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HREM of 〈c〉-type dislocations in high-dose neutron-irradiated and postirradiation annealed Zircaloy-2
Journal of Nuclear Materials 205 (1993) 284-292 North-Holland
HREM of ( c) -type dislocations in high-dose neutron-irradiated and postirradiation annealed Zircaloy-2 Rodney A. Herring and T. Tanji Tonomura Electron Wavejiont Project, Exploratory Research for Advanced Technobm, JmC, Hatoyama, Saitama 350-03, Japan
c/o
ARL Hitachi Ltd.,
Derek 0. Northwood Engineering Materi& Group MechanicalEngineering Department, Universityof W&or,
W&or,
Ontario, Canada N9B 3P4
High resolution electron microscopy (HFtEM) was performed on (c)-type dislocations in annealed Zircaloy-4 which bad been neutron irradiated to 20 dpa at 690 K Lattice images of the (c)-type. dislocations, which lie on the basal plane, reveal unfaulted dislocation structure. Conventional TEM analysis of these dislocations in postirradiation annealed Zircaloy-2 shows that the (c)-component dislocations have Burgers vectors of i(1123) and [OOOl].Large bends in the (c)-component dislocation network are associated with multiple (a)-type dislocations. Initial results by electron holography of the (c)-type dislocation core region produces phase shifts having a length of m 3 basal planes in isolated regions along the length of the dislocations. The shift may occur due to large solute concentrations.
1. Introduction
Zirconium alloys are used in the nuclear industry because of their much lower neutron absorption per unit strength than other commercially available structural materials. Irradiation-induced growth and irradiation-enhanced creep are of concern both for fuel cladding and nuclear reactor structural components such as pressure tubes and calandria tubes. To understand growth and creep it is necessary to understand the nature of the irradiation-induced microstructural changes. This includes the altering of the zirconium alloy’s chemistry. Our understanding of the evolution of microstructure in zirconium alloys during irradiation has improved substantially over the last 10 years. In his 1977 review of irradiation damage in zirconium, Northwood [l] summarized the interpretation of the microstructure developed during irradiation as follows: (1) Only dislocation loops having Burgers vetior of b = f(ll2o), lying on or close to prism {ioio)planes, were formed during irradiation. They were: mostly aligned in rows parallel with the trace of the basal plane WJO1). (2) Both interstitial and vacancy loops were formed during neutron irradiation, the relative proportions 0022-3115/93/$06.00
varying with irradiation temperature. At high temperatures (2 675 IO the majority were vacancy in nature and at low temperatures (5 575 K) the majority were interstitial in nature. Interstitial loops were markedly elliptical. As pointed out by Griffiths [2] this simple interpretation is invalid. The microstructure developed in Zralloys during irradiation is more complicated as indicated by specimens of varying composition and thermo-mechanical history that have been irradiated to higher fluences. In particular, Griffiths [2] noted that, with the development of techniques for microchemical analysis, it has become increasingly obvious that alloying elements or impurities are important in affecting and contributing to the evolving microstructure. Also attention has been focused on c-component dislocations and loops and their contribution to accelerated, or “breakaway growth” [3]. Generally, whenever large growth strains have been observed, c-component dislocations were also present [4,5]. cxomponent dislocations can be introduced into the material during neutron irradiation [2,4-111 and during cold-working [1216]. In our earlier work [17] some data were given on the analysis of a dislocation network in neutron-irradiated Zircaloy-2. More work is presented in this paper.
8 1993 - Elsevier Science Publishers B.V. All rights reserved
R.A. Herring et al. / HREM of (c)-type &locations in Zircaloy-2
285
The comparatively new technique of electron holography [18] is used to explore dislocation cores. Excellent reviews on holography have been provided by Tonomura 1191 and Hanszen [201. Holograms are produced from the interference between a reference wave and an object wave. Contained within the hologram is the amplitude and phase information of the object. Conventional electron images only contain the amplitude or intensity distribution of the object.
2. Experimental details Two thin foil Zircaloy-2 specimens that had been neutron irradiated (E > 1 MeV) as bulk material in the EBR-II reactor, Argonne National Laboratory West, at 690 K to a fluence of 20 dpa were supplied to us by Chalk River Laboratories of AECL Research. Composition was 1.2-1.7 wt% Sn, 0.07-0.20 wt% Fe, 0.05-0.15 wt% Cr, 0.03-0.08 wt% Ni, an oxygen concentration of 1500 ppm, with the balance being Zr. One thin foil was subjected to an in-situ annealing experiment in an AEI EM7 high voltage electron microscope. The foil was held at 1075 K for over 20 min. A JEM lOO-CX TEM, capable of large two-axis tilt angles, was used to examine the dislocation network. Burgers vectors were determined by the g - b = 0 condition, i.e., zero contrast, plus the diffraction behaviour of the dislocation when they were in contrast, such as when g-b = 1, 2 or higher. Line directions, U, and habit planes, R, were determined with the aid of a stereogram and noting the projected directions of the dislocations. A dedicated high resolution microscope, Hitachi HF 2000, was used for the lattice imaging. This microscope has been fully characterized for its high resolution potential, having a well-controlled environment with respect to temperature fluctuations and noise, a spherical aberration of the objective pole piece of 0.65 mm and a chromatic aberration of the cold field emission electron gun of 0.3 eV for an acceleration energy of 200 keV. The point to point resolution is 0.21 nm, and thus the line widths (which are much easier to see than point to point images) of (0001) Zr at 0.51 pm were quite easily seen in our micrographs. The electron gun had a very high emission, 30 uA, which enabled very low beam divergence to be used N 0.3 to 0.5 mrad. The lattice images were taken using the standard photographic method for lattice images which starts from about exact focus to about the Scherzer focus (48.5 nm) in 5 to 7 steps. Multiparallel-beam lattice images were taken using the (0OOl)g family of
Fig. 1. The hologram is formed by illuminating the object (the dislocatioti core region) with a coherent electron beam wider than the object. The half of the beam that bypasses the object (the Zr matrix) serves as the reference wave. Electron lenses focus both halves of the wave, and the biprism superposes them in the image plane, where the interference pattern is recorded on film.
systematic reflections at orientations parallel to the basal plane between the 1100 and 1210 zone axes. This condition was used to clearly see the c-component dislocations without any image contribution by a-component dislocations. To clearly see the dislocation core region, the Zircaloy-2 crystal was tilted _ 15 to 20 mrad from the basal plane, i.e., the main beam approximately occupied the [0008] reflection position. Fig. 1 illustrates the procedure for formation of an off-axis hologram image. In one-half of the specimen plane, the electron wave illuminates the object or region of interest. The other half of the plane is reserved for a reference wave. The reference wave can pass through the vacuum or a part of the specimen with a constant inner potential and thickness [21]. A positive potential of about 100 V is applied to a centrally located Cr-coated glass-filament of the electron biprism beyond the objective lens which makes the image and the reference wave overlap to form a hologram or interference pattern. Magnification of the interference pattern by the projector lenses enable the fringes to be easily seen on the recording media, i.e., film. Contained within the hologram is both the phase and amplitude information of the object.
3. Results 3.1. HREM 3.1.1. As-irradiated
material The DOO21-type diffraction
vector was used to see the c-component dislocations in the as-irradiated material. One region was found with a particularly high
R.A. Hemhg et al. / HREhi of (c)-type dislocationsin Zircaloy-2
286
density of c-component dislocations having line directions substantially deviated from the basal plane, whereas most c-component dislocations had line directions lying in the basal plane. A large number of lattice images were taken of the c-component dislocations using weak diffracting conditions, i.e., [00081. Fig. 2 shows two examples of the lattice images taken at many places between the (liOO> and (l%O> zone axes: one is along a straight section, i.e., in the basal plane direction, and the other is at a bend, i.e., away from the basal plane direction, where the c-component dislocations lie between the bracketed regions. These lattice images show that all of the basal planes are intact, i.e., there are no missing planes of atoms, and no stacking faults in the basal plane are present in the as-irradiated material. 3.1.2. Postirradiation annealed material The analysis of a dislocation network which was adjacent to a grain boundary and contained c-component dislocations is given in fig. 3. Only a few of the micrographs used to characterize the dislocation network are given and these show the two-beam electron diffraction vector and the crystal orientation (given within the square brackets). To fully determine the diffraction behaviour of the dislocations, pictures were taken of the dislocations at a range of deviations from the exact Bragg condition. A summary of the contrast given by the dislocation segments is given in table 1, with an asterisk (*) for the dislocation segments in contrast and a cross (XI for the dislocation segments out of contrast, given for their respective electron diffraction conditions. Fig. 4 is a schematic illustration of the dislocation network. In fig. 3 and fig. 4 the
Fig. 2. Lattice images of as-irradiated material; (a) along a straight section, i.e., in the basal plane direction; (b) at a bend, i.e., away from the basal plane direction. The c-component dislocations lie between the bracketed regions.
c-component dislocations are contained within the dislocation segments marked B, C, D, E, K, M and possibly 0 and are clearly in contrast for the
Table 1 Diffraction contrast of dislocation segments given in fig. 4 Diffraction vector and orientation
x :
Contrast of dislocation segment (* : in,
out)
ABCDEFGHIJKLMNOPQR
rtooo4[oiioI x *Zl12[oilo] -
*
*
*
*
*
x
x
x
x
*
x
*
x
-
x
x
x
*
x
*
*
*
*
*
*
*
*
*
*
x
*
*
*
*
*
*
x
*
x
*
x
*
*
*
*
*
*
*
*
*
*
*
*2ii~oiloi * ~lioi[ll2o] *
*
*
*
*
x
-
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
x
*
x
*
x
*
*
*
*
*
*
*liol[oili] *
*
*
*
*
*
*
x
*
x
*
x
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
x
*
*
*
*
*
*
*
*
*
*
*
*
*Zllqoilo]
*loii[l2lo] *
*
*
*
*20zq12101
*
-
*
x*-****-*******
*2ii2@3231
*
*
*
*
*
R.A. Herring et al. / HREM of (c)-type didocations in Zircakq-2
Fig. 3. Analysis of dislocation network in postirradiation annealed material. The network was adjacent to a grain boundary and contained c-component dislocations; (a) c-component dislocations in contrast and a-component dislocations out of contrast; (b) some of the dislocation network in contrast; (cl and (d) dislocations having b = [Oool] out of contrast, i.e., segments B and D; (e) dislocation haying b = f[%lz] out of contrast; (0 and (h) dislocation haying b = f[Z113] out of contrast. Micrographs for (d), (s> and (h) orientations were used to determine the line direction, u, for the dislocation segments, D, b = [OOOl]and C, b = #113].
288
RA. Herring et al. / HREIU of (c >-type dislocations in Zircaloy-2
whereas segment F is also out of contrast for + 2ii2g[OilO] and + 2ii2g[2ri23] and has b = $1131, i.e., the opposite c-component to that of segment C. The pure c and (c + a)-dislocation segments are joined to the pure a-component dislocations having b = $llO] which are all out of contrast for the 0004g[OllO] and given by the symbols A, G, I, L, P, Q, and R. The pure u-component dislocations are consistent with the Burgers vectors of the c and (a + c>dislocations. The bend in the c-component dislocation, given by segment E, is a complex junction between c and u-component dislocations. Note that there is a high density of a-component dislocations associated with segment E. The Burgers vectors of the dislocation segments of the dislocation network illustrated in fig. 4 are listed in table 2. There are segments of the dislocation network which could not be identified with high confidence including segments E, K, M, N, and 0. The line directions and habit planes for the dislocations labelled B, C, and D in fig. 3, which had Burgers vectors b = [OOOl]and b = $[2113], were found to have line directions close to the prism directions and lie within the (0001) habit planes. Thus, the c-component dislocations lie on non-glide planes. Fig. 5 is a series of micrographs and schematic diagram showing a c-component dislocation network
Table 2 Burgers vectors of dislocation segments given in fig. 4 Fig. 3 (continued).
0004g[O’ilO]. The segments given by B, D and F are pure c-component and are out of contrast with +211Og[OilO]. There appears to be an overlap of the pure c and (a + c)-dislocation, especially in the region of the F segment. Segment C is out of contrast with +2112g[OilO] and has a Burgers vector of b = $!llJ],
Dislocation segment A B C D E F
;[zllol
DJoo11 j[Z113]
mw ? [OCKU] and f[?113]
G
;[2110]
H
+[llzo]
I
f[ZllO]
J K L M N 0 P
;r11zo1 ? f[ll20] ? ? ? ;[zllo]
Q
~[sllol ;[zllo]
R Fig. 4. Schematic diagram of dislocation network in fig. 3.
Burgers vectors b
?: uncertain of analysis.
R.A. Herring et al / HREM of (c)-type disbcatiom
running through a precipitate. Fig. 5a is the OOO& diffracting vector and figs. 5b and 5c are bright and dark field micrographs using long. These micrographs show that the dislocation segment A (see fig. 5d) is out of contrast (shadowy with g close to Bragg) and that segment B has interacted with an a-type dislocation to form a bend in the c-component segment and a new c-component dislocation with a different Burgers vector has formed. All of the i(2ti3) dislocations are in contrast with
in Zircaloy-2
289
the lOi1 diffraction vectors and weak beam imaging with this condition showed the presence of some +(20?3) dislocations associated with stacking faults. This agrees with Griffiths’ I21 suggestion that edge sections of [cl or (c + a)-dislocations climb by splitting into partials during neutron irradiation of Zircaloy-2 at 690 K. The reactions being [OOOl] + ;(2023)
+ +(2023),
+(1123) --) *(2023) + ;(0223),
(1) (2)
d
Fig. 5. Series of micrcgraphs and schematic diagram showing a c-component dislocation network running through a precipitate in postirradiation annealed material; (a) 0006g diffracting vector; (b) and Cc) bright and dark field micrographs using 1Oug; (d) schematic diagram.
290
RA. Herring et al. / HREM of ( c>-typedislocations in Zircaloy-2
with the result that both k(2023) dislocations and stacking faults would be present. However, the number of c-component dislocations having stacking faults were at a very much lower density than the c-component dislocations. An example is given in fig. 6 which are dark field and bright field images taken using the Olilg at the [1213lZA. Many of the stacking faults observed were missing their :(20?3) partial dislocations (out of the foil) and faults are visible only in the weak-beam micrographs.
Fig. 7. Electron holograms of (a) as-irradiated and (b) postirradiation annealed material showing fringe shifts which occur at isolated regions along the c-component dislocations and the dislocation line direction. In summary, the c-component dislocation network in postirradiation annealed material consists of [OOOl] and i(1213) dislocations and faults on (0001) having $(2023) partial dislocations. 3.2. Electron holography
Fig. 6. Dark field (a) and bright field (b) images of postirradiation annealed material taken using the Ol’Ilg at the [1213lZA showing stacking faults.
Electron holography was performed on both asirradiated and postirradiation annealed specimens. Typical holograms for as-irradiated and postirradiation annealed material are given in fig. 7a and 7h, respectively. The common feature of these holograms are the fringe shifts which occur at isolated regions along the c-component dislocations and the dislocation line direction. In fig. 7a, as-irradiated material, the phase shifted region is along the @001) planes and can be seen to have a width of 3-4 basal planes, whereas in fig. 7b, postirradiation annealed material, the phase shifted region is at a bend in the dislocation where there should be intersections from other dislocations.
R.A. Herring et al. / HREM of (c >-typedislocationsin Zircaloy-2
These fringe shifts, which were present for both as-irradiated and postirradiation annealed material, are indicative of small changes in the mean inner potential believed to be due to solute segregation to the dislocation core region.
4. Discussion The dislocation networks analyzed in the as-irradiated Zircaloy-2 specimens are essentially the same as that reported by Holt and Gilbert [4]. High density of u-component dislocations obscure the visibility of the c-component dislocation at any diffracting condition other than the (0002)g-type. The fact that the c-component dislocations are not faulted suggests that they are either a different type of c-component dislocation than those previously identified, which were faulted defects on the @OOl), [2,11,22l or the defects have changed their nature during the course of irradiation (these specimens were at a fairly high fluence: 20 dpa = 6.2 X lO= nm-* in EBR II). Analysis of the dislocation networks in the postirradiation annealed material provided evidence of pure u-component dislocations reacting with c-component dislocations. Possible dislocation reactions would be as follows +(2023) + +(2113) --) ;(2203),
(3)
;(2023)
(4)
+ ;(zllO>
+ +(2203),
+(2205) +;(2203)-[oooi].
(5)
Thus,combining reactions (31, (4) and (5) 2. {(2023) + f(2113)
+ #llO>
+ [OOOi].
formation of unfaulted [OOOll dislocations by this mechanism is highly energetically favourable by 4u*. Eq. (7) has already been proposed by the present authors 1171and is also energetically favourable. Further work, principally electron holography, is required both to ‘quantify’ the extent of solute segregation to the dislocation cores and to ensure that the phase shifts are not caused by hydrides.
5. Conclusions The main conclusions from the present study are as follows: (i) HREM of annealed Zircaloy-2 neutron irradiated at 690 K to 20 dpa shows a high density of c-component dislocations with line directions lying in, or substantially deviated from, the basal plane. Lattice images of the (c)-type dislocations which lie in the basal plane, reveal an unfaulted dislocation structure. (ii) Conventional TEM analysis of the dislocations in postirradiation annealed material shows ( c )-component dislocations with Burgers vectors of the type +(llB) and [OOOll. Large bends in the (c)component dislocation network are associated with multiple (a)-type dislocations. (iii) Electron holography of the (c)-type dislocation core region for both as-irradiated and postirradiation annealed material produces phase shifts having a length of - 3-4 basal planes in isolated regions along the lengths of the dislocations. A possible interpretation is solute segregation to the dislocation core.
(6)
As well, +(1123)+
291
Acknowledgements
~(ii2o)+[oooi]
(7)
Eq.(3)is equivalent to eq. (2) proposed by Griffiths [2]. As far as the authors know, eq. (4) was not proposed but it may be responsible for some of the bends seen in the dislocations. Both eqs. (3) and (4) are highly energetically favourable, by $f’ and a*, respectively, since the energy of a i(2023) dislocation is proportional to a* = cc*/4 + a*/3). Griffith’s [2]eq. (2)is highly energetically unfavourable, by $a*, even if splitting were to occur by climb. If we consider eq. (51, it is only slightly energetically unfavourable, by $a*, and is the opposite to eq. (1) given by Griffiths [21, so eq. (5) is questionable. However, if the overall dislocation reaction is considered, given by eq. (6), then the
The authors are grateful to Dr. R.A. Fleck, Ontario Hydro Research Laboratories, for the use of their electron microscope and to R.A. Holt and R.W. Gilbert for the provision of the neutron-irradiated specimens. One of the authors (D.O.N.) would also like to acknowledge the continuing financial support of the Natural Sciences and Engineering Research Council of Canada through Research Grant A4391.
References [1] D.O. Northwood, At. Energy Rev. 15 (1977) 547. [2] M. Griffiths, J. Nucl. Mater. 159 (1988) 190.
292
B.A. Herring et al. / HREM of (c)-type dislocations in Zircaby-2
[3] A. Rogemon and R.A. Murgatroyd, J. Nucl. Mater. 113
[4] [5] [6] [7] [8] [9] [lo] [ll] [12]
(1983) 256. R.A. Holt and R.W. Gilbert, J. Nucl. Mater. 116 (1983) 127. R.A. Holt and R.W. Gilbert, J. Nucl. Mater. 137 (1986) 185. M. Griffiths, R.W. Gilbert and C.E. Coleman, J. Nucl. Mater. 159 (1988) 405.. A. Jo&sons, P.M. Kelly, R.G. Blake and K. Farrell, ASTM-STP 683 (1979) 46. M. Griffiths and R.W. Gilbert, J. Nucl. Mater. 150 (1987) 169. A. Jostsons, R.G. Blake, J.G. Napier, P.M. Kelly and K_ Farrell, J. Nucl. Mater. 68 (1977) 267. R.G. Fleck, R.A. Holt, V. Perovic and J. Todros, J. Nucl. Mater. 159 (1988) 75. M. Griffiths, R.W. Gilbert and V. Fidleris, ASTM-STP 1023 (1989) 658. V. Perovic, G.C. Weatherley and R.G. Fleck, Can. Metah. Quart. 24 (1985) 253.
[13] J.A. Jensen and W.A. Backofen, Can. Metall. Quart. 11 (1972) 39. 1141 O.T. Woo, G.J.C. Carpenter and S.R. MacEwen, J. Nucl. Mater. 87 (1979) 70. 1151 R.A. Holt, M. Griffiths and R.W. Gilbert, J. Nucl. Mater. 149 (1987) 51. [16] P. Merle, J. Nucl. Mater. 144 (1987) 275. [17] R.A. Herring and D.O. Northwood, J. Nucl. Mater. 159 (1988) 386. 1181 D. Gabor, Proc. R. Sot. London, Ser. A 497 (1949) 45. 1191A. Tonomura, Rev. Mod. Phys. 59 (1987) 639. [20] K.J. Hanszen, in Advances in Electronics and Electron Physics, ed. L. Marton, vol. 59 (Academic Press, New York, 1982) p. 1. 1211 K.J. Hanszen, J. Phys. D19 (1986) 373. [22] O.T. Woo and G.J.C. Carpenter, J. Nucl. Mater. 159 (1988) 397.
HREM of ( c) -type dislocations in high-dose neutron-irradiated and postirradiation annealed Zircaloy-2 Rodney A. Herring and T. Tanji Tonomura Electron Wavejiont Project, Exploratory Research for Advanced Technobm, JmC, Hatoyama, Saitama 350-03, Japan
c/o
ARL Hitachi Ltd.,
Derek 0. Northwood Engineering Materi& Group MechanicalEngineering Department, Universityof W&or,
W&or,
Ontario, Canada N9B 3P4
High resolution electron microscopy (HFtEM) was performed on (c)-type dislocations in annealed Zircaloy-4 which bad been neutron irradiated to 20 dpa at 690 K Lattice images of the (c)-type. dislocations, which lie on the basal plane, reveal unfaulted dislocation structure. Conventional TEM analysis of these dislocations in postirradiation annealed Zircaloy-2 shows that the (c)-component dislocations have Burgers vectors of i(1123) and [OOOl].Large bends in the (c)-component dislocation network are associated with multiple (a)-type dislocations. Initial results by electron holography of the (c)-type dislocation core region produces phase shifts having a length of m 3 basal planes in isolated regions along the length of the dislocations. The shift may occur due to large solute concentrations.
1. Introduction
Zirconium alloys are used in the nuclear industry because of their much lower neutron absorption per unit strength than other commercially available structural materials. Irradiation-induced growth and irradiation-enhanced creep are of concern both for fuel cladding and nuclear reactor structural components such as pressure tubes and calandria tubes. To understand growth and creep it is necessary to understand the nature of the irradiation-induced microstructural changes. This includes the altering of the zirconium alloy’s chemistry. Our understanding of the evolution of microstructure in zirconium alloys during irradiation has improved substantially over the last 10 years. In his 1977 review of irradiation damage in zirconium, Northwood [l] summarized the interpretation of the microstructure developed during irradiation as follows: (1) Only dislocation loops having Burgers vetior of b = f(ll2o), lying on or close to prism {ioio)planes, were formed during irradiation. They were: mostly aligned in rows parallel with the trace of the basal plane WJO1). (2) Both interstitial and vacancy loops were formed during neutron irradiation, the relative proportions 0022-3115/93/$06.00
varying with irradiation temperature. At high temperatures (2 675 IO the majority were vacancy in nature and at low temperatures (5 575 K) the majority were interstitial in nature. Interstitial loops were markedly elliptical. As pointed out by Griffiths [2] this simple interpretation is invalid. The microstructure developed in Zralloys during irradiation is more complicated as indicated by specimens of varying composition and thermo-mechanical history that have been irradiated to higher fluences. In particular, Griffiths [2] noted that, with the development of techniques for microchemical analysis, it has become increasingly obvious that alloying elements or impurities are important in affecting and contributing to the evolving microstructure. Also attention has been focused on c-component dislocations and loops and their contribution to accelerated, or “breakaway growth” [3]. Generally, whenever large growth strains have been observed, c-component dislocations were also present [4,5]. cxomponent dislocations can be introduced into the material during neutron irradiation [2,4-111 and during cold-working [1216]. In our earlier work [17] some data were given on the analysis of a dislocation network in neutron-irradiated Zircaloy-2. More work is presented in this paper.
8 1993 - Elsevier Science Publishers B.V. All rights reserved
R.A. Herring et al. / HREM of (c)-type &locations in Zircaloy-2
285
The comparatively new technique of electron holography [18] is used to explore dislocation cores. Excellent reviews on holography have been provided by Tonomura 1191 and Hanszen [201. Holograms are produced from the interference between a reference wave and an object wave. Contained within the hologram is the amplitude and phase information of the object. Conventional electron images only contain the amplitude or intensity distribution of the object.
2. Experimental details Two thin foil Zircaloy-2 specimens that had been neutron irradiated (E > 1 MeV) as bulk material in the EBR-II reactor, Argonne National Laboratory West, at 690 K to a fluence of 20 dpa were supplied to us by Chalk River Laboratories of AECL Research. Composition was 1.2-1.7 wt% Sn, 0.07-0.20 wt% Fe, 0.05-0.15 wt% Cr, 0.03-0.08 wt% Ni, an oxygen concentration of 1500 ppm, with the balance being Zr. One thin foil was subjected to an in-situ annealing experiment in an AEI EM7 high voltage electron microscope. The foil was held at 1075 K for over 20 min. A JEM lOO-CX TEM, capable of large two-axis tilt angles, was used to examine the dislocation network. Burgers vectors were determined by the g - b = 0 condition, i.e., zero contrast, plus the diffraction behaviour of the dislocation when they were in contrast, such as when g-b = 1, 2 or higher. Line directions, U, and habit planes, R, were determined with the aid of a stereogram and noting the projected directions of the dislocations. A dedicated high resolution microscope, Hitachi HF 2000, was used for the lattice imaging. This microscope has been fully characterized for its high resolution potential, having a well-controlled environment with respect to temperature fluctuations and noise, a spherical aberration of the objective pole piece of 0.65 mm and a chromatic aberration of the cold field emission electron gun of 0.3 eV for an acceleration energy of 200 keV. The point to point resolution is 0.21 nm, and thus the line widths (which are much easier to see than point to point images) of (0001) Zr at 0.51 pm were quite easily seen in our micrographs. The electron gun had a very high emission, 30 uA, which enabled very low beam divergence to be used N 0.3 to 0.5 mrad. The lattice images were taken using the standard photographic method for lattice images which starts from about exact focus to about the Scherzer focus (48.5 nm) in 5 to 7 steps. Multiparallel-beam lattice images were taken using the (0OOl)g family of
Fig. 1. The hologram is formed by illuminating the object (the dislocatioti core region) with a coherent electron beam wider than the object. The half of the beam that bypasses the object (the Zr matrix) serves as the reference wave. Electron lenses focus both halves of the wave, and the biprism superposes them in the image plane, where the interference pattern is recorded on film.
systematic reflections at orientations parallel to the basal plane between the 1100 and 1210 zone axes. This condition was used to clearly see the c-component dislocations without any image contribution by a-component dislocations. To clearly see the dislocation core region, the Zircaloy-2 crystal was tilted _ 15 to 20 mrad from the basal plane, i.e., the main beam approximately occupied the [0008] reflection position. Fig. 1 illustrates the procedure for formation of an off-axis hologram image. In one-half of the specimen plane, the electron wave illuminates the object or region of interest. The other half of the plane is reserved for a reference wave. The reference wave can pass through the vacuum or a part of the specimen with a constant inner potential and thickness [21]. A positive potential of about 100 V is applied to a centrally located Cr-coated glass-filament of the electron biprism beyond the objective lens which makes the image and the reference wave overlap to form a hologram or interference pattern. Magnification of the interference pattern by the projector lenses enable the fringes to be easily seen on the recording media, i.e., film. Contained within the hologram is both the phase and amplitude information of the object.
3. Results 3.1. HREM 3.1.1. As-irradiated
material The DOO21-type diffraction
vector was used to see the c-component dislocations in the as-irradiated material. One region was found with a particularly high
R.A. Hemhg et al. / HREhi of (c)-type dislocationsin Zircaloy-2
286
density of c-component dislocations having line directions substantially deviated from the basal plane, whereas most c-component dislocations had line directions lying in the basal plane. A large number of lattice images were taken of the c-component dislocations using weak diffracting conditions, i.e., [00081. Fig. 2 shows two examples of the lattice images taken at many places between the (liOO> and (l%O> zone axes: one is along a straight section, i.e., in the basal plane direction, and the other is at a bend, i.e., away from the basal plane direction, where the c-component dislocations lie between the bracketed regions. These lattice images show that all of the basal planes are intact, i.e., there are no missing planes of atoms, and no stacking faults in the basal plane are present in the as-irradiated material. 3.1.2. Postirradiation annealed material The analysis of a dislocation network which was adjacent to a grain boundary and contained c-component dislocations is given in fig. 3. Only a few of the micrographs used to characterize the dislocation network are given and these show the two-beam electron diffraction vector and the crystal orientation (given within the square brackets). To fully determine the diffraction behaviour of the dislocations, pictures were taken of the dislocations at a range of deviations from the exact Bragg condition. A summary of the contrast given by the dislocation segments is given in table 1, with an asterisk (*) for the dislocation segments in contrast and a cross (XI for the dislocation segments out of contrast, given for their respective electron diffraction conditions. Fig. 4 is a schematic illustration of the dislocation network. In fig. 3 and fig. 4 the
Fig. 2. Lattice images of as-irradiated material; (a) along a straight section, i.e., in the basal plane direction; (b) at a bend, i.e., away from the basal plane direction. The c-component dislocations lie between the bracketed regions.
c-component dislocations are contained within the dislocation segments marked B, C, D, E, K, M and possibly 0 and are clearly in contrast for the
Table 1 Diffraction contrast of dislocation segments given in fig. 4 Diffraction vector and orientation
x :
Contrast of dislocation segment (* : in,
out)
ABCDEFGHIJKLMNOPQR
rtooo4[oiioI x *Zl12[oilo] -
*
*
*
*
*
x
x
x
x
*
x
*
x
-
x
x
x
*
x
*
*
*
*
*
*
*
*
*
*
x
*
*
*
*
*
*
x
*
x
*
x
*
*
*
*
*
*
*
*
*
*
*
*2ii~oiloi * ~lioi[ll2o] *
*
*
*
*
x
-
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
x
*
x
*
x
*
*
*
*
*
*
*liol[oili] *
*
*
*
*
*
*
x
*
x
*
x
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
*
x
*
*
*
*
*
*
*
*
*
*
*
*
*Zllqoilo]
*loii[l2lo] *
*
*
*
*20zq12101
*
-
*
x*-****-*******
*2ii2@3231
*
*
*
*
*
R.A. Herring et al. / HREM of (c)-type didocations in Zircakq-2
Fig. 3. Analysis of dislocation network in postirradiation annealed material. The network was adjacent to a grain boundary and contained c-component dislocations; (a) c-component dislocations in contrast and a-component dislocations out of contrast; (b) some of the dislocation network in contrast; (cl and (d) dislocations having b = [Oool] out of contrast, i.e., segments B and D; (e) dislocation haying b = f[%lz] out of contrast; (0 and (h) dislocation haying b = f[Z113] out of contrast. Micrographs for (d), (s> and (h) orientations were used to determine the line direction, u, for the dislocation segments, D, b = [OOOl]and C, b = #113].
288
RA. Herring et al. / HREIU of (c >-type dislocations in Zircaloy-2
whereas segment F is also out of contrast for + 2ii2g[OilO] and + 2ii2g[2ri23] and has b = $1131, i.e., the opposite c-component to that of segment C. The pure c and (c + a)-dislocation segments are joined to the pure a-component dislocations having b = $llO] which are all out of contrast for the 0004g[OllO] and given by the symbols A, G, I, L, P, Q, and R. The pure u-component dislocations are consistent with the Burgers vectors of the c and (a + c>dislocations. The bend in the c-component dislocation, given by segment E, is a complex junction between c and u-component dislocations. Note that there is a high density of a-component dislocations associated with segment E. The Burgers vectors of the dislocation segments of the dislocation network illustrated in fig. 4 are listed in table 2. There are segments of the dislocation network which could not be identified with high confidence including segments E, K, M, N, and 0. The line directions and habit planes for the dislocations labelled B, C, and D in fig. 3, which had Burgers vectors b = [OOOl]and b = $[2113], were found to have line directions close to the prism directions and lie within the (0001) habit planes. Thus, the c-component dislocations lie on non-glide planes. Fig. 5 is a series of micrographs and schematic diagram showing a c-component dislocation network
Table 2 Burgers vectors of dislocation segments given in fig. 4 Fig. 3 (continued).
0004g[O’ilO]. The segments given by B, D and F are pure c-component and are out of contrast with +211Og[OilO]. There appears to be an overlap of the pure c and (a + c)-dislocation, especially in the region of the F segment. Segment C is out of contrast with +2112g[OilO] and has a Burgers vector of b = $!llJ],
Dislocation segment A B C D E F
;[zllol
DJoo11 j[Z113]
mw ? [OCKU] and f[?113]
G
;[2110]
H
+[llzo]
I
f[ZllO]
J K L M N 0 P
;r11zo1 ? f[ll20] ? ? ? ;[zllo]
Q
~[sllol ;[zllo]
R Fig. 4. Schematic diagram of dislocation network in fig. 3.
Burgers vectors b
?: uncertain of analysis.
R.A. Herring et al / HREM of (c)-type disbcatiom
running through a precipitate. Fig. 5a is the OOO& diffracting vector and figs. 5b and 5c are bright and dark field micrographs using long. These micrographs show that the dislocation segment A (see fig. 5d) is out of contrast (shadowy with g close to Bragg) and that segment B has interacted with an a-type dislocation to form a bend in the c-component segment and a new c-component dislocation with a different Burgers vector has formed. All of the i(2ti3) dislocations are in contrast with
in Zircaloy-2
289
the lOi1 diffraction vectors and weak beam imaging with this condition showed the presence of some +(20?3) dislocations associated with stacking faults. This agrees with Griffiths’ I21 suggestion that edge sections of [cl or (c + a)-dislocations climb by splitting into partials during neutron irradiation of Zircaloy-2 at 690 K. The reactions being [OOOl] + ;(2023)
+ +(2023),
+(1123) --) *(2023) + ;(0223),
(1) (2)
d
Fig. 5. Series of micrcgraphs and schematic diagram showing a c-component dislocation network running through a precipitate in postirradiation annealed material; (a) 0006g diffracting vector; (b) and Cc) bright and dark field micrographs using 1Oug; (d) schematic diagram.
290
RA. Herring et al. / HREM of ( c>-typedislocations in Zircaloy-2
with the result that both k(2023) dislocations and stacking faults would be present. However, the number of c-component dislocations having stacking faults were at a very much lower density than the c-component dislocations. An example is given in fig. 6 which are dark field and bright field images taken using the Olilg at the [1213lZA. Many of the stacking faults observed were missing their :(20?3) partial dislocations (out of the foil) and faults are visible only in the weak-beam micrographs.
Fig. 7. Electron holograms of (a) as-irradiated and (b) postirradiation annealed material showing fringe shifts which occur at isolated regions along the c-component dislocations and the dislocation line direction. In summary, the c-component dislocation network in postirradiation annealed material consists of [OOOl] and i(1213) dislocations and faults on (0001) having $(2023) partial dislocations. 3.2. Electron holography
Fig. 6. Dark field (a) and bright field (b) images of postirradiation annealed material taken using the Ol’Ilg at the [1213lZA showing stacking faults.
Electron holography was performed on both asirradiated and postirradiation annealed specimens. Typical holograms for as-irradiated and postirradiation annealed material are given in fig. 7a and 7h, respectively. The common feature of these holograms are the fringe shifts which occur at isolated regions along the c-component dislocations and the dislocation line direction. In fig. 7a, as-irradiated material, the phase shifted region is along the @001) planes and can be seen to have a width of 3-4 basal planes, whereas in fig. 7b, postirradiation annealed material, the phase shifted region is at a bend in the dislocation where there should be intersections from other dislocations.
R.A. Herring et al. / HREM of (c >-typedislocationsin Zircaloy-2
These fringe shifts, which were present for both as-irradiated and postirradiation annealed material, are indicative of small changes in the mean inner potential believed to be due to solute segregation to the dislocation core region.
4. Discussion The dislocation networks analyzed in the as-irradiated Zircaloy-2 specimens are essentially the same as that reported by Holt and Gilbert [4]. High density of u-component dislocations obscure the visibility of the c-component dislocation at any diffracting condition other than the (0002)g-type. The fact that the c-component dislocations are not faulted suggests that they are either a different type of c-component dislocation than those previously identified, which were faulted defects on the @OOl), [2,11,22l or the defects have changed their nature during the course of irradiation (these specimens were at a fairly high fluence: 20 dpa = 6.2 X lO= nm-* in EBR II). Analysis of the dislocation networks in the postirradiation annealed material provided evidence of pure u-component dislocations reacting with c-component dislocations. Possible dislocation reactions would be as follows +(2023) + +(2113) --) ;(2203),
(3)
;(2023)
(4)
+ ;(zllO>
+ +(2203),
+(2205) +;(2203)-[oooi].
(5)
Thus,combining reactions (31, (4) and (5) 2. {(2023) + f(2113)
+ #llO>
+ [OOOi].
formation of unfaulted [OOOll dislocations by this mechanism is highly energetically favourable by 4u*. Eq. (7) has already been proposed by the present authors 1171and is also energetically favourable. Further work, principally electron holography, is required both to ‘quantify’ the extent of solute segregation to the dislocation cores and to ensure that the phase shifts are not caused by hydrides.
5. Conclusions The main conclusions from the present study are as follows: (i) HREM of annealed Zircaloy-2 neutron irradiated at 690 K to 20 dpa shows a high density of c-component dislocations with line directions lying in, or substantially deviated from, the basal plane. Lattice images of the (c)-type dislocations which lie in the basal plane, reveal an unfaulted dislocation structure. (ii) Conventional TEM analysis of the dislocations in postirradiation annealed material shows ( c )-component dislocations with Burgers vectors of the type +(llB) and [OOOll. Large bends in the (c)component dislocation network are associated with multiple (a)-type dislocations. (iii) Electron holography of the (c)-type dislocation core region for both as-irradiated and postirradiation annealed material produces phase shifts having a length of - 3-4 basal planes in isolated regions along the lengths of the dislocations. A possible interpretation is solute segregation to the dislocation core.
(6)
As well, +(1123)+
291
Acknowledgements
~(ii2o)+[oooi]
(7)
Eq.(3)is equivalent to eq. (2) proposed by Griffiths [2]. As far as the authors know, eq. (4) was not proposed but it may be responsible for some of the bends seen in the dislocations. Both eqs. (3) and (4) are highly energetically favourable, by $f’ and a*, respectively, since the energy of a i(2023) dislocation is proportional to a* = cc*/4 + a*/3). Griffith’s [2]eq. (2)is highly energetically unfavourable, by $a*, even if splitting were to occur by climb. If we consider eq. (51, it is only slightly energetically unfavourable, by $a*, and is the opposite to eq. (1) given by Griffiths [21, so eq. (5) is questionable. However, if the overall dislocation reaction is considered, given by eq. (6), then the
The authors are grateful to Dr. R.A. Fleck, Ontario Hydro Research Laboratories, for the use of their electron microscope and to R.A. Holt and R.W. Gilbert for the provision of the neutron-irradiated specimens. One of the authors (D.O.N.) would also like to acknowledge the continuing financial support of the Natural Sciences and Engineering Research Council of Canada through Research Grant A4391.
References [1] D.O. Northwood, At. Energy Rev. 15 (1977) 547. [2] M. Griffiths, J. Nucl. Mater. 159 (1988) 190.
292
B.A. Herring et al. / HREM of (c)-type dislocations in Zircaby-2
[3] A. Rogemon and R.A. Murgatroyd, J. Nucl. Mater. 113
[4] [5] [6] [7] [8] [9] [lo] [ll] [12]
(1983) 256. R.A. Holt and R.W. Gilbert, J. Nucl. Mater. 116 (1983) 127. R.A. Holt and R.W. Gilbert, J. Nucl. Mater. 137 (1986) 185. M. Griffiths, R.W. Gilbert and C.E. Coleman, J. Nucl. Mater. 159 (1988) 405.. A. Jo&sons, P.M. Kelly, R.G. Blake and K. Farrell, ASTM-STP 683 (1979) 46. M. Griffiths and R.W. Gilbert, J. Nucl. Mater. 150 (1987) 169. A. Jostsons, R.G. Blake, J.G. Napier, P.M. Kelly and K_ Farrell, J. Nucl. Mater. 68 (1977) 267. R.G. Fleck, R.A. Holt, V. Perovic and J. Todros, J. Nucl. Mater. 159 (1988) 75. M. Griffiths, R.W. Gilbert and V. Fidleris, ASTM-STP 1023 (1989) 658. V. Perovic, G.C. Weatherley and R.G. Fleck, Can. Metah. Quart. 24 (1985) 253.
[13] J.A. Jensen and W.A. Backofen, Can. Metall. Quart. 11 (1972) 39. 1141 O.T. Woo, G.J.C. Carpenter and S.R. MacEwen, J. Nucl. Mater. 87 (1979) 70. 1151 R.A. Holt, M. Griffiths and R.W. Gilbert, J. Nucl. Mater. 149 (1987) 51. [16] P. Merle, J. Nucl. Mater. 144 (1987) 275. [17] R.A. Herring and D.O. Northwood, J. Nucl. Mater. 159 (1988) 386. 1181 D. Gabor, Proc. R. Sot. London, Ser. A 497 (1949) 45. 1191A. Tonomura, Rev. Mod. Phys. 59 (1987) 639. [20] K.J. Hanszen, in Advances in Electronics and Electron Physics, ed. L. Marton, vol. 59 (Academic Press, New York, 1982) p. 1. 1211 K.J. Hanszen, J. Phys. D19 (1986) 373. [22] O.T. Woo and G.J.C. Carpenter, J. Nucl. Mater. 159 (1988) 397.