Microstructure and Strength of Al₂O₃ and Carbon Fiber Reinforced 2024 Aluminum Alloy Composites

The re-melting of delivered extruded Al2024 alloy and its subsequent squeeze casting resulted in development of cast materials microstructure with characteristic segregation of precipitates at the grain boundaries, prevailing formation of η₂-AlCu phase, Al₇₇Cu₁₇Mn₂Fe₃ intermetallic compounds, and α/S-Al₂CuMg eutectic mixture (Fig. 1). Generally, the precipitates in composite material are evidently refined due to the presence of large number of places available for heterogeneous nucleation. A lower volume of molten metal due to the presence of solid reinforcement released less latent heat during solidification. It causes a faster solidification of infiltrating metal inside preform leading to higher matrix saturation with copper. Therefore, even without homogenization heat treatment, the composite material may show a small positive strengthening response to aging similar to T6 treatment for its Al2024 aluminum alloy matrix.

The homogenization treatment at slightly higher parameters as compared with standard T6 conditions for Al2024 (increased temperature to 515 °C and longer time −1.5 h) enables to dissolve more Cu2Al intermetallic phases segregated at grain boundaries. Only the largest particles in the vicinity of reinforcement remain un-dissolved probably due to its stabilization with small amounts of oxygen transferred from the fibers (Fig. 2a). Outside of this area at the binder/matrix interface, MgO forms according to reaction SiO₂ + 2 Mg → 2MgO + Si (−64 kcal/mol). In places where fibers were covered with thinner layer of binder, the latter was completely dissolved and fine MgO crystalline precipitated directly on the fiber surface (Fig. 2b). However, such prolonged heat treatment results in substituting of most of the amorphous SiO₂ bridges and connections between fibers with a mixture of dense nanocrystalline materials including the Θ′ phase and almost pure silicon (see Fig. 3).

Furthermore, density of intermetallic phases was found higher near the matrix/fiber interfaces. Thermal mismatch of the components induces thermal-misfit dislocations at the fiber/matrix interface which can act as heterogeneous nucleation sites for precipitates. Summarizing, the amorphous SiO₂ binder was partly dissolved and acted as an oxygen source substituted partially by MgO, but the more brittle MgAl₂O₄ spinels were not detected. The subsequent age hardening caused the formation of the Θ platelet phase from saturated matrix at the {100} aluminum planes. It should be also noted that dissolution of amorphous silica leads to formation of microvoids, which can serve as crack nucleation sites.

To improve wear properties at dry conditions, preform was supplemented with small amount of C fibers. In such case, carbon fibers should act as lubricant pockets providing a thin carbon film separating working parts and protecting them against seizure. During infiltration of carbon fiber with aluminum matrix, they may react leading to formation of Al₄C₃. The latter could cause fast deterioration of the fiber/matrix interface and induce a high stresses within a composite. In the case of composite reinforced with fiber matting, a large gap between matrix and the fiber bundles was observed (Fig. 4a). Usually, delamination and cracking of C fiber occurred already shortly after casting or during subsequent machining of the specimens. Thermal mismatch of components caused large stresses which in the case of C fibers relates on their direction. Depending on C fiber type, the mean transverse CTE varies from 5 × 10⁻⁶ to 10 × 10⁻⁶/K and the longitudinal one from 1.6 × 10⁻⁶ to 2.1 × 10⁻⁶/K (Ref 6). Therefore, after cooling, fibers are under compression (CTE for the matrix is ca. 23 × 10⁻⁶/K), with large shear forces produced at the interface. Additionally, in the case of fiber mat, separate bundles, arranged normal to each other, act alternately in tension and compression. Similar deterioration of microstructure and mechanical properties caused by the high residual porosity and cracks originated from shrinkage was observed in carbon fiber multilayer composites (Ref 7).

Infiltration of fiber preform requires preheating of reinforcement almost to temperature of molten metal. In squeeze casting method placed, preform in die is exposed to surrounding atmosphere. C fibers, being in contact with surrounding atmosphere, start to burn at temperature of over 560 °C. Thus, in contact with molten aluminum alloy, CO or CO₂ gases can be produced and trapped at the interface. This phenomenon was occasionally observed during carbon bundle infiltration. Usually complete solidification lasts from 15 to 20 s, so during that time, the remnants of oxygen can react and produce indicated bubbles. Near the C-fiber/aluminum matrix interface within large bubbles, a presence of Al₃C₄ carbides was occasionally noticed (Fig. 4b). It looks that at the beginning of squeeze infiltration process, molten metal came into contact with C fiber, reacted producing aluminum carbide which was next detached by newly formed gasses. It was determined (Ref 8) that in the case of carbon, wires infiltrated and exposed to the liquid aluminum already after one second carbides may be formed in rapidly increasing quantity. Fortunately, the above phenomenon was observed rather rarely. Mapping of composites reinforced additionally with carbon fiber has shown similar intermetallc distribution, though reaction of SiO₂ binder with Mg can also result in precipitation of Mg₂Si, see Fig. 5

Solidification of metal under pressure increases liquidus temperature by 6-8 K/100 MPa (Ref 9), and as a result of undercooling, a finer microstructure is produced. Furthermore, during infiltration of preform at the temperature of 500-600 °C, metal quickly nucleates on fiber surfaces. Therefore, crystals inside the strengthened region are 5-10 times smaller than in the unreinforced area. Preform acts as a rigid skeleton restraining metal flow and after solidification contributes stresses due to thermal mismatch between reinforcement and Al alloy matrix. Preform coefficient of thermal expansion, being comparable with that characteristic for strengthening fibers, amounted 7-10 × 10⁻⁶/K.

During cooling, it shrinks much less than the matrix, what causes tension stresses up to 35÷57 MPa as determined by Ref 10. During subsequent heating, the matrix expands and deforms, annihilating most of the strain. Therefore, tensile stress in the matrix and compressive stress in fibers at about 80-100 °C are fully reduced. Further heating, reverses tendency and expansion of the matrix with higher expansion coefficient (20-24 × 10⁻⁶/K) will be restrained by the fibers. As a result, at about 200-250 °C, local stress can achieve matrix yield point. Then the weakest areas located in the vicinity of the fibers bundle or at the fiber-matrix interface can crack even under low load during strength examinations.

Beside thermal stresses, lattice deformation with high density dislocations induces precipitation of Θ′ phase even in the as-cast matrix, what was also observed in composite materials AlCu4MgAg/SiC (Ref 5). Thus, hardness of the as-cast composite is relatively high and subsequent heat treatment only slightly enhanced its value, see Table 3. Similarly small reinforcing effect was observed in the case of Young modulus and the elongation. Alumina fibers slightly reduced the ductility of aluminum matrix presumably due to grain refinement and the strong stresses induced by the mismatch in thermal expansion between the matrix and fibers

Strength investigations performed on age hardened composite materials show slight increase of Young modulus and distinct reinforcing effect for bending and tensile strength properties. In relation to the unreinforced alloy, the bending strength of composite materials strengthened with 10% of Al₂O₃ fiber increased by ca. 100 MPa at the 20-300 °C temperature range, see Fig. 6. Higher, i.e., 20% volume content of fiber only slightly enhances this effect.

The increasing fiber volume from 10% to 20% resulted in increase of UTS up to ~200 MPa, see Fig. 7. Strength of squeeze cast 2024 alloy was ca 230 MPa, comparable with UTS = 250 MPa obtained by Hajjari and Divandari (Ref 11). The best result was achieved at the temperature of 300 °C, at which strength of composite with 20% of fiber was ~100% higher than that of the matrix. This corresponds to 384 MPa calculated from the rule of mixture with efficiency factor related to fiber orientation and length η = 0.5. This confirmed proper shear strength of the interface approximately equal to the matrix strength.

Observations of specimen fractures after tensile and bending examinations revealed significant amount of cracks in the matrix, adjacent to the fibers, or even propagating transversely or along fiber axes. Thus, process can be intensified by interdendritic porosities which, in strongly intertwined fiber network, are difficult to feed with liquid alloy, see Fig. 8(a). Existed in Saffil fiber α-Al₂O₃ plates occurred close to microvoids, meaning that they frequently initiate crack and premature material fracture, see Fig. 8(b).

Basing on the shear-lag theory, the ratio of the maximum tensile stress in a fiber, σ_max, and the maximum shear stress at the end of fiber, τ_max, is described by equation (Ref 12):where G _m is the matrix shear modulus, E _f is the fiber axial modulus, and V _f is the volume fraction of fibers. If the shear strength of the interface is higher, and tensile strength of fiber is smaller than the maximum stresses, then interface should withstand the local loading. However, crack may develop in the vicinity of a next fiber. On the other hand, using appropriate data for the experiment, i.e., G _m = 28 GPa, E _f = 300 GPa, V _f = 10%, the maximum shear stress at the interface can be estimated for τ_max = 0.201, σ_max = 402 MPa, what if compared with shear strength of the matrix of 283 MPa, seems unlikely. Therefore, it should be assumed that fiber cracking in semi-ordered preform is caused by accompanying bending and low matrix strength at grain boundaries.

Usually strengthening fibers present at the fracture surface are strongly segmented. Even in such conditions, the “pull-out” effect or detaching of fiber ends, where the highest shear stresses arise, is rarely observed. Fractured fibers, from the one side, joint and anchor grains between them and on the other hand, when oriented perpendicular to direction of tension, initiate crack, and induce fracture development. Generally, in composite materials with high fiber volume fraction, the intercrystalline fracture is observed, whereas those with low reinforcement volume fraction or in matrix alloy, with larger grains, cracking in the transcrystalline way take place, see Fig. 9(b). Cracks in unreinforced volume propagate through grain boundaries whereas in the composite matrix grains are split in area with broken fibers. Reinforcement beside transferring the load and restraining stresses significantly refines microstructure. Dimension of the grains in unreinforced 2024 casting is about 50-100 μm whereas in the matrix, inside of the reinforced volume, decreases to 10-30 μm, see Fig. 9(a). Thus, this could be the predominant reinforcing effect causing increase of strength properties.

In the necking area subjected to higher load, fibers oriented parallel or at a slight angle to direction of tension were cracking at the beginning. It is because in such condition, they bear most of the acting stresses. Analyzed by Kang et al. (Ref 13), stress transfer in randomly oriented δ-Al₂O₃ fibers in composite materials confirmed that axial stresses for tilted fibers meets its minimum at 60° of orientation, whereas the highest interfacial shear stresses occurred at the angle of 45°. Thus, the most adverse fiber orientation is 60°, instead of 90°. Observations of the microstructure in the vicinity of the sample fracture, after tensile test, allowed determining the depth of cracking zone as 300-400 μm. In the vicinity of grain boundaries where lower elastic modulus and high plastic deformation occurred, fibers cracks were often observed, see Fig. 9(b). Thus, inside the grains, modulus of the matrix is higher than at the interface, and the matrix is effectively reinforced, whereas between boundaries, the most loads are transferred by strengthening fibers. According to elastic-plastic analysis (Ref 13, 14), fiber axial stresses can reach even 4000-7000 MPa, i.e., higher than their tensile strength (~2000 MPa). All these indicate that most of investigated stress transfer models considered homogeneous matrix properties, and prediction of strength properties can significantly differ from those determined experimentally.

Wear resistance of composite materials reinforced with Al₂O₃ fibers is better than matrix especially at higher load (see Fig. 10). During friction hard fragments of alumina, Saffil fibers were transferred between counterparts and composite material. Accelerated wear rate of both parts increased friction coefficient from 0.36 for unreinforced alloy to 0.45 for 2024/10% Al₂O₃ composite. Strengthening the composite material with addition of 3 vol.% of graphite fibers significantly improved wear resistance at dry condition. Usually carbon fibers, similarly as alumina fibers, restrain plastic deformation simultaneously piling up stresses and arresting dislocation (Ref 15, 16). With progress of abrasion at the wear surface, carbon fibers primarily provide lubricant film which protect against adhesion and seizure. At the wear surface, fragmented and plastically deformed matrix components can be found. Thus, graphite fibers slightly restrain deformation in subsurface, but first of all, they create solid lubricant pocket, see Fig. 11.