Effect of alloying elements on thermal stability of Aluminium-Cerium based alloys

The effect of alloying elements on thermal stability of ve different near eutectic Al-Ce based alloys, i.e., Al-12Ce, Al-12Ce-4Si, Al-12Ce-0.4Mg, Al-12Ce-4Si-0.4Mg and Al-12Ce-4Si-0.4Mg-0.25Sr alloys were investigated. The alloys were heat treated at three different temperatures, i.e., 200 °C, 300 °C and 400 °C to establish the thermal stability. Binary Al-12Ce alloy consisting of α-Al and showed higher resistance to coarsening at all temperatures. Addition of Si decreases the thermal stability above 200 °C while combined addition of Si and Mg increases the thermal stability up to 300 °C. Addition of Sr to quaternary alloy was not benecial to enhance the thermal stability.


Introduction
Aluminium alloys are desirable for automotive and aerospace industries due to their lightweight, outstanding castability and excellent mechanical properties. These alloys can be used as replacement in the automotive sector with currently used iron based alloys and titanium alloys wherever it is suitable. The high strength to weight ratio of aluminium alloys lead to better fuel economy and help reduce greenhouse gases. Besides, aluminium alloys tend to make a robust passivizing protective oxide layer that makes it corrosion resistant. One of the major drawbacks of aluminium alloys is that it does not retain their mechanical properties at higher temperatures. The addition of rare earth elements to aluminium alloys can ll this gap. In this context, recent studies on Al-Ce alloys show improved design e ciency for the transport industry and better mechanical properties for high temperature applications. At the hyper eutectic point, Al-12Ce alloy contains a eutectic mixture of Al and Al 11 Ce 3 intermetallic. This Al 11 Ce 3 present in aluminium matrix is metastable and does not respond to heat treatment. In lanthanide series, earlier lanthanides up to Sm form Al 11 RE 3 (RE refers to rare earth) intermetallic and later lanthanides from Eu form Al 3 RE phase that is stable in aluminium matrix. In Al-RE system, thermodynamically stable Al 3 RE phase precipitates out during solidi cation and no further heat treatment is needed for Al 3 RE precipitation [1]. Al-Ce alloys are castable for a wide range of Ce contents and compatible with available casting infrastructure.
Addition of Mg to Al-Ce binary alloy allows precipitation of intermetallic phase during heat treatment and ensures better properties than previously developed binary Al-Ce alloys [2]. The addition of Si in Al-Ce-Mg alloys has a negative impact on castability of Al-Ce alloys and creates a aky type CeAlSi intermetallic that adversely affects the mechanical properties.
Zachary C. Sims et al. have found that a lower amount of Si addition to Al-Ce based alloys does not degrade the castability of alloys and provides better wear resistance at room temperature [2]. It was found that the addition of Ce up to 12 wt. % into Al increases the mechanical properties of binary system (Figure1) [3]. To develop reasonable room temperature and high temperature (400 °C) strength, other alloying elements (Si, Mg and Cu) were added to binary Al-Ce alloys. Figures 1 and 2 represent the summary of mechanical properties of Al-Ce based alloys available in literature. Previous studies on binary Al-Ce alloys show improved thermal stability at high temperatures for various applications. Therefore, the main objective of the present work is to understand the effect of alloying elements and correlate the microstructure-hardness-thermal stability of Al-Ce based alloys like Al-12Ce, Al-12Ce-4Si, Al-12Ce-0.4Mg, Al-12Ce-4Si-0.4Mg and Al-12Ce-4Si-0.4Mg-0.25Sr (wt. %) alloys.

Materials And Methods
Experimental details

Alloy preparation
Near-eutectic Al-Ce based alloys with composition of Al-12Ce, Al-12Ce-4Si, Al-12Ce-0.4Mg, Al-12Ce-4Si-0.4Mg and Al-12Ce-4Si-0.4Mg-0.25Sr (wt. %) were cast for the present study. Commercial pure aluminium (CPAl) ingot and the following master alloys having purity level 99.7% were used as feedstock: Al-50Si, Al-20Ce, Al-20Mg, and Al-10Sr (wt. %). All the alloys were prepared using resistance type furnace (KANTHAL, SN 1015) and clay bonded graphite crucible. Master alloys of Al-Ce, Al-Si and CPAl together were placed in a crucible and allowed to melt in furnace held at 785 °C. Al-Mg and Al-Sr master alloy were added to the melt at the end with intermittent stirring to ensure chemical homogeneity and minimize oxidation loss. After complete melting, the alloy was degassed with a commercially available degasser, C 2 Cl 6 . The top oxide layer of the melt was skimmed off before pouring the molten alloy into mild steel mould (1.5 cm diameter and 15 cm height), which was preheated at 300 °C in an air circulated oven. After casting, the cylindrical alloy bars were sectioned using the abrasive cutter to prepare a sample of 2-3 mm thickness for metallography, heat treatment and subsequent studies.

Metallography and microstructure characterization
For microstructural analysis, ASTM E3 standard metallographic practice was adopted, where nal polishing was carried out using 0.01 μm colloidal silica.
Keller's reagent was used for chemical etching, whenever required. The FESEM (Field Emission Scanning Electron Microscope) was equipped with a backscattered electron (BSE) and a secondary electron (SE) detector. The attached energy dispersive spectroscopy (EDS) was used for phase speci c chemical analysis with 5 kV and 20 kV acceleration voltages. Phase speci c EDS analysis was carried out on three random locations for each phase and average composition calculated. For heat treatment studies, samples were polished using alumina slurry up to 12-13 μm surface roughness. Image analysis was carried out using Image J software (version 1.53b) and the systematic manual point count method (ASTM E562) was used for calculation of volume fraction of phases. X-ray diffractometer was equipped with a solid-state detector and operated using Cu-K α radiation. Quanti cation of phases and XRD pattern was analysed using Xpert high score (3.0.0) software.

Results And Discussion
3.1 Al-12Ce alloy Binary Al-12Ce (Figure 3a) alloy shows the presence of Al (light grey in the matrix) and intermetallic Al 11 Ce 3 . Al 11 Ce 3 is also referred to as Al 4 Ce by some authors [1].The alloy shows a very ne faceted eutectic mixture of Al and lathe like intermetallic Al 11 Ce 3 . Figure 3b shows the XRD pattern of as cast and selected heat treated alloys. The XRD pattern con rms the presence of Al and Al 11 Ce 3 phases. From the heat treatment studies (Figure 4), few critical points (encircled) were selected based on high hardness variation compared to as cast alloy. The XRD pattern shows that no new phases form on heat treatment up to 100 hours. The calculated lattice parameters for Al (cubic) and Al 11 Ce 3 (orthorhombic) was found to be a=4.07Å and a=4.37Å b=16.63Å c=8.24Å, respectively which are in agreement with reported a=4.39Å b=13.2Å c=10.09Å [5]. Figure 4a and b show heat-treated (300 °C for 10 hours) and as cast microstructure of Al-12Ce alloy respectively. The heat-treated alloy showed coarsening of phases. Heat treatment at 300°C for 10 hours results in a eutectic microstructure that has undergone minor morphological changes. TheAl 11 Ce 3 phase seems to have spheroidized at many regions and become less intertwined as compared to small and entangled laths in as cast alloy. This suggests a localized minimization of micro constituent surface energy at the eutectic through surface diffusion within the intermetallic and accompanying spheroidization, rather than bulk diffusion through the matrix. In as cast condition, volume fraction of Al 11 Ce 3 was found to be 15.5 ± 1.2 % as determined using the systematic manual point count method. Theoretical volume fraction of Al 11 Ce 3 from equilibrium phase diagram was found to be 9 .45 % [6,7]. This disparity in volume fraction is due to the non-uniform distribution of intermetallic particle into the Al matrix. The volume fraction of Al 11 Ce 3 increased from 0.155 to 0.185 upon heat treatment at 300 °C for 10 hours as compared to as-cast alloy. This is in agreement with XRD pattern (Figure 3b). It shows that the ratio of area integrated intensity ( ~13.9 to 13.6) slightly changed for a xed plane (( , ) and 2 interval with heat treatment [8]. This indicates the gradual change in phase fraction of intermetallic with time and temperature. Strong vacancy binding of Al to Ce atoms decreases the degradation of Al 11 Ce 3 intermetallic and therefore reduces vacancy diffusion (the dominant transport mechanism for solute atoms within the matrix) [9,10]. The intermetallic is trapped by the zero solubility of Ce in aluminium matrix. This trapping prevents the system from minimizing surface energy through diffusion which in turn prevents the alloys from coarsening. So, low solubility and large atomic size difference between Ce (1.81 Å) and Al (1.43 Å) result in a low diffusion coe cient when compared to other alloying elements. As an illustration, the diffusion data for Al, Ce and other solutes in Al were calculated using the data available in literature (Table1).  [13]. The increased hardness from as-cast to heat treated condition for 300 °C up to 10 hours can be approximated by Orowan strengthening.
An attempt was made to understand the strengthening mechanism in Al-Ce based alloys. Solid solution strengthening, Hall-Petch hardening and precipitation hardening contribute to signi cant hardening in aluminium alloys. In the presence of precipitates, precipitation hardening can dominate all other hardening mechanisms. Orowan described that precipitation strengthening is a result of precipitate-dislocation interaction in the matrix leading to formation of dislocation loop around the precipitate and increase in yield strength of the material is given by [14] Where f: precipitate volume fraction The parameters for aluminium in Equation 1are as follows: M=3.06 [18], ν=0.345 [18], b=0.286 nm [19] and G=25.4 GPa [19].
approximated as the increment in strength and de ned as the difference in microhardness values of as-received alloy and pure Al [20]. Z. C Sims et al. [21] explained the disparity in strength. The neutron diffraction study showed that the load transfer mechanism played signi cant role in improving the strength.
Orowan strengthening mechanism and load transfer mechanism are expected to be active at higher temperatures, although less e cient than at ambient temperature, as dislocation can climb to bypass Al 11 Ce 3 precipitate and the fast creeping Al matrix transfers less load to Al 11 Ce 3 precipitate. An increase in the mean diameter of Al 11 Ce 3 from 142 ± 26 nm in as cast alloy to 175 ± 21 nm in heat treated alloy at 300 °C for 10 hours was observed (Figure 4a and b).
However, Eric T. Stromme et al. [22] observed the ageless behaviour in Al-Ce alloys. Although there was coarsening after heat treatment at 300 °C for 10 hours (Figure 3d), Orowan strengthening still dominated in the heat-treated alloy due to increase in volume fraction of intermetallic [23]. The hardness values show a peak for all temperatures studied and then stabilized on prolonged heat treatment for up to 100 hours. This demonstrates that both the hardening mechanisms were active at room and elevated. This study shows the thermal stability of intermetallic Al 11 Ce 3, after heat treatment. The thermal stability of Al 11 Ce 3 can also be ascribed to its high melting point above 1200 °C [24] Figures 6a, b). This shows the high thermal stability of intermetallic phase in aluminium [28][29][30]. This result can be justi ed with the higher melting point (more than 1400 °C) of Al 2 Ce [31].
The effect of Si on yield strength appears to be inconsistent and affects the ultimate tensile strength and work hardening [22].Therefore, Si addition in Al-12Ce alloy results in marginal change in hardness as compared to binary Al-12Ce alloy. Figure 6c shows the heat treatment studies for Al-12Ce-4Si alloy. After heat treatment at 200 °C for 10 hours, there is a signi cant increase in hardness. This could be due to the combined effect of dispersion strengthening, solid solution strengthening and load transfer mechanism activated at this temperature. Diffusion data (Table 1) shows that diffusion coe cient of Si at 400 °C is 10 4 times higher than 200 °C. Thus, diffusion time for Si at high temperature is much lower and strengthening due to solid solution could be lowered. Though microstructural changes observed at 200 and 400 °C were not signi cant (Figure 6a and b) but decrease in hardness at 400 °C after 10 hours of aging time was signi cant suggesting poor thermal stability of the alloy.
3.3 Al-12Ce-0.4Mg alloy Figure 7a shows the as cast microstructure of Al-12Ce-0.4Mg alloy containing eutectic mixture of Al and Al 11 Ce 3 in which Al 11 Ce 3 lathes are rmly interlinked.
The XRD pattern also shows Al andAl 11 Ce 3 in as cast alloy. However, upon heat treatment at higher temperatures two new phases, Al 3 Mg 2 and Mg 3 Al, appear ( Figure 7b).This could be due to precipitation of phases in the alloy assisted by higher diffusion coe cient of Mg in Al ( Table 1).
The corresponding microstructures of the alloy are shown in Figure 8a and c. Lathe-like interconnected Al 11 Ce 3 transforms into discrete particles after heat treatment at 400 °C for 10 hours and becomes more globular as compared to heat treatment at 200 °C for 100 hours (Figure 8c).These alloys showed a 23% increase in hardness at 200 °C for 100 hours, while at 400 °C, there is a 10% decrease in hardness. The variation in the hardness can be inferred from solid solution strengthening, Orowan strengthening and load transfer mechanism. In order to nd the contribution of Orowan strengthening microstructural study was conducted on the critical point (Table 3).  (Table 3). This suggests that the strength increase was mostly due to solid solution strengthening and a mechanism of load transfer.
The decrease in hardness at 400 °C after 10 hours of heat treatment can be inferred to loss of solid solution strengthening due to the release of strain energy, inactivation of load transfer mechanism due to fragmentation of Al 11 1Ce 3 lathes and softening of phases low melting point phases like Al 3 Mg 2 (~447 °C) [33]).
It was di cult to quantify the individual contribution of solid solution strengthening and load transfer mechanisms due to the formation of new phases during heat treatment. Increasing the heat treatment time beyond 50 hours shows a rapid decrease in hardness value accompanied by softening of phases 3.4 Al-12Ce-4Si-0.4Mg alloy Addition of 4 wt. % Si and 0.4 wt. % Mg to binary Al-12Ce alloy leads to complete suppression of intermetallic Al 11 Ce 3 due to formation of Ce(Si 1-X Al X ) 2 with x=0.1 to 0.9, which has lower formation energy (-0.585 eV) than Al 11 Ce 3 (-0.349 eV) [25,27]. Figure 9a show ne eutectic mixture of Al 9 Siand Al matrix with dispersed Ce(Al X Si 1-X ) 2 .CeAl 1.2 Si 0.8 phase is tetragonal with lattice parameters of a = b = 4.24, c = 14.538 A° [34].XRD pattern shows that no phase transformation occurs on heat treatment as compared to as-cast condition (Figure 9b).After the heat treatment, based on high hardness variation, some critical points(encircled) were selected for microstructural study. Figure 10a and b show the micrograph of as cast Al-12Ce-4Si-0.4Mg alloy at 200 °C and 400°C respectively after 10 hours of heat treatment. Figure 10c shows the heat treatment analysis of Al-12Ce-4Si-0.4Mg alloy. Heat treatment at high temperature (400 °C) results in fragmentation of Al 9 Si intermetallic lathe and transformation into particle-like morphology. Hardness improves by 33 % at 200 ° C for 10 hours of aging time compared to as-cast condition (Figure 10c). This increase in hardness was associated with the combined effect of Si and Mg. The decrease in hardness at 400 °C could be expected due to the loss in solid solution strengthening at high temperature and inactivation of load carrying capacity of the microstructure due to fragmentation (Figure 10b).
3.5 Al-12Ce-4Si-0.4Mg-0.25Sr alloy Figure 11a demonstrates that further addition of Sr to quaternary alloy re nes intermetallic Ce(Si 1-X Al X ) 2 .Based on hardness values, some critical points (encircled) were selected and XRD study was performed (Figure 11b).Al 0 . 9 Mg 3.1 phase was observed in the as-cast condition and disappears after heat treatment possibly due to its lower melting point. Si intermetallic, as observed in the alloy without Sr, is not observed in the alloy with Sr. Figure 11c reported the heat treatment study up to 100 hours for quinary alloy. It shows that Sr addition to quaternary alloy enhances the quinary alloy's room temperature strength. This may be due to increment in solid solution strengthening by addition of Sr. After 10 hours of heat treatment, a signi cant reduction in hardness was observed at 200 °C. Quinary alloy is characterised by presence of multiple phases and thus making the analysis challenging. The decrease in the hardness can be correlated with less thermal stability and softening and possibly dissolution of Mg 3 Al phase which is con rmed by the XRD pattern ( Figure   11b). For a prolonged heat treatment period, Mg 3 Al phase disappears.

Conclusion
Microstructure stability of ve different Al-Ce based alloys were studied after heat treatment at three different temperatures, i.e., 200 °C, 300 °C and 400 °C.
Microhardness measurements were used to ascertain microstructure stability and strengthening mechanisms were discussed. The major conclusions from the work are as follows: 1. Al-Ce binary alloys show thermal stability at all the temperatures studied (200 -400 ° C) due to higher stability of 2. Al-12Ce-4Si is thermally stable at 200 °C but there is progressive decrease in stability at higher temperatures. Mechanical properties of Al-Ce alloys at room temperature [2][3][4]. In binary alloys, Al-12Ce shows better mechanical properties as compared to other binary alloys. In ternary alloys, Mg increases UTS and YS but simultaneously decreases % elongation. 3D printed alloys (manufactured by selective laser melting) show considerable improvement in mechanical properties compared to as cast ternary alloys.