TECHNICAL FIELD

[0001] The present disclosure is directed generally to powder processing and more particularly to the preparation of quaternary silicon composites from powder mixtures.

BACKGROUND

[0002] Toughened silicon-rich alloys are being developed to compete with conventional and engineered ceramics. It has recently been recognized, for example, that silicon eutectic alloys can be fabricated by melting and casting processes to have properties competitive with technical ceramics. A challenge has been fabricating such alloys with sufficient control over the melting and casting process to achieve a uniform microstructure that exhibits the desired property set.

[0003] It has been hypothesized that amorphous silicon-rich materials could display enhanced mechanical properties compared to their crystalline counterparts due to the lack of long range crystalline order. Such possible improvements could include enhanced fracture toughness, high yield strength, high hardness, and improved ductility. Similarly, nanostructured materials may be advantageous in comparison to their microcrystalline counterparts, as improvements in strength and toughness have been achieved for some compositions.

BRIEF SUMMARY

[0004] A precursor for forming a quaternary silicon composite comprises a mixture of nanocrystalline powders comprising Si, Mn, one of Ti and Zr, and B.

[0005] A quaternary silicon composite prepared by powder consolidation comprises a compact including Si, Mn, one of Ti and Zr, and B. The compact has a relative density of at least about 95%.

[0006] A method of forming a quaternary silicon composite comprises mixing together powders comprising Si, Mn, B, and one of Ti and Zr, and heating the powders to a hold temperature of at least about 700°C. While at the hold

temperature, an external load is applied to the powders, and the powders are consolidated to form a dense compact, thereby forming a quaternary silicon composite. BRIEF DESCRIPTION OF THE DRAWINGS

[0007] FIG. 1 shows X-ray d iff ractog rams of powder blends corresponding to Si- 20Mn-15Zr-15B (top) and Si-20Mn-15i-15B (bottom) milled for 14 hours with overlays of the reflections for Mn ₂ 7Si47, ni ₅ Si26, MnnSiig, and Mn ₄ Si ₇ .

[0008] FIGs. 2A and 2B show thermal profiles acquired by high temperature differential scanning calorimetry (HTDSC) at 10°C/min for the Si-20Mn-15Zr-15B and Si-20Mn-15Ti-15B milled powders, respectively.

The upper line represents the first cycle and the lower line represents the second cycle, and exothermic is up.

[0009] FIG. 3 shows HTDSC isothermal scans of Si-20Mn-15Ti-15B acquired at 500°C, 550X, and 600°C. The temperature ramp rate was 20°C/min until the hold temperature was reached.

[0010] FIG. 4A shows an XRD pattern of Si-20Mn-15Zr-15B following HTDSC analysis up to 800°C.

[0011] FIG. 4B shows an XRD pattern of Si-20Mn-15Ti-15B following HTDSC analysis up to 800°C.

[0012] FIG. 5 is a schematic of an exemplary spark plasma sintering (SPS) system.

[0013] FIG. 6 is a graph of exemplary temperature and load profiles for SPS.

[0014] FIG. 7 shows XRD patterns obtained from milled Si-20 n-15Ti-15B powders after HTDSC analysis (top profile) and after SPS processing into a compacted specimen at 900°C (bottom profile).

[0015] FIG. 8 shows XRD patterns obtained from milled Si-20Mn- 5Zr-15B powders after HTDSC analysis (top profile) and after SPS processing into a compacted specimen at 900°C (bottom profile).

[0016] FIGs. 9A and 9C-9E show EDS maps of a Si-20Mn-15Ti-15B sample consolidated using SPS at a hold temperature of 950°C, and FIG. 9B shows the corresponding SEM image.

FIGs. 10A and 10C-10E show EDS maps of Si-20Mn-15Zr-15B consolidated using SPS at a hold temperature of 950°C, and FIG. 10B shows the corresponding SEM image. DETAILED DESCRIPTION

[0017] Fueled by the recognition that dense compacts comprising silicon (Si), manganese (Mn), titanium (Ti) or zirconium (Zr), and boron (B) may be formed from nanocrystalline precursor powders using suitable consolidation conditions, the present inventors have developed quaternary silicon composites that may exhibit high hardness, high stiffness, and at least moderate fracture toughness when compared with various engineered ceramics and cermets.

[0018] The present disclosure describes these quaternary silicon composites, which are based on the Si-Mn-Ti-B and Si-Mn-Zr-B systems, as well as the precursor powders from which they may be prepared. Also described herein is a consolidation method that may be employed to prepare dense compacts of the quaternary silicon composites, which are shown to attain relative densities in excess of 95% as well as promising mechanical properties.

[0019] It is noted that the terms "comprising," "including" and "having" are used interchangeably throughout the specification and claims as open-ended transitional terms that cover the expressly recited subject matter alone or in combination with un recited subject matter.

Quaternary Silicon Composite

[0020] A quaternary silicon composite prepared by powder consolidation

comprises a compact including Si, Mn, one of Ti and Zr, and B and having a relative density of at least about 95%.

[0021] The compact may have a phase composition that includes one or more silicide compounds comprising Si and a boride compound comprising B. The one or more silicide compounds may include a disilicide selected from the group consisting of ZrSi ₂ and MnTiSi ₂ . The boride compound may comprise Ti or Zr. For example, the boride compound may be a diboride selected from the group consisting of ZrB ₂ and TiB ₂ .

[0022] The one or more silicide compounds may also include a manganese silicide having a formula Mn _x Si _y , where 1.71< ~≤1.75. It is believed that, during consolidation, the manganese silicide diffuses more readily than the other silicide and boride phases, and thus may act as a sintering aid, as discussed in greater detail below. Accordingly, the densified compact may have a polycrystalline structure including grains separated by grain boundary regions that include the manganese silicide. In one example, the manganese silicide may comprise Mn ₄ Si ₇ , where x is 4 and y is 7.

[0023] The elements Si, Mn, one of Ti and Zr, and B may be present in the compact at the following concentrations: from about 40 at.% to about 60 at.% Si; from about 10 at.% to about 30 at.% Mn; from about 5 at.% to about 25 at.% Ti or Zr; and from about 5 at.% to about 25 at.% B.

[0024] For example, the concentration of Si may be about 50 at.%; the

concentration of Mn may be about 20 at.%; the concentration of the Ti or Zr may be about 15 at.%, and the concentration of B may be about 15 at.%. The composition of such a quaternary silicon composite may be abbreviated as Si-20Mn-15Zr-15B or Si- 20Mn-15Ti-15B, depending on whether Ti or Zr is present in the compact.

[0025] The microstructure and properties of the quaternary silicon composite are described in greater detail below, after a discussion of Si-Mn-Ti/Zr-B precursor powders and a method of making the composite.

Precursor Powders for the Quaternary Silicon Composite

[0026] A precursor for forming a quaternary silicon composite comprises a mixture of nanocrystalline powders that includes Si, Mn, one of Ti and Zr (Ti or Zr), and B. The Si may be present in the mixture at a concentration of from about 40 at.% to about 60 at.%; the Mn may be present at a concentration of from about 10 at.% to about 30 at.%; the one of Ti and Zr may be present at a concentration of from about 5 at.% to about 25 at.%; and the B at a concentration of from about 5 at.% to about 25 at.%.

[0027] For example, the concentration of Si in the mixture of nanocrystalline powders may be about 50 at.%; the concentration of Mn may be about 20 at.%; the concentration of Ti or Zr may be about 15 at.%; and the concentration of B may be about 15 at.%. Powders having this composition may be referred to as Si-20Mn-15Zr- 15B or Si-20Mn-15Ti-15B powders, depending on whether Ti or Zr is used in the mixture.

[0028] The mixture of nanocrystalline powders may be formed by mechanical alloying, which is sometimes referred to as mechanical milling. Mechanical alloying is a solid-state materials processing technique that involves repeated particle fracturing and welding events in a mechanical milling apparatus, such as a high energy ball mill. Blends of elemental and/or alloy powders may be mixed together and milled to create a milled powder of the desired composition. The mixture of nanocrystalline powders resulting from milling may be partially or fully nanocrystalline.

[0029] The nanocrystalline powders may have an average primary particle size ranging from about 1 nm to about 100 nm. The average primary particle size may also be about 50 nm or less, about 20 nm or less, or about 10 nm or less, and is typically at least about 5 nm. For example, the average primary particle size may be from about 5 nm to about 8 nm

[0030] For example, the nanocrystalline powders may be produced from elemental powder blends corresponding to Si-20Mn-15Zr-15B and Si-20Mn-15Ti-15B. The powder blends may be milled continuously for 14 hours using a SPEX 8000D mill with tungsten carbide milling vials and bearings. FIG. 1 shows x-ray diffraction (XRD) profiles for exemplary Si-20 n-15Zr-15B (top) and Si-20Mn-15Ti-15B (bottom) milled powders prepared in this manner. Both XRD patterns display obvious reflections (peaks) with broadening indicative of a nanocrystalline particle size. Defects and lattice strain may also contribute to peak broadening, however, and thus there are challenges with making a precise quantitative determination of an average primary particle size (or grain size) from the XRD data. '

[0031] Along with XRD, the milled powders may be analyzed using high

temperature differential scanning calorimetry (HTDSC). Calorimetry can be useful to distinguish between amorphous and nanocrystalline materials. FIGs. 2A and 2B display thermal profiles for each quaternary system. From the thermograms, which were acquired at 10°C/min up to 800°C, a striking difference between the two systems can be observed. Si-20Mn-15Ti-15B undergoes a small exothermic transition at about 530°C while Si-20Mn-15Zr-15B undergoes no apparent enthalpic changes. The lack of any change in the zirconium-containing powder sample immediately corroborates the nanocrystalline nature of the material, assuming all transitions occur below 800°C.

[0032] As for the titanium-containing sample, the possibility that the exothermic event might be representative of a recrystallization event was explored; however, the enthalpic change is quite weak at just 5.8 J/g and there is no obvious or subtle indication of a glass transition prior to the event/Typical glass forming alloys display exothermic recrystallizations on the order of 25 J/g or greater and are accompanied by subtle if not obvious glass transitions. Additionally, isothermal HTDSC scans acquired at temperatures below, close to, and above the exothermic transition of the Si-20Mn- 15ΤΊ-15Β compound did not yield any evidence of an enthaipic change. FIG. 3 displays the collected isothermal scan profiles. The lack of an exothermic enthaipic signal indicates that grain growth is occurring without any nucleation event. This was further indication of the nanocrystallinity of the milled powders. Thus, HTDSC corroborated the nanocrystalline nature of both of the milled blends.

[0033] The most likely dominant crystalline species in either sample has been identified as a higher order manganese silicide (Mn _x Si _y , where 1 .71< ^ <1.75). FIG. 1 shows the XRD patterns of the milled powders along with the expected reflections of Mn Si ₇ , MniiSi-i 9, Mni ₅ Si26, and Mn ₂ 7Si ₄₇ . There is significant overlap of the major crystallographic reflections of the various higher order manganese silicides, which leads to some ambiguity in positively identifying the phase(s). Apparent peak shifts in the XRD data relative to the expected reflections may be attributed to crystalline defects and lattice strain introduced during milling.

[0034] Following HTDSC, the heated samples were analyzed via XRD in order to obtain additional compositional information. FIGs. 4A and 4B depict the obtained diffraction profiles for both Si-20Mn-15Zr-15B and Si-20Mn-15Ti- 5B heated to 800°C, respectively. Based on the patterns, it is fairly evident that significant grain growth or annealing occurred in both specimens. Peaks are much sharper and there are significantly more reflections than were detected prior to heating. In the case of the zirconium-containing sample, identified species include Mn ₄ Si ₇ , ZrSi ₂ , and ZrB ₂ . As for Si-20 n-15Ti-15B, Mn ₄ Si ₇ , TiB ₂ , MnTiSi ₂ , and residual silicon are observed. It is possible the newly identified species were present prior to heating but due to crystallite size or defect density they were not detectable, with the exception of the manganese silicide, which is believed to be Mn ₄ Si ₇ . Only after annealing, which allowed defects to migrate and grains to grow, could the additional phases be observed by XRD. Thus, it is believed the milled powders include microcrystalline grains of Mn ₄ Si ₇ and nanocrystalline grains of the various transition metal diborides and disilicides.

[0035] The presence of zirconium and titanium refractory compounds as well as residual silicon indicates that consolidation of the milled powders may be challenging. These constituents are hard and thermally stable materials with melting points above 1400°C and hardness values greater than 10GPa. Table 1 shows the melting points and hardness values of these species. Sintering typically requires temperatures of 0.5-0.75 times the melting point of the material being studied in order to activate the diffusion process and achieve successful consolidation. Hard materials present added difficulties such as minimal creep behavior which may reduce the effectiveness of pressure assisted sintering. In the case of T1B ₂ and ZrB ₂ , the difficulties associated with fabricating bulk materials via powder metallurgical routes have been recognized. It is because of their high thermal stability and hardness values that sintering aids, materials with lower melting points, are often used to consolidate the powders at more realistically achievable temperatures and loads.

Table 1. Melting Points of Zirconium and Titanium

Diborides and Disilicides and Silicon

Method of Preparing the Quaternary- Silicon Composite

[0036] A method of forming a quaternary silicon composite may include mixing together powders comprising Si, Mn, B, and one of Ti and Zr, and heating the powders to a hold temperature of at least about 700°C; while at the hold temperature, an external load may be applied to the powders, thereby consolidating the powders to form a dense compact.

[0037] The powders may be contained in an electrically conductive die for the consolidation, and the heating may comprise applying a pulsed DC current to the die, as shown schematically in FIG. 5. When a combination of rapid joule heating and an external load pressure are used to consolidate the powders, the method may be referred to as spark plasma sintering (SPS), field assisted sintering technique (FAST) and/or pulsed electric current sintering (PECS). Higher temperatures and greater thermal uniformity may be achieved in shorter amounts of time using this approach. Heating rates in excess of 100°C/min are achievable with SPS, whereas conventional hot pressing techniques typically achieve maximum rates of only 50°C/min or less.

[0038] The external load applied to the powders at the hold temperature may be about 100 MPa or less, or about 75 MPa or less, or about 50 MPa or less. Typically, the external load is at least about 30 MPa.

[0039] The resulting compact may have a relative density of at least about 95%, or at least about 97%. The relative density may be determined by comparing a density measured for the compact (compact density) with the theoretical or actual density of a fully dense reference sample comprising the same elements and/or having the same composition (reference density), where the relative density is equivalent to the ratio of the compact density to the reference density. For example, for compacts having compositions of Si-20Mn-15Zr-15B and Si-20Mn-15Ti-15B, reference density values of 4.954 g/mL and 4.501 g/mL, respectively, may be used.

[0040] To avoid melting of any phases during sintering, the hold temperature is generally kept at less than 1 100°C. For example, the hold temperature may be no more than about 1050°C, or no more than about 1000°C. To ensure that sufficient diffusion occurs for successful consolidation, the hold temperature may also be least about 750°C, at least about 800°C, at least about 850°C, or at least about 900°C. For example, the hold temperature may be in the range of from about 700°C to about 950°C.

[0041] The powders may be heated to the hold temperature at a rapid rate of at least about 100°C/min, at least about 130°C/min, or at least about 150°C/min.

Typically, the heating rate is no higher than about 200°C/min. For example, the powders may be heated to the hold temperature at a heating rate of from about 130°C/min to about 170°C/min. In one example, the heating rate is about 150°C/min.

[0042] The silicon powders being consolidated may include Si at a concentration of from about 40 at.% to about 60 at.%; Mn at a concentration of from about 10 at.% to about 30 at.%; one of Ti and Zr at a concentration of from about 5 at.% to about 25 at.%; and B at a concentration of from about 5 at.% to about 25 at.%. For example, the powders may have a composition of Si-20Mn-15Ti-15B or Si-20Mn-15Zr-15B. The compact formed by the above method may have the same composition as set forth for the powders.

[0043] A set of experiments to consolidate Si-20Mn-15Zr-15B and Si-20Mn-15Ti- 15B milled powders was carried out on a commercially available SPS device (Thermal Technology LLC, Santa Rosa, CA). 16 samples were compacted under vacuum at a heating rate of 150°C/min and a load of 50MPa. Hold temperatures of 700°C, 900°C, 950°C, and 1 100°C were used. The duration of time at the hold temperature was 5 minutes. FIG. 6 depicts an exemplary processing profile obtained at 900°C. All powders were sieved to 63 pm (230mesh) or less before compaction. Compaction of the quaternary milled samples at hold temperatures in the range of 700-950°C resulted in the greatest degree of densification. Relative densities of 95+% for the titanium-containing samples and 96+% for the zirconium-containing samples were achieved under these conditions. Table 2 lists the relative densities and associated SPS conditions for the consolidated specimens. A hold temperature of 1 100°C resulted in liquid material leaking from the die and punch which indicated melting under those conditions.

Table 2. Consolidation Conditions and Relative Density Values for Powder Compacts

Si-20Mn-15Zr-15B Si-20Mn-15Ti-15B

Temperature (°C) 700 900 950 700 900 950

Load (MPa) 50 50 50 50 50 50

Density (g/mL) 4.7621 4.869 4.9732 4.296 ' 4.3476 _t 4.3542

Relative Density

(%) 96.13 98.28 100.39 95.45 96.59 96.74

Characterization and Testing of Quaternary Silicon Composites

[0044] Following densification trials, compacts prepared at 950°C were

characterized via scanning electron microscopy (SEM) and XRD. FIGs. 7 and 8 show the XRD patterns of the quaternary Si-20Mn-15Ti-15B and Si-20Mn-15Zr-15B consolidated specimens (bottom) in comparison with the XRD patterns obtained from the milled powders after HTDSC analysis (top). FIGs. 9A-9E and 10A-10E show respective SEM (B) and energy dispersive spectroscopy (EDS) images (A.C-E) of the compacts.

[0045] Similar to the post-HTDSC samples, the compacted samples are

significantly more crystalline than the milled material. In fact, the x-ray patterns for the consolidated and HTDSC analyzed specimens, as shown in FIGs. 7 and 8, are almost identical for each of the respective compositions. As discussed above, the resolved, sharp reflections detected in each sample are indicative of either or both annealing and grain growth that occurred during heating; consequently, the nanocrystalline structure of the milled powders was not maintained.

[0046] From the EDS maps, it appears that Mn ₄ Si ₇ aggregates at the boundaries of larger grains which are comprised of a mixture of the various disilicide and diboride phases. Due to equipment limitations, boron could not be detected via EDS, and accurate compositional assignment of the larger grains could not be completed using the technique. The aggregation of Mn ₄ Si ₇ suggests the manganese silicide is the most thermally active material and diffuses quickly under the studied conditions. The more thermally stable components contained within the larger grains diffuse much more slowly and consequently these regions remain more intimately mixed, larger in size, and irregularly shaped.

[0047] Following compositional and microstructurai characterization, several mechanical properties of the compacts were evaluated using a variety of techniques. Table 3 summarizes the assessed physical properties and their associated values. The Si-20Mn-15Zr-15B compact was moderately harder and more wear resistant than the titanium-containing compact; however, the titanium-containing compact exhibited a slightly higher fracture toughness. Both compacts displayed mechanical properties superior to silicon and many engineered ceramics.

Table 3. Assessed mechanical properties of Si-20Mn-15Zr-15B and Si-20Mn-15Ti- 15B compacts consolidated at 950X

precisely as the ASTM standard requires. Due to sample size, a smaller diameter wear track had to be used which prevents direct comparison of the assessed values with other known materials.

[0048] From nanoindentation, the modulus of these compacts was found to be 350-400 GPa. Nanoindentation analysis using a Berkovich indenter revealed hardness values of 10-14 GPa; whereas, the bulk hardness examined via Vicker's indentation was found to be 12-14 GPa. In these specific cases, the bulk hardness value is probably more representative of these materials due to their composite-like structure. Nanoindentation samples only a very small volume of material which may not necessarily be representative of the bulk.

[0049] Wear volumes of 8.36x10 ⁷ pm ³ and 1 .04x10 ⁸ m ³ and coefficients of friction of 0.7828 and 0.7631 were found for the zirconium- and titanium-containing compacts, respectively. These values were determined following ASTM G99-05 and using a pin on disk testing set-up with a tungsten carbide pin; however, the overall travel path diameter was narrower than required by the standard and thus a direct comparison with other materials is not presented.

[0050] The fracture toughness of the compacts was determined using the Vicker's indentation fracture toughness (VI F) test developed by Anstis, as described in Anstis, G.R., et al., "A Critical Evaluation of Indentation Techniques for Measuring Fracture Toughness: I, Direct Crack Measurements, " Journal of the American Ceramic Society, 1981. 64 (9): p. 533-538. Although the VI F is designed for homogeneous materials and not composites, the indent sample volume is significantly larger than the size scale of the composite microstructure; thus, it is believed that a fairly representative VI F test can be conducted. The fracture toughness of the compacts was found to be about 2.5-3 MPaVm.

[0051] In summary, SPS was successfully employed to consolidate the thermally stable, nanocrystalline materials prepared from milling elemental blends

corresponding to Si-20Mn-15Zr-15B and Si-20Mn-15Ti-15B. Relative densities of 95+% were achieved using the conditions described above, and it is believed that sintering was assisted by the increased diffusion of Mn ₄ Si ₇ . The final compacts did not retain the nanocrystalline structure of their powder precursors due, presumably, to the elevated temperature conditions required to achieve densification. In addition to successful consolidation, several mechanical properties of the compacts were assessed, as summarized in Table 2. Each of the compacts out performed silicon in all examined categories except for wear (which could not be accurately analyzed per the ASTM standard as explained above). When compared to other engineering materials, the quaternary silicon composites display high hardness, high stiffness, and moderate fracture toughness. The bulk hardness of these materials (12-14 GPa) is comparable to tungsten carbide cermets, stabilized zirconia, alumina, as well as other engineered ceramics. The stiffness of the quaternary silicon composites (350-400 GPa) is on the high end of the class of engineered ceramics and cermets. The fracture toughness of the composites (2.5-3 MPaVm) is not as competitive, but still exceeds values for silicon.

Experimental Details

Mechanical Alloying (MA)

[0052] All milling experiments, including loading and collection, were conducted inside in an inert argon environment inside a glove box. The oxygen concentration inside the glove box was less than 10 ppm. Mechanical alloying was performed using a SPEX 8000D (double container mill). Studies on the multi-component systems were carried out in tungsten carbide lined vials (SPEX part# 8004) with tungsten carbide ball bearings (diameter = 7/16 inches). The ball bearing to powder weight ratio was kept constant throughout experimentation at -10:1. In all case ~3 g of powder/starting materials was combined with ~31 g of WC ball bearings. All milling experiments were performed continuously over 14 hours.

Spark Plasma Sintering (SPS)

[0053] Powder compaction was conducted using a Pulsed Electric Current

Sintering (PECS) system, Model SPS 10-3, manufactured by Thermal Technology, LLC, and installed at Michigan State University. All powders were sieved to 63 μιη or less prior to compaction. Tests were conducted under vacuum (2-3x10 ^"3 Torr), using a16mm graphite die and punch set. 0.005 inch thick graph-foil was used to line punch and dies. A heating rate of 150°C/min was applied with a hold time of 5 min at the specified maximum temperature (hold temperature). Hold temperatures were 700, 900, 950, and 1 100°C. The maximum load applied was 50 MPa with a hold time of 4 min. The load was ramped at a rate of 45 MPa/min.

X-Ray Diffraction (XRD)

[0054] Powder diffraction patterns were collected in Bragg-Brentano geometry from 20 to 80° 2d in 0.02° increments at 0.57minute with a Cu anode operating at 40 kV and 44 mA. A 10 mm height limiting slit, 1/2° divergence slit, 1/2° scattering slit, and 0.6 mm receiving slit were used, and intensity data were collected with a scintillation counter equipped with a diffracted beam monochromator

High Temperature Differential Scanning Calohmetry (HTDSC)

[0055] DSC thermograms were collected using isothermal and non-isothermal heating rates separately. Non-isothermal studies were conducted from 50°C to 800°C at heating rates of 10°C/min. Isothermal analyses were conducted using a heat rate of

20°C/min to reach hold temperature. All studies were conducted under an Ar flow of

70 mL/min and used sample masses of 25-35 mg.

Density

[0056] Density analyses were performed using an AccuPyc II 1340 gas

pycnometer manufactured by Micromeritics. Samples were analyzed in a 0.1 cm ³ sample cup. At least 10 replicates per sample were conducted. Measurements were taken at 24.05°C.

Nanoindentation

[0057] Nanoindentation was performed using a Hysitron Triboindentor equipped with a Berkovich identor tip. Information on hardness and reduced modulus was obtained from the load vs. displacement data acquired from each indentation.

Samples consisted of 1 1 mm diameter, thin cylinders. Indents were preformed on the Triboindentor under Load Control. The Load function was linear and consisted of a 10 second loading segment and a 10 second unloading segment. An indentation depth of 100 microns was used for hardness calculations. The indentor was calibrated using a fused quartz standard. The standards' experimental modulus value was 70.83 GPa (69.6 = theoretical value).

Wear Testing

[0058] Wear testing was performed using the Bruker UMT-3 Wear Tester equipped with a ball-on-disk fixture. Tests were conducted for a specified sliding distance at a set load and speed. Samples consisted of small, 1 1 mm diameter, thin cylinders. Test conditions used are shown in Table 2. Wear information (wear volume, coefficient of friction, etc.) was obtained using a modified version of ASTM G99-05, Standard Test Method for Wear Testing with a Pin-on-Disk Apparatus.

Table 4. Wear testing conditions and material Wear Test Conditions

Relative Humidity ambient

Temperature ambient

3/8 inch Tungsten

Pin Material

Carbide Bearing

Force ION

Sliding Distance 100m

Velocity 100 re /min

Interferometry

[0059] After wear testing, the samples were analyzed using a Zygo NewView 7300 interferometer. Images were collected using a 5x objective coupled with a 0.5x optical zoom, which resulted in an image with dimensions of 2.83 x 2.12 mm. The images were fitted as needed to flatten the data. Data fill was ON to fill in missing data points. Indentation

[0060] Indentation studies were conducted with a Leco LM247AT microhardness tester using a Vickers pyramidal diamond indenter. Optical micrographs were captured through a 50x objective with a 2.0 megapixel digital microscope camera (PAXcam2). Hardness values were assessed from post-indentation micrographs using Confident Hardness Tester software (v. 2.5.2). Indent crack profiles were analyzed using Stream Motion image analysis software from Olympus.

Optical Microscopy

[0061] Optical micrographs were acquired using an Olympus BX51 or GX51 microscope equipped with 2.5x, 5. Ox, 10x, 2 Ox, 50x and 100x objectives. Images were captured with a 5.0 megapixel digital microscope camera (Olympus UC50).

Micrographs were analyzed using Stream Motion image analysis software from

Olympus.

Scanning Electron Microscopy/ Energy Dispersive Spectroscopy

[0062] Samples were analyzed without coating in the JEOL JSM-0601 OLA SEM. Images and spectra were collected using a 15-20 kV accelerating voltage. EDS maps were collected for 6 minutes per location. Spectra were acquired in spot mode for 10 seconds per location.

[0063] Although the present invention has been described in considerable detail with reference to certain embodiments thereof, other embodiments are possible without departing from the present invention. The spirit and scope of the appended claims should not be limited, therefore, to the description of the preferred

embodiments included here. All embodiments that come within the meaning of the claims, either literally or by equivalence, are intended to be embraced therein.

[0064] Furthermore, the advantages described above are not necessarily the only advantages of the invention, and it is not necessarily expected that all of the described advantages will be achieved with every embodiment of the invention.