Characterization of surface layers formed on DU10Mo ingots after processing steps and high humidity exposure

Journal of Nuclear Materials 514 (2019) 28e39

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Characterization of surface layers formed on DU10Mo ingots after processing steps and high humidity exposure Tiffany C. Kaspar a, *, Christina L. Arendt b, Derek L. Neal c, Shawn L. Riechers a, Crystal Rutherford c, Alan Schemer-Kohrn c, Steven R. Spurgeon c, Lucas E. Sweet b, Vineet V. Joshi c, Curt A. Lavender c, Rick W. Shimskey c a b c

Physical and Computational Sciences Directorate, Pacific Northwest National Laboratory, Richland, WA, 99354, USA National Security Directorate, Pacific Northwest National Laboratory, Richland, WA, 99354, USA Energy and Environment Directorate, Pacific Northwest National Laboratory, Richland, WA, 99354, USA

h i g h l i g h t s
Surface characterization of monolithic UeMo fuel elements.
Efficacy of acid etching vs. electropolishing to remove surface oxide.
Characterized by x-ray photoelectron spectroscopy and scanning electron microscopy.
Surface roughness enhances oxidation in high humidity environment.

a r t i c l e i n f o

a b s t r a c t

Article history: Received 3 August 2018 Received in revised form 11 October 2018 Accepted 13 November 2018 Available online 14 November 2018

The design of monolithic UeMo fuel elements fabricated from low-enriched uranium for use in highpower research reactors requires bonding of the fuel foil to either Al cladding or a Zr barrier layer. Processing of the UeMo ingot to final foil form has the potential to generate surface layers on the foil that differ from the bulk, metallic UeMo. The interfacial properties between the UeMo and Zr or Al cladding layers will then be determined by these surface layers. We use x-ray photoelectron spectroscopy, crosssectional scanning electron microscopy, and atomic force microscopy to characterize the composition, oxidation state, and morphology of the surface layers that form after hot rolling and cold rolling depleted Ue10 wt% Mo alloy (DU10Mo). A thick uranium nitride layer is observed after hot rolling, although its origin is likely from a previous processing step. The efficacy of acid etching in HNO3 is compared to that of electropolishing in H2SO4 to remove surface nitride and oxide layers, and both methods are found to be similarly effective. Both laboratory (low humidity) air exposure and longer rinse times in water are shown to promote the formation of surface oxide layers. Exposure of both acid-etched and electropolished DU10Mo foils to humid air (97% relative humidity) for six weeks results in formation of a thick oxide layer due to corrosion. The oxide layer on the acid-etched foil is thicker and more highly oxidized than the oxide layer that forms on the electropolished foil, and these differences in oxidation behavior are attributed to higher surface roughness on the acid-etched foil. In general, Mo is found to play a role as a sacrificial element, typically exhibiting a larger ratio of Mo6þ/Mo4þ than U6þ/U4þ. This is unexpected, given the greater thermodynamic driving force to form U oxides than Mo oxides. © 2018 Elsevier B.V. All rights reserved.

Keywords: U-Mo USHPRR XPS Surface oxide Acid etch Electropolish Corrosion

1. Introduction Alternative fuel compositions and forms are being explored to

* Corresponding author. E-mail address: [email protected] (T.C. Kaspar). https://doi.org/10.1016/j.jnucmat.2018.11.022 0022-3115/© 2018 Elsevier B.V. All rights reserved.

meet the U.S [1,2]. and European [3] goals to convert civilian research nuclear reactors from highly enriched uranium (235U typically >90 wt%) to low-enriched uranium (LEU, 235U < 20 wt%) fuel. For these high-power applications, the required 235U density necessitates the use of metallic LEU fuels. Alloying U with 6e10% Mo stabilizes the radiation-tolerant g-U phase as the material is quenched from high temperature [4,5].

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Extensive studies of g-UeMo alloys (4e12 wt% Mo) as dispersed particles embedded in Al 6061 have been undertaken [1,3,6e9], and it was found that at higher operating temperature and higher burnup, interfacial reactions occur between the UeMo particles and the Al alloy matrix to form uranium aluminides. This leads to undesirable swelling or “pillowing” of the dispersion fuel plates. As an alternative design that reduces the contact area between UeMo and Al, and simultaneously increases the 235U density, monolithic g-UeMo fuel plates are being developed that employ a Zr barrier between the Al and UeMo [2,10e14]. Fabrication of a monolithic fuel element requires UeMo ingots to undergo several processing steps such as heat treatments (homogenization) [2,15] and mechanical deformation (hot rolling, cold rolling) [16e19] to the final foil thickness. Each of these steps has the potential to generate surface layers on the foil that differ from the bulk, metallic UeMo [11]. These surface layers will then be the surface that is bonded to the Zr barrier layer, and the interfacial properties [14] will be determined in part by these surface layers [11]. It is expected that thick oxide, nitride, and/or carbide layers on the UeMo surface will be detrimental to the final properties of the clad fuel element. In addition, changes in composition of the surface region during processing can occur. For example, a recent study of U0.8Zr0.2 observed Zr segregation to the surface during heat treatments [20]; Mo segregation or depletion at the surface of UeMo will likely affect the bonding properties. Thus, it is necessary to understand the surface layers that form at each processing step, the efficacy of methods such as acid etching [11,12,21e23] or electropolishing [22] to remove them, and the effect of air and humidity exposure between steps or during storage, in order to

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develop material- and time-efficient fabrication protocols. In this work, we characterize the surface layers that form on depleted uraniume10 wt% Mo alloys (DU10Mo, as a simulant for LEU-Mo) as it undergoes processing steps analogous to those expected in the production of LEU-Mo fuel plates. A key component of this study is the use of x-ray photoelectron spectroscopy (XPS), which is a surface-sensitive spectroscopic technique that provides information on composition and valence charge states of the nearsurface region (approximately the top 5e10 nm) of a material. Using XPS, scanning electron microscopy (SEM), and atomic force microscopy (AFM), we explore both the effect of mechanical processing steps and subsequent cleaning procedures, and the effects of corrosion in air and in a high-humidity environment such as might be found in some fabrication facilities, on the surface of DU10Mo. 2. Experimental methods Nitric acid, HNO3, is well established as a safe and efficient etchant to dissolve metallic U [22]. Note that when U is alloyed with Mo, the dissolution process in high concentrations of HNO3 can lead to the precipitation of poorly soluble MoO3 and uranyl polymolybdate species; to avoid this precipitation when etching U10Mo alloys, the HNO3 concentration should be less than 10.7 M (55.7 wt %) [22,23]. As an alternative to dissolution via etching, electropolishing can remove both surface layers and surface protrusions. Electropolishing U in 6 M sulfuric acid, H2SO4, at 6 V against a Pt cathode results in a clean, polished surface [22]. Developmental DU10Mo ingots were fabricated at the Y12

Table 1 DU10Mo processing and treatment steps, and the surface composition after each processing step (as measured by XPS). Carbon was present on all surfaces as a result of air exposure. Sample

Treatment

Mo/(U þ Mo), at%

Contaminants

A B

Hot roll Hot roll, HNO3 etch, DI rinse 2 s Hot roll, HNO3 etch, DI rinse 30 min Cold roll to 0.635 mm Cold roll to 0.635 mm, HNO3 etch Cold roll to 0.203 mm Cold roll to 0.203 mm, Electropolish Cold roll to 0.203 mm, Electropolish, Corrosion study

12.3% 25.6%

Er None

24.1%

None

3.9% 29.6% 25.9% 36.7% Spot 0: Spot 1: Spot 2: Spot 3: Spot 0: Spot 1: Spot 2:

Ca, Zr Ca, Zr Ca S S, Y, Si

C D E F G H

I

Cold roll to 0.203 mm, HNO3 etch, Corrosion study

48.1% 36.1% 36.6% 35.2% 35.7% 40.3% 27.0%

Ca, K, S, Y, Si

Fig. 1. Processing scheme for DU10Mo material. Red letters indicate where XPS data were collected, and correspond to the samples listed in Table 1. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

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National Security Complex (Oak Ridge, TN), then cold rolled to 5.08 mm (0.200 ) and homogenized at 900  C for 48 h by BWX Technologies (BWXT, Lynchburg, VA). This material as received from BWXT is referred to below as “ingot” material. Ingots were then processed in several steps, as detailed in Table 1 and Fig. 1. Hot rolling was accomplished by wrapping the ingot in Zr foil (to act as an oxygen getter) and heating in a furnace at 700  C for 15e20 min in air, then rolling to a final thickness of 1.02 mm (0.04000 ). After removing the Zr foil getter, the DU10Mo foil was cleaned to remove surface layers by etching in 35 wt% HNO3 for 5 min, followed by rinsing in deionized (DI) water for 2e5 s (unless noted otherwise). Cleaned foils were cold rolled to a final thickness of 0.635 mm (0.02500 ) or 0.203 mm (0.00800 ), followed by annealing at 700  C in an inert Ar gas atmosphere. A second cleaning step was then employed, either another acid etch in 35 wt% HNO3, or electropolishing in 30 wt% H2SO4 at an applied current of 1e2 A; both cleaning treatments were followed by rinsing in DI water for 2e5 s. DU10Mo specimens were exposed to controlled temperature and humidity conditions to determine the effects of cleaning methods and relative humidity on surface layer formation (i.e., corrosion) at room temperature. Two specimens of ingot material and four foil specimens cold rolled to 0.635 mm were prepared, each with a surface area of about 1e2 cm2. Half of the specimens were etched in 35 wt% HNO3 for 5 min, while the other half were electropolished in 30 wt% H2SO4 at 1e2 A for 5 min. The humidity was controlled using saturated potassium sulfate (K2SO4) salt solutions in sealed vessels at a fixed temperature to maintain a stable relative humidity. Potassium sulfate provides a relative humidity of 97.0% at 30  C (86  F). Approximately 2 g of saturated salt solution was added to each of six vessels (22 mL 4700 series Parr vessels). Specimens and temperature/humidity data loggers were suspended above the solution using platinum wire. The Parr pressure vessels were sealed at atmospheric pressure and placed in a water bath at 30  C for 1008 h (6 weeks). The specimens were weighed using a balance accurate to ±0.0001 g and photographed before and after the corrosion study. For XPS analysis, samples (1  1 cm2 or smaller) were cut from the larger coupons and foils. All samples were measured asreceived, with no additional cleaning or degreasing procedure. In all cases, the samples were transferred through air to an oxygenfree glove box attached to the XPS chamber. For the cleaned (etched or electropolished) foils, the transfer occurred within 10e15 min of the final rinse step to minimize air exposure. In the glove box, samples were mounted on copper sample stubs with carbon tape, then loaded into the vacuum system. XPS data were collected with a Kratos Axis Ultra DLD spectrometer, using monochromated Al Ka radiation (l ¼ 1486.6 eV). The spectrometer was operated in hybrid mode, with pass energies of 160 eV for survey spectra and 40 eV for high resolution spectra. The analysis spot size under these conditions is approximately 700  300 mm. A lowenergy electron gun was employed as a charge neutralizer during all acquisitions. XPS spectra were typically collected from two or more regions on each sample. Unless noted otherwise, similar results were obtained from each region, and only one region from each sample is discussed below. Low resolution survey spectra (collected over a wide energy range) were used to identify the presence of contaminants, and to semi-quantitatively determine the relative amounts of U, Mo, O, and other identified elements on the surface. For analysis, CasaXPS software was used to subtract a Shirley background from each spectrum before peak areas were calculated. Sensitivity factors [24] were used to quantify the atomic fractions presented here. These sensitivity factors are likely not strictly correct for our particular spectrometer setup and material form (oxides, not metals), and thus the calculated atomic fractions are not highly accurate. However, the calculated values are

representative of trends in the Mo/U ratio. In some cases, it was necessary to shift peaks by 0.1e0.2 eV to align. High resolution spectra (collected over selected, narrow energy ranges) were used to qualitatively identify the oxidation state of U, Mo, O, C, Er, and S. SEM images and compositional x-ray data on the hot-rolled and cold-rolled samples were collected on a JEOL JSM-7600F SEM fitted with an Oxford X-Max 80 mm2 detector. Energy-dispersive spectroscopy (EDS) data were analyzed using Oxford AZtec nanoanalysis software version 3.2. Most data collection occurred at 30 keV beam voltage and a nominal beam current of 6 nA. Crosssectional secondary electron imaging of the corroded coupons was performed with an FEI Helios 660 DualBeam FIB/SEM using a 5 kV accelerating voltage and 3.2 nA current. Cross-section cuts were performed using an ion beam accelerating voltage of 30 kV and a current of 47 nA, followed up by cleanup cuts at 2.7 nA. X-ray diffraction (XRD) data was collected on the corroded samples using a Rigaku Ultima IV diffractometer equipped with a Cu sealed tube X-ray source operating at 1.6 kW, and a linear position sensitive detector covering 5 on a 285 mm radius goniometer. q-q scans were conducted with a 0.25 divergence slit, 2.5 incident and receiving side solar slits, a Ni foil filter to reduce contributions from Kb, 0.02 count binning steps and a 1 /min scan rate. The zero error and instrument line profiles were obtained from diffraction data collected on the NIST SRM 660c (LaB6). The data was analyzed using the TOPAS version 6 (Bruker) software. Surface topography was measured with AFM using an Oxford Instruments Asylum MFP-3D Infinity. Bruker MSLN tips with a nominal spring constant of 0.03 N/m were used in lateral force (contact) mode at forces between 0.2 and 1.8 nN. Gwyddion (http:// gwyddion.net/, V 2.38) was used for image processing and analysis. Plane subtraction and line correction were used with masking where appropriate. Some obvious line spikes (noise) were removed from 3D representations using the Remove Individual Grains tool. 3. Results Exposure of bare, metallic DU10Mo to air at room temperature will result in the immediate formation of an oxide layer on the exposed surfaces. Oxidation of U to U4þ (as in UO2) and Mo to Mo4þ (as in MoO2) are the expected reactions under these conditions, given their standard enthalpies of formation (Table 2). Under more oxidizing conditions and/or elevated temperature, further oxidation to U6þ (as in g-UO3) and Mo6þ (as in MoO3) is likely to occur. Note that MoO3 is volatile, with a vapor pressure of 4.69  104 atm at 700  C and a comparatively low melting temperature of ~1150  C. High temperature oxidation of Mo thus can result in significant mass loss through the volatilization of MoO3 [25]. Characterizing the relative fractions of U4þ/U6þ and Mo4þ/ Mo6þ provides information on the oxidizing strength of a given process step and the efficacy of a cleaning technique to remove the oxide layer thus formed. 3.1. Hot rolling The expected Mo concentration in DU10Mo is 21.6 at% (10 wt%).

Table 2 Standard enthalpies of formation of various U and Mo compounds. Compound

DfHm , kJ mol1

Reference

UO2 UO3 UN MoO2 MoO3

1085 1223.8 290.3 ~ 590 745.1

[38] [38] [39,40] [41,42] [42]

T.C. Kaspar et al. / Journal of Nuclear Materials 514 (2019) 28e39

As discussed above, the sensitivity factors [24] used to quantify the elemental compositions measured by XPS were not benchmarked against a known UeMo metallic or oxide standard, and thus the compositions reported here are used only to identify trends between datasets. After hot rolling, the DU10Mo surface is enriched in U and depleted in Mo, as indicated in Table 1, Sample A. The Mo content is only 12.3 at% (note that, although this is not an absolute quantification, the value is significantly lower than those observed on the surfaces of etched samples discussed below). A small amount of Er is present on the surface, and arises because Er2O3 is

31

used as a release agent in the crucibles used in casting. The large chemical shifts between the 3d5/2 peaks of Mo0 (228.0 eV), Mo4þ (MoO2, 230.0 eV [26]), and Mo6þ (MoO3, 232.5 eV [26]) allow these valence states to be readily distinguished as individual peaks. However, peak overlap leads to a characteristic three-peak structure in mixed-valence Mo oxides, where the central peak at ~233 eV is a convolution of the Mo 3d5/2 peak from Mo6þ and the Mo 3d3/2 peak from Mo4þ [27]. In addition, the presence of both Mo5þ (231.1 eV [26]) and Modþ (4 < d < 0, ~228.7e229.8 eV [27,28]) can complicate the spectra. As shown for Sample A in Fig. 2(a), after

Fig. 2. XPS Mo 3d (a) and U 4f (b) spectra from DU10Mo coupons after various processing steps (see Table 1). A Shirley background has been subtracted from each spectrum; the peak areas have not been normalized. The positions of the black dashed lines representing various valence states of Mo were taken from Bhosle et al. [26]. The black dashed U 4f spectrum is from a single crystal of UO2, and represents the peak shape for U4þ. The inset in (b) compares this UO2 spectrum to that of Sample G (blue solid line). (c) Peak fitting of U 4f/N 1s region of Sample A, hot-rolled DU10Mo. The thick black line is the experimental data, the Shirley background and fitted peaks are shown as thin lines, and the overall fit is shown in red. Tentative peak assignments are indicated; unlabeled peaks are assigned as U 4f satellites. (d) XPS spectra from the S 2p/Er 4d/Y 3d/Si 2s region of Samples A, F, G, and I. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

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hot rolling, Mo is present primarily as Mo6þ, with some metallic Mo0 also observed. The U 4f7/2 peak shift between U4þ (UO2, 380.0 eV [29]) and U6þ (UO3, 382.0 eV [29]) is large enough to be distinguishable, but is not resolved into separate peaks in mixed-valence compounds because of the inherent width of the peaks. Qualitatively, the presence of U6þ broadens the 4f7/2 peak on the high binding energy (BE) side of the U4þ peak when compared to the reference spectrum from a single crystal of UO2 [30], and the 4f7/2 and 4f5/2 satellite peaks associated with U4þ at ~386.8 eV and 397.5 eV, respectively, are reduced. In Fig. 2(b), Sample A, no metallic U is observed; U is fully oxidized to U4þ, with a fraction further oxidized to U6þ. The U 4f spectrum from Sample A is presented in more detail in Fig. 2(c). The broadening to the high BE side of the 4f7/2 peak is apparent, although additional broadening to the low BE side is also evident. The U4þ satellite peaks are reduced. Deconvolution of this peak envelope reveals three peaks located at 379.6 eV, 380.6 eV, and 382.1 eV. Confirming the qualitative assignments from Fig. 2(b), the peaks at 380.6 eV and 382.1 eV are assigned to U4þ and U6þ, respectively. The low-BE peak at 379.6 eV is not U0 (377.2e377.5 eV [20,29]). Instead, we assign it as arising from UNx. Justification for this assignment is obtained by observation of the small but sharp peak appearing at 395.8 eV, which we assign to N in UNx: in oxygen-free UNx, the U 4f7/2 peak appears at 379.0e379.4 eV in conjunction with the N 1s peak at 396.2e396.4 eV [31,32]. In addition to this signature of N in oxygen-free UNx, the broad peak at 399.6 eV indicates that the majority of N is incorporated into uranium oxy-nitride [31]. EDS maps and a line scan associated with cross-sectional SEM images of Sample A are presented in Fig. 3, and reveal the complex structure of the surface layers present on DU10Mo after hot rolling. In the near-surface region, an oxide (or oxy-nitride) layer is observed in the O Ka1 map; the oxygen signal appears strongest at the surface, and extends approximately 5 mm deep. This surface oxidation is likely a result of air exposure and is consistent with the XPS results from Sample A in Fig. 2(b and c). Correlation of the N Ka1,2 and O Ka1 maps reveals a thick layer of UNx extending approximately 10 mm below this surface layer. The average composition of Mo in this layer (calculated as Mo/(U þ Mo) from region #1 in Fig. 3(a)) is slightly Mo-deficient (19.5 at%), in qualitative agreement with the reduced Mo concentration measured by XPS at the sample surface (see Table 1, Sample A). A thin (1e2 mm), highly oxidized layer is present at the interface between the UNx surface layer and pristine DU10Mo. This layer exhibits strong Mo depletion, with an average composition of 11.2 at% as calculated from the line scan data in Fig. 3(b)), although in some regions (not shown), small clusters (~1 mm) enriched in Mo are observed. In the pristine DU10Mo region, low N Ka1,2 counts indicate that the N composition is zero within the statistical error of the EDS measurement. The average Mo composition in the bulk of the pristine DU10Mo is 20.6 at% (Fig. 3(a), region #2), which matches well (within the error of the EDS quantification) to the nominal composition of 21.6 at% Mo. It is possible that this nitride layer formed during the hotrolling process. Specifically, it might have formed during the 700  C heat treatment in air; the ingot was wrapped in Zr foil to getter oxygen, and the ingot may have reacted with nitrogen in the resulting low-oxygen environment. Mo loss due to volatilization of MoO3 may have also occurred during this heat treatment; the reaction of Mo with O would have further reduced the partial pressure of oxygen between the foil and the ingot. However, it is also possible that the Mo-depleted nitride layer was formed during an earlier processing step. More characterization of the as-received ingot material would be necessary to confirm the origin of the nitride layer.

Fig. 3. (a) Cross-sectional EDS maps of Sample A, hot-rolled DU10Mo. Map color scale intensities are on arbitrary scales; brighter colors indicate a higher concentration. Positions of EDS line scan data and compositional measurement regions are indicated by a line and rectangles, respectively, in the Mo La1 map. (b) EDS line scan data as a function of distance along the scan direction. Concentration was calculated using only U, Mo, O, and N data. (c) Cross-sectional SEM image collected with a low-angle backscattered electron detector of Sample D, cold-rolled DU10Mo. Three measurements of the thickness of the apparent surface layer in nanometers are shown. The black arrow is an artifact of the measurement routine. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

3.2. Acid etching hot-rolled DU10Mo To remove the uranium nitride and oxide layers from hot-rolled DU10Mo, the foil was immersed in HNO3 for 5 min. After a quick, 2e5 s rinse in DI water, XPS measurements (Fig. 4, Sample B spectrum) reveal that the N 1s signal has disappeared, all the U6þ is reduced to U4þ, and a significant fraction of the Mo6þ is reduced to Mo4þ. In addition, the concentration of Mo at the surface increases to 25.6 at% (Table 1, Sample B), which is near the nominal value for pristine DU10Mo. A longer DI water rinse time (30 min, Sample C) has little effect on the oxidation of U compared to the short rinse; only a very slight fraction of U is oxidized to U6þ. In contrast, a significant fraction of Mo4þ is oxidized back to Mo6þ after the 30 min DI water rinse. Exposing the coupon that had been rinsed for 2e5 s to air produces a similar effect: the oxidation state of U remains U4þ, while some Mo4þ oxidizes to Mo6þ, as shown in Fig. 4. This oxidation is detected after 24 h of air exposure, and becomes more pronounced after 96 h. Estimating the ratio of the intensities of the Mo 3d5/2 peaks at ~230 eV (Mo4þ) and ~233 eV (Mo6þ þ Mo4þ) as a guide, the amount of Mo6þ after air exposure

T.C. Kaspar et al. / Journal of Nuclear Materials 514 (2019) 28e39

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Fig. 4. XPS Mo 3d (a) and U 4f (b) spectra of Sample B immediately after HNO3 etching and 2 s rinse in DI water, and after exposure to laboratory air for 24 h and 96 h. Spectra from Sample C after HNO3 etching and 30 min rinse in DI water are also presented.

for 96 h is still somewhat less than the amount observed after a 30 min rinse in DI water. 3.3. Cold rolling and etching Cold rolling the hot-rolled and cleaned DU10Mo coupons to 0.635 mm (0.02500 ) does not produce an easily identifiable surface layer in SEM images, as shown in Fig. 3(c). However, the DU10Mo surface is strongly depleted in Mo relative to U, with a Mo composition of 3.9% (Table 1, Sample D). As seen in the Sample D spectra in Fig. 2, a large fraction of Mo remaining on the surface is present as Mo6þ. The valence state of U is primarily U4þ, although a small fraction has oxidized to U6þ. The XPS results indicate that a surface layer has formed that is too thin to be clearly resolved in cross-sectional SEM. Close inspection of the SEM images in Fig. 3(c) reveals a region at the surface that exhibits slightly different Zcontrast than the bulk. The thickness of this region is roughly estimated to be 150e300 nm, and thus this puts an upper bound of approximately 300 nm on the surface layer thickness. The mechanism responsible for the observed significant Mo depletion after cold-rolling is not apparent. Both Ca and Zr are present on the surface (as observed by XPS) after cold rolling. Zr is likely present as residue from the Zr foil used during the heating and annealing steps. Ca is an impurity in DU10Mo. After etching in HNO3 for 5 min, all the U6þ is reduced to U4þ, and the fraction of Mo6þ decreases relative to Mo4þ, as seen in the Sample E spectrum in Fig. 2. In addition, the fraction of Mo present on the surface increases to be somewhat greater than the Mo concentration on similarly etched surfaces of Samples B and C. Zr is removed, but Ca remains on the surface. 3.4. Cold rolling and electropolishing In contrast to the first cold-rolling step to reduce the thickness to 0.635 mm (0.02500 ), further cold rolling DU10Mo to 0.203 mm (0.00800 ) does not lead to surface depletion of Mo, as indicated in Sample F of Table 1. This may indicate that the Mo depletion

observed in Sample D is due to sample-to-sample variation, not an intrinsic phenomenon related to cold-rolling. As seen in the Sample F spectra in Fig. 2, the charge states of Mo (primarily Mo6þ) and U (primarily U4þ, fraction of U6þ) are very similar to those after cold rolling to 0.02500 (Sample D, Fig. 2). Zr is not detected on the surface, but Ca is. After electropolishing in H2SO4 at 1e2 A until the surface layer appeared by eye to be removed (Sample G), all U6þ is reduced to U4þ, and a very small shoulder is present on the low BE side of the U 4f7/2 peak (approximately 377 eV; see inset to Fig. 2(b)) that might indicate the detection of U0. The presence of U0 is likely an indication that the surface oxide layer is very thin (on the order of 5e10 nm thick), and metallic U is detected below this layer. Some Mo is reduced from Mo6þ to Mo4þ, although the fraction of Mo6þ remains somewhat higher than for the foil etched in HNO3 (Sample E); some Mo5þ may also be present. The Ca contamination is removed by electropolishing. A trace of S may be present before electropolishing, as shown in Fig. 2(d), Sample F, but the overlap of the S 2p and Er 4d core level peak positions at ~168 eV makes it difficult to identify the origin of the weak peak. However, after electropolishing, the peak increases significantly as shown in Fig. 2(d), Sample G, and a second peak appears at lower BE, both of which arise from S. S is present primarily as sulfate, SO2 4 (high BE peak at ~169 eV), with a smaller signal from sulfide, S2 (low BE peak at ~162 eV) [33]. For comparison, the Er 4d signal from hotrolled DU10Mo (Sample A) is also shown in Fig. 2(d). The asymmetry in both of the peaks assigned as S 2p from Sample G indicates that the spin-orbit splitting (SOS) of the S 2p3/2 and S 2p1/2 peaks is weakly resolved (SOS ¼ 1.18 eV [34]). In contrast, the broad peak assigned to Er 4d from Sample A is symmetric, consistent with the negligible SOS of the Er 4d5/2 and Er 4d3/2 peaks [34]. The concentration of Mo on the surface of Sample G is 36.7%, which is significantly higher than that measured for any other rolled or cleaned sample. This may be due in part to an artificial increase in the Mo 3d peak area arising from overlap between the S 2s and Mo 3d peaks at ~228 eV (as indicated in Fig. 2(a)). It is also feasible that U was preferentially etched into solution during the electropolishing process, leaving behind a Mo-enriched surface.

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T.C. Kaspar et al. / Journal of Nuclear Materials 514 (2019) 28e39

Fig. 5. Digital photographs of the front and back of cold-rolled and etched DU10Mo ingot (a) and foil (b, Sample I) before and after exposure to 97% humidity at 30  C for six weeks. Coupon dimensions are approximately 1 cm  1 cm. XPS analysis areas are indicated in yellow in (b). (c) XPS Mo 3d (left) and U 4f (right) spectra from Sample I. Dashed U 4f spectrum is from a single crystal of UO2. (d) XRD pattern of Sample I. Asterisks indicate reflections from g-U. (e) Cross-sectional SEM image of Sample I. Surface oxide layer is visible as a dark band at the top of the DU10Mo. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

3.5. Corrosion behavior of etched and electropolished DU10Mo Both ingot material and cold-rolled foils of DU10Mo were cleaned by either rinsing in HNO3 or electropolishing for 5 min. Samples were then sealed at atmospheric pressure in Parr pressure

vessels at 97% humidity and heated to 30  C for six weeks. Due to the small size of the samples, no significant weight gain was measured for any sample after the corrosion treatment. Photographs from selected samples before and after the corrosion treatment are presented in Figs. 5(a,b) and 6(a,b) and show that all specimens experienced some color change, indicating surface oxidation. Etching the ingot material revealed a surface topography of dense, small features, as shown in the left image of Fig. 5(a). These surface features became more pronounced after the corrosion treatment (right image of Fig. 5(a)), turning an orange/brown color, while the surface area between the features became darker. In contrast, the ingot appeared relatively smooth and uniform after electropolishing (Fig. 6(a), left image). However, after the corrosion treatment (Fig. 6(a), right image), surface features similar in size to those from the etched ingot were revealed. The color after corrosion is primarily brown, with only a few small areas of darker color. The color differences between etched and electropolished surfaces are seen more distinctly for the cold-rolled foils. The coldrolled specimen in Fig. 6(b) appeared clean and polished after electropolishing. After corrosion, it developed a thin layer of monochromatic brown color on the surface, with spots of the original surface still visible. In comparison, the cold-rolled foil retained some brown color variations after etching (Fig. 5(b)), and after corrosion the coloration was more varied and intense than for the electropolished foil, with regions of pale blue, dark blue, and red/orange. Color on the surface of metals is typically caused by the interference of light refracting through the surface oxide layer(s) and reflecting outward. While a qualitative measure, this indicates that the HNO3-etched samples had developed an oxide layer of different thickness and/or optical properties (e.g., transparency to visible light) compared to the electropolished samples. The etched and electropolished specimens were examined with XPS after corrosion. The locations where spectra were collected are indicated in Figs. 5(b) and 6(b). The etched foil (Sample I) had a pattern of red and blue oxide film, with a small region of “staining” from the etching process (this region appears pale blue in the photographs in Fig. 6(b)). One XPS spectrum was taken from each of the distinctive regions: “Spot 0” and “Spot 2” were the red and blue regions, respectively, while “Spot 1” was from the pale blue area of “staining.” As shown in the spectra presented in Fig. 5(c), a significant fraction of the uranium has been oxidized to the U6þ state in all three regions. Likewise, the Mo 3d spectra reveal that nearly all the Mo in both Spot 0 and Spot 2 has been oxidized to the Mo6þ state, although Spot 1 (the “stained” region) has remained a mixture of Mo4þ and Mo6þ. The Mo composition of Spot 2 (blue region) is approximately the nominal value, while Spot 0 (red region) is somewhat enriched and Spot 1 (pale blue) is highly enriched in Mo, as presented in Table 1, Sample I. The XRD pattern from the etched foil after corrosion (Sample I) is presented in Fig. 5(d). The x-ray penetration depth in U is ~10 mm, making the XRD data fairly surface-sensitive. In addition to strong diffraction peaks arising from g-U, several broad peaks could be indexed to UO2 with a small crystallite size. Rietveld refinement of the pattern resulted in a composition of approximately 86 wt% g-U and 14 wt% UO2. Although the XPS data revealed significant fractions of U6þ and Mo6þ, no higher oxides were detected in the diffraction pattern. This may indicate that the highly-oxidized species only occur at the sample surface, and do not constitute sufficient volume fraction to be detectable by XRD. Note that no Mo oxides or mixed UeMoeO oxides [35] were detected. In contrast to the etched sample, the electropolished foil (Fig. 6(b)) was a uniform brown color after corrosion (Sample H), albeit with some intensity differences across the surface. Two of the corners were shinier in appearance than the rest of the foil. Two spectra (“Spot 0” and “Spot 1“) were collected from separate, small

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Fig. 7. AFM surface topography micrographs (90  90 mm scan region; 3D representations are cropped to 40  40 mm) of DU10Mo coupons cold rolled to 0.203 mm and (a) etched in HNO3 (Sample I), with an rms roughness of 0.598 mm; (b) electropolished in H2SO4 (Sample H), with an rms roughness of 0.185 mm. Line profiles are presented in (c) that correspond to the white lines in the associated topography micrographs.

Fig. 6. Digital photographs of the front and back of cold-rolled and electropolished DU10Mo ingot (a) and foil (b, Sample H) before and after exposure to 97% humidity at 30  C for six weeks. Coupon dimensions are approximately 1 cm  1 cm. XPS analysis areas are indicated in yellow in (b); analysis areas 0 and 1 were collected from two different spots within the dashed circle. (c) XPS Mo 3d (left) and U 4f (right) spectra from Sample H. Dashed U 4f spectrum is from a single crystal of UO2. (d) XRD pattern of Sample H. Asterisks indicate reflections from g-U. (e) Cross-sectional SEM image of Sample H. No oxide layer is resolved at the surface of the DU10Mo. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

areas of the central, uniform region. “Spot 2” and “Spot 3” were collected from the shiny corners. As shown in Fig. 6(c), in all four regions the uranium was primarily in the U4þ state, with a small fraction of U6þ. Molybdenum was primarily in the Mo6þ state, with a small fraction of Mo4þ. Spot 0 is slightly more oxidized (more U6þ, more Mo6þ) than the other regions. The Mo enrichment in Spots 1,

2, and 3 is nearly identical (Table 1, Sample I), and very similar to that for Sample G. Again, the Mo content of these areas may be somewhat overestimated due to the overlap between the Mo 3d and S 2s regions (see Fig. 2(d), Sample H). Notably, Spot 0 is highly enriched in Mo, although its visual appearance was not strikingly different from that of Spot 1. The XRD pattern of the electropolished foil after corrosion (Sample H) revealed a surface oxide layer of UO2 with a small crystallite size, as shown in Fig. 6(d). In contrast to the etched foil, however, the electropolished sample also exhibited peaks corresponding to UC and UC2. Rietveld refinement of the pattern resulted in a composition of approximately 79 wt% g-U, 13 wt% UO2, 6 wt % UC, and 1 wt% UC2. Uranium carbides are well-known inclusions in U10Mo [18], and the carbides observed by XRD are assumed to have been present before the electropolishing and subsequent corrosion processes. The presence of carbides near the surface of the electropolished sample, but not the etched sample, indicates that etching in HNO3 was more successful at dissolving the carbide inclusion than electropolishing in H2SO4. The dramatic color difference between the etched and the electropolished foils after corrosion may be due, in part, to differences in thickness of the oxide layer that formed on the two

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samples. To quantify the oxide layer thickness, cross-sections of the near-surface region of each corroded coupon were removed by FIB milling and imaged with SEM. The cross-sectional image for the etched and corroded foil (Sample I) is shown in Fig. 5(e). Between the pristine DU10Mo and the rough surface, the surface oxide layer is visible as a dark band. The approximate thickness of this band is 100e150 nm. A cross-section of the electropolished and corroded foil (Sample H) is presented in Fig. 6(e). Unlike the etched foil, there is no dark band visible at the top of the pristine DU10Mo. Thus, the oxide layer on the electropolished and corroded surface is too thin to resolve in these cross-sectional SEM images. It is considerably thinner than that on the etched and corroded surface (i.e., its thickness is considerably less than 100 nm). This is consistent with the light brown color of the corroded surface of the electropolished sample, with the bare DU10Mo surface visible underneath. To better understand the differences in corrosion behavior between etched and electropolished DU10Mo, the surface topography of selected specimens was characterized with AFM. Fig. 7 presents the surface topography of cold-rolled DU-10Mo foils after etching in HNO3 (Fig. 7(a), analogous to Sample E) or electropolishing in H2SO4 (Fig. 7(b), analogous to Sample G). In both cases, the cleaned foils were exposed to air for a few hours before and during the AFM measurements, and thus the topography is a convolution of the bare metal surface after treatment and the beginnings of surface oxide formation. Nonetheless, clear differences in topography and roughness are observed when comparing the etched and electropolished foils. The foil etched in HNO3 exhibits a rough surface with numerous protrusions that are on the order of half a micron tall and a few microns wide. Large pits are also observed; the pit imaged in Fig. 7(a) is approximately 2e4 mm deep, with a flat bottom. The root mean square (rms) roughness of the surface (excluding the deep

pits) is 0.598 mm. In contrast, the surface topography of the electropolished sample (Fig. 7(b)) is much smoother, and lacks large protrusions. Pits are also observed on this surface, but their depth is shallower; the pit in Fig. 7(b) is < 2 mm deep. The rms roughness of this surface (excluding the deep pits) is 0.185 mm. A comparison of the surface roughness differences between the etched and the electropolished foils is given by an overlay of line profiles from the surface topography images of each sample in Fig. 7(c). The smoother surface morphology and shallower pits characteristic of the electropolished surface produce a smoother line profile than that resulting from the protrusions and deep pits present on the etched surface. SEM and AFM images of the surface topography of the etched and electropolished foils exposed to humidity for six weeks are presented in Figs. 8 and 9, respectively. At lower magnification, SEM images of the etched and corroded foil (Sample I) show a surface that is relatively flat and uniform, as seen in Fig. 8(a). As the magnification is increased, the surface is revealed to exhibit a “cobblestone” morphology that is characteristic of oxide formation. This morphology is visible in both SEM (Fig. 8(b)) and AFM (Fig. 8(c)) images. The surface has been nearly entirely covered by an oxide layer, and the original morphology of the metal after etching cannot be clearly discerned. The SEM image in Fig. 8(b) shows pits with small particles located inside. These pits (but not the particles inside) are also clearly resolved in AFM images of similar magnification (Fig. 8(c)). Lines are visible in the AFM deflection image in Fig. 8(c) (red arrows) which may correspond to grain boundaries in the underlying metal. Alternatively, the lines may represent cracks formed during oxidation. As shown in Fig. 8(d), the height and diameter of the oxide grains on the surface is on the order of 0.1e0.3 mm. In several regions, flat-bottomed pits are observed that are similar to those seen on the freshly etched

Fig. 8. (a, b) Plan-view SEM images collected with a backscattered electron detector of DU10Mo coupon cold rolled to 0.203 mm, etched, and exposed to humidity for six weeks (Sample I). AFM surface topography (top, color) and deflection ((c) only, bottom, grayscale) micrographs of Sample I. Red arrows in (c) indicate underlying grain boundaries or cracks in the oxide layer. Line profiles are presented that correspond to the white lines in the associated topography micrographs. (c) 40  40 mm scan region, (d) 2  2 mm scan region, (e) 16  16 mm scan region. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

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Fig. 9. (a, b) Plan-view SEM images of DU10Mo coupon cold rolled to 0.203 mm, electropolished, and exposed to humidity for six weeks (Sample I). (a) Image collected with a secondary electron detector. (b) Image collected with a backscattered electron detector. (cee) AFM surface topography (top, color) and deflection ((c) only, bottom, grayscale) micrographs of Sample I. Red arrows in (c) indicate underlying grain boundaries or cracks in the oxide layer. Line profiles are presented that correspond to the white lines in the associated topography micrographs. (c) 40  40 mm scan region, (d) 2  2 mm scan region, (e) 16  16 mm scan region. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

surface (Fig. 7(a)). One such pit is imaged in Fig. 8(e) with a wall height of 3.2 mm. The surface morphology of the electropolished and corroded foil is very different from that of the etched and corroded sample. Lower-magnification SEM images of the electropolished and corroded foil (Sample H), such as those shown in Fig. 9(a) and (b), reveal a rough surface morphology, as well as significant organic contamination (not shown). Between the protrusions, the surface looks relatively smooth (seen more clearly in Fig. 9(b)). AFM images collected over a smaller area and shown in Fig. 9(c) confirm that a majority of the surface does not show the characteristic “cobblestone” morphology indicative of a thick oxide layer. Instead, the majority of the surface at this scale is relatively smooth, with grain boundaries clearly visible (red arrows in Fig. 9(c)). In addition to the smooth morphology, two characteristic features of the surface can be identified in the AFM images in Fig. 9: (1) a thin surface covering with pitted edges, and (2) a rough morphology with small grains (~40e80 nm diameter, Fig. 9(e)). Feature (2) likely corresponds to patches of thicker oxide formation. The origin of Feature (1) is not understood, but may be surface contamination arising from the electropolishing process. Although it may serve as a protective layer (no rough features analogous to Feature (2) are observed within Feature (1) regions in Fig. 9), its presence alone cannot account for the dramatically improved corrosion resistance of electropolished DU10Mo compared to etched specimens, because it appears to cover only a small fraction of the surface. The differences in surface topography of the etched sample compared to the electropolished sample after corrosion are highlighted in Fig. 10. After exposure to humidity for six weeks, the electropolished sample retains the smooth surface exhibited soon after cleaning, while the etched surface remains rough.

4. Discussion No matter the mechanism of oxide formation (mechanical processing, air exposure, humidity), the Mo component of the DU10Mo alloy exhibited a higher average oxidation state than the U component. This “protective” role of Mo is consistent with the oxidation results of Kang et al. [36], who observed a slower oxidation rate for DU10Mo than for U at elevated temperatures (200e400  C) in air. A similar “protective” role of the alloying element on the oxidation of U was observed in U0.8Zr0.2 alloys [20]. In this case, it was hypothesized that the observed Zr segregation to the surface during heat treatment to 500e600 K formed a UZr2þx phase that both resisted subsequent oxidation and partially protected the underlying U from oxidation. A similar explanation is not likely in the case of DU10Mo because substantial Mo enrichment was not observed by XPS. The protective role of Mo is all the more surprising because, as shown in Table 2, the standard enthalpies of formation of MoO2 and MoO3 are significantly less than the equivalent enthalpies of formation of UO2 and g-UO3, which indicates that the thermodynamic driving force for U oxidation is higher than that for Mo oxidation. In addition, the oxidation of Mo in water-vapor-containing atmospheres at elevated temperatures (500e1200  C) was found to be significantly retarded compared to oxidation in dry air [25]. Further work is required to determine the role of Mo in the oxidation of DU10Mo. The effectiveness of acid etching in HNO3, compared to electropolishing in H2SO4, can be determined by comparing Samples B (HNO3 etched), E (HNO3 etched), and G (electropolished). In all three cases, the treatment reduced a significant fraction of Mo6þ to Mo4þ, and reduced the small fraction of U6þ that was present to U4þ. These reductions in valence state indicate the removal of the

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Fig. 10. AFM surface topography micrographs of cold-rolled DU10Mo cold-rolled to 0.203 mm and (a) etched in HNO3, followed by exposure to humidity for six weeks (Sample I), with an rms roughness of 0.437 mm, or (b) electropolished, followed by exposure to humidity for six weeks (Sample H), with an rms roughness of 0.332 mm. (c) Overlay of a portion of the line profiles from the etched sample (from white line in Fig. 8(c)) and the electropolished sample (from white line in Fig. 9(c)) after humidity exposure. (a,b) 90  90 mm scan region.

oxide layers that formed during processing. Cross-sectional SEM images after a later processing step (Fig. 3(c)) confirm that the thick UNx layer that was present after hot rolling has also been removed. Because the samples were then exposed to air after treatment, a new, thin oxide layer consisting of U4þ and Mo4þ formed on (and covered) any bare metal surfaces. The differences observed by XPS between the two treatments were relatively minor: electropolishing resulted in S contamination, somewhat more Mo6þ than after etching, and some Mo surface enrichment. Subsequent reoxidation occurred as the etched U10Mo surfaces were soaked in water for long periods. This is evident when comparing etched Samples B (rinsed 2 s) and C (rinsed 30 min). After 30 min of rinsing, Sample C has more Mo6þ than does Sample B. Likewise, minimizing air exposure after treatment can reduce surface oxide formation. Sample B was exposed to air for 96 h, and afterwards exhibited nearly the same fraction of Mo6þ as did Sample C (rinsed 30 min). This illustrates the faster kinetics of Mo oxidation in water

than in air. Holding the etched and electropolished DU10Mo at 30  C in 97% humidity for six weeks simulates an aggressive ambient corrosion environment that will result in accelerated surface oxidation. The photographs presented in Figs. 5(a,b) and 6(a,b) confirm that surface corrosion occurs under these conditions. A qualitative assessment of the visibility of the bare metal surface below the brown oxide layer on the electropolished sample, compared to the blue and red oxide layer on the etched sample that fully obscures the underlying metal surface, indicate that the oxide layer on the electropolished surface is thinner than that on the etched surface. The cross-sectional SEM images in Figs. 5(e) and 6(e) confirm this assessment: the oxide layer on the etched and corroded foil is 100e200 nm thick, while the oxide layer on the electropolished and corroded foil is too thin to be visible in these images. The thin oxide on the electropolished sample is expected to be easier to remove with subsequent cleaning steps, resulting in less loss of material, than the thicker oxide on the etched sample. From the XPS data in Figs. 5(c) and 6(c), it is apparent that the deep blue and red colors of the etched sample correspond to the presence of significant fractions of U6þ and Mo6þ, while the light brown color of the electropolished sample corresponds to less U6þ and Mo6þ (i.e., higher fractions of U4þ and Mo4þ). Note that the XRD patterns of both foils (Figs. 5(d) and 6(d)) indicated the oxide layer was composed of UO2; no higher oxides such as U3O8 or UO3 were detected. Because the probe depth of XPS is only 5e10 nm, compared to ~10 mm for XRD, the higher-valence U and Mo species may be limited to the surface region of the oxide layer. As discussed above, the U and Mo valence charges were similar for the electropolished and etched samples before the corrosion treatment, so this cannot contribute to the differences observed after corrosion. Two factors may play a role in the difference in oxidation behavior: first, the electropolished sample exhibited a higher concentration of Mo on the surface, which may act as a protective layer. Second, surface topography measurements by AFM (Fig. 7) reveal significantly less surface roughness after electropolishing compared to etching in HNO3, which provides less surface area over which oxidation will occur. A similar correlation between low surface roughness and reduced surface corrosion was recently found for 304 stainless steel exposed to a highly oxidizing CO2 environment [37]. 5. Conclusions The surface layers that form on DU10Mo during processing and storage, and their removal via acid etching or electropolishing, were characterized with XPS, cross-sectional SEM/EDS, and AFM. The thick, Mo-poor UNx layer present on the foil surface after hot rolling was effectively removed by etching in HNO3 and quickly rinsing in DI water, leaving only a thin surface oxide composed primarily of U4þ and Mo4þ. Extending the DI water rinse time, or increasing the air exposure time, led to further oxidation that increased the fractions of U6þ and Mo6þ. After cold rolling, most of the Mo was present as Mo6þ (Mo depletion was also observed, but could not be definitively attributed to result from the cold-rolling process), while the U component remained U4þ. Both etching and electropolishing the cold-rolled DU10Mo reduced the surface oxide layer, although somewhat more Mo6þ remained after electropolishing than after etching. There was also some evidence of Mo enrichment on the surface of the electropolished sample, which may indicate preferential dissolution of U during the electropolishing process. The electropolished foil was found to be more resistant to oxidation in a humid environment than was the etched foil. This difference is not likely due to the relatively small differences in surface composition or oxidation state of the foils after

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etching or electropolishing, but instead is attributed in part to the smoother surface obtained after electropolishing, which reduced the surface area available for oxidation compared to the etched foil. In all cases, Mo appeared to oxidize more readily than U, although the thermodynamic driving force for oxidation favors U oxidation over Mo oxidation. This effect is particularly observed for the electropolished DU10Mo foil exposed to high humidity, where the Mo enrichment on the surface appeared to act as a protective layer that prevented significant oxidation of U4þ to U6þ. These insights can inform the processing and storage decisions for LEU-Mo alloys as they are fabricated in preparation for subsequent Zr or Al cladding. Acknowledgments The authors thank Dr. Du for providing the single crystal UO2 XPS spectra. This work was funded by the U.S. Department of Energy National Nuclear Security Administration's Office of Material Management and Minimization and performed at Pacific Northwest National Laboratory (PNNL) under contract DE-AC0576RL01830. A portion of this work was performed using EMSL, a national science user facility sponsored by the Department of Energy’s Office of Biological and Environmental Research and located at PNNL. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jnucmat.2018.11.022. References [1] D.D. Keiser, S.L. Hayes, M.K. Meyer, C.R. Clark, High-density, low-enriched uranium fuel for nuclear research reactors, JOM-J. Miner. Met. Mater. Soc. 55 (9) (2003) 55e58. [2] A.J. Clarke, K.D. Clarke, R.J. McCabe, C.T. Necker, P.A. Papin, R.D. Field, A.M. Kelly, T.J. Tucker, R.T. Forsyth, P.O. Dickerson, J.C. Foley, H. Swenson, R.M. Aikin, D.E. Dombrowski, Microstructural evolution of a uranium-10 wt.% molybdenum alloy for nuclear reactor fuels, J. Nucl. Mater. 465 (2015) 784e792. [3] S. Van den Berghe, P. Lemoine, Review of 15 years of high-density lowenriched UMo dispersion fuel development for research reactors in Europe, Nucl. Eng. Technol. 46 (2) (2014) 125e146. [4] V.P. Sinha, P.V. Hegde, G.J. Prasad, G.K. Dey, H.S. Kamath, Phase transformation of metastable cubic gamma-phase in U-Mo alloys, J. Alloy. Comp. 506 (1) (2010) 253e262. [5] B.W. Howlett, A.J. Eycott, I.K. Kang, D.R.F. West, The kinetics of the isothermal decomposition of a gamma-phase uranium-6 atomic percent molybdenum alloy, J. Nucl. Mater. 9 (2) (1963) 143e154. [6] G. Champion, R. Belin, H. Palancher, X. Iltis, H. Rouquette, M. Pasturel, V. Demange, P. Castany, V. Dorcet, O. Tougait, Development of characterisation methods on U(Mo) powders for material testing reactors (MTRs), Powder Technol. 255 (2014) 29e35. [7] Y.S. Kim, G.Y. Jeong, D.S. Sohn, L.M. Jamison, Pore growth in U-Mo/Al dispersion fuel, J. Nucl. Mater. 478 (2016) 275e286. [8] A. Leenaers, W. Van Renterghem, S. Van den Berghe, High burn-up structure of U(Mo) dispersion fuel, J. Nucl. Mater. 476 (2016) 218e230. [9] D.E. Burkes, A.J. Casella, A.M. Casella, Measurement of fission gas release from irradiated U-Mo dispersion fuel samples, J. Nucl. Mater. 478 (2016) 365e374. [10] D.E. Burkes, A.J. Casella, A.M. Casella, The influence of cladding on fission gas release from irradiated U-Mo monolithic fuel, J. Nucl. Mater. 486 (2017) 222e233. [11] Y. Park, N. Eriksson, D.D. Keiser, J.F. Jue, B. Rabin, G. Moore, Y.H. Sohn, Microstructural anomalies in hot-isostatic pressed U-10 wt.% Mo fuel plates with Zr diffusion barrier, Mater. Char. 103 (2015) 50e57. [12] Y. Park, N. Eriksson, R. Newell, D.D. Keiser, Y.H. Sohn, Phase decomposition of gamma-U (bcc) in U-10 wt% Mo fuel alloy during hot isostatic pressing of monolithic fuel plate, J. Nucl. Mater. 480 (2016) 271e280. [13] J.F. Jue, D.D. Keiser, C.R. Breckenridge, G.A. Moore, M.K. Meyer, Microstructural characteristics of HIP-bonded monolithic nuclear fuels with a diffusion barrier, J. Nucl. Mater. 448 (1e3) (2014) 250e258.

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