DEFORMABLE HIGH-STRENGTH ALUMINUM ALLOY COMPOSITIONS AND METHODS OF MAKING THE SAME

CROSS-REFERENCE TO RELATED APPLICATIONS

[0001]The present U.S. patent application is related to and claims the priority benefit of U.S. Provisional Patent Application Ser. No. 63/612,129, filed Dec. 19, 2023, the contents of which are hereby incorporated by reference in their entirety into the present disclosure.

STATEMENT REGARDING GOVERNMENT FUNDING

[0002]This invention was made with government support under DMR 2210152 awarded by the National Science Foundation, under N00014-17-1-2921 and N00014-20-1-2043 awarded by Office of Naval Research. The government has certain rights in the invention.

TECHNICAL FIELD

[0003]The present disclosure generally relates to aluminum alloy compositions of high strength and high deformability and containing medium entropy metallic lamella. Additive manufacturing methods for fabrication of the alloy are also disclosed.

BACKGROUND

[0004]This section introduces aspects that may help facilitate a better understanding of the disclosure. Accordingly, these statements are to be read in this light and are not to be understood as admissions about what is or is not prior art.

[0005]Aluminum (Al) alloys are widely utilized as structural materials in aerospace and automobile industries. To fulfill the complex geometrical constraints for industrial applications, selective laser melting (SLM) has been increasingly used to fabricate parts of Al alloys with a high degree of freedom for complex design and integration. Most existing studies have been conducted mainly for near-eutectic Al—Si and Al—Si—Mg alloys. These alloys exhibit medium strength but great hot-tearing resistance, making them good candidates for 3D printing2,4,5. In contrast, high-strength Al alloys, such as Al60616 and Al70757, are inherently susceptible to hot cracking during additive manufacturing process.

[0006]One method to alleviate hot cracking during additive manufacturing of high-strength Al alloys is to introduce fine and hard particles. These particles can be introduced via external inoculation, e.g. TiN, TiC, TiB₂ or aging, e.g. Al₃Zr, Al₃Sc, Al₂Cu. They can strengthen the Al alloy by impeding dislocation movements. Meanwhile, these particles promote heterogeneous nucleation, and break down columnar grains where intergranular cracks are prone to initiate and propagate. In spite of these studies, the highest strength achieved in additively manufactured (AM) Al alloys remain in the range of 300-500 MPa. There is scattered success in producing high strength Al alloys via severe plastic deformation, such as high-pressure torsion and accumulative roll-bonding, or cryo-milling followed by powder consolidation. The high strength in these cases arises from significant grain refinement to nanoscales. Ultra-strong AM Al alloys with high flow strength and deformability remain to be discovered.

[0007]Transition metal (TM) intermetallics, such as Al—Fe, Al—Co and Al—Ni are largely avoided in AM Al alloys as prior experience in casting shows that the addition of TM elements often lead to large and brittle intermetallics. These intermetallics, such as Al₉Co₂, Al₁₃Fe₄ have crystal structures with low symmetry (monoclinic) and thus are known to be brittle materials at room temperature.

[0008]Due to the factors mentioned above, a need exists for ultrastrong deformable aluminum alloys.

SUMMARY

[0009]An alloy containing 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel is disclosed. A technical effect of alloys of this disclosure is that they possess high strength and high deformability.

[0010]A method of making an alloy is disclosed. The method includes the steps of: (1) providing particles wherein each particle has a composition 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel; (2) utilizing a selective leaser melting (SLM) apparatus producing a first layer of the particles on a substrate and melting and solidifying a first group selected areas of the layer of particles, wherein the melting and the solidification results in an alloy of composition 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel, containing intermetallic lamellae of compositions Al₉(Fe,Co,Ni)₂ and Al₃Ti wherein the alloy formed has thickness equal to thickness of the first layer; (3) repeating utilization of SLM apparatus to produce a second layer of the particles and laser melting and solidification of a second group of selected areas of the second layer of the particles, wherein the second group of selected areas is coincident or in contact with the first group of selected areas wherein the melting and the solidification results in an alloy of composition 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel, containing intermetallic lamellae of compositions Al₉(Fe,Co,Ni)₂ and Al₃Ti wherein the alloy formed has thickness equal to thickness of the first layer; and (4) repeating the utilization of SLM apparatus to produce an object of specified thickness and shape containing the melted and solidified areas.

BRIEF DESCRIPTION OF DRAWINGS

[0011]While some of the figures shown herein may have been generated from scaled drawings or from photographs that are scalable, it is understood that such relative scaling within a figure are by way of example, and are not to be construed as limiting.

[0012]FIG. 1A shows the back scattered SEM image showing the horizontal (XY) view of microstructure of the as-printed Al₉₂Ti₂Fe₂Co₂Ni₂ alloy with 300 W laser.

[0013]FIG. 1B shows back scattered SEM image showing the vertical (XZ) planes view of the microstructure of the as-printed Al₉₂Ti₂Fe₂Co₂Ni₂ alloy with 300 W laser, with the melt pool (MP) boundaries outlined.

[0014]FIG. 1C is a micrograph of a representative melt pool with the outlined boundary and striated fine rosettes.

[0015]FIG. 1D shows microstructure across the melt pool boundaries showing distinctive features in the coarse rosettes region and fine rosettes region.

[0016]FIGS. 1E and 1F show Magnified micrographs of the coarse and fine rosette regions respectively with arrows indicating cellular precipitates.

[0017]FIGS. 2A and 2B show STEM HAADF images of representative morphologies of fine rosettes region and coarse rosettes region respectively and corresponding EDS composition maps of various elements.

[0018]FIGS. 2C and 2D are line scans showing relevant phase constituents of FIGS. 2A and 2B respectively.

[0019]FIG. 3A shows an overview bright field (BF) TEM image showing three major phases in the alloy, Al matrix, Al₉(Fe,Co,Ni)₂ and Al₃Ti.

[0020]FIG. 3B shows Inverse pole figure mapping of Al/Al₃Ti phases by ASTAR (nano-EBSD).

[0021]FIG. 3C shows High-resolution TEM (HRTEM) with the corresponding Fast Fourier Transform (FFT) of the image of Al₃Ti intermetallics showing its highly defective nature.

[0022]FIG. 3D shows an inverse FFT of the image with (002) plane filtered indicating the presence of abundant dislocations in the nanoscale Al₃Ti intermetallics.

[0023]FIGS. 3E and 3F show Identical structures obtained from VESTA crystallographic visualization of its prototype Al₉Co₂ with monolithic crystal structure.

[0024]FIG. 3G shows HRTEM micrograph of the interface between two genres of intermetallics.

[0025]FIG. 3H shows selected area electron diffraction (SAED) pattern corresponding to FIG. 3G indicating the orientation relationships.

[0026]FIG. 4A shows TEM micrograph of the atom probe tip taken prior to APT acquisition demonstrating three Al grains.

[0027]FIG. 4B shows atomic maps demonstrating uneven saturation of solute elements.

[0028]FIG. 4C shows 1D concentration profile capturing the behavior of the solute elements across two grains: Ni atoms demonstrate segregation at the grain boundary, while Ti atoms show uneven distribution between grains.

[0029]FIG. 4D shows atomic maps of an interphase plane between the Al grain and the monoclinic Al₉(Fe,Co,Ni)₂ phase.

[0030]FIG. 4E shows radial Distribution Function taken from the inset in (d) highlighting the lack of clustering in the bulk monoclinic phase.

[0031]FIG. 4F shows 1D concentration profile taken across the interphase boundary. (A spike in the middle highlights some segregation near the boundary prior to a near equiatomic composition of Fe, Co, and Ni in the monoclinic phase.)

[0032]FIG. 5A shows a back scattered SEM micrograph on the as-printed Al₉₂Ti₂Fe₂Co₂Ni₂ after nanoindentation hardness measurements.

[0033]FIGS. 5B and 5C show hardness and Young's modulus contour maps respectively reconstructed from series of nanoindentation, revealing a relatively high hardness near the melt pool boundaries with finer microstructure.

[0034]FIG. 6A shows engineering stress-engineering strain curves for bulk compression tests performed on the as-printed Al₉₂Ti₂Fe₂Co₂Ni₂ pillars with schematic diagrams showing the geometry of specimen and the typical barreling phenomenon after compression for one of the deformed specimens (200 W). Numbers in the legend denote laser power.

[0035]FIG. 6B shows true stress-true strain curves superimposed with working hardening rates showing 7% uniform compression in the early stages of deformation.

[0036]FIG. 7A shows true stress-true strain curves for in situ micropillar compression tests on both fine and coarse rosettes regions of samples printed with 300 W laser tested at room temperature. The flow stress on fine rosette region could reach 1 GPa.

[0037]FIG. 7B shows screenshots of morphological evolutions for pillars under compression. Arrows indicate the formation of shear planes.

[0038]FIG. 7C shows back stress measurements vs. strain curves for micropillars with coarse and fine rosettes.

[0039]FIG. 7D shows a stress-strain hysteresis loop for the pillar in fine rosettes region at 14% strain. Back stress, yield points for unloading and loading curves are labeled for illustration.

[0040]FIGS. 8A, 8B, 8C and 8D show STEM and selective EDS mapping on the post-deformation pillar in the coarse rosette region.

[0041]FIGS. 8E and 8F show BF and weak-beam DF TEM micrographs respectively of the Al grain revealing high-density dislocations in the vicinity of cellular walls.

[0042]FIG. 9A shows an overview TEM micrograph on the post-deformation pillar in the fine rosette region. (Arrows indicate the deformation bands).

[0043]FIG. 9B shows a TEM micrograph demonstrating the morphology of fine intermetallic rosettes. Phases are labelled Al₃Ti and Al₉(Fe,Co,Ni)₂, and EDS maps were presented. Arrows indicate the abundant stacking faults in the monoclinic Al₉(Fe,Co,Ni)₂. The insert shows SAED pattern along Al₉(Fe,Co,Ni)₂ [120] zone with the significant crystal rotation by plastic strain.

[0044]FIG. 9C shows a TEM image for Al₉(Fe,Co,Ni)₂ outlined near the pillar top in FIG. 9B with the corresponding FFT pattern. Stacking faults and crystal rotation are observed.

[0045]FIG. 9D shows a HRTEM micrograph for the Al₉(Fe,Co,Ni)₂ particle in FIG. 9B with planar defects. The corresponding FFT pattern reveals stacking faults habit plane is (100).

[0046]FIG. 9E shows an inverse FFT image masking diffraction spots circled shows lattice distortion around stacking faults along with some additional half planes, indicating the existence of dislocations.

[0047]FIG. 9F shows an HRTEM micrograph for the Al₉(Fe,Co,Ni)₂ particle in FIG. 9B slightly off the pillar top with planar defects. The indexed FFT pattern along [110] zone shows diffraction spots originating from defects circled in pink. Solid line segments help to visualize the direction change of crystal planes.

[0048]FIG. 9G shows an HRTEM micrograph on the end of a stacking fault ribbon.

[0049]FIG. 9H shows an inverse FFT image masking the extra spots (in the inserted FFT) confirming they result from the stacking fault ribbon.

[0050]FIG. 9I shows an HRTEM image for Al₉(Fe,Co,Ni)₂ showing a low-angle grain boundary (LAGB) ˜5° close to (110) plane where some atomic disorder is observed.

DETAILED DESCRIPTION

[0051]For the purposes of promoting an understanding of the principles of the disclosure, reference will now be made to the embodiments illustrated in the drawings and specific language will be used to describe the same. It will nevertheless be understood that no limitation of the scope of the disclosure is thereby intended, such alterations and further modifications in the illustrated device, and such further applications of the principles of the disclosure as illustrated therein being contemplated as would normally occur to one skilled in the art to which the disclosure relates.

[0052]In this disclosure are presented compositions and methods are presented produce intermetallics-strengthened Additive Manufactured (AM) Al alloys utilizing transition metals including Co, Fe, Ni and Ti. Colonies of nanoscale intermetallics lamellae aggregate into fine rosettes and give rise to a high strength, exceeding 900 MPa, with prominent plastic deformability under compression. Heterogeneous microstructure also introduced significant back stress. Surprisingly, complex dislocation structures and stacking faults were present in the sandwiched monoclinic brittle Al₉(Fe,Co,Ni)₂ phase. This study demonstrates an effective strategy to develop ultra-high strength AM Al alloys via nanoscale laminated deformable intermetallics. In this disclosure AM is used to represent Additive Manufacturing or Additively Manufactured based on the context. Further high-strength and ultra-high-strength are phrases and adjectives to denote strength higher than previously obtained with similar deformability.

[0053]The following methods, structural characterization and mechanical testing were employed in experiments leading to this disclosure.

[0054]Methods: Powder processing and manufacturing. Spherical powder with a nominal composition of Al₉₂Ti₂Fe₂Co₂Ni₂ (at. %) satisfying −53+15 μm were gas atomized by Atlantic Equipment Engineering, Inc. Additive manufacturing was performed by using a laser powder bed fusion (LPBF) instrument, SLM 125 HL metal 3D printer in Argon atmosphere with the oxygen level below 1000 PPM. Printing was conducted by utilizing a 400 W IPG fiber laser (λ=1070 nm) with a laser power of 200-300 W, a scan speed of 1200 mm/s, a hatch space of 100 μm, a layer thickness of 30 μm and a laser spot of 70 μm in diameter. Build plate was preheated to 200° C. and each layer rotated by 67°. Cylindrical samples with height 12 mm and diameter 6 mm were fabricated for bulk compression tests. Cubic samples with dimensions 10×10×5 mm were printed for microstructure characterization, nanoindentation and micropillar compression tests.

[0055]Structural characterization: The microstructure of Al alloy was investigated by X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM) and atom probe tomography (APT). Samples were mechanically grinded and polished down to 1 μm diamond paste. XRD was performed on a PANalytical Empyrean X'pert PRO MRD diffractometer with a 2×Ge (220) hybrid monochromator to select Cu K_α1 in the 2θ-ω geometrical configuration. Scanning electron microscopy (SEM) experiments were performed by using a Thermo Fischer Quanta™ 3D and Teneo™ high-resolution Field Emission SEM microscopes with a back scattering detector operated at 30 kV. A Thermo Fisher Talos 200×TEM microscope with an acceleration voltage of 200 kV was utilized to capture bright field (BF), dark field (DF), scanning transmission electron microscopy (STEM) images, and Energy dispersive spectrometry (EDS) maps. Crystal orientation mapping was performed by using a NanoMEGAS detector. APT Samples were prepared using standard focused ion beam (FIB) lift-out procedures on a Scios 2 DualBeam FIB/SEM, followed by a series of annular milling steps with decreasing radii to achieve a tip radius of approximately 50 nm. Atom probe data were collected on a CAMECA LEAP 5000×S APT, using both voltage and laser mode acquisition. For the former, a pulse fraction of 20%, temperature of 50 K, and a pulse rate of 200 kHZ were employed. For the latter, similar values for temperature and pulse rate were employed, with a laser pulse energy of 80 pJ to ensure complete field ion evaporation. Data reconstruction and analyses were conducted using AP Suite 6.1 software.

[0056]Mechanical testing. Nanoindentation experiments were performed with a Bruker's Hysitron TI Premier nanoindenter with a Berkovich tip under displacement-control mode at 800 nm depth on well-polished samples. Hardness information was assessed from an area of 100×100 μm² covering representative microscale features with 121 indents with 10 μm spacing in both dimensions. Progressive indentation with multiple continuous loading-unloading segments at incremental penetration depths were conducted for each indentation. Hardness and Young's modulus were determined from an average of 10 measurements. Bulk compression tests were performed on an MTS framework with a 30 kN load cell and a strain rate of 10⁻³ s⁻¹ after polishing and leveling the top and bottom surfaces of as-printed cylindrical samples for better alignment. In situ micropillar compression tests were performed in the Quanta™ 3D SEM microscope equipped with a Hysitron PI 88×R PicoIndenter and a real-time video recorder. Micropillars were produced by FIB, with the height of 10 μm, the diameter of 5 μm, and an aspect ratio of 2:1. Both 10 and 20 μm diamond flat-punch tips were used and strain rate was set as 5×10⁻³ s⁻¹. An average drift rate of 0.2-0.6 nm/s was determined for displacement correction.

[0057]Microstructural characterization. Back scattered SEM images reveal the microstructure of the as-printed Al₉₂Ti₂Fe₂Co₂Ni₂ fabricated with 300 W laser power in FIGS. 1A through 1F. Morphologies of inter-woven laser tracks and inverse-parabolic melt pool cross sections typical in SLM-processed alloys are evident on horizontal XY plane (FIG. 1A)) and vertical XZ plane (FIG. 1B), where Z axis is the build direction. Melt pools outlined by dash lines are around 120 μm in width and 80 μm in depth, with some inherent variation due to layer rotations. FIG. 1C shows a gradient heterogeneous microstructure in the melt pools. Colonies of layered aggregates (referred to as rosettes) shown in light contrast are intermingled with the Al-rich matrix (in dark contrast). Referring to FIG. 1D, rosettes with finer laminate spacing (fine rosettes) are the dominant features near the melt pool boundaries, while rosettes with thicker lamellae (coarse rosettes) are abundant in the melt pool center. In addition, there are also some fine rosettes in the melt pool center arranged in a striated fashion. Referring to FIGS. 1E and 1F, high-magnification SEM micrographs in FIGS. 1E and 1F show that the melt pool center contains coarse rosette region (36 vol. %), nanoscale cellular precipitates (4 vol. %) denoted by arrows and Al rich matrix (60 vol. %). In contrast, near the melt pool boundaries, fine rosettes (97 vol. %) are separated by thin layers of Al matrix (3 vol. %).

[0058]Representative fine rosettes and cellular precipitates in coarse rosette region were characterized by STEM and EDS in FIGS. 2A through 2C. The fine rosettes as shown in FIGS. 2A and 2C have Al₃Ti cores surrounded by two alternating intermetallics laminates, Al₃Ti and medium entropy Al₉(Fe,Co,Ni)₂ intermetallics, with a laminate thickness of 20-60 nm. The Al₃Ti layers are thinner than Al₉(Fe,Co,Ni)₂ in the laminates. The chemical compositions for coarse rosettes are nearly identical to fine rosettes with a laminate thickness of 150-300 nm. The coarse-rosette region also contains cellular boundaries enriched in Al₉(Fe,Co,Ni)₂ as shown in FIGS. 2B and 2D.

[0059]Detailed microstructure examinations of the same (300 W) specimen using TEM and STEM are summarized in FIGS. 3A through 3H. The BF TEM image if FIG. 3A shows the coexistence of three major phases, Al matrix, Al₉(Fe,Co,Ni)₂, and Al₃Ti. XRD patter and selected area electron diffraction (SAED) of the TEM image suggest the Al₃Ti exists mainly in D0₂₂ phase (space group 14/mmm, a=0.384 nm, c=0.860 nm) and partially in L1₂ phase (space group Pm3m, a=3.98 nm³⁰), and Al₉(Fe,Co,Ni)₂ has monoclinic structure with the prototype of Al₉Co₂ (space group P2₁/c, a=0.622 nm, b=0.629 nm, c=8.559 nm, β=94.8°³¹) or Al₉FeNi. Besides, it could be seen that plenty of defects exist in Al₃Ti (FIG. 3a & FIG. S3a). The inverse pole figure (IPF) map FIG. 3B acquired from high-resolution ASTAR orientation mapping demonstrates the crystallographic orientation of Al and Al₃Ti. The polycrystalline composite has colonies with dimension of ˜1 μm. HRTEM image (FIG. 3C) along Al₃Ti D0₂₂ [100] zone axis demonstrates the lattice distortion induced by the 2 nm scale scattered patches. An inverse Fast Fourier Transform (FFT) image based on (002) diffraction shows plenty of dislocations residing in the disordered lattice patches as shown in FIG. 3D. In comparison, the Al₉(Fe,Co,Ni)₂ phase possesses few defects as shown in FIG. 3A. HRSTEM (FIG. 3E) shows the atomic arrangements of the medium-entropy Al₉(Fe,Co,Ni)₂ along [110] zone axis with blue dots representing sites of Co atoms in its prototype. The micrograph is consistent with the simulated 3×3×3 cells by VESTA (FIG. 3F) It is difficult to distinguish Fe, Co, Ni atoms in the current HRSTEM micrograph due to their close proximity in atomic number. FIG. 3G depicts the typical well-defined interface between Al₉(Fe,Co,Ni)₂ and Al₃Ti. The SAED pattern (FIG. 3H) illustrates the crystallographic orientation relationships between the two phases, where [130]_Al3Ti/[100]_{Al9(Fe,Co,Ni)2}, (002)_Al3Ti/(001)_{Al9(Fe,Co,Ni)2}, (310_{) Al3Ti}/(010)_{Al9(Fe,Co,Ni)2} and their planar spacing has the following match ratio:

$d_{{\mathrm{Al}}_3\mathrm{Ti}}^{\left(002\right)}:d_{{{\mathrm{Al}}_9\left(\mathrm{Fe},\mathrm{Co},\mathrm{Ni}\right)}_2}^{\left(001\right)}\approx 2:1\mathrm{and}{d}_{{\mathrm{Al}}_3\mathrm{Ti}}^{\left(310\right)}:d_{{{\mathrm{Al}}_9\left(\mathrm{Fe},\mathrm{Co},\mathrm{Ni}\right)}_2}^{\left(010\right)}\approx 5:1.$

[0060]FIG. 4A corresponds to a TEM image of three Al matrix grains from a fine rosette region to gauge the microstructure before atom probe acquisition. Referring to FIG. 4B, atomic maps for individual elements show that there are some inhomogeneities in the distribution of the TM solute elements. This variation is highlighted in the 1D composition distribution FIG. 4C where the profile follows the white arrow in FIG. 4A. Ti concentration increases (from 0.1 to 0.5 at % towards the top grain with a decrease in Ni concentration concomitantly. Grain boundaries appear to enrich in Fe, Co and Ni. FIG. 4D corresponds to a dual-phase interface between Al and Al₉(Fe,Co,Ni)₂ taken from another area. A relatively small amount of TM solutes are present in the matrix, specifically, 0.20% Ti, 0.30% Fe, 0.03% Co and 0.09% Ni (at %). Ti is highly rejected by the intermetallic phase and present in the matrix. The 1D composition distribution profile in FIG. 4F following arrow FIG. 4C shows minute TM solutes in matrix, vs. nearly equiatomic distribution of Fe, Co and Ni (78.06% Al, 0.06% Ti, 5.45% Fe, 7.53% Co and 8.90% Ni (at %)) in the intermetallic phase. To understand the interactions between the alloying elements in the brittle intermetallic region of the tip, partial radial distribution functions (FIG. 4E) were computed (in the area highlighted by the rectangle). All profiles remain close to 1.0, the value for random distribution, signifying well-mixed TM solutes in lattice.

[0061]Mechanical properties. In order to assess the mechanical properties of heterogenous AM Al alloys, nanoindentation experiments were conducted over a representative melt pool (FIG. 5A). Hardness contour reconstructed from nanoindentation map (FIG. 5B) shows a majority of the region has a high hardness ranging from 2.5 to 4.5 GPa. Melt pool boundaries often have a higher hardness (3.5-4.5 GPa), whereas the melt pool interior has a lower hardness (2.5-3.0 GPa). Referring to FIG. 5C, a similar trend for Young's modulus is observed. Hard melt pool boundaries are associated with a relatively high Young's modulus (140-150 GPa), while melt pool interiors have a relatively low Young's modulus (130-140 GPa).

[0062]Bulk compression tests were performed on cylinders with dimensions of 6×12 mm, fabricated at various laser power. FIG. 6A shows engineering stress-engineering strain curves for bulk compression tests performed on the as-printed Al₉₂Ti₂Fe₂Co₂Ni₂ pillars with schematic diagrams showing the geometry of specimen and the typical barreling phenomenon after compression for one of the deformed specimens (200 W). Numbers in the legend denote laser power. Specimens printed with 200 W laser exhibit an ultra-high stress exceeding 900 MPa concurring with substantial plastic deformability around 20%. The inserted optical micrographs in FIG. 6A reflect the typical barreling phenomenon for ductile metallic materials. At higher laser energy, 250 W and 300 W, the flow stress of pillar decreases to 550 MPa with compressive strain of 5-20%. FIG. 6B presents analyses on plastic instability by using Considère's criterion. The superimpositions of selective work hardening curves over true stress-true strain curves imply uniform deformation strain at 7%.

[0063]To probe the influence of heterogeneous microstructures on mechanical behavior of the AM Al alloys, micropillar compression tests were carried out over coarse and fine rosettes regions. As shown in FIG. 7A, fine rosette region can sometimes possess an exceptionally high strength around 1 GPa, while coarse rosette region has a flow stress of 500 MPa. In situ SEM micrographs in FIG. 7B (collected from supplementary videos) reveal the morphological evolution for representative pillars from the two distinct regions. Particles were readily visible on the pillar surfaces in the coarse rosette region, while a wrinkled and wavy surface shows less apparent protrusion in the fine rosette region. For the coarse rosettes, the evident particle extrusions to the outer surfaces manifest the preferential plastic flow in the soft aluminum matrix and insignificant plasticity in the hard intermetallics precipitates. Whereas the retention of cylindrical shape of pillars in the fine rosettes implies uniform co-deformation of matrix and precipitates to constrain the generation of localized shear bands. FIG. 7C shows back stress measurements vs. strain curves for micropillars with coarse and fine rosettes. FIG. 7D shows a stress-strain hysteresis loop for the pillar in fine rosettes region at 14% strain. Back stress, yield points for unloading and loading curves are labeled for illustration. The multiple loading-unloading experiments display prominent hysteresis loops evident from the stress-strain curves. The implication of such loops on deformation mechanisms will be discussed later in this description.

[0064]Post-deformation microstructure analyses: To better understand the deformation mechanism of AM Al alloys with intermetallic rosettes, cross-section TEM (XTEM) samples from post-deformation micropillars were investigated. Referring to FIG. 8A, in the coarse rosette region, some microcracks are witnessed in intermetallics in STEM image. EDS elemental maps are provided in FIGS. 8B through 8D to identify composition in these intermetallics precipitates. Microcracks (labeled in dash circles) appear in both Al₃Ti and Al₉(Fe,Co,Ni)₂ intermetallics. It is seen that these microcracks are typically constrained in single intermetallic phase, not extending into the neighboring phases. FIGS. 8E through 8F show abundant dislocation activities in the vicinity of cellular walls in Al matrix.

[0065]Referring to FIG. 9A, in the fine rosette region, the dilation of the pillar top and shear bands were observed. Apart from dislocation activities and microcracks, significant crystallographic rotation (as evidenced by the inserted SAED pattern) and some planar defects were observed in the monoclinic Al₉(Fe,Co,Ni)₂ phase (FIG. 9B). FIG. 9C shows the Al₉(Fe,Co,Ni)₂ precipitate has an evident curvature with planar defects and grain rotation on the pillar top, where the deformation is the most intense. HRTEM micrograph of the Al₉(Fe,Co,Ni)₂ show abundant stacking faults (FIG. 9D and streaks in the inserted FFT pattern suggest the habit plane for these SFs is (001). The related inverse FFT image masking two brightest diffraction spots reveals dislocations aligned along the SFs (FIG. 9E). Another HRTEM micrograph (FIG. 9F) is acquired for Al₉(Fe,Co,Ni)₂ at 400 nm underneath the top surface, where the deformation is less severe. The lattice arrangement is disturbed by (001) SFs along [110] zone. Streaks in the indexed FFT pattern are pronounced and extra spots circled in pink are identified. FIG. 9G depicts the ends of a SF ribbon. The corresponding inverse FFT image in FIG. 9H confirms the additional spots originate from the faulted region, which might suggest a possible phase transformation, as the interplanar spacing has no matching with the parent phase. FIG. 9I demonstrates a low-angle grain boundary (˜5° within Al₉(Fe,Co,Ni)₂. The boundary is roughly on (110) where some local disorder is present.

[0066]A highly heterogeneous microstructure composed of coarse rosettes, fine rosettes and cellular Al matrix was observed in the AM Al alloy. To further understand the formation of rosettes, equilibrium phase constitution was calculated by Thermo-calc using TCAL8 database The calculation suggests that Al₃Ti forms the cores of intermetallic rosettes due to its high melting temperature, providing nucleation sites for co-precipitation of Al₉(Fe,Co,Ni)₂. The rapid cooling rate significantly refines precipitates, whereas traditional casting of transition-metal-bearing Al alloys often leads to overgrown large precipitates and hence, embrittlement. The morphology distinctions for coarse and fine rosettes are attributed to the complex and location-specific thermal history with respect to melt pools. Sufficient supplies of TM solutes and a higher quenching rate adjacent to melt pool boundaries can enable the precipitation of a greater volume fraction of intermetallics with finer lamellae, compared to the coarse rosettes that dominate melt pool center. The arrangement of alternating fine and coarse rosettes region in this alloy results from periodic thermal cycles during layer-wise construction. Additional refinement is realized by the striated precipitation possibly due to the turbulent Marangoni flow. Marangoni flow stirs fine rosettes owing to the complex thermal gradient, varying surface tension and dynamic hydromechanics. Besides, Ti has been reported to be effective to refine microstructures of Al alloys. The rosette structure was reported in other Al alloys where Ce and Mn were introduced for precipitate strengthening, yet the resultant deformation mechanisms remain unexplored.

[0067]The high quenching rate characteristic of laser fusion not only refines the microstructure but also has profound impact on the formation of various non-equilibrium phases. First, L1₂ Al₃Ti is often unstable and will spontaneously transform to equilibrium D0₂₂ Al₃Ti³⁶. But the rapid solidification process retains some L1₂ Al₃Ti by not allowing atoms sufficient time for complete ordering. The cruciform geometry of Al₃Ti core in FIG. 2A correlates well with literature report on typical morphology characteristics of trialuminides. L1₂ Al₃Ti can also be fabricated via mechanical alloying with or without a ternary element. It is generally accepted that L1₂ Al₃Ti shall be more deformable than its D0₂₂ counterpart as the former has more independent slip systems rendered by a cubic crystal structure. It is also speculated that the high cooling rate establishes significant defects in both phases. Second, a partitioned medium-entropy intermetallic phase Al₉(Fe,Co,Ni)₂ was maintained, while at equilibrium it shall decompose into two isomorphic monoclinic phases (Al₉Co₂+Al₉FeNi). The supersaturated Fe, Ni atoms in Al₉(Fe,Co,Ni)₂ distort the lattice locally and thus change its mechanical behavior. Preservation of these metastable phases would play a significant role in deformation mechanisms of the AM Al alloys.

[0068]Mechanical behavior of high strength AM Alloys: High strength of the current AM Al alloy is confirmed by multiple experiments. This alloy exhibits over 900 MPa engineering stress from macroscale compression tests. Micropillar compression tests show that the fine rosette region could reach 1 GPa true flow stress or an engineering stress of 1.18 GPa. An estimation based on the rule of mixture is attempted. Hardness assessments from nanoindentation mapping show values of 2.5-4.5 GPa. The flow stress derived from the empirical Tabor's relation,

$\sigma =\frac{\mathrm{Hardness}}{3},$

varies from 800 MPa to 1.5 GPa. The variations of hardness values across melt pools arise from the heterogenous microstructures containing coarse rosettes in melt pool center and fine rosettes near melt pool boundaries. The current Al₉₂Ti₂Fe₂Co₂Ni₂ has an excellent combination of mechanical strength and plastic strain under compression, compared with other AM Al alloys.

[0069]Next, we consider the related strengthening mechanisms leading to the ultrahigh mechanical strength in our AM Al alloys, including solid solution strengthening and Orowan strengthening, Hall-Patch strengthening, dislocation strengthening, and hetero-deformation induced (HDI) strengthening. Solid solution strengthening can be ignored as the accumulative solubility of TM solutes in Al (though in supersaturated state) is very low, <1 at % based on EDS measurements (FIG. 2) and APT studies (FIG. 4C)). Dislocation strengthening from statistically stored dislocations (SSDs) is insignificant as dislocation density in Al matrix is minute in the as-printed state (Build plate heating at 200° C. would remove a good number of dislocations). Orowan strengthening is also secondary as the morphology of dislocations are not those typical to Orowan strengthening, which would be manifested by dislocation loops encircling intra-granular precipitates. In the current AM Al alloy, dislocation entanglements are primarily identified against inter-granular Al₉(Fe,Co,Ni)₂ cellular precipitates in the post-deformation TEM samples (FIGS. 8E and 8F).

[0070]HDI strengthening has been observed in heterogenous materials and can provide back stress and work hardening ability in metallic materials. Significant stress-strain hysteresis loops were observed during micropillar compression tests (FIG. 7A). The evolution of accumulated back stress with progressing strain in FIG. 7C (calculated using method shown in FIG. 7D reveals that the fine rosette regions with a higher flow strength carries a very high back stress, ˜600 MPa, compared to the back stress of ˜250 MPa in the coarse rosette regions. Back stress is the long-range stress component typically ascribed to the pileups of geometrically necessary dislocations (GNDs) in materials with heterogeneity or gradient structure. These GNDs could arise from mismatch of coefficients of thermal expansion (CTE) between Al and intermetallics during solidification, and strain incompatibility across interfaces of hard and soft phases during compression. It's worth mentioning that HDI strengthening is the major component leading to Bauschinger effect as GND pileups near interfaces have reversible dislocation configuration during loading-unloading processes. TEM study reveals ample GNDs in Al matrix shown in FIGS. 8E and 8F. Rigid intermetallics tend to remain elastically deformed while Al matrix sustains high-density dislocations near the interfaces to accommodate strain field gradient. Existing GNDs will hamper further dislocation motion and also trigger forward stress in the hard intermetallic phases across interfaces. Back stress and forward stress will collectively harden the material. As the deformation progresses, back stress rises quickly and reaches a plateau for both regions. This trend can be explained by the absorption of GNDs by the interfaces after straining to a critical value. At certain strain levels, dynamic generation and annihilation of GNDs reach an equilibrium state, leading to a saturated strengthening contribution from HDI stress. The strain gradient carried by absorbed GNDs, though decoupled from strengthening, will trigger Al-intermetallics interface debonding in microscale and eventually fracture at macroscale.

[0071]Apart from HDI stress from Al-intermetallic interface, there might be HDI stress originating from the interfaces between two genres of intermetallics Al₉(Fe,Co,Ni)₂ and D0₂₂-Al₃Ti. Prior study suggests both intermetallics phases are brittle at room temperature. This assertion is especially applicable to Al₉(Fe,Co,Ni)₂ inferred from its monoclinic crystal structure with low symmetry. However, abundant SFs and dislocations were surprisingly observed in the deformed nanoscale monolithic Al₉(Fe,Co,Ni)₂ (FIGS. 9G through 9I), suggesting unique plastic deformation mechanism in the often-brittle intermetallics. This study presents what may be the first experimental evidence of plasticity in monoclinic Al₉(Fe,Co,Ni)₂ medium entropy intermetallics. Scattered prior study suggests that complex metallic compound could deform by introducing metadislocations, which rarely exist in high symmetry metallic materials, and metadislocations have been reported in Al₁₃Co₄ with monoclinic crystal structure. There are dislocations in Al₃Ti in the as-printed state (FIGS. 3a, 3B and 3C). Dislocations can also carry out plastic flow in Al₃Ti to ensure co-deformation between the two types of intermetallics across the laminated intermetallics interfaces in the fine rosette region (FIG. 2A and FIGS. 9A through 9!). Deformability and deformation mechanisms for nanoscale sandwiched intermetallic branch could differ from their bulk counterpart, due to the discrepancy in scale and confined loading state, as corroborated by studies on laminated nanolayers. Comparing to transient and reversible GND pileups adjacent to Al-intermetallic interfaces, temporary defects configuration existing in intermetallic-intermetallic interface could generate back stress as well. For nanolaminated intermetallics, strain transfer into the monolithic Al₉(Fe,Co,Ni)₂ could be very challenging, which necessitates a large back stress to drive defect activity. Under such a large back stress, SFs or other defects may be activated in intermetallic phases. Due to the limited plasticity of intermetallics, HDI stress stemming from intermetallic interfaces would increase rapidly during the initial loading process and remain saturated after plastic relaxation, which is consistent with experiment observations of back stress saturation for both regions (FIG. 7C)

[0072]In the context of strength-ductility paradox for most metallic materials, it is surprising to achieve 20% plasticity in these high-strength AM alloys as shown from both macropillar and micropillar compression tests. First, the Al matrix accommodates a majority of plastic strain as verified by dislocations in Al in the deformed pillars. Second, the back stress from heterogeneous interfaces sustains significant work hardening. As discussed earlier, SFs and other defects have been observed in deformed nanoscale intermetallics to accommodate plasticity under high stresses. Third, the interfaces between the two nanoscale intermetallic phases may have increased the fracture strength in the fine rosettes region, so that plastic yielding can occur before fracture. The improved fracture toughness of intermetallic nanolaminates is witnessed by microcracks restrained within lamellae in fine rosettes, as shown in FIGS. 8A through 8F. This crack inhibition effect will release local stress concentration and delay catastrophic fracture.

[0073]It is clear from the above description that in this disclosure, a custom-made Al₉₂Ti₂Fe₂Co₂Ni₂ alloy was fabricated by LPBF. This alloy has rosettes of nanoscale intermetallics and a macroscopic engineering compressive strength exceeding 900 MPa and 20% plasticity. Micropillar compression tests reveal that the fine rosette regions can achieve a microscopic compressive strength of nearly 1.0 GPa and at least 15% plasticity. The simultaneous achievement of high strength and plasticity arises from the large back stress accommodated through heterogenous intermetallic nanolaminate interfaces. Significant plasticity was also observed in the medium entropy monoclinic Al₉(Fe,Co,Ni)₂ intermetallic phases. The mechanisms that trigger the formation of abundant stacking faults in monolithic Al₉(Fe,Co,Ni)₂ remain to be illuminated by future modeling investigations. Our results shed light on incorporation of nanoscale intermetallics rosettes in the design of ultra-strong Al alloys with prominent plasticity.

[0074]Based on the above description it is an objective of this disclosure to describe an alloy composition containing An alloy comprising 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel. In some embodiments of the alloy, the contains medium-entropy intermetallic lamellae. In some embodiments of the alloy containing medium-entropy intermetallic lamellae, the medium-entropy intermetallic lamellae are of nanoscale. In some embodiments of the alloy containing medium entropy intermetallic lamellae, medium entropy intermetallic lamellae are either Al₉(Fe,Co,Ni)₂ or Al₃Ti. In some embodiments of the alloys of this disclosure, the compressive strength of the alloy ranges from 600-900 MPa. In some embodiments of the alloys of this disclosure, wherein the compressive strain of the alloy ranges from 10-20%.

[0075]Based on the above description, it is another objective of this disclosure to describe a method of making an alloy. The method contains the steps of 1) providing particles wherein each particle has a composition 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel; 2) utilizing a selective leaser melting (SLM) apparatus producing a first layer of the particles on a substrate and melting and solidifying a first group selected areas of the layer of particles, wherein the melting and the solidification results in an alloy of composition 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel, containing intermetallic lamellae of compositions Al₉(Fe,Co,Ni)₂ and Al₃Ti wherein the alloy formed has thickness equal to thickness of the first layer; 3) repeating utilization of SLM apparatus to produce a second layer of the particles and laser melting and solidification of a second group of selected areas of the second layer of the particles, wherein the second group of selected areas is coincident or in contact with the first group of selected areas wherein the melting and the solidification results in an alloy of composition 92 at % aluminum, 2 at % titanium, 2 at % iron, 2 at % cobalt, and 2 at % nickel, containing intermetallic lamellae of compositions Al₉(Fe,Co,Ni)₂ and Al₃Ti wherein the alloy formed has thickness equal to thickness of the first layer; and 4) repeating the utilization of SLM apparatus to produce an object of specified thickness and shape containing the melted and solidified areas. In some embodiments of the method of this disclosure, the power of laser used in SLM is in the range of 200-300 Watts. In some embodiments of the method, the particles are nearly spherical. In some embodiments of the method the particles are produced by gas atomization.

[0076]Additional disclosure is found in Appendix A attached to this specification. The contents of Appendix A are herein incorporated be reference in their entirety into this specification.

[0077]While the present disclosure has been described with reference to certain embodiments, it will be apparent to those of ordinary skill in the art that other embodiments and implementations are possible that are within the scope of the present disclosure without departing from the spirit and scope of the present disclosure. Thus, the implementations should not be limited to the particular limitations described. Other implementations may be possible. It is therefore intended that the foregoing detailed description be regarded as illustrative rather than limiting. Thus, this disclosure is limited only by the following claims.