Tailoring the formation of twins in Al by introducing epitaxial layer interfaces
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Materials Science↗
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Iodine capture with mechanically robust heat-treated Ag-Al-Si-O xerogel sorbents
Silver-loaded heat-treated aluminosilicate xerogels (Ag-HTX) were evaluated as sorbents for iodine [I2(g)] capture. The heat treatment step was performed to help increase the mechanical integrity of the gels. The synthesized xerogels were characterized using powder X-ray diffraction, scanning electron microscopy, energy-dispersive X-ray spectroscopy, transmission electron microscopy, Brunauer-Emmett-Teller analysis, gravimetric iodine loading, nanoindentation, and dynamic mechanical analysis. The structural and chemical analyses of Ag-HTX showed uniform distribution of Ag throughout the gel network after Ag-exchange. After I2(g) capture, the AgI crystallites were observed in the sorbent verifying chemisorption as the iodine capture mechanism. Iodine loading of this xerogel was 0.43 g g-1 at 150°C over 1 d and 0.52 g g-1 at 22°C over 33 d. The specific surface area of Ag-HTX was 202 m2 g-1 and decreased to 87 m2 g-1 after iodine loading. The hardness of the heat-treated xerogel was 160 times higher than heat-treated aerogel of the same composition. The heat-treatment process increased the modulus value to 40.77 MPa from 6.99 MPa of as-made xerogel, demonstrating the need for this added step in the synthesis process. These results show that Ag-HTX is a promising sorbent for I2(g) capture with good loading capacity and mechanical stability.
xerogels, aerogels, iodine capture, hardness testi↗
In situ characterization of tensile behavior of laser rapid solidified Al–Si heterogeneous microstructures
Not provided.
Materials Science↗
BrzostekEcologyLab/InSitu_Incubation: Ridgeway et al., 2023 Data
This data is from an labelled litter decomposition experiment in root ingrowth cores in Miscanthus x giganteus
Ridgeway, Joanna↗
Materials Data on Al(SiO3)2 by Materials Project
Al(SiO3)2 is Low Tridymite-derived structured and crystallizes in the triclinic P1 space group. The structure is three-dimensional. there are four inequivalent Al sites. In the first Al site, Al is bonded to four O atoms to form AlO4 tetrahedra that share corners with four SiO4 tetrahedra. There are a spread of Al–O bond distances ranging from 1.74–1.76 Å. In the second Al site, Al is bonded to four O atoms to form AlO4 tetrahedra that share corners with four SiO4 tetrahedra. There is two shorter (1.74 Å) and two longer (1.75 Å) Al–O bond length. In the third Al site, Al is bonded to four O atoms to form AlO4 tetrahedra that share corners with four SiO4 tetrahedra. There is two shorter (1.74 Å) and two longer (1.75 Å) Al–O bond length. In the fourth Al site, Al is bonded to four O atoms to form AlO4 tetrahedra that share corners with four SiO4 tetrahedra. There are a spread of Al–O bond distances ranging from 1.73–1.76 Å. There are eight inequivalent Si sites. In the first Si site, Si is bonded to four O atoms to form SiO4 tetrahedra that share corners with two AlO4 tetrahedra and corners with two SiO4 tetrahedra. There is two shorter (1.61 Å) and two longer (1.65 Å) Si–O bond length. In the second Si site, Si is bonded to four O atoms to form SiO4 tetrahedra that share a cornercorner with one SiO4 tetrahedra and corners with three AlO4 tetrahedra. There are a spread of Si–O bond distances ranging from 1.61–1.65 Å. In the third Si site, Si is bonded to four O atoms to form SiO4 tetrahedra that share a cornercorner with one SiO4 tetrahedra and corners with three AlO4 tetrahedra. There are a spread of Si–O bond distances ranging from 1.61–1.64 Å. In the fourth Si site, Si is bonded to four O atoms to form SiO4 tetrahedra that share a cornercorner with one SiO4 tetrahedra and corners with three AlO4 tetrahedra. There are a spread of Si–O bond distances ranging from 1.61–1.64 Å. In the fifth Si site, Si is bonded to four O atoms to form corner-sharing SiO4 tetrahedra. There is one shorter (1.62 Å) and three longer (1.63 Å) Si–O bond length. In the sixth Si site, Si is bonded to four O atoms to form SiO4 tetrahedra that share a cornercorner with one SiO4 tetrahedra and corners with three AlO4 tetrahedra. There is three shorter (1.62 Å) and one longer (1.64 Å) Si–O bond length. In the seventh Si site, Si is bonded to four O atoms to form corner-sharing SiO4 tetrahedra. There is one shorter (1.62 Å) and three longer (1.63 Å) Si–O bond length. In the eighth Si site, Si is bonded to four O atoms to form SiO4 tetrahedra that share corners with two AlO4 tetrahedra and corners with two SiO4 tetrahedra. There are a spread of Si–O bond distances ranging from 1.61–1.64 Å. There are twenty-four inequivalent O sites. In the first O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the second O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the third O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the fourth O site, O is bonded in a bent 150 degrees geometry to two Si atoms. In the fifth O site, O is bonded in a bent 150 degrees geometry to two Si atoms. In the sixth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the seventh O site, O is bonded in a bent 150 degrees geometry to two Si atoms. In the eighth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the ninth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the tenth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the eleventh O site, O is bonded in a bent 150 degrees geometry to two Si atoms. In the twelfth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the thirteenth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the fourteenth O site, O is bonded in a bent 150 degrees geometry to two Si atoms. In the fifteenth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the sixteenth O site, O is bonded in a bent 150 degrees geometry to two Si atoms. In the seventeenth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the eighteenth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the nineteenth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the twentieth O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the twenty-first O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the twenty-second O site, O is bonded in a bent 150 degrees geometry to one Al and one Si atom. In the twenty-third O site, O is bonded in a bent 150 degrees geometry to two Si atoms. In the twenty-fourth O site, O is bonded in a bent 150 degrees geometry to two Si atoms.
36 MATERIALS SCIENCE↗
Microstructure and creep properties of cast near-eutectic Al–Ce–Ni alloys
This study investigates the as-cast and aged microstructures, thermal stability, ambient temperature strengthening, and creep resistance of three ternary Al–Ce–Ni alloys (wt%): near-eutectic Al–10Ce–5Ni (with both eutectic and hypoeutectic regions), hypoeutectic Al-7.5Ce-3.75Ni (with numerous primary Al dendrites), and hypereutectic Al-12.5Ce-6.25Ni (with coarse, blocky primary Al3Ni and Al11Ce3 precipitates and some primary Al dendrites). Depending on the alloy composition and local solidification conditions, the following eutectic morphologies are found: (i) coarse Al–Ce eutectic colonies where Al 11 Ce 3 is in the form of “Chinese script”, (ii) intermingled regions of binary Al–Ce and Al–Ni eutectic colonies, with finer Al 3 Ni and Al 11 Ce 3 fibers, (iii) large ternary eutectic colonies, where the binary Al 3 Ni and Al 11 Ce 3 phases are alternating or intertwining within the individual, fine fibers (diameters of ~60–170 nm, depending on solidification rates), and (iv) ternary eutectic zones (between primary Al dendrites), where fine Al 3 Ni and Al 11 Ce 3 build up a 3D-interconnected network. The high volume fraction of intermetallic phases and extremely fine eutectic spacing/fiber diameter both contribute to high ambient strengthening (higher as-cast microhardness than binary Al–Ce or Al–Ni), and also provide enhanced creep resistance at 300 and 350 °C. Additionally, the alloys are coarsening-resistant up to 425 °C for extended periods, with a gradual decrease in microhardness. The alloys aged at 400 °C to 1050 h show fiber fragmentation and coarsening of the resulting particles, with the faster-diffusing Ni driving more rapid coarsening of the Al 3 Ni particles which engulf finer, more stable Al 11 Ce 3 particles. Severe overaging (performed at 590 °C for 24 h) leads to Al 3 Ni and Al 11 Ce 3 spheroids which remain submicron-sized in eutectic colonies, but micron-sized at colony boundary and at Al dendrite-eutectic interface. Creep resistance at 300 °C of overaged Al–10Ce–5Ni remains substantial, consistent with load-transfer based composite strengthening being an important strengthening mechanism in these alloys, making them excellent candidates for replacement of heavier steel or titanium parts operating under stress up to 300 °C.
36 MATERIALS SCIENCE↗
Thermodynamics of Tritium Trapping by Point Defects in Intermetallic Al 12 (TM) 2.35 Aluminide Coating Phase
Density functional theory simulations have been carried out to investigate the potential for tritium trapping by metal vacancies in intermetallic Al 12 (TM) 2.35 phase (TM = Fe, Cr, and Ni) as function of temperature and tritium partial pressure. It was found that tritium could be favorably trapped by Fe and Ni vacancies and not favorably trapped by Al and Cr vacancies. However, due to the presence of partially occupied Al sites in bulk Al 12 (TM) 2.35 , leading to the approximate number of ~255 Al atoms in the unit cell, 86 sites were found energetically favorable to the creation of an Al vacancy. While adding a tritium atom in an Al vacancy is not energetically favorable, the tritiated defect still has a negative Gibbs free energy because the energy gain for creating an Al vacancy overcome the energy cost of adding the tritium species. Based on the calculated Gibbs free energy, the first tritiation of a metal vacancy, at conditions relevant to in-reactor operations, should be more favorable for Al, followed Fe, Ni, and Cr vacancies. By comparing the behavior of tritium in Al 12 (TM) 2.35 with previously studied Fe-Al coating phases (i.e., FeNiAl 5 , Fe 4 Al 13 , and Fe 2 Al 5.6 ), we found that there is a correlation between interstitial tritium solubility and the potential for vacancy trapping. The current trend suggests that if the insertion of an interstitial tritium cost more than 0.3 eV, then trapping by metal vacancies should be preferred. By combining the simulations results obtained to date, we noticed different trapping mechanisms of tritium in the Al coating. Tritium is mostly trapped by Fe and Ni vacancies in the outer Fe-Al coating phase Al 12 (TM) 2.35 while tritium should be preferentially trapped by Al and Fe vacancies for the inner Fe-Al coating phases (FeNiAl 5 , Fe 4 Al 13 , Fe 2 Al 5.6 ). Altogether, these studies show that tritium interacts differently with the various Fe-Al aluminide phases, they also suggest that tritium trapping and retention could be more efficient if metal defects are present and if the solubility of interstitial tritium in the different phases is low.
36 MATERIALS SCIENCE↗
Heterogeneous phase transformation pathways in additively manufactured Al-Ce-Mn alloys
Heat treatment of additively manufactured Al-Ce based multicomponent alloys leads to complex microstructure evolution. In this research, the ability to extend the phase transformation theories involving nucleation of a product phase from a heterogeneous multi-phase microstructure typical to that of additively manufactured samples is explored. The Al-10Ce-8Mn (wt%) was used as a model alloy system. Under additive manufacturing conditions different solidification microstructures were obtained due to spatial and temporal variations of thermal gradients (G) and liquid-solid interface velocities (R) within a given melt pool. Near the melt pool boundary (high G and low R, referred as MPB region), initially, Al 20 Mn 2 Ce forms from the liquid followed by a eutectic of FCC Al and Al 11 Ce 3 . In the melt pool interiors (low G and high R referred as ES region) a eutectic structure between FCC Al and Al 20 Mn 2 Ce is observed. During subsequent heat treatments, the MPB and ES regions transform into different sets of microstructures. In the MPB region, a fine globular microstructure containing FCC Al, Al 11 Ce 3 , Al 6 Mn, and Al 12 Mn results from the decomposition of Al 20 Mn 2 Ce. In the ES region a faceted Al 51 Mn 7 Ce 4 plate phase results from the decomposition of Al 20 Mn 2 Ce. The formation of the Al 51 Mn 7 Ce 4 phase within the eutectic microstructure at the boundaries of FCC Al and Al 20 Mn 2 Ce has not been reported in the literature. Further, these two distinct phase transformation pathways are rationalized based on the role of driving force on the nucleation of (Al 6 Mn) and/or metastable intermetallic (Al 51 Mn 7 Ce 4 ) phases at the interface of aluminum (FCC) and the non-equilibrium intermetallic (Al 20 Mn 2 Ce) phases.
36 MATERIALS SCIENCE↗
The CAI Database: 26 Al– 26 Mg Isotope Systematics
We present a publicly available calcium–aluminum-rich inclusion (CAI) database that focuses on the initial 26 Al/ 27 Al 0 ratio in CAIs, designed in a way that researchers in cosmochemistry and astrophysics may find useful. To date, the database contains 497 CAIs from 75 peer-reviewed papers. The CAIs are from all chondrite groups and cover different CAI types, textures, and sizes. The database includes the paper; the host meteorite; the CAI name and type; the 26 Al/ 27 Al 0 , δ 26 Mg$^*_0$, and δ 25 Mg values and their uncertainties; the number of regression points; the maximum 27 Al/ 24 Mg; the mean-squared weighted deviation; the CAI size; and CAI descriptions. We grouped the CAIs in different ways to discuss 26 Al/ 27 Al 0 ratio distributions with implications for the CAI formation timeline. Overall, we agree with previous authors that CAIs have a bimodal 26 Al distribution: CAIs with robust isochrons (n = 151) have a median 26 Al/ 27 Al 0 = 4.8 × 10 −5 (with a 1σ standard error of 0.1), while those with isotopic anomalies (n = 87) have a median 26 Al/ 27 Al 0 = 0.3 × 10 −5 (with a 1σ standard error of 0.2). However, the large standard deviation of both groups (1.3 and 2.3, respectively) indicates that the 26 Al/ 27 Al 0 values scatter significantly within each population. CAI types and groups can have distinct 26 Al/ 27 Al 0 and δ 26 Mg$^*_0$, but the unmelted inclusions (n = 33) have the highest median 26 Al/ 27 Al 0 = 5.1 × 10 −5 and a low median δ 26 Mg$^*_0$ = −0.05‰. We find slightly different 26 Al/ 27 Al 0 distributions between CAI chondrite types, but no differences between petrographic types or sizes. These observations can help us to understand CAI formation in the context of astrophysical models.
Astronomy and AstroPhysics↗
Combining solution-, precipitation- and load-transfer strengthening in a cast Al-Ce-Mn- Sc -Zr alloy
Here, a cast Al-9Ce-0.75Mn-0.18Sc-0.12Zr (wt%) alloy is designed to combine three strengthening phases: (i) micron-scale Al 11 Ce 3 platelets formed during eutectic solidification, (ii) nano-scale L1 2 -Al 3 (Sc,Zr) precipitates formed during aging, and (iii) Mn in solid solution in the α-Al matrix. Microstructural analyses by SEM, TEM, and atom-probe tomography reveal that Mn remains in solid solution in the as-cast alloy, providing solution strengthening with no influence on the eutectic Al-Al 11 Ce 3 microstructure, which provides precipitation- and load-transfer strengthening. During long-term over-aging at 400 °C, Mn-rich precipitates grow at the Al-Al 11 Ce 3 interface, with no effect on the microhardness. However, after short aging at 350 °C, a high number density of fine L1 2 -Al 3 (Sc,Zr) nanoprecipitates form in the Al matrix (with a coarser size at the Al-Al 11 Ce 3 interface), providing precipitation strengthening. The synergistic combination of the three strengthening mechanisms (solution, precipitation, and load transfer) in our Al-Ce-Mn-Sc-Zr alloy results in higher microhardness after aging at 350 and 400 °C, and higher creep resistance at 300 °C, as compared to alloys with two strengthening mechanisms: an Al-10Ce-0.93Mn control alloy (without precipitation strengthening from Sc and Zr), Al-Ce-Sc-Zr (without solution strengthening from Mn), and Al-Mn-Zr-Er (without load-transfer strengthening from Ce). Furthermore, these dual-strengthened alloys are more creep resistant than alloys with a single strengthening mechanism (Al-Ce, Al-Mn, and Al-Sc-Zr), confirming that the three mechanisms can be combined in pairs or all together.
36 MATERIALS SCIENCE↗
Atomic-Scale Insights into Carbon Dissolution in α-, γ-, and θ-Al 2 O 3 : Phase-dependent Transport Dynamics from First-Principles Calculations
α-Al 2 O 3 exhibits superior carburizing corrosion resistance compared to metastable γ-Al 2 O 3 and θ-Al 2 O 3 phases in high-temperature CO 2 environments, yet its atomic-scale origins remain unclear. Using first-principles density functional theory, we systematically investigate carbon dissolution and diffusion in α-Al 2 O 3 , γ-Al 2 O 3 , and θ-Al 2 O 3 , including the effects of oxygen (O) and aluminum (Al) vacancies. Our results show that α-Al 2 O 3 consistently exhibits higher carbon solution enthalpies than γ-Al 2 O 3 and θ-Al 2 O 3 in both pristine and defective structures, indicating lower intrinsic carbon solubility in α-Al 2 O 3 . Vacancies significantly enhance carbon incorporation: O vacancies reduce solution enthalpy, while Al vacancies further amplify this effect, with a strong preference for carbon at Al vacancy sites. Carbon diffusion barriers are also highest in α-Al 2 O 3 , reflecting slower carbon mobility. Al vacancies increase diffusion barriers across all phases, while O vacancies raise barriers in α- and γ-Al 2 O 3 but slightly lower them in θ-Al 2 O 3 . These results reveal a dual mechanism behind the carburizing resistance of α-Al 2 O 3 : reduced carbon solubility and elevated diffusion barriers. Furthermore, this work provides atomic-scale insights to guide the design of alumina-based materials with improved carburizing resistance through phase selection and defect engineering.
36 MATERIALS SCIENCE↗
Thermodynamic Modeling of the Al-Ce-Cu-Mg-Si System and Its Application to Aluminum-Cerium Alloy Design
Recently discovered AlCe alloys have shown promise in a number of applications, but the propensity of Ce to react with Al and other alloying elements can complicate the phase equilibria and design approach. To solve this, the CALPHAD method is used to explore an alloy within the quinary Al-Ce-Cu-Mg-Si system by developing a thermodynamic database with self-consistent parameters. The database includes a description of all 10 binary systems and 8 ternary systems consisting of: (i) 6 Al-containing ternaries (Al-Ce-Cu, Al-Ce-Mg, Al-Ce-Si, Al-Cu-Mg, Al-Cu-Si and Al-Mg-Si); and (ii) 2 additional ternaries that include Mg and Si (i.e., Ce-Mg-Si and Cu-Mg-Si). The thermodynamic description for the Al-Ce-Mg and Al-Mg-Si systems were reassessed to ensure consistency with the binary systems and the Ce-Mg-Si system is presented for the first time and compared to theoretical data from DFT (Density Functional Theory). In addition to the ternary interactions, the quaternary compound Al3Cu2Mg9Si7 and solid solution extending from the ternary Al2CuMg phase (Al,Si)2CuMg are incorporated. The CALPHAD method is employed and leveraged through the use of a Materials Design Simulator (MDS) to accelerate the design of novel aluminum-cerium-based alloys. The combination of a CALPHAD-based framework with experimental efforts and industrial insight permits the development of three new Al-Ce alloys: Al-3.5Ce-0.4Mg-7Si (Ce-modified A356), Al-5Ce-1Cu-0.5Mg-10Si and Al-19Ce-0.9Mg-1.1Si.
36 MATERIALS SCIENCE↗
Ab Initio Simulations of Tritium Diffusion in Al-Rich Iron Aluminide Coating Phases
Surface coatings of steels used in extreme conditions and corrosive environments generally aim to provide protection and increased durability. In the case of tritium-producing burnable absorber rods (TPBARs) used in nuclear reactors, a 316 stainless steel has been coated with Al. Scanning transmission electron microscopy (STEM) characterization of the coating found three Al-rich (>60 atom % Al) iron aluminide alloys identified as hexagonal FeNiAl 5 , monoclinic Fe 4 Al 13 , and orthorhombic Fe 2 Al 5 . Density functional theory simulations using nudged elastic band have been performed to investigate the diffusion of interstitial tritium in each Al-rich iron aluminide phase. While FeNiAl 5 and Fe 4 Al 13 can be viewed as the stacking of two layers, the structural peculiarity of Fe 2 Al 5 is that channels of variable Al vacancy content are present along the c-axis. Therefore, three stoichiometries for Fe 2 Al x phase, namely, Fe 2 Al 4 , Fe 2 Al 5 , and Fe 2 Al 6 , have been considered to evaluate the impact of Al vacancy concentration on tritium diffusion behavior. Altogether, we found that at 600 K, tritium diffusion decreases from a faster rate in the channels of Fe 2 Al x phases (D T ≤ 10 –11 m 2 ·s –1 ) to Fe 4 Al 13 (D T ≈ 10–12 m2·s–1), and finally in FeNiAl5 (D T ≈ 10 –13 m 2 ·s –1 ). Here, we also find that interstitial tritium generally diffuses faster in Fe–Al coating phases than in the tritium breeding material γ-LiAlO 2 (D T ≈ 10 –14 m 2 ·s –1 ) but slightly slower than in 316 stainless steel (D T ≈ 10 –10 m 2 ·s –1 ).
37 INORGANIC, ORGANIC, PHYSICAL, AND ANALYTICAL CH↗
Improving creep resistance of Al-0.27Zr-0.08Sn (wt%) via cold swaging while maintaining high electrical conductivity
This study investigates the effect of mechanical cold work (0, 50, and 90% cross-sectional area reduction via swaging) on the creep properties of cast aluminum alloys with high electrical conductivity including high-purity Al (HP-Al), Al-0.15Zr (Al-Zr), and Al-0.27Zr-0.08Sn (Al-Zr-Sn, wt%). The Al-Zr alloy is strengthened by Zr in solid solution while the Al-Zr-Sn alloy is additionally strengthened by Al3Zr nanoprecipitates (8 nm diameter). After 90% cold swaging, the alloys exhibit grains elongated along the swaging direction with their widths ranging between 8 and 152 μm perpendicular to the swaging direction. Creep testing at 200 °C shows that the minimum creep rate of all 90% swaged alloys shows high sensitivity to stress, which can be modeled via a threshold stress σth. Increasing swaging magnitude from 0 to 50 to 90% in Al-Zr-Sn improves the creep resistance without reducing electrical conductivity. The formation of subgrains, increased dislocation density, and columnar grain structure in swaged alloys all enhance creep resistance. HP-Al swaged to 90% exhibits much lower creep resistance (σth = 18 MPa) due to subgrain coarsening and the absence of precipitates compared to Al-Zr (σth = 40 MPa) and Al-Zr-Sn (σth = 40 MPa), which maintain stable subgrain structures and benefit from solid solution strengthening and Al3Zr nanoprecipitates, respectively. Although Al3Zr precipitates are known to stabilize the deformed microstructure by pinning grain-boundaries, this work demonstrates that Zr in solid solution alone can effectively stabilize the deformed microstructure resulting in enhanced creep resistance. This effect of Zr solute on stabilizing grain boundaries for creep performance was not previously well-defined in the literature. Despite similar creep performances of Al-Zr and Al-Zr-Sn, the latter shows 70% higher room-temperature microhardness and slightly improved electrical conductivity (56%IACS (International Annealed Copper Standard) vs. 57%IACS, respectively). This work provides new insights into creep mechanisms of cold-worked Al–Zr alloys that will guide the design of heat-resistant Al conductors for high-demand applications.
Coello ramirez, Ismael [Northwestern University,]↗
Structure-Composition Subtleties in NaZn 13 -type Derivatives of Sr/Ca(Au x Al 1– x ) 12–13
In the ternary (Sr/Ca)-Au-Al phase space for 5–8 at % Sr/Ca and nearly equimolar mixture of Au and Al, five new NaZn 13 -derivatives Sr/Ca(Au x Al 1–x ) 12–13 were discovered and their atomic site preferences and electronic structures were studied: (1) at higher Au content, cubic SrAuxAl 13–x [7.24(2) ≥ x ≥ 6.68(2)] shows no obvious atomic site preferences; (2) at lower Au content, tetragonal SrAu x Al 13–x [6.59(1) ≥ x ≥ 6.35(3)] shows icosahedra exclusively centered by Al and preferential atomic arrangements to maximize the number of Au–Al (or Al-rich) shortest distances; (3) a monoclinic SrAu 6.10(3) Al 6.40 compound was uncovered from studies of a single crystal specimen with problematic refinement issues but more detailed powder X-ray diffraction analysis which suggest even more pronounced geometric distortion from the NaZn 13 -type structure as well as preferential arrangements for Au–Al nearest contacts; (4) tetragonal SrAu 5.75(2) Al 6.25 with all empty icosahedra; and (5) orthorhombic CaAu 6.09(2) Al 6.01(1) with icosahedra that are half empty and half partially occupied by Al. Electronic DOS and COHP curves were used to rationalize the structural depiction of Sr/Ca(Au x Al 1–x ) 12–13 as icosahedra packing of (Au x Al 1–x) 12–13 with voids filled by the electropositive Sr/Ca.
37 INORGANIC, ORGANIC, PHYSICAL, AND ANALYTICAL CH↗