**1. Introduction**

Currently, light materials with a given set of mechanical and physical properties are extremely in demand for manufacturing engineering parts [1–3]. Specifically, high-load pistons operated at up to 300 ◦C require not only specific mechanical properties (e.g., strength, hardness and ductility) but also low Coefficient of Thermal Expansion (CTE) and high thermal conductivity. Among numerous materials, hypereutectic Al-Si alloys (e.g., A390.0 and FM180 alloys), have been established due to acceptable performance in the aforementioned properties [1,4,5]. In reality, they are natural metal matrix composites (NMMC) consisting of hard (Si) crystals distributed throughout a eutectic matrix. However, they show some pivotal drawbacks such as brittleness and the need for modification operation for refining primary and eutectic (Si) phase [1,5].

In the present study, alternative Al-Ca alloys are proposed which were recently reported with excellent processability through shape casting and metal forming [6–10]. When considering hypoeutectic Al-Ca alloys, they have low density, an appropriate combination

**Citation:** Naumova, E.; Doroshenko, V.; Barykin, M.; Sviridova, T.; Lyasnikova, A.; Shurkin, P. Hypereutectic Al-Ca-Mn-(Ni) Alloys as Natural Eutectic Composites. *Metals* **2021**, *11*, 890. https://doi.org/ 10.3390/met11060890

Academic Editor: Tilmann Beck

Received: 27 April 2021 Accepted: 26 May 2021 Published: 29 May 2021

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of mechanical properties and corrosion resistance [6,8]. Meanwhile, their phase composition is complicated and available thermodynamic data are relatively poor. Specifically, there is a lack in the ternary Al-Ca-X phase diagrams not to mention the quaternary systems Al-Ca-X-Y. In our previous studies, many ternary Ca-rich compounds were found including those previously unknown [7,9,11,12]. Their primary crystallisation depending on the composition matters for the design of the hypereutectic alloys. Initially, we find it expedient to develop the multicomponent phase diagrams in the regions for hypereutectic alloys is of profound scientific importance for designing novel materials.

It should be noted that a remarkably fine as-cast eutectic structure in the Al-Ca alloys can be achieved without any refining agents. Moreover, multicomponent eutectic can appear to be even thinner in comparison to a binary [(Al) + Al4Ca] eutectic. The work [13] reported Al-Ca-Mn-Fe alloys with a set of superfine multiphase eutectics (ternary and quaternary) that provides a high-tech performance in casting and metal forming. It should be noted that the best metal forming processability was shown for Al-Ca-Mn alloys [14]. Moreover, Mn acts both as a eutectic-forming element and a solid-solution strengthening agen<sup>t</sup> [1,15]. As for Al-Ca-Ni alloys, an appearance of the ternary Al9NiCa phase together with Al4Ca and Al3Ni phases was found in equilibrium with (Al) solid solution [9]. Being a part of the eutectic mixture [(Al) + Al4Ca + Al9NiCa], the ternary compound has a submicron size and demonstrates a response to spheroidising annealing. Primary Al9NiCa crystals exhibit a polyhedron shape that is similar to the primary (Si) phase in the hypereutectic Al-Si alloys [4]. When considering a detailed investigation of the primary crystals' shape in the multicomponent Al-Ca alloys, there has been virtually no relevant research so far. Within the range of various compounds, some may primarily solidify acquiring a compact shape without compromising the ductility but increasing both strength and the CTE value. It is known that hypereutectic Al-Si alloys are manufactured by either squeeze casting or isothermal hot stamping or hot extrusion [4]. However, massive (Si) primary crystals deteriorate the plasticity of alloys. For this reason, special approaches are being developed for refining primary crystals [1,4,5]. For ensuring processability of hypereutectic alloys through deformation, small size (50–70 μm) and uniformity in distribution of primary crystals must be maintained as well as a globular shape of eutecticorigin intermetallics as illustrated in Figure 1. Based on the preliminary results, we reasonably find hypereutectic multicomponent Al-Ca alloys very appropriate for designing NMMC materials with special physical and mechanical properties competitive to those of the Al-Si piston alloys.

**Figure 1.** Schematic representation of the target NMMC microstructure: (**a**) as-cast condition; (**b**) after spheroidising annealing.

The aim of this work is to explore promising hypereutectic Al-Ca alloys and study their structure and properties with Mn and Ni alloying.

#### **2. Materials and Methods**

The experimental alloys (Table 1) were prepared in an electric resistance furnace LAC (LAC, s.r.o., Židlochovice, Czech Republic) using pure aluminium (99.99% Al), calcium (99.99% Ca), silicon (99.99% Si) and Al—20% Ni and Al—20% Mn master alloys. For the sake of comparison, a commercial AlSi18Cu1Mg1Ni1 (FM180) piston alloy was prepared. The molten alloys were poured at 730–750 ◦C into a graphite mould (without preheating) with an inner size of 15 mm × 30 mm × 180 mm. The cooling rate under that conditions was about 10 K/s.


**Table 1.** Chemical composition of the experimental alloy.

Heat treatment of ingots was carried out at 500 ◦C in a muffle electric furnace SNOL 8.2/1100 (Umega Group, AB, Utena, Lithuania) with a temperature deviation of about 3 ◦C. This procedure is necessary for increasing the plasticity of ingots by fragmentation and spheroidisation of eutectic intermetallics. Hot rolling of ingots was carried out at a speed of 0.2 m/s on a reversible laboratory duo-mill with a maximum rolling width of 250 mm. An initial ingot's thickness was 15 mm. We conducted rolling with a 10% reduction for each pass and an average total reduction ratio of 80%.

The microstructure was examined by scanning electron microscopy (SEM, TESCAN VEGA3, Tescan Orsay Holding, Brno, Czech Republic) equipped with an electron microprobe analysis (EMPA, Oxford Instruments plc, Abingdon, UK), and the Aztec software (Oxford Instruments plc, Abingdon, UK). Thermo-Calc software (TTAL5 database) was used to calculate the phase composition of the alloys [16]. Actual phase transformation temperatures were determined using a DSC 404 F1 Pegasus differential scanning calorimeter (Netzsch Group, Selb, Germany). X-ray analysis was carried out on a Bede D1 System (Bruker, Karlsruhe, Germany) with Cu K α radiation (λ = 0.15406 nm) and treated by software package [17]. Electrical conductivity (EC) was determined using the eddy current method with a VE-26NP eddy structures instrument (CJSC Research institute of introscopy SPEKTR, Moscow, Russia). A high purity Al was employed for comparison purposes. Vickers hardness (HV) was determined using a Wilson Wolpert 930 N setup (Wilson Instruments, Instron Company, Norwood, MA, USA) at 50 N load and 10 s dwell time. Room temperature tensile tests for the sheet samples of 150 mm × 12 mm in dimensions were carried out using a Zwick/Z 250 setup (Zwick Roell AG, Ulm, Germany).

#### **3. Results and Discussion**

### *3.1. Al-Ca-Mn Alloys*

Liquidus projection of the Al-Ca-Mn system (Figure 2) was calculated for substantiating the choice of the experimental hypereutectic alloys solidified under different phase fields. Experimental investigations were next conducted for revision of the isothermal boundaries. Since Al4Ca primary crystals were reported with coarse platelet shape [7], most of the selected compositions encountering crystallisation of the primary Al6Mn phase. Approximately 30–40 wt.% intermetallics in the structure are expected to provide essential physical and mechanical properties.

**Figure 2.** The predicted view of the Al-Ca-Mn phase diagram in the Al-rich region (solid line— calculation, dotted line—experiment).

Al-3Ca-2Mn, Al-6Ca-1.5Mn, Al-8Ca-0.5Mn and Al-10Ca-1.5Mn alloys are the closest to the phase-field boundaries. Microstructural observations and EMPA analysis of these alloys showed that Al-3Ca-2Mn alloy (Figure 3a) almost totally includes the colonies of the binary eutectic [(Al) + Al4Ca]. Some insufficient Al6Mn particles (including some calcium as solute) are also observed in the structure, while according to the equilibrium conditions their amount should have been far larger. A very fine eutectic structure in the Al-6Ca-1.5Mn alloy is decorated by equiaxed crystals of the ternary compounds (Figure 3b), namely Al11CaMn2, as the latter was determined according to EMPA. It should be noted that this ternary compound is not counted in the Thermo-Calc calculation but its existence indicates one more phase field in the phase diagram. Furthermore, the Al-6Ca-1.5Mn alloy displays some primary crystals of the Al6Mn phase. The structure of the Al-8Ca-0.5Mn alloy (Figure 3c) includes eutectic [(Al) + Al4Ca] and the primary (Al) dendrites of a small fraction. The Al-10Ca-1.5Mn alloy exhibits primary Al4Ca crystals and a minor amount of the Al11CaMn2 phase. When considering the Al-10Ca-2.5Mn alloy (Figure 3e), primary Al4Ca and Al11CaMn2 phases co-exist in equal fractions due to apparent phasefield boundary composition. Ultimately, primary Al4Ca crystals and [(Al) + Al4Ca] fine binary eutectic are observed in the structure of the Al-10Ca alloy (Figure 3f).

**Figure 3.** SEM images showing the microstructure of the experimental Al-Ca-Mn alloys in the as-cast condition: (**a**) Al-3Ca-2Mn; (**b**) Al-6Ca-1.5Mn; (**c**) Al-8Ca-0.5Mn; (**d**) Al-10Ca-1.5Mn; (**e**) Al-10Ca-2.5Mn; (**f**) Al-10Ca.

According to the aforementioned investigations, we can reasonably propose a shift of isothermal lines and extension in the (Al) existence region as it is shown in Figure 2 (dotted line). Thus, an Al-rich angle liquidus projection has been predicted. In this regard, two nonvariant reactions including peritectic L + Al6Mn→(Al) + Al10Mn2Ca (point P) and eutectic L→(Al) + Al4Ca + Al10Mn2Ca (point E) are expected.

The Al-6Ca-3Mn and Al-8Ca-2Mn alloys appeared to have the most favourable NMMC microstructure. Their solidification starts with the formation of the Al10CaMn2 phase (Figure 1a). In the Al-8Ca-2Mn alloy, a few compact angular crystals of Al6(Mn, Fe) appeared likely due to the presence of Fe impurity. According to the calculation and structural observations, these alloys manufactured under graphite mould gravity casting (cooling rate 10 K/s) exhibit approximately 40% intermetallics (Table 2). They consist of fine eutectic (containing approximately 8%Ca and 1%Mn) and the Al10CaMn2 primary crystals ranged between 10 and 20 μm in linear size (Figure 4a,b). The FM180 alloy counterpart prepared under the same conditions contains less than 20% (Si) crystals of more than 40–50 μm (Figure 4c). As the eutectic structure of the Al-6Ca-3Mn and Al-8Ca-2Mn alloys is sufficiently thinner as compared to the FM180 alloy, spheroidising annealing may be utilised for improving ductility [1,10]. Some physical and mechanical properties are presented in Table 2.

**Figure 4.** SEM images showing the microstructure of the experimental as-cast alloys showing a natural composite structure and FM180 alloy: (**a**) Al-6Ca-3Mn; (**b**) Al-8Ca-2Mn; (**c**) FM180.

**Table 2.** Calculated phase transformation temperatures, hardness and electrical conductivity of the as-cast Al-Ca-Mn alloys.


\* hereinafter electrical conductivity is designated as Ω.

Instead of the Al11CaMn2 compound roughly estimated by EMPA, the Al10Mn2Ca compound with a precise atomic composition was determined by X-ray analysis. The latter was utilised for the Al-6Ca-3Mn alloy solidified in furnace ambience as it shows a far larger fraction (12.4 vol.%) of the large ternary phase in the structure (Figure 5a). While the X-ray database does not contain the foregoing phase, the latter is known for the related systems CaCr2Al10 and YMn2Al10 [18,19]. A crystal lattice belongs to the tP52/2 space group and has the parameters a = 1.2845 nm, c = 0.5134 nm (Figure 5b).

**Figure 5.** SEM images showing the Al10Mn2Ca crystals appeared in the Al-6Ca-3Mn alloy after furnace cooling (**a**) and X-ray analysis data identified the Al10Mn2Ca compound (**b**).
