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Der Artikel geht auf die Synthesemethoden für Mg-Al-Ca-Legierungen ein und hebt die Vorteile dieser leichten Werkstoffe für die Luftfahrt- und Automobilindustrie hervor. Es werden die Herausforderungen industrieller Produktionsmethoden wie Guss und die Vorteile von Knetlegierungen diskutiert. Der Schwerpunkt liegt auf dem Einfluss von Al- und Ca-Gehalten auf die intermetallische Phasenbildung und die mechanischen Eigenschaften. Verschiedene Synthesetechniken, darunter Induktionsschmelzen, Diffusionspaare, Flusswachstum und Bridgman-Methoden, werden im Detail untersucht. Die Studie betont auch die Bedeutung hochauflösender Charakterisierungstechniken für das Verständnis der Mikrostruktur und der Eigenschaften dieser Legierungen. Die Ergebnisse bieten wertvolle Einblicke in die Optimierung der Synthese und Verarbeitung von Mg-Al-Ca-Legierungen zur Leistungssteigerung.
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Abstract
Further improving the mechanical and corrosion properties of Mg-based alloys requires the investigation of underlying causalities between synthesis, processing, and microstructure, ideally down to the near atomic scale. However, such investigations with high resolution characterisation techniques require model materials of exceptionally controlled chemical composition and microstructure. In this context, the Mg-Al-Ca system, being attractive for offering significant improvement of the inherent brittleness of Mg alloys, is exceptionally challenging from a metallurgical perspective due to the inclusion of components with high reactivity and vapor pressures. Here, we demonstrate the applicability of various synthesis methods from arc-melting over solution growth to diffusion couples, extending to parameters for thermomechanical processing. Suitable pathways to overcome the specific challenges of the Mg-Al-Ca system are demonstrated, as well as the persistent limitations of the current state-of-the-art metallurgical laboratory technology. Solid solution sheet or wrought materials are ideally produced by induction melting in steel crucibles and casting under elevated pressure, with intermittent annealing steps between hot rolling passes to ensure homogenous chemical composition without excessive grain growth. Composite microstructures strongly depend on the solidification rate, and special care needs to be taken with regards to crucible reactions with increasing Al concentrations. Binary intermetallic compounds can be successfully produced in bulk by a variety of techniques, but the interplay between reactivity and evaporation of the melt when both Ca and Al are present required more complex approaches such as the Bridgman method for the synthesis of ternary compounds.
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1 Introduction
The application of lightweight components in the aeronautic and automotive industry has both environmental and economic benefits.[1] Mg-based alloys are especially attractive due to their low density compared to steels, aluminium, and even some polymer-based materials. The most common industrial production method for structural Mg materials is casting[2] after electric melting (either inductive or conductive) using steel crucibles and Cl-F-based flux protection. While fast and cost effective once the required tooling is set-up, the presence of typical casting artefacts such as inhomogeneous grain size distribution and porosity is often resulting in comparatively low mechanical performance, leading to the application of Al-Zn alloyed compositions or even composite structures optimized for effective processing. Wrought alloys on the other hand,[3] often based on Mn and rare earth element additions, generally offer an improved property profile, and are typically shaped by hot extrusion techniques—sometimes coupled with additional cold-rolling and ageing treatments—into profiles and shaped slabs for further machining operations. However, most of the commercial Mg alloys typically still suffer from limited strength and low ductility inherent to the hexagonal crystal structure of Mg.[4] Recent studies show that Mg-based solid solution alloys with a small amount (typically below 1 wt pct) of Ca and Al exhibit improved formability and ductility due to the activation of \(\langle c + a\rangle\) slip.[5] With increasing the Al and Ca content, different types of intermetallic phases precipitate in the Mg-Al-Ca alloys, resulting in a composite structure.[6] The precipitates change from Mg17Al12 to the C15-CaAl2, C36-Ca(Mg,Al)2 and C14-CaMg2 Laves phases with increasing Ca/Al ratio.[7,8] A continuous Laves phase network reinforces the soft Mg-matrix and improve the creep resistance of the Mg-Al-Ca alloys.[9‐11] However, the knowledge of the impact on mechanical[12‐14] and corrosion properties[15‐17] of the Laves phase is still limited, motivating further investigations based on suitable samples and model materials. However, such fundamental investigations require the application of high resolution characterisation techniques such as transmission electron microscopy, micropillar compression testing and atom probe tomography, which in turn necessitate model materials of exceptionally controlled chemical composition and microstructure, e.g., with respect to the grain size, morphology, and chemical homogeneity, as well as the content and size of oxide inclusions.
The main challenges in the production and processing of Mg-Al-Ca alloys are related to the high reactivity[17] and the high vapor pressures[3,18] especially of Mg and Ca. Although the synthesis of solid solutions and Mg-Al-Ca composites on an industrial scale is feasible,[19,20], it is more challenging to prepare the Mg-Al-Ca alloys for specific scientific purposes, i.e., achieving a desired chemical composition and precisely controlling the microstructure with regards to constitution, grain size and morphology of the intermetallic phases. Induction melting for example under a protective atmosphere with elevated pressure can reduce the evaporation of Mg and Ca, while steel crucibles can be utilised by exploiting the low mutual solubility and difference in melting temperature between Mg and Ca with Fe, respectively. However, this does not hold true when Al is added to the melt.[21] Ceramic crucibles, on the other hand, are typically not suitable as they are severely attacked by the highly reactive Mg and Ca. While a limited lifespan of crucibles due to chemical interactions with Mg-Al-Ca melts is not a direct problem for synthesizing scientific model materials (albeit increasing the cost of sample production substantially), it does lead to undesired contamination of the melt, for example with oxide inclusions, and additional elements affecting microstructure and properties thereafter, for example by the formation of Fe-aluminides. Furthermore, independent of the utilized melting technology and crucible material, phenomena such as grain boundary precipitation and intergranular segregation need to be addressed by subsequent thermomechanical processing for as-cast Mg-based solid solution materials. However, the reduced ductility of composites—induced by the intermetallic compounds which in parallel increase the strength—typically renders such procedures unfeasible.[22]
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The underlying phenomena of plasticity as well as the co-deformation processes with the surrounding Mg-matrix illustrate the motivation for the investigation of binary and ternary[23‐27] intermetallic phases of the Mg-Al-Ca system. Similar basic considerations apply as those for the synthesis of solid solution and composite materials: while induction melting using a steel crucible under Ar atmosphere is applicable to the synthesis of C14-CaMg2 Laves phase, it is not suitable for the synthesis of the C15-CaAl2 Laves phase due to the Al-induced contamination with Fe. On the other hand, the absence of Mg with its low boiling point now allows for considering arc-melting without an expected severe evaporation by its extreme local temperatures. However, the inherent brittleness of the Laves phases[28] needs to be considered as well for all liquid metallurgical techniques, as the stresses from solidification and further cooling might lead to their cracking up to total disintegration of the solidified ingot, even after the melting procedure itself could be performed successfully. Casting into a pre-heated alumina crucible followed by slow cooling for example was reported to reduce thermal stress during cooling.[29] However, alumina crucibles cannot be used in synthesis of the C14-CaMg2 and C15-CaAl2 Laves phases synthesis due to the high reactivity of Ca. It has also been reported that the Mg17Al12 intermetallic phase with a small amount of a second phase can be produced using an electrical resistance melting furnace in a graphite crucible under protective Ar atmosphere, but the potential formation of porosity and Al-carbides needs to be considered.[28] The majority of these challenges, i.e., excessive crucible reactions and unwanted solidification phenomena, might be avoided using solid state synthesis techniques. Powder metallurgy, however, appears not favourable for the specific alloy system here, as the huge surface area of the Mg, Al and Ca powder particles substantially increases the risk of oxide inclusions and also poses a considerable health and safety risk by spontaneous combustion. More interesting are solid state diffusion couples, which are comparatively straightforward to fabricate for binary compounds, but appear more challenging for the most complex intermetallic compound of this system from a metallurgical perspective, namely the C36-Ca(Mg,Al)2 Laves phase, due to its high Mg and Ca contents coupled with a narrow compositional and temperature range.[25,29]
Together with the published overview of respective metallographic preparation methods,[30] we explore the challenges and pitfalls in the corresponding synthesis procedures—ranging from induction melting, diffusion couples, flux-growth and Bridgman methods—of Mg-Al-Ca bulk materials in solid solution, composite and purely intermetallic configuration in the current study.
2 Materials and Methods
The chemical compositions of the investigated materials in conjunction with the deployed synthesis techniques are summarized in Table I. The Mg-Al-Ca alloys were prepared from pure Mg (99.95 wt pct), Ca (98.8 wt pct) and Al (99.999 wt pct). The chemical analyses of the starting materials and the synthesized alloys were performed by inductively coupled plasma optical emission spectroscopy (ICP-OES). The metallographic preparation methods for characterizing the obtained Mg-Al-Ca alloys are explained in details elsewhere.[30] Selected samples were colour-etched with a picric acid-based solution for optical microscopy (OM; Leica DMR, Leica AG) observations. Microstructures of the alloys were characterized using scanning electron microscopy (SEM; Helios Nanolab 600i, FEI Inc.). The compositions of the phases in the alloys were determined by energy dispersive X-ray spectroscopy (EDS; EDAX Inc.). Phase analyses were carried out by electron backscatter diffraction (EBSD; Hikari, EDAX Inc.).
Table I
Summary of the Metallurgical Synthesis Methods for the Mg-Al-Ca Solid Solutions, Composites and Intermetallics
Materials
Composition (Wt Pct)
Targeted Phases
Synthesis
Processing
Solid Solutions
Mg-1Al-0.5Ca
Mg
induction melting (60 kW furnace, Ar atmosphere of 0.8 and 15 bar, steel crucible) and casting (Cu moulds of 30x60 mm internal dimensions and wall thickness of 15 mm)
hot rolling at 450 °C, 5 passes with 10 pct reduction per pass, annealing at 500 °C for 24 h
Mg-1Al-0.05Ca
Mg-1Al-0.1Ca
Mg-1Al-0.2Ca
Mg-2Al-0.05Ca
Mg-2Al-0.1Ca
Mg-2Al-0.2Ca
Composites
Mg-6Al-2Ca
Mg + laves phase
induction melting (16 kW furnace, Ar atmosphere at 0.8 bar, steel crucible) and casting (Cu moulds of 25x65 mm internal dimensions and a wall thickness of 15 mm)
annealing (Ar atmosphere at 500 °C for 48 h, furnace cooling)
Mg-5Al-3Ca
Mg-4Al-4Ca
Intermetallic Compounds
Mg-45Ca
C14-CaMg2
as composite materials
as-cast
Al-42Ca
C15-CaAl2
arc-melting on a water-cooled copper hearth under Ar atmosphere of 800 mbar
annealing (glass tube furnace at 600 °C for 24 h under Ar atmosphere)
Mg-28Al-46Ca
C36-Ca(Mg,Al)2
bridgman apparatus (cylindrical tantalum crucible sealed with 0.6 bar Ar, placed on a water-cooled cold finger heated to 900 °C, lowering with 5 mm/h)
as-cast
Mg-53Al-36Ca
C15-CaAl2
3 Results and Discussion
3.1 Mg-Al-Ca Solid Solutions
The limited amount of Al within the solid solution materials (max. 2 wt pct) led to very low Fe contamination of the melt (< 0.0068 wt pct) despite the use of steel crucibles, which offered the additional benefit of being heated by the inductive field of the furnace coil, facilitating rapid heating and melting of the base materials. The elevated pressure of 15 bar was found to be beneficial for reducing the Mg evaporation by local overheating and simultaneously minimising shrinkage porosity during solidification after casting. As shown in Figure 1(a), the Mg-1Al-0.05 Ca alloy cast under high pressure Ar exhibit only a small amount of pores, which translated into a higher strength as well as ductility compared to materials prepared under conventional lower pressure (0.8 bar) conditions. A detailed analysis on the pore formation, size and dispersion—for example by X-ray tomography—as a function of atmospheric pressure during solidification of Mg alloys would be worth of future investigations. Microstructure and resulting mechanical properties of solid solution materials can be substantially improved by thermomechanical treatments (TMTs) such as hot rolling and subsequent annealing procedures, as they generally minimize porosity remaining in as-cast materials, reduce or at least homogenize the grain size, and reduce segregations by recrystallization and diffusion processes. The corresponding parameters, especially the temperature, has to be chosen based on the Mg-Al-Ca phase diagram.[25,29] On the one hand, the temperature should be as high as possible to facilitate recrystallization and diffusion processes, as well as dissolving any (here unwanted) intermetallic phases, which becomes especially for alloys close to the solubility limit of Mg. On the other hand, too high temperatures can cause grain growth, localized melting of segregated areas, and generally enhance surface oxidation when the TMT is performed in a non-protective atmosphere.
Fig. 1
(a) SEM overview of an Mg-1Al-0.05Ca (wt pct) alloy cast under Ar atmosphere with pressure of 15 bar. (b) Corresponding tensile testing results, illustrating the beneficial effects of limiting porosity by increasing the pressure within the melting and casting furnace. (c) Porosity in fractured tensile samples from material cast under 0.8 bar, causing the observed reduction in ductility. Parts of this Figure are being reproduced from Ref. [37] under license CC-BY-4.0
We found that a temperature of 450 °C ensured an optimum balance for the investigated alloy compositions. As with all other metallic structural materials, the other parameters such as reduction per pass, rolling diameter and speed, as well as surface preparation including the use of lubricants need to be carefully adjusted for the given geometries and targets of the TMT processing. Figure 2 outlines the effect of a final annealing treatment at 500 °C after hot rolling the Mg-2Al-0.2Ca (wt pct) alloy at 450 °C in five passes with 10 pct thickness reduction per pass, with reheating for 10 minutes between the individual passes. After the final pass the alloy was annealed at 450 °C for 15 minutes followed by quenching in water [Figure 2(a)]. The corresponding hot-rolled microstructure [Figure 2(b)] shows well recrystallised grains with an average size of 74 ± 6 μm, but with traces of intragranular segregations and most probably intermetallic precipitates visible in the EDS mappings [Figure 2(c)]. These detrimental phenomena—both from a mechanical as well as a corrosion perspective—could be successfully mitigated by additional annealing at 500 °C for 24 h followed by water quenching [Figures 2(d) through (f)]. However, this subsequent annealing also caused a significantly enlarged grain size. These finding suggest that a combined strategy consisting of (i) initial hot rolling to break up the as-cast microstructure, reducing the diffusion length and recrystallizing coarse dendritic structures followed by (ii) a prolonged homogenization annealing step, and (iii) final hot rolling—maybe even with additional cold-rolling beforehand—to refine the grain size would produce optimal mechanical properties. For that purpose, the initial dimensions of the cast blocks must be of sufficient size to allow for the required rolling reduction, while not getting too thin for both later testing procedures as well as too strong heat loss to the cold forming tools (rolls) during the TMT procedure.
Fig. 2
Thermomechanical processing of a Mg-2Al-0.2Ca (wt pct) alloy. (a) hot rolling sequence and (b) corresponding OM overview image and (c) SEM-EDS map of Al (the areas in red and yellow indicate high concentrations of Al) showing traces of remaining segregations as well as Al-rich precipitates (TEM insert). After additional annealing at 500 °C for 24 h (d) significant grain growth can be observed (e) as well as successful improvement of chemical homogeneity (f). Parts of this Figure are being reproduced from Ref. [37] under license CC-BY-4.0
Similar to the solid solution materials, Mg-Al-Ca composites consisting of a Mg solid solution matrix containing a network of Laves phases were successfully synthesized by induction melting using a steel crucible under protective Ar atmosphere. While Fe-Al intermetallic layers forming at the crucible wall can hinder the diffusion of Fe into the Al-containing melt,[31] the Fe contamination remained negligible (0.002–0.018 wt pct Fe) when the Al content was below 6 wt pct. The types of intermetallic phases formed within the composites can be modified by adjusting the alloy composition.[4‐7,21]. With increasing Ca/Al in the range of 0 to 1, the intermetallic phases change from Mg17Al12 to the C15-CaAl2, C36-Ca(Mg,Al)2 and C14-CaMg2 Laves phases.[5,6] Moreover, the precipitates become more interconnected and form a skeleton structure.[6,7] Representative images of the as-cast Mg-6Al-2Ca, Mg-5Al-3Ca and Mg-4Al-4Ca (wt pct) alloys are shown in Figures 3(a) through (c), respectively. The respective Ca/Al ratios of the Mg-6Al-2Ca, Mg-5Al-3Ca and Mg-4Al-4Ca (wt pct) alloys are 0.21, 0.41 and 0.65. It has been confirmed that their main types of intermetallic phases are C15-CaAl2, C36-Ca(Mg,Al)2 and C14-CaMg2, respectively. After heat treatment at 500 °C for 48 h, the interconnectivity decreases and the Laves phase morphology changes to more isolated and spherical particles.[32] Interestingly, the porosity observed within the solid solution materials (Figure 1) was strongly reduced in the composite materials despite the use of sub-atmospheric pressure during melting and casting. This effect is most probably linked to the closer positioning of the alloy composition to eutectic ridges as the alloying content is increased, which reduces the solidification interval.
Fig. 3
Image of as-cast composite materials (a) Mg-6Al-2Ca, (b) Mg-5Al-3Ca and (c) Mg-4Al-4Ca (wt pct) alloys synthesized by induction melting. The corresponding intermetallic compounds were identified as C15-CaAl2, C36-Ca(Mg,Al)2 and C14-CaMg2 Laves phases, respectively. ImageJ analysis gave the following volume fractions of the intermetallic compounds: Mg-6Al-2Ca 5.5 vol pct, Mg-5Al-3Ca 6.8 vol pct, Mg-4Al-4Ca 8.7 vol pct. This Figure is reproduced from Ref. [37] under license CC-BY-4.0
The Mg-Al-Ca ternary system as well as the Mg-Al, Mg-Ca and Al-Ca binary systems show various types of intermetallic phases. The ternary C36-Ca(Mg,Al)2 Laves phase and the binary C14-CaMg2, C15-CaAl2 and Mg17Al12 intermetallic phases are in focus of the present work. In order to synthesize the C14-CaMg2 Laves phase, Mg-45Ca (wt pct) alloys were successfully produced from pure Mg and pure Ca by induction melting using a steel crucible. In order to avoid increasing Fe-contaminations, the Al-42Ca (wt pct) alloys were prepared by arc-melting on a water-cooled copper hearth to synthesize the C15-CaAl2 Laves phase. After arc-melting, the Al-42Ca (wt pct) alloy was heat treated in a glass tube furnace at 600 °C for 24 h under Ar protection. A Mg-Al diffusion couple was prepared from two blocks of pure Mg and Al, targeting the formation of the Mg17Al12 intermetallic phase. The two 5 × 5 × 5 mm3 blocks of Mg and Al were ground up to #4000 grit using SiC paper and polished to 1 μm using diamond paste and ethanol lubricant. The blocks were subsequently cleaned in an ultrasonic bath in acetone, dried and lightly pressed together using a molybdenum clamp. Thereafter, the clamped blocks were placed in the furnace, evacuated and purged with Ar, and then annealed at 400 °C for 1 week, followed by furnace cooling.
The obtained microstructures of the C14-CaMg2 and C15-CaAl2 Laves phases are shown in Figures 4(a) and (b), respectively. While a high density of pores can be observed in the C14-CaMg2 Laves phase [Figure 4(a)], the C15-CaAl2 Laves phases [Figure 4(b)] exhibits a more homogeneous microstructure and a comparatively coarse grain size. The synthesis of Mg17Al12 intermetallic phase using the diffusion couple technique was not successful. As shown in Figure 4(c), neither a concentration gradient indicating measurable interdiffusion nor the formation of intermetallic phases could be achieved despite numerous trials. This is most probably due to rapid re-oxidation of the polished Mg surface despite only few minutes between the last polishing step and evacuation of the furnace. This result highlights the requirement of materials preparation in an oxygen-free atmosphere (such as a glovebox) and transport of the samples in vacuum containers to the diffusion bonding furnace.
Fig. 4
Synthesis of intermetallic phases. SEM images of the (a) C14-CaMg2 Laves phase synthesized by induction melting showing high number of porosity, (b) C15-CaAl2 Laves phase synthesized by arc-melting, and (c) an unsuccessful attempt at fabricating an Mg-Al diffusion couple after heat treatment at 400 °C for 1 week, showing no sign of interdiffusion or intermetallic phase formation. This Figure is reproduced from Ref. [37] under license CC-BY-4.0
However, several other pathways remain for the synthesis of intermetallic compounds of the Mg-Al-Ca system, especially for the ternary C36-Ca(Mg,Al)2 Laves phase, which combines all three elements and thus also their specific limiting factors of reactivity (Ca), vapor pressure (Mg) and solubility (Al). While it has been reported that the Mg-Al-Ca intermetallic phases can be synthesized by combinatorial sputtering, desorption of volatile film-forming species during sputtering at elevated temperatures demonstrates the challenges in this process routine for the C36-Ca(Mg,Al)2 Laves phase.[33]
We therefore investigated the feasibility of alternative approaches: induction melting of a Mg-30Al-44Ca (wt pct) using a graphite crucible with boron nitride spray proved unfeasible to rapid disintegration of the crucible, while both sintered boron nitride and aluminumtitanite crucibles repeatedly fractured during heating. As it allows for more gentle heating of the base materials and does not require any additional crucible material, manual induction melting on a water-cooled Cu finger within an Ar filled quartz tube was performed (“copper boat” furnace by Edmund Bühler GmbH, 12 kW power). However, Mg evaporation could not be avoided entirely, and the resultant microstructure [Figure 5(a)] predominantly consists of the here undesired C15-CaAl2 Laves phase with small amounts of Mg. Utilizing a Mg-Al-Ca diffusion multiple with a 50 µm thin Ca foil between two polished blocks of pure Mg and Al failed as the oxidation of Ca—even more rapid than that of Mg—precluded any interdiffusion in a similar manner as for the Mg-Al binary diffusion couple. Cold rolling Ca foils in between Al and Mg did not lead to a successful bond due to the vastly different rates of deformation of the three materials involved. Another approach relied on positioning a block of Mg on top of a C15-CaAl2 Laves phase alloy, placing the stack in an Ar filled annealing furnace and bringing it to 700 °C (above the melting point of Mg). Although liquid/solid interdiffusion could be achieved, the desired formation of the C36 Laves phase could not be observed [Figure 5(b)].
Fig. 5
SEM images of the trials for fabricating bulk C36 ternary Laves phase: (a) Mg-30Al-44Ca (wt pct) alloy synthesized by manual induction melting, and (b) reaction zone of a Mg-CaAl2 liquid-solid diffusion couple. This Figure is reproduced from Ref. [37] under license CC-BY-4.0
In view of these unsuccessful attempts, the flux solution growth method[34] was used to synthesize the targeted C36-Ca(Mg,Al)2 Laves phase. Figures 6(a) and (b) show a schematic illustration of the utilized setup. The starting materials of a Mg-28.4Al-43.2Ca (wt pct) alloy were mixed and placed in an alumina crucible with a mesh fixed at a position above the alloy for decanting the melt at the end of crystal growth. The entire crucible was encapsulated in a Quartz tube filled with Ar [Figure 6(c)]. The quartz tube was heated up to 870 °C and annealed for 2 h to melt the starting materials. After melting, the quartz tube was cooled down to 850 °C at 10 oC/h and then slowly cooled down to 750 °C at 1°/h, followed by centrifugation to separate the remaining liquid phase from the solidified material. According to the phase diagram,[25,29,35], at the used composition the C36-Ca(Mg,Al)2 Laves phase is in equilibrium with the liquid phase at 750 °C. Therefore, after removing the liquid phase, the crystal of the C36-Ca(Mg,Al)2 Laves phase should be obtained.
Fig. 6
(a) A schematic illustration of the set-up of the solution growth method. (b) After heat treatment, the primary phase was separated from the remaining liquid phase by centrifugation. (c) An alumina crucible encapsulated in a quartz tube filled with Ar. (d) After cooling the quartz tube was blackened on the inside and showed radial cracks. (e) Only one ingot was obtained in the crucible and no materials were left on the grid. This Figure is reproduced from Ref. [37] under license CC-BY-4.0
However, after cooling the quartz tube was blackened on the inside and showed radial cracking at the position where the crucibles met [Figure 6(d)]. After opening the tube one ingot was found in the crucible and no material was left on the grid [Figure 6(e)]. This suggests that the entire melt solidified above 750 °C, with no liquid left to be centrifuged into the upper crucible. The fact that no liquid phase was present at 750 °C could either hint at inconsistency with the existing phase diagram, be linked to a significant change in melt composition due to evaporation losses, or caused by a temperature offset between the thermocouple and the actual temperature in the crucible. Lastly, the feasibility of synthesizing the C36-Ca(Mg,Al)2 Laves phase by the Bridgman method was investigated. Pure elements were positioned in a cylindrical Ta crucible of 14 mm in diameter and 70 mm in length. The crucible was sealed containing an Ar atmosphere of 0.6 bar, and placed on a water-cooled cold finger in a Bridgman apparatus After temperature equilibration at 900 °C, the growth was carried out by lowering the crucible out of the hot zone at a velocity of 5 mm/h. As shown in Figure 7(a), the obtained alloy contained two parts and a large void of yet unknown origin in the center. The microstructures of the top and the bottom parts of the sample are shown in Figures 7(b) through (d), with average compositions (at. pct) of Mg-45.37Al-38.63Ca and Mg-2.27Al-39.40Ca, respectively. At least three phases are present in both top and bottom regions. EBSD analyses showed that the right top regions consists predominantly of the targeted C36-Ca(Mg,Al)2 phase [Figure 7(b)], but with inclusions of eutectic regions containing the C15-CaAl2 Laves together with Mg solid solution. On the left top, the main region revealed the C15-CaAl2 Laves phase and small eutectic regions [Figure 7(c)]. The bottom is mainly consisting of the C14-CaMg2 phase, with traces of other components of the compositions of Mg-53Al-36Ca (at. pct) as indicated by the white arrow in Figure 7(d). While the resultant size of the C36-Ca(Mg,Al)2 domains in the right top of the ingot are not large enough for bulk mechanical testing, they allow for micromechanical testing such as micropillar compression or nanoindentation testing after f.e. focussed ion beam preparations. While surely representing substantially increased effort compared to other laboratory scale synthesis methods, the Bridgeman technique thereby offers at last the possibility to probe the mechanical properties of the C36 Laves phase, thereby elucidating its contribution to Mg-Al-Ca composite materials. Another potential strategy to overcome the outlined difficulties would be to follow a powder metallurgical route, blending either elemental or pre-alloyed binary powders and sintering them—preferably under isostatic pressure—at elevated temperature in the solid state into a bulk specimen. However, close attention would have to paid to the formation of oxide inclusions potentially originating from the large surface area of the powder particles. Larger samples with surface areas in the order of 100 mm2 to facilitate corrosion studies are as of yet not feasible with current metallurgical laboratory techniques.
Fig. 7
Results from synthesising the C36-Ca(Mg,Al)2 Laves phase using the Bridgman method. (a) overview of the obtained ingot with SEM images of the (b) top right of the sample. EBSD allowed us to identify the grain orientation of the cubic CaAl2 phase (c) as well as the hexagonal CaMg2 phase (d). Parts of this Figure are being reproduced from Ref. [37] under license CC-BY-4.0
Due to the high reactivity and high vapor pressures of Mg and Ca, and the mutual solubility of Al with Fe, the synthesis of bulk materials of Mg-Al-Ca solid solutions, composites and intermetallic phases is a complex metallurgical challenge. Different bulk metallurgical methods were investigated, demonstrating their specific challenges and pitfalls, and the following conclusions can be drawn:
(1)
Solid solutions can be synthesized by induction melting using a steel crucible under Ar atmosphere. Elevated pressure was found to minimize evaporation and porosity within the solidified sample. Intermetallic precipitations and intragranular segregation can be significantly reduced by subsequent hot rolling and homogenization annealing, with careful selection of temperatures and processing sequence to minimise grain growth.
(2)
Composites can also be synthesized by induction melting under Ar atmosphere, ideally using Mg-Ca master alloys instead of pure elements to minimize evaporation and oxidation. When the Al content is low (≤ 6 wt pct), the Fe contamination in the Mg-Al-Ca composites is negligible, and thus steel crucibles can be used. Porosity formation and thus the need for high pressure solidification seems to be less pronounced than for solid solution materials.
(3)
The intermetallic C14-CaMg2 Laves phase can be readily synthesized by induction melting, while the C15-CaAl2 Laves phase can be prepared by arc-melting. The Bridgman method was identified as the most promising pathway for synthesizing of the C36-Ca(Mg,Al)2 Laves phase, albeit in small amounts only suitable for micromechanical testing.
(4)
Solid state diffusion couples require extreme care to avoid oxidation of Mg and especially Ca, which otherwise inhibits interdiffusion and interfacial phase formation.
Acknowledgments
We thank Mr. Jürgen Wichert and Mr. Michael Kulse for their technical supports in alloy synthesis. This work was supported by the German research foundation (DFG) within the Collaborative Research Centre SFB 1394 “Structural and Chemical Atomic Complexity—From Defect Phase Diagrams to Materials Properties” (Project ID 409476157). H. Springer wishes to acknowledge funding through the Heisenberg-program of the Deutsche Forschungsgemeinschaft (Project ID 416498847).
Conflicts of interest
The authors declare no conflict of interest.
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