Chemical Engineering Aspects of Advanced ... - ACS Publications

Ind. Eng. Chem. Res. 1996, 35, 349-377

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REVIEWS Chemical Engineering Aspects of Advanced Ceramic Materials Vladimir Hlavacek* Laboratory for Ceramic and Reaction Engineering, Chemical Engineering Department, State University of New York at Buffalo, Amherst, New York 14260

Jan A. Puszynski Chemistry and Chemical Engineering Department, South Dakota School of Mines and Technology, Rapid City, South Dakota 57701

Advanced ceramic materials can be synthesized and processed using a large variety of different techniques. Synthesis and processing methods of oxide and non-oxide ceramic materials is reviewed with emphasis on solution techniques and high-temperature gas phase and condensed phase syntheses. Net-shape methods such as chemical vapor deposition, reaction sintering, joining, and liquid infiltration are presented. Critical aspects of materials synthesis and future research needs are also addressed. 1. Introduction

2. Solution Synthesis of Ceramic Powders

Materials have always been an important part of human culture and civilization. Likewise, today’s advanced technologies involve sophisticated materials, since all of them utilize devices, products, and systems that must consist of various advanced materials. The current technical development is strongly dependent on new materials with particular mechanical, chemical, electrical, magnetic, or optical properties. The market for advanced ceramics is generally considered to consist of several potential and existing applications as heat engines, heat-dissipating microelectronic devices, cutting tools, superconductors, ceramic armors, bioceramics, optoelectric devices, etc. The significance of this market sector is best demonstrated by the substantial amount of government funding that has been committed to the research and development of advanced ceramic materials. If high performance is demanded on a large scale, the chemical industry is capable of supplying the precursors in the form of powders, monomers, or organometallic compounds. However, current delays in the market development represent the reason why the chemical engineering aspects have not been adequately addressed. Reaction engineering methodology is uniquely suitable to cope with the synthesis, manufacturing, and optimization of advanced ceramic processes. This review includes chapters where the authors made contributions in the past. Because of the space limitations, an idea of a complete review of the recent advances in the field must be abandoned. Two excellent reviews written by chemical engineers are available, namely, “Reaction Engineering of Advanced Ceramic Materials” by Luss (1990) and “Combustion Synthesis of Advanced Materials” by Varma and Lebrat (1992). The objective of our survey is to indicate areas where progress has been made and to delineate other aspects where considerable work remains to be performed.

The oxide ceramics are still the most used materials for technical applications, particularly in electronic and structural areas. The purity of these materials is extremely important, especially for electronic and hightemperature applications. For instance, alumina ceramic components, which are used in electronic devices, are very sensitive to alkaline metallic ions because of their negative influence on the insulating property. Alumina is the most widely used of the synthetic raw materials for ceramics. The production technologies for this material have been developed and used for many years. Although alumina can be represented by the chemical formula Al2O3, its nature varies considerably depending upon its crystalline form, the impurities present, and the average particle size diameter. The major source of alumina is the Bayer process; however, for applications requiring low-impurity content other techniques are used (Somiya (1984)): (a) pyrolysis of ammonium alum, (b) hydrolysis of organoaluminum compounds, (c) pyrolysis of ammonium dawsonite, (d) electric discharge of aluminum in water. Zirconia (ZrO2) based ceramics are becoming more popular in recent years. The main production method of zirconia is the high-temperature decomposition of zircon sand, ZrSiO4, which is melted in an electric arc furnace, with vaporization of silica. Novel solution techniques used for synthesis of oxide and non-oxide ceramic powders are discussed in detail in the next sections. 2.1. Sol-Gel Synthesis of Ceramic Materials. Sol-gel processing has proved to be a very effective method for producing nuclear ceramics such as oxides of thorium, uranium, and plutonium (Wilson and Heathcote (1990)) and structural as well as electronic materials (Yan (1981); Gherardi and Matijevic (1988)). This method allows one to synthesize finely dispersed powders, consisting of particles uniform in size and shape. The microstructure of the fabricated ceramics depends strongly on the characteristics of precursor powders (Blendell et al. (1983); Matijevic (1985)), including

* Author to whom correspondence should be addressed.

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© 1996 American Chemical Society

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Ind. Eng. Chem. Res., Vol. 35, No. 2, 1996

particle shape, size, and distribution, agglomerate strength, adsorbed species on the particle surface, impurities, and the homogeneity of the constituents. These powder properties can be controlled by careful selection of processing parameters such as concentration of reactants, ionic strength, pH, temperature, addition sequence, adsorption of added ions, etc. In spite of considerable effort made in this area, only empirical correlations between process variables and powder properties have been established. So far, kinetic factors and mechanisms controlling the nucleation process, growth, and nature of the aggregation phase are not well documented or understood. A sol-gel process can be described as a synthesis of ceramic particles or polymer type structures from a solution (Schmidt (1991)). It can be distinguished from precipitation by its specific property to stabilize a finely dispersed (mostly colloidal) phase in solution by surface chemistry. The size of solid agglomerates is usually in the range 1-100 nm. The overall sol-gel process can be represented by the following sequence of transformations:

precursors w sol w gel w oxide A very important first step is homogeneous precipitation which should take place in a kinetically controlled regime without subsequent secondary nucleation. In such a procedure the decomposition of complexes is one of the convenient ways to release the necessary precursors to particle precipitation. In principle, the simplest process is the deprotonation of hydrated metal ions to form hydrolyzed intermediates in the production of hydrous metal oxides. This can be controlled by adjusting the temperature, concentration, and pH of the solution. The hydrolysis process can be also promoted by addition into the reacting solution compounds (e.g., urea) which release hydroxide ions in a controlled manner. There exist a large class of organometallic compounds which are suitable for preparation of colloidal metal oxides because of their easy hydrolyzability. Matijevic and his co-workers have synthesized numerous oxide powders from metal inorganic salts or alkoxides using the sol method (Catone and Matijevic (1974); Matijevic et al. (1977); Visca and Matijevic (1979); Matijevic (1985)). Titanium dioxide sols consisting of spherical particles of narrow size distribution were obtained when a highly acidic titanium tetrachloride solution containing sulfate ions was aged at elevated temperatures for extended periods of time. Titanium dioxide can be similarly synthesized by hydrolysis of titanium ethoxide. The list of apparent advantages and disadvantages of the sol-gel method for the synthesis of glasses is presented in Tables 1 and 2 (Mackenzie (1984)). In recent years, a significant effort has been paid to the synthesis of microcomposite powders and electronic and superconductive submicron powders (Gai et al. (1989); Murakami et al. (1990); Kani et al. (1991)). Fine powders of Y-Ba-Cu-O were synthesized using yttrium, barium, and copper alkoxides with stoichiometric composition Y:Ba:Cu ) 1:2:3 by their dissolution in a mixture of toluene and xylene. The scheme of YBa2Cu3Ox powder synthesis by the sol-gel method is shown in Figure 1. One of the major advantages of sol-gel processing is that it permits intimate mixing of components. Addition of certain metal ions (e.g., transition metals) to oxide sol (e.g., alumina sol) followed by vigorous stirring

Figure 1. Schematics of YBa2Cu3Ox powder synthesis by a solgel method. Table 1. Advantages of the Sol-Gel Method over Conventional Melting of Glasses 1. better homogeneity from raw materials 2. better purity-from raw materials 3. lower temperature of preparation a. save energy b. minimize evaporation losses c. minimize air pollution d. no reaction with container e. bypass phase separation f. bypass crystallization 4. new noncrystalline solids outside the range of normal glass formation 5. new crystalline phases from new noncrystalline solids 6. better glass products from special properties of gel 7. special products, such as powders, films, and coatings Table 2. Some Disadvantages of the Sol-Gel Method 1. 2. 3. 4. 5. 6. 7.

high cost of raw materials large shrinkage during processing residual fine pores residual hydroxyl residual carbon health hazards of organic solutions long processing times

prior to the gelation results in a molecular level mixing of the two materials. After gelation and subsequent pyrolysis, the transition element is homogeneously dispersed throughout the alumina matrix on a scale difficult to achieve with standard doping procedures (Johnson (1985)). The full potential of sol-gel processed ceramic powders is, as yet, underexploited, particularly in the field of microcomposite synthesis, formation of non-oxide ceramic powders, and little possibility to a priori predict the actual size and shape of formed particles. The reason for the inability to fully anticipate the results is in the complexity of chemical processes underlying the precipitation with subsequent particle growth and surface interactions. Silicon nitride powder is synthesized by the liquidinterfacial reaction method developed by UBE Industries. The major steps of this synthesis are shown in Figure 2. Submicron powder synthesized by this technique contains a minimum amount of metallic impurities and a high content of the desired R-phase. The typical properties of silicon nitride produced by this technique are listed in Table 3. The uniform particle size distribution and powder morphology makes this material a prime candidate as a high-quality raw

Ind. Eng. Chem. Res., Vol. 35, No. 2, 1996 351

Figure 2. Synthesis of silicon nitride by reactions in a liquid phase. Table 3. Typical Properties of Silicon Nitride Powder: UBE-SN-E-10 purity total nitrogen total oxygen total chlorine total iron total calcium total aluminum bet surface area bulk density after tapping phase composition degree of crystallinity β/(R + β)

38.7 wt% 1.3 wt % 50.0 ppm 50.0 ppm 1200 K), and therefore other types of precursors are frequently considered. The list of metal carbonyls and organometallic compounds commonly used in the industry with selected physicochemical properties is presented in Table 8. Product purity, particle size distribution, and powder morphology are a strong function of (a) thermodynamics, (b) mass transport and reaction kinetics, and (c) hydrodynamics and reactor configuration. Homogeneous gas phase reaction led to the formation of species with low-equilibrium vapor pressures. Once the concentration exceeds the equilibrium level, the vapors undergo a spontaneous phase transition. Ho-

reaction

temp range (K)

(C2H5)2Te(g) + Cd(g) + H2(g) f CdTe(s) + 2C2H6(g) H2Se(g) + Zn(g) f ZnSe(s) + H2(g) TiCl4(g) + 2Mg(s) f Ti(s) + 2MgCl2(g) SiCl4(g) + Na(g) f Si(l) + 4NaCl(g) NbCl5(g) + (5/2)Mg(g) f Nb(l) + (5/2)MgCl2(g) TaCl5(g) + 5Na(g) + (1/2)C2H2(g) f TaC(s) + 5NaCl(g) + (1/2)H2(g)

598-633 873-1073 1073-1273 2000-2400 2600-2800 2900-3200

Table 7. Synthesis of Oxides by Oxidation and Hydrolysis of Gaseous Halide Precursors reaction

temp range (K)

UF6(g) +2H2O(g) + H2(g) f UO2(s) + 6HF(g) TiCl4(g) + 2H2O(g) f TiO2(s) + 4HCl(g) SiH4(g) + 2O2(g) f SiO2(s) + 2H2O(g) 2AlCl3(g) + 3H2O(g) f Al2O3(s) + 6HCl(g) SiCl4(g) + 2H2O(g) f SiO2(s) + 4HCl(g)

1200-1500 1273-1573 473-723 973-1373 1273-1773

mogeneous nucleation is a critical step in the synthesis of powders from the gas phase. There are numerous theories of condensed phase nucleation from vapors. The comprehensive review of these theories is given by Seinfeld (1986). However, these theories still become quite controversial due to many assumptions made in their derivation. Once particles have been formed by homogeneous nucleation, they grow due to condensation, surface reactions, or coalescence. There are a large number of publications related to the mechanism of particle growth and aerosol dynamics (Fucks (1964); Friedlander (1977); Flagan and Seinfeld (1988)). The homogeneous and surface reactions are quite complex and might involve tens or hundreds of elemental reaction steps. Allendorf and Kee (1991) have developed a detailed mechanism of silicon carbide formation by gaseous reaction of silane and propane. This mechanism involves 83 fundamental homogeneous reaction steps and 36 surface reactions. The majority of the gaseous reactions can be categorized as endothermic in their nature. Therefore, an external source of energy is normally required. This

Ind. Eng. Chem. Res., Vol. 35, No. 2, 1996 353 Table 8. Selected Properties of Metal Carbonyls and Organometallic Precursors compound

formula

melting point, °C

boiling point, °C

trimethylaluminum triethylaluminum diethylberillium dimethylcadmium trimethylgallium triethylgallium trimethylindium chromium tert-butoxide germanium ethoxide niobium ethoxide tantalum ethoxide titanium isopropoxide titanium n-butoxide chromium hexacarbonyl iron pentacarbonyl nickel tetracarbonyl molybdenum hexacarbonyl tungsten hexacarbonyl

(CH3)3Al (C2H5)3Al (C2H5)2Be (CH3)2Cd (CH3)3Ga (C2H5)3Ga (CH3)3In Cr(OC4H9)3 Ge(OC2H5)4 Nb(OC2H5)5 Ta(OC2H5)5 Ti(OC3H7)4 Ti(OC4H9)4 Cr(CO)6 Fe(CO)5 Ni(CO)4 Mo(CO)6 W(CO)6

15 -58 12 4 -15 -82 88 36 49 6 21 20 -55 164 -20 -25 150 169

126 194 194 105 5 143 134 66 (2.7 Tr) 185.5 156 (0.05 Tr) 146 (0.15 Tr) 58 (1 Tr) 312 180 (dec) 103 43 180 (dec) (dec)

energy may be provided by (i) flame, (ii) rf or dc plasma, and (iii) laser. 3.1. Flame Synthesis. Flame synthesis has proven to be very effective and the most dominant hightemperature synthesis of ultrafine oxide ceramic powders from chloride precursors (Sato et al. (1983); Kimura et al. (1988); Nickel et al. (1989); Zachariah et al. (1989); Zhao et al. (1990)). Premixed, concurrent, and countercurrent diffusion flames have been explored for the synthesis of both oxide and non-oxide ceramic powders. The ultrafine silica is prepared on a commercial scale by injecting vapor of silicon tetrachloride in a concurrent hydrogen-oxygen diffusion flame (Spalter and MacKenzie (1952)). This process was developed before World War II in Germany by Deutsche Gold Silberscheidenanstalt (Vormals Ro¨ssler), and later it was commercialized in the United States. In 1952 the G. L. Cabot Company under licensing arrangement from the German company offered the product to the American market. Particle size and surface characteristics of the silica prepared via the flame technology depend mainly on the reaction temperature. After combustion the silica is agglomerated, separated from the hydrogen chloride vapor in cyclone separators, and finally calcined in order to remove traces of residual hydrogen chloride. During the combustion, excess air is used to ensure a complete combustion of the hydrogen. Since the presence of even minute quantities of moisture would upset the reaction, air used in the furnace must be dried. During the startup, the hydrogen is burned until the furnace reaches its operating temperature. Then excess air and silicon tetrachloride are fed into the burner. The reaction temperature varies from 1000 to 1100 °C. Within this range, the flame temperature is controlled by adequate ratios of air, hydrogen, and silicon tetrachloride. Fluctuations in the flame temperature can strongly affect the physical properties of the resulting silica product. When the temperature is at the lower bound (∼1000 °C), the product formed is characterized by very small average particle size. Silica made under these conditions is particularly well suited as a thickening agent. At higher temperatures, a product with lower specific surface area is formed. This material is better suited as a thixotropic agent. The discharge from the reactor contains a silica fume which is too fine to be removed from the gas stream by any conventional method of separation. Therefore, ultrafine silica particles must be agglomerated to form particles large enough to be taken out by a cyclone separator, followed

Figure 4. Schematics of various flame configurations.

by a bag filter. Recovered silica is calcined at about 340 °C in a gas-fired screw conveyer. The extremely fluffy nature of silica prepared via the flame route makes it difficult to bag. A special bagging machine is used in which bags are filled under vacuum. Premixed flames are formed by mixing the fuel and oxidizer gases and ignition mixture at a location outside the mixing zone. Since the synthesis of ceramic particles requires very high temperatures (and therefore high concentrations of fuel and oxidizer) the premixed flames have limited applications due to the safety requirements. Diffusion flames, on the other hand, are produced when the fuel and oxidizer are brought into contact at the position of the flame. The combustion front is in the form of a very thin sheet, and the rate of the process is determined by the diffusion of fuel and oxidizer. There are two kinds of diffusion flames: (i) concurrent and (ii) countercurrent (Tsui (1982); Katz and Hung (1982)). The concurrent flame is generated when fuel and oxidizer are fed separately into the center tube and annulus of a concentric tube burner and ignited at the burner mouth. The result is a conical, candle-type flame. The countercurrent flame is obtained when fuel and oxidizer are fed into the reaction zone through two separate and opposing jets. The result is a flat (laminar) flame front. Schematics of various kinds of flames are shown in Figure 4. Different types of flames have been used for syntheses of ultrafine powders from halide precursors: (a) hydrogen-oxygen flame (Formenti (1972)); (b) hydrogenfluorine flame (Smiley (1965)); (c) hydrogen-chlorine flame (Sato et al. (1983)). Titanium dioxide and other metallic oxides can be prepared by a similar route of high-temperature oxidation from titanium tetrachloride or other volatile metallic chlorides, respectively (Frey (1961)). Examples of chlorides which can be used include those of zirconium,

354


hafnium, aluminum, tin, chromium, and iron. Metal chlorides, in the vapor phase, can be burnt in oxygen at high temperatures, e.g., above 1000 °C, to give an aerosol of metal oxide particles in a chlorine-containing gas. The fine divided oxide particles thus produced make the oxide especially suitable for severe commercial applications, e.g., as pigments. The titanium dioxide has a very uniform grain size, with an average diameter of less than 0.5 µm. Chung and Katz (1985) described the development of the counterflow H2-O2 diffusion flame burner as a tool for the study of nucleation of refractory oxides. Light scattering measurement were used to study the nucleation process of SiO2 particles. This study was extended to TiO2-SiO2, SiO2-GeO2, Al2O3-TiO2, and V2O5-TiO2 systems (Hung and Katz (1992); Hung et al. (1992); Miquel et al. (1993)). The authors have proposed a mechanism for the formation of mixed oxides. The nucleation and growth processes vary from system to system, and therefore the theoretical prediction of particle composition and morphology without sophisticated experiments is very difficult. Zhao et al. (1990) and Revankar et al. (1991) performed a detailed experimental study and theoretical analysis of H2-Cl2 and its application to tungsten and tungsten carbide synthesis. They have concluded that submicron powders can be synthesized; however, the control of stoichiometric composition in tungsten carbide powders is very difficult. It can be concluded that reaction kinetics of combustion reaction(s), chemical kinetics of halide reduction, nucleation of the solid phase, and the purity and phase composition of the final products are still not well understood, and additional research is needed to answer these questions. Fine particles can also be synthesized by pyrolysis of inorganic salts. Spray pyrolysis, in its simplest representation, is a process in which a liquid solution of various compounds is sprayed and exposed to high temperatures. As a result of solvent vaporization and thermal decomposition, powders with complex composition can be synthesized. Several oxide powders including MgO, ZrO2, Al2O3, and SiO2 can be synthesized using this technique (Albin and Risbud (1987)). Zachariah and Huzarewicz (1991) used this technique to synthesize YBa2Cu3O7-δ powders. An aqueous solution of yttrium, barium, and copper nitrite salts with a 1:2:3 molar ratio was atomized (0.5 µm droplets), entrained in a stream of oxygen, passed through a diffusion drier, and introduced into the reactor. The pyrolysis of nitrites was carried out in a coannular H2-O2 diffusion flame. The oxidizer was introduced in the inner stream. This type of a flame geometry seems to be the most suitable for controlling particle residence time and time-temperature history. The proposed synthesis appears to be very promising for the production of superconducting particles with well-controlled particle size distribution. A glass optical waveguide preform can be produced by generating fine soot of silica in a flame (Siegfried (1980); Bailey and Morrow (1981)). The soot is attracted to the surface of a mandrel by thermophoresis. The mandrel rotates and moves axially in a furnace equipped with multiple flame soot generators. Heretofore, the mandrel is usually removed, and the soot, high porous preform, is sintered to form a consolidated, clear glass draw blank. The resulting tubular draw blank is heated to a temperature at which the material has a low enough viscosity, and a fiber of the desired dimension is drawn.

Figure 5. dc plasma system for the production of SiC powder: 1, plasma torch; 2, reactor; 3, thermocouple; 4, collector; 5, cyclone separator; 6, bag filter; 7, feeding system; 8, evaporator (taken from Zhu et al. (1992)).

3.2. Plasma Synthesis. The development of plasma chemical synthesis for the preparation of ultrafine powders represents the broadest application of thermal plasma technology to chemical vapor deposition (CVD). Plasma chemical synthesis offers the advantage of a single-step continuous process. Ceramic, intermetallic, and metallic powders are formed in a vapor phase at extremely high temperatures (T > 5000 K). Powders synthesized by this technique are usually very fine ( 2000 °C

(2) (3)

T > 2000 °C (4)

The reactivity of the silicon source increases in the following sequence: Si(CH3)4 > SiH4 > CH3SiCl3 > SiCl4. The equilibrium constants of the above reaction systems are presented in Figure 7 (Zhu et al. (1993)). Reactions 3 and 4 have low equilibrium constants at low temperatures, and therefore high temperatures must be used to obtain significant conversions. The same authors have presented experimental results of silicon carbide synthesis in a 15-kW dc plasma reactor. They investigated the formation of SiC by the vapor reaction SiCl4-CH4-H2-Ar and CH3SiCl3-CH4-H2Ar systems. They demonstrated that high-quality, nanosize silicon carbide can be synthesized in a plasma environment. 3.3. Laser Synthesis. Development of powerful CO2 lasers, which have now reached the market of applications for industrial processes, from both technical and economical points of view, stimulated research in the area of laser-assisted photodecomposition of gaseous precursors in order to form high-purity ceramic powders. Photochemical reactors can be used for the production of ultrapure materials, since the laser-assisted processes take place along the laser beam and no contamination from hot surfaces occurs. Fantoni et al. (1990) have synthesized SiC powder through chemical vapor deposition in a flow reactor in which a continuous-wave CO2 laser was focused. The reaction between silane and acetylene

2SiH4 + C2H2 f 2SiC + 5H2 was initiated by the IR laser, which caused the dissociation of SiH4 to SiH2. The schematics of the experimental setup are shown in Figure 8. They demonstrated that β-SiC with uniform sizes of microcrystallites was formed. The growth mechanism seems

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Figure 8. Laser reactor for the synthesis of SiC (taken from Fantoni et al. (1990)).

to occur via coalescence of β-SiC microcrystallites, and the surface area of powders is in the range of 50-100 m2/g. Haggerty (1984) and Haggerty and Flint (1990) presented conditions for growth of precisely controlled Si, Si3N4, and SiC powders from laser-heated gases. The authors have projected that the manufacturing cost of these powders might be lower than those synthesized by conventional methods. It has been proposed that ceramic powders of the type prepared via the laser pyrolysis may be ideal for the low-temperature processing of parts exhibiting superior microstructure and possessing enhanced physical properties. To date, however, characterization of the densification and physical properties of parts derived directly from lasersynthesized powders is incomplete. Knudsen (1987) has reported the synthesis of boron carbide powders and their densification characteristics. The product consisted of equiaxed, nonagglomerated, ultrafine powders (10-100 nm) with well-defined stoichiometry. The physical properties as well as the microstructure of the densified laser-derived B4C compare favorably to the best described in the literature. The microhardness and elastic modules of 31.4 and 492 GPa, respectively, are greater than those typically reported for dense B4C (29.4 and 450 GPa). The fracture toughness (3.7 MN/m3/2) and transverse rapture strength (455 MPa) are equivalent to the values for high-quality dense B4C. It can be concluded that the nature of the particles generated by flame, plasma, laser, and other thermal techniques is determined by chemical kinetics, particle inception or nucleation, coagulation to form agglomerate particles, and coalescence of those agglomerate particles. The dynamics of refractory aerosols are still not well understood. In order to understand the formation of powders by aerosol processes, the particle properties must be measured while the particles are still entrained in the carrier gas. A detailed description of the dynam-

ics of hard agglomerate formation by sintering or vapor deposition and their fractal structure is also needed. 4. High-Temperature Condensed Phase Synthesis 4.1. Calcination and Reactive Calcination. Traditionally ceramic powders have been obtained by calcination. In general, calcination refers to the heating of inorganic solid materials to remove volatile components (e.g., decomposition of calcium or magnesium carbonates). Many oxide powders of high purity are produced by calcination of crystalline or amorphous hydroxides, carbonates, or other inorganic or organometallic salts. Numerous studies of the mechanisms of thermal decomposition of hydroxides and carbonates have established that the overall reaction process is topotactic (Brett et al. (1970)). Calcination of these materials results in the formation of porous oxide powders consisting of small crystallites. The size of these crystallites strongly depends on calcination temperature, and it may vary from a few to hundreds of nanometers. The calcination process is also strongly influenced by the surrounding gas atmosphere. In many cases the effect of the gas atmosphere on powder morphology is experimentally found; however, the overall mechanism is not well understood. Mixed oxides with perovskite or spinel structure (e.g., titanates, zirconates, and ferrites) are also produced in a calcination process. The solid-state reactions between oxides to form various mixed solid compounds involve solid-state diffusion of metallic ions. Sometimes this process might be accompanied by solid-state or gaseous transport of oxygen (Schmalzried (1974)). Precise control of powder formation and its morphology may be difficult (Samsonov and Upadkhaya (1971)). The ratedetermining step in the reaction rate is the solid-state

Ind. Eng. Chem. Res., Vol. 35, No. 2, 1996 357 Table 9. Carbothermal Synthesis of Ceramic Powders TiO2 + 3C f TiC + 2CO Ta2O5 + 7C f 2TaC + 5CO TiO2 + B2O3 + 5C f TiB2 + 5CO 2B2O3 + 7C f B4C + 6CO SiO2 + 3C f SiC + 2CO Al2O3 + 3C + N2 f 2AlN + 3CO 3SiO2 + 6C + 2N2 f Si3N4 + 6CO

diffusion. Examples of the reactive calcination are as follows:

BaO + TiO2 f BaTiO3 TiB2 + TiC f TiB2‚TiC TiC + TiN f TiC‚TiN In order to obtain fine powders, lower calcining temperatures are recommended. However, if the complete conversion is desired, then proper mixing of precursors and their ultimate contact are essential. This can be accomplished by using coprecipitated gels and spray-dried or freeze-dried mixed solutions of salts. Even though the calcination technique belongs to the most important in ceramic powder manufacturing, the chemical engineering principles of the design of calcining reactors are still lacking. The current design of such reacting systems is entirely based on the rule-of-thumb. 4.2. Carbothermal Reduction of Oxides. Carbothermal reduction of oxides is one of the oldest techniques to manufacture non-oxide ceramics. This technique has been successfully used to synthesize carbides (e.g., TiC, TaC, ZrC, WC, NbC, Cr3C2, Mo2C, MoC, VC, HfC, B4C), borides (e.g., TiB2, ZrB2, HfB2, LaB6), nitrides (e.g., Si3N4, TiN, ZrN, TaN, AlN), and solid solutions based on carbide-carbide, carbide-nitride, and boridenitride materials. Carbothermal reduction of oxides is an endothermic process which usually requires high temperatures (1500-2300 K) to be completed within a reasonable period of time. Some carbides might be formed at significantly lower temperatures (e.g., Mo2C). During carbothermal synthesis of carbides, borides, nitrides, composite powders, and solid solutions, a very large amount of carbon monoxide is formed as a byproduct. Typical carbothermal reduction reactions are listed in Table 9. Commercial production of boron and silicon carbides is mainly done by the electrochemical reaction of highgrade silica sand or boric oxide and carbon (pitch coke or anthracite coal) in an electric batch furnace. This type of process is energy inefficient and environmentally difficult to control. In addition, enormous temperature gradients inside a reaction vessel cause the product’s particle size distribution and chemical composition to be highly nonuniform. There have been many attempts and modifications to produce carbides, borides, and nitrides more efficiently (Smudski (1966); Kim and McMurtry (1985); Shaffer and Blakely (1987)). Resistively or induction heated graphite reactors have been tested and successfully used. This type of equipment has a relatively high capital cost and is subjected to corrosion of heating elements by reactants and byproducts. However, products formed in these reactors have more uniform composition and particle size distribution, and the process is continuous. Many transition-metal carbides and borides are commercially produced in this type of equipment. The carbothermal synthesis of aluminum and silicon nitrides is more complicated because it involves a reduction of

Figure 9. Continuous graphite transport reactor (taken from Weimer et al. (1991)).

alumina or silica by carbon and subsequent nitridation of intermediate products. Thermodynamic calculations show that the reduction of alumina by carbon at temperatures below 1900 K is not possible without nitrogen presence. In nitrogen, the alumina reduction begins at 1400 K. Formation of aluminum oxynitride phases is a very critical step in this process. The method, used by the Japanese company Tokuyama Soda to manufacture high-purity AlN powders, consists of reducing alumina with carbon in a nitrogen atmosphere. The resulting product is later heated in air to remove unconverted carbon. Processing of aluminum nitride powders has to be done in a low-humidity environment due to their high sensitivity to moisture. Any substantial hydrolysis prior to a final sintering results in a poor thermal conductivity of fully dense articles (40-80 W/m‚K). However, when AlN powders with a low oxygen content are sintered, the material exhibits significantly higher thermal conductivity (170-220 W/m‚K). Protection of AlN powders (polymer coating, hydrophobic agents) during a processing prior to a final sintering is extremely important; therefore, more research in this area is needed. Silicon nitride exists in two hexagonal phases (R and β). In general, R-Si3N4 is considered to be a metastable low--temperature phase which seems to be the more desirable phase because of its high-temperature transformation to the β-phase during a densification process. This high-temperature transformation is important to fabricate parts with a very high fracture toughness. Carbothermal reduction of colloidal silica by carbon in the presence of nitrogen at temperatures below 1700 K results in the formation of an almost pure R-Si3N4 phase (>97 wt %). However, at higher temperatures, the product contains a significantly higher content of β-Si3N4 and other complex solid solutions and silicon carbide which are thermodynamically more stable. Recently, Weimer et al. (1991) propose a novel graphite transport reactor which allows a rapid carbothermal reduction. The schematics of this reactor are shown in Figure 9. An intimate carbon and boron oxide mixture was continuously fed under an argon atmosphere into the reactor where the temperature was kept between 2073 and 2573 K. The residence time within the reactor was controlled by the flow of argon, the quantity of CO generated by the reaction, the temperature, and the

358


Table 10. Properties of Ceramic Whiskers (Richerson (1992))

whiskers

density (g/cm3)

tensile strength (MPa)

Young’s modulus (GPa)

thermal expansion (10-6/°C)

SiC Al2O3 BeO B4C Si3N4

3.2 3.9 2.9 2.5 3.2

21 000 21 000 13 000 14 000 14 000

840 580 345 480 380

4.8 5.5 9.5 4.5 2.8

cross-sectional area of the reaction chamber. Chemically homogeneous, uniformly sized submicron boron carbide was obtained. This concept of transport reactor has been also used for the synthesis of aluminum nitride (Weimer et al. (1994)). Carbothermal reduction processes are also commonly used to synthesize non-oxide materials in the form of whiskers. Whiskers are filamentary single crystals that have ultrahigh strength (on the order of 1 million psi) and they range in size from about 0.5 to 10 µm in diameter and a few microns to a few centimeters in length. Analytical and experimental studies indicate that low-density, high-modulus, high-melting-point whiskers such as B, B4C, C, Al2O3, SiC, and Si3N4 have great potential as reinforcing agents for metal, ceramics, and plastics. With the development of emerging technologies such as aerospace vehicles, structural temperatures exceeding 1100 °C are not uncommon. To accommodate these applications, high-temperature, oxidation-resistant materials with high strength and fracture toughness are required. Major efforts have been conducted with SiC whiskers in Al2O3, Si3N4, and MoSi2. A comparison of the properties of some whisker materials and potential matrix materials is presented in Table 10. At the present time over one-hundred different compounds have been grown in whisker form. A variety of methods have been developed for whisker growth (Levitt (1970)). Sublimation or evaporation of a source material, mass transport through the vapor phase, and condensation at the growth site cause whisker growth under low supersaturations. The growth rate can be controlled by the proper adjustment of temperature gradients. The application of this method was primarily to the growth of metallic whiskers. Other techniques which include a formation of whiskers from gaseous phase include the following: 1. Reaction between the source material and an ambient gas phase to produce a volatile species, vapor transport to the growth site, and condensation result in whisker growth under low supersaturations (Webb and Forgeng (1957)) 2. Reaction between gaseous species and subsequent condensation at growth sites causes deposition or crystal growth under suitable conditions and supersaturations (Price (1961)). Several basic requirements must be satisfied to utilize a vapor-phase reaction for whisker growth. These include thermodynamic calculations, kinetics of the reaction involved, and control of supersaturation conditions. For whisker growth, the control of supersaturation is a prime consideration because there is a good evidence that the degree of supersaturation determines the whisker morphology. Whiskers can also be grown by a different so-called vapor-liquid-solid (VLS) mechanism. This process occurs in two steps. First, there is deposition from the vapor directly on a liquid solution in a vapor-liquid

Figure 10. Idealized growth of SiC whiskers by the VLS mechanism. Table 11. Whisker-Reinforced Ceramics (Richerson (1992))

material hot-pressed Al2O3 hot-pressed Al2O3-15 vol % SiCwhiskers hot-pressed Al2O3-30 vol % SiCwhiskers hot-pressed Si3N4 hot-pressed Si3N4-30 vol % SiCwhiskers reaction-bonded/hot-pressed Si3N4-5 wt % SiCwhiskers hot-pressed MoSi2 hot-pressed MoSi2-20 vol % SiCwhiskers

flexural strength (MPa)

fracture toughness (MPa‚m1/2)

480 570

3.8 5.8

660

6.9

780 970

4.7 6.4

871

10.4

150 310

5.3 8.2

system. The second step occurs in a liquid-solid system and consists of precipitation from the supersaturated liquid solution at the liquid-solid interface. For the VLS mechanism to operate, the liquid solution must be stable. An idealized representation of SiC whisker growth is shown in Figure 10. There exists another technique of silicon carbide whisker formation. Rice hulls are a unique waste product of agriculture. Rice hulls contain silica (1520 wt %) and cellulose which yield carbon when thermally decomposed. With the very high surface area and intimate contact available from reactants in the rice hulls, it is possible to form silicon carbide readily and economically. Lee and Cutler (1975) were the first to explore this technique. They found that addition of fine iron particulates is required to obtain a preferential growth of whiskers. This indicates that the VLS mechanism is dominant in this process. Silicon carbide whiskers have been the focus of much research during the past several years, being studied as a possible reinforcement for ceramic matrices such Al2O3 and Si3N4. Significant improvement in fracture toughness has been widely reported in the literature (Karasek et al. (1991)). Examples of whisker-reinforced ceramics are presented in Table 11. 4.3. Combustion Synthesis. Self-propagating hightemperature synthesis (SHS), also called combustion synthesis, is a novel and simple method for making certain advanced ceramic and intermetallic materials. This method has received considerable attention as an alternative to conventional furnace technology. The principle concept of this technique is that, once initiated, a highly exothermic reaction can become self- sustaining and will propagate through the reactant mixture in the

Ind. Eng. Chem. Res., Vol. 35, No. 2, 1996 359 Table 12. Materials Synthesized in the Self-Propagating High-Temperature Regime TiC, ZrC, TaC, Ta2C, Mo2C, SiC, VC, HfC, NbC, Nb2C borides TiB, TiB2, ZrB2, HfB2, NbB2, TaB2, CrB2, WB, W2B, LaB6 nitrides TiN, ZrN, HfN, VN, Nb2N, NbN, Ta2N, TaN, Si3N4, BN, AlN, UN silicides Ti5Si3, TiSi, TiSi2, ZrSi2, NbSi2, MoSi2, Mo3Si, Mo5Si3, WSi2 sulfides CdS, CeS, MnS, MoS2, WS2 phosphides Ni2P, Co2P, AlP, Mg3P2 intermetallics Ni3Al, NiAl, NiAl3, TiAl, TiAl3, NiTi, CoAl, ZrAl, ZrAl3, FeAl composites Si3N4-SiC, TiC-Al2O3, TiB2-Al2O3, TiC-TiB2, MoSi2-Al2O3 solid solutions TiC-TiN, WC-TiC, R- and β-sialons

carbides

form of a combustion wave. For this to happen a chemical reaction must have a relatively high activation energy and must also generate a sufficient amount of heat. The synthesis of refractory high-temperature materials (e.g., borides, carbides, nitrides, and silicides) from elemental constituents exhibits both of the characteristics mentioned above. The historical development of combustion technology is well documented by Hlavacek (1991). Combustion synthesis of inorganic materials is a technology which has been in existence for more than a century. However, a systematic study of advanced ceramic material synthesis was initiated by Merzhanov and his associates in 1967. Since then, many theoretical and experimental papers have been published on this subject. The details of this development can be found in review papers by Merzhanov (1981, 1990), Frankhauser et al. (1985), Munir and Anselmi-Tamburini (1989), McCauley (1990), Luss (1990), and Varma and Lebrat (1992). A wide variety of refractory materials, ranging from single compounds to multicomponent composites and solid solutions, have been synthesized by this method. The partial list of various compounds synthesized by SHS is shown in Table 12. In general, the combustion synthesis can be categorized as (a) gasless combustion and (b) gas-solid combustion. In the case of the gasless combustion, solid reactants are usually dry mixed in a blender and the resulting mixture is pelletized to increase an intimate contact between individual particles. The reactant mixture is placed in a refractory container, degassed, and ignited in a vacuum or an inert atmosphere. Purities of inert gas and solid reactants, completeness of the degasification process, and an inertness of the container’s material affect the oxygen content in the final product. In some cases the purity of the resulting product might be even higher than that of the starting reactants due to evaporation of volatile species and suboxides in the combustion zone. The gasless combustion process is schematically shown in Figure 11. The resulting combustion product might be in the form of a powder or a dense compact. In order to accomplish a simultaneous combustion synthesis and densification, the use of uniaxial or isostatic force is required. Various aspects of in-situ densification are described by McCauley (1990), Merzhanov and Yukhvid (1990), and Odawara (1990). One of the most promising features of simultaneous combustion synthesis and densification is a formation of dense, refractory coatings inside metallic pipes. A mixture of iron oxide, aluminum, and aluminum oxide as a diluent is wet coated inside a metal pipe when it rotates. Next, this mixture is dried and subsequently ignited by an external flame. High tem-

Figure 11. Gasless combustion synthesis.

Figure 12. Centrifugal densification of in-situ synthesized materials.

peratures (>2500 K) generated during aluminothermal reduction cause a melting of both products (Fe and Al2O3). Due to a significant difference in the densities of both product and applied centrifugal force, a complete separation of phases takes place. A schematic representation of this process is shown in Figure 12. A direct reaction between a solid reactant (aluminum, silicon, boron, transition metals) and gaseous nitrogen can also be self-sustaining; however, in many cases high nitrogen pressures (Pnitrogen > 5 MPa) are required to complete the nitridation process. Even under very high gas pressures (hundreds of atmospheres), the amount of nitrogen gas present in the voids of the porous metallic specimen is insufficient for a complete conversion. However, as the gaseous reactant gets depleted at the reaction front, a pressure gradient results and nitrogen is supplied to the reaction front from the surrounding atmosphere by “filtration” through the porous sample. As a result, high degrees of conversion can be achieved, without application of ultrahigh nitrogen pressures. The term “filtration combustion” is thus prevailing for the description of gas-solid combustion reactions. The gas-solid reactions can be carried out using various reactor configurations shown in Figure 13. In the first two configurations depicted in this figure the metal powder is contained in a refractory crucible and the nitrogen flow to the reaction zone is cocurrent (Figure 13a) or countercurrent (Figure 13b) to the direction of the propagating combustion front. In the cocurrent configuration, the gas has to be filtered through the hot product layer to reach the unreacted

360


Figure 14. Typical structure of a combustion wave. Table 13. Modeling Equations for Gas-Solid Combustion Systems (Dandekar et al. (1990)) mass balance on the gas phase:

∂(Fgνx) ∂(Fgνy) ∂F )- µWg ∂t ∂x ∂y energy balance on the system:

∂(Fgcg + (1 - )Fgcg)T ) ∂T Figure 13. Reactor configurations for gas-solid combustion reactions.

∂(FgνxcgT) ∂(FgνycgT) ∂ ∂T ∂ ∂T + (-∆H)W λ + λ ∂x ∂x ∂y ∂y ∂x ∂y

( )

( )

equation of state:

metal. For systems that exhibit melting or sintering of the product, permeability is significantly reduced and an incomplete conversion or extinction of the combustion wave may occur. The countercurrent configuration is more suitable in such cases. Combustion of pressed metal pellets in nitrogen is governed by a combination of co- and countercurrent filtration (Figure 13c). However, at low nitrogen pressures, the nitridation reaction usually takes place only in a very thin layer on the outer surface of the pellet due to the relatively low initial porosity of the metal specimen. Complete nitridation inside the pellet can be accomplished by the application of very high nitrogen pressures (Hirao et al. (1987)) or by use of a solid nitrogen source (e.g., NaN3) (Holt and Kingman (1984)). The configuration depicted in Figure 13d is essentially equivalent to that shown in Figure 13b. However, in this case, due to the significantly shorter filtration path, less sintered and more uniform product in a composition can be synthesized. This type of arrangement is called a cross-flow configuration. The final product composition and its morphology in the case of both gasless and gas-solid reactions depend on (a) reactant particle size distribution, their morphology, and purity, (b) the initial density of the reacting mixture, (c) the initial temperature and pressure of the reacting system, (d) dilution of the reacting mixture with the final product, and (e) the size of the specimen and reactor configuration. The combustion reaction is initiated locally by supplying a short intense burst of energy (e.g., electrical spark, laser beam, resistively heated wire, chemical ignition source, etc.). A combustion wave then propagates through the unreacted material. A typical structure of the combustion wave is shown in Figure 14. However, this structure might be more complicated when melting of a product and/or formation of new phases take place. Combustion characteristics include the rapid propagation of a combustion wave (1.0 × 10-4-1.5 × 10-1 m/s), generation of high temperatures (1500-3500 °C), and the rapid heating of the product (103-106 K/s). Due to relatively fast cooling rates, a

P)

FgRT Mg

Darcy’s law:

νx ) -k

∂p ∂p ; νy ) -k ∂x ∂y

reaction rate expression:

∂FR ) -Wg ) -FR°k0f(η)e-E/RTpν ∂t

solid product consisting of nonequilibrium phases might be formed (Borovinskaya (1974)). A direct reaction between aluminum, silicon, boron, and transition metals with nitrogen at elevated N2 pressures is more complicated than a gasless synthesis, and therefore a detailed understanding of concentration and temperature fields inside a porous solid structure is needed to predict reaction conditions which would guarantee a complete conversion. There are certain limitations of the counterflow and concurrent configurations which are depicted in Figure 13a,b. These configurations are rather unsuitable for industrial application. A novel configuration shown in Figure 13d allows better gas accessibility inside a porous structure. Gaseous reactant is transported from surroundings not only by longitudinal flow from above but also by transverse (countercurrent) flow from regions ahead of the front. As a result, the nitridation reaction can be completed at lower pressures without excessive sintering of the product. Furthermore, the configuration provides easy product removal. For nitridation on the industrial scale, nitrogen pressures should always be higher than 2-5 atm to guarantee a high degree of nitridation of large batches. Dandekar et al. (1990a) performed detailed numerical studies of combustion synthesis in a cross-flow configuration. The system of equations describing such reacting systems is shown in Table 13. These equations were solved using a sophisticated numerical program with multidimensional adaptive grids (Degreve et al. (1987)). As can be seen from Figure 15, temperature,


Figure 15. Three-dimensional perspective view of a gas-solid combustion reaction (taken from Dandekar et al. (1990)).

conversion, gas density, and pressure profiles are quite complex, and they significantly change when the combustion front progresses. A more detailed description of profile developments and propagation characteristics might be found elsewhere (Dandekar et al. (1990a,b); Matkowsky and Shivasinsky (1978); Shkadinsky et al. (1992)). One of the most important features of SHS technology is a possibility of in-situ formation of composites and solid solutions. Lis et al. (1991) have demonstrated insitu synthesis of various β- and R/β-sialons by a combustion technique. They found that a full range of different β-sialons can be synthesized through combustion nitridation. The properties of sintered parts derived from combustion-synthesized sialon powders are comparable to others reported in the literature. 4.3.1. Selected Applications of the SHS Technology. Manufacture of Refractory High-Temperature Ceramics. High-temperature ceramic materials can be manufactured by a combustion technique. Typical reactions belonging to this category are as follows: (a) Mo + Si f MoSi2: The reaction occurs without any major gas expulsion, and propagation velocity is constant in a wide range of initial concentrations. Molybdenum disilicide is an important material for high-temperature heating elements with significant resistance to oxidation. (b) Ti + 2B f TiB2: The reaction occurs violently with gas generation (vaporization of impurities and reactants). The ignition characteristics depend on the particle size of metallic titanium. Titanium diboride has major applications as wear- and corrosion-resistant material. (c) Ti + C f TiC: The reaction occurs with moderate gas release. Titanium carbide is an important material in the cutting tool industry. (d) Si + C f SiC: The reaction is strongly exothermic but not self-sustaining without a significant preheating of the reacting mixture. Silicon carbide has major application as a wear-resistant refractory material. (e) 3Si + 2N2 f Si3N4: The self-propagating combustion of silicon in nitrogen does not occur at atmospheric pressure, and approximately 40 atm of nitrogen is

required to make a nitridation process self-sustaining (Munir and Holt (1987)). Silicon nitride is one of the most important high-performance ceramics and has important application in a new generation of heat engines. Isolation of High-Level Nuclear Waste by Combustion Synthesis. A potential application of SHS technology is to isolate highly active nuclear waste in an inert matrix. The problem of immobilizing radioactive waste by melting a radioactive product with borosilicate glass is not the final solution. The SHS technology is a simple approach for immobilization. Transition-metal borides, carbides, and nitrides have some unique properties that are more responsive than oxides to the high-level waste immobilization requirements. With selection of specific compositions of transitionmetal borides, carbides, or nitrides, thermal stability in essentially all media (air, acid, water, molten metals) can be guaranteed up to very high temperatures. With the SHS technology, the immobilization can be performed without a melting unit. The energy requirements are also lower. Fabrication of Advanced Nuclear Fuel. The SHS technology can be successfully used for fabrication of advanced nuclear fuel for the space reactor. The SHS technology provides unique capabilities: advanced nuclear fuel, based on uranium nitride, can be readily synthesized without an application of high-temperature furnaces. Direct nitridation of uranium metal powder by nitrogen is straightforward, and any major chemical contaminations are eliminated. Composition adjustments, i.e., additions of transition-metal nitrides, carbides, and/or carbonitrides, can be accomplished without a significant process modification. Metallic Hydrides as Shield Materials. There are significant needs of high-efficiency radiation shields, where reduced weight and volume are major design constraints. For example, in space applications, reactor shields for neutron radiation must be more efficient per unit volume by several orders of magnitude than shields normally placed around nuclear reactors at central power stations. Emphasis in the search for such shield materials has been on transition-metal hydrides (e.g., ZrH2, TiH2, LiH). These materials can be synthesized within seconds using the SHS technology. Cutting Tool Materials. Nickel-based alloys and titanium alloys are the materials featuring rapid tool wear. Therefore, the application of cutting tools in highspeed machining has its limitation. The development of Al2O3-TiC cutting tools enabled an increase in the cutting speed from 100 to 600 ft/min for nickel-based superalloys. The synthesis of alumina-based composites by SHS technology is one of the main trends in this field. Thick Coatings Technology. Presently U.S. technology for making ceramic coatings is not energy efficient. A typical process requires use of deposition techniques, e.g., plasma spraying, which require a very high energy input per square foot covered. A new process for making continuous thick ceramic coatings using SHS technology has been developed and successfully used for the past several years. 4.3.2. Scaleup of Combustion Synthesis Operations. The design of reactors for SHS systems is a very complex reactor design problem. As such, combustion characteristics must be carefully controlled if highquality ceramic powders are to be synthesized. The scaleup of the SHS operation is still in its infancy. There are a wide variety of SHS reactions with proper-

362


ties of the phases changing dramatically during the reaction process. Certain solid-solid reactions are accompanied by an explosive expulsion of gases during combustion and by melting of intermediate or final products. In some cases, the large heat released during the reaction might cause partial or complete vaporization of precursors (e.g., during the synthesis of sulfides, sulfur might completely vaporize at low ambient pressures; a similar situation takes place during magnesium thermal reduction of oxides). For gas-solid SHS reactions, a similar variety of situations may exist. In addition, complex flow patterns of a gaseous reactant inside a porous structure might affect the overall conversion in the system. Because of the rich spectrum of physical properties and associated phenomena, there is no trivial answer to the scale-up problems. Thus, each particular situation requires the design of a specific reactor.

Figure 16. Reaction bonding of silicon nitride (taken from Moulson (1979)).

5. Reaction Sintering and Net Shape Processes 5.1. Reaction Sintering: Gas-Solid. The densification of ceramic powder compacts by sintering or hot pressing involves usually linear shrinkage of 10-25%, causing such problems as warping and shape distortion, cracking, and the occurrence of harmful residual stresses. These problems are usually eliminated with melt oxidation derived materials (Lanxide type of processes) or with reaction-formed oxide and non-oxide ceramics, such as, reaction-bonded aluminum oxide (RBAO) or reaction-bonded silicon nitride (RBSN). Reaction forming of ceramics results in a number of advantages when compared with conventional fabrication. Among the most important are low raw material costs, avoidance of glassy phases produced by sintering aids, and netshape capabilities. Especially, low shrinkage of reaction sintering techniques makes this process suitable for the manufacture of complex shapes produced by either slip casting, injection molding, or machining of green bodies while maintaining very close dimensional control. 5.1.1. Reaction Bonding of Silicon Nitride (RBSN). Reaction-bonded silicon nitride can be produced by the reaction of Si green body with a nitrogencontaining atmosphere according to the reaction scheme (Moulson (1979)):

3Si(s) + 2N2(g) f Si3N4(s) A large amount of heat is produced during the reaction, and reaction control is very important in producing high-quality RBSN. The reaction rate of nitridation can be lowered by nitrogen dilution, e.g., by hydrogen. The use of a N2/H2 nitriding atmosphere results in significant improvement of the creep resistance of RBSN. Since RBSN is carried out at temperatures below the melting point of silicon (1420 °C), the noncatalytic reaction is of the type solid + gas f solid. The schematics of the nitridation process are shown in Figure 16. The produced film of silicon nitride results in a decrease of the reaction rate. Following each temperature increase, a rapid nitridation occurs. Consequently, a multistep nitriding cycle was proposed. More details about the process can be found in the patent literature (Washburn (1978)). Mangels (1980, 1981) showed that the strength of RBSN can be 50% higher if nitridation is accomplished in a multistep temperature cycle. A good microstructure can be observed if the silicon melting point is not exceeded.

Figure 17. Schematic diagram of RBAO (taken from Claussen (1989)).

Because of the large amount of heat produced during the reaction, a huge overheating of the silicon green body can be observed. This overheating can amount to several hundred degrees Celsius, and with large bodies of silicon preform, a careful temperature control must be accomplished. The RBSN bodies can be produced to near net shape with good tolerance since little shrinkage occurs during the reaction. Typically, the density of RBSN is in the range 75-82% of theoretical. As a result, most mechanical properties of RBSN are inferior to sintered or hot-pressed silicon nitride. Supplying sintering aids to RBSN can increase the density to 95% of theoretical. Two techniques were recommended in the literature to introduce sintering aids into the silicon nitride body: (i) Impregnation of the RBSN body with an alcoholsalt solution, with the salt eventually decomposing into the oxide. Repeated impregnations are required to reach the desired concentration. As an example, 2-4% MgO can be achieved by multiple impregnation of RBSN by a methanol solution of MgCl2, followed by calcination. (ii) Introduction of sintering aids to the Si powder prior to nitridation. For example, Y2O3 can be added to the silicon powder. 5.1.2. Reaction Bonding of Aluminum Oxide (RBAO). The RBAO process follows principally the same strategy as RBSN, namely, a mixture of Al powder (usually 30-60% by volume), and Al2O3 is oxidized in an air atmosphere to produce sintered Al2O3 (Claussen (1989), Wu (1993)). The schematic diagram of the RBAO process is presented in Figure 17. The mixture of Al + Al2O3 is vigorously mixed in an attrition mill. Green compact can be easily produced by conventional compaction units. Al particles are plastically deformed during the compaction process, the green body density is high, and green body strength reaches values more than an order of magnitude higher (20-50 MPa) than those of conventional ceramic green bodies. Evidently, green body machining can be used to produce complex shapes.


After heating the compact in an oxidizing atmosphere at temperatures above 350 °C, Al is converted to Al2O3, providing a net volume expansion of 28%. During the sintering stage (T > 1200 °C) a shrinkage occurs which compensates the expansion. Nevertheless, in the Al + Al2O3 systems, 15% shrinkage may remain when the final density exceeds 95% of theoretical. As a result, the RBAO process can be modified by incorporating other metals that exhibit larger volume expansion during the oxidation step. For example, volume expansion is 49% for Zr, 76% for Ti, 102% for Cr, and 174% for Nb. At temperatures below 450 °C, the Al2O3 film covering individual Al particles grows by Al3+ ions diffusing outward through the Al2O3 skin and reacting with oxygen. At temperatures T > 450 °C, Al is oxidized directly to γ-Al2O3. The metal oxidation takes place because of oxygen diffusion through the grain boundary and the occurrence of microcracks in the film. Above the melting temperature (660 °C) oxidation is accelerated because of volumetric expansion of liquid Al which penetrates through pores and microcracks. This process continues until all Al is converted. Comparison of the bending strength of conventional Al2O3 and RBAO has shown that RBAO ceramics feature higher strength than conventional Al2O3 and ZrO2-toughened Al2O3. Porous Al2O3 has many important technical applications, among them are catalytic convertors, filters, membranes, etc. Evidently, RBAO represents a promising fabrication route for highstrength porous ceramics. 5.1.3. Reaction Sintering of Silicon Carbide. Silicon carbide-silicon nitride materials are important for their high strength at high temperatures, high fracture toughness, excellent thermal shock and oxidation resistance, high thermal conductivity, and superior resistance to metal penetration. The reaction-bonded silicon carbide bricks are widely used in aluminum melting and holding furnaces and in aluminum reduction cells (Maltseva (1966)). The reaction sintering of the SiC body is based on strong bonding of SiC by Si3N4 which is prepared in situ by direct nitridation of Si. Typically, SiC is mixed with 10-15% of the Si powder, pressed into a green body, and sintered in the nitrogen stream at 1350-1450 °C (Reddy (1991)). 5.1.4. Reaction-Bonded TiC, TiN, and Al2O3. Various ceramic powders, if mixed with fine silicon powder, will sinter in nitrogen atmosphere at 13501450 °C to corresponding composite materials (Yasutomi et al. (1991a)). The porosity of the TiC-Si3N4 composite is very low because of partial conversion of TiC to TiCNx and β-SiC. Carbon, in particular, seems to have a remarkable effect on reducing the number of pores in the sintered body. For reaction sintering of the Al2O3 material, for samples containing 20-70% of Al2O3, only a very low shrinkage of about 0.3% was observed. An electroconductive TiN-Si3N4 composite can also be produced by reaction sintering. Low shrinkage of ∼0.3% was observed. 5.2. Reaction Sintering: Solid-Solid. Usually under the term reaction sintering the ceramic community understands reaction sintering of the type solid-gas in general and bonding by silicon nitride or aluminum oxide in particular. Nevertheless, there is a broad class of reactions of the type solid-solid which can also be used for reaction

Figure 18. Schematics of the formation of a ceramic composite by the outward growth of ceramic/metal reaction produced matrix through a mass of filler material. Table 14. Examples of Reaction Bonding (Solid-Solid Systems) Ti + 2B + Cu f TiB2 + Cu Ti + Zr + 4B + Al f TiB2 + ZrB2 + Al Ti + C + Ni f TiC + Ni Mo +2Si +Al2O3 f MoSi2 + Al2O3 Ti + 2B +Mg f TiB2 + Mg Al + Ti f TiAl Mo + 2Ge + Al2O3 f MoGe2 + Al2O3

bonding. Generally speaking, these reactions are of exothermic nature. For strongly exothermic reactions (as, e.g., Ti + 2B f TiB2) the reaction is self-propagating; other reactions (as, e.g., 3TiH2 + AlB6 f 3TiB2 + Al) are only weakly exothermic, and an auxiliary source of energy (furnace) must be used to keep them going. The process was apparently first described by Krapf (1964) and then rediscovered several times (Williams (1967); Maksimov et al. (1984); De Angelis (1986); Brupbacher (1987)). For the self-propagating type of reactions, the combustion process is accompanied by violent desorption of adsorbed gases and, as a result, the reaction-bonded body is highly porous. On the other hand, for weakly exothermic reactions occurring in a die (sometimes under pressure in a hot press die), the reaction-bonded bodies can be fully densified. A few important reactions belonging to this class are reported in Table 14. 5.3. Reaction Sintering: Liquid-Gas. This type of reaction bonding is known in ceramic literature as the Lanxide process (Newkirk (1986); Creber (1988)). The Lanxide process involves formation of ceramicceramic or ceramic-metal composites by catalytic reaction of “molten metal-reacting gas”. Typically, the molten metal is aluminum (or its alloys); the reacting gas is oxygen or nitrogen. The Lanxide process is displayed in Figure 18. The process is carried out in a crucible, by direct reaction of the molten metal with gas. Reinforcing materials (e.g., fibers) are incorporated in the matrix by placing them in a layer adjacent to the metal, in the path of the penetrating gas once started under appropriate conditions; the process generates a composite of uniform microstructure as long as (i) molten metal and gas are available to sustain the process, and (ii) the process temperature is maintained. The reaction is sustained by the wicking of the molten alloy through microscopic channels in the product layer. Composites up to 10-15 cm thick can be produced by this technology. Since growth occurs from the surface of the molten metal, one-piece ceramic bodies of virtually any size can be produced. The volume fraction of the metal and ceramic, the type of microstructure, the porosity, and the degree of interconnectivity of the metallic phase in the product can be

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controlled by varying the processing temperature, time of reaction, and the dopants used. Lanxide materials exhibit little or no ductility. They can be produced with a variety of densities, strength, hardness, fracture toughness, thermal conductivity, and thermal expansion coefficients. The high-temperature strength of Al2O3Al Lanxide material can be compared to that of sintered Al2O3. 6. Joining of Ceramics A major restriction on the wider use of ceramics in the industry is the lack of suitable techniques for joining ceramics to other materials, particularly metals. Ceramic-to-metal sealing started a few decades ago when development of electron tubes required insulating material superior to glass in electrical and thermomechanical properties. The traditionally used joining techniques include organic adhesives and cements. The methods have serious limitations especially in engineering applications including a hot, dirty, corrosive high-wear environment, where, for ceramics to be used to their full potential, joining techniques must be made available. Ceramic bonding involves creation of interfaces between dissimilar materials. Part of this interface can be ceramic, glassy, or metallic material. Evidently, there is a discontinuity in the material properties at the interface. The abrupt changes of properties at the interface may belong to one of the following categories (Loehman and Tomsia (1988)): (a) crystallographic discontinuitiesslattice mismatch between ceramic and metallic parts; (b) electronic property changesschange in an electronic structure and bonding from ionic or mixed ionic-covalent in ceramics or glasses to metallic in metals; (c) thermodynamic differencessthe materials close to the interface will be not in thermodynamic equilibrium, and they react to generate interfacial products; (d) mechanical property gradientssdifferences in elastic modulus across the interface; (e) thermomechanical changessat the interface there may be a large difference of coefficients of thermal expansion. A good bond between two dissimilar materials depends critically on how these differences are reconciled. For example, an elastic modulus match is important if the joint is going to be exposed to cyclic mechanical stresses. Strong joining of two materials requires good interfacial bonds and a match of thermal expansion coefficients. If the interface is at chemical equilibrium, long-term stability at high temperatures can be expected. Various bonding methods can be classified as follows (Klomp (1971)): (i) sintered metal powder processes, (ii) active metal brazing process, (iii) ceramic frit process, and (iv) solid-state processes. Generally speaking, the bonding of ceramics is not a trivial process, and the “art of bonding” requires a lot of “know-how”. 6.1. Sintered Metal Powder Process. This process is based on the discovery by Pulfrich and Vatter (1936) on sintering a mixture of tungsten and iron powder on a MgO-SiO2 ceramic surface. Later this process was modified, a mixture of molybdenum and manganese powder is sintered on ceramic surfaces. The metallic powder is suspended in a nitrocellulose binder and mixed in a ball mill. The slurry is then applied to ceramic surfaces by dipping, rolling, or brushing. After drying, the coated ceramic layer is fired in a wet hydrogen at temperatures 1300-1550 °C. The ceramic body is then ready to be brazed to a metallic body if

brazing filler materials are used which wet the molybdenum layer, e.g., Ag-Cu. Addition of a reactive component is necessary if ceramics containing more than 97% Al2O3 have to be metalized since the molybdenum-manganese process fails to metalize ceramic surfaces with a high concentration of alumina. For metalization of pure Al2O3 ceramic bodies, a metal-metal oxide mixture was developed, e.g., Mo-Al2O3/CaO/MgO/SiO2. The reactions that may occur between alumina and the coating containing molybdenum with oxide additions are based on the fact that alumina dissolves in the oxide additives and a strong bond of the oxides to the alumina and of the oxides to the molybdenum is formed. 6.2. Active Metal Brazing Process. This method involves application of certain transition metals with high melting point (titanium, zirconium, tantalum, and niobium) to a ceramic surface. Brazing to a metallic part occurs with a brazing filler in the argon atmosphere at temperatures ∼1000 °C. During the bonding operation the active metal dissolves in the brazing alloy and active metal can react in the liquid phase with the ceramic interface. This technique has only limited use since a very careful control of the brazing temperature and amount of the active metal is necessary. An excessive amount of the active metal or the brazing alloy results in spreading of the braze over areas of the ceramic that are not supposed to be metalized. The seals, unfortunately, do not withstand a hightemperature environment since the reaction of active metal with the ceramic will continue. 6.3. Ceramic Frit Process. This method represents a one-step technique which can bond ceramic parts or ceramic-to-metal parts. Two important oxide mixtures can be used: (i) Al2O3-CaO-MgO-SiO2 and (ii) Al2O3-MnO-SiO2. The mixture of type i can be used for bonding refractory metals to ceramics. Depending on the composition of the bonding mixture, the seal can withstand temperatures ∼1200 °C. The sealing mixture ii can be used for bonding iron alloys to ceramics and can withstand 800 °C or higher. Important properties of these seals are strength, vacuum tightness, and good temperature resistance. 6.4. Solid-State Reaction Processes. The abovementioned sealing processes take advantage of the liquid phase at a certain stage of the sealing process. There are, however, processes where the bonding occurs in the solid phase. At the reacting interface a certain type of cement is created via the interface reaction. A few examples are mentioned below. (i) Joining of Oxides to Metals through Intermetallic Compounds. The technique of reaction hot pressing of intermetallic components can be used to produce strong ceramic-to-metal seals. When a powder mixture of elemental components, which react to refractory aluminides or silicides, is heated rapidly, an exothermic reaction is initiated between them at relatively low temperatures, usually about 600-800 °C for aluminum reactions and 1000-1200 °C for silicon reactions (Faulkner and Williams (1964)). The reaction bonding process consists of formation of a thin compact between the surfaces of the components to be joined, the whole assembly being contained in a hot press die. Bonding between the metal surface and the sealing layer was found to take place by rapid intermetallic diffusion during the overall heating cycle in the hot


press die. Bonding between a ceramic surface and the sealing layer takes place by very short-range interaction. (ii) Joining Two Ceramic Bodies by the Ti-CNi System. The method is described for joining two SiC elements. The principle is similar to that described above for intermetallic compounds. An exothermic reaction between Ti + C in Ni environment ignites below 1200 °C and forms a TiC-Ni bonding material. The bonding process occurs in a hot press die which is heated electrically. The joining surfaces, coated by the Ti + C + Ni mixture, are mated together. Pressure ranging from 5 to 20 MPa is applied in the die and maintained during heating beyond the temperature required to initiate the exothermic reaction. 6.5. Application of Ceramic Joining. Nb-Al2O3 System. Transparent Al2O3 is used as the arc tube material in high-pressure sodium lamps because it is corrosion resistant to sodium vapors at temperatures above 600 °C. The end cap and the electrode structure are fabricated from Nb because of very close thermal expansion coefficients. This combination is important for minimization of mechanical stress produced during thermal cycling. The bond is produced by a mixture of CaO-MgO-Al2O3 frit which may contain a small amount of BaO, SiO2, B2O3, and V2O3. Bonding of Non-Oxide Ceramics. Many applications of high-performance non-oxide ceramics require the development of suitable sealing of joining techniques. A silicon nitride heat engine part must be connected at some point to metal, for example, a ceramic turbocharger rotor to a metal shaft. Hermetic packaging of electronic devices in an AlN substrate requires methods for metal-AlN sealing. Non-oxide ceramics can be joined using reactive metals or braze alloys that contain a reactive metal. Commercially available braze Cu-Ag-Ti exhibits very high values of flexure strength when used to braze silicon nitride pieces. The same alloy can also be effectively used for welding of SiC bodies. The bonding of two non-oxide ceramic bodies can also be accomplished by a polymeric precursor; Yajima et al. (1981) welded two pieces of SiC by polyborosiloxane. Bonding of Non-Oxide Ceramic and Metal. Welding of ceramic and metal is strongly affected by the values of the thermal expansion coefficient. In the case of bonding larger ceramic parts, an interlayer material with a thermal expansion coefficient that matches that of the ceramic part must be used. Welding of SiC and stainless steel can be accomplished by using a Ag-CuTi braze. The application of ceramic-to-metal welding plays an important role in the design of high-performance turbochargers, gas turbines, and diesel engines. 7. Chemical Vapor Deposition Methods 7.1. Application and Modeling. In the last years chemical vapor deposition (CVD) has been gaining in popularity, becoming a widely used process in modern technology. The current application of CVD includes the coating of particles and extended surfaces, vaporformed free-standing shapes, high-purity materials, vapor-impregnated parts, filaments, fibers, whiskers, crucibles, optical windows, and tubes among others. The scientific community has been impressed at large by the unique structures attainable from the deposition of solid products via homogeneous and heterogeneous gas reactions promoted by heated solid surfaces. A large number of chemical reactions, including the thermal

Figure 19. Physicochemical processes in a CVD system.

decomposition or reduction of fluorides, chlorides, bromides, hydrides, organometallics, carbonyls, and hydrocarbons, are available for the deposition of solid films or coatings. Despite widespread and routine application, the level of understanding on the interaction between the fluid dynamics of the process and the mass and energy transfer phenomena was rather limited until recent scientific contributions were published. The design and operation of chemical reactions has received much attention in the past decade and is rapidly developing from an empirical art toward synthetic and rational activity. A CVD system is simply a chemical reactor. As such, flow rates and flow patterns of reactant vapors along with substrate temperature must be carefully controlled if uniform film layers are to be obtained. In turn, prediction of the deposition rates and uniformity requires detailed understanding of thermodynamics, kinetics, fluid flow, and mass transport phenomena for the appropriate reactions and reactor designs. An exact analysis of the phenomena occurring in CVD systems would certainly have to include a complex description of the different processes involved: from a mechanistic description of the reaction scheme to accurate equations of state to describe the transport properties of the reacting flow. All these components blended with the governing equations for the mass, energy, and momentum transport phenomena would constitute a “complete” model of the reaction system (cf. Figure 19). This approach has been, with justice, usually categorized as a formidable task and even in the advent of the supercomputers is still seen as a real challenge. Broadly defined, chemical vapor deposition, or CVD, is the formation of solid products via chemical reactions of gaseous precursors. A typical CVD process would involve a dynamic flow system in which gaseous reactants pass over a heated substrate. The gases react chemically to produce a condensed coating on the substrate plus product gases, which, in turn, together with any remaining reactant gases, exit from the hot reaction zone. The application of CVD processes centers around the production of thin films for semiconductors and other solid-state electronic devices (Hess et al. (1985); Dupuis (1984); Green and Levy (1985)), coating of cutting tools and surfaces needing erosion and/or corrosion protection (Yee (1978)), coating of fibers used in forming composite materials (Di Carlo (1985); Guinn and Middleman

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(1989)), and containment coatings for nuclear fuel and nuclear waste particles (Spear (1982)). CVD processes are also used in the manufacturing of objects with complex shapes (e.g., refractory crucibles) out of materials such as tungsten, molybdenum, and rhenium which resist conventional machining and fabrication techniques (Iwasa et al. (1987); Spear (1982); Shinko and Lennartz (1987)). A vast number of chemical compounds are considered for these applications (Blocher et al. (1984); Campbell et al. (1949); Stinton et al. (1988)). Important materials being prepared by CVD for electronic and optic devices include Si, Ge, Si3N4, SiO2, GaAs, CdS, ZnSe, and related compounds (Spear (1982); Green and Levy (1985); Gentilman et al. (1981); Hess et al. (1985)). Superconducting materials such as Nb3Sn, Nb3Ge, and NbCxNy (Wahl and Schmanderer (1989)) and YBa2Cu3O7-δ (Kim et al. (1991)). Erosion- and corrosionresistant coatings commonly applied by CVD to tooling materials are TiC, TiN, and Al2O3. In addition, an appreciable amount of CVD research is being performed on coatings of TaC, TaN, W2C, TiB2, SiC, Si3N4, B4C, BN, and B for use on cutting tools, tubes, crucibles, fibers, and applications in which the materials are subjected to high-temperature, corrosive and erosive environments (Spear (1982); Iwasa et al. (1987); Powell et al. (1966); Pierson (1981); Wachtman and Haber (1986); McCandless and Withers (1970)). In the manufacturing of the mentioned films and coatings the performance of the deposition process is governed by the underlying thermodynamic, kinetic, and transport processes. In the design and control of a CVD system, these processes have to be addressed and analyzed. 7.2. CVD and Chemical Equilibrium. In CVD, a nonequilibrium inlet state approaches an equilibrium state by chemical reaction and mass transfer. The chemical reaction and phase equilibria relationships provide a description of the reactor performance attainable. Since operating temperatures are typically far above room temperature and molar flow rates are generally low, near-equilibrium conditions are often approached in various segments of the reactor (Rosner et al. (1987)). Calculation of compositions in phase and reaction equilibrium is useful for several reasons. First, computation of chemical equilibrium in complex systems is simple relative to the determination of the actual (experiment) and thus permits the feasibility of a proposed process to be evaluated efficiently (Randich and Gerlach (1983); Schlichting (1980); Carlton et al. (1970); Walsh (1973)). In addition, knowledge of the equilibrium state can bound the values of operating parameters necessary for successful film growth and provide information on the response of the process to changes in operating conditions. Finally, the computation of equilibrium compositions with intentionallylimited reactants can assist in the assessment of reaction mechanisms. The attractiveness of an equilibrium model for describing CVD reactor operation is the simplicity of the calculation. The results will, however, be semiquantitative at best. 7.3. CVD Kinetics. CVD chemistry is complex and usually involves both gas phase and surface reactions. The role of gas phase reaction expands with increasing temperature and partial pressure of the reactants and might even lead to particle formation (van den Brekel and Bollen (1981)). CVD kinetic data have traditionally been reported in terms of growth rates and their

dependence on temperature. Such data are, however, only marginally useful in reactor design. The Langmuir-Hinshelwood or Langmuir-Rideal rate expression for CVD reactions can, however, be formulated for CVD systems in a manner similar to the way it is done for catalytic reactors. The analysis is, however, complicated by gas phase reactions that form precursors for surface reactions. Nevertheless, several studies propose Langmuir-Hinshelwood rate expressions for CVD reactions. Unfortunately, kinetic parameters are rarely estimated (Farrow (1974); Carlton et al. (1970); Gruber (1970); Michaelidis and Pollard (1984); Roegnick and Jensen (1987)). 7.4. CVD Reactors. Two important CVD systems have been designed in the past: horizontal and vertical reactors. In vertical reactors the main flow direction is parallel to the gravitational field, in contrast to horizontal reactors in which the forced flow of reactants is normal to the gravitational field. Both horizontal and vertical arrangements can be utilized as cold- or hotwall reactors. Typical cold-wall reactors include the horizontal, barrel, and pancake reactors where the reactor walls are cooled while the susceptors are heated by several means. In fiber applications the substrate is often heated resistively (e.g., Krukonis (1977)). rf induction coils and quartz radiant heaters are often used in horizontal, barrel, and pancake reactors (e.g., Hess et al. (1985)). Less traditional methods include heating by microwave, laser, and plasma techniques (Blocher et al. (1984); Randich and Gerlach (1983); Allen (1981); Stinton et al. (1988); Hess et al. (1985)). Although coldwall reactors minimize deposition on the walls, the large temperature gradients induce secondary flow phenomena with associated difficulties in obtaining uniform thickness and composition. In horizontal and barrel reactors, susceptors are often tilted relative to the main flow direction to improve film thickness and composition uniformity. Film uniformity is further increased in barrel and pancake reactors by a rotational movement of the substrate. Large substrate packing densities can be realized in hot-wall low-pressure (LPCVD) reactors. Such reactors are used for the production of multiple Si, SiO2, and Si3N4 wafers (Hitchman (1979); Meyerson et al. (1986)). As these reactors are operated at low pressure (about 100 Pa), diffusion coefficients are nearly 3 orders of magnitude larger than at atmospheric conditions. This makes it easier to work in a kinetically controlled regime which ensures a uniform deposition rate. In spite of the low pressures, rates of film growth in LPCVD systems are also only slightly less than those possible at atmospheric pressure processes as smaller dilution ratios can be used for low-pressure operation. 7.5. Fluid Flow Phenomena. Fluid flow patterns can have a pronounced influence on the performance of CVD reactors. Flow disturbances resulting from entrance effects and natural convection can result in nonuniform deposition rates. Experimentally complex flow structures, including back flow the longitudinal and transversal flows, have been visualized by TiO2 smoke tests (Takahashi et al. (1972); Wang et al. (1986)) and inferred from laser interference holography observations of perturbations (Giling (1982)). Flow visualization is, however, limited by the thermophoresis effect on seed particles which are rejected from the hot susceptor (Talbot et al. (1980); Rosner and Kim (1984)). The holography technique, on the other hand, gives spatially averaged density gradients (Williams and Peterson


(1986)). This might, however, not reflect the actual flow of species in very low concentrations when reactants have transport properties which differ substantially from the transport properties of the carrier gas (e.g., H2 vs SiCl4). Several theoretical investigations of the classical Rayleigh-Benard problem relating flow of a Boussinesq fluid in enclosures with conducting, insulating, and/or heated walls have been performed (e.g., Chiu and Rosenberger (1987); Gatica et al. (1989)). Due to the large temperature gradients present in cold-wall CVD reactors, none of these studies are, however, directly applicable to the problem at hand. Several studies, combining momentum heat and mass transfer, have been performed as well. In these studies the governing equations, conserving momentum, energy, and mass, are solved numerically over a domain that includes the entire reactor. This kind of analysis gives insight regarding the macroscopic behavior of the reactor and has been performed for several microelectronic applications (e.g., Evans and Greif (1987); Hess et al. (1985); Holstein et al. (1989); Kee et al. (1987); Moffat and Jensen (1986, 1988); Roegnick and Jensen (1987); Rosenberger (1987)). 7.6. Microsctructure and Surface Phenomena. As a global analysis of a CVD reactor, as described above, stretches over space scales much larger than the microstructural morphology of the deposited surface itself. It is not possible to capture all the transport phenomena that occur over the various space scales at the same time, especially in a numerical analysis in which the discretized elements are orders of magnitude larger than irregularities in the coating itself, due to the fact that this transport phenomenon occurring on a microlevel is lost. To optimize the mechanical properties of a material, it is, however, necessary to control the microstructural evolution of the product during the deposition process. Nonuniformities in the film surface seem to be a real life problem in CVD and have been reported repeatedly (see, for example, Campbell et al. (1949); Smith et al. (1987); Armas and Combescure (1987); Messier et al. (1984)). Nonuniformities can occur on three different length scales namely, the near atomic or nanoscale (3-30 nm), the microscale (0.1-1 µm), and the macrolevel (>10 µm). The characteristics of the surface are, however, often quite similar over the various levels (Messier et al. (1984)). Qualitatively it seems as if these nonuniformities can be blamed on two major factors. First, they can be caused by the interaction between the adatom effect and surface or bulk diffusion in the film. Second, mass and heat transfer limitations at the gas-solid interface can lead to a nonuniform growth of the film. Nonuniformities on the nano- and microlevels are most likely caused by the surface effects, while larger nonuniformities are more likely related to nonuniformities in the boundary layer. The morphology of a depositing film depends on the way deposited particles are incorporated into the existing structure. On a smooth surface, preferential growth of one area over another may result from varying surface mobilities of the depositing particles, usually as a result of grain orientation. Preferential growth leads to a dominant grain orientation and surface roughening with thickness. As the surface gets rough, geometrical shadowing can lead to preferential growth of the elevated region, giving a columnar morphology to the deposit. A preexisting surface roughness or preferential

Figure 20. Structure zone model of deposited material.

nucleation will also lead to shadowing and a columnar morphology. An elevated substrate temperature will effect the morphology by increasing surface mobility and allowing recrystallization to occur. When atoms or molecules impinge on a surface, they lose energy to the surface and finally condense by forming stable nuclei. During condensation, the particles have a degree of mobility on the surface which is determined by their kinetic energy and the strength and type of interaction between the particle and the surface. A strong surface particle interaction will give a high density of nuclei, while a weak interaction will result in a widely spaced nuclei. If the depositing particle/substrate interaction strength is low, the deposited particle will nucleate on surface discontinuities or nucleate by collision with adsorbed atoms or molecules or other particles migrating on the surface. These effects led to the structure zone model of deposited materials. The model consists of the formation of three zones which depend on the ratio of the surface temperature (Ts) to the melting point of the deposited material (Tm) (see Figure 20). Zone I results when particle diffusion is insufficient to overcome the effect of shadowing, giving a columnar structure with low-density boundaries between the columns. The individual columns are polycrystalline and usually highly defected with small grain size. The surface morphology is generally rounded. Zone I deposition usually occurs if Ts/Tm < 0.3 for metals and if Ts/ Tm < 0.24 for nonmetals. Zone II is defined in the range of Ts/Tm, where the particle surface diffusion dominates and the column structure consists of less defected and larger grains with higher density boundaries between the columns. The surface morphology of zone II material is generally more angular than that of the zone I material. Zone III is the Ts/Tm region where surface diffusion, bulk diffusion, and recrystallization dominate. The material is more equiaxed with high-density grain boundaries and large grains. For zone III deposition Ts/Tm will typically be higher than 0.45. Rough surfaces, high gas pressures, and reactive gases will tend to lower the Ts/Tm values for the zone boundaries. The microstructure is also influenced by the supply of reactants and the deposition rate. A higher concentration of reactants enhances the reaction rate in the gas phase. As a result, the gas may become supersaturated with product, leading to the formation of “snow” which results in a refinement of grain size (Blocher (1981); Stinton et al. (1980)). The effect of increased supersaturation can, therefore, be seen as opposite to that of increased temperature. At a higher deposition rate there is also less time for microvoids to be filled by surface diffusion. The described three zones are reported by several authors (Messier et al. (1984); Stinton et al. (1988); Mazor et al. (1988); Muller (1985)). Although the three-zone classification is used traditionally to classify nonuniformities on the microscale, it seems to be applicable to the nano- and macroscales as well (Messier et al. (1984)).

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Several models have been proposed to simulate the surface evolution as it occurs in zones I and II. These models are based on either a discrete or continuous approach. The Eden model (Meakin (1986)) was developed to simulate the growth of cell colonies. In the original version of this model, the growth process is started with a single occupied lattice site and unoccupied surface sites are occupied randomly with a probability which is proportional to the number of occupied nearest neighbors. In the ballistic aggregation model the growth process is started with a single stationary particle, and other particles are allowed to follow random ballistic (linear) trajectories in the vicinity of the stationary particle. If a mobile particle contacts the stationary particle, it sticks at that point and a stationary cluster is formed. Additional particles are added in the same way, one at a time, until a large aggregate has grown. Random ballistic aggregation of particles onto a surface under low particle surface mobility conditions leads to a columnar morphology observed in a wide range of vapordeposited thin films (Messier et al. (1984)). Diffusion-limited aggregation (Witten and Sander (1983)) is an example of how totally random motion can give rise to self-similar clusters. The simulation is started with a square lattice with one occupied site. A particle is then released from the perimeter of a large circle whose center coincides with the seed particle. The particle executes a random walk until it either leaves the circle or reaches a neighboring site of the seed particle. In the latter case, it becomes a part of the growing cluster. Muller (1985) developed a discrete model to simulate vapor-deposited thin film growth in two dimensions. In this model it is assumed that atoms travel in straight lines, which are inclined at an angle from the normal to the film surface. The points of intersection of these lines with the substrate are chosen randomly. When atoms impinge on the surface, they are assumed to relax instantaneously into the nearest-lying cradle formed by at least two other atoms. To simulate the migration of film atoms, the model allows atomic jumping if certain criteria are satisfied. Muller (1985) performed numerical calculations for the deposition of Ni, Ti, and W. From his simulations he determined the transition temperature for which the surface morphology changes from zone I to II as a function of deposition rate. His results compared well with transition temperatures determined experimentally for these materials. Mazor et al. (1988) proposed a continuous model to describe zone I and zone II deposition. His model assumes a constant uniform deposition rate J of atoms of radius δ and a finite surface diffusivity. By analyzing the derived equation and solving it numerically, Mazor et al. (1988) were able to predict under which circumstances small perturbations to the surface will grow or disappear. To estimate the transition temperature between zones I and II, he equated the critical temperature to that temperature at which perturbations of 5δ become stable. Mazor illustrated that this method gave results which were in good agreement with experimental results. 7.7. Instabilities Due to Interfacial Mass and Heat Transfer. An irregular surface morphology can also be expected when the growth rate is limited by the transport of reactants to the growing surface (Smith et al. (1987); Blocher et al. (1984)). When mass or thermal

boundary layers do not follow the curvature of perturbations on the surface closely, such perturbations can become stable and grow into dendrites. To determine whether such perturbations are stable or not, methods of linear stability analysis can be applied to the boundary layer equations for mass and heat transfer (van den Brekel (1978)). From the study van den Brekel concluded that the dimensionless Sh number will under most realistic growth conditions determine the smoothness of the deposition (van den Brekel (1981)). 7.8. Reactor Configuration and Mathematical Modeling. Historically, the development of CVD systems has proceeded along two directions: horizontal and vertical reactors. In the vertical reaction chamber, a gaseous precursor carrying the reactants is directed either toward the susceptor in a perpendicular direction or parallel to the deposition surface; the main characteristic is that the main flow direction is parallel to the gravitational field. The horizontal systems, on the other hand, are configured in such a way that the flow of reactants occurs parallel to the surface of the substrate and the gravitational field is normal to the main flow direction. Vertical systems have appeared as more amenable to scaleup than horizontal ones. Besides, the depletion of the reacting species as the gases flow over the substrate can adversely affect uniformity of concentration, doping, and thickness. The reactor design plays a decisive role on the properties of the materials deposited. The control and elimination of recirculating gas patterns over the substrate as well as the improvement in the species transport to the reaction surface are essential goals pursued when designing the chambers to carry out CVD processes. In simple geometries, as the horizontal planar channel reactor, one way to attain a homogeneous distribution of the reactant concentration gradients has been to tilt the susceptor in order to obtain a constant mass transfer driving force. In vertical reactors used in the microelectronic industry, the substrate is commonly rotated to expose the substrate wafer uniformly to the reacting gas stream. An innovative approach has been favored recently in vertical reactors: a very high speed rotating disk as a means of controlling the flow pattern in the reactor. In more complex configurations or applications other than microelectronics, the alternatives are not so apparent. The homogeneous distribution of reactants is a design problem without a definite answer so far. For instance, in the manufacture of crucibles the flow design of the CVD system precludes an even distribution of the reactants, resulting in nonuniform thickness depositions (Shinko and Lennartz (1987)). The thermal boundary conditions are also an important factor as sources of instabilities. Basically two distinct configurations are available: the cold-wall and the hot-wall reactors. The distance between the susceptor and the walls of the chamber is usually small with temperatures that can differ in hundreds of degrees; as a result, large-temperature gradients are observed in these reactors. A different set of thermal conditions is encountered in the hot-wall reactors, usually required for exothermic reactions. This suppresses the large transversal temperature gradients typical for cold-wall reactors, but gradients in the main direction of the flow can still produce sizable deviations from the traditional fluid dynamics scenarios. As was mentioned previously, the application of CVD


processes is also used to manufacture materials for the aerospace industry and a host of other applications as, for example, infinite inorganic fibers for metal and ceramic matrix reinforcement, optical windows for CO2 high-power laser systems, domes for high-speed missiles, windows and domes for passive imaging in guidance systems, refractory tungsten tubes for aerospace applications, etc. While the design and simulations of simple geometries in horizontal and vertical arrangements as well as the deposition of thin films are going to mature in the near future, the modeling of reactor configurations for applications mentioned above is in its early stages of development. Furthermore, the literature available on the latter problem is certainly scarce. Historically, the analysis of these systems did not draw much attention from traditional fluid dynamics: thus, most of the early works in CVD reactor design tried to obtain guidance from more specialized information available about the complex fluid dynamics in similar geometries. These efforts have been thoroughly discussed (Wahl (1977); Juza and Cermak (1982); Ludowise (1985)), and they are summarized in recent reviews by Hess et al. (1985) and Jensen (1987). In general, these models, based on oversimplified descriptions of the transport phenomena, were unable to provide a detailed description of the concentration, temperature, and velocity fields and the consequent spatial variations of the deposition thickness and growth rate. However, the good agreement with experimental average growth rates encouraged their use. The increasing interest in CVD processes and their broad industrial applications generated a tremendous feedback for the theoretical analyses. Thus, during the last years it has been found that most of the physical credence granted to the simplified models was undeserved and was just a result of the fortuitous cancellation of some of the consequences of the simplifications involved in the modeling step (Rosenberger (1987)). The advent of the supercomputers with their high speed and large memory capacity has produced the renaissance of the ambitious goal of an “accurate” modeling of CVD processes. A detailed description of two-dimensional or “incomplete” three-dimensional structures (assuming no breaking of symmetry or parabolized Navier-Stokes equations) provides a valuable insight in the parametric influence and a reliable guidance for the design of CVD reactors. Unlike in combustion systems the thermal effects associated with the chemical reactions are small in CVD systems. Due to the often very small concentrations of the reactants, the effect of the chemical reactions on the temperature and velocity fields is considered negligible. This allows a major simplification in both the analysis and simulation of transport phenomena in CVD systems: The mass conservation equation can be decoupled from the energy and momentum balances. This assumption is commonly adopted in CVD modeling and simulation studies. However, the disparate molecular weights between reactants and precursor (for instance, H2 and SiCl4) might generate considerable density gradients associated with the concentration field, giving rise to destabilizing effects of the same order as those due to the thermal expansion of the fluid. Thus, even for low-concentration reactants, solutal effects might have to be accounted for, and the coupling of the mass transport with the energy and momentum equations will be restored. One way to increase the rate of production in CVD

processes has been to reduce the operating pressure, increasing the mass transfer rate to the susceptor. Due to these features the flow field appears to be of less importance while multicomponent diffusion and chemical reaction are crucial elements in LPCVD models (Roegnik and Jensen (1987)). Although low, the typical pressures in LPCVD still permit the continuum hypothesis; however, recent efforts are being done in the direction of operating epitaxy reactors in the transition to free molecular regime (Meyerson et al. (1986)). The combination of large temperature gradients and large differences in molecular weights prevalent in CVD reactors adds the Soret (thermal diffusion) effect as a component of the mass transfer. Early modeling efforts did not considered this effect to be important; however, more recent publications have shown that the importance of this effect cannot be overlooked (Jenkinson and Pollard (1984)). Due to the strong temperature dependence of the physical properties of gases, the large temperature changes are accompanied by drastic changes in the physical properties, the transport coefficients, and, consequently, the dimensionless groups in the mass, energy, and momentum governing equations. Thus, one might expect that the modeling of CVD processes resorting to average properties and assuming an incompressible fluid will yield totally distorted results. Certainly, the predicted temperature and flow fields depend strongly on whether the physical properties are considered temperature dependent or not. However, the thickness of the deposited films, apparently due to compensating effects, does not show major variations when these temperature dependencies are neglected. Typically for the treatment of flow in ducts, the parabolic form of the Navier-Stokes equations is used. Such an approximation is strictly valid for flows dominated by convection in one direction, becoming more accurate as the Reynolds number increases. The main assumption is to neglect the axial diffusion of momentum as compared with the transport of momentum by convection. Also, pressure variations in the axial direction are assumed to be larger than variations of pressure in the transverse direction; thus, the axial pressure gradient is decoupled to form the transverse pressure gradients. The pressure field is represented by

P(x,y,z) ) P(x) + p(y,z) where P(x) represents the average pressure and is used only in the axial momentum balance; p(y,z), instead, represents small transverse pressure variations and is used in the transverse momentum equations. For the case of horizontal reactors, comparison between flow profiles determined with the parabolic approximation and with the full Navier-Stokes treatment has shown that the approximation is justified even for low values as used in CVD conditions. However, consideration should be given to its validity in the entrance region of the horizontal reactor where the gas expands rapidly as it contacts the hot susceptor. When an isothermal fluid is hydrodynamically developed, the velocity profiles will be given by the solution to the Navier-Stokes subject to the pertinent boundary conditions. Upon entering the heated region of the susceptor in an adiabatic chamber, the fluid starts to heat up and it will evolve to a new temperature profile within its thermal entry length. It can be shown that thermal instabilities will be present in this development region for all values of the Rayleigh number different

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than zero. Indeed, two-dimensional numerical results for a GaAs hot-wall reactor showed that rather small temperature differences between successive isothermal zones can cause backflow against the imposed forced flow. Depending on the kind of reactor, secondary flows induced by these conditions will be confined to the entrance region or will persist all through the reactor (Rosenberger (1987)). Thermal entry lengths determined in the absence of natural convection are expected to be conservative estimates of regions for transversal flow instabilities in the entrance of CVD reactors (Gatica et al. (1989)). The temperature distribution in the reactor, besides governing transport and physical fluid properties, will be responsible for the distribution of homogeneous gas phase reactions. Thus, the importance of an accurate description of this region becomes apparent. From a mathematical point of view it is important to know where the linear temperature profile is established, natural convection can still be present, and this new profile will determine the critical conditions for the onset of instabilities in the developed region. 7.9. Plasma-Enhanced Chemical Vapor Deposition (PECVD). Plasma-enhanced CVD (PECVD) has gained prominence as an important material processing tool since its introduction in the mid-1960s. In the PECVD process, gases are ionized by an electric energy source to form a plasma. Through the intermediation of the highly energized electrons, radical species are formed from the precursor molecular gases. These radicals are chemically very active, and their stability in the gas phase is determined by the rate of recombination and disproportionation reactions. The more stable radicals diffuse to the reactor walls where they are depleted by rapid insertion reactions. One outstanding feature of the PECVD process is the ability to energize electrons to very high levels, without any significant rise in the gas temperature. Processing at lower substrate temperature holds several advantages; e.g., temperature-sensitive material can be used, and a reduction in residual stress results, due to thermal mismatch. Whereas traditional CVD systems employed thermal energy to assist the deposition reaction, PECVD uses electrons produced by the glow discharge. The advantages of PECVD are obvious, but the understanding of and the ability to control this system are considerably more complex than thermal CVD processes. Historically, PECVD was developed for microelectronic applications. Passivation layers, diffusion masks, and interlayer dielectrics can be deposited at low temperatures. But the major advantage of PECVD was realized when macroelectronic devices were fabricated by this method. Examples are photovoltaic cells, large area display panels, linear arrays, and other thin film based technology. Recently, PECVD was used to deposit thin films of polycrystalline diamond. Microwave plasmas are normally used for diamond deposition, but they have certain limitations. If a PECVD process can be devised that produces a large concentration of atomic hydrogen and sufficient ion bombardment to remove any pyrolytic carbon, it has a good chance to deposit a diamond-like film at a fast rate. This brings us to the philosophy of a generic reactor design. Different objectives will lead to different designs. For example, if it is desirable to have a large concentration of atomic hydrogen, gap space and electric field strength are the crucial factors;

if a uniform deposition rate is the objective, then the hydrodynamic design will require more attention, and in industrial reactors, conversion and yield are the decisive factors in the design process. The glow discharge decomposition of silane to produce amorphous hydrogenated silicon (a-Si:H) is the most well-known PECVD process. It is important because of its use in solar cells and thin film transistors (Joannoupolos and Lucovsky (1984)). The Japanese have also embarked on a program to scale up plasma processes to achieve time and cost reduction in device fabrication through mass production. Another crucial development was the discovery in 1975 that the electronic properties of the flow discharge deposited material could be controlled very effectively by substitutional doping from the gas phase (Thomas (1985)). This possibility has opened a rapidly growing new field of research and applications. Orlicki et al. (1992) developed a model of a Si:H deposition in a dc glow discharge (Ar-SiH4) reactor. The parallel-plate configuration is used in this study. Electron and positive ion densities have been calculated in a self-consistent way. A macroscopic description that is based on the Boltzmann equation with forward scattering is used to calculate the ionization rate. The dissociation rate constant of SiH4 requires knowledge about the electron energy distribution function. Maxwell and Druyvesteyn distributions are compared, and the numerical results show that the deposition rate is lower for the Druyvesteyn distribution. The plasma chemistry model includes silane, silyl, silylene, disilane, hydrogen, and atomic hydrogen. The sensitivity of the deposition rate toward the branching ratios SiH3 and SiH2 as well as H2 and H during silyl dissociation is examined. Further parameters that are considered in the sensitivity analysis include anode/ cathode temperatures, pressure, applied voltage, gap distance, gap length, molar fraction of SiH4, and flow speed. This work offers insight into the effects of all design and control variables. 7.10. CVD Selected Applications. 7.10.1. Coating of Inorganic Fibers. Over the past decade the advanced composites industry has experienced a dramatic growth, in particular with respect to technological developments. This growth has been primarily driven by high-performance requirements in the aerospace industry, both commercial and military. The development of fiber-reinforced ceramic and metal matrix composites has been a key component. These lowdensity materials exhibit high elastic moduli and tensile strengths which makes them attractive in a host of engineering applications (Watts (1980); Suzuki and Umehara (1987); Di Carlo (1985)). Several fibers have been considered for the composite industry. Tungsten fibers coated with boron were the first commercially available (Morley (1987)) and have been used for the reinforcement of both polymeric materials and metals, particularly aluminum. These fibers are typically grown at a rate of 14 µm/s in monofilament reactors at a temperature 1260 °C (Watts (1980)). Below 1000 °C very low deposition is obtained, while large crystallites are obtained above 1300 °C (Morley (1987)). Typical final diameters vary from 100 to 236 µm (Morley (1987)). Other inorganic coated fibers commonly considered are SiC, Si3N4, BN, B4C, Al2O3, and TiB2 (Morley (1987); Galasso (1969); Watts (1980)). Examples of deposition reactions that can be performed are shown in Table 15 (cf. Blocher et al. (1984); Campbell et al. (1949); Green and Levy (1985); Stinton


Figure 21. Schematics of a chemical vapor deposition apparatus for fiber coating. Table 15. Typical Deposition Reactions no.

reaction

temp, °C

pressure, atm

1 2

2Ta + N2 f 2TaN (1/x)W + BCl3 + (3/2)H2 f (1/x) WBx + 3HCl TiCl4 + 2BCl3 + 5H2 f TiB2 + 10HCl 2BCl3 + 3H2 f 2B + 6HCl CH3SiCl3 f SiC + 3HCl

1000 1800-2000

1 1

1100-1300

1

1100-1300 1100-1400

1 1

3 4 5

et al. (1988); Wachtman and Haber (1986)). Examples 1 and 2 in this table demonstrate how a coating can be formed via a reaction between the substrate and a constituent in the gas phase. Alternatively, the desired coating can be obtained if two or more gas phase compounds react with each other (examples 3 and 4). Ceramic coatings can also be obtained from the decomposition of a single gas phase species (example 5). A versatile reactor configuration suitable to carry out these and many other reactions is shown schematically in Figure 21 (Galasso (1969); Krukonis (1977); Morley (1987)). The fiber core is mounted in a Pyrex tube between two electrodes. Depending on whether the reactor is operated in a batch or continuous mode, these electrodes can either be copper blocks or pools of mercury. The reactive gases flow over the substrate and react near or on the fiber surface to form the desired inorganic coating. In the synthesis of inorganic fibers, in contrast to other CVD applications, the substrate plays a decisive role in the performance of the reactor. As the deposition thickness can be of length scales similar to that of the original core, overall physicochemical properties of the substrate can change substantially in space and time. Furthermore, chemical changes can occur in the core itself. An illustrative example of this is the deposition of boron on a tungsten core (Krukonis (1977); Pierson (1981); Gallasso (1969)). The evolution of several distinctive layers can be noticed during the deposition process. Initially, a thin layer of WB forms adjacent to the surface. A second distinctive layer containing W2B5 appears then beneath the WB layer. As the process continues, a B mantle develops external to the original core, the WB layer disappears, and the core is converted completely to W2B5. 7.10.2. Manufacturing of Infrared Windows. The CVD process offers several advantages over other techniques for preparing infrared-transmitting materials (Gentilman et al. (1981)). First, the resulting material is usually theoretically dense and also very pure; thus, light scattering and impurity absorptions, respectively, are minimized. Furthermore, CVD systems have no inherent limitations in the size, thickness, or shape of the deposit which can be produced. Optical systems operating in the longer wavelength bands tend to have large apertures. Therefore, large

windows are required with linear dimensions up to 0.5 m or more. Free-standing windows of this size need to be 0.01-0.02 m thick in order to have sufficient mechanical strength. For the long infrared wavelength applications, ZnSe with its extremely low absorption coefficient and ZnS with its lower cost and higher strength have become the leading materials for highpower CO2 laser windows and various passive imaging systems. 7.10.3. Manufacturing of Crucibles for Use in the Czochralski Method. Silicon single crystals for semiconductor devices are usually prepared by the Czochralski (CZ) method. The CZ crystals are pulled from high-purity fused quarts crucibles. The highly corrosive power of the molten silicon invariably produces erosion in the crucible, and the silicon crystals are contaminated with oxygen. An alternate way is to produce silicon nitride crucibles by CVD instead of quartz crucibles (Shinko (1987)). The CVD silicon nitride crucibles have been shown to permit growth of nearly oxygen-free silicon single crystals. The chemical and thermal stability of crystalline silicon nitride makes it well-suited to the molten silicon environment. Another application of crucibles manufactured by CVD processes is the pyrolitic boron nitride crucibles used in the LEC (liquid encapsulated Czochralski) of III-V compound semiconductors including GaAs and InP (Iwasa et al. (1987)). The crystals are pulled through a B2O3 melt that suppresses the escape of As vapors from the molten GaAs and maintains the stoichiometry of the crystal. The molten B2O3 remains in the crucible after the crystals have been grown; during the cooling step the melt shrinks faster than the crucible itself and stresses develop in the crucible. These stresses are the source of delamination of the crucibles. The CVD process determines the crystalline structure of the crucible, and thus the cycle life of the crucibles can be improved. 7.10.4. Manufacturing of Tubes. CVD is also used in the manufacturing of tungsten (Yang et al. (1973)) and silicon carbide (Zhao et al. (1993)) tubes. The direct process can be used to make tube shells for further working or tubing in its final form. The as-deposited structure of the chemical vapor deposited tungsten may be sufficient and/or desirable for some applications. Simple shapes such as tubing in lengths up to 1.2 m can be made easily and are less expensive than equivalent tubing made by the conventional extrusion-drawing method. These tubes can be made in various diameters with very uniform wall thickness. The method has proven to be very attractive for the manufacture of short-length tubes of higher purity (less than 25 ppm total metal impurities). The manufacture of these tubes is performed by arranging a resistance-heated mandrel which is surrounded by a water-cooled chamber fitted with insulated electrodes. 8. Inorganic Fibers A new generation of composite materials is revolutionizing today’s aircraft and automotive industries. In these high-technology applications, weight, strength and stiffness are critical parameters which can be achieved by using high-performance ceramic material composites (CMC) and metal matrix composites (MMC). These composite materials offer many outstanding properties for high-temperature applications.

372


The much higher thrust/weight performance requirements for improved engines for 21st century supersonic aircraft demand a major increase in engine operating temperatures and a major reduction in engine weight. This improved performance will be dependent upon the development of lightweight materials capable of withstanding much higher turbine operating temperatures, higher pressure ratios, and increased rotational speeds. Conventional materials are limited in the improvements that they can offer, so in order to meet these goals, new materials such as intermetallic matrix composites (IMC), ceramic matrix composites (CMC), and carbon/carbon composites (C/C) must be developed. This trend led to the discovery of fibers with excellent high-temperature strength and corrosion resistance, good thermal shock resistance, low density, and low coefficient of thermal expansion. High-performance fibers are available only at high prices and, as a result, tend to serve specialized and restricted markets. Basically, three factors effect this price: raw material, manufacturing, and most importantly new market development cost. Second, the use is partly limited because of their brittle nature and large degree of scatter in strengths which often lead to catastrophic failure. The ideal fiber for reinforcement of ceramic and metallic matrix composites would be chemically compatible with the matrix, would have a coefficient of thermal expansion (CTE) that matches that of the matrix, and would have a low density and a high strength and modulus at elevated temperatures. Unfortunately, there are no commercially available fibers that meet exactly all these criteria. The key factor in the development of advanced composite materials is the development of advanced fibers to reinforce these composites. The fiber provides strength and stiffness to the matrix, so the key to the development of superior advanced composites is the development of superior reinforcing fibers. The most important goal is that the fiber should have a CTE close to that of the matrix, so that the thermal expansion mismatch between the fiber and the matrix can be minimized. A fiber with a CTE of about 10 ppm/ °C would be needed to match Titanium aluminides, while a CTE of up to 16 ppm/°C would be needed for matching NiAl matrices. Matching thermal expansions reduces thermal strains and stresses that could cause fiber or matrix cracking during thermal cycling. The fiber must also have a modulus that is higher than that of the matrix, so that the fiber can carry a greater portion of the elastic loads imposed on the composite. The fiber must also be chemically compatible with potential matrix materials, so that the properties of the composite are not severely degraded during fabrication or service. For aerospace applications currently available fibers do not meet all these fiber goals. Tungsten and molybdenum fibers have a high density, while boron fibers do not have the required strength at elevated temperatures. SiC fibers have strength at elevated temperatures but are reactive with most matrices and have a low CTE. Graphite fibers, available as small-diameter multifiber yarns, are difficult to fabricate into hightemperature intermetallic matrix composites. Single crystal Al2O3 and polycrystalline TiB2 and TiC fibers are under development for aerospace applications but have not yet demonstrated the required elevated temperature properties and CTE needed to reinforce nickel aluminide matrices.

Over the past several years, numerous patents and publications have revealed the use of simple organometallic compounds, metal halides, and polymer precursors for routes to ceramic fibers. Essentially, two different types of technologies have been studied extensively, i.e., chemical vapor deposition (CVD) and polymer pyrolysis (Yajima et al. (1978)). This paper will address the two most widely evaluated continuous silicon carbide fibers: Nippon Carbon’s Nicalon fiber marketed by Dow Corning in the U.S. and Textron Specialty Materials Division’s fiber. These two fibers are different in several respects, including fabrication, chemistry, size, and properties. One of the major differences in Nicalon and Textron’s fibers is their method of preparation. Nicalon is fabricated by polymer pyrolysis, while Textron’s fiber is made via chemical vapor deposition. The process of making silicon carbide fibers from polycarbosilane was developed in Japan (Yajima (1983)). The reaction of dimethyldichlorosilane with Na metal yields poly(dimethylsilane). This will polymerize in an autoclave to obtain a product known as polycarbosilane. A distillation step removes a low molecular weight component from the polycarbosilane, which can be melt spun at 350 °C to produce a fiber. The fiber is heated in air to render it infusible and then pyrolyzed at 12001300 °C in an inert gas or vacuum to yield the final 1015 µm diameter SiC fiber. Textron’s fiber is manufactured by chemical vapor deposition, in a process very similar to the production process for boron fiber. A 33 µm carbon monofilament substrate is fed through a glass reactor and resistively heated through a mercury contact. Pyrolytic graphite (∼1 µm) is first deposited onto the substrate to improve the surface for silicon carbide deposition. Following this, hydrogen, argon, and silane vapors are fed into the reactor at approximately 1300 °C, resulting in the deposition of β-SiC in a columnar fashion onto the substrate. If the fiber was removed from the reactor at this point and the strength tested, its ultimate tensile strength would be approximately 300 Ksi. This relatively low strength is due to the exposed SiC grain boundaries, which can act as stress concentrators and lower the potential strength of the fiber. By adding a thin coating of amorphous carbon (∼1 µm), these grain boundaries can be sealed and fiber strength is nearly doubled. By varying the ratio of Si to C, a final thin coating (1-6 µm) can be tailored to the specific matrix material for optimal fiber-matrix strength. This is particularly critical for ceramic composites, in that it is generally conceded that the interfacial bond strength is one of the most important factors in determining toughening. Theoretically, the stoichiometric SiC fiber should retain the strength up to 1700 °C. One of the reasons for reduction in strength is due to surface pitting which is associated with metallic impurities and the conversion of excess carbon to carbon monoxide. Due to the low solubility of Si and C in silicon carbide, a laminar structure is formed with alternating thin Si/SiC layers or C/SiC layers. The free Si and C and morphology of the fiber entirely depends on deposition parameters. The thermal instability of these ceramic composites, which is often a result of fiber-matrix interfacial degradation, has prevented prolonged application of these composites at high temperature. The composition and properties of different SiC fibers are given in Table 16. Hence,

Ind. Eng. Chem. Res., Vol. 35, No. 2, 1996 373 Table 16. Properties of SiC Fibers (Sheppard (1990))

fiber Nicalon (Nippon Carbon Co.) SCS-6 (Textron) Tyranno (Ube Industries Corp.) HPZ (Dow Corning Corp.)

composition

tensile elastic density strength modulus diameter (g/cm3) (GPa) (GPa) (µm)

Si, C, O

2.5

2.8

180

10-20

SiC on carbon core Si, C, O, Ti

3.0

3.9

406

143

2.4

3.2

180

10-12

Si, N, C, O

2.4

2.8

180

10-12

future development will be concentrated on stoichiometric deposition of ultraclean silicon carbide. 9. Conclusions If high-performance ceramics should reach the massive level of production needed for, e.g., car industry, it seems inevitable that integration of chemical routes back to the source ores will occur for economical reasons. Significant reduction of processing costs is necessary to expand the current limited markets. A heavy involvement of the chemical industry is necessary; however, the markets, despite generous funding from several government agencies, have not been developed. Currently, we feel that with the end of the Cold War the economic forecasts were too optimistic and the current growth does not agree with the published predictions in the past. Nevertheless, we believe that a larger number of chemical engineers should be involved in the ceramic area. This field represents a challenging world of new opportunities and requires skills and expertise from several disciplines. Mathematical modeling of complex reaction processes involving a formation of solid phases and their high-temperature densification in the presence of sintering aids is still in the virgin state. Advances in this theoretical approach should help to solve many complex problems in the advanced materials synthesis and processing. Based on our past experience and knowledge of the current state-of-the-art, we are listing below problems which have not been adequately studied and solved or which require new input: (1) A theoretical foundation of the sol-gel method is missing. (2) The sol-gel method should expand more aggressively into the non-oxide ceramic area. Attention should be paid to B-C-N complex ceramics. (3) The development of new inexpensive organometallic precursors is needed. (4) The engineering of production, handling, and processing of ultrafine powders is still not developed. Special attention should be paid to minimization of oxygen surface concentration which is usually very high. (5) Reactivity of solid materials was historically neglected by the chemical engineering community. Reaction engineering principles developed for catalytic processes must be extended to several important classes of noncatalytic solid-solid (liquid) and gas-solid reactions. Among the most important issues is the development of a systematic approach of measurement and evaluation of reaction rate data. (6) Carbon reactivity and elimination of carbonaceous residuum in products formed in carbothermal synthesis need more investigation.

(7) The theoretical foundation of reaction sintering is fragmentary and theoretical work is imperative for future progress in the field. (8) The same conclusion is true for other net shape processes; chemical engineering ingenuity may be a catalyst in a future development of this field, (9) Joining of ceramics is currently at the level of alchemistic exploration. Theoretical modeling based on reactivity, diffusivity, and stress distribution is needed. (10) CVD is a mature field. New methods of fast deposition of films in a stable regime and synthesis of “thermodynamically unstable” composites should be developed. Diamond deposition techniques should be generalized to diamond composite materials and thick films. (11) Recent advances in modeling of “classical CVD arrangements” ought to be extended to plasma and microwave-assisted CVD. (12) A field of advanced composites needs an inexpensive source of SiC fiber. Methods of a fast deposition of SiC in a stable regime must be developed so that a residence time in CVD reactors will be reduced by an order of magnitude. (13) Systematic application of CVD in a nonelectronic area is rare. Use of CVD to grow thick films should be pursued; new development can be expected. (14) The global modeling approach, including stress distribution, surface phenomena, sintering, effect of electromagnetic fields, etc., is a challenge for people trained in chemical engineering analysis of reacting systems. Nomenclature c ) specific heat, J‚K-1‚kg-1 E ) activation energy, J‚mol-1 J ) deposition rate, kg‚mole‚m-2‚s-1 k ) filtration coefficient, m3‚s‚kg-1 k0 ) frequency factor Mg ) molecular weight, kg‚kmol-1 p ) pressure, N‚m-2 R ) universal gas constant, J‚mol-1‚K-1 t ) time, s T ) temperature, K v ) gas velocity, m‚s-1 W ) reaction rate, kg‚m-3‚s-1 x ) length coordinate, m y ) height coordinate, m Greek Letters ∆H ) heat of reaction, J‚kg-1 δ ) atomic radius, m ) void fraction λ ) effective thermal conductivity, W‚m-1‚K-1 µ ) stoiciometric coefficient ν ) order of reaction with respect to a gas reactant F ) density, kg‚m-3 Subscripts g ) gas s ) solid R ) solid reactant

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Received for review February 7, 1995 Revised manuscript received June 29, 1995 Accepted October 13, 1995X IE9501034

X Abstract published in Advance ACS Abstracts, December 15, 1995.

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