Chemical compatibility of uranium based metallic fuels with T91 cladding

Nuclear Engineering and Design 250 (2012) 267–276

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Chemical compatibility of uranium based metallic fuels with T91 cladding Santu Kaity a , T.R.G. Kutty a,∗ , Renu Agarwal b , Arijit Laik c , Arun Kumar a a

Radiometallurgy Division, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India Product Development Division, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India c Materials Science Division, Bhabha Atomic Research Centre, Trombay, Mumbai 400 085, India b

h i g h l i g h t s  Performance of Zr as FCCI barrier layer was evaluated by diffusion experiments.  Rate constant for reaction at U/Zr interface was 2.07 × 10−8 m s−1/2 at 973 K.  Rate constant for reaction at Zr/T91 interface was 1.95 × 10−8 m s−1/2 at 973 K.  Activation energy for reaction at Zr/T91 interface was found to be 54.7 kJ mole−1 .  Interdiffusion between U–6Zr and T91 resulted in formation of three layers.

a r t i c l e

i n f o

Article history: Received 26 December 2011 Received in revised form 26 April 2012 Accepted 27 April 2012

a b s t r a c t Studies related to development of fast reactor fuels based on ternary U–Pu–Zr and binary U–Pu alloys has been initiated in India for building a data base on thermo-physical and thermodynamic properties, fuel-clad compatibility etc. which are very useful to the fuel-designer to optimize the design feature and to predict the in-reactor fuel behaviour. Fuel-clad chemical compatibility is considered as one of the major concerns for metallic fuels. In the present investigation, the performance of Zr as fuel-clad chemical interaction (FCCI) barrier layer between U and T91 was evaluated by diffusion couple experiments. The growth kinetics of reaction layers at U/Zr and Zr/T91 interfaces were established. The growth kinetics of the reaction zone at both the U/Zr and Zr/T91 interfaces were determined at 973 K from the plot of log (width) versus log (time). The value of reaction index n was found to be around 2 at both the U/Zr and Zr/T91 interfaces. The reaction constant (k) for the growth of reaction layer at the U/Zr interface was determined to be 2.07 × 10−8 m s−1/2 at 973 K. Similarly, the rate constant at the Zr/T91 interface was found to be 1.95 × 10−8 m s−1/2 at 973 K. The activation energy Q for the reaction at the Zr/T91 interface was determined and was found to be 54.7 kJ mole−1 . The fuel-clad chemical compatibility between U–6Zr alloy and T91 steel was also investigated in the present study by diffusion couple experiments. The interdiffusion between U–6Zr and T91 at 973 K resulted in the formation of three different layers at the interface. The mechanism of formation of these layers was analysed in detail. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Metallic fuel is considered for future fast breeder reactors (FBR) due to its high breeding potential, high thermal conductivity, high fissile and fertile atom densities, low doubling time and ease of fabrication compared to other ceramic fuels (Burkes et al., 2009; Crawford et al., 2007; Hofman and Walters, 1994; Kittel et al., 1993; Lahm et al., 1993; Nevitt, 1989; Riyas and Mohanakrishnan, 2008; Walters, 1999). One of the key features of metallic fast reactor fuels is improved safety through lower stored energy and enhancements in the negative reactivity feedback during off-normal events. As a consequence of these characteristics, it is possible to avoid core damage even for circumstances where

∗ Corresponding author. Tel.: +91 22 2559 5361; fax: +91 22 2550 5151. E-mail address: [email protected] (T.R.G. Kutty). 0029-5493/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.nucengdes.2012.04.028

the automatic scram system fails to operate or heat removal systems are severely degraded (Kim et al., 2006; Kramer et al., 1992; Kwon et al., 2003; Lahm et al., 1993; Lehto et al., 1988; Nam and Hwang, 1998; Peppler and Will, 1988; Tsai, 1995). However, a few shortcomings of metallic fuels such as, low solidus temperature, high swelling rate and susceptibility to chemical and mechanical interaction with cladding materials prevent it from achieving its full potential. The fuel-clad chemical interaction (FCCI), which is considered as a potential problem area in the application of the metallic fuel in liquid-metal cooled fast reactors, can be avoided in two ways i.e. either by alloying of uranium or by using a barrier layer between fuel and clad. Zr is known to increase the solidus temperature and to improve the chemical compatibility between fuel and steel cladding material by suppressing the interdiffusion between the fuel and cladding components (Hofman and Walters, 1994). Hence, Zr was chosen as an alloying element as well as a barrier layer.

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Fig. 1. Cross-sections of conceptual fuel pin design for metallic fuels; (a) U–Pu–Zr ternary and (b) U–Pu binary alloy (not to scale).

Primarily, two design concepts have been proposed for the metallic fuel development programme for FBR’s in India (Devan et al., 2011). Two concepts are being explored: (a) mechanically bonded pin with U–15Pu alloy as fuel, and (b) sodium bonded pin with U–15Pu–6Zr alloy as fuel (composition in wt.%). In the case of mechanically bonded binary alloy fuel, Zr liner is proposed to be used inside the clad to prevent fuel–clad chemical interaction. T91 grade steel has been used as the cladding material in these designs. The cross-sections of the conceptual fuel pin designs of the metallic fuel, with and without Zr barrier, are shown in Fig. 1. Uranium is the proposed blanket materials for mechanically bonded binary U–Pu alloy fuel, whereas, U–6 wt.% Zr is the proposed blanket materials for that of sodium bonded ternary U–Pu–Zr fuel. U–6 wt.% Zr (hereafter, referred to as U–6Zr) is also a subsystem of U–15 wt.% Pu–6 wt.% Zr alloy. T91 grade steel is a 9Cr–1MoVNb type steel, having ferriticmartensitic structure. In contrast to austenitic alloy steels, the ferritic steels have the advantage of a higher resistance to void swelling. However, a rise of the ductile-brittle transition temperature (DBTT) due to fast neutron irradiation is a concern in these types of steels. The above steel exhibits high mechanical strength and combines low thermal expansion with high heat conductivity (Kaity et al., 2010a; Klueh and Nelson, 2007). The integrity of the cladding is of fundamental concern for designers since it provides the primary barrier to the release of radionuclides (Keiser and Dayananda, 1993, 1994; Keiser and Petri, 1996). For reliable operation of a fast breeder reactor, the fuel elements must be resistant to breaching even in case of overpower transients. From the standpoint of accident transients, the most important factor is the accelerated rate of cladding attack, once eutectic liquefaction forms at the interface. The chemical compatibility between the fuel and clad material also known as fuel-clad chemical interaction (FCCI) is of prime importance because of the formation of low melting eutectic which may often limit the life of the fuel pin in a reactor. The temperature of the eutectic reaction between the fuel and the cladding is considered as a critical parameter for the design of the metallic fuel pin. The fuel-clad interdiffusion behaviours involving U–10Zr, U–Pu–10Zr, U–Pu alloy fuels and Fe, Fe–Cr, Fe–Cr–Ni, SS316, D9 type of cladding materials have been reported in the literature (Cohen et al., 1993; Hofman and Walters, 1994; Keiser and Dayananda, 1993, 1994; Keiser and Petri, 1996; Lee et al., 2009; Nakamura et al., 1999, 2001; Ogata et al., 1997). The FCCI issues with

austenitic cladding might be largely academic, as swelling alone renders these steels unacceptable for use. High burn-up operation requires a low-swelling type material such as HT9 and T91 grade steels. Therefore, FCCI is more important in these types of steels. The interdiffusion between U–10Zr alloy and HT9 at 973 K was studied by Keiser and Dayananda (1994), Keiser and Petri (1996) and Lee et al. (2009). The diffusion behavior between U–10Zr with T91 was investigated by Ryu et al. (2009). Studies on the fuel-clad chemical compatibility or diffusion behaviour between U–Zr alloys and feritic/martensitic steels have been reported in the literature, with emphasis on the U–10Zr alloy fuel and HT9 cladding steel. However, it appears that the work on fuel–clad chemical interaction between U–6Zr alloy and T91 steel was very limited in the literature and this requires further investigation. Uranium with low concentration of Zr (∼6 wt.%) in the fuel is preferred in the Indian nuclear power programme with the objective of achieving a higher breeding ratio (Devan et al., 2011). On account of the above, the determination of the fuel-clad chemical compatibility of this particular fuel composition is significant. The performances of different FCCI barrier layers between fuel and clad, such as V, Cr, Zr, Nb, Ti and Mo, have been reported in the literature (Keiser and Cole, 2007; Kim et al., 2009; Ryu et al., 2009; Yang et al., 2010). The present investigation mainly emphasizes on Zr as FCCI barrier layer. However, detailed report on the effectiveness of Zr as a barrier layer between U and T91 steel is limited. The Zr barrier layer between U and T91 will interact with both U as well as T91. The kinetics of interdiffusion of Zr and U has been reported in the literature (Mash and Disselhorst, 1954; Rough, 1955). However, the reaction kinetics between Zr and T91 is not available in the open literature. The main objectives of the present study are: (a) evaluation of performance of Zr as FCCI barrier layer between U and T91 by diffusion couple experiments, (b) determination of growth kinetics of reaction layers at U/Zr as well as Zr/T91 interfaces, and (c) studies on the fuel-clad chemical compatibility between U–6Zr and T91 by diffusion couple experiments. 2. Experimental procedure Slugs of uranium and U–6Zr alloy were prepared using 99.9% pure uranium and 99.95% pure crystal bar of zirconium by following injection casting route. The typical diameter of the alloy slugs was about 5 mm. The trace metallic constituents were totaling to only 70 ppm as revealed by the inductively coupled plasma-atomic

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269

Table 1 Chemical composition of T91 cladding steel (wt.%). Cr

Mo

V

Nb

Al

Ti

Ni

Cu

Mn

Si

C

N

P

S

Fe

9.161

0.882

0.207

0.079

0.008

0.003

0.197

0.068

0.368

0.209

0.099

0.0457

0.015

0.0013

balance

emission spectroscopic (ICP-AES) analysis. T91 grade steel was used in the standard normalized and tempered condition with a hardness of 220 kg mm−2 . The heat treatment of T91 consists of austenization at 1323 K and air quenching, followed by tempering at 1023 K for 1 h. The chemical composition of T91 steel is given in Table 1. The uranium and U–6Zr slugs were cut into 3 mm thick discs and T91 steel rod of the same diameter was cut into discs about 0.5 mm thick. The surfaces of all these discs were metallographically polished to 1 ␮m surface finish. Two types of diffusion couples were prepared as described below: (a) couples between U and T91 discs with a Zr foil of thickness ∼200 ␮m between them, would be referred to as U/Zr/T91 couples, and (b) couples of a disc of U–6Zr alloy between two T91 steel discs, referred to as U–6Zr/T91 couples. The components of the diffusion couples were kept in fixtures made of Inconel 600, to ensure intimate contact during annealing. Ta foil was used to prevent any chemical reaction between couples and fixture. The diffusion couple-fixture assembly is shown schematically in Fig. 2. The fixtures containing these couples were encapsulated in quartz tube in helium atmosphere and annealed in a resistance heating furnaces maintained at 923, 973 and 1023 K for duration up to 1500 h. The heat treatment schedule of the couples is given in Table 2. Subsequent to annealing, the couples were sectioned using a slow speed diamond cutting wheel. The exposed cross sections were metallographically polished to 1 ␮m surface finish. The extent of reaction and phases formed at the interface were characterised using a scanning electron microscope (SEM) with energy dispersive spectroscope (EDS) and an electron probe microanalyzer (EPMA) (CAMECA SX100) equipped with three wavelength dispersive spectroscopes (WDS). The X-ray line scans of U, Fe, Cr and Zr were acquired across the interfaces of the diffusion couples to determine the distribution of each element.

3. Results 3.1. Interfacial reactions in U/Zr/T91 couples The microstructure of the U/Zr/T91 diffusion couple annealed at 923 K for 1500 h is shown in Fig. 3(a). The intensity profiles of U-M␣, Zr-L␣, Fe-K␣ and Cr-K␣ X-ray lines, as shown in Fig. 3(b), were recorded along a line AB marked in Fig. 3(a). The microstructure and the intensity profiles across the interface of the U/Zr/T91 couple annealed at 973 K for 1500 h are shown in Fig. 4(a) and (b) respectively. Similarly, Fig. 5(a) and (b) shows the microstructure and the intensity profiles of the diffusion couple annealed at 1023 K for 1500 h. The thicknesses of the interdiffusion layers at U/Zr and Zr/T91 interfaces for U/Zr/T91 diffusion couples were summarized in Table 3 . The microstructural analysis of U/Zr/T91 diffusion couples revealed an excellent bonding formation at both the U/Zr and Zr/T91 interfaces after annealing. The thicknesses of interdiffusion layers at the U/Zr interface at 923 and 973 K were found to be in the range of 10–20 ␮m even after annealing up to 1500 h (Table 3). The microstructure of U/Zr/T91 diffusion couple at 1023 K for 1500 h (Fig. 5(a)) shows that U had diffused into 200 ␮m thick Zr barrier layer and reached very close to the Zr/T91 interface. A reasonable amount of Zr was also diffused into U metal. However, no eutectic melting between U–Zr interdiffusion layer and T91 has been observed at Zr/T91 interface. The magnified micrographs of the U/Zr interdiffusion layer are shown in Fig. 5(c). The interdiffusion layer is having Widmanstatten morphology. The EDS spectra of the phases formed on the U/Zr interdiffusion layer are shown in Fig. 5(d) and (e). These clearly indicate that the U/Zr interdiffusion layer does not contain any Fe or Cr. On the other hand, the interdiffusion layer at Zr/T91 interface is found to be only 20 ␮m and formation of any intermetallic compound was not found at the Zr/T91 interface, even at even at 1023 K. As was evident from the intensity profiles of U-M␣, Zr-L␣, Fe-K␣ and Cr-K␣ X-ray lines, the depth of penetration of Cr into the Zr metal is smaller than that of Fe. 3.2. Interfacial reactions in U–6Zr/T91 couples Fig. 6(a) shows the microstructure of the interdiffusion layer formed at the interface of the U–6Zr/T91 diffusion couple after annealing at 973 K for 500 h. The magnified micrograph of the interface is shown in Fig. 6(b). The intensity profiles of U-M␣, ZrL␣, Fe-K␣, Cr-K␣ X-ray lines, as shown in Fig. 6(c) were recorded across the interface along the line GH marked in Fig. 6(a). The interdiffusion between the two specimens resulted in the formation of three different layers at the interface: (U,Zr)(Fe,Cr)2 phase,

Table 2 Annealing time and temperature of the diffusion couple experiments.

Fig. 2. Schematic diagram of the diffusion couple-fixture assembly for U/Zr/T91diffusion couple experiment.

No.

Capsule

Temperature (K)

Time (h)

1. 2. 3. 4. 5. 6. 7.

U/Zr/T91 U/Zr/T91 U/Zr/T91 U/Zr/T91 U/Zr/T91 U–6Zr/T91 U–6Zr/T91

923 973 973 973 1023 973 1023

1500 500 1000 1500 1500 500 100

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Table 3 Widths of the reaction layers at U/Zr and Zr/T91 interfaces for U/Zr/T91 diffusion couples, under different heat treatment conditions. No.

Capsule

Temperature (K)

Time (h)

1. 2. 3. 4. 5.

U/Zr/T91 U/Zr/T91 U/Zr/T91 U/Zr/T91 U/Zr/T91

923 973 973 973 1023

1500 500 1000 1500 1500

a Zr-rich layer and a Zr-depleted layer as shown in Fig. 6(b). The (U,Zr)(Fe,Cr)2 layer of thickness around 10 ␮m was formed on the T91 side. The formation of 15–20 ␮m thick Zr-depleted layer (containing ∼94 at.% U, 4 at.% Zr, 2 at.% Fe), was found on the U–Zr alloy side. The Zr-rich layer (containing ∼9 at.% U, 87 at.% Zr, 4 at.% Fe) of thickness 2–3 ␮m was formed in between the above two. This Zrrich layer at the interface acts as fuel-clad diffusion barrier. It may be noted here, that no melting has been observed at this temperature. The (U,Zr)(Fe,Cr)2 , Zr-rich and Zr-depleted layers are marked as ‘X’, ‘Y’ and ‘Z’, respectively in Fig. 6(b). The U–6Zr/T91 diffusion couple was melted when it was annealed at 1023 K for 100 h. The microstructure of the reaction zone, as shown in Fig. 7, revealed that three different phases were formed due to the interdiffusion reaction between U–6Zr and T91 at 1023 K. It has been confirmed by the EDS analysis that the

Fig. 3. (a) Secondary electron micrograph of U/Zr/T91 diffusion couple annealed at 923 K for 1500 h. (b) Intensity profiles of U-M␣, Zr-L␣, Fe-K␣ and Cr-K␣ X-ray lines along the line AB marked in (a).

Width of reaction layer (␮m) U/Zr

Zr/T91

9 12 14 20 200

10 10 14 16 20

bright phase was U6 Fe, the grey and dark phases are U(Fe,Cr)2 and Zr(Fe,Cr)2 , respectively. A limited solubility of Zr in U(Fe,Cr)2 phase and U in Zr(Fe,Cr)2 phase have also been observed. 4. Discussion 4.1. Interfacial reactions in U/Zr/T91 couples The results of the U/Zr/T91 diffusion couples indicate that the Zr liner was effective in preventing fuel-clad chemical interaction up to 973 K. Also, Zr liner helps in achieving a good thermal bond between fuel and clad, facilitating excellent heat transfer from fuel to clad. The thickness of the interdiffusion layer at the U/Zr interface was increased enormously when the annealing temperature was increased from 973 K to 1023 K for the same annealing time of

Fig. 4. (a) Secondary electron micrograph of U/Zr/T91 diffusion couple annealed at 973 K for 1500 h. (b) Intensity profiles of U-M␣, Zr-L␣, Fe-K␣ and Cr-K␣ X-ray lines along the line CD marked in (a).

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271

Fig. 5. (a) Microstructure of U/Zr/T91 diffusion couple annealed at 1023 K for 1500 h, (b) Intensity profiles of U-M␣, Zr-L␣, Fe-K␣ and Cr-K␣ X-ray lines along the line EF marked in (a). (c) Magnified micrographs of the U/Zr interdiffusion layer showing Widmanstatten morphology. (d) EDS spectrum on acicular phase on U/Zr interdiffusion layer. (e) EDS spectrum on bright matrix phase on U/Zr interdiffusion layer.

1500 h. However, the increase in thickness of the interdiffusion layer at the Zr/T91 interface was very insignificant. Therefore, it is evident that the rate of interdiffusion and penetration is sluggish in nature at the Zr/T91 interface compared to the U/Zr interface. At 1023 K, uranium and zirconium exist as ␤-U and ␣-Zr forms, respectively. The interdiffusion between them results in formation of a solid solution (␥-U, ␤-Zr). The higher rate of diffusion at the U/Zr interface compared to the Zr/T91 interface at 1023 K may be attributed to the formation of (␥-U, ␤-Zr) solid solution between U and Zr over a wider range of composition. The microstructure of the U–Zr solid solution shows a typical Widmanstatten morphology, which contains acicular Zr structure on (␥-U, ␤-Zr) matrix as evident from the EDS spectra (Fig. 5(d) and (e)). The solubility of U in

the acicular Zr phase was found to be negligibly small (Fig. 5(d)). The (␥-U, ␤-Zr) matrix phase contains ∼17.5 wt.% (7.5 at.%) U (Fig. 5(e)). This typical structure was formed during cooling from 1023 K after diffusion annealing. The formation of similar acicular alpha Zr in (␥U, ␤-Zr) matrix was also reported by Rough (1955) for Zr–19 wt.% U alloy when it was air cooled after annealing at 1123 K for 1 h. Any evidence, related to formation of liquid phase was not observed in the U/Zr/T91 diffusion couple even after annealing at 1023 K for 1500 h. The eutectic reaction temperature between U and T91 was found to be 995 K by the present authors (Kaity et al., 2010b). Under accidental condition, the fuel-clad interface may reach 1023 K. Therefore, the present observation indicate that even holding at temperatures higher than the eutectic temperature for

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Fig. 6. (a) Microstructure of interdiffusion layer formed at the interface of the U–6Zr/T91 diffusion couple after annealing at 973 K for 500 h. (b) The magnified micrograph of the interface showing the formation of three diffusion layers. (c) Intensity profiles of U-M␣, Zr-L␣, Fe-K␣, Cr-K␣ X-ray lines recorded along the line GH marked in (a).

more than a month, does not cause any damage to fuel pin because of the presence of Zr barrier. This fuel is expected to reach such high temperatures under accidental condition, for very short duration. Hence, it can be concluded that the performance of a Zr barrier layer

of about 200 ␮m would suffice to prevent the formation of liquid phase. At the Zr/T91 interface, no intermetallic compounds were found to form by diffusion reaction. Fe being the primary constituent of T91, this interface may be compared with the diffusion behaviour of Zr/Fe couple. The Zr–Fe phase diagram shows the existence of three intermetallic compounds, viz. ZrFe2 , Zr2 Fe and Zr3 Fe (Arias and Abriata, 1988). However, the diffusion reaction experiments by Bhanumurthy et al. (1991) showed that none of these compounds formed at the Zr/Fe interface up to 1073 K. The first phase to form at this interface was Zr3 Fe and the interdiffusion coefficient of this compound (DZr3 Fe ) was estimated to be 5.5 × 10−16 m2 s−1 at 1134 K (Bhanumurthy et al., 1991). The sluggish nature of diffusion reaction at the Zr/T91 interface may be attributed to such low value of the interdiffusion coefficient. The depth of penetration of Cr, which is the next important component of T91, to the Zr metal is smaller than that of Fe. Pande et al. (1968) reported that diffusivity of Cr in Zr is much smaller than the diffusivity of Fe in Zr. 4.1.1. Kinetics of the interfacial reaction at U/Zr and Zr/T91 interfaces The growth kinetics of a reaction layer formed at the interface can be represented by the following relation: w = kt 1/n

Fig. 7. The microstructure of U–6Zr/T91 diffusion couple annealed at 1023 K for 100 h showing the formation of U6 Fe, U(Fe,Cr)2 and Zr(Fe,Cr)2 phases.

(1)

where, w is the width of the reaction layer, t is the time of the reaction, k is the reaction rate constant, and n is the reaction index. The value of the index n depends upon the mechanism of the growth

S. Kaity et al. / Nuclear Engineering and Design 250 (2012) 267–276

 −Q 

(2)

RT

where, Q is the activation energy required for the growth of the layer. Combining Eqs. (1) and (2), w = k0 t 1/n exp

 −Q 

(3)

RT

The value of activation energy, Q can be evaluated using a log w versus 1/T plot. The diffusion behaviour of U in Zr was studied by Mash and Disselhorst (1954) over the temperature range 573–1323 K for time interval up to 1812 h. According to them, the temperature dependence of k follows the following relations:



Above 873 K,

k = 7.26 × 10−3 exp

 Below 873 K,

k = 1.37 × 10

−7

exp

−92.728 kJ mol−1 RT



− 18.543 kJ mol−1 RT



−8

D = 3.54 × 10

exp

−128.030 kJ mol−1 RT

In ˛–Zr(867 to 1053 K),



D = 1.08 × 10

−8

exp

−133.888 kJ mol−1 RT

2.00x10 -5

1.50x10 -5

-5

2x10

6

3x10

6

4x10

6

5x10

6

6x10

6

Time (sec)

b

2.00x10

-5

1.50x10

-5

1.00x10

-5

(4)

 (5)

where, k is the rate constant in m s−1/2 , R is the universal gas constant (8.314 J mol−1 K−1 ), T is temperature in the absolute scale. The slow kinetics of the diffusion reaction below 873 K is attributed to the sluggish nature of formation of ␦-UZr2 (Mash and Disselhorst, 1954). The temperature dependence rate constant (k) above 873 K (Eq. (4)) was determined by carrying out the diffusion reaction at 1013, 1123, 1273 and 1323 K where the U–Zr system forms a single ␥-solid solution, almost in the whole range of composition. This indicates that the rate constants at 923 and 973 K were determined by them by extrapolating these high temperature data. The rate of interdiffusion of U in the ␤-Zr zirconium phase is much faster than the rate in the ␣-Zr phase (Rough, 1955). The rate of interdiffusion of U and Zr is reproduced as follows (Rough, 1955): In ˇ–Zr(1053 to 1303 K),

2.50x10 -5

1.00x10 6 1.5x10

 (6)

 (7)

where, D is the interdiffusion coefficient in m2 s−1 . In the present investigation, the growth kinetics of the reaction zone at both the U/Zr and Zr/T91 interfaces were determined at 973 K. The plot of log (width) versus log (time) for the U/Zr interface at 973 K was shown in Fig. 8(a). The value of n was determined from the slope of the plot and was found to be around 2. The reaction constant (k) was calculated from the intercept of the plot. The rate constant for the growth reaction layer formed at the U/Zr interface was determined to be 2.07 × 10−8 m s−1/2 at 973 K. The rate constant or the penetration coefficient (k) is known to be proportional to the square root of the diffusion coefficient (Mash and Disselhorst, 1954). The square root of the interdiffusion coefficient (i.e. D1/2 ) of U in Zr at 973 K calculated by Eq. (7) was found to be 2.65 × 10−8 m s−1/2 . This indicates that the present data of rate constant at U/Zr interface at 973 K (i.e. 2.07 × 10−8 m s−1/2 ) and the square root of the diffusion coefficient (i.e. D1/2 ) of U in Zr at 973 K

Width of reaction zone ( µm)

k = k0 exp

a Width of reaction zone ( µm)

of the reaction layer. In case of solid state reactions, the growth is diffusion controlled and the value n becomes 2 (Sengupta et al., 2006). The reaction constant k is temperature dependent parameter and increases exponentially with temperature as

273

-6

9.00x10 6 1.5x10

2x10

6

3x10

6

6

4x10

6

5x10

6

6x10

Time (sec) Fig. 8. (a) Plot of log (width) versus log (time) for the U/Zr interface at 973 K. (b) Plot of log (width) versus log (time) for the Zr/T91 interface at 973 K.

calculated by Eq. (7) (i.e. 2.65 × 10−8 m s−1/2 ) are very close to each other. Similarly, the plot of log (width) versus log (time) for the Zr/T91 interface at 973 K was shown in Fig. 8(b). The reaction index n was found to around 2 and the rate constant for the growth reaction layer at the Zr/T91 interface was found to be 1.95 × 10−8 m s−1/2 at 973 K. The activation energy Q for the reaction at Zr/T91 interface could be estimated from the data for the width of the reaction layer formed at different temperature for equal length of time. The widths of the reaction layers were plotted against the inverse temperature on a semi-log plot for the samples annealed for 1500 h (Fig. 9). The activation energy Q for the reaction at the Zr/T91 interface was determined from the slope of straight line and was found to be 54.7 kJ mole−1 . The pre-exponential factor k0 was calculated from the intercept of the plot and was found to be 5.54 × 10−6 m s−1/2 . The growth of interdiffusion layer at the Zr/T91 interface can be written as:



−6

k = 5.54 × 10

exp

−54.7 kJ mol−1 RT



(923 ≤ T/K ≤ 1023) (8)

Eq. (1) can also be expressed as:



w = 5.54 × 10−6 t 1/2 exp (923 ≤ T/K ≤ 1023)

−54.7 kJ mol−1 RT



(9)

274

S. Kaity et al. / Nuclear Engineering and Design 250 (2012) 267–276

-5

Width of reaction layer ( µm)

2.00x10

-5

1.50x10

-5

1.00x10

-6

9.00x10

0.95

1.00

1.05

1.10

3

1/T (x10 ) Fig. 9. Plot of log (width) of the reaction layers versus 1/T for the Zr/T91 interfaces of U/Zr/T91 couples annealed for 1500 h.

4.2. Interfacial reactions in U-6Zr/T91 couples The interdiffusion between U–6Zr alloy and T91 steel resulted in formation of parallel layers of (U,Zr)(Fe,Cr)2 phase, a Zr-rich layer and a Zr-depleted layer at the interface. The formation of similar type of diffusion layers was reported by Hofman and Walters (1994) in the diffusion couple experiments between U–19Pu–10Zr and HT9 cladding. According to them, a Zr rich layer containing interstitial nitrogen is first formed at the fuel-clad interface of U–19Pu–10Zr and HT9. Then U and Pu from fuel diffuse through the Zr layer and interact with the elements in the cladding material. This leads to the formation of UFe2 and U6 Fe type phases in a single zone on the cladding side of the Zr rich layer. The formation of U(Fe,Cr)2 phase and a Zr-rich layer in the diffusion reaction between the U–10Zr and HT9 cladding at 973 K was reported by Keiser and Dayananda (1994). According to them, Cr plays crucial role in the development of Zr rich layer. It was expected that Cr affects the activity of Zr in the ␦-UZr2 precipitates and cause them to become unstable. The decomposition of ␦ precipitates causes formation of a Zr rich layer, which can reduce interdiffusion of fuel and cladding elements. The formation of (U,Zr)(Fe,Cr)2 , Zr-rich layer and Zr-depleted layer was also observed by Lee et al. (2009) in the diffusion couple experiment between U–10Zr and HT9 at 973 K. According to them, the ␦-UZr2 phase in the vicinity of the boundary layers between two specimens would decompose into U and Zr elements. The U in the matrix would diffuse to the interface between the two specimens forming (U,Zr)(Fe,Cr)2 phase, whereas the decomposed Zr is saturated to form a Zr-rich layer. This reaction results in the formation of a Zr-depleted layer between Zr-rich layer and U–Zr matrix. Recently, the interaction of the elements in the clad (Fe, Ni, Cr) with U–Zr alloy has been discussed with the help of computational atomistic modeling approach by Bozzolo et al. (2010, 2011). The segregation in U–Zr solid solution and fuel–clad interaction has been interpreted by means of the concepts of strain and chemical energy underlying the BFS method (Bozzolo et al., 1992). The segregation in U–Zr solid solution is strongly dependent on lattice parameter, temperature, concentration of Zr and surface energy. According to them, at low Zr concentration range (0 < XZr < 22.45 (at.%) (U–10 wt.% Zr)), a slightly favorable Zr surface energy and lattice contraction leads to the Zr segregation. On the other hand, as the concentration of Zr increases, the lattice parameter shows a rapid growth, fast approaching bulk Zr values, thus reducing the BFS strain energy of Zr atom and U finds itself in

a high BFS strain environment which favors U segregation to the surface. On the basis of the above observation, let us analyze the results of the present study. The U–Zr phase diagram (Sheldon and Peterson, 1989) shows that the ␦ phase is not stable above 890 K. If any ␦ phase is present in the sample, it should decompose above 890 K. The fact, that a Zr-rich layer forms at the interface of U–10Zr/Fe couple (Keiser and Dayananda, 1993) where Cr is not present, eliminates the requirement for Cr in decomposition of the ␦-UZr2 phase. The U–Zr alloy used by both Keiser and Dayananda (1994) and Lee et al. (2009) was homogenized at 1173 K prior to the diffusion couple experiment. Since the peritectoid reaction, ␥ + ␣ → ␦ is extremely sluggish in nature, their samples may not have contained any ␦-phase at room temperature (Basak et al., 2009a, 2009b). In fact, Hills et al. (1965) have found that up to 20 at.% Zr (∼10 wt.% Zr), supersaturated ␣ phase is present at room temperature. Hence, the formation of the Zr-rich layer cannot be explained on the basis of decomposition of ␦-UZr2 in the U–Zr alloy. The formation of such a diffusion layer can easily be explained with the help of atomistic modeling of fuel-clad interaction using BFS method of alloys given by Bozzolo et al. (2010, 2011). The deviation of the lattice parameter of U–Zr solid solution with respect to their average values calculated by Vegard’s law, as a function of concentration of Zr (in at.%), was determined by Bozzolo et al. (2010). According to them, the lattice parameter of U–6Zr solid solution, in the temperature range of 1000–1400 K, shows a minor contraction with respect to their average values calculated by Vegard’s law and the lattice parameters exhibit a negligible qualitative and quantitative dependence on temperature. U–6Zr alloy exists as solid solution (␥-U, ␤-Zr) at 973 K and therefore, it will follow same trend of contraction of lattice parameter at 973 K. The surface energy of Zr was found to be slightly lower than that of U in the solid solution of U–6Zr alloy (Bozzolo et al., 2010). The lower Zr surface energy as well as lattice contraction lead to the segregation of Zr toward the surface. The Zr enrichment at the surface creates an initial barrier for Fe, Cr diffusion. Due to the sluggish nature of diffusion reaction between Zr and Fe, as discussed in Section 4.1, this Zr cannot react with T91 steel. Thus, there is a Zr-rich layer (layer Y, Fig. 6b) on the surface and a Zr-depleted inner zone (layer Z) to maintain mass balance of Zr element. However, due to strong interaction of U with Fe of T91 steel, some U diffuses though this Zr rich layer and comes in contact with Fe of the steel. This results in formation of a U–Fe layer (layer X) closer to T91 and outside Zr rich layer. This outermost layer (layer X) has small amount of Cr and Zr, forming (U,Zr)(Fe,Cr)2 with ∼10 at.% Zr and ∼5 at.% Cr. The literature data on diffusion couple experiment between U–10Zr and HT9 showed that the thickness of the (U,Zr)(Fe,Cr)2 layer gradually increases with annealing time (Lee et al., 2009). This is a direct evidence of U diffusion through the Zr rich layer. The Zr rich layer acts as a barrier by reducing the penetration of Fe and Cr elements through it towards fuel lattice. Therefore, the amounts of Fe and Cr in layer Z are negligible. This is also evident from our diffusion couple experiments with a Zr liner between U and T91 cladding, as discussed in Section 3.1. Considerable U had diffused into Zr barrier layer and reasonable amount of Zr diffused into U metal, but no detectable Fe or Cr was found in either Zr barrier layer or U lattice (Fig. 5a). From the intensity profiles of U M␣, Zr L␣, Fe K␣ and Cr K␣ X-ray lines of U–6Zr/T91 couple (Fig. 6c) it was observed that the diffusion of Cr into the U–Zr alloy is slower than that of Fe. U does not form any intermetallic compound with Cr and the mutual solubilities of Cr and U are almost negligible. Hence, Cr shows weak interaction with U–Zr alloy. When the annealing temperature was increased to 1023 K, eutectic-melted microstructure was observed as shown in Fig. 7

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and it was much different from the layered structures formed below the eutectic temperature. The authors have measured that the eutectic temperature between U–6Zr and T91 as 995 K. At 1023 K, the Zr-rich barrier layer formed at the interface is not effective to prevent the interdiffusion and both the U and Zr completely react cladding components which lead to the eutectic melting. The result indicates the measured eutectic temperature between U–6Zr and T91 is very crucial from reactor safety point of view. In any case, if the clad temperature reaches above the eutectic temperature, the direct contact of U–6Zr and T91 cladding would cause eutectic melting as observed in this diffusion couple experiment. 5. Summary and conclusions The performance of Zr as FCCI barrier layer between U and T91 was evaluated by diffusion couple experiments. The growth kinetics of reaction layers at U/Zr as well as Zr/T91 interfaces was established. The fuel–clad chemical compatibility between U–6Zr and T91 was also studied by diffusion couple experiments. The following conclusions are drawn from this study: (a) The thicknesses of interdiffusion layers at the U/Zr and Zr/T91 interfaces at 923 and 973 K were found to be in the range of 10–20 ␮m even after annealing up to 1500 h. (b) The microstructure of U/Zr/T91 diffusion couple at 1023 K for 1500 h showed that U had diffused into 200 ␮m thick Zr barrier layer and reached very close to the Zr/T91 interface. However, no eutectic melting between U–Zr interdiffusion layer and T91 has been observed at Zr/T91 interface. (c) The growth kinetics of the reaction zone at both the U/Zr and Zr/T91 interfaces were determined at 973 K from the plot of log (width) versus log (time). The value of reaction index n was found to be around 2 at both the U/Zr and Zr/T91 interfaces. The reaction rate constant (k) for the growth reaction layer at the U/Zr interface was determined to be 2.07 × 10−8 m s−1/2 at 973 K. (d) Similarly, the rate constant for the growth reaction layer at the Zr/T91 interface was found to be 1.95 × 10−8 m s−1/2 at 973 K. (e) The activation energy Q for the reaction at the Zr/T91 interface was determined from the slope of straight line and was found to be 54.7 kJ mole−1 . (f) The interdiffusion between U–6Zr and T91 at 973 K resulted in the formation of three different layers at the interface: (U,Zr)(Fe,Cr)2 phase, a Zr-rich layer and a Zr-depleted layer. The Zr-rich layer at the interface acts as diffusion barrier. (g) Fuel-clad interdiffusion was strongly retarded by the addition of Zr in the fuel and this is due to formation of Zr-rich barrier layer at the interface. (h) At 1023 K, U–6Zr/T91 couple reacted completely with each other causing a complete melt down of the clad in 100 h. This contrary to what observed for U/Zr/T91 couple at the same temperature. This is a very important observation in support of a Zr barrier layer concept compared to that when Zr is present as a constituent element in the fuel slug. References Arias, D., Abriata, J.P., 1988. The Fe–Zr (iron–zirconium) system. Bull. Alloy Phase Diagrams 9 (5), 597–604. Basak, C., Keswani, R., Prasad, G.J., Kamath, H.S., Prabhu, N., 2009a. Phase transformations in U–2 wt% Zr alloy. J. Alloys Compd. 471, 544–552. Basak, C., Prasad, G.J., Kamath, H.S., Prabhu, N., 2009b. An evaluation of the properties of as-cast U-rich U–Zr alloys. J. Alloys Compd. 480, 857–862. Bhanumurthy, K., Kale, G.B., Khera, S.K., 1991. Reaction diffusion in the zirconium–iron system. J. Nucl. Mater. 185, 208–213. Bozzolo, G., Ferrante, J., Smith, J.R., 1992. Method for calculating alloy energetics. Phys. Rev. B 45, 493–496.

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