The properties of the aluminum alloys, titanium alloys, nickel-based superalloys, polymer-matrix composites, and ceramic-matrix composites that are candidate materials for HSCT (High-Speed Civil Transport) structures and engines may degrade with time at the elevated temperatures associated with the operation of the aircraft. The general degradation mechanisms (i.e., the physical event or chain of events that underlie observed degradation effects) that must be considered include:
microstructural and compositional changes,
time-dependent deformation and resultant damage accumulation,
environmental attack and the accelerating effects of elevated temperature, and
synergistic effects among the above.
The lack of prior service experience with the candidate materials makes the design of advanced high-speed vehicles dependent on predictive models to characterize long-term behavior. The development of such predictive models and their application with some level of confidence requires a thorough understanding of the degradation mechanisms of interest.
Deformation and failure mechanism maps can provide a framework for understanding the time-dependent damage (Ashby, 1983). Mechanisms of property degradation and damage accumulation under complex thermomechanical and environmental conditions, especially those uniquely associated with the operation of HSCT, are discussed in this chapter.
Historically, the chief damage mechanisms for aluminum alloys in aircraft applications are corrosion and fatigue, mechanisms generally associated with an aging fleet (Bucci and Konish, 1994). There is currently a great deal of work being supported by the Federal Aviation Administration, NASA, and the materials and aircraft industries to characterize the effects of corrosion and multiple site fatigue accumulation. Information is sparse concerning damage mechanisms associated with high-temperature applications of aluminum alloys. Potential damage mechanisms include microstructural changes, fatigue, creep, and environmental effects.
Elevated-temperature exposure under applied stress can introduce a number of microstructural changes, including nucleation and growth of new phases; formation of subgrains; changes in dislocation density, configuration, and distribution; and nucleation, growth, and coalescence of microcracks. Many of these changes will occur in the absence of stress, while others can be exacerbated by stress. Cyclic stresses may induce accelerated local overaging (coarsening) of the microstructure (Martin and Doherty, 1976), as well as promote the nucleation of microstructural damage (e.g., microcracks, voids, reinforcement fracture in composites) at various sites in the material. In monolithic materials, grain boundaries, inclusions, and other second phases and microstructural inhomogeneities may provide nucleation sites, while in metalmatrix composites, additional damage sites are provided by the presence of the reinforcement (Singh and Lewandowski, 1993a).
Determination of the most critical degradation mechanisms depends on what properties are important in a particular application. For example, if strength is critical, coarsening of the matrix precipitates during elevated-temperature service will be important; if toughness is critical, grain-boundary precipitation or the development of a precipitate-free zone will be important; and if creep or fatigue are critical, the nucleation, growth, and coalescence of microcracks will be important.
The level of residual stresses in aluminum alloys and aluminum-matrix composites are strongly dependent on processing conditions, heat treatment, and thermal excursions. In monolithic aluminum alloys, changes in the processing conditions (e.g., hot and warm working, cold working, rolling, etc.) significantly affect the mechanically induced residual stresses, while quenching and heat treatment produce thermal residual stresses and the precipitation of second phases which possess lattice parameters different than that of the matrix. The mismatch in lattice parameters produces local residual
stresses that may be reflected in the macroscopic mechanical properties. The level of residual stresses are affected by subsequent thermal exposures and prestraining.
In addition to the residual stresses produced via processing, heat treatment, and straining in the monolithic materials, the introduction of reinforcements with different coefficients of thermal expansion than the matrix alloy induces additional thermal residual stresses as well as the possibility of increased mechanical residual stresses because the reinforcements do not deform in a plastically deforming matrix (Bourke et al., 1993; Liu et al., 1993; Withers and Clyne, 1993). As with monolithic alloys, the levels of residual stress may be affected by subsequent thermal exposures and prestraining (Liu et al., 1993).
Coarsening of matrix precipitates is expected to be the major cause of the loss of yield strength in the 2XXX-series ingot alloys. In addition to many theoretical treatments of the effects of applied stress on precipitate growth and coarsening, there are experimental observations of particular relevance (Weatherly and Nicholson, 1968; Singer and Blum, 1977). For example, Singer and Blum (1977) looked at the effects of various thermomechanical processes on the creep behavior of 2618. They showed that the loss of precipitation strengthening during overaging was accelerated by the application of stress. The early work of Weatherly and Nicholson (1968) showed that aging with applied stress produced a significant acceleration in the rate of coarsening of precipitates in an Al-Cu-Mg system. Extended exposures at the intended service temperatures, even without stress, may degrade the yield strength of the ingot 2XXX alloys. Recent data show degradation of tensile yield strength of ingot 2XXX alloys in the T8-type tempers after 1,000 hours at 135°C (275°F; Angers, 1994). Experience on materials used in the Concorde showed a degradation in such properties in early-generation aluminum alloys tested prior to utilization on the Concorde (Peel, 1994). Coarsening of the dispersed phase in dispersionstrengthened alloys such as X8019 is not expected to be a problem at the temperatures of interest (see Figure 3-1 ).
Various researchers have highlighted the importance of microstructural features on the degradation of toughness because of overaging in both monolithic and metal-matrix composite systems (Lewandowski et al., 1987, 1989; Lewandowski, 1989; Klimowicz and Vecchio, 1990; Manoharan and Lewandowski, 1990). In conventional aluminum alloys such as 2XXX-, 6XXX-, and 7XXX-series alloys, age hardening generally produces increases in strength and an accompanying decrease in fracture toughness, while overaging subsequently recovers the toughness while decreasing the strength (Garrett and Knott, 1978; Chen and Knott, 1981; Lewandowski and Knott, 1985). However, microstructural changes such as gain-boundary precipitation, the development of precipitate-free zones, and changes in the deformation modes have been shown to significantly affect the strength/toughness relationship primarily through a reduction in the fracture-critical properties (Vasudevan and Doherty, 1987). For example, the toughness is not recovered upon overaging in a variety of metal-matrix composites (Lewandowski et al., 1987, 1989; Lewandowski, 1989; Klimowicz and Vecchio, 1990; Manoharan and Lewandowski, 1990) and A1-Li alloys (Vasudevan and Doherty, 1987; Blankenship and Starke, 1991).
Deleterious effects of elevated-temperature exposure have been noted in experimental ingot 2XXX-series alloys aimed at high-speed aircraft applications. As shown in Figure 4-1 , some degradation in both strength and toughness were observed after 1,000 hours at 135°C (275°F). Creep deformation of 2618-T6 has been shown to degrade fracture toughness and increase the amount of intergranular fracture (Bobrow et al., 1991).
Time-Dependent Deformation and Damage Accumulation
Fatigue resistance is degraded by the nucleation, growth, and coalescence of voids or microcracks. The level of residual stresses discussed earlier can reduce the level of applied stresses that can be accommodated in service without the nucleation and growth of microcracks. High-cycle fatigue resistance is sensitive to the nucleation of microcracks at
microstructural inhomogeneities, while fatigue crack growth thresholds are affected by the level of residual stresses and crack tip shielding produced by variations in the microstructure. For example, in Al-Li alloys, microstructural features such as the planarity of slip deformation (Blankenship and Starke, 1991), as well as the presence of grain-boundary precipitates and precipitate-free zones (Vasudevan and Doherty, 1987), have been shown to affect the fatigue properties. In the threshold regime, the intense slip localization and the resulting crack tip bifurcation produces nonplanar fracture surfaces and reduces the local stress intensity for crack growth (Blankenship and Starke, 1991). It has been demonstrated that the reinforcements in metal-matrix composites improve high-cycle fatigue resistance while decreasing both the low-cycle fatigue resistance and fatigue crack growth resistance (Allison and Jones, 1993).
Although a great deal is known about the relationships between creep and microstructure, most authors have restricted their theoretical and experimental studies to secondary (“steady-state”) creep. While there are no theoretical treatments that describe primary creep rates in terms of microstructural features, the effects of alloying, grain size, percent stretch, and extent of aging have been investigated for 2618 and other precipitation-strengthened aluminum alloys.
Creep resistance appears to improve with increasing grain size in precipitation-strengthened alloys. For example, the creep resistance of Hiduminium-RR58 (alloy 2618) improved progressively with increases in grain size (Doyle, 1969a, b). As shown in Figure 4-2 , the creep resistance of 2519-T87 at small strains improved with increases in grain size, in the range from 15 to 250 microns (Alcoa, 1995).
The effect of existing strain (pre-strain) on subsequent creep is different in pure metals and alloys than in particlehardened materials. In pure metals and alloys, cold work improves the creep strength by removing part or all of the primary creep strain. The opposite trend has been observed in the precipitation-strengthened aluminum alloys. For example, Wilson (1973) and Peel (1994) showed that pre-strain in 2618 variants, either before or after artificial aging, has a deleterious effect on creep resistance. Doyle (1969a, b) also showed that while tensile properties of Hiduminium-RR58 were increased by cold work after quenching, the creep resistance was reduced. For microstructures that were composed of a fine precipitate and any of three types of dislocation substructure, such as a uniform dislocation distribution, a cellular distribution, or well-defined subgrains (e.g., in commercial and modified 2024, 7075, and 2618 and an Al-Cr alloy), creep strengths were decreased relative to a condition where only fine precipitates were present (Clauer and Wilcox, 1974).
A similar effect of pre-strain was observed in the dispersion-strengthened powder metallurgy Al-Fe-Ce alloys (Alcoa, 1995). Increasing rolling reductions tended to increase the rate and extent of primary creep, presumably because a nonequilibrium substructure was created. Reheats between rolling passes served to anneal the substructure and reduce primary creep rates.
In general, while applying creep deformation (pre-creeping) at the temperature and stress of interest may reduce the extent and rate of primary creep, pre-straining to large strains at different temperatures and strain rates may have the opposite effect. This may be particularly true in metal-matrix composites where it has been demonstrated that pre-straining produces damage (Liu and Lewandowski, 1988; Brechet et al., 1991; Lloyd, 1991; Singh and Lewandowski, 1993a), the extent of which is very dependent on the reinforcement size, matrix microstructure, and nature of the reinforcement-matrix interfaces (Singh and Lewandowski, 1993a).
Degradation mechanisms due to service environmental interactions that need to be considered for aluminum alloys include corrosion, stress corrosion cracking, hydrogen embrittlement, solid-metal embrittlement, and liquid-metal embrittlement. Susceptible sites within the material, such as grain-boundary regions, interfaces, and precipitates are favored sites for embrittlement by all of the above means. The microstructural features that affect the performance of various monolithic aluminum alloys in such environments were summarized by Holroyd (1989). Grain-boundary precipitation or precipitate-free zones are favored sites for corrosion due to the differences in composition and the resulting electrochemical behavior. A number of the other embrittlement
mechanisms also depend on the matrix temper, making thermal exposure effects on the resulting microstructure relevant. Although very little published work exists on the behavior of metal-matrix composites under various environmental exposures, research indicates that susceptibility to corrosion, stress corrosion, and hydrogen embrittlement are extremely sensitive to the electrochemical behavior of the reinforcement and reinforcement-matrix interfaces, as well as to the microstructural features outlined above (Singh et al., 1992a, b).
Heavy-metal impurities, introduced through the use of recycled scrap containing lead or bismuth, may significantly accelerate the evolution of creep damage and sustained load cracking in aluminum alloys at temperatures ranging from −4–80°C (25–176°F). Sustained load cracking can occur at rates in excess of 100 mm/yr at high stress intensities in 6XXX-series alloys containing levels of lead or bismuth that are within the specification limits for those alloys (Lewandowski et al., 1987, 1992). Of concern to the long-term behavior of those systems is that crack initiation and growth near the threshold value of crack initiation is very dependent on time, temperature, and heavy-metal impurity level, as well as the distribution of the heavy-metal impurities in the microstructure. Such problems have been reported in the electric power industry overseas, the domestic automotive industry, and in the pressure vessel industry.
Summary: Aluminum Degradation Mechanisms
Potential damage mechanisms associated with high-temperature applications of aluminum alloys include microstructural changes, fatigue, creep, and environmental effects.
Elevated-temperature exposure under applied stress can introduce a number of microstructural changes including coarsening of the matrix precipitates (important in strength-critical applications) and grain-boundary precipitation or the development of a precipitate-free zone (important in toughness-critical applications)
Fatigue resistance is degraded by the nucleation, growth, and coalescence of voids or microcracks. High-cycle fatigue resistance is sensitive to the nucleation of microcracks at microstructural inhomogeneities, while fatigue crack growth thresholds are affected by the level of residual stresses and crack tip shielding produced by variations in microstructure.
Creep resistance appears to improve with increasing grain size. Also, cold work reduces creep resistance in precipitation- and dispersion-strengthened aluminum alloys.
Degradation mechanisms due to service environmental interactions that need to be considered for aluminum alloys include corrosion, stress corrosion cracking, hydrogen embrittlement, solid-metal embrittlement, and liquid-metal embrittlement.
The primary results of degradation due to aging in titanium alloys in the temperature range of 200°C (392°F) or less would be possible loss of strength, fracture toughness, and fatigue crack growth resistance at these temperatures. Changes in microstructure (e.g., precipitation of chromium and iron intermetallic phases) during long-term exposure at the service temperature and their effect on these properties is of concern. Creep deformation and rupture at temperatures of 200 °C (392°F) or less is not likely to be relevant. However, establishing the highest recommended use temperature for the selected alloys from the point of avoiding creep deformation and crack growth is a worthwhile goal and will help the designers of supersonic aircraft to assess safety margins and also to optimize the use of these alloys.
Three interrelated problems associated with the complexity of fracture processes in titanium alloys, and the possibility of time-temperature-hydrogen-dependent failure modes, could hinder development of high-strength and high-toughness titanium alloys for HSCT applications. These problems include the effects of microstructural variations on deformation and local fracture, uncertainties in intermediate temperature deformation behavior, and hydrogen embrittlement.
Fundamental fracture relationships—particularly among microstructure, deformation mode, and local fracture resistance—are not sufficiently understood for beta titanium alloys. Because of the large number of different microstructures that can be developed by thermomechanical processing (e.g., variations in beta grain size, alpha volume fraction, grain-boundary alpha formation, alpha morphology on aging, and metastable phases), a better understanding of the relationship between processing conditions, resulting microstructure, and properties in this class of alloys, analogous to what has been developed for aluminum alloys, needs to be established. In addition, the effects of crystallographic texture on properties and property anisotropy need to be better understood.
For example, consider the relationships among tensile ductility, fracture toughness, and grain size. Ductility depends on grain size, and the strength/ductility balance can normally be improved by thermomechanical treatments that refine grain structure. Fracture toughness appears to be less sensitive to changes in grain size; consequently, the strength/toughness balance may be less affected by grain refinement (Kawabe and Muneki, 1993). Initiation and growth fracture toughness
are expected to depend on tensile strength, ductility, and modulus. The relationships between basic mechanical properties and fracture toughness have been represented by strain-controlled micromechanical models (Gangloff, 1994). A detailed analysis of the strength/toughness balance must consider such relationships, including basic understanding of ductility-microstructure relationships, and must involve careful measurement of both initiation and growth fracture toughness. Grain-structure-dependent crack path tortuosity will affect crack growth resistance, but not plane-strain initiation toughness, if the latter is carefully measured.
Studies on deformation and fracture of titanium have focused on either near-ambient or relatively high temperatures. Little data are available concerning the effects of moderate service temperatures and loading rates on alloy behavior; the assumption is that ductility and toughness will increase with increasing temperature within the HSCT range. Concerns include the possibility for dynamic strain aging from solute such as oxygen or carbon to reduce tensile ductility and toughness by affecting microvoid growth and coalescence, the possibility for thermally activated slip localization in locally soft regions of the microstructure, time-temperature effects on hydrogen embrittlement, and deformation and fracture behavior at cryogenic temperatures.
The dissolved hydrogen content of HSCT-candidate titanium alloys could increase during processing, component fabrication, or elevated-temperature service in aggressive environments, such as airplane hydraulic fluid, and subsequently degrade tensile ductility and fracture toughness.
Although beta titanium alloys are believed to be tolerant of dissolved hydrogen (Eylon et al., 1994), additional study is required to verify their resistance to hydrogen effects. Brittle hydrides do not form in body-centered-cubic beta alloys due to exposure in hydrogen-producing environments such as modest pressure gases and electrolytes. The diffusivity of hydrogen in beta is relatively rapid (similar to that in high-strength martensitic steels), even at 25°C (77°F; Schutz, 1993). In contrast, the alpha phase is susceptible to both internal hydrogen embrittlement (IHE) and hydrogen environment embrittlement (HEE), as demonstrated by a variety of studies on alpha-beta alloys such as Ti-6Al-4V and Ti-6Al-6V-2Sn (Nelson, 1974; Lucas, 1990; Moody and Costa, 1991). Hydrogen is relatively insoluble in alpha, with a hydride phase possible, and is slow-diffusing compared with diffusion in beta (Schutz, 1993). Hydrogen effects are likely to be microstructure sensitive; either primary or precipitated alpha in the beta matrix may reduce hydrogen tolerance, particularly if this phase is arrayed in a continuous path along beta-grain boundaries or if this phase promotes locally intense slip.
HSCT-candidate alloys may be prone to hydrogen embrittlement. Such alloys are generally higher strength than the alpha-beta and beta titanium alloys that have been examined to date; high strength exacerbates hydrogen embrittlement (Gangloff, 1988). Based on results for alpha-beta alloys with lower concentrations of beta-stabilizing elements, Ti-6-22-22 could be prone to hydrogen embrittlement. Metastable beta alloys such as Beta-21S, Ti-15-3, and Beta-C are embrittled by predissolved hydrogen during subsequent, relatively rapid-rate loading at 25° C (77°F; Young and Scully, 1993, 1994). The fracture strain of both peak-aged and lower-strength solution-treated alloys was degraded by dissolved hydrogen. Intergranular and a variety of transgranular cleavage and slip-band cracking modes were promoted by hydrogen, without evidence for hydriding of either the alpha phase or the continuous beta matrix. Several factors exacerbated IHE, including higher hydrogen concentrations (between 100 and 1,000 ppm by weight), triaxial constraint, as well as grain-boundary alpha colonies or locally intense planar slip. IHE and HEE susceptibility is likely coupled (Nelson, 1974). Peak-aged Beta-21S and Beta-C are susceptible to intergranular environmental crack propagation in aqueous chloride (or chloride stress corrosion cracking) when subjected to actively rising stress intensity levels at about one-half of the planestrain fracture toughness (Grandle et al., 1994; Young et al., 1995). This behavior mirrors the well-known brittle cracking of alpha-beta alloys such as Ti-6Al-4V and parallels the IHE behavior (Nelson, 1974).
High-strength or high-toughness HSCT-candidate titanium alloys may be prone to HEE in typical aircraft environments. Low-cost beta and Beta-CEZ alloys have not been examined to assess their susceptibility to hydrogen embrittlement, including the effects of temperature, loading rate, and the effects of moderate hydrogen concentration. The HEE resistances of low-cost beta, Beta-CEZ, and Ti-6-22-22 are not characterized. While most IHE and HEE studies emphasize brittle cracking under plane-strain conditions, the range of HSCT component geometries dictates consideration of adverse hydrogen effects on both plane-strain and plane-stress deformation and fracture processes.
Summary: Titanium Degradation Mechanisms
Three primary factors need to be considered in long-term, elevated-temperature applications of high-strength and high-toughness titanium alloys for HSCT applications, including:
the effect of microstructural variations (e.g., variations in beta grain size, alpha volume fraction, grain-
boundary alpha formation, alpha morphology on aging, and metastable phases) should be developed, analogous to what has been developed for aluminum alloys,
the effects of moderate service temperatures and loading rates on alloy behavior, including the possibility for dynamic strain aging, thermally activated slip localization, time-temperature effects on hydrogen embrittlement, and deformation and fracture behavior at cryogenic temperatures, and
the degradation of tensile ductility and fracture toughness due to dissolved hydrogen from aggressive environments such as airplane hydraulic fluid.
The degradation mechanisms that affect nickel-based superalloys include overaging of microstructure, fatigue, creep, and oxidation. The trade-off that needs to be made among the competing properties is difficult, as whatever is done to improve one property usually affects all of the other properties. An example of this is shown in Figure 4-3 , in which several key properties are shown to vary with one of the controllable parameters in the microstructure grain size. There are many controllable parameters such as grain size, microstructure, and concentration of major and minor alloying elements that can be optimized to yield the balance of properties.
Overaging of the superalloy microstructure depends on time, temperature, and stress. All of the nickel-based
superalloys used in gas turbine engines are metastable under use conditions. Characterization of selected alloys must be performed to define the safe-use limits of the alloy for the engine-operating environment. Creep experiments are the primary mechanical characterization test to determine the limits for an alloy.
The strength of nickel alloys is derived primarily from the particles that precipitate in the alloys on cooling. These particles are generally ordered arrays of Ni3A1, although cobalt can substitute for nickel and niobium, and titanium and tantalum can substitute for aluminum (the general chemistry of these particles is given as [Ni, Co]3[Al, Ti, Nb, Ta]). The alloys undergo heat treatment to obtain a uniform array of particles of a specific size and spacing to achieve the maximum strength and stability. However, at high temperatures and stresses the larger particles grow while the smaller ones dissolve —a process known as Ostwald Ripening—which in time change both the structure and behavior of the grains. This process can be slowed by altering the balance of elements that form particles and adding large atoms that diffuse slowly in the matrix, slowing down the entire coarsening process. The problem with this approach is that new flat and brittle phases, called topologically close-packed (TCP) phases, can form and cause strength reduction at grain boundaries. A primary goal of most alloy development is to avoid the formation of TCP and to develop as stable a microstructure as possible with the many constraints imposed on the process.
Deformation and Damage Accumulation
Superalloys develop damage from cyclic applications of stress and strain. The most prevalent sources of mechanical damage are low-cycle fatigue and creep deformation. Both processes are at work in most turbine applications.
Creep failure is generally not a problem in turbine or compressor disks, but creep deformation can be a tremendous problem. As the disk creeps during operation it can redistribute stresses, allowing accelerated fatigue in other locations. Also, creep deformation of an engine disk can cause compressor or turbine blades to rub against the outer case. The allowable creep in a turbine disk is very small, generally less than 0.2 percent creep in the life of the disk. Creep, in conjunction with stress and oxygen effects at the tip of a crack, can result in dramatic accelerations in crack growth rates. This is referred to as dwell crack growth, creep crack growth, or time-dependent crack growth.
Oxidation in nickel-based alloys or structural castings is more than simple surface oxidation, which in many ways would be an easier problem to solve. Since the oxidation rates
for grain boundaries are far greater than the rate for bulk oxidation, grain boundaries usually serve as sites for oxidative degradation.
Over the years the reduction or elimination of grain boundaries via directional solidification and single-crystal processing has greatly extended reliability of blade alloys. However, single-crystal and directionally solidified alloys are high in cost, have directional properties, and exhibit a large reduction in low-temperature strength and fatigue properties compared with conventional alloys.
Another approach to improve oxidation resistance is to include alloying additions that inhibit the embrittling effects of oxidation on the boundaries to the basic alloy chemistry of elements. The alloying elements that have the greatest effect are those that partition primarily to the grain boundaries, including boron, carbon, hafnium, and zirconium. These elements are generally added in very small quantities, but assuming that they partition solely to the grain boundaries, they may be the primary elements locally at the boundary. The challenge in reducing grain-boundary oxidation is to add just enough of the correct alloying elements to inhibit oxidation while maintaining mechanical properties.
Despite the myriad of composite materials and structural arrangements, there are a limited number of mechanical degradation mechanisms (NRC, 1991). These include:
local compressive instability, and
Any one of these damage modes may be small as regards the representative volume element of the structure in question and therefore viewed as subcritical. However, the combined effects of several damage modes may radically alter this picture. The complex coupling of any number of these damage modes is poorly understood. Furthermore, the entire process will be strongly influenced both by material types and structural lay-up.
For elevated-temperature applications such as the HSCT, the effects of various environmental factors must also be accounted for in addition to the mechanical degradation modes. Continuing damage accumulation is induced and driven by combined cyclic loads, high-temperature exposure, oxidative attack, solvent infusion, moisture, and other factors. The coupling process linking the growth of various damage modes and the external environmental drivers will undoubtedly prove complex.
Damage mechanisms to consider for elevated-temperature composite applications include thermal oxidation, hygrothermal (combined moisture and temperature) effects, matrix cracking, and microstructural changes. The evaluation of the degradation of polymer-matrix composites is complex, requiring not only an integration of many contributing factors, but also an assessment of poorly understood synergistic accelerations in damage accumulation as driven by external factors. Moisture absorption, high-temperature exposure, heating (and cooling) rates, and loading rates represent a number of those factors that affect basic composite properties such as toughness, glass transition temperature, and strength. The combined influence of such factors on the failure mechanism may be pivotal in deriving any reliable modeling process.
High-temperature thermosetting polymers such as bismaleimides commonly incorporate discrete toughening phases, typically 25–35-weight percent of a soluble thermoplastic polyimide in granular or particulate form, to improve impact resistance. Isothermal air aging studies at temperatures above 150°C (302°F) have shown a tendency for these systems to undergo phase separation that is clearly distinguishable under examination in the scanning electron microscope. Using a common optical microscope, distinct color differences can be seen between the two phases, with both phases showing considerable darkening as exposure time increases. This phase separation progresses from the surface to the interior and is assumed to be driven by differences in oxidation rates of both the bismaleimide and the added toughening agent. Phase separation has a deleterious effect not only on matrix toughness, but can also lead to matrix cracking since phase boundaries represent discontinuities with associated stress concentrations. This combined degradation phenomenon can be seen generally after 5,000 hours of exposure and represents a valid concern for long-term HSCT applications.
Damage Accumulation: Matrix Cracking
Perhaps the most critical damage mechanism operating in high-temperature polymeric composites is the formation of transverse ply cracks (shown schematically in Figure 4-4 ) and in-plane microcracks in matrix polymers of multiaxial composites. Matrix cracking can result from initial laminate processing, mechanical static, and fatigue loading (Reifsnider and Giacco, 1990), residual stresses resulting from hygrothermal exposures, thermal cycling, and combined effects of mechanical and environmental cycles (Sensmeier, 1994). As the cycles (mechanical or thermal) advance, matrix cracks become more pronounced and increase in density until a saturation level is reached (NRC, 1991).
Fiber anisotropy 1 and differences between thermal expansion coefficients of the matrix and fiber can result in residual thermal stresses during processing or exposure to temperature. These stresses can cause fiber-matrix interfacial failure or radial cracking (in the matrix radiating from fiber surfaces).
The fundamental cause of thermally induced transverse matrix cracking in a laminate is the residual stresses resulting from differences in thermal expansion of the lamina in longitudinal (parallel to the fiber orientation) and transverse (perpendicular to fiber orientation) directions (αL ≠ αT). Hygrothermal (combined moisture and temperature) cycling of composite laminates can produce transverse matrix cracks that initiate in surface plies and progress deeper into the laminate with accumulating cycles. The rate and severity of transverse matrix cracking is dependent on several conditions, including matrix properties, fiber properties (especially thermal expansion and stiffness), processing conditions, service conditions (including temperature cycle and humidity), ply thickness, and ply stacking sequence.
Composite strength, stiffness, and thermal properties as well as failure modes can be affected by transverse matrix cracking. Figure 4-5 shows a reduction in stiffness as a function of crack density for a cross-plied carbon/epoxy laminate. Glass/epoxy cross-plied and pseudoisotropic laminates made from nonwoven unidirectional fabric and laminates made from 181-style cloth had stiffness reductions of 15–20 percent due to microcracking caused by subjecting the laminates to cyclic tension loading (McGarry and Wilner, 1968).
In addition to damage-induced reduction of matrix-dominated mechanical properties, the advancing microcracking promoted higher uptake of moisture deeper in the laminate (Carey, 1957; Anderson and Healey, 1958; Epstein and Bandaruk, 1964).
Thermal Degradation and Oxidation
Exposure to high temperatures can cause chemical degradation of composite-matrix polymers. Thermal degradation usually implies chemical reactions associated with polymer chain scission as a result of temperature or diffusion of small molecules (e.g., O2). Degradation of composite-matrix polymers will more likely be due to thermal instability and the accelerating effects of oxidative attack.
Polymer degradation in an inert atmosphere is the result of the breaking of covalent bonds in the polymer network. The result is a reduction in molecular weight and eventually volatilization of low-molecular-weight fragments (void growth). This process is important and is used positively in the production of polymer-precursor carbon and ceramic materials. The susceptibility of a polymer to thermal degradation is a function of the bond strength (Rodriguez, 1970). The thermal stability of high-performance polymers in an inert atmosphere are generally very good due to their highly crosslinked and aromatic nature.
Unfortunately, most high-temperature aircraft applications involve exposure in oxidizing environments (e.g., ambient air). While both carbon fibers and matrix polymers are susceptible to oxidative degradation, the degradation of the fibers is negligible at the service temperatures envisioned for application of polymeric composites (Magendie et al., 1990). However, oxidation of high-temperature matrix
Coefficients of thermal expansion of fibers (such as carbon or polymeric fibers) are different in the longitudinal and transverse directions.
polymers must be considered in composite design and characterization.
HSCT-candidate systems have exhibited varied susceptibility to oxidation (Brunner, 1994). Matrix systems that have been evaluated include polyimides (Bowles et al., 1993), bismaleimides (Stenzenberger et al., 1976; Magendie et al., 1990), cyanate esters, benzocyclobutene, and thermoplastics such as polyarylenes (Arnold and Maskell, 1991). Oxidation is usually characterized in terms of weight loss. In general the candidate materials can be ranked in terms of oxidation resistance from most stable to least stable: polyimides ≈ polyarylenes ≫ bismaleimides ≈ cyanate esters ≫ epoxies (Hipp et al., 1993).
Microstructural studies have shown that oxidation of high-temperature polymers is largely a surface phenomena (Nam and Seferis, 1992). Oxidation reactions in the interior relies on mass-transport processes, including diffusion of oxygen inward and subsequent diffusion of reaction products outward.
In carbon-reinforced composites, oxidation of matrix polymers is strongly influenced by geometry and ply orientation. Matrix oxidation is dominated by the exposed surface and the fiber-matrix interface, progressing inward from
exposed edges along the fiber direction as shown in Figure 4-6. The nature of the fiber-matrix interphase region and the influence of fiber coatings or sizings are critical in limiting oxidation of the composite. The presence of significant matrix cracking makes more surface available and can greatly accelerate oxidative degradation throughout the laminate.
Composite mechanical properties, especially interlaminar and interfacial shear, and flexural have been shown to be seriously degraded due to oxidative damage (Bowles, 1993; Hipp et al., 1993). Much more work needs to be done to understand threshold damage conditions and the effects of low levels of oxidation on composite performance.
The detrimental effects of moisture, particularly at high temperature, are often underestimated. Since the projected flight cycle for the HSCT reaches hot, dry conditions during the supersonic cruise phase of each flight, many believe that HSCT applications can be considered to be dry structure for evaluation purposes. However, although the time to moisture saturation is long, all structures will carry a moisture load when entering flight. On a global basis the moisture level may be low, but the surface plies could carry significant moisture levels. Concentration gradients (Figure 4-7 )—from high moisture levels at the surface to near dryness at the midplies —along with subsequent desorption in the near-surface plies and redistribution of moisture during flight can lead to significant residual stresses or outer-ply delamination or blistering during rapid heating. The effects of moisture are dependent on many factors, including thickness, material type, prevailing humidity, flight temperature profile, time and conditions between flight times, and presence of matrix cracking. Potential long-term hygrothermal effects include microvoid
generation and reduction in intrinsic mechanical properties due to matrix plasticization. Potential short-term hygrothermal effects include matrix cracking, outer-ply delamination, or surface blistering during rapid heating.
Although it is well known that moisture plasticizes matrix polymers, with an apparent increase in toughness, studies have shown that repeated cyclic exposures designed to simulate flight histories can result in growing microcracking initiated at the surface as shown in the micrograph in Figure 4-7. Moisture has been shown to increase the residual stresses that can lead to matrix cracking in many systems, including carbon/polyimide, carbon/bismaleimide, and aramid/epoxy. As described in previous sections, matrix cracking can increase the composite's vulnerability to other damage accumulation mechanisms, including fatigue loading, chemical attack by fluids, and changes in mechanical or chemical properties of the matrix or interface. McGarry and Wilner (1968) showed that for cross-plied glass/epoxy laminates the weight gain resulting from one day water immersion increased significantly as a result of microcracking caused by tension fatigue (50–75 percent weight gain increase for 30 percent cyclic stress and 100–150 percent weight gain increase for 50 percent cyclic stress).
Simple hygrothermal cycling, with no externally applied loads, has been shown to cause a multiplicity of damage mechanisms for composites including (1) alteration or loss of intrinsic properties, (2) microcracking, (3) microvoid generation, (4) delamination, and (5) surface blistering.
For example, hygrothermal exposure of thermoplastic polyimide resins —AFR700B and Avimid®N—intended for very high-temperature service (up to 371°C [700°F]) leads to significant loss in glass transition temperature (Tg) (>50°C [90°F]). Contrary to previous expectations, neither matrix system offers a stable Tg nor 371°C (700°F) performance capability. Adopting a strict adherence to MIL-HDBK-17B standards, the safe material-operating limit would be less than 260 °C (500°F) in spite of the fact that the initial, dry Tg was greater than 400°C (752°F). As the wet Tg decreases, approaching the upper exposure cycle temperature, the elevated internal moisture vapor pressures can be sufficient to generate both interior microvoids and small surface blistering.
Amorphous, high-temperature thermoplastic polyimides, such as Avimid N and AFR700B, that are subjected to a post-cure thermal treatment to attain very high Tg are glassy polymers in an energy state above thermodynamic equilibrium. The basic, intrinsic properties of both resins can change significantly under hygrothermal exposures. Exposures at temperatures well below the dry Tg results in a cumulative, permanent reduction in Tg. Matrix-dominant mechanical properties (e.g., shear modulus, compressive strength, and transverse tensile strength) that are closely coupled to the Tg, particularly at elevated temperatures, can be significantly affected (Cornelia, 1994). High-temperature humidity exposures effectively plasticize the matrix, accelerating polymer relaxation processes toward thermodynamic equilibrium. As the Tg decreases, the acceptable upper limit of the service temperature also decreases. Under similar conditions another thermoplastic polyimide, Avimid K3B, shows no alteration of intrinsic properties. Although bismaleimide systems have not exhibited this behavior, highly crosslinked systems were not totally immune; the polyimide PMR-15 shows significant loss of Tg (20–25°C [36–45°F]) after a modest number of hygrothermal cycles. Morphological and crystalline character have a strong influence on matrix-dominant properties. The effects of hygrothermal exposures on matrix morphology must be considered when determining acceptable performance.
Summary: Polymer-Matrix Composite Degradation Mechanisms
Damage mechanisms to consider for elevated-temperature composite applications include thermal oxidation, hygrothermal (combined moisture and temperature) effects, matrix cracking, and microstructural changes.
High-temperature thermosetting polymers commonly incorporate discrete toughening phases to improve impact resistance. These systems can undergo phase separation that has a deleterious effect, not only matrix toughness, but it can also lead to matrix cracking.
Perhaps the most critical damage mechanism operating in high-temperature polymeric composites is the formation of transverse ply cracks and in-plane microcracks within the matrix of multiaxial composites. Matrix cracking can result from initial laminate processing, mechanical static and fatigue loading, residual stresses resulting from hygrothermal exposures, thermal cycling, and combined effects of mechanical and environmental cycles. Strength, stiffness, and thermal properties as well as failure modes are affected by the mechanical and chemical degradation of composites (Greszczuk, 1988). Generally, stiffness, strength, and coefficients of thermal expansion decrease with increases in microcrack density.
Exposure to high temperatures can cause chemical degradation of composite-matrix polymers. Degradation of composite-matrix polymers are generally due to thermal instability and the accelerating effects of oxidative attack. In carbon-reinforced composites, oxidation of matrix polymers is strongly influenced by geometry and ply orientation.
Hygrothermal effects can lead to significant residual stresses, outer-ply delamination, or blistering during rapid heating. The effects of moisture are dependent on many factors including thickness, material type, prevailing humidity, flight temperature profile, time and conditions between flight times, and presence of matrix cracking. Absorbed moisture can increase the residual stresses and lead to matrix cracking, whereas microcracking can accelerate moisture absorption and desorption.
Since the candidate ceramic-matrix composite materials are just now being developed, there is very little direct information about their durability. Therefore, the durability of the constituent materials is considered here to identify potential degradation mechanisms for these composites. This section includes a review of the degradation of the monolithic materials, a discussion of the degradation of the high-temperature fibers, and a discussion of the interphase (material at the matrix-fiber interface) durability.
Recent reviews have summarized the high-temperature stability of non-oxide structural ceramics (Jacobson, 1993; Tressler, 1993). The review by Jacobson concentrated on the corrosive degradation of silicon-based (i.e., monolithic silicon carbide [SiC] and silicon nitride [SiN]) ceramics in combustion environments. Figure 4-8 shows the major types of corrosive attack as a function of reciprocal temperature. The passive region refers to the formation of a protective SiO2 film that grows by diffusion of oxidants through the growing scale. Active oxidation refers to material removal by the formation of gaseous SiO(g). Reaction of the SiO2 with the substrate to form volatile SiO(g) only occurs at the very highest temperatures. Molten salt deposition (e.g., Na2SO4) can occur at lower temperatures and accelerate the passive oxidation by dissolving the SiO2 film, leading to rapid rates of material removal and often to selective pitting of the substrate (Fox et al., 1990). Aerodynamic conditions can affect these rates if the corrosive product is swept off the surface into the gas stream, leaving a fresh, reactive surface or a very thin protective film.
Current research suggests three major areas of concern with silica films:
Influence of impurities (Zheng et al., 1992; Opila, 1994). It appears that small amounts of such elements as sodium and aluminum accelerate the growth rate of silica by an order of magnitude or more.
Ease of reduction of SiO2 to SiO(g). The formation of SiO(g) via a classical, active oxidation mechanism only occurs under extreme conditions. However, there are many situations where both oxidizing and reducing gases are present (e.g., fuel-rich combustion in the HSCT). Under these conditions, SiO2 can form on a silicon-based component, but it is readily reduced to SiO(g). Recent work at NASA Lewis has demonstrated this both in furnace and burner tests (Opila et al., 1994; Opila and Jacobson, in press).
Formation of silicon hydroxy species. It is well known that in steam atmospheres SiO2 readily hydrates to form
Si(OH)4(g) and Si2O(OH)6(g) (Brady, 1953; Kashimoto, 1992). High-pressure combustion environments contain substantial amounts of water vapor, and this volatilization has recently been observed at NASA Lewis in combustion atmospheres (Opila et al., 1994).
These three issues have major implications in the application of SiO2-protected materials to the HSCT. Only limited data are available to delineate the extent of each phenomena.
Because of the above effects, as well as the susceptibility of silica to basic molten salt attack, there has been an effort by several groups to develop refractory oxide coatings for silicon-based ceramics (Lee et al., 1994, 1995). These systems show the promise of combining the desirable physical and mechanical properties of a silicon-based system with the better durability of a refractory oxide. Mullite coatings on SiC show improved corrosion resistance. However, mullite contains silica at a fairly high activity. Recent developmental work suggests it may be possible to go from mullite to an even more refractory coating.
Non-oxide fibers are also subject to the same thermochemical degradative processes if they are exposed to the environment. In addition, the polymer-derived fibers undergo microstructural changes at elevated temperature which lead to altered mechanical properties (Bodet et al., 1993). Specifically, in the Nicalon fiber the Si-O-C phase decomposes to yield additional SiC particles and SiC grain growth with CO being a gaseous decomposition product. Encapsulation in an impermeable matrix or exposure at high temperature in a CO ambient suppresses this behavior, but in general the fiber degrades under long-term exposure.
The most critical chemical durability issue involves the interphase material. The carbon interphase can be actively gasified in an oxidizing ambient atmosphere, leaving
ultimately a SiC/SiC interface that is too strong. Under the highest temperature conditions the interphase channel can fill with SiO2 so that the removal of the interphase is limited to the near-surface region (Filipuzzi and Naslain, 1994). In service, the matrix will be cracked, providing points of interphase degradation.
The most common fiber coatings—carbon and boron nitride—are currently unacceptable because of their susceptibility to oxidation. The need for alternate fiber coatings is the major issue in the development of fiber-reinforced ceramicmatrix composites in their present form. Many efforts have been directed to this end, but the requirements of fiber pull-out (to provide toughness) and inertness with silicon (which is one of the most reactive elements) limit the possibilities. The issue of a microstructurally stable fiber is being pursued with the development of fully crystalline SiC fibers.
Mechanical Degradation: Effects of Corrosive Reactions
In addition to the purely thermodynamic and kinetic aspects of the corrosion reactions, a recognition is emerging that the selective attack of certain microstructural features can lead to nucleation of strength-limiting flaws. The time-dependent properties of slow crack growth and creep leading to creep damage are altered because of grain-boundary effects caused by reactions with the environment. Examples are the effects of oxidation in the high-temperature creep and fracture of βSi3N4 ceramics and the effects of oxidation on crack growth in sintered SiC. For temperatures below 1300–1400°C (2372–2552°F), crack growth is the key issue in non-oxide ceramics, particularly sintered and chemical-vapor-deposited SiC that are extremely creep resistant.
The primary parameter of concern for long-term reliability is the threshold stress intensity, Kth, above which crack growth will occur. Very short lifetimes of tens of hours are predicted from crack growth rate (V) versus stress intensity factor (K) data for components stressed above the threshold K value. Thus, environmental effects on the threshold K value are of critical importance in engineering design. For sintered α-SiC with boron and carbon-sintering additives, either preoxidation or testing in air caused a substantial reduction in the estimated value of Kth at 1200°C (2192°F) and 1400°C (2552°F; Minford and Tressler, 1983; Minford et al., 1983). In general, oxidation has a deleterious effect on crack growth in SiC materials. Recent studies of the effects of interphase oxidation and embrittlement of SiC/SiC composites with a carbon interface (e.g., Filipuzzi and Naslain, 1994) suggest that a carbon interphase is simply not going to be acceptable for long-lived composites.
Fibers that exhibit excessive creep or undergo slow crack growth can also be life limiting as observed for monolithic SiC (Tressler and DiCarlo, 1993). Rugg and Tressler (n.d.) have demonstrated the effect of crack growth in the stressrupture behavior of a developmental sintered SiC fiber. The limiting aspect of the c-axis sapphire fibers as load-bearing members is slow crack growth which leads to unacceptable stress-rupture characteristics above ~ 1000°C (1832°F; Newcomb and Tressler, 1993).
Thus, although the phenomena that lead to degradation of the individual constituents of current ceramic-matrix composites are beginning to be understood, the study of the durability of any given composite system is just now beginning. Accelerated aging evaluations can be based only on degradation mechanisms for the constituent materials, not a detailed understanding of a composite system.
Summary: Ceramic-Matrix Composite Degradation Mechanisms
Based on the durability of the constituent materials, the most important degradation mechanism for ceramic-matrix composites for HSCT applications is thermochemical degradation (including corrosion and oxidation reactions) in combustion environments. The selective attack of certain microstructural features of the composite can lead to the initiation of flaws that affect slow crack growth and creep deformation processes.
The development and application of aging characterization tests and predictive models require a thorough understanding of the important basic materials properties and degradation and damage accumulation mechanisms to provide confidence in their ability to accurately assess service life. Also, an understanding of the damage mechanisms and affected mechanical properties can provide important criteria for the developers of high temperature materials so that material candidates can be improved. The evaluation of basic materials properties and definition of critical degradation mechanisms for these materials is, at present, incomplete. The scarcity of available property data makes the analysis of subsequent test results to determine aging responses challenging.
The committee recommends that NASA emphasize fundamental work to characterize materials properties and, most important, to determine critical degradation and damage accumulation mechanisms over a broad range of environmental conditions. This work should allow determination of the mechanisms that are likely to have the greatest influence on component durability.