Alloy and Microstructural Design by John K. Tien

By John K. Tien

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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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Alloy and Microstructural Design by John K. Tien
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