Patent Application
20080093779
High toughness SiC may be obtained using applied pressure, such as by hot pressing. For recent publications see D. A. Ray, R. M. Flinders, A. Anderson, and R. A. Cutler, “Hardness/Toughness Relationship for SiC Armor,” Ceram. Eng. Sci. Proc., 24[3], 401-10 (2005); M. Flinders, D. Ray and R. A. Cutler, “Toughness-Hardness Trade-Off in Advanced SiC Armor,” Ceram. Trans., 151, 37-48 (2003); M. Flinders, D. Ray, A. Anderson, and R. A. Cutler, “High-Toughness Silicon Carbide as Armor,” J. Am. Ceram. Soc., 88[8], 2217-26 (2005), and R. Marc Flinders, D. Ray, A. Anderson and R. A. Cutler, “Microstructural Engineering of the Si—C—Al—O—N System,” Ceram. Trans., 178, 63-78 (2005), all of which are hereby incorporated by reference. However, it is challenging to pressureless sinter these same systems to high density due to loss of volatile species, changes in fracture modes, and sluggish sintering kinetics, among other problems. Using applied pressure to obtain high density with oxygen at grain boundaries and triple junctions may result in intergranular fracture and enhanced fracture toughness.
Pressureless sintering with AlN as the only additive is also difficult. Hot pressing materials using an active powder with a surface area of about 15 m2/g (e.g. Starck UF-15) and the addition of 2.5 wt. % AlN resulted in with densities of 3.21 g/cc and fracture toughness of 3.5 MPa·m1/2 as measured by the SEBP technique. When the same material was pressureless sintered, it had a density of 1.84 g/cc after sintering in Ar at 2000° C. for one hour. Since the green density was 1.75 g/cc, little densification occurred. When 0.5 wt. % B4C was added in addition to the 2.5 wt. % AlN, the density increased to 2.82 g/cc. As shown in
The addition of Al together with the AlN, C, and B4C, as shown by the SiC-1.5 wt. % AlN-1 wt. % Al-1 wt. % C-0.25 wt. % B4C composition, also allowed for high density to be achieved over an even wider range of temperatures. The amount of residual O was 0.11 wt. % and N was 0.55 wt. % after sintering at 2000° C., but the toughness was still below about 3.5 MPa·m1/2 as measured by the SEBP technique. This was consistent with the mostly transgranular fracture surface for this material.
When Al contents increased above about 2 wt. %, it was necessary to heat rapidly (33° C./min above 1500° C.) in order to achieve high density. This permitted compositions to be processed without carbon additions, as shown by the density data in
In one embodiment of the present invention, higher fracture toughness is obtained by increasing the amount of intergranular fracture without decreasing the ability to get high density by pressureless sintering. Traditional thinking precludes adding yttria or lanthanide oxides together with boron, since yttrium borides form at high temperatures along with boron oxides (see T. Mori, F. Zhang, and T. Tanaka, “Synthesis and Magnetic Properties of Binary Boride REB25 Compounds,” J. Phys: Condens Matter: 13[20], L423-30 (2001) incorporated herein by reference). However, Pujari, et al. (U.S. Pat. No. 6,762,140 and U.S. Pat. No. 6,680,267), as mentioned above, used boron and lanthanide oxides in liquid phase sintering SiC with C, AlN, and/or Al2O3, by adding at least 2 wt. % Y2O3. Their objective was to induce liquid phase sintering, and this resulted in large amounts of additives, such that the SiC content was below about 94 wt. %. The resulting microstructures showed liquid phase at grain boundaries and required powder beds or control of temperature to limit volatility.
One surprising result of the present invention is the ability to create pressureless sintered SiC with additives in addition to B and C, in minimal quantities to achieve high fracture toughness. In one embodiment, using Al in place of Al2O3 or Al(OH)3, in SiC compositions containing a source of nitrogen (AlN, Si3N4, or BN) and boron (B, B4C, or BN), very little yttria is required to change the fracture mode and increase the fracture toughness. Due to the low amount of additives, weight loss is not an issue, and packing powders or gaseous overpressures are not required even at high processing temperatures. It is expected that metallic aluminum reacts with the surface silica on the silicon carbide to form alumina, such that the Al is present only as a transient liquid. Regardless of the mechanisms involved, substituting Al for Al2O3 or Al(OH)3 changes the sintering behaviour and allows densification at lower additive levels than would have been expected.
Thus, in the present invention, pressureless sintered ceramic bodies containing more than about 94% by weight of SiC can be achieved which appears to be an improvement over the prior art. According to the prior art, it does not appear to have been possible to obtain a pressureless sintered silicon carbide body with a density above 3.1 g/cc, a SEPB fracture toughness above about 4 MPa·m1/2 as measured by the single-edged precracked beam technique, when the SiC content was above 94 wt. % without means for controlling volatile species.
In the present invention, density is measured by the Archimedes' method of water displacement, helium pyenometry, or with any accurate assessment of mass and volume. Fracture toughness is measured on 3 mm×4 mm×45 mm SEPB bars as described in ASTM C 1421-99 (Standard Test Methods for Determination of Fracture Toughness of Advanced Ceramics at Ambient Temperature, pp. 641-672 in 1999 Annual Book of Standards. ASTM, Philadelphia, Pa., 1999). A dye, such as any black ink used for ink jet printers, can be vacuum infiltrated into the precrack for 10 minutes and then dried overnight at 60° C. in order to help locate the precrack. The bars are cooled to room temperature and fractured as described in 1421-99. This marking of the precrack allows an easy measurement of the crack length and negates the need for the tedious strain gage measurement described in 1421-99. The fracture toughness values measured this way are conservative since they do not take into account any stable crack growth prior to fracture. Testing hundreds of samples has shown that this method is a good assessment of long-crack fracture toughness. In one embodiment, the pressureless sintered ceramic body may include a density greater than about 3.1 g/cc. In another embodiment, the pressureless sintered ceramic body includes a density of above about 3.15 g/cc. The pressureless sintered ceramic body may also have a fracture toughness of at least 4 MPa·m1/2.
Oxygen and nitrogen contents in sintered samples are measured by inert gas fusion/thermal conductivity detection techniques with commercially available analyzers, such as model TC500 from Leco Corporation (St. Joseph, Mich.). A weighed sample, placed in a high-purity graphite crucible, is heated in He at high temperatures where the oxygen, released as CO, and N2 are released. The CO is converted to CO2 and measured with an IR detector, while the N2 is measured due to the thermal conductivity decrease for the gas.
An intergranular fracture mode or mechanism means that at least one-third (33%) of the fracture surface fractures in between grains. The percent of intergranular fracture can be assessed by polishing a SEPB bar, precrackling the bar, etching the precracked bar, and viewing the crack mode over a distance of 25-250 μm, depending on the grain size, it is important to measure a length equivalent to at least 100 grains.
X-ray fluorescence (XRF) or inductivity coupled plasma (ICP) chemical analysis, routine methods employed in many laboratories, can be used to assess the Al, Y (or lanthanide elements), and B contents. SiC content can be assessed with x-ray diffraction using Rietveld fitting to quantify secondary phases present. Alternatively, Si content can be measured by XRF and combined C by combustion with an IR detector in order to quantify SiC content. Lanthanide elements are defined as La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu.
The SiC starting powder can be primarily alpha, beta, or mixed polytypes. In some applications, alpha powder is less expensive. The starting SiC powder can have impurities, such as oxygen, carbon, and iron, as is typical of material currently used in the industry. An inexpensive SiC grade can be milled to increase its surface area or a premilled powder can be used. The milled SiC powder should have a surface area above 10 m2/g, and in one embodiment above 15 m2/g. Powders between about 10 m2/g and about 20 m2/g have higher sintering activity. In some embodiments, surface areas above 20 m2/g can be used for this invention.
Boron can be added as elemental boron, as boron carbide, boron nitride, aluminum boride (e.g. AlB2 or AlB12), boron oxide, boric acid, or any other form. In one embodiment, boron carbide (B4C) is used for adding boron since fine-grained powders are readily available. In one embodiment boron carbide powder having a mean particle size between 0.5 and 8 μm is used. Since boron is relatively mobile at high temperatures, some boron may be provided within the sintering environment. The amount of boron used in one embodiment is less than 1% by weight of the pressureless sintered ceramic body. In other embodiments, the amount of boron used is between about 0.1% and about 0.5% by weight of the pressureless sintered ceramic body.
Carbon, while note necessarily present in the sintered sample, aids in densification and may be added. Since carbon has a slow diffusion coefficient, it is distributed prior to green densification. Phenolic resins may be used for distributing the carbon. The carbon may be utilized to remove oxygen from the SiC and it is therefore distributed uniformly on the SiC particles by dissolving the carbon and leaving it on the surfaces through evaporation or sublimation. In other embodiments, oxygen in removed by other methods.
Yttrium may be added in the form of yttrium oxide (Y2O3) since this is readily available and inexpensive due to its use in other industrial processes. In one embodiment, In one embodiment, added yttria is less than less than 20 μm in mean diameter. In other embodiments, the added yttria is less than 5 μm in size. Due to the low volume fraction added, it can be less than 1 μm in size without affecting the green density substantially. Yttrium can also be added as a hydroxide Y(OH)3, a nitrate (Y(NO3)3,xH2O), acetate, carbide, nitride, or as elemental Y. The yttrium, in one embodiment is distributed well in the green state and it is generally milled with the starting SiC powder to aid in this process. In one embodiment, the yttrium accounts for less than about 1.5% by weight of the pressureless sintered ceramic body. In other embodiments, the pressureless sintered ceramic body includes between about 0.01% and about 1.0% yttrium by weight. In other embodiments, other lanthanide elements can be used alone or in combination to replace some or all of the yttrium. One surprising result of the present invention is that only minor additions of yttrium and/or lanthanide elements are required to achieve comparatively high levels of fracture toughness and density, which allows sintering without the use of powder beds or overpressures to control volatilization.
Nitrogen may be added as Si3N4, AlN, or BN. In one embodiment, silicon nitride is used when milling in water since it does not have the same tendency for hydroxide formation as does aluminum nitride. The nitrogen-containing powders may be added as fine powders. In other embodiments a nitrogen-containing powder is added having a particle size of less than about 20 μm in mean diameter. In other embodiments, the particle size is equal to or less than about 5 μm in size. The amount of nitrogen in the pressureless sintered ceramic body may be less than about 2% by weight. In other embodiments, the amount of nitrogen in the pressureless sintered ceramic body is between about 0.15% and 1% by weight. The nitrogen is well distributed in the green state. These powders may also contain oxygen due to their fine particle size. Additional carbon is added to compensate for the free silicon formed as the silicon nitride dissociates at elevated temperatures.
Aluminum may be added as a pure metal. In one embodiment, aluminum is added as a spherical powder. In another embodiment, the aluminum is added in flake form. It may be coated or uncoated, as a multitude of fine aluminum powders are available. If processing occurs in water, then a coarser powder may be used to limit the amount of aluminum hydroxide formation. The amount of aluminum hydroxide formation can also be controlled by mixing or low-energy milling the Al into the SiC slip rather than high-energy milling. The aluminum powder may contain a surface layer of alumina and some additional oxygen may be introduced during processing. In one embodiment, the aluminum particle size is less than 100 μm. In other embodiments, the particle size is less than 50 μm. Powders less than 10 μm in size can also be used. Distribution of the Al may occur in form of liquid upon heating. In one embodiment, the amount of aluminum in the pressureless sintered ceramic body is less than about 5% by weight. In another embodiment, between about 0.25% and about 3% by weight of the pressureless sintered ceramic body is aluminum.
It will be appreciated by those of skill in the art that various amounts of the above mentioned additives in various particle sizes may be used in a variety of combinations to achieve a pressureless sintered ceramic body having a fracture toughness of at least about 4 MPa·m1/2 and a density of at least about 3.1 g/cc while maintaining a high percentage of silicon carbide.
The choice of solvent or carrier is dependent on the processing route chosen. This invention can be practiced using traditional ceramic mixing processes including ball milling, vibratory milling, attrition milling, jet milling, and high shear mixing. In one embodiment, the powder is mixed dry. In another embodiment, it is preferable to mix wet in order to distribute the carbon and other additives more uniformly. Non-aqueous solvents limit oxidation of the Al, while water is preferred for production batches for environmental reasons. Milling can occur in liquid nitrogen, although this is a relatively expensive milling approach. Forming methods include tape casting, tape calendaring, slip casing, extrusion, injection molding, pressing, or other means for achieving high green density. Dispersants, binders, and plasticizers are specific to the milling and forming methods, but are similar to those used by those skilled in the art of SiC component manufacturing. The organics can be removed in air, inert gas, or vacuum. If removed in air, care should be taken to limit oxidation of the aluminum and removal of all of the carbon. The organics should be removed in a manner to prevent pore formation, cracks, or other defects from being introduced during this step. The green density after binder removal in one embodiment is at least 1.5 g/cc. In another embodiment, the green density is about 1.7 g/cc. The parts may be bisqued or green machined prior to sintering. Bisquing may be performed in nitrogen, argon, helium, or vacuum at temperatures up to 1600° C.
Sintering can occur in graphite, tungsten, or any furnace capable of reaching temperatures above 1800° C. in reducing environments. In one embodiment, graphite resistance-heated furnaces are used in such applications. The atmosphere may be vacuum below 1500° C. and helium or argon above 1500° C. A stagnant Ar may be used with a slight positive pressure to avoid introduction of air into the furnace chamber. The parts may be placed on graphite, graphite foil, BN, coarse SiC powder, or any non-reactive surface during sintering. It is convenient to put the parts inside a graphite box or enclosure to enable loading of the furnace.
The sintering cycle is best designed around the parts being manufactured, but is typical of those already used for making SiC sintered with boron and carbon, with the exception that the isothermal hold temperature is 50-150° C. lower. If larger grains are desired either temperature or time can be increased. Sintering need not be performed in a batch kiln and the use of a continuous kiln is acceptable. A pressureless sintered density above 3.1 g/cc ensures closed pores, which can be removed with an optional hot isostatic pressing step without cladding the part. Hot isostatic pressing, if used, will typically occur at temperatures above 1800° C. in Ar pressures above 100 MPa.
Finishing operations using grit blasting or diamond tools is optional, since the surface finish of the parts are acceptable for many applications and volatility, which is less than 3 wt. %, and preferably less than 1 wt. %, is low.
The following examples further illustrate this invention.
Eight compositions were prepared by using alpha silicon carbide (H. C. Starck grade UF-15), boron carbide (H. C. Starck grade HS), phenolic resin (Durez Corporation grade 7716), aluminum (Valimet grade 1413), aluminum nitride (Tokuyama Soda grade F), yttrium oxide (Molycorp grade 5600), and reagent grade acetone. The compositions shown in Table 1 were prepared by making slurries of 350 grams of acetone and appropriate amounts of powdered phenolic resin (with a C content of 50% after pyrolysis), and powders so that each composition contained 300 g of powder based on a post-pyrolysis weight. The slurries were ball milled in one-liter high density polyethylene (HDPE) bottles filled with 1 kg of SiC (solid state sintered with boron and carbon additives) media for 62 hours. The slurry was then stir-dried and screened through an 80-mesh sieve. Following stir-drying, the powders were prepared for pressing by lubing with 2 wt. % polyethylene glycol (Union Carbide grade 8000) and 0.25 wt. % polyvinyl butyral (Solutia grade B-79). Plates (approximately 6 mm×60 mm×60 mm) were uniaxially pressed at about 70 MPa and then isostatically pressed to about 200 MPa. The binders were removed from the plates and the carbon pyrolyzed in an atmosphere of flowing N2 at 600° C. for one hour. Green densities were about 1.75 g/cc for all compositions after this process. The plates were placed on BN-coated graphite foil inside of a graphite crucible without packing powder. Pressureless sintering was performed by evacuating and backfilling the chamber with argon and heating to 1500° C. in 90 minutes. The parts were then heated to the final temperature (2000° C. or 2100° C.) in 30 minutes. Weight loss was less than 2% for all compositions, with linear shrinkages ranging between 17 and 18%. The plates were machined to make bars (3 mm×4 mm×45 mm) used for measurements of SEBP fracture toughness. Knoop hardness (1 kg load), grain size, and aspect ratio as listed in Table 2.
Plasma etched microstructures of the materials sintered at 2000° C./1 hr are shown in
Example 1-8 show high toughness and density, as well as the specified chemistry at both sintering temperatures.
It is clearly demonstrated that small yttria additions permit increased intergranular fracture to result in improvement of fracture toughness by a factor of two. More importantly, by limiting the amount of yttria added, volatility is suppressed and there is no need to use packing powders or gas overpressures.
A composition similar to Example 8 but with 0.5 wt. $ Y2O3 was prepared in larger quantities for pressing into large tiles. The powder (95.25wt. % SiC, 2.0 wt. % Al, 1.0 wt. % AlN, 1.0 wt. % C, 0.25 wt. % B4C, and 0.5 wt. % Y2O3) was made by batching slurries of 1500 grams using the same raw materials described in Examples 1-8. The slurries were ball milled in 4 liter HDPE bottle with 4 kg of SiC media for 16-20 hours. The slurries were then stir-dried and screened through a 20-mesh screen. Large plates (˜20 cmט30 cm) were pressed uniaxially at about 35 MPa and then isostatically pressed to 207 MPa. Pressed parts were delubed in an atmosphere of flowing N2 to 600° C. The green density of the plates (˜1.73 g/cc) was nearly identical that of the smaller tiles in Examples 1-8.
Large tiles were sintered between BN coated graphite foil, supported between graphite plates and placed nearly upright in a large graphite crucible with a lid. Due to size constraints, the tiles had to be tilted slightly in order to fit into the crucible. Sintering was done by evacuating and backfilling the chamber with argon and heating to 1500° C. in 90 minutes. The parts were then heated to 2000° C. in 30 minutes and held for one hour. Densities of the eight tiles sintered were 3.13±0.01 g/cc. Weight change was 0.5-0.6% for the plates which were approximately 3 cm×15 cm×32 cm after pressureless sintering. Creep flattening of the plates occurred in stagnant Ar by heating to 2000° C. for 0.5 hour under approximately 5 kPa of pressure and was performed in a separate operation without change in density and a 0.2% loss in mass. Bars were ground from one tile and compared with a commercially available pressureless sintered material (Morgan Advanced Ceramic's PS 5000) prepared in an identical manner. Table 3 compares some properties of the two materials. It is very evident that there is a trade-off between hardness and toughness between the two materials. Example 9 is more than twice the SEPB fracture toughness but is not nearly as hard. This trade-off is inherent in the change in fracture mode which occurs between the two materials. The significance of Example 9 is that is demonstrates that large components can be fabricated by pressureless sintering while maintaining high SEPB fracture toughness without the use of powder beds
The compositions shown in Table 4 were prepared in the same manner as Examples 1-8. For Example 16, Si3N4 (Ube, grade E-10) was substituted for the AlN at the same level of nitrogen as in Example 7, with additional Al added so that the Al contents were fixed for the two compositions. Tiles were pressed, debinderized, and sintered at 1925° C. for 1 hour in an identical manner as described for Examples 1-8. Weight loss was less than 2 wt. % for all samples. The samples were subsequently hot isostatically pressed (HIPed) by heating to 1950° C. and holding for one hour under 200
MPa Ar. Table 5 displays the sintered density, as well as the increased density due to HIPing, with Example 9 included for comparison. These examples demonstrate that post-HIPing can be used to increased density, as well as demonstrate that the nitrogen source can be supplied by silicon nitride instead of aluminum nitride. HIPing coarsened the microstructures and increased the toughness. This is not a consequence of HIPing but rather due to exposure to elevated temperature. Microstructural control is necessary to control toughness regardless of whether the materials go through a post-sintering HIPing cycle.
Example 17, which is identical in composition to Example 12 was prepared by water-based processing along with Example 18, which was identical in composition to Example 17 except that Al(OH)3 (Alfa Asear grade 12366) was substituted for the Al and Si3N3 (Ube Si3N4) for the nitrogen in the AlN. For both examples, 400 grams of deionized water was added to a one-liter HDPE jar filled with one kilogram of SiC media. The pH was adjusted to about 9.5 with ammonium hydroxide and 15 grams of phenolic resin was added (Capitol Resin Corporation grade 720) and the pH was again adjusted to about 9.5 with ammonium hydroxide. Three grams of a lignosulfunate dispersant (Borregaard Industries grade Vanisperse CB) was added prior to adding 574.5 grams of SiC (Starck grade UF-15), 3.0 grams of B4C (Starck grad HS), and 1.5 grams of Y2O3 (Molycorp grade 5600). Example 17 was milled for 14 hours with 18 grams of PEG (Union Carbide grade 8000) while Example 18 had 26.0 grams of Al(OH)3 and 3.08 grams of Si3N4 added along with the same amount of PEG before milling for 14 hours. AlN (6.0 grams of Tokuyama Soda grade F) and Al (9.0 grams of Valimet grade I13) were added to Example 17 and both mills continued to roll for four additional hours. The apparent viscosity of Example 17 ranged from 2470 mPa-s at 10 rpm to 450 mPa-s at 100 rpm while Example 18 had a lower viscosity (1700 mPa-s at 10 rpm to 300 mPa-s at 100 rpm). These slips, which were about 30 vol. % solids, were dried by spraying the slurry into a bath of liquid nitrogen and then freeze drying the powders. The dried powders were screened −80 mesh and processed similar to Examples 1-8. The green density of Example 17 was 1.78 g/cc, which was slightly higher than Example 18 (1.73 g/cc). Sintering occurred as for Examples 1-8 with densities as shown in Table 6. Example 17 sintered to a density greater than 3.1 g/cc but Example 18 did not sinter well as evidenced by the low density and high open porosity. Al is much more effective than Al(OH)3 in achieving high density at these low levels of sintering additives. The fracture toughness for Example 17 was 4.8±0.4 MPa·m1/2 when sintered at 2000° C. for one hour.
The compositions shown in Table 7 were made using water based processing by adding phenolic resin (CRC 720) to 400 grams deionized water, adjusting the pH by adding 3.0 grams NH4OH, adding B4C (Starck HS), Y2O3 (Molycorp 5600), Si3N4 (Ube E-10), and SiC (Starck UF-15) and milling with 1 kg SiC media inside a one-liter HDPE jar for 4 hours. Al (Valimet H-3), AlN (Tokuyama Soda F), and PEG (Carbowax 8000) were then added and an additional two hours of milling was performed before freeze drying. The materials were pressed to a green density of about 1.75 g/cc and sintered in stagnant Ar for one hour at 2000° C. with results as shown in Table 8. A comparison of Examples 19 and 20 clearly shows that Si3N4 can be substituted for AlN in water based processing. Examples 19 and 21 compare aluminum at the same concentration and clearly show the benefits of metallic aluminum as compared to aluminum oxide. Not only is densification and fracture toughness improved but the surface segregation of yttria that occurs with the oxide is minimal when using the metal. A comparison of Examples 19 and 22 clearly shows that there is no advantage in using excess yttria, which migrates to the surface at higher concentration.
To further illustrate the uniqueness of this invention comparative compositions (see Table 9) were made using examples from three previous patents. Examples 23 and 24 are batches 2 and 7, respectively from U.S. Pat. No. 6,762,140 to Pujari, et al. Examples 25 and 26 are Examples 3 and 25, respectively, from U.S. Pat. No. 5,298,470 to Chia, et al. Example 27 is Example 5 from U.S. Pat. No. 4,829,027 to Cutler, et. al. Example 28 is from the present invention.
All of the samples were processed in an identical manner using the same raw materials whenever possible. The compositions shown in Table 9 were made using water based processing by adding phenolic resin (CRC 720) to 400 grams deionized water, adjusting the pH to 9.5 using NH4OH, adding B4C (Starck HS), Y2O3 (Molycorp 5600), Al2O3 (Ceralox SPA-0.5), Nd2O3 (PIDC grade 1030), Al (Valimet H-3), AlN (Tokuyama Soda grade H), PEG (Carbowax 8000) and SiC (Starck UF-15) before milling with 1 kg SiC media inside a one-liter HDPE jar for 24 hours. The compositions were freeze dried, pressed at 100 MPa uniaxially and 200 MPa isostatically before sintering all examples together at either 1900° C. for three hours, 1950° C. for three hours, or 2000° C. for one hour in stagnant Ar after delubing at 600° C. in N2. Results are shown in Table 10.
Fracture surfaces from samples sintered at 1950° C. for three hours are shown in
The fact that little yttria is necessary to increase fracture toughness is shown by Examples 29-33 made at the same time as Example 28. The compositions are shown in Table 11, with Example 28 added for convenience in seeing its yttria level. Processing of Examples 29-33 was identical to Examples 23-28, with results as shown in Table 12. The level and amount of yttria needed to increase the fracture toughness is likely dependent on the method of distributing the yttrium or lanthanide oxide materials prior to sintering. Note that increased yttria in Examples 29-33 had little effect on density, but the smallest addition of 0.125% was enough to allow increased intergranular fracture, which nearly doubled the fracture toughness.
The importance of intergranular fracture is illustrated further in Examples 34-36. Table 13 lists the compositions, which were processed as 2.4 kg batches in four-liter HDPE jars in a manner similar to Examples 28-33. Examples 34 and 36 are identical in composition in Examples 29 and 28, respectively. Table 14 lists the properties after sintering at 2000 C. for one hour. HK1 and HV1 are the Knoop and Vicker's hardness using one-kilogram loads. It is clear that the ability to fracture intergranularly influences the fracture toughness and is a function of composition.
Selected examples were sent for chemical analysis, with results as shown in Table 15. ICP was used to determine Al, Y, Nd, Si, and B contents. Oxygen and nitrogen were determined using inert gas fusion/thermal conductivity techniques. Example 24, which is not part of the present invention, clearly has less than 94 wt. % SiC, whereas the other examples clearly have greater than 94 wt. % SiC. Substantial oxygen and nitrogen are retained in the sintered structure
The above examples are only illustrative to teach how to pressureless sinter SiC to high density with a SEPB toughness greater than 4 MPa·m1/2 with density greater than 3.1 g/cc and without the need for powder beds or gaseous overpressures for SiC contents greater than 94 wt. %. While the invention has been described in terms of several specific embodiments, it must be appreciated that other embodiments could readily be adapted by one skilled in the art. For example, one could add other secondary phases to increase toughness, change crystallinity by annealing during cooling, or seed the microstructures to increase the tendency to form elongated grains.
The important point is that low levels of additives limit volatility and when added in the proper manner can double the fracture toughness compared to silicon carbide sintered conventionally with boron and carbon. Process steps such a hot isostatic pressing may be used in addition to or in place of pressureless sintering. It should also be noted that the heating step will result in a change of chemical composition from that of the initial starting materials. Accordingly, the scope of the invention is to be limited only by the following claims.
This application claims the benefit of U.S. Provisional Application No. 60/787,717 filed Mar. 29, 2006.
| Number | Date | Country | |
|---|---|---|---|
| 60787717 | Mar 2006 | US |
