STRUCTURAL INTEGRITY CHANGES IN THE Zr-Nb ALLOY DURING HYDRIDE PHASE TRANSFORMATION UNDER THE THERMAL SHOCK. 1. Introduction

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1 STRUCTURAL INTEGRITY CHANGES IN THE Zr-Nb ALLOY DURING HYDRIDE PHASE TRANSFORMATION UNDER THE THERMAL SHOCK R.Levinskas 1, A.Grybenas 1, V.Makarevicius 1, José Rodríguez 2, Diego Martínez 2 1 Lithuanian Energy Institute, Breslaujos 3, LT-3035 Kaunas, Lithuania 2 Plataforma Solar de Almeria, CIEMAT, PO Box 22, Tabernas (Almeria), Spain Abstract: Structural changes and properties of zirconium alloy subjected to the thermal shock in the solar furnace were investigated. Specimens were prepared from unirradiated cold-worked RBMK Zr-2.5Nb fuel channel pressure tube material after adding hydrogen to the sections of the tubes up to 150 ppm. Tests were done at temperatures up to 1460 o C and heating rates o C per second. After thermal treatment changes were estimated by optical and electron microscopy as well as tensile and microhardness testing. Treatment in the solar furnace revealed, that the influence of hydrogen under thermal shock conditions results in structural degradation and loss of ductility of an alloy. Keywords: zirconium, hydrogen, thermal shock, solar furnace 1. Introduction Zirconium alloys are used as a constructional material for manufacturing of cladding of fuel assemblies and fuel channels of NPP. Zr alloys can pick up hydrogen during in-service corrosion. Hydrogen has very limited solubility in zirconium alloys; when the terminal solid solubility (TSS) (Pan et al, 1996) is exceeded in a component such as a pressure tube that is highly stressed for long periods of time, hydride cracking failures may occur (Grybenas et al, 2000). Hydride induced cracking has been recognized as the potential cause of failure of pressure tubes in nuclear reactors (Cheadle et al, 1998). With increase of temperature hydrides dissolves and atomic hydrogen easily redistributes in the alloy (Sagat et al, 2000). At cooling, when hydrogen concentration exceeds terminal solid solubility, hydrides preferentially precipitate on a certain crystallographic planes of Zr α-phase. Consequently, their orientation depends on orientation of this phase. The precipitation of zirconium hydride embrittles Zr alloy and leads to the cracking failures of the pressure tube. The mechanism of Zr interaction with hydrogen is complex and insufficiently investigated. Influence of both hydrogen and hydrides to the properties and structure of Zr-Nb alloy can be different depending on a range of factors. Usually zirconium matrix properties remains unchanged during repetitive hydride precipitation and dissolution. After eliminating hydrogen properties of the alloy recovers. When hydrogen-containing Zr-Nb alloy is heated over 1000 o C and rapidly cooled, continuous structure of Zr α-phase forms in a shape of sharp pellets or needles. Grains consisting of α-phase forms as cooling rate decreases. Hydrides formed on cooling are distributed between the pellets of α-phase and on the grain boundaries, their size increases with slower cooling rate. Generally hydrides are in a shape of thin pellets of various sizes which dislocation depends on orientation of Zr α-phase. (Sung Soo Kim et al, 1999) In this work an effort was made to investigate influence of relatively low amounts of hydrogen (150 ppm) to the structure of Zr-2.5Nb alloy subjected to the thermal shock up to 1000 o C and over, and also to explore structural as well as material property changes, when temperature increase rate is larger than hydride dissolution rate.

2 2. Experimental Specimens were prepared from unirradiated cold-worked RBMK Zr-2.5Nb fuel channel pressure tube material manufactured using the heat treatment technology TMO-1. Hydrogen was added to the sections of the fuel channel tubes up to 150 ppm by electrolytically depositing a layer of hydride on the surface of the pressure tube material followed by dissolving hydride layer by diffusion annealing at elevated temperature (Makarevicius et al, 2001). From hydrided material two types of rectangular specimens were made as shown in Fig. 1. Type I specimen thickness was 4 mm. Some of the tests were done again with type II specimens of a half thickness t B R6 3x2,5 Fig 1. Test specimen drawing Fig 2. Tensile test specimen Specimens to be tested were placed horizontally on the mobile test table at the centre of the chamber in which argon at 1 bar pressure was flowing. Chromel-alumel or platinum-rhodium thermocouples (for temperatures above 1200 o C) were mounted in the drilled hole at the centre of the bottom end of the specimen. The temperatures for each treatment were recorded every 1 sec. All tests have been done at two different heat increase rates with shutters opened 30 and 60 %, maximum test temperatures were o C. The end of experiment was determined by reaching required temperature. After the end of the treatment, the samples were left in the chamber until temperature reached 200 o C and then were cooled in the ambient air (9 o C /min). Table 1. Testing conditions in the solar furnace Specimen type Specimen thickness, mm Max. Shutter Aperture, % Average heating rate, o C/s I II After thermal treatment in the solar furnace from the axial orientation of the pressure tube test samples (type I), compact tension test specimens (Fig. 2) were cut. Tensile tests were performed at the ambient temperature at strain rate mm/s. Strain was defind as changes (%) of the value B (6 mm) compared to the original length. Microhardness (Vickers, using a 0.49N load,) measurements were done on samples prepared for metallographical analysis. The samples were polished and etched in the mixture of H 2 SO 4 : HNO 3 : HF: H 2 O (3: 3: 1: 3) for 10 s. Microstructure was studied by both optical microscopy and SEM (JEOL JSM-5600) 3. Results and discussion Experiments were performed under thermal shock conditions in the solar furnace in a certain sequence according to the temperatures, corresponding the following phase transformations: α-zr +β-nb +Zr[H] α-zr +β-zr β-zr

3 After thermal treatment, obtained specimens were used for investigation of microstructural and mechanical properties. A number of metallographic specimens were made, which enabled to find out microstructural changes depending on the thermal shock temperature. In this paper are presented some characteristic samples of microstructure, which in the best way illustrate microstructural changes after thermal treatment. In Fig. 3 is presented an example of microstructure of original Zr-2.5Nb material, containing 150 ppm hydrogen. Hydrides are formed on slow cooling and are arranged in lines along the tube texture. Alongside hydrides α-zr pellets are visible (SEM photograph). Fig. 3. Microstructure of the original Zr-2.5Nb material, containing 150 ppm hydrogen After thermal shock up to 586 o C, close to eutectoid transformation temperature, microstructure, hydride size and distribution remains almost unchanged. Obviously, hydrides cannot dissolve so far; otherwise, under a more rapid cooling rate than that was during hydride formation in the original specimens, it must result in formation of even smaller hydrides. Above eutectoid transformation temperature ( o C), in a α-zr+β-zr region, primarily only small hydride redistribution as well as decrease in size occurs and starts formation of 5-15 µm hydride areas, without apparent boundaries. However, because only part of hydrides dissolves or due to structural memory, previous hydride allocation remains unchanged and no significant changes in α-zr phase occurs (Fig. 4a). In the upper region of α-zr+β-zr phase ( o C) as a result of recrystalization process orientation of α-phase disappears and grain formation begins. In this case structure augmentation of hydride origin areas can be observed, part of zirconium hydride distributes around the grains. Consequently, integrity of grain boundaries is lost. Raising thermal shock temperature up to o C, transformation to β-zr phase region occurs and after recrystallization on cooling, structural degradation results in broken integrity of grain structure and boundaries. The fragments of such structure are shown in Fig. 4a, b. Inside the grains and on the grain boundaries flaws obviously can be seen. Some of them are similar to the microcracks, formed in a place of hydride dislocations. As the testing revealed, namely these structural regions have significant microhardness increase. Moreover, microhardness testing indicated that most impact of thermal shock is on the surface layer of the specimens. Especially noticeable structural changes occurred in type II specimens, in which temperature increase rate was the largest. Changes are different across the specimen; large scatter of microhardness values decreases moving away from the thermal shock affected surface. As can be seen in the Fig. 5a, in the samples without hydrogen scatter of microhardness values is less, as well as increase of microhardness near the surface.

4 a b c Fig. 4. Fragments of microstructure after the thermal shock, hydrogen content 150 ppm.: a o C (type I); b o C (type II); c o C (type II). As a comparison presented is microstructure of the original alloy without hydrogen, after the same thermal treatment (Fig. 5b). In the furnace on cooling typical Zr-Nb structure forms, with characteristics large grains. Grain boundaries are uniform. Inside grains are located platelets of sharp needle-shaped α-zr phase and large α-quasimonotectoides. Tensile strength and strain behavior was evaluated depending on the thermal shock temperature. Changes of ultimate tensile stress and strain are shown in Fig. 6. It can be observed that structural changes of the Zr alloy have considerable impact on stress and strain properties. Compared to the original alloy, increase of tensile strength and significant decrease of strain in hydrided alloy indicates that ultimate strain essentially decreases. Microhardness, HV (0.49N) Depth from surface, µm a Fig. 5. a - Changes in microhardness at different depth (1223 O C, type II, [H] = 150 ppm); b - microstructure of Zr-2.5Nb alloy after the thermal shock (1200 o C type I, without hydrogen) In some works (Douglass, 1972; Klepfer et al, 1960; Forscher, 1956) similar investigations devoted to the influence of strain rate and hydrogen content on zirconium plastic properties after different thermal treatment operations. It is shown that under low strain rate, when ductile properties diminish due to hydride deformation, cavities are forming and developing. Formation of cavities is observed also under a relatively large strain rate, when deformation occurs according to bicrystal formation mechanism. Cavities are forming as a result of low bond strength between matrix and hydride. Obtained results show that structural transformations under the thermal shock conditions in hydrogen containing Zr alloy also determines fracture mechanism changes. In the fractured tensile specimens near the strained fracture zone formation of intrinsic cavities can be observed. Such fracture mechanism is characteristic for hydrogen-containing zirconium. Cavities b

5 are forming because of low bond strength between hydride and matrix. Hydrides disorder deformation continuity and as a result microcracks can form, subsequent deformation induces cavities formation. Ultimate tensile stress, MPa Temperature, o C Strain, % Temperature, o C [H]=150 ppm Without hydrogen Fig. 6. Ultimate tensile stress and strain as a function of temperature As temperature increases up to transition to α-zr+β-zr and β-zr phase region, fracture mode changes from ductile to brittle. In this case near the fracture zone there are no cavities and ductility decreases. If in the case of ductile fracture surface was fibrous, in this case it has grainy structure. Besides, microscopic observations indicated, that rupture surface is not typical transgrannular fracture, it pass through grains as well as through the grain boundaries. Such fracture mode corresponds to the observed structural changes. Flaws inside grains or on the grain boundaries caused by hydride formation, which shape and allocation are similar to microcracks, actually results in reduction of zirconium matrix ductility. Such flaws, under certain stress level initially do not cause catastrophic failure, but prevents migrating of dislocations (Ericson et al, 1964). Under such structure matrix couldn t distort to the extent that plastic deformation of microcracks could induce formation of cavities. As a result ductility is lost along with temporal increase of strength. To estimate more accurately influence of structural changes on the loss of ductility, microhardness changes were analyzed. For that purpose average microhardness values were determined farther from specimen edges and hydride locations. Average microhardness values before and after testing of hydrided non-hydrided specimens were similar (HV=186). If microhardness of hydrided specimens after the thermal shock only slightly changes, a microhardness hydrided specimen significantly increases as well as the scatter of microhardness values. The influence of the thermal shock temperature on the average microhardness values is demonstrated on the Fig. 7a. It shows that the alloy structure after thermal treatment becomes more heterogeneous. Microphotographs of the surface after prolonged etching also reveal that surface areas of the hydrided alloy differs in chemical reactivity compared with non-hydrided material (Fig. 8). Similar results can be observed from the SEM photographs (Fig. 7b). It is not fully clear what is the reason for such changes. It can be possibly related with dispersive hydride distribution among α-zr phase or due to hydrogen interaction with Zr-2.5Nb structural constituents.

6 Microhardness, HV (0.49N) [H]=150 ppm a Temperature, o C Without hydrogen Fig. 7. a - relationship between microhardness and annealing temperature; b - flaws of microstructure in Zr alloy after thermal shock up to 1145 o C (type II). b Fig. 8. Microstructure of zirconium alloy after prolonged etching: a-without hydrogen (1223 O C); b-[h] = 150 ppm (1200 O C). Hydride dissolution time calculated according to the following expression of diffusion model [Kearns, 1968] t= 1.65h 2 exp(24900/rt), where h is the hydride thickness (cm), demonstrate that during the thermal shock for both specimen types hydrides cannot dissolve up to eutectoid temperature. Under certain conditions, when thermal shock temperature goes above terminal hydride solubility limit temperature, ordinary state can be retained only under the some equilibrium internal pressure (LaGrange et al, 1959; Libowitz, 1962), which increases with temperature until all hydrides dissolve. Thus can be explained formation of hydride zones, which, as referred (Briant et al, 1983) can arise as a consequence of local stresses. Formation of such zones after thermal treatment can be seen in Fig. 4a, when texture and previous hydride arrangement remains unchanged. It is known that in some metals, which contains dissolved hydrogen, under elevated temperatures and internal pressure heterogeneous reactions can occur (Briant et al, 1983) and as a consequence hydrogen compounds can form and alloy brittleness increases. It was found that during thermal shock hydride layer on the specimen surface diffuses instantly into the alloy increasing hydrogen concentration (Fig. 9). Microhardness profile revealed higher hardness values up to 0.5 mm from the surface.

7 a Microhardness,HV (0.49N) Depth from surface, µm Fig. 9. a - increase of hydrogen concentration in the surface layer; b - changes in microhardness at different depth (586 O C, type II) Testing in the solar furnace revealed, that the influence of hydrogen under thermal shock conditions on the zirconium alloy in some cases is different compared with the slow rate heat treatment. In order to find out the reasons of increased brittleness after thermal shock and determine factors causing failure mechanism changes is necessary to carry out additional experiments. 4. Conclusions Testing in the solar furnace have demonstrated that heat treatment under thermal shock conditions results in decrease of ductility of hydrided Zr-2.5Nb alloy. Below eutectoid temperature structural changes and hydride zone formation can be observed. In addition, hydride layer on the specimen surface diffuses instantly into the alloy increasing hydrogen concentration. More investigations are required to explore conditions and temperature range for this critical process. Above eutectoid temperature hydrogen has time to dissolve in β-zr phase of the alloy and after recristallization on cooling structural degradation takes place, which results in loss of ductility and changes in the fracture mode. It is related not only with hydride dislocation but also with structural changes of the alloy, the reasons however are not obvious. Acknowledgements This work was supported by EU Improving Human Potential DGXII programme. We are grateful to the staff of Plataforma Solar de Almeria assistance and support. References Briant C. L. Banerji. Treatise on materials science and technology: Embrittlement of engineering alloys, 1983, v. 25. Cheadle, B.A., Coleman, C.E., Rodgers, D.K., Davies, P.H., Chow, C.K. and Griffits, M. Examination of Core Components Removed from CANDU Reactors International Conference on CANDU Maintenance, Canadian Nuclear Society, November, p. Douglass D.L. The metallurgy of zirconium; International Atomic Energy Agency-Viena, Ericson W. H., Hardie D. The role of hydride precipitate in the fracture of zirconium and its alloys.- J. Inst. Metals, 1964, v. 93, p Forscher F. Strain-induced porosity and hydrogen embrittlement in zirconium.-trans AIME, 1956, v. 206, p. 536 b

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