure and Properties After Thermo-Mechanical Treatment

After the analysis of the previous results, the authors decided to focus on the two most promising alloys (according to best mechanical parameters)—2A and 3B. The main difference between them is copper content. The tests were carried out on the same 500 T extrusion press as mentioned in the previous chapter, and billets were extruded in the same process parameters. All extruded profiles were cooled on a press runway with water. The next step was thermo-mechanical treatment according to the following scheme:

• F—Fabricated (extruded). Profiles intended as reference.
• T1—On-line cooled with water wave, and naturally aged.
• T2 1%—On-line cooled with water wave, elongated by 1% and naturally aged.
• T2 2%—On-line cooled with water wave, elongated by 2% and naturally aged.
• T2 3%—On-line cooled with water wave, elongated by 3% and naturally aged.
• T5—On-line cooled with water wave, and artificially aged (175 °C/8 h).
• T5510 1%—On-line cooled with water wave, elongated by 1% and artificially aged (175 °C/8 h).
• T5510 2%—On-line cooled with water wave, elongated by 2% and artificially aged (175 °C/8 h).
• T5510 3%—On-line cooled with water wave, elongated by 3% and artificially aged (175 °C/8 h).

The elongation of samples was conducted on the standard tensile machine Instron 5582—max load 100 kN at room temperature. Additionally, profiles in F temper were made (slowly cooled after the plastic forming process, with no additional heat treatment). The mentioned profiles were used as a reference to the profiles that were the subject of the research.

Heat-treated samples were prepared for SEM/EBSD and the TEM microstructure. Figure 6 shows the crystallographic orientation maps of the 2A and 3B alloys after extrusion from the cross-section in the F condition. In both alloys, the crystallographic orientation of grains was dominant in the directions [111]. Under extrusion conditions, the textures of the profiles were mainly in <100>//ED and <111>//ED orientations and the texture intensity in the centre of the profiles was higher than at the surface. However, at higher Cu, Si, and Mg contents, <111>//ED was found to be the dominant orientation [37]. The average grain size for alloy 2A was about 23 µm, while for alloy 3B, it was about 18 µm. The difference in grain size between alloys 2A and 3B is not large, but the differences may be due to the casting process of the material and the contribution of the TiB2 grinder.

Microstructural studies using transmission microscopy were carried out on samples 2A and 3B at T1, T5, and T5510 3% (Figure 7, Figure 8, Figure 9 and Figure 10). The bright-field (BF) visible structures of the alloys are characterised by a high number of precipitates due to the excess of elements such as Cu and Cr in the analysed alloys. The phases visible in Figure 7 are shaped like elongated “sticks” with rhombohedral corners oriented along the plastic process in the microstructure image. When cut at different angles, they give an image with rounded shapes. The size of the fragmented phases observed in the TEM ranged from 100 nm to about 700 nm. In the alloys studied, their size and shape were comparable. In sample 3A, the phases were denser in some areas than in sample 2A due to the increased Cu and Cr content. The detailed phase analysis shown in Figure 8 revealed the presence of Al_93.38Cu_6.02Fe₂₄Si_16.27 or Al_17.1Fe_3.2Mn_0.8Si_1.9 phases in the structure of the alloys studied. The chemical composition analysis (Figure 8A) for point A did not reveal the presence of Fe, but elements such as Cr and Mn have a similar atomic radius and can perform exchanges at junctional or interjunctional positions, replacing Fe atoms. The differences in the phases are mainly related to the chemical composition, which has minimal influence on the lattice as the interatomic distances are 12,480 Å in one case and 12,560 Å in the other (Figure 8B). The interatomic distances in these compounds are large, which may favour atom exchange and the lack of exact stoichiometry of the compound. A similar study was presented in [38] where α phases in a 6xxx series alloy were analysed on a scanning microscope using Kikuchi line diffraction. On the other hand, Zr was found at point B (Figure 8A) for sample 2A. Zr was added to sample 2A with slightly lower Cu and Cr contents. Zr in aluminium alloys forms Al₃Zr phases [39]. Figure 8 shows Al₃Zr phases surrounded by α phase precipitates and Cu-containing phases.

For [110], a solid solution decomposition was analysed where different heat treatment variants resulted in different microstructures (Figure 9). The microstructure of alloys 2A and 3B after T1 heat treatment (Figure 9A,D) does not show any changes. There is no clear evidence of a solid solution breakdown in the grains. After T5 treatment, in which water wave press supersaturation and artificial ageing at 175 °C for 8 h were carried out, finely dispersed strengthening phases and a free zone in the grain boundary were visible. The T5 treatment revealed differences in the microstructure of alloys 2A and 3B. In the case of alloy 3B, fine oval phases of about 10 nm are visible in addition to elongated finely dispersed phases.

These phases were also observed at grain boundaries (Figure 9F) and at interfaces with the α phase (Figure 10A). EDX analysis results showed higher Cu contents (Figure 10). Due to the size and shape of the phases, these may be stable Q phases [40]. The oval phases for this sample are also visible after T5510 treatment and 3% deformation. Interestingly, the dislocation boundaries appear inside the grains where finely dispersed phases are additionally released. In sample 3B, finely dispersed oval-shaped phases, which are not oriented according to the planes of the densest alignment matrix, appear in dislocation boundaries inside the grains. Dislocation boundaries were also observed in sample 2A, but no phases similar to those in sample 3B were observed. After heat treatment with T5510 and a strain of 3% for both samples 2A and 3A, a high density of dislocations was seen in the microstructure (Figure 7C,F and Figure 9C,F) compared to the T5 heat treatment.

Local grain disorientation (KAM) analysis using EBSD was performed from the longitudinal section of the profiles. In the analysis, the value of the KAM parameter for recrystallised grains was assumed to be below 0.5°. In the KAM maps (Figure 11 and Figure 12), areas with values lower than 0.5° were characterised in blue. The KAM parameter values show that the dislocations were probably stopped at the grain boundaries during dislocation movement (in green).

The lowest value of the KAM parameter was observed for sample 2A in the T2 condition deformed by 1% and 3%. The average local disorientation was less than 0.3. On the other hand, for the T2 treatment, an increase in disorientation of 0.06 was observed for specimen 2A, while for specimen 3A, although the average local disorientation was above 0.3, the increase associated with deforming the specimen to 3% did not increase the average disorientation. When analysing the T5510 treatment, it is possible to observe a higher average disorientation in the microstructure of the samples studied than in the T2 treatment. The dislocation density was significantly higher for sample 2A in the T5510 condition and did not change significantly despite an increase in strain to 3%. Sample 3B deformed to 3% in the T5510 condition had the highest mean disorientation of 0.57 and, as can be seen in Figure 12d, the highest dislocation density of the samples and states analysed. Sample 3A also showed the highest increase in the mean local disorientation (0.25) between the 1% and 3% deformation profiles.

In graphs in Figure 13, Figure 14 and Figure 15, the hardness of each variant of the profiles is shown. When analysing the hardness of both alloys, it should be noted that the natural ageing process results in proper mechanical parameters. Natural ageing, as well as artificial processes, can be used for the precipitation hardening process. For the 2A alloy, the natural ageing time should be a minimum of 7–14 days to obtain a maximum hardness level—95 HB in the case of T1 temper and 99 HB for T2 (3% elongated) temper. Alloy 3B showed noticeably higher hardness after natural ageing with a maximum hardness level—110 HB in the case of T1 temper and 115 HB for T2 (3% elongated) temper. When it comes to artificially aged tempers, both materials showed around a 20% increase in hardness, up to 140 HB for 3B alloy, and 123 HB for 2A. The results clearly showed that T5 and T5510 (prolongated before ageing) had very similar hardness levels (differences were lower than 3% for each investigated alloy).

The results of the study of 2A and 3B alloy samples taken from extruded and heat-treated profiles showed significant differences in structure and properties. In the F condition, the profiles differed in grain size (Figure 6). Alloy 3B had a finer grain size and a slightly higher hardness and tensile strength of approximately 35 MPa compared to sample 2A (Figure 15 and Figure 16). The thermal and thermo-mechanical treatments carried out did not affect the modification and grain growth. TEM observations did not confirm the presence of stable strengthening phases in the T1-treated samples (Figure 7A,D and Figure 9A,D). Probably metastable GP phases [41] were present in the microstructure, resulting in an increase in strengthening (Figure 13, Figure 14 and Figure 15) due to natural ageing. In the two alloys studied, 2A and 3B, a large number of precipitates containing phases Al_93.38Cu_6.02Fe₂₄Si_16.27 or Al_17.1Fe_3.2Mn_0.8Si_1.9 representing a typical α phase [38,42] were observed (Figure 7, Figure 8 and Figure 10). The orientation of these phases was along the direction of the extrusion process. Their abundant presence was associated with the increased Cu, Mn, and Cr contents in the alloys studied. Their presence can block strong strain localisation and lead to increased ductility. Observation of the microstructure of the alloys after T5 heat treatment enabled the identification of finely dispersed strengthening phases within the grains (Figure 9B,E). Due to the differences in the chemical composition of alloys 2A and 3B, and due to the Mg/Si ratio < 1, it was found in the microstructure of alloy 2A that Q′ phases could form in addition to the β′ phase, and in the case of alloy 3B additionally Q and θ phases (Figure 9E) [23,43]. In addition, the presence of minor precipitates in areas of the α phases was observed. Zr-containing phases were observed in alloy 2A and Cu-containing phases in alloy 3B. The alloys studied have identical mechanical properties in the F and T1 states. The formation of phases during both plastic deformation and T5 heat treatment contributes to an increase in strength properties compared to the F and T1 states. It was found (Figure 16,

Table 3. Mechanical properties of extruded 30 × 9 mm profiles in as-extruded condition, and after different thermo-mechanical treatments for T1 and T2 tempers with 1%, 2%, and 3% elongation and T5 temper and T5510 with 1%, 2%, and 3% elongation.

Alloy Designation	Treatment	Rp0.2 [MPa]		Rm [MPa]		A10 [%]
		Mean Value	Standard Deviation	Mean Value	Standard Deviation	Mean Value	Standard Deviation
Alloy 2A	F	211	2	380	1.5	20.8	0.8
	T1	210	1	378	1	19.9	0.05
	T2 1%	253	0.5	383	0.5	19.6	0.35
	T2 2%	275	0.5	381	0	19.1	0.85
	T2 3%	296	0	384	1	16.9	0.5
	T5	362	1.5	413	1	15.6	0.05
	T5510 1%	373	0.5	422	0.5	16.5	0.4
	T5510 2%	377	1	419	0	15.8	0.45
	T5510 3%	384	1	420	0.5	14.9	0.05
Alloy 3B	F	238	2.5	416	0	21.7	0.6
	T1	242	2	422	1	21.5	0.2
	T2 1%	290	0.5	429	0	21.3	0.6
	T2 2%	316	1	434	1.5	19.7	0.2
	T2 3%	380	0.5	445	0	17.3	0.35
	T5	399	2	452	1	16.6	0.4
	T5510 1%	407	0.5	454	0	16.6	0.15
	T5510 2%	413	0.5	460	0.5	16.1	0.1
	T5510 3%	416	1	457	0.5	15.6	0.2

) that the T5 heat treatment resulted in an increase in tensile strength Rm of 35 MPa for alloy 2A and 30 MPa for alloy 3B and a decrease in ductility of about 5% compared to the T1 heat treatment. Artificial ageing performed immediately after extrusion and on-line cooling on the press induces higher lattice stresses than the natural solid solution disintegration that occurs with T1 treatment [44]. The effect of increasing the Cu and Cr content in alloy 3B results in an increase in strength properties by approximately 8% compared to sample 2A as a result of the T5 heat treatment.

Thermo-mechanical treatments improve mechanical properties. Analysis of the hardness results and mechanical properties showed that the T2 treatment, together with deformation in the range of 1% to 3% and natural ageing, resulted in an increase in strength, yield stress, and tensile strength. The yield strength for alloy 2A increased by 14% and for alloy 3B by 23%. In contrast, the increase in tensile strength was much lower for the alloys tested. The increase in tensile strength was 0.2% for alloy 2A and 3% for alloy 3B. The KAM images with EBSD confirm an increase in dislocation density between the deformation of 1% and 3% for the T2 treatment for the alloys tested.

Analogous results were obtained by heat treatment of T5510 with tensile strains of 1%, 2% and 3% and artificial ageing at 175 °C/8 h. Based on the hardness results, there was a slight increase in strengthening compared to the T5-treated samples in both alloys tested. When comparing the mechanical properties between the T2 and T5510 heat treatments, an increase was observed, but differences were found between the treatments between R_m and R_p0.2 with regard to individual deformations. In the case of the T5 treatment, the differences in values were relatively large, while in the case of the T5510 treatment, the increase in R_p0.2 was about 2%, while Rm did not increase significantly in the alloys studied. The differences in the local dislocation density images show a significant increase in the dislocation density in sample 3B deformed by 3%. In specimen 2A, the differences in the dislocation density between a strain of 1 and 3% are not significant. The significant increase in the dislocation density after 3% strain in specimen 3B may be due to the appearance of internal stresses after ageing, where the growth of dispersoids across dislocation boundaries prevented the dislocation migration and annihilation and blocked natural healing processes [45,46]. The increase in the dislocation density observed in the T5510 treatment translates into the highest yield stress values, R_p0.2, for both alloys tested. The analyses indicated higher mechanical properties as a result of the 1.4% Cu addition compared to the 0.8% Cu AlMgSi alloy. This is due to the thermo-mechanical processing associated with the decomposition of the solid solution and the appearance of additional strengthening phases such as Q′, Q, and θ′ within and at grain boundaries and subgrains. The appearance of a high density of dislocations and the separation of finely dispersed phases on them, as well as the low influence of Cr on the formation of alpha phases, made it possible to obtain high-strength properties comparable to those of 2xxx alloys.

4. Conclusions

• The research confirmed that by using a high copper content in the 6xxx series alloys, it is possible to increase the strength properties without reducing the ductility. As a result of the research, optimum thermo-mechanical processing parameters were selected. Structural studies confirmed the importance of the α and β phases and Q finely dispersed phases formed by solid solution decomposition in shaping the properties.
• Extruded and heat-treated sections of AlMgSi(Cu) alloys with increased copper content in the range of 0.8–1.4% show a significant increase in strength properties compared to typical 6xxx series alloys, in which Cu content typically does not exceed 0.4%. This is also the average level of strength properties for the 2xxx series. These results for the investigated alloys compared to the 2xxx series alloys, due to economic (less Cu means cheaper alloys) and technological (less Cu means better extrudability) indicators, pave the way for their future wide commercialisation.
• The influence of the extrusion process parameters on the tensile properties of on-line supersaturated AlMgSi(Cu) alloys is typical of other 6xxx series aluminium alloys. Increasing the temperature of the extrusion process (material temperature) and increasing the extrusion speed promotes a better supersaturation of the profiles in the press run and results in higher strength properties after ageing. For AlMgSi(Cu)-type alloys with higher copper contents, in the range of 0.8–1.4%, the extrusion speed should be controlled and, if necessary, reduced, since too of a high speed will cause hot cracking. The addition of copper to 6xxx series aluminium alloys reduces their susceptibility to extrusion.
• The thermo-mechanical treatment in the form of stretching with controlled deformation of the extruded profiles after supersaturation had the greatest effect on the yield strength of the materials that were subsequently naturally aged (T2 condition). For the profiles stretched by 3%, an increase in yield strength of approximately 100 MPa was found compared to the unstretched profiles (T1 condition). For the artificially aged specimens, controlled desaturation deformation (condition T5510) had a much smaller effect on the yield strength increase (an increase of about 1–2% for the subsequent permanent strain value) compared to the heat-treated-only specimens (condition T5). The effect of the thermo-mechanical treatment on the yield stress was much smaller, ranging from 7 to 23 MPa for alloys 2A and 3B in both T2 and T5510 conditions.
• The higher content of elements such as Cu and Cr in the alloy leads to the formation of a large number of α phase Al_93.38Cu_6.02Fe₂₄Si_16.27 or Al_17.1Fe_3.2Mn_0.8Si_1.9 precipitates with sizes between 100 and 700 nm after the extrusion process. The contribution of these phases as a result of deformation mechanisms helps to maintain the high ductility of the alloy together with high yield and tensile strength.
• In the alloys studied, depending on the heat or thermo-mechanical treatment, differences were found in the redistribution of the finely dispersed strengthening phases formed by the disintegration of the solid solution. The introduced additional deformation and the resulting increase in the dislocation density catalyses the growth of dispersoids. This results, especially in alloys with higher Cu content and T5 treatments deformed from 1 to 3%, in high yield stress values.
• High strength and good elongation (R_p0.2—397 MPa; R_m—473 MPa; A—13.1%) may successfully predispose the tested alloys to the automotive and transportation industry for any construction. where the weighting factor is of critical importance. As examples, all parts designed as extruded profiles can be mentioned (crash energy absorbers, frame structure parts, brackets, etc.). The alloys can be used in a variety of applications, but it is recommended that due to the elevated copper content, they will require additional surface protection (e.g., painting, anodising).