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temperatures, so the coated fibers can be used to fabricate copper matrix composites (ref. 41).
In the past 10 years, substantial effort has been made to develop carbonfiber-reinforced carbon and ceramic matrix composites that can be used for high-temperature structural applications in oxidizing environments. Although the development of external coatings to protect the composites has been vigorously pursued, little work has been conducted in the United States on fiber coatings as a method of internal oxidation protection. However, studies reported by Soviet investigators on CVD SiC- and ZrC-coated carbon fibers have been encouraging (ref. 42). This work has shown an order-of-magnitude decrease in the oxidation rate at 600°C and 800°C in air compared with the uncoated fibers. Increasing the thickness of the SiC coatings from 0.02 pm to 0.10 /mi resulted in an additional fourfold decrease in the oxidation rate. Tests in CO2 revealed very low oxidation rates for the coated fibers, even at 1000°C. It is clear that the Soviet tests were conducted so that oxidation from the ends of the fibers did not occur; similar tests have demonstrated rapid oxidation of SiC-coated fibers at temperatures above 500°C because of attack from the unprotected cut ends (ref. 20).
The effects of coatings on fiber mechanical properties are well documented. Generally, studies show that without major strength degradation, adherent coatings change the elastic modulus according to the simple rule of mixtures (refs. 15, 26, and 43). The tensile strengths of high-performance carbon fibers are rarely enhanced by the presence of coatings, although electroplated ductile metal coatings have relatively little effect on fiber strength because of the high failure strain of the coatings (refs. 24, 26, and 43). In contrast, brittle coatings that are well bonded to the fibers tend to decrease fiber strengths because cracks that propagate in the coatings at low strains extend into the fibers and result in premature failure (refs. 15, 17, and 18). This physical effect combined with possible chemical damage to the fiber surfaces during coating can produce substantial fiber strength losses (refs. 19 and 21). Strength reductions can be minimal, however, if ceramic coatings are kept thin and are made under conditions that reduce the tendency for chemical attack.
Maximum coating defect sizes decrease with decreasing coating thickness, so coating failure strains increase. Thin coatings produce a more flexible yam and are less likely to join fibers together, so fiber damage during weaving and handling is less of a factor (ref. 20). Coating conditions such as lower temperatures and carbonrich environments can be anticipated to reduce the effects of aggressive chemicals. Under the best conditions, strength reductions of 10 percent to 15 percent can be expected when 0.1 -pm- to 0.2-/xm-thick brittle coatings are deposited on 7-/xm- to 8-/xm-diameter, high-performance carbon fibers (refs. 14, 15, 17, 18, and 21).
High-Temperature Coatings on CC Composites
In the late 1950's, converting carbon fiber-resin matrix composites into carbon fiber-carbon matrix composites by slow pyrolysis of the resin was found to produce useful high-temperature materials (ref. 44). Densification of these materials by repeated cycles of resin or pitch impregnation and pyrolysis or by infiltration and decomposition of hydrocarbon gases produced bodies with thermal shock and fracture toughness properties far superior to those of earlier synthetic graphite materials (ref. 45).
To this day, CC composites are principally known as lightweight, veryhigh-temperature materials with superior thermal shock, toughness, and ablation properties. These attributes and specific tribological properties are emphasized in present aerospace and military applications for CC composites such as rocket nozzles, heat shields, reentry vehicle nosetips, and aircraft brakes (refs. 46 to 50).
It is not because of a lack of structural capability that current applications do not involve long-term operation under substantia] primary loads. When made with high-performance carbon fibers, CC composites can have strength and stiffness characteristics comparable to those of many metal alloys. Because a variety of fiber weaves are available, properties can be tailored for specific component applications.
Carbon-carbon composites with superior mechanical properties contain highly oriented fibers derived from mesophase pitch or PAN resin. Fibers of this type have been thermally stabilized for normal CC processing and have strengths in the 2000 MPa to 2400 MPa range and elastic moduli between 350 GPa and 400 GPa. Composite elastic moduli in the principal fiber directions usually reflect the fiber modulus according to the rule of mixtures. Strength utilization of the fibers is 25 percent to 50 percent of rule-of-mixtures predictions, depending on the fiber architecture and processing specifics (refs. 51 to 53).
Similar to many fiber composites, shear and cross-fiber tensile properties are often a major issue and frequently have a strong influence on design and materials selection. In general, the shear strengths and moduli of CC composites are low compared with those of more conventional materials. The same limitations exist for tensile properties in directions normal to the fiber directions because of processing that creates less than optimum bonding between the matrix and fibers.
One of the reasons CC composites are attractive is because they can exhibit relatively high values of fracture toughness. These materials belong to a special group of composites in which the failure strain of the matrix is much lower than that of the fibers. As a result, strong bonding between the matrix and fibers produces composites that are both brittle and low in strength because strong bonds promote failure of the fibers at the low strains where the matrix fails. Conversely, relatively weak bonds allow matrix cracking to occur without crack propagation throughout the fibers.
Matrix cracking can occur during CC processing as a result of matrix shrinkage during pyrolysis, the thermal expansion coefficient (CTE) mismatch between the fibers and matrix, and when the composite is loaded in service. Strong bonding in carbonization can result in fiber damage and, combined with strong bonding that persists in the final product, can result in weak and brittle composites. In contrast, weak bonds leave fibers undamaged and allow the matrix to crack in service; however, they continue to transfer load to the fibers so the composite can utilize the high strength of the fibers. Shear displacement of the matrix relative to the fibers and eventual fracture of the fibers at points away from the primary zone of matrix failure create the fibrous pullout fracture surfaces indicative of toughness.
A good method of comparing the fracture toughness of two materials is the notched beam work-of-fracture test. This simple test measures the energy required to move a tensile crack through the material in a controlled fashion. Work-offracture tests on strong, fully densified CC specimens containing high-performance fibers showed fracture energy values of approximately 2 x 104 Jm-2 for a unidirectional material and 3 x 10 Jm for a fabric laminate when cracks were propagated across the fibers (refs. 54 and 55). By comparison, plastics give values of 103 to 10 Jm , technical-grade ceramics and premium-grade graphites show values less than 102 Jm-2, wood exhibits values of about 104 Jm , and ductile metals have values of more than 105 Jm-2 (ref. 56).
Carbon materials are unique in their retention of mechanical properties at high temperatures. High-performance fibers exhibit modest tensile strength increases as temperatures increase to about 2200°C, while the tensile elastic modulus gradually decreases to about 50 percent of the room temperature value at 2200°C (ref. 57). Carbon-carbon composites show similar behavior in the principal fiber directions. The matrix and matrix bonding to the fibers, however, play a large role in determining the effect of temperature on the shear, cross-fiber tensile, and compressive strengths. One reason for matrix cracking is the CTE mismatch with the highly oriented fibers. Cracking is extensive at room temperature. Cracks close as the temperature increases, providing a more coherent and, perhaps, adherent matrix. This phenomenon accounts for increases in shear and crossfiber tensile strengths that are observed with increasing temperature and, together with increasing fiber strengths, is consistent with observed increases in compressive strength. Moduli either decrease or increase with temperature, depending on the influence of matrix cracking relative to the basic thermal reductions in modulus of the constituents. This change appears to depend on fiber architecture and specifics of the CC processing.
The high-temperature time-dependent deformation of carbon fibers and CC composites is currently an active area of investigation (refs. 58 to 60). Although comprehensive scientific and engineering studies are needed to define creep mechanisms and the effects of microstructure, fabrication processing, and thermal and mechanical history on deformation behavior, the available data clearly demonstrate the superior structural capability of CC composites at very high temperatures. For example, conservative interpretation of the data projects creep rates on the order of 10~3 percent per hour in the 2000°C to 2200°C range at stresses of 70 MPa to 140 MPa for unidirectionally reinforced CC composites.
The primary reason coatings are applied to CC composites is to provide oxidation or erosion protection at high temperatures. Until recently, CC composites were considered to be nonstructural thermal protection materials based on their excellent performance in a number of rocket propulsion and atmospheric reentry applications. These applications involve extremely high temperatures but require very short lifetimes and take advantage of the unmatched ablative properties of the all-carbon composite. Although ceramic and refractory metal external coatings were considered and evaluated for these applications, they are generally found to be of little or no benefit.
The first real need for a coated CC composite was identified during the development of the Space Shuttle orbiter vehicles. Again, CC was used as a thermal protection material, but in this case, the temperatures were much lower than those for earlier applications, and the Space Shuttle material had to be reusable. The success that was achieved with coated CC in the Space Shuttle program provided the impetus for the current development of oxidation-protected CC composites for structural applications in aircraft and limited life of turbine engines, and as airframe components for multiuse hypersonic vehicles. A renewed interest also exists in the use of coated CC composites for rocket propulsion and missile aerodynamic heating applications because of the requirements of advanced higher performance systems.
Space Shuttle Orbiter Thermal Protection
The first coated CC was developed in the early 1970's for use on the National Aeronautics and Space Administration (NASA) Space Shuttle orbiter vehicles. Termed reusable carbon carbon (RCC). this CC is a resin-densified fabric laminate made from low-modulus, rayon-precursor carbon fibers (ref. 44). The RCC material is protected from oxidation by a two-layer coating system consisting of a porous SiC inner layer sealed with SiC<2 and an outer coating of alkali silicate glass filled with SiC particulate. A pack process is used to convert the CC surface to SiC (refs. 44, 61. and 62). followed by repeated cycles of impregnating the porous converted surface with acid-activated tetraethoxysilane (TEOS) liquid that is gelled and then heated to produce SiCh (ref. 63). The outer coating on the carbon structure is made from a paste of commercial alkali silicate bonding liquid filled with SiC powder. The coating is applied to the structure by brushing and then it is heat-cured (ref. 63).
The RCC has been used to construct the nose cap and leading edge components for all of the Space Shuttle vehicles. It has performed well as a reusable thermal protection material through multiple rapid heating and cooling cycles, with shortterm exposures at temperatures of more than 1300°C (ref. 44). The durability of the RCC coating system can be attributed to several factors. First, the RCC has a relatively high CTE and a low elastic modulus in the plane of the fibers. Second, the inner SiC conversion layer is a strongly bonded, integral part of the CC surface. Finally, the outer coating is a fluxed SiO_> glaze that can plastically deform at high temperatures to accommodate differential strains and is fluid enough at temperatures of more than 900°C to be self-healing. A problem with the RCC outer coating is that the rate of alkali volatilization from SiO2 becomes significant in the 1200°C to 1300°C range, allowing the loss of the fluxing agent over periods of hours to compromise the compliancy and sealing characteristics of the glaze.
By the late 1970's, the encouraging test results that had been obtained with the RCC material were well known. NASA's success with CC, combined with disappointments in a Department of Defense (DoD) sponsored high-performance monolithic ceramics program, generated strong DoD support for the development of oxidation-protected CC materials for very-high-temperature structures in advanced military systems. Although it was recognized that major differences exist between RCC and structural CC materials, it is clear that the amount of development required to meet a number of very optimistic design and performance goals has only recently been fully appreciated.
Coated RCC has very modest mechanical properties and a relatively high CTE in the plane of the fibers. A tensile strength of 62.5 MPa, a tensile elastic modulus of 14.5 GPa, and a CTE of 2.6 x 1(T6 °C-1 are average RCC in-plane properties (ref. 44). In contrast, the projected structural applications required tensile strengths in the 207-MPa to 345-MPa range. This dictated the use of high-performance carbon fibers that produce composites with elastic moduli in the 90-GPa to 117GPa range and CTE values of approximately 1.8 x 10~6 °C-1.
Early systematic oxidation testing of the first structural oxidation-protected CC was conducted by two U.S. Air Force contractors in 1981. The CC was a twodimensional laminate made from high-performance PAN-precursor fibers and was designated advanced carbon carbon (ACC). The best results were obtained with the RCC coating system in which the SiC conversion layer was modified with boron (ref. 64). Burner rig testing that involved rapid heating and cooling between 150°C and 1371°C with half-hour holds at the maximum temperature resulted in several early failures; this testing also resulted in a number of specimens that