After the March 2011 events at the Fukushima site, the U.S. congress directed the Department of Energy (DOE) to focus efforts on the development of fuels with enhanced accident tolerance. In comparison with the standard UO2–Zircaloy system, the new fuels need to tolerate loss of active cooling in the core for a considerably longer time period while maintaining or improving the fuel performance during normal operations. A GE Global Research led program is investigating the behavior of advanced steels both under normal operation conditions in high temperature water (e.g. 288°C) and under accident conditions for reaction with steam. Results show that, under accident conditions, the advanced steels: (1) have low reactivity with steam, (2) would generate less hydrogen than the current zirconium based alloys, (3) offer higher mechanical properties, and (4) offer excellent retention of fission products. These advanced steels are also highly resistant to environmental cracking under normal operation conditions and they are not susceptible to suffer irradiation damage such as void swelling.
Keywords: light water reactors, fuel cladding, advanced steels, steam, crack propagation, mechanical properties
The commercial nuclear energy in the United States had its origin in the nuclear navy. The navy originally adopted zirconium based alloys over stainless steels for the fuel cladding mainly because of the higher transparency to neutrons of the former making the reactors more compact for submarine applications [Terrani et al. 2013]. In spite of the thermal neutron cross section of stainless steels being approximately 12–16 times higher than for zirconium alloys it is now understood that the fuel enrichment penalty incurred by the use of stainless steel cladding can be partially overcome by using thinner wall advanced stainless cladding because they are stronger than the zirconium alloys [Terrani et al 2013].
At the beginning of the nuclear navy program, the susceptibility to cracking of sensitized stainless steels in high temperature water initially relegated the stainless steels in research and development as compared to the zirconium alloys. Six decades later, it is now understood that high strength ferritic stainless steels are resistant to environmental cracking and irradiation damage, which would not limit their application as fuel cladding in light water reactors. Moreover, current significant progress in steelmaking practices shows that the chemical purity in modern steels can be highly controlled. Similarly, there is an increased ability in the fabrication of these advanced steels into thin walled tubes, including readily joining (welding) by several techniques.
Terrani et al. [2013] cite several uses in the industry of non-zirconium alloys as fuel cladding, including type 304, 316 and 347 austenitic stainless steels and austenitic nickel based alloys such as Inconel 600 and Incoloy 800. Type 304 SS fuel cladding was used for some time in US commercial light water reactors, for example at the Connecticut Yankee and San Onofre 1 power stations [Rivera and Meyer 1980]. The early cracking of type of 304 SS was linked to sensitization due to welding of high carbon alloys. The cracking phenomenon of sensitized stainless steels is now well understood and controlled and it is not a current concern in light water reactors.
In the post Fukushima Daiichi scenario there is an increased effort by the international community to find a material or materials that would an alternative to the zirconium alloy for fuel cladding. This material should have
One of these alternatives could be advanced steels since earlier concerns about stainless steels such as cracking, radiation damage and neutron economy can be now retired by using advanced ferritic steels (e.g. Fe-Cr-Al alloys). The advanced steels are also far superior to zirconium based alloys regarding higher strength, higher retention of fission products, and higher resistance to reaction with steam.
Austenitic stainless steels such as types 304 and 316 are highly susceptible to stress corrosion cracking in chloride containing environments, especially at temperatures higher than 60°C [Sedriks 1996]. The most common tests used in the industry to determine susceptibility to chloride cracking are immersion of U-bend specimens (ASTM G 30) into hot solutions of chloride salts including magnesium chloride or sodium chloride (ASTM G 36 and G 123). Ferritic stainless steels such as types 405 and 430 are highly resistant to SCC in hot chloride solutions [Sedriks 1996, Bond et al 1967].
Austenitic stainless steels such as type 304, 308, 316, 321 and 347 are used worldwide as construction materials for light water power reactors [Andresen 2012]. In the USA the most common austenitic alloy may be type 304 SS (UNS S30403) and in Japan the preferred stainless steel is type 316 (S31603). European countries such as Germany may prefer to use titanium (Ti) or niobium (Nb) stabilized types of stainless steel such as type 321 (S32100) and 347 (S34700). Austenitic stainless steels are susceptible to stress corrosion cracking (SCC) in boiler water reactor (BWR) service and in a lesser extent in pressurized water reactor (PWR) service [Andresen 2012].
Austenitic stainless steel (SS) core internals components are susceptible to irradiation assisted stress corrosion cracking (IASCC) during service in nuclear power plants light water reactors [Chung and Shack 2006, Cookson et al 1993, Jacobs et al 1993]. One of the effects of irradiation is the hardening of the SS due to modifications in the dislocation distribution in the alloy [Was 2003, Bruemmer 2002]. Irradiation also alters the local chemistry of these austenitic alloys, for example in the vicinity of grain boundaries by a mechanism of radiation induced segregation (RIS). The segregation or depletion phenomena at or near grain boundaries may enhance the susceptibility of these irradiated alloys to stress corrosion cracking (SCC) [Was et al 2006, Yonezawa et al 2000]. The effect of the IASCC on austenitic stainless steels may impact the life extension of currently operating light water reactors due to the progressive dose accumulation [Hojná 2012].
In nuclear power plant applications, ferritic steels have superior void swelling resistance because they experience delayed void nucleation and they sustain less than 2% swelling even at irradiation levels close to 200 dpa [Raj and Vijayalakshmi 2010]. On the other hand, austenitic stainless steels such as type 304 undergo the onset of significant void swelling and possible embrittlement at dose rates in the order of 20 dpa [Was et al 2006]. Besides the higher resistance of ferritic steels to radiation damage, other benefits that could make these steels more attractive than the austenitic stainless steels in nuclear applications include: (1) Ferritic materials have lower cost since they do not contain nickel (Ni), and generally contain lower chromium (Cr), (2) They do not contain Ni or cobalt (Co) that could be become activated in commercial reactors, (3) They offer a lower coefficient of thermal expansion (CTE) that matches the CTE of pressure vessel ferritic alloys such as type A508, A516, or A533 [Ren et al 2008], and (4) Ferritic steels have higher thermal conductivity for heat transfer capabilities (Table 1).
Table 1 – Physical Properties of Ferritic and Austenitic Steels
| Steel | CTE (0-538°C) μm/m/°C | Thermal Conductivity at 100°C (W/m.K) |
| Zircaloy-2 | 8.32 & 15.7 (orientation dependent) | 13.8 |
| Ferritic type 430 (16% Cr) | 11.4 | 23.9 |
| Austenitic type 304L (18% Cr) | 18.4 | 16.2 |
In the case of a loss of coolant accident (LOCA), such as in the Fukushima Daiichi situation, the cladding of the fuel will be exposed to steam. The zirconium alloy plus steam reaction has been widely studied under loss of coolant accident scenarios (Whitmarsh 1962, Baker and Just 1962, Leistikow and Schanz 1987, IAEA 1992, Grandjean and Hache 2008, Terrani et al 2013). Zircaloy oxidizes in presence of steam to form zirconia and hydrogen following an exothermic reaction:
According to Baker and Just the chemical heat generated by the reaction of zirconium and steam in (Equation 1) could exceed the nuclear heat generation during a destructive nuclear transient. Moreover, the hydrogen generated by the reaction could give rise to a pressure surge and might subsequently react explosively with air [Baker and Just 1962].
All the ferrous materials listed in Table 2, including ferritic steel T91, have lower reaction kinetics with steam than Zircaloy-2. The oxidation behavior of iron based alloys in steam was recently reviewed and updated comparing to the behavior of zirconium alloys [Pint et al. 2012, Terrani et al 2013] At 1200°C, the degradation of APMT was practically nil (no mass change) after 8 h exposure at 1000°C while the degradation of Zircaloy-2 was complete for the same period of time [Pint et al. 2012] APMT offers extraordinary resistance to reaction with steam at temperatures higher than 1000°C because it allows first for the formation of a protective Cr2O3 scale which subsequently allows for the formation of a continuous protective Al2O3 scale between the metal and the Cr2O3 scale. It is this Al2O3 scale what protects the alloy against further oxidation in steam [Opila 2004, Pint et al 2012]. Cheng et. al., [2012] studied the oxidation behavior of several new cladding candidates (SiC, stainless steels 304 and 317, alloy PM 2000 and iron-based alloys with 15–25% Cr) at 800, 1000 and 1200°C. They concluded that, aluminum in the PM 2000 alloy formed a very protective alumina layer which significantly reduced the mass loss compared to the other materials under superheated steam conditions [Cheng et al 2012]. Pint et al. [2012] also showed the effect of the content of Cr on the degradation of ferrous alloys. It was reported that the content of Cr is important and that, in the absence of Al, at least a 25% of Cr may be required in the iron alloy to offer protection against steam [Pint et al. 2012]. However, it is likely that if the alloy also contains approximately 5% Al, lower amount of Cr may be needed to offer a similar resistance to oxidation.
Table 2 – Studied Alloys for Accident Tolerant Cladding in Light Water Reactors
| Alloy | Nominal Composition |
| Zircaloy-2 | Zr + 1.2–1.7 Sn + 0.07–0.2 Fe + 0.05–0.15 Cr + 0.03–0.08 Ni |
| Ferritic steel T91 | Fe + 9 Cr + 1 Mo + 0.2 V |
| Ferritic steel HT9 | Fe + 12 Cr + 1 Mo + 0.5 Ni + 0.5 W + 0.3 V |
| Nanostructured ferritic alloys NFA | e. g. 14YWT; Fe + 14 Cr + 0.4 Ti + 3 W + 0.25 Y2O3 |
| MA956 | Fe + 18.5–21.5 Cr + 3.75–5.75 Al + 0.2–0.6 Ti + 0.3–0.7 Y2O3 |
| APMT | Fe + 22 Cr + 5 Al + 3 Mo |
| E-BRITE – S44627, | Fe + 25-27.5 Cr + 1 Mo + 0.17 (Ni + Cu) |
| Alloy 33 – R20033 | 33 Cr + 32 Fe + 31 Ni + 1.6 Mo + 0.6Cu + 0.4N |
The objective of the present work is to compare the environmental and mechanical behavior of advanced steels to the behavior of the currently used zirconium based alloys. The comparison will include behavior under (a) normal operation conditions and (b) under accident conditions such as loss of coolant accident (LOCA).
The aim of the current research is to characterize the behavior of advanced steels as candidate cladding materials in comparison to the behavior of the current zirconium alloys cladding. Table 2 shows the list of alloys that currently are being studied. The characterization studies are being performed both under (a) normal operation conditions (e.g. water at 288°C), and (b) accident conditions (e.g. steam at 800°C).
Under normal operation conditions the cladding may not breach releasing fission products into the water. That is, similarly to the actual zirconium based alloys, the candidate replacement alloys should not corrode excessively in water at ~300°C nor suffer environmentally assisted cracking under similar conditions. The life of a fuel bundle in a commercial reactor is generally less than 10 years, that is, under normal operation conditions the cladding should be able to survive for this period of time.
It can be anticipated that general corrosion of the advanced steels materials would be in the acceptable group of alloys for cladding applications. Nevertheless, coupons of the candidate materials are being tested for general corrosion under laboratory simulated normal operation conditions of commercial light water reactors. Four sets of autoclaves are being used (Table 3). The degradation of the immersion coupons is being evaluated by weight (mass) change, standard metallographic procedures and surface analysis techniques. In all the autoclave systems (S-2, S-5, S-6, and S-14) the water is re-circulated at a flow rate of 100 cm3/min and reconditioned (filtered) before entering again the autoclaves.
Table 3 – Immersion Tests under Simulated Normal Operation Conditions
| Autoclave | Test Conditions at GE GRC | Alloys under Testing |
| 2584 S-2 | Simulated PWR, High Purity Water, 330°C | T91, 14YWT, APMT |
| 2584 S-5 | Simulated BWR, Hydrogen Water Chemistry (63 ppb H2), 288°C | T91, 14YWT, APMT |
| 2584 S-6 | Simulated BWR, Normal Water Chemistry (2000 ppb O2), 288°C | T91, 14YWT, APMT |
| 2520 S-14 | Simulated BWR, Normal Water Chemistry (2000 ppb O2), 288° | Zircaloy-2, T91, HT9, 14YWT NFA, MA956, APMT, Ebrite, Alloy 33 |
Table 4 shows the non-irradiated materials that are being tested for resistance to environmental cracking [Andresen et al 2014, Rebak 2013b].
Table 4 – Non-Irradiated CT Specimens under Testing at GE GRC
Most of the 0.5T CT specimens were machined from the plates so that the notch in the specimen was placed in the SL direction of the plate (ASTM E399). Most of the materials were cold forged (CF) by 20%, which increases their yield strength The 20% cold reduction was performed to compare crack propagation results with literature data for austenitic steels, which are normally cold reduced. (Cold reduction is used to accelerate crack propagation rate and, therefore, minimize the time required for testing). For a plate material, the SL direction should be the most susceptible to environmentally assisted cracking, since the crack will have a tendency to separate the material along pre-existing rolling lines.
0.5T compact type specimens were machined with 5% side grooves on each side. The CT specimens were instrumented with platinum current and potential probe leads for dc potential drop crack length measurements. In this technique, current flow through the sample was reversed about once per second primarily to reduce measurement errors associated with thermocouple effects and amplifier offsets. The test was computer controlled using inputs from the relationship between the measured potential and crack length. Data were stored in a permanent disk file typically once every 0.69 hours. In addition to the data record number, total elapsed and incremental time, and crack length, the system measured and stored temperature, current, corrosion potential, dissolved gases, influent and effluent conductivity, load and time/date. Additionally, both operator and automated program messages describing changes in the test conditions and test status were a permanent part of the data record.
The CT specimens were electrically insulated from the loading pins using zirconia sleeves, and within the autoclave a zirconia washer also isolated the upper pull rod from the internal load frame. The lower pull rod was electrically isolated from the autoclave using a pressure seal and from the loading actuator using an insulating washer. Ground isolated instrumentation was used for the platinum current and potential probe attachments to the specimen. Testing was performed using servo-electric testing machines equipped with a single stage, low flow servo-valve to ensure optimal (non-noisy) response. Crack growth rates can be considered statistically meaningful when the crack growth increment is at least 10 times the resolution of the technique, which was typically 1 to 5 μm. Thus, crack length increments were typically > 50 urn, although for very low growth rate conditions, smaller increments were occasionally used to reduce testing time from several months per datum to several weeks. Generally, the lowest test time for each combination of variables (e.g. stress intensity) was in the order of 2 weeks or 3 mils (76 μm) of crack growth. The R2 correlation coefficients from linear regression analyses of the crack lengths vs. time data from which growth rates are calculated were typically >0.90.
Deaerated, de-mineralized water was drawn through a demineralizer and submicron filter to ensure ultra-high purity (0.055 μS/cm or 18 MΩ-cm) and then into a glass column (6.4-cm diameter by 183 cm long). The volume of the column is approximately 4 L, which added to the volume of the autoclave and the piping results on a total volume of solution in the order of 7–8 L. A high pressure pump recirculated the water from the column to the autoclave and back to the column at a rate of approximately 100–200 cm3/min (which represents two volumes replenishment of the autoclave each hour). The autoclave effluent was back-pressure regulated, then continuously monitored for conductivity and dissolved oxygen. The oxygen concentration was controlled by bubbling gas mixtures blended by mass flow controllers. Impurities of interest (such as 30 ppb sulfate ions as sulfuric acid) were added to the glass column using a metering pump, which was controlled via a preset value in the conductivity meter. The crack tests were performed in a 4-L (1 gallon) stainless steel autoclave at 550°F (288°C) and 1500 psia (10.3 MPa). The corrosion potentials of the CT specimen and a Pt coupon were continuously monitored using a zirconia membrane reference electrode containing copper and copper oxide, whose reference potential is 273 mV higher than the standard hydrogen electrode (SHE) in pure water at 550°F (288°C).
Once the specimen was loaded in the autoclave and connected to the leads, water recirculation started and the temperature and pressure were raised to 550°F (288°C) and 1500 psi (10.3 MPa), respectively. Cyclic fatigue started at a stress intensity of 25 ksi√in (27.3 MPa√m) using a trapezoidal wave at a frequency of 0.001 Hz, a load ratio (Kmin/Kmax) R = 0.6 and zero holding time. Ideally, once the crack front propagated 3 mils (76 μm) in the first step a holding time of 9000 s is applied for each cycle at the highest value of the stress intensity in that cycle. Again, ideally, after the crack advanced another 3 mils (76 μm), the stress intensity is kept constant at the highest value (or a static load, R = 1) and the crack advance is typically monitored for a minimum time of 2 weeks or a growth of 3 mils (76 μm). Current results show that these ideal situations cannot be fully applied for the alloys in Table 4 because they are so resistant to environmental cracking and cracking generally stops growing once the high frequency loading is transitioned to low frequency loading [Andresen et al 2012, Andresen et al 2014].
Changes in the crack growth rate were also monitored when the water chemistry was changed from pure water to water contaminated with sulfuric acid to give 30 ppb concentration of sulfate ions or chloride ions. Similarly the crack propagation rate was also monitored under oxygenated conditions or normal water conditions (NWC or containing 2 ppm of dissolved oxygen) and under hydrogen water conditions (HWC or containing 63 ppb of dissolved hydrogen). The presence of oxygen or hydrogen controls the corrosion potential of the specimen under test. The crack propagation rate is generally lowered when the corrosion potential is lowered.
The specimens listed in Table 4 are currently under testing. After the CT test in the autoclave is considered completed, the system will cooled down, the CT is removed and then fatigued cracked in air at ambient temperature for observation of the crack patterns both optically and under a scanning electron microscope.
Stress corrosion cracking (SCC) testing was initiated on a nano-ferritic alloy (NFA) (specimen c642 in Table 4), containing 14% Cr and oxide dispersion hardened with Y2O3 in the as-received condition (not cold worked) [Andresen et al 2014]. The in-situ fatigue pre-cracking proceeded as anticipated. As the frequency was slowed to transition to intergranular SCC conditions, cracking slowed and ceased [Andresen et al 2014]. The stress intensity and frequency had to be increased to re-initiate the cracking; however, as the loading frequency decreased, crack growth stalled. This crack growth cessation behavior at low frequencies was observed before for other ferritic steels containing 5, 9 and 13% Cr [Andresen et al 2013, Rebak et al 2013, Rebak 2013, Rebak 2013b].
Testing was also initiated on HT-9 ferritic steel with 20% cold work (specimen c647 in Table 4). The in-situ fatigue pre-cracking proceeded well, but after decreasing the frequency to 0.004 Hz at 260 hours, the crack growth rate slowed and stopped [Andresen 2014]. Testing on the APMT ferritic steel with 20% cold work (c648 in Table 4) for APMT are very low compared to known reference materials such as cold worked nickel based alloy 600 or type 316 stainless steels. The last specimen mentioned in Table 4 is c649 made using T91 ferritic steel with 23% cold work. T91 behaved similarly as the other ferritic alloys in the sense that when the conditions were made less aggressive in terms of applied stress intensity and loading frequency, cracking slowed down below measurable rates [Andresen et al 2014].
All of the ferritic alloys being evaluated for SCC response in this program have excellent resistance to stress corrosion cracking, even under quite aggressive conditions of elevated oxidants (2 ppm dissolved O2) and 30 ppb sulfate or chloride (well above that allowed by the BWR water chemistry guidelines).
All crack propagation reported here under cyclic loading condition can be considered fatigue cracking. Only under constant load conditions (R=1) the crack propagation may be recognized as environmentally assisted cracking or stress corrosion cracking (SCC). Current results show that ferritic steels containing Cr are extremely resistant to cracking in high temperature water.
Tests are being conducted to determine the relative steam oxidation resistance of advanced steels to compare their behavior with the current zirconium alloys. Table 5 shows the test conditions for the steam tests. Initially all the tests will be performed in 100% steam at a flow rate of 2.5 g/min, and eventually, tests will be performed in mixtures of steam plus hydrogen gas, and steam plus air. Experimental alloys, currently under fabrication, will also be tested for resistance to oxidation in steam and compared to the results from the baseline alloys in Table 5. The specimens were flat rectangular measuring approximately 25 × 8 × 2.3 mm with a total exposed area of 5–5.3 cm2. The surfaces were ground on wet SiC paper up to 320 grid finish. All specimens were washed with solvents and dried. The weight (mass) of the specimens was measured with direct reading microbalance at room temperature before and after each test (3 readings), and the mass change due to exposure to steam was calculated.
Table 5 – Current Steam Test Conditions at GE GRC
Figure 1 shows the superheated steam system (SSS). Some of main components include a vertical alumina retort where the five specimens were exposed to steam hanging vertically from a tree. Five thermocouples monitored continuously the temperature next to each specimen. The retort was connected to a steam generator where water was pumped at a rate of 2.5 g/min. The ultra-high purity (UHP or 18 MΩ) water was deaerated with argon before it was injected into the steam generator using a metering pump with a reciprocating piston design. The steam was forced to flow through four alumina diffusers (Figure 2) to allow for the preheating and homogenization of the steam. The temperature of the retort was controlled by a three-zone furnace. The steam exited the retort through another set of alumina diffusers to avoid back convection of steam onto the specimens. The steam was condensed at the exit of the system and the volume of water collected was comparable to the amount of water injected into the steam generator.
Figure 1 – Superheated Steam System (SSS)
Figure 2 – Set of Four Alumina Diffusers to Allow for Steam Homogenization in the Retort
Once the coupons are inserted in the retort and the system is sealed, the entire system is purged using a constant flow of pure argon (30 cm3/min) for 1–2 hours and the gas flow is maintained while heating the chamber (from room temperature to the testing temperature). When the testing temperature is reached, the argon gas flow is stopped and the argon deaerated water injection to the steam generator is started. The top and bottom stainless steel caps of the retort are maintained approximately 150–180°C to avoid steam condensation. At the end of the test, the steam injection is stopped and the testing chamber is cooled down using a flow of dry argon (30 cm3/min).
Figure 3 shows the appearance of the coupons before and after 8 hours exposure in steam at 800°C with a 2.5 g/min flow rate [Rebak 2013b]. After the test, the Zircaloy-2 specimen was slightly bent, perhaps due to the growth of oxides on the surface. The Zircaloy-2 oxide layer presented a white snake-like skin appearance and signs of spallation. Alloy T91, did not show any sign of deformation but small oxide spallation was evident. The nano ferritic alloy (NFA) exhibited a uniform black oxide scale whereas APMT and Alloy 33 showed minimal pink and light green discoloration, respectively [Rebak 2013b].
Figure 3 – Specimens before (top) and after (bottom) exposure to 100% steam at 800°C for 8 h
Figure 4 shows the mass gain per surface area as a function of testing time for the five alloys in Table 5. The Zirc-2 coupon was consumed at the 48 h testing time; therefore there is no data in the plot. Figure 4 shows that there were two evident groups of alloys regarding resistance to degradation in steam. Zircaloy-2 and T91 are in Group 1, with a higher oxidation rate, and Group 2 included NFA, APMT and Alloy 33 with an oxidation rate that was approximately two orders of magnitude lower than for Group 1. Overall, the highest oxidation rate was for Zircaloy-2 and the lowest for APMT.
Figure 5 shows a plot for the weight change as a function of testing time for the three alloys in Group 2 and the respective fitting according to a power law. The fitting equation and parameters are given in Table 6. The oxidation of the iron alloys seems to follow a parabolic law with an exponent coefficient close to 0.5. The coefficient is higher for the zirconium alloy, suggesting that oxidation was not controlled by diffusion through a protective oxide film on the surface.
Figure 4 – Mass change as a function of test time in 100% steam at 800°C
Figure 5 – Mass change as a function of test time in 100% steam at 800°C
Table 6 – Fitting parameters mass change as a function of time in pure steam at 800°C (Y = mass change in mg/cm2 and X = time in h)
| Alloy | Equation | |
| Zirc-2 | ln(Y) = 1.315 * ln(X) + 0.081 | R2 = 0.584 |
| T91 | ln(Y) = 0.491 * ln(X) + 0.789 | R2 = 0.708 |
| NFA | ln(Y) = 0.511 * ln(X) - 3.342 | R2 = 0.996 |
| APMT | ln(Y) = 0.467 * ln(X) - 4.117 | R2 = 0.937 |
| Alloy 33 | ln(Y) = 0.616 * ln(X) - 4.472 | R2 = 0.863 |
Figure 6 shows the appearance of the coupons in a scanning electron microscope after the exposure to 100% steam at 800°C for 8 h. The magnification for both the alloys in Group 1 (Zirc-2 and T91) is approximately 10 times lower than for the alloys in Group 2. For the APMT specimen there are two images, one to the X2000 magnification. Figure 6 shows the evident difference in oxidation behavior between these five materials. APMT has the lowest oxidation susceptibility in steam since marks from sample preparation are clearly discernible. On the other hand Zircaloy-2 is completely covered by a partially cracked oxide film.
Figure 6 – Appearance of the coupons after testing in 100% steam at 800°C for 8 h
As compared with the current zirconium based cladding, the advanced steel materials should offer higher resistance to reaction with steam and hydrogen gas generation, resistance to stress corrosion cracking in high temperature water, higher structural integrity, higher mechanical properties as the temperature increases, lower coefficient of thermal expansion, higher thermal conductivity and enhanced resistance to radiation damage such as material hardening or embrittlement and maintain dimensional instability caused by void and helium driven swelling. Oxide Dispersion Strengthened steels such as MA956 are resistant to radiation-induced swelling and have improved creep strength and oxidation/corrosion resistance at elevated temperatures compared to conventional steels [Hsiung 2010). It is generally accepted that an irradiated component in a nuclear power plant undergoes three stages during irradiation damage [Odette et al 2008]: (1) Primary defect production, (2) Long-range diffusion of the primary defect, and (3) Changes in the properties and dimensions of the component as a consequence of the “new” microstructure. Ferritic/martensitic steels are more resistant to irradiation damage than their cousins the austenitic type 316/304 stainless materials. This is especially true at the lower helium levels or lower dpa ratios [Odette at al 2008]. The main explanation of this higher resistance of the ferritic/martensitic materials is due to their bcc structure, which offers a higher self-diffusion coefficient, more traps for helium bubbles and a lower energy for dislocation climb. Another explanation is the lack of nickel in the ferritic/martensitic alloys (thermal neutrons may transmute Ni into helium) [Odette et al 2008].
One of the attributes of an improved cladding material is that it will also retain fission products under normal operation conditions (as the current zirconium alloys normally do). Moreover, the advanced steel cladding would also have improved retention of fission products under accident conditions over the current zirconium alloys. That is, the candidate material should not breach releasing fission products to the external environment. Breaching of the cladding may be produced from the OD or water side via a mechanism of stress corrosion cracking or environmentally assisted cracking or from the ID of the cladding or fuel side also via a mechanism of stress corrosion cracking. Zirconium alloys are resistant to cracking from the water side under normal operation conditions [Lemaignan 2006] but may suffer cracking from the fuel side due to a combined effect of hoop stresses and iodine [Cox 1976]. Zirconium alloys are also notorious for reacting with hydrogen to form internal hydrides, which render the cladding brittle and subject to enhanced cracking susceptibility [Cox 1976].
Current zirconium based alloys sometimes suffer from failure from the pellet side of the cladding (ID of the tubing) due to the phenomenon of pellet cladding interaction (PCI) [Cox 1990]. Zirconium based alloys are susceptible to cracking when there is fuel swelling that imparts hoop stresses in the ID of the cladding, which with the presence of fission products such as iodine, may develop cracks from the ID of the zirconium alloy cladding [Cox 1990].
The great advantage of the advanced steels (such as Fe-Cr-Al alloy) is their outstanding resistance to reaction with steam under accident conditions [Terrani et al 2013]. Previous work from the 1960s showed compatibility between this type of steel and the UO2 fuel, especially if the cladding material is pre-oxidized [Edwards and Bohlander 1969].
The great advantage of zirconium based alloys for fuel cladding is their transparency to neutrons. The advanced steels could overcome this barrier by making the walls of the cladding thinner (since at a given temperature the proposed advanced steel is stronger than Zircaloy) and, if necessary, by incrementally increasing the enrichment of the fuel.
The technical expertise of J. F. Flores, R. J. Blair, F. Wagenbaugh, and P. J. Martiniano is gratefully acknowledged.
This material is based upon work supported by the Dept. of Energy [National Nuclear Security Administration] under Award Number DE-NE0000568.
This report was prepared as an account of work sponsored by an agency of the United States Government. Neither the United States Government nor any agency thereof, nor any of their employees makes any warranty, express or implied, or assumes any legal liability or responsibility for the accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately owned rights. Reference herein to any specific commercial product, process or service by trade name, trademark, manufacturer, or otherwise does not necessarily constitute or imply its endorsement, recommendation, or favoring by the United States Government or any agency thereof. The views and opinions of authors expressed herein do not necessarily state or reflect those of the United States Government or any agency thereof.
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