Current status of materials development of nuclear fuel cladding tubes for light water reactors

Nuclear Engineering and Design 316 (2017) 131–150

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Current status of materials development of nuclear fuel cladding tubes for light water reactors Zhengang Duan a, Huilong Yang b, Yuhki Satoh c, Kenta Murakami b, Sho Kano b, Zishou Zhao b, Jingjie Shen b, Hiroaki Abe b,⇑ a b c

Department of Quantum Science and Energy Engineering, Graduate School of Engineering, Tohoku University, Sendai, Miyagi 980-8577, Japan Department of Nuclear Engineering, School of Engineering, The University of Tokyo, Nakagun, Ibaraki 319-1188, Japan Institute for Materials Research, Tohoku University, Sendai, Miyagi 980-8577, Japan

a r t i c l e

i n f o

Article history: Received 7 October 2016 Received in revised form 15 February 2017 Accepted 26 February 2017

Keywords: LWR BWR PWR Fuel cladding Zr-based alloy Coating SiC

a b s t r a c t Zirconium-based (Zr-based) alloys have been widely used as materials for the key components in light water reactors (LWRs), such as fuel claddings which suffer from waterside corrosion, hydrogen uptakes and strength loss at elevated temperature, especially during accident scenarios like the lost-of-coolant accident (LOCA). For the purpose of providing a safer, nuclear leakage resistant and economically viable LWRs, three general approaches have been proposed so far to develop the accident tolerant fuel (ATF) claddings: optimization of metallurgical composition and processing of Zr-based alloys, coatings on existing Zr-based alloys and replacement of current Zr-based alloys. In this manuscript, an attempt has been made to systematically present the historic development of Zr-based cladding, including the impacts of alloying elements on the material properties. Subsequently, the research investigations on coating layer on the surface of Zr-based claddings, mainly referring coating materials and fabrication methods, have been broadly reviewed. The last section of this review provides the introduction to alternative materials (Non-Zr) to Zr-based alloys for LWRs, such as advanced steels, Mo-based, and SiC-based materials. Ó 2017 Elsevier B.V. All rights reserved.

Contents 1. 2.

3. 4.

5.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Development of Zr-based alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. The characteristics of zirconium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. Effects of alloying elements on Zr-based alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. The development of Zr-based claddings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.1. Zircaloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.2. E110, E125 and E635 alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.3. ZIRLO (Zr low oxygen) series alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.4. M5 alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.5. X5A alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.6. Other Zr-based alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Coatings on Zr-based alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Alternative materials to Zr-based alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. Advanced steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. Mo-based alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3. SiC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

132 132 132 132 134 134 135 136 137 137 138 139 142 142 143 144 147 147 147

⇑ Corresponding author at: Department of Nuclear Engineering, School of Engineering, The University of Tokyo, Shirokata Shirone2-22, Tokai-mura, Naka-gun, Ibaraki 319-1188, Japan. E-mail addresses: [email protected] (Z. Duan), [email protected] (H. Abe). http://dx.doi.org/10.1016/j.nucengdes.2017.02.031 0029-5493/Ó 2017 Elsevier B.V. All rights reserved.

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1. Introduction Generally, fission products are confined primarily by means of four successive physical barriers: the fuel matrix, the fuel cladding, the boundary of the reactor coolant system and the containment system, of which integrities are protected against internal and external hazards by the implementation of the defence in depth (DiD) concept (IAEA, 1997). As the second security barrier, the selection of materials and corresponding fabrication process for the cladding is of importance. Due to combinations with a low thermal neutron capture cross section, excellent corrosion resistance in high-temperature water and adequate mechanical properties, Zr-based alloys are widely used as materials for all commercial LWR fuel cladding. The motivations to achieve acceptable safety margin and higher burnup are driving the evolution in the reliability of Zr-based cladding, thereby contributing to the birth of advanced Zr-based alloys like E635, ZIRLO, M5, MDA, X5A and J-alloys, etc., which strode a major step forward in improvement of corrosion resistance and mechanical performances. However, the well-known inherent demerits of Zr-based alloys such as rapid oxidation and hydrogen production at hightemperature steam, cannot be altered by optimizing the Zr-alloy chemical composition and/or their manufactory processes, which was highlighted further by the nuclear accident happened at the Fukushima Daiichi plants. As a result, the implementation of DiD concept has been reinforced and the term ATF has become popular (NEA/OECD, 2016). Comparison with the current UO2/Zr system, the enhanced ATF, described by US Department of Energy (US DoE), can tolerate loss of active cooling in the core for a considerably longer time period while maintaining or improving the fuel performance during normal operations (Goldner, 2012). Therefore, the ATF cladding concept is inevitably a hot topic against the background of the deployment of the ATF rods, resulting in innovative designs, such as Zr alloys with coating, advanced steels, and ceramic cladding. Coating technology has been attempted to enhance the waterside corrosion resistance and wear resistance without modifying the existing Zr-based alloys, which enable this coated cladding to be put into commercial application in a short term; additionally, the alternative materials (Non-Zr) have been under development to substitute existing Zr-based claddings. Whereas, it is worth noting that the long time together with the high cost is needed to achieve a reasonable return due to the involvement of modifications on fuel enrichment, size of assemblies associated with the corresponding alternative materials. In this paper, an introduction to over past 50-year evolution in Zr-based cladding for LWR (primarily BWR, PWR and SCPWR) fuel is presented, in which the impacts of alloying elements on the performances of Zr-based alloys are briefly described. Then, the up-todate ATF claddings, coating technology and alternative materials, have been summarized successively in this review.

2. Development of Zr-based alloys 2.1. The characteristics of zirconium Cladding tubes are the vital part of a nuclear reactor because they not only provide an enclosure to the highly radioactive fuel but also remain in direct contact with the coolant during reactor operation which makes it vulnerable to corrosion (Alam et al., 2011). Hence, the material used for cladding tubes must have the important characteristics as follows; low thermal neutron capture cross-sections, high thermal conductivity, high strength, and high corrosion resistance. From considerations of neutron economy, the suitability of materials for water-cooled thermal power reactors is restricted

Table 1 Calculated effective neutron absorption cross section (neutron absorption cross section per unit of yield strength) for pure elements in comparison to Zr (Azevedo, 2011). Elements

Neutron absorption cross section (Barns)

Yield strength (MPa)

Relative effective neutron absorption cross section in relation to Zr

Be C Mg Si Zr Al Mo Cr Nb Fe Ni V Sn

0.009 0.004 0.063 0.16 0.185 0.231 2.48 3.05 1.15 2.55 4.43 5.04 0.63

200–350 24–28 65–100 165–180 135–310 30–40 170–350 185–280 75–95 110–165 80–280 125–180 7–15

0.04 0.2 1 1 1 8 10 15 15 20 30 40 70

to metals like aluminum (Al), magnesium (Mg), and Zr or their alloys (see Table 1 (Azevedo, 2011)). However, severe blistering and accelerated corrosion precluded application of aluminum for power reactors at the temperature above 150 °C. Beryllium (Be), manganese (Mn) and their alloys also were not considered for cladding due to their failure to meet the mechanical properties and corrosion resistance at water-cooled reactor operating temperature (200–300 °C) (IAEA, 1987). At normal operating reactor temperature (300 °C), Zr is an exceptional substance and has been employed as fuel cladding tubes since the early 1950’s because it has much lower neutron absorption cross section as well as adequate mechanical properties than the other commercially available structural materials (Baczynski, 2014). It was initially thought that the poor corrosion resistance of the unalloyed zirconium produced by the van Arkel process was a result of stray impurities, but it was found that improving the purity by Kroll process could not eliminate the problems. The oxidation rate of pure Zr decides the oxide grain orientation in the oxide films, which led to severe growth stresses and cracking of the oxide and oxide spalling (Cox, 2005). Therefore, to get additional properties for better operation in water cooled reactors, Zr was alloyed with constituents which have low nuclear impacts like tin (Sn), oxygen (O), and niobium (Nb), while other transition metals (iron (Fe), chromium (Cr), nickel (Ni), etc.) can be accepted up to limited concentrations (below 0.5 wt% total). It is worthwhile to discuss the effects of the alloying elements and impurities on the properties of Zr-based alloys. 2.2. Effects of alloying elements on Zr-based alloys Pure Zr has two distinct type of crystal structure, hexagonal close-packed (HCP) in a-phase, and body-centered cubic (BCC) in b-phase and the a-to-b transformation temperature is 850 °C (Alam et al., 2011) (other different value, 865 °C, also was reported (Banerjee and Banerjee, 2016; Lemaignan, 2012)). Further, a twophase (a + b) regime exists in the temperature range of 850– 950 °C in Zr-based alloys. Zr-based alloys remain in a phase at the normal plant operation temperature (below 400 °C) (Azevedo, 2011). However, it is possible that phase transient conditions occur in the accident and deteriorates the accident continuously. The alloying elements could be divided into a stabilizer and b stabilizer. O and Sn, having high solubility in a phase and stabilize it at high temperature (Fig. 1), are able to raise the a-tob transformation temperature, while b phase is stabilized by the addition of Fe, Cr, Ni and Nb (Fig. 2) (Lemaignan, 2012). As an alloying element, O was added to Zr-based alloys to increase the yield strength by solution strengthening and its level in the range of 0.11–0.16% (all wt% in this review) is recommended

Z. Duan et al. / Nuclear Engineering and Design 316 (2017) 131–150

133

Fig. 1. Zr-O and Zr-Sn binary phase diagram (Lemaignan, 2012).

Fig. 2. Zr-Fe (a), Zr-Cr (b), Zr-Ni (c) and Zr-Nb (d) binary phase diagram (Lemaignan, 2012).

in most of Zr-based alloys due to the detrimental effects on the ductility of the Zr-based alloys, while nitrogen (N) is suggested to remove as much as possible because it deteriorates the corrosion resistance if its level exceeds about 0.07% (Banerjee and Banerjee, 2016). Sn is a good corrosion inhibitor, and most effective in improving corrosion resistance in N-containing Zr alloys, especially by mitigating the adverse effects of N in high temperature, without seriously affecting the neutron economy. Additionally, Sn also has some impacts on the mechanical properties by increasing the yield strength and maintaining good creep properties of the Zr-based

alloys (Lemaignan, 2012). The addition of Sn makes contributions to the increase of creep resistance by lowering the stacking fault energy, thereby increasing interaction between solid solution alloying elements and extending dislocations (Pahutová et al., 1985), while another different mechanism was proposed that Sn addition reduced the self-diffusion coefficient of Zr matrix which resulted in the increment in creep strengthening (Cheol et al., 2000). Nevertheless, the high Sn content (higher than 1.5%) was observed to deteriorate the corrosion resistance of zircaloy-2 (Zry-2), thus, the Sn content is reduced to the amount just necessary to counteract the N levels(Krishnan and Asundi, 1981).

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The important effects of Fe, Cr, and Ni additives on the corrosion behavior of Zr-based alloys were found in a fact that the zircaloy-1 (Zry-1) was contaminated by stainless steel during the fabrication process, which proved these impurities are beneficial to the corrosion resistance at high temperature. Fe and Cr, possessing atomic sizes considerably different from that of Zr, have very low solid solubility and for the same level of solute content the strengthening effects of these solutes are much larger than that of Sn (Banerjee and Banerjee, 2016). The addition of Cr is not only used to improve corrosion resistance, it was also reported Cr addition enhanced the creep resistance and greatly improved the tensile properties of Zrbased alloys (Jung et al., 2010; Zhang et al., 2016). In addition, it seems that the presence of Zr(Fe,Cr)2 precipitates act as local electric shortcuts favoring electronic conduction, thus promoting hydrogen evolution at the oxide–water interface and resulting in a low hydrogen pickup fraction while Ni has a detrimental effect on hydrogen pickup so its amount is kept within the permissible limit (Couet et al., 2015). The sum of Fe, Cr and Ni must be controlled between 0.18 and 0.38% (Alam et al., 2011). Nb is highly soluble in b-Zr (see Fig. 2) and has a reasonable solubility in a-Zr, which is about 1.1% at 600 °C (Bethune and Williams, 1969). Therefore, it is favorable to mechanical treatment, and thus, the parameters of the alloys can be varied widely by heat treatment. Additionally, Nb has beneficial effects similar to Sn, for instance, it also offsets the adverse effects of N, C (carbon), Al, Ti (titanium) and other impurities, and more importantly, it is far better than Sn as a hardening agent, inferior only to Al and Mo (molybdenum), and substantially reduces the absorption of hydrogen (Zaimovskii, 1978). Further, Nb addition in Zr-based alloys improves the mechanical properties as well as retains a relatively low thermal neutron capture cross-section (see Table 1) (Charit and Murty, 2008). As a b-stabilizer, the effects of Mo addition on the mechanical properties and corrosion resistance of Zr-based alloys have also been reported widely (Cheol et al., 2000; Chun et al., 1999; Isobe and Matsuo, 1991; Lee et al., 2001; Pahutová et al., 1985; Yang et al., 2015). The creep rate of Zr-based alloys can be lowered by Mo addition (Pahutová et al., 1985), but its creep strengthening effect is lower than Nb, Sn and Cr (Cheol et al., 2000). Moreover, it was pointed out that the addition of Mo enhanced the yield strength and the ultimate tensile strength (UTS) both at room temperature and at high temperature (Chun et al., 1999), and that Mo addition is more effective than Sn, Nb, and Cr in strengthening the Zr-based alloys (Yang et al., 2015). On the other hand, it is worth noting that a higher strength increment is gained accompanied by the sacrifice of ductility by minor Mo addition (Yang et al., 2015). In addition, since Mo is generally harmful to corrosion resistance and has relatively higher thermal neutron absorption crosssection, the addition was previously recommended to be limited to 0.1% (Isobe and Matsuo, 1991). On the contrary, Lee reported that the corrosion resistance of Mo-containing Zr-Sn-Nb-Fe alloy was comparable to the Mo-free alloy if the alloy was properly heat treated (Lee and Hwang, 2003). Si (silicon) also has low thermal neutron absorption crosssection, and trace Si addition can provide an increase in the oxidation resistance (Chen et al., 2015a,b). Si could easily precipitate in the form of second phase particles (SPP) in Zr-based alloys because of its considerably low solubility in a- or b-Zr. Accordingly, the tensile strength of Zr-based alloys increased with the increment in Si content, meanwhile, the alloys’ resistance to oxidation could be reduced if the Si content is lower or higher than the optimum value (the value turned out to be 0.01% in Ref. (Hong et al., 2002) or 0.008%
0.014% in Ref. (Sell et al., 2006)). The decrease in oxidation resistance induced by higher Si content is probably ascribed to the formation of Zr3Si precipitates in the b range (Sell et al., 2006). Additionally, it was also reported that Si addition (0.5%

and 1.5% Si) enhanced the creep resistance by the formation of the eutectic phase which retarded the dislocation motion at grain boundaries (Matsunaga et al., 2014). However, the effects of the Si addition to the Zr-based alloys are quite complicated, and have not yet been understood clearly. Al is a very effective a-solution strengthener due to its high solubility in a-Zr and also possesses low thermal neutron absorption cross-section, while Al addition in a-Zr is not acceptable because of its detrimental to the corrosion resistance of a-Zr in water (Banerjee and Banerjee, 2016; Krishnan and Asundi, 1981; Northwood, 1985; Williams et al., 1972). Except for these alloying elements mentioned above, some other minor alloying elements, such as Cu (copper) (Hong et al., 2000, 2001), V (vanadium) (Sabol et al., 2000), Ge (germanium) (Xie et al., 2014), etc. have also been examined to improve the performance of Zr-based alloys to be capable in advanced nuclear reactors with higher coolant temperature, longer fuel cycle and higher fuel burnup. 2.3. The development of Zr-based claddings Up to date, although many Zr-based alloys have been studied for potential usage in nuclear reactors, only a few of them are of commercial importance, and new Zr-based alloys are continuously developed and then are recommended to be candidates for the advanced nuclear power plant, as shown in Table 2. However, the development of advanced fuel claddings is a persistent issue in the nuclear power area and many investigators have been developing new alloy compositions or exploring new processes for improving performances of Zr-based alloys. 2.3.1. Zircaloys Unalloyed Zr shows poor corrosion resistance and deteriorates further with the increment of the N content. 2.5% Sn, as a good compromise between corrosion resistance, strength, and manufacturability, was added to unalloyed Zr to mitigate the adverse effects of N. This alloy was designated as Zry-1 (Krishnan and Asundi, 1981). However, Sn addition could not exert its full beneficial effect unless Fe, Ni, or Cr were present. Furthermore, the corrosion rate of Zr-Sn alloys decreases with decreasing Sn content (Kass, 1962). Accordingly, the Sn content was reduced to 1.8% on the basis of Zry-1 and Fe, Ni, and Cr were added to enhance the corrosion resistance of the binary alloy, which resulted in the development of Zry-2. The optimum values and finally fixed values of Sn, Fe, Cr and Ni in Zry-2 are shown in Table 3. In order to further improve the corrosion resistance or the manufacturability of Zry2, three prototype alloys, namely Zry-3A, 3B, and 3C (chemical compositions are listed in Table 2), were developed. The Zry-3 was not accepted by the commercial power plants due to their lower strength (about 75% of that of Zry-2) and it is much more complexed to minimize the stringer effect in Zry-3. But the application of vacuum arc-melting technique developed during this period promoted to minimize the occurrence of stringers, thereby increasing the corrosion resistance of Zry-2. Hydrogen absorption during corrosion in aqueous environments was one of the earliest observations made during the study of Zr alloy corrosion (Cox, 2005). The hydrogen ions produced by the oxidation of Zr with water can permeate into the Zr metal. What’s more, during the service period, a significant amount of hydrogen gas was produced, especially in the LOCA by the following exothermic reaction.

Zr þ 2H2 O ! ZrO2 þ 2H2

ð1Þ

This reaction becomes significant at the temperature above 1200 °C, and could be the dominant heat source because it is a heat production process. Hydrogen precipitates out as hydrides because of the low solubility of hydrogen (80–100 wt.ppm) at light water

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Z. Duan et al. / Nuclear Engineering and Design 316 (2017) 131–150 Table 2 Composition* (weight%), material condition of various Zr-based alloys.

* **

Alloys

Sn

Fe

Cr

Nb

Others

Condition**

Remarks

Zry-1 Zry-2 Zry-3A Zry-3B Zry-3C Zry-4 E110 E125 E635 E110 opt. E110M E635M E125 opt. ZIRLO OPT ZIRLO M5 MDA M-MDA NDA S2 HiFi

2.5 1.5 0.25 0.5 0.5 1.5 – – 1.10–1.30 – – 0.70–0.90 – 1.02 0.60–0.80 – 0.80 0.5 1 0.8 1.5

0.12 0.25 0.4 0.2 0.2 – – 0.3–0.4 0.025–0.07 0.07–0.15 0.30–0.40 0.025–0.05 0.1 0.11 0.05 0.2 0.4 0.3 0.3 0.4

– 0.1 – – – 0.1 – – – – – – – – – 0.015 0.10 0.3 0.2 0.1 0.1

– – – – – – 0.95–1.05 2.20–2.60 0.95–1.05 0.90–1.10 0.90–1.10 0.70–0.90 2.4–2.7 1.01 1.02 1 0.5 0.5 0.1 0.1 0.08

– 0.12 O,0.05 Ni – – 0.2 Ni 0.09–0.13 O 0.10 O 2.20–2.60 O 0.05–0.12 O 0.06–0.099 O 0.10–0.15 O 0.04–0.10 O 0.06–0.099 O 0.09–0.15 O 1.04 O 0.09–0.13 O 0.12 O – 0.12 O, 0.01 Ni – –

– – – – – – RX – RX RX – RX – SR/RX PR RX SR SR/RX SR RX RX

Not suitable BWRs

AXIOM alloys X1 X2 X4 X5 X5A

0.3 – – 0.3 0.45

0.05 0.06 0.06 0.35 0.35

– – 0.25 0.25 0.25

0.7–1 1 1 0.7 0.3

0.12Cu, 0.2 V

80%PR RX 80%PR 50%PR 50%PR

PWR

J-Alloys J1 J2 J3

– – –

– – –

– 0.1 –

1.8 1.6 2.5

–

PWR

–

RX RX RX

Quaternary Alloys (Zr1NbSnFe) by AREVA, abbreviated as QA1, QA2, QA3 for simplicity QA1 0.5 0.1 – 1 – QA2 0.3 0.1 – 1 – QA3 0.3 0.2 – 1 –

RX RX RX

PWR

HANA alloys HANA-3 HNAN-4 HANA-5 HANA-6

PR PR PR PR

PWR

Used since

PWR PWR/RBMK/VVER 1997 PWR/RBMK/VVER PWR/RBMK/VVER

2009

PWR PWR PWR PWR PWR PWR PWR BWR

1991 2008 1998 2004 2004 2004

1996

0.08 Cu 0.05 Ni

1997 0.4 0.4 0.8 –

0.1 0.2 0.35 –

– 0.1 0.15 –

1.5 1.5 0.4 1.1

0.1 Cu 0.12Cu, 0.2 V 0.1 Cu 0.05Cu

Remainder zirconium. SR: Stress relieved, PR: Partially recrystallized, RX: fully recrystallized.

Table 3 The optimum and finally fixed values of the alloying element contents in Zry-2 (Hong et al., 2000). Alloying Elements

Optimum Values wt %

Fixed Values wt%

The reasons for fixed values

Sn

1.5

1.5

Fe

0.22

0.12

Cr

0


0.1

Ni

0.25

0.05

Be adequate to counteract the deleterious effect of nitrogen The total alloy content should be kept to a minimum to ensure easy fabrication Avoid the excessive increase in the hardness of the alloy The total alloy content should be kept to a minimum to ensure easy fabrication

reactor (LWR), which causes local swelling in the Zr-based alloys (Barrett et al., 2012). Zry-1 and Zry-2 absorb hydrogen during corrosion and the quantity of hydrogen uptakes is proportional to the amount of corrosion. Subsequently, the Ni-free Zry-2 were investigated, and the Zry-4 also was developed by increasing the Fe content on the basis of Ni-free Zry-2 to compensate for the removal of Ni (Kass, 1962). Zry-2 exhibits superior corrosion resistance at elevated temperatures in steam because Fe, Cr, and Ni, singly or in combination,

enhance steam corrosion resistance, which makes it best suit for the boiling water reactor (BWR) conditions. Meanwhile, Zry-4 is recommended in the pressured water reactor (PWR) conditions because of the facts that Zry-4 provides almost similar mechanical properties but the absorption of hydrogen by Zry-4 is lower than that of Zry-2 (Whitmarsh, 1962). Consequently, Zry-2 and Zry-4 had become the standard fuel cladding materials for BWRs and PWRs, respectively from the early 1970s. To improve U (uranium) utilization and reduce nuclear fuel cycle cost, low Sn Zry-4 (1.32– 1.44%) was found to possess superior corrosion resistance in-PWR without significant influence on hydrogen uptake (Garde et al., 1994). Subsequently, Zry-4 and low Sn Zry-4 have become the basis and reference for other Zr-based claddings. 2.3.2. E110, E125 and E635 alloys Instead of Sn addition, Russian focused more on the effects of Nb addition, which has beneficial effects similar to Sn addition. The E110, E125, and E635 alloys, whose chemical compositions are listed in Table 2, have been designed successfully in Russia for VVER (Vodo-Vodyanoi Energetichesky Reactor) (also denoted as WWER) and RBMK (Reaktory Bolshoy Moshchnosti Kanalniy) reactors in 1958, 1958, and 1971 successively (Novikov et al., 2011a). It was observed that both of E110 and E635 had been found to have high strength, creep resistance, and resistance to

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Fig. 3. Creep and irradiation growth of E110 and E635 (Markov et al., 2005).

radiation growth except for the differences in their characteristics that E110 had a high corrosion resistance in high-temperature pressurized water but corroded intensely in boiling water medium in the presence of Li (lithium), while, E635 provided good mechanical properties and corrosion resistance in boiling water whereas inferior corrosion resistance than E110 in water under pressure is concerned (Nikulina, 2004). The optimal solution for the VVER fuel assembly was proposed that E110 is used as a material for fuel rod claddings and E635 as a material of the skeleton components since E110 has higher uniform corrosion resistance but insufficient geometrical stability, which could be compensated by the advantages of E635 (Markov et al., 2005). However, E635 are also adequately competent as materials for claddings of VVER-1000, proved by superior strength and resistance to irradiation induced growth than E110 (Fig 3) (Markelov et al., 2005; Nikulina et al., 1996). E125 is employed for shroud pipes of VVER-440 and pressure pipes of RBMK (Nikulina, 2004). In comparison to the Zry-4, E110 seems to be more susceptible to breakaway which leads to enhanced oxidation and hydrogen uptake, and thus absorbs considerably higher hydrogen and exhibits a faster embrittlement (Hózer et al., 2008; Steinbrück et al., 2010). Additionally, the creep and growth behaviors of E110 may accelerate as the fuel burnup increased, hence, the current properties of E110 could not allow a guarantee to be given with full assurance of the reliability of fuel claddings for a new generation of VVER with higher fuel burnup (Shebaldov et al., 2000). E110 opt., E110M, E125 opt., and E635M were created for WWER-1200 (AES-2006), which has longer refueling interval (18–24 months) and higher unit power provided by both the changed core geometry and increased operation parameters than existing reactor WWER-1000 (Novikov et al., 2011b). To ensure the high serviceability of the items made from E110, E125, and E635 during conditions of extended burnup, all of them were updated by perfecting their chemical composition, respectively. As mentioned above, O could increase the strength of binary alloys. In addition, it was found that the presence of iron in O-bearing Zr-based alloys increases the resistance to nodule corrosion characterized by the formation of white pustules thicker than the uniform oxide film in steam and the resistance to radiation-induced growth (Nikulina et al., 1996). In terms of E110, Fe addition was employed which resulted firstly in the birth of E110 opt., and subsequently, E110 opt. was updated to E110M via optimizing the permissible ranges of the Fe and O content. E125 was modified to E125 opt. by optimizing the contents of O and Fe, and E125opt. features

higher strength characteristics and resistance to shape changes (Novikov et al., 2011a). On the basis of the E635, the E635M was designed having lower contents of Sn and Nb and the optimized Fe/Nb ratio (shown in Table 2) with a view of the enhancement of uniform corrosion resistance without degrading the mechanical performance such as strength and creep resistance (Markelov et al., 2005). Even though each of the modified alloys, E110M, E635M, and E125 opt. is superior to other in different areas, for example, E635M was found to have the highest thermal hoop creep resistance while the lowest corrosion resistance in steam at the temperature of 400 °C, all of them were considered to ensure the higher characteristics for the resistance to shape changes in comparison to E110 (Novikov et al., 2011a). 2.3.3. ZIRLO (Zr low oxygen) series alloys In response to the desire to obtain higher burnup by increasing the coolant temperature and refueling interval, thereby reducing the radioactive wastes, the Zr-Sn-Nb-Fe alloy with annealing treatment at 873 K for 8 h, called ZIRLO (Table 2), was developed on the basis of Russian E635 for PWR by Sabol et al. in 1987 (Sabol et al., 1989). By the means of long-term autoclave corrosion test and high burnup irradiation test in PWR environment, ZIRLO showed lower autoclave corrosion rate at high temperature, lower sensitivity against lithia attack, lower in-reactor corrosion rate and higher creep resistance as compared to Zry-4 (Sabol et al., 1989). Subsequently, the merits were confirmed further, which the in-reactor creep and irradiation growth of ZIRLO were approximately 80% and 50%, respectively lower than that of Zry-4 (Sabol et al., 1994). Additionally, except Zry-4, ZIRLO possesses better corrosion resistance, lower irradiation growth and superior creep properties for high burnup fuels than other Zr-Nb alloys, and has been in use in almost all Westinghouse-designed plants in the United States, as well as in several European countries by the late 1990s (Sabol, 2005). As pointed out previously, even though a decrease in Sn concentration increases uniform corrosion resistance, the Sn addition also is needed to provide mechanical property, such as increment in strength, and creep resistance for Zr-based alloys. Therefore, similar to the work undertaken to improve the in-PWR performance of Zry-4 by a reduction in Sn content, the efforts also have been done to improve the performance of ZIRLO since about 20 years ago through comparison the variations in the thermal heat treatments and chemical compositions, especially a decrease of Sn content (Comstock et al., 1996; Yueh et al., 2005). The impacts, on corro-

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Fig. 4. Variation in normalized (to 1%Sn-1%Nb-0.1%Fe) corrosion rate at 633 K pure water and 700 K steam environments as a function of chemistry variable: (a) Sn, (b) Nb, and (c) Fe (Comstock et al., 1996).

sion rate and SPP size, of alloying element content in normalized to 1%Sn-1%Nb-0.1%Fe associated to the changes of heat treatments were studied by Yueh et al. (2005). It has been found that corrosion rate decreased with decreasing Sn content while no significant correlations were observed with Fe and Nb variations (Fig. 4), and a threshold of about 0.6–0.7% Sn plus Fe is beneficial in significantly reducing the lithium-accelerated corrosion (Fig. 5). What’s more, cold work in addition to thermal aging was able to obviously enhance the precipitation of SPPs. Soon afterward, a new material, named optimized ZIRLO (denoted as OPT ZIRLO) with a lesser amount of Sn (0.66% or 0.77%) and partially recrystallized microstructure, was designed by Foster et al. (2008). They concluded that compared with ZIRLO, OPT ZIRLO showed higher corrosion and creep resistance by reducing Sn content in parallel to fabrication changes like final tube area reduction and final annealing temperature to compensate the decrease in in-reactor creep strength. 2.3.4. M5 alloys As has been described previously, a reduction of Sn content in Zr-based alloys results in a reduction of corrosion rate, therefore, modern trend to optimize existing alloys or to develop new alloys is to reduce Sn content. New Zr-based alloys with lower Sn content (1.3%) have been produced by AREVA and its industrial partners: quaternary materials (Zr, Sn, Fe and V, denoted as M4), (Zr, Sn, Fe and Nb, including Alloy 3, Alloy 1), and ternary materials (Zr, Nb,

Fig. 5. Impact of Sn and Fe content on corrosion rate at 633 K 77 ppm lithiated water environment (Comstock et al., 1996).

and O, including Alloy 2 and Alloy 5 denoted as M5). But anyway, only M5 has been applied as a cladding material in a wide range of PWR environments at burnups up to 78 GWd/Metric Ton of Uranium (MTU) (Doriot et al., 2005). M5, developed on the basis of Russian material E110 in the 1980s, is a Zr-Nb alloy devoid of Sn with a controlled O, Fe and S (sulfur) content and is distinguished by a homogenous highly refined dispersion of b-Nb precipitates in fully recrystallized microstructure which is primarily responsible for the advantageous creep resistance (Mardon et al., 2010). The absence of Sn, the controlled alloy chemistry, and the optimized heat treatment result in a very high corrosion resistance and superior mechanical properties in high burnup and high dose irradiation conditions. The O shows the beneficial effect on thermal creep in the range of 900–1800 ppm, while C must be held at a relatively low level due to its adverse influence on corrosion in steam at 400 °C. Similarly, the content of S is optimized to be lower than10 ppm, since deliberate addition of S is able to provide an improved high-temperature mechanical strength of M5 cladding without degrading the corrosion resistance (Mardon et al., 2000). What’s more, a significant improvement in creep resistance could be obtained by a low temperature (600 °C) process (Gilbon et al., 2000). Variations in chemical composition and heat treatment made M5 different from other Zr-based alloys supplying much more superior properties for M5. Many literatures (Brachet et al., 2002; Le Saux et al., 2013; Thomazet et al., 2005; Garat et al., 2012) reported that M5 had been considered a better choice than Zry-4 thanks to its impressive improvements in corrosion resistance, creep resistance and lower hydrogen pickups out of the pile and in the pile. It also has been pointed out that M5 shows a very good behavior in LOCA and RIA (reactivity insertion accident) conditions and outperforms Zry-4 and other Zr-based alloys (even Zr-1%Nb alloys) at the same burnup level due to its lower hydrogen pickup (Mardon and Dunn, 2007). Consequently, M5 has been widely granted by US, UK, Chinese etc., safety authorities and by the end of 2011, over 4.5 million M5 claddings had been utilized in 94 commercial PWRs to fuel rod burnup of 80 GWd/MTU (Garat et al., 2012). 2.3.5. X5A alloys Garde et al. concluded the Sn reduction showed beneficial effects on improvement of corrosion resistance by the investigation on the effects of variation in Sn content, which resulted in the low Sn Zry-4 (OPTIN) (Garde et al., 1994). On the basis and reference of low Sn Zry-4, they found that the Zr-based alloys with dilute tin (about 0.5%) possessed excellent uniform waterside cor-

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rosion resistance necessary for the PWR fuel application, since corrosion rate was reduced by decreasing Sn content with the concurrent adjustments of other components, such as the addition of Fe, Cr, Nb, and O (Garde et al., 2002). In this way, Alloys A (Zr-0.35%Fe0.25%Cr-0.30%Nb-0.5%Sn-O), Alloy C (Zr-0.50%Fe-0.24%Cr-0.4%SnO), and Alloy E (Zr-0.45%Fe-0.24%Cr-40%Nb-0.7%Sn-0.03%Ni-O) were designed, and results revealed that Alloy C had superior corrosion resistance than Alloy A whereas its drawbacks, referring tube fabrication difficulties and higher diametrical strain, made Alloy A much more outstanding (Garde et al., 2002). Consequently, Alloy A was the most promising material with integrated properties such as improved in-PWR corrosion resistance, good tube manufacturability, and then was renamed as X5A and included as one of the five candidate AXIOM alloys (referred as X1, X2, X4, X5, and X5A) which are being developed by Westinghouse for both high-burnup and high-duty PWR (Garde et al., 2010). The fabrication process of original X5A fabricated by high temperature processing (HTP X5A) was optimized through the application of a low-temperature processing appropriate (LTP X5A) for Niobiumbearing alloys, and X5A has been considered as a promising alloy for PWR with burnup about 50 GWd/MTU by the comparison with ZIRLO (Garde et al., 2010). 2.3.6. Other Zr-based alloys In Japan, MDA (Mitsubishi developed Alloy) and NDA (New Developed Alloy) with burnup limit up to 55GWd/MTU were created by Mitsubishi and Nuclear Fuel Industries (NFI) for reloads since 2004. MDA was developed by Mitsubishi via keeping chemical composition as close to Zry-4 as possible while reducing Sn content to improve corrosion resistance, as well as the Nb addition to compensate the decrease of mechanical strength and to reduce the hydrogen pickups. On the other hand, according to the discover that the Nb content in the range of 0.1–0.2% results in lowest corrosion rate for Zr-Sn-Nb-Fe-Cr alloys, NFI and Sumitomo Metal Industry (SMI) designed the NDA having 0.1% Nb, 1% Sn, and 0.4% (Fe + Cr) (Garzarolli et al., 2011). For the future high duty fuel in Japanese PWRs with burnups exceeding 55 GWd/MTU, the new fuel cladding materials Modified- Mitsubishi Developed Alloy (MMDA) and S2 alloys have been developed on the basis of MDA and NDA by reducing Sn content similar to OPT ZIRLO. M-MDA, with two type of manufacturing processes, Stress-Relieved (SR) and Recrystallized (RX), was found to inherit an excellent performance from MDA meanwhile corrosion resistance were improved by optimization of alloying elements, and will be further evaluated in commercial reactors (Fujii et al., 2009). Moreover, J-Alloys (referred as J1, J2, J3 (Motta et al., 2015), chemical compositions are shown in Table 2) are being developed for burnups up to 70 GWd/MTU, and the irradiation of test rods has started since 2006 (Garzarolli et al., 2011). Higher hydrogen uptake at high burnup strongly affects the embrittlement of BWR fuel cladding, therefore, for the development of BWR fuel claddings, a new BWR fuel cladding material named HiFi (High corrosion resistance and high Fe Zr alloy), a Zry-2-based alloy with Fe content increased up to 0.4% (shown in Table 2 (Adamson et al., 2003)), is being developed by NFI (Kakiuchi et al., 2006a). Out-of-pile and in-pile tests have been carried out in Japanese and a European commercial BWRs for screening candidate materials, results of which confirmed that HiFi showed lower hydrogen absorption in both out-of-pile and in-pile tests than Zry-2, and did not exhibit accelerated hydrogen pickup in the high burnup region (Kataoka et al., 2012). It was reported that the superiority of HiFi in hydrogen uptake resistance is contributed to the higher Fe addition which decreased the electrical potential gradient over the oxide film, thereby retarded the proton diffusion in the oxide film (Kakiuchi et al., 2006b). As an evolutionary development of M5, ultra-low Sn quaternary alloys Zr1NbSnFe were developed by AREVA with the addition of

small amounts of Sn (0–0.65%) and Fe (0.03–0.35%), and controlled S and O contents, respectively of 20 and 1400 ppm. Sn addition is to increase the irradiation creep strength, but the limited content is employed to mitigate corrosion resistance degradation. Fe content was limited to 0.35% to improve corrosion resistance (Chabretou et al., 2012). The chemical compositions of three typical Zr1NbSnFe were described in Table 2. Irradiation of Quaternary Alloys (Zr1Nb0-0.5Sn0.1-0.2Fe) was started in 2003, and a second irradiation campaign was launched in 2008 in a French and a German PWRs (Trapp Pritsching et al., 2012). The preliminary result revealed the addition of Sn improved diametrical irradiation creep resistance while inducing a degradation of oxidation resistance when Sn content was higher than 0.3% (Chabretou et al., 2012). The research and development of the advanced fuel to meet the global demand for the extension of the fuel discharge burnup to more than 70GWd/ MTU was started in 1997 in South Korea (Jeong et al., 2006), By the systematical investigation on the effects of alloying elements (Nb, Sn, Fe, Cr, Cu, Mn, V, Sb, Ag, Te, and Ge) and heat treatment on the corrosion of the alloys, the Zr-Nb alloys named HANA (High-performance Alloy for Nuclear Application) were being developed, of which HANA-3, 4, 5 and 6 were evaluated in and out of pile environments and their chemical compositions were listed in Table 2 (Kim et al., 2012). Compared with other Zr-based alloys, Cu addition to Zr-based alloy with a high Nb content was a new attempt, and what’s more, it revealed that Cu addition is of benefit to the fine distribution of b-Nb, which promoted the stability of tetragonal zirconia (t-ZrO2), thereby increasing the corrosion resistance of Zr-Nb alloys (Park et al., 2005). Some delighted improvements over Zry-4 have been reported, such as superior corrosion resistance, higher strength and elongation, lower hydrogen absorption in out-of-pile and in-pile environment (Jeong et al., 2006; Kim et al., 2014). Consequently, HANA claddings have been regarded as promising fuel claddings to improve the safety and the operation economy of the nuclear fuel. In China, new Zr-based alloys for PWR fuel cladding also have been under development with the purpose to increase the discharge burnup up to 60 GWd/MTU. Compared with Zry-4, the superior performance of two new alloys (N18 and N36) containing Sn, Nb, Fe and Cr, in the out-of-pile tests were reported (Zhao et al., 2005). Overall, the element, Zr, was discovered in 1789 by Klaproth, isolated as a metal by Berzelius in 1824, highly purified by Van Arkel in 1925, selected firstly as structural material for nuclear reactors of submarines in 1949, and then, became the main cladding materials for water-cooled reactors replacing stainless steels in the middle of 1960s (Northwood, 1985), followed by the development of Zr-based alloys for nuclear reactors. In the history of Zr alloy development, Sn, Fe, Cr, and Ni were firstly used by U.S. as alloying elements to improve the corrosion resistance and mechanical properties of Zr-based alloys, resulting in Zry-2 and 4. Subsequently, E110 alloy was created by Union of Soviet Socialist Republics by Nb addition. As the trend of nuclear reactor development, continues improvements of Zr-based alloys are urgently needed to meet the requirements of growing burnup, which have been driving the optimization of existing alloys and developments of new alloys by modifying alloying elements (e.g., Sn, Nb, Fe, Cr, Ni, Cu, V, Si, etc.) and heat treatments. All candidate materials should be verified through out-of-pile tests, research reactor tests, and in-pile pathfinder rods before putting into service. As described above, almost evaluation tests revealed an improvement on the modified or developed alloys, however, it is considered unlikely that the relatively small alloying additions in Zr-based alloy can be further modified beyond the impressive improvements already achieved in order to achieve the needed additional
100fold reduction in high-temperature oxidation resistance (Zinkle et al., 2014). Consequently, on the other hand, the development

FeCrAl Zhong et al. (2016)

 



Hot steam oxidation at 700 °C and normal BWR condition

Ni, glass lubricant Valeeva et al. (2012)

Zry-2

Magnetron sputtering

0.3–1.3

   

Air annealing at 700-1000 °C E125

Cr Brach et al. (2015)

Chemical deposition

30–40

  

Zry-4

Cr Kim et al. (2015)

Physical vapor deposition

3D laser cladding

Ti Baczynski (2014)

Zry-4


20

Autoclave corrosion test, steam test at 1000-1200 °C

No peeling phenomena were observed in mechanical and corrosion tests. Cracks were formed on coatings in the tensile and compression tests. Coated alloy showed better corrosion resistance than non-coated alloy. The Cr coating was characterized by a fully dense microstructure without defects. Cr-coated alloys exhibited significantly higher post-quenching residual strength and ductility due to their much lower oxidation kinetics. Hydrogen uptake was decreased drastically by Cr coating in the steam test. Some important issues, especially neutron irradiation, should be investigated. Single Ni-layer coating and Ni-layer with glass lubricant as outer layer coating were investigated. The highest efficient protection of Ni-layer coating could be kept up to 800 °C. Oxide and defective layers were not formed in double-layer coated alloys. Four FeCrAl film compositions were investigated. FeCrAl with higher Al content promoted alumina formation, inhibiting oxidation of base metal and the formation of zirconia which is an effective barrier to cation transport but ineffective to impede oxygen diffusion. Alumina growth was accompanied by a counter flow of vacancies into the FeCrAl, driving porosity. No FeCrAl film degradation was observed in autoclave test. A thicker coating is required to promote oxidation resistance of Zry-2.      Adhesion tests, immersion at 360 °C, 18.9 MPa.

 Ti began to have a positive effect on preventing oxidation at high temperature when the thickness of coating layer was increased to between 42 nm and 64.5 nm. Hot steam oxidation at 700 °C

0.0215, 0.043, 0.0645 80 Magnetron sputtering

Experimental conditions Thickness (lm) Fabricating methodology Cladding material Coating materials Author (year)

Table 4 Studies on coatings on Zr-based alloys.

During the development of advanced fuel claddings for LWRs, the waterside corrosion of the fuel claddings is generally recognized as one of the main limitations to burnup extension of nuclear fuels, reducing operator refueling downtime (Zhao et al., 2005). Moreover, accelerated hydrogen uptake in the claddings was observed at high burnups, which is also one of the most important issues limiting high burnup fuel performance from the viewpoint of cladding integrity (Une et al., 2011). As introduced above, many new Zr-based alloys with higher resistance to uniform corrosion and hydrogen uptake have been created by optimizing the metallurgical composition and/or manufactural process of the Zr-based alloys, which, however, was considered that the likelihood of further improvements making a significant difference was limited (Barrett et al., 2012). Currently, the coating technology has been widely applied in cladding to increase water corrosion and wear resistance due to its outstanding profits. The major benefit is the economics as the resistances can be improved using a coating on existing Zr-based claddings without the necessity to modify the base materials, contributing to the possibility for commercial application in the very near term (5 years or less). Additionally, coated Zr-based alloys are characterized by superior properties i.e. higher melting point, lower hydrogen absorption and generation at elevated temperature, thereby mitigating severe accident consequences (Barrett et al., 2012). However, there are also significant disadvantages: the increment on the mechanical strength at higher temperature could not be obtained by coating techniques, as well as that the resolution of the adhesion properties of Zr-based alloy, phase stability up to high temperature, thermal expansion coefficient, neutron economy, thermal conductivity, irradiation susceptibility and tube manufacturability are hindering the developments of coated Zr-based cladding (Kim et al., 2016). The relevant coating studies (Alat et al., 2015; Ashcheulov et al., 2015; Baczynski, 2014; Brachet et al., 2015; Daub et al., 2015; Jin et al., 2016; Khatkhatay et al., 2014; Kim et al., 2015a; Kuprin et al., 2015; Maier et al., 2015; Rezaee et al., 2013; Valeeva et al., 2012; Wiklund et al., 1996; Zhong et al., 2016) on Zr-based claddings were summarized in Table 4. According to Table 4, various materials were selected for coatings, whose thicknesses differed each other significantly, and were deposited on Zr-based alloys by different coating techniques. The fabricating methodologies were generally determined by the materials chosen for coatings in order to obtain high adhesion to the substrate. Therefore, both of them confine the choices of potential materials and fabrication in coating technology for nuclear fuel claddings. In respect to the coating techniques, thermal spray, physical and chemical vapor deposition (PVD and CVD) are widely investigated for cladding coatings. Thermal spray is a coating process in which melted or heated particles (1–50 lm) of materials (feedstock) by electrical (plasma or arc) or combustion flame means, are sprayed onto the substrate and condensed to form a coating layer, which is pictured as Fig. 6 (Herman et al., 2000). There are several thermal spray methods, such as cold spraying (CS), high-velocity oxy-fuel spraying (HVOF), two-wire electric-arc spray and plasma spray, etc. of which the coatings fabricated by CS (Maier et al., 2015) and HVOF (Jin et al., 2016) have been evaluated for cladding coating. PVD covers a broad class of vacuum coating processes, in which materials are physically removed from a source by evaporation or sputtering, transported through

Remarks

3. Coatings on Zr-based alloys

Zry-2

of new oxidation-resistant alloys with a markedly different chemical composition and application of coatings have been employed as approaches to reduce oxidation rate.

(continued on next page)

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140

Table 4 (continued) Coating materials

Cladding material

Fabricating methodology

Thickness (lm)

Experimental conditions

Maier et al. (2015)

Ti2AlC

Zry-4

Cold spray


90

Scratch test, air oxidation tests at 700 and 1005 °C

Alat et al. (2015)

Ti1-xAlxN (0.54 < x < 0.67)/ TiN

ZIRLO

CA-PVD

4–12

Immersion, 360 °C, saturation pressure.

Khatkhatay et al. (2014)

TiN/Ti0.35Al0.65N

Zry-4

Pulsed laser deposition

1

500 °C, 25 MPa dearated water

Daub et al. (2015)

CrN, TiAlN, AlCrN

Zry-4

Physical vapor deposition

2–4

Autoclave corrosion test, 1100 °C steam test and hydrogen ingress test.

Remarks              

Rezaee et al. (2013)

YSZ

Zry-4

Sol-gel process

0.97

Heat treatment temperature, electrochemical measurements, oxidation test.

Jin et al. (2016)

Cr3C2-NiCr

Zr-2.5Nb

HVOF

250

Autoclave corrosion test, air and steam test from700 to 1000 °C

Ashcheulov et al. (2015) Wiklund et al. (1996)

Polycrystalline diamond

Zry-2

0.3

Hot steam oxidation at 950 °C, impedance spectroscopy

Multilayered Ti/TiN

Zry-4

Microwave plasma enhanced linear antenna chemical vapor deposition apparatus Physical vapor deposition

1.0, 2.0, 3.5, 3.7

Autoclave corrosion test, hydrogen ingress test.

Kuprin et al. (2015)

Multilayered Cr-Zr/Cr/Cr-N

E110/ Zr1Nb

Vacuum-arc plasma separation system.

7

Air tests at 660, 770, 900, 1020, 1100 °C.

     

        

YSZ: Yttria Stabilized Zirconia. RT: room temperature. HT: high temperature. HVOF: high-velocity oxygen-fuel spraying.

The coating possessed a higher hardness (800 Hk) than Zry-4 (
180 Hk). The coating exhibited high density and was well adhered to the substrate. Well protection of coating to the substrate was observed at elevated temperature. Thinner coating (
30 lm), autoclave test is needed to be further undertaken. Ti layer as bond coating exists between the ceramic coating and ZIRLO. Better corrosion resistance and well adhesion were obtained. Al depletion occurred in the TiAlN coated samples. A TiAlN layer with an outer TiN layer was recommended. Non-coated alloy undergone most severe oxidation TiN coating performed better than the Ti0.35Al0.65 N coating. Further optimization of the film deposition parameters is needed to reduce the susceptibility of Ti0.35Al0.65 N to oxidation at high-temperature water. AlCrN was dissolved mostly in autoclave test and cracked in steam test due to its poor adhesion to the substrate. Without a breakdown, CrN coating provided the strongest resistance to corrosion and hydrogen ingress. Located aggressive corrosion beneath the coatings was observed in steam test caused by the coating cracking. Further work was proposed to be done on the study of scratched or damaged coatings. The surface quality and corrosion properties of the coatings strongly depended on the heat treatment temperature. Cracks occurred on coatings dried at lower temperatures. The higher corrosion protection was achieved with drying at 700 °C. The coating with many micropores was well-boned to the substrate. The coated alloys exhibited lower weight gain in HT air while higher in autoclave test than non-coated alloys due to the accelerated corrosion of substrate by the coating in HT water. Cracks were observed at the interface in the autoclave test. The coating was found to be of high quality without defects and contaminations. A higher resistance and thinner oxide layer was formed on coated alloy. Coating layer blocked the hydrogen diffusion into Zry-2. Multilayered Ti/TiN and a TiN single layer were investigated. No peeling occurred on any multilayered coating, but the single layer coating showed minor peeling. All coated alloys showed much less hydriding than the non-coated alloy. Zirconium alloy layer, Cr2Zr layer, Cr layer and CrN layer were deposited on the substrate successively. The coating was considered to provide high wear resistance due to its high hardness (27 GPa) resulting from the absence of defects and pores.

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Author (year)

Z. Duan et al. / Nuclear Engineering and Design 316 (2017) 131–150

141

Fig. 6. Schematic overview of the thermal spray process (updated from Rf. (Herman et al., 2000)).

a vacuum or partial vacuum and then condensed as a film (Moarrefzadeh and Branch, 2012). According to the two methods: thermal evaporation and sputtering, several variations of PVD techniques exist, for instance, cathodic arc deposition, electron beam PVD, evaporative deposition, pulsed laser deposition, and sputter deposition (magnetron sputtering) (Shahidi et al., 2015). In comparison to the PVD, CVD involves the chemical reactions in the vapor phase and on the substrate, and the reactions can be promoted or initiated by heat, higher frequency radiation or a plasma, which contributes to the variety of CVD, such as plasmaenhanced CVD (PECVD), metal-organic CVD (MOCVD) (Jones and Hitchman, 2009). The mature experience has existed in thermal spray, PVD, CVD and other coating techniques (sol-gel, hot dipping, impregnation, etc.), however, the appropriate method and parameters suitable for specific coating materials still need further investigation and optimization to achieve the desired coating. Therefore, the priority should be given to the candidate materials for cladding coatings. From the viewpoint of materials for Zr-based cladding coatings, as same as the standards of the selection for alloying elements, Fe, Cr and Ni have been added to Zr-based alloys due to their relatively low thermal neutron cross section absorption, and beneficial effects on corrosion resistant properties at elevated temperature, they are the candidate materials for the coatings to reinforce the corrosion resistance of Zr-based cladding by coating. Cr has been deposited on Zry-4 by the techniques such as laser beam scanning (Kim et al., 2013), three-dimensional (3D) laser coating (Kim et al., 2015a) and PVD (Brachet et al., 2015). All results confirmed that Cr was well-boned to the substrate without peeling or spalling, and the Cr coating contributed to improve the corrosion resistance at elevated temperature, although the brittle characteristic of the Cr element as well as the influence of the laser treatment resulted in the formation of cracks on Cr coating in tension and compression (Kim et al., 2015a). Monolithic FeAlCr coating layer showed no susceptibility to substantial oxidation or corrosion at high temperature contributed from the formation of a protective alumina oxide film in addition to the chromia formation which is effective as well as alumina (Zhong et al., 2016). Ti (Titanium) was considered to have overall advantages in thermal neutron absorption cross-section, melting temperature, thermal conductivity and phase stability at high temperature, which led to the selection of Ti as the materials for coatings on Zr-based alloys (Baczynski, 2014). Titanium nitride (TiN), a refractory metal nitride ceramic with high hardness, high melting temperature, high thermal conductivity, and excellent corrosion and erosion resistant properties, has been demonstrated to reduce the fretting damage of Zry-4 in

PWR (Khatkhatay et al., 2014). However, the fracture toughness and oxidation resistance of TiN coatings were not able to satisfy the requirements of many advanced engineering applications, and were improved with Al incorporation resulting in TiAlN, whilst it has been noted that the presence of Ti would perturb the formation of the protective alumina film at high temperature (>1000 °C), weakening the coating oxidation resistance (Chim et al., 2009). Compared with TiN and TiAlN, another nitride ceramic CrN is the most corrosion resistant coating due to the highly protective properties of the formed oxide scale at temperature above 700 °C water following as the chemical reaction below (Korablov et al., 2005):

2TiAlN þ 7H2 O ! TiO2 þ Al2 O3 þ N2 " þ7H2 "

ð2Þ

2CrN þ 3H2 O ! Cr2 O3 þ 2N2 " þ3H2 "

ð3Þ

Similar to the TiAlN coating system, the incorporation of Al into cubic CrN crystalline structure was found to greatly enhance the hardness and oxidation resistance of AlCrN coating system at the elevated temperature air (Chim et al., 2009). Nevertheless, AlCrN coating was dissolved mostly in autoclave test and cracked in steam test due to its poor adhesion to the substrate, which did not occur in the case of CrN and TiAlN coatings (Daub et al., 2015). Moreover, Mn+1AXn (where n = 1, 2, 3) phase materials, composed of hexagonal structure ternary carbides and nitrides (referred as ‘‘X”) of a transition metal (referred as ‘‘M”) and an Agroup element, are promising for advanced coatings due to the combinability of metallic and ceramic properties (the strong metallic-covalent M-X bonds in combination with the relatively weak M-A bonds), good thermal and electrical conductors, superb machinability, and lightweight (Eklund et al., 2010). The outstanding high-temperature oxidation resistance of the Al-containing MAX phase compounds were reported to be related to the formation of the protective Al2O3 layer, like Ti2AlC (Maier et al., 2015). In addition, Ti3SiC2 and Ti3SiC2 also have the same favorable properties as other MAX-phase materials which makes them a possible candidate for use as a coating against high-temperature oxidation onto Zr-based cladding (Barrett et al., 2012). Like the nitride compounds, carbide compounds, especially silicon carbide (SiC), also recently have been considered as a candidate material for coatings. SiC mainly used as the alternative material for Zr-based alloys would be introduced in the subsequent parts. From the viewpoint of improvement in adhesion properties of coatings, Ti layer as bond coating exists between the ceramic coating and substrate (Alat et al., 2015). Moreover, the Zr oxide film formed on the Zr-based alloys would provide a high chemical stability, thereby reducing the oxidation rate of Zr-based alloys, on basis of which, zirconia

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coating, consisting of t-ZrO2 and monoclinic zirconia (m-ZrO2), was formed by plasma electrolyte oxidation (PEO) to improve the corrosion and wear resistance of the substrate (Cheng et al., 2012b; Xue et al., 2010). The zirconia coatings formed in silicate electrolyte with lower t-ZrO2 content than m-ZrO2 content, comprised a porous inner layer and a more compact outer layer, the latter being rich in t-ZrO2 (Cheng et al., 2012b; Xue et al., 2010). However, it should be noted that chemically stabilized zirconia ceramics undergo degradation from t-ZrO2 to m-ZrO2 when exposed to moisture even at low temperature (
100 °C) (Cox, 2005) . To recap briefly, some coatings have been fabricated by common coating techniques and desired chemical and physical properties of the coated Zr-based cladding also have been confirmed in the autoclave tests and steam or air exposure tests at elevated temperature. Concurrently, investigators also proposed their further work to put their productions into service in the nuclear power plants (summarized in Table 4). There are almost no thorough experiments having been done to acquire the performances of the coatings, such as the penalty to thermal conductivity, neutron economy and mechanical properties, of which oxidation resistance, and the irradiation-introduced degradation resistance, should be evaluated in a long duration exposure tests. That is why it is not clear that if the coatings on Zr-based alloys will survive for a sufficiently long duration in a service environment or even under accident scenarios. Consequently, the coating techniques are still under development to increase the fuel burnup and improve safety margin but the feasibility of each proposal need to be further evaluated comprehensively, for example the long term in-pile tests. 4. Alternative materials to Zr-based alloys The waterside oxidation and hydrogen pickup serve to limit the service of Zr-based alloys by degrading the mechanical properties, although today’s Zr-based alloy claddings (such as M5, ZIRLO, and X5) exhibit optimized behavior under normal operating conditions via decades of active research and development. The waterside oxidation and hydrogen pickup could be hopefully addressed well by the coating techniques as described above part, however, the likelihood exists that the coatings are peeled off in the accident scenarios, and as a result, the matrix (Zr-based alloy) is exposed to the corrosive environment. It is an inherent defect for the Zr-based alloy that hydrogen production and ductility degradation of the fuel cladding are accelerated by the reaction between Zr and steam in the case of LOCAs, which was underlined recently by the Fukushima Daiichi accident. Consequently, another approach to achieve larger safety margins against severe accident scenarios is to replace Zr-based alloy with substantially more oxidationresistant cladding and/or cladding with less heat and hydrogen production. It is understood that the alternative materials should meet or even exceed the performance metric of existing fuel designs while providing additional benefits/relief during severe accidents as well as alleviate some normal operational performance issues (Yueh and Terrani, 2014). Currently, candidate alternative cladding materials to Zr-based alloys have included: advanced steels, refractory metal (primarily molybdenum alloy) and SiC. 4.1. Advanced steels Earlier in the 1950s, stainless steel was used as cladding for the fuel elements but it was replaced by Zr alloy due to its corrosive nature at the higher temperature in the 1960s (Alam et al., 2011). However, currently, Advanced steels are being reexamined for their potential application as nuclear fuel cladding owing to the specific limitations with Zr alloy under both design-basis and

beyond-design-basis accident scenarios, and particularly the significant progress over the past five decades in advanced steel (Terrani et al., 2014b). Oxidation-resistant structural materials, such as Ni-based alloys and austenitic stainless steels, offer the improved strength and oxidation resistance compared to Zrbased alloys. Ni-based alloys are limited in reactor service ascribed to helium (He) embrittlement, the susceptibility to stress corrosion cracking (SCC) in LWR and high neutronic penalty of Ni (Pint et al., 2015a)although they have higher temperature strength than either austenitic stainless steels or ferritic/martensitic (F/M) steels (Allen et al., 2009). Austenitic stainless steels (containing 16–25%Cr with face center cubic (FCC) structure) have higher strength and better corrosion resistance than F/M steels but higher susceptibility to SCC and void swelling under radiation, which resulted in He embrittlement and high temperature creep rupture (Allen et al., 2009; Kimura et al., 2007). Neutron irradiation can cause the yield strength to quadruple for austenitic stainless steels and double for F/M steels compared to their unirradiated value at 300 °C, and F/M steels also experience embrittlement for LWR cladding conditions because of dislocation loops and precipitates (Terrani et al., 2014b). Moreover, 18Cr-8Ni stainless steels (e.g. type 304L) cannot form a protective scale in steam and instead oxidizes to form FeOx at <1473 K (1200 °C) in steam which does not meet the accident tolerance criteria (a 100X improvement in oxidation resistance for Zr-based alloys to achieve larger safety margin) (Pint et al., 2015a). It was reported that Cr content higher 20% was needed to form a protective layer for ferritic and conventional 300-series stainless steels at 1200 °C steam, while ferritic FeCrAl alloys are the most promising alloys for severe accident tolerance claddings because an additional protective a-Al2O3 scale, higher stable compared to Cr2O3 in the presence of H2O, can be formed, generally at temperature above 800 °C (Cheng et al., 2012a; Pint et al., 2013; Terrani et al., 2014b). Commercial FeCrAl alloys (melting point is 1200–1500 °C), MA956 (Fe-20Cr-4.5Al), PM2000 (Fe-20Cr-5Al-0.01Mo), and APMT (Fe-21.5Cr-5.0Al-3.1Mo) have been assessed for application in LWR claddings. They showed more excellent mechanical properties and oxidation resistance at 1200 °Csteam or at 360 °C pressured water compared to Zr-based alloys and 304, 310 stainless steels (Park et al., 2015; Pint et al., 2015b; Terrani et al., 2014b). Nevertheless, the typically around 20% Cr-containing alloys are not suitable for use at the temperature around 500°C, since the microstructure of high Cr-content (
20%) Fe-Cr and Fe-Cr-Al alloys decompose into a Fe-rich (a) phase and a Cr-rich (a0 ) phase, called 475 °C embrittlement or a-a0 phase separation, degrading the mechanical performance during thermal aging (Ejenstam et al., 2015). Additionally, above 14% Cr content will result in a large number density of a0 precipitates leading to potential embrittlement under neutron irradiation (Field et al., 2015), while
17.5% Cr is needed with lower Al contents but for the preferred lower content alloys,
4.5% Al was needed to form a protective alumina scale at 1200 °C steam (Pint et al., 2015b). In addition to the benefit that Al can perturb irradiation-induced a0 precipitation under a dose of 7 displacements per atom (dpa) at 320 °C (Edmondson et al., 2016), it was worth noting that Al addition may raise the ductile-brittle transition temperature (DBTT) and the work hardenability (Yamamoto et al., 2015). Oak ridge national laboratory (ORNL) (Snead et al., 2014; Yamamoto et al., 2015) is developing nuclear-grade ATF FeCrAl alloys aiming to a new, metal-base structural material for LWR cladding, that exhibits greatly improved accident tolerance including good mechanical properties in a wide temperature range as well as oxidation and irradiation resistance under normal and transient operating conditions; The results are segmented into two phases: determination of the major alloying constituents (namely Cr and Al) and further optimization of the alloys with minor alloying addition, such as Mo, Si, Nb, and/or C;

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Fe-13Cr-4.5Al+Y was down-selected as a base alloy and optimized further in phase II, as a result, Fe-13Cr-4.5Al+Y+Mo+Si+Nb with hot-rolled and annealed processes is most attractive for further development. Meanwhile, FeCrAl alloys with BCC structure show roughly two orders of magnitude higher tritium permeation compared with Zr-based alloys and even higher than 304 stainless steel, leading to costly disposal of tritium-contaminated water and higher radiation risk for the plant personnel. For Zr-based alloys, the formation of ZrO2 on both inner and outer cladding surface impedes the tritium permeation effectively. Although an alumina layer on the surface of FeCrAl alloys could dramatically reduce tritium diffusion, alumina does not form under normal operating condition of LWRs (Hu et al., 2015). Undoubtedly, FeCrAl alloys have been the candidate ATF cladding materials for their superior performance in accidents, however, more effective, reliable tritium barrier is necessary to facilitate application of FeCrAl alloys in LWRs. Oxide dispersion strengthened (ODS) F/M steels, through adding thermally stable oxide particles dispersed in the F/M matrix, have been developed for the application to fuel cladding material to Generation IV nuclear energy systems such as supercritical pressurized water reactor (SCPWR), fast breeder reactor (FBR), and it also can be easily applied to the cladding material for high burnup operation of LWR (Cho et al., 2004; Ukai et al., 1993). Martensitic 9Cr-ODS steels (9Cr-0.13C-2W-0.2Ti-0.35Y2O3) and ferritic 12CrODS steels (12Cr-0.03C-2W-0.3Ti-0.23Y2O3) have been developed, and both of them can retain their tensile strength and ductility properties after 15 dpa at the temperature ranging from 400 to 500 °C, in addition to keeping complex oxide particles substantially stable after 150 dpa with Ni ion bombardment (Allen et al., 2009). However, the (9–12) Cr-ODS F/M steels are not suitable for SCPWR owing to an insufficient corrosion resistance which is influenced by Cr and Al content, therefore, high Cr-content (>13%) ODS F/M steels with Al addition have been under investigation (Cho and Kimura, 2007; Cho et al., 2004, 2006; Je and Kimura, 2014; Kimura et al., 2007, 2011). It was observed that high Cr (>13%) and Al (4.5%)containing ODS steels exhibited higher resistance to corrosion than SUS316L in SCPW and that Al addition decreased the DBTT and the thermal embrittlement (Cho and Kimura, 2007; Cho et al., 2004). 14–16% Cr content was recommended on the balance between a merit of corrosion resistance and a demerit of aging embrittlement with maintaining strength at elevated temperatures, in parallel, small amount of Zr or Hf addition can significantly increase creep strength at 973 K in Al-added ODS steels (Kimura et al., 2011), which was further confirmed by the results from ORNL that 12Cr-ODS containing 5%Al with Zr and Hf performed very well in creep at 800 °C and had a maximum using temperature of 1475°Cin steam (Pint et al., 2014). 14Cr- and 16Cr-ODS steels with 4%Al addition did not suffer from aging embrittlement at 773 K for 1000 h and SCC at strain rates from 1  104 to 3  107s1 in simulated BWR conditions(Cho et al., 2006), while it was also confirmed that ODS steel (15Cr-4Al-2W) showed no susceptibility to SCC in SCPW with dissolved hydrogen (DH) and dissolved oxygen (DO) regardless of the strain rate (Je and Kimura, 2014). Reduced cladding thickness, and/or increased fuel enrichment are required to substitute the Zr-based claddings with FeCrAl cladding, in order to overcome the neutron penalty and decrease in refueling cycle length. Accordingly, reducing cladding thickness from 571.5 lm to 350 lm, and increasing fuel enrichment from 4.9% to 5.06% were proposed to match lifetime requirements of standard U2O for PWR, resulting in a rough 15% increment in the fuel pellet production cost for FeCrAl cladding (George et al., 2015). Subsequently, the feasibility of the design (reducing the cladding thickness and slightly increasing the fuel enrichment) was confirmed, demonstrating that FeCrAl cladding is a promising candidate to alternate Zr-based alloys, despite of the insignificant

Table 5 Ultimate Tensile Strength of candidate cladding materials (MPa) (Cheng, 2012). Materials

300 °C

1000 °C

Zry-4 Stainless steel 304 F/M steel SiCf/SiCm Composite Mo alloys

270 475 480 300 400–570

Nil <10 <10 300 200–300

demerits that fission gas release and fuel temperature (<50 K even at high burnup) is increased (Wu et al., 2015). 4.2. Mo-based alloys Refractory metals (such as Mo, Nb, Ta, W, etc.) have melting temperatures in excess of 2000 °C and possess good creep/swelling resistance up to high burnups, whereas showing poor oxidation resistance coupled with low-temperature radiation embrittlement and fabrication (joining) difficulties (Murty and Charit, 2008). Among these metals, only Mo was considered for LWR fuel cladding application based on its high tensile strength (Table 5), absence of hydride forming, acceptable thermal neutron capture cross section (2.6 bar), high melting temperature (2623 °C) and excellent thermal conductivity (138 Wm1K1), while Nb was considered as an alloying element for Mo alloys (Cheng, 2012). Mo and its alloys react readily with oxygen or steam at high temperature via below reactions but slower than observed for optimized Zr-based alloys (e.g. M5, ZIRLO, E110) above 1000 °C (Nelson et al., 2014):

2Mo þ 3O2 ! 2MoO3

ð4Þ

Mo þ 3H2 O ! MoO3 þ 3H2

ð5Þ

Oxidized Mo is susceptible to losses from volatile Mo trioxide, MoO3, in air and hydroxide, MoO2(OH)2, in water, but Mo oxides (MoO2/MoO3) become less protective with increasing the temperature and subsequently evaporate rapidly at elevated temperature. Therefore, Zr-based alloys or FeCrAl alloys were selected as coating materials by PVD on thin walled Mo alloys to provide good corrosion resistance at high water reactor operation temperatures via forming an in-situ oxide layer of ZrO2 or Al2O3 (Kim et al., 2015c). An ATF cladding design developed with coated Mo-alloy cladding was proposed by Electric Power Research Institute (EPRI) as shown in Fig. 7, consisting of duplex or triplex layers with the diameter of Mo-alloy tube being 9–10 mm: the outer surface is coated with a thin, metallurgically bonded Zr-based alloys or Alcontaining stainless steel (FeCrAl) to provide corrosion resistance in LWR coolants during normal operation and oxidation resistance during severe accidents; optional Zr- or Nb-alloy on the inner sur-

Fig. 7. Schematic of coated Mo-alloy cladding (Cheng et al., 2016).

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Fig. 8. High-temperature strength of various SiC VS zirconium alloys (Hallstadius et al., 2012).

4.3. SiC As a ceramic, SiC has been identified as a potential candidate material to replace Zr-based alloys thanks to a higher melting temperature, reduced hydrogen generation under off-normal conditions, low chemical activity, and a lower neutron absorption cross-section (Younker and Fratoni, 2016). Compared with Zrbased alloys, SiC can retain its strength to very high temperature, as shown in Fig. 8; no balloon occur until its failure temperature (>2000 °C) (Hallstadius et al., 2012). It has been reported there are different types of monolithic SiC, such as chemical vapor deposited SiC (CVD SiC), sintered SiC (SSiC) (Kim et al., 2002), reaction-bonded SiC (RBSC) (Kim et al., 2003), and liquid-phase sintered SiC (LPS SiC) (Kazunari Okonogi et al., 2015). As common ceramic consumer products, monolithic SiC fares brittle failure when subjected to a sudden mechanical shock, which not only leads to the loss of use of the component but also generates a host of foreign debris issues that could cause fuel failures during operation (Yueh and Terrani, 2014). To overcome the brittle behavior, the continuous SiC fiber-reinforced SiC matrix (SiCf/SiCm) structural composites were developed with multiple industrial processes: chemical vapor infiltration (CVI), liquid silicon infiltration (LSI), melt infiltration (MI), polymer impregnation and pyrolysis (PIP) and nano-infiltration and transient eutectic-phase (NITE) (Katoh et al., 2014b). However, the challenge exists in achieving

very low porosity for SiCf/SiCm, indicating that SiCf/SiCm alone may be insufficient to enclose fission gasses within the fuel cladding. As a consequence, the SiCf/SiCm cladding was optimized to SiC-SiC cladding consisting of a SiCf/SiCm layer outer layer for strength and a dense monolithic SiC inner layer for impermeability by Industry Advisory Committee, meanwhile the thermal conductivity and neutronic economic were reduced resulting from more complex structure (containing porosity and interface) and thick wall (Barrett et al., 2012). A novel SiC-based fuel cladding for PWR (referred to as Triplex SiC cladding) is under deployment by Massachusetts Institute of Technology (MIT), consisting of three layers: an inner monolith, a central SiCf/SiCm, and an outer environmental barrier coating (EBC); the inner layer is a high-density SiC to hold fission gases, the intermediate layer provides the required strength, and the EBC is a dense SiC layer for corrosion protection, which is pictured in Fig. 9 (Carpenter, 2010; Stempien et al., 2013). The Triplex SiC cladding was found to have acceptably low irradiation-enhanced corrosion rates and predictable swelling behavior (Carpenter, 2010), whereas a significant reduction in hoop strength and thermal diffusivities was also observed (Stempien et al., 2013). The properties and fabrication of

1000

SiC-Cg SiC-Type-S SiC-Hi-nic SiC-Tyranno Zircaloy-4 Zircaloy-2

900

Ultimate Tensile Strength (MPa)

face provides additional resistance to attack by fission products or higher toughness; the thickness of the Mo-alloy cladding is limited to 0.25 mm to reduce its impact on neutron absorption meanwhile the coating thickness of Zr-alloy or FeCrAl is nominally
0.05– 0.125 mm and
0.05 mm, respectively (Cheng et al., 2015, 2016). Currently, the feasibility of the fabrication for thin-walled Moalloy tube has been proved and welding of Mo cladding to endcaps have been successfully developed using electron beam welding and several other techniques (Cheng et al., 2015). Moreover, Zralloy-and FeCrAl-coated Mo-alloy cladding by PVD also has been demonstrated to be capable of partially or completely surviving in 1200 °C for 24 h; In order to further confirm the feasibility of Mo-alloy as cladding for LWR, alternative coating techniques (hipping, coextrusion, pilgering, drawing, and their combinations) for commercial production, evaluations on properties of Mo-Nb alloys, irradiation impacts, and steam oxidation (>1200 °C) have been initiated (Cheng et al., 2016).

800 700 600 500 400 300 200 100 0

0

500

1000

Temperature( C) Fig. 9. Cross section of the Triplex SiC design (Carpenter, 2010).

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145

Table 6 Studies on the properties and fabrications of SiC-based claddings. Authors

Materials

Tests

Kim et al. (2002, 2003)

RBSC, CVD SiC and SSiC

HT and pressured water

Okonogi et al. (2015)

CVD SiC and LPS SiC

High-temperature steam and pressured water

Stempien et al. (2013)

Triplex SiC tube

Post-irradiation examination

Barringer et al. (2007) Yueh et al. (2012)

CVD SiC

500°C supercritical water

SiCf/SiCm

Feasibility and neutronic evaluation

Terrani et al. (2012, 2014)

CVD SiC, SiC-SiC duplex tube, and NITE SiC

Steam and steam-hydrogen environments

Kim et al. (2015)

CVD SiC and SiC triplex composite

Immersion test and mechanical test

Lorrette et al. (2015)

CVI SiC

Autoclave tests and steam, air or O2 mixture exposures

Remarks  CVD SiC exhibited superior oxidation resistance to RBSC and SSiC.  Free Si existing in RBSC was preferentially hydrolyzed, which is faster than SiC in aqueous conditions.  SSiC suffered readily intergranular corrosion due to the impurities, eg. Boron and carbon.  The corrosion rate of SiC in steam was less than 1/1000 of Zrcaloy-2.  The corrosion rate of LPS SiC depended on the additive content relative to CVD in the water environment.  Ion effect on the swelling of SiC is insignificant while increasing the corrosion rate in the aqueous condition.  Corrosion behavior of SiC strongly depended on the fabrication process and the environment.  Three layers: a monolith SiC layer/a SiC/SiC composite layer/ a monolith SiC layer.  10%-60% of hoop strength was lost after varying durations of coolant and neutron exposure.  Irradiation-assisted corrosion was observed via weight loss measurements.  The thermal diffusivities of the triplex tube is roughly three times lower than that of single layer tube and decreased significantly with irradiation due to the accumulation of phonon-scattering defects.  CVD SiC was preferentially attacked at the train boundaries.  SiC at the surface was hydrolyzed to form Si(OH)4 species and subsequently dissolved into water.  SiC is susceptible to irradiation-induced swelling and needed to be improved in irradiation resistance.  SiC/SiC composite is remarkably resistant to fragmentation.  The utilization of SiC increased reactivity, indicating the fuel enrichment may be decreased by about 0.1%.  Two-step recession process: silica scale formation and silica volatilization.  CVD SiC has similar oxidation behavior as SiC-SiC while NITE SiC exhibited more stable due to the presence of yttria and alumina in the surface layer.  Micro-cracks in SiC/SiC composites exhibited self-healing behaviors in atmospheric pressure steam.  The oxidation kinetics of SiC materials was at least two orders of magnitude of that of Zrbased alloys.  SiC triplex tubes: a monolith SiC layer/a SiC/SiC composite layer/ a monolith SiC layer.  Hydrogen injection retarded effectively the formation of the oxide or hydroxide layer, reducing the corrosion rate of SiC in simulated-PWR water.  Hoop strength of SiC triplex tubes was proportional to the fiber volume fraction.  Weight loss was led by the silica dissolution which was independent of water chemistry.  Both the integrity and geometry of SiC/SiC composite were fully retained after oxidation at the high-temperature water and steam conditions while fracture surfaces were observed at 1400 °C air/H2O.

SiC-based cladding have been investigated widely, some of which are listed in Table 6 (Barringer et al., 2007; Lorrette et al., 2015; Kazunari Okonogi et al., 2015; Kim et al., 2002, 2003, 2015b; Stempien et al., 2013; Terrani et al., 2014a, 2012; Yueh et al., 2012). In spite of the potential benefits of SiC-based cladding, there are a lot of the technical issues required to be settled, such as the fabrication of thin-walled long tubes, capability of fission products, corrosion under the normal operating condition of PWRs, and the hermetic joining of end-cap seals (Kim et al., 2015b). Fig. 10 shows the effects of the temperature and irradiation on the thermal conductivity of SiC (Snead et al., 2007). Below 300 K, the thermal conductivity increases monotonically with grain size, while the phono-phono scattering effect becomes domination above 300 K. It is obvious that the thermal conductivity of SiC degrades significantly due to the accumulation of the irradiation-induced defects. In addition, the brittle property is one of the main obstacles preventing the application of SiC as the cladding material, as it cannot dissipate the accumulated stresses after pellet-cladding mechanical interaction(PCMI) through the plastic deformation and creep. The gap width and interfacial pressure between the fuel and cladding were calculated by the computer code FRAPCON, and were plotted in Fig. 11 (Li, 2013). Gap width increases initially due to the densification the fuel pellet. Unlike to Zircaloy, SiC cladding could not creep inward to the fuel, the reduction of pelletcladding gap is only contributed by fuel swelling and thermal

expansion. Therefore, the interfacial pressure for SiC cladding increase rapidly after the higher burnup, thereby causing the accumulation of hoop stress on SiC cladding until it finally fails. Consequently, suitable the fuel-cladding gap or an optimized fuel design is needed to avoid PCMI for the integrity of SiC cladding under high burnup. In the normal operating LWR conditions, the oxidation reaction, which is independent of the water chemistry (Lorrette et al., 2015), of the free Si in the SiC matrix and SiC are considered to progress as follows (Kim et al., 2002):

SiðsÞ þ 2H2 OðlÞ ! SiO2ðsÞ þ 2H2ðgÞ

ð6Þ

SiCðsÞ þ 2H2 OðlÞ ! SiO2ðsÞ þ CH4ðgÞ

ð7Þ

At the high-temperature and high-pressure water, the SiO2 layer is believed to be unstable and dissolved into water through the following reaction (Kim et al., 2002):

SiO2ðsÞ þ 2H2 OðlÞ ! SiðOHÞ4ðaqÞ

ð8Þ

With the appearance of dissolved oxygen at the hightemperature and high-pressure water, Reaction (9) can take place, and the presence of lithium in the solution may accelerate the corrosion under oxygenated, alkaline conditions (Carpenter, 2010).

SiCðsÞ þ 2O2 þ 2H2 O ! SiðOHÞ4ðaqÞ þ CO2ðgÞ

ð9Þ

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Fig. 10. Effect of radiation and temperature on the thermal conductivity of SiC (Snead et al., 2007).

Fig. 11. Pellet-cladding gap width and interfacial pressure vs. burnup (Li, 2013).

Given that water vapor is the dominant oxidizing agent for SiC at high temperature, the recession of SiC is governed by a two-step process of silica formation on the surface of SiC by the reaction (10) followed by the volatilization via the reaction (10) (Terrani et al., 2014a) .

SiCðsÞ þ 3H2 OðgÞ ! SiO2ðsÞ þ 3H2ðgÞ þ COðgÞ

ð10Þ

SiO2ðsÞ þ 2H2 OðgÞ ! SiðOHÞ4ðgÞ

ð11Þ

It was reported that UO2 and SiC underwent a chemical reaction at temperature above 1377 °C in the following formula (Lippmann et al., 2001).

2SiCðsÞ þ UO2ðsÞ ! USi2 þ 2COðgÞ

ð12Þ

The deployment of SiC-based materials for LWR necessitates the development of a joining method and joint material for the SiCbased claddings. Various methods and materials have either already been established or are considered promising for joining, including solid state diffusion bonding, transient eutectic-phase joining, MAX-phase joining etc. methods with Ti3SiC2, SiC, and CaO-Al2O3, etc. materials (Katoh et al., 2014a). Moreover, the mechanical properties of four kinds of joint geometries (as shown in Fig. 12: butt, butted lap, scarf, and butted scarf) were evaluated, results showed that a butted scarf endplug was capable to provide the required out-of-pile strength and permeability performance for

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147

Fig. 12. Schematic representation of cylindrical joint geometries (Khalifa et al., 2015).

LWR (Khalifa et al., 2015). However, the technology for manufacturing high temperature resistant joints of ceramic materials has not yet been adequately developed and the lack of a widely accepted standard method to assess the SiC joints also hinders the development of joining for SiC-based cladding (Gentile and Abram, 2015). Additionally, the manufacturing costs, modification of the rule system (e.g. the emergency core cooling system), and long time together with the high cost needed for verification are also impeding the revolutionary change from Zr-based cladding to SiC-based cladding for LWR (Hallstadius et al., 2012).

5. Conclusions The efforts to improve the performance of existing fuel claddings or to deploy novel claddings for LWR (primarily BWR, PWR, SCPWR) made by international researchers have been extensively reviewed in the present paper. The highlighted points are listed as below: A) Zry-2 and Zry-4 have been used as fuel cladding materials since the 1960 s after substituting stainless steels. Advanced nuclear reactor design such as higher fuel burnup and zero accident, is driving the optimization of the existing Zrbased alloys and deployments of new alloys by modifying their chemical compositions and/or processes. As a result, various advanced cladding materials like E635, ZIRLO, M5, MDA, HiFi, and X5A, etc., have been developed, some of which (E635, ZIRLO, M5, etc.) have been commercialized for LWR fuel application. B) Coating techniques have been utilized to address the waterside corrosion of the current Zr alloy claddings which is the main limitation to burnup extension, so as to achieve a higher accident tolerance. Currently, the candidate materials for coatings on Zr alloy cladding surface mainly composed of Cr, Fe, Al containing metallic materials and C-or Ncontaining ceramic, such as TiN, CrN, MAX-phase materials, via various methods: thermal spray, PVD, and CVD, etc. Coating on Zr alloy cladding surface is considered as a promising pathway to realize the extension in burnup and enhancement in safety margin within a short term (less than 5 years), although the thorough test data are still in shortage. C) Recently, SiC-based materials, Al-containing steels (Fe-CrAl) and Mo-based alloys have been investigated to substitute Zr-based alloys as accident tolerant fuel cladding for LWRs. The utilization of alternative materials (Non-Zr) for LWR fuel cladding is the most effective pathway to address the inherent demerits of Zr-based alloys i.e.: strength loss at high temperature and rapid oxidization at high-temperature steam. Nevertheless, to accomplish this purpose requires a long time together with a highly-increased cost, which need to take into consideration for achieving a reasonable return.

Acknowledgement A part of this study is the result of ‘‘The development of selfhealing intelligence on nuclear fuel cladding” carried out under the Center of World Intelligence Project for Nuclear S&T and Human Resource Development by the Ministry of Education, Culture, Sports, Science and Technology of Japan.

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