Iron-chromium-aluminum (FeCrAl) alloys such as APMT (advanced powder metallurgy tubing) are candidate materials to replace zirconium alloys for the light water reactor fuel cladding. This alloy meets the requirements to be a material more tolerant of high-temperature accidents than the current zirconium alloys. One concern is that the use of FeCrAl may result in an increase in tritium presence in the coolant compared to the current design. The aim of the current research was to obtain effective diffusion coefficients (Deff) for hydrogen through APMT using the Devanathan-Stachurski cell. Results showed that at 30°C the Deff value was 2.8 × 10−8 cm2/s. Results also showed a linear relationship between the permeated hydrogen flux and the inverse of the test specimen thickness, which demonstrated the validity of the permeation measurements.

The concept of accident tolerant fuels (ATF) was born in the aftermath of the Fukushima Daiichi nuclear power stations accident in March 2011 due to a 15 m tall tsunami wave generated by the 9.0 magnitude Tohoku earthquake.1  Four reactors at the site suffered severe damage after the lack of active cooling increased the temperatures above 1,000°C, which led to zirconium-based components in the reactor core and in the used fuel cooling pools to react exothermically with water and steam to produce ignitable hydrogen gas (Equation [1]).
formula

Even though there was no loss of human life as a consequence of this event, the explosion of the accumulated hydrogen in the external buildings caused concern about the safe use of nuclear energy to generate electricity. After the accident, the global nuclear materials community acknowledged the need to develop and implement more accident tolerant materials. One of the most developed ATF concepts is to use monolithic iron-chromium-aluminum (FeCrAl) alloys to replace zirconium alloys for the cladding of the fuel.2-3  Ferritic FeCrAl alloys are suitable cladding materials because they have a low coefficient of thermal expansion, high thermal conductivity, good mechanical properties at normal operation temperatures, and outstanding resistance to attack by steam at a loss of coolant accident conditions (T > 1,000°C). The two less desirable attributes of FeCrAl alloys are their higher than zirconium capture of thermal neutrons and higher than zirconium permeation of tritium.1-4 

Mitigation Measures for Increased Tritium Release into the Coolant

The proposed use of FeCrAl for cladding of light water reactor fuels is being currently investigated in the international materials community. However, there are two attributes of FeCrAl that may not make them ideal for cladding application: (1) a higher neutron absorption cross section than the current zirconium alloys; and (2) the possibility of a higher than current tritium release into the coolant.2,5-6  Regarding the parasitic absorption of neutrons, the FeCrAl cladding wall was made 0.3 mm thick or approximately half the thickness of the current Zr alloy to make it cost neutral. The wall thickness reduction is possible due to the superior mechanical properties of FeCrAl compared to zirconium alloys at temperatures in the order of 400°C.7  Also, it is likely that the tritium release into the coolant will decrease as a function of time because of the development of oxide films on the surface of the cladding, i.e., alumina from the fuel cavity side and chromia from the coolant side. Both of these oxides would make tritium less likely to migrate from the fuel cavity to the coolant (Figure 1).2,6 

The Electric Power Research Institute (EPRI) reported that when austenitic stainless steel claddings were used for power generation the amount of tritium in the coolant water was approximately 10 times higher than when zirconium claddings were used.8  As FeCrAl are ferritic (bcc) in nature, the diffusion of tritium through the cladding wall into the coolant could be even higher than when the austenitic (fcc) iron-based material was used.6  It is known that hydrogen intake into iron alloys decreases in the presence of oxide films on the surface of the iron-based alloys.9-11 

Recently, tritium permeation studies were conducted on FeCrAl oxide dispersion strengthened (ODS) alloys.12-13  Sakamoto, et al., studied the permeation of hydrogen as a function of the temperature on noncorroded FeCrAl specimen and on an FeCrAl ODS specimen immersed for 30 d in water with 8 ppm of dissolved oxygen at 290°C. They reported a modest decrease, by a factor of less than 10, in the tritium permeation through the preoxidized specimen as compared to the nonoxidized FeCrAl ODS specimen.13 

One of the common failure mechanisms of the current cladding of zirconium alloys is the embrittlement by hydrides. Atomic hydrogen can form on the surface of the zirconium cladding and diffuse into the cladding wall and react with the metal to form stable zirconium hydrides. The hydrides tend to develop near the outer diameter (OD) of the zirconium alloy fuel rod due to the temperature gradient across the cladding wall. Hydrogen is less soluble in the zirconium alloy matrix near the cooler OD and may react with zirconium to form hydrides. By using FeCrAl alloys for the cladding, the hydride in the cladding issue is eliminated because none of the elements in FeCrAl (Fe, Cr, Al, or Mo) react with hydrogen (or tritium) to precipitate stable metal hydrides in the manner that Zr does.14 

The literature contains limited data related to hydrogen or tritium permeation through ferritic alloys of FeCr and of some FeCrAl, but not advanced powder metallurgy tubing (APMT).6  The purpose of the current research was to use the Devanathan-Stachurski cell (captured in standards ASTM G148 and ISO 17081)15-17  to measure the permeation of hydrogen at 30°C through clean (freshly polished) FeCrAl APMT alloy. The use of this technique, where atomic hydrogen is generated on the charging side by applying a cathodic potential, has never been reported in the literature for FeCrAl alloy.

Materials

The tested alloy was FeCrAl APMT, which has a nominal weight percent composition of 21Cr + 5Al + 3Mo + Fe (balance). The APMT base material was made using yttrium-free powder, which was pilgered to a 0.875-in (22.2 mm) round bar at Sandvik in Sweden. The bar was then reduced in height (forged) by 30% at 875°C and then recrystallized at 850°C at GE Research. Specimens for hydrogen permeation were made by slicing this upset bar sample at three different thicknesses (0.44 mm, 1.10 mm, and 1.99 mm). The fully recrystallized microstructures of the APMT specimens had an average grain size that was below 10 µm (Figure 2). This microstructure will not be the same as in the actual cladding tubes, but the grain size is in the same range. Figure 2 showed the typical APMT microstructure of small grains and second-phase precipitates (as white nano-sized dots) which provide strength to the material.7  Energy dispersive spectra (not included here) indicate that the white nano particles are oxides, carbides, or nitrides. The coarse bright particles are refractory elements. The gray particles at grain boundaries are Cr-, Mo-, and C-rich while the black particles are Al-, O-, Y-, Zr-, and Si-rich.
FIGURE 1.

Fuel rod configuration showing hydrogen permeating from the fuel cavity to the coolant across the cladding wall. Surface oxides hinder H permeation flux.

FIGURE 1.

Fuel rod configuration showing hydrogen permeating from the fuel cavity to the coolant across the cladding wall. Surface oxides hinder H permeation flux.

Close modal
FIGURE 2.

Microstructure of the APMT specimens used for the electrochemical behavior and the hydrogen permeation tests.

FIGURE 2.

Microstructure of the APMT specimens used for the electrochemical behavior and the hydrogen permeation tests.

Close modal

Electrochemical Evaluations

Before running the permeation tests, the electrochemical behavior of APMT was characterized in relevant electrolytes. Four types of electrochemical tests were conducted: (a) linear polarization resistance (LPR) tests, which involved scanning the potential from −20 mV below the open-circuit potential (OCP) to +20 mV above OCP at a rate of 0.167 mV/s; (b) electrochemical impedance spectroscopy (EIS) by applying 10 mV (rms) sine wave with respect to OCP and measuring the impedance of the system from 100 kHz to 10 mHz; (c) cyclic potentiodynamic polarization (CPP) tests, which consisted of scanning the potential at a rate of 0.167 mV/s from −150 mVOCP until reaching a current density of 1 mA/cm2 and then reversing the scan direction until reaching −100 mVOCP; and (d) potentiostatic test by applying a constant potential of +300 mVSCE or +400 mVSCE. Prior to conducting any of the polarization measurements, the OCP was monitored for 24 h while test conditions stabilized.

The electrochemical evaluations were performed under deaerated conditions (high-purity N2) at 30°C in the following four electrolytes: (1) 3.5% NaCl; (2) 0.1 N NaOH; (3) 1 N H2SO4; and (4) 0.5 M Na2SO4. These electrolytes were selected because they spanned a wide pH range and involved typical anions such as Cl and SO42–. However, these electrolytes are not intended to represent the water chemistry in contact with the fuel rods in a reactor. A conventional three-electrode setup was used for conducting the measurements. A saturated calomel electrode (SCE) was used as a reference electrode, platinized-niobium served as a counter electrode, and the APMT specimen was the working electrode. The specimens were wet-polished with 600-grit silicon carbide paper within 1 h of starting the experiments. The exposed area of the specimens was approximately 1 cm2. The density of APMT is 7.25 g/cm3 and its equivalent weight is 22.7 (unitless).

Hydrogen Permeation Tests

Hydrogen diffusion through APMT specimens was determined using the Devanathan and Stachurski15  cell and according to the ASTM G148 and ISO 17081 standards.16-17  The permeation setup consisted of a thin APMT specimen (membrane) that was clamped between a hydrogen charging cell and an oxidation cell to create a hydrogen entry and a hydrogen exit side, respectively. The charging cell was filled with 3.5% NaCl and a cathodic potential was applied to the entry side of the specimen to generate hydrogen. The oxidation cell was filled with 0.5 M Na2SO4 solution and an anodic potential was applied to the exit side of the specimen to detect hydrogen. Atomic hydrogen (H) was generated on the cathodic side and it diffused through the specimen to the anodic side where it was oxidized to H+, which produced an oxidation current. This oxidation current is proportional to the amount of flux of hydrogen that permeates through the specimen.

The APMT surface that was exposed to the detection side was not electroplated with palladium, which is a common practice to reduce the background current and enhance oxidation kinetics.18  Instead, the specimens were tested with a bare surface finish condition (600-grit), which still allowed for very low background currents to be achieved based on the careful selection of the oxidation solution. The exposed area of the APMT specimens was approximately 1 cm2 on each side of the permeation cell.

Figure 3 is a schematic of the experimental setup that was used to conduct the hydrogen permeation tests. The APMT specimen was secured between two jacketed glass cells. Each cell included a gas inlet and outlet, an SCE reference electrode, and platinized-niobium counter electrode. In the charging cell, the SCE was housed in a right-angle Luggin capillary to position the reference electrode within a few millimeters of the specimen. The solution temperature in each cell was kept at 30°C using a thermostatically controlled recirculating bath. Both sides of the APMT surface were controlled by an independent potentiostat to allow for cathodic polarization in the charging cell and anodic polarization in the oxidation cell. The test procedure consisted of the following steps:
  • Introduce deaerated 0.5 M Na2SO4 solution to oxidation cell and then monitor OCP of the APMT specimen for 30 min.

  • Apply an oxidizing potential of +400 mVSCE and establish the oxidation/background current for at least 24 h. Ensure that the background current density is sufficiently low (less than 100 nA/cm2) prior to loading the charging cell with test solution.

  • Introduce deaerated 3.5% NaCl electrolyte to the charging cell and record OCP for 24 h.

  • Apply cathodic potential of −1 VSCE to charging cell for 48 h.

  • Remove applied potential from a charging cell for 48 h.

  • Repeat steps 4 and 5 to produce additional transients.

FIGURE 3.

Schematic of the experimental setup used for hydrogen permeation tests.

FIGURE 3.

Schematic of the experimental setup used for hydrogen permeation tests.

Close modal

Electrochemical Behavior of Advanced Powder Metallurgy Tubing at 30°C

As APMT is an alloy that was designed for high-temperature applications, little or no data exist on the electrochemical and corrosion behavior of this alloy in wet conditions. The objective of the current electrochemical tests was to select the most appropriate electrolyte to be used in the charging cell and oxidation cell of the permeation tests. Figure 4 shows the comparative CPP curves for APMT in 3.5% NaCl, 0.5 M Na2SO4, and 0.1 N NaOH solutions (reverse scans are omitted for clarity). The CPP curves in the figure show that the corrosion current density (icorr) of APMT was very low (approximately 1 × 10−8 A/cm2) for all three solutions, which indicated stable passivity around the corrosion potential. However, once the APMT was polarized to more positive potentials, stable passivity was not maintained in some cases. In the 3.5% NaCl solution, the polarization curve exhibited a sharp increase in current density that occurred near +300 mV, which was due to crevice corrosion initiation. In the 0.1 N NaOH solution, the polarization curve displayed current densities that were on the order of 10 µA/cm2 once the potential went above +150 mV, which suggested that passivity was being compromised at elevated potentials. The best range of passivity was found in the 0.5 M Na2SO4 solution; the polarization curve maintained low passive current densities (approximately 1 µA/cm2) until the potential was scanned above +900 mVSCE.
FIGURE 4.

Potentiodynamic polarization of APMT in NaCl, Na2SO4, and NaOH electrolytes.

FIGURE 4.

Potentiodynamic polarization of APMT in NaCl, Na2SO4, and NaOH electrolytes.

Close modal

Based on the CPP data that was obtained on APMT, it was decided that 3.5% NaCl solution could be used in the charging cell because there would be no concern of crevice corrosion under the applied cathodic potential (−1 VSCE). However, 3.5% NaCl solution would be a poor choice for the oxidation cell because of the threat of crevice corrosion under an oxidizing potential of +300 mVSCE. In the available standards for hydrogen permeation testing, it is recommended that an alkaline solution, such as 0.1 N NaOH, be used in the oxidation cell because it is conducive to passivating low-alloy steels. But for APMT alloy, the 0.1 N NaOH solution did not appear to be a good option for the oxidizing solution because of the higher “passive” current densities that were observed (Figure 4). Ultimately, the 0.5 M Na2SO4 solution was considered the most suitable electrolyte for the oxidation cell because it led to low passive currents over a wide potential range.

Constant potential tests were conducted on APMT to corroborate that 0.5 M Na2SO4 would be the ideal electrolyte for the oxidation cell. Figure 5 contains the results from the potentiostatic tests, which show that the current density for APMT decreased as the time increased. After 96 h (4 d) of constant potential treatment at +300 mVSCE, the current density in the NaOH electrolyte declined to around 200 nA/cm2, which was still too high of a background current for detecting hydrogen. For the same period (96 h) of constant potential treatment, the current density of APMT in the Na2SO4 electrolyte decreased to around 3 nA/cm2, which was almost two orders of magnitude lower than in the NaOH electrolyte. In addition, the APMT specimen showed substantial discoloration after the potentiostatic treatment in the NaOH electrolyte while no evidence of film damage was found after potentiostatic treatment in the Na2SO4 solution (Figure 6). These results confirmed that the most suitable electrolyte for the oxidation cell was 0.5 M Na2SO4.
FIGURE 5.

Potentiostatic polarization of APMT in (a) 0.5 M Na2SO4 (+0.4 VSCE) and (b) 0.1 N NaOH (+0.3 VSCE) electrolytes.

FIGURE 5.

Potentiostatic polarization of APMT in (a) 0.5 M Na2SO4 (+0.4 VSCE) and (b) 0.1 N NaOH (+0.3 VSCE) electrolytes.

Close modal
FIGURE 6.

Appearance of APMT specimen after potentiostatic polarization tests in (a) 0.1 N NaOH and (b) 0.5 M Na2SO4.

FIGURE 6.

Appearance of APMT specimen after potentiostatic polarization tests in (a) 0.1 N NaOH and (b) 0.5 M Na2SO4.

Close modal

Hydrogen Permeation Tests Through Advanced Powder Metallurgy Tubing at 30°C

The hydrogen permeation tests were performed on APMT specimens at 30°C and under deaerated conditions. The specimen side in the charging cell was exposed to 3.5% NaCl solution and was cycled between cathodic (−1 VSCE) and OCP conditions. The specimen side in the oxidation cell was exposed to 0.5 M Na2SO4 and was maintained at a potential of +0.4 VSCE. The specimens were bare APMT sheets of three different thicknesses (0.44 mm, 1.10 mm, and 1.99 mm). Three transients of hydrogen charging and discharging were conducted for each specimen thickness.

Figure 7 shows the permeation transients as a function of time for three periods of charging and discharging of hydrogen through the 0.44-mm-thick APMT specimen. Figure 7(a) displays the potential of the APMT specimen in the charging cell, which was cycled between −1 V and OCP for 48 h periods to produce each transient. Figure 7(b) contains the cathodic current density that was measured during the −1 V hold periods in the charging cell. Figure 7(c) shows the current density that was measured in the oxidation cell, which was very low (approximately 30 nA/cm2) prior to each hydrogen charging step. The results show that when the cathodic potential of −1 VSCE was applied, it generated a current density of approximately −5 µA/cm2 to −10 µA/cm2 (Figure 7[b]). The oxidation current due to hydrogen diffusion through the APMT specimen was in the range of 600 nA/cm2 to 650 nA/cm2 according to Figure 7(c). Similar results were obtained for the tests involving specimen thicknesses of 1.10 mm and 1.99 mm.
FIGURE 7.

Hydrogen charging and discharging transients at 30°C in 0.44-mm thick APMT specimen: (a) APMT potential in charging cell, (b) CD = current density or current generated on the charging cell, and (c) current generated on the oxidation or hydrogen discharging cell.

FIGURE 7.

Hydrogen charging and discharging transients at 30°C in 0.44-mm thick APMT specimen: (a) APMT potential in charging cell, (b) CD = current density or current generated on the charging cell, and (c) current generated on the oxidation or hydrogen discharging cell.

Close modal

Estimations of Effective Diffusivities Through Advanced Powder Metallurgy Tubing at 30°C

The current transients that were measured in the permeation tests were used to calculate the effective hydrogen diffusivity (Deff), which was the main objective of this work. The Deff of hydrogen was calculated using two different methods that are prescribed in ASTM G148 [2018].16  The first method is based on the break-through time (tb) and is calculated according to the following equation:
formula
where Deff is the effective diffusivity (cm2/s), L is the specimen thickness (cm), and tb is the break-through time (s). The breakthrough time was obtained from the intersecting point of the approximated baseline current and a line that was extrapolated from the slope of the permeation rise (Figure 7[c]).
The second method that was used for calculating the effective diffusivity was based on the elapsed time (tlag) and is shown in the equation below:
formula
where tlag is defined as the elapsed time at J(t)/Jss = 0.63, or the time required for the permeation flux (J) to reach 0.63 times the steady-state value (Jss). The value of Jss was determined by subtracting the background or baseline flux from the approximate steady-state flux (Figure 7[c]).

Both methods can be subject to sources of error in the approximation of Deff. For example, an unstable or inconsistent baseline current can lead to poor estimations of tb in the break-through method while the inability of a system to achieve “steady-state” permeation current (Jss) often confounds the application of the tlag method. In the literature, it seems that errors are more frequently encountered in the tlag method because there are a number of systems that can lead to surface films such as metal sulfides, corrosion layers, etc., which manifest as peaks in the permeation current rather than steady-state flux. The formation of voids in a material can also interfere with the attainment of steady permeation current. In this work, these issues were not encountered in the permeation data, which allowed for both methods to be utilized without any special caveats for the reported Deff values. The calculated values of Deff for the three transients that were obtained for each membrane thickness are summarized in Table 1. The Deff values were estimated using both the tb and tlag methods, which are shown in separate columns in the table. Both methods for calculating Deff yielded the same average value of 2.8 × 10−8 cm2/s, with only small differences in the standard deviation.

Table 1.

Effective Diffusivities for Hydrogen in APMT Membranes Generated by Applying a Cathodic Potential in 3.5% NaCl cell at 30°C

Effective Diffusivities for Hydrogen in APMT Membranes Generated by Applying a Cathodic Potential in 3.5% NaCl cell at 30°C
Effective Diffusivities for Hydrogen in APMT Membranes Generated by Applying a Cathodic Potential in 3.5% NaCl cell at 30°C

Comparison of Permeation Transients Through Advanced Powder Metallurgy Tubing at 30°C

The current transients that were measured in the permeation tests were also plotted in the form of normalized flux (J/Jss) vs. dimensionless time (Dt/L2) to produce rising transients, which are shown in Figure 8. Where D is the lattice diffusion coefficient of pure iron, which was 7.2 × 10−5 cm2/s.19  The plots in Figure 8 are organized by membrane thickness and each plot includes a theoretical transient that represents the lattice diffusion of pure iron. The normalized transients associated with the 0.44-mm-thick membrane (top plot) exhibited shallower slopes compared to the transients of the thicker membranes (1.10 mm and 1.99 mm). It was unclear whether this lower steepness in the rising slope was related to lower trap occupancy of the thinner membrane or a difference in surface conditions.16  The first permeation transient took slightly longer than the second and third transient for the tests involving membrane thicknesses of 0.44 mm and 1.99 mm, but in general, there was not a significant displacement of the permeation transients. The displacement of normalized transients to longer times indicates the existence of trap sites for hydrogen, which was outside the scope of this work to try to quantify.
FIGURE 8.

Rising hydrogen permeation transients for APMT in 3.5% NaCl at 30°C.

FIGURE 8.

Rising hydrogen permeation transients for APMT in 3.5% NaCl at 30°C.

Close modal

Fitting of Normalized Transients

An analytical solution to Fick’s second law of diffusion was used to analyze the permeation rises associated with the 0.44-mm thick membrane to evaluate whether the data conformed to Fick’s second law and to validate the Deff values that were determined from the manual methods (tb and tlag). The analytical solution that was used for analyzing the permeation rises can be expressed as follows:20-21 
formula
where it is the permeation current at time t, io is the initial steady-state (baseline) permeation current prior to the rising transient, i is the steady-state permeation current as t → ∞, and D is the diffusion coefficient, which is the same as effective diffusivity in this case. MATLAB was used to perform a nonlinear least-squares (NNLS) fit of the experimental permeation transients using Equation (4). The curve-fit of the first and second transients for the 0.44-mm thick membrane are shown in Figures 9(a) and (b), respectively. The curve fits are shown as solid lines (red curves) along with the experimental data points, which are plotted in the form of the normalized permeation current vs. time. In general, the MATLAB curve-fits correlated well with the experimental transients from both data sets, which indicated that H diffusion for this system was in accordance with Fick’s second law. Furthermore, the effective diffusivities that were determined from the curve-fits showed good agreement with the diffusivities that were obtained from the manual methods. Table 2 compares the Deff values for the three transients that were performed on the 0.44-mm thick membrane. For all three transients, the Deff values that were obtained from the MATLAB curve fitting were very close to the values obtained from the tlag method. The Deff values that were estimated from the break-through method (tb) were consistently higher by comparison but still similar in magnitude.
FIGURE 9.

MATLAB fit (red curves) of rising transients obtained on 0.44-mm-thick APMT membrane in 3.5% NaCl at 30°C: (a) first transient and (b) second transient.

FIGURE 9.

MATLAB fit (red curves) of rising transients obtained on 0.44-mm-thick APMT membrane in 3.5% NaCl at 30°C: (a) first transient and (b) second transient.

Close modal
Table 2.

Comparison of Methods for Determining Effective Diffusivities for Hydrogen in APMT Membrane (0.44-mm thick) in 3.5% NaCl at 30°C

Comparison of Methods for Determining Effective Diffusivities for Hydrogen in APMT Membrane (0.44-mm thick) in 3.5% NaCl at 30°C
Comparison of Methods for Determining Effective Diffusivities for Hydrogen in APMT Membrane (0.44-mm thick) in 3.5% NaCl at 30°C

Assumption of Volume-Controlled Transport

The estimated Deff values that are reported in Tables 1 and 2 are intended to represent volume-controlled transport (bulk diffusion) of hydrogen through the APMT alloy. Bulk diffusion implies that surface reactions such as transport through oxide films and absorption kinetics are not influencing the estimated diffusivities.16  This becomes especially important for thin membrane materials with high bulk diffusivities. One method of confirming that bulk diffusion is being characterized is to assess whether a linear relationship exists between the steady-state hydrogen flux and the inverse of the specimen thickness as given below:
formula
where Jss is the atomic H permeation flux at steady state (mol/s/cm2), Dl is the lattice diffusion coefficient (cm2/s), Co is the subsurface concentration of atomic H at the charging side (mol/cm3), and L is the specimen thickness (cm).
Figure 10 is a plot of Jss vs. L−1 for the first and second permeation transients that were obtained for the three different APMT membrane thicknesses (0.44 mm, 1.10 mm, and 1.99 mm). As evident from the figure, a linear relationship was found, which indicates that the Deff values in Table 1 correspond to volume-controlled transport. This approach is valid as long as Co is consistent between tests because it can affect diffussivity.22-23  The data in Figure 10 were obtained under the same charging conditions (i.e., 3.5% NaCl at −1 VSCE), therefore, it is expected that Co did not vary significantly between tests.
FIGURE 10.

Steady-state hydrogen flux (Jss) as a function of the inverse of the specimen thickness (L) for APMT in 3.5% NaCl at 30°C.

FIGURE 10.

Steady-state hydrogen flux (Jss) as a function of the inverse of the specimen thickness (L) for APMT in 3.5% NaCl at 30°C.

Close modal

Implications of the Findings

This is the first time that the Devanathan-Stachurski cell has been used to measure effective hydrogen diffusion coefficients through APMT alloy. Current findings show that the effective diffusion of hydrogen through ferritic APMT at 30°C is 2.8 × 10−8 cm2/s (Table 1), which is approximately three orders of magnitude lower than the Deff values reported for pure iron (8 to 9.5 × 10−5 cm2/s).24-26  The presence of alloying elements in APMT (Cr, Al, and Mo) would decrease the movement of hydrogen through the metal.6  The reduction in the effective diffusion of H in steel-containing alloying elements compared to pure iron is an established fact.6  The steady-state flux of hydrogen through APMT in Figure 10 is for cathodically generated atomic hydrogen on a clean surface. It has been shown many times that the presence of surface oxides on the charging side may limit the amount of atomic hydrogen that is generated and absorbed by the metal, which would affect migration to the other side of the membrane.6  Figure 6 shows that the membrane specimen did not show obvious oxide formation on the discharge or oxidation side of the membrane when the sodium sulfate electrolyte was used. Even though the current findings are not intended to actually embody the behavior of a nuclear fuel rod, it is a simplistic analogy. The APMT specimen that was used in this study is a representation of the cladding barrier between the fuel cavity (where the atomic hydrogen may be charged into the cladding wall) and the water side (where the permeated atomic hydrogen through the cladding wall would be oxidized to water). A recent review article6  showed that the permeation rates of hydrogen through FeCrAl such as APMT is between those of ferritic FeCr steels, which have faster diffusivities, and those of austenitic FeCrNi steels, which have slower diffusivities compared to APMT. As stated in the introduction, one of the concerns of using FeCrAl alloys such as APMT for the cladding of the fuel is the risk of an increase in the tritium concentration in the water. This risk of tritium increase in the coolant will be reduced once the cladding develops an aluminum oxide in the fuel cavity (charging side) and a chromium oxide on the water side. It has been reported that oxides can reduce the flux rates of hydrogen by a factor of 1,000.9  Thermodynamic calculations have been conducted which show that the dissociation of uranium dioxide in the fuel cavity may generate enough partial pressure of oxygen to allow for the formation of alumina on the inner diameter of the cladding. These predictions still need to be corroborated by performing post irradiation examinations on FeCrAl rods that are currently under irradiation in test and commercial reactors.

  • The Devanathan-Stachurski cell technique, which was the foundation for the ASTM G148 and ISO 17081 methods, was used to measure the effective hydrogen diffusion coefficient through APMT at 30°C.

  • For evaluating APMT specimens, it was determined that the most suitable electrolyte to use in the oxidation/detection cell was 0.5 M Na2SO4 and not the traditionally used NaOH electrolyte.

  • The average Deff for APMT was found to be 2.8 × 10−8 cm2/s for the tested conditions.

  • A linear relationship was found between the hydrogen flux and the inverse of the APMT specimen thickness, which indicates that the Deff values that were determined corresponded to bulk diffusion of hydrogen.

Trade name.

The authors would like to express their gratitude to Dr. Feng Gui (DNV, USA) for his help in analyzing the permeation data, Steve J. Buresh at GE Research for the thermal-mechanical processing of the APMT material, and Dr. Shenyan Huang (GE Research) for performing the metallographic studies of APMT membranes. This material is based on work supported by the Department of Energy [National Nuclear Security Administration] under Award No. DE-NE0009047. The financial support of Global Nuclear Fuel and GE Research is gratefully acknowledged. This report was prepared as an account of work sponsored by an agency of the United States Government. Neither the United States Government nor any agency thereof, nor any of their employees makes any warranty, express or implied, or assumes any legal liability or responsibility for the accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately owned rights. Reference herein to any specific commercial product, process, or service by trade name, trademark, manufacturer, or otherwise does not necessarily constitute or imply its endorsement, recommendation, or favoring by the United States Government or any agency thereof. The views and opinions of the authors expressed herein do not necessarily state or reflect those of the United States Government or any agency thereof.

1.
Rebak
R.B.
,
in
Accident Tolerant Materials for Light Water Reactor Fuels
(
Amsterdam, Netherlands
:
Elsevier
,
2020
),
p
.
45
47
.
2.
Rebak
R.B.
,
EPJ Nucl. Sci. Technol.
3
(
2017
):
article id 34
.
3.
Nagase
F.
,
Sakamoto
K.
,
Yamashita
S.
,
Corros. Rev.
35
,
3
(
2017
):
p
.
129
140
.
4.
Rebak
R.B.
,
Yin
L.
,
Andresen
P.L.
,
Corrosion
76
,
11
(
2020
), https://doi.org/10.5006/3632.
5.
Goodson
C.E.
,
Geelhood
K.J.
, “
Degradation and Failure Phenomena of Accident Tolerant Fuel Concepts
,”
Pacific Northwest National Laboratory, Richland, Washington, D.C., report no. PNNL-30445 (prepared for the Nuclear Regulatory Commission, Sept.
2020
).
6.
Garud
Y.S.
,
Hoffman
A.K.
,
Rebak
R.B.
,
Metall. Mater. Trans. A
53
(
2022
):
p
.
773
793
.
7.
Huang
S.
,
Dolley
E.
,
An
K.
,
Yu
D.
,
Crawford
C.
,
Othon
M.A.
,
Spinelli
I.
,
Knussman
M.P.
,
Rebak
R.B.
,
J. Nucl. Mater.
560
(
2022
):
p
.
153524
.
8.
Strasser
A.
,
Santucci
J.
,
Lindquist
K.
,
Yario
W.
,
Stern
G.
,
Goldstein
L.
,
Joseph
L.
, “
An Evaluation of Stainless Steel Cladding for Use in Current Design LWRs
,”
Electric Power Research Institute, Palo Alto, CA, report no. EPRI NP-2642, Dec.
1982
.
9.
Levchuk
D.
,
Bolt
H.
,
Döbeli
M.
,
Eggenberger
S.
,
Widrig
B.
,
Ramm
J.
,
Surf. Coat. Technol.
202
(
2008
):
p
.
5043
5047
.
10.
Strehlow
R.A.
,
Savage
H.C.
,
Nucl. Technol.
22
(
1974
):
p
.
127
137
.
11.
Van Deventer
E.H.
,
Maroni
V. A.
,
J. Nucl. Mater.
113
(
1983
):
p
.
65
70
.
12.
Sakamoto
K.
,
Hirai
M.
,
Ukai
S.
,
Kimura
A.
,
Yamaji
A.
,
Kusagaya
K.
,
Kondo
T.
,
Yamashita
S.
, “
Overview of Japanese Development of Accident Tolerant FeCrAl-ODS Fuel Claddings for BWRs
,”
in
WRFPM 2017 Conference (
Daejeon, South Korea
:
Korea Atomic Energy Research Institute
,
2017
).
13.
Sakamoto
K.
,
Miura
Y.
,
Ukai
S
,
Kimura
A.
,
Yamaji
A.
,
Kusagaya
K.
,
Kondo
T.
,
Yamashita
S.
, “
Progress on Japanese Development of Accident Tolerant FeCrAl-ODS Fuel Cladding for BWRs
,”
TopFuel 2018
,
paper no. A0011
(
Brussels, Belgium: European Nuclear Society
,
2018
).
14.
Rebak
R.B.
,
Kim
Y.J.
, “
Hydrogen Diffusion in FeCrAl Alloys for Light Water Reactors Cladding Applications
,”
2016 ASME PVP Conference, paper no. PVP2016-63164
(
Vancouver, BC
:
American Society of Mechanical Engineers
,
2016
).
15.
Devanathan
M.A.V.
,
Stachurski
Z.
,
Proc. R. Soc. A: Math. Phys. Eng. Sci. London
270
(
1962
):
p
.
90
102
.
16.
ASTM G148
, “
Evaluation of Hydrogen Uptake, Permeation, and Transport in Metals by an Electrochemical Technique
” (
West Conshohocken, PA
:
ASTM International
,
2018
).
17.
ISO International Standard 17081
, “
Method of Measurement of Hydrogen Permeation and Determination of Hydrogen Uptake and Transport in Metals by an Electrochemical Technique
” (
Geneva, Switzerland
:
ISO International
,
2014
).
18.
Manolatos
P.
,
Jerome
M.
,
Galland
J.
,
Electrochim. Acta
40
(
1995
):
p
.
867
871
.
19.
Turnbull
A.
,
Carroll
M.W.
,
Corros. Sci.
30
,
6/7
(
1990
):
p
.
667
679
.
20.
Fan
D.
,
White
R.E.
,
Gruberger
N.
,
J. Appl. Electrochem.
22
(
1992
):
p
.
770
772
.
21.
Owczarek
E.
,
Zakroczymski
T.
,
Acta Mater.
48
(
2000
):
p
.
3059
3070
.
22.
Turnbull
A.
,
Wright
L.
, “
Hydrogen Permeation Modelling with Generalised Boundary Conditions at the Charging Interface
,”
National Physical Laboratory, NPL report no. MAT 69
,
2014
.
23.
Griffiths
A.J.
,
Turnbull
A.
,
Corros. Sci.
37
(
1995
):
p
.
1879
1881
.
24.
Brass
A.M.
,
Collet-Lacoste
J.R.
,
Acta Mater.
46
(
1998
):
p
.
869
879
.
25.
Kiuchi
K.
,
McLellan
R.B.
,
Acta Metall.
31
,
7
(
1983
):
p
.
961
984
.
26.
Zakroczymski
T.
,
Scr. Metall.
19
,
4
(
1985
):
p
.
521
524
.