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Electron damage in zirconium
313
Journal of Nuclear Materials 115 (1983) 313-322 North-Holland RubIishing Company
ELECTRON DAMAGE IN ZIRCONIUM - I. Defect structure and loop character M. GRIFFITHS,
M.H. LORETTO
and R.E. SMALLMAN
Department of Metallurgy and Materials, University of Birmingham, P. 0. Box 363, Birmingham Bl5 2TT, UK Received 18 October 1982; accepted 8 November 1982
Zirconium (499.95% Zr) has been irradiated using 1 MV electrons at temperatures between 273 and 773 K in a high-voltage electron microscope and the nature of the clustered damage assessed. It has been shown that both interstitial and vacancy loops of 6 z f (11%) nucleate and grow during electron irradiation within this temperature range. Some loops of b = 4 (1123) and
%(2023) are also observed. The significance of the observation of vacancy lcxqs is discussed in terms of current theories of defect clustering during irradiation and it is concluded that the presence of imposed stress influences the clustering.
1. Introduction
2. Experimental
There is now incont~vertible evidence that large stable vacancy loops are formed during neutron irradiation of Zr at temperatures in excess of 0.3 r, [1,2]. The loops are far too large to have been formed merely by cascade collapse and it is apparent that their existence is related to the resistance of Zr to void swelling. Cases of void growth during neutron and electron irradiation of Zr have been reported but these have been shown to be largely dependent both on specimen purity and on the presence of implanted or fission-generated inert gases such as He [1,3]. Recent work by Carpenter and Watters [4] has shown that voids are not formed during HVEM irradiation of relatively pure undoped Zr for temperat~es up to 725 K and doses up to 23 dpa. This is consistent with much of the earlier work on this material but not with the results of Buckley and Manthorpe [5] who reported voids during electron irradiation of doped and undoped Zr. Carpenter suggests that the difference is probably related to impurity effects. Vacancy loops have not been reported following electron irradiation in any previous study. This work is part of a general programme aimed at understanding the behaviour of point defects in cph metals and alloys. The present paper assesses the relevance of basic observations of loop character and clustering behaviour to void-swelling. Paper II (see p. 323) deals primarily with the c-component loop which is considered in relation to radiation growth phenomena (dimensional changes in the absence of swelling).
Single crystal, zone-refined iodide rod and Marx grade Zr of ~99.95% purity were used for this work. Discs of 3 mm diameter and - 0.25 mm thick were prepared by spark-machining and were subsequently wrapped in Ta sheet and annealed in vacua for 8 h at 1073 K. Microscope specimens were prepared by electropolishing the discs in a Struers twin-jet Tenupol using 6% perchloric acid in ethanol at a temperature below 233 K and at 18 volts. Foils having normals - 4.5’ from [OOOl]were prepared from the single crystal in order to facilitate loop analysis. Foils for which a significant number of grains had a selected orientation with respect to the foil plane could be prepared from the iodide rod. Electron irradiation was carried out at 1 MV in the temperature range of 273-773 K using a f 25O doubletilt heating stage in an EM7 at a dose rate of - lOmu dpa/s. In general, specimens were cooled to room temperature and transferred to a rt 4S’ double-tilt stage for analysis at 400 keV in the EM7 or at 120 keV in a Philips EM4OOT. Si8nificant oxidation occurs during irradiation at temperatures in excess of 623 K and the stresses caused by the oxide give rise to slip in thick regions and bending in thin regions of the foil. Oxygen depth profiles have been obtained using an Auger spectrometer and comparison between a freshly prepared and irradiated sample shows that the depth of significant oxygen permeation from the foil surface increases by a factor of - 10, and is typically - 500 A following in situ irradia-
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M. Griffiths et al. / Electron damage in zirconium
1
tion at 675 K. Contrast from the oxide complicates analysis, especially when examining small defects, and care must be taken to reduce the contrast by use of appropriate diffraction conditions. One other adverse effect of oxidation is an asymmetric imaging condition in which defect contrast is obscured for one sense of virtually all diffracting vectors and this makes loop analysis difficult. An explanation of this effect has been reported elsewhere [6].
3. Results 3.1. Point defect clustering Dislocation loops nucleate and grow during the irradiation of Zr with 1 MV electrons. No other point defect clusters were observed over a range of irradiation temperatures between 273-773 K (0.13-0.36 T,) and for doses up to - 10 dpa. Loop structures do not generally reach a near steady-state at high doses; loops coarsen to a dislocation network and loop nucleation continues. Increasing the irradiation temperature leads to a decrease in loop density and an increase in the size of the loops for a given dose rate. Damage alignment was observed parallel to the trace of the basal plane when the electron beam direction B was close to prism orientations (i.e. (lOjO) and (1120)) as illustrated in fig. 1. Tilting from basal to prism orientations shows that most of the damage is clustered in layers parallel with the basal plane. The spacing between the layers increases with increase of irradiation temperature and is comparable with the size of the loops at moderate doses ( - 0.1 dpa). The loop distribution within the layers is non-uniform and consists mainly of small strings of, typically, 4-6 loops. These loops are non-overlapping, have identical Burgers vectors and have approximately the same orientation but are not necessarily coplanar. Loop clustering in this manner is observed from the early stages of irradiation suggesting that their nucleation is influenced by a localised stress field probably from their mutual interaction. During continued irradiation adjacent loops sometimes coalesce, forming larger elliptical loops of complex shape, as illustrated in fig. 2. Burgers vector analysis showed that loops having b = 4 (112b), -f (1123) and more uncommonly b = d (2023) are present in electron-irradiated foils of Zr. All loops clustering in layers were analysed as having Burgers vectors of the type b = 4( 1120), alternatively referred to as (B) type. The (u) type loops can be distinguished from the c-component damage using a fOOO4 diffracting vector. This puts dislocations of the type b = i (1120) out of contrast and
Fig. 1. Alignment of damage in Zr in layers parallel with the basal plane following electron irradiation at - 675 K to - 0.15 dpa, g = 1120, B - (lioo].
shows clearly those having a c-component. Loops having b = f (1123) or (c + a) type loops were observed between the layers of (a) type damage. The ellipticity and sizes of the (u) type loops varied considerably because of the absence of steady-state conditions and as a result of the proximity of the foil surfaces. Sometimes a bimodal distribution of loops was observed, as shown in fig. 1 where large numbers of small loops cluster close to the foil edge separated, by a denuded zone, from the bulk of the damage within the foil. Although it is difficult to determine accurately the plane of nucleation of the loops, from observations made with B- (1213), (l~lO), or [OOOl] there is evidence to suggest that (u) type loops do not nucleate exclusively on (1210) planes and c-component loops do not nucleate exclusively on basal planes. The most probable nucleation planes for (u) type damage are then (lOi0) and (lOi1) planes. Figs. 3(a) and 3(b) show loop nucleation during irradiation using B - [OOOl] and B - (1213), respectively. The traces of relevant planes are marked. There may be some nucleation on (lzlO} planes but the (10’50) and (lOi1) planes appear to be more likely candidates and this is hardly surprising considering that they are more widely spaced than the (1210) planes even though they are corrugated. Measurements on (c + u) type loops when viewed along
M. Griffiths et al. / Electron damage in zirconium I
315
Fig. 3. Loop nucleation during
Fig. 2. Loop coalescence in Zr during electron irradiation at - 750 K. Many beam imaging conditions, B - (11~0); (a) 0.08 dpa (b) 0.1 dpa.
(l?lO) indicate that they are tilted slightly away from (lOi1) towards (lOi0) by about 10’ during the early stages of irradiation. This is consistent with loop growth on a combination of (lOi1) and (1OiO) planes. It is
electron irradiation of Zr at - 675 K. Traces of important planes are shown; (a) g = 2ii0, B - [OOOI], (b) g = ioio,B - [1213].
difficult to say whether this is also the case for (a) type damage. Evidence for c-component loops nucleating on basal planes in prism foils is sparse and nucleation on planes inclined to basal appears to be dominant. For discussion purposes the nucleation planes of (a) type loops will be considered to be { lOiO> and (107 1) whereas
the planes of nucleation for (c + (0 type loops will be considered to be {lOit>. All loops subsequently rotate during growth to less well-defined orientations. The long dislocations seen in fig. 3(b) are formed during heating of the specimen by slip of edge dislocations through the foil, leaving dislocations trapped at the coherent metal/oxide interface. Only the (a) type loops will be analysed in more detail as we are mainly concerned with the interstitial/vacancy nature of these loops at present. A more detailed account of the c-component damage is presented in paper II. 3.2. Dislocation ioop characterisation Because of the variability in experimental conditions for thin foil irradiations such as foil thickness, specimen purity (dependent also on the microscope vacuum), degree of oxidation, absence of a steady-state damage structure etc., no attempt will be made to give a statistical account of the experimental observations. Some quantitative results are being processed, however, and will be presented in a future publication. Adverse imaging conditions associated with asymmetric contrast tends to bias loop analysis especially for loops giving g 1b = 1 images. Thus the loop marked with an A in fig. 4 shows clear outside contrast with g =I 1120 compared with g = ii20. Loop B, however, can be seen clearly with g = 1120 but is obscured with g = 1720 and it is difficult to say whether it exhibits inside or outside contrast. The net effect is that, dependin8 on the direction of contrast asymmetry, analysis of the loops can be biased in favour of interstitial or vacancy loops. Both interstitial and vacancy loops have been analysed using the Maher and Eyre 171 and Fiill and Wilkens [S] techniques. As a result of the adverse conditions already mentioned only a small percentage ( - 10%) of the total loop population could be analysed with confidence. Many of the loops had a large screw component and this made it imperative that “safe” regions for analysis be determined for each particular loop and that the loop plane should be determined wherever possible. It should be emphasised therefore that two different methods have been used to determine the nature of the loops and the loop plane normal has been obtained to within & 20” for all the loops which have been analysed. No assumption has been made concerning the maximum shear component. At temperatures below 575 K the loops were generally too small to analyse with confidence. Examples of
Fig. 4. Asymmetric (f g) contrast from a Zr foil electron irradiated at -675K.LoopAhasb=+~[ll~O]and~oopB has b= *If[lilO]; (a)g=llTO, B-[IiOIj,(b)g=ii20,B-
[IiOlJ.
loop characterisation in three higher temperature mes will now be described in detail.
regi-
M. Griffiths et al. / Electron damage in zirconium I 3.3. irradiation
at - 585 K, foil normal N - [If131
The loops 1, 2, 3 and 4 in fig. 5 are out of contrast with g = Oil0 and g= OLil and thus have Burgers vectors of b = & f[21 lo]. They are almost edge-on close to B - [ 121 l] and micrographs taken over a large tilt range show that they have habit planes with n - [i103]. They show outside contrast for g = iOl0 at B - [iZiS] and therefore have 6 = $1 lo], indicating that they are interstitial in nature. No loops in thicker regions of the foil could be analysed because of adverse imaging conditions. The majority of large loops having b = #l lo] lie close to the foil edge and the partition between these and the rest of the loop population is quite clear. No vacancy loops were identified.
oiio
ioio tooa
317
3.4. rrrud~~tion at - 675 K, foil normaI N - [Jill]
A mixed population of vacancy and interstitial loops was observed following irradiations at this temperature. The loops marked 1, 2, 3 and 4 in fig, 6 are out of contrast with g= TO11 and g= 0004 and have b = &$f[l210]. They show outside contrast with g= ilO1 and have habit planes of n - [Oi 1l] and are therefore vacancy in nature with b = #sio]. Loop 5 is out of contrast with g = ?Oli, has n - [OilO] and exhibits outside contrast with g = 1iOi. It therefore has b = +[I2101 and is interstitial in nature. The majority of damage has b = f $[iZ‘iO] and most of this is comprised of vacancy loops. Only a small percentage of vacancy loops ( < 20%) have Burgers vectors other than b =f[i2iO]. The major-
ioio
Fig. 5. Chara~te~sation of dislocation loops in Zr electron irradiated at - 585 K to - 0.1 dpa. Loops 1,2,3 and 4 all have R - [i103], and are interstitial in nature, g and B are indicated for each figure.
b = f[Sl lo]
318
M Griffiths et al. / Electron damage in zirconium
I
iioi
To0
[oili]
[?213]
p201
toil
ooo&
Fig. 6. Characterisation of dislocation loops in Zr electron irradiated at - 675 K to - 0.15 dpa. Loops 1,2, 3 and 4 haven - [OilI], character. Loops 6 and 7 have b = f[i2iO] and are vacancy in nature. Loop 5 has n - [OTlO], b = f[l?lO] and has interstitial n - [iOl I], b = f[ii20] and are interstitial in nature. Loop 8 has n - [ 1 IOO], b = 4[ZiiO] and is also interstitial in nature, g and B are indicated for each figure.
319
pit
1120
iJ
iori
[liOl]
Fig. 7. Character of dislocation loops in Zr electron irradiated at - 730 K to -0.1 dpa.ioops 1 andzhaven-[IiOl],b=~~~iO] and are vacancy in nature. Loops 3 and 4 have n - [Oli2], 6 = $1201 and are interstitial in nature. Loops 5 and 6 have II - [1230], b = $I?lO] and are also interstitial in nature, g and 3 are indicated for each figure.
ity of interstitial loops have Burgers vectors of b =j(ll~O) or +(ZiiO). Loops 6 and 7 are out of contrast with g = ilO but show inside contrast with g = IOil. They have n - [iOl lf which shows that they
are interstitial in nature having b = f[ ii20]. Loop 8 is out of contrast with g = Olil and shows outside contrast with g = IOil. It has n - [IiOO] and therefore is also interstitial in nature having b = j[2iiO].
It is apparent that interstitial loops are more common close to the edge of the foil, although this may be a result of the difficulty in analysis of similar loops within the dense dislocation network found in thicker regions of the foil. Even so, interstitial loops account for perhaps only 20-30% of the total loop population. 3.5. ~rra~ia~io~at - 730 K, N - [Oj23] Vacancy and interstitial loops were observed following irradiation at this temperature. Loops 1 and 2 in fig. 7 are out of contrast with g = iOli, have habit planes n - f 1101J and exhibit outside contrast with g = i2i2 and therefore have Burgers vectors of b = +[iZiO] and are vacancy in nature. Loops 3 and 4 are out of contrast withg = 1101, have habit planes n - [Ol]Zj and exhibit outside contrast with g= i2i2, They have Burgers vectors b = f[ 11201 and are interstitial in nature. Loops 5 and 6 are out of contrast with g =i lOjO, have n - []230] and exhibit outside contrast with g = 1212 and therefore have b = b[lzlO] and are also interstitial in nature. Clearly, as with the previous example of fig. 6, there is a partition of interstitial loops close to the edge of the foil and in this example these loops are markedly elliptical. There is no predominance of one particular
Fig. 8. Climb of dislocations at the metal/oxide interface during irradiation of Zr at 675 K, g =ilOl, B - 112131.
Burgers vectors in contrast to the previous example. The majority of loops appear to be vacancy in nature (provided that they have a safe orien~tion) since they exhibit outside contrast for g = i2i2. Unfortunately the majority of loop habit planes could not be ascertained accurately enough for confident loop analysis and therefore it is not possible to estimate the number of interstitial and vacancy loops. It is apparent, however, that no vacancy loops were observed amongst those anaiysed close to the edge of the foil. Also it is highly probable, on the basis of the observations, that the majority of the loops (say > 70%) in the bulk of the foil are vacancy in nature.
4. Discussion Vacancy loops are not expected to nucleate and grow during electron irradiation except under very specific conditions e.g., either in very thin areas [9], where the more mobile interstitials reach the free surface leaving a supersaturation of vacancies, or in the compression region of a dislocation, for example close to the edge of a large interstitial loop /lo]. One material in which vacancy loops have been reported following electron irradiation in the absence of the above conditions is vanadium. Shiraishi et al. [ 111 observed vacancy loop growth during electron irra~ation of vanadium at temperatures > 0.3 T, and they attributed this result to the effect of impurities in the matrix (0.15%). Their results suggest that unfaulted vacancy loops are favoured energetically over voids in irradiated vanadium although under normal circumstances any vacancy loops are unstable, attracting interstitials as a result of the dislocation bias term [ 121. Although the abundant interstitial impurities such as 0 and C may play a role in pinning self-interstitial atoms in V, it seems unlikely that this in itself is enough to stabilise the vacancy loops. Vacancy loop stability may be the result of impurity segregation. Impurities can change the interstitial bias for vacancy or interstitial loops depending on whether the solute segregates to the compression or dilational side of the dislocation line in a manner similar to that proposed by Bentley et al. for oversized solutes in neutron irradiated TZM 1131 or Tjhia et al. for undersized solutes in electron irradiated Al-Si alloys [14]. The results presented here show beyond doubt that vacancy and interstitial loops of b = 4( 11~0) nucleate and grow during electron irradiation of Zr. It therefore follows that cascades are not essential for the nucleation of vacancy loops - although of course the vacancy loops
observed in neutron irradiated Zr may be formed in cascades. Two questions need to be answered concerning vacancy loop growth: first, how do vacancy loops nucleate in electron-irradiated samples and second, how do vacancy loops grow during electron and neutron irradiation? 4.1. Nucieatidn
of vacancy
loops during electron irradia-
tion
A consideration of planar geometry and the observations made during electron irradiation of Zr suggests that the most common nucleation plane is (lOiO}. The {lOiO} planes are the most widely spaced planes because of the low c/a ratio of 1.59 and a hard sphere model suggests that coincident collapse and shear is reasonably likely on these planes, giving loops _of b = 4 (1120). Nucleation is also possible on the (1011) planes which have a slightly lower packing density than {IOiO} planes. Since vacancy loops have been identified in electron irradiated samples it is clear that a sufficient vacancy supersaturation exists locally to nucleate the loops. This perhaps suggests that there is a rapid, efficient removal of interstitials and this may be associated with the presence of dislocations generated by slip (cf. figs. 6 and 7). Indeed fig. 8 shows that these dislocations can climb significantly during irradiation. It should be noted, however, that such climb is not noticeable for heavily oxidised samples where they are strongly pinned at the interface, and there are insufficient dislocations to account for widespread vacancy clustering. A high vacancy supersaturation may result from pinning of interstitial atoms by impurities and the impurities themselves may enhance the stability of voids or loops. A small vacancy cluster is more likely to collapse to give a loop because most gases, such as 0 and N, are highly soluble in Zrand cannot therefore act as stabilisers for void nuclei. It is significant that such gases are very soluble in V in which vacancy loops are also observed, however it is not obvious that gas solubility is a sufficient prerequisite for vacancy loop stability. One possible factor contributing to loop stability is the segregation of impurities to the loops. There is some evidence that there is considerable impurity segregation to loops in electron- and protonirradiated Ti for a given purity and irradiation temperature [ 151. Such loops have irrational Burgers vectors but the same cannot be said for loops in Zr which appear to have rational Burgers vectors and for which segregation has so far been undetectable. In the absence of reliable evidence for impurity segregation to the loops in Zr alternative explanations for vacancy loop stability must
be sought, although impurities may indeed be important. The fact that vacancy loops nucleate may imply that there is an energy gain to the system and this would be the case if the nucleation (and subsequent growth) were to relieve an appropriate stress. There is clearly a stress induced by the growing oxide (viz., the dislocations in figs. 6 and 7). This stress arises because of the volume increase during oxidation (Pilling-Bedworth ratio = 1.5) and because of the lattice expansion close to the oxide-metal interface caused by the solution of interstitial oxygen. The stress may be expected to be tensile within the plane of the foil and compressive in the oxide. It is notable that interstitial loops only have been observed in electron irradiated ruthenium at 0.39 T, [16]. It has a similar c/a ratio to Zr but does not suffer from oxidation to any appreciable extent and is therefore likely to be relatively stress-free. Si~ficantly, gases such as 0 and N are quite insoluble in this material and it is possible that the vacancies cluster to form sub-microscopic, gas-stabilised voids. Any vacancy cluster may be unstable, however, at such a high temperature. 4.2. Vacancy loop growth during electron and neutron irradiation
The growth of vacancy loops during electron (and neutron) irradiation of Zr is not compatible either with the theories developed to explain void-swelling [ 171 or with the large body of experimental evidence on many other metals and alloys - the notable exception being Ti [IS], another hcp metal. The theories developed to account for void-swelling are all based on the premise that the dislocation strain field interacts more strongly with interstitials than with vacancies so that any dislocation attracts more interstitials than vacancies. On this basis, vacancy loops are unstable during irradiation and it is only interstitial loops that can grow. In order to accommodate the neutron-irradiation observations on Zr and Ti [ 1, 181 into the theories of void-swelling, Bullough et al. [19] have put forward a theory based on the fact that small dislocation loops have a smaller interstitial bias than do large loops or network dislocations. This theory thus relies on the network dislocations to remove the interstitials, so allowing a vacancy supersaturation to develop and hence vacancy loops to grow. This is an adequate explanation for a heavily cold-worked material but neither the present observations nor the observations of Jostsons et al. [l] seem compatible with this model since for annealed material the network dislocation density is low (lOs/cmz). Indeed on this basis prior cold-working of any metal before
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M. Griffiths et al. / Electron damage in zirconium I
irradiation should lead to the growth of vacancy loops formed during cascade collapse and there is no evidence in the literature to suggest that this is true. Jostsons et al. [l] have suggested that the existence of vacancy loops in neutron irradiated Zr and Ti depends on impurities in these materials. Indeed they have shown that the point defect clustering behaviour does vary from specimen to specimen but they have not been able to show which specific impurity is responsible. More recently, Woo [20] has suggested that it is not inconceivable for vacancy or interstitial loops to have a preference for vacancies in some material because of the anisotropy in the point defect saddle-point configuration. It certainly seems necessary to call on a specific property of Zr and Ti to account for the observations rather than a general argument of the type advanced by Bullough et al. [19]. One feature of loop-clustering to support this view is the banding of loops in layers parallel with the basal plane which appears to be linked with the crystallographic anisotropy of Zr and the tendency for autocatalytic nucleation of the loops [4]. Clusters of vacancy loops which are separated from interstitials loops are, by the very nature of this segregation, stabilised. Comparing the present work with previous studies, the main discrepancy lies in the observation of vacancy loops. Carpenter and Watters [4] and Buckley et al. [21] reported that interstitial loops only were formed during electron irradiation of Zr. No loop analysis was carried out by Buckley et al. but Carpenter and Watters carried out detailed loop analysis. Their results are very similar to the present observations made in relatively thin regions. For example, fig. 7 shows loops near the edge of a foil, all of which were found to be elliptical - as was the case in Carpenter’s work at the same temperature. All loops analysed in the present work, which were in such thin regions, were found to be interstitial. It seems likely that in thin regions, where the oxide-induced stress is very large, the vacancies diffuse to the metal-oxide interface to relieve the stress whereas the interstitials are repelled by the compressive stress in the oxide. These interstitials condense and grow as loops, with the loop shape governed by the stress field which tends to produce elliptical loops. The role of stresses on the clustering of point defects during electron and neutron irradiation is discussed more fully in paper II (see next pages). 5. Conclusions
(1) Vacancy and interstitial loops of b = f (1120) are observed following electron irradiation. The loops prob-
ably nucleate on (1010) and (1011) planes. (2) Theories based on the loop-size dependence of the bias term do not seem able to explain these observations. (3) The solubility of gases and the presence of stresses are likely to be important in determining whether voids or vacancy loops are formed during irradiation.
References [l] A. Jostsons, P.M. Kelly and R.G. Blake, 9th ASTM Intern. Symp. on Effects of Radiation Damage in Structural Materials, ASTM-STP-683 (1978) 46. [2] D.O. Northwood, R.W. Gilbert, L.E. Bahen, P.M. Kelly, R.G. Blake, A. Jostsons, P.K. Madden, D. Faulkner, W. Bell and R.B. Adamson, J. Nucl. Mater. 79 (1979) 379. [3] D. Faulkner and C.H. Woo, J. Nucl. Mater. 90 (1980) 307. [4] G.J.C. Carpenter and J.F. Watters, J. Nucl. Mater. 96 (1981) 213. [5] S.N. Buckley and S.A. Manthorpe, Proc. Intern. Conf. on Physical Metallurgy of Reactor Fuel Elements, Berkeley Nuclear Laboratories (1973) (The Metals Society, London) p. 127. [6] M. Griffiths, J. White, R.E. Smallman, M.H. Loretto and I.P. Jones, Proc. 10th Intern. Congress on Electron Microscopy, Hamburg (1982). [7] D.M. Maher and B.L. Eyre, Phil. Mag. 23 (1971) 409. [8] H. Foil and M. Wilkens, Phys. Status Solidi (a)31 (1975) 519. [9] M. Kiritani and H. Takata, J. Nucl. Mater. 69/70 (1978) 277. [lo] K. Urban, Phys. Status Solidi (a)4 (1971) 761. [l l] K. Shiraishi, A. Hashinuma, Y. Katano and T. Taoka. Proc. 3rd Intern. Conf. on HVEM, Kyoto (1974) 365. [12] P.T. Heald and M.V. Speight, Acta Met. 23 (1975) 1389. 1131 J. Bentley, B.L. Eyre and M.H. Loretto, Proc. Intern. Conf. on Radiation Effects and Tritium Technology for Fusion Reactors, Gatlinburg, Tenn. (1975) p. 297. [ 141 E. Tjhia, P. Wilkes and G.L. Kulcinski, Radiation Effects 51 (1980) 49. [15] M. Griffiths, J. White, M.H. Loretto and R.E. Smallman, Proc. Yamada Conf. V on Point Defects and Defect Interactions in Metals, Kyoto (1981). [16] M. Griffiths, I.G. Salisbury and R.E. Smallman, unpublished work. [ 171 R. Bullough, B.L. Eyre and K. Krishan, Proc. Reg. Sot. A346 (1975) 81. [18] A. Jostsons, R.G. Blake and P.M. Kelly, Phil. Mag. A41 ( 1980) 903. [ 191 R. Bullough, M.R. Hayns and C.H. Woo, J. Nucl. Mater. 84 (1979) 93. [20] C.H. Woo, J. Nucl. Mater. 107 (1982) 20. [21] S.N. Buckley, R. Bullough and M.R. Hayns, J. Nucl. Mater. 89 (1980) 283.
Journal of Nuclear Materials 115 (1983) 313-322 North-Holland RubIishing Company
ELECTRON DAMAGE IN ZIRCONIUM - I. Defect structure and loop character M. GRIFFITHS,
M.H. LORETTO
and R.E. SMALLMAN
Department of Metallurgy and Materials, University of Birmingham, P. 0. Box 363, Birmingham Bl5 2TT, UK Received 18 October 1982; accepted 8 November 1982
Zirconium (499.95% Zr) has been irradiated using 1 MV electrons at temperatures between 273 and 773 K in a high-voltage electron microscope and the nature of the clustered damage assessed. It has been shown that both interstitial and vacancy loops of 6 z f (11%) nucleate and grow during electron irradiation within this temperature range. Some loops of b = 4 (1123) and
%(2023) are also observed. The significance of the observation of vacancy lcxqs is discussed in terms of current theories of defect clustering during irradiation and it is concluded that the presence of imposed stress influences the clustering.
1. Introduction
2. Experimental
There is now incont~vertible evidence that large stable vacancy loops are formed during neutron irradiation of Zr at temperatures in excess of 0.3 r, [1,2]. The loops are far too large to have been formed merely by cascade collapse and it is apparent that their existence is related to the resistance of Zr to void swelling. Cases of void growth during neutron and electron irradiation of Zr have been reported but these have been shown to be largely dependent both on specimen purity and on the presence of implanted or fission-generated inert gases such as He [1,3]. Recent work by Carpenter and Watters [4] has shown that voids are not formed during HVEM irradiation of relatively pure undoped Zr for temperat~es up to 725 K and doses up to 23 dpa. This is consistent with much of the earlier work on this material but not with the results of Buckley and Manthorpe [5] who reported voids during electron irradiation of doped and undoped Zr. Carpenter suggests that the difference is probably related to impurity effects. Vacancy loops have not been reported following electron irradiation in any previous study. This work is part of a general programme aimed at understanding the behaviour of point defects in cph metals and alloys. The present paper assesses the relevance of basic observations of loop character and clustering behaviour to void-swelling. Paper II (see p. 323) deals primarily with the c-component loop which is considered in relation to radiation growth phenomena (dimensional changes in the absence of swelling).
Single crystal, zone-refined iodide rod and Marx grade Zr of ~99.95% purity were used for this work. Discs of 3 mm diameter and - 0.25 mm thick were prepared by spark-machining and were subsequently wrapped in Ta sheet and annealed in vacua for 8 h at 1073 K. Microscope specimens were prepared by electropolishing the discs in a Struers twin-jet Tenupol using 6% perchloric acid in ethanol at a temperature below 233 K and at 18 volts. Foils having normals - 4.5’ from [OOOl]were prepared from the single crystal in order to facilitate loop analysis. Foils for which a significant number of grains had a selected orientation with respect to the foil plane could be prepared from the iodide rod. Electron irradiation was carried out at 1 MV in the temperature range of 273-773 K using a f 25O doubletilt heating stage in an EM7 at a dose rate of - lOmu dpa/s. In general, specimens were cooled to room temperature and transferred to a rt 4S’ double-tilt stage for analysis at 400 keV in the EM7 or at 120 keV in a Philips EM4OOT. Si8nificant oxidation occurs during irradiation at temperatures in excess of 623 K and the stresses caused by the oxide give rise to slip in thick regions and bending in thin regions of the foil. Oxygen depth profiles have been obtained using an Auger spectrometer and comparison between a freshly prepared and irradiated sample shows that the depth of significant oxygen permeation from the foil surface increases by a factor of - 10, and is typically - 500 A following in situ irradia-
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M. Griffiths et al. / Electron damage in zirconium
1
tion at 675 K. Contrast from the oxide complicates analysis, especially when examining small defects, and care must be taken to reduce the contrast by use of appropriate diffraction conditions. One other adverse effect of oxidation is an asymmetric imaging condition in which defect contrast is obscured for one sense of virtually all diffracting vectors and this makes loop analysis difficult. An explanation of this effect has been reported elsewhere [6].
3. Results 3.1. Point defect clustering Dislocation loops nucleate and grow during the irradiation of Zr with 1 MV electrons. No other point defect clusters were observed over a range of irradiation temperatures between 273-773 K (0.13-0.36 T,) and for doses up to - 10 dpa. Loop structures do not generally reach a near steady-state at high doses; loops coarsen to a dislocation network and loop nucleation continues. Increasing the irradiation temperature leads to a decrease in loop density and an increase in the size of the loops for a given dose rate. Damage alignment was observed parallel to the trace of the basal plane when the electron beam direction B was close to prism orientations (i.e. (lOjO) and (1120)) as illustrated in fig. 1. Tilting from basal to prism orientations shows that most of the damage is clustered in layers parallel with the basal plane. The spacing between the layers increases with increase of irradiation temperature and is comparable with the size of the loops at moderate doses ( - 0.1 dpa). The loop distribution within the layers is non-uniform and consists mainly of small strings of, typically, 4-6 loops. These loops are non-overlapping, have identical Burgers vectors and have approximately the same orientation but are not necessarily coplanar. Loop clustering in this manner is observed from the early stages of irradiation suggesting that their nucleation is influenced by a localised stress field probably from their mutual interaction. During continued irradiation adjacent loops sometimes coalesce, forming larger elliptical loops of complex shape, as illustrated in fig. 2. Burgers vector analysis showed that loops having b = 4 (112b), -f (1123) and more uncommonly b = d (2023) are present in electron-irradiated foils of Zr. All loops clustering in layers were analysed as having Burgers vectors of the type b = 4( 1120), alternatively referred to as (B) type. The (u) type loops can be distinguished from the c-component damage using a fOOO4 diffracting vector. This puts dislocations of the type b = i (1120) out of contrast and
Fig. 1. Alignment of damage in Zr in layers parallel with the basal plane following electron irradiation at - 675 K to - 0.15 dpa, g = 1120, B - (lioo].
shows clearly those having a c-component. Loops having b = f (1123) or (c + a) type loops were observed between the layers of (a) type damage. The ellipticity and sizes of the (u) type loops varied considerably because of the absence of steady-state conditions and as a result of the proximity of the foil surfaces. Sometimes a bimodal distribution of loops was observed, as shown in fig. 1 where large numbers of small loops cluster close to the foil edge separated, by a denuded zone, from the bulk of the damage within the foil. Although it is difficult to determine accurately the plane of nucleation of the loops, from observations made with B- (1213), (l~lO), or [OOOl] there is evidence to suggest that (u) type loops do not nucleate exclusively on (1210) planes and c-component loops do not nucleate exclusively on basal planes. The most probable nucleation planes for (u) type damage are then (lOi0) and (lOi1) planes. Figs. 3(a) and 3(b) show loop nucleation during irradiation using B - [OOOl] and B - (1213), respectively. The traces of relevant planes are marked. There may be some nucleation on (lzlO} planes but the (10’50) and (lOi1) planes appear to be more likely candidates and this is hardly surprising considering that they are more widely spaced than the (1210) planes even though they are corrugated. Measurements on (c + u) type loops when viewed along
M. Griffiths et al. / Electron damage in zirconium I
315
Fig. 3. Loop nucleation during
Fig. 2. Loop coalescence in Zr during electron irradiation at - 750 K. Many beam imaging conditions, B - (11~0); (a) 0.08 dpa (b) 0.1 dpa.
(l?lO) indicate that they are tilted slightly away from (lOi1) towards (lOi0) by about 10’ during the early stages of irradiation. This is consistent with loop growth on a combination of (lOi1) and (1OiO) planes. It is
electron irradiation of Zr at - 675 K. Traces of important planes are shown; (a) g = 2ii0, B - [OOOI], (b) g = ioio,B - [1213].
difficult to say whether this is also the case for (a) type damage. Evidence for c-component loops nucleating on basal planes in prism foils is sparse and nucleation on planes inclined to basal appears to be dominant. For discussion purposes the nucleation planes of (a) type loops will be considered to be { lOiO> and (107 1) whereas
the planes of nucleation for (c + (0 type loops will be considered to be {lOit>. All loops subsequently rotate during growth to less well-defined orientations. The long dislocations seen in fig. 3(b) are formed during heating of the specimen by slip of edge dislocations through the foil, leaving dislocations trapped at the coherent metal/oxide interface. Only the (a) type loops will be analysed in more detail as we are mainly concerned with the interstitial/vacancy nature of these loops at present. A more detailed account of the c-component damage is presented in paper II. 3.2. Dislocation ioop characterisation Because of the variability in experimental conditions for thin foil irradiations such as foil thickness, specimen purity (dependent also on the microscope vacuum), degree of oxidation, absence of a steady-state damage structure etc., no attempt will be made to give a statistical account of the experimental observations. Some quantitative results are being processed, however, and will be presented in a future publication. Adverse imaging conditions associated with asymmetric contrast tends to bias loop analysis especially for loops giving g 1b = 1 images. Thus the loop marked with an A in fig. 4 shows clear outside contrast with g =I 1120 compared with g = ii20. Loop B, however, can be seen clearly with g = 1120 but is obscured with g = 1720 and it is difficult to say whether it exhibits inside or outside contrast. The net effect is that, dependin8 on the direction of contrast asymmetry, analysis of the loops can be biased in favour of interstitial or vacancy loops. Both interstitial and vacancy loops have been analysed using the Maher and Eyre 171 and Fiill and Wilkens [S] techniques. As a result of the adverse conditions already mentioned only a small percentage ( - 10%) of the total loop population could be analysed with confidence. Many of the loops had a large screw component and this made it imperative that “safe” regions for analysis be determined for each particular loop and that the loop plane should be determined wherever possible. It should be emphasised therefore that two different methods have been used to determine the nature of the loops and the loop plane normal has been obtained to within & 20” for all the loops which have been analysed. No assumption has been made concerning the maximum shear component. At temperatures below 575 K the loops were generally too small to analyse with confidence. Examples of
Fig. 4. Asymmetric (f g) contrast from a Zr foil electron irradiated at -675K.LoopAhasb=+~[ll~O]and~oopB has b= *If[lilO]; (a)g=llTO, B-[IiOIj,(b)g=ii20,B-
[IiOlJ.
loop characterisation in three higher temperature mes will now be described in detail.
regi-
M. Griffiths et al. / Electron damage in zirconium I 3.3. irradiation
at - 585 K, foil normal N - [If131
The loops 1, 2, 3 and 4 in fig. 5 are out of contrast with g = Oil0 and g= OLil and thus have Burgers vectors of b = & f[21 lo]. They are almost edge-on close to B - [ 121 l] and micrographs taken over a large tilt range show that they have habit planes with n - [i103]. They show outside contrast for g = iOl0 at B - [iZiS] and therefore have 6 = $1 lo], indicating that they are interstitial in nature. No loops in thicker regions of the foil could be analysed because of adverse imaging conditions. The majority of large loops having b = #l lo] lie close to the foil edge and the partition between these and the rest of the loop population is quite clear. No vacancy loops were identified.
oiio
ioio tooa
317
3.4. rrrud~~tion at - 675 K, foil normaI N - [Jill]
A mixed population of vacancy and interstitial loops was observed following irradiations at this temperature. The loops marked 1, 2, 3 and 4 in fig, 6 are out of contrast with g= TO11 and g= 0004 and have b = &$f[l210]. They show outside contrast with g= ilO1 and have habit planes of n - [Oi 1l] and are therefore vacancy in nature with b = #sio]. Loop 5 is out of contrast with g = ?Oli, has n - [OilO] and exhibits outside contrast with g = 1iOi. It therefore has b = +[I2101 and is interstitial in nature. The majority of damage has b = f $[iZ‘iO] and most of this is comprised of vacancy loops. Only a small percentage of vacancy loops ( < 20%) have Burgers vectors other than b =f[i2iO]. The major-
ioio
Fig. 5. Chara~te~sation of dislocation loops in Zr electron irradiated at - 585 K to - 0.1 dpa. Loops 1,2,3 and 4 all have R - [i103], and are interstitial in nature, g and B are indicated for each figure.
b = f[Sl lo]
318
M Griffiths et al. / Electron damage in zirconium
I
iioi
To0
[oili]
[?213]
p201
toil
ooo&
Fig. 6. Characterisation of dislocation loops in Zr electron irradiated at - 675 K to - 0.15 dpa. Loops 1,2, 3 and 4 haven - [OilI], character. Loops 6 and 7 have b = f[i2iO] and are vacancy in nature. Loop 5 has n - [OTlO], b = f[l?lO] and has interstitial n - [iOl I], b = f[ii20] and are interstitial in nature. Loop 8 has n - [ 1 IOO], b = 4[ZiiO] and is also interstitial in nature, g and B are indicated for each figure.
319
pit
1120
iJ
iori
[liOl]
Fig. 7. Character of dislocation loops in Zr electron irradiated at - 730 K to -0.1 dpa.ioops 1 andzhaven-[IiOl],b=~~~iO] and are vacancy in nature. Loops 3 and 4 have n - [Oli2], 6 = $1201 and are interstitial in nature. Loops 5 and 6 have II - [1230], b = $I?lO] and are also interstitial in nature, g and 3 are indicated for each figure.
ity of interstitial loops have Burgers vectors of b =j(ll~O) or +(ZiiO). Loops 6 and 7 are out of contrast with g = ilO but show inside contrast with g = IOil. They have n - [iOl lf which shows that they
are interstitial in nature having b = f[ ii20]. Loop 8 is out of contrast with g = Olil and shows outside contrast with g = IOil. It has n - [IiOO] and therefore is also interstitial in nature having b = j[2iiO].
It is apparent that interstitial loops are more common close to the edge of the foil, although this may be a result of the difficulty in analysis of similar loops within the dense dislocation network found in thicker regions of the foil. Even so, interstitial loops account for perhaps only 20-30% of the total loop population. 3.5. ~rra~ia~io~at - 730 K, N - [Oj23] Vacancy and interstitial loops were observed following irradiation at this temperature. Loops 1 and 2 in fig. 7 are out of contrast with g = iOli, have habit planes n - f 1101J and exhibit outside contrast with g = i2i2 and therefore have Burgers vectors of b = +[iZiO] and are vacancy in nature. Loops 3 and 4 are out of contrast withg = 1101, have habit planes n - [Ol]Zj and exhibit outside contrast with g= i2i2, They have Burgers vectors b = f[ 11201 and are interstitial in nature. Loops 5 and 6 are out of contrast with g =i lOjO, have n - []230] and exhibit outside contrast with g = 1212 and therefore have b = b[lzlO] and are also interstitial in nature. Clearly, as with the previous example of fig. 6, there is a partition of interstitial loops close to the edge of the foil and in this example these loops are markedly elliptical. There is no predominance of one particular
Fig. 8. Climb of dislocations at the metal/oxide interface during irradiation of Zr at 675 K, g =ilOl, B - 112131.
Burgers vectors in contrast to the previous example. The majority of loops appear to be vacancy in nature (provided that they have a safe orien~tion) since they exhibit outside contrast for g = i2i2. Unfortunately the majority of loop habit planes could not be ascertained accurately enough for confident loop analysis and therefore it is not possible to estimate the number of interstitial and vacancy loops. It is apparent, however, that no vacancy loops were observed amongst those anaiysed close to the edge of the foil. Also it is highly probable, on the basis of the observations, that the majority of the loops (say > 70%) in the bulk of the foil are vacancy in nature.
4. Discussion Vacancy loops are not expected to nucleate and grow during electron irradiation except under very specific conditions e.g., either in very thin areas [9], where the more mobile interstitials reach the free surface leaving a supersaturation of vacancies, or in the compression region of a dislocation, for example close to the edge of a large interstitial loop /lo]. One material in which vacancy loops have been reported following electron irradiation in the absence of the above conditions is vanadium. Shiraishi et al. [ 111 observed vacancy loop growth during electron irra~ation of vanadium at temperatures > 0.3 T, and they attributed this result to the effect of impurities in the matrix (0.15%). Their results suggest that unfaulted vacancy loops are favoured energetically over voids in irradiated vanadium although under normal circumstances any vacancy loops are unstable, attracting interstitials as a result of the dislocation bias term [ 121. Although the abundant interstitial impurities such as 0 and C may play a role in pinning self-interstitial atoms in V, it seems unlikely that this in itself is enough to stabilise the vacancy loops. Vacancy loop stability may be the result of impurity segregation. Impurities can change the interstitial bias for vacancy or interstitial loops depending on whether the solute segregates to the compression or dilational side of the dislocation line in a manner similar to that proposed by Bentley et al. for oversized solutes in neutron irradiated TZM 1131 or Tjhia et al. for undersized solutes in electron irradiated Al-Si alloys [14]. The results presented here show beyond doubt that vacancy and interstitial loops of b = 4( 11~0) nucleate and grow during electron irradiation of Zr. It therefore follows that cascades are not essential for the nucleation of vacancy loops - although of course the vacancy loops
observed in neutron irradiated Zr may be formed in cascades. Two questions need to be answered concerning vacancy loop growth: first, how do vacancy loops nucleate in electron-irradiated samples and second, how do vacancy loops grow during electron and neutron irradiation? 4.1. Nucieatidn
of vacancy
loops during electron irradia-
tion
A consideration of planar geometry and the observations made during electron irradiation of Zr suggests that the most common nucleation plane is (lOiO}. The {lOiO} planes are the most widely spaced planes because of the low c/a ratio of 1.59 and a hard sphere model suggests that coincident collapse and shear is reasonably likely on these planes, giving loops _of b = 4 (1120). Nucleation is also possible on the (1011) planes which have a slightly lower packing density than {IOiO} planes. Since vacancy loops have been identified in electron irradiated samples it is clear that a sufficient vacancy supersaturation exists locally to nucleate the loops. This perhaps suggests that there is a rapid, efficient removal of interstitials and this may be associated with the presence of dislocations generated by slip (cf. figs. 6 and 7). Indeed fig. 8 shows that these dislocations can climb significantly during irradiation. It should be noted, however, that such climb is not noticeable for heavily oxidised samples where they are strongly pinned at the interface, and there are insufficient dislocations to account for widespread vacancy clustering. A high vacancy supersaturation may result from pinning of interstitial atoms by impurities and the impurities themselves may enhance the stability of voids or loops. A small vacancy cluster is more likely to collapse to give a loop because most gases, such as 0 and N, are highly soluble in Zrand cannot therefore act as stabilisers for void nuclei. It is significant that such gases are very soluble in V in which vacancy loops are also observed, however it is not obvious that gas solubility is a sufficient prerequisite for vacancy loop stability. One possible factor contributing to loop stability is the segregation of impurities to the loops. There is some evidence that there is considerable impurity segregation to loops in electron- and protonirradiated Ti for a given purity and irradiation temperature [ 151. Such loops have irrational Burgers vectors but the same cannot be said for loops in Zr which appear to have rational Burgers vectors and for which segregation has so far been undetectable. In the absence of reliable evidence for impurity segregation to the loops in Zr alternative explanations for vacancy loop stability must
be sought, although impurities may indeed be important. The fact that vacancy loops nucleate may imply that there is an energy gain to the system and this would be the case if the nucleation (and subsequent growth) were to relieve an appropriate stress. There is clearly a stress induced by the growing oxide (viz., the dislocations in figs. 6 and 7). This stress arises because of the volume increase during oxidation (Pilling-Bedworth ratio = 1.5) and because of the lattice expansion close to the oxide-metal interface caused by the solution of interstitial oxygen. The stress may be expected to be tensile within the plane of the foil and compressive in the oxide. It is notable that interstitial loops only have been observed in electron irradiated ruthenium at 0.39 T, [16]. It has a similar c/a ratio to Zr but does not suffer from oxidation to any appreciable extent and is therefore likely to be relatively stress-free. Si~ficantly, gases such as 0 and N are quite insoluble in this material and it is possible that the vacancies cluster to form sub-microscopic, gas-stabilised voids. Any vacancy cluster may be unstable, however, at such a high temperature. 4.2. Vacancy loop growth during electron and neutron irradiation
The growth of vacancy loops during electron (and neutron) irradiation of Zr is not compatible either with the theories developed to explain void-swelling [ 171 or with the large body of experimental evidence on many other metals and alloys - the notable exception being Ti [IS], another hcp metal. The theories developed to account for void-swelling are all based on the premise that the dislocation strain field interacts more strongly with interstitials than with vacancies so that any dislocation attracts more interstitials than vacancies. On this basis, vacancy loops are unstable during irradiation and it is only interstitial loops that can grow. In order to accommodate the neutron-irradiation observations on Zr and Ti [ 1, 181 into the theories of void-swelling, Bullough et al. [19] have put forward a theory based on the fact that small dislocation loops have a smaller interstitial bias than do large loops or network dislocations. This theory thus relies on the network dislocations to remove the interstitials, so allowing a vacancy supersaturation to develop and hence vacancy loops to grow. This is an adequate explanation for a heavily cold-worked material but neither the present observations nor the observations of Jostsons et al. [l] seem compatible with this model since for annealed material the network dislocation density is low (lOs/cmz). Indeed on this basis prior cold-working of any metal before
322
M. Griffiths et al. / Electron damage in zirconium I
irradiation should lead to the growth of vacancy loops formed during cascade collapse and there is no evidence in the literature to suggest that this is true. Jostsons et al. [l] have suggested that the existence of vacancy loops in neutron irradiated Zr and Ti depends on impurities in these materials. Indeed they have shown that the point defect clustering behaviour does vary from specimen to specimen but they have not been able to show which specific impurity is responsible. More recently, Woo [20] has suggested that it is not inconceivable for vacancy or interstitial loops to have a preference for vacancies in some material because of the anisotropy in the point defect saddle-point configuration. It certainly seems necessary to call on a specific property of Zr and Ti to account for the observations rather than a general argument of the type advanced by Bullough et al. [19]. One feature of loop-clustering to support this view is the banding of loops in layers parallel with the basal plane which appears to be linked with the crystallographic anisotropy of Zr and the tendency for autocatalytic nucleation of the loops [4]. Clusters of vacancy loops which are separated from interstitials loops are, by the very nature of this segregation, stabilised. Comparing the present work with previous studies, the main discrepancy lies in the observation of vacancy loops. Carpenter and Watters [4] and Buckley et al. [21] reported that interstitial loops only were formed during electron irradiation of Zr. No loop analysis was carried out by Buckley et al. but Carpenter and Watters carried out detailed loop analysis. Their results are very similar to the present observations made in relatively thin regions. For example, fig. 7 shows loops near the edge of a foil, all of which were found to be elliptical - as was the case in Carpenter’s work at the same temperature. All loops analysed in the present work, which were in such thin regions, were found to be interstitial. It seems likely that in thin regions, where the oxide-induced stress is very large, the vacancies diffuse to the metal-oxide interface to relieve the stress whereas the interstitials are repelled by the compressive stress in the oxide. These interstitials condense and grow as loops, with the loop shape governed by the stress field which tends to produce elliptical loops. The role of stresses on the clustering of point defects during electron and neutron irradiation is discussed more fully in paper II (see next pages). 5. Conclusions
(1) Vacancy and interstitial loops of b = f (1120) are observed following electron irradiation. The loops prob-
ably nucleate on (1010) and (1011) planes. (2) Theories based on the loop-size dependence of the bias term do not seem able to explain these observations. (3) The solubility of gases and the presence of stresses are likely to be important in determining whether voids or vacancy loops are formed during irradiation.
References [l] A. Jostsons, P.M. Kelly and R.G. Blake, 9th ASTM Intern. Symp. on Effects of Radiation Damage in Structural Materials, ASTM-STP-683 (1978) 46. [2] D.O. Northwood, R.W. Gilbert, L.E. Bahen, P.M. Kelly, R.G. Blake, A. Jostsons, P.K. Madden, D. Faulkner, W. Bell and R.B. Adamson, J. Nucl. Mater. 79 (1979) 379. [3] D. Faulkner and C.H. Woo, J. Nucl. Mater. 90 (1980) 307. [4] G.J.C. Carpenter and J.F. Watters, J. Nucl. Mater. 96 (1981) 213. [5] S.N. Buckley and S.A. Manthorpe, Proc. Intern. Conf. on Physical Metallurgy of Reactor Fuel Elements, Berkeley Nuclear Laboratories (1973) (The Metals Society, London) p. 127. [6] M. Griffiths, J. White, R.E. Smallman, M.H. Loretto and I.P. Jones, Proc. 10th Intern. Congress on Electron Microscopy, Hamburg (1982). [7] D.M. Maher and B.L. Eyre, Phil. Mag. 23 (1971) 409. [8] H. Foil and M. Wilkens, Phys. Status Solidi (a)31 (1975) 519. [9] M. Kiritani and H. Takata, J. Nucl. Mater. 69/70 (1978) 277. [lo] K. Urban, Phys. Status Solidi (a)4 (1971) 761. [l l] K. Shiraishi, A. Hashinuma, Y. Katano and T. Taoka. Proc. 3rd Intern. Conf. on HVEM, Kyoto (1974) 365. [12] P.T. Heald and M.V. Speight, Acta Met. 23 (1975) 1389. 1131 J. Bentley, B.L. Eyre and M.H. Loretto, Proc. Intern. Conf. on Radiation Effects and Tritium Technology for Fusion Reactors, Gatlinburg, Tenn. (1975) p. 297. [ 141 E. Tjhia, P. Wilkes and G.L. Kulcinski, Radiation Effects 51 (1980) 49. [15] M. Griffiths, J. White, M.H. Loretto and R.E. Smallman, Proc. Yamada Conf. V on Point Defects and Defect Interactions in Metals, Kyoto (1981). [16] M. Griffiths, I.G. Salisbury and R.E. Smallman, unpublished work. [ 171 R. Bullough, B.L. Eyre and K. Krishan, Proc. Reg. Sot. A346 (1975) 81. [18] A. Jostsons, R.G. Blake and P.M. Kelly, Phil. Mag. A41 ( 1980) 903. [ 191 R. Bullough, M.R. Hayns and C.H. Woo, J. Nucl. Mater. 84 (1979) 93. [20] C.H. Woo, J. Nucl. Mater. 107 (1982) 20. [21] S.N. Buckley, R. Bullough and M.R. Hayns, J. Nucl. Mater. 89 (1980) 283.