The microstructure of fast neutron irradiated Nimonic PE16

N

ELSEVIER

Journal of Nuclear Materials 230 (1996) 148-157

The microstructure of fast neutron irradiated Nimonic PE 16 R.M. B o o t h b y * AEA Technology, Harwell, Oxfordshire OXI I ORA, United Kingdom

Received 1 June 1995; accepted 24 January 1996

Abstract The microstructures of a number of conditions of Nimonic PE16, irradiated in EBR-II at temperatures from 394 to 636°C to displacement doses up to 74 dpa, have been examined. Volume changes, determined from density measurements, were less than 1% in the solution treated and aged (STA) and overaged (OA) conditions. Specimens which were cold worked by 20% or more prior to aging exhibited no significant void swelling. Redistribution of ~/' to point defect sinks occurred, particularly at higher temperatures, resulting in a skeletal ~/' form in both the STA and OA conditions. In cold worked conditions the ~/' precipitates tended to retain a more equi-axed shape. Continuous grain boundary layers of ~' were produced in STA PE16 at intermediate irradiation temperatures. Grain boundary precipitation of M6C and Laves phase occurred at higher temperatures and was more prevalent in a high-boron alloy. Mass spectrometry measurements confirmed that the helium generation rate in PE16 was about 1.2 appm per dpa, and assessments of the total helium contents of gas bubbles yielded a value for the bubble surface energy of 3 J m -2. The mechanism of irradiation embrittlement in PE16 is discussed in the light of these observations.

1. Introduction The high creep strength [1] and relatively low void swelling rate (see e.g. Refs. [2-4]) of Nimonic PE16 led to its gradual emergence from back-up status [5] to being the UK's favoured alloy for fuel element cladding in PFR and in conceptual advanced fast reactor designs [6-9]. Some experimental PEl6-clad fuel pins attained over 21% bumup (up to 150 dpa) in PFR [8,9]. Nevertheless, this type of material is widely regarded as being susceptible to irradiation embrittlement and continuing development work is necessary to ensure that the integrity of PE16 cladding can be routinely maintained to the high burn-ups which will be required in commercial fast reactors. Detailed microstructural examinations of irradiated PE16 are an essential part of this work. In this paper the effects of heat treatment, cold working and boron content on precipitation effects, void swelling behaviour and helium bubble formation in EBR-II irradiated PE16 are examined. Previous work on

* Corresponding author. Tel.: +44-1235 434 356; fax: +441235 435 941; e-mail: [email protected].

irradiation embrittlement in PE16 is briefly reviewed, and the embrittlement mechanism is discussed in the light of the microstructural observations made.

2. Experimental procedure Alloy compositions and heat treatments are shown in Tables 1 and 2, respectively. Cast DAA766 was irradiated in solution treated and aged (STA), overaged (OA), and three cold worked and aged (CWA) conditions. Cast Z260D was irradiated in the STA condition only. The alloys were irradiated in row 2 of the EBR-II fast reactor throughout runs 1-3 of the UK-1 experiment. Specimens were irradiated in the form of 3 mm diameter discs of thickness 0.38 mm (for subsequent density measurements and helium determinations) and 0.13 mm (for preparation of electrontransparent foils). Atomic displacement doses (NRT-Fe) were calculated from fast neutron fluences using the following averaged (along axial position) relationship for row 2 irradiations: 10 26

n / m 2 ( E > 0.1 MeV) = 4.72 dpa.

0022-3115/96/$15.00 Copyright © 1996 Elsevier Science B.V. All rights reserved PII S0022-3 115(96)00169-9

(1)

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Table 1 Nimonic PE16 cast analyses (wt%) Cast

Ni

Cr

Mo

Ti

A1

Si

Mn

C

N

B

Fe

DAA766 Z260D

43,5 45.8

16.5 16.8

3.17 3.50

1.18 1.35

1.15 1.20

0.15 0.26

0.05 na

0.05 0.07

0.02 na

0.0018 0.0070

bal. bal.

(na: not analysed).

The displacement doses ranged from 44 to 74 dpa at dose-weighted temperatures from 394 to 636°C. Thin foils were prepared by jet electropolishing in a solution of 5% perchloric acid in methanol cooled to - 50°C, and were generally examined in a 300 keV Philips EM430 transmission electron microscope (TEM). Some additional examinations using energy dispersive X-ray analysis (EDX) to identify precipitated phases were carried out using 100 keV Philips EM400 or VG HB501 microscopes. The EM400 TEM utilised an Edax SW9100/60 EDX system, whilst the HB501 S(scanning)TEM used a Link analysis system. Foil thicknesses, for evaluations of void or helium bubble concentrations, were determined by stereo microscopy. Voi~d sizes were determined by measuring their inscribed diaraeter d. The mean void volume V is then given by V =f(~,d3/n)

=fd3v,

(2)

and the swelling S by S = V N v / ( 1 - VNv),

(3)

where f is a shape factor, n is the number of voids measured and N v is the void concentration. The voids observed in PE16 were generally cuboidal with truncated corners/edges, and the shape factor f, which would be I r / 6 for perfectly spherical voids and 1 for perfectly cuboidal ones, was therefore estimated to be 0.9. Errors in swelling values determined from TEM examinations are estimated to be about +_25%. Density determinations were made using Archimedes' principle, weighing in air and in water using a Mettler ME22 balance positioned in a vibration-free, temperature controlled environment. Specimens were electropolished to remove any scale prior to weighing, The fractional density change (which may include contributions from precipitation as well as void sveelling) is defined as ( Po - P i ) / P i , where Po and Pi are the unirradiated and irradiated densities respectively. Defined in this way, the density change is, like the TEM swelling data, equivalent to the volume change relative to the initial volume. The uncertainty in the density change data is estimated to be ±0.2 percentage points. The 4He content of irradiated PE16 was measured by mass spectrometry using added 3He as an internal standard. Gas release was achieved by dissolving the samples in a molten mixture of vanadium pentoxide and sodium metavanadate held at 990°C for 2 h. Evolved gases were

scrubbed using a zirconium getter pump and a charcoal trap cooled to liquid nitrogen temperature before being admitted to the mass-spectrometer. Unirradiated PE16, in which no significant 4He was found, was used to set up the apparatus prior to measurements being made on irradiated samples.

3. Results

3.1. Precipitate structures

Prior to irradiation the precipitate structures of the two STA casts of PE16 and the three CWA conditions of cast DAA766 (all of which were aged for 4 h at 750°C) were similar, the main features being intragranular y' of mean diameter 9 nm and grain boundary M 2 3 C 6 . The more complex OA treatment applied to cast DAA766 resulted in a coarser 3" dispersion (mean diameter 33 nm), with rows of MC precipitates located intragranularly at the sites of prior grain boundaries but with the current boundaries largely carbide-free. In addition the alloys contained primary particles, predominantly (Ti, Mo)C and TiN in cast DAA766 but also (Mo, Cr)3B 2 in the high boron cast Z260D. As expected from previous microstructural examinations of fast reactor exposed PEI6 (see e.g. [10-13]), redistribution of the y' phase to point defect sinks occurred during irradiation. The extent of the "y' redistribution depended on the thermo-mechanical condition of PE16 as well as the irradiation temperature. Fig. 1 shows typical -,/' structures in STA and CWA cast DAA766 irradiated to

Table 2 Alloy heat treatments Designation Treatment STA CWA OA

20 min/1020°C + 4 h /750°C 20 min/1020°C + cold work (10, 20 or 40%)+4 h /750°C 1 h/1100°C + 50% cold work + 1 h/900°C + 30% cold work + 1 h/900°C, cool at 40°C/h to 750°C, hold 16 h at 750~C, cool at 40°C/h to 650°C then air cool.

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?

l

Fig. 1. ~/' structures in PE16 (cast DAA766) irradiated to 54.8 dpa at 438°C. (a) STA, (b) 10%CWA.

54.8 dpa at 438°C. Little or no coarsening of the original ~/' dispersion occurred in the STA alloys at this temperature, but irradiation-induced ~,' was precipitated at dislocations and void surfaces. Comparison of the ",/' structures in the STA and CWA conditions showed a generally coarser but more uniform precipitate dispersion in the latter. Figs. 2 and 3 show examples of ",/ dispersions in STA PE16 irradiated at higher temperatures. Virtually continuous layers of "V' were formed at grain boundaries in the STA condition irradiated to 44.4 dpa at 477°C. EDX analysis of one sample (STA Z260D) using the HB501 STEM indicated that the approximate composition of the grain boundary ",/ phase was Ni3(Al0.6, Ti0.25, Si0.15 ). Intragranular ~,' structures remained similar to those at 438°C (Fig. 1) for irradiations at temperatures up to and including 513°C. Above this temperature the ~,' phase was

Fig. 2. ~' structure in STA DAA766 i~adiated to 44.4 dpa at 477°C.

Fig. 3. "y' structures in STA Z260D irradiated to 69.4 dpa at 570°C.

almost entirely redistributed to point defect sinks. The spherical ~/' precipitates formed during thermal aging which were retained at low irradiation temperatures were replaced by a 'skeletal' form decorating dislocations and voids at high temperatures. Skeletal ~/' precipitates in Z260D, irradiated to 69.4 dpa at 570°C, are shown in Fig. 3. Although skeletal precipitates were the dominant form of ~/' at this temperature (Fig. 3(a) and (c)), spherical precipitates were retained in regions close to mobile grain boundaries (Fig. 3(b)). Cellular ~/' was generally left in the wake of migrating grain boundaries (Fig. 3(b)) and ~/' layers again coated static boundaries (Fig. 3(c)). Increased grain boundary mobility at higher irradiation temperatures restricted the range over which continuous boundary 3,' layers were formed to 477 to 570°C. Note that ~' coated grain boundaries were observed only in the STA alloys and not in the CWA or OA conditions. Typical ~' dispersions in CWA DAA766 at high irradiation temperatures are compared with the skeletal form found in the STA condition in Fig. 4. At 570°C (Fig. 4(a))

R.M. Boothby / Journal of Nuclear Materials 230 (1996) 148-157

151

and 600°C (Fig. 4(b)) ,dislocations in CWA samples were decorated by groups of irregularly shaped ~/' precipitates. Once again the original spherical ~/' dispersion was retained near grain boundaries (Fig. 4(a)). At the highest irradiation temperature of 636°C there was a marked contrast between the coarse skeletal "y' structure in the STA sample (Fig. 4(c)) and the rounded precipitates present in CWA specimens (Fig. 5(d)). In DAA766 in the OA condition the thermally aged "y' structure was largely retained, with some secondary precipitation at intragranular point defect sinks, at irradiation temperatures up to 570°C. At 600°C and above a skeletal ~/' structure developed. In DAA766 (all conditions) grain boundary Cr-rich M23C 6 precipitates were accompanied by the M o - S i - C r Ni phase M6C at temperatures of 513°C and above. Fig. 5(a) shows the two intergranular carbide phases in a CWA sample irradiated at 636°C. As noted previously [14], the composition of M6C in PE16 was dependent on the irradiation temperature, being more enriched in Mo but less so in Cr at higher temperatures. A small amount of grain boundary Laves phase - - another Mo-Si enriched phase, which like M6C also forms during prolonged thermal ageing [15] - - precipi~Eatedin CWA DAA766 irradiated at 636°C. Laves phase occurred more abundantly in CWA

Fig. 5. Carbide and intermetallicphases in PE16 irradiated to 74.1 dpa at 636°C. (a) Grain boundaryM6C and M23C6 in 10%CWA DAA766; (b) intragranularLaves phase in 10%CWA DAA766; (c) grain boundaryM6C and Laves phase in STA Z260D.

Fig. 4. ~/' structures in PEI6 cast DAA766; (a) 10%CWA, 69.4 dpa/570°C; (b) 40%CWA, 72.7 dpa/600°C; (c) STA, 74.1 dpa/636°C; (d) 40%CWA, 74.1 dpa/636°C.

(and occasionally in STA and OA) DAA766 irradiated at 600 and 636°C in an intragranular form nucleated at primary MC particles (Fig. 5(b)). MrC and Laves phase were more prevalent in STA Z260D than in similarly heat treated and irradiated DAA766. Thus, in Z260D irradiated at 636°C, grain boundary Laves phase and M6C precipitates had largely replaced M23C 6 (Fig. 5(c)). An earlier investigation of grain boundary segregation in PE16 during heat treatment and thermal neutron irradiation showed that molybdenum levels at the boundary are related to the boron content and degree of boron segregation in the alloy [16]. The increased tendency for Mo-rich phases (M6C and Laves) to form at grain boundaries in EBR-II irradiated Z260D (70 ppm B) compared with DAA766 (18 ppm B) is consistent with this effect. Primary boride particles, (Mo, Cr)3B 2 , tended to occur in widely dispersed clusters in Z260D and were only encountered in foils irradiated at 477, 540 and 636°C.

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Table 3 Density change (%) in irradiated Nimonic PE16 Temp.

dpa

DAA766 STA

OA

10%CWA

20%CWA

40%CWA

44.4 54.8 44.4 58.1 66.6 69.4 72.7 74.1

0.44 0.36 0.39 0.61 0.41 0.22 0.30 0.30

0.22 0.56 0.48 0.98 0.95 0.41 0.70 0.21

-0.18 - 0.09 0.04 0.09 0.14 -0.12

-0.15 - 0.17 - 0.20

-0.12 - 0.13 - 0.12 -0.14 -0.15 -0.07

(°C) 394 438 477 513 540 570 600 636

-0.16 -0.21 - 0.18 - 0.14

- 0.22

Transmutation of boron by the l°B(n, ot)7Li reaction gave rise to localised damage regions around the boride particles. Fig. 6(a) illustrates such damage around a primary boride particle in PE16 irradiated at 5400C. The ranges of the lithium ions and a-particles formed in the transmutation reaction are about 1.3 and 2.7 I~m respectively, giving rise to two concentric damage shells around a boride particle [17,18]. In practice these two damage regions may overlap if the boride particle (i.e. the volume within which the transmutation occurs) is sufficiently large. In previous work on an austenitic steel the bright-contrast features evident in the inner damage region corresponding to the range of the lithium ions were thought to be voids [18]. More recently, however, Dumbill and Bishop [19] have examined the region around boride particles in thermal

- 0.02

Z260D STA 0.44 0.33 0.33 0.46 0.27 0.19 0.34 0.04

neutron irradiated PE16 (cast Z260D) using high resolution secondary ion mass spectrometry and have identified lithium precipitates in the inner damage shell. Moreover, TEM examination showed that the bright-contrast features in this region were retained after a high temperature (800°C) annealing treatment that eliminated voids [19]. It is therefore considered that the large (typically 50 to 70 nm diameter) bright-contrast features evident near the boride particle in Fig. 6(a) are lithium precipitates rather than voids. As expected, a high density of small helium bubbles (about 5 to 10 nm diameter) is visible in regions corresponding to the range of a-particles. An additional feature of the specimen irradiated at 540°C, though not those irradiated at 477 or 636°(2, was the presence of small faulted precipitates near the primary

Table 4 TEM estimates of void swelling Temp. (°C)

dpa

Void parameter

394

44.4

438

54.8

477

44.4

513

58.1

540

66.6

570

69.4

Nv (m -3) d v (rim) S (%) Nv (m -3) d v (nm) S (%) Nv (m -3) d v (nm) S (%) Nv (m -3) d v (rim) S (%) Nv (m -3) d v (nm) S (%) Nv (m -3) d v (nm) S (%)

600

72.7

Nv (m -3)

dv (nm) S (%)

DAA766 STA

2.2 X 1020 29.7 0.51 1.6 X 10 20 40.2 0.95

7.2 X 1019 52.2 0.93

2.9 X 1019 50.3 0.33

DAA766 OA

DAA766 10%CWA

Z260D STA

2.5 × 1020 23.3 0.29 1.6 X 1020 35.8 0.67

2.1 X 1020 17.3 0.10 1.9 X 1020 23.8 0.23 5.5 X 1019 32.2 0.17

5.1 × 23.2 0.58 2.2 × 33.4 0.74 1.1 X 43.3 0.81 8.2 X 44.8 0.67 6.0 X 36.5 0.26

1.2 X 1020 51.8 1.62 6.4 X 1019 58.6 1.31 5.3 X 1019 57.9 0.94 3.0 X 1019

43.5 0.22

6.0 X 1019 45.0 0.50

1020

102°

1020

1019

1019

R.M. Boothby/ Journal of Nuclear Materials 230 (1996) 148-157

153

indicated by the density measurements, particularly when the latter are small or negative, by about 0.2 to 0.5 percentage points. Thus, a general densification effect (attributable to precipitation) of this magnitude is suggested by the data. Examples of the void structures observed in DAA766 are shown in Fig. 7. Voids in PE16 tended to be cuboidal with truncated comers and/or edges. Void concentrations and sizes in the STA and OA conditions were similar, although in the latter some larger voids, which could account for marginally higher swelling, were associated with the rows of carbide particles which delineated prior grain boundaries as shown in Fig. 7(b) and (c). Surprisingly, EDX and diffraction analyses of these carbides in irradiated OA material revealed them to be M23C 6 (or occasionally M6C at the higher temperatures) - - not in a cube/cube orientation relationship with the matrix - rather than the MC particles present in the unirradiated condition. TEM examinations of CWA DAA766 revealed

Fig. 6. (a) Primary M3B2 boride particle and damage region and (b) M2B precipitates in STA Z260D irradiated to 66.6 dpa at 540°C.

borides. Such precipitates, which are evident in Fig. 6(a) and shown in more detail in Fig. 6(b), were identified from EDX analyses and electron diffraction patterns to be the Cr-rich phase M2B (tetragonal structure with a = 0.52 nm and c = 0.43 nm). 3.2. Void swelling Volume changes determined by density measurements are given in Table 3. Density changes of less than 1% were measured in these samples. Comparison of the results for DAA766 in the STA and OA conditions revealed only a slight dependence of void swelling on heat treatment. Swelling tended to be marginally higher in the OA condition at temperatures from 438 to 600°C but, in view of experimental uncertainty, the increase was barely significant except perhaps at 513 and 540°C. The CWA samples of DAA766 generally exhibited slight densification on irradiation, though marginal volume increases were registered in 10%CWA specimens irradiated in the temperature range 477 to 540°C. Swelling in STA Z260D was generally low and comparable to that in similarly heat treated DAA766. Void swelling values determined from TEM measurements are given in "Fable 4. Agreement with the trends shown by the density data is good in view of the large errors inherent in the TEM measurements. The TEM estimates of void swelling tend to exceed the volume changes

Fig. 7. Void structures in PEt6 cast DAA766; (a) STA, 66.6 dpa/540°C; (b) OA, 66.6 dpa/540°C; (c) OA, 58.1 dpa/513°C.

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Table 5 Helium generation rate in Nimonic PE16 (cast DAA766) irradiated in EBR-II Element

Atomic conc. (%)

Reaction rate (appm He per 10 26 n / m 2)

Helium generated (appm per dpa)

Ni Fe Cr AI Ti Mo Si C N I°B Total

41.5 34.4 17.8 2.39 1.38 1.85 0.30 0.24 0.08 0.00174

9.63 0.675 0.262 1.57 0.337 0.170 6.23 4.45 231 9710

0.98855 0.05744 0.01154 0.00928 0.00115 0.00078 0.00462 0.00264 0.04571 0.04179 1.1635

appreciable void concentrations, and a peak swelling value of 0.5% at 540°C/66.6 dpa, at the 10% cold work level, but only irregular dispersions of voids and insignificant swelling in the more heavily worked samples. In practice, heavily cold worked conditions of PEI6 may not be suitable for use as fuel element cladding because a moderate degree of void swelling can be helpful in reducing fuel/clad interaction stresses. However, for sub-assembly wrappers, where dimensional stability is a primary concern, the minimal swelling rate of CWA PE16 could be beneficial. 3.3. Helium bubbles

Helium bubble dispersions were examined in CWA samples irradiated at the three highest temperatures. Irradiations at these temperatures were carded out in near centreline positions in row 2 of EBR-II. Helium generation rates for the centreline position are given by Gabriel et al. [20] for a reactor power of 62.5 MW and can be used to estimate the amount of helium produced in PE16. The amount of helium generated from each component of the alloy is given by multiplying the appropriate reaction rate by the total neutron fluence and by the atomic fraction of the element. The fast neutron fluence ( E > 0.1 MeV), which was measured in the current experiment, was taken to be 85.65% of the total neutron fluence in EBR-II. Using the relationship between fast neutron fluence and dpa, Eq. (1), it may be deduced that a total neutron fluence of

2.474 X 10 25 n / m 2 is equivalent to 1 dpa. As indicated in Table 5, the calculated helium generation rate for PE16 cast DAA766 is 1.1635 appm per dpa, about 85% of which is generated from nickel transmutation. The total amount of helium generated is therefore estimated to have been 80.7 appm at 570°C/69.4 dpa, 84.6 appm at 600°C/72.7 dpa and 86.2 appm at 636°C/74.1 dpa. Variations in reactor power and deviations from the ideal centreline position give rise to some uncertainty in these calculations. As a check, therefore, actual helium concentrations were measured using mass spectrometry. The measured helium contents of 75.1 appm for the 570°C irradiation, 87.0 appm at 600°C, and 82.6 appm at 636°C showed reasonable agreement with the calculated values. Fig. 8 shows examples of the intragranular helium bubble dispersions in CWA PE16 irradiated at 600 and 636°C. Bubbles were imaged in an underfocused condition and sized by measuring the inner diameter of the dark Fresnel fringe. Number densities were determined by stereomicroscopy. Mean bubble diameters and number densities are given in Table 6. Assuming the measured diameters to be those of bubbles in which the gas pressure was in equilibrium with the surface tension at the irradiation temperature, the number of helium atoms per bubble can be calculated. Total helium contents, calculated using the equation of state for helium given by Mills et al. [21], are shown in Table 6 for various assumed values of the bubble surface energy, Ys • Comparison of these data with the helium contents determined by mass spectrometry a n d / o r

Table 6 TEM estimates of helium contents in CWA Nimonic PE16 (cast DAA766) for different values of the surface energy Ys (J m--•) Temp. (°C)

dpa

570 600 636 636

69.4 72.7 74.1 74.1

Condition 40%CWA 40%CWA 10%CWA 20%CWA

Bubble conc. (m -3)

Mean diam. (nm)

Helium content (appm) %=1

3(,=2

%=3

2.49 X 10 21 1.85 X 1021 1.02 X 10 21 1.08 X 1021

4.05 4.72 6.04 5.72

43.8 46.2 45.6 42.2

63.8 68.7 70.3 64.6

77.1 84.0 87.4 80.0

R.M. Boothby / Journal of Nuclear Materials 230 (1996) 148-157

155

Fig. 10. Grain boundary migration in 10%CWA PEI6 irradiated to 74.1 dpa at 636°C.

Fig. 8. Helium bubble dispersions in irradiated PE16 (cast DAA766); (a)40%CWA, 72.7 dpa/600°C; (b) 10%CWA, 74.1 dpa/636°C.

calculated from gas generation rates indicates a value of about 3 J m -2 for T~ in the given temperature range. A similar value would be predicted by extrapolating Murret al.'s high temperature surface energy measurements in type 304 stainless steel to 600°C [22]. If the helium bubbles in irradiated PE16 were not in equilibrium as assumed, but instead were underpressurised owing to the onset of void swelling, then a higher surface energy than 3 J m -2 would be required to account for the generated helium. Grain boundary helium bubbles, illustrated in Fig. 9, ranged up to about 15 nm diameter in PE16 irradiated at 570 and 600°C, and 20 nm at 636°C. There was no apparent effect of heat treatment or cold working, or of intergranular precipitates, on the size of grain boundary helium bubbles. Grain boundary migration was not uncommon at these temperatures and resulted in precipitates and helium bubbles originally sited at the boundary becoming detached from it. An example of this phenomenon is shown in Fig. 10, in which the original position of the grain boundary is indicated by a row of M23C 6 precipitates. The boundary itself appears to have moved by about 0.5 Ixm, with most of the larger helium bubbles being deposited mid-way between the original and current positions of the boundary. Pinning of the grain boundary at ",/ particles is evident, and it appeared that the ",/ in the adjacent grain dissolved at the boundary and re-precipitated (in the usual cube-cube orientation relationship with the matrix) in its wake.

4. Discussion

Fig. 9. Grain boundary helium bubbles in irradiated PE16 (cast DAA766); (a) OA, 72.7 dpa/600°C; (b) 10%CWA, 72.7 dpa/600°C; (c) STA, 74.1 dpa/636°C; (d) 20%CWA, 74.1 dpa/636°C.

Although the low void swelling behaviour of PE16 is widely recognised and is a prime consideration for fast reactor applications, this type of precipitation-hardened high-nickel alloy is also perceived as being particularly susceptible to irradiation embrittlement. Post-irradiation tests measuring conventional tensile ductility [14,23,24] or ring ductility (i.e. the strain to initiate cracking in a tube compression test) [25,26] have shown severe embrittlement in fast neutron irradiated PE16 and other ~'- or ~'/'y"-

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R.M. Boothby / Journal of Nuclear Materials 230 (1996) 148-157

hardened alloys (e.g. Inconel 706) at temperatures of about 450°(? and above. In the case of Inconel 706 irradiation embrittlement has been attributed to a combination of matrix hardening due to 3" precipitation and grain boundary weakening due to the formation of intergranular -q phase (hexagonal structured Ni3(Ti, Nb)) [27-29]. Similarly, Yang [23] and Vaidyanathan et al. [25] have suggested that the formation of grain boundary 3" layers is responsible for the severe irradiation embrittlement found in PE16 and similar alloys. However, it should be noted that these observations were made on material irradiated in the solution treated condition, which therefore had more 3,'-forming solutes available to segregate and precipitate at grain boundaries than would have been the case if the material had been aged. Vaidyanathan et al. [25] showed that the ductility of irradiated PE16 was partially recovered by a thermal soak of 4 h at 785°C and argued that this result was not consistent with a helium embrittlement mechanism. However, the possibility that helium bubbles could have become detached from migrating grain boundaries during the thermal treatment was not considered. In the present work it has been shown that grain boundary layers of 3,' form in STA PE16 over a restricted range (477 to 570°C) of irradiation temperatures. Such layers were not produced in OA or CWA material, and these conditions might therefore be expected to be less susceptible to irradiation embrittlement. Ductility data are not available for the irradiated conditions examined in the current work. However, earlier work on STA and OA PE16 irradiated to about 20 dpa at 465 to 635°C showed severe embrittlement in both conditions [14]. Embrittlement was most marked in specimens irradiated at 535°C and tested under low strain-rate conditions at 550°C, where the ductility of both STA and OA material was < 0.3%. No indications of grain boundary 3,' formation were found in either STA or OA PE16 in this lower displacement dose experiment, and the ductility loss was interpreted as being due to helium embrittlement. Although the possibility remains that the formation of grain boundary 3,' layers, or other grain boundary segregation effects, may exacerbate irradiation embrittlement, it is therefore considered that the embrittlement is primarily due to helium. It may be of interest to note that, in a proton irradiated nickel-silicon alloy, Packan et al. [30] found that the radiation induced formation of the Ni3Si 3" phase at grain boundaries did not lead to low ductility intergranular failures unless helium was preimplanted into the specimens. In practice the normal operating stresses in a reactor core component are low compared to the tensile/creep strength of the material. It has been suggested that the rupture life of a stressed core component will be controlled by the gas driven growth rate of sub-critical grain boundary helium bubbles [31]. Once the critical size is reached, the bubbles become unstable and grow by vacancy absorption - - leading to cavity coalescence and low ductility intergranular failure. In PE16, unstable bubble growth has

been observed in helium-implanted specimens which were tested under low strain-rate tensile conditions at 650°C [32]. The instability criterion for grain boundary gas bubbles was derived by Hyam and Sumner [33] assuming ideal gas behaviour. For the ideal gas case, a bubble with equilibrium radius r o becomes unstable once the stress exceeds a critical value o"c given by o"c = 0.77ys/r o . Using more realistic equations of state for helium, it can be shown that the critical stress for small bubbles is somewhat higher than indicated by the ideal gas assumption [34]. For example, for a surface energy of 3 J m - 2 and grain boundary bubbles 20 nm in diameter (i.e. the largest bubble size observed in the present experiment), o'c is about 230 MPa in the ideal gas case but 280 MPa using Mills et ai.'s equation of state. Continued irradiation would of course result in further bubble growth and a reduction in o-~ . The extent of this effect can be estimated if it is assumed that the number of helium atoms per bubble increases linearly with neutron dose (i.e. there is no bubble coalescence). For example, an increase in dose from 74 dpa (the highest exposure level in the current experiment) to 180 dpa (the target dose for commercial fast reactors [8]) would, using Mills et al.'s equation of state for helium, result in the growth of a 20 nm diameter bubble to about 30 nm with a corresponding reduction in trc to 180 MPa. The build-up of pressure from fission gases gives rise to a hoop stress in fuel element cladding which increases with time but is not expected to exceed 70 MPa throughout the element life [7]. Thus, there would appear to be an adequate margin between the applied stress and ~c, and fuel element failure due to unstable growth of grain boundary helium bubbles would not generally be expected to occur in PE16 cladding exposed to high neutron doses.

5. Summary The microstructures of a number of conditions of Nimonic PE16, irradiated in EBR-II at temperatures from 394 to 636°C to displacement doses in the range 44 to 74 dpa, have been examined. The following observations were made. (i) Volume changes derived from density measurements were less than 1% in all specimens. The swelling behaviour of cast DAA766 in the STA and OA conditions showed little dependence on heat treatment. Marginally higher swelling, which could possibly be accounted for by the presence of larger voids attached to intragranular carbide particles, occurred in the OA condition. Cold worked and aged specimens of cast DAA766 exhibited low (10% CWA) or insignificant (20 and 40% CWA) void swelling. Swelling in STA cast Z260D was similar to that in similarly heat treated cast DAA766. (ii) The original 3" dispersion was largely retained in STA PE16 at temperatures up to 438°C though irradiation-induced precipitation occurred at dislocations

R.M. Boothby / Journal of Nuclear Materials 230 (1996) 148-157 and void surfaces. At higher irradiation temperatures the ~/' phase in the STA condition was almost completely redistributed to point defect sinks, giving rise to a skeletal -/' form. The coarser "y' dispersion produced in the OA condition was relativety stable at irradiation temperatures up to 570°C but was again replaced by a skeletal structure at 600°C and above. In C W A PE16 ~/' redistribution occurred at all temperatures, but the precipitates tended to retain a more equi-axed shape. Continuous grain boundary layers of ~/' were formed only in STA material at intermediate irradiation temperatures (477 to 570°C). (iii) Grain boundary precipitation of the molybdenumrich phases M6C and Laves at high irradiation temperatures was more prevalent in the high boron cast Z260D than in DAA766. (iv) Mass spectrometry measurements of the helium content in specimens of cast DAA766 which were irradiated in near centreline positions confirmed the calculated helium generation rate of between 1.1 and 1.2 appm per dpa. Assessments of the helium bubble sizes and concentrations in C W A samples in which void swelling was negligible yielded a value for the bubble surface energy of 3 J m -2. Consideration of previously published work on irradiation embrittlement in PE16 in the light of the current microstructural obserwations suggests that helium embrittlement, rather than the formation of grain boundary 7' layers, is primarily responsible for low ductility failures in post-irradiation tests. However, unstable growth of grain boundary helium bubbles would not generally be expected to occur at the relatively low stresses to which fuel element cladding is subjected.

Acknowledgements Thanks are due to M.R. Newman for carrying out the mass spectrometry measurements of helium content, W.C. Fuller for the preparation of TEM foils, and S. Dumbill for STEM examinations.

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