Hydrogen uptake during oxidation of zirconium alloys

Jo~Jrnel o!

A~OYS A~qDCO~FO'~ND~ ELSEVIER

Journal of Alloys and Compounds 256 (1997) 244-246

Letter

Hydrogen uptake during oxidation of zirconium alloys B. C o x Centre fi~r Nuclear Enghleering. Uuiversity of Toronto. Tor~mt,. Ontario M.qS 3E4. (~tltada

Received 26 July 1996: revised 27 September 1996

Abstract By analysing hydrogen uptake data for Zircaloy specimens with the same chemistry, but different thicknesses, it is found that hydrogen uptake rates declrease when TSS in the metal is exceeded Hydride precipitation is usually uniform in the bulk metal, and could only affect hydrogen uptake if the tendency to precipitate in the bt,',k is overcome and precipitation occurs at the point of ingress. This requires very high local ingre,~s rates and a very local uptake proce,;s ( 10 "* of the surface area) to achieve. © 1997 Elsevier Science S.A. Kewvordx: Hydrogen diffusion; Zr alloys; Steam oxygen and hydrogen exposure

Investigators trying to deduce the mechanism by which hydrogen enters zirconium alloys during oxidation in water or steam are faced with a dilemma. While there is much macroscopic evidence on the quantity of hydrogen entering the zirconium alloy as a function of oxidation time, there is no evidence to justify the assumption that this hydrogen migrates through the oxide either through the ZrO 2 lattice or crystallite boundaries in the ZrO 2 film I11. Obviously. the hydrogen has to get through the oxide film somehow, however, there are a number of other heterogeneities in these films that might provide the route. There is equally little informatilon on the chemical state of the hydrogen that enters the metal, although at one time or another all the possible ferms of hydrogen have been suggested. The hydrogen is generated at the oxide/environment interface by the disch~trge of a proton by an electron migrating through the oxide and, hence, represents the cathodic half-cell reactilon [ll. Thus, the hydrogen that enters the metal starts as a hydrogen atom at whatever point on the surface the ca,~hodic half-cell reaction occurs. Because of the highly reducing conditions inside a protective ZrO, film. it is unlikely that it could ever return to a higher oxidation state. Attempts to measure diffusion coefficients in ZrO z films by exposure of preformed (in HzO) oxides to D~O have shown that the D profiles measured by SIMS only follow Fick's law for a few hours and then stop [21. This suggests that D that is being measured in these experiments is merely exchanging for H in some structural and immobtle constituent of the oxide (e,g., O H - groups) and that once o925-8388/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved PII S0925- 8388(96)02852-6

the exchange is complete no further migration occurs. These O H - groups may be associated with porosity in the oxide, and the hydrogen in them is probably not the hydrogen referred to above that is generated by the cathodic half-cell reaction. Exposure to hydrogen atmospheres causes Iocalised breakdowns of the oxide, in the form of through-oxide porosity, that are well established in zirconium and Zircaloys Ill, and are suspected to occur in cold-worked Z r - 2 . 5 % Nb alloys as well. Thus, there is no reason to believe that any measurements of uptake in hydrogen gas [I,3,41 give values for H diffusion coefficients in ZrO 2. Even the results obtained for hydrogen diffusion in ZrO 2 films by implanting hydrogen and measurement with nuclear reaction techniques 15,61 are suspect because of possible irradiation damage during implantation and the unknown form of the hydrogen after implantation which may be different from that generated by the cathodic process. Thus there is no reason to believe that implantation leaves the hydrogen in a chemical form that is similar to that by which it normally migrates into the zirconium, and it is possible that if implanted hydrogen atoms recombined (e.g., at irradiation induced dislocations) and formed small hydrogen bubble nuclei they may behave very differently from those migrating through the oxide during the normal oxidation process. It has been known for a long time that the initial uptake of hydrogen by different zirconium alloys is very dependent on the chemical nature of the intermetallics present [ 1,7-9l. Fig. ! is a composite of early data for zirconium

B. Cox I Jot, nlal of AIh~vs and Compozmds 256 (1997) 244-246

245

ZrH --* io,ooo

-/ oJo

/

/~

.J/

, .J"~'"

ppm U

lua,,,f,.~.,,;~ Fig. I. Hydrogen uptake of various Zr alloys in steam at .%'79°C. I arm.

alloys containing different intermetallic phases. Unalloyed zirconium, unless it is of very high purity usually contains ZGFe particles, the Zr-1.0% Fe alloy will have contained more of the same phase: Zircaioy4 contains tin in solid solution and the iron in the form of a 7_.r(Fe/Crl: phase; Zircaloy-2 in addition to the Zr(Fe/Cr), phase and the tin, contains a second intermetallic phase containing nickel, Zr:(Fe/Ni). It can be seen that it is the presence of this Ni containing phase that causes an increase in the hydrogen uptake during the early period of oxidation. A big increase in uptake was obtained with a binary Zr-1.0% Ni alloy, but hydrogen uptake was so high as to be off-scale on this figure. Thus, there is every reason to believe that hydrogen uptake is a Iocalised phenomenon occurring at anomalous areas of oxide film at the sites of second phase panicles. This may also he true for the Zr-2.5% Nb specimens reported here which were from a batch that contained 600 ppm wt. Fe, some of which may be presem as a second phase after a 580 °C anneal. Confirmation of such an hypothesis comes from the observation that, when a single Zircaloy-2 sheet was coldworked and annealed to give a number of different thickness sheets from 0.125-2.5 mm the oxidation curves in high temperature, atmospheric pressure steam lay within the same scatter band. However, the hydrogen uptake curves in the post-transition region that showed 100% uptake for the thickest sheet, departed from this straight line at progressively thinner oxide thicknesses for the thinner sheets [81. The hydrogen content in the metal at which the curves departed from the straight line at different temperatures equalled the terminal solid solubility (TSS) at each temperature (Fig. 2). The first hydride particles m precipitate when TSS is exceeded during the oxidation of zirconium alloys usually form in the bulk of the specimen rather than at the oxide/metal interface, because of the strong tendency for precipitates to nucleate at preferred sites, which are often at metal grain boundaries [I]. Thus, the solubility of hydrogen in zirconium is ¢ffect,.'ve!y position dependent, and significant supersaturation of the metal matrix is needed to force hydride to

0 he,~ialtou poiutl ro¢ Zircllloy 2,

ss0

,.2

|

500

45oL

4oo

3~,0

~

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I ~ (K"~) Fig. 2. Points of departure of hydrogen uptake from a linear posttransition rate for shccls of Zircaloy-2 of two different thicknesses, compared with the hydrogen solubility at the same temperature.

precipitate at sites other than the preferred ones. It is difficult to see how such precipitates could influence the rate of hydrogen uptake by the specimens. Only if the hydrogen uptake process was localized and the hydride precipitated at the oxide/metal interface at these local sites is it easy to see how the first hydride to precipitate could affect the subsequent hydrogen uptake. Hydrogen diffusion in the Zr matrix at elevated temperatures is so high compared with average ingress rates during corrosion that hydrogen in solid solution in the metal is effectively uniformly distributed, and in the absence of a heat flux, the relatively uniform distribution of preferred sites for hydride nucleation gives rise to a relatively homogeneous precipitation of hydrides once the local TSS at these points is exceeded. Only if the local hydrogen uptake rates are very high, as when corroding in concentrated LiOH solutions, is it possible to form solid hydride layers on the specimen surface during corrosion at high temperatures. Only under such conditions can steep enough hydrogen concentration gradients be maintained at the oxide/metal interface to overcome the tendency for precipitation to occur within the bulk of the metal rather than at the metal/oxide interface. Slow hydrogen diffusion through such a layer might then inhibit the rate of uptake of further hydrogen. The rate of uptake needed to form such surface layers can be plotted from experimental data obtained after tests in concentrated lithium hydroxide

246

B. Co.r / Journal of Alloys and C, mq~oumls 256 (Iq97) 244-246

solution where very high uptake rates were observed (Fig. 3). Then, if wet take the actual hydrogen uptake rates at the points on the oxidation curves where TSS was exceeded in Fig. 2, and compare them with the rate needed to precipitate solid hydride at the o: "de/metal interface from Fig. 3, it is possible i1:ocalculate the fractional area o f the o x i d e / metal interface that was providing active sites for hydrogen absorption. The numbers so obtained are ~ 2 . 5 × 1 0 - * at 400 °C and ~ 1 . 5 × 10 -'~ at 350 °C; the only temperatures for which both types o f data are avai!able well beyond the oxidation rate transition. It is interesting to note that these values are abciut a factor o f 100 lower than the fractional area o f the intermetallics in Zircaloy-2 ( < 1 % ) , and a similar factor below the observed fractional areas o f the porosity in the post-transition oxide films ( ~ i - 2 % ) . Thus, we cannot distinguish between the intermetallics themselves providiag the local " w i n d o w s " for hydrogen uptake, a,:d pores that penetrate to the o x i d e / m e t a l interface in the post-transition oxidation region from these results [1]. Similar changes in hydrogen uptake rate, however, IO"E" f,

Hydride lllyen form

N~

N~

.......

\~,

--

400 t~',,

corrosion temperature.

O UNWOR,MHYDRIDE

~ \

!0 a .

Zircaloy-2Ipc:imenl,corroded inLiOH,withoutthermal ¢yd¢~, to hydrogen contentl above the solubilityat the

1,~

appear to occur for very thin specimens where the TSS is exceeded during pre-transition oxidaticr,, when extensive porosity will not be present. Therefore, we can conclude that only a fraction o f either feature o f the system can be operating as a " w i n d o w " at any moment. In the thin oxide growth region it has been observed that, in water, small pores form at an early stage o f oxidation in the heavily doped oxide formed on the intermetallics, and that such sites ~ersist into the posttransition oxidation region [i,10]. Electron microscope studies [ I l ] have shown the presence o f small particles o f nearly pure iron oxide in the oxide formed around such iron containing intermetallics. The dissolution o f this phase in high temperature water may provide the observed pores, and a possible route for hydrogen into the metal. With the accumulated evidence that hydrogen uptake is a very Iocalised phenomenon, it is surprising that EImoselhi includes little or no discussion o f his own work [2] that disagrees with the hypothesis now espoused [3], that a relatively uniform diffusion process tL, ough the ZrO 2 (although perhaps at the crystallite boundaries) is operating, whether the specimens were exposed to steam, oxygen or hydrogen.

SURFAC~

"NiX

36O 316 300 I ,, . . . . .h . . 1

250 "C

,a

t,,..

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10J/TOK.t) Fig. 3. Critical hydrogen fluxes for surface hydride layer formation (fi'om Chemistry and Materials Division, CRNL, Progress Report, JanuaryMarch 1979. AECL-6538).

I I ] Corrosion of Zirconimn Alloys h~ Nuclear Power Plants° I A E A -

TECDOC-684. International Atomic Energy Agency, Vienna, 1993, pp. 38-56 and 129-147. [2] M.B. Elmoselhi, B.D. Wart and N.S. Mclmyrc. A study of the hydrogen uptake mechanism in zirconium alloys, in A.M. Garde and E.R. Bradley (eds.). Proc. IOth hn. S.xwtp on Zirconium i, the Nuclear InduMr3", ASTM-STP-1245, 1994. pp. 62-79. 13] M.B. EImoselhi, J. AIho's Comp., 231 (1995) 716-721. 141 T. Smith, d. NucL Mater., 18 (1966) 323-336. 15] D. Khatamian and D. Manchester, J. NucL Mater.. 166 (1989) 300-306. 161 D. Khatamian, Z. Phys. Clwmie, 181 (1993) 435-440. 171 B. Cox, J. Electrochem. Sot., 109 (1962) 6-12. 181 B. Cox, Some factors which affect the rate of oxidation and hydrogen absorption of Zircaloy-2 in steam, Report AERE-R434B, U.K.A.E.A. Harwell, 1963. 191 D.W. Freer, D.R. Silvester and J.N. Wanklyn, Corrosion, 21 (1065) 137-142. 1101 B. Cox, M. Ungurelu, Y-M, Wong and C. Wu, Mechanisms of LiOH degradation and H~BO, repair of ZrO., films, in R.R. Bradley and G.P. Sabol (eds.), Proc. 12th htt. Syrup. on Zr in the Nuclear huhtsrry, ASTM-STP-1295, 1996, pp. 114-136. [I I1 H.I. Sheikh, M.A.Sc. Tlwsis, Dept. of Metallurgy and Materials Science, University of Tornnto, 1996.