A comparison of gaseous hydrogen embrittlement, slow-strain-rate hydrogen embrittlement, and stress-corrosion cracking I
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IT is well established that near-alpha alloys such as Ti-8A 1-1Mo-IV undergo various forms of hydrogen embrittlement, including both gaseous hydrogen embrittlement (GHE) and slow-strain-rate hydrogen embrittlement (SSRHE), and stress-corrosion cracking (SCC) in aqueous media. While several details remain to be clarified, the mechanism of GHE appears to be satisfactorily explained in terms of the hydride-rupture model proposed by Nelson and his colleagues?,z It has been suggested that hydride formation also plays a key role in SSRHE, 3-5 but the evidence for this is less direct. Hydride models have been advocated for SCC of this alloy, but there is considerable opposition to this view) The present work was undertaken to resolve several of the remaining issues in the case of GHE, and, in particular, to make a direct comparison with this form of failure and SSRHE and SCC, thereby testing the hydride model in these cases. Nelson, Williams, and Stein 1studied the roomtemperature behavior of Ti-6A1-4V heat treated to produce a variety of microstructures, and concluded that the occurrence of significant GHE requires the existence of continuous fi paths of the type present in Widmanstatten structures. This finding, together with the fact that the hydrogen solubility and diffusivity are greater for the/3 than the a phase, led them to suggest that GHE involves diffusion of hydrogen through the/3 The authors were formerly with the Department of Metallurgy and Mining Engineering and Materials Research Laborary, University of Illinois, Urbana, Illinois. G. H. KOCH is now with Battelle's Columbus Laboratories, Columbus, Ohio 43201, A. J. BURSLE is now with Nabalco, Northern Territory, Australia, R. LIU is now with the Bell Laboratories, Murray Hill, New Jersey 07974, and E. N. PUGH is now with the National Bureau of Standards, Washington, D.C. 20234. Manuscript submitted October 27, 1980.
METALLURGICAL TRANSACTIONS A
to the a-fl interface, where it causes embrittlement of the neighboring a phase. Specifically, Nelson et al 1,2 proposed that embrittlement resulted from precipitation of titanium hydride in the a-phase, cracking proceeding by the repeated formation and rupture of the brittle hydride phase. Meyn6 provided some evidence for this hydride-rupture model by reporting the existence of the hydride in studies of GHE of Ti-8A1-1 Mo-1V with the Widmanst~itten structure. The fracture surfaces were reported to be covered by a black deposit whose X-ray diffraction pattern indicated a mixture of bct, fcc, and bcc titanium hydrides. The bct and bcc structures seem to be at variance with the literature on titanium hydrides.7,8The equilibrium hydride phase is 3' titanium hydride, whose composition can range from approximately Till15 to Till2; the 3' phase has a fluorite structure (a fcc arrangement of Ti atoms with H at the tetrahedral sites) over much of the range of stoichiometry, but has been observed to become fct as the hydrogen to metal atomic ratio exceeds 1.9. There is also some confusion concerning the path of GHE. Nelson
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