Introduction

Present day's industries are continuously focusing on improving the properties of the originally used materials. One direction of improving the properties of Al alloys is the application of alloying elements with specific properties. More sophisticated alloys with increased mechanical and physical properties allow, for example in the automotive industry, increased material savings, improved environmental aspects in the form of reduced fuel consumption and increased performance due to the increased durability of Al castings.

Al–Si–Mg–Cu-based alloys are nowadays very often used in various industries, where the automotive and aerospace industries play a priority role. Castings made from these alloys are characterized by high strength characteristics as well as excellent foundry properties and good corrosion resistance.1,2 An important property of Al–Si–Mg–Cu-based alloys is the possibility of increasing the strength characteristics through precipitation hardening by the formation of Al2Cu or Mg2Si precipitates.3,4,5,6 In the automotive industry, there is a great effort to increase the properties of aluminum alloys in order to reduce the weight of vehicles. Weight reduction effectively contributes to fuel economy, emission reduction and also increases vehicle performance.7,8 At the same time, aluminum alloys, even after the application of precipitation hardening, are characterized by a limited working temperature of 150–170 °C.9

One of the possibilities to increase the working temperature range (i.e., increasing tensile strength and hardness of the alloy) is alloying with a transition element—Zr. Ten stable phases can be formed in Al–Zr binary systems such as AlZr, AlZr2, AlZr3, Al2Zr, Al3Zr, or Al2Zr3, Al3Zr2.9 Crystallization of Zr-rich phases occurs during the peritectic reaction at Zr contents above 0.1 wt% with long needle or plate morphologies. The resulting phases may crystallize as the semi-coherent tetragonal equilibrium DO22 and DO23 phases, or in the coherent cubic L12 form. Both of these forms are characterized by the formation of a low-energy interface with the matrix.10 Zr in aluminum alloys causes the presence of a thermally stable dispersoids. Zr particles are resistant to coarsening since its interface energy with the matrix is low, and the diffusivity and solubility of the rate-controlling element are small.11 In general, the addition of Zr to an Al-Si based alloy leads to an increase in strength properties and hardness.12 Even a small addition of Zr increases the hardness of the castings, as a consequence of the precipitation of the Al3Zr phase, which not only induces an increase in the hardness of the alloy, but also increases the wear resistance of the material.12,13 The Al–Zr binary phase diagram states that the precipitation of dispersoids due to heat treatment can be thermodynamically possible only if the concentration of Zr is higher than 0.08 wt% at temperatures approaching 500°C. Precipitation occurs at the cores of dendrites, and if the concentration of Zr in the alloy is higher, it will continue until the concentration is reduced to 0.08 wt%, which is the minimum required limit.12,14

The addition of Zr and Ti to Al–Si–Mg–Cu based alloys forms (Al,Si)2(Zr,Ti) and (Al,Si)3(Zr,Ti) intermetallic phases in the microstructure. The (Al,Si)2(Zr,Ti) phase is characterized by an increased Si content reaching an increased Al content in the (Al,Si)3(Zr,Ti) phase.15,16,17 The aging produces Al3Zr (L12)- or Al3(Zr1-xTix) (L12)-based precipitates in the microstructure of solute-enriched dendrites, reflecting the microsegregation of Zr and Ti atoms during solidification. The initial inhomogeneous saturation of the solute atoms during solidification affects not only the spatial distribution of the precipitates, but also the size and morphology of the precipitate through the nucleation mechanism. The increased Ti (inoculant) content in the alloy also promotes increased crystallization of Zr-rich phases such as (Al,Si)3(Ti,Zr). Ti atoms are able to form clusters which act as nucleation sites.18 The synergistic interaction of Zr and Sr in Al-Si alloys can form Al3(Sr1-x Zrx)-based phases. In general, Sr (as an Si eutectic modifier) is added to the melt to modify the morphology of the eutectic. Higher content of Sr may encourage the formation of SrO or Al2SrO3 oxides. However, several investigations describe that in the presence of Zr, the small size SrO particles can act as nucleation sites for the formation of Zr-rich phases.18,19,20,21,22

This paper deals with the analysis of the influence of Zr and the synergistic effect of Zr with Ti or Zr with Sr on the aluminum alloy AlSi7Mg0.3Cu0.5. Alloying by Zr combined with Ti has already been addressed by several authors,5,10,11,12,13,18 but no clear conclusion have been reached. The alloying of aluminum alloys through addition of Zr and Sr and their synergistic effect has not been comprehensively studied. So the main purpose of this paper is to investigate the effect of Zr, Zr + Ti, Zr + Sr on the changes of microstructure, substructure, microhardness of the formed phases and their effects on the mechanical properties of the synthesized alloys. The appropriate amount of alloying, grain refinement and modifying elements in the alloying process is a key factor to achieve suitable and more sophisticated aluminum alloys.

Experiment Procedures

AlSi7Mg0.3Cu0.5 Alloy

To evaluate the microstructure, substructure and mechanical properties, we used the hypoeutectic AlSi7Mg0.3Cu0.5 aluminum alloy, which is based on the normalized alloy AlSi7Mg0.3. The main use of the AlSi7Mg0.3 alloy is for the production of components for the automotive, aerospace, space or defense industries, where the emphasis is on good casting properties, corrosion resistance, pressure tightness and weldability. The AlSi7Mg0.3Cu0.5 alloy was supplied by the supplier and used in the experimental process as the so-called reference alloy—R. Its chemical composition is shown in Table 1.

Table 1 Chemical Composition of Reference Experimental Alloys (wt%)
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AlSi7Mg0.3Cu0.5 Alloys with Addition of Zr, Ti and Sr

In the first step of the experimental part, a "new" AlSi7Mg0.3Cu0.5 alloy was created with an addition of 0.15 wt% Zr (alloy AlSi7Mg0.3Cu0.5Zr0.15 with designation—N). AlZr15 master alloy was used as an alloying material. The chemical composition of the alloy AlSi7Mg0.3Cu0.5Zr0.15—N is shown in Table 2.

Table 2 Chemical Composition of AlSi7Mg0.3Cu0.5Zr0.15—N Alloys (wt%)
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In the second step of the experimental part, another six material variants of the AlSi7Mg0.3Cu0.5 were created, combining the alloying elements Zr + Ti and Zr + Sr, respectively. The content of alloying elements Ti and Sr was graded into three levels, namely 0.1; 0.2; and 0.3 wt%, while the Zr content was in all cases constant—at the level of 0.15 wt%. Master alloys AlZr15, AlTi5B and AlSr10 were used as alloying material. The chemical composition of the experimental alloys AlSi7Mg0.3Cu0.5 based on the combination of Zr + Ti and Zr + Sr and their designation are given in Table 3. The chemical composition was determined by arc spark spectroscopy (Bunker - Q2 ION).

Table 3 Chemical Composition of AlSi7Mg0.3Cu0.5 Alloys Based on the Combination of Zr + Ti and Zr + Sr (wt%)
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Melt Preparation and Casting Process

The melting and preparation of all experimental alloys were carried out in an electric resistance furnace using a graphite crucible with a protective coating. The total weight of one batch was 8 kg, and the melt was deliberately not degassed. The individual experimental variants were cast in ceramic molds using the investment casting technology (Figure 1). The ceramic mold consisted of a contact, insulating and strengthening layer. The ceramic molds were preheated for 1.5 hours to 750 ± 10 °C. The surface temperature of the ceramic molds prior to casting ranged from 510 to 540 °C, and the casting temperature was set at 750 ± 10 °C. The casting height from the pouring cup was 200 mm, and the casting rate was 0.6 kg.s-1. After casting, the ceramic mold was getting cooled on air for 1 hour. The above casting parameters were constant for all material variants.

Figure 1
figure 1

(a) Ceramic mold with 10 test samples, (b) Scheme of test specimen.

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The microstructures of the samples were evaluated by light optical microscopy on a Nikon Epiphot 200 optical microscope, by scanning electron microscopy (SEM) on an Oxford UltimMax 65 electron microscope. Samples for microstructural analysis were prepared by standard metallographic procedures (coarse and fine wet grinding, polishing on an automatic machine using diamond emulsion, etching). To highlight the Zr-rich phases, the samples were etched with 20 ml H2SO4 and 100 ml of distilled water under light microscopy. Samples intended for SEM and energy dispersive X-ray spectroscopy (EDS) analysis were etched with 0.5% HF solution.

Samples for substructural analysis were prepared by grinding on P320 to P2500 sandpaper. They were prepared in the form of a 200-nm-thick “film”, from which the evaluated sample was cut out in a circular shape with a diameter of ø 3 mm. Before the individual samples were evaluated, they were finalized to a thickness of 50–70 nm using a Gatan Disc Grider. Subsequently, the samples were ion-polished using the Gatan PIPs device. Samples were observed using a Jeol 2200FS transmission electron microscopy (TEM) equipped with an Oxford Instruments EDS analyzer. EDS analysis was performed in scanning transmission electron microscopy (STEM) mode with a local size of 1 nm. As part of the substructural analysis, energy dispersive X-ray spectroscopy (EDS/elemental) mapping and high-angle annular dark-field (HAADF) mode were also used. The microhardness of the structural components was evaluated on a Hahnemann type Mod 32 at a load of 10 p. The device uses the Vickers method, whereby the measuring body was embedded in the front lens of the objective. This is a direct measurement of the created impression. The resulting value represents the average value of ten measurements.

Mechanical tests were carried out on the Inspekt table 50 kN device according to the EN ISO 6892-1 standard. The hardness tester CV 3000LDB was used to evaluate the hardness, under conditions of HBW 5/250/15 (pressing a carbide bead with a diameter of 5 mm under a test load of 2,542 kN for 15 seconds). All reported values of mechanical characteristics are average values from five measurements.

Heat Treatment

From each experimental variant, half of the samples (5 pcs) were heat treated by T6 precipitation hardening. The temperature mode consisted of solution treatment (540 ± 5°C/12 hours), quenching (water 65 ± 2°C), and artificial aging (150 ± 5°C/5 hours).

Results

Microstructure Investigation

The microstructure of the as-cast reference alloy R is characterized by the presence of eutectic Si in the morphology of fine, imperfectly round grains that crystallized between the primary dendrites of the α phase (Figure 2a). The content of Mn in the alloy together with the presence of Fe led to the crystallization of Al15(FeMn)3Si2 intermetallic phases. These Mn- and Fe-based phases occur in the form of primary Chinese script, which can be seen in Figure 2b together with its chemical composition from EDX analysis. Applying heat treatment caused a significant change in the spheroidization of eutectic Si, which is in the form of round grains with minimal coarsening (Figure 2b).

Figure 2
figure 2

Microstructure of reference alloy R. (a) As-cast state (Mag. 100× and 500×); (b) EDS analysis of the as-cast state Fe-rich phase; (c) After heat treatment (Mag. 100× and 500×).

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In the microstructure of alloy N in the as-cast state, intermetallic Zr-rich phases are present due to the addition of 0.15 wt% Zr. The eutectic crystallized in the same morphology as in the reference alloy R. The intermetallic phases rich in Zr are most often observed in the form of long needles, with a smooth surface and split ends. Based on EDS analysis, they can be identified as Al3Zr- or AlSiZr-type phases with a higher content of Ti (Figure 3). Similar results of chemical analysis are reported by Hernandez-Sandoval18, and the recommended formula is (AlSi)3(TiZr).13,23,24,25,26 As a result of heat treatment of alloy N, spheroidization of eutectic Si occurred, similar to the case of reference alloy R. The morphology of Zr phases remained almost unchanged, and a partial change was observed only by the local breakdown of large acicular phases into smaller plate-like structures. A comparison of the effect of the application of heat treatment on alloy N is shown in Figure 4.

Figure 3
figure 3

EDS analysis of (AlSi)3(TiZr) phases present in alloy N, as-cast state.

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Figure 4
figure 4

Microstructure of alloy N (Mag. 350×). (a) As-cast state; (b) After heat treatment.

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By observing the microstructures of group T experimental alloys (with a constant Zr content of 0.15 wt% and a graded amount of Ti), it can be concluded that the addition of Ti to the alloy led to an increase in the nucleation of Zr-rich phases. It can also be observed that with the increase in Ti content from 0.1 wt% (Figure 5a) to 0.3 wt% (Figure 5b), the number of intermetallic phases based on Zr evenly distributed in the microstructure of the alloys also increases. The combination of Zr and Ti led to the crystallization of Zr-rich phases in the morphology of thinner longer needles or in thicker plate-like formations without split ends. As a result of the heat treatment, there was a partial disintegration of the larger Zr phases of the acicular morphology, but the overall characteristics of the Zr-rich shapes remains unchanged (Figure 5c).

Figure 5
figure 5

Microstructure of T alloys. (a) T1 alloy, as-cast state (Mag. 100x and 1000×); (b) T3 alloy, as-cast state (Mag. 100× and 200×); (c) T2 alloy, after heat treatment (Mag. 100× and 350×).

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Figure 6a shows a detailed view of the Zr-rich intermetallic phase in the form of a thicker plate without split ends. EDX analysis points to the presence of Al, Si, Zr, and similarly to alloy N (Figure 3), the Ti content is about 16.3 wt%, so it is a (Al,Si)3(Ti,Zr) phase. From the EDX analysis, it follows that Ti favorably participates in the increased number of nucleation sites as well as in the crystallization of nuclei of Zr-rich phases. The increased content of Ti in intermetallic Zr-rich phases is confirmed by the mapping shown in Figure 6c. The mapping also shows the distribution of Zr in the intermetallic phase and in the aluminum matrix.

Figure 6
figure 6

T2 alloy as-cast state. (a) Detail of (Al,Si)3(Ti,Zr) phase; (b) EDS phase analysis; (c) Mapping.

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In the microstructures of material variant S (constant Zr content of 0.15 wt% and graded amount of Sr), a reduced number of formed Zr-rich phases was observed compared to material variant T. The size and morphology of Zr-rich phases was identical to alloy N (more split-end acicular formations) (Figures 7a and 8). An increase in the Sr content in the alloy did not lead to an increase in the number of nucleation sites of phases rich in Zr. The precipitation of Zr-rich phases was observed especially near the eutectic, which is confirmed by EDX analysis (Figure 8b) and mapping (Figure 8c). EDX analysis of the Zr-rich phase shows a high Si content (56 wt%), approximately the same Ti content (20 wt%) but a lower Zr content (approx. 15 wt%) compared to alloy N and variant T. Change in chemical composition of phases based on Al–Si–Zr–Ti means that these are (Al,Si)2(Zr,Ti) phases with a crystallization temperature of around 660°C.18,20 With the increasing content of Sr in the alloys (S2 and S3), the local elongation of the excluded phases rich in Zr and their transformation into a square form was observed. After the application of heat treatment, a more significant breakdown of phases rich in Zr in the form of needles was observed (Figure 7b— detail in the red rectangle).

Figure 7
figure 7

Microstructure of S alloys. (a) S1 alloy, as-cast state (Mag. 500× and 5000×); (b) S3 alloy, after heat treatment (Mag. 500× and 5000×).

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Figure 8
figure 8

Alloy S2 after heat treatment. (a) Detail of (Al,Si)2(Zr,Ti) phase; (b) EDS phase analysis; (c) Mapping.

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Microhardness Evaluation

It is believed that the individual phases formed can have a significant effect on the hardness and strength of the alloy and can also significantly affect the increase in operating temperatures of the researched alloys. For the above reasons, microhardness measurements of the precipitated phases were performed. Figure 9 shows the measured values of the microhardness of the α phase, eutectic and Zr-rich phases for all material variants after heat treatment. The microhardness of the α phase of the reference alloy R reached 145 HVM5. Similar values of the microhardness of the α phase were also measured for alloys N, S1, S2 and S3. The combination of Zr and Ti (variant T) led to an increase in the microhardness of the α phase, while the highest value was achieved by alloy T3 (159 HVM5). The measured microhardness values of eutectic Si ranged from 201 to 211 HVM5 for all material variants. The exception is alloy N with the lowest microhardness of eutectic Si (191 HVM5). The highest microhardness of the Zr-rich phases was measured in the T variant alloys (260 HVM5—T1 alloy to 266 HVM5—T2 alloy). Alloying by Zr and Sr (variant S) resulted in a decrease in the microhardness of the Zr-rich phases and the lowest value was 238 HVM5—alloy S3).

Figure 9
figure 9

Microhardness values of individual structural components depending on the material variant in the cast state.

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Substructure Investigation

Coarser particles are not observed in the observed substructure of the reference alloy R after heat treatment (Figure 10a). The detailed view (red rectangle in Figure 10a) shows the presence of a structural disorder, probably clusters of dislocations. The substructure of alloy N is characterized by the presence of precipitates with the morphology of thicker rods with lengths of 100 to 700 nm and a width of 20 to 60 nm. Figure 10b (red rectangle) shows a detail of a part of the precipitate with a total length of 700 nm.

Figure 10
figure 10

Substructure after heat treatment. (a) Reference alloy R; (b) Alloy N.

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As can be seen in Figure 11, the substructure of the R and N alloys is formed by an Al matrix with Si structures. The application of EDS mapping showed that in the substructure of the reference alloy R, there are particles of both round and sharp-edged morphology, primarily formed by the elements Si, Fe and Mn (Figure 11a). In detail and using the HAADF mode (Figure 11b—gray image), variously oriented precipitates in rod morphology were observed, mainly formed by the elements Zr, Ti and Si (Figure 11b—mapping).

Figure 11
figure 11

Substructure detail and STEM/HAADF mapping. (a) Reference alloy R; (b) Alloy N.

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From the material variant T, sample T3 with 0.15 wt% Zr and 0.3 wt% Ti was selected for transmission microscopy. In the observed substructure of the T3 alloy after heat treatment, it is possible to observe a large number of perpendicularly oriented particles of acicular morphology or slightly thicker rods with an approximate length of 200 nm (Figure 12). Perpendicularly oriented particles represent θ'—Al2Cu phases, and together with them Si-rich phases in a granular morphology are visible. Figure 12a, in the right-hand part, shows structural disturbances (clusters of dislocations). EDS mapping in Figure 12b also shows the presence of Zr and Ti elements in the matrix.

Figure 12
figure 12

T3 alloy after heat treatment. (a) Substructure, (b) Detail and STEM/HAADF mapping.

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From the material variant S, sample S3 with 0.15 wt% Zr and 0.3 wt% Sr was selected for transmission microscopy. The observed substructure of the S3 alloy after heat treatment is characterized by a lower number of rod formations with a more irregular distribution when compared to the T3 alloy. The precipitates in the form of sticks reached a size in the range of 100–500 nm and a width of 20–80 nm (Figure 13b—gray image). EDS mapping showed that the precipitates present are mostly formed by an increased content of Ti and precipitates based on Cu (θ′—Al2Cu) are present only partially. EDS mapping also shows the distribution of Zr and Sr elements in the matrix (Figure 13b).

Figure 13
figure 13

S3 alloy after heat treatment. (a) Substructure, (b) Detail and STEM/HAADF mapping.

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Point EDS analysis of the chemical composition of the precipitates in the S3 alloy showed a high content of Zr, which was up to 64 wt% at a content of 5 wt% Ti and 30 wt% Al (precipitates marked as spectrum 44 and 46) (Figure 14). In another case, the measured Zr content was 34 wt% Zr together with a low presence of Ti and Cu (spectrum 45).

Figure 14
figure 14

EDS analysis of precipitates in S3 alloy after heat treatment

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Mechanical Properties

Figure 15 shows the course of tensile strength values depending on the experimental material variant. The reference alloy R reached a strength of Rm259 MPa. Similar Rm strength values were also measured for alloys N, T1 and T2. A decrease in strength Rm strength occurred in the T3 alloy (246 MPa), which, together with the S1 and S3 alloys reached the lowest value of Rm strength. Conversely, the highest Rm strength of 270 MPa was measured for the S2 alloy.26

Figure 15
figure 15

Relation between average tensile strength and alloy.

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The reference alloy R reached the lowest yield strength value of Rp0.2 of 198 MPa. Addition of 0.15 wt% Zr in alloy N increased the agreed yield strength value of Rp0.2 to a value of 207 MPa (Figure 16). The agreed yield strength value of Rp0.2 of alloys of material variant T ranged from 232 to 237 MPa. The addition of Zr and Sr (variant S) led to an increase in the agreed yield strength Rp0.2, and the highest value was measured for alloy S2 (264 MPa).26

Figure 16
figure 16

Relation between average yield strength and alloy.

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The addition of Zr and its combination with Ti and Sr had a negative effect on the elongation values A5 (Figure 17). The highest A5 elongation value of 3.2% was measured for the reference alloy R. Alloying by Zr, Ti and Sr resulted in a significant reduction of the A5 elongation. The lowest A5 elongation values of 0.2% were achieved by alloys S1 and S2.26

Figure 17
figure 17

Relation between average elongation and alloy.

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The best hardness results according to Brinell HBW were achieved by material variant T (Figure 18), of which the T3 alloy had the highest value of 102 HBW. Material variant S achieved lower values when compared to variant T. The lowest Brinell HBW hardness values were measured for the reference alloy R, namely 89 HBW.26

Figure 18
figure 18

Relation between average Brinell hardness and alloy.

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Discussion

By adding the AlZr15 master alloy and increasing the Zr content in the N alloy to the level of 0.15 wt%, phases of the (Al,Si)3(Ti,Zr) type were present in the microstructure in acicular or thicker rod morphology. Zr-rich phases are characterized by the smallest diffusion toward the α matrix and therefore crystallize preferentially as Zr-rich phases at a temperature of around 670 °C.18 Similar results are also published in the work by Vočin.29 In the N alloy, in addition to the spheroidization of Si, the partial disintegration of the longer acicular phases rich in Zr into smaller plate-like formations was also observed due to heat treatment.

The increase in the number of Zr-rich phases with the gradual addition of Ti to the alloy in material variants T1, T2 and T3 can be explained by the fact that the increased precipitation of Al–Zr–Ti–Si particles accompanying the addition of Ti (or Si) may be the result of the fact that the given element affects either the interphase energy, the deformation energy of the nucleation process, or the interaction with vacancies and thus form heterogeneous nucleation sites. The strong binding energy between Ti and aluminum vacancies could explain the fact that “quenched” excess vacancies and Ti atoms are able to form clusters that act as nucleation sites for Al–Si–Zr–Ti phases. The simultaneous addition of titanium would tend to increase the formation of Al-, Si-, Zr- and Ti-based phases such as the (Al,Si)3(Ti,Zr) phase. These phases are then present in the structure as metastable phases with a DO2218,28,29 crystal structure. An increase in Ti content also led to an improvement in the compactness of the observed needles or rods (without split ends). As a result of the heat treatment, there was no significant change in the Zr-rich phases in the alloys of material variant T. The Zr-rich phases were mostly observed in the original morphology, which we can attribute to the fact that the Al-Si-Zr-Ti-based phase was formed during solidification and could not dissolve by heat treatment due to its high stability at elevated temperatures.16 Decay or the breaking of longer needles of phases based on Al-Si-Zr-Ti due to heat treatment was observed only minimally when compared to alloy N, which is related to the more compact morphology of the given phases.

The addition of a graded amount of Sr instead of Ti led to a decrease in the number of nucleation sites for the Al–Si–Zr–Ti phases, and their number was comparable to alloy N. Increasing the amount of Sr in the alloys of variant S also led to a change in the chemical composition of phases rich in Zr (predominant phase was (AlSi)2(ZrTi)) in a less compact (split) form. Similar to alloy N, the heat treatment resulted in a partial breakdown of the longer acicular Zr-rich phases into smaller plate-like morphologies.

Due to the addition of 0.15 wt% Zr, a slight increase in Al–Si–Ti–Zr in rod-like morphology precipitates was observed in the N alloy. Al–Si–Ti–Zr type precipitates show low diffusivity in the Al matrix at elevated temperatures, which can largely improve the mechanical properties of Al–Si–Mg–Cu alloys at high temperatures.18 Increasing the Ti content in the material variant T resulted in an increase in the number of perpendicularly oriented precipitates based on Cu θ′—Al2Cu. θ′—Al2Cu precipitates are relatively stable at elevated temperatures,30 as a result of which they may benefit in terms of strength.31 With the last material variant S (combination of Zr and Sr), precipitates with a high Zr content (Al3Zr or Al3(ZrTi)) with an unoriented distribution in the substructure were mostly observed. In alloys of variant S during aging, the decomposition of supersaturated Al–Zr solid solutions occurred by the nucleation of Al3Zr and Al3(ZrTi) precipitates, which are characterized by the metastable cubic structure L12 and thermal stability.32

The agreement of the results of the mechanical characteristics and the microstructure together with the substructure was expected to be manifested only at the agreed yield strength and hardness. The improvement of the agreed yield strength is mainly caused by the formation of precipitates. An increase in the volume fraction of precipitates in the substructure in the alloys of variants T and S, when compared to the reference alloy R and alloy N, led to a decrease in the distance between the precipitates.33,34 Based on the results of the agreed yield strength, we can conclude that the Al3(ZrTi) Zr-based precipitates, which were present in the substructure of the S variant alloys, prevented the dislocations from sliding more effectively, thereby increasing the resistance to further deformation under loading than the Cu-based precipitates (θ' - Al2Cu) present in T variant alloys. Several sources state that the results of the tensile strength and the agreed yield strength are also positively influenced by grain refinement due to the addition of Zr, or as the case may be the combined effect of Zr and Ti.10,33,34,35,36,37 Based on the comparison of the images, it can be concluded that in our case there was no significant refinement of the grains of the α phase due to the addition of Zr and its combination with the elements Ti and Sr, similar to the paper.1 The increase in hardness in T1, T2, T3 and S3 alloys is caused by a higher number of Zr-based phases (Al,Si)2(Ti,Zr) and (Al,Si)3(Ti,Zr), which are characterized by higher hardness.

The significant decrease in elongation due to the addition of Zr and its combination with the elements Ti and Sr can be explained by the presence of a large number of Zr-based phases in the long acicular or plate-like morphologies, which crystallize already at a temperature of 660–670°C, which leads to the impeding of the flow of the melt into the interdendritic spaces and the creation of an easier path to the formation of shrinkage.38 The second explanation is directly related to the acicular morphology of the Zr-rich phases. After the solidification of the aluminum matrix, there is an increase in the total stresses in the system. A critical point of microtear formation or as a stress concentrator is then the sharp ends of needles of Zr-rich phases oriented perpendicular to the load axis.39

Conclusion

The paper investigated the influence of alloying with the transition element Zr and its synergistic effect together with the elements Ti and Sr on the microstructure, substructure and mechanical properties of the AlSi7Mg0.3Cu0.5 alloy. Based on the results, the following results were obtained:

  1. 1.

    The effect of Zr alloying in the microstructure of N—AlSi7Mg0.3Cu0.5Zr0.15 alloy resulted in the formation of Zr and Ti(AlSi)3(TiZr)-based phases with DO22type crystal structure. The (AlSi)3(TiZr) phases crystallized in a less compact needle or thicker rod morphology.

  2. 2.

    The addition of Ti clearly increased the number of nucleation sites of Zr-rich phases in alloys of the T variant. By gradually increasing the Ti content (0.1; 0.2 and 0.3 wt%), the number of (Al,Si)3(Ti,Zr) phases in compact (unsplit) needles and rods also increased.

  3. 3.

    As a result of the addition of graded amounts of Sr, Zr-based phases with a changed chemical composition, namely (Al,Si)2(Zr,Ti) phases in a less compact (split-ends) form, were formed in the microstructure of the S variant alloys.

  4. 4.

    The effect of T6 heat treatment only minimally affected the characteristics of the microstructures of the experimental alloys. There was a local breakdown of Zr-based phases from the morphology of long needles to smaller plate-like formations and spheroidization of eutectic Si.

  5. 5.

    The substructure of the experimental alloys was characterized by the presence of precipitates, while in alloy N, these were strengthening precipitates based on Al–Si–Ti–Zr. In T alloys (combination of Zr and graded Ti content), precipitates based on Cu θ'—Al2Cu clearly prevailed. The substructure of the alloys of variant S (combination of Zr and graded Sr content) preferentially formed precipitates with a high Zr content Al3Zr and Al3(ZrTi) with a metastable cubic structure L12.

  6. 6.

    Zr together with the elements Ti and Sr led to an increase in the yield strength and hardness of all experimental alloys. Tensile strength and elongation in particular were negatively affected to a large extent, and the desired values were not reached for tensile strength in alloys T3, S1 and S3 and ductility in all experimental alloys except the reference alloy.

  7. 7.

    From the evaluated spectrum of experimental alloys, the AlSi7Mg0.3Cu0.5 alloy with 0.15 wt% Zr and 0.2 wt% Ti (designated T2) can be named an alloy with an optimal combination of the results of microstructural, substructural analysis and mechanical properties.

  8. 8.

    The AlSi7Mg0.3Cu0.5 alloy with 0.15 wt% Zr and 0.2 wt% Ti appears to be a possible direction for the development of new, more sophisticated alloys for specific uses. However, it is necessary to take steps to increase the elongation of the castings in question.