NANOSTRUCTURE ASSISTED CASTING OF THERMALLY STABLE, ULTRAFINE GRAINED, NANOCRYSTALLINE METALS

BACKGROUND

Ultrafine grained (UFG), nanocrystalline metals, exhibiting extraordinary properties, are highly demanded in the fields of aerospace, electronics and transportation, amongst others. However, other fabrication methods for UFG, nanocrystalline metals, such as mechanical alloying, severe plastic deformation, fast cooling and thin film deposition, are challenging for economical mass production of bulk samples with complex geometries. Moreover, the unsatisfied thermal stability also constrains the application of UFG, nanocrystalline metals.

It is against this background that a need arose to develop the embodiments described herein.

SUMMARY

Some embodiments include a nanocrystalline material comprising: a matrix including one or more metals; and nanostructures dispersed in the matrix, wherein the matrix is polycrystalline and includes grains having an average size of about 1 μm or less. In some embodiments, the average size of the grains is about 600 nm or less. In some embodiments, the average size of the grains is about 400 nm or less. In some embodiments, the nanostructures are dispersed in the matrix at a volume fraction of about 5% or greater of the nanocrystalline material. In some embodiments, the volume fraction of the nanostructures in the nanocrystalline material is about 10% or greater. In some embodiments, the volume fraction of the nanostructures in the nanocrystalline material is about 15% or greater. In some embodiments, the matrix includes copper, and the nanostructures include a transition metal or a transition metal carbide. In some embodiments, the matrix includes zinc, and the nanostructures include a transition metal or a transition metal carbide. In some embodiments, the matrix includes aluminum, and the nanostructures include a transition metal carbide or a transition metal boride.

Some embodiments include a manufacturing method of a nanocrystalline material, comprising: heating a matrix material including one or more metals to form a melt; loading a mixture including a salt and reinforcing nanostructures over a surface of the melt, such that the salt is heated to form a molten salt including the nanostructures dispersed therein; agitating the melt to incorporate the nanostructures from the molten salt into the melt; delivering the melt to a mold defining a hollow space with a requisite shape; and cooling and solidifying the melt to form a metal part including the nanocrystalline material and having the requisite shape.

Some embodiments include a manufacturing method of a nanocrystalline material, comprising: mixing a powder of a matrix material and reinforcing nanostructures to form a powder mixture; compacting the powder mixture to form a preform; heating the preform under compression to form a melt including the nanostructures dispersed therein; and cooling the melt including the nanostructures dispersed therein to form a master material. Some embodiments further comprise subjecting the master material to casting to form a metal part. In some embodiments, subjecting the master material to casting includes heating the master material to form a master material melt, delivering the master material melt to a mold defining a hollow space with a requisite shape, and cooling and solidifying the master material melt to form the metal part including the nanocrystalline material and having the requisite shape. In some embodiments, cooling the master material melt is performed at a rate of less than about 100 K/s.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows microstructure of bulk UFG, nanocrystalline Cu containing distributed WC nanoparticles. (FIG. 1A) SEM image of Cu-5 vol. % WC (by a cooling rate of about 4 K/s) acquired at about 52° showing well dispersed WC nanoparticles in the Cu matrix. Inset is the image of a typical as-cast bulk Cu-5 vol. % WC ingot with a diameter of about 50 mm. (FIG. 1B) Magnified SEM image of Cu-5 vol. % WC (about 4 K/s) showing the ultrafine and nanoscale Cu grains. (FIG. 1C) SEM image of the cross section showing ultrafine-grained Cu matrix and WC nanoparticles present beneath surface of the sample. (FIG. 1D) A typical FIB image of Cu-5 vol. % WC (about 4 K/s) showing the UFG, nanocrystalline microstructure. (FIG. 1E) FIB image of pure Cu cast under the same condition showing coarse Cu grains. (FIG. 1F) EBSD image of Cu-5 vol. % WC (about 4 K/s) with grain size color code. Black phases are WC nanoparticles, red grains are smaller than about 100 nm, yellow and orange grains are smaller than about 1.0 μm. (FIG. 1G) Summary of the average Cu grain sizes for different volume fractions of nanoparticles under different cooling rates. Error bars show the standard deviation.

FIG. 2 shows nucleation and grain growth control mechanisms by nanoparticles. (FIG. 2A) Typical DSC scanning result during the cooling of substantially pure Cu and Cu-10 vol. % WC. (FIG. 2B) A typical TEM image of Cu-13 WC interface showing the interface between Cu and WC nanoparticle. (FIG. 2C) Fourier-filtered high resolution TEM image of the marked red rectangle area in (FIG. 2B) showing a characteristic interface between WC nanoparticle and Cu matrix. Insets are the fast Fourier transformation of the Cu matrix (top right) and WC nanoparticle (bottom left). (FIG. 2D) Undercooling to overcome Gibbs-Thompson pinning effect for Cu, Al and Zn. (FIG. 2E) Schematic illustration of the nanoparticle pinning effects. (FIG. 2F) SEM image of a Cu grain refined by WC nanoparticles. (FIG. 2G) Nanoparticles break the fundamental constraint in other grain refinement methods. (FIG. 2H) and (FIG. 21) Schematic illustrations of phase evolution during solidification of substantially pure metal (FIG. 2H) and metal with nanoparticles (FIG. 21).

FIG. 3 shows nanoparticle assisted grain refinement in other materials systems. (FIGS. 3A-B) FIB images of Al-10 vol. % TiB2 cast by furnace cooling (about 0.7 K/s) showing the distribution of TiB2 nanoparticles and ultrafine AL grains. (FIG. 3C) TEM image of Al-10 vol. % TiB2 (about 0.7 K/s) showing one ultrafine Al grain surrounded by TiB2 nanoparticles. (FIG. 3D) Al Grain size distribution of Al-10 vol. % TiB2 (about 0.7 K/s). (FIGS. 3E-F) SEM image of Zn-5 vol. % WC by air cooling (about 3.7 K/s). (FIG. 3G) FIB image of Zn-5 vol. % WC (about 3.7 K/s). (FIG. 3H) Zn Grain size distribution of Zn-5 vol. % WC (about 3.7 K/s).

FIG. 4 shows thermal stability of ultrafine, nanocrystalline Cu containing WC nanoparticles. (FIGS. 4A-D) STEM images of an area with a high percentage of WC nanoparticles at room temperature, about 400° C., about 600° C. and about 850° C., respectively. (FIGS. 4E-H) STEM image of an area with a relative low percentage of WC nanoparticles at room temperature, about 400° C., about 600° C. and about 850° C., respectively. (FIG. 41) SEM image of Cu-34 vol. % WC after heat treatment (about 750° C. for about 2.0 hours). (FIG. 4J) EBSD image corresponds to the marked white rectangle in (FIG. 41). (FIG. 4K) Cu grain size distribution of the heat treated Cu-34 vol. % WC sample.

FIG. 5 shows fabrication of Cu containing WC nanoparticles. (FIG. 5A) Schematic illustration of the salt-assisted self-incorporation for Cu-13 WC before casting bulk ingots. (FIG. 5B) Schematic illustration of the powder melting method to cast Cu containing WC nanoparticles.

FIG. 6 shows cooling curves for furnace cooling, air cooling and water quenching of Cu-13 WC samples.

FIG. 7 shoes the size distribution of WC nanoparticles in as-solidified Cu-13 WC sample.

FIG. 8 shoes the structure of bulk UFG, nanocrystalline Cu containing WC nanoparticles. (A-C) FIB image of Cu-5 vol. % WC, Cu-10 vol. % WC and Cu-20 vol. % WC, respectively.

FIG. 9 shows a STEM image of the WC nanoparticle rich area showing that Cu grain size is correlated with WC inter-particle spacing.

FIG. 10 shows a cooling curve during the DSC tests at a cooling rate of about 5° C./min.

FIG. 11 shows mechanical properties of UFG, nanocrystalline Cu containing WC nanoparticles. (FIG. 11A) SEM image of a Cu-34 vol. % WC micropillar machined by FIB. (FIG. 11B) FIB image of the micropillar showing polycrystalline Cu matrix with WC nanoparticles. (FIG. 11C) Engineering stress-strain curves of as-solidified pure Cu samples (blue), with nanoparticles (black) and heat treated sample (red). (FIGS. 11D-E) SEM images showing the morphology of post-deformed samples with (FIG. 11D) and without (FIG. 11E) nanoparticles. (FIG. 11F) Young's modulus of pure Cu, Cu-19 vol. % WC and Cu-34 vol. % WC.

FIG. 12 shows an undercooling profile with respect to solid fraction. The solid curves correspond with the constitutional undercooling for Al-Ti alloys. Constrained by the phase diagram and the maximum solubility of Ti in Al, the growth restriction factor reaches a maximum value of Q. in the A1-0.15Ti alloy. However, the undercooling profile of the nanoparticle assisted phase growth indicates an almost infinitely large Qn_p.

DETAILED DESCRIPTION

Embodiments of this disclosure are directed to an improved, cost effective method to fabricate bulk, thermally stable, UFG, nanocrystalline metals through a casting process by the addition of nanostructures (e.g., nanoparticles). UFG, nanocrystalline metals, encompassing metals with a grain size smaller than one micrometer (μm), can exhibit extraordinary mechanical, physical, and chemical properties. However, it has been considered impractical to synthesize bulk UFG, nanocrystalline metals by casting, due to its slow cooling (e.g., less than about 100 K/s). Moreover, other UFG, nanocrystalline metals are generally not thermally stable, and nanocrystalline metals (such as Al, Cu, Zn, and Mg) can exhibit extensive grain growth even at room temperatures. Embodiments of this disclosure can pave a path towards scalable manufacturing of bulk, thermally stable, UFG, nanocrystalline metals via casting. More specifically, some embodiments demonstrate casting of UFG, nanocrystalline copper (Cu), aluminum (Al) and zinc (Zn) by incorporation of nanostructures.

Aspects of some embodiments include the following:

1. Nanostructure Selection

Nanostructure selection is one consideration factor. In some embodiments, tungsten carbide (WC) nanoparticles are used for Cu and Zn, and titanium diboride (TiB2) nanoparticles and titanium carbide (TiC) nanoparticles are used for Al. The general selection standard is: (a) Nanoparticles should be thermally/chemically stable in the molten metal; (b) small molten metal-nanoparticle wetting angle (e.g., <about 90°); and (c) close lattice matching between metal matrix and nanoparticles.

2. Incorporation of Nanoparticles in Molten Metal

In some embodiments, there are at least two methods for incorporation:

(a) Melt pressing of metal powders with nanoparticles. For example: Cu powder (Sigma Aldrich, <about 10 μm) is first mixed with a designed volume fraction of WC nanoparticles (US Research Nanomaterials, about 150-200 nm) for about 1.0 hour by a mechanical shaker (SK-O330-Pro). The mixed powders are then compressed into disks of about 2.0 cm in diameter under about 250 MPa by a hydraulic machine. Then these disks are compression-melted at about 1250° C. by an induction heater while under a pressure of about 7-10 MPa under a graphite piston. After solidification in different cooling rates (furnace cooling, air cooling and water quenching), Cu ingots with WC nanoparticles are obtained.

(b) Salt-assisted self-incorporation of nanoparticles into molten metal: nanoparticles are mixed with a salt (potassium aluminum fluoride, borax, calcium fluoride, and so forth) by mechanical mixing. A metal is melted above its melting temperature. Inert argon gas is purged on the surface of the molten metal to avoid severe oxidation. The mixed salt-nanoparticles powder is loaded on the surface of the molten metal. Then a propeller is located below the molten metal-salt interface and operated to stir the melt. Finally, the molten melt is allowed to solidify to obtain ingots.

3. A Theoretical Model of Nanostructure Assisted Continuous Nucleation and Controlled Grain Growth

A theoretical model is successfully established to explain the casting of UFG, nanocrystalline metals. Several mechanisms can contribute: (a) Large amount of nanostructures can serve as potential nucleation sites during solidification; (b) continuous nucleation during solidification; (c) nanostructures impede the grain growth by Gibbs-Thompson effects; and (d) nanostructures tune the thermal properties (e.g., thermal conductivity) of the melt.

4. Extraordinary Termal Sability of UFG, Nanocrystalline Metals Assisted by Nanostructures

As-cast UFG, nanocrystalline Cu with WC nanoparticles shows outstanding thermal stability at high temperatures. The UFG structure is stable at about 750° C. and has little grain growth at about 850° C.

Advantages of embodiments of this disclosure include:

Casting process can be used to achieve casting of bulk UFG, nanocrystalline metals.

Scalable manufacturing process.

Excellent thermal stability up to about 750° C. or greater.

Excellent mechanical properties.

EXAMPLE EMBODIMENTS

In some embodiments, a nanocrystalline material includes a matrix including one or more metals, along with reinforcing nanostructures dispersed in the matrix. Examples of suitable matrix materials include aluminum (Al), magnesium (Mg), iron (Fe), silver (Ag), copper (Cu), manganese (Mn), nickel (Ni), titanium (Ti), chromium (Cr), cobalt (Co), zinc (Zn), alloys, mixtures, or other combinations of two or more of the foregoing metals, and alloys, mixtures, or other combinations of one or more of the foregoing metals with other elements. In some embodiments, the matrix is polycrystalline and includes grains having an average size (or an average dimension) of about 2 μm or less, about 1.5 μm or less, about 1.2 μm or less, about 1 μm or less, about 800 nm or less, about 600 nm or less, about 400 nm or less, or about 300 nm or less, and down to about 200 nm or less, or down to about 150 nm or less.

In some embodiments, the nanostructures can have at least one dimension in a range of about 1 nm to about 1000 nm, such as about 1 nm to about 500 nm, about 1 nm to about 400 nm, about 1 nm to about 300 nm, about 1 nm to about 200 nm, or about 1 nm to about 100 nm. In some embodiments, the nanostructures can have at least one average or median dimension in a range of about 1 nm to about 500 nm, about 1 nm to about 400 nm, about 1 nm to about 300 nm, about 1 nm to about 200 nm, or about 1 nm to about 100 nm. In some embodiments, the nanostructures can include nanoparticles having an aspect ratio of about 5 or less, or about 3 or less, or about 2 or less and having generally spherical or spheroidal shapes, although other shapes and configurations of nanostructures are contemplated, such as nanofibers and nanoplatelets. In the case of nanoparticles of some embodiments, the nanoparticles can have at least one dimension (e.g., an effective diameter which is twice an effective radius) or at least one average or median dimension (e.g., an average effective diameter which is twice an average effective radius) in a range of about 1 nm to about 1000 nm, such as about 1 nm to about 500 nm, about 1 nm to about 400 nm, about 1 nm to about 300 nm, about 1 nm to about 200 nm, or about 1 nm to about 100 nm.

In some embodiments, the nanostructures can include one or more ceramics, although other nanostructure materials are contemplated, such as metals. Examples of suitable nanostructure materials include metal oxides (e.g., alkaline earth metal oxides, post-transition metal oxides, and transition metal oxides, such as aluminum oxide (Al₂O₃), magnesium oxide (MgO), titanium oxide (TiO₂), and zirconium oxide (ZrO₂)), non-metal oxides (e.g., metalloid oxides such as silicon oxide (SiO₂)), metal carbides (e.g., transition metal carbides, such as titanium carbide (TiC), niobium carbide (NbC), chromium carbide (Cr₃C₂), nickel carbide (NiC), hafnium carbide (HfC), vanadium carbide (VC), tungsten carbide (WC), and zirconium carbide (ZrC)), non-metal carbides (e.g., metalloid carbides such as silicon carbide (SiC)), metal silicides (e.g., transition metal silicides, such as titanium silicide (TiSi)), metal borides (e.g., transition metal borides, such as titanium boride (TiB₂), zirconium boride (ZrB₂), hafnium boride (HfB₂), vanadium boride (VB₂), and tungsten boride (W₂B₅)), metal nitrides (e.g., transition metal nitrides), non-metal nitrides (e.g., metalloid nitrides such as silicon nitride), alloys, mixtures, or other combinations of two or more of the foregoing. Particular examples of suitable nanostructure materials include transition metal borides (e.g., TiB₂) and transition metal carbides (e.g., TiC and WC), amongst other transition metal-containing ceramics.

In some embodiments, the nanocrystalline material can include the nanostructures at a high volume percentage of, for example, greater than about 3%, such as about 5% or greater, about 6% or greater, about 7% or greater, about 8% or greater, about 9% or greater, about 10% or greater, about 15% or greater, about 20% or greater, about 25% or greater, or about 30% or greater, and up to about 40% or greater.

In some embodiments, the matrix includes Cu, and the nanostructures include a transition metal (e.g., W) or a transition metal carbide (e.g., WC).

In some embodiments, the matrix includes Zn, and the nanostructures include a transition metal (e.g., W) or a transition metal carbide (e.g., WC).

In some embodiments, the matrix includes Al, and the nanostructures include a transition metal carbide (e.g., TiC) or a transition metal boride (e.g., TiB₂).

In some embodiments, a manufacturing method of a nanocrystalline material includes: (1) heating a matrix material including one or more metals to form a melt; (2) loading a mixture including a salt and reinforcing nanostructures over a surface of the melt, such that the salt is heated to form a molten salt including the nanostructures dispersed therein; (3) agitating the melt to incorporate the nanostructures from the molten salt into the melt; and (4) cooling the melt including the nanostructures dispersed therein to form a master material.

In some embodiments of the method, the method further includes subjecting the master material to casting to form a metal part. In some embodiments, subjecting the master material to casting includes heating the master material to form a melt, delivering the melt to a mold defining a hollow space with a requisite shape, and cooling and solidifying the melt to form the metal part including the nanocrystalline material and having the requisite shape. In some embodiments, cooling the melt is performed at a rate of less than about 100 K/s, such as about 90 K/s or less, about 70 K/s or less, about 50 K/s or less, about 30 K/s or less, or about 10 K/s or less.

In some embodiments of the method, features of the matrix material, the nanostructures, and the nanocrystalline material are as described for the foregoing embodiments of the nanocrystalline material.

In some embodiments, a manufacturing method of a nanocrystalline material includes: (1) mixing a powder of a matrix material and reinforcing nanostructures to form a powder mixture; (2) compacting the powder mixture to form a preform, such as in a hydraulic press and under a pressure of about 100 MPa to about 300 MPa, or about 250 MPa; (3) heating the preform under compression, such as under a pressure of about 7 MPa to about 10 MPa and to a temperature up to or below (e.g., up to about 100° C. below, or up to about 200° C. below) a melting temperature of the matrix material, to form a melt including the nanostructures dispersed therein; and (4) cooling the melt including the nanostructures dispersed therein to form a master material.

EXAMPLES

The following example describes specific aspects of some embodiments of this disclosure to illustrate and provide a description for those of ordinary skill in the art. The example should not be construed as limiting this disclosure, as the example merely provides specific methodology useful in understanding and practicing some embodiments of this disclosure.

Bulk Ultrafine Grained, Nanocrystalline Metals via Slow Cooling
Overview:

Cooling, nucleation, and phase growth are ubiquitous processes in nature. Effective control of nucleation and phase growth is of significance to yield refined microstructures with enhanced performance for materials. Ultrafine grained (UFG), nanocrystalline metals can exhibit extraordinary properties. However, other microstructure refinement methods, such as fast cooling and inoculation, have reached certain fundamental constraints. It has been considered impractical to fabricate bulk UFG, nanocrystalline metals via slow cooling. Here this example reports that nanostructures can refine metal grains to ultrafine nanoscale by instilling a continuous nucleation and growth control mechanism during slow cooling. The bulk UFG, nanocrystalline metal with nanostructures also reveals an unprecedented thermal stability. This method overcomes the grain refinement constraints and can be extended to any other processes that involve cooling, nucleation and phase growth for widespread applications.

Introduction:

Cooling, nucleation, and phase growth are ubiquitous processes of significance in various aspects, such as cloud formation, ice nucleation, and volcanic rock evolution. It is established that effective control of nucleation and phase growth will yield refined microstructures with enhanced performance for materials, and hence vital to numerous broad fields, including materials science, climate and atmospheric sciences, biomedicine and chemistry. Widely used technologies that involve cooling, such as casting, are of significance for the mass production of complex materials and components. Ultrafine grained (UFG), nanocrystalline metals can exhibit extraordinary properties. However, it has been considered impractical to fabricate bulk ultrafine grained, nanocrystalline metals by casting, partly due to its slow cooling (e.g., less than about 100 K/s). Other microstructure refinement methods, such as fast cooling (thousands to millions K/s) and inoculation, have reached certain fundamental or technical constraints. Fast cooling substantially restricts the size and complexity of as-solidified materials. The minimum grain size generally achievable by inoculation in casting falls in the range of tens of micrometers.

Grain refinement in metals during solidification is of great interest due to the enhanced mechanical properties, more homogeneous microstructure and improved processability of refined microstructures. Various approaches such as inoculation, growth restriction by adding alloy elements and fast cooling (up to about 10⁷K/s) are investigated through both theoretical and experimental pathways in an effort to attain optimal grain refinement effects. However, these approaches have failed to demonstrate whether it is possible to cast UFG, nanocrystalline metals via a casting process, which would represent a revolutionary approach given the pervasiveness of casting in manufacturing. As a consequence of the inability to implement solidification processes to fabricate UFG, nanocrystalline metals, various methods emerge for the fabrication of UFG, nanocrystalline metals, including: mechanical alloying, severe plastic deformation and thin film deposition and although some degree of success has been achieved using these processes they remain constrained to the solid state and present challenging issues for economical mass production of bulk samples with complex geometries. Solidification behavior of metals can be controlled by the addition of nanostructures. Nanostructure-controlled solidification can provide a pathway for casting metals with refined microstructures.

Here this example reports that nanostructures can refine metal grains down to ultrafine or even nanoscale by instilling a continuous nucleation and growth control mechanism during slow solidification. When casting substantially pure copper (Cu) with tungsten carbide (WC) nanoparticles, the grain sizes of Cu are refined substantially down to ultrafine and even nanoscale. The as-solidified bulk ultrafine, nanocrystalline Cu reveals an unprecedented thermal stability up to about 1023 K (0.75 melting point of Cu) and high mechanical properties. Furthermore, this revealed grain control mechanism is successfully applied in other materials systems such as aluminum-titanium boride (Al-TiB₂) and zinc-tungsten carbide (Zn-WC) for ultrafine grains via slow cooling. This method paves a pathway for the mass production of bulk stable UFG, nanocrystalline materials that can be readily extended to any other processes that involve cooling, nucleation and phase growth for widespread applications.

Results and Discussion:
Nanoparticles Incorporation and Dispersion in Bulk Samples

Preparation is made of bulk Cu ingots with WC nanoparticles by two different methods. The first method is a salt-assisted self-incorporation of nanoparticles into molten metal (see Methods). As shown in FIG. 5A, molten salt (Borax+about 5% CaF₂) could dissolve the oxide layer at the top of molten metal and provide a clean interface between Cu melt and nanoparticles. In addition, the wetting angle between Cu and WC at about 1250° C. is below about 10°, which indicates a good wettability between Cu and WC so that WC nanoparticles prefer to transport from the molten salt to the Cu melt to reduce the system energy. Combined with mechanical mixing, WC nanoparticles can be readily incorporated into molten Cu. This salt-assisted incorporation method opens up a scalable manufacturing method to fabricate metals containing different volume percentages of nanoparticles. In this example, bulk Cu ingots with about 5%, about 10% and about 20% volume fraction of WC nanoparticles were cast. Another method is a powder melting process (see Methods) especially suitable for high volume percent of nanoparticles as shown in FIG. 5B. A cold compacted Cu-13 WC powder preform was melted by an induction heater under a pressure of about 7-10 MPa. Cu ingots with about 19% and about 34% volume fraction of WC nanoparticles were fabricated by this powder melting method. It should be noted that the second method may not be as scalable as the first method, but is more suitable for a high volume loading of nanoparticles if desired.

To evaluate the effects of the cooling rate under a same initial condition, all these Cu-13 WC samples were then melted again and cast under different cooling rates, namely furnace cooling (about 2-4 K/s), air cooling under the protection of argon gas (about 7-12 K/s) and water quenching (about 70-100 K/s). The typical cooling curves are shown in FIG. 6, and the thermal arrest caused by the solidification of Cu is clearly identified in the furnace and air cooling curve, while not so noticeable in the water quenching curve due to the rapid heat dissipation.

Grain structure characterization

Characterization is made of the distribution and dispersion of WC nanoparticles and grain structures in cast bulk Cu samples via scanning electron microscope (SEM), transmission electron microscope (TEM), focused ion beam (FIB) imaging, scanning transmission electron microscope (STEM) and electron backscatter diffraction (EBSD). SEM samples were sectioned along a cross-section of the sample. To clearly reveal the nanoparticles, the mechanically ground and polished samples were further polished by low-angle ion milling at about 4° for about 1.5 hours.

The inset in FIG. 1A shows one typical bulk Cu-13 WC ingot cast after the salt-assisted incorporation method. FIG. 1A shows the typical SEM microstructure of Cu-5 vol. % WC by furnace cooling acquired at about 52° . WC nanoparticles are uniformly dispersed in the Cu matrix. A theoretical analysis, detailed in the subsection below entitled ‘Nanoparticle dispersion and self-stabilization mechanism’, attributes that the excellent wettability between Cu and WC can stabilize the dispersed nanoparticles in the metal melt. The average size of WC nanoparticles was measured to be about 200 nm in diameter as shown in FIG. 7. The UFG microstructure of Cu-5 vol. % WC (about 4 K/s) is shown in the SEM image of FIG. 1B (Cu grains marked by white dash lines). Grains in this SEM image are smaller than about 1000 nm. It should be noted that some areas without nanoparticles from a top view under SEM may have nanoparticles beneath the surface as shown in the SEM image of FIG. 1C, showing the cross-section cut by FIB. Cu grains are pinned by nanoparticles beneath the surface. Pt coating is used to protect the Cu surface during the FIB cutting. The channeling contrast of different grains induced by ion beam makes FIB a powerful tool to characterize grain structures. FIG. 1D and FIG. 1E are the typical ion beam micrographs of Cu-5 vol. % WC and substantially pure Cu cast under the same casting condition (about 2-4 K/s), respectively. The dark phase in FIG. 1D corresponds to WC nanoparticles whereas the white or grayish phases indicate the Cu grains. More FIB images of as-cast Cu-13 WC samples with about 5, about 10 and about 20 vol. % of WC nanoparticles are shown in FIGS. 8A-C. With the addition of WC nanoparticles into Cu, the grain sizes are readily refined to ultrafine and even nanoscale. From the FIB micrographs, it can be observed that a higher percentage of nanoparticles yield more dense nanoparticles distribution and thus more refined grain microstructures. For comparison, the as-solidified pure Cu sample, however, has an average grain size of 270±132 μm under the same cooling rate.

EBSD analysis was used to further investigate the grain sizes. FIG. 1F shows the typical EBSD micrograph of Cu-5 vol. % WC cast by a cooling rate of about 4 K/s. The black phases are corresponding to the WC nanoparticles. Given that the surfaces of nanocomposite samples were not perfectly flat (nanoparticles stick out from the matrix) after ion milling, some regions (marked as greyish region) were not identifiable during the EBSD scanning. The colors of different grains correspond to different sizes, as shown in the legend. Red grains are smaller than about 100 nm, yellow and orange grains are smaller than about 1.0 μm, while green and deep blue grains are smaller than about 2.0 μm. The majority of the Cu grains are smaller than about 1.0 μm and significant number of nanosized Cu grains (marked with red color in the EBSD micrograph) can be observed in the nanoparticle-rich areas. The average grain sizes of different areas under different cooling rates are summarized in FIG. 1G. The average grain size of Cu-13 WC samples with different fraction of WC nanoparticles ranges from about 236-434 nm, about 227-384 nm and about 193-347 nm for furnace cooling, air cooling and water quenching, respectively. The error bars indicate the standard deviation of the grain size measurements. It is shown that a higher percentage of nanoparticles yields more refined grains. FIG. 1G also indicates that the grain size slightly decreases with an increased cooling rate although the differences from the cooling rates are not as significant. The EBSD scanning also supports that the addition of WC nanoparticles can refine Cu grains to ultrafine/nanoscale by regular casting.

If it is assumed that nanoparticles are spherical and homogeneously distributed in the matrix, the theoretical inter-particle spacing between WC nanoparticles can be calculated by:

$d = {r (\frac{4 }{3 f_{v}})}^{\frac{1}{3}}$

where d is the theoretical inter-particle spacing, r is the radius of the particles (e.g., about 100 nm in this case), and f_vis the volume fraction of particles. The green dotted line in FIG. 1G corresponds to the theoretical inter-particle spacing of the WC particles at different volume fractions. It is noted that the experimental results for Cu grain sizes are close to the theoretical inter-particle spacing between nanoparticles. FIG. 9 shows the scanning transmission electron microscope (STEM) image of the nanoparticle rich area, which indicates that Cu grain size is correlated to the WC inter-particle spacing. The equation also indicates that a smaller nanoparticle would allow a lower volume fraction to yield smaller grains as long as the nanoparticles are not engulfed. Lower cooling rates actually favor the non-engulfment (pushing) of nanoparticles during solidification.

Nucleation and grain growth control

To further provide fundamental insight into the underlying nucleation and grain growth phenomena, differential scanning calorimetry (DSC) analyses were conducted and the typical cooling curves at a cooling rate of about 5° C./min of substantially pure Cu and Cu with WC nanoparticles (samples have the same mass, about 50 mg) are shown in FIG. 2A (see more details in Methods and FIG. 10). The bump on the cooling curve of substantially pure Cu sample corresponds to the exothermal peak from the solidification of Cu started at about 1033° C., which indicates an about 51° C. undercooling to activate major nucleation of Cu grains. However, the starting point of the exothermal peak of Cu-10 vol. % WC is about 1078° C., which means about 6 ° C. undercooling results by the addition of WC nanoparticles. The undercooling difference indicates that WC nanoparticle is a relatively potent nucleation particle for Cu. For effective nucleation, the crystallographic lattice discrepancy between particles and matrix is of significance for the nucleation; thus the interface between matrix and nanoparticles were evaluated at atomic scale by high resolution transmission electron microscope (HRTEM). FIG. 2B shows the typical interface between WC nanoparticles and Cu matrix at atomic scale. FIG. 2C is the Fourier-filtered atomic resolution TEM image at the marked area of FIG. 2B. The atomic structure indicates a clean and well-matched Cu-13 WC interface with no intermediate phases present. The top-right and bottom-left insets are the fast Fourier transformation of the Cu matrix and WC nanoparticles, respectively. It is identified that the (1011) planes of WC nanoparticles are parallel with (200) planes of Cu matrix. The plane distance of (200) Cu and (1011) WC are about 0.1806 nm and about 0.1881 nm, respectively. Thus the misfit is calculated to be about 4.1%, which implies a coherent lattice matching at this specific interface. The presence of a clean and coherent interface indicates a strong interfacial bond. The coherent interface indicates that WC nanoparticles could serve as potent nucleation sites for Cu grains.

Moreover, the number of potential nucleation sites (WC nanoparticles) in Cu-13 WC sample is significantly higher than that in an inoculation method. Nanoparticles can serve as effective nucleation sites due to their high number density. For example, the population density of particles for 5 parts per thousand (ppt) Al-5Ti-1B refiner is about 5×10¹⁴m⁻³. In comparison, the population densities of WC nanoparticles are calculated to be about 1.2×10¹⁹m⁻³, about 2.4×10¹⁹m⁻³and about 4.8×10¹⁹m⁻³(assuming that WC nanoparticles are spherical) for Cu-5 vol. % WC, Cu-10 vol. % WC and Cu-20 vol. % WC samples, respectively. The population densities of nanoparticles are five orders of magnitude higher than that with the Al—Ti—B grain refiner.

Following the initial nucleation of grains from the melt, the nucleated grains grow rapidly and release latent heat, which usually impedes the nucleation of additional grains nearby. In contrast, with nanoparticles in the molten metal, nanoparticles can rapidly assemble/adhere to the solid-liquid interface to effectively restrict the grain growth, preventing other potential nucleation sites from being suppressed by the latent heat release. Therefore, later nucleation events can occur continuously during the solidification, which is significant for this grain refinement and growth control mechanism. This growth restriction and continuous nucleation can be validated by comparing the width of the exothermal peaks of substantially pure Cu and Cu with nanoparticles in DSC analyses. FIG. 2A indicates that the Cu-10 vol. % WC sample took about 83% longer time to complete the solidification. The heat flow curves during DSC tests are shown in FIG. 10. The exothermal peaks correspond to the latent heat release from the solidification. The width of the peak in substantially pure Cu is about 6.9° C., while it is about 12.6° C. in the Cu-10 vol. % WC sample, which means exothermal peak in substantially pure Cu is sharp and intensive, and it is gradual and wide in Cu-13 WC samples. These results indicate that the nanoparticles are able to slow down solidification and allow a continuous nucleation and an effective grain growth control during solidification.

After nucleation, grain growth control during solidification is of significance to achieve refined grain structures. It is proposed here that nanoparticles will impede grain growth during solidification by forcing the solidification front to grow with a non-linear geometry (e.g., curved solidification front). When Cu grains form curvatures with a small radius during solidification, the free energy (Gibbs-Thomson effect) increases. The molar free energy increase can be calculated by:

$Δ G_{γ} = \frac{2 γ V_{m}}{r}$

where ΔG_γ is the molar free energy increase, γ is the interfacial energy of the interface, V_mis the molar volume of the phase, and r is the radius of curvature of the interface. If this increment of free energy is large enough (thus larger undercooling is involved), the growth of Cu grains may be inhibited. The undercooling to overcome a curved interface generated by nanoparticles pinning can be described by the Gibbs-Thompson effects and can be calculated by:

$Δ T = \frac{2 γ T_{m}}{r H_{f}}$

where y is the interface energy of solid-liquid interfaces (Cu: about 0.185 J/m², Al: about 0.116 J/m²and Zn: about 0.09 J/m²), T_mis melting of the metal, r is the radius of curvature of the solidification front and H_fis the enthalpy of fusion. The undercooling to overcome the grain front curvatures when pinned by nanoparticles for Cu, Al, and Zn are shown in FIG. 2D.

As shown in the schematic illustration (FIG. 2E) of the nanoparticle pinning effects, there are mainly two types of curvature generated by nanoparticles: (1) Marked as r₁and r₃in FIG. 2E, when a solidification front meets with dispersed nanoparticles, the Cu phase can grow through the micro/nano-channels between nanoparticles and form a curvature with a small radius that results in an increase of free energy (Gibbs-Thomson effect). (2) Marked as r₂in FIG. 2E, nanoparticles have curved surfaces and when Cu solidification front meet nanoparticles, the curvature at the interface between Cu and WC nanoparticles will involve extra undercooling to remain solid. The range of the undercooling specified by the curvatures generated by nanoparticle pinning effects are approximately in the range of tens degrees of Kelvin as marked by r₁, r₂and r₃in FIG. 2D. The curvatures generated by nanoparticles are observed in SEM image as shown in FIG. 2F. Cu grain is surrounded by several WC nanoparticles and shows both types of curvatures as mentioned above.

On the basis of theories for nucleation and grain growth in alloys, grain growth velocity can be reduced by the solute atoms at the solid/liquid interface during alloy solidification. An empirical relationship (as shown in FIG. 2G) can be revealed between average grain size (d) and the restriction factor Q for alloys:

$d = a + \frac{b}{Q}$

where a is a constant related to the number of particles that actually nucleate grains at infinite values of Q. In the ideal case, when Q is infinity large, the grain size could be refined to a which is dictated by inter-particle spacing. b is a constant related to the potency of the nucleation particles. Q is the restriction factor, which is inversely proportional to the constitutional undercooling, traditionally specified as:

Q=mC
₀(k−1)

where m is the liquidus slope in a linear phase diagram, C₀is the solute content in the alloy and k is the equilibrium solute partition coefficient. From the equation of Q, it can be seen that the restriction factor is determined by the phase diagram and the specific chemical composition of the specific alloy. Thus the chemical restriction from solute atoms is constrained. The maximum value of Q (Q_max) in some practical alloys can be about 50 K. As shown in FIG. 2G, with a constrained Qmax, the smallest grain size achievable by inoculation and solute atoms is d_limit, which is approximately tens of micrometers.

In contrast, here it is proposed that nanoparticle pinning can induce a restriction factor, Q_np, which can break the fundamental constraint set by Q that depends on constitutional undercooling. The nanoparticle assisted grain growth restriction can extend the chemical restriction factor Q to a physical restriction factor Qn_np, which could be readily increased to a significant large number, if not infinity (see below section entitled ‘Growth restriction factor by nanoparticles’ and FIG. 12). Moreover, populous nanoparticles (five orders of magnitude more) could serve as potent nucleation sites in the continuous nucleation; thus a (related to the number of effective nucleation sites) is significantly reduced. The average grain size can be reduced to dmin when 1/Q_npapproaches to zero. It is proposed that d_minachievable is determined by the theoretical inter-particle spacing. The trend line of theoretical inter-particle spacing in FIG. 1G is consistent with the experimental data of average grain sizes in Cu-13 WC samples. This indicates that nanoparticle assisted continuous nucleation and grain growth control can refine grains down to the inter-particle spacing under slow cooling. Therefore, nanoparticles effectively break the fundamental constraints set by other solidification processes and readily refine grains down to ultrafine or even nanoscale by regular casting process. Based on the experimental results and theoretical analysis, the mechanisms of nanoparticles assisted grain refinement and control, in comparison to other grain refinement, during regular casting are schematically illustrated in FIG. 2H and 2I.

Other less dominant mechanisms may also play minor roles in phase growth control, such as blocking of the diffusion of atoms to the surface of the growing phase and modification of the local temperature field by nanoparticles. Nanoparticles can remain at the solidification front and block the transportation of the atoms, thus slowing down the grain growth. When the density of nanoparticles is low, this effect will not play a major role. Moreover, when WC nanoparticles are close to the solidification front, the nanoparticles could affect the local thermal fields. The lower thermal conductivity of ceramic nanoparticles could slow down the transportation of latent heat from the solidification front and protect potential nucleation sites.

Other Materials Systems

The nanoparticle assisted grain refinement has also been validated for other materials systems. Al—TiB₂and Zn—WC were also evaluated in order to determine whether it is possible to achieve UFG, nanocrystalline microstructures via slow cooling for different metals. The salt-assisted self-incorporation method (see Methods) was used to fabricate Al containing about 10 vol. % TiB₂nanoparticles (with an average size of about 100 nm) by furnace cooling (about 0.7 K/s as measured). FIG. 3A shows the FIB image of Al-10 vol. % TiB₂, where the dark phases are TiB₂nanoparticles while the grey matrix is Al. TiB₂nanoparticles are reasonably well dispersed in the Al matrix. The higher magnified image is shown in FIG. 3B, and one can clearly identify Al grains (marked by white dash lines and pink arrows) smaller than about 1.0 μm surrounded by TiB₂nanoparticles (marked by yellow arrows). TEM image of FIG. 3C demonstrates one ultrafine Al grain (marked by white dash lines) surrounded by several TiB₂nanoparticles. The grain size distribution of Al is shown in FIG. 3D with an average grain size of 460±220 nm. Zn-5 vol. % WC samples were also fabricated by casting with a cooling rate of about 3.7 K/s (see Materials and Methods). As shown in FIG. 3E and 3F (SEM images of Zn-5 vol. % WC), WC nanoparticles are distributed and dispersed in Zn matrix. Zn grains (marked by white dash lines and pink arrows) close to about 1.0 μm are clearly identified. The FIB image of FIG. 3G shows a better contrast from the Zn grains. The size distribution of Zn grains is shown in FIG. 3H with an average grain size of 991±746 nm. It is thus validated that Zn grains were refined to about 1.0 μm by the addition of 5 vol. % WC nanoparticles, although not as effective as in the Cu and Al cases. The reason is mainly due to the smaller undercooling to overcome the Gibbs-Thompson pinning effects for Zn than for Al and Cu, as shown in FIG. 2C. This nanoparticle assisted grain refinement approach provides an additional general pathway to produce UFG, nanocrystalline metals by casting.

Thermal Stability

The thermal stability of UFG, nanocrystalline metals is another grand challenge that constrains their widespread use in many applications. Pure nanocrystalline metals (e.g., Al, Cu, Sn, Pb, Zn, and Mg) can exhibit extensive grain growth even at room temperature (about 300 K). Nanocrystalline metals with higher melting points (e.g., about 1700 K), such as Co, Ni, and Fe, can exhibit a rapid grain growth over a moderate temperature range of about 220-450° C. (about 493-723 K), resulting in micrometer-sized grains at a temperatures less than a half of their melting temperatures. Approaches that involve adding solute atoms to stabilize nanocrystalline metals by pinning the grain boundaries or lowering the grain boundary energy have yielded some successful results. However, these approaches remain inherently constrained by the thermodynamic properties of the particular alloy systems. It is proposed that dispersed nanoparticles could stabilize the UFG, nanocrystalline metals at elevated temperatures. To evaluate the nanoparticle assisted stabilization in the as-cast Cu-13 WC nanocomposites at elevated temperatures, the grain growth was further investigated by STEM equipped with in-situ heating capability. The in-situ heating path is staying at about 200, about 400, and about 600° C. for about 10 min and about 850° C. for about 30 min. FIG. 4A and 4E show the local microstructures of as solidified Cu-13 WC samples with different local volume percentages of nanoparticles. The local nanoparticle concentration in FIG. 4A is higher than that in FIG. 4E. The relatively bright phases are WC nanoparticles and the black phase is the Cu matrix. The grain structures at about 400, about 600 and about 850° C. shown in FIG. 4(B-D) and FIG. 4(F-H) are acquired at longer camera lengths than FIG. 4A and 4E to enhance the diffraction contrast of the image in order to distinguish individual Cu grains. In the nanoparticle rich local area, FIG. 4(A-D) show that nanoscale grains were thermally stable at temperatures up to about 850° C. (FIG. 4D). In the area with fewer nanoparticles (FIG. 4(E-H)), ultrafine/nanoscale grains were thermally stable up to about 600° C. However, when the temperature was raised to about 850° C., ultrafine/nanoscale grains started to grow and became larger grains (FIG. 4H). Although it is observed that the grains started to grow at about 850° C., this growth was restricted by nanoparticles close to this area and therefore the grain sizes did not exceed the ultrafine grain size. This restriction indicates that nanoparticles effectively inhibit grain growth in solid state up to a temperature of about 850° C. (about 1123 K), which significantly corresponds to about 0.83 of the melting temperature of Cu (about 1353 K).

To further validate the thermal stability of the refined Cu grains and evaluate the effects on mechanical properties by heat treatment, heating is performed on the Cu-34 vol. % WC sample obtained by casting under air cooling (cooling rate of about 7 K/s), which had an average grain size of 208 ±94 nm, to about 750° C. (about 1023 K, about 0.75 of the melting point of Cu) and the sample is held at this temperature for about 2.0 hours under the protection of argon. FIG. 4I shows a typical SEM image of the Cu-34 vol. % WC sample after the high temperature exposure. EBSD was utilized to evaluate the potential grain growth in the nanoparticle-poor zone such as the marked white rectangle area in FIG. 41. FIG. 4J clearly shows that most Cu grains are still smaller than about 1.0 μm while a significant number of nanoscale grains remained after the heat treatment. The grain size distribution in FIG. 4K confirms that the average Cu grain size is 248±135 nm after the heat treatment, which is close to the average grain size, 208±94 nm, before the heat treatment.

The thermal stability may be attributable to a Zener pinning effect derived from the presence of dispersed nanoparticles. A maximum mean grain size can be theoretically estimated by:

$D_{\max} = \frac{4 r}{3 f_{v}}$

where r is the nanoparticle radius and f_vis the volume fraction of the nanoparticles. The D_maxfor Cu-34% WC is estimated to be about 392 nm, which is comparable with the experimental data. When compared with other precipitate particles from alloy systems, the ex-situ WC nanoparticles will not dissolve in the Cu matrix at high temperatures, offering a superior thermal stability than alloys.

Mechanical Properties

The presence of fine, well-dispersed nanoparticles in a metal matrix can significantly enhance mechanical responses. To gain fundamental insights into the property enhancement induced by the refined grains and nanoparticles, micropillar compression tests are conducted as shown in FIG. 11. The typical micropillar of Cu-13 WC sample with diameter of about 4 μm and length of about 9 μm is shown FIG. 11A, demonstrating that dense nanoparticles are dispersed in the micropillar. The polycrystalline nature of the micropillar is shown in FIG. 11B. The typical results from the micropillar compression tests are shown in FIG. 11C. The yield point increased from 180±8 MPa to 827±74 for substantially pure Cu and as-solidified Cu-34 vol. % WC, respectively. Moreover, micropillars from Cu-34 vol. % WC sustained a gradually increasing load up to about 1490 MPa and strain over about 25% without catastrophic failure. In comparison, the substantially pure Cu sample experienced extensive slip as the strain increases discontinuously. Furthermore, it is observed that the mechanical properties did not change much (Yield point drops from 827±74 to 780±9) after a thermal exposure at about 750° C. for about 2.0 hours, which further confirms that the Cu samples with WC nanoparticles exhibit excellent thermal stability. The detailed strengthening mechanisms are analyzed in the supplementary subsection ‘Strengthening mechanisms’. The Young's modulus of substantially pure Cu, Cu-19 vol. % WC, and Cu-34 vol. % WC are measured by the nanoindentation method with a Berkovich tip. The Young's modulus of pure Cu, Cu-19 vol. % WC, and Cu-34 vol. % WC are 139±8 GPa, 192±13 GPa, and 223±16 GPa, respectively. It is hypothesized that the enhancement of Young's modulus is due to the high Young's modulus of WC (about 530-700 GPa) and the effective load transfer by the nanoparticles.

CONCLUSIONS

In summary, an approach to effectively control nucleation and grain growth down to ultrafine/nanoscale during slow solidification is proposed by use of dispersed nanostructures in molten metals. The nanostructures assisted grain refinement mechanisms, which combine continuous nucleation and grain growth control, and break the fundamental constraint of other grain refinement methods. This approach provides a pathway to directly cast bulk UFG, nanocrystalline metals under slow cooling rates (e.g., less than about 100 K/s) for large scale production. An unprecedented thermal stability up to about 1023 K, about 0.75 T_mof Cu, for the as-cast bulk UFG, nanocrystalline Cu is demonstrated. As-cast bulk UFG, nanocrystalline metals with nanoparticles also show exceptional strengths and Young's modulus enhancement. Furthermore, this general approach is applicable for different materials systems such as Al-TiB₂and Zn—WC. This approach can have a significant impact on the solidification processes for metals, and also on numerous applications such as biomedical, chemical, and atmospheric sciences.

Materials and Methods:
Materials Fabrication
1. Salt Assisted Incorporation Method

WC nanoparticles were mixed with Borax (Na₂B₄O₇)-5 wt.% CaF₂salt powders by a mechanical shaker (SK-0330-Pro) for about 1.0 h. The volume fraction of nanoparticles in the salt mixture is designed as about 10%. As shown in FIG. 5A, substantially pure oxygen free Cu (about 99.99%, Rotometals, Inc.) ingots were melted at about 1250° C. in a graphite crucible by induction heater. Inert argon (Ar) gas was purged on the molten Cu to avoid severe oxidation. The mixed Na₂B₄O₇-5 wt.% CaF₂-WC nanoparticles were manually loaded on the surface of the molten Cu. A graphite propeller was located below the Cu-salt interface and stirred at a speed of about 400 rpm for about 20 min to incorporate WC nanoparticles into Cu melt. Then the melt was allowed to cool down to about 900° C. to allow Cu solidify first while the salt mixture was still in a liquid state. Liquid salt was poured out from the crucible and leave a Cu-13 WC ingot. The volume fraction of WC nanoparticles in Cu was designed to be 0, about 5, about 10, and about 20 vol. %.

Surface clean TiB₂nanoparticles were synthesized by the magnesiothermic reduction of TiO₂nanoparticles and B₂O₃powders in molten salt. The synthesized TiB₂nanoparticles and KAl₄flux were then mechanically mixed in solid state for about 3.0 h. Mixed powders were dehydrated at about 120° C. for about 1.0 hour in a vacuum oven. An electrical resistance furnace was used to melt the Al ingots at about 820° C. under Ar gas protection. Then, the mixed powders were added to the melt surface and the melt was mechanically stirred at about 200 rpm for about 10 min with titanium (Ti) mixing blade. The designed volume fraction of TiB₂in Al is about 10 vol. %. The melt was taken out from the furnace and naturally cooled down to room temperature under Ar gas protection.

2. Powder Melting Method

Cu powders (Sigma Aldrich,<about 10 μm) and Zn powders (Alea Aesa, about 150 μm) were first mixed with a designed volume fraction of WC nanoparticles (US Research Nanomaterials, about 150-200 nm) for about 1.0 hour by a mechanical shaker (SK-0330-Pro). The mixed powders were then compressed into disks with about 2.0 cm in diameter under about 250 MPa by a hydraulic machine. As shown in FIG. 5B, Cu-13 WC cold compacted disks were compression-melted at about 1250° C. by an induction heater while under a pressure of about 7-10 MPa under a graphite disk and alumina piston. After solidification in air, Cu ingots with WC nanoparticles were obtained. Zn—WC cold compacted billets were then melted at about 500° C. and ultrasonic processed by a Niobium (Nb) probe for about 10 min to disperse WC nanoparticles in Zn melt. Then the melt was taken out from the furnace and cooled down under Ar gas protection.

Cooling Rates Measurement

To evaluate the effects of the cooling rate, Cu-13 WC samples were then melted again at about 1250° C. under a protection of Argon gas and then cooled down in furnace, air and water. The cooling rate was measured by a K-type thermocouple connected to an Arduino UNO board for recording the cooling curve. The measured cooling rates of furnace cooling, air cooling and water quenching are about 2-4 K/s, about 7-12 K/s and about 70-100 K/s, respectively. The typical cooling curves are shown in FIG. 6, and the thermal arrest caused by the solidification of Cu is clearly identified in the furnace and air cooling curve, while not so noticeable in the water quenching curve due to the rapid heat dissipation.

Structure Characterization

The microstructure, distribution and dispersion of nanoparticles in metals were evaluated by scanning electron microscope (SEM), focus ion beam (FIB) imaging, electron backscatter diffraction (EBSD) and transmission electron microscope (TEM). To clean the surface and reveal the nanoparticles in metal matrix, the mechanically ground and polished as-cast samples were further polished by low-angle ion milling (Model PIPS 691, Gatan). SEM images were acquired at 0° and a tilted about 52° by ZEISS Supra 40VP and FEI Nova 600, respectively. The composition of the material was characterized by energy-dispersive X-ray spectroscopy (EDS).

The volume fraction of nanoparticles was estimated based on the atomic fraction of the major element in the base metal and nanoparticles. Taking advantages of the channeling contrast of different grains induced from the ion beam, FIB imaging was used to reveal the grain structure. Grain size and orientation were evaluated by EBSD (FEI Quanta 3D) at about 30 kV with a current of about 12 nA and FIB with a channel detection electron multiplier (CDEM) detector. The interfaces between matrix and nanoparticles were evaluated by a FEI Titan TEM at about 300 kV. The thin film TEM samples were machined by FIB.

Solidification Behavior

To evaluate the solidification behaviors during cooling, Cu-13 WC and substantially pure Cu samples with the same mass (about 50 mg) were analyzed by Differential Scanning calorimetry (DSC) (TA Instruments, Q600). Samples were heated up to about 1250° C. with a heating rate of about 40° C./min and then cooled down with a cooling rate of about 5° C./min.

High Temperature Stability

To evaluate the UFG, nanocrystalline structures stability at high temperatures, in situ heating Scanning Transmission Electron Microscope (STEM) was conducted at about 300 kV with a convergence angle of about 4.3 mrad and geometric aberrations in the probe corrected to third order. The small convergence angle was chosen to enhance diffraction contrast for grain identification. An FIB milled TEM sample was quickly heated to about 200° C., about 400° C., about 600° C., about 700° C. and about 850° C. by a Gatan in-situ heating holder and held for about 10 min at each step for imaging. The diffraction camera length was increased at higher temperatures in order to enhance diffraction contrast, emphasizing the difference between individual Cu grains for robust identification. The collection semi-angle was about 48-240 mrad at room temperature, about 15-75 mrad at about 400° C., and about 12-60 mrad at about 600° C. and above. Moreover, the bulk solidified Cu-34 vol. % WC (air cooling, about 7 K/s) samples were heated up to about 750° C. (about 1023K, about 0.75 of Tm) and stayed at that temperature for about 2.0 hours under the protection of Ar gas. The grain structure was then evaluated by FIB imaging and EBSD. Mechanical properties were evaluated by microcompression test.

Mechanical characterization and properties

An MTS nanoindenter with a flat punch tip was used for microcompression tests at a strain rate of about 5×10⁻²s⁻¹under room temperature. Micropillars of about 4 μm in diameter and about 9 μm in length were machined by FIB. To evaluate the elastic modulus, microindentation tests with an indent depth of about 2 μm were performed by the same MTS nanoindenter with a Berkovich tip. For each sample, at least ten points were measured. Accordingly micropillars with a diameter and a length of about 4 μm and about 9 μm, respectively, were machined by FIB from the as-solidified (air cooling, cooling rate of about 7 K/s) and heat treated samples (about 750° C. for about 2.0 hours).

Nanoparticle Dispersion and Self-Stabilization Mechanism

Silicon carbide nanoparticles can be dispersed in magnesium matrix by a self-dispersion and self-stabilization mechanism. There are three major factors contributing in this mechanism: (1) Good wetting between molten metal and nanoparticles creates an energy barrier to prevent atomic contact and sintering of nanoparticles in the melt. In the Cu-13 WC system, the energy barrier can be calculated by the following equation:

W_barrier=Sσ_Cucos θ

where S is the effective area and can be calculated by S=πRD₀(D₀=0.2 nm), _σCuis the surface energy of Cu at the processing temperature (about 1.27 J m⁻²), and θ is the wetting angle, e.g. at about 1250° C. (about)10° . In the Cu-13 WC system, an energy barrier of 7×10⁴zJ is obtained; (2) Thermal energy that allow nanoparticles to move randomly in the molten melt by overcoming the attractive van der Waals potential between nanoparticles. A higher thermal energy is desired. The processing temperature for Cu-13 WC is about 1250° C., which provide a thermal energy of E=k_bT=21.0 zJ. (3) A van der Waals potential between nanoparticles to lessen the attraction of nanoparticles from each other to form nanoparticles clusters in molten metals. A small attractive van der Waals potential is desired. It can be calculated by the following equation:

$W_{v d w} = - \frac{{(\sqrt{A_{C u}} - \sqrt{A_{W C}})}^{2}}{6 D} \frac{R}{2}$

where A_Cu=410 zJ and A_WCare the Hamaker constants, D is the distance between two nanoparticles that can be as small as two atomic layer thick (about 0.4 nm), R is the radius of the nanoparticle (about 100 nm). Although the data of A_wcis not available, a 7×10⁴zJ energy barrier is several orders of magnitude higher than thermal energy. Since WC is a conductive ceramic material, A_wcis estimated to be in the range of 200 to 500 zJ, indicating that W_barrierwould always be much higher than W_vdwfor stabilization of dispersed WC nanoparticles in Cu melt.

Growth Restriction Factor by Nanoparticles

For growth restriction by extra solute atoms, the constitutional undercooling can be described by:

$Δ T c = m C_{0} (1 - \frac{1}{{(1 - f_{s})}^{(1 - k)}})$

where ΔT_cis the constitutional undercooling, m is the liquidus slope in a linear phase diagram, C₀is the solute content in the alloy and k is the equilibrium solute partition coefficient, and f_sis the solid fraction solidified. By taking the derivative of this equation respect to f_s, Q can be expressed as:

$Q = m C_{0} (k - 1) = {(\frac{\partial (Δ T_{c})}{\partial f_{s}})}_{f s \to 0}$

This equation indicates that the physical meaning of growth restriction factor is the initial rate of development of constitutional undercooling. Taking Al—Ti as one example, the constitutional undercooling of Al—Ti alloy system is shown in FIG. 12. Ti is the one of the most effective atoms to restrict grain growth in Al. The growth restriction factor is described as Q₁, Q₂and Q₃for different Ti concentration of 0.05, 0.10 and 0.15, respectively. With higher concentration of Ti atoms in Al melt, the growth restriction factor increases. As marked by the black arrow in the FIG. 12, a steeper slope provides larger growth restriction factor. However, the maximum solubility of Ti atoms in Al is 0.15 which constrains Q to reach a maximum value as Q_max.

In contrast, the growth restriction factor, Q_np, introduced by nanoparticle assisted phase control break the fundamental constraint set by Q that depends on constitutional undercooling. As far as the solidification front touches with a nanoparticle with curved shape or a nanoscale channel between nanoparticles, unlike the constitutional undercooling built gradually ahead of the solidification front, an undercooling is immediately established by the Gibbs-Thompson effect. As shown in FIG. 12, the undercooling profile established by Gibbs-Thompson effect is a step function with an infinitely large slope at the initial point. Thus Q_npwhich is the initial rate of the undercooling development by nanoparticles, can be readily increased to a significantly large number, if not infinity.

Strengthening Mechanism

The strengthening from nanoparticles and refined grain structures in the as-solidified samples of Cu-34 vol. % WC is about 647 MPa. The major strengthening mechanisms include Orowan strengthening induced from the populous and dispersed nanoparticles, Hall-Petch effect, and load bearing transfer.

The contribution from Orowan strengthening (Δσ_Orowan) from well-dispersed nanoparticles can be calculated by the following equation:

$Δ σ_{Orowan} = \frac{0.1 3 G_{m} b}{d_{p} [{(\frac{1}{2 V_{p}})}^{\frac{1}{3}} - 1]} \ln \frac{d_{p}}{2 b}$

where G_m, b, V_pand d_pare the shear modulus of the matrix, the Burger vector, the volume fraction and the size of the nanoparticles, respectively. In this example, G_m=46 GPa, b=0.256 nm, V_p=0.34 and d_p=200 nm, and the calculated Δσ_Orowanis 333 MPa. It should be noted that this value is estimated based on ideal dispersion.

The contribution from Hall-Petch effect can be calculated by the following equation:

Δσ_y=kd^−1/2

where d is the grain size and k is a material constant. For Cu, k=0.11 MPa·m, and d=208 nm, and the calculated Δσ_y=246 MPa.

The load bearing strengthening can be calculated by the following equation:

Δσ_load=1.5V_pσ_i

where σ_iis the interfacial bonding strength between nanoparticles and metal matrix. The strong interfacial bonding between Cu and WC nanoparticles contributes in the load bearing mechanism. It can be difficult to estimate the interfacial bonding strength.

Definitions

As used herein, the singular terms “a,” “an,” and “the” may include plural referents unless the context clearly dictates otherwise. Thus, for example, reference to an object may include multiple objects unless the context clearly dictates otherwise.

As used herein, the term “set” refers to a collection of one or more objects. Thus, for example, a set of objects can include a single object or multiple objects.

As used herein, the terms “connect,” “connected,” and “connection” refer to an operational coupling or linking. Connected objects can be directly coupled to one another or can be indirectly coupled to one another, such as via one or more other objects.

As used herein, the terms “substantially” and “about” are used to describe and account for small variations. When used in conjunction with an event or circumstance, the terms can refer to instances in which the event or circumstance occurs precisely as well as instances in which the event or circumstance occurs to a close approximation. When used in conjunction with a numerical value, the terms can refer to a range of variation of less than or equal to ±10% of that numerical value, such as less than or equal to ±5%, less than or equal to ±4%, less than or equal to ±3%, less than or equal to ±2%, less than or equal to ±1%, less than or equal to ±0.5%, less than or equal to ±0.1%, or less than or equal to ±0.05%. For example, a first numerical value can be “substantially” or “about” the same as or equal to a second numerical value if the first numerical value is within a range of variation of less than or equal to ±10% of the second numerical value, such as less than or equal to ±5%, less than or equal to ±4%, less than or equal to ±3%, less than or equal to ±2%, less than or equal to ±1%, less than or equal to ±0.5%, less than or equal to ±0.1%, or less than or equal to ±0.05%.

As used herein, the term “size” refers to a characteristic dimension of an object. Thus, for example, a size of an object that is spherical or spheroidal can refer to a diameter of the object. In the case of an object that is non-spherical or non-spheroidal, a size of the object can refer to a diameter of a corresponding spherical or spheroidal object, where the corresponding spherical or spheroidal object exhibits or has a particular set of derivable or measurable properties that are substantially the same as those of the non-spherical or non-spheroidal object. When referring to a set of objects as having a particular size, it is contemplated that the objects can have a distribution of sizes around the particular size. Thus, as used herein, a size of a set of objects can refer to a typical size of a distribution of sizes, such as an average size, a median size, or a peak size.

As used herein, the term “nanostructure” refers to an object that has at least one dimension in a range of about 1 nm to about 1000 nm. A nanostructure can have any of a wide variety of shapes, and can be formed of a wide variety of materials. Examples of nanostructures include nanofibers, nanoplatelets, and nanoparticles.

As used herein, the term “nanoparticle” refers to a nanostructure that is generally or substantially spherical or spheroidal. Typically, each dimension of a nanoparticle is in a range of about 1 nm to about 1000 nm, and the nanoparticle has an aspect ratio of about 5 or less, such as about 3 or less, about 2 or less, or about 1.

As used herein, the term “nanofiber” refers to an elongated nanostructure. Typically, a nanofiber has a lateral dimension (e.g., a width) in a range of about 1 nm to about 1000 nm, a longitudinal dimension (e.g., a length) in a range of about 1 nm to about 1000 nm or greater than about 1000 nm, and an aspect ratio that is greater than about 5, such as about 10 or greater.

As used herein, the term “nanoplatelet” refers to a planar-like, nanostructure.

Additionally, amounts, ratios, and other numerical values are sometimes presented herein in a range format. It is to be understood that such range format is used for convenience and brevity and should be understood flexibly to include numerical values explicitly specified as limits of a range, but also to include all individual numerical values or sub-ranges encompassed within that range as if each numerical value and sub-range is explicitly specified. For example, a ratio in the range of about 1 to about 200 should be understood to include the explicitly recited limits of about 1 and about 200, but also to include individual ratios such as about 2, about 3, and about 4, and sub-ranges such as about 10 to about 50, about 20 to about 100, and so forth.

While the disclosure has been described with reference to the specific embodiments thereof, it should be understood by those skilled in the art that various changes may be made and equivalents may be substituted without departing from the true spirit and scope of the disclosure as defined by the appended claim(s). In addition, many modifications may be made to adapt a particular situation, material, composition of matter, method, operation or operations, to the objective, spirit and scope of the disclosure. All such modifications are intended to be within the scope of the claim(s) appended hereto. In particular, while certain methods may have been described with reference to particular operations performed in a particular order, it will be understood that these operations may be combined, sub-divided, or re-ordered to form an equivalent method without departing from the teachings of the disclosure. Accordingly, unless specifically indicated herein, the order and grouping of the operations are not a limitation of the disclosure.

NANOSTRUCTURE ASSISTED CASTING OF THERMALLY STABLE, ULTRAFINE GRAINED, NANOCRYSTALLINE METALS

Information

Publication Number

Date Filed

Date Published

Inventors

Original Assignees

CPC

International Classifications

Abstract

Description

Claims

CROSS-REFERENCE TO RELATED APPLICATIONS

Provisional Applications (1)