METHOD OF LASER TREATING AN ALUMINA SURFACE

Abstract

The method of laser treating an alumina surface includes applying a coating of a phenolic resin including particles of titanium carbide (TiC) and boron carbide (B4C) to an alumina (Al2O3) surface to form a resin-coated alumina surface, heating the resin-coated alumina surface to form a carbon-coated alumina surface, and scanning the carbon-coated alumina surface a nitrogen gas-assisted CO2 laser beam to form a laser-treated surface. The particles of titanium carbide (TiC) and boron carbide (B4C) each have a diameter of about 350 nm.

Description

CROSS-REFERENCE TO RELATED APPLICATION

This application claims the benefit of U.S. Provisional Patent Application Ser. No. 62/137,213, filed on Mar. 23, 2015.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to laser treatment of metal surfaces, and particularly to a method of laser treating an alumina surface, thus yielding a hardened wear-resistant surface.

2. Description of the Related Art

Alumina (Al₂O₃) tiles find wide application in industry due to tribological and thermal properties, such as corrosion resistance and thermal stability at high temperatures. Alumina is widely used in lined process piping, chutes, cyclones, lined metal fabrication, and grinding mill components. This is due to the fact that all of these applications require a unique combination of hardness, extremely high abrasion resistance, and high strength over a broad range of temperatures.

Alumina tiles are formed from fine alumina powders, e.g., microsized alumina powders, through sintering. The tiles have high porosity, particularly in the surface region. Depending on the powder size, in some cases, structural non-homogeneity and abnormalities, such as scattered small voids, are formed in the tiles. These abnormalities in the structure, particularly in the surface region, can be minimized through controlled laser melting.

Nitriding is the process of adding nitrogen to the surface of metals for improved hardness and wear resistance. A nitride coating may be applied to many different metals, including steel and other ferrous metals, titanium, molybdenum, and aluminum. Conventional gas nitriding generally uses ammonia, which dissociates to form nitrogen and hydrogen when brought into contact with a heated metal workpiece. Laser gas assisted nitriding involves exposing a workpiece to laser radiation for comparatively brief periods in the presence of nitrogen. Nitrogen is supplied under pressure. Laser nitriding has several advantages, including the ability to apply a very thin coating, the ability to apply the coating in a very narrow beam, if desired, relatively low temperatures and pressures to avoid deformation of the metal, and quick processing times, with exposure to radiation often being less than a second.

Thus, a method of laser treating an alumina surface addressing the aforementioned problems is desired.

SUMMARY OF THE INVENTION

The method of laser treating an alumina surface includes forming a carbon film on the alumina surface and laser treating the carbon-coated alumina surface. The carbon film can include particles of titanium carbide (TiC) and boron carbide (B₄C). The alumina surface with the carbon film formed thereon is then scanned with a nitrogen gas-assisted CO₂ laser beam in order to form a laser-treated surface. The laser-treated surface can include aluminum nitride (AlN) and/or aluminum oxynitride (AlON) compounds. The resultant laser-treated surface has increased hardness.

These and other features of the present invention will become readily apparent upon further review of the following specification and drawings.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a graph showing X-ray diffraction results of an alumina surface prepared by a method of laser treating an alumina surface according to the present invention, specifically showing d spacing thereof as a function of sin² ψ, where ψ is the tilt angle.

FIG. 2A is an optical photograph of an alumina surface prepared by the method of laser treating an alumina surface.

FIG. 2B is an optical photograph of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing laser scanning tracks.

FIG. 2C is a scanning electron micrograph of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing overlapping of laser irradiated spots.

FIG. 2D is a scanning electron micrograph of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing the formation of a small void and a pair of closely positioned hard particles at the surface.

FIG. 2E is a scanning electron micrograph of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing the formation of a semi-embedded hard particles and fine grains at the surface.

FIG. 3A is a scanning electron micrograph showing a cross-section of a laser-treated layer of the alumina surface prepared by the method of laser treating an alumina surface.

FIG. 3B is a scanning electron micrograph showing a cross-section of a laser-treated layer of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing the formation of a dense layer with small voids formed therein.

FIG. 3C is a scanning electron micrograph showing a cross-section of a laser-treated layer of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing the formation of small columnar structures.

FIG. 3D is a scanning electron micrograph showing a cross-section of a laser-treated layer of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing the formation of large columnar structures.

FIG. 3E is a scanning electron micrograph showing a cross-section of a laser-treated layer of the alumina surface prepared by the method of laser treating an alumina surface, specifically showing the formation of large grains.

FIG. 4 is an X-ray diffractogram of the alumina surface prepared by the method of laser treating an alumina surface compared against the X-ray diffractogram of an untreated alumina surface.

Similar reference characters denote corresponding features consistently throughout the attached drawings.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

A method of laser treating an alumina surface includes forming a carbon film on the alumina surface to provide a carbon-coated surface and laser treating the carbon-coated surface to provide a laser-treated surface. The alumina surface can be a surface of an alumina tile. The carbon film can include at least two chemically different hard particle types. The hard particles can include titanium carbide (TiC) and boron carbide (B₄C), for example. The hard particles can have a particle size of about 350 nm. The carbon film can include titanium carbide (TiC) and boron carbide (B₄C) in equal proportions. The carbon film can be about 40 μm thick. The carbon-coated surface can then be scanned with a nitrogen gas-assisted CO₂ laser beam, thereby converting alumina at the surface to aluminum nitride (AlN) and/or aluminum oxynitride (AlON), with carbon dioxide escaping from the irradiated surface. The present method enhances the microhardness and lowers the fracture toughness of the tile surface.

Laser treating the coated alumina surface melts the surface of the alumina tile to provide thermal integration of the powders forming the sintered tiles. During the laser melting process, use of an assisting gas, e.g., nitrogen, can prevent excessive oxidation reactions taking place in the irradiated region. The use of nitrogen in the laser melting process can also prevent surface defects that would otherwise result from high temperature oxidation reactions.

A phenolic resin and hard particle mixture can be provided to form the carbon film at the alumina tile surface. The phenolic resin and hard particle mixture can be prepared by adding a mixture of at least two chemically different hard particle types, e.g., titanium carbide (TiC) and boron carbide (B₄C), to water dissolved phenolic resin. Hard particles can include any suitable powder with high melting points and high hardness. Other examples of hard particles include ZrO₂, TiN, and VC. The phenolic resin and hard particle mixture is applied to the alumina tile surface and heated under pressure, e.g., 175° C. for 2 hours under 8 bar, then 400° C. in an argon environment for 6 hours, to form the carbon film.

In experiment, 3 mm thick alumina (Al₂O₃) tiles were used. The TiC and B₄C powders were mixed in the ratio of 3 wt % of TiC and 3 wt % B₄C, with homogeneous mixing, to form a hard particle mixture. The TiC and B₄C powders had particle sizes of about 350 nm. A phenolic resin was dissolved in water then mixed with the hard particle mixture to form a phenolic resin and hard particle mixture. The phenolic resin and hard particle mixture was applied uniformly on the alumina tile workpiece surface. The workpiece was then placed in a control chamber set at 8 bar pressure and 175° C., for two hours. The use of high pressure ensured evaporation of water and drying of the phenolic-particle mixture at the workpiece surface. The workpieces were then heated to 400° C. in an argon environment for six hours to ensure the conversion of the phenolic resin into carbon, and form a 40 μm thick carbon film at the alumina tile surface.

A CO₂ laser delivering a nominal output power of 2 kW was used to irradiate the sample surface including the carbon film. The delivery of the laser output power was in the form of repeated pulses and the maximum frequency of the laser pulse repetition was 1500 Hz. The proper combination of the laser beam scanning speed and the pulse repetition rate provides controlled melting of the surface. A carbon film, uniformly hosting the hard particles, was formed prior to the laser scanning, as described above. This enhanced the absorption of the incident beam at the workpiece surface. It should be noted that high reflection of the incident beam from the untreated surface resulted in poor melting characteristics in the irradiated region. This was because of the wavelength of the laser beam, which was 10.6 μm. In addition, the operational cost of the CO₂ laser was lower than those of the other high power laser sources. Therefore, the CO₂ laser was selected to irradiate the pre-prepared workpiece surfaces.

The nominal focal length of the focusing lens was 127 mm. The laser beam diameter focused at the workpiece surface was ˜0.25 mm. Nitrogen assisting gas emerged from a conical nozzle, co-axially with the laser beam. It should be noted that the laser surface ablation/melting process was carried out with a variety of laser parameters. During the treatment process, laser parameters resulting in a minimum of surface defects, such as very small cavities with no cracks or crack networks, were selected. Optimal laser parameters resulting in minimum surface defects, e.g., very small cavities with no cracks, included a scanning speed of 10 cm/s, a peak power of 2 kW, a frequency of 1.5 kHz, a nozzle gap of 1.5 mm, a nozzle diameter of 1.5 mm, a focus setting of 127 mm, and an N₂ pressure of 600 kPa.

Material characterization of the laser treated surfaces was conducted using an optical microscope, a scanning electron microscope (SEM), energy-dispersive spectroscopy (EDS), and X-ray diffraction (XRD) analysis. Copper tapes were used to ensure contact between the workpiece surface and the sample holder during the SEM micro-graphing. X-ray diffraction (XRD) analysis was carried out (Cu—Kα; λ=1.5406 Å) using XRD equipment with a Bragg-Brentano geometry arrangement. A typical setting of the XRD equipment was 40 kV and 30 mA, and the scanning angle (2θ) ranged from 20° to 80°. Surface roughness measurement of the laser-melted surfaces was performed using a scanning probe microscope in contact mode. The tip was made of silicon nitride probes (r=20-60 nm) with a force constant, k, of 0.12 N/m.

A microphotonics digital microhardness tester was used to obtain Vickers micro-indentation hardness values at the ablated surface. The standard test method for Vickers indentation hardness of advanced ceramics (ASTM C1327-99) was adopted. Microhardness was measured at the workpiece surface after the laser ablation/melting process. The measurements were repeated five times at each location for consistency of the results.

The XRD technique was used to measure the residual stresses in the surface region of the laser ablated surface. The XRD technique provided data in the surface region of the specimens due to the low penetration depth of Cu—Kα radiation into the treated layer; i.e., the penetration depth was on the order of a few μm. Measurements relied on the stresses in fine grained polycrystalline structure, and the position of the diffraction peaks exhibited a shift as the specimen was rotated by an angle ψ. The relationship between the peak shift and the residual stress, σ, is given by

$\sigma =\frac{E}{\left(1+v\right){\sin }^2\psi }·\frac{\left(d_n-d_0\right)}{d_0},$

where E is Young's modulus, v is Poisson's ratio, ψ is the tilt angle, and d_n are the d spacing measured at each tilt angle.

If shear strains are not present in the specimen, the d spacing changes linearly with sin² ψ. d₀ is the inter-planar spacing when the substrate material is free from stresses. FIG. 1 shows the linear dependence of d(311) on sin² ψ in the region of the laser cut surface. The γ-Al₂O₃ peak (ICDD 29-1486) is obtained at 37.5°, which corresponds to the (311) plane with an inter-planar spacing of 0.239 nm. The linear dependence of d(311) has a slope of −2.4761×10⁻¹² m/deg and an intercept of 0.239 nm, as shown in FIG. 1. The residual stress determined by the XRD technique in the vicinity of the surface is on the order of −3.1±0.09 GPa. XRD measurements were repeated three times and the error related to the measurements was on the order of 3%.

The fracture toughness of the surface was measured using the indenter test data for microhardness (Vickers) and crack inhibition. In this case, microhardness (in HV) and the crack length generated due to indentation at the surface were measured. The length of the cracks measured, l, corresponded to the distance from the crack tip to the indent. The crack lengths were individually summed to obtain Σl. The crack length, c, from the center of the indent is the sum of individual crack lengths, Σl, and half the indent diagonal length, 2a. Therefore, c=a+Σl. However, depending on the ratio of c/a, various equations were developed to estimate the fracture toughness, K. The following equation was used to determine the fracture toughness, K, and is applicable for 0.6≦c/a≦4.5:

$K_c=0.079{\left(\frac{P}{a}\right)}^{1.5}\log \left(4.5P·\frac{a}{c}\right),$

where P is the applied load on indenter, c is the crack length, and a is the half indent diagonal length. Table 1 below gives fracture toughness data.

TABLE 1
Fracture Toughness for Treated and Untreated Surfaces
	Fracture Toughness	P	a	c
	(MPa {square root over (m)})	(N)	(μm)	(μm)
Untreated surface	3.53 ± 0.2	5	45	60
Treated surface	3.05 ± 0.4	5	45	90

FIGS. 2A-2E show optical photographs and SEM micrographs of the laser treated surface. The laser treated surface is free from asperities, such as large scale cracks and voids. However, a few locally scattered small cavities were observed at the surface, which is attributed to the local evaporation of the surface during the laser treatment process. Although laser power intensity was set at a level to avoid evaporation at the surface, laser peak intensity at the irradiated spot was high and the contribution of local exothermic reactions to the energy density available at the surface caused local evaporation. It should be noted that laser power intensity distribution is Gaussian at the surface, which in turn results in peak intensity at the center of the irradiated spot. Nevertheless, localized surface evaporation occurs only in a relatively small area and the locations are scattered on the surface.

Laser scanning tracks are visible on the treated surface and they are formed by the repetition of laser pulses during the laser scanning of the workpiece surface (as seen in FIG. 2B). It should be noted that the pulse repetition rate of laser irradiation is 1500 Hz, which results in a 76% overlapping ratio of the irradiated pulses. A high overlapping ratio of laser pulses causes continuous melting at the surface, as shown in FIG. 2C. In addition, no melt over-flow between the scanning tracks was observed, which indicates controlled melting of the surface without excessive heating.

The presence of partially embedded hard particles at the surface, as shown in FIGS. 2D and 2E, is attributed to the high melting temperatures of TiC and B₄C particles, which remain mainly in the solid phase during the laser treatment process. It should be noted that the melting temperatures of TiC and B₄C are 3160° C. and 2763° C., respectively, which are higher than the melting temperature of alumina (2040° C.). Although the laser treated surface is subjected to high cooling rates due to the impinging assisting gas and the high rate of melting and solidification, thermally induced cracks are not formed at the surface, except for a few micro-cracks which occur around the hard particles. It should be noted that estimation of the exact cooling rates at the laser irradiated surface is difficult, however numerical simulations carried out previously indicated that the heat transfer coefficient is on the order of 3000 W/m² K at the laser treated surface.

Large crack formation is suppressed by the self-annealing effect of subsequently scanned tracks on the initially treated regions. In this case, heat conduction from the subsequently scanned sites towards the initially formed tracks modifies the cooling rates in the surface region of the workpiece. This suppresses the high cooling rates while lowering the thermally induced strains in the surface region of the laser treated layer. As the chemically distinct hard particles (TiC and B₄C) are located close to each other in the surface region, micro-stress levels increase further. This is because of the contribution of the mismatch of thermal expansion coefficients of TiC and B₄C particles to the stress levels. Therefore, in the neighborhood of these particles, micro-stress levels increase significantly when these particles are located close to each other at the surface during the solidification cycle. Consequently, micro-cracks are formed in the vicinity of these particles. It should be noted that once a crack is formed, the strain energy stored due to the contraction of these particles during solidification is released and crack extension ceases at the surface. Thus, a few scattered micro-cracks are observed rather than large size crack formation at the surface.

The SEM micrographs in FIGS. 3A-3E show a cross-section of the laser treated layer. As best seen in FIG. 3B, the laser treated layer is composed of a dense layer at the surface, with a columnar structure below the dense layer (see FIGS. 3C and 3D). In addition to the hard particles, the dense layer consists of fine grains, which are formed in the surface region of the laser treated layer due to the high cooling rates at the surface. The fine grains cause volume shrinkage while resulting in high stress levels in the dense layer. However, the self-annealing effect of the subsequently formed laser scanning tracks suppresses the stress levels in the dense layer. Thus, conduction heat transfer from the subsequently formed tracks towards the initially treated regions lowers the strain energy in the dense layer. Hard particles remain in the solid phase in the surface region because of their high melting temperatures.

Due to the differences between the thermal expansion coefficients of alumina and the hard particles, the volume shrinkage causes some micro-scale void formation in the vicinity of the hard particles. However, as the micro-size voids are few and localized, they are scattered on the surface. In addition, the wetting state of the particles, due to the particle surface texture, contributes to the void formation around the hard particles. Micro-size cracks are also visible in the vicinity of the hard particles, which are attributed to the differences between the thermal expansion coefficients of the chemically distinct hard particles. Therefore, the TiC and B₄C hard particles located close to each other cause micro-crack formation in their neighborhood. In the region below the dense layer, the cooling rates are low, which cause the formation of fine columnar structures (FIG. 3C). As the cooling rates are non-uniform, the size of the columnar structures varies. As the depth below the dense layer increases and approaches the solid bulk, small columnar structures are replaced with large columnar structures, as shown in FIG. 3D. This is attributed to the decrease of the cooling rates in this region.

As shown in FIG. 3E, large grains are observed in the region adjacent to the large columnar structures, which is associated with a further decrease of the cooling rates in this region. As the thermal conductivity of alumina is low, a clear demarcation line is observed between the laser treated layer and the untreated substrate material. However, the heat affected zone is not visible in the region adjacent to the demarcation line because of the low thermal conductivity.

FIG. 4 shows X-ray diffractograms for laser treated and “as received” (i.e., untreated) surfaces. The high pressure nitrogen assisting gas used forms nitride species at the surface, which is indicated by the diffractogram peaks; i.e., the formation of AlN compounds during the laser processing of the surface is evident. The formation of aluminum nitride may take place in two steps. In the first step, Al₂O₃+C→Al₂O+CO₂, and in the second step, the presence of high pressure nitrogen assisting gas causes Al₂O+CO+N₂→2AlN+CO₂. Carbonic gas (carbon dioxide) formed during the process escapes from the irradiated surface. An AlON compound peak also appears in the XRD diffractogram, the formation of which is associated with high temperature processing. When the temperature reaches the melting temperature, Al₂O₃→2AlO+½O₂ is formed. The oxygen atom remains in the alumina structure and high pressure N₂ causes the formation of AlON compounds through the reaction 2AlO+N₂→2AlON. The transformation of γ-Al₂O₃ phase into thermodynamically stable α-Al₂O₃ is also evident when diffractogram peaks corresponding to the laser treated and “as received” workpiece surfaces are compared. Moreover, TiC and B₄C peaks in the diffractogram indicate the presence of hard particles in the solid phase, which are located in the vicinity of the surface of the laser treated layer.

Table 2 below presents the energy dispersive spectroscopic data at three different locations on the laser treated surface. Close examination of the energy dispersive spectroscopic data reveals that elemental composition does not alter significantly at the treated surface. Although a large error is associated with the quantification of the light elements, such as nitrogen, with energy dispersive spectroscopy data it can be concluded that nitrogen is present at the treated surface which is attributed to the formation of nitride species at the surface during the laser treatment process. The elemental composition of Table 2 comes from the EDS data (wt %). Spectrum corresponds to a location at the surface and the EDS data represent weight percentile of constituting elements.

TABLE 2
Elemental Composition of Laser Treated Surface
	Spectrum	N	O	B	Ti	Al
	Spectrum 1	7.3	37.2	2.9	3.1	Balance
	Spectrum 2	6.2	39.4	3.1	2.9	Balance
	Spectrum 3	7.5	38.2	3.1	2.8	Balance

Table 3 below presents microhardness and residual stress determined by the X-ray diffraction technique for the laser treated surface, as well as “as received” (i.e., untreated) alumina surface, and samples which were laser nitrided at the surface, laser treated with solely TiC particles at the surface, laser treated with solely B₄C particles at the surface, and laser treated with B₄C and TiC at the surface. The use of the two chemically distinct hard particles increases microhardness when compared to that corresponding to laser nitrided and laser treated surfaces with hard particles of only a given carbide. This can be attributed to the formation of micro-stress at the surface because of the differences between the thermal expansion coefficients of the constituting hard particles. The residual stress developed is on the order of −3.1±0.09 GPa, which is compressive, for the laser treated surface with two types of hard particles. The formation of a dense layer consisting of fine grains due to high cooling rates, nitride phases, and differences between the thermal expansion coefficients of the constituting hard particles is responsible for the high residual stress at the surface of the treated workpiece.

TABLE 3
Microhardness and Residual Stress Data
	Hardness	Residual
	(HV)	Stress (GPa)
Untreated surface	1150 ± 50	—
Laser nitrided surface	1650 ± 50	2.1 ± 0.06
Laser treated with TiC at surface	1950 ± 50	2.4 ± 0.06
Laser treated with B4C at surface	1800 ± 50	2.9 ± 0.06
Laser treated with TiC and B4C at surface	2050 ± 50	3.1 ± 0.06

It is to be understood that the present invention is not limited to the embodiments described above, but encompasses any and all embodiments within the scope of the following claims.

Claims

1. A method of laser treating an alumina surface, comprising the steps of:

providing a phenolic resin and hard particle mixture, the phenolic resin and hard particle mixture including a phenolic resin and a mixture of at least two chemically different hard particles;

applying the phenolic resin and hard particle mixture to the alumina (Al2O3) surface to form a resin-coated alumina surface;

heating the resin-coated alumina surface to form a carbon-coated alumina surface, the carbon-coated alumina surface including a carbon film; and

scanning the carbon-coated alumina surface with a nitrogen gas-assisted CO2 laser beam to provide a laser-treated surface, the laser-treated surface including AlN and/or AlON.

2. The method of laser treating an alumina surface as recited in claim 1, wherein the at least two chemically different hard particles include titanium carbide (TiC) and boron carbide (B4C).

3. The method of laser treating an alumina surface as recited in claim 1, wherein the hard particle mixture includes titanium carbide (TiC) and boron carbide (B4C) in a ratio of about 3 wt % of TiC and 3 wt % of B4C.

4. The method of laser treating an alumina surface as recited in claim 1, wherein the carbon film has a thickness of about 40 μm.

5. The method of laser treating an alumina surface as recited in claim 4, wherein the step of heating the resin-coated alumina surface comprises heating the resin-coated alumina surface at a temperature of about 175° C. and a pressure of about 8 bar.

6. The method of laser treating an alumina surface as recited in claim 5, wherein the step of heating the resin-coated alumina surface further comprises heating the resin-coated alumina surface at a temperature of about 400° C. in an inert gas atmosphere.

7. A method of laser treating an alumina surface, comprising the steps of:

applying a coating of a phenolic resin and hard particle mixture to an alumina (Al2O3) surface to provide a resin-coated alumina surface, the phenolic resin and hard particle mixture including a phenolic resin, titanium carbide (TiC), and boron carbide (B4C);

heating the resin-coated alumina surface to form a carbon-coated alumina surface, the carbon-coated alumina surface including a carbon film having a thickness of about 40 μm; and

scanning the carbon-coated alumina surface with an inert gas-assisted CO2 laser beam to provide a laser-treated surface.

8. The method of laser treating an alumina surface as recited in claim 7, wherein particles of titanium carbide (TiC) and boron carbide (B4C) in the phenolic resin and hard particle mixture each have a diameter of about 350 nm.

9. The method of laser treating an alumina surface as recited in claim 7, wherein the titanium carbide (TiC) and the boron carbide (B4C) are present in a ratio of about 3 wt % of TiC and 3 wt % of B4C.

10. The method of laser treating an alumina surface as recited in claim 7, wherein the step of heating the resin-coated alumina surface comprises heating the coated alumina surface at a temperature of about 175° C. and a pressure of about 8 bar.

11. The method of laser treating an alumina surface as recited in claim 10, wherein the step of heating the resin-coated alumina surface further comprises heating the resin-coated alumina surface at a temperature of about 400° C. in an inert gas atmosphere.

12. A method of laser treating an alumina surface, comprising the steps of:

applying a phenolic resin and hard particle mixture to an alumina (Al2O3) surface to provide a resin-coated alumina surface, the phenolic resin and hard particle mixture including a phenolic resin, titanium carbide (TiC), and boron carbide (B4C), the titanium carbide (TiC) and the boron carbide (B4C) being present in a ratio of about 3 wt % of TiC and 3 wt % of B4C;

heating the resin-coated alumina surface to form a carbon-coated alumina surface; and

scanning the carbon-coated alumina surface with an inert gas-assisted CO2 laser beam to provide a laser-treated surface.

13. The method of laser treating an alumina surface as recited in claim 12, wherein particles of titanium carbide (TiC) and boron carbide (B4C) in the phenolic resin and hard particle mixture each have a diameter of about 350 nm.

14. The method of laser treating an alumina surface as recited in claim 12, wherein the step of heating the resin-coated alumina surface comprises heating the resin coated alumina surface at a temperature of about 175° C. and a pressure of about 8 bar.

15. The method of laser treating an alumina surface as recited in claim 14, wherein the step of heating the resin-coated alumina surface further comprises heating the resin coated alumina surface at a temperature of about 400° C. in an inert gas atmosphere.