With respect to alloy M3S, the improved strength values following T5- and T6-heat treatments can be ascribed to (i) the presence of [alpha]-[Al.sub.15][(Fe, Mn).sub.3][Si.sub.2] in the form of blocky hard particles, (ii) the precipitated strengthening Cu- and Mg-containing dispersoids, and (iii) the probable formation of [Al.sub.6]Mn fine dispersoids in the presence of a high Mn content of 0.75 wt.% [39, 40].
Thus, by applying artificial aging directly after solidification (i.e., T5 temper), large proportions of the dissolved Cu and Mg in the solid solution will form strengthening dispersoids. The T6 heat treatment, on the other hand, involves solutionizing the as-cast structure at a sufficiently high temperature in order to dissolve higher amounts of Cu and Mg in solid solution in order to form a supersaturated solid solution upon quenching and to achieve further modification of the eutectic Si particles.
The enhanced strength values of alloy M4S after T6 heat treatment may be attributed to the presence of [delta]-[Al.sub.3]CuNi and eutectic Al-[Al.sub.3]Ni phases that prove to contribute effectively to the elevated-temperature strength of alloy M4S, in spite of a considerable amount of Cu that is consumed in forming the [delta]-[Al.sub.3]CuNi phase, which will certainly affect the amount of fine [Al.sub.2]Cu dispersoids formed, which is consistent with the findings reported in references [38, 49].
Theory suggests that, if prefabricated dispersoids
in the form of nanoballoons were introduced into a pure aluminum melt, the resulting alloy could not only be stronger, but lighter as well.
 also observed short rod-like dispersoid particles of the Al-Cu-Mn compound in the solutionized Al-Si-Cu-Mn alloy and identified it as [T.sub.Mn] ([Al.sub.20][Cu.sub.2][Mn.sub.3]) by calculating the SAD patterns.
It is just in the Cu-rich regions that precipitation of [T.sub.Mn] and Al[Cu.sub.3][Mn.sub.2] dispersoid particles occurs.
After solutionizing, a great number of fine [T.sub.Mn] and Al[Cu.sub.3][Mn.sub.2] dispersoid particles are precipitated in the [A.sub.2] alloy except for the redissolving of part of the eutectic 6 phase (Figures 1-3 and 8).
The extent to which C was converted to [Al.sub.4][C.sub.3] dispersoids
was analysed by gas chromatographic measurements in .
Several morphological forms exist that include colonies of 1-5 micron dispersoids
The dispersion strengthening can be described as dislocations inhibited by Ni or/and Sn dispersoid in the slipping planes.
It shows the prevalence of Ni-rich dispersoids particles within the matrix.
The fully dense materials were composed of a nanostructured ferrite matrix having an average grain size of 150 nm and cementite [Fe.sub.3]C dispersoids
a few nanometers in size distributed in the ferrite grains and along the grain boundaries.