Rowcliffe and Horak86 investigated the tensile prop­erties of Inconel 706 (in a multistep ‘fully aged’ condition) and Inconel 718 (ST condition) following irradiation in EBR-II to fluences of 4-5 x 1026nm~2 (E > 0.1 MeV). Irradiation temperatures (Ti) ranged from 450 to 735 °C, with tensile tests being per­formed at a strain rate of 4 x 10~4 s~3 at temperatures corresponding to Ti and to Ti + 110 °C. Yield stresses and total elongation data for Inconel 706 are shown in Figure 12 and for Inconel 718 in Figure 13. Data for Inconel 706 showed very high (>1000 MPa) yield stresses and ultimate tensile strengths (UTS) in
specimens irradiated at temperatures up to and including 500 °C. This high tensile strength was maintained in a specimen irradiated at 500 °C but tested at 610 ° C. Although there was some reduction in strength in specimens irradiated at 560 °C and above, the UTS remained above 650 MPa in speci­mens irradiated at 625 ° C. The very high tensile strengths exhibited at the lower irradiation tempera­tures were attributed to the instability of the (ordered body-centered tetragonal) g» phase below 525 °C and its consequent dissolution, leading to the reprecipita­tion of nickel and niobium as (ordered face-centered cubic) g’ on dislocation loops. At higher irradiation temperatures, both g’ and g" were stable, but

200 —

800

Temperature (°C)

precipitate coarsening resulted in lower tensile strength. Elongations to failure for tests carried out at the irradiation temperature were between 1.5% and 3% up to 625 °C, compared to >8% in unirradi­ated material. Irradiation embrittlement was gener­ally more severe in tests at T + 110 °C, particularly at 610-735 °C where the lowest recorded ductility was 0.2%. Fractures in irradiated Inconel 706 were predominantly intergranular, with failure believed to be facilitated by the decohesion of Z phase (hexago­nal Ni3(Ti, Nb)) platelets which were formed at grain boundaries during the initial heat treatment.

Rowcliffe and Horak’s data for ST Inconel 718 showed similar trends to Inconel 706. Precipitation of the g and g" phases occurred during the irradia­tion of Inconel 718, resulting in yield strengths in excess of 1000 MPa at irradiation temperatures up to 560 °C and above 800 MPa at 625 °C. The ductility of Inconel 718 was reduced from more than 30% in the unirradiated condition to 0.2% or less in specimens which were irradiated at 500-560 °C and tested at Ti + 110 °C. In contrast to Inconel 706, failures in irradiated Inconel 718 were reported to be predomi­nantly transgranular. Crack propagation in Inconel 718 appeared to have been via a ‘channel’ fracture mechanism, that is, with deformation occurring by highly localized planar slip and consequent linkage of radiation-induced voids.

Bajaj et al.87 examined the tensile properties of Nimonic PE16 irradiated in EBR-II to neutron fluences up to a maximum of ^7 x 1026nm 2 (E > 0.1 MeV), at temperatures in the range of 450­735 °C. The alloy was in a STA (1 h at 900 °C plus 8 h at 750 °C) condition, and appears to have been the same low-Si heat of PE16 that was subsequently used in the AA-1 swelling experiment described by Garner and Gelles.22 Tensile tests were carried out at 232 °C (to simulate refueling conditions), at the irradiation temperature Ti and at Ti + 110 ° C (to simulate reac­tor transients), at a strain rate of 4 x 10~4s~ and with a small number of tests at 4 x 10~3s~ . Irra­diated specimens tested at 232 °C generally showed a substantial increase in yield stress and a small increase in UTS over the unirradiated values (although samples irradiated at the highest tempera­ture of 735 °C exhibited some softening), and retained good levels of ductility with total elonga­tion values above 10%. Yield stress and total elon­gation data for PE16 at higher test temperatures are shown in Figure 14 for specimens irradiated to a fast neutron fluence of 4.3 x 1026nm~2 (enabling direct comparison with the data for the similarly irradiated Inconel alloys shown in Figures 12 and 13). Speci­mens tested at the irradiation temperature again showed strengthening at temperatures in the range of 450-625 °C and softening at 735 °C, with good

ductility at 450 °C but total elongations reduced to ^3% at 560-625 °C. Tests at T + 110 °C showed fUrther increases in tensile strength (consistent with the greater hardening expected from irradiation at a lower temperature) and more severe embrittlement with ductility levels at 670-735 °C reduced to 0.3% at a fluence of 4.3 x 1026nm~2 and to zero (i. e., failure before yield) in higher dose samples (7.1 x 1026nm~2). Tests at Ti at the higher strain rate resulted in an improvement in ductility by a factor of two or three. Examination of fracture surfaces showed that failures were predominantly intergranular in irradiated sam­ples tested above ^550 °C, transgranular at 232 °C, and mixed mode at 450-550 °C. Bajaj et al. considered that the irradiation embrittlement of PE16 evident at high temperatures could simply be explained by matrix hardening with little or no change in the grain bound­ary fracture strength — evidenced by increases in yield strength but no significant changes in true (as opposed to engineering) UTS values — so that mechanisms relying on the weakening of grain boundaries could be discounted for the test conditions studied.

Sklad et al50 reported tensile data for two aged conditions of Nimonic PE16 which were irradiated in EBR-II to 1.2 x 1026nm~2 (E >0.1MeV) at 500 °C and tested at strain rates from ^3 x 10~5 to 3 x 10~3s~ . There was no significant difference in the postirradiation properties of the two differently aged conditions, although one aging treatment (2 h at 800 °C plus 16 h at 700 °C) resulted in an unirradiated yield stress ^25% higher than the other condition (1 h at 900 °C plus 8 h at 750 °C). No effect of strain rate on tensile properties was evident in tests at the irradiation temperature, where total elongations remained above 10%. Tests at higher temperatures were made only at the lowest strain rate, with failure elongations being reduced to 1.6% at 600 °C and 0.5% at 700 °C. The low ductility failures were associated with an increased tendency toward intergranular fracture, and additional tests, in which samples irradiated to 4 x 1026nm~2 at 500 °C were fractured in situ in an Auger spectrometer, revealed helium release from samples which fractured intergranularly as well as the segregation of Ni, P, and S to grain boundaries. Helium release was estimated at ^0.03 He atoms per grain boundary atom. No grain boundary helium bubbles were observable by TEM, and it was therefore considered that helium either remained in solution as a partial monolayer or was present in unresolved bubbles less than 1-2 nm in diameter.

The presence of grain boundary helium bubbles in Nimonic PE16 was reported by Fisher et a/.88 in sections of AGR (advanced gas-cooled reactor) tie bars irradiated at 512 °C and above. AGR tie bars, which are approximately 10 m long and are under load only during charging and discharging of the fuel element stringers, operate at temperatures from 325 to 650 °C from bottom to top, with peak doses of ^3 dpa occurring at around the 4 m position. Stress-rupture testing at 600 °C at an applied stress of 500 MPa showed a trough in properties (i. e., a minimum in failure times) and intergranular failures in sections of some tie bars which were irradiated at temperatures in the range of 350-400 °C where grain boundary helium bubbles were not generally observed. Even so, grain boundary cavitation was observed in a fractured tie bar section which was irradiated at 360 °C, with the cavities appearing to be nucleated (possibly at submicroscopic helium bubbles) at the intersection of slip bands with the boundary. The trough in stress-rupture properties occurred in tie bar sections which exhibited both high yield strengths (attributable to high concentra­tions of dislocation loops and small voids) and high levels of grain boundary segregation. EDX (energy dispersive X-ray) analyses showed a significant enrichment of Ni and Si, and a depletion of Fe, Cr, and Mo, at the grain boundaries ofsections irradiated at 335-585 °C. In addition, high levels of Si were detected in sections irradiated at 335-512 ° C in the g phase that precipitated at the surface of voids, with the Si content increasing with decreasing irradiation temperature. Although the presence of Si-enriched g phase at grain boundaries could not be confirmed, it was suggested that its formation may have contribu­ted to the minimum in stress-rupture life, which was thought to result from the weakening of the bound­aries relative to the matrix.

Grain boundary helium bubbles were also observed by Boothby and Harries89 and Boothby28 in PE16 irradiated in DFR and EBR-II at 535 °C and above. Tensile testing of DFR-irradiated PE16, exposed to ^20dpa at 465-635 °C, and strained at a rate of 2.5 x 10~6s_1 at temperatures approximating those of irradiation, revealed severe embrittlement with mini­mum elongations of ^0.2% at 550 °C; TEM examina­tion of strained specimens provided evidence of intergranular cavitation, and the ductility data were interpreted using a model for the diffusion-induced growth of cavities nucleated at grain boundary helium

bubbles.89

The postirradiation tensile properties and micro­structure of developmental g (D21, D25, and D66) and g/g’ (D68) strengthened alloys were discussed by Yang et al. The alloys were all irradiated in a ST condition; additionally, D25 was irradiated in an aged (24h at 700°C) condition (STA), and D66 in a 30% cold-worked plus aged (11 h at 800 °C plus 2 h at 700 °C) condition (CWA). Specimens were irradiated at 450-735 °C to a fast neutron fluence of 4 x 1026nm~2 (E > 0.1 MeV) in EBR-II, and were tested at Ti, Ti + 110 °C and 232 °C. Severe irradiation embrittlement was evident in the ST alloys and STA D25, particularly in tests at Ti + 110 °C. Zero ductility was recorded in the lower-Ni alloy D21 (25Ni-8Cr) irradiated and tested at 550 and 600 °C. Severe ductil­ity losses were associated with intergranular failures, which were attributed to irradiation-induced solute segregation and consequent precipitation of brittle g0 layers at grain boundaries. However, reasonable levels of ductility, ranging from 2 to 6%, coupled with trans­granular failures, were obtained at all temperatures in irradiated CWA D66 (45Ni-12Cr). The preirradiation grain boundary structure of this material, comprising a ‘necklace’ of small recrystallized subgrains plus large g0 particles and discrete Laves particles, remained stable with no indication of irradiation-induced g0 layers. Yang etal. considered that the radiation-induced segregation of g0 forming solutes to grain boundaries was inhibited by the introduction of a high density of dislocation sinks by cold working.

Vaidyanathan et al.90 and Huang and Fish91 examined the embrittlement of EBR-II-irradiated, precipitation-hardened alloys, using ring ductility tests. In this test, small sections of tubing are com­pressed and the ductility, defined as the strain at the initiation of cracking, is deduced from the change in the sample radius of curvature at maximum load. Both experiments included Inconel 706 and Nimonic PE16 in ST conditions, while Vaidyanathan et al. also examined the developmental alloys D25 and D68 in ST and STA conditions. Peak fluences in these experiments were around 6-7 x 1026nm~2 (E > 0.1 MeV) and irradiation temperatures were in the range 460-616 ° C. All the materials exhibited low ductility failures at high test temperatures, particu­larly in tests at about Ti + 110 °C where ductilities were generally below 0.1%, though Vaidyanathan et al. found that postirradiation heat treatments (typi­cally of 4 h at 785 °C) produced a moderate recovery in ductility. Based largely on observations reported by Yang81 for irradiated ST PE16, Vaidyanathan et al. and Huang and Fish considered that the irradiation — induced embrittlement of precipitation-hardened alloys could generally be attributed to the formation of brittle g0 layers at grain boundaries. However, the arguments presented were far from conclusive — microstructural examinations of the developmental alloys which were reported by Vaidyanathan et al. showed only weak indications of g0 precipitation in D25 even within the grains, and evidence for g0 pre­cipitation at grain boundaries in D68 was not found in the low ductility tested samples but only in material irradiated to a higher fluence. Yang81 examined the microstructure of a low Si (0.01%) heat of ST PE16, which was irradiated in EBR-II to doses of about 30 and 50 dpa at temperatures from 425 to 650 °C. Grain boundary g0 layers were observed in ST PE16 samples which were irradiated at 510 °C or above but not at 425 °C, and helium bubbles were detected at bound­aries in samples irradiated at 600-650 °C. It was considered by Yang that the irradiation-induced embrittlement of ST PE16 was mainly attributable to the cleavage fracture of grain boundary g0 layers and that any effects of helium were of secondary impor­tance. However, although grain boundary precipitation of g0 was observed by Boothby28 in PE16 irradiated to relatively high doses in EBR-II, there was no evidence for the formation of intergranular g0 layers in the aged conditions of PE16 which exhibited low ductility fail­ures following irradiation in DFR to ^20 dpa.89 Thus, although it remains possible that the formation of grain boundary g0 layers may aggravate the embrittlement, it was considered by Boothby28 that the irradiation embrittlement of PE16 is primarily due to helium.

A breach in solution-annealed Inconel 706 fuel pin cladding, irradiated to 5% burn-up in EBR-II, was reported by Yang and Makenas.92 The rupture occurred from 12.7 to 18.4 cm from the bottom of the pin, corresponding to irradiation at 447-526 °C at a fluence of 6 x 1026nm~2 (E > 0.1 MeV). Microstruc­tural examinations revealed a brittle intergranular fracture, with failure being attributed to a combina­tion of matrix hardening due to g0 precipitation and grain boundary weakening due to the formation of interconnected Ni3(Ti, Nb) Z phase particles. In con­trast to the work of Rowcliffe and Horak86 where grain boundary Z phase was precipitated during a preirradiation aging treatment, this phase formed during the irradiation period in the solution — annealed cladding. Precipitation of Z was considered to be irradiation enhanced because it was not formed in long-term thermal annealing experiments at 480­540 °C. Grain boundary precipitation of Z phase was also observed at the hot (650 °C) end of the fuel pin cladding, with both g0 and g00 in the matrix.

Cauvin et al.93 and Le Naour et al.94 also attributed irradiation embrittlement effects in Inconel 706 cladding to the combined effects of matrix hardening and the precipitation of Z at grain boundaries. Inconel 706 fuel pin cladding, fabricated from four heats with Nb contents varying from 1 to 3% and in two heat-treated conditions (solution annealed or aged), was irradiated in the Phenix fast reactor up to a maximum of 100 dpa. Tensile tests on cladd­ing sections were carried out at a strain rate of 3 x 10-4 s-1. Tensile tests performed at ambient tem­perature showed high UTS (>1000 MPa) along the full length of the pins with peak values of 1500 MPa in sections irradiated near 500 °C; ductility values (uniform elongations only were given) remained low (<2%) for irradiation temperatures up to 550 °C and then increased sharply. In tests carried out at the irradiation temperatures, however, prema­ture failures occurred above about 500 °C, with ten­sile strengths reduced to 300 MPa and ductilities close to zero. All of the materials examined by Cau — vin and Le Naour et al. showed similar properties, with no systematic influence of composition or heat treatment. Plitz et a/.57 listed three fuel pin failures in Inconel 706 cladding in the Phenix reactor, with a further 12 failures in austenitic steel cladding; two of the Inconel 706 failures were associated with long periods of low power operation followed by a rise to full power conditions, resulting in mechanical inter­action between the fuel and the low ductility cladding material.