**3. Results**

## *3.1. Grain Refinement*

The results of the grain structure investigation showed very large dendritic grains, the sizes of which were several millimeters (Figure 1a), in the Al0.3Mg1Si alloy without the addition of any transition elements. This is generally characteristic of low-alloyed alloys [35]. Upon the addition of Sc (Figure 1b) and Zr (Figure 1c), the size of the grain structure continues to decrease but not significantly. The scandium concentration is 0.6% which is required to achieve a hypereutectic composition and the formation of primary Al3Sc particle in the liquid [21]. Therefore, in Al0.3Mg1Si0.3Sc alloy grain refinement occurs due to supercooling at the boundaries of the nuclei and the liquid phase [36]. The concentration of zirconium Al0.3Mg1Si0.15Zr reaches a peretectic point of 0.11% [36]. However, significant grain refinement is not observed. It is explained by the fact, that according to [36] at the given concentration, the primary particles of Al3Zr only begin to form in the liquid and a significant grain-refinement effect is observed at a concentration of 0.2%.The concentration of 0.3 Sc and 0.15 Zr is sufficient for the hypereutectic composition and the formation of primary Al3Sc particles in the liquid phase, which leads to a significant appearance of a fine-grained equiaxed structure during casting (Figure 1d).

**Figure 1.** Grain structure after casting (**a**) Al0.3Mg1Si, (**b**) Al0.3Mg1Si0.3Sc, (**c**) Al0.3Mg1Si0.15Zr, (**d**) Al0.3Mg1Si0.3Sc0.15Zr.

A dendritic structure is observed in the Al0.5Mg1.3Si alloy (Figure 2a), as well as in Al0.3Mg1Si. However, the grain size in Al0.5Mg1.3Si is two times smaller. This effect is caused by an increase in the content of Mg and Si, which also leads to supercooling between the nuclei and the liquid phase. Additions of scandium (Figure 2b) or zirconium (Figure 2c) alone only slightly affect the grain size. Their combined use leads to the appearance of fine equiaxed grains which, however, only partially replace the dendritic structure (Figure 2d). The decrease in alloying efficiency will be explained later.

**Figure 2.** Grain structure after casting (**a**) Al0.5Mg1.3Si (**b**) Al0.5Mg1.3Si0.3Sc, (**c**) Al0.5Mg1.3Si0.3Zr, (**d**) Al0.5Mg1.3Si0.3Sc0.15Zr.

Thus, a gradual grain refinement occurs during complex alloying with small additions of zirconium and scandium. The decrease in the grain size is explained, first, by the fact that the growing content of scandium and zirconium causes supercooling between the liquid and the surface of the nuclei formed during solidification thereby facilitating the passing of the latter [36]. This also explains the fact that the grain in Al0.5Mg1.3Si (Figure 3a) is finer than in Al0.3Mg1Si (Figure 3b) after solidification, since the total number of alloying elements increases. With a further increase in the Sc and Zr contents, an almost 10-fold decrease in grain size occurs, mainly because the eutectic composition for Al-Sc-Zr system is reached [37]. As already mentioned, due to the high concentration of Zr and Sc in the Al0.5Mg1.3Si alloys, the grain refinement is more effective than in Al0.3Mg1Si. An explanation of this fact will be given later. Note, however, that no appreciable refinement occurs with an increase in Zr or Sc alone. This is due to the fact that a scandium content of 0.6% is required for the hypereutectic composition and significant grain refinement. At the same time, an increase in Zr content to 0.15% reduces the scandium content which is necessary for the hypereutectic composition to 0.2%. It should be noted that the amount of zirconium required for the overflow reaction of the zirconium content should be about

0.26%. Therefore, it is not efficient for a significant refinement of the grain structure during casting in the concentration range studied here.

**Figure 3.** Grain size at complex alloying with small Sc and Zr additives (**a**) in alloys based on Al0.3Mg1Si, (**b**) on AL0.5Mg1.3Si.

### *3.2. Microhardness Behavior*

The results of microhardness measurements are very interesting. It is generally higher (Figure 4a) in an alloy with a higher silicon content than in an alloy with a low silicon content (Figure 4b). This can be explained by the fact that alloy Al0.3Mg1Si contains more particles of the type (AlSi)3ScZr and/or AlSc2Si2 than Al0.5Mg1.3Si. They can appear due to the fact that the cooling process during casting requires a certain time, which is sufficient for the precipitation of a certain amount of these particles from the supersaturated solid solution [26]. An increased ratio of magnesium and silicon will accelerate the precipitation of these particles from the supersaturated solid solution and thereby increase the microhardness. At the same time, an increase in the zirconium content will reduce the ability to form them upon cooling, since this element reduces the diffusion of scandium [19]. As for the Al0.5Mg1.3Si alloy (Figure 4b), it can be assumed that grain refinement, associated with an increase in the proportion of zirconium and scandium, will have a stronger effect than a decrease in the diffusion of the latter. As mentioned earlier, due to the larger amount of Mg and Si, the refinement of the grain structure will be much stronger. Therefore, the length of grain boundaries is longer. The latter, in turn, serve as an additional source in the formation of secondary metastable particles, contributing to an increase in their number [38]. Note that the effect of the grain structure size can be excluded, since this indicator begins to exert a significant effect on the microhardness when reaching submicron sizes. Even in alloys with the finest grains its size is several microns.

#### *3.3. Phase Diagram Calculation Result*

Figure 5a,b shows polythermal sections of the systems Al-0.3%Mg-1%Si-0.15%Zr- (0–0.4)%Sc (wt. %) and Al-0.5%Mg-1.3%Si-0.15%Zr-(0–0.4)%Sc. Note that alloys with a combined content of scandium and zirconium seem to be the most promising. Firstly, they allow to refine the grain structure which will positively affect the mechanical properties and, secondly, slow down the diffusion of scandium and foster the formation of metastable nanoparticles (AlSi)3ScZr coherent to aluminum matrix, instead of harmful AlSc2Si2. Therefore, polythermal sections with the highest zirconium content are considered. On the whole, phase transformations in both groups of alloys (with different Mg/Si ratios) have common laws. Slightly below 700 ◦C, primary Al3Zr particles begin to precipitate from the liquid. Three phases L + Al3Zr + (Al) are simultaneously present in a narrow range between 653–650 ◦C in Al0.3Mg1Si (Figure 5a) and 650–647 ◦C in Al0.5Mg1.3Si (Figure 6b). After that, Al3Zr dissolves and the ZrSi phase appears instead. With Sc content of up to 0.23%

in Al0.3Mg1Si (Figure 5a) and up to 0.26% in Al0.5Mg1.3Si (Figure 5b), the ZrSi phase is transformed into Zr2Si. At zero scandium content, upon reaching a temperature of 601 ◦C in Al0.3Mg1Si (Figure 5a) and 578 ◦C in Al0.5Mg1.3Si (Figure 5b), the metal finally solidifies. After that an aluminum solid solution and the ZrSi2 phase can be observed in the metal. With an increased concentration of scandium, the solidus temperature gradually decreases up to 0.14% and 0.07% of Sc for Al0.3Mg1Si and Al0.5Mg1.3Si alloys respectively and at higher concentration of Sc the solidus temperatures is slightly increased. When the Sc concentration reaches 0.21% in Al0.3Mg1Si (Figure 5a) and 0.25% Al0.5Mg1.3Si (Figure 5b), L + (Al) + ZrSi2 first transforms into L + (Al) + ZrSi2 + Al3Sc and then finally is solidified as (Al) + ZrSi2 + Al3Sc. When scandium reaches 0.24% in the Al0.3Mg1Si alloys (Figure 6a) and 0.27% in Al0.5Mg1.3Si (Figure 6b), solidification occurs according to the following scheme: L + (Al) + ZrSi2 → L + (Al) + ZrSi2 + Al3Sc → L + (Al) + ZrSi2 + SiSc → (Al) + ZrSi2 + Al3Sc. Upon reaching Sc 0.24% and 0.28% in the alloys Al0.3Mg1Si (Figure 5a) and Al0.5Mg1.3Si (Figure 5b), solidification occurs along an even more complicated path: L + (Al) + ZrSi → L + (Al) + ZrSi + Al3Sc → L + (Al) + ZrSi2 + Al3Sc → L + (Al) + ZrSi2 + SiSc → (Al) + ZrSi2 + SiSc. In addition, in a short interval between 0.23 ÷ 0.24% for Al0.3Mg1Si and 0.27 ÷ 0.28%, solidification occurs according to the following scheme: L + (Al) + ZrSi → L + (Al) + ZrSi2 + Al3Sc → L + (Al) + ZrSi2 + SiSc → (Al) + ZrSi2 + SiSc.

**Figure 4.** Microhardness with complex alloying with small additives Sc and Zr (**a**) in alloys based on Al0.3Mg1Si, (**b**) on Al0.5Mg1.3Si.

It should be noted that the appearance of primary particles in the liquid phase will facilitate grain refinement but not all of them will be equally effective. Primary Al3Zr, and especially Al3Sc, are effective inoculators due to the good correspondence between them and the aluminum matrix. There is no such information about ZrSi2 and SiSc particles. However, it can be assumed that their effectiveness is rather low.

Note that in the Al0.3Mg1Si0.3Sc0.15Zr alloy the combined grain refinement effect of zirconium and scandium does not differ from that observed in pure aluminum and alloys with a high magnesium content [37]. However, in the Al0.5Mg1.3Si0.3Sc0.15Zr alloy at concentrations sufficient to obtain a fine grain structure in pure aluminum and alloys with high magnesium content, a large number of dendritic grains are still observed. A first explanation for this is that Al3Sc and Al3Zr particles are not present at all stages of solidification. They will have a slightly weaker effect than in high-magnesium alloys. It should also be noted that Al3Sc particles, which contribute to efficient grain refinement, act even more efficiently in Al0.3Mg1Si alloys because their formation requires a lower scandium content than in Al0.5Mg1.3Si. This explains the fact that the grain is larger in this group of alloys with a low content of Zr and Sc than in Al0.5Mg1.3Si with a similar content of these elements. However, the grain in Al0.3Mg1Si is refined better with an increase in their concentration. At a low concentration of Sc and Zr, the reason for the finer grain in the

case of Al0.5Mg1.3Si is the aforementioned supercooling at the interface between the nuclei and the liquid phase (arising from the dissolved Mg and Si). However, due to the lower Sc concentration required for the formation of Al3Sc particles during metal solidification, the refinement of the grain structure at high scandium concentrations is more efficient in Al0.3Mg1Si.

**Figure 5.** Polythermal sections (**a**) Al-0.3%Mg-1%Si-0.15%Zr-(0–0.4)% Sc (wt. %) and (**b**) Al-0.5%Mg-1.3%Si-0.15%Zr-(0– 0.4)%Sc (wt. %) of the Al–Mg–Si–Zr–Sc phase diagram calculated using Thermo-Calc software.

**Figure 6.** Al0.3Mg1Si0.3Sc main types of intermetallic compounds (**a**); small dispersoids (**b**); large particles (**c**); result of the energy-dispersive spectroscopy (EDS) investigation of chemical composition of the intermetallic compounds (**d**).

After solidification in both types of alloy, a similar chain of phase transformations is observed. At a low concentration of scandium, the following phase transformations will occur: (Al) + ZrSi2 → (Al) + ZrSi2 + (Si) → (Al) + ZrSi2 + (Si) + Mg2Si. At a higher

concentration of scandium, the following transformations will be observed: (Al) + ZrSi2 + SiSc, (Al) + ZrSi2 + SiSc + (Si); (Al) + ZrSi2 + SiSc + (Si) + Mg2Si. It may seem that the Al0.5Mg1.3Si system is more promising from the point of view of further use, since it has a larger Mg/Si relationship than Al0.3Mg1Si, which decreases the risk of the appearance of harmful particles of the AlSc2Si2-type. However, a lower dissolution temperature of Mg2Si particles in Al0.3Mg1 Si alloys allows one to avoid additional heat treatment. To date, multistage heat treatment has been required to obtain two hardening phases (AlSi)3ZrSc and β" (Mg5Si6) [26]. It consists in homogenization in order to maximize dissolution of Sc, Mg, Si and Zr into a supersaturated solid solution which includes eliminating large non-equilibrium particles. Three-stage annealing is applied after homogenization: The first step is the formation of (AlSi)3Sc particles. Then (AlSi)3ScZr particles are formed on their basis. The third stage consists of the dissolution of phases close to Mg2Si, which at the fourth stage can be isolated in the form of a strengthening β" (Mg5Si6) phase. However, at a low temperature of Mg2Si solubility, the fourth stage of annealing can be avoided, since all the necessary processes will occur at the third stage. Note that, according to the calculations performed in the current Thermo-Calc version, only equilibrium primary particles Al3Zr are present in the alloys studied. The appearance of Al3Zr as solid state equilibrium particle which has D023-structure was not predicted by this calculation. This, however, does not exclude the emerging of metastable particles which have L12 structure and composition close to (AlSi)3Zr.

#### *3.4. Investigation of the Intermetallic Participles by Scanning Electron Microscopy (SEM) Method*

The Al0.3Mg1Si0.3Sc alloy contains two types of particle; the first one can be described as being like Al3Fe2Si2. The second can be interpreted as being like quaternary π phase Fe2Mg7Si10Al18 [39] which in this alloy can form due to there being no equilibrium solidification. Primary intermetallic compounds have both close to round (smaller) and needle-like shapes. Their average size is 12.4 microns. This alloy has an increased amount of fine particles which, among other things, explains its high microhardness (Figure 4a). Considering that no traces of scandium were found in large intermetallic particles with a fraction of its content in the solid solution of only 0.11%, it can be assumed with a high degree of probability that it is in the form of fine particles, the fraction of which is 0.14%. It should be noted that there is a high proportion of magnesium in the supersaturated solid solution. Therefore, it is insufficient for the formation of large intermetallic particles of the Mg2Si type. This chemical composition, as predicted in [25], is extremely unfavorable for the production of strengthening Al3Sc particles, since scandium actively interacts with silicon already during cooling after casting. In this case, Mg is used rather rarely. This means, firstly, the need for a prolonged homogenization (completed by quenching) to dissolve excess Sc and Si at high temperatures. In addition, with a high degree of probability the subsequent heat treatment and aging at 300 ◦C will not be accompanied by the precipitation of Mg2Si-type particles. As a result, silicon will interact only with scandium, which will favor the predominant growth of unwanted AlSc2Si2 particles.

Figure 7 shows the results of electron microscopy for the alloy Al0.3Mg1Si0.3Sc0.15Zr. It contains a number of large primary intermetallic such as Al9Fe2Si2, (Al, Si)3Sc, Al3Sc0.6Zr0.4 with an average size of 11.6 microns. It also contains particles with a composition close to quaternary π phase Fe2Mg7Si10Al18. In this alloy, a fairly large amount of Si and Sc is in solid solution. This suggests that, for a given chemical composition, Si and Sc do not interact so intensely with each other. The presence of particles of both the (Al, Si)3Sc and Al3Sc0.6Zr0.4 types indicates that Sc, upon further heat treatment, will react both with silicon, forming AlSc2Si2 particles, and with zirconium, forming Al3Sc0.6Zr0.4 particles. In this case magnesium also "absorbs" part of the silicon, contributing to the formation of Mg2Si-type particles. Taking into account the results of Thermo-Calc calculations (Figure 5a), it can also be assumed that various Zr and Si compounds will occur during heat treatment. Thus, this alloy seems to be very promising for further study, since excessive silicon will interact both with magnesium and zirconium, as well as various impurities. With proper heat treatment,

this will allow the required amount of Al3ScZr particles which have a L12 structure to be obtained.

**Figure 7.** Al0.3Mg1Si0.3Sc0.15Zr main types of intermetallic compounds (**a**); small dispersoids (**b**); large particles (**c**); result of the EDS investigation of chemical composition of the intermetallic compounds (**d**).

The Al0.5Mg1.3Si0.3Sc alloy contains the main types of large primary non-equilibrium intermetallic compounds Al9Fe2Si2, Fe2Mg7Si10Al18, and (Al, Si)3Sc. Their average size is 10.6 microns. In this case, there are spherical and needle-shaped particles, as well as large elongated eutectic particles arising along the boundaries of dendritic grains. Besides, there are many dispersoids in this alloy occupying 0.2% of the cross-sectional area, but with a rather large size.

In the Al0.5Mg1.3Si0.3Sc0.15Zr alloy, large intermetallic compounds are represented by Fe2Mg7Si10Al18 and (Al, Si)3(Sc) particles with an average diameter of 12.8 microns. There are not so many dispersoids in the alloy, however, they have a size of about 30 nm. Smaller dispersoids with an average size of 20–15 nm were detected using transmission microscopy. Considering also the fact that zirconium can practically not be observed in a supersaturated solid solution and in large intermetallic particles, it can be assumed that the main type of fine particles will be Al3ScZr. Moreover, in this alloy Si reacts sufficiently with magnesium which makes it very promising for further research.

The results of electron microscopy are presented in:


**Figure 8.** Al0.5Mg1.3Si0.3Sc main types of intermetallic compounds (**a**); small dispersoids (**b**); large particles (**c**); result of the EDS investigation of chemical composition of the intermetallic compounds (**d**).

**Figure 9.** Al0.5Mg1.3Si0.3Sc0.15Zr main types of intermetallic compounds (**a**); small dispersoids (**b**); large particles (**c**); result the EDS investigation of chemical composition of the intermetallic compounds (**d**).


**Table 2.** Size and amount of dispersoids in the investigated alloys.

**Table 3.** Size and amount of intermetallic compounds in the investigated alloys.


**Table 4.** Chemical composition of intermetallic particles in the studied alloys.


It must be mentioned that most of the phases listed in Tables 2–4 are non-equilibrium, as result they cannot be predicted by the diagram. The reason for this is the non-equilibrium conditions of solidification and the inevitable presence of iron even in the high frequency batch. However, the diagram on Figure 6 can be very useful for predicting phase transformations after homogenization annealing (for example, combined with quenching) when the composition is close to equilibrium.

#### *3.5. Investigation of the Despersoid by Transmission Electron Microscopy (TEM) Method*

A sufficiently large number of Al3Sc particles is observed in the alloy Al0.3Mg1Si0.3Sc. The size of these particles is 20 nm on average (Figure 10a,b). Particles of a given size are at the coherence loss thresholds. Therefore, some of them are completely coherent (Figure 10d), while others only partially (Figure 10d). The presented data are in good agreemen<sup>t</sup> with the data of scanning microscopy, according to which most of the scandium is not in a supersaturated solid solution, and at the same time, it is not observed in large particles of eutectic origin, such as Al0.3Mg1Si0.3Sc0.15Zr. In addition, it is the particles of this type that serve as the main reason for the growth of microhardness (see Figure 4a). It should

also be noted that these particles contain unwanted silicon. However, considering that these particles are fully or partially coherent and considering the increase in microhardness, it is possible with a high degree of probability that they are metastable (AlSi)3Sc, rather than equilibrium harmful AlSc2Si2. The appearance of such particles during solidification in the 6XXX series alloys with small additions of scandium and zirconium as a result of discontinuous precipitation has already been described in many investigations [40,41]. However, it has been described mainly in alloys that do not have such a high excess of silicon. The presence of metastable coherent and semi-coherent particles which precipitate during the cooling process after casting of the Al0.3Mg1Si0.3Sc alloy indicates that they can also be obtained as a result of multistage heat treatment. Thus, the cooling rate achieved in a steel chill mold without zirconium is insufficient so that (AlSi)3Sc-type particles are not precipitated during cooling. It should be noted that in spite the short cooling time of the ingot, a large number of (AlSi)3ScZr dispersoids precipitate in this and other alloys investigated by TEM. This can be explained by the fact that silicon accelerates the kinetics of the precipitation of (AlSi)3ScZr [42].

A different picture of the distribution and chemical composition of fine particles is observed in the Al0.5Mg1.3Si0.3Sc alloy. It also contains the (AlSi)3Sc-type particles but their number is somewhat less than in the Al0.3Mg1Si0.3Sc alloy (Figure 11a). However, in general, they have the same size and the same chemical composition. In addition, needle-shaped particles similar to β" (Mg5Si6) appear (Figure 11b,c).

It should be noted that in spite of the rather short cooling time of the ingot, a large number of Al3ScZr particles precipitate in this (and other alloys investigated by TEM). This can be explained by the fact that silicon accelerates the kinetics of the precipitation of these particles.

Note that emerging of β" actually became possible due to the fact that additional magnesium appeared for the reaction with silicon. Thus, an increase in the Mg/Si ratio will contribute to obtaining an additional synergistic effect between the (AlSi)3Sc and β" (Mg5Si6) particles. In general, the cooling rates after casting in a steel chill mold are also insufficient in this alloy to contain the main elements in a solid solution, and they form hardening phases (AlSi)3Sc and β" (Mg5Si6).

The Al0.3Mg1Si0.3Sc0.15Zr alloy contains a fairly large number of fine coherent or semicoherent (AlSi)3ScZr particles with the L12-structure (Figure 12b–d). At the same time, there is a second type of elongated particles (Figure 12a,c), which can be confused with β" (Mg5Si6), given the rather high content of magnesium as well as silicon in the surrounding solid solution. However, these particles are much longer than β" (Mg5Si6) (from 0.5 to 1 μm). Therefore, a chemical analysis (Figure 12e) shows that they are also discontinuous precipitation of (AlSi)3ScZr particles. This fact is also confirmed by the comparison with the literature data [43], in which discontinuous Al3Sc precipitates are practically identical in shape, size and location of relatively high-angle boundaries compared to the dispersoids observed in this work. It should be noted that, as in Ref. [44], high-angle boundaries serve as a source of discontinuous precipitation of the (AlSi)3ScZr-type dispersoids. When these particles are precipitated, energy is released and the high-angle boundaries begin to move, leaving behind a large number of the (AlSi)3ScZr-type dispersoids. It is noteworthy that dendritic grains, in which intermittent precipitates are found, do not contain L12-nanoparticles (Figure 12c) and vice versa. A possible explanation for this is that in discontinuous precipitation, as in recrystallization [45], there are orientations more favorable for the motion of high-angle boundaries. In grains with an intermittent release of (AlSi)3ScZr, scandium and zirconium become insufficient for the appearance of more coherent particles which have L12-structure. It should be noted that, in the case of the Al0.3Mg1Si0.3Sc alloy, β" (Mg5Si6) particles can be observed together with (AlSi)3Sc particles, in contrast to the L12 and needle shape (AlSi)3ScZr dispersoids. It can also be noted that for the discontinuous and simultaneous precipitation both needle-shaped and L1-structured dispersoids the concentration of scandium in the Al0.3Mg1Si0.3Sc and Al0.5Mg1.3Si0.3Sc alloys is insufficient.

**Figure 10.** (**a**) (AlSi)3Sc coherent and semi coherent particles, (**b**) EDS profile line scan; (**<sup>c</sup>**,**d**) direct resolution of the crystal lattice with (AlSi)3Sc particles of various sizes.

**Figure 11.** (**a**) Al3Sc coherent and semi coherent particles particles; (**b**) Al3Sc coherent and semi coherent particles + Mg2Si (rod-like) particles (**c**) EDS profile line scan.

**Figure 12.** (**<sup>a</sup>**,**c**,**d**) Discontinuous precipitation of (AlSi)3ScZr needle-shaped particles; (**b**) (AlSi)3ScZr particles with the L12-structure, (**c**) discontinuous precipitation of (AlSi)3ScZr needle-shaped particles near high-angle boundaries. (**e**) EDS profile line scan.

Mainly coherent or semi-coherent (AlSi)3ScZr-particles, which have the L12-structure (Figure 13a,b), are formed during discontinuous precipitation and are observed in the Al0.5Mg1.3Si0.3Sc0.15Zr alloy. The lack of needle-shape precipitates is most likely due to the limitations of TEM associated with the small area observed. The average grain size in this alloy is 260 μm, while the survey area is 4 μm2, at best. For the alloy Al0.3Mg1.3Si0.3Sc0.15Zr, where the grain is finer, it is easier to find grain boundaries around which intermittent precipitates are concentrated. In the case of Al0.3Mg1Si0.3Sc0.15Zr, the β" (Mg5Si6)-phase is not detected in the alloy. Its absence can be explained by the fact that the bulk of magnesium and silicon form primary Mg2Si-particles at this concentration of zirconium and scandium (see Figure 6).

 **Figure 13.** (**<sup>a</sup>**,**b**) (AlSi)3ScZr particles with the L12 structure, (**c**) EDS profile line scan.
