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two principal reinforcements used in current CC composites, and the two most promising candidates for the CC applications of tomorrow.

Carbon Fiber Processes

Carbon fibers are produced by the thermal decomposition of various organic fiber precursors, such as PAN, rayon, and mesophase pitch. As mentioned in chapter 2, rayon and PAN are normally extruded into filaments using solutionspinning techniques. These precursor fibers then are thermoset by oxidizing them in air for extended periods. For PAN precursor fiber, this results in the formation of an infusible cyclized network of hexagonal, carbon-nitrogen rings. For rayon precursor fiber, some depolymerization occurs, and a similar aromatization of the remaining structure begins. The stabilized rayon and PAN fibers finally are carbonized by heating them in an inert atmosphere to approximately 1500°C (ref. 1). Subsequent graphitization at temperatures ranging from 2500°C to 3000°C can improve the mechanical properties of both fibers. The mechanical properties of rayon are enhanced considerably if the rayon is stretched during this graphitization. If the thermoset rayon precursor fiber is graphitized without stretching at 2800°C, it develops a tensile modulus of only 69 GPa. The same fiber can develop a modulus of 690 GPa if it is stretched 100 percent during graphitization at the same temperature (ref. 2). However, this stress-graphitization step is extremely expensive, and the resulting high-modulus fiber is brittle, making it difficult to handle.

The starting material used to produce pitch-based carbon fibers is usually a highly aromatic coal or petroleum pitch. For the production of high-performance fibers, the pitch first must be converted to an oriented liquid-crystalline material—a mesophase. Heating the raw pitch to temperatures ranging from 400°C to 410°C for periods of up to 40 hr (ref. 3) usually accomplishes this conversion. Solvent extraction (ref. 4) is another method for preparing mesophase, which involves first separating the high-molecular weight fraction of the pitch. The extraction step removes the smaller disordering molecules and concentrates the higher molecular weight portion of the pitch in the insoluble portion. This insoluble fraction then can be polymerized to a 100-percent anisotropic phase by heating it to between 230°C and 400°C for only 10 min.

This liquid-crystalline product, or mesophase, is then melt-spun into a precursor fiber. The viscosity of commercial mesophase precursors is extremely temperaturedependent, making precise temperature control during melt-spinning vital. Because mesophase is actually a lyotropic liquid-crystalline solution, it becomes highly oriented during extrusion and subsequent drawdown (ref. 5). Therefore, unlike PAN and rayon, as-spun pitch precursor fiber has an extremely high degree of molecular orientation.

Similar to PAN and rayon, the melt-spun mesophase pitch fibers must be thermoset to render them infusible and then carbonized to develop their final properties. Because mesophase pitch fibers develop an extremely high degree of molecular orientation during fiber formation, the object of subsequent heat treatment is not to develop, but instead to preserve, the molecular orientation. Because of this, the conditions used in the thermosetting and carbonization steps are usually mild compared with those used for PAN and rayon precursor fibers. Typically, temperatures of approximately 300°C and exposure to air for a period ranging from 30 min to 2 hr are required to fully thermoset mesophase fibers. Subsequent heat treatment in an inert atmosphere, such as nitrogen, allows the thermoset mesophase precursor fiber to develop its final properties. Because the precursor fiber is already highly oriented, mesophase pitch precursor fibers develop tensile moduli of approximately 690 GPa when heat-treated at 2800°C for a few minutes, even when no tension is applied (ref. 1).

Effect of Graphite Structure on Fiber Properties

Carbon fibers are composed of 99.9 percent pure carbon, most of which is arranged into graphite crystallites. These graphite crystallites organize into layer planes, and the mechanical properties of carbon fibers are a result of this structure. This layer-plane structure and the corresponding lattice dimensions of graphite, as well as its approximate orientation within carbon fibers, are shown in figure 1 (ref. 6). Within these layers, chemical bonds link the carbon atoms in the crystallographic a-direction. These anisotropic, nonpolar, a-bonds, created by the sp2 hybridization of the electron orbitals in the carbon layers, make the graphite structure extremely strong in the crystallographic a-direction (ref. 7). This makes the theoretical tensile modulus and the ultimate tensile strength of graphite extremely high, approximately 1060 GPa (ref. 8) and 106 GPa (refs. 9 and 10), respectively, for a load applied along this crystallographic axis. On the other hand, very weak van der Waals bonding exists between atoms in adjoining planes. These weak bonds make the mechanical properties of the graphite crystal in this direction, the crystallographic c-direction. quite low. As a result, the theoretical tensile modulus for a load applied in this direction, normal to the crystallographic a-axis. is only 35 GPa (ref. 1).

The structure of the graphite crystallites in a carbon fiber is not perfect, as shown in figure I. The layer planes of carbon are slightly offset, creating what is termed turbostratic graphite. This offset also results in a slight increase in the interlayer spacing (3.37 A to 3.45 A), compared with the 3.35 A found in a perfect crystal. Even though the interlayer spacing is not perfect, the orientation of these graphite layers is more or less parallel to the fiber axis, and it is the strong covalent bonding within these layer planes that accounts for the longitudinal properties in carbon fibers. Thus, as one would expect, the layer planes in high-modulus carbon fibers are almost perfectly aligned with the fiber axis.

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Figure I. Structure of graphite and approximate orientation in carbon fibers (ref. 6).

Today commercial pitch-based carbon fiber is available with a tensile modulus that is 85 percent of the theoretical maximum for the perfect graphite crystal. By comparison, the tensile strength of this fiber is only 5 percent of that theoretically possible. Although increasing the layer plane alignment directly improves the tensile modulus of carbon fibers, flaws limit the tensile strength of these brittle fibers. A brief review of the interaction between crystallite size and brittle fracture follows to help understand how this limitation relates to the microstructure of the carbon fiber.

Brittle Failure Mechanism

Because flaws limit the tensile strength, considerable effort has been made to minimize contamination in both PAN-based and pitch-based carbon fibers (refs. 11 and 12). Yet even when flaw-inducing particles are minimized, structural flaws within carbon fibers still limit the tensile strength. Johnson (ref. 13) recently presented an excellent explanation of these structural limitations of carbon fibers and of the interaction between microstructure and the flaw sensitivity of these fibers, based on the mechanism for tensile failure proposed by Reynolds and Sharp (see ref. 13). As mentioned previously, the crystallites and carbon layer planes are not aligned perfectly in carbon fibers, and misoriented crystallites are relatively common. When a stress is applied parallel to the fiber axis, these crystallites align until movement is restricted by a disclination in the structure (as shown in fig. 2). If the stress is sufficient, a basal-plane rupture of the misoriented crystallite can relieve the stress within the fiber. When the size of this ruptured crystallite (perpendicular to the fiber axis) is larger than the critical flaw size, a catastrophic failure occurs, breaking the entire fiber.

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Figure 2. Reynolds and Sharp mechanism for tensile failure (ref 13). Misoriented crystallite linking two crystallites parallel to fiber a.\is (left), tensile stress causing basal plane rupture in Ln direction (center), and catastrophic failure occurring if crystallite size is greater than flaw size (right).

Even if the ruptured crystallite is smaller than the critical flaw size, catastrophic failure can still occur if the crystallites surrounding the disclination are continuous enough to allow a crack to propagate into neighboring crystallites. Johnson (ref. 13) indicates that this failure mode, created by extended three-dimensional (3-D) crystalline order, explains the difference in flaw sensitivity between PANbased carbon fibers and the more highly graphitic, pitch-based carbon fibers.

The strong covalent bonding within the layer planes thus causes the tensile modulus of carbon fiber to increase as the crystallite orientation parallel to the fiber axis increases. However, when a stress is applied parallel to the fiber axis, misoriented crystallites must rupture to relieve the stress. As the size of the crystallites increases, it becomes more likely that some will exceed the critical flaw size. Because of this, the final balance in the physical properties of carbon fibers depends on the size, the extent of 3-D order, and the orientation of the graphite crystallites.

Microstructure of Carbon Fibers

Because the precursor fiber for PAN-based carbon fibers is formed by precipitating a linear polymer from solution, whereas the precursor fiber for pitch-based carbon fibers is formed by melt-spinning a cyclic organic liquid crystal, it is not surprising that the arrangement of the graphite crystallites (or microstructure) in these carbon fibers differs considerably. As explained in the previous section, the physical properties of carbon fibers are a direct result of their microstructure. Therefore, the microstructures of both PAN-based and pitch-based carbon fibers must be understood before the potential, as well as the limitations, of both types of carbon fiber can be appreciated.

Microstructure of PAN-Based Carbon Fibers

Early work by Diefendorf and Tokarsky (ref. 14) indicated that the macrostructure of PAN-based carbon fibers, similar to that of many other synthetic precursor fibers, is fibrillar. Although the exact nature of this microstructure has been refined in more recent studies, the ribbon-like undulations (characteristic of a fibrillar structure) are an excellent visualization of the microstructure of PAN-based carbon fibers. Diefendorf and Tokarsky (ref. 14) also showed that the amplitude of the undulation in the fibrillar structure was highest in the center and lowest near the surface of PAN-based carbon fibers. This indicates that the modulus of a PANbased carbon fiber varies throughout the cross section of the fiber. Figure 3 shows a model of this microstructure.

Johnson (ref. 13) studied various types of PAN-based carbon fibers in an effort to elucidate the relationship between the microstructure of these fibers and their properties. Using wide-angle X-ray diffraction, he found that the layer planes of PAN-based carbon fibers have no regular 3-D order. In the skin region of the fiber, small-angle X-ray diffraction and transmission electron microscope (TEM) analysis of longitudinal and transverse fiber sections revealed needle-shaped voids between crystallites and layer planes that essentially were parallel to the surface. However, in the core region, Johnson found that the layer planes were folded extensively, often through angles of 180°. He proposed that misoriented crystallites interlink with other oriented and misoriented crystallites, accounting for the voids, shown in figure 4. Based on these results, Johnson developed the 3-D schematic representation of the microstructure of PAN-based carbon fiber shown in figure 5.

Endo (ref. 15) also conducted a crystallographic analysis of a PAN-based carbon fiber, Torayca M46, manufactured by the Toray Company. X-ray diffraction studies revealed no separation of the 100 and 101 peaks and the absence of a 112 peak. He also concluded that these fibers have little 3-D order. Endo found that the average crystallite thickness Lc was approximately 6.2 nm and the average layer spacing (dno2) was 0.3434 nm.

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