Secondary phase increases the elastic modulus of a cast aluminum-cerium alloy | Communications Materials

Communications Materials volume 5, Article number: 185 (2024) Cite this article

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Alloying in metal castings is one of the principal methods of strengthening an alloy for various structural and functional applications, but very rarely does it modify an alloy’s elastic modulus. We report a methodology of combining isostructural Laves phases to form a multi-component, high symmetry, isotropic phase that was discovered to enhance the elastic modulus of a cast aluminum alloy to 91.5 ± 7.4 GPa. Flux grown single crystals of the rhombicuboctahedron phase (RCO), so named for the observed morphology, were used to enhance understanding of the structure and mechanical properties of the phase. The pure RCO phase’s structure and site occupancies were co-refined using x-ray and neutron diffraction. Dynamic nanomechanical testing of the cast alloy shows the primary RCO phase has a high, relatively isotropic, elastic modulus. This RCO containing aluminum alloy is found to have a specific modulus that exceeds that of the leading Al, Mg, Steel, and Ti alloys.

Structural and functional materials engineering heavily relies on metals that are prized for their tunable properties by the addition of alloying elements and various processing techniques. Alloying in multi-component metals with significant substitutional solubility, known as high entropy alloys, have expanded the potential of alloying to improve material properties1. Intermetallics and precipitate phases can significantly impact strength, but they typically have little effect on the material's elastic modulus, which is often considered an intrinsic property2 of specific alloy series/classes i.e., Mg3, Fe4, 5, Ti6, Cu7, and Al8,9,10,11.

Titanium alloys, for example, are selected for many applications due to their specific modulus. When combined with its other properties, i.e., specific strength, high melting point, and corrosion resistance, there are no complete substitutes12, 13. Titanium alloys are used in aerospace applications for 80% of their total use. The remaining 20% is used in armor, chemical processing, marine applications, medical implants, power generation, and other applications14. Despite being relatively costly to extract and limited in production (with only 5% of total titanium extraction, or 210,000 metric tons in 202115, going directly to the creation of titanium alloys), titanium’s high specific modulus has allowed it to maintain its position in various applications. This is because it cannot be easily replaced by more readily available and less expensive metal alloys of iron and aluminum13.

Modulus modification of many alloy systems has been achieved through combining metals with ceramic systems into metal matrix composites (MMCs)16. MMCs typically rely on a residual stress field caused by a coefficient of thermal expansion mismatch for strain transfer but are limited by poor interfacial coherency17. Casting MMCs can be difficult and costly, resulting in defects and uneven composite distribution that negatively affects properties18, 19. The shape of composite or precipitate particles can have an important impact on the modification of the alloy stiffness due to how stress fields interact with composite particles20. Elastic modulus is typically measured through two primary approaches: static mechanical testing and dynamic methods. Dynamic techniques, including resonant ultrasound spectroscopy (RUS), impulse excitation (IE), and ultrasound point analysis (UPA), are particularly effective for analyzing materials with simple geometries and high homogeneity. These dynamic methods are ideal for non-destructive, in-situ testing and material qualification21. However, their application is limited in materials with inherent heterogeneities, such as castings characterized by porosity and preferential grain orientations22. In contrast, while static mechanical testing methods are not immune to inaccuracies caused by casting defects, they are less affected by particle interfaces or matrix defects in metal matrix composites (MMCs), offering a more robust measurement in such cases23.

Recently, the development of aluminum rare earth element (REE) alloys has demonstrated high room temperature yield and UTS with high strength retention at temperature24, 25. These alloys, however, have exhibited no increase in modulus using REE additions. Al-Ce alloys, for example, have shown enhanced properties by taking advantage of the high reactivity of Ce to form strengthening intermetallics25,26,27. Specifically, these alloys have demonstrated castability similar to Si containing alloys28, 29 with a reduced reliance on heat treatment30. Al-Ce alloys also exhibit enhanced stability at elevated temperatures due to limited coarsening of Ce containing intermetallics31, 32. The Al-Ce intermetallics reported to date have anisotropic crystal structures and tend to precipitate into high aspect ratio growth morphologies27, 32. Additionally, the intermetallics complicate traditional Al alloying due to the large number of stable ternary and quaternary Al-Ce-X phases. Thermodynamic modeling has been utilized to control phase fraction and morphology of these detrimental phases to improve mechanical properties33.

Generally, the performance of alloys containing needle-like and plate-like, low-symmetry, intermetallic morphologies are hindered by weak interfaces, low dislocation densities, and insufficient slips systems34. Additive manufacturing has been shown to improve mechanical properties via rapid solidification to control phase morphology and refine grain structures35, 36. Sims et al. showed that the yield strength of Al-8Ce-10Mg (wt%) alloy as-cast and hot isostatic pressed (HIP) was 186.5 MPa and 280.4 MPa, respectively27 whereas powder bed fabrication of the same alloy composition has a yield strength of 350 MPa37. In our investigation, high-symmetry cubic phases that are isotropic and exhibit high mechanical properties were sought to strengthen an Al alloy without the need for secondary processing. It was discovered that this approach also leads to a substantial increase in the elastic modulus of the alloy. In the following we first describe our approach to design an isotropic intermetallic phase and then report in detail the observed improvements of the resulting Al alloy.

An alloy design approach combining individual components of isostructural Fd-3m compounds in the CeNi, CuMg, and AlMnCe alloy systems was used to create an isotropic Al-REE based high symmetry phase. Compounds of matching crystal structures from these systems were chosen based on the Hume-Rothery rules38 to achieve substitutional solubility. Cu2Mg and CeNi2 are a cubic Laves structure, C15, with a Fd-3m space group and a lattice parameter of 7.034 Å and 7.248 Å, respectively39,40,41. One ternary study by Coury et al. showed that the rapid solidification of Al-Mn-Ce compositions forms a nanocrystalline Al20Mn2Ce intermetallic with the same space group as cubic Laves (Fd-3m) with a lattice parameter of ~14.49 Å42. Zn and Mg were added to enhance alloy strength, using a method established by Sigworth et al, which has proven to be effective in strengthening the Al matrix phase43. Here, we demonstrate these alloy design criteria can be easily incorporated with traditional casting techniques to produce an isotropic, low aspect ratio precipitate phase with high substitutional solubility.

Figure 1a presents a backscatter micrograph from a scanning electron microscope (SEM) of the as-cast microstructure produced in a gravity cast B108 permanent mold of the composition Al-8.3Cu-7.65Ce-3.3Ni-1Mn-4.5Zn-1.5Mg (wt%). The microstructure is dominated (21.8 ± 2.0 area%) by a light gray euhedral phase with an average diameter of 28.1 ± 3.0 µm (Supplementary Fig. 1). The septenary phase’s composition is estimated to be Al-14.5Cu-17.5Ce-4.8Ni-2.3Mn-4.0Zn-1.2Mg (wt%), using energy-dispersive X-ray spectroscopy (EDS) (Supplementary Table 1). We utilize flux grown single crystals to enhance our understanding of the structure and mechanical properties of the polyhedral phase44. Detailed structural evaluations were performed on high purity 0.5–5 mm single crystals grown in an Al rich flux. Figure 1b depicts one of these crystals exhibiting the characteristic Archimedean solid rhombicuboctahedron (RCO) shape (Fig. 1c). These crystals epitomize the equiaxial habit of this compound observed in the alloy, so we name this phase the RCO.

a A backscatter scanning electron microscopy (SEM) image, revealing an evenly distributed polyhedral phase ranging in size from 10 μm to 30 μm. b A single crystal of the RCO phase with colored reflections to show facets more clearly and (c) a schematic figure of the RCO single crystal showing its high symmetry with red {100}, green {110}, and blue {111} planes. d Neutron and X-Ray diffraction patterns with the plotted Rietveld co-refinement. e The resulting crystal structure from the co-refinement of ground single crystals showing a high symmetry primitive structure that has varying site occupancies for Al, Cu, and Ni.

Single crystal diffraction revealed the structure of the RCO phase (Supplementary Fig. 2a) to be the cubic space group \({Pm}\bar{3}m\) (No. 221) and analogous to a AlCuCeMn phase45; a detailed table of single crystal diffraction results can be seen in Supplementary Materials Table 2. Additionally, the Al8Ce3Cu reported by Perrin et al. has a matching space group in both cast and AM versions of the alloy46. As mentioned before, ternary and quaternary strengthening intermetallics for Al-Ce alloys have only been reported to have high aspect ratio morphologies and presumably anisotropic properties. One example of this is in Al-Ce-Si-X casting’s where large plate-like lamella of CeAlSi intermetallics readily form as primaries during solidification32. A slowly precipitated single crystal of CeAlSi reflects this tabular morphology (Supplementary Fig. 2b). We will compare the anisotropic mechanical properties of this phase to the RCO phase. An idealized model of the CeAlSi crystal shape is rendered in Supplementary Fig. 2c consistent with the refined tetragonal structure depicted in Supplementary Fig. 2d (Supplementary Table 3). X-ray single crystal diffraction data is limited in accuracy due to atomic scattering factor overlap between Cu and Ni in the RCO phase, the content of which was confirmed with EDS measurements in Supplementary Table 1. Figure 1d presents powder x-ray and powder neutron diffraction measurement of ground RCO single crystals. The refined structure using XRD alone shows a cubic crystal structure with a lattice parameter of 8.39513(2) Å, close to the single crystal diffraction result (Supplementary Table 4). X-ray diffraction of the cast alloy (Supplementary Fig. 3) shows three main phases Al-FCC at 74(2)wt%, Al8Cu4Ce at 6(1)wt%, and the RCO phase with a = 8.38 Å at 21(1)%. The small volume fraction phases in the SEM micrograph were not observed with XRD. A co-refinement of the x-ray and neutron data allows us to refine the site occupancies in the RCO phase by leveraging the distinctly different scattering factors for the elements with each technique. A common crystal structure was used for both data sets, but the lattice parameters were allowed to differ (XRD a = 8.39516(7) and Neutron a = 8.3935(2) Å). Data from two of the NOMAD detector banks were used to extend the q-range covered as can be seen in the upper half of Fig. 1d. The 12j Wyckoff site was observed to have significant compositional complexity due to substitutional occupancy by Al, Ni, and Cu taking up one third of the total available sites in the crystal structure. First principles calculations could be used to validate that the ratio measured in these crystals is the most stable ratio of occupancies for a given composition. Based on the co-refinement the composition of the RCO phase was estimated to be Al25.3Ce3Cu3.6Ni3.1Mn. One interesting note for the RCO phase’s structure is that the lattice parameter is 2.07 times the lattice parameter for pure Al and both phases have a cubic structure offering an opportunity for an epitaxial interface, but no evidence of this has been observed.

Figure 2a shows a 200 kV scanning transmission electron microscopy (STEM) high-angle annular dark field (HAADF) image for the [110] projection of the RCO phase taken from the cast alloy. The inset image shows the same zone axis imaged at 60 kV with a high current to measure atomic resolution EDS as seen in Fig. 2b (Supplementary Note 1). Atomic spacing and elemental mapping are in good agreement with the calculated structure using XRD and neutron diffraction as projected on the same [110] zone axis in Fig. 2c. STEM and EDS measurements confirm that the refined structure of the single crystal matches the structure of the RCO phase in the alloy.

a A 200 kV STEM HAADF of the RCO phase in the alloy taken in the [110] zone axis with an elevated level of atomic clarity. An in set of a region imaged at 60 kV showing the (b) EDS mapping results with layers of Ce and Mn atoms with Ce and Al, Cu, Ni shared sites. c A projection along the [110] direction of the RCO crystal structure matching the refined structure in Fig. 1e and the measured structure in a. d 200 kV STEM HAADF and BF images of the RCO and FCC-Al with an inset SAED image of the interface.

Characterization of the RCO’s interface with the Al matrix was performed to elucidate the RCO phases’ effect on the mechanical properties of the alloy. The interface between the RCO phase and the Al matrix could affect the elastic modulus of the alloy by changing the way that strain is distributed within the material. TEM of the RCO and the FCC-Al matrix interface was performed to better understand the interfacial structure and an orientation relationship. Figure 2d shows STEM HAADF and brightfield (BF) images of the RCO phase and the Al matrix with an inset of a selected area diffraction (SAED) pattern. The tilt between the [110]RCO and the [110]Al zone axes on a double tilt holder was calculated to be 43.3 degrees. In this interface, the tilt between the {100}RCO and {100}Al planes was measured to be 11.5 degrees. Plane matching of the {100} plane families is seen to be relatively periodic at the interface. The measured lattice mismatch between the (400)RCO and (200)Al planes was found to be relatively low at 2%, see Supplementary Fig. 4 for detailed qualitative analysis. Our images have not, however, revealed an orientation relationship or a preferred interfacial structure. Secondary phases that effectively enhance the elastic modulus of aluminum alloys often involve coherent or semi-coherent bonding, allowing efficient load transfer between the matrix and the phase. The physical nature and thermodynamic stability of composite interfaces significantly influence the mechanical properties of MMCs. Minimizing deviations from ideal planar geometries and impurity segregation improves load transfer and performance17, 47. Although more work is needed to determine the exact nature and variability of the interfacial relationship between the phases, fractography in Supplementary Fig. 6 shows RCO particle interfaces are strong enough to prevent premature failure or pull-out. The isotropic nature of the RCO phase, in both mechanical properties and shape, helps to minimize stress concentrations as is discussed further.

Mechanical property testing was performed on various permanent mold castings with extensometery. Figure 3a shows the Young’s Modulus [GPa] vs Density [g/cm3] for our cast alloy Al-20%RCO and various metal alloy systems for comparison such as Mg3, Fe4, 5, Ti6, Cu7, and Al8,9,10,11. The variance in modulus values is small for metal alloys especially when compared to the variance seen in yield strength. As stated previously, it is difficult to increase a cast aluminum’s elastic modulus through alloying additions alone. The cast aluminum alloy containing the RCO phase exhibits an elastic modulus averaging 91.5 ± 7.4 GPa measured in tensile tests seen in Supplementary Fig. 5 and Table 1.

a A summary of Young’s Modulus vs. Density for select cast alloy systems including our work. b Compliance corrected compression testing at temperature for the Al-20%RCO tested at RT, 100 °C, 150 °C, and 200 °C showing strength retention of 60% at 150 °C. c Load sharing behavior of the RCO and Al matrix with lattice strains of both phases along and perpendicular to the loading axis plotted as a function of stress. d Phase load-sharing for Al–20%RCO from tensile testing that shows a higher load carried by the RCO phase during elastic loading.

Young’s modulus was determined using a variety of sample geometries, including both flat and round subsized bars compliant with ASTM E8 standards, as well as an as-cast bar following ASTM B108 guidelines. The modulus measurements were conducted using knife-edge extensometers, with the modulus itself being derived from the initial linear portion of the stress-strain curves obtained during testing. Notably, all samples displayed a predominantly brittle failure with limited ductility. Fractographic analysis, illustrated in Supplementary Fig. 6, reveals a distinct failure pattern in the tensile bars. The aluminum matrix exhibited ductile dimpling, indicative of plastic deformation, whereas the RCO grains showed signs of brittle fracture, namely radial marks. This behavior underscores the effective stress transfer and crack propagation through RCO crystals, demonstrating a robust interface between the phases. This is in contrast to some MMCs, where failure often involves the pull-out of composite particles, highlighting a weaker interfacial bond48, 49. The modulus value for the Al-20%RCO alloy is in the range of Ti and Cu alloys, however, it has a density of 3.202 g/cm3, 30% lower than Ti-6Al-4V. For direct comparison, the specific Young’s modulus [GPa cm3/g] of the Al-20%RCO alloy was found to be 27.86 ± 2.9 which is higher than A356 by 5%, 1050 Steel by 6%, AZ91C by 12%, and Ti-6Al-4V by 16% (Supplementary Fig. 7). To our knowledge, modulus modification using the RCO phase is the highest achieved in a cast aluminum alloy without using beryllium, MMC’s, or powder processing techniques50, 51.

Figure 3b shows compression testing results for thermally exposed samples tested at room temperature, 100 °C, 150 °C, and 200 °C. The room temperature compressive yield was measured to be 325 MPa while yield strength at 100 °C, 150 °C, and 200 °C is 307 MPa, 274 MPa, and 194 MPa. Compressive strength retention at 100 C and 200 C was measured at 94% and 60%, respectively. A comparison to a literature measurement of aluminum A356 can be seen in Supplementary Fig. 8 showing Al-20%RCO is stronger but exhibits a lower compressive ductility. The RCO phase in Fig. 4e, f demonstrates a lack of coarsening due to thermal exposure, aiding strength retention at temperature. The reduction of strength observed at 200 °C is associated with Cu precipitate coarsening (Supplementary Fig. 9) allowing the Al-FCC matrix to flow more freely around the RCO phase at temperatures above 200 C. Precipitates formed in the alloy during solidification due to excess Cu added based on Henderson et al.’s work to introduce Cu precipitates in Al-Ce-Cu alloys52 and ultimately increase the alloys strength.

a CSM nanoindentation results for the RCO and CeAlSi single crystals measured on different crystallographic planes showing isotropic and anisotropic behavior, respectively. b AFM measurements of a 50mN indent on the RCO single crystal in the [100] direction with (c) line profiles showing pile-up around the edges of the indent. d SEM images of a 1000mN indent on the RCO single crystal showing deformation extending past the indent edge. e A backscatter SEM image of a RCO containing alloy in the as cast condition. f A backscatter SEM image of a RCO containing alloy heat treated at 500 °C for 4 h. g Energy-dispersive X-ray spectroscopy (EDS) mapping from the same location as e showing the elements contained within the RCO phase. h EDS mapping from the same location as f showing the elements contained within the RCO phase and an increased relative Zn content in the thermally exposed condition. i CSM mapping results in the as cast alloy and j thermally exposed sample both showing a high modulus for the low aspect ratio RCO phase. k Cluster analysis of the as cast sample’s modulus mapping and l cluster analysis of the thermally exposed sample’s modulus mapping.

To determine what affect the RCO phase has on the stress-strain response in the alloy we measured the elastic strain accumulation rate for RCO and the Al matrix during mechanical loading using a methodology described in Choo et al. and applied to cast alloys in Sims et al. and Henderson et al.17, 32, 37. In-situ diffraction during tensile testing was performed on a recast Al-20%RCO alloy (Supplementary Fig. 10) using a synchrotron X-ray source, resulting in the lattice strain vs. applied tensile stress plot that can be seen in Fig. 3c. Diffraction similarities as seen in Supplementary Fig. 11 makes peak deconvolution difficult, but the lattice strain vs stress slopes for the RCO phase and FCC-Al are nearly the same as the system deforms elastically. Slope matching in this region indicates strain is effectively transferred across the phase interfaces. The RCO phase exhibits elastic load sharing up to 300 MPa. In contrast, Sims et al. measured the load partitioning of the Al11Ce3 intermetallic and found elastic load sharing only present up to 50 MPa32. During yielding of Sims et al.’s alloy, Al began to take on far less of the strain and the strain response is dominated by the Al11Ce3 phase32. Figure 3d shows how the load is shared in our alloy between the RCO phase and the FCC-Al matrix, more explicitly the RCO phase begins to take on strain directly at the onset of loading. Henderson et al. found that by using laser powder bed manufacturing the Al11Ce3 was nanostructured, and the load partitioning showed elastic strain transfer up to ~300 MPa due to a new low aspect ratio nano scale morphology37. The same elastic strain transfer is seen up to 300 MPa in the Al-20%RCO alloy with no nano structuring of an intermetallic. By directly bulk casting the alloy we formed a uniformly distributed RCO phase that load shares directly with the Al matrix.

The RCO’s cubic symmetry could mean its intrinsic mechanical properties, like elastic modulus, are nearly isotropic53. Figure 4a shows continuous stiffness measurement (CSM) nanoindentation results for the RCO and CeAlSi single crystals in two different crystallographic directions. The RCO was indented normal to the (100) and (110) facets with average modulus values of 174.9 ± 2.3 GPa and 172.2 ± 2.6 GPa, respectively, showing variation that is within measurement errors. In contrast, the CeAlSi crystal was measured on the (001) and (101) facets with values of 167.5 ± 1.5 GPa and 113.9 ± 0.8 GPa, respectively. This 47% difference of the stiffness of CeAlSi along two directions exemplifies the anisotropic nature of previously reported Al-Ce-X intermetallics.

Figure 4b shows a topographical analysis of a 50mN indent that exhibits pileup of around ~50 nm around the indent’s edges. Figure 4c shows profiles along each indent vertex that more clearly displays the pile up around the indents edge. Figure 4d shows a 1000mN indent of the RCO single crystal with cracking at the vertices and a crack 5 \({{\rm{\mu }}}{{\rm{m}}}\) away from the indent edge. The RCO phase exhibits a brittle fracture similar to the CeAlSi crystal in Supplementary Fig. 12, but to a lesser extent.

Phase specific property measurements, namely nanomechanical mapping, can establish structure-property relationships at microstructural length scales. Figure 4e, f shows backscatter SEM images of the as cast Al-20%RCO and the same alloy that has undergone a thermal exposure at 500 °C for 4 h with a water quench. Under these conditions, no change is seen in the intermetallic phases present, including the RCO phase of interest. The RCO phase displayed a lack of coarsening due to the limited solubility of Ce into the Al matrix54, as well as its low aspect ratio polyhedral shape. Formed as a primary phase, the RCO has a low surface area to volume ratio which is a lower energy state compared to lamellar plates, reducing the thermal diffusion driving force. Figure 4g, h shows EDS mapping for the same regions as the SEM micrographs displaying the elements measured inside the RCO phase. Full EDS maps can be seen in Supplementary Fig. 13. Ce, Cu, Ni, Mn, Zn, and Mg are all present in the cast RCO phase (Supplementary Note 2). Figure 4i, j shows nanomechanical mapping of the same locations as the SEM images and EDS maps, tying composition and phase morphology to phase specific properties. Modulus distributions for the nano-mapping can be seen in Supplementary Fig. 14, but a cluster analysis was done on each map and can be seen in Fig. 4k, l. The cluster analysis shows that the modulus for the RCO exceeds the modulus measurements of both the matrix and the other Al-Ce-X intermetallics present. The average modulus of the RCO phase was measured to be 165 ± 1.2 GPa for the as cast alloy and 169 ± 1.2 GPa for the thermally exposed sample (Supplementary Note 2).

A general rule of mixtures can be used to determine a MMC’s elastic modulus but has not been shown to work with secondary phases in metal castings55. The upper and lower bounds of composite elastic moduli can be calculated using a general rule of mixtures (ROM) and inverse rule of mixtures (IROM) formulas for elastic modulus:

And,

Where E is the phases elastic modulus and V is the volume fraction of the phase present. Using the averaged values after clustering values from nanomechanical mapping reveals an upper bound modulus of approximately 89.0 GPa, only 5% different from the measured modulus of the alloy using tensile testing. The lower bound is calculated to 77.9 GPa and no samples measured showed this value. Additional volume fractions of Al-8%RCO, Al-12%RCO, and Al-14%RCO were cast to validate the change in modulus following the ROM and IROM models. Supplementary Fig. 15 shows the modulus vs area fraction for each sample, where the elastic modulus of the Al-RCO system increases with the RCO fraction, ranging from 87.1 ± 5.5 GPa for Al-8%RCO to 91.5 ± 7.4 GPa for Al-20%RCO. This trend is consistent with theoretical predictions based on the ROM and IROM. Although the measured values are not fully consistent with the nano-indentation mapping results, the enhancement in elastic modulus aligns with the rule of mixtures when considering solid solution and nano-precipitate contributions of Cu, Mg, Ni, and Mn, raising the matrix’s modulus to the calculated 78.2 GPa, and fitting the rule of mixtures with the measured modulus of the single crystal RCO phases average modulus, 173.6 ± 2.5.

Based on the evidence presented the Al-20%RCO alloys modulus is what is expected from an MMC containing spherical, isotropic, high modulus particles with a strong, strain-sharing interface with the matrix. One measured interface, which is not enough to completely define an orientation relationship, shows no voids and few defects. The consistent interface could contribute to the phases ability to transfer strain between phases, as seen in both fractography and validated using in-situ tensile testing. The advantage to other MMC systems is the possibility of controlling the phase size, morphology, and volume fraction using casting and solidification techniques.

Material selection and design require a comprehensive evaluation of properties, including strength and modulus. The elastic modulus of an alloy is usually determined by its primary alloying elements bonding energy, which can make it challenging to optimize its design through metallurgy. In the most advanced Al alloys high strengths and strength retention at temperature can be achieved, but little modification is made to their elastic modulus. Using traditional bulk casting techniques, we have developed a new class of Al alloys with an increased elastic modulus. Multi-component, isostructural alloy design criteria were used to cast a seven-element alloy to form an Al-rich, cubic phase that forms rhombicuboctahedron (RCO) crystals during primary solidification. Nanomechanical testing of the single crystal and the Al-20%RCO alloy shows the RCO phase has a high modulus that is also isotropic. Additionally, thermal exposures have almost no coarsening effect on the RCO phase, and strength is retained at elevated temperature. Modulus modification is achieved through load sharing during elastic deformation through a strong, strain-sharing interface between the RCO phase and the FCC-Al matrix.

Development of isotropic, high symmetry precipitate phases that allow for stiffness modification like our RCO phase presents a new alloy design methodology for lightweight applications. Ti alloys are commonly selected for their specific modulus but are less accessible compared to other alloys due to primary and secondary availability. The alloy that features the RCO phase exhibits a modulus comparable to that of Ti and Cu alloys, and a density that is lower, surpassing the specific modulus of the leading Al, Mg, Steel, and Ti alloys.

All samples were initially melted and mixed in an open-air resistive Baker furnace in 22.7 kg batches. Samples were casting in an ASTM B108 steel permanent mold preheated to 400 °C. Commercially pure elements were melted in a two-stage process to ensure good mixing. A master alloy of Al, Ni, and Cu were melted together first, and then the remaining elements, Ce, Mn, Zn, and Mg were added and mixed further. The thermally exposed sample was heated to 500 °C for 4 h followed by a water quench in 80 °C water.

Samples for compression and tensile testing were remelted in an Arc200 arc melter and mixed using an electromagnetic stirring function. Material was cast into a 0.5”x3.25” cylindrical mold to be machined into sub-sized E8 tensile bars. Material was also cast into a 1”x0.35”x1.5” rectangular mold to be machined into square subsized ASTM E8 tensile bars and compression cylinders.

Atomic Force Microscopy (AFM) was performed on an Asylum Research Cypher S AFM using an OTESPA silicon probe with a resonant frequency of 300 kHz and spring constant of 26 N/m. Scans consisted of 256 lines at a scan rate of 0.1 Hz. Data collected was analyzed using Gwyddion software to quantify pileups around nanomechanical testing sites56.

The SEM samples were ground using SiC papers and polished using water-based diamond polish up to 1 μm. All images were collected using backscatter on a Helios 5 Hydra DualBeam. Energy-dispersive X-ray spectroscopy (EDS) point scans and maps were collected using an AMETEK EDAX Octane Elite Super.

Single crystals of cubic Al-Ce-Ni-Cu-Mn phase were grown by slowly cooling a high temperature aluminum melt. A master alloy was generated by arc melting together elemental metals with the atomic ratio of Al53.7%-Ce18.14%-Cu19.08%-Ni7.44%-Mn2.86%. Pieces of this master alloy and additional aluminum metal yielding a composition of Al72.38%-Ce10.74%-Cu10.75%-Ni4.43%-Mn1.70% were loaded into an alumina 5 mL Canfield crucible set57. This was sealed in a fused silica ampoule filled with argon to exclude air. This assembly was heated to 1050 °C in 6 h and then held for 12 h to completely melt everything. Crystals were grown by slowly cooling to 700 °C over 100 h. The remaining liquid was removed from the crystal by removing the ampoule from the furnace and immediately centrifuging.

Crystals of CeAlSi were also grown from an aluminum rich melt. The elements were loaded into a 5 mL alumina Canfield crucible set57 with an atomic fraction of: Al86.99%-Ce5.01%-Si8.00%. This was sealed in a fused silica ampoule and the assembly was placed in a box furnace. The furnace was heated to 1100 °C over 9 h and held for 12 h to form a homogenous melt. Crystals were grown during a slow, 100 h cool from 1100 °C to 700 °C. Finally, the assembly was removed from the furnace and centrifuged to separate the remaining liquid from the crystals. Both crystals were modeled in JCrystal and rendered in in POV-Ray.

The measurements of single crystal X-ray diffraction were performed on Rigaku XtaLAB mini II diffractometer at 0.6 kW (50 kV and 12 mA) power using graphite-mono-chromated Mo-Ka radiation (λ = 0.71073 Å). The data was collected with an exposure time of 1.5 s and the scan width of 1°. Cell refinements were carried out using CrysAlisPro software. The measured crystal structure was solved by direct methods and refined using the full-matrix least-squares against F2 using ShelXT58 and ShelXL59 programs as implemented in OLEX2 software60. The non-hydrogen atoms were assigned with isotropic (or anisotropic) displacement parameters.

Single crystals of the RCO phase were ground into a fine powder. The powder was dispersed on a flat “zero-background” silicon sample holder with acetone. The XRD spectra were collected by a Malvern Panalytical Empyrean using a Cu Kα source (wavelength, 0.15406 nm) operated at 40 kV and 40 mA. The diffraction patterns were recorded over a 2θ range from 10° to 140° at a step size of 0.0065°. The refinement of the XRD patterns was obtained using FullProf software. This same process was used to measure the XRD pattern for the cast alloy on a polished sample.

Neutron powder diffraction was performed using the Nanoscale Ordered Materials Diffractometer (NOMAD) at the Spallation Neutron Source at Oak Ridge National Laboratory61. Samples were loaded into quartz glass capillaries with 2.8 mm inner and 3 mm outer diameter. The samples were aligned in the beam by translating it through the beam while monitoring the scattered intensity and held at constant temperature by immersion into an Ar gas stream kept at 300 K. Neutron diffraction data was collected for an accelerator proton charge of 8 C, corresponding to about 90 min data collection at 1.4 MW accelerator power. An empty capillary as well as the empty diffractometer was measured as well, scattering from a Vanadium rod was used for normalization and scattering from diamond powder for calibration. The co-refinement of the XRD and Neutron patterns was calculated using Fullprof software.

A transmission electron microscopy (TEM) lamella was prepared by pulling a sample of the RCO phase in a cast specimen using a plasma focused ion beam (PFIB) with a xenon source on a ThermoFisher Helios 5 Hydra DualBeam. Scanning transmission electron microscopy (STEM) high-angle annular dark field (HAADF) and bright field (BF) images were taken using a ThermoFisher FEI Spectra 300 MC. at 200 kV. Selected area diffraction of the phase interfaces was performed using the same machine at 200 kV. For STEM EDS scans, the voltage was dropped to 60 kV with a high current to increase spectroscopy signal. Image processing was performed in the ThermoFisher Velox software. Simulated diffraction was performed in SingleCrystal® by CrystalMaker Software Ltd using the structure refined from the neutron diffraction and XRD co-refinement.

Values from direction specific modulus measurements on single crystals were averaged from 300 microns forward due to increased errors below 300 microns.

Nanomechanical properties were investigated via continuous stiffness measurement (CSM) nanoindentation method following ISO 14577 standards. A KLA iMicro Pro equipped with an InForce 50mN and 1000mN actuator both with a diamond Berkovich indenter tip was used for all nanoindentation methods. All methods used the dynamic CSM test method with a target frequency of 110 Hz and a displacement amplitude of 2 nm. For the specific mechanical property maps target loads of 5mN were used on a region with a resolution of 0.833 indents per μm (i.e., 300 μm X 300 μm with 250 × 250 indents). Outliers were removed from the data for heat map clarity based on them being placed on defects on the surface. After performing an area function calibration on fused silica samples and correcting for load-frame compliance, the hardness and elastic modulus for every indent was calculated using the standard Oliver-Pharr method. Single crystal samples were ground up to 1200 grit on SiC paper and then polished with 3 µm, 1 µm, and 0.05 µm diamond suspension polishing fluid.

A TA Instruments SDT-Q600 was used to perform a DSC measurement with a heating rate of 10 C per minute. A 10 mg sample of the Al-20%RCO alloy was heated to 700 C, then cooled to 30 C and heated again to 700 C inside an alumina pan.

Subsized E8 cylindrical and flat tensile bars were tensile tested with an Instron 5566 with a 10kN load cell and knife extensometer attached. B108 cylindrical tensile bars were tested on a UNITED STM-20 with a 100kN load cell and a knife extensometer. Samples of differing geometries and tensile frames showed similar elastic modulus measurements using graphical analysis of the tested samples. In-situ high-energy synchrotron X-ray diffraction experiments during tensile loading were performed at the FAST beamline (ID-3A), Cornell High Energy Synchrotron Source (CHESS), US. The X-ray beam energy was 41.991 keV (i.e., wavelength 0.29526 Å) and the beam was collimated to a size of 0.25 × 0.25 mm (W × H). Uniaxial tensile loading was performed in displacement control with a strain rate of \(1\times {10}^{-4}{s}^{-1}\) using the Rotation and Axial Motion System (RAMS2) load frame. The diffraction data acquisition was performed using an area detector (GE Detector 2048 × 2048 pixels, 200 × 200 μm^2 per pixel). The far field detector was centered on the incoming X-ray beam and sat 800 mm behind the sample. A calibration measurement on a cerium dioxide (CeO2) powder specimen was performed and the software package GSAS-II was used for the detector calibration62. For calculating the lattice strain, the diffraction patterns were integrated in 10° azimuthal bins from 85° to 95°. Single-peak fitting was performed using GSAS-II software62 to obtain the d-spacing of (320) RCO and (200) Al phases. The lattice strains were then calculated using

Where, \({d}_{0,{hkl}}\) and \({d}_{{hkl}}\) are the measured interplanar spacings for the stress-free and the stressed states, respectively. The strain was corrected using SEM imaging of untested samples machined using the same method to measure the actual gauge length which was found to be 0.85 mm instead of the test methods assumed 1 mm.

Compression testing was performed on machined cylinders from the recast alloy with dimensions of 3 mm in diameter and 6 mm in length using an Instron 5566 with a 10kN load cell. Inconel 718 rods with Inconel bearing blocks with a diameter of 12.5 mm were used to ensure an even distribution of initial force. The sample was placed between the bearing blocks and the sample was brought up to test temperature within 15 min and held for 15 min with a thermocouple placed on the sample during the hold to ensure temperature accuracy. Samples were compressed using crosshead speed control with rates equivalent to 0.005 (mm/mm-min) in the elastic region for the room temperature test. Compliance corrections were made by compressing the two bearing blocks together at each given temperature and modeling the extension vs. load to a 4th order polynomial. The sample’s measured load is then used to calculate extension contributions from the test setup and are then subtracted from the extension test data.

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We thank Michael Koehler and Andreas Kreyssig for their assistance with structural characterization. Microscopy and XRD instrument access provided by the Microscopy and Diffraction facilities at the Institute for Advanced Materials & Manufacturing (IAMM) at the University of Tennessee, Knoxville. Single crystal growth and structural characterization by WRM is supported by the Gordon and Betty Moore Foundation’s EPiQS Initiative through Grant No. GBMF9069. This research received funding from the DEVCOM Army Research Laboratory and was accomplished under Cooperative Agreement Number W911NF2220007. The views and conclusions contained in this document are those of the authors and should not be interpreted as representing the official policies, either expressed or implied, of the Army Research Laboratory of the U.S. Government. The U.S. Government is authorized to reproduce and distribute reprints for Government purposes notwithstanding any copyright notation herein. In-situ tensile testing work is based on research conducted at the Center for High-Energy X-ray Sciences (CHEXS), which is supported by the National Science Foundation (BIO, ENG and MPS Directorates) under award DMR-1829070.

Nitish Bibhanshu

Present address: Indian Institute of Technology–Ropar, Bara Phool, Punjab, India

Materials Science and Engineering, The University of Tennessee Knoxville, Knoxville, TN, USA

Max L. Neveau, William R. Meier, Hyojin Park, Michael J. Thompson, Nitish Bibhanshu, Matthew F. Chisholm, Orlando Rios & Gerd Duscher

Institute of Environmental Sciences, University of Leiden, Leiden, The Netherlands

Catrin Böcher & Tomer Fishman

Eck Industries, Inc, Manitowoc, WI, USA

David Weiss

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Max Neveau: conceptualization, methodology, investigation, formal analysis, visualization, writing – original draft, visualization, writing – review & editing. William R. Meier: methodology, investigation, formal analysis,, writing – original draft, writing – review & editing. Hyojin Park: investigation, formal analysis. Michael J. Thompson: investigation, formal analysis. Nitish Bibhanshu: investigation, visualization. Catrin Böcher: investigation, writing – original draft, writing – review & editing. Tomer Fishman: investigation, writing – review & editing. David Weiss: conceptualization, methodology, resources, supervision. Matthew F. Chisholm: methodology, writing – original draft, writing – review & editing, supervision. Gerd Duscher: conceptualization, methodology, writing – review & editing, project administration, funding acquisition. Orlando R. Rios: conceptualization, project administration, funding acquisition.

Correspondence to Gerd Duscher.

The authors declare no competing interests.

Communications Materials thanks Qingsong Pan and the other, anonymous, reviewer(s) for their contribution to the peer review of this work. Primary Handling Editors: Xiaoyan Li and John Plummer. A peer review file is available.

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Neveau, M.L., Meier, W.R., Park, H. et al. Secondary phase increases the elastic modulus of a cast aluminum-cerium alloy. Commun Mater 5, 185 (2024). https://doi.org/10.1038/s43246-024-00611-3

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DOI: https://doi.org/10.1038/s43246-024-00611-3

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