Early models ofprecipitate stability under irradiation were based on the ideas of Nelson et a/.,76 who suggested that precipitates would evolve to an equi­librium size determined by competing processes affecting their growth, via the radiation enhanced and/or thermal diffusion of solutes, and their simul­taneous dissolution due to damage arising in collision cascades. Two dissolution mechanisms were sug­gested: recoil dissolution due to the displacement of atoms from the precipitate into the matrix, and dis­ordering dissolution of ordered phases such as g, with the latter predicted to be the more effective. The model predicted that fine precipitates would continue to grow to some equilibrium size (depen­dent on temperature, dose rate, and solute levels), but that precipitates greater than this size would shrink. Experimental evidence for the dissolution of large preexisting Ni3Al g precipitates in heavy-ion- irradiated Ni-Al alloys was shown by Nelson et a/.76

These ideas were developed further and applied to g precipitates in ion and neutron-irradiated alloys by Baron et a/.77 The model developed by Baron et a/. indicated that, at a given particle size, a higher solute supersaturation was required under irradiation than in a purely thermal environment. The model appeared to be consistent with the observed coarsen­ing behavior of g0 precipitates during irradiation, though no evidence for the shrinkage oflarge particles was presented. For example, data for PE16 irradiated at fluences up to 7.5 x 1026nm—2 at 560 °C, which were reported by Chang and Baron,78 only examined the growth of g0 particles up to a maximum radius of — 15 nm under conditions where the predicted maximum equilibrium radius was — 35 nm.

However, detailed examinations of g structures in neutron-irradiated Nimonic PE16 which were made by Gelles79 found no evidence to indicate that irra­diation-induced dissolution mechanisms limited the particle size. Microstructural examination of PE16, originally in ST, STA, and OA conditions, irradiated in EBR-II to -27 dpa (5.4 x 1026nm—2, E> 0.1 MeV) at 600 ° C, revealed that preexisting g dispersions in aged material were maintained but continued to coarsen even in the OA condition, and that a fine dispersion formed in ST material. Coarsening of the g0 particles in the OA material was accompanied by the formation of fine background precipitates in some regions. Further in-reactor precipitation of g0 also occurred at point defect sinks, including void surfaces and dislocations, in all the heat-treated con­ditions. Additional examinations by Gelles80 of ST PE16, irradiated to —30-50 dpa at temperatures in the range of 430-650 °C, indicated that g coarsening was controlled by radiation-enhanced diffusion below 600 °C with an activation energy that (in agreement with theoretical predictions for a process governed by point defect recombination) was approx­imately a quarter of that for thermal diffusion.

As described in Section 4.04.5.1 in relation to irradiation embrittlement effects, Yang81 examined an identically irradiated set of ST PE16 samples as Gelles, focusing on the precipitation of g at grain boundaries. Similar g structures to those described by Gelles and Yang were also observed by Boothby2 in the aged conditions of EBR-II-irradiated PE16, though at higher irradiation temperatures (>540 °C for the STA condition, and >600 °C for the OA condition), where doses were in the range 66-74 dpa, the spherical g precipitates which formed during thermal aging were almost entirely replaced by ‘skel­etal’ forms nucleated at point defect sinks. Figure 11 shows an example of the g distribution, imaged in dark field, in STA PE16 irradiated to 69 dpa at 570 °C; although small spherical precipitates were retained in a narrow region adjacent to the grain boundary, a much coarser dispersion is evident at the boundary itself and within the bulk of the grain.