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Effect of carbides on the hydriding and oxidation behavior of a Zr-2.5Nb alloy
jeu.rnalof nuclear ELSEVIER
Journal of Nuclear Materials 218 (1994) 18-29
Effect of carbides on the hydriding and oxidation behavior of a Zr-2.5Nb alloy G. Vigna a, L. Lanzani a, G. Domizzi a, S.E. Bermfidez a, j. Ovejero-Garcla a, R. Piotrkowski b a Dto. Materiales, Comisidn Nacional de Energfa Atdmica, Av. del Libertador 8250, 1429 Buenos Aires, Argentina b Dto. Combustibles Nucleares, Comisidn Nacional de Energfa Atdmica, Av. del Libertador 8250, 1429 Buenos Aires, Argentina
Received 18 May 1994; accepted 25 August 1994
Abstract
The surface carbide layer, interior chain-like carbide arrangements and the c a r b i d e / m a t r i x interfaces present in an initial Zr-2.5Nb alloy were studied in hydriding and oxidation experiments. Hydriding performed by cathodic charge in 0.1 M H2SO 4 solution showed that neither the inner carbide chains nor the corresponding interfaces offered short circuit paths for hydrogen diffusion, while the surface carbide layer acted as a protective layer. The oxidation tests performed in autoclaves during 24 h in deaerated steam at 400°C and 105 bar showed weight gains higher than one order of magnitude with respect to the samples free of carbides. The surface carbides promoted the accelerated oxidation of the surrounding matrix and were stripped during the experiments. Evidence was found of the oxidation of the inner carbides, probably due to the transformation strain energy produced by the former oxidation of the surrounding matrix. Moreover, the carbide/matrix interfaces acted as short circuit paths for oxygen diffusion.
I. Introduction
Carbide particles were detected in some samples of the Zr-2.5Nb alloy used in the production of pressure tubes for C A N D U nuclear reactors. These particles can play a deleterious role in the tube performance; in particular, it is necessary to know the influence of the carbide phase during the oxidation of Zr-2.5Nb under reactor operation conditions and with respect to the hydrogen uptake. In previous work [1,2] results were obtained concerning the nature and growth kinetics of carbide particles in a Zr-2.5Nb alloy. Carbides precipitated in solid phases at 555~C (or-phase), 716°C (ct + 13 two-phase structure) and in the liquid phase at 1850°C. The obtained carbides corresponded mainly to the cubic MCI_ ~ phase, M for metal. At 1850°C, carbides similar in size and composition to those observed in pressure tubes were obtained. The experiment consisted of melting the Zr-2.5Nb alloy in
a graphite crucible, where isothermal carbon diffusion in the Z r - N b melt took place. As a result of the diffusion couple (liquid Z r - 2 . 5 N b / s o l i d graphite) a carbide layer 100 I~m thick grew attached to the crucible wall, together with carbide particles whose size ranged between 1 and some tens of Ixm. The smallest particles were arranged in rows or chains determined by the prior 13-phase grains. The detrimental effect of carbides on fracture of cold-worked Zr-2.5Nb pressure tubes was established by Aitchison and Davies [3]. They demonstrated that a variation in toughness for specimens fractured on the radial-axial plane is related to decohesion or fissuring on the transverse-axial plane as a result of microsegregation of C as carbide and of C1. Much earlier, Cox [4] reported a lower corrosion resistance of the material surrounding the carbide particles, in oxidation treatments in steam in an autoclave at 300°C and atmospheric pressure. The present work deals with the effect of the car-
0022-3115/94/$09.50 © 1994 Elsevier Science B.V. All rights reserved SSDI 0022-3115(94)00371-8
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29 bide phase and the carbide/matrix interfaces both in hydriding (in cathodic: charge essays) and oxidation (by autoclaving in deaerated steam at 400°C and 105 bar) of an initial Zr-2.5Nb alloy. The selected oxidation condition corresponds to that which gives place to a protective oxide layer on the pressure tube material.
19
Table 1 Thermal treatments prior to the hydriding and oxidation experiments Sample Treatment 1 Treatment 2 Cooling (°C, h) (°C, h) M1 M2
892+1, 1 900+1, 1
852+1, 3 850+1, 3
Treatment 3 (°C, h)
in furnace 400+1, 24 in furnace -
2. Experimental 2.1. Material preparat;on and characterization The Z r - N b material enriched in carbide phase was obtained by melting a commercial Zr-2.5Nb alloy, provided by Wah Chang, in a graphite crucible at 1850°C under high wlcuum conditions, in a Brew furnace. The obtained Z r - N b - C alloy was characterized by X-ray diffraction (XRD), electron probe microanalysis (EPMA), optical microscopy (OM) and scanning electron microscopy (SEM). The experimental details are displayed in Ref. [2]. A disc-shaped sample with a height of 2.6 mm and a diameter of 14 nun was extracted from the ingot by cutting the crucible through planes perpendicular to its axis. The metallic disc was easily removed from the surrounding graphite, as the bonding at the graphite/carbide interface was extremely loose. The di:~cwas then cut in two equal parts through a plane perpendicular to its faces, giving place to two samples, hereafter named M1 (later hydrided) and M2 (later oxidized), of which shape and features C, D and T are shown in Fig. 1. Both samples were submitted to therm~Ll treatments destined to obtain carbide particles intmersed in a two-phase ct + 13 (Zr, Nb) matrix, similar to that of the pressure tubes. The samples were wrapped in high purity Ta foils and vacuum sealed in quartz tubes with a slight overpressure of high purity Ar. The thermal treatment conditions are described in Table 1. The purpose of treatment 1 was to achieve a homogeneous 13-phase. Treatment 2, designed to obtain the (ct + 13) phase structure,
FACE D :tBIDE LAYER
FACE T
--
FACE C
Fig. 1. Schematic repret;entation of samples M1 and M2. The surface carbide layer is on face C.
was performed lowering the furnace temperature from the temperature corresponding to treatment 1. Treatment 3, to which the M1 sample was submitted, was performed to simulate the aging by the pressure tube during the autoclaving process [5]. The cooling following treatment 2 took place in the furnace, in order to allow the retention of the 13-Zr phase in the two-phase structure. After finishing the heat treatments the samples were characterized by XRD, EPMA, OM and SEM. Z2. Hydriding The M1 sample was cathodically charged under the conditions described in Table 2 after pickling in a 47 ml H20, 47 ml HNO3, 6 ml HF solution. The conditions were selected to obtain a hydride surface layer with a thickness of more than 10 p,m, advantageous for SEM observations. From the equations displayed in Refs. [6,7] we assume the same diffusion coefficient valid for hydrogen through a hydride layer on Zry-2, D = 3 . 2 3 × 10 -5 e x p ( - 4 0 5 8 5 / R T ) cm 2 s -1, where (and hereafter) the activation energy is in J / m o l and R = 8.314 J / K mol. Then, when T = 368 K and the charging time t = 9 h, a hydride layer with a thickness of more than 10 ~m is expected. For the given temperature it is not expected that H could significantly enter the solution in the adjacent metal. The surface hydride formed on face D was characterized by XRD. After this the hydride layer was removed (down to emery paper grit 1000) and then the sample was attacked in a (45 ml lactic acid, 45 ml HNO3, 10 ml HF) solution. In this way, by observing face D, the thickness of the hydride layer grown on face T could be measured, and the modifications in the hydride layer due to the presence of carbides could be evaluated. At the same time, the role of the surface
Table 2 Hydrogen cathodic charge conditions Current dens. (mA/cm 2)
Chem.sol.
Time (h)
Temperature (°C)
150
0.1 M H2SO 4
9
95
G. ~gna et al. //Journal of Nuclear Materials 218 (1994) 18-29
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G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
21
carbide layer (on face C) in the hydriding process could also be studied. 2.3. Oxidation The M2 sample w~ts cut through a plane parallel to face D and the two szmples obtained in this way, M2 A and M2B, were prepared in different ways in order to observe the possible effects of the surface preparation on the ulterior autoclaving process. All faces, except face C, of both samples were mechanically abraded, polished with emery paper down to grit 1000 and then polishing was continued with diamond paste down to 1 p.m. The M2 B sample was then pickled with a (47 ml HNO 3, 6 ml HF, 47 ml H 2 0 ) solution and both samples were washed in bi-distilled water at 80°C. Finally, before the autoclaving process, the samples were weighed and measure,d in order to calculate later the weight gain. The autoclaving was performed in H 2 0 steam at 400°C and 105 bar during 24 h, together with Zr-2.5Nb reference samples with the nominal C content. The oxide layer developed on face D was analyzed by XRD and observed with SEM. In order to investigate the behavior of the inner carbides, in particula,r the oxygen-carbide interaction at different depths, lhe surface oxide layer on face T was removed and successive abrading sectionings parallel to that face were performed. To evaluate the behavior of the oxide layer developed on the Z r - N b ~lloy with high C content, cathodic charge essays were performed on samples extracted from the autoclaved M2 A and M2~ samples under conditions similar to those performed with the nonoxidized samples.
3. Results
3.1. Starting material The disc extracted from the molten little ingot presented an annular layer of (Zr, Nb)Ca_ x, nearly 35 ~m thick, and inner carbides with some I~m in size, arranged in chains, some of them with dendritic appearance. The XRD results of this starting material are presented in Table 3. The peaks were assigned according to data taken from the literature [8]. The major phases were c~-Zr and ZrC. Some peaks corresponding to 13-Zr and one corresponding to to-Zr were also detected. Fig. 2 shows the microstructure of sample M1 after submission to treatraents 1, 2 and 3. We can observe the carbide layer surrounding the lateral face of the
Fig. 2. Microstructure of sample M1 after T1 + T2 + T3 (Table 1). (a) The surface carbide layer on face C can be observed on the left side. (b) Interior carbides merged in the a + 13(Zr, Nb) structure. Darker phase is ct-(Zr,Nb).
half-disc, the interior carbides immersed in the metallic matrix and the (Zr, Nb) a + B-phase structure (the darker plates are a-Zr). This a-phase began to nucleate in the grain boundaries of the B-phase when the temperature was lowered from 900 to 850°C and continued to form and also grew during cooling inside the furnace, giving place to a Widmanst~itten type structure. The B-phase can be seen surrounding the a-plates. Fig. 3 shows the Nb concentration profile with a Nb content higher than 12 wt% in the B-phase; as the B-phase thickness is lower than 1 txm, the quantitative result obtained by EPMA has to be considered as a lower limit for the Nb content. The XRD results obtained after treatment 3 are summarized in Table 4 and verify the presence of the a-Zr, 13-Zr and ZrC phases. The to-Zr phase was not detected in the studied plane of this sample, contrary to the results reported in Ref. [5] for aged material.
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
22
Fig. 4 shows a S E M m i c r o g r a p h of face D, after the surface hydride layer h a d b e e n removed. T h e hydride layer grown o n face T of the half-disc can b e observed. Figs. 5a a n d 5b show t h e protective c h a r a c t e r against hydriding of the carbide layer.
3.3. Oxidized material T h e samples M 2 A a n d M2B, s u b m i t t e d to the previously described oxidation t r e a t m e n t u n d e r w e n t the weight gain indicated in T a b l e 6. This gain is extremely high c o m p a r e d with the 13 m g / d m z m e a n value corres p o n d i n g to various Z r - 2 . 5 N b r e f e r e n c e samples without carbides. However, the thickness of t h e surface oxide layer developed o n the M 2 a a n d M 2 B samples, less t h a n 3 ixm, was similar to t h a t o b t a i n e d o n t h e r e f e r e n c e sample. T a b l e 7 shows the X R D results from face D of the M 2 B sample. C o m p a r i n g with data from t h e literature [8,11], evidence arises of two oxide phases, monoclinic a n d tetragonal, t h e latter stable in general above 1200°C, b u t stabilized in this case by stresses a n d / o r dissolved C atoms [2,12]. Peaks from the ct-Zr a n d 13-Zr p h a s e s also appear. T h e p r e s e n c e of t h e Z r C p h a s e c a n n o t b e g u a r a n t e e d because t h e p e a k s c o r r e s p o n d i n g to the carbide p h a s e e i t h e r d i s a p p e a r e d or t h e i r n u m b e r a n d height r e m a r k a b l y diminished, a n d in a prop o r t i o n t h a t c a n n o t b e explained by t h e a b s o r p t i o n of the X-rays by a n oxide layer some p,m thick. Similar results were o b t a i n e d with sample M2 A.
Fig. 3. Nb concentration profile. An increase of Nb content can be observed on the [~-(Zr, Nb) phase that appears as light plates less than 1 Cm thick.
3.2. Hydrided material T a b l e 5 shows t h e X R D results o b t a i n e d from face D of t h e h y d r i d e d sample t o g e t h e r with literature data [8,9]. T h e o b t a i n e d diagram essentially indicated the p r e s e n c e of a significant quantity of hydride phases. T h e r e were d e t e c t e d e, 5 a n d possibly ~/ phases. T h e p r e s e n c e of ~ could not b e certainly e s t a b l i s h e d because its peaks would overlap with those coming from o t h e r phases. T h e cubic carbide p h a s e was also detected. T h e r e were n o peaks from a c a r b o h y d r i d e p h a s e r e p o r t e d by o t h e r a u t h o r s [10].
Table 4 X-ray diffraction patterns before hydriding Experiment
a-Zr
d (pm)
I (a.u.)
d (pm)
280.1 270.6 257.8 248.5 246.0 234.6 189.9 175.6 165.8 141.3 140.0 136.9 135.2 117.1 111.1 108.4 104.8 103.7
14 13 18 10 100 6 18 8 6 7 3 90 3 9 9 3 3 6
279.8 . 257.3 245.9 . 189.4 . . 139.9 136.8 135.0 116.9 108.4 . 103.6
ft-Zr
I (a.u.)
hkl
33 .
. 32 100
.
. 17 -
. .
. . 3 18 12 3 4
.
. 6
100 . 002 101 . 102 . . 200 112 201 104 203 . 211
ZrC
I (a.u.)
d (pm) . . . 250.7 . . . 177.3 . . . . 112.1 . . .
hkl
.
.
.
. .
. 100
.
.
.
.
. 110 .
. 17
. 200
. . . .
. .
.
.
. . 310 .
.
.
.
15
d (pro)
.
I (a.u.)
hM
100
111
-
-
82
200
62 50
220 311
19 10 -
222 400 -
23
420
. 270.9 . . 234.6 . 165.9 141.5 . . 135.5 117.4 . 104.9 .
. . .
. .
. .
G. ~gna et al. ~Journal o f Nuclear Materials 218 (1994) 18-29
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G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
Fig. 4. SEM micrograph from face D of sample MI. The hydride layer grown adjacent to face T can be observed. I: resin; II: hydride layer.
Fig. 6 shows SEM micrographs corresponding to face D of the M2 A and M2 B samples immediately after the autoclaving treatment. The thin surface oxide layer allows the observation of the underlying phase structure. Thick branches associated with the prior carbide arrangement can be seen. Fig. 6b shows these branches with higher magnification: they consist of cavities left by the carbides detached during the autoclaving, covered by a thick oxide. After oxidation the surface carbide phase was not detected by SEM, XRD and EPMA. Particles detected in the inner regions of some branches like those shown in Fig. 6b resulted in ZrO 2 by EPMA and they probably originated from the spalling of the oxide layers. In the left bottom corner of Fig. 6c a portion of face C of sample M2 B can also be observed. The structure, shown in detail in Figs. 7a and 7b consists of an arrangement of ceils that correspond to the oxide layer grown at the carbide/matrix interface. The shape and the average size of the cells are similar to those corresponding to the carbide grains in the surface layer on face C prior to the oxidation process. These observations prove that the surface carbides were destroyed during the autoclaving treatment. The volume change associated with the oxide growth around the carbide particles would lead to the development of a stress field that would promote the ejection of the particles. Moreover, the cohesion between the oxide and carbide phases could be lower than that between carbide and matrix. Similar results were obtained with the M2 A sample. Fig. 8 shows a SEM micrograph from a plane parallel to the T face, after extraction of the surface oxide layer, putting in evidence the inner structure. The appearing thick and long branches resemble in their distribution the previous carbide chains. The branches correspond in general to an oxide phase as revealed by EPMA and in some cases they present small holes or cracks. Carbon segregation was observed in some of these cracks. We can observe that the branches that touch the surface oxide phase are perfectly connected with it. Some clean carbides can also be observed and they are distinguished because they look like much smaller and smoother particles. Fig. 9 shows with higher magnification the border corresponding to a D face and the enhanced oxidation of the 13-phase with respect to a-phase, as was stated by other authors [12].
Table 6 Weight gain of oxidized samples Fig. 5. Evidence of the protective character of the surface carbide layer against hydriding. (a) The carbide layer has suppressed the formation of the hydride. (b) The accidental partial failure of the carbide layer permitted the formation of the hydride.
Sample
Area (dm2)
Initial weight (mg)
Final weight (mg)
Weight gain (mg/dm2)
M2A M2 B
0.015725 447.01 0.013562 272.36
453.3 275.58
400 237.43
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
I
25
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I
i
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26
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
Fig. 7. SEM micrograph of face C of sample M2 B after autoclaving. (a) Cell type microstructure of the oxide phase that remains after the carbide phase detachment, (b) Micrograph showing the walls of one of the cells shelled and cracked.
Fig. 6. SEM micrographs from faces D of samples M2 A and M2 B after autoclaving. The thickness of the oxide layer, less than 3 ixm, allows to see the underlying structure. The surface carbides have been detached and thick branches of the oxide phase appear in the corresponding locations. Micrographs (a) and (b): sample M2 A mechanically polished before autoclaving; (c): sample M2 B mechanically polished and chemically pickled before autoclaving. It must be e m p h a s i z e d that the same general features were f o u n d at different d e p t h s both in planes parallel to face T or to face D and the same results were o b t a i n e d with b o t h samples.
Fig. 8. SEM micrograph of a plane parallel to face T corresponding to sample M2 a after extraction of the surface oxide. The oxide layer grown on face D can be seen at the upper right border. Thick interior oxide branches are distributed resembling the previous carbide arrangement. Some few clean aligned carbide particles are shown.
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
Fig. 9. Electron image showing the oxide layer grown on face D of the M2A sample. The preferential oxidation of the 13-(Zr,Nb) phase is evidenced.
XRD and SEM studies performed on the samples that were submitted to hydrogen cathodic charge after having been oxidized showed no hydride phase formation.
4. Discussion For the starting material the thickness of the surface carbide layer is in reasonable accord with the extrapolated values from those obtained by other authors in Z r - C reactions performed at higher temperatures [13], where it was supposed that the reaction rate was governed by C diffusion in the carbide layer. In our case, considering that X 2 = kpt,
where
k o = 1.83 exp -352711/RT cm 2 s -1,
with T = 2123 K and t = 1440 s, the carbide thickness X should be 23.5 ~ra, which is not too far from the obtained value. In Ref. [14] it was shown that the growth rate of the inner carbides was also controlled by C diffusion through the external carbide during the melting process. The appearance of some 13-Zr peaks and an to-Zr peak in the molten material would indicate that cooling after melting was sufficiently slow as to allow the partial retention of the 13-phase. A certain quantity of it transformed by aging at lower temperatures to the o~-phase. After treatment 1 a homogeneous 13-(Zr, Nb) phase was obtained ~nd after treatment 2 the 13-phase resulted with a Nb content higher than 12 wt%. This Nb enrichment took place specially during slow cooling after treatment 2; this allowed the 13-phase to be retained down to room temperature, since this is possible only for Nb contents higher than 7 wt% [15]. Other authors [5] pointed out that the I~-phase obtained after treatment similar to that performed in this work contained approximately 20 wt% Nb. For the hydrided material, it could be observed that the carbide layer had a protective character. The fact that no hydride phase was detected in the region adjacent to the carbide layer when it was undamaged,
27
evidenced that hydrogen diffusion should be very slow in that phase. Another explanation could be the formation of a carbohydride phase, that was not detected in our study; this formation could be very sluggish. On the other hand, the thickness of the hydride layer was not modified when it was intersected by a carbide now on chain. Neither the hydrides precipitated preferentially along the carbide/matrix interfaces. These interfaces did not show up as short-circuit paths for hydrogen diffusion. For the oxidized material some results agree with those obtained by Cox dealing with surface carbides in Zr and Zry-2 [4]. It was stated in that paper that a lower oxidation resistance in the vicinity of the surface carbides led to a rapid development of a thick oxide, and to the appearance of stresses due to the high Pilling-Bedworth relation for the Zr oxidation ( r = 1.56). This ends in the appearance of cracks in the oxide and finally the carbides were pressed out. The oxidation of the surface carbides in steam is not expected under the conditions of our experiment, as was shown in Ref. [4]. The surface carbides would play, because of their good electric conductivity, the same role as the surface intermetallic particles in the oxidation of Zry-2 and Zry-4. In that case it was established by Cox [16] that the high temperature corrosion is mainly controlled by the transport of electrons through the intermetallic phases Zr2(Fe, Ni) and Zr(Fe, Cr) 2 which partially or totally short-circuit the oxide layer in these alloys. The electrons that leave the particles reduce the H 2 0 producing 0 2- ions, which diffuse through the oxide crystallite boundaries and react with the metal, generating ZrO e and free electrons. The electrons flow to the surface through the intermetallic particles closing in this way the corrosion circuit. This model predicts that a nonhomogeneous distribution of particles would produce an enhanced local corrosion. In a similar way, carbides in Z r - N b would facilitate the transfer of electrons, promoting in the first stages of the autoclaving the rapid oxidation of the surrounding metallic matrix. As corrosion advances, the appearance of stresses leads to the separation of the surface carbides and in this way the proposed mechanism is not more active. With respect to the interaction of oxygen with the inner carbides, it can be deduced from the described observations that the carbide/matrix interfaces act as short circuit paths for oxygen diffusion. This allowed oxygen, produced after the H 2 0 decomposition at the surface, to enter deeply into the samples and in this way the oxidation of the inner carbides could occur. They were submitted both to high oxygen pressure and to stresses due to the prior oxidation of the surrounding matrix. Zr oxidation is a phase transformation which in-
28
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
volves the formation of a new phase with a molar volume different from that of the matrix from which it precipitates. During this type of reactions a Gibbs free energy loss occurs to produce the strain field around the precipitated phase [17,18]. In our case this happened during the precipitation of the oxide phase at the carbide/matrix interface. The inner carbides submitted to stresses would be in a higher Gibbs free energy state for the oxidation process because the activation energy for oxidation would have been reduced. The oxygen atoms would gradually have replaced the carbon atoms until the oxygen concentration favored the phase transformation from carbide to oxide, pushing the C atoms to the outer regions of the oxide phase where they could segregate or enter the surrounding metal. The high-temperature oxidation of carbides is a particular topic in the wider field of carbon-oxygen combination in interstitial alloys of transition metals. In fact, all refractory mono- and semicarbides of the groups IV and V are capable to dissolve oxygen, initially by C / O substitution but in oxygen excess by higher oxide formation [19]. The weight gain difference between samples M2 A (mechanically polished) and M2 B (chemically attacked) cannot be exclusively assigned to the different surface preparation. Cox [4] observed that samples pickled in hydrofluoric-nitric acid solutions present a lower surface oxidation and explained this result by the fact that the region surrounding the carbides, with lower corrosion resistance, had been eliminated during the pickling. In our case the main reason for the higher weight gain of sample M2 A is due to its greater volume. If the weight gain due to inner oxidation is directly related to the volume of the carbide phase, which is also directly related to the volume of the sample, a factor of 1.6 is expected between both values. The cathodic charge essay performed on the oxidized sample showed that the oxide layer maintained its protective character against hydrogen entrance in spite of the special features occurring due to fhe number and distribution of the surface carbide particles.
lized by dissolved carbon atoms. The surface carbides seem to play the same role in the oxidation of Z r - N b as the surface intermetallic particles in the oxidation of Zry-2 and Zry-4. They facilitate the transfer of electrons, promoting the rapid oxidation of the surrounding matrix in the first stages of autoclaving. On the other hand, a non-homogeneous distribution of particles results in an enhanced local corrosion. The formation of oxide phases around the carbide particles led to the development of stresses, the appearance of cracks and the detachment of the surface carbides. The carbide/matrix interfaces proved to act as short circuit paths for oxygen diffusion; this permitted that oxygen entered deeply into the samples. A high number of the inner carbides were oxidized, since they were submitted both to high oxygen pressure and to stresses due to the prior oxidation of the surrounding matrix. The weight gain obtained after oxidation is more than one order of magnitude higher than that obtained in normal Zr-2.5Nb samples. However, an oxide layer of only 3 ixm grew on the samples. The main reason for the enhanced weight gain was the oxidation occurring both around and at the inner carbides. The surface oxide layer remained protective against the hydrogen cathodic charge in spite of the special features produced by the surface carbides. The inner carbide chains connected with the surface could play a deleterious role in the oxidation of pressure tubes in service, because they could promote an internal oxidation process.
Acknowledgements The present work was performed under Contract No. 6249 with the International Atomic Energy Agency. The authors wish to thank R. Bordoni and M. Villegas from the Water Chemistry Department for performing of the autoclave treatments and for interesting discussions.
References 5. Conclusions After hydriding, e, 8 and possibly 3' hydride phases were obtained. The cubic carbide phase was also detected, but no peaks from a carbohydride phase were observed. The surface carbide showed a protective character against hydriding. The carbide/matrix interfaces did not act as short circuit paths for hydrogen diffusion. After the oxidation treatments the monoclinic and tetragonal oxide phases were detected, the latter stabi-
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G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29 [8] Joint Committee on Powder Diffraction Standards: 5-665; 34-657; 26-1399; 35-784. [9] R.L. Beck, Trans. ASM 55 (1962) 542. [10] R.A. Andrievskii, E.B. Boiko and V.P. Kalinin, Soviet Physics, Crystallography 12 (6) (1968) 930. (Translated from Kristallografiya 12 (6) (1967) 1068). [11] Joint Committee on Powder Diffraction Standards: 371484; 17-923. [12] A.J.G. Maroto, R. Bordoni, M. Villegas, M.A. Blesa, A.M. Olmedo, A. Iglesias and G. Rigotti, Growth and characterization of oxide layers on zirconium alloys, paper presented in the IAEA Technical Committee Meeting on Influence of Water Chemistry on Fuel Cladding Behavior, Prague, October 4-8, 1993.
29
[13] L.M. Adelsberg, L.H. Cadoff and J.M. Tobin, Trans. MetaU. Soc. AIME 236 (1966) 972. [14] R. Piotrkowski and G. Vigna, Mater. Sci. Forum 126-128 (1993) 571. [15] J.P. Abriata and J.C. Bolcich, Bull. Alloy Phase Diagr. 3 (1982) 34. [16] B. Cox, in Advances in Corrosion Science and Technology, eds. M.G. Fontana and R.W. Stahele, vol. 5 (Plenum, New York, 1970) p. 173. [17] R. Balasubramaniam, Acta Metall. Mater. 41 (1993) 3341. [18] J.K. Lee, Y.Y. Earmme, H.I. Aaronson and K.C. Russell, Metall. Trans. l l A (1980) 1837. [19] H.J. Goldschmidt, Interstitial Alloys (Butterworth, London, 1967) p. 564.
Journal of Nuclear Materials 218 (1994) 18-29
Effect of carbides on the hydriding and oxidation behavior of a Zr-2.5Nb alloy G. Vigna a, L. Lanzani a, G. Domizzi a, S.E. Bermfidez a, j. Ovejero-Garcla a, R. Piotrkowski b a Dto. Materiales, Comisidn Nacional de Energfa Atdmica, Av. del Libertador 8250, 1429 Buenos Aires, Argentina b Dto. Combustibles Nucleares, Comisidn Nacional de Energfa Atdmica, Av. del Libertador 8250, 1429 Buenos Aires, Argentina
Received 18 May 1994; accepted 25 August 1994
Abstract
The surface carbide layer, interior chain-like carbide arrangements and the c a r b i d e / m a t r i x interfaces present in an initial Zr-2.5Nb alloy were studied in hydriding and oxidation experiments. Hydriding performed by cathodic charge in 0.1 M H2SO 4 solution showed that neither the inner carbide chains nor the corresponding interfaces offered short circuit paths for hydrogen diffusion, while the surface carbide layer acted as a protective layer. The oxidation tests performed in autoclaves during 24 h in deaerated steam at 400°C and 105 bar showed weight gains higher than one order of magnitude with respect to the samples free of carbides. The surface carbides promoted the accelerated oxidation of the surrounding matrix and were stripped during the experiments. Evidence was found of the oxidation of the inner carbides, probably due to the transformation strain energy produced by the former oxidation of the surrounding matrix. Moreover, the carbide/matrix interfaces acted as short circuit paths for oxygen diffusion.
I. Introduction
Carbide particles were detected in some samples of the Zr-2.5Nb alloy used in the production of pressure tubes for C A N D U nuclear reactors. These particles can play a deleterious role in the tube performance; in particular, it is necessary to know the influence of the carbide phase during the oxidation of Zr-2.5Nb under reactor operation conditions and with respect to the hydrogen uptake. In previous work [1,2] results were obtained concerning the nature and growth kinetics of carbide particles in a Zr-2.5Nb alloy. Carbides precipitated in solid phases at 555~C (or-phase), 716°C (ct + 13 two-phase structure) and in the liquid phase at 1850°C. The obtained carbides corresponded mainly to the cubic MCI_ ~ phase, M for metal. At 1850°C, carbides similar in size and composition to those observed in pressure tubes were obtained. The experiment consisted of melting the Zr-2.5Nb alloy in
a graphite crucible, where isothermal carbon diffusion in the Z r - N b melt took place. As a result of the diffusion couple (liquid Z r - 2 . 5 N b / s o l i d graphite) a carbide layer 100 I~m thick grew attached to the crucible wall, together with carbide particles whose size ranged between 1 and some tens of Ixm. The smallest particles were arranged in rows or chains determined by the prior 13-phase grains. The detrimental effect of carbides on fracture of cold-worked Zr-2.5Nb pressure tubes was established by Aitchison and Davies [3]. They demonstrated that a variation in toughness for specimens fractured on the radial-axial plane is related to decohesion or fissuring on the transverse-axial plane as a result of microsegregation of C as carbide and of C1. Much earlier, Cox [4] reported a lower corrosion resistance of the material surrounding the carbide particles, in oxidation treatments in steam in an autoclave at 300°C and atmospheric pressure. The present work deals with the effect of the car-
0022-3115/94/$09.50 © 1994 Elsevier Science B.V. All rights reserved SSDI 0022-3115(94)00371-8
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29 bide phase and the carbide/matrix interfaces both in hydriding (in cathodic: charge essays) and oxidation (by autoclaving in deaerated steam at 400°C and 105 bar) of an initial Zr-2.5Nb alloy. The selected oxidation condition corresponds to that which gives place to a protective oxide layer on the pressure tube material.
19
Table 1 Thermal treatments prior to the hydriding and oxidation experiments Sample Treatment 1 Treatment 2 Cooling (°C, h) (°C, h) M1 M2
892+1, 1 900+1, 1
852+1, 3 850+1, 3
Treatment 3 (°C, h)
in furnace 400+1, 24 in furnace -
2. Experimental 2.1. Material preparat;on and characterization The Z r - N b material enriched in carbide phase was obtained by melting a commercial Zr-2.5Nb alloy, provided by Wah Chang, in a graphite crucible at 1850°C under high wlcuum conditions, in a Brew furnace. The obtained Z r - N b - C alloy was characterized by X-ray diffraction (XRD), electron probe microanalysis (EPMA), optical microscopy (OM) and scanning electron microscopy (SEM). The experimental details are displayed in Ref. [2]. A disc-shaped sample with a height of 2.6 mm and a diameter of 14 nun was extracted from the ingot by cutting the crucible through planes perpendicular to its axis. The metallic disc was easily removed from the surrounding graphite, as the bonding at the graphite/carbide interface was extremely loose. The di:~cwas then cut in two equal parts through a plane perpendicular to its faces, giving place to two samples, hereafter named M1 (later hydrided) and M2 (later oxidized), of which shape and features C, D and T are shown in Fig. 1. Both samples were submitted to therm~Ll treatments destined to obtain carbide particles intmersed in a two-phase ct + 13 (Zr, Nb) matrix, similar to that of the pressure tubes. The samples were wrapped in high purity Ta foils and vacuum sealed in quartz tubes with a slight overpressure of high purity Ar. The thermal treatment conditions are described in Table 1. The purpose of treatment 1 was to achieve a homogeneous 13-phase. Treatment 2, designed to obtain the (ct + 13) phase structure,
FACE D :tBIDE LAYER
FACE T
--
FACE C
Fig. 1. Schematic repret;entation of samples M1 and M2. The surface carbide layer is on face C.
was performed lowering the furnace temperature from the temperature corresponding to treatment 1. Treatment 3, to which the M1 sample was submitted, was performed to simulate the aging by the pressure tube during the autoclaving process [5]. The cooling following treatment 2 took place in the furnace, in order to allow the retention of the 13-Zr phase in the two-phase structure. After finishing the heat treatments the samples were characterized by XRD, EPMA, OM and SEM. Z2. Hydriding The M1 sample was cathodically charged under the conditions described in Table 2 after pickling in a 47 ml H20, 47 ml HNO3, 6 ml HF solution. The conditions were selected to obtain a hydride surface layer with a thickness of more than 10 p,m, advantageous for SEM observations. From the equations displayed in Refs. [6,7] we assume the same diffusion coefficient valid for hydrogen through a hydride layer on Zry-2, D = 3 . 2 3 × 10 -5 e x p ( - 4 0 5 8 5 / R T ) cm 2 s -1, where (and hereafter) the activation energy is in J / m o l and R = 8.314 J / K mol. Then, when T = 368 K and the charging time t = 9 h, a hydride layer with a thickness of more than 10 ~m is expected. For the given temperature it is not expected that H could significantly enter the solution in the adjacent metal. The surface hydride formed on face D was characterized by XRD. After this the hydride layer was removed (down to emery paper grit 1000) and then the sample was attacked in a (45 ml lactic acid, 45 ml HNO3, 10 ml HF) solution. In this way, by observing face D, the thickness of the hydride layer grown on face T could be measured, and the modifications in the hydride layer due to the presence of carbides could be evaluated. At the same time, the role of the surface
Table 2 Hydrogen cathodic charge conditions Current dens. (mA/cm 2)
Chem.sol.
Time (h)
Temperature (°C)
150
0.1 M H2SO 4
9
95
G. ~gna et al. //Journal of Nuclear Materials 218 (1994) 18-29
20
¢q
¢q
t~
o~
¢q
¢q
¢¢)
t~ ¢q
o)
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
21
carbide layer (on face C) in the hydriding process could also be studied. 2.3. Oxidation The M2 sample w~ts cut through a plane parallel to face D and the two szmples obtained in this way, M2 A and M2B, were prepared in different ways in order to observe the possible effects of the surface preparation on the ulterior autoclaving process. All faces, except face C, of both samples were mechanically abraded, polished with emery paper down to grit 1000 and then polishing was continued with diamond paste down to 1 p.m. The M2 B sample was then pickled with a (47 ml HNO 3, 6 ml HF, 47 ml H 2 0 ) solution and both samples were washed in bi-distilled water at 80°C. Finally, before the autoclaving process, the samples were weighed and measure,d in order to calculate later the weight gain. The autoclaving was performed in H 2 0 steam at 400°C and 105 bar during 24 h, together with Zr-2.5Nb reference samples with the nominal C content. The oxide layer developed on face D was analyzed by XRD and observed with SEM. In order to investigate the behavior of the inner carbides, in particula,r the oxygen-carbide interaction at different depths, lhe surface oxide layer on face T was removed and successive abrading sectionings parallel to that face were performed. To evaluate the behavior of the oxide layer developed on the Z r - N b ~lloy with high C content, cathodic charge essays were performed on samples extracted from the autoclaved M2 A and M2~ samples under conditions similar to those performed with the nonoxidized samples.
3. Results
3.1. Starting material The disc extracted from the molten little ingot presented an annular layer of (Zr, Nb)Ca_ x, nearly 35 ~m thick, and inner carbides with some I~m in size, arranged in chains, some of them with dendritic appearance. The XRD results of this starting material are presented in Table 3. The peaks were assigned according to data taken from the literature [8]. The major phases were c~-Zr and ZrC. Some peaks corresponding to 13-Zr and one corresponding to to-Zr were also detected. Fig. 2 shows the microstructure of sample M1 after submission to treatraents 1, 2 and 3. We can observe the carbide layer surrounding the lateral face of the
Fig. 2. Microstructure of sample M1 after T1 + T2 + T3 (Table 1). (a) The surface carbide layer on face C can be observed on the left side. (b) Interior carbides merged in the a + 13(Zr, Nb) structure. Darker phase is ct-(Zr,Nb).
half-disc, the interior carbides immersed in the metallic matrix and the (Zr, Nb) a + B-phase structure (the darker plates are a-Zr). This a-phase began to nucleate in the grain boundaries of the B-phase when the temperature was lowered from 900 to 850°C and continued to form and also grew during cooling inside the furnace, giving place to a Widmanst~itten type structure. The B-phase can be seen surrounding the a-plates. Fig. 3 shows the Nb concentration profile with a Nb content higher than 12 wt% in the B-phase; as the B-phase thickness is lower than 1 txm, the quantitative result obtained by EPMA has to be considered as a lower limit for the Nb content. The XRD results obtained after treatment 3 are summarized in Table 4 and verify the presence of the a-Zr, 13-Zr and ZrC phases. The to-Zr phase was not detected in the studied plane of this sample, contrary to the results reported in Ref. [5] for aged material.
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
22
Fig. 4 shows a S E M m i c r o g r a p h of face D, after the surface hydride layer h a d b e e n removed. T h e hydride layer grown o n face T of the half-disc can b e observed. Figs. 5a a n d 5b show t h e protective c h a r a c t e r against hydriding of the carbide layer.
3.3. Oxidized material T h e samples M 2 A a n d M2B, s u b m i t t e d to the previously described oxidation t r e a t m e n t u n d e r w e n t the weight gain indicated in T a b l e 6. This gain is extremely high c o m p a r e d with the 13 m g / d m z m e a n value corres p o n d i n g to various Z r - 2 . 5 N b r e f e r e n c e samples without carbides. However, the thickness of t h e surface oxide layer developed o n the M 2 a a n d M 2 B samples, less t h a n 3 ixm, was similar to t h a t o b t a i n e d o n t h e r e f e r e n c e sample. T a b l e 7 shows the X R D results from face D of the M 2 B sample. C o m p a r i n g with data from t h e literature [8,11], evidence arises of two oxide phases, monoclinic a n d tetragonal, t h e latter stable in general above 1200°C, b u t stabilized in this case by stresses a n d / o r dissolved C atoms [2,12]. Peaks from the ct-Zr a n d 13-Zr p h a s e s also appear. T h e p r e s e n c e of t h e Z r C p h a s e c a n n o t b e g u a r a n t e e d because t h e p e a k s c o r r e s p o n d i n g to the carbide p h a s e e i t h e r d i s a p p e a r e d or t h e i r n u m b e r a n d height r e m a r k a b l y diminished, a n d in a prop o r t i o n t h a t c a n n o t b e explained by t h e a b s o r p t i o n of the X-rays by a n oxide layer some p,m thick. Similar results were o b t a i n e d with sample M2 A.
Fig. 3. Nb concentration profile. An increase of Nb content can be observed on the [~-(Zr, Nb) phase that appears as light plates less than 1 Cm thick.
3.2. Hydrided material T a b l e 5 shows t h e X R D results o b t a i n e d from face D of t h e h y d r i d e d sample t o g e t h e r with literature data [8,9]. T h e o b t a i n e d diagram essentially indicated the p r e s e n c e of a significant quantity of hydride phases. T h e r e were d e t e c t e d e, 5 a n d possibly ~/ phases. T h e p r e s e n c e of ~ could not b e certainly e s t a b l i s h e d because its peaks would overlap with those coming from o t h e r phases. T h e cubic carbide p h a s e was also detected. T h e r e were n o peaks from a c a r b o h y d r i d e p h a s e r e p o r t e d by o t h e r a u t h o r s [10].
Table 4 X-ray diffraction patterns before hydriding Experiment
a-Zr
d (pm)
I (a.u.)
d (pm)
280.1 270.6 257.8 248.5 246.0 234.6 189.9 175.6 165.8 141.3 140.0 136.9 135.2 117.1 111.1 108.4 104.8 103.7
14 13 18 10 100 6 18 8 6 7 3 90 3 9 9 3 3 6
279.8 . 257.3 245.9 . 189.4 . . 139.9 136.8 135.0 116.9 108.4 . 103.6
ft-Zr
I (a.u.)
hkl
33 .
. 32 100
.
. 17 -
. .
. . 3 18 12 3 4
.
. 6
100 . 002 101 . 102 . . 200 112 201 104 203 . 211
ZrC
I (a.u.)
d (pm) . . . 250.7 . . . 177.3 . . . . 112.1 . . .
hkl
.
.
.
. .
. 100
.
.
.
.
. 110 .
. 17
. 200
. . . .
. .
.
.
. . 310 .
.
.
.
15
d (pro)
.
I (a.u.)
hM
100
111
-
-
82
200
62 50
220 311
19 10 -
222 400 -
23
420
. 270.9 . . 234.6 . 165.9 141.5 . . 135.5 117.4 . 104.9 .
. . .
. .
. .
G. ~gna et al. ~Journal o f Nuclear Materials 218 (1994) 18-29
,.do6 I
I
I
I
I
~q
¢q
¢',I
23
¢q
¢q
~'~ o ~
~.eq. ¢q
¢q
~
n ¢q
¢q
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¢.q
¢q
t'q ¢q
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~.
tO.
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N
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24
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
Fig. 4. SEM micrograph from face D of sample MI. The hydride layer grown adjacent to face T can be observed. I: resin; II: hydride layer.
Fig. 6 shows SEM micrographs corresponding to face D of the M2 A and M2 B samples immediately after the autoclaving treatment. The thin surface oxide layer allows the observation of the underlying phase structure. Thick branches associated with the prior carbide arrangement can be seen. Fig. 6b shows these branches with higher magnification: they consist of cavities left by the carbides detached during the autoclaving, covered by a thick oxide. After oxidation the surface carbide phase was not detected by SEM, XRD and EPMA. Particles detected in the inner regions of some branches like those shown in Fig. 6b resulted in ZrO 2 by EPMA and they probably originated from the spalling of the oxide layers. In the left bottom corner of Fig. 6c a portion of face C of sample M2 B can also be observed. The structure, shown in detail in Figs. 7a and 7b consists of an arrangement of ceils that correspond to the oxide layer grown at the carbide/matrix interface. The shape and the average size of the cells are similar to those corresponding to the carbide grains in the surface layer on face C prior to the oxidation process. These observations prove that the surface carbides were destroyed during the autoclaving treatment. The volume change associated with the oxide growth around the carbide particles would lead to the development of a stress field that would promote the ejection of the particles. Moreover, the cohesion between the oxide and carbide phases could be lower than that between carbide and matrix. Similar results were obtained with the M2 A sample. Fig. 8 shows a SEM micrograph from a plane parallel to the T face, after extraction of the surface oxide layer, putting in evidence the inner structure. The appearing thick and long branches resemble in their distribution the previous carbide chains. The branches correspond in general to an oxide phase as revealed by EPMA and in some cases they present small holes or cracks. Carbon segregation was observed in some of these cracks. We can observe that the branches that touch the surface oxide phase are perfectly connected with it. Some clean carbides can also be observed and they are distinguished because they look like much smaller and smoother particles. Fig. 9 shows with higher magnification the border corresponding to a D face and the enhanced oxidation of the 13-phase with respect to a-phase, as was stated by other authors [12].
Table 6 Weight gain of oxidized samples Fig. 5. Evidence of the protective character of the surface carbide layer against hydriding. (a) The carbide layer has suppressed the formation of the hydride. (b) The accidental partial failure of the carbide layer permitted the formation of the hydride.
Sample
Area (dm2)
Initial weight (mg)
Final weight (mg)
Weight gain (mg/dm2)
M2A M2 B
0.015725 447.01 0.013562 272.36
453.3 275.58
400 237.43
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
I
25
I t'N
I
i
I
I
m.
,=.i
t',,I,-.~t",lt'xl
0
0
m.~.
It",l
0
0
0
,..~t'xl,~
26
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
Fig. 7. SEM micrograph of face C of sample M2 B after autoclaving. (a) Cell type microstructure of the oxide phase that remains after the carbide phase detachment, (b) Micrograph showing the walls of one of the cells shelled and cracked.
Fig. 6. SEM micrographs from faces D of samples M2 A and M2 B after autoclaving. The thickness of the oxide layer, less than 3 ixm, allows to see the underlying structure. The surface carbides have been detached and thick branches of the oxide phase appear in the corresponding locations. Micrographs (a) and (b): sample M2 A mechanically polished before autoclaving; (c): sample M2 B mechanically polished and chemically pickled before autoclaving. It must be e m p h a s i z e d that the same general features were f o u n d at different d e p t h s both in planes parallel to face T or to face D and the same results were o b t a i n e d with b o t h samples.
Fig. 8. SEM micrograph of a plane parallel to face T corresponding to sample M2 a after extraction of the surface oxide. The oxide layer grown on face D can be seen at the upper right border. Thick interior oxide branches are distributed resembling the previous carbide arrangement. Some few clean aligned carbide particles are shown.
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
Fig. 9. Electron image showing the oxide layer grown on face D of the M2A sample. The preferential oxidation of the 13-(Zr,Nb) phase is evidenced.
XRD and SEM studies performed on the samples that were submitted to hydrogen cathodic charge after having been oxidized showed no hydride phase formation.
4. Discussion For the starting material the thickness of the surface carbide layer is in reasonable accord with the extrapolated values from those obtained by other authors in Z r - C reactions performed at higher temperatures [13], where it was supposed that the reaction rate was governed by C diffusion in the carbide layer. In our case, considering that X 2 = kpt,
where
k o = 1.83 exp -352711/RT cm 2 s -1,
with T = 2123 K and t = 1440 s, the carbide thickness X should be 23.5 ~ra, which is not too far from the obtained value. In Ref. [14] it was shown that the growth rate of the inner carbides was also controlled by C diffusion through the external carbide during the melting process. The appearance of some 13-Zr peaks and an to-Zr peak in the molten material would indicate that cooling after melting was sufficiently slow as to allow the partial retention of the 13-phase. A certain quantity of it transformed by aging at lower temperatures to the o~-phase. After treatment 1 a homogeneous 13-(Zr, Nb) phase was obtained ~nd after treatment 2 the 13-phase resulted with a Nb content higher than 12 wt%. This Nb enrichment took place specially during slow cooling after treatment 2; this allowed the 13-phase to be retained down to room temperature, since this is possible only for Nb contents higher than 7 wt% [15]. Other authors [5] pointed out that the I~-phase obtained after treatment similar to that performed in this work contained approximately 20 wt% Nb. For the hydrided material, it could be observed that the carbide layer had a protective character. The fact that no hydride phase was detected in the region adjacent to the carbide layer when it was undamaged,
27
evidenced that hydrogen diffusion should be very slow in that phase. Another explanation could be the formation of a carbohydride phase, that was not detected in our study; this formation could be very sluggish. On the other hand, the thickness of the hydride layer was not modified when it was intersected by a carbide now on chain. Neither the hydrides precipitated preferentially along the carbide/matrix interfaces. These interfaces did not show up as short-circuit paths for hydrogen diffusion. For the oxidized material some results agree with those obtained by Cox dealing with surface carbides in Zr and Zry-2 [4]. It was stated in that paper that a lower oxidation resistance in the vicinity of the surface carbides led to a rapid development of a thick oxide, and to the appearance of stresses due to the high Pilling-Bedworth relation for the Zr oxidation ( r = 1.56). This ends in the appearance of cracks in the oxide and finally the carbides were pressed out. The oxidation of the surface carbides in steam is not expected under the conditions of our experiment, as was shown in Ref. [4]. The surface carbides would play, because of their good electric conductivity, the same role as the surface intermetallic particles in the oxidation of Zry-2 and Zry-4. In that case it was established by Cox [16] that the high temperature corrosion is mainly controlled by the transport of electrons through the intermetallic phases Zr2(Fe, Ni) and Zr(Fe, Cr) 2 which partially or totally short-circuit the oxide layer in these alloys. The electrons that leave the particles reduce the H 2 0 producing 0 2- ions, which diffuse through the oxide crystallite boundaries and react with the metal, generating ZrO e and free electrons. The electrons flow to the surface through the intermetallic particles closing in this way the corrosion circuit. This model predicts that a nonhomogeneous distribution of particles would produce an enhanced local corrosion. In a similar way, carbides in Z r - N b would facilitate the transfer of electrons, promoting in the first stages of the autoclaving the rapid oxidation of the surrounding metallic matrix. As corrosion advances, the appearance of stresses leads to the separation of the surface carbides and in this way the proposed mechanism is not more active. With respect to the interaction of oxygen with the inner carbides, it can be deduced from the described observations that the carbide/matrix interfaces act as short circuit paths for oxygen diffusion. This allowed oxygen, produced after the H 2 0 decomposition at the surface, to enter deeply into the samples and in this way the oxidation of the inner carbides could occur. They were submitted both to high oxygen pressure and to stresses due to the prior oxidation of the surrounding matrix. Zr oxidation is a phase transformation which in-
28
G. Vigna et al. /Journal of Nuclear Materials 218 (1994) 18-29
volves the formation of a new phase with a molar volume different from that of the matrix from which it precipitates. During this type of reactions a Gibbs free energy loss occurs to produce the strain field around the precipitated phase [17,18]. In our case this happened during the precipitation of the oxide phase at the carbide/matrix interface. The inner carbides submitted to stresses would be in a higher Gibbs free energy state for the oxidation process because the activation energy for oxidation would have been reduced. The oxygen atoms would gradually have replaced the carbon atoms until the oxygen concentration favored the phase transformation from carbide to oxide, pushing the C atoms to the outer regions of the oxide phase where they could segregate or enter the surrounding metal. The high-temperature oxidation of carbides is a particular topic in the wider field of carbon-oxygen combination in interstitial alloys of transition metals. In fact, all refractory mono- and semicarbides of the groups IV and V are capable to dissolve oxygen, initially by C / O substitution but in oxygen excess by higher oxide formation [19]. The weight gain difference between samples M2 A (mechanically polished) and M2 B (chemically attacked) cannot be exclusively assigned to the different surface preparation. Cox [4] observed that samples pickled in hydrofluoric-nitric acid solutions present a lower surface oxidation and explained this result by the fact that the region surrounding the carbides, with lower corrosion resistance, had been eliminated during the pickling. In our case the main reason for the higher weight gain of sample M2 A is due to its greater volume. If the weight gain due to inner oxidation is directly related to the volume of the carbide phase, which is also directly related to the volume of the sample, a factor of 1.6 is expected between both values. The cathodic charge essay performed on the oxidized sample showed that the oxide layer maintained its protective character against hydrogen entrance in spite of the special features occurring due to fhe number and distribution of the surface carbide particles.
lized by dissolved carbon atoms. The surface carbides seem to play the same role in the oxidation of Z r - N b as the surface intermetallic particles in the oxidation of Zry-2 and Zry-4. They facilitate the transfer of electrons, promoting the rapid oxidation of the surrounding matrix in the first stages of autoclaving. On the other hand, a non-homogeneous distribution of particles results in an enhanced local corrosion. The formation of oxide phases around the carbide particles led to the development of stresses, the appearance of cracks and the detachment of the surface carbides. The carbide/matrix interfaces proved to act as short circuit paths for oxygen diffusion; this permitted that oxygen entered deeply into the samples. A high number of the inner carbides were oxidized, since they were submitted both to high oxygen pressure and to stresses due to the prior oxidation of the surrounding matrix. The weight gain obtained after oxidation is more than one order of magnitude higher than that obtained in normal Zr-2.5Nb samples. However, an oxide layer of only 3 ixm grew on the samples. The main reason for the enhanced weight gain was the oxidation occurring both around and at the inner carbides. The surface oxide layer remained protective against the hydrogen cathodic charge in spite of the special features produced by the surface carbides. The inner carbide chains connected with the surface could play a deleterious role in the oxidation of pressure tubes in service, because they could promote an internal oxidation process.
Acknowledgements The present work was performed under Contract No. 6249 with the International Atomic Energy Agency. The authors wish to thank R. Bordoni and M. Villegas from the Water Chemistry Department for performing of the autoclave treatments and for interesting discussions.
References 5. Conclusions After hydriding, e, 8 and possibly 3' hydride phases were obtained. The cubic carbide phase was also detected, but no peaks from a carbohydride phase were observed. The surface carbide showed a protective character against hydriding. The carbide/matrix interfaces did not act as short circuit paths for hydrogen diffusion. After the oxidation treatments the monoclinic and tetragonal oxide phases were detected, the latter stabi-
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