1 Introduction

The total aluminium production has increased more than 2.5 times over the last 20 years [1]. This is due to, amongst others, the recent growth in demand for light, high-strength materials, mainly in construction and transport [2]. The aluminium alloys most commonly used today are those of the 6xxx series, or AlMgSi. The PN-EN 573–3 Standard [3] specifies more than 40 alloys, whilst a dozen or so of them are used in everyday life. Depending on the required mechanical properties, alloys that are used have either a lower content of the main alloying elements (Si and Mg), which results in their lower mechanical properties, or higher content of these elements, which allows achieving higher mechanical properties. For more demanding applications, such as aviation, aluminium alloys of the 2xxx series, i.e. AlCu, are used. Amongst aluminium alloys, these are characterised by the highest strength properties. Currently, many types of AlCu alloys with different mechanical and technological properties are being produced. The most important factors shaping these properties include the content of alloying elements (mainly the addition of copper), as well as parameters of plastic forming and heat treatment [4, 5]. In most of the 2xxx series alloys currently used, e.g. EN AW-2014, EN AW-2017A, EN AW-2024, copper content is in the range of 3.5 ÷ 5 wt% [3]. There are single alloys with copper content below 3.5 wt%, but they are usually designated for special applications and also contain other additives, such as Fe, Ni or Li [3]. In the 2xxx series alloys, the content of Cu is of key importance, as it is precisely the variable solubility of copper in aluminium that enables conducting the heat treatment which, through the precipitation hardening mechanism, enables control and maximisation of mechanical properties. The higher the Cu content in the alloy, the more intense the hardening effect is [4,5,6]. In the 2xxx series alloys, the “θ” Al2Cu phases play a key role in the process of precipitation strengthening [4,5,6,7]. As regards Al–Cu alloys containing also silicon and magnesium as alloying additives, numerous studies and publications describe complex phases of the AlCuMgSi type with different stoichiometric ratios [7,8,9,10,11,12]. In the 2xxx series alloys, the “β” Mg2Si phases are involved in the strengthening process to a lesser extent [13]. The formation of “β” phases in Al–Cu alloys is favoured by the situation when the Mg/Si ratio is over 1.

As already mentioned, aluminium–copper alloys belong to the family of high-strength materials, but unfortunately, compared to aluminium alloys of the 6xxx series, they offer much lower corrosion resistance. Numerous studies show that copper as an alloying element has a detrimental effect on the corrosive properties of aluminium, it especially reduces the resistance to solutions containing chlorides. The research results show that with increasing Cu content, the susceptibility to pitting corrosion of Al–Cu alloys also increases [14]. Another obstacle in a more widespread use of 2xxx series aluminium alloys is their high manufacturing cost. It is higher than the production cost of 6xxx series alloys. This is mainly due to the fact that these alloys are hardly deformable—they are characterised by high deformation resistance and the highest coefficient of friction in plastic working processes. As a consequence, the operation of extruding semi-finished products made from these alloys is carried out at low speeds, which significantly reduces the process efficiency [15, 16]. Added to this is the high price of copper, which is the alloying addition with the highest content. All these factors taken together ultimately lead to the situation when, for economic reasons, AlCu alloys with higher mechanical properties are often abandoned in favour of cheaper, however weaker, AlMgSi alloys. Therefore, despite the rich bibliography and a vast amount of research carried out on aluminium alloys of the 2xxx and 6xxx series, it is reasonable to continue the research, mainly to gain a better insight into the subject and allow for optimisation of the chemical composition in terms of the widely understood functional properties. These assumptions have also accompanied the present research work carried out as part of the project, the aim of which is to develop optimal materials/aluminium alloys for the new design of a suspension system dedicated to the so-called light rail and tram overhead contact lines. In the cited project, the main alloys selected for the overhead contact line suspension structures are the 6xxx series alloys. However, alloys from this series, even after modification of their chemical composition and application of the manufacturing process based on improved heat treatment parameters, allow obtaining the tensile strength in the range of 360–380 MPa. Therefore, attempts were made to examine the functional properties of selected aluminium alloys whose chemical composition is partly based on the 6xxx series (mainly Si and Mg content), but the addition of copper qualifies them to be included into the 2xxx series. The research was focussed on modified aluminium alloys based on the 2xxx series with different copper content, i.e. from 2.5 to 4.5 wt%, the additions of Mg and Si maintained at the same average level as for the 6xxx series alloys, and with the addition of Zr as a recrystallization inhibitor. The key issue in these materials is to obtain a fine-grained structure, as this is the structure that can significantly increase the mechanical properties [17]. The results obtained in the research into the introduction of zirconium in this type of alloys to maximise the structure fragmentation and effectively inhibit recrystallization at elevated temperatures of plastic working and heat treatment have confirmed that alloys with zirconium content up to about 0.2 wt% are characterised by a definitely lower tendency for grain growth [18, 19].

The aim of the studies was to initially develop a new alloy with the tensile strength higher than the tensile strength obtained in typical 6xxx series alloys, i.e. over 400 MPa, characterised moreover by satisfactory hot workability.

2 Test materials and methods

The materials selected for tests were alloys whose reference base is the EN AW-2017A alloy. Three variants of the AlCuMgSiMn alloys with different Cu content: 2.5, 3.5 and 4.5 wt% were tested. The addition of Mg and Si was selected based on the average content of these elements in a typical EN AW-6063A alloy. The Si to Mg ratio was chosen to be around 0.7, i.e. with an excess of Mg to Si. The opposite of typical 6xxx series aluminium alloys is where this ratio is usually above 1 or with an excess of Si. When determining the content of Mn, the average content adopted in the 2xxx series alloys was taken into account, bearing in mind that the addition of this element to aluminium alloys enhances the hardening effect in materials subjected to plastic working and heat treatment carried out directly from the temperature of plastic forming. The last alloying element was Zr added in a content typical for aluminium alloys to obtain the recrystallisation inhibiting effect described in the Introduction. All alloying additives for each variant were maintained at the same level and the only variable factor was the copper content. The alloys were melted in electric crucible furnaces and then ingots with a diameter of 100 mm were cast by the vertical semi-continuous method. Chemical compositions and the designation of their individual variants are presented in Table 1.

Table 1 Chemical composition of EN AW-2017A and EN AW-6063A base alloys (according to PN-EN 573–3:2013) and selected alloy variants (W1-W3) for the study
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Test billets were subjected to the process of homogenising annealing (homogenisation heat treatment) carried out at a temperature of 490 °C (heating for 3 h, holding for 6 h, rapid cooling in air). Then the billets were rolled (peeled) to obtain the final diameter of ø96 mm for extrusion.

Tests of direct extrusion were carried out on an experimental line consisting of a horizontal press (maximum force of 5 MN) with a run-out table and a device for cooling the profiles directly on the run-out table (“on-line” solution heat treatment). Cooling was carried out in the so-called “water wave” ensuring the highest cooling rate. Before extrusion, the stock was heated in a chamber furnace. Dies with an I cross-section and with a rectangular 70 × 6 mm cross-section were used for the process of extrusion (Figs. 1, 2). The processability of the extrusion process was assessed on the basis of extruded model profiles: a flat bar and an I-beam. The dies for these profiles were selected in such a way that it was possible to perform a preliminary assessment of the processability of the extrusion process by macroscopically evaluating the surface of the extruded profiles. The correct surface (without overheating, hot cracks, tears, blisters, etc.) in different places of these profiles (flat surfaces, edges, wall joints) is classified as sufficient processability at this stage of research. The direct extrusion was carried out at a temperature of 460 °C with a ram speed of 1 mm/s for an I-beam, whilst for a flat bar, the speed was 0.5 mm/s. The extruded profiles were tested in as-extruded condition (F), after heat treatment which consisted of solution treatment on the press run-out table and natural ageing (T1), and after heat treatment which consisted of solution treatment in furnace and natural ageing (T4). The heat carried out by the following route: heating in furnace at 500 °C/2 h and cooling in water, natural ageing for 4 days.

Fig. 1
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Die for extrusion I-beam and extruded profiles

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Fig. 2
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Die for extruding 70 × 6 mm flat bars and extruded profiles

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The structure and mechanical properties (hardness and static tensile strength) were examined on the manufactured sections. The structure study mainly focussed on grain size analysis. The analysed samples were cold-included in KEM 605 resin. Then the samples were ground with a sandpaper of various gradations and polished with a diamond polishing suspension of 3 µm granulation. The final stage of the preparation process was electrochemical etching using Barker reagent. Structure studies were also realised using an EBSD detector on a Hitachi Su70 scanning electron microscope (SEM) at 20 kV. The average grain size in the studied materials was determined. To further analyse the structure and phase precipitates in the studied alloy variants, studies were carried out using a transmission electron microscope.

As part of the mechanical property research, hardness tests were carried out using Vickers HV0.2 method on a hardness tester Falcon 500 and Brinell HB according to PN-EN 6506–1:2014. Brinell hardness testing was performed using a 2.5 mm diameter carbide ball with a load of 613N on a hardness tester Duramin 2500. Mechanical properties were also determined in a static tensile test. The experiments were conducted using an MTS 810 hydraulic testing machine with a maximum force of 120 kN, using a strain rate of 10–3 1/s.

3 Test results and discussion

The test extrusion of sections was carried out whilst maintaining constant the dimensions of the stock material and process parameters such as temperature (container, die and extruded material) and ram speed. Table 2 gives the parameters of the extrusion process and the registered process force.

Table 2 Examples of process parameters for extrusion of 70 × 6 mm flat bars and I-beams
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All extrusion tests were successful. The extruded profiles had the correct surface without defects such as cracks, tears, blisters, etc. The registered extrusion force for the flat bar was increasing along with the increasing Cu content in the alloy. For I-beam, the situation was comparable, but these differences were not so distinctly visible because, due to the more developed surface of the profile, and thus the greater impact of friction in the extrusion process, it was carried out at maximum press force. Representative samples were taken from the extruded profiles after heat treatment for structure studies. The study of the structure focussed on the middle part of the profiles, from where samples for mechanical properties were also taken. Possible inhomogeneities of the structure at the surface of the profiles were not thoroughly analysed. The obtained images of the structure of the extruded profiles studied and the subgrain sizes determined from them are presented in the form of histograms and the values of the average subgrain size (d2) are included in Figs. 3, 4, 5, 6, 7 and 8.

Fig. 3
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EBSD structure with subgrain size distribution histogram for alloy W1 (2.5 wt% Cu) at T1 temperature

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Fig. 4
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EBSD structure with subgrain size distribution histogram for alloy W1 (2.5 wt% Cu) at T4 temperature

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Fig. 5
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EBSD structure with subgrain size distribution histogram for alloy W2 (3.5 wt% Cu) at T1 temperature

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Fig. 6
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EBSD structure with subgrain size distribution histogram for alloy W2 (3.5 wt% Cu) at T4 temperature

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Fig. 7
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EBSD structure with subgrain size distribution histogram for alloy W3 (4.5 wt% Cu) at T1 temperature

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Fig. 8
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EBSD structure with subgrain size distribution histogram for alloy W3 (4.5 wt% Cu) at T4 temperature

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In the posted structures, Figs. 3, 4, 5 and 6 from the tested profiles for alloy variants W1, W2 after heat treatment for the T1 condition, (solution heat-treated on the press run-out table) and for T4 condition (after additional heat treatment in furnace), the grain size was homogeneous, fine and comparable for these alloys (variants with 2.5 wt% and 3.5 wt% copper content). In the structures shown in the pictures 7–8, from the extruded profiles from alloy variants W3 (of 4.5 Cu percent weight content) after heat treatment (T1 and T4 conditions), the grains were very fine and homogeneous. The medium subgrain size for W3 variant (~ 3.6 μm) is approximately 25% lower in comparison to W1 and W2 variants (4.6 μm on average). The largest average subgrain size of 5.3 µm was observed for the W2 variant sample in the T1 heat treatment condition. The structure of this sample also shows elongated grains in the direction of material deformation. Similar areas of grain were observed near the surface of the extruded profile where grains differ, and there are areas with grains of a slightly larger size. This was due to the occurrence of slight growth of the grains on the surface of the profiles. Coarse-grain zones were identified in each tested profile. There are no noticeable differences between various heat treatment conditions for a specific alloy variant. To further analyse the structure of the materials studied and identify the phases that strengthen these alloys, studies were realised using a transmission electron microscope. Referring to the attached EBSD key, which is located in the upper right corner of each EBSD map, it can be easily noticed that most of the subgrain orientation is near < 101 > crystal direction. Their smaller number arranged according to the < 001 > direction, which in the case of the fcc materials is a typical orientation for recrystallization.

In AlCu alloys, precipitation strengthening takes place with the intermetallic phases Q, (Al2Cu) and (Mg2Si) [9]. According to various sources, the Q phases are: Al5Cu2Mg8Si6 [7, 8], Al4CuMg5Si4 [7], Al4Cu2Mg8Si7 [10] or Al3Cu2Mg9Si7 [13]. A high Cu content, above 4% by weight, guarantees the presence of phase θ. This phase is also observed at low Cu content of 0.2–0.5 wt%. At high Cu content and at Mg/Si > 1 ratio, in addition to the phase θ, the phase β may also appear, and at Mg/Si < 1 ratio, Q or S phases may form, depending on the amount of Si. Very low Si content favours the formation of S phase, whilst higher Si content favours the formation of Q phase and Si particles. In the studied alloys, the Mg/Si ratio is about 1.4, which would indicate the presence of phases θ and β in the structure. The actual order of formation of precipitates in the 2xxx series alloys strongly depends on the amount of alloying elements, as well as on the history of the material, including production methods (rolling, forging, etc.), heat treatment or natural ageing. Figures. 9, 10, 11, 12, 13 and 14 show the TEM images of selected microstructures of the tested alloys. Two types of precipitates are present—very fine precipitates of the β phase (Mg2Si) and slightly larger θ (Al2Cu) (red arrows for β precipitations and green one for θ) [20]. In the T1 temper, the presence of β (Mg2Si) precipitates is dominant for all variants, higher than in the T4 temper at lower Cu contents.

Fig. 9
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Microstructure of alloy in variant W1 (2.5 wt% Cu) in T1 temperature (a) with indication of β and θ precipitates (b)

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Fig. 10
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Microstructure of alloy in variant W1 (2.5 wt% Cu) in T4 temperature (a) with indication of β and θ precipitates (b)

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Fig. 11
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Microstructure of alloy in variant W2 (3.5 wt% Cu) in T1 temperature (a) with indication of β and θ precipitates (b)

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Fig. 12
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Microstructure of alloy in variant W2 (3.5 wt% Cu) in T4 temperature (a) with presence of β and θ precipitates (b)

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Fig. 13
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Microstructure of alloy in variant W3 (4.5 wt% Cu) in T1 temperature (a) with indication of β and θ precipitates (b)

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Fig. 14
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Microstructure of alloy in variant W3 (4.5 wt% Cu) in T4 temperature (a and b)

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At the highest Cu content, the trend is reversed (Figs. 13 and 14). The presence of primary Si particles was not observed in any of the tested alloys (tempers). The higher the Cu content, the smaller the size of the θ type (Al2Cu) precipitates. These precipitates are also smaller after T4 treatment compared to T1. The electron diffraction images showed the coherence of the observed precipitations θ with the matrix. Small β (Mg2Si) precipitates are evenly distributed throughout the tested volume—no places with a privileged distribution were observed. It is worth noting that there are no clusters of precipitates along the grain boundaries. A difference in the amount of β precipitates depending on the type of treatment used was observed. In low Cu content, specimen on-line solution treatment and natural ageing result in the formation of higher amount of β phase, whilst at higher Cu contents, the trend is reversed. After solution treatment in the furnace and natural ageing, higher amount of β precipitates was noted. The variable amount of precipitates of both phases β and θ does not allow to show their significant influence on the mechanical properties depending on the T1 or T4 treatment applied. However, there is a clear influence in increasing of these properties after the heat treatment.

Microhardness measurements were taken in areas a, b, c (Fig. 15). Based on the obtained results, the average hardness was determined in individual areas along with the standard deviation. The determined values are presented in respective diagrams (Fig. 16). In areas a and c, five measurements were taken, whilst in area b, three measurements were taken, due to the smaller area of the sample.

Fig. 15
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Scheme of the test samples with indication of measurement points

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Fig. 16
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Mean HV0.02 hardness in specific areas for samples in F-temper (raw material after processing) and T4 temperature (after heat treatment)

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The microhardness of samples in F-temper is characterised by high heterogeneity, showing differences both within individual areas (reflected in high values of the standard deviation) and between these areas. The highest microhardness was obtained for the alloy with composition W1. As a result of heat treatment, the heterogeneity in individual areas decreases, but the effect of strain hardening in area b is still visible in all samples. Moreover, compared with as-extruded condition (F-temper), all alloys show an increase in microhardness, the highest for the alloy with composition W3.

Microhardness measurements taken on alloys with different copper content were completed with measurements made by Brinell hardness. Measurement points were marked on the prepared surfaces of I-beam according to the scheme in Fig. 15. The aim was to obtain more averaged results, which this method of measurement could ensure. Figure 17 shows the average value of eight hardness measurements with standard deviation marked in the drawing. In the as-extruded condition, variant W1 was again characterised by the highest hardness. After heat treatment to the T4 condition, the value of hardness increased for each variant and the increase was parallel to the increasing Cu content in the alloy. For variant W3, the recorded increase in hardness was 36%.

Fig. 17
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Hardness on the cross-section of I-beam in F and T4 tempers

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Static tensile tests were the next step in the research. The values of yield strength (Rp0.2), tensile strength (Rm) and elongation (A) were determined by the recorded data. The obtained results are summarised in the diagrams in Fig. 18, including the effect of heat treatment.

Fig. 18
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Comparison of mechanical properties determined by static tensile test depending on the heat treatment state of the material

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The results of static tensile tests confirm the results of hardness tests and for the samples in as-extruded condition (F), the highest strength properties are definitely obtained in sample W1, whilst for samples W2 and W3, they are lower and similar to each other. The applied heat treatments to the T1 and T4 conditions produced an increase in all determined parameters, namely, an increase in yield strength, tensile strength and elongation. These results confirm the relationship, already well established in other numerous research works, that the higher the copper content in the alloy is, the higher are the mechanical properties obtained [4, 6, 7, 11]. Higher Cu content releases more of the “θ” phase, which strengthens the alloy. The most intense growth was observed in the alloy with composition W3 and, amongst the samples after heat treatment, this alloy had the highest strength properties. The T1 heat treatment caused a slightly lower increase in strength properties than T4, meaning 13 MPa. However, the elongation results show that the heat treatment to T1 conditions gives slightly higher values than T4 for W1 and W3.

4 Summary

Analysing the obtained test results in terms of the scientific goal adopted in the research, it can be stated that variant W1 of the tested alloy with 2.5% Cu content complies with the requirements set at the beginning of the research. This alloy, being on the borderline between the 6xxx series alloys (because of the Si and Mg content) and 2xxx series alloys (because of the Cu and Mn content), achieves a strength of over 400 MPa and satisfactory hot workability in processes such as, e.g. extrusion. It requires higher extrusion force than the 6xxx series alloys, but the extruded profiles, even with sharp edges, are of high quality. Properly selected extrusion parameters (temperature, extrusion ratio, stock length) and design of the technological process enabled successful manufacture of profiles adopted for studies.

5 Conclusion

  1. 1.

    Based on the surface of the extruded profiles of the three alloy variants, on which no defects such as overheating, cracks, tears or blisters were found, it was assessed that the selected alloys have a good ductility for the hot extrusion process.

  2. 2.

    The analysis of the EBSD structure and subgrain size measurements made on the extruded profiles have not revealed any significant differences between the alloy variants with 2.5% Cu (W1) and 3.5% Cu (W2). The exception is the W2 variant in the T1 temper resulting from the greater effect of grain growth. For variant W3—4.5% Cu, the structure is finer with the average subgrain size smaller by about 20% (3.6 ÷ 3.9 μm).

  3. 3.

    The results of microhardness and Brinell hardness show that variant W1 without heat treatment (temper F) obtained the highest hardness. The material with the lowest content of alloying additives is easier to supersaturate and harden without rapid cooling. For the remaining variants, the microhardness was comparable, with only greater dispersion of results. Brinell hardness measurements, for all heat-treated variants (T4 condition), indicate that hardness of the alloy after heat treatment increases with the increasing Cu content.

  4. 4.

    The mechanical properties show a dependence similar to hardness: for the extruded temper F, variant W1 has the highest properties, and for variants heat treated (T1 and T4 temper), the mechanical properties increase with the increasing Cu content.

  5. 5.

    In case of the extruded profiles, variant W1 of the alloy with 2.5 wt.% Cu content meets the strength requirements over 400 MPa, it achieves strength of over 450 MPa in the T4 condition. This variant meets the present goal and satisfies the requirements set at the beginning of the research. For full use and implementation of the new alloy, additional tests of its corrosion resistance and fatigue strength are required.