Langmuir 1987, 3, 867-873
867
Stress Corrosion Cracking of Zirconium Alloyst B. Cox Reactor Materials Division, Atomic Energy of Canada Limited, Chalk River Nuclear Laboratories, Chalk River, Ontario KOJ 1J0, Canada Received June 20,1986 The environments in which zirconium alloys suffer stress corrosion cracking are summarized and the conditions leading to delayed hydride cracking identified. The two processes are clearly distinguishable fractographically,and these differences are illustrated. The mechanisms of stress corrosion (under conditions where hydride cracking is absent) are discussed in terms of the important steps in the process. These rate-controlling steps are examined for their relevance to the pellet-clad interaction (PCI) phenomenon that can lead to the failure of nuclear reactor fuel cladding. The active fission product species causing PCI failures are discussed, and it is concluded that between I2 and Cs/Cd as culprits, the choice is iodine.
Introduction Zirconium alloys are very similar, both chemically and metallurgically, to titanium alloys. This similarity was used, when we were first investigating stress corrosion cracking (SCC) in zirconium alloys,l to decide on those environments that would be worthy of study (Table I), based on the already extensive studies of SCC in titanium alloy~.~J At that stage little was published about the SCC of zirconium alloys, and only some work in methanolic solutions4 and in ferric chloride solutions5p6and hydrochloric acid7was in the open literature. However, General Electric had already experienced some brittle failures of experimental fuel irradiated at high power.8 In this instance they suspected SCC by fission products, thought to be iodine, and had proceeded to demonstrate SCC in iodine vapor in the laborat~ry.~ Our interest in this field was stimulated by further brittle failures of irradiated fuel element cladding that did not appear to have originated purely from mechanical causes, or from the secondary hydriding of the sheaths that often followed a primary defect, and in particular by the onset of an epidemic of fuel failures that followed the first on-load refueling of the Douglas Point and Pickering 1reactors.lOJ1 Although this problem was essentially eliminated for CANDU fuel by modifications to the refueling sequence and the introduction of the CANLUB graphite coating in Dec. 1972, pellet-clad interaction failures that are basically SCC by fission products have continued to present a significant economic penalty in other reactor systems. Solutions for BWR fuel now seem to be centering on “zirconium-liner” cladding,12J3although graphite coatings seem to be almost equally as g00d.l~ This problem has spawned extensive studies of the detailed mechanism of the transgranular cracking p r o ~ e s s ’ ~and J ~ much discussion about the important fission product species and its form.16J7 The remainder of this paper will be devoted to recent work in these two areas. SCC Mechanism The environments in which SCC has been observed in zirconium alloys are listed in Table I in comparison with those found to crack titanium alloys. In the zirconium alloy column an indication of the propagation mode is given and the earliest reference to cracking of zirconium alloys in this environment (or class of environments). One observation about the entries in this table would be the prevalence of mixed intergranular (IG) and transgranular (TG) cracking. ‘Presented at the symposium on “Corrosion”, 191st National Meeting of the American Chemical Society, New York, NY, April 13-18, 1986. 0743-7463f 87 f 2403-0867$01.50 f 0
The mechanism of intergranular cracking is relatively easily established. For cracking in solutions it seems to be usually a dissolution process (Figure 1)and weight losses are generally observed.lJ8 The corresponding process in halogen vapors is argued to be the removal of zirconium from the grain boundaries as a volatile halide.19,20 One (1) Cox, B. “Environmentally Induced Cracking of Zirconium Alloys”; Atomic Energy of Canada Ltd., Reports (PtI) AECL-3551, (PtII) AECL-3612. (PtIII) AECL-3799: 1970/1971. (2) Papers on titanium: Proceedings of the Conference on the Fundamental Aspects of Stress Corrosion Cracking, Staehle, R. W., Forty, A. J., van Rooyen, D., Eds; Columbus, OH, Sept. 1967; National Association of Corrosion Engineers: Houston. (3) Blackburn, M. J.; Feeney, J. A,; Beck, T. R. “Stress Corrosion Cracking of Titanium Alloys”; Boeing Scientific Research Laboratory, Report DI-82-1054, 1970; Adv. Corros. Sci. Technol. 1973, 3, 67-292. (4)Mori, K.; Takamura, A.; Shimose, T. Corrosion 1966, 22, 29. (5) Dunham, J. T.; Kato,H. ‘Stress Corrosion Cracking Susceptibility of Zirconium in Ferric Chloride Solution”; US. Bureau of Mines Report, BM-R1-5784; 1961. (6) Thomas, K. C.; Allio, R. J. Nucl. Appl. 1966, I , 252. (7) Payne, B. S.; Priest, D. K. Corrosion 1961, 17,196t. (8) Lyons, M. F.; Coplin, D. H.; Jones, G. G. “High Performance UOz Program”; Quarterly Progress Reports; GEAP-3771-10, -11, -12; Weidenbaum, B., Ed.; Oct. 1963-April 1964. (9) Rosenbaum, H. S.; Davies, J. H.; Pon, J. Q. “Interaction of Iodine with Zircaloy-2”;G.E. Report GEAP-51005; 1966; Electrochem. Technol. 1966, 4, 153. (10) Robertson, J. A. L. “Nuclear Fuel Failures, their Causes and Remedies”, Proceedings of the Joint ANSICNS Topical Meeting on Commercial Nuclear Fuel Technology Today, Toronto, April 1975; CNS-ISSN 0068-8517 (75-CNA/ANS-100) p 2-2. (11)Fanjoy, G. R.; Bodie, L. L.; Page, R. D.; Tarasuk, W. R.; Debnam, H. R. “Canadian Nuclear Fuel Operating Experience”, ref 10; pp, 1-66. (12) Rosenbaum, H. S.; Davies, J. H.; Adamson, R. B.; Tucker, R. P.; Rowland, T. C.; Paustian, H. H.; Thompson, J. R.; O’Boyle, D. R. “large Scale’Demonstration of Barrier Fuel”, Proceedings of ANS Meeting on LWR Fuel Performance, Orlando, April 1985; Vol. 2, p 7-63, DOE/ NE/ 34130-1. (13) Inoue, K.; Suzuki, K.; Maki, H.; Yasuda, T.; Oi, N.; Hayashi, Y.; Wakashima, Y.; Ogata, K.; Junkrans, S.; Vesterlund, G.; Lysell, G.; Ronnberg, G.“An Overview of the Joint Development Work on PCI Remedy Fuel”, ref 12; p 6-1. (14)Cox, B.; Wood, J. C. “Iodine-Induced Cracking of Zircaloy FuelCladding-A Review”, Presented at ECS Fall Meeting, New York, Oct. 1974. Proceedings volume: Corrosion Problems in Energy Conversion and Generation; Tedmon, Jr., C. S., Ed.; Electrochemical Society: New York; p 275. (15) Haddad, R.; Cox, B. ‘Investigations of Crack Initiation During Environmentally-Induced Cracking of the Zircaloys”, Atomic Energy of Canada Ltd., Report AECL-8104; Oct. 1983. (16) Cox, B.; Surette, B. A,; Wood, J. C. “Pellet-Clad Interaction Failures: Stress-Corrosion Cracking by Iodine or Metal-Vapour Embrittlement by Cesium/Cadmium Vapours”, Proceedings of the Conference on Environmental Degradation of Engineering Materials by Aggressive Enuironments, Blacksburg, VA, Sept. 1981; V.P.I. and S.U.,p 293. (17) Cox, B.; Surette, B. A.; Wood, J. C. J.Nucl. Mater. 1986,138, 89. (18) Cox, B. Corrosion 1972, 28, 207. (19) Wood, J. C. J. Nucl. Mater. 1972/73, 45, 105. (20) Cox, B.; Wood, J. C. “The Mechanism of SCC of Zirconium Alloys in Halogens”, Proceedings of the International Conference on Mechanisms of Environment Semitive Cracking of Materials, Guildford, U.K., April 1977; Swann, P. R., Ford, F. P., Westwood, A. R. C., Eds; Metals Society: London; p 520.
0 1987 American Chemical Society
868 Langmuir, Vol. 3, No. 6, 1987
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T a b l e 1. Environments Causing Cracking ofT i t a n i u m a n d Zirconium Alloysa SCC in Zr allow environment organic liquids methanol and solutions with I,, Br,, NaCI. HCI higher alcohols with similar additives chlorinated hydrocarbons (e.&CCI., CHCIJ Freon8 ethylene glycol aqueous solutions HCI NaCI, KCI, etc. FeC18, CuCI, hot and fused halidea NaCI, KCI. CsCI. etc. KNO,/NaNO, KBr/CsI KNO,/NaNO, NaCl halogen and halogen acid vapors
+ +
metala (solid. liquid, vapor) Hg,Cd, Cs, Zr oxides and oxy acids of N2 N,?, nitric acid hydrogen gas sulfuric acid
temp,' OC
first O M ,
SCC in Ti alloys
yes/no
cracking mode'
ref
RT
Yea
IG initiation TG propagation
1,4
RT
Yes NT yes yes only in I, solution
mixed IC/TG
4
IC IC
51 51 51
3 ~ e l m RT RT RT
25-150 25-154 RT
Yes yes (needs polarization) yes
300-450 300-400
TG (DCB)
IC
7
mixed IG/TG
52 5,6
mixed IG/TG mixed IG/TC mixed IT/TC
1 18 21 51 9
TG
1
T
53, 54 55
IC
300400
Tk (DCB)
RT 250-500 RT-
Yes
RT RT-lo0 RT-350 RT
Yes Yea NT
NT
TG
.A more detailed description of the various environments causing cracking is given in ref 56. 'RT = r m m temperature. (DCB indica& fractures only in precracked double cantilever beam specimens.
Figure 3. Etched section of specimen cracked in KN03/ NaNOs/NaCI a t 300 "C showing intergranular internal oxide (20-rm bar).
Figure 1. Weight lcea by zirconiumdoya m methanol/l%iodine solution.
f.
Y
Figure 4. Higher magnification view showing 'double-layer" nature of the internal oxide layers (IO-pm bar).
Figure 2. Intergranular cracking of Zircaloy fuel cladding in KNOJ/NaNO,/NaCI a t 300 "C (IO-" bar).
exception to this simple explanation of IG cracking is the phenomenon observed in fused nitrate/chloride mixtureaZ1
Here the specimen undergoes a weight gain, and although the fracture surface (Figure 2) is apparently fully IG, etching of metallographic sections (Figure 3) shows that the grain boundaries are penetrated by a form of intemal oxidation process. Higher resolution scanning electron microscope observations show that an open channel for (21)Cor, B. Oxid. Met. 1971,3.399.
Stress Corrosion Cracking of Zirconium Alloys
Langmuir, Vol. 3, No.6, 1987 869
GOLD
GREY
BLACK
IHITE
Figure 7. Bands of differentdored oxide on DHC fracture face in a Zr-2.5% Nh alloy pressure tube (2." bar).= - p~ q& p?R*e??q ,.;., , ,>".#.>'.''-
..
Pigum 8. End-cap weld cracking in Zr-2.5% Nb clad fuel: (a) hydride at crack tip; (b) general view of crack starting at weldupset notch."
Figure 5. Crack tip showing absence of hydride hut possibly a small void ahead of the furthest oxide penetration (10-Mmbar).
1mm1m1. .c
Figure 6. Electron m i c r m p p replica of oxidid fracture surface formed by DHC in a Zr-2.3 w Yo Sb pressure t u b e (l-pm bar). passage of the fused salt mixture is maintained (Figure 4) between the two oxidized grain boundary faces. Thus, despite the high Pilling-Bedworth ratio (- 1.56) for conversion of metal to oxide, and the undoubtedly large stresses that this generates in the metal, the two faces of the crack do not sinter together, but maintain an open channel. This is possible because the growth of zirconia films proceeds entirely by the inward migration of oxygen anions and because of the very high comprensive strength of zirconia at low temperatures. The precise method by which the intergranular crack is formed in the first place has not been established. It may be that a small amount of zirconium is removed a t the crack tip as a volatile chloride before the crack walls are oxidized. The highly oxidizing nature of the salt, the absence of moisture or of absorbed hydrogen in the metal, and the absence of any precipitated hydride at crack tips (Figure 5) show that this is not a form of delayed hydride cracking. In zirconium alloys delayed hydride cracking (DHC) is almost entirely transgranular in f ~ r m ? ~although .~' grain boundary precipitation of hydrides in unstressed specimens (22) Ells, C. E. Can. IML Metall. Annu. Vol. 1978. 17.32. 123) Coleman. C. E.: Ambler. J. F. R. 'S-atihilitv of Zirconium
AU&tn Delayed Hyd-en
".+ , I..
,
,
,
,
,
,
II
I.,
1.0
1 2
2.1
I I
.
1.31,
.
..I
Figure 9. Solubility of hydrogen in zirconium measured from the temperature dependence of incubation time for DHC and compared with other techniques for measuring the solubility?* is quite common.s Contrary to the situation in titanium alloy^,^ DHC in zirconium alloys is only possible in the presence of precipitated zirconium hydride particles." The velocities of DHC cracks in zirconium alloys can be modeled well by the stress-induced growth of hydrides at a notch, or the crack tip, and their fracture once a critical size is reached.n The cracked hydrides produced in this manner may then oxidize to produce a transgranular "oxide-replica" of the original fracture face that preserves the details of the original fractured hydride (Figure 6), even
I~~~~~~~ I ~~
Crickiig", h c e e d i n g s of the 3rd Interm-
tioml Conference on Zzreonrum m the Nuclear Industry. Quebec City, 1976; ASTM-SPT-633, p 589. (24) Cox. B. 'Hydride Cia& e Initiators for Streas Corrosion
Cracking of Zirdoya", Rmeedings of the 4th h t e r m t i o m l Conference on Zirconium in the Nuclear Industry, Stratford-upon-Avo". U.K..June 1978: ASTM-STI-681. p 306.
(25) Bradbrook, J. S.;Lorimar. G. W.;Ridley. N.J.Nwl. Meter. 1972, 42. 142. '(26) William, D. N . J. Imt. Met. 1962,91, 147. (27)Simpson. L. A. Met. TIOM. A 1981,12A,2113.
870 Longmuir, Vol. 3,No.6, 1987
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Figure 10. Appearance of steps on pseudocleavage faces formed on Zircaloy-2 by delayed hydride cracking (note -90° angles) (1-pm bar).
though the remains of this hydride may have diffused away when the temperature was raised to oxidize the crack fa=. I t was this sequence of cracking and oxidation that gave rise to the pressure tube failures (Figure 7) in CANDU and similar cracking has been seen in the end-caps (Figure 8) of fuel claddingJO In titanium alloys there has been much disagreement as to whether or not transgranular SCC is a form of DHC.S1 In this alloy system it is difficult to reach an unequivocal conclusion because DHC seems to be possible in the absence of precipitated hydrides and the fraetography of the two processes seems to he identical. In zirconium alloys DHC is not possible in the absence of precipitated hydrides, and in fact the onset of DHC and the known hydrogen content of the zirconium alloy sample have been used3* to measure the hydrogen solubility (Figure 9) in zirconium alloys at lower temperatures than are possible by most other techniques. The results of the two methods are in excellent agreement a t comparable temperatures. Because the DHC crack in a zirconium alloy is passing through a hydride phase, which is either tetragonal or cubic, the fine detail on the fracture surface shows angles of 90' in the steps (Figure 10). The DHC cracks also proceed through hydrides that occur on a number of preferred planes for precipitation or on grain boundaries 80 that a constant relationship to the position of the basal pole would not be expected. By contrast, the transgranular SCC process in zirconium alloys proceeds on the basal plane. Although a few sketchy reports in the literature33suggested that it might deviate significantly from the basal plane, as appears to be the case for SCC of titanium alloys?% recent work on the cracking of large-grained zirconium alloy specimens (oriented by a micro-back-reflection Laue technique) has shown that TG cracks propagate within *lo of the basal plane.% This is within the present accuracy of the technique. If these ~
(28) Perry",
E. C. W. Nuel. Energy (Br. N u l . Emrgy Sof.) 1978,
I?, 95. (29) Cox, B. 'Failure Analysis in the Canadjan Nudem Indwfry', Pmeeedmna 01 the htWMtiOMl Metdlonmnhthle Society . . . Svmposrum. . . Bostnn. A i g . l9W. Chapvr 6. p VI-3. (301Simpon. C. J.: Ells. C . E. J . Nurl. Mote,. IY74.62.289. 1311 Wanhill. R. J. H. Rr. Corms. J . 1976. 10.69. (32) Coleman, C. E.; Ambler, J. F. R S&: Meroll. 19.%3,17,71. (33) Peehs. M.; Svhle. H.; Steinberg. E 'Outof-Pile Testingofldiw Streas Corrosion Cracking in Zirealoy Tubing in Relation to the Pellet Cladding Interadion Phenomenon", Proceedings of the 4th I n t e m t i o m l Conlerence on Zirconium in the Nuek-ar Industry, Stratford-upon-Avo", U.K., June 1978; ASTM-STP-681, p 244. (34) Wanhill. R. J. H. "Fraetogrsphie lnterpmiation of Subcritical Cracking in a High-Strength Titanium Alloy (IMI-550)", NLR Report MP72025U. National Aerospace Laboratory. Amsterdam. 1972. (35) Hsddsd. R.. CNEA Argentina. unpublished work.
Figure 11. Appearance of steps on pseudocleavage faces formed
on Zircaloy-2 during cracking in mercury (1-pm bar).
r
-.-
1
1
Figure 12. Crack formed in MeOH/HCI solution showing similarity to DHC (5-pm bar).
P i 13. Stress corrosion crack in iodine initiating at an already cracked surface hydride platelet.
cracks are propagating along the basal plane in a hexagonal a-Zr crystal, then the edges of any small steps in the pseudocleavage plane should be produced by slip on the prism planes. These features should have the hexagonal symmetry typical of a-Zr (Figure 11)and not the cubic/ tetragonal symmetry of fractured hydride, and such is found to he the case.% Thus, transgranular SCC in zirconium alloys can be positively distinguished from DHC by the examination of the shape of microfaceting of the pseudocleavage faces. In most instances of transgranular SCC, such examinations have shown that hydride cracking is not implicated in the mechanism;" there remain some ambiguous results (36)Cox. B. "Identifying Failure Mechanisms of Zirconium AUoys hom Rsctqmphie Studtea-. Pmceedings of the Joint NACElASMllMS Symposium on Motollogmphy and Carmsion. 16th Annu Tech. Meeting of Int. Metallag. Sa..Calgary, Aug. 1983. National h o e . of Cor,. Eng., p 153.
Stress Corrosion Crocking of Zirconium Alloys
I.D.
Langmuir, Vol. 3. No.6, 1987 871
QD.
Figure 14. Through-wall crack in a failed fuel element sheath showing mixed intergranular and transgranular mode of crack propagation (50-wm bar).
Figure 16. Crack produced in Cs/5% Cd vapor at 300 OC in a Zircaloy fuel sheath (IO-pm bar).
m , '8
b!
Figum 15. Crack produced in Zircaloy fuel sheath in iodine vapor at 300 OC (10-pm bar). in alcoholie/HCI solutions where hydride cracking may be impli~ated.~*~ In@these instances, although detailed fractography was not done, the circumstantial evidence implicates DHC in the mechanism although the metallography was not able to demonstrate the presence of hydride at the crack tip or along the crack faces. We have examined some of our specimenscracked in methanol/HCl solution ( F i i 12)and find that the fractures bear a close similarity to DHC, although unequivocal microfacet shapes were not identified. Apart from this example, however, SCC in other halogen and halide environments does not appear to involve hydride cracking in the transgranular process. The economicallyimportant PCI cracking of fuel cladding in-reador is a true SCC proceaa and does not involve hydride cracking in its propagation, although the SCC cracks may initiate (Figure 13) at cracked hydrides." However, there has continued to be much w e n t about whether iodine or cesium/cadmium vapoPf" are the important fmion product environments. This disagreement has arisen bemuse of thermodynamic arguments'* that iodine should be combined as CsI inside a fuel pin and that CsI would not cause SCC of Zircaloy,* but the predicted excess of cesium in the fuel might cause cracking. Frac(37) Majumdar. P.: seullv. J. C. C o m . Sei. 1979.19.141.
(39) Galozar. M. A,; S d y ,J. 6.Corros. Sei. 1982,ZZ. 1015. (40) Gmbb, W. T.; Morgan, M. H., 111 'Cadmium Embrittlemeat of Zircalov-2". Roeeeditus of the ANS Tooicd Meetinc on Water Reactor f i e 1 Pbrfohonee, Si: Charles, IL.May 1977; p 2 9 i (41) GNbb, W. T.: Morgan. M. H., 111 "A Survey of Chemical Envimnments for Activity in the Embrittlement of Zircaloy-2".b e e d i n g * 01 the 4th International Conference on Zirconium in the Nuclear Induatry. Stratford-upon-Avon,'U.K..June 1978; ASTM-STF-681. p 145. (42) Olander. D. R. J . Nucl. Mater. 1982, 110, 343. (43) Shann, S. H.; Olander, D. R. J. Nuel. Mater. 19SJ. 113, 234.
tographic evidence supports the argument that iodine c a w the cra~king.'~ PCI cracks in CANDU fuel cladding are usually mixed IG and TG (Figure 14),as are cracks produced in iodine vapor (Figure 15). whereas cracks produced in Cs/Cd vapors are always entirely TG (Figure 16),as is typical for liquid metal embrittlement in zirconium alloys.' This is not unreasonable if removal of metal is necessary for IG cracking since zirconium has a very low solubility in most liquid metals (perhaps because of the thin zirconia film almost invariably present) and no mechanism for removing Zr as a volatile compound is known in Cs/Cd vapor. Thus,since the fractographicevidence is h e a d y in favor of iodine as the important species,we must conclude that thermodynamic equilibrium does not control the processes inside the fuel cladding. Indeed there are other exceptions, since, as far as the inner surface of the cladding is concerned, thermodynamics that at equilibrium the Zucaloy cladding would be oxidized to a solid solution of oxygen in zirconium by reduction of the UOp. While most of the long-lived iodine will have had time to react with cesium, iodine is continually produced in the fission pras a precursor to eesium, and a short Yunditioning" irradiation is known to be necessary if failure of a preirradiated fuel pin is to be obtained in a ramp test once the short-lived iodine species have been allowed to decay. It cannot be determined on present evidence whether this 'conditioning" is necessary to generate and release further short-lived iodine species into the gas gap or to permit (44)Komank, K. L; Silver. M. Thsrmodyoamic Pmprtiss of zirmnium-Oxygen. Titanium-en and Hafnium-Orygcn AUoys'. Pmceeditua ol the 1.AAB.A. Conference on Thermodvmmics of Rcoclor Mote.&b,'Vienna. 1962; p 749. (45) Bauresu. G.; Gerdanian, P. J. P h y . Chem. Solids 1984,45141.
872 Langmuir, Vol. 3, No. 6,1987
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Table 11. Cracking of Zircaloy Fuel Cladding Batches by CsI
experiment 1
2 3 4
5 6
CsI form/ dried slurry
laboratory
(Rb,Cs,Sr,Ba)NOS" (Rb,Cs)2Mo04/Mo20~" powder a
powdern,e powder" b C
d
temp, "C
location
-
laboratory laboratory gammacell gammacell gammacell laboratory u-5 loop u-5 loop u-5 loop u-5 loop u-5 loop
batch no. (MLI-)
350
---
340 400 300 300 300 300
300 300 300 300 300
exposure time, h
788 801 3034 787 787 788 788 788 788 788 788 788 788 788
---
crack free
0 0 0 0 4 3 4 2 6 0 1 0 0 0
144 216 216 72 72 1000 2000 3000 2300 2400 1400 1400 1400 1400
incidence of cracking SCC cracks broken inciDient 2
0 0 3 0 0 0 0
0 0 0 1 0 2
DHC cracks
0 2 2 0 0 3 0 2 0 0 2 2 3 1
0 0 2 0 0 2 4 (2)
0 6
0 0
0 0
Precracked, K, 2.5 MPa m1I2. Precracked and graphite coated. dPrecracked and siloxane coated. e Included two a Not precracked. graphite-coated and one siloxane-coated specimen. Specimens in parentheses showed mixed-mode failure of both DHC and SCC on the same fracture.
Table 111. Effect of Preheating on the Apparent Incubation Time for Cracking at 300 "C" number of specimens environment Cs/Cd
specimen type sealed tube sealed tube sealed tube sealed tube sealed tube sealed tube sealed tube sealed tube sealed tube sealed tube split ring split ring split ring split ring sealed tube sealed tube sealed tube sealed tube sealed tube
preheat of apparatus no
hot unstressed time, min 60 30 15 10 8 7 5 5 3 2 30 30 30 30 60 30 5 3 2
total 2 4 2 1 1 1
5 1 2 3 1 1 1 1 1 1 2
intact 0 0
0 0 0 0 5 0 0 2 0
0 0 1 1 1 0
with incipient cracks 0
0 0 0 1 0 0 0 1 1 1 1 1 0
7
1
0 0 2 5
1
0
1
failed completely 2 4
2 1
0 1 0 1 1 0 0 0 0 0 0 0 0 1 0
" Batch MLI-786 was used. irradiation-induced dissociation of CsI. However, experiments have shown that radiolysis of CsI in-reactor is capable of releasing enough iodine to cause SCC (Table 11). If the chemical environment in contact with the specimen is carefully equilibrated, the application of stress to the system causes the immediate nucleation (Table 111)of the TG cracks.& Furthermore, cracking of Zircaloy in CsI is possible (Figure 17) at 300-350 "C if a rapid transport process for I- through flaws in the oxide on the Zircaloy is provided.17 This transport process might take the form of a low-melting mixture of other fission products, such as the species that give rise to the small globular fission product particles often seen on the inside of the cladding.47g48 It would appear, therefore, that there are several processes that could generate sufficient iodine species at the inner surface of the cladding in chemical forms that can ~~~~~
(46) Haddad, R.; Cox, B. J.Nucl. Mater. 1986, 138, 81. (47) Cubicciotti, D.; Sanecki, J. E.; Strain, R. V.; Greenberg, S.; Neimark, L. A.; Johnson, C. E. J.Nucl. Mater. 1978, 78, 96. (48) Pasupathi, V.; Perrin, J. G.; Roberts, E.; Piker, E. H.; Metcalf, N. R.; Schmidt, G. R.; Rosenbaum, H. S.; Wolff, U. E.; Bell, W. L.; Mattas, R. F.; Sanecki, J. E.; Neimark, L. A. "Determination and Microscopic Study of Incipient Defects in Irradiated Power Reactor Fuel Rods", US. Report, EPRI-NP-812, July 1978.
cause SCC of Zircaloy. Therefore, we have no reason to doubt the fractographic evidence that also points to iodine as the important fission product species leading to PCI failures of fuel cladding. The observations that nucleation times (following the imposition of a stress on a system in chemical equilibrium) can be as short as 1-2 min is also in line with recent observations of incipient cracks in fuel pins given very short power ramps i n - r e a ~ t o r . ~The ~,~~ observation that initiation sites produced by loading an (49) LaVake, 3. C.; Gaertner, M. 'Ramp Test Behavior of High Burnup PWR Fuel Rods", Proceedings of ANS Meeting on L WR Fuel Performance, Orlando, April 1985; Vol. 2, p 6-53, DOE/NE/34130-1. (50) Kjaer-Pederson, N.; Woods, K. N. 'Ramp Tests on Fuel Rods Containing Solid Pellet, Annular Pellet and Sphere-Pac Particle Fuel", ref 49; p 6-35, (51) Cox, B. "Stress-Corrosion Cracking of Zircaloys in Iodine Containing Environments", Proceedings of the Symposium on Zirconium in Nuclear Applications, Portlaud, OR, Aug. 1973; ASTM-STP-551,p 419. (52) Cox, B. Corrosion 1973, 29, 157. (53) Beavers, J. A.; Grieas, J. C. Boyd, W. K. Corrosion 1981,36, 292. (54) Yau, T. L. "SCC of Zirconium and its Alloys in Nitric Acid", Preprint No. 102, Corrosion/81, National Association of Corrosion Engineers, Houston. (55) Nelson, H. G.; Wachob, H. F. "Stress Corrosion Cracking of Zircaloys", Electric Power Research Institute Report EPRI-NP-717. Palo Alto, CA, March 1978, Section 6. (56) Cox, B. Reu. Coat. Corros. 1975, I , 366.
Langmuir 1987,3,873-885 already chemically and thermally equilibrated system are identical with cleavage facets produced during propagation suggests that there is basically no difference between the initiation and the propagation of transgranular cracks. Because (given the right local stress state) no incubation time we arethe mechanism to an adsorption-enhanced process for TG cracking in iodine.*' Thus, there is essentially no mechanistic difference between the SCC process in iodine vapor and the liquid metal embrittlement or metal vapor embrittlement mechanisms. However, because in iodine there is an al-
a73
ternative grain-boundary attack process, the SCC in iodine can be distinguished from MVE in Cs/Cd, and iodine is concluded to be the operative species in causing PCI cracking. Registry No. 12, 7553-56-2;KN03,7757-79-1;NaN03, 763199-4; NaC1,7647-145;Nb2.5, Zrg8,50813-12-2;H2,1333-74-0;HC1,
7647-01-0; Cs, 7440-46-2; Cd, 7440-43-9; &I, 7789-17-5; Brz, 7726-95-6;ccb, 56-23-5; CHC13,67-66-3;FeC13,7705-08-0;cUcl2, 7447-39-4; H ~7439-97-6; , a,7440-67-7; N ~ O10544-72-6; ~, HNO~, 7697-37-2; C U ~ .Nb2.5, ~, Zr9,, 12644-43-8; HzS04, 7664-93-9; methanol,67-56-1;zircaloy-2,11068-94-3;ethylene glycol, 107-21-1.
Corrosion, Layer Formation, and Oxide Reduction of Passive Iron in Alkaline Solution: A Combined Electrochemical and Surface Analytical Study? S. Haupt and H.-H. Strehblow* Institut for Physikalische Chemie, Universitat Dusseldorf, 0-4000 Dusseldorf 1, Federal Republic of Germany Received November 6, 1986. I n Final Form: December 29, 1986 The formation and reduction of passive films on iron in 1 M NaOH are examined by combined electrochemical and surface analytical studies. The electrochemical examination involves the study of POtentiodynamic polarization curves and potentiostatic transients starting with oxide-free surfaces and well-defined preformed oxide films. The pretreatment of the metal surface especially to yield an oxide-free situation was carefully controlled in order to obtain reliable and reproducible results. The use of a hydrodynamically modulated rotating ring disk electrode (HMRRD) permits the determination of soluble corrosion products although the corrosion current densities are extremely small in 1 M NaOH even for unstationary conditions of the passive layer. For this purpose the application of rectangular hydrodynamic pulses was developed as an analytical method to subtract the analytical ring current of the corrosion products from the electronic offset and from the contribution of the impurities of the bulk electrolyte. X-ray photoelectron spectroscopy (XPS) was applied to get information about the composition as, e.g., the Fe(III)/Fe(II) ratio and the hydroxide or oxide content of the passive films. This work examines the dependence of the formation and reduction of the passive layer on potential and time, including very negative potentials and very short oxidation times in the millisecond region. For this purpose a specimen transfer within a closed system was necessary. Iron(II1) oxide grows linearly with the potential and log t, as expected for a barrier-type film. The formation of an inner iron(I1) oxide film and its oxidation to Fe(II1) with increasing potential and with time during potentiatatic transients were followed by X P S and electrochemical methods. The parameters for the growth kinetics of oxide were determined and are in good agreement with literature data at different pH values. The formation of soluble corrosion products has a minor influence and amounts to ca. 5% at maximum. The reduction of a passive layer leads to Fe(OH), and Fe metal for pure iron and to Fe(OH)2only for Fe5%Cr. These studies give a close coincidence for the XPS and electrochemicalresults and provide a deeper insight to the passive behavior by the application of the two disciplines.
Introduction There has been a strong interest in passivity of metals and especially of iron for many years. Numerous works have been devoted to the natural structure and properties of the passive film, although they are still areas of deElectrochemical methods provide good insight to the formation and properties of passive layers, but a better understanding and a reliable interpretation of the results require the application of surface analysis. These methods, working in UHV, provide a variety of valuable information. However, ex situ methods require the extraction of the electrode from ita environment which may introduce unwanted changes after a well-defined prepa-
* Author to whom correspondence should be addressed. 'Presented at the symposium on "Corrosion", 191st National Meeting of the American Chemical Society, New York, NY, April 13-18, 1986.
ration. It has been shown previously that a specimen transfer from the electrolyte via a protecting gas atmosphere into the UHV within a closed system prevents changes which may be introduced by exposure to oxygen." Therefore, the examination of a specimen prepared carefully with regard to the electrochemical behavior of the (1) Passivity of Metals; Frankenthal, R. P., Kruger, J., Eds.; Electrochemical Society: New Jersey, 1978; and references therein. (2) Paasiuity of Metals and Semiconductors; Froment, M., Ed.; Elsevier: 1983, and references therein. (3) Calinski, C.; Collisi, U.; Haupt, S.;Hoppe, H. W.; Strehblow, H.-H.; Speckmann, H.-D. Abstracts of Papers, 191st National Meeting of the American Chemical Society, New York, Ny, American Chemical Society: Washington, DC, 1986; COLL 140; Strehblow, H.-H., COLL 230. (4) Haupt, S.; Collisi, U.; Speckmann, H. D.; Strehblow, H.-H. J . Electroanal. Chem. 1986, 194, 179. (5) Haupt, S.; Calinski, C.; Collisi, U.; Hoppe, H. W.; Speckmann, H. D.; Strehblow, H.-H. Surf. Interface Anal. 1986, 9, 357. (6) Haupt, S.; Calinski, C.; Collisi, U.; Hoppe, H. W.; Speckmann, H. D.; Strehblow, H.-H. DECHEMA-Monogr. 1986, 102, 33.
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