The Growth of Quaternary Divorced Eutectic in Al–Ce–Si–Mg Alloys

The pattern formation through advancing solidification front in a variety of organic, metallic, or ceramic materials is of special interest in metallurgy for eutectic systems, being the essential components of cast alloys. The eutectics crystallize into well-ordered patterns, often of high complexity and their control through the alloy chemistry and crystallization conditions is utilized to create the intricate structures across multiple length scales.^[1] For lightweight aluminum alloys, the pattern control through converting the Al–X binary systems towards ternary or higher-order Al–X₁…X_n eutectics with fine phases of higher hardness was identified as the effective strategy of the strengthening improvement.^[2] This concept was recently applied to the Al–Al₁₁Ce₃ eutectic in Al–Ce alloys, known of very high coarsening resistance but insufficient strength^[3]; an addition of 3 wt pct Si to the Al5Ce0.5Mg (wt pct) base increased the room temperature yield stress almost three times, from 47 to 135 MPa at an expense of ductility through the formation of ternary Al–Ce–Si and quaternary Al–Ce–Si–Mg eutectics.^[4] The latter one is of special interest since it was not anticipated based on existing data^[5] and our FactSage computational thermodynamic calculations. In addition to ultrafine morphology, a unique feature of the new quaternary eutectic is its divorced nature. The objective of this study was to assess the growth of the quaternary eutectic in Al–Ce–Si alloys with minor additions of Mg by a combination of in-situ thermal analysis and ex-situ microstructural characterizations. To determine the role of alloying elements, eutectics in related ternary Al–Ce–Mg and binary Al–Ce systems were examined as a reference.

The Al5Ce3Si0.5Mg, Al5Ce0.5Mg, and Al5Ce (wt pct) alloys with compositions listed in Table I were cast using elemental ingredients of 99.9 wt pct purity in a resistance furnace coupled with a steel crucible coated with a mica wash under argon cover gas. For plates with dimensions of 305 × 152 × 25 mm, the average cooling rate from casting temperature of 730 °C to 500 °C, as measured in the plate center, was 4.5 °C/s. The solidification characteristics of alloys were assessed using the Computer-Aided Cooling Curves Thermal Analysis (CCTA) associated with measurements by a Universal Metallurgical Simulator and Analyzer (UMSA).^[6] The solidified structures were examined on polished sections using optical and scanning electron microscopy, SEM Nova NanoSEM650, and on thin foils using transmission electron microscopy, TEM FEI’s Tecnai Osiris equipped with X-FEG gun at 200 keV. The energy-dispersive spectrometry EDS with ESPRIT software was used for TEM-STEM mapping the elemental distribution. The X-ray diffraction was measured using PANalytical X’Pert Pro X-ray diffractometer, equipped with copper-sealed tube source radiation, λ_CuKα = 1.5418 Å. The computational thermodynamic calculations of phase composition under non-equilibrium (Scheil) cooling conditions were conducted using the FactSage (The Integrated Thermodynamic Databank System) with the FTlite database.

The experimental Al5Ce3Si0.5Mg (wt pct) alloy of hypoeutectic nature in regard to both the Ce and Si contents was designed based on the commercial A356 (Al–7Si–0.3Mg, wt pct) cast high-temperature grade by substituting a larger portion of the silicon content with cerium. A thermodynamic description of the Al–Ce–Si–Mg system requires database for six binary systems with appropriate ternary configurations; although some compositions of the Al–Ce–Si–Mg system were a subject of thermodynamic modeling, the results published cannot directly be applied to the alloy of this study.^[7−10] According to FactSage Scheil’s solidification calculation, the Al5Ce3Si0.5Mg alloy should contain 93.4 vol pct of aluminum solid solution with 0.51 at. pct Si and 0.3 at. pct Mg. A content of 0.0019 at. pct Ce indicates practically a lack of solid-state solubility of Ce in Al. The minority phases predicted include 5.59 vol pct of AlCeSi₂ along with 0.72 vol pct Mg₂Si and 0.24 vol pct of Si. Although FactSage is a reliable tool in understanding the solidification paths, as proven especially for many lower order systems, experimental verifications of the Al5Ce3Si0.5Mg alloy revealed differences. It should be emphasized that the Scheil method does not take into account either the effect of solute back-diffusion or the effect of dendritic coarsening in the final stage. Therefore, the Scheil approach is generally only acceptable up to the solid fraction (f_s) of about 0.9 but beyond f_s = 0.9, i.e., in the last stage of solidification, the Scheil predictions are increasingly off the mark. There are models that consider both the solute back-diffusion and dendrite coarsening but their application to the solidification of multicomponent alloys is not straightforward.

Solidification characteristics of Al5Ce, Al5Ce0.5Mg, and Al5Ce3Si0.5Mg alloys, expressed through solid fraction vs temperature plots, obtained from CCTA, allowed determining the influence of individual alloying elements (Figure 1). For the binary Al–5Ce alloy, the proeutectic Al solidification that started at liquidus temperature of 651 °C continued until solid fraction of 52 pct at 644 °C, where the isothermal eutectic reaction took place. The last 4 pct liquid solidified under non-equilibrium conditions until temperature decreased to 618 °C. For the Al5Ce0.5Mg composition, nucleation of the proeutectic Al started at lower liquidus temperature of 648 °C and continued until the eutectic temperature of 638 °C when the solid fraction reached 54 pct. For this alloy, the last 10 pct of liquid solidified under non-equilibrium conditions until 612 °C. An addition of 3 wt pct Si led to essential changes in solidification characteristics. The first derivative dT/dt (T—temperature, t—time) plotted as a function of temperature in the inset of Figure 1 clarifies the location of thermal events, marked then on solid fraction vs temperature plot. At the liquidus temperature of 630 °C, the primary Al initiated the solidification that was followed by the univariant eutectic reaction L → α(Al) + AlCeSi₂ + L that started at 618 °C. It is understood that univariant growth takes place along the tie-line combining the binary eutectic point with ternary eutectic point, thus along this line, the liquid remains in equilibrium between two solid phases. The univariant reaction was followed by crystallization of mostly cored compounds with AlCe₂Si core and Al₂CeSi₂ shell. Some bulky compounds, however, exhibited the Al₂CeSi₂ composition only, most likely due to local melt inhomogeneities. A similar type of silicides in Al–Ce–Si system with varying Ce content, where Ce(Si_1−xAl_x)₂ were surrounded by AlCeSi₂ observed in,^[11] were suggested to be formed by a peritectic (transition)-type reaction. Since phases formed at this stage contained Ce but did not contain Mg, the moving solidification front led to depriving the molten alloy in Ce but enriching in Mg. It is understood that changes of melt composition led to continuous reduction in the otherwise constant eutectic temperature, similarly as it is seen during solidification of the primary Al. As the final solidification event, the quaternary eutectic transformation L → αAl + Si + Al₂MgSi/Al₂Mg₂Si + Al_78.9Ce_5.9Si_12.7Mg_2.6 took place isothermally at 539 °C with the last 1 pct portion of the alloy continuing freezing under non-equilibrium conditions until 535 °C.

An influence of alloy chemical composition on the eutectic morphology is explained in Figure 2. For the Al5Ce alloy, the binary Al–Al₁₁Ce₃ eutectic is clearly lamellar (Figure 2(a)) with further details described earlier.^[12] Additions of 0.5 pct Mg did not change the eutectic phase, which was still Al₁₁Ce₃ with all Mg being present in αAl solid solution. Instead, Mg modified the eutectic morphology towards a mixture of lamellae and Chinese script features; the structure became faceted, suggesting an interface boundary between two solids that has changed in nature (Figure 2(b)). For the Al5Ce3Si0.5Mg alloy, the microstructure became more complex. In addition to lamellae type ternary two-phase eutectic αAl(Ce, Si) + AlCeSi₂ that was generated during the univariant eutectic reaction and mostly cored AlCe₂Si/Al₂CeSi₂ compounds that follow, islands of ultrafine quaternary eutectic were present, as marked in Figure 2(c). The quaternary eutectic that solidified as last liquid portion of the entire alloy volume is located within interdendritic regions, mainly at triple junctions and its volume fraction, according to thermal analysis, reached 7 pct (Figure 2(d)). As seen under higher magnification, the eutectic is composed of near-spheroidal phases, exhibiting evident contrast differences during optical microscopy imaging (Figure 2(e)). Individual phases within the quaternary eutectic can also be distinguished under a backscattered electron contrast in SEM, where high-magnification imaging revealed an additional phase, which is smaller by an order of magnitude than remaining phases and it is seen as bright spheroids in Figure 2(f). As emphasized in the inset of Figure 2(f), the quaternary eutectic is adjacent to bulky compounds, marked there as BC. This feature, detailed through TEM imaging in Figure 2(g), allowed to determine the solidification sequence; as pointed there, solidification of the fine quaternary eutectic followed the solidification of bulky compounds. This observation confirms conclusions reached through thermal analysis that the conditions for the quaternary eutectic transformation to proceed, in terms of the melt composition, were reached after prior solidification of the bulky compounds was completed. Moreover, the TEM imaging confirmed a presence of the ultrafine phase within the eutectic, although the artifacts caused by the thin foil thickness distorted determination of its precise distribution in regard to other eutectic phases (Figure 2(h)). It should be noted that the TEM image contrast did not clearly reveal the geometry of interfaces between the eutectic phases.

Due to low volume fraction of quaternary eutectic phases, X-ray diffraction had limited capabilities in their identification and therefore STEM-EDS microchemical analysis was employed for this purpose. As portrayed in Figures 3(a) and (b), the bulk phase corresponds to Al₂CeSi₂ with particles of Si (with some Al and Mg dissolved) and Mg₂Si. Within the quaternary eutectic, four phases were identified: αAl(Ce,Si), Si, Al₂MgSi/Al₂Mg₂Si, and Al_78.9Si_12.7Mg_2.6Ce_5.9 (Figures 3(c) and (d)). No literature equivalent was found for the Ce-containing latter phase, which formed clusters of very fine compounds with a diameter below 50 nm. The remaining eutectic phases were of the order of 500 nm in diameter.

To determine the role of Mg minor additions in the quaternary eutectic transformation, its content in Al solid solution was analyzed. First, in the Al5Ce0.5Mg alloy, in the absence of Si, Al–Al₁₁Ce₃ binary eutectic was formed, the same one as in the Al5Ce alloy. Then, after addition of 3 wt pct Si, the same minor amount of 0.5 wt pct Mg led to the formation of two eutectic phases. The EDS microchemical measurements indicated the reduced content of Mg in Al solid solutions after Si additions from roughly 1.2 to 0.8 wt pct. Since EDS results at such low Mg contents may involve an error, X-ray measurement of Al lattice constant was used to confirm the trend detected by EDS. The interplanar spacings d_hkl for selected planes and lattice constants a of Al, calculated from X-ray diffraction data, are listed in Table II. As summarized in Figure 4, a presence of 0.5 wt pct Mg in the Al5Ce0.5Mg alloy led to an increase in lattice constant as compared to pure Al. However, a presence of Si in the Al5Ce3Si0.5Mg alloy caused a reduction of the lattice constant to the level below that of pure Al. This means that Si limited the Mg content dissolved in Al, thus made it available for the quaternary eutectic phases. An influence of minor alloying additions on eutectic transformations is well documented for binary systems where small amounts of a ternary element can modify the growth mode of binary eutectic or lead to the formation of a ternary eutectic, with three different phases. The latter effect was shown, e.g., for Ag–Cu eutectic and Sb addition.^[13] It should be noted that the reduced Mg content in Al solid solution alone was not sufficient to generate new phases; Mg did not influence the ternary eutectic reaction in this study. The quaternary reaction was triggered, when the moving solidification front further increased the Mg concentration in the molten alloy. Such accumulation of the modifying element creates necessary conditions of constitutional undercooling and supersaturation for the complex compounds to form, as it was proposed for lower order systems, e.g., Eu additions to Al–Si system led to the formation of clusters of ternary compound Al₂Si₂Eu.^[14]

A major drawback of eutectics in conventionally cast Al–Ce-based alloys is their coarse morphology.^[12] For other eutectic systems, there are reports of refining effects of Mg and Ce. For example, a synergistic effect of Ce and Mg in Al–7Si–0.3Mg–0.2Fe (wt pct) alloy was described, where the combined addition of 0.3 wt pct Mg and 0.1 wt pct Ce refined α-Al, Fe-intermetallics, and modified the eutectic Si from plate-like to fibrous morphology.^[15] In another example, in the Al–10Si (wt pct) alloy, additions up to 2 wt pct Mg refined the primary αAl dendrites, modified Si to coral-like shape and reduced the size of Mg₂Si.^[16] This study shows that in the Al5Ce3Si0.5Mg alloy, the extraordinary refining effect applied mainly to the quaternary eutectic. The ternary eutectic remained generally coarse with intermetallic lamellae of the same order of magnitude as seen in the binary Al–Al₁₁Ce₃ eutectic of Al5Ce and Al5Ce0.5Mg alloys.

For binary eutectic systems, the scaling law, analytically developed by Hillert^[17] and Jackson and Hunt,^[18] allows to determine the microstructure of regular binary eutectics during directional, steady-state solidification. In contrast, the growth of ternary or higher-order eutectics, being of importance in commercial alloys, is more complex, their structure cannot be described by a simple geometrical approach and the analogous theory for multicomponent alloys still does not exist.^[19] Their experimental assessment is, therefore, paramount to fill the present knowledge gap and to provide validation data for numerical modeling. The experimental evidence is even more important in the case of the unique nature of the quaternary eutectic in this study, termed as divorced, in contrast to normal and anomalous growths observed for other systems. The term divorced eutectic implies that during growth the liquid and all four phases do not have triple points and phases do not have direct contact but are separated (divorced) through the liquid.^[20] In contrast to coupled growth, showed schematically for the ternary two-phase eutectic in Figure 5(a), where growth takes place by exchanging the solute atoms between two solid phases, in divorced eutectic the phases grow independently but in a successive manner and as a result there is no connection or mutual relationship between them. This is schematically depicted in Figure 5(b) for the quaternary eutectic of this study α(Al) + Si + Al₂MgSi/Al₂Mg₂Si + Al_78.9Ce_5.9Si_12.7Mg_2.6. It should be emphasized that the schematics show an idealized directional solidification front with a two-phase system and plane isotherms what will differ when a sample involves a curvature of the solid–liquid interface that must play a role in the morphology observed. The divorced eutectic shows no coupling; in fact, the two phases attempt to minimize their area of contact and to form separate crystals. When considering the growth mechanism of divorced eutectic the question arises on how the eutectic phases nucleate. According to Reference 21 in the Al5Mg2Si0.4Mn0.7Fe (wt pct) alloy, the heterogeneous nucleation of the eutectic leading phase during solidification controlled the eutectic growth with examples of binary eutectic of (Al₁₅(Fe,Mn)₃Si₂ + α-Al) nucleated and grew from the primary Al₁₅(Fe,Mn)₃Si₂ phase while ternary eutectic of (FIMCs + Mg₂Si + α-Al) nucleated on the surrounding primary α-Al (FIMCs—Fe containing intermetallic compounds). During the study of divorced eutectic growth L → αAl + α(AlMnSi) + L in ternary Al1Mn3Si (wt pct) alloys under microgravity of the Spacehab mission STS-95, the eutectic silicide phase nucleated on TiB₂ at the advancing solid/liquid interface of αAl.^[20] It is claimed that the symbiotic growth of primary αAl and α(AlMnSi) particles took place with simultaneous pushing of the particles, while solute was exchanged across the liquid gap that existed between the particle and the αAl interface. Although detailed verification is still required, it is likely that the mechanism, depicted in the inset of Figure 5(b), was also valid during growth of the quaternary divorced eutectic of this study.

This research shows that despite its minor content in the Al5Ce3Si0.5Mg (wt pct) alloy, magnesium became the major ingredient of quaternary eutectic phases, thus directly contributing to the quaternary reaction. A reduction of magnesium content in aluminum solid solution due to the addition of silicon along with an enrichment of molten alloy in magnesium, caused by moving solidification front of univariant two-phase coupled ternary eutectic L → α(Al) + AlCeSi₂ + L, followed by crystallization of mostly cored AlCe₂Si/Al₂CeSi₂ compounds, both led to the formation of magnesium-containing eutectic phases. The invariant transformation at 539 °C, L → α(Al) + Si + Al₂MgSi/Al₂Mg₂Si + Al_78.9Ce_5.9Si_12.7Mg_2.6, generated a fine morphology of near-spheroidal phases with 50–500 nm in diameter, located individually without the ordered arrangement, being characteristic for the non-coupled growth of divorced nature. As revealed by preliminary experiments, although the quaternary eutectic represents only 7 pct of the Al5Ce3Si0.5Mg alloy volume, due to its melting temperature and strategic distribution within interdendritic spaces, it effectively controls the alloy thermal stability.