Flow Behavior of the L1 2 (Al,Fe) 3 Ti Single Crystals
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FLOW BEHAVIOR OF THE LI 2 (AI,Fe) 3 Ti SINGLE CRYSTALS Z.L. Wu, D.P.Pope and V.Vitek Department of Materials Science and Engineering, University of Pennsylvania, Philadelphia, PA 19104-6202.
ABSTRACT The compressive flow behavior of single crystalline L12 A16 7Fe 8 Ti 2 5 was investigated as a function of temperature and orientation at temperatures from 77K to about 1250K, using specimens with compressive axes orientated near [001], [1 13], [011], [122] and [ 111]. The operating slip systems seen in these specimens after 0.4% plastic deformation are predominantly of the octahedral type at all temperatures, even in near-[122] and [I 11] specimens in which the Schmid factors for the primary cube slip system are larger than that for the primary octahedral slip system. The yield stress increases rapidly with decreasing temperature at low temperatures, while it decreases gradually from room temperature to higher temperatures. The critical resolved shear stress (CRSS) on the [I01](111) slip system does not seem to be orientation-dependent over a wide range of temperatures, except at temperatures from 1050K to 1250K where the CRSS exhibits a mild orientation-dependence. Fracture tests at room temperature were also conducted. No special orientation-dependence of the ductility was observed.
INTRODUCTION The flow behavior of various L1 2 titanium trialuminides has been studied by many groups, however, the results often differ. For instance, Kumar and Pickens observed an anomalous flow stress peak at around 700K in polycrystalline AI 2 2 Fe3 Ti 8 [1]. Similar peaks were also reported in A16 7 Mn 8Ti 2 5 [2], A16 7 Cr 8Ti 2 5 [3], and AI 5 Cu 3 Ti2 [4]. Wu et al, on the other hand, did not observe an anomalous peak in Fe-modified alloys [5-7]. Instead, they observed a flow behavior which is similar to that of L1 2 Pt3Al-type alloys [8], and the flow stress vs temperature curve at intermediate temperatures is plateau-like. The Pt3AI-type flow behavior has also been recently seen in A16 7 Cr8 Ti 2 5 and A16 6 Mn 9 Ti2 5 [9], results that, however, differ from the ones reported earlier in [2, 3], even though the compositions of the alloys used in these studies were about the same. The structure of [110] superdislocations in deformed alloys has also been investigated by several groups using weak beam TEM. Interestingly, some of the studies showed that the superdislocations are dissociated into antiphase boundary (APB)-associated superpartials, and lie on octahedral slip planes at low and intermediate temperatures [10-12], and mainly on cube slip planes at high temperatures [13]. Such a transition in slip plane from low temperatures to high temperatures has been seen in L1 2 Ni3AI-type alloys and is believed to be resulted from an anisotropy of the APB-energy. According to the so-called PPV model [14], the anisotropy of the APB-energy on the two different slip planes provides a driving force for the cross-slip pinning of { 111) superdislocations onto the cube planes (with the aid of thermal activation). The cross-slip pinning in
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