The effects of hydride precipitates on the mechanical properties of zirconium-hydrogen alloys

Journal of the Less-Common Metals, 40 (1975) 129 - 138 0 Elsevier Sequoia S.A., Lausanne - Printed in the Netherlands

THE EFFECTS OF HYDRIDE PRECIPITATES PROPERTIES OF ZIRCONIUM-HYDROGEN

129

ON THE MECHANICAL ALLOYS

B. J. GILL N~~~o~~i Gas Turbine Est~btishme~t, Pyestock, F~rnboro~gh, (Gt. Britain) J. E. BAILEY

and P. COTTERILL

Department of Metallurgy and Materials Technology, fGt. Britain) (Received

Hunts. GU14 OLS

September

University of Surrey, Guildford

28, 1974)

Summary The mechanical behaviour of zirconium-hydrogen alloys of low hydrogen content, (100 - 6OO.ppm) has been examined at room temperature under tensile-testing conditions at absolute strain rates in the range lo-* lo- ’ s-l. Alloy microstructure and constitution have been modified means of variations in the cooling rate from the CYphase field and it has been shown that the reduction in elongation with increasing hydrogen content depends on the volume fraction of hydride present. Failures at strains of -1% have been related to the presence of continuous networks of precipitate throughout the specimens and, under these conditions, it has been shown that the hydrides 6 and y are equally effective embrittling agents. The strains to failure of zirconium and dilute zirconium-hydrogen alloys have been shown to be markedly reduced in the presence of notches.

Introduction There have been many investigations of the hydrogen embrittlement of zirconium and its alloys since the low-notched impact-energy values of certain specimens were first related to an appreciable hydrogen content, and a precipitate observed metallo~aphically was identified as a hydride [ 1 - 61. The subject is of technological importance because dilute zirconium alloys are used for fuel element cladding and in-core structural components in water-cooled, water-moderated nuclear reactors, and hydride precipitates are formed under operating conditions. The presence of hydride precipitates in zirconium may result in sudden failures at very low strains (<0.5%). Following an extensive review of the subject, Coleman and Hardie [ 73 were able to state that the embrittling effect of hydrogen in zirconium was undoubtedly due to precipitated hydride. Moreover, it has long been appre-

130

ciated that there are two hydride phases which can co-exist with zirconium under conditions of metastable equilibrium below the eutectoid temperature, and there has been much debate as to which of these hydrides is the embrittling agent [ 81. The object of the present study was an examination of the effects of the nature and distribution of the hydride on the room-temperature tensile properties of zirconium’hydrogen alloys of low hydrogen content. Phase relationships in the system The zirconium-hydrogen system consists of two solid solutions, a(c.p.h.) and P(b.c.c.), and three hydride phases, 6 (f.c.c.), T(f.c.t., c/a > 1) and E(f.c.t., c/u < 1). There is a eutectoid reaction at 550 “C, with the eutectoid composition at 33 at.% hydrogen, at which the reaction 0 2 1y+ S occurs. Relationships below the eutectoid temperature, where the phases CY, 6 and y co-exist, have been clarified by the work of Mishra et al. [9] who proposed a peritectoid reaction at about 255 “C in which the reaction 01+ 6 f y occurs. Experiment Alloys of different microstructure and constitution were produced by cooling solid solutions of hydrogen in a-zirconium to room temperature at different rates. Tensile specimens (to BSS 18,1962) were pressed from annealed, reactor-grade zirconium and gaseously hydrided to hydrogen concentrations of 100,250 and 600 ppm in a modified Sieverts’ apparatus. Groups of specimens were cooled to room temperature at rates of - lo- 2, - lo- ’ and lo2 deg C s-l and examined metallographically. Phase analysis was carried out by means of X-ray diffraction in a Nonius Guinier focussing camera with transmitted CL&I X-radiation. The relative concentrations of the hydride phases present in these alloys were estimated from comparisons of the intensities of the X-ray lines 6 I 1I and y1 11. Specimens were tested in tension at room temperature at absolute strain rates in the range lo-* s-l to 10V2 s-l and the fractured regions were examined metallographically and by scanning electron microscopy. Samples were also single-edge notched and tested at a strain rate of lo- 2 s- ‘. Results Metallographic examination showed the presence of increasing volume fractions of precipitate with increasing hydrogen concentration, and a transition from coarse-grain boundary precipi~tion to ~idm~st~tten structures as the cooling rate was increased. Microstructures of the alloy Zr-600 ppm hydrogen cooled at -10e2 and lo2 “C s-l are presented as Figs. 1 and 2. Since the specimens were only 0.5 mm in thickness, the grain

Fig. 1. Alloy Zr-600 ppm hydrogen films of hydride. (x400)

cooled

at - lo-’

Fig. 2. Alloy Zr-600 hydrides. (X400)

ppm hydrogen

cooled

at -lo2

Fig. 3. Alloy precipitates.

ppm hydrogen

cooled

at - lop2

Zr-100

“C s-l “C s-l

“C s-l

showing showing

showing

grain boundary fine, needle-like

discrete

hydride

(X 100)

size was measured to ensure that no single grain boundary traversed the section. The mean linear intercept was 52 pm and the largest observed was 165 pm. In alloys of lower hydrogen concentration, the hydride was present as discrete precipitates, as shown in Fig. 3. X-ray phase analysis showed the presence of 6 and y hydrides in all of the alloys produced and an overall increase in hydride concentration with hydrogen content. However, the alloy constitution was markedly dependent on the cooling rate. In alloys cooled at -lo- 2 “C s-l the hydride was

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0

100

200

300

400

Hydrogen

content

(ppm)

Fig. 4. Failure strain us. hydrogen

0

100

200 Hydrogen

500

content

300

400

content

(pprn1

Fig. 5. Failure strain us. hydrogen

content

600

for alloys cooled

500

at - lOa

“C!s-l.

at -10-l

“C s-l.

600

for alloys cooled

predominantly 6 but the y :6 ratio increased with cooling rate. Thus, in alloys cooled at the highest rate, - lo2 “C s- ‘, the hydride consisted almost excfusively of y. The results of tensile testing, which are summarized in Figs. 4 - 6, show that the values of elongation to failure were markedly affected by the hydrogen concentration and not significantly affected by the cooling rate from the hydriding temperature. The effect of increasing the hydrogen content was to lower the elongation to failure in a progressive manner from about 30% for unhydrided material to about 5% for alloys which contained 600 ppm hydrogen. These data also show tkat the elongation to failure was little affected b:~ changes of strain rate in the range 10e4 s-l to 10e2 s-l. The values of the 0.1% proof stress, and the ultimate and nominal fracture stresses were not affected by variations in the hydrogen content, cooling rate, and strain rate, used in the present work. In the case of the single-edge notched specimens, the only parameter

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0

150

200 Hydrogen

305 content

400

500

600

(ppm)

Fig. 6. Failure strain VS.hydrogen content for alloys cooled at - lo2 “C s-l.

Fig. 7. Necked region of alloy ZrfOO ppm hydrogen cooled at -lo-” voids and precipitate realignment. (Tensile axis $ X100)

“C s-l showing

Fig. 8. Fractured region of alloy Zr-600 ppm hydrogen cooled at - lo2 “C s-l showing fracture path along grain boundary precipitates. (Tensile axis $ X400)

significantly affected by experimental variables was, again, the elongation to failure. The progressive reduction in elongation with increasing hydrogen content is typified by Fig. 4. Comparison of the curves for notched and unnotched specimens in Figs. 4 - 6 shows that the presence of an edge notch lowered the values of elongation progressively as before, but at lower values overall. Thus, for unnotched specimens the elongation fell from -30% to -10% while, for notched specimens, it fell from -10% to - 1%. For unhydrided zirconium, the effect of a notch in reducing the elongation from -30% to - 10% is clearly shown in these figures. The deformed and fractured regions of the specimens were examined

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met~lo~aphically and some of the results are presented as Figs. 7 - 12. Figure ‘7 shows the presence of voids in the matrix in the necked region of the alloy Zr-100 ppm hydrogen cooled at lo-’ “C s-l and that the hydrides in this area were aligned in the direction of testing. In constrast, the structure of the alloy Zr-600 ppm hydrogen, cooled at the same rate, contained a continuous grain-boundary network of hydride throughout the specimen, and Fig. 8 shows that cracking occurred exclusively in these regions. Figure 9, a micrograph of the fractured region of this alloy at X 1000 magnification, shows that not only was cracking confined to the grain-boundary precipitates but that it occurred within them and not at the matrix /precipi~te interface. The hydride network which was so clearly controlling the alloy properties was shown by X-ray analysis to consist almost entirely of 6 phase.

Fig. 9. Fractured region of alloy Zr-600 fracture path within hydrides. (X 1000) Fig. 10. Scanning electron necked region. (X 2500)

micrograph

ppm hydrogen of unhydrided

cooled zirconium

at - 10P2 “C s-l showing

porosity

showing in

The fracture surfaces of unhydrided zirconium, and of alloys which contained 100 and 600 ppm hydrogen, are presented in Figs. 10 - 12. The presence of rupture peaks and porosity in the necked regions of the unhydrided material and of the alloy Zr-100 ppm hydrogen is shown in Figs. 10 and 11. The fractograph of the alloy Zr-600 ppm hydrogen (Fig. 12) shows the occurrence of grain-boundary cracking and may be related to the optical micrographs presented as Figs. 8 and 9. The elongations to failure of the alloys shown in Figs. 10 - 12 are -3O%, -20% and -l%, respectively.

Fig. 11. Scanning electron micrograph of alloy Zr-100 “C s-l showing peaks and porosity in necked region.

ppm hydrogen cooled at -lo-’

Fig. 12. Scanning electron micrograph of alloy Zr600 “C s-l showing fracture at grain boundaries.

ppm hydrogen cooled at - 10e2

Discussion The effects of increases in cooling rate on microstructure were consistent with increases in the degree of supersaturation and the rate of nucleation. Under conditions of slow cooling where the rate of nucleation was relatively low, precipitation occurred on the most favourable sites, namely, the grain boundaries (Fig. 1). Increases in nucleation rate by way of increases in cooling rate from the solid-solution region led to the formation of Widmanstatten structures (Fig. 2), with the assumption of plate- or needle-like shape by the precipitates in order to minimise the strain energy effects [lo] . Growth of hydride precipitates in this way in zirconiumhydrogen alloys was observed previously by Bailey [ 11) . The effects of cooling rate on the constitution of these alloys were demonstrated and it was shown that increases in cooling rate from the solid-solution phase field favoured an increase in the y:6 ratio. Thus, in alloys cooled at - lo- 2 “C s-i the precipitate consisted almost entirely of 6 hydride, whilst in those cooled at -lo2 “C s-l it was almost entirely y hydride. The formation of y hydride from quenched solid solutions is in good agreement with the earlier results of Bailey [ 111 and Bradbrook et al. [ 121, but Mishra et al. [ 91 showed that it could also form by means of a peritectoid reaction, QI+ 6 P y. The mechanism of formation of the y hydride in zirconium alloys of low hydrogen content is still not completely understood. As regards the mechanical properties and the embrittlement effect, the alloys studied in the present work were divided into two classes, one of

136

which showed appreciable elongation to failure and the other which showed failures at N 1% strain. The unhydrided zirconium and the alloys which contained 100 and 250 ppm hydrogen fell into the first class. These alloys were characterised microstructurally by the presence of hydride as discrete particles either at grain boundaries or within the grains, and by the presence of voids and rupture peaks in the fractured regions (Figs. 3, 7, 11). Alloys in the second class which failed at low strains contained a high concentration of hydrogen (600 ppm), and were characterised microstructurd@ by the presence of continuous networks of hydride precipitate either at grain boundaries or within the grains (Fig. 1). The fractured regions of these alloys showed cracking which was confined to the hydride precipitate (Figs. 8,9). The first class of alloys which exhibited ductile behaviour also showed a progressive reduction in elongation with increasing volume fraction of hydride precipitate. The presence of peaks and porosity in the fractured regions of these alloys, together with the alignment of the precipitates in the testing direction, suggested that the matrix had flowed before fracture (Figs. 7, 11). Comparison of the fractured regions of the unhydrided metal with those of the low-hydrogen-content alloys (Figs. 10, ll), showed similar microstructural features and suggested that the matrix controlled the failure in both cases. This also suggested that the progressive reduction in failure strain with increasing hydride precipitate concentration was due to a reduction of the load bearing area of the matrix. The failure of unhydrided zirconium and of the low-hydrogen-con~nt alloys which contained discrete particles of hydride appeared to occur by the nucleation and coalescence of voids and by a matrix shear process. The nucleation of voids at matrix/precipitate interfaces has already been reported [13,14] and a theoretical treatment by Tetelman and McEvily [ 151 showed that elongation could be reduced by an increase in void concentration. An increase in the number of precipitates and, hence, an increase in the number of void nucleation sites can be used to explain the reduction in elongation with increasing hydrogen concentration. The results of the present study support the conclusion of Coleman and Hardie [16] that, in these cases, failure occurs by ductile parting of the matrix between discrete hydride precipitates. The present study showed that the reduction in elongation with increasing hydrogen content was not affected by the initial cooiing rate from the hydriding temperature and was therefore not dependent on alloy constitution. The role of the hydride precipitate in the embrittlement of zirconium has been the subject of many investigations, and although Mueller et al. [S] realized that either phase could be the embrittling agent, they concluded in view of the work of Whitwam [17] that it was probably y hydride. Recently, Barraclough and Beevers [ 181 also suggested that the embrittling agent in alloys of low hydrogen content was y hydride. In the present study, particular attention was given to the mechanical properties of the alloy Zr-600 ppm Hz cooled at -lop2 “C s-l which was

137

shown to have a continuous network of hydride at the grain boundaries. In this alloy the hydride viras shown by X-ray analysis to be composed almost entirely of 6 hydride. The fracture profile of this alloy, which failed at -1% elongation, showed no evidence of necking, rupture or porosity, and this suggested little or no contribution from the matrix. Metallography showed that cracking was confined to the precipitate itself and did not occur at the matrix/ precipitate interface, Fig. 9. Together with electron fractography (Fig. 12), the evidence showed that specimen failure occurred by cleavage of the 6 hydride which was present as an envelope around almost every grain. Alloys cooled at -lo2 “C s-r which failed at strains -1% were also shown to have continuous paths of hydride through their microst~~tures, but in these cases the phase was found to be mainly y hydride. Since there was a precipitate at the grain boundaries in these alloys and the fraction of 6 hydride was found to be very low, it was concluded that y hydride could also precipitate at grain boundaries. This result is substantiated by the work of Bradbrook et al. [12]. It was therefore concluded that brittle behaviour in zirconiumhydrogen alloys under the conditions of the present study was related to the attainment of continuous paths of hydride throughout the specimens, irrespective of whether the hydride was 6 or y. The effects of continuous paths of brittle hydride in zirconiumhydrogen alloys have been discussed by several authors [7,19 - 211. Coleman and Hardie [ 71 stated that once an intergranular film became continuous, the hydride precipitates provided a continuous path for unstable fracture, and more recently, Hardie ]21] summarized the effect of brittle grain-boundary films by stating that if a crack could propagate continuously through a brittle intergranular phase then the ultimate failure wnuld itself be brittle. The observations and conclusions of the present study clearly support the statement. Single-edge notched specimens were tested in order to examine the effects of triaxiality in the stress system and to achieve a condition of localised deformation before general yielding. The overall lowering of elongation values for notched specimens observed in the present study was consistent with the reduction in the magnitude of the shear stresses which are responsible for plastic deformation [ 221. The susceptibility of zirconium with hydrogen contents of up to 600 ppm to low-strain failure in the presence of a notch was demonstrated. Conclusions (1) The progressive reduction in elongation to failure in zirconiumhydrogen alloys is due to an increasing fraction of the load-bearing area being covered by hydride. (2) Brittle failures in these alloys are associated with the presence of

138

continuous hydride networks in the specimens either in or around the grams. (3) In the form of continuous networks, the hydride phases, 6 and yt are equally effective embrittling agents. (4) The presence of discrete hydride precipitates does not appear to alter the mode of failure of the zirconium matrix. (5) The strains to failure in these alloys are markedly reduced by the presence of a triaxial stress system. Acknowledgments This work was carried out in the Dep~tment of ~et~lur~ and laterals Technology, University of Surrey, Guildford, and financial support from both the United Kingdom Atomic Energy Authority and the University is gratefully acknowledged. References 1 H. A. Saller and R. F. Dickerson, US. At. Energy Comm. Rept. Batt. Mem. Inst., 72, 1951. 2 G. T. Muehlenkamp and A. D. Schwope, US. At. Energy Comm. Rept. Batt. Mem. Inst., 845, 1953. 3 F. Forscher, Trans. AIME, 206 (1956) 536. 4 D. G. We&lake, Trans. Amer. Sot. Metals, 56 (1963) 1. 5 C. J. Beevers, Trans. AIME, 233 (1965) 780. 6 C. E. Coleman and D. Hardie, J. Less-Common Metals, 10 (1966) 12. 7 C. E. Coleman and D. Hardie, J. Less-Common Metals, 11 (1966) 168. 8 W. M. Mueller, J. P. Blackledge and G. G. Libowitz, Metal Hydrides, Academic Press, New York, 1968. 9 S. Mishra, K. S. Sivaramakrishnan and M. K. Asundi, J. Nucl. Mater., 45 (1972/73) 235. 10 A. Kelly and R. B. Nicholson, Progr. Mater. Sci., 10 (1963) 149. 11 J. E. Bailey, Acta Met., 11 (1963) 267. 12 J. S. Bradbrook, G. W. Lorimer and N. Ridley, J. Nucl. Mater., 42 (1972) 142. 13 K. E. Puttick, Phil. Mag., 4 (1959) 964. 14 A. H. Cottrell, Fracture, Wiley, London, 1959. 15 A. S. Tetelman and A. J. McEvily Jr., Fracture of Structural Materials, Wiley, Chichester, 1967. 16 C. E. Coleman and D. Hardie, J. Less-Common Metals, 10 (1966) 12. 17 D. Whitwam, Mem. Sci. Rev. Met., 57 (1960) 1. 18 K. G. Barraclough and C. J. Beevers, J. Less-Common Metals, 35 (1974) 177. 19 G. P. Walters, Electrochem. Tech., 4 (1966) 216. 20 C. E. Coleman and D. Hardie, J. Nucl. Mater., 19 (1966) 201. 21 D. Hardie, L’Hydrogene dans les Metaux, 2 (1972) 497. 22 S. P. Timoshenko and J. N. Goodier, Theory of Elasticity, McGraw-Hill, New York, 1970.