Modern nanostructured ferritic alloys: A compelling and viable choice for sodium fast reactor fuel cladding applications

Journal of Nuclear Materials 529 (2020) 151928

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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat

Modern nanostructured ferritic alloys: A compelling and viable choice for sodium fast reactor fuel cladding applications* D.T. Hoelzer a, *, C.P. Massey a, S.J. Zinkle a, b, D.C. Crawford c, K.A. Terrani a a

Oak Ridge National Laboratory, Oak Ridge, TN, 37831, USA University of Tennessee, Knoxville, TN, 37886, USA c Idaho National Laboratory, Idaho Falls, ID, 83415, USA b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 17 June 2019 Received in revised form 8 November 2019 Accepted 25 November 2019 Available online 28 November 2019

Sodium fast reactor cores present a truly challenging environment for the fuel cladding and core structural materials that limit achievable fuel burnup in these systems. Modern nanostructured ferritic alloys have the potential to deliver the desired high performance characteristics where historic austenitic and ferritic/martensitic alloys fall short. In this paper, a new nanostructured ferritic alloy, OFRAC (Oak Ridge Fast Reactor Advanced Fuel Cladding), is developed and demonstrated in cladding geometry with a length > 1 m. The alloy composition and microstructure are tailored to deliver dramatically improved strength and creep resistance. At 600
C the ultimate tensile strength is ~600 MPa and the stress to induce a steady state strain rate of 106 s1 is 500 MPa vs. 200 MPa for traditional ferritic/martensitic alloys. A very high defect sink density was engineered in these alloys that is expected to further enhance the traditional good swelling and irradiation creep resistance of ferritic/martensitic steels. These significant improvements along with the demonstrated viability for seamless cladding production indicate that this alloy is a compelling and viable choice for sodium fast reactor fuel cladding applications. © 2019 Elsevier B.V. All rights reserved.

Keywords: Sodium fast reactor Nanostructured ferritic alloy OFRAC Thin wall tubing Tensile and creep properties

1. Introduction Innovative nuclear engineering designs for Generation IV fission reactor concepts have been proposed that, if successfully built and operated, would provide marked improvements compared to current Generation II water cooled reactors in terms of economics, fuel utilization and safety [1]. However, these innovative designs impose extreme operating conditions on structural materials such as fuel cladding in terms of high operating temperatures, high applied stress, and high neutron displacement damage levels [2e6]. In most cases, the technical feasibility of proposed

* This manuscript has been authored by UT-Battelle, LLC under Contract No. DEAC05-00OR22725 with the U.S. Department of Energy. The United States Government retains and the publisher, by accepting the article for publication, acknowledges that the United States Government retains a non-exclusive, paid-up, irrevocable, world-wide license to publish or reproduce the published form of this manuscript, or allow others to do so, for United States Government purposes. The Department of Energy will provide public access to these results of federally sponsored research in accordance with the DOE Public Access Plan (http://energy. gov/downloads/doe-public-access-plan). * Corresponding author. 1 Bethel Valley Road, USA E-mail address: [email protected] (D.T. Hoelzer).

https://doi.org/10.1016/j.jnucmat.2019.151928 0022-3115/© 2019 Elsevier B.V. All rights reserved.

Generation IV reactor concepts hinges on the identification or development of new ultra-high-performance structural materials that will allow the proposed reactor to achieve the necessary high operating temperatures and fuel burnup levels. In general, the materials used in current light water reactors (LWRs) are unsuitable candidates for structural applications in proposed Generation IV reactors due to the large difference in operating temperatures and displacement damage exposure levels. Even the cladding and wrapper/duct materials that were developed during major international liquid metal fast breeder reactor research programs in the 1970s and 1980s do not have sufficient thermal creep strength and void swelling resistance to meet the design requirements for sodium cooled fast reactors (SFRs) [4,7e9]. Although SFR fuel burnup design requirements depend in part on the reactor concept (transmuter, converter or breeder with respective fuel cycle conversion ratios of <1, ~1 or >1) [10], in general the required SFR fuel burnups exceed 10% and can be as high as 40% [11,12]. This corresponds to cladding and wrapper/duct damage levels in excess of 150 displacements per atom (dpa) and typical cladding stress levels of >100 MPa during the latter half of the multi-year in-core exposure due to fission gas release into the fuel rod plena. These high burnup levels are necessary to reduce the

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number of recycles needed (in recycle systems), reducing overall system cost and proliferation concerns and to enable significant reductions in long-lived, high-level waste [9,11]. For conventional SFR designs, the maximum operating temperature for fuel cladding can reach 550-600
C. Higher cladding temperatures approaching 700
C would be desirable to increase system efficiency or to enable more compact and thermodynamically efficient nitrogen or supercritical CO2 Brayton power conversion systems with reduced capital construction costs and enhanced safety (by eliminating the potential for highly reactive Na-steam reactions). Steady progress has been documented since the 1950s in the development of high performance steels for a variety of technologies with progressively higher upper use temperatures and creep stress capability [5,13]. Using the nomenclature in Refs. [13e15] for modern ferritic steels and their maximum use temperature, 1stgeneration steels were introduced in the 1960s and included HT-9 (Fee12%Cre1%MoVW). Modified 9%Cre1Mo (Alloy 91) is a 2ndgeneration steel developed in the late 1970s. Recently developed nanostructured ferritic alloys (NFAs) [1e4] that represent a marked advance over prior oxide dispersion strengthened (ODS) steels could be considered 5th-generation steels. During the initial commercialization of Generation I nuclear energy systems in the 1950s and 1960s, nuclear power was an aggressive adopter of advanced high-performance structural materials. However, relatively incremental changes to the structural materials in LWRs have occurred in the intervening ~60 years [9]. This relative stagnation with respect to introduction of new materials is somewhat understandable for LWRs since existing structural materials have a proven adequate performance and do not necessarily need to be replaced with higher performance materials. The worldwide liquid metal fast breeder reactor research and development programs in the 1970s similarly selected some of the most advanced steels for initial reference designs of fuel cladding, including HT-9 and modified 9Cre1Mo steels, and in parallel developed new advanced Ti-modified austenitic stainless steels that provided significant improvements in void swelling radiation resistance and thermal creep strength compared to Type 316 stainless steel [16e19]. However, these 2nd-generation steels in general do not simultaneously provide sufficient void swelling resistance at doses >150 dpa and requisite thermal creep strength. In summary, successful development of advanced Generation IV reactor concepts must utilize advanced high performance (4thgeneration or better) structural materials. In the following sections, a review of SFR fuel pin design philosophy and historic material limits are discussed. This is followed by the recent progress in the development and properties of an advanced NFA for cladding and wrapper/duct applications. The alloy is dubbed OFRAC (Oak Ridge Fast Reactor Advanced Fuel Cladding). An important milestone, demonstrating the ability to manufacture a cladding geometry from this material is then showcased. Finally, a discussion on key focus areas to enable full deployment of OFRAC as cladding and core structure materials in SFRs is offered. 2. SFR cladding design philosophy and historic limits In order to make clear the benefits of modern NFAs as SFR cladding, it is useful to briefly review the operational history, lessons learned, and the design philosophy of this application. Given the well-documented benefits and the long-term experience with metal fuels [20e24], they remain the current focus for further development and application as driver fuels for the upcoming SFRs in the US, including the recently proposed Versatile Test Reactor (VTR) [25e27]. In the meantime, oxide and other ceramic fuels were developed and used in SFRs both inside [28,29] and outside

the U.S [30]. and are used today to power these systems [31]. The reader should note that many of these desired cladding properties discussed here are applicable to both metal and ceramic fuels in SFRs. The choice of cladding material has had a great influence on the course of fuel pin development for SFRs [32]. Experimental Breeder Reactor-I (EBR-I) that went into operation in 1951 in Idaho was a sodium-potassium cooled system. It used uranium metal clad in 347 austenitic stainless steel (SS) as its Mark I fuel. This fuel system proved useful to a burnup of only ~0.3 at% FIMA (fissions per initial metal atom), reaching the maximum of what was achieved with other fuel variants tested in this reactor [20]. Type 347 SS, and a more-tightly-specified variant, was selected as the material for incore structures, including Mark-I and Mark-II fuel and blanket cladding, because of its corrosion resistance with NaK alloy coolant [33]. However, issues with welding and welded joints and with embrittlement from sigma phase formation led early reactor designers to favor types 304 SS and 316 SS over 347 SS for application to early SFRs [34]. EBR-II started its operation in 1964 with its 304L SS clad Mark I fuel that proved useful to ~1.2 at% burnup. In the ensuing decades significant progress was made to qualify fuels that robustly operated to burnups of 10 at% in EBR-II [35]. Mark II fuel with solution annealed 316 SS, introduced into EBR-II in 1973, was fully qualified and operated robustly up to at least 8 at% [20], while the Mark-IIIA fuel employed 20% cold-worked 316 SS cladding for an increased burnup limit of 10 at.% [21]. Cold-working incorporated high dislocation densities into the microstructure of Type 316 SS cladding that in turn served to delay the onset of void swelling [4]. This thirty-fold increase in achievable fuel burnup without failure was made possible by elucidating the governing phenomena that resulted in fuel failure and development of engineering solutions to remedy them. Simply put, fuel failure in SFRs results from the stress exerted on the cladding as the sum result of fuel-cladding mechanical interaction (FCMI), driven by fuel swelling resulting from accumulation of gaseous and solid fission products, and pin internal pressure increased by release of fission gas to the fuel rod plenum. Fuel-cladding chemical interaction (FCCI) simultaneously exacerbates this by corroding the inner surface of the cladding, reducing its thickness and limiting its load-bearing capability. Fuel rod deformation due to void swelling in the cladding and cladding outward creep under the applied load leads to damage accumulation in the fuel cladding and increased stresses to other fuel rods and components caused by interference; under more severe conditions of fuel rod deformation, flow through fuel assemblies is impeded by reductions to coolant channel flow area. Remedies to address these phenomena and enable higher-burnup (higherexposure) capability include: 1. Reducing fuel smeared density (i.e., the area fraction inside the fuel rod occupied by solid, as-fabricated fuel material), such as by incorporating a large fuel-cladding gap, to allow for sufficient swelling. For metal fuels, accommodation of fuel swelling of around 30 vol % allows interconnection of fission gas bubbles [36] that in turn results in fission gas release into the plenum [37,38] (relieving the highest contribution to FCMI force). 2. Incorporating into the fuel rod a large gas plenum to accommodate released fission gas. 3. Utilizing high-strength cladding material with relatively-high creep and swelling resistance to withstand FCMI stress and fission gas pressure. The high strength cladding in turn exerts a constraining force on the fuel, pressing it to creep instead. 4. Incorporating barriers to limit FCCI. This paper presents recent work to improve on the third

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engineering solution. While cladding plastic deformation to accommodate fuel strain is allowable (if not desirable) in light water reactors - except in cases of excessive growth that may result in pin-assembly interaction and bowing, pin dimensional stability is of greater importance in SFRs. This stems from the fact that many of the reactivity feedback mechanisms in these fast-spectrum systems are associated with location and density of the fuel. Furthermore, small dimensional changes affect heat transfer phenomena in the tightly packed pins in SFR fuel assemblies [39]. Finally, thermal creep-induced dimensional changes are much more pronounced at SFR operating temperatures compared to the relatively low LWR operating temperatures [15,40] and the irradiation degradation phenomena of void swelling [6,41] and irradiation creep [41] are maximized in steels near SFR operating temperatures (versus minimal concerns at the lower LWR temperatures). Therefore, the principle phenomena that govern cladding deformation, namely thermal creep, irradiation creep, and swelling, need to be carefully considered and minimized to the extent possible. Since the latter two phenomena are dose-dependent, high displacement damage dose in SFRs, 150 dpa, places a serious demand on the cladding material. Although we treat irradiationinduced swelling and creep in our constitutive models as separate terms, in reality they stem from the same defect generation and transport phenomena [42] and therefore they are strongly coupled [41]. This implies that alloys susceptible to excessive swelling are bound to suffer from significant creep as well. Finally, the radiation damage dose also results in formation of defects in the material that reduce its ductility and make cladding susceptible to brittle failure during handling [43]. Except for EBR-I’s Mark-III & IV cores that used Zircaloy-2 cladding [44], Fe-based alloys have been and remain the cladding of choice for these systems. A complete summary of various fuel systems used in two of the most important U.S. SFRs is available in Ref. [21]. Austenitic (Fe,Cr,Ni) alloys have been extensively used while a sizeable database on ferritic/martensitic (Fe,Cr) alloys also exists. Austenitic alloys offer high strength and creep resistance at relevant temperatures. However, owing to their face centered cubic (fcc) crystal structure, they are susceptible to significant void swelling after an incubation period. In order to delay the onset of this high steady-state swelling rate (as high as 1%/dpa [41,45]), cold-worked or Ti-modified (D9) grade 316SS was used and they continue to remain as design options today [27]. Ferritic/martensitic (F/M) body centered cubic/tetragonal (bcc/bct) alloys on the other hand offer better resistance to swelling compared to the austenitic alloys (longer incubation period and slower steady state swelling rate afterwards, ~0.2%/dpa [41]) while suffering from insufficient thermal creep resistance and strength at relevant temperatures. Embrittlement due to increase in ductile to brittle transition temperature after low dose irradiation in F/M steels is another issue, particularly for operation temperatures below ~400500
C [3]. HT-9 steel (12%Cr) with a high density of lath boundaries and fine precipitates was the alloy of choice in this category for the US in the 1980s [4,21], whereas EM10 and EM12 steel (9%Cr) emerged as the preferred options in French SFR research [4,7]. As discussed in the previous paragraphs the desirable properties of SFR cladding are:    

high strength resistance to thermal and irradiation creep resistance to swelling at high dose (>150 dpa) ability to retain some ductility at high dose

Both the historic austenitic and ferritic alloys with tailored composition and microstructure failed to meet all of these criteria. The modern NFAs described in this paper leverage modern

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materials processing techniques and a sound understanding of radiation damage processes to tailor cladding properties and microstructure that meets this challenge and offer a viable solution to meet the demanding environments of SFRs. Specifically, high strength is made possible by grain size refinement (Hall-Petch [46,47]) that remains remarkably stable to high temperatures owing to grain boundary pinning by precipitates. Excellent thermal creep resistance is achieved by incorporation of large number density of fine (<5 nm) precipitates. These same precipitates provide a high density of sinks [48] for radiation defects that is expected to suppress irradiation creep, swelling, and embrittlement. 3. Modern NFA development and properties OFRAC is a modern nanostructured ferritic alloy that was derived from extensive experiences gained on mechanical alloying (MA) and alloy composition during the development of the NFA 14YWT [49e52]. Modern NFA’s have rapidly evolved during the past ~15 years and exhibit marked property improvements over earlier ODS alloys developed in the 1960s. The improvement in the NFA properties is mainly due to refinement in the microstructure compared to conventional ODS alloys, in particular creation of an ultra-fine grain structure and high concentration of nano-size (~2e5 nm instead of 10e100 nm for earlier ODS alloys) oxide particles. The nominal composition of OFRAC is Fee12Cre1Mo-0.3Ti0.3Nb-0.3Y2O3 (wt.%), which differs from that of 14YWT (Fee14Cre3W-0.4Ti-0.3Y2O3). The Cr level of OFRAC was lowered to 12%Cr to explore the partial transformation of bcc-ferrite matrix to fcc-austenite grains above at elevated temperatures. This differs from the excellent phase stability of the matrix in 14YWT which remains fully bcc ferritic to the melting point due to the 14%Cr level. However, as shown later in section 4, the transformation of bccferrite to fcc-austenite in OFRAC is sluggish based on tensile test results obtained after annealing at 1150
C for 8 h. The contribution of the 1%Mo addition is mainly for solid solution strengthening and may also help in retarding recrystallization and reducing the prior austenite grain size assuming that the bcc-ferrite to fcc-austenite occurs in OFRAC at high temperatures [53]. The effect of partial phase transformation in ODS alloys was shown to be beneficial to fabrication of thin wall tubing for advanced fast reactors [54e56]. The addition of 0.3%Nb allows for improved sequestration of impurity C and N so that they cannot segregate to and potentially lower the grain boundary cohesion energy that was suggested as the caused for improved high temperature fracture toughness of 14YWT [52]. Nevertheless, the arduous task of fabricating complexshaped components from ODS alloys has historically prevented their use in high performance applications. This challenge is successfully tackled here. 3.1. Processing The OFRAC alloy was produced using standard mechanical alloying techniques [57]. Powder was prepared by blending 99.7% pre-alloyed powder (Fee12Cre1Mo-0.3Ti-0.3Nb: 150/þ45 mm size range, produced by Ar gas atomization by ATI Powder Metals), with 0.3% Y2O3 powder (17e31 nm crystallite size, produced by Nanophase, Inc.,) and ball milling using the high kinetic energy Zoz Simoloyer CM08 in static Ar gas atmosphere. A steel can was filled with the ball milled powder and degassed at 300
C in vacuum and sealed. The sealed can was extruded at 850
C through a rectangular shaped die to form a bar (20.3 long  6.4 wide  3.5 cm high) for assessing the microstructure and mechanical properties or through a circular shaped die to form a master rod (20 long x 2.2 cm diameter) for fabricating thin wall tubes (section 4).

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3.2. Microstructure The important microstructural features of OFRAC are observed in Fig. 1. The grain structure observed in the backscattered electron (BSE) image obtained by scanning electron microscopy (SEM) consists of an average grain size along the extrusion (e) direction of 0.643 ± 0.067 mm and 0.468 ± 0.034 mm normal (n) to the extrusion direction for a grain aspect ratio (e/n) of 1.373. Within intragranular regions of the microstructure is a high concentration of fairly uniform distribution of nano-size oxide particles and a low number density of coarser size particles identified by element mapping using scanning transmission electron microscopy (S/TEM) x-ray energy dispersion spectroscopy (XEDS) to be enriched in Nb, Ti and C [57]. The size, spatial distribution and XEDS maps of the Nb,Ticarbides are shown in Figs. 8 and 9 of reference [57]. The Local Electrode Atom Probe (LEAP) analysis reveals that the oxide particles are enriched in YeTieO solute atoms and have an average size of 3 nm diameter (d) and number density of N~1  1024 m3 result in a large sink strength for defects (S ¼ 2pdN z 2  1016 m2). The combination of ultra-fine grains and a high concentration of nanosize oxide particles have been shown to contribute a significant magnitude of strengthening as well as high sink strength (>1016 m2) responsible for the attractive radiation tolerance of the NFA 14YWT [40,58]. 3.3. Mechanical properties The tensile properties of the OFRAC are shown in Figs. 2 and 3. Stress-strain curves of OFRAC in the as-extruded condition (Fig. 2) were obtained from tensile tests of type SS-3 miniature sheet tensile specimens (overall length of 25 mm with gage dimensions of 7.62 mm long  1.524 mm wide  0.762 mm thick) that were conducted from room temperature (25
C) to 800
C in air using a strain rate of 103 s1. The data shows a balance of high strength and good ductility over the entire temperature range. Although high work hardening was not observed after yielding, the alloy exhibited good ductility at all test temperatures, including at room temperature where the yield and ultimate strength exceeded 1.2 GPa. Fig. 3 demonstrates the high-temperature stability of the tensile properties for OFRAC. The figure compares the tensile properties of the as-extruded alloy and after subsequent annealing in vacuum (7  105 mbar) at relatively extreme conditions of 1050
C and 1150
C for 8 h. Although the yield stress (YS) is decreased after the 1150
C annealing (e.g. by ~100 MPa at room

Fig. 2. Stress-strain curves over the temperature range from 25
C to 800
C.

temperature), there are only minor variations in ultimate tensile strength (UTS) for the as-extruded vs. high temperature annealed specimens over the 20-800
C tensile test temperature range (Fig. 2a). Plus, the YS values for the annealed and extruded conditions converge at 600
C and 800
C. The post-extrusion annealing treatments influenced the ductility in a more complicated manner, but generally lowered the uniform elongation (UE) above ~200
C while slightly improving in most cases the total elongation (TE). Overall, the tensile results on the high temperature annealed specimens suggest that the as-extruded specimens were relatively resistant to property deterioration for annealing conditions as high as 1150
C (~0.80 TM, where TM is the absolute melting temperature). The thermal creep properties of OFRAC were determined using the strain-rate jump (SRJ) method. This SRJ performance metric was compared with that of a late 1960s-era conventional HT-9 F/M steel (Fee12%Cre1%MoVW) and three austenitic stainless steels, which are historic candidate materials for fast reactor fuel cladding [21]. In the SRJ test, the specimen is deformed by applying a constant strain rate that causes the flow stress to increase until it reaches a steady-state followed by increasing the strain rate and repeating the deformation process. These tests were conducted on an MTS servohydraulic uniaxial load frame in displacement control mode using SS-3 tensile specimens in the as-extruded condition. The SRJ tests were performed in air at 550
C, 600
C, 700
C and 800
C by increasing the strain rate by one order of magnitude for each SRJ test, with strain rate sensitivity examined over a total of four orders of magnitude. The SRJ tests started at 1.5  107 s1 and ended at 1.5  103 s1; each SRJ test ended when the stress no longer changed with time. The stress exponent (n) is determined from the following relationship:

 n¼

Fig. 1. Microstructure characteristics of the extruded OFRAC. (a) BSE (SEM) micrograph of the grain structure and (b) LEAP chemical map of the distribution of nano-size Y-, Tiand O-enriched oxide particles. The LEAP chemical map shows the Ti þ O isoconcentration surface.

 log_εm log s

and is determined by plotting the values of the log minimum creep rate (_εm), or strain rate from the SRJ test, against log stress (s) and measuring the slope from the exponential line fit. The derived creep properties of OFRAC are compared to published results on HT-9 [59] and three stainless steels [60,61] in Fig. 4. The calculated values of the stress exponents for OFRAC and

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Fig. 3. High-temperature stability of the tensile properties. (a) strength and (b) ductility.

Fig. 4. Comparison of thermal creep behavior of OFRAC (determined from SRJ tests) with (a) HT-9 [59] and (b) three austenitic stainless steels [60,61]. The calculated stress exponent (n) is shown for each set of creep data.

the published values for HT-9 and stainless steels are shown in the figure legends. The data shows similar creep stress exponent behavior for OFRAC over the temperature range of 550
C-800
C. The stress exponent values of OFRAC were comparable to that of HT-9 (Fig. 4a). However, the test temperature required to induce a given creep rate at the same applied stress was much higher for OFRAC compared to HT-9. Equivalently, the stresses required to induce a comparable creep strain rate at a given test temperature were markedly higher in OFRAC. For example, the stress to induce a strain rate of 106 s1 at ~600
C in OFRAC was ~300 MPa higher compared to HT-9 (500 vs. 200 MPa). For a comparable applied stress of 150 MPa, a similar creep rate of 107 s1 was achieved in HT-9 at ~600
C and in OFRAC at ~800
C. Similarly, at higher applied stress levels a creep rate of 107 s1 was achieved in HT-9 at

540
C and in OFRAC at 700
C; collectively these results imply an improvement of ~150-200
C operating temperature limit for OFRAC compared to HT-9. OFRAC also exhibited superior thermal creep performance compared to Type 316 and 316LN austenitic stainless steel (Fig. 4b). For example, a creep rate of 106 s1 was achieved in the stainless steels at 650
C for applied stress levels of 200e250 MPa and in OFRAC similar creep rates occurred for applied stress levels of 150e250 MPa at much higher test temperatures of 800 and 700
C, respectively. The thermal creep stress exponent values for 316 SS and 316LN SS at 650
C were a factor of two to three lower than that of OFRAC. The lower values of stress exponent suggest that diffusional creep mechanisms may be occurring in the stainless steels at 650
C while dislocation (power law) creep mechanisms may be occurring in OFRAC, even up to

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Longer term ([10,000 h) thermal creep testing of OFRAC is needed to confirm the short-term SRJ predictions. However, published long term thermal creep tests on 14YWT and other similar 9e14%Cr NFA’s [40,53,62] have consistently observed superior high temperature thermal creep strength. 4. Modern NFA cladding demonstration There are several challenges that in the past have hindered wide acceptance of ODS alloys in demanding high-temperature nuclear applications that also apply to NFA. First and foremost is the ability to fabricate geometries that are relevant to functional components from these materials and to weld the components without degrading the important microstructural features that make these materials attractive for high performance applications. For advanced SFRs with high fuel burnup targets, demonstration of the ability to fabricate and weld thin wall tubing for fuel cladding without property degradation to retain high fission gas pressures at end-of-life is paramount. Historically, the scale-up production, or scalability, and the costs associated with commercial production of ODS alloys have been prohibitive for many applications. Thus, the question of whether high performance thin wall tubing can be fabricated from modern NFAs and the issues influencing the use of modern NFA as advanced fuel cladding for sodium fast reactors need to be addressed. 4.1. Cladding production The fabrication of a thin wall tube from OFRAC was explored in collaboration with Nippon Nuclear Fuel Development Co., Ltd (NFD), Japan. The general procedure involved producing a master rod of OFRAC at ORNL that was shipped to NFD, who then produced a master tube with an initial dimension of 18 OD  12 ID  200 L, where units are in mm, OD and ID are outside and inside diameters, respectively, and L is length. The initial master tube was pilgerrolled at room temperature, i.e. cold-pilger, in four passes, with each pass resulting in 50% reduction in area. Following each pilger pass, an optimum annealing condition was determined using a small sample that was cut from the end of the tube and annealed at temperatures ranging from 700
C to 1,100
C for 30 min followed by characterization using optical microscopy and micro-hardness measurements. . The annealing conditions used between each pilger pass run were at temperatures that did not cause recrystallization or coarsening of the Y-, Ti- and O-enriched nanoclusters. Inspection of the OFRAC tube following each of the four cold-pilger passes showed that no cracking occurred. The total reduction in tube wall thickness was 84% (3.04e0.51 mm) with an increase in length of 790%. The final length of the cladding was 1.78 m, almost twice the length of reference U.S. SFR fuel pin designs [21]. The resulting OD of the tube was 8.51 mm. Fig. 5 shows the 1.78 m long tube and a composite micrograph of light microscopy images of the thin wall tube cross-section. The precision of the wall thickness was ±1 mm. 4.2. Cladding properties Fig. 5. Optical images of the OFRAC thin wall tube (top) and the cross-section view of the wall thickness (bottom).

800
C. The high dependence of minimum creep rate with changes in stress indicates that quite low creep rates will occur for OFRAC at higher temperatures and stresses compared to HT-9 and stainless steel, which is essential for achieving thermal creep performance requirements (while simultaneously achieving high radiation resistance to void swelling) of fuel cladding in sodium fast reactors.

Although no cracks formed during fabrication of the OFRAC thin wall tube by cold-pilger rolling, the large magnitude of deformation associated with the wall thickness reduction and tube length increase severely altered the grain structure in terms of size, morphology and texture. Fig. 6 shows S/TEM micrographs of the grain structures observed in the master rod after annealing at 1150
C for 8 h (Fig. 6a) and after cold-pilger rolling into the thin wall tube. The grains in the annealed master rod are mostly equiaxed in shape and contain network dislocation and sub-grain

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Fig. 6. S/TEM BF micrographs showing the elongated grains that formed during fabrication of the OFRAC thin wall tube. (a) extruded condition and (b) cold-pilger rolled thin wall tube.

Fig. 7. EBSD pole figure maps of the (a) annealed (1150
C/8 h) master rod and (b) cold-pilger thin wall tube. Both EBSD maps were obtained from the longitudinal (L) orientation, i.e. along rod/tube axis.

structures. Several dark particles are observed within the grains that are consistent with the Nb,Ti-carbides [57]. After the cold pilger rolling, the grains become highly elongated are oriented in the cold working direction of the thin wall tube. The average width of the elongated grains is ~107 nm and are typically longer than 3 mm in length, or grain aspect ratio of >30 (l/w). Fig. 7 shows electron backscattered diffraction (EBSD) pole figure maps comparing the microstructure of the annealed master rod with that of the cold-pilger rolled thin wall tube. Both EBSD pole figure maps were obtained from the longitudinal (L) orientation to show the grain structure along the axis of the rod (Fig. 7a) and tube (Fig. 7b). The normal direction to the grains is ascertained using the color inverse pole figure (IPF) inset in Fig. 7a. The EBSD pole figure map of the annealed master rod (Fig. 7b) shows bimodal grain size distribution. The coarse grains possess significant sub-grains denoted by low misorientations <15
[57]. These coarse grains containing subgrain structures viewed by EBSD are not as evident as observed in the BSE micrograph (Fig. 1a). Research is in progress to determine if the size and number of the YeTieO particles are different between the coarse and fine grains and the conclusion will be published in the future. After severe cold working by pilger rolling, some remnant of the bimodal grain size is present in addition to the

significant elongation of the grains that occurred. For the annealed master rod (Fig. 7a), a large number of grains are oriented in the <100> and <111> directions in the L orientation, which is consistent with a strong <110> a-fiber texture parallel to the extrusion axis. The severe cold working by pilger rolling does not appear to alter the strong <110> a-fiber texture significantly since many of the grains still are oriented in the <100> and <111> directions (Fig. 7b). For a better understanding of the texture changes induced by deformation during cold-piler rolling are shown in Fig. 8 for the annealed master rod and the cold-pilger rolled thin wall tube. For the annealed master rod (Fig. 8a), the grains are broadly oriented around <110> in the T (normal) orientation (Fig. 8b) and between <111> and <100> in the L orientation (Fig. 8a). For the cold-pilger rolled tube, the deformation caused the grains to become more aligned to <110> for the T orientation (Fig. 8d) and split between <100> and a branch from <111> to <110> for the L orientation (Fig. 8c). Thus, these results suggest that the severe cold working intensifies the <110> a-fiber texture. The tensile properties of the OFRAC pilger rolled tube in the hoop and axial orientations were obtained using ring- and tubetype specimens, respectively, that are illustrated in Fig. 8. The room temperature stress-strain curves for both orientations show a

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Fig. 8. IPF showing the grain texture of the annealed master rod (a) and (b) and cold-pilger thin wall tube (c) and (d). Figures (a) and (c) show the L orientation and (b) and (d) show the T orientation.

highly anisotropic microstructure (Figs. 6 and 7) during cold pilger rolling. The ring specimen possessed a smaller elongation compared to the tube specimen since the gage is normal to the elongated grains in the microstructure. The increase in strength for the ring specimen may not be as great a problem as is the decrease in ductility since internal stresses on the fuel cladding during reactor operation will exert higher stresses in the circumferential direction than along the thin wall tube direction that may increase the propensity for fracture. Although no cracks formed during cold working by pilger rolling, studies are in progress to determine if post pilger annealing can recover the ductility with a concomitant decrease in strength. In addition, further research of the fabrication method of cold pilger rolling is required for attempting to decrease the high degree of anisotropy that develops in the microstructure of OFRAC by severe cold working. 5. Considerations towards full scale deployment 5.1. Scalability

Fig. 9. Comparison of tensile properties between the pilger rolled tube and the extruded bar of OFRAC.

large increase in strength and decrease in elongation compared to the as-extruded condition of OFRAC, which was obtained using a dog-bone shaped SS-J3 specimen (overall length of 16 mm with gage dimensions of 5.00 mm long  1.20 mm wide  0.75 mm thick). These detrimental changes in tensile properties of the OFRAC tube are due to the significant cold working that led to the

The commercial production of ODS alloys in the past by the International Nickel Company (INCO) and Plansee demonstrates their scalability potential [63e65]. INCO developed ODS alloys based on Ni (MA754 and MA6000), Fe (MA956 and MA957) and Al (IN-9052) with annual production capabilities of about 350 tons in past years [66]. Many of the applications were for hightemperature gas turbine engines and components such as blades and vanes except for MA957, which was developed for fuel cladding for fast neutron reactors [63,64,66]. Plansee developed the PM2000 alloy for overcoming some of the technical problems associated with INCO MA956 in applications requiring improved stressrupture creep properties and corrosion resistance in gaseous environments at very high temperatures [65]. Commercial products

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of PM2000 that were produced by Plansee included sheet, bar, tube and wire. Thus, scale-up production of modern NFAs including OFRAC is feasible based on these examples. However, there are some technical issues with the mechanical alloying (MA) approach for producing large quantities of ODS alloys that pertains to NFA OFRAC. It is generally accepted that the composition of powders and the processing and fabrication parameters used in MA, including ball milling conditions and consolidation method, affect the microstructure of ODS/NFA alloys which in turn influences the performance in terms of strength, ductility, fracture toughness and degree of anisotropy in mechanical properties. It is also known that inconsistencies in processing and fabrication of ODS/NFA alloys often results in heat-to-heat variations in mechanical properties, which is a concern. This problem may be reduced by carefully controlling the composition of the powders during production, such as by gas atomization, and minimizing the contamination of the powders during high energy ball milling. The most likely source of contamination during high energy ball milling is the unwelcome pickup of interstitial C, N and O atoms. As an example, high levels of O and N were present in the NFA 14YWT (SM10 heat) that most likely occurred via contamination of the powder by ingress of air in the mill chamber during ball milling; this led to very high strength properties but poor fracture toughness properties at elevated temperatures for this heat [67]. It is also important that once the optimum MA parameters for producing consistent heats of ODS/NFA alloys are obtained at the research and development (R&D) level, these parameters must be scalable to larger ball mills and to potentially different powder consolidation methods than was used at the R&D level. Ultimately, a robust quality assurance program will be necessary to guarantee uniform properties across batches during large scale manufacturing of NFAs. 5.2. Joining There are many techniques available for welding conventional alloys that are not suitable for joining ODS alloys. For joining ODS alloys, methods that do not destroy the dispersion of oxide particles and significantly alter the grain size are required. This requirement is paramount for modern NFA due to the refinement in grain and oxide particle sizes to meet the demands for radiation tolerance in fast reactor applications. A recent review of joining techniques for a wide range of ODS alloys revealed solid-state approaches were the least disruptive on the microstructure and ensuing mechanical properties [68]. This includes diffusion bonding, resistance welding, explosive bonding, friction-based welding and magnetic impulse welding. Of the different solid-state joining techniques, diffusion bonding, resistive welding and friction-based welding have been investigated the most for joining ODS/NFA alloys. Studies on diffusion bonding were conducted on MA956 and PM2000 using transient liquid phase (TLP) bonding [69] and on 9Cr ODS alloys [70,71] by applying uniaxial pressure on components at elevated temperatures. For the TLP technique, a thin boron film was deposited on polished surfaces of MA956 and PM2000 and heated to an elevated temperature below the melting point of the Fe-base ODS alloys under compressive stress to form a liquid eutectic with the Fe-base substrate that solidifies to form the bond. In general, fine grain microstructures of MA956 and PM2000 improved the microstructural continuity across the bond layer [69]. The diffusion bonding technique was applied to 9Cr ODS alloys, including ODS Eurofer, to induce transformation between bcc a-Fe and fcc g-Fe that causes the bonding interface to migrate by diffusional processes. The results from these studies reported little change in tensile properties and microstructure between the diffusion bond

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layer and base metal [70,71]. However, diffusion bonding approaches require smooth non-oxidized surfaces for the components to be bonded and require relatively long times at elevated temperatures. Resistive welding, which combines pressure with resistive heating to cause deformation at the interface between two components, has been investigated for joining PM2000 [72] and MA957 [73]. Heating occurs by the Joule effect as current is applied and bonding results from application of the contact force between the components to be bonded. These studies showed that the electrical energy dissipation (ED) in the components will affect the weld quality. At low ED, incomplete bonding can occur while high ED may cause recrystallization and coarsening of the oxide particles that degrades the mechanical properties. Nevertheless, welds exhibiting small changes in grain structure and mechanical properties can be achieved using an ED parameter between the low and high ED extremes. Recent investigations using friction stir welding (FSW) have been conducted on NFA MA957 [74,75] and NFA 14YWT [76,77]. This solid-state joining method involves high speed rotation of a pin and shoulder tool that travels along the joint separating two components, generating heat by friction that softens the components allowing material to plastically deform and flow around the spinning tool to form the weld. A related friction-based joining method involves rotating one component against another at high speeds under an applied load that generates heat that softens the components until they deform and bond together; this technique is particularly suitable for joining end caps on fuel cladding. In the MA957 studies [73,74], advanced microstructural characterization methods such as TEM, SANS and APT were used to show that welds produced by FSW consisted of uniform, fine equiaxed ferrite grain structure, a high density of dislocations and slightly coarsened nano-size oxides that resulted in a modest reduction in hardness compared to the base metal. The FSW investigation on 14YWT showed that the ultra-small grains in the weld zone had increased by ~4 times and defects such as a wormhole on the advancing side of the weld zone and small pores associated with other joints and interfaces were observed [76]. However, TEM and APT showed no statistical difference in the size and number density of the nanosize oxides between the weld zone and base metal [76,77]. Similar to the MA957 study, the hardness of the FSW weld zone decreased by ~20% compared to the base metal. These results indicate FSW may not degrade the nano-size oxide dispersion in the welds of ODS/NFA alloys, but further studies are required for optimizing the FSW parameters to reduce defects such as voids and wormholes. In conclusion, several solid-state welding approaches show much promise for welding modern NFA without degrading the important microstructural features. 5.3. Cost The costs for producing ODS/NFA alloys will always be higher than that for producing conventional alloys by the ingot metallurgy approach. The higher costs in production are primarily because of the mechanical alloying approach that incorporates both powder metallurgy principles and high energy ball milling of powder, which was quintessential when MA was developed and patented by INCO [63,64]. Thus, additional costs compared to conventional ingot metallurgy approaches are incurred for producing the powders, i.e. gas atomization, and for processing the powders into solid products by MA. To illustrate this point, the estimated cost for producing PM2000 extruded stock is ~$345/kg, which is considerably higher than that associated with several conventional steels and stainless steels that range from $5.50 kg to $28.20/kg [78]. The total cost for producing the two master rods of OFRAC in the

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present study was ~$460/kg, but this value is considered to be more than two times higher than what would be a commercial largerscale master rod cost due to the high costs for exploratory oneof-a-kind alloy research and development and the relatively small quantity of powder (~2.5 kg per rod) that was processed by mechanical alloying in this study. Therefore, if OFRAC was produced commercially, then the estimated cost per kilogram is expected to be similar to that of PM2000. The cost for fabricating Zr alloy fuel cladding for LWRs was estimated to be $20-$30 per meter corresponding to $210-$310 per kg for LWR cladding geometry, with a 50% increase in cost to $30-$45 per meter for new technology criteria being considered by the LWR Sustainability Program [79]. As mentioned previously, the 1.78 m long thin wall tube of OFRAC was fabricated in collaboration with Nippon Nuclear Fuel Development Co., Ltd (NFD), Japan who contracted Zirco Products, which commercially produces Zircaloy cladding tubes for nuclear power plants. Since the same pilger rolling equipment and procedures were used, the costs for fabricating the thin wall tube of OFRAC are known to be similar to that incurred for fabricating Zircaloy cladding tube. Thus, the major cost involved for producing advanced fuel cladding from OFRAC is the higher cost associated with producing the alloy by mechanical alloying. There are other factors to consider besides the high cost of production for ODS/NFA alloys, especially when considering the modern NFA OFRAC for advanced fuel cladding in sodium fast reactors. As pointed out by Busby [78], fast reactors will require similar economics and operation costs for power generation to that of current generation light water reactors while also improving safety and reducing the radioactive decay period of actinide and long-lived fission isotopes through transmutation. Despite the high production costs, the improved strength and creep properties of NFA alloys and the high radiation resistance (due to high point defect sink strength associated with the ultra-fine features of their microstructure) can potentially reduce the levelized cost of electricity (LCOE) of SFRs compared to conventional stainless steel or ferritic steel cladding designs, by allowing higher-burnup (higherexposure) operation. The LCOE considers the cost for building and operating a reactor that generates power over a period per total energy output of the reactor over the same period. Higher-burnup capability reduces fuel cycle costs by improving fuel utilization in once-through fuel systems and by reducing the number of recycle passes needed in a fuel-recycling energy system. Either improvement reduces the amount of fuel handling, fuel fabrication cost, and fuel hardware cost per kW-h of energy generated. (Currently, fuel costs in commercial LWRs are roughly 5% of the LCOE, compared to operations and maintenance (O&M) cost of about 15% LCOE [80,81]).However, fuel costs are projected to be a higher proportion of the LCOE in some Generation IV energy systems with simpler plant designs that reduce operating cost associated with staffing and with more complex fuel designs that seem likely to cost more than LWR fuel designs; for example, a recent study projected TRISO fuel costs in a high-temperature gas reactor to be nearly 15% of LCOE [81]. The improved creep properties of NFA OFRAC compared to that of HT-9 and austenitic stainless steels may allow increased operating temperatures up to ~700
C, i.e., close to the estimated upper operating temperature limit for steels due to corrosion in lowoxygen sodium [82]. In addition to allowing increased efficiency with a steam Rankine cycle, higher-temperature operation could also enable use of relatively compact (lower capital cost) supercritical CO2 or other Brayton power conversion systems rather than steam Rankine systems, which would lower construction costs, improve power generation efficiency, and improve accident safety scenarios by replacing highly exothermic Na-water reactions with less aggressive NaeCO2 reactions [83]. Long-term thermal creep

tests in Na are needed to confirm the upper use temperature of advanced steels such as OFRAC, since decarburization could potentially cause significant strength loss [82]. The high sink strength associated with the ultra-fine microstructural features of NFA OFRAC will improve the radiation resistance behavior for the cladding and thereby may enable neutron irradiation doses greater than 200 dpa to be attained, which is highly valuable for achieving the SFR fuel deep burnups for optimized actinide conversion and transmutation. 6. Summary A new nanostructured ferritic alloy, OFRAC, has been developed to address the challenging material requirements for advanced sodium fast reactor cladding and core internals. Leveraging decades of experience, the OFRAC’s composition and microstructure have been tailored to meet strength as well as creep and swelling resistance requirements where historic austenitic and ferritic SFR cladding materials (i.e. D9 and HT9) fell short. Excellent strength and thermal creep resistance were confirmed with experiments as a part of this study. While irradiation of OFRAC to high dose is currently ongoing, based on the current state of knowledge regarding radiation damage processes the alloy is expected to exhibit excellent swelling and irradiation creep resistance. Finally, and of key significance, production of long (>1 m) cladding segments with ~1 mm diametral tolerance from OFRAC was demonstrated in this study. By addressing these key areas, the authors intend to emphasize to the reactor design community that material system performance targets previously unattainable may now be in reach for SFR cladding. Key focus areas towards full scale commercial deployment of this alloy are scalability, joining, and production cost management, none of which are deemed insurmountable. It is perceived that a focused government- or industry-led development program may address all of these to enable full scale deployment of OFRAC. 7. Data availability statement Data will be made available upon request. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements The authors are grateful to Kan Sakamoto (NFD) and Kory Linton (ORNL) in facilitating the tube manufacturing. Maxim Gussev (ORNL) performed tube mechanical testing. Kevin Field provided useful comments on the manuscript. This work was supported by the US Department of Energy, Office of Nuclear Energy, Advanced Fuels Campaign. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jnucmat.2019.151928. References [1] A technology roadmap for generation IV nuclear energy systems, generation IV international forum. http://www.gen-4.org/PDFs/GenIVRoadmap.pdf, 2002. [2] W.R. Corwin, US Generation IV reactor integrated materials technology

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