Effect of Heat Treatment on Interfacial Strengthening Mechanisms of Second Phase Particulate Reinforced Aluminium Alloy

Transcription

1 , Hradec nad Moravici Effect of Heat Treatment on Interfacial Strengthening Mechanisms of Second Phase Particulate Reinforced Aluminium Alloy S.T. Hasan Faculty of Arts, Computing, Engineering and Sciences, Sheffield Hallam University, Pond Street, Sheffield S1 1WB, UK; Abstract The satisfactory performance of reinforced aluminium alloys depends critically on their integrity, the heart of which is the quality of the matrix-reinforcement interface. The nature of the interface depends in turn on the processing of the metal matrix composite component. At the micro-level the development of local stress concentration gradients around the reinforcement, as the metal matrix attempts to deform during processing, can be very different to the nominal conditions and play a crucial role in important microstructural events such as segregation and precipitation at the matrix-reinforcement interface. These events dominate the cohesive strength and subsequent mechanical properties of the interface. The aim of this study is to investigate the features, which significantly affect the microstructural developments when subjected to fully hardened heat treatments and their effects on precipitate distribution within the aluminium matrix. The compositional variations at the matrix-reinforcement interface of a metal matrix composite are reported, with emphasis on the interfacial strengthening mechanisms during thermo-mechanical processing. A method of calculation has been applied to predict the interfacial fracture strength of aluminium and SiC interface, in the presence of magnesium segregation. Keywords: MMC, Segregation, Precipitation, Interfacial-strength 1. INTRODUCTION Metal matrix composites are rapidly becoming strong candidates as structural materials for many high temperature and aerospace applications. Metal matrix composites combine metallic properties with ceramic properties, leading to high stiffness and strength with a reduction in structural weight. The main objective of using a metal matrix composite system is to increase service temperature or specific mechanical properties of structural components by replacing existing superalloys. At present the relationship between the strength properties of metal matrix composites and the details of the thermo mechanical forming processes is not well understood. The purpose of this study is to define the features which significantly affect the interfacial strength of a practical aluminium alloy/silicon carbide system and which are directly related to the forming processes currently being used by the industry. 2. MATERIALS The metal matrix composite studied was an aluminium-zinc-magnesium-copper (AlZnMgCu) alloy matrix (N707) reinforced with varying amounts of silicon carbide particles of F600 grit, which has a mean diameter of approximately 10 µm. AlZnMgCu alloy is heat treatable and is produced by spray deposition technology and in the peak aged temper T6 condition shows very high strength at room temperature without a significant loss in ductility. The AlZnMgCu alloy obtains its excellent properties from a very

2 fine and homogeneous microstructure and extended alloying possibilities allows maximum hardening as well as a stabilisation of the desired microstructure. For this investigation two types of material were used, 1) monolithic AlZnMgCu alloy tubes and 2) metal matrix composite tubes with varying amounts of silicon carbide particles as reinforcement. Table 1 contains the details of the chemical composition of the matrix alloy and the amount of silicon carbide particles in the metal matrix composites. The material was supplied as 17 mm diameter tubes with 1mm wall thickness following hot extrusion alone and also after hot extrusion, annealing and then cold drawing. 3. HEAT TREATMENT Table 1 Chemical Composition (wt %) Code Zn Cu Mg Zr SiC AZC-I AZC-II AZC-III HT1 produces the highest value of mechanical properties and is achieved by solution treating the material at 450 o C for 1 hour per 2mm section thickness and water quench to room temperature. This is followed by precipitation treatment (artificial ageing) for 24 hours at 170 o C and allowed to air cool. 4. MICROSCOPIC EXAMINATION The samples were examined using a Philips XL40 analytical scanning electron microscope equipped with EDAX energy dispersive X-Ray analyser and Jeol 100 CX TEM system with Link 860 energy dispersive spectrometer (100kV) to perform microanalysis. Etching of the samples was achieved using a two stage process. The first stage involved the use of modified Keller's reagent, with the HF replaced with Sodium Flouride on health and safety grounds. This etchant highlighted the secondary phases of the material when viewed in secondary electron mode. The second stage involved the use of Sodium Hydroxide Potassium Ferricyanide etchant, which highlighted the grain boundaries. The samples were also examined in the as polished conditions, which produced excellent results when viewed in back scattered electron mode. To study the matrix reinforcement interface in greater detail, an ion beam thinning technique was applied to achieve better interface detail [1]. 5. RESULTS Scanning electron microscope observations of the microstructure of the monolithic N707 alloy (AZC-I), with or without reinforcement, after heat treatment show large differences in the size and number of precipitates.

3 Figure 1: The monolithic alloy AZC-I in T6 condition Figure 2: SiC p reinforced alloy AZC-II in T6 condition Figure 1 shows the monolithic alloy T6 conditions, the equiaxed grain structure is clearly visible, with the dark areas being MgZn 2 precipitates; the bright areas are Cu-MgZn 2 precipitates. Compare with Figure 2 which shows the alloy reinforced with 7.8 wt% SiC p, the size and distribution of Cu-MgZn 2 precipitates has changed significantly. The majority of the precipitates are clustered around the reinforcing particles with some areas denuded of precipitates altogether. Figure 3: SiCp reinforced Aluminium alloy (AZC-III) in T6 heat treated condition Figure 3 shows the alloy reinforced with 11.5wt% SiC p in the T6 condition. There is a marked increase in the number of Cu-MgZn 2 precipitates when compared with 7.8wt% SiC p reinforced alloy. There are fewer and smaller precipitates around the reinforcing particles compared with the 7.8wt% SiC p alloy. This leads us to conclude that the dislocation network promoted by the reinforcing particles is more evenly spread throughout the alloy and is not mainly clustered around the particles, which appears to be the case in the 7.8wt% SiC p reinforced alloy. Figure 4 shows the interface micrograph, revealed by ion beam thinning of the disc specimens. STEM analysis was performed on the interface boundary, as marked on the micrograph.

4 Figure 4: Thin foil micrograph (AZC-III) showing the matrix-reinforcement interface. Figure 5 shows the plot of zinc and magnesium concentrations at the matrix-reinforcement interface analysed for longer times (up to 300s) to pick up the traces of zinc or magnesium segregation at the interface boundary. Figure 5: Compositional variations of Zn and Mg at AlZnMgCu - SiC interface. The analysis was carefully repeated by stopping after every 50s to adjust the spot target in order to verify the magnesium content at the interface boundary and not to be confused with MgZn 2 precipitates appearing in the matrix. However, the zinc content shows no such indication of being segregated at the interface boundary. 6. DISCUSSION Lim and Watanabe [2] and Shvindlerman and Faulkner [3] have recognised that the interface structure is important in determining the amount of predicted segregation and hence the change of the interfacial energy caused by the segregation. Equations have been developed to forecast the energy change in terms of the coincidence site stress (σ a) value describing the boundary, and the formation energies of impurities on the boundary. The predictions fit well

5 with observations of specific grain boundaries in molybdenum and aluminium. The model indicates that low angle boundaries and those with low σ coincidence site values will be less susceptible to fracture than high angle boundaries. This indication is supported by experimental measurements made by Lim and Watanabe [2]. The work of intergranular fracture, G k is given by [3]; G k = A σ p exp(n ln (σ/σ o )) (1) where, A is the dislocation pile up term describing the effectiveness of dislocations in providing stress concentration at the advancing intergranular crack tip (=100) [4], n is the strain rate hardening exponent, σ p is the energy involved with creating the fracture surface = 2σ s - σ gb (= σ o ), σ s is the surface energy, σ gb is the grain boundary energy and σ a is the new interfacial energy caused by segregation, given by; σ a = σ o - ZRT ln (1- c + Bc) (2) where, Z is the term describing the density of interface sites which are disordered enough to act as segregation sites, R is the gas constant, T is the absolute temperature, c is the segregate concentration needed to cause embrittlement (=0.1), B is a term describing the modification of the boundary energy by impurities using the Zuchovitsky equations [5], given by; B = exp ((σ gb /Z) (Hv/kT)) (3) where; Hv is the enthalpy of formation of the impurity atom in the bulk (in ev), k is Boltzmann s constant, Z is assumed constant for all boundaries in equation 3. Values of Hv can be estimated for metallic solids using elasticity theory, based on misfit arguments [4]. These calculations have been applied to the case for reinforcement-matrix interfaces in aluminium based MMCs by assuming that magnesium is the main segregation species. Clearly at this stage of research there will be differences between the approach used for grain boundaries and interphase interfaces. In this study it has been assumed that the SiC interface can be treated as a random grain boundary (Z=1), and, in the absence of any reliable data on Al-SiC interfacial energy, property values for pure aluminium be used. The results are summarised, together with the data used, in Table 2.

6 Table 2. Interface Fracture Strength r AL r Mg E Al σ s σ F x10-10 m x10-10 m Nm -2 Jm -2 MPa x x * * unsegregated 7. CONCLUSIONS It can be seen that Mg reduces the interface fracture strength of Al. If the segregation mechanism is the non-equilibrium type then this source of weakness can be removed by the appropriate heat treatment. If it is equilibrium segregation then it will be impossible to remove and the problem can only be overcome by removing magnesium or by introducing another, more strongly segregating element. The results demonstrate that compositional variations over 100 to 300 nm can be reliably measured. The presence of zinc depletion and magnesium segregation at a matrixreinforcement interface has been measured. A method of calculation has been applied to predict the interfacial fracture strength of aluminium, in the presence of magnesium segregation. The model shows success in making prediction possible of trends in relation to segregation and intergranular fracture strength behaviour in metallic solids. Further work is in progress to give a more quantitative forecast of the effects of second phase particles/ reinforcement on mechanical properties of matrix-reinforcement interface in metal matrix composites. REFERENCES [1] S.T. Hasan, J.H. Beynon and R.G. Faulkner, 2004, J. of Materials Processing Technology, , pp [2] L.C. Lim and T. Watanabe, 1990, Acta Met., 38, [3] R.G. Faulkner and L.S. Shvindlerman, 1996, Materials Science Forum, Vol , No. pt 1, pp [4] R.G. Faulkner, 1985, Materials Science and Technology, 1, 442. [5] A.A. Zuchovitsky, 1944, J. of Phys. Chemistry, Vol. 18, No. 3.