US 5,000,986 AGrant
Metallized Coatings on Ceramics for High-Temperature Uses
Issue Date:1991-03-19
•44 Claims
•1 Drawing Sheets
Abstract
A method for forming metallized coatings on ceramics for high-temperature uses above about 630.degree. C. comprising the steps of: preparing a metallizing composition of mixed ingredients of differing sizes, proportioning the differing sizes to have nonsegregating qualities when applied onto the ceramics, coating the metallizing composition on the ceramics; and heating to form the desired metallized layer.
Metadata
Inventor
- Chou H. Li
Application Information
Application Number:US 2776723
Filing Date:1988-12-14
Priority Date:1988-12-14
Art Unit:139
Classifications
IPC:
B05D 302B23K 3100B23K 3522
Field of Search:
42742822862;63;229;383.5;376.7;226;376.6446;432;704;698263.12;122;120;124
Patent Drawings (1 sheets)
Drawing sheets are not available for this patent.
Description
Background
1. Field
This invention relates to ceramic-metal joining, and more particularly relates to ceramic-metal joining with uniform ceramic metallizing compositions and specially graded seals to produce reproducibly strong and thermomechanically shock-resistant and substantially defect-free joints.
By ceramic I mean not only the usual ceramics such as alumina, zirconia, beryllia, mullite; but also quartz, intermetallics, diamond, boron, graphite, carbon, silicon, and various other carbides, nitrides, aluminides, or borides; glasses, machinable glasses, Corning's Vision glass; and the surface of many metals, particularly reactive metals such as aluminum, magnesium, chromium, silicon, titanium, or zirconium which always have oxides or other compounds of reactions of the metal with the environment.
2. Prior Art
Many problems exist with present ceramic metallizing methods. A serious problem is the difficulty of achieving uniform metallized layers formed on the ceramic. Take, for example, the commonly used heavy metal processes, such as W-yttria (W-Y.sub.2 O.sub.3), W-Fe, or Mo-Mn. In these and many similar joining methods, segregation of the mixed metal or other powder takes place due to their differing specific gravities, shapes, sizes, porosities, and surface smoothness. These segregations occur at all times; during the mixing of the powders, storing of the power suspensions, application of the suspensions, settling of the suspended power particles, and drying of the suspension. Further, these segregations occur so fast as to be practically uncontrollable, as will be shown shortly.
In general, spherical, heavy, large, smooth, and dense particles settle first and early in the binder or suspension medium. Upon settling, these particles tend to roll or move sidewise or downward toward the corners or boundaries faster and further than odd-shaped, light, small, rough, and porous particles of otherwise identical characteristics.
Take the W-Y.sub.2 O.sub.3 mixed powders in an organic binder of nitrocellulose in butyl carbitol acetate with specific gravities of 19.3, 4.5, and 0.98, respectively. Such a suspension, even if perfectly mixed up by shaking, stirring, roller-milling, or otherwise, will immediately tend to segregate. More specifically, the initial settling acceleration due to gravitational minus buoyancy forces on W powders is 980.6.times.(19.3-0.98)/19.3=930.8 cmxcm/sec, while that of Y.sub.2 O.sub.3 powders is only 767.0 cmxcm/sec.
In a mixing, storing, or carrying bottle 10 cm high and containing a perfectly mixed suspension of these metallizing powders, the time to completely settle out is only 147 ms (milliseconds) for W powders, if uniform acceleration is assumed. At the tip of a paint brush having a suspension drop 0.3 cm in diameter, the complete settling time of these same W powders is merely 25.4 ms, while on a horizontally painted or sprayed layer 0.1 cm thick, the same settling time is only 14.7 ms. In all these cases, the complete settling time for the Y.sub.2 O.sub.3 powders is always the square root of 930.8/767.0=1.21, or 21% longer.
Note in particular that the powder segregations with uniform accelerations may be completed within 147 to 14.7 ms. Such short times indicate that the W-Y.sub.2 O.sub.3 powder segregations are beyond human controls. Painted or sprayed mixed powder layers are thus always not uniform.
In metallizing onto a horizontal ceramic surface to be metallized, most of the W powders immediately settles out. The first layers are therefore always very rich in W, and correspondingly very poor in Y.sub.2 O.sub.3. These first layers are too refractory for the preset metallizing temperature (up to about 1550 C) so that the ceramic surfaces are not sufficiently metallized, or not at all. The last settling layers, on the other hand, are too rich in the fluxing Y.sub.2 O.sub.3. Again, the ceramic surfaces are improperly metallized, with only a glassy layer being formed which is very weak in strength and thermal or thermal shock resistance.
Thus, common metallizing results on ceramics are often erratic and uncontrollable. The metallized surface may contain loose and unmetallized spots with high refractory metal content, as well as non-wettable spots due to the high flux content. The entire process is critical and involved, and yet nonuniform. The resultant ceramic-metal joints or ceramic coatings on metals are weak, costly, nonreproducible, and usually not vacuum-tight, or temperature-resistant.
Painting or spraying onto vertical or inclined surfaces results in vertical and additional lateral segregations and gradations, and gives added poor uniformity, reproducibility, and bonding strength.
While only the effect of gravitational density segregation has been considered in some detail, the other segregation variables such as powder shape, size, porosity, and surface roughness are also important.
A second important problem with common joining processes is the lack of control, or even understanding, or dynamic mismatches of temperatures, stresses, and strain profiles in the joint region, and their variations with time. Another aspect of this invention is therefore to describe such dynamic mismatch phenomena, and to specially tailor-grade the composition and/or physical property profiles of the joint region so that the maximum or critical transient mismatch stresses never exceed the local material strength at any point inside the joint region, at any time during the heating or cooling of such joints in processing or service.
A third problem results from our poor understanding of the required microstructural, chemical, and physical properties of the interfacial regions in the ceramic-metal joints.
Accordingly, an object of this invention is to provide improved ceramic-metal joints and joining methods;
A further object of this invention is to provide improved ceramic metallizing methods for these joints;
A broad object of this invention is to minimize gravitational segregations of the components in the metallizing methods during or prior to the joining;
Another broad object of the invention is to specially tailor-grade, both in and normal to the joining plane, the composition and/or property profiles in the joint regions to ensure that the maximum dynamic or transient stresses do not exceed the local material strengths at any point and time;
A further object of the invention is to provide a specially microengineered interfacial region of the optimum characteristics to achieve defect-free, tough, and very strong joints;
Another object of the invention is to flawlessly coat metals or ceramics with protective materials;
A yet another object of the invention is to provide substantially flawlessly coated reinforcements for the manufacture of tough, strong, thermochemically stable, and thermomechanically shock-resistant composites;
Further objects and advantages of my invention will appears as the specification proceeds.
Summary
A method for improving the strength of a ceramic-metal bond comprising: providing a uniform metallizing composition; and
forming by liquid diffusion between the ceramic and metal a shock-absorbing, interfacial region whose microstructure is free of voids, inclusions, microcracks, and excessive dynamic mismatch stresses/strains and stress gradients.
Brief Description
The invention and its further objects and features will be more clearly understood from the following detailed description taken in conjunction with the drawings in which:
FIG. 1 shows a system for real-time monitoring of mixed settling powders;
FIG. 2 shows nodular bonding spots on reinforcing carbon fibers in carbon composites;
FIG. 3 shows a multi-purpose bonding method for high temperature ceramic superconductors;
FIG. 4 shows newly microengineered microstructures of the bonding interfacial regions; and
FIG. 5 shows a bonding method for mounting diamond or other gem stones.
Detailed Description
It will be understood that the specific embodiments described herein are merely illustrative of the general principles of the invention and that various modifications are feasible without departing from the spirit and scope of the invention. That is, the invention is of general applicability for improving the quality of the ceramic-metal joints or joining methods, or coatings of ceramics on ceramics, or on metals. It is also evident that materials, structures, and methods other than those especially described can be used to practice the invention.
Stokes in 1851 first considered the resistance R which a fluid medium of density d.sub.m and V viscosity n offers to the movement of a spherical particle of velocity V diameter D and density d.sub.p suspended in it, and arrived at the equation R=3.pi.Dvn.
The small sphere settling in the fluid (i.e., gaseous or liquid) suspension medium is acted on by the force of gravity with gravitational constant g, .pi.D.sup.3 d.sub.p g/6 acting downward; and by the buoyant force of the fluid, .pi.D.sup.3 d.sub.m g/6, given by Archimedes' principle and acting upward. The resultant net gravitational force G is .pi.D.sup.3 (d.sub.p -d.sub.m)g/6 acting downward, producing a downward acceleration, a.
When the resistance R exactly equals this net gravitational force G, the acceleration reduces to zero; the final velocity, v.sub.f, becomes constant. There than results:
Hence, the final velocity is: v.sub.f =(d.sub.p -d.sub.m) g D.sub.2 /18 n, the equation of Stokes' law which has been shown to be widely valid.
For a given fluid density (d.sub.m) at a specific temperature (viscosity n) and a given sphere (of density d.sub.p and mass M), the Stokes' equation gives a velocity constant:
Also, the velocity at any time starting from rest, t, is:
while the settling distance at time t is:
The velocity equation shows that the exact v.sub.f is not reached until after infinitely long time when the exponential term in the equation turns to zero and then the velocity reduces to v=v.sub.f, as it should.
With the Stokes' law, one can calculate the velocity constants, v.sub.c in 1/cm-sec, for the settling in water at 20 C (d.sub.m =1.0 and n=0.010) of various metal or oxide powders, with densities in g/cc in parentheses, as follows: W (19.35) 100,000, Y.sub.2 O.sub.3 (5.01) 21,900, Fe (7.87) 37,400, Mo (10.2) 50,100, Mn (7.2) 33,800, WO.sub.3 (7.16) 33,600, Fe.sub.2 O.sub.3 (5.24) 23,100, MoO.sub.3 (4.692) 20,100, and MnO.sub.2 (5.026) 21,900.
Thus, in the W-Y.sub.2 O.sub.3 metallizing process, because the W powders are 3.9 (19.35/5.01) times heavier than Y.sub.2 O.sub.3, the velocity constants v.sub.c 's of the two components differ by a factor of 100,000/21,900=4.6 times. That is, for a given powder size D, the final constant settling velocity v.sub.f of W spheres is 4.6 times greater than that of Y.sub.2 O.sub.3 spheres. As discussed above, this wide difference in velocities results in severe gravitational segregation and early depletion of W particles in the settling mixtures and, therefore, poor metallizing results.
It can also be seen that the powders in the mixed oxide processes, e.g., WO.sub.3 -Fe.sub.2 O.sub.3, are much more uniform, or less varying, in densities, d.sub.p, than mixed particles of the same metals, e.g., W-Fe. This, the WO.sub.3 -Fe.sub.2 O.sub.3 process shows density and velocity constant ratios of 1.366 and 1.455, vs 2.459 and 2.674, respectively, for the W-Fe process.
Similarly, in the Mo-Mn process, replacing the metal powders by their respective oxides reduces the differences in the ratios of velocity constants, v.sub.c, and final velocities, v.sub.f, from 48.2% to only 9.0% and 19.2% to 4.2%, respectively. In addition, the metal particles, i.e., W, Fe, Mo, and Mn when reduced during metallizing from their respective oxides are smaller than the initial oxide powders. These smaller sizes further promote homogenizations and metallizing results.
Hence, if we select and mix the WO.sub.3 and Fe.sub.2 O.sub.3 spherical powders in the size (diameter D) ratio of the square root of (33,600/23,100=1.455) i.e., 1.206, the final settling velocities of both these size-ratioed powders will be exactly the same. That is, by simply making the Fe.sub.2 O.sub.3 powders 20.6% larger than the WO.sub.3 powders, the mixed particles will finally settle in water at 20 C at exactly the same velocity. This condition leads to metallizing uniformity due to the uniform composition of the finally deposited layers.
The final settling velocities of the two mixed powders, v.sub.s 's, however, come only after some settling time, t.sub.s, when a specific amount, Q, of the mixed powders has already settled out at differing velocities. From this settling time, t.sub.s, for the specific combination of component powders, the settled amount Q and material use efficiency can be computed from the materials remaining after the settling time, t.sub.s. The materials already settled before t.sub.s is the presettled distances, s.sub.t, multiplied by the initial material densities. But the already settled materials are not lost, since they can be recirculated and reused in subsequent metallizing runs.
In this way, gravitational segregations between, for example, cosettling W and Fe, Mo and Mn, WO.sub.3 and Fe.sub.2 O.sub.3, or MoO.sub.3 and MnO.sub.2 powders, are minimized. Naturally, the smaller the percentage of velocity or useful powder size differences, .DELTA.v and .DELTA.D, respectively, the lower the material use efficiency on a particular mixed-powder combination. An engineering compromise must, therefore, be struck.
It can be seen that the fluid suspension medium may be either a gaseous or liquid medium. The liquid may be water, alcohol, other inorganic or organic liquids of fairly constant viscosity at room temperature. A varying viscosity liquid may also be used, for example, a polymerizing organic substance containing a polymer and a hardener, a nitrocellulose in an evaporating solvent such as butyl carbitol acetate, or Duco cement diluted with rapidly evaporating acetone, to achieve rapidly increasing viscosity, n. The velocity constant of the settling powders is, as shown above, inversely proportional to this viscosity. In all cases, the starting time for achieving nearly equal settling velocities is shortened by the increasing viscosity due to polymerization or solvent evaporation. With increasing viscosities, the absolute difference is centimeters per second between the settling velocities of the two mixed powders of differing densities then become less, and nearly equal settling conditions are more easily achieved. The real-time monitoring system to be described in FIG. 1 is also useful, but the nearly equally settling mixed powders must be quickly used before much further polymerization or evaporation takes place.
Apparently, the above technique for minimizing gravitational segregation through minimized settling differences can be used to handle more than two types of powders of differing densities.
In practice, we specify that the two settling velocities of the mixed particles are within a certain prespecified percentage, e.g., 20 or 10%, of each other. Still, gravitational segregations are minimized.
By repeated iteration or computer simulation, the best mixed-powder metallizing process for optimal combined metallizing uniformity and material use efficiency can be systematically determined. Based on these principles, method and equipment can be developed for controlling the turn-on time for starting to deposit the mixed powder at nearly equal final settling velocity, v.sub.f, into metallizing layers with the size-ratioed powders.
FIG. 1 shows that a system for real-time monitoring of the settling particles is employed to determine the starting time for collecting the residual or still unsettled mixed particles to be used for metallization. This system has a vertical settling cylinder 10. Near the bottom of the cylinder 10, two pairs of light emitters 11 and detectors 12 are located at two different heights with emitters on one side and detectors on the opposite side of a vertical cylinder 10, to sense the settling particles. The times for the particles to pass the top or bottom emitter/detector pair determine the particle size or type being monitored, while the times for the particles to transverse through the vertical distance d between the heights give their velocities. When the settling velocities of the two types (and sizes) of the powders are within a specified percentages, a slide shuttle 14 is moved to catch on the shuttle the residual or unsettled mixed powder of nearly equal settling velocities. These equal-settling mixed powders in suspension are separated for immediate metallizing use while the already settled powders are drained through the valve 15 for subsequent reuse.
Useful metallizing compositions include the commonly used W:Fe or Mo:Mn system containing 10 to 30 weight percent of Fe or Mn, or their derivatives WO.sub.3 :Fe.sub.2 O.sub.3, MoO.sub.3 :MnO.sub.2, or other non-oxide systems. From the atomic or molecular weights of the elements W, Mo, Fe, Mn, O, Cl, F, I, Br, . . . or radicals NO.sub.3, SO.sub.4, . . . , the weight percentage of the heavy metal W or Mo and the other braze and melting temperature-lowering metals such as Cu, Zn, Pb, Sn, Bi, Fe, Mn, Ag, Au, In, . . . used for the past, suspension, or solution metallizing compositions can be readily determined. Generally, I maintain the same ratio of 10 to 40 weight percent of braze metal to the 90 to 60 percent of heavy metal in these compositions.
There are other ways to insure a substantially constant chemical composition consisting of at least two types of metallizing materials having different densities and carried in a fluid suspension medium. One way is to cause the two types of materials to come out of the suspension medium in a substantially constant chemical composition thereby ensuring uniformity and reproducibility of the metallizing results. For example, the two types of materials may be integrated into physically integral and inseparable forms, such as by alloying the materials into integrated alloy form, or coating the internal and/or exterior surface of one type of material particles with the other material to form integrated coated powders.
Thus, tungsten particles may be alloyed or coated with iron to form integral or inseparable W-Fe powders. Similarly molybdenum powders may be alloyed or coated with manganese to form integral Mo-Mn powders that will not segregate.
Another method to minimize segregation of a single fluxing (e.g., MgO, Y.sub.2 O.sub.3) or brazing (e.g., Cu, CuO, Zn, ZnO), cometallizing (e.g., Mn or MnO2 with Mo or Fe or Fe.sub.2 O.sub.3 with W) material is the use of an aqueous or other solution of W and/or Mo compounds such as sodium molybdate or tungstate which is soluble in water, or MoO.sub.3 or WO.sub.3 which is soluble is hot water particularly in the presence of NH.sub.4 OH. Here, the solution is the settling medium itself and suspended powders being of a single type cannot segregate. Solutions of compounds of Cu, Zn, Fe, Mo, . . . used with powders of W, Mo, Wo.sub.3, or MoO.sub.3 achieve the same results.
to completely eliminate gravitational segregations, solution metallizing is the ideal process. Many molybdenum and tungsten compounds are soluble in water, alcohol, acid, or bases. MoO.sub.3, for example, is soluble in hot or ammoniated water. Oxide, chloride, nitrate, sulfate, halogen, and other compounds of iron, manganese, nickel, antimony, lead, tin, copper, zinc, and bismuth are similarly soluble. Mixtures of W/Mo and the other solutions may be compounded into proper compositions for the metallization of various ceramics. The use of solutions of compounds, e.g., halides, of nickel, lead, tin, zinc, and copper allow these metal compounds to be reduced in a hydrogen or nitrogen/hydrogen atmosphere to supply the braxe metal. In a single processing step, then, complete metallizing, brazing, and bonding is possible.
One difficulty of metallizing MACOR, Corning Glass's machinable glass ceramic, by the solution method is the relatively low, allowable metallizing temperature of about 950.degree. C. The solubilities of the metallizing compounds are also restricting factors. Still, many potential metallizing compounds are soluble or at least partly soluble. Zinc chloride and sodium molybdate, for example, are soluble up to 432 and 65 grams, respectively, per 100 cc of cold water. Such a composite solution may be filtered to remove solid particles and used for metallizing various ceramics.
Useful W/Mo-based metallizing compounds include: X(X=W or Mo), XO.sub.3, Na.sub.2 XO.sub.4, K.sub.2 XO.sub.4, Li.sub.2XO 4, and XH (H=F.sub.2, Br.sub.2, Cl.sub.2, and I.sub.2). Useful braze metal compounds include: many YNO.sub.3, YH (Y=Cu, Ag, Au, Zn, In, Fe, Ni, Mn, Ga, Sn, Pb, Cd, Tl, . . . , and H=F, Br, Cl, and I). Many of these compounds are soluble in water, alcohol, or solvents and can, therefore, be used to prepare metallizing solutions. Knowing the elemental atomic weights, one can readily compute the weight of metallizing w or Mo or braze metal in each gram of these chemical compounds.
Another important consideration in making joints between dissimilar materials relates to thermal mismatch stresses and strains. In any ceramic-metal joints, or for that matter, any joining of two dissimilar materials, the matching or mismatch of their thermomechanical characteristics in general, and thermal expansion coefficients in particular, is extremely important. From this mismatch of their thermal expansions, thermal stresses are generated.
Mismatches in other thermomechanical characteristics also result in other thermomechanical mismatch stresses and strains. The magnitude of these mismatch stresses and strains determines the failure probability of the joint.
generally, the thermal expansion mismatch differentials of within 100 ppm (parts per million) are considered as allowable, according to Hagy and Ritland's paper on "Viscosity Flow in Glass-to-Metal Seals," J. Amer. Ceram. Soc., vol. 40, pp. 58-62, 1957. Such thermal expansion coefficients and differentials relate only to the static or equilibrium case, and may not truly represent dynamic or transient conditions when the joint is being heated up or cooled down. Yet such transient conditions often exist during the services of the joint.
Unlike the commonly used static thermal expansion mismatch, the dynamic mismatch in thermal expansion coefficients is not constant but varies with the bonded materials shapes and sizes, physical and surface properties, and heating or cooling conditions and times.
As can be shown, the dynamic expansion strain mismatch may exceed the yield point of the ceramic materials, while the dynamic mismatch stress often exceeds the flexure or even comprehensive strengths of these same materials. What fails most ceramic-metal joints, or cause most coatings to crack, peel, flake, or spall, is thus the dynamic, rather than the static, thermal expansion mismatch.
Using this dynamic mismatch technique, we have been able to determine the location, magnitude, and occurrence time of the maximum or critical mismatch stresses, and take measures to reduce the dynamic mismatch stresses on the relatively weaker ceramic so that the ceramic is no longer failing from the high stresses.
Dynamic mismatches result partly from the fact that metals and ceramics have widely different thermal conductivities. The conductivities for metals range from 0.014 cal/sq. cm/cm/degree C/sec for tellurium, to 1.0 for silver (same unit), while those for ceramics are from 0.0018 for glass to 1.8 for beryllia.
During heating of a ceramic-metal joint, the ceramic temperature lags behind that of the metal, often markedly so; while under cooling the opposite is true. This produces different temperature profiles in the metal and ceramic at a particular time instant on either heating or cooling. Dynamic mismatches in temperatures, strains (i.e., expansions on heating or shrinkages on cooling), and stresses (strains multiplied by Young's moduli) then result.
Take the special example of the case of a long ceramic rod joined end-to-end to a similarly sized metal rod. The ceramic may be, for example, Corning Glass's machinable glass ceramic (MACOR), while the metal may be SAE 1010 carbon steel. The joint is brazed at 950.degree. C. and is, for the worst-case condition, suddenly air quenched in a room-temperature (20.degree. C.) ambient.
The Fourier equation for independent radial heat conduction in long ceramic and metal cylinders is well known. The solution of the cylindrical heat conduction problem consists of an infinite series, each term of which is a product of a Bessel's function and an exponential function, as given in various textbooks on heat conduction. Data tables and master charts for cylindrical heat diffusion have been compiled. See, e.g., 1961 Gebhart's "Heat Transfer," McGraw-Hill, New York). With these equations, data tables, and master charts, one can determine the temperature profiles at different locations (e.e., radial positions, r, in a cylindrical end-to-end joint) at various time instants. From these profiles the critical temperature profile with the associated, maximum transient mismatch stresses and strains can be calculated.
The cooling down of a MACOR-metal joint from the brazing to room temperatures represents one of the most severe thermal changes, because of the wide temperature range involved. The step-by-step temperature changes, i.e., u.sub.m and u.sub.s for the temperatures of MACOR and steel, respectively, at cooling time t in seconds, at the center, (r=0) of the interfacial regions of a two-inch diameter, rod-type MACOR-steel joint are given in Table 1. Other tables have also been prepared for rods of different diameters.
The data used in the computations for Table 1 are: rod diameter D=2 in=5.08 cm, rod radius r=2.54 cm, surface heat transfer coefficient=0.1 per inch for both steel and MACOR, thermal diffusivities=0.108 cm.sub.2 /sec for steel and 0.0054 for MACOR, initial temperature of both MACOR and steel=950.degree. C., and room temperature=20.degree. C.
The computed data in Table 1 show, for the particular case treated, the maximum temperature differential between MACOR and steel at the axial center point, (or r=0), i.e., .DELTA.u=u.sub.m -u.sub.s, at different cooling times t in seconds. Thus immediately upon cooling after brazing (t=0), this differential is zero because both the MACOR are at the same brazing temperature of 950.degree. C.
Subsequently, faster cooling of the steel rod relative to MACOR causes this differential to increase with time t, until both rods are significantly cooled when the temperature differential decreases. After 29,900 seconds (8.3 hours), for example, both rods are within a few degrees of the rooms temperature at 20.degree. C. The maximum temperature differential reaches 775.degree. C. at about 1,000 seconds after the air cooling, giving rise to the maximum or critical dynamic mismatch stress and strain. Table 1 also shows that the temperature differential T=u.sub.m -u.sub.s reaches 113.degree., 144.degree., 217.degree., 282.degree., 478.degree., and 703.degree. C. at 47.8, 59.8, 89.6, 119, 239, and 358 seconds, respectively, after the same cooling from 950.degree. C.
By comparison, the maximum temperature differential of 727.degree. C. at the axial center point of a one-inch (or r=1.27 cm) diameter MACOR-steel joint is reached sooner, at about 440 seconds after cooling.
The linear thermal expansion coefficients, f, are defined as the thermal expansion per unit length per unit degree Centigrade. As given in the literature, they refer only to the static case. These coefficients are constants, at least for the respective temperature ranges. Within these ranges, they are, therefore, independent of the specimen temperature, cooling or heating rates. In addition, these coefficients do not depend on the specimen geometries, sizes, diffusivities, surface characteristics, heating or cooling conditions, and initial and final temperatures. Each material thus has a singular, unique static expansion coefficient, at least for a given temperature range.
The static thermal shrinkage (or negative expansion) strain, e, for a given material is, by definition, the static thermal expansion coefficient, f, multiplied by the temperature range of cooling, u, i.e., e=f.times..DELTA.u. Thus, for the steel rod, this strain is: e.sub.s =f.sub.s .times..DELTA.u.sub.s, and for the MACOR rod, it is: e.sub.m .times..DELTA.u.sub.m.
Under equilibrium conditions, the materials of the joint, i.e., MACOR and steel, are supposed to be in constant thermal equilibrium. That is, U.sub.m =u.sub.s. Both materials are thus at the same brazing temperature of u.sub.0 at the beginning of cooling (t=0). Also, at any time t during the cooling, the cooling temperature range for MACOR and steel are always the same in the static case. Thus:
and the static expansion mismatch strain between steel and MACOR is:
On the other hand, dynamic thermal expansion coefficients, f*, and the resultant dynamic mismatch strains, e*, and stresses, s*, strongly depend on the joint materials, geometries, sizes, physical and surface properties, and heating or cooling conditions.
Starting with zero strain on cooling from the brazing temperature of 950.degree. C., the dynamic strain in the steel rod is: e*.sub.s =f.sub.s .times..DELTA.u.sub.s where .DELTA.u.sub.s =950-u.sub.s, while that in the MACOR rod is:
The difference in dynamic mismatch strain, i.e.,
The dynamic mismatch strain reaches a maximum of about 0.0123 at t=1,000 seconds. Such high strains exceed even the yield point of steel joined to the rigid MACOR ceramic.
The dynamic thermal expansion coefficient mismatch, .DELTA.f*, can be computed by dividing the dynamic mismatch strain, e*.sub.s -e*.sub.m, by the average cooling temperature range, i.e., .DELTA.u.sub.v =950-(u.sub.s +u.sub.m)/2. This dynamic coefficient mismatch, for the 2-inch MACOR-steel rod joint cooling from 950.degree. C. to 20.degree. C., still depends greatly on the cooling time t. It reaches a maximum rate of about 29.6 ppm/degree C. at a cooling time of about 90 seconds, but continuously drops down to less than 5.6 pp/degree C. at t=1,000 seconds, as can be computed from the data in Table 1. The total dynamic coefficient mismatch over the temperature range of 930.degree. C. far exceeds the maximum of 100 ppm considered allowable by Hagy and Ritland.
In can also be shown that the dynamic expansion coefficient mismatch, .DELTA.f*=(f*.sub.s -f*.sub.m)av, for the 2-inch MACOR-steel rod joint cooling from 950.degree. C. to 20.degree. C., is two to five times greater than the corresponding mismatches for the static or equilibrium case, for cooling time t of 10 to 10,000 seconds.
Statically, MACOR only marginally "matches" with a few low-expansion metals such as Sylvania #4, Dumet, 50% nickel alloys, chrome-iron stainless, platinum, Sealmet, and titanium, according to Corning Glass. Because of this two to five times greater dynamic expansion coefficient mismatch relative to the static coefficient mismatch, we must conclude that, dynamically, MACOR and steel joints now become totally "mismatched", at least in so far as the specimen configuration, size, and brazing conditions given above are concerned.
To approximately compute the dynamic mismatch stresses, one may further neglect the presence of the braze and the metallized layers, and use a Timoshenko approach as follows. Consider a portion of the steel specimen of unit length and unit cross-sectional area, brazed together with a MACOR specimen of equal length and cross-sectional area. At time t=t after cooling from the brazing temperature of 950.degree. C., the temperature of the steel is u.sub.s and .DELTA.u.sub.s =950-u.sub.s, while the temperature of MACOR is u.sub.m and .DELTA.u.sub.m =950-u.sub.m. The steel specimen has thus shrunk from unit length to 1-f.sub.s .times..DELTA.u.sub.s, while the MACOR has shrunk to 1-f.sub.m .times..DELTA.u.sub.m. The steel has shrunk more than MACOR, since both f.sub.s and .DELTA.u.sub.s are greater than f.sub.m and u.sub.m. To maintain joint integrity particularly at the ends the originally stress-free but overshrunk steel must be stretched with dynamic tensile stress s.sub.s * by the adjoining MACOR, to length y from length 1-f.sub.s .times..DELTA.u.sub.s, while the undershrunk MACOR must be compressed with dynamic compressive stress s.sub.m * by the steel, to the same length y from length of 1-f.sub.m .times..DELTA.u.sub.m.
Hence, the tensile stress in the steel, s.sub.s *, is
where e.sub.s is the Yound's modulus of steel, i.e., 30,000,000 psi;
while the compressive stress in MACOR, s.sub.m, is
where E.sub.m is the Young's modulus of MACOR, i.e., 5,000,000 psi.
Apparently, s.sub.s *=s.sub.m. Hence,
From these equations, we compute the equal dynamic mismatch stresses in MACOR and Steel, s.sub.m *=s.sub.m, to reach over 52,800 psi, well above MACOR's flexural strength of 15,000 psi or even its comprehensive strength of 50,000 psi.
Similarly, dynamic or transient differences in temperatures, thermal expansion coefficients, thermal expansion strains, and thermal mismatch stresses have been computed for differently sized cylindrical MACOR-steel joints, at various radial locations and cooling time instants. The dynamic mismatch stresses and strains are all unexpectedly high. Measures must therefore be taken to reduce the dynamic mismatch stresses on the relatively weak ceramic so that the ceramic is no longer subjected to the high stresses. This reduction can be achieved by, e.g., absorbing a major portion of the dynamic mismatch stresses normally present in the ceramic through the use of a soft, yieldable metallic braze. These measures prevent the braxed joint failures particularly from these dynamic mismatch stresses, because residual or actual mismatch stress between the two joined materials is the theoretical mismatch stress with a portion thereof absorbed in the metallized or brazed layer.
Specifically, this invention also describes the following methods, for uses singly or in combination, to minimize or neutralize these high mismatch stresses and strains:
(1) Using a soft, yieldable metal layer to braze the metallized ceramic to the metal, and to absorb within the braze layer a large or major portion of these mismatch stresses so that the relatively weak MACOR or other ceramic is no longer subjected to high stresses thereby preventing fractures;
(2) Radially grading, or controllably and gradually changing in (i.e., parallel rather than perpendicular to) the joining plane or bonding interfacial region, the thermal conductivity conductivity (or reciprocal of thermal resistivity), expansion coefficient, and tensile strength of the braze metal, to ensure that the maximum residual mismatch stress, after absorption in the braze or the shock-absorbing interfacial region to be described below, will not exceed the local material strength in the ceramic at any point and time;
(3) Axially grading, or controllably changing normally of or perpendicular to the joining plane or bonding interfacial region, from the ceramic side toward the metal side, the thermal expansion coefficient of the braxe layer to minimize direct mechanical interaction between the steel and ceramic members;
(4) A toughened and strengthened microengineered interfacial region between the ceramic and metallized layer to absorb thermomechanical shocks; and
(5) A new method to achieve flawless bonding regions.
The first two objectives are achieved by providing a novel composite metallic braze layer of disc 10. This composite metallic disc joins together a ceramic body 14 and a metal body 15, as shown in FIGS. 1a and 1b. This disc, lying parallel to and forming part of the bonding interfacial region, has a central copper core 11 inside an outer copper alloy ring or washer 12 made of, e.g., 70:30 Catridge brass. The linear thermal expansion coefficient of pure copper is 16.5 ppm.degree C, while that of, for example, 70 Cu:30 Zn Catridge brass is 19.9 ppm/degree C. Also, the tensile strength of the brazing-annealed, soft pure copper is only 15,000 psi, while that of the 70:30 Catridge brass is over 40,000 psi, or about three times greater.
Hence, the tensile strength and thermal expansion coefficient of the peripheral region in my composite brake disc is 2.67 times and 1.21 times, respectively, those of the central pure copper core. The thermal conductivity of pure copper at 0 degrees C. is 0.920, while those of 11% and 32% Zn:Cu are 0.275 and 0.260, respectively. Hence, the thermal conductivity of 30% Zn:Cu Catridge brass is about 0.261. That is, the thermal conductivity of the peripheral catridge brass in our composite braze disc is only, 0.261/0.920=28.4%, or much less than 50% or 70% of, that of the central pure copper core.
These combinations of linear thermal expansion and tensile properties achieve the required results. In a ceramic-steel joint, the maximum or critical transient mismatch temperatures, dynamic expansion mismatch, and thermal strains and stresses occur in the axial centers of the interfacial regions. I therefore have dead soft, brazing-annealed, pure copper at the core regions. This copper, with a yield strength less the fracture strength of the ceramic, is highly and readily yieldable to absorb much or most of the dynamic mismatch thermal strains and, therefor, stresses. Pure copper also has relatively low thermal expansion to reduce these mismatch effects in the first place. In addition, the pure copper is a good thermal conductor, equalizing the temperature between the centers, as well as their outer and regions, of the steel-MACOR joint, and thus further minimizes mismatch strains and stresses.
On the other hand, the outer peripheral regions of the braze disc is made of relatively highly expansive but the low thermal-conducting brass. At these peripheral regions, the mismatch temperature differentials are relatively smaller. The higher tensile strength is even desirable at the large-area peripheral regions to enhance the joint strength.
This composite braze disc design will thus provide the required radially tailor-graded profiles of braze composition, thermal expansion coefficient, braze strength, and thermal conductivity needed to overcome or minimize the critical dynamic mismatch stresses.
The composite braze discs can be made by, for example, metallurgically cladding or mechanical press-forming a sphere or disc inside a washer, at least two concentric tubes, or other combinations of different materials.
Elemental interdiffusion during the braze manufacture, brazing operation, or special pre- or post-brazing heat-treatment can modify or provide any reasonable composition profiling in the braze discs for achieving the desired tailor-grading results.
If all these measures are still insufficient to prevent dynamic thermal mismatch failures axial elemental grading or sudden composition changes may be added. One method consists of providing a disc of low-expansive metals such as Sylvania #4, Dumet, 50% nickel alloy, chrome-iron stainless, platinum, Sealmet, and titanium placed intermediately between the steel and the copper braze. In this way, the ceramic MACOR is mechanically isolated from the highly expansive steel. The desired elemental profiling can also be achieved through controlled diffusion.
Skilled persons can, of course, select other yieldable metals such as gold, silver, tin, lead, indium, zinc, or even iron or nickel, or other materials to replace copper, and select other chemical elements to replace the copper-strengthening zinc. The resultant new alloys will, of course, be different in compositions, strengths, diffusivities, thermal conductivities, melting or softening points, and other properties.
In addition to achieve metallizing uniformity and minimal mismatch stresses, I have also found it desirable to microengineer the chemical compositions, microstructures, and mechanical properties of the bonding interfacial regions between the ceramic and metallized layer. Merely perfecting the interfaces surfaces alone, as is commonly done, is inadequate to produce strong and reliable joints for withstanding the unavoidable, severe mismatches stresses and strains as shown above.
Different physical, chemical, and electrical metallizing or film-forming methods have been developed. Each has its unique advantages. Some, for example, are atomically precise. Others thoroughly clean the substate surfaces for better adhesion. Some other result in crystalline epitaxy, which is necessary for semiconductor or other devices. Others produce splat cooling and superfine grains, with resultant enhanced mechanical properties, for example, increased Young's modulus. Still others are done at low temperatures to avoid unwanted thermal effects. But non deal effectively with the critical problem of thermal mismatch stresses and strains.
For extremely shock-resistant joints or metallized layers, I have found it absolutely necessary to have a carefully microengineered interfacial layer between the ceramic and the metallized layer. this layer is designed to absorb the major portion of the always present mismatch stresses and strains. Many of my ceramic metallizing processes typically last more than 20 minutes and involve liquid-forming layers containing, directly or indirectly, MoO.sub.3 which melts at 801.degree. C., and WO.sub.3 which melts at about 1550.degree. C. but can be further reduced by alloying with other compounds of metals such as ZnO or PbO. Liquid diffusion is rapid with diffusion coefficient D.sub.1 =1 E-4 to 1 E-5 cm.sup.2 /sec. Processing for t=20 minutes gives a diffusion length of up to the square root of D.sub.1 .times.t=0.35 to 0.11 cm. In addition, a diffused interfacial layer of graded composition, microstructures, and mechanical properties is formed which can be highly shock-absorbing.
In contrast, most conventional coating processes involve only solid-state diffusion. Solid diffusion is slow with diffusion coefficient D.sub.s =1 E-10 to 1 E-20. Even for the same processing or diffusion time t, which these processes do not have, the diffusion length is only 3.2 microns to 3.2 A, or several orders of magnitude shorter than that in my liquid diffusion case. The mismatch stress gradient is thus proportionately steeper.
Plasma spraying does involve liquid droplets in rapid transit. These extremely high-temperature droplets impact the substrate at very high velocities resulting in splat cooling with millsecond liquid dwell times. The resultant diffusion is thus also over three orders of magnitude shorter than my metallizing or metallizing-brazing case. Splat cooling gives very fine grains with high Elastic moduli which actually increase the mismatch stresses. Also, the superheated liquid particles form oxides, mitrides, or other surface layers during transit preventing perfect bonding between the particles themselves. Laser, electron, and some other energetic beam enhanced coating processes also give splat cooling and solid-duffusion conditions.
Without applying any external pressure to force the joining members together, I have used metallizing and bonding processes described above to join various ceramics to metals with pure copper brazes. A typical metallizing process comprises using a mixture of metallizing composition such as WO.sub.3 -Fe.sub.2 O.sub.3 or MoO.sub.2 -MnO.sub.2 in suspension or paste form and applied onto the ceramic, heating for 5 to 25 minutes the coated ceramic to about 800.degree. C.-1500.degree. C. but under no applied pressure. The ratio of heavy metal W or Mo to Fe or Mn after reduction from the compounds is between 9:1 to 6:4. This metallizing may be followed by or simultaneously done with brazing with, e.g., copper or its alloys. Hydrogen or forming gases of 10 to 40 volume % of hydrogen is the desirable metallizing atmospheres.
The metallizing temperatures and times depend on other factors. Diamond, for example, should not be metallized above about 900.degree. C., to minimize chemical reactions nor should graphite fibers be treated above about 750.degree. C. In both cases, a carburizing atmosphere, such as one containing CH.sub.4 or propane, may be useful to prevent too much loss of carbon. The ceramic I have already bonded with my W/Mo-based metallizing methods described here include: alumina, zirconia, silicon carbide, beryllia, yttria, graphite, quartz, silicon, mullite, cordierite, Corning's MACOR and Vision glass, diamond, peizoelectric ceramics, and 123 high-temperature superconductors. Useful structural metals for the joints include copper, nickel, stainless steel, high-nickel or cobalt iron alloys, or even highly "mismatched" ordinary cold-rolled SAE 1010 carbon steel. These joining metals can, therefore, be used as the braze metals themselves for more refractor metals in the joints. Even with the "mismatch" between ceramic and carbon steel, structural joints brazes with pure copper can be repeatedly thermal cycled without fractures between 980.degree. C. and ice water followed by mechanical shocks including 8 to 10foot drop tests onto carpeted, wood, or even marble floors.
These results show that with my improved precesses, low-cost "mismatched" ceramic/metal, carbon-metal, ceramic-ceramic, or ceramic-graphite joints can be made; that these joints can be mechanically strong and thermally shock resistent; that the bonding processes, being ceramic material-limited, need no further improvement for the particular material combinations and thermal shock requirements; and that these joints are, after bonding and thermomechanical shocks, free of pores, microcracks, inclusions, inhomogeneities, and other defects at which fractures originate. Each of these shocks would multiply the number of defects exponentially and have failed the joints. These joints, including particularly the metallized layers, thus compare favorably with, e.g., certain ceramic-metal jonts or ceramic materials developed at great cost, as reported in the literature.
Note that our new joints use only thin lyeres, not bulks, of tungsten/molybdenum; and generally contain no other strategic and expensive metals such as nickels, cobalt, or chromium. The metallized layer adheretly joins to the ceramic. Upon the metallized layer, tenacious, protective metal or ceramic layers can be brazed or formed which resist spalling, peeling, and thermomechanical shocks. Improved corrosion, wear, or frictional properties on these coatings are also possible by suitable selection of the coating materials. A solid lubricant system may be made, e.g., comprising graphite, talc, or MoS.sub.2 powders chemically bonded in copper, bronze, nickel, steel, or cast iron. Also, carbon-carbon composites with improved strength and oxidation resistance are possible. Advanced chemically bonded intermetallic compounds and materials (titanium or hafnium carbide, and titanium or nickel aluminides) are also made available. The same W/Mo-based metallizing compositions are even usefusl as almost universal high-temperature adhesives or sealants for ceramics or metals.
It is even possible to leave only the metallizing molybdenum and/or tungsten between the materials to be joined without any braze metal layer, the operating temperature of the joint is then generally limited by the melting point of the metallized layer.
The flawless and defect-free quality of my ceramic-metal joints or metallized layers on ceramics or graphite are particularly important for tough, fatigue-resistant, protective, easily wettable, and thermochemically stable coatings on, or joints between, ceramic, metals, or graphite, A metallized or coated graphite fiber, for example, cannot tolerate a single pinhole or microcrack that allows oxygen to penetrate and to destroy fiber. Ceramic coatings on metals also cannot have defects when exposed to chemically reactive, high-intensity ion or plasma, high temperature, or other extreme environments. High-melting precious metals such as Pt, Os, and Pd and oxidation resistant metals such as Cr, Al, and Ni are therefore beneficially applied onto the metallized layer, or be formed simultaneously with a metallizing-brazing composition in a single-step metallizing-coating process. Less protective metals such as gold, copper, magnesium, titaniu, or zirconium may also be applied onto, formed simultaneously with, the metallizing layer, followed by coating by electrolytic, electroless, or spraying methods, of the above-mentioned oxidation resistant metals for oxidation protection.
In addition, the metallized or metallized/brazed layers have good wetting characteristics. Further, the metallizing or metallized/brazed layer penetrates and seals all surface pinholes, microcracks, or other defects to strengthen the ceramic at the bonding region. A thick (over 100 microns thick) metal layer further toughness the brittle ceramic. Graphite cor carbon fibers or particles may thus not only be oxidation resistnat but surface toughened and non-brittle.
Coated with my metallized/brazed films up to 20 microns thick, ceramics, boron, graphite, diamond, or glass powders 0.5 through 50 to 200 microns in microns in diameters, are also suited for specific particulate reinforced composites. Upon compacting and sintering these metal coated particles to proper densities and mechanical properties, special acoustic or otherwise damping materials are obtained.
Sapphite, quartz, alumina, or zirconia tubes sealed vacuum tightly to nibobium tantanum, or other ceramic tubes make useful electronic cavity or optical windows for services to or over 1300.degree. C. or 1500.degree. C. My bonding method will also avoid the usual frits seals which are weak, contaminating, short-lived, deteriorating to electrooptical characteristics of the component, and otherwise unreliable in operations.
Defect-free or flawless coatings or bonding are also necessary to contain dangerous materials and should be used to replace weldments which almost always have bubbles, oxides inclusions, segregation patterns, severe residual stresses, weak grain boundaries, or other defects.
The strong, defect-free, and thermomechanically shock-resistant quality of the metallized layers on ceramics, graphite, diamond, and reactive metals such as titanium zirconium, aluminum, or stainless is especially important in the manufacture of advanced composites. Here, the reingorcing fibers, particulates, sheets, or two- or three-dimensional weaves of the ceramics, graphite, boron, oxides of aluminum or zirconium; and carbides or nitrides of Ti, Zr, Hf, V, Nb, Ta, Cr, No, or W; borides of carbon or nitrogen; silicides, aluminides, other intermetallics; diamond; and metals are then perfectly not only wetted by, but bonded to, the matrix of metals, ceramics, carbon, borides, carbides, diamond, . . . Good interfacial bond strengths in, e.g., about 20 volume % graphite, SiC, or Si.sub.3 N.sub.4 fibers or particles in Type 6061 alumina, or zirconia, allow load transfer to occur between matrix and reinforcing particulates, fibers, or weaves thereby giving maximum specific moduli and strengths. These defect-free bondings at the interfaces prevent debondings and allow ideal load transfer between, within, and along the reinforcing members thereby achieving maximum strength, prodcution yield, and productivity at minimum costs.
By replacing the soft, yieldable braze metal pure copper (with melting point 1083.degree. C.), silver (961.9.degree. C., gold (1064.4.degree. C.), tin (232.0.degree. C.), zinc (419.6.degree. C.), lead (327.5.degree. C.), antimony (630.degree. C.), cadmium (320.degree. C.), aluminum (660.4.degree. C.), magnesium (648.8.degree. C.), gallium (29.8.degree. C.), indium (156.4.degree. C.), thallium (303.5.degree. C.), bismuth (271.3.degree. C., . . . , and their alloys with higher-melting metals such as beryllium, chromium, cobalt, hafnium iridium; iron, manganese, nickel, niobium, osmium; palladium, platinum, protoactiniu, rhenium, rhodium; ruthenium, samarium, scandium, silicon, tantalum; thorium, titanium, uranium, vanadium, yttrium, and zirconium, the allowable operating temperatures of the joints are raised to near their respective melting points of 1278, 1857, 1495, 2227, 2410; 1535, 1244, 1455, 2468, 2700; 1554, 1772, 3000, 3180, 1966; 2310, 1300, 1541, 1430, 2996; 1800, 1660, 1130, 1890, 1522, and 1852, Centigrade, respectively.
When molybdenum is used as the metallized layer together with an osmium, rhenim, platinum, protoactinium, rhenium, and tantalum braze layer, the melting point of molybdenum, i.e., 2810.degree. c., rather than that of the braze layer, generally limits the useful temperature of the joint. Similarly, when tungsten (melting point 3410.degree. C.) and carbon (melting point 3650.degree. C.) are used as the metallized and brazed layers for more refractary materials, respectively, the lower tungsten melting point dominates. A variety of new, metallized fibers or particulates of, e.g., SiC, Si3N4, Al203, ZrO2, mullite, cordierite, diamond, glass, quartz, and other ceramics can thus be produced that can be used as reinforcement in composites for temperatures over 1500, 2000, 2500, 3000.degree. C., or higher.
Matrix-reinforcement chemical reactions are serious problems in composites. In graphite-aluminum composites, for example, the graphite reinforcement may react with matrix aluminum to form brittle aluminum carbide. At a given service, processing, or other operating temperature over about 800.degree. C., the graphite-aluminum interfacial reactions may thus be intolerable. High-melting metals given above and used as the metallized/brazed layers on the graphite slow down the elemental diffusion rates and, therefore, graphite particulate--or fiber-matrix interfacial reactions. The heavy metals W or Mo and refractory metals slow down even further. This is because the elemental diffusion rates are functions of the ratio of the operating temperature to the absolute melting temperature. At the same operating temperature of, e.g., 550.degree. C., this ratio for aluminum directly contacting graphite is (550+273.1)/(660.4+273.1)=0.882. With nickel braze on the graphite fibers according to my invention, the interfacial reaction is now between nickel and graphite, and the same ratio is reduced to 823.1/(1455+273.1)=0.476. When the graphite fibers are metallized with Mo or W, the same ratios are further reduced to 0.267 or 0.223, respectively. With a wide available variety of metallizing alloys (e.g., W-Fe, Mo-Mn, Cr-Ni Cu-Zn, . . . ) and coated layers on ceramic reinforcing fibers and particulates, these ratios can be selectively chosen to be less than, e.g., 0.6, 0.5, 0.4, 0.3, 0.22, or even less. The matrix-reinforcement interfacial chemical reactions are thereby reduced, weakening of composite strength is minimized and embrittlement of reinforcement or destruction of composite avoided.
Interfacial chemical reactivity between, e.g., ceramic reinforcement and the metal matrix, can be further suppressed or totally eliminated by coating the metallized/brazed layer with chromium or aluminum. Chromium, aluminum, and their alloys form adherent, dense oxides that resist further oxygen penetration to, e.g., the underneath graphite fibers. These specially metallized/coated graphite or carbon fibers are thermochemically stable in oxygen or other oxidizing atmospheres.
Even mismatch ceramic-metal joints made according to my invention refused to fail under repeated, rapid and severe thermomechanical shocks. Further, the final forced fractures occur away from the bonding regions. This shows that the bonds are free of flaws, microcracks, inclusions, and other defects. In addition, the bond is actually stronger than the weaker ceramic member. This is because the liquid layer formed on the ceramic surface during the metallizing step, generally from 5 to 50 microns thick, actually seals surface notches and other flaws. The metallizing W/Mo ingredients, as will be shown, also strengthen the ceramic at the interfacial region through solution strengthening, or formation of microcomposite reinforcement in the form of precipitated particulates and reinforcing roots, branches, or networks. In many composites, weight is a critical consideration. A very thin W/Mo-based metallized/brazed layer, down to several atomic layers in thickness, may be used with or without any copper, nickel, or other braze metal. The formation of a surface liquid diffusion layer 3 to 30 atomic layers (about 10 to 100 A) takes only 10E-9 to 10E-7 seconds, if a liquid diffusion coefficient of 10E-5 cm.times.cm/sec is used. The control of such extremely thin layer can still be achieved by applying a thin layer of the metallizing solution containing the appropriate amount of molybdate or tungstate compounds.
Another problem with composites is that ceramic, graphite, and carbon fibers are very difficult to be perfectly wetted by, or bonded to, metals, other ceramics, or even the epoxy. Because of this difficulty, an airplane or other vehicle made of these composites often structurally fails under cyclic environmental heat-moisture conditions. Under capillary attraction forces, rain or condensed moisture on the composite surface deeply penetrates, or is sucked in, along the tiny passageways in the unbonded or poorly bonded interfacial regions between the graphite or other ceramic fibers and the epoxy, metal, or ceramic matrices. This penetration is facilitated by air release in, for example, an improperly oriented one-dimensional reinforcement where water enters from the outside skin and move freely along the entire length of the fibers, with entrapped air being forced to leave out of the inner surfaces. This fills the composite structure with water. When the environment turns cold, the filled water expands on freezing, disruptively enlarging the passageways and further debonding the reinforcement from the matrices. Repeated filling-expanding cycles may destroy the composites. When a high-altitude airplane lands in a hot humid weather, moisture automatically condenses onto the very cold composite skin and similarly fill the passageways. The vehicle may take off again into the same freezing attitude where the filled water also expands on freezing with disruptive forces. Multiple cycles of landing and high-altitude flying also also destroy the composite.
By uniformly covering these fibers with flawless metallized W/M0-based coatings, with or without brazing materials, the bonding between these coatings and the matrix will also be flawless. Water penetration is then impossible. Periodic coating of all the strands of these fibers 21 along their lengths with nodular metallized spots 22 at a specific distance d apart breaks up the passageways into small compartments of length d (FIG. 2a). Water can now penetrate to no more than the same distance d below the composite surface. Dipping a two-dimensional or three-dimensional fiber weave into a W/Mo-based metallizing solution or paste, again with or without braze, preferentially coats only the intersections of the fibers with the metallizing compound to thereby form the required nodules for stopping deep water penetration (FIG. 2b). The size of the nodular metallized spots can be controlled by adjusting the viscosity and/or solid content of the solution or paste. Wetting control with the addition of acetone, alcohol, house detergent (e.g., Wisk) also helps.
The reinforcing graphite or other ceramic fibers selectively but perfectly bonded at the nodulated or coated spots in the composites achieve excellent load transfer between fibers, or even along the fibers in metal-matrix composites, but allow systematically and controlably unbonded or weakly bonded regions between the nodules, lending to excellent toughness as well as heat and shock resistances.
The ceramic metallizing processes described in this invention also allow the brazing or coating of the internal or external surfaces of ceramics of controlled densities or porosities. More specifically, porous alumina, zirconia, silicon carbide, mullite, and cordierite have already been metallized on either the internal pores, external surfaces, or both. Substantially 100% of the internal surfaces of the porous ceramic can be metallized by my processes. Ceramic filters for, for example, molten steel, aluminum, or other metals or materials are already in wide uses. But the difficulty of perfectly bonding these weak and porous filter ceramic medium to each other or the metal make their uses difficult, tricky, unreliable, and often dangerous. By bonding these ceramic filters to steel wires or plates, as I have don, these handling problems are minimized.
Multi-stage ceramic filters of alumina, zirconia, silicon carbide, yttria, mullite, cordierite, glass, or other ceramics strongly bonded to the same or different ceramic of the same or increasingly finer pore sizes can now be joined together, one on top of the other. Metal-reinforced multi-stage filters can also be made for, e.g., added strength through metal strengthening; multiple-purpose separations of gases, liquids, or solids from one another through physical means due to size differences; absorption by carbon; catalytic reaction by platinum; liberation or desorption of gases such as oxygen, nitrogen, carbon oxides, or hyrogen from the bonded oxides, nitrides, carbides; hydride for doping or addition to the molten metals or other materials; separation of substances of the same gas, liquid, or solid phases; and other special features functions.
Ceramic filters for air, gas, oil, transmission fluids, and cooling water on automobiles, diesels, power generating equipment, and other machineries are already available. Similar filters for various other fluids including molten metals such as steel or aluminum, or catalytic reactors can, with my bonding method, be strongly attached to internal or external carbon steel or stainless steel containers, other metallic, carbon, or ceramic hooks, knobs, holders, fasteners, protrusions, strengtheners, friction contacts, or springy devices for easy handling or to form fluid-tight enclosures without fluid by-passings.
Catalytic materials such as platinum alloys may also be coated on the metallized layer via diffuison coating, brazing, electrolytic or electroless plating. Reactive materials such as yttria or CaO can also be made porous by sol gel, or by controlled powder packing and sintering, to achieve any desired powder sizes and packing or sintered densities. Such reactive ceramic filters, properly bonded to metal structures, may be used, for example, to remove weakening sulfur in high-quality tool steel poured through these filters.
An electric heater may surround, or be embedded in, the porous ceramic filter for periodical activation with electric ohmic heating to burn to ashes or gases the materials remaining on the ceramic filtering medium. This achieves reusable or self-cleaning results.
Many other uses in differing industries of my bonding methods are readily seen. These include ceramic composites, graphite composites, intermetallic composites, metal-matrix composites, coatings on ceramics, graphite, or metals, high-strength chemically bonded ceramics, and self-lubricating materials containing, e.g., lubricating talc, MoS.sub.2, or graphite particles in iron, steel, copper, or nickel. The composites may involve reinforcing fibers or particulates of ceramics, intermetallics, graphite, or metals in a matrix of ceramic, intermetallic, graphite, or metal.
Using my metallizing methods described above, metallized refractory metallic compounds can be formed for uses as the matrix or reinforcement for composites. These compounds include: oxides of Al, Ba, Be, Ca, Cr, Eu, Gd, La, Mg, Mi, Pu, Ru, Sm, Sc, Si, Th, Ti, U, V, Y, and Zr; carbides of Al, B, Ba, Ve, Ca, Hf, Mo, Nb, Si, Ta, Th, Ti, U, V, W, and Zr; borides of Ba, Ca, Ce, Hf, Mo, Ni, Sr, Ta, Th, Ti, U, V, and Zr; Sulfides of Ca, Gd, Sr, U. and Y; nitrides of Al, Hf, La, Nb, Nd, Sc, Si, Pr, Pu, Ta, Th, Ti, U, V, Y, and Zr; and aluminides of Fe, Ni, Pt, Be, and Ti. Particularly attractive among these compounds are: Si.sub.3 N.sub.4, SiAlon, Sic, Al.sub.2 O.sub.3, mullite, AlN, B.sub.4 C, TiB.sub.2, and BN.
Light, strong, tough, and reliable structural Al, Mg, Be, Ti alloys in composite forms can thus be made with metallized graphite, SiC, or other ceramic reinforcement that will operate over 480.degree. C.
Powder of a ceramic, carbon, intermetallics, or reactive metal may be similarly metallized to achieve flawless and perfectly wetting surface characteristic so that the sintered powder compacts or liquid metal infiltrated compostes will form that have unusually high strengths, densities, and thermal conductivites. Such metallized powders can also be cast as particulate reinforcements or strengtheners. These same powders can be cast (by, e.g, hot squeeze method) to achieve near net shape or net shape into complex structures or components.
A multi-purpose procedure for bonding, sintering, purifying, densifying, strengthening, and otherwise improving the high temperature 123 ceramic superconductor is shown in FIG. 3. High temperature superconductors are superconductors which superconducts at above 90 degrees K (Kelvin). In this multi-purpose procedure, a layer of a suitable MoO.sub.3 -based mixture 31 is formed at selected spots on the copper substrate 30, as shown in FIG. 3a. MoO.sub.3 is the key ingredient in many Mo-based metallizing operations. It melts at 801.degree. C. but the melting point can be lowered or raised to selectable temperatures by forming eutectics or compounds with, e.g., CuO, BaO, and Y.sub.2 O.sub.3, and other oxides such as AgO, CaO, or TlO (Thallium oxide), or even flourides, chlorides, or iodides in view of Ovshinsky's promising results on superconducting and particularly current-carrying capabilities, upon this MoO.sub.3 -based layer is spread the YBa.sub.2 Cu.sub.3 O.sub.7-x powders 32. A vertical temperature gradient is applied to the composite so that the top of the superconductor powders is at least 20.degree. to 50.degree. C. below its melting point, while the bottom of the MoO.sub.3 -based layer is above the melting point of this mixture. This mixture layer will melt and sweep upward (FIG. 3a) to achieve the following highly desirable results:
1. Metallizing and bonding of the bottom layer of 123 superconductor to the copper substrate;
2. Temperature gradient zone-melting to purify the superconductor boundaries according to Pfann (See: Zone Melting, Wiley, 1966);
3. Vertically oriented, upward coloumnar grain growth 34;
4. Grain boundary scavenging, oxgenation, or halogen doping;
5. Liquid phase sintering of the superconductor particles for improved sintering speed, density, mechanical strength, and material stabilities partly also due to the purified or doped grain boundaries;
6. High critical current density of the purified, thinner, and oriented grain boundaries;
7. Cushioning or shock-absorbing qualities of the liquid-diffused, chemically and mechanically graded interfacial layer 33 between the superconductor film and substrate; and
8. Simple, low-cost, single-stop and mass-producing but potentially high-yielding flim-making operation.
After this special temperature-gradient multi-purpose operation, most of the impurities will be dissolved in the sweeping zone. This zone eventually comes up to the surface to be frozen into a highly impure layer 35. This impure layer can be removed by, e.g., grinding or chemical etching with mineral acids. See FIG. 3b.
Other high-temperature ceramic superconductors such as Tl.sub.2 Ba.sub.2 Ca.sub.2 Cu.sub.3 O.sub.10 and TlCa.sub.2 Ba.sub.3 Cu.sub.4 O.sub.x can be similarly bonded or treated for properties improvement with the above method. The substrate does not have to be pure copper, but can be other metals such as aluminum, nickel, oriron, glasses, graphite, or diamond. In addition, other ceramics such as Al203, ZrO2, SiC, carbon glasses, diamond, or even metals powders or filaments, may be similarly bonded onto metallic, ceramic, glass, or carbon substrates. The ceramic layer 34 with thinned, purified, oriented grain boundaries have improved physicochemical properties including thermal and electrical conductivities since grain boundaries are well-known to contribute to resistivity.
In ceramic-metal joints other than for superconductor application, however, the above zone-melting procedure is harmful from the bond strength viewpoint. This is because the last-solidifying layer, usually of complex ceramic eutectic compounds, is weak and brittle and reduces the joint strength. The proper cooling direction after the metallizing here should, therefore, not be vertical but horizontal. In this way, the last-forming layer is lateraly swept out of the joint region without harmfully affecting the joint strength.
According to the above disclosures, I microengineer the ceramic-metal, ceramic-metallizing layer, and/or metallizing-braze layers by substantial thickness and, more important, graded composition, thermoconductivity, and mechanical properties. The W/Mo-based metallized layer may be, for example, 10 20 or 30 microns containing a graded interfacial layer up to 5 or 10 microns. The effective liuqid diffusion length described above may range from 5 to the entire 30 microns. These layers are obtained by liquid diffusion, generally over five minutes but up to one hour. The Cu, Ni, or alloy braze layers may also be chemically, mechanical, and physically graded, as described above.
Another important grading of the interfacial layer relates to the microstructure. Many conventional joints rely on superficial adhesion, weak and defective chemical bonding, or mechanical anchoring wit roughened surfaces. Rough surfaces increase surface area by about 41.4% with 45-degree slopes or valleys (FIG. 4a). An important feature of my invention is the principle of rooting (FIG. 4b), branching (FIG. 4c), and networking (FIG. 4d). Straight roots of the metallizing materials 41 penetrate, during the metallizing or rapid liquid diffusion period, deep along the ceramic grain boundaries 40, These roots may be in the form of fibers located at the intersections of the multiple boundaries, or in the form of sheets each located between two adjacent ceramic grains. These fibers and sheets may be straight, extending generally perpendicularly to the ceramic-metal interface (FIG. 4a). They may form branches following the grain boundaries (FIG. 4b). These roots may even flow deeply into the grain boundaries and turn or curve around to form a partial or complete network (FIG. 4c). The formation of these fibers or sheets depend on the surface energies of the metallizing compounds relative to those of the ceramic grains at the metallizing temperature. The depth of penetration also depends on these energies, but primarily on the metallizing temperature and time.
Preferably, these penetrating metallizing material form reinforcement in a matrix of the ceramic material at the interfacial region. This can be achieved by selecting a W/Mo-based metallizing composition which, with the ceramic at the metallizing temperature, forms hare (Mohr hardness over 8 or 9 versus less then 7 or 6 for the matrix), tough, and strong compounds. Useful compounds include PbMoO.sub.4, MgWO.sub.4, CaMoO.sub.4, MnWO.sub.4, MnMoO.sub.4 and the like. In practice, I simply use pure starting materials such as MoO3, WO3, PbO, CaO, . . . , prepare the exact or near stoichiometric compositions for the metallizing compositions, and metallize at a temperature 50.degree. to 200.degree. C. above the melting points of these compounds. By varying the metallizing time, the grain-boundary reinforcing compounds penetrate to different depths, according to the square root of time diffusion law. For example, for a liquid diffusion case with a diffusion coefficient of 10E-5 cm.times.cm/sec, metallizing for 5 to 60 minutes gives a diffusion length or penetration depth of about 0.055 to 0.19 cm. I also achieved moderately different penetrations of reinforcing particles, fibers, or sheets of different penetration depths by changing the metallizing compositions, e.g., from the W-based type to the Mo-based type.
Thus, with my new ceramic-ceramic or ceramic-metal joining methods, new structural joints, coatings, or surfaces can be produced that have wide uses due to their hardnesses (diamond, alumina, zirconia), hardness and resistances to wear (diamond, zirconia) or corrosion (diamond, carbon, alumina), electrical or thermal conductivity/insulation (zirconia, beryllia, diamond, silver, stainless steel), catalytic activity (platinum), and other properties or appearances.
Tool bits of silicon carbon or nitride, alumina, diamond, and other cutting or abrasive materials can, for example, be metallized with my methods and joined to steel holders to form cutting, drilling, milling, or other machining tools. Particles of the same materials, mixed with the W/Mo metallizing compounds together with copper or nickel brazing alloys, can be spread onto inexpensive carbon steel sheets 0.010 to 0.250 mils thick. Upon heating in a reducing atmosphere, a steel sanding sheet or block is formed. The braze metal may be very thin and merely joins the abrasive particles to the steel plate. The same braze metal may have a thickness up to 95% of the size of the particles, to support fully and hold strongly these particles while still allowing their sharp cutting edges to perform.
Gem stones such as diamond, sapphire, quartz, and the like can be mounted onto metal holders. Because of the excellent strength of the bond, minimum contact with holding metals is needed. As shown in FIG. 5, diamond 51 can now be mounted on the tip of a fine wire 52 so that practically its entire back surface can be illuminated. Also, different back characteristics (color, texture, and reflectivity) can now be instantly changed.
The invention, as described above, is not to be construed as limited to the particular forms disclosed herein, since these are to be regarded as illustrative rather than restrictive. Various combinations, equivalent substitutions, or other modifications of the preferred embodiments described herein are obviously possible in light of the description, without departing from the spirit of the invention. In particular, other ceramics such as alumina or zirconia may be used instead of MACOR with the same or a modified metallizing composition. Accordingly, the invention is to be limited only as indicated by the scope of the following appended claims:
Claims
I claim:
1. A method for coating a ceramic with a strong, adherent, substantially defect-free, and thermomechanically shock resistant metallized layer, said layer in its solid form being practically useful at temperatures over about 630.degree. C., comprising: selecting a ceramic metallizing composition having a plurality of mixed powdered metallizing ingredients of differing sizes, said composition when molten causing reactions between the ingredients and also with the ceramic to form the metallized layer thereon; preparing the composition by proportioning the differing sizes of the mixed ingredients to have gravitationally substantially nonsegregating qualities when applied onto said ceramic; coating onto a selected surface of the ceramic a layer of the metallizing composition, heating the coated ceramic surface to a temperature at which the metallizing composition melts to cause said reactions between the ingredients and with the ceramic thereby achieving ceramic metallization; and keeping the composition molten for a sufficiently long time to thereby form on the ceramic by liquid diffusion the strong, adherent, substantially defect-free, and thermomechanically shock resistant metallized layer including a controlled interfacial region of substantial thickness whose microstructure is substantially free of voids, inclusions, and microcracks.
2. A method as in claim 1 wherein the metallizing composition is of the W/Mo-based type; and including providing a metallic shock-absorbing layer at least five microns in thickness at the interfacial region.
3. A method as in claim 2 wherein said metallic layer is made of an annealed metal selected from the group consisting of Cu, including thermal and electrical conductivities since grain boundaries are well-known to contribute to resistivity.
4. A method as in claim 1 for application with carbon reinforcing fibers for the manufacture of carbon composites each fiber having a plurality of strands, and wherein said coating, heating, and keeping steps coat each strand of every fiber and form flawless, nodular metallized spots along the length of each fiber at a prespecified periodic distance apart.
5. A method as in claim 1 for metallizing a ceramic wherein said coating step coats the selected surface of the ceramic with a uniform layer of the metallizing composition comprising an element selected from the group consisting of tungsten and molybdenum, said layer in its entirety having a substantially constant chemical composition; and said heating and keeping steps keep the coated ceramic heated for at least five minutes at a temperature at least 50 degrees Centigrade above a temperature at which the composition melts to thereby form a metal-containing shock-absorbing liquid-diffused surface layer of graded chemical composition.
6. A method as in claim 5 wherein said ceramic is a material selected from the group consisting of diamond and graphite and wherein said heating and keeping steps are done an atmosphere selected to minimize the loss of carbon from the ceramic.
7. A method as in claim 5 for metallizing a ceramic in the form of a graphite fiber and wherein said heating and keeping steps coat the fiber with a liquid-solidified, pinhole-free and microcrack-free layer of a metal selected from the group consisting of Cu, Ag, Au, Sn, Zn, Pb, Sb, Cd, Al, Mg, Ga, In, Th, Bi, Cr, Co, Fe, Mn, Ni, Nb, Pt, Pd, Rh, Ir, Os, and Ru; and including providing the thus coated surface with an ambient-resistant, top surface metal layer at least 100 A thick.
8. A method as in claim 5 for use in forming a powder metallurgy product with powders selected from the group consisting of ceramics, boron, graphite, diamond, or glass in the range of 0.5 to 200 microns in diameter and wherein said heating keeping steps provide metallized films of up to 20 microns thick on each powder, and including the additional step of compacting and sintering the thus surface-metallized powders to prespecified densities mechanical, and other physical properties.
9. A method as in claim 5 for use in a ceramic fiber-reinforced composite subjected to cyclic environmental heat-moisture conditions and wherein said coating step periodically coats all strands of the ceramic fibers along their lengths with nodular metallized spots at a specific distance apart to break up the moisture passageways into small compartments between the nodular metallized spots.
10. A method as in claim 9 wherein said ceramic fibers are in the form of multi-dimensional weave and including: providing the metallizing composition in a liquid or paste form; and dipping the multi-dimensional ceramic fiber weave into the metallizing liquid or paste to preferentially coat only the intersections of the fibers with the metallizing composition to thereby form the nodules for stopping deep water penetration; and controlling the size of the nodular metallized spots by adjusting at least one parameter of the metallizing liquid or paste composition selected from the group consisting of viscosity, solid content, and wettability of the composition.
11. A method as in claim 5 wherein said ceramic is a porous ceramic of controlled density or porosity and wherein said coating, heating, and keeping steps coat and mettalize substantially the entire internal surface of all the pores in the porous ceramic.
12. A method as in claim 1 wherein said selecting step selects a ceramic metallizing composition of two metallizing ingredients; and including the additional step of providing one of the ingredients in a solution form while supplying the other ingredient in powders of a single substance suspended in said solution thereby ensuring substantially uniform and reproducible metallizing results.
13. A method as in claim 1 wherein said selecting step selects a ceramic metallizing composition of at least two metallizing ingredients; and including the additional step of integrating the at least two ingredients into a physically inseparable form so that the ingredients will provide a substantially constant chemical composition throughout the coated layer thereby ensuring substantially uniform and reproducible metallizing results.
14. A method as in claim 13 wherein said integrating step consists of alloying the at least two ingredients into a single substantially homogeneous alloy form.
15. A method as in claim 13 wherein one of said at least two ingredients is a solid powder while the other ingredient or ingredients are solid, and said integrating step consists of coating the surface of said one ingredient powder with the other ingredient or ingredients to form physically integrated coated solid powders.
16. A method as in claim 1 wherein said selecting step comprises selecting the ceramic metallizing composition having two powdered ingredients of densities and sizes d.sub.1 and D.sub.1, and d.sub.2 and D.sub.2, respectively, d.sub.1 and D.sub.2 being respectively greater than d.sub.2 and D.sub.1, both powders being suspended in a common suspension medium of density d.sub.m, and including the additional step of selecting the ratio of the powder sizes D.sub.1 /D.sub.2 to be at least equal to square root of (d.sub.1 -d.sub.m)/(d.sub.2 -d.sub.m) final settling velocities of the two powders in said medium are nearly the same thereby ensuring substantially uniform and to improve the uniformity and reproducibility of the metallizing results.
17. A method as in claim 1 wherein said selecting step comprises selecting the ceramic metallizing composition having a plurality of powdered ingredients of densities and sizes d.sub.1 and D.sub.1, d.sub.2 and D.sub.2, . . . , d.sub.i and D.sub.i, . . . , respectively, all powders being suspended in a common suspension medium of density d.sub.m, and including the additional step of selecting the powder sizes D.sub.1, D.sub.2, . . . , D.sub.i, to make the D.sub.i.sup.2 /(d.sub.i -d.sub.m) substantially constant so that the final settling velocities of all the powders in said medium are nearly the same thereby ensuring substantially uniform and reproducible metallizing results.
18. A method as in claim 1 wherein said coating step coats only a single point on the ceramic while the heating and keeping steps provides a metallized coating on the ceramic at the single point thereby leaving substantially the entire back surface of the ceramic uncoated and fully exposed to the ambient.
19. A method as in claim 18 wherein said ceramic is a gem stone selected from the group consisting of diamond, sapphire, and quartz.
20. A method of strengthening the surface of a ceramic with a metallized coated layer, said layer in its solid form being practically useful at a temperature about 630.degree. C. comprising: preparing a highly penetrating and wettable ceramic metallizing composition selected from the group consisting of refractory metals W and Mo which when molten, is highly penetrating and wettable in the ceramic surface region particularly relative to the surface defects therein; coating the composition onto a selected surface of the ceramic; heating the coated ceramic to a temperature at which the metallizing composition becomes molten; and keeping the coated ceramic thus heated for a sufficiently long time so that the molten metallizing composition becomes highly wettable to thereby penetrate and seal in the coated ceramic surface region defects at which fractures normally originate, said defects being selected from the group consisting of microcracks, pinholes, porosities, and notches.
21. A method as in claim 20 including the additional step of additionally strengthening the ceramic surface layer at the interfacial region through formation in the ceramic of microcomposite reinforcement in the form of precipitated particulates, reinforcing roots, branches, or networks.
22. A method as in claim 20 including the additional step of adding a metal layer over 100 microns thick to further toughen the ceramic.
23. A method as in claim 20 including the additional step of further strengthening the ceramic surface layer at the interfacial region through solution hardening of the ceramic.
24. A method as in claim 20 wherein the penetrating metallizing composition additionally forms in the ceramic microscopic reinforcement in a matrix of the ceramic material in the interfacial region.
25. A method in claim 20 wherein said preparing step comprises selecting a composition which, with the ceramic at the metallizing temperature, forms a substance which is harder than the ceramic.
26. A method as in claim 25 wherein said selecting step selects the composition which forms the substance having a Mohr hardness of at least 8.
27. A method as in claim 25 wherein said selecting step selects the composition which forms the hard substance with the ceramic at the metallizing temperature. because of the surface energies of the substance relative to those of the ceramic grains at the metallizing temperature, said hard substance, being in liquid form and highly penetrating in the ceramic, forming fibers of the hard substance located at the intersections of the multiple ceramic grain boundaries.
28. A method as in claim 25 wherein said selecting step selects the composition which forms the hard substance with the ceramic at the metallizing temperature, said substance penetrating substantially into the ceramic to form a sheet of the hard substance located along the boundary between two neighboring ceramic grains.
29. A method as in claim 25 wherein said selecting step selects the composition which forms the hard substance with the ceramic at the metallizing temperature, said hard substance penetrating substantially into the ceramic to form branches following the ceramic grain boundaries.
30. A method as in claim 25 wherein said selecting step selects the composition which forms the hard substance with the ceramic at the metallizing temperature, said substance alloying substantially into the ceramic and forming precipitated reinforcing particles inside the ceramic grains.
31. A method as in claim 25 wherein said selecting step selects the composition which forms the hard substance with the ceramic at the metallizing temperature, said hard substance penetrating substantially into the ceramic to form roots which flow deeply into the grain boundaries and turn or curve around to form at least a partial network of the hard substance around a multitude of the ceramic grains.
32. A method as in claim 25 wherein said selecting step comprises selecting a W/Mo-based metallizing composition which forms the hard substance in the form of a hard compound of W/Mo, and including: preparing the metallizing composition at the exact or near stoichiometric compositions for the W/Mo hard compound; metallizing at a temperature at least 50.degree. C. above the melting point of the hard compound; and varying the metallizing time from five minutes up to allow the grain-boundary reinforcing W/Mo compound to penetrate to a depth of a fractional centimeter.
33. A method for coating a high-quality layer of powders of a ceramic onto a selected substrate, said ceramic being selected from the group consisting of alumina, zirconia, beryllia, mullite, quartz, intermetallics, diamond, boron, graphite, carbon, silicon, various other carbides, nitrides, aluminides, or borides; glasses, machinable glasses, and oxidized surfaces of reactive metals, comprising: supplying the substrate onto which the powdered ceramic particles are to be coated; providing a layer of a metallizing composition on the substrate, said composition at a sintering temperature being capable of sintering said powdered ceramic particles together into a strong ceramic layer and simultaneously, of coating the sintered ceramic powders onto the substrate; providing a layer of the powdered ceramic onto the metallizing-composition layer; applying at the substrate surface a first temperature at least about 50.degree. to 150.degree. C. above the melting point of the metallizing composition; and applying onto the top surface of the ceramic layer a second temperature which is above both the first temperature and the sintering temperature of the ceramic so as to form a temperature gradient across the ceramic layer and to cause the sintering of the ceramic powders and regrowth of the ceramic grains; the first and second temperature being sufficiently high to form a liquid layer of the metallizing composition, said liquid layer sweeping under the influence of the temperature gradient across the sintered ceramic particulate layer from the substrate upward and carrying therewith undesirable impurities in the ceramic so as to purify the ceramic layer, to facilitate the sintering, densification, and strengthening of this ceramic layer, and to orient generally normally of the substrate surface or along the temperature gradient the sintered and regrown ceramic grains in columnar forms thereby forming the high-quality coated ceramic layer.
34. A method as in claim 33 wherein said two applying steps sinter, purify, densify, strengthen, and otherwise improve at least a specified physicochemical property of the sintered ceramic layer.
35. A method as in claim 34 wherein said improved physicochemical property is the enhanced electrical conductivity of the resultant ceramic layer because of the thinned, reduced, purified, and oriented ceramic grain boundaries.
36. A method as in claim 33 wherein said providing step provides a layer of a W/Mo-based at selected spots on the substrate.
37. A method as in claim 36 wherein said providing step comprises providing a layer of a metallizing composition containing MoO.sub.3.
38. A method as in claim 37 wherein said two applying steps comprises applying at the bottom of the MoO.sub.3 -based layer the first temperature which is above the melting point of this MoO.sub.3 -based layer while simultaneously applying on the top of the ceramic powders the second temperature which is at least 20.degree. C. below the melting point of the ceramic powders, thereby causing the MoO.sub.3 -based layer to melt and sweep upward to achieve the desired results.
39. A method as in claim 33 wherein said two applying steps cause most of the impurities associated with the ceramic powders to be dissolved in the upward sweeping liquid layer, said liquid layer eventually coming up to the surface of the ceramic layer to be frozen into a highly impure layer.
40. A method as in claim 39 including the additional step of removing the highly impure layer swept up on top of the ceramic layer.
41. A method of coating a ceramic with a refractory metal which in solid form is practically useful over about 630.degree. C. comprising: selecting a metallizing compound comprising a metallizing substance selected from the group consisting of refractory metals W and Mo, said compound at a metallizing or reduction temperature being reducible to the corresponding metallic substance, said metallic substance having a melting point above both the reduction or metallizing temperature and the melting temperature of the metallizing compound; preparing the metallizing compound composition; coating the compound of the metallizing substance onto a selected portion of the ceramic surface; providing a reducing environment around the coated ceramic surface; and heating the coated surface to a metallizing temperature above both the melting and reduction temperatures so that the metallizing compound becomes molten and is reduced to the metallic substance to thereby form a refractory metallized layer on the coated surface for practical use at temperature above the reduction temperature but below the melting point of the metallizing substance.
42. A method as in claim 41 wherein said selecting step comprises selecting MoO.sub.3 and said heating step heats to above about 801.degree. C. both to melt the MoO.sub.3 and to reduce it to Mo so that the metallized ceramic is thereby useful above 801.degree. C. but below the melting point of Mo at 2810.degree. C.
43. A method as in claim 41 wherein said selecting step comprises selecting WO.sub.3 and said heating step heats to above about 1550.degree. C. both to melt the WO.sub.3 and to reduce it to W so that the metallized ceramic is thereby useful above the melting point of WO.sub.3 but below the melting point of W at 3410.degree. C.
44. A method as in claim 41 including laterally cooling the molten metallizing layer to sweep the impurities therein laterally outward from the central area of the coated portion.
Patent Citations (8)
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