Friction stir processing (FSP) was applied to produce aluminum based in situ composites from powder mixtures of Al-Fe, Al-Mo, and Al-Fe2O3. Billet of powder mixtures was prepared by the use of conventional pressing and sintering route. The sintered billet was then subjected to multiple passages of FSP. During FSP, the material has experienced both high temperature and very large plastic strain. The basic idea for fabricating the composites is to combine the hot working nature of friction stir processing (FSP) and the exothermic reaction between aluminum and transition metals (Al-Fe, Al-Mo) or metal oxides (Al-Fe2O3). In the Al-Fe alloy, in situ Al¡VFe reaction can be induced during FSP and form Al-Al13Fe4 composite. The size of reinforcing particles formed by the in-situ reaction is ~100 nm. In Al-Mo alloys, fine Al-Mo intermetallic particles with an average size of ~200 nm were formed and uniformly dispersed in the aluminum matrix by FSP. The Al-Mo intermetallic particles were identified mainly as Al12Mo with minor amount of Al5Mo. The exothermic reaction could result in local melting of Al at the Al/TM interface, and the liquid Al may accelerate the reaction. In addition, it is suggested that the critical mechanism responsible for the rapid reaction and the formation of nanometer sized particles in FSP is the effective removal of the Al-TM intermetallic phase from the Al-TM interface, maintaining an intimate contact between TM and Al. In the Al-Fe2O3 system, the reactions taking place during FSP includes the thermite reaction (2Al +Fe2O3 ¡÷ Al2O3 + 2Fe), and the reaction between the reduced Fe and Al to form Al13Fe4. In the FSPed Al-Fe2O3 specimens, there are two types of second phase particles, Al13Fe4 and Al2O3. The Al2O3 particles (about 10 nm in size) usually appear as a cluster of 100-200 nm in diameter. There are two types of Al2O3 phases existed in the Al matrix after FSP passes, depending on the content of Fe2O3. One is £^-Al2O3 in Al-2Fe2O3 specimens, and the other is £\-Al2O3 in Al-4Fe2O3 specimens. It is suggested that the formation of different type of Al2O3 particles in the Al-Fe2O3 composites may be attributed to different heat release in each system. The lower heat release in Al-2Fe2O3 sample favors the formation of the while the higher heat release in Al-4Fe2O3 sample results in the £\-Al2O3. The Al-Al13Fe4/Al2O3 composite produced by FSP exhibits both high strength and good tensile ductility. The higher strength in Al-Fe2O3 specimen may be due to the presence of fine Al2O3 particles. The flow stress of the Al-4Fe2O3 composite can maintain at 100 MPa even at 773 K. The good thermal stability and high temperature strength of Al-Al13Fe4/Al2O3 composites could be attributed to the fine dispersion of second phase particles in the aluminum matrix, especially the nanometric Al2O3 particles. These Al2O3 particles are very stable at elevated temperatures, even after long time exposure at 873 K. The temperature excursion in FSP is determined by both the FSP parameters and the exothermic reaction involved. The peak temperature in Al-Fe or Al-Fe2O3 system during FSP was calculated as a function of the fraction of Fe or Fe2O3 reacted. Based on calculated results, it is noted that with the in situ reaction, the value of can easily reach the melting point of Al, especially for the Al-Fe2O3 system. The reaction mechanism and microstructure evolution during FSP are discussed.

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We have studied precipitate growth and coarsening in ternary alloys using two different phase held models. The first one is a ternary extension of the classical Cahn-Hilliard (C-H) model in which both the phases are characterized using conserved held variables i.e. composition (cB; cC ); mobility matrix and gradient energy efficient are the other input parameters in this model. In the second model, each phase is treated as separate, and phase identify cation is through a (non-conserved) phase held variable ; we have used a grand potential-based (GP) formulation, due to Plapp [1], Choudhury and Nestler [2], where interfacial energy and interface width, as well as free energy and diffusivity matrix for the relevant phases are the input parameters. The first model i.e. the Cahn-Hilliard (C-H) type model is conceptually simple. The model for ternary is a straight forward extension of the binary. The grand potential (GP) formulation has the advantage of being able to incorporate thermodynamic database like Thermocalc in it. We present below a summary of the findings of our research on (a) precipitate growth, precipitate coarsening, and (c) a critical comparison between results from phase held simulations and those from experiments on an Ni-Al-Mo alloy Precipitate growth In our study of precipitate growth in ternary alloys, we end that when both the solute elements have the same diffusivity, precipitate growth behaviour in ternary alloys is identical to that binary alloys; specifically, we recover the temporal power law r2 = kgt relating the particle radius to time, and the growth kg depends only on supersaturation (i.e., equilibrium volume fraction of the precipitate phase), and is independent of the slope of the tie line. However, when one solute element, (say, C) di uses slower than the other (i.e. (DCC =DBB) < 1,(where DBB, DCC are intertie suavities’ in the lab frame of reference), the ux of C at the interface is smaller than that of species B, causing the precipitate to become depleted in C and enriched in B; this process continues until the growth phase enters a scaling regime where we recover the temporal law for growth: r2 = kgt. In this regime, the tie line selected by the precipitate and matrix interfacial compositions is different from the thermodynamic tie line containing the alloy, a result first reported by Coates [3]. After validating our phase held model quantitatively through a critical comparison with Coates' theory of tie line selection, we have characterized the growth behaviour: specifically, we end that growth kg drops with decreasing value of DCC ; the magnitude of this drop is stronger for alloys which (a) are on higher-C tie lines (i.e., the slope of the tie line is higher), and (b) have smaller precipitate volume fractions. Precipitate coarsening In our simulations, we end that precipitate coarsening does indeed enter a scaling regime where the temporal power law r3 = kt (which relates the average precipitate radius r to (b) time t) is valid; the coarsening rate k depends, as expected, not only on precipitate volume fraction, but also on the slope of the tie line and diffusivity ratio (DCC =DBB). (c) (d) When the solutes have equal diffusivity (i.e., (DCC =DBB) = 1), the coarsening behaviour is essentially the same as that in a binary alloy. However, when solute C (say) is the slower di using species, the coarsening rate k drops, with a deeper drop in alloys on higher-C tie lines. Both these conclusions are similar to those from our study of precipitate growth. (e) (f) However, there is a crucial difference between precipitate growth and coarsening in ternary alloys: The suppression in coarsening rate (for DCC < DBB) in ternary alloys is accompanied by another e ect: larger (and growing) precipitates are richer in the faster di using species B, while the smaller and shrinking precipitates are richer in the slower di using C. In other words, during coarsening in ternary alloys, the tie line selected by precipitate and matrix interfacial components depends on precipitate size; during growth, however, the scaling regime is characterized by the same tie line, independent of precipitate size. (g) (h) (i) Critical comparison between theory and experiment (j) (k) (l) We have used the grand potential based phase held model [1] [2] to study coarsening in Ni-Al-Mo alloys. This model has the advantage of ease with which we can incorporate the thermodynamic and kinetic data on real alloys. (m) (n) A comparison of coarsening rate from our 3D simulations with the experimentally observed rate reveals that diffusivity of the faster di using species (which, in Ni-Al-Mo alloys, is aluminium) from our simulations is within an order of magnitude from the experimental value. However the dominant term in the (@ =@c) matrix is underestimated by 2 to 3 orders of magnitude (compared to its value computed from CALPHAD-based thermodynamic data).