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Extra info for Alloy and Microstructural Design
25. Light micrograph showing heavy precipitation of a phase at β grain bounda ries in Ti-8Mo-8V-2Fe-3Al quenched from 1175°K and aged 4 hr at 775°K. 200 ι 0 1 1 1—ι ι ι ι ι ι 1 10 1 1—ι ι ι ι ι | AGING TIME ( h r ) 100 Fig. 26. Showing variations in yield strength and α-phase volume fraction in Ti-8Mo-8V-2Fe-3Al as a function of aging time at 775°K after quenching from 1175°K. Fig. 27. Thin foil electron micrographs showing the reduction of a precipitate size in tempered Ti-8Mo-8V-2Fe-3Al. (a) Quenched from 1175°K and aged 2 hr at 775°K; (b) quenched from 1175°K, deformed 30% at 300°K, and aged 2 hr at 775°K.
At Al contents > 6 wt % the ordered a phase can be formed during slow cooling or during prolonged elevated temperature exposure. The a phase forms as a uniform dispersion of coherent precipitates. The presence of a serves to intensify the planar slip and further reduces ductility. Alloys that contain between 8 and 9 wt % Al, when aged to produce a maximum volume fraction of a , be come completely embrittled (Crossley and Carew, 1957) and fracture by cleavage (Blackburn and Williams, 1969a). Thus it is clear that a practical upper limit for solid solution strength ening exists for α-phase alloys because of the intervening brittleness problem.
Yielding in such alloys is believed to occur at the stress where dislocations move past the par ticles by Orowan looping (Ashby, 1964). Indeed, the yield stress as measured on dispersion-hardened single crystals at small plastic strains (~10~ ) is found to vary with the reciprocal of the planar interparticle 4 STRENGTHENED STRAIN Fig. 14. Typical yielding behavior of solution-hardened, precipitation-hardened, and dispersion-hardened alloys. 30 Stephen Μ. Copley and James C. Williams spacing in agreement with Eq.
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