In situ ceramic layer growth on coated fuel particles dispersed in a zirconium metal matrix

Journal of Nuclear Materials 437 (2013) 171–177

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In situ ceramic layer growth on coated fuel particles dispersed in a zirconium metal matrix K.A. Terrani a,⇑, C.M. Silva a, J.O. Kiggans b, Z. Cai c, D. Shin b, L.L. Snead b a

Fuel Cycle and Isotopes Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA c Advanced Photon Source, Argonne National Laboratory, Argonne, IL 60439, USA b

a r t i c l e

i n f o

Article history: Received 4 January 2013 Accepted 18 February 2013 Available online 27 February 2013

a b s t r a c t The extent and nature of the chemical interaction between the outermost coating layer of coated fuel particles embedded in zirconium metal during fabrication of metal matrix microencapsulated fuels were examined. Various particles with outermost coating layers of pyrocarbon, SiC, and ZrC have been investigated in this study. ZrC–Zr interaction was the least substantial, while the PyC–Zr reaction can be exploited to produce a ZrC layer at the interface in an in situ manner. The thickness of the ZrC layer in the latter case can be controlled by adjusting the time and temperature during processing. The kinetics of ZrC layer growth is significantly faster from what is predicted using literature carbon diffusivity data in ZrC. SiC–Zr interaction is more complex and results in formation of various chemical phases in a layered aggregate morphology at the interface. Published by Elsevier B.V.

1. Introduction Metal matrix microencapsulated (M3) fuels were recently proposed as an advanced fuel form for light water reactors (LWRs) with potential for utilization in other advanced platforms or target applications [1]. The fuel consists of a zirconium alloy metal matrix that hosts microencapsulated (or coated) fuel particles and has been described in detail elsewhere [2]. The coated fuel particle in turn is adopted from the technology originally developed for high-temperature gas-cooled reactors (HTGRs) [3,4]. At its center, the coated fuel particle contains a spherical fuel kernel that is a fissile-material-bearing ceramic. The fuel kernel is then coated with successive coating layers that are designed to contain fission products. Consistent among this family of fuels, the first coating layer is a porous graphite layer dubbed the buffer layer that absorbs the damage deposited by the fission recoils and provides the volume to contain fission gases. On the buffer is layered an isotropic pyrocarbon coating (PyC). The PyC, given it remains intact, is impermeable to fission gases and a majority of metallic fission products. This two-layer structure is what is commonly referred to as a bistructural isotropic (BISO) particle. However, through the addition of an additional SiC or ZrC shell and an outer PyC layer the tristructural isotropic (TRISO) fuel particle is formed. In adding these additional layers, fission products are held inside a SiC micropressure vessel of remarkable stability. The details of the fabrica⇑ Corresponding author. Tel.: +1 865 576 0264. E-mail address: [email protected] (K.A. Terrani). 0022-3115/$ - see front matter Published by Elsevier B.V. http://dx.doi.org/10.1016/j.jnucmat.2013.02.042

tion, properties, and irradiation behavior of coated fuel particles are described in detail elsewhere [3]. The interaction between the outermost coating layers of various microencapsulated fuel particles and the zirconium metal matrix during fabrication of M3 fuels was reported previously [2]. The focus of this paper is to provide a detailed characterization of the extent and nature of this interaction under specific fuel fabrication conditions. The characterization has been performed via micro X-ray diffractometry using a synchrotron source and scanning electron microscopy. It is recognized that this interaction could be potentially exploited to produce a ZrC coating layer on the surface of BISO particles once they are incorporated into the metal matrix to transform them into TRISO particles. This is of interest given that BISO particles are produced via a much more rapid fluidized-bed chemical vapor deposition process. A kinetic model for the growth of ZrC layer at the PyC–Zr interface has also been developed and presented here.

2. Surrogate fuel fabrication Fuel specimens for this study were fabricated using a graphite element hot-press. The metal matrix was produced using 300 lm zirconium metal powder feedstock (American Elements, Los Angeles, CA, lot 1901392447-TYP). The coated fuel particles used during this study are described in Table 1. TRISO particles that are missing the outer PyC layer and their outermost layer consist of SiC and ZrC are designated by S and Z in Table 1, respectively. Full TRISO particles with a SiC shell and an outer PyC are designated by F in Table

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Table 1 Description of the constituents of the coated fuel particles used during M3 fuel fabrication. Measurement (lm)

BISO

TRISO-S

TRISO-F

TRISO-Z

ZrO2 Kernel diameter Buffer thickness Inner PyC thickness SiC thickness ZrC thickness Outer PyC thickness

521.3 ± 16.9 92.9 ± 6.4 29.0 ± 1.6 N/A N/A N/A

552.8 ± 20.1 57.1 ± 4.8 61.9 ± 4.1 43.5 ± 2.0 N/A N/A

530 85 41 31 N/A 36

525 N/A 50 N/A 13 N/A

2.2

1300

2.0

Temperature [°C]

1.8

1100

1.6

1000

1.4 1.2

900

1.0 Punch Movemnet Profiles

800

0.8

700

0.6 0.4

600 10 MPa applied

0.2

500 0

10

20

30

40

50

60

70

80

Fractional Punch Movement

1200

0.0 90 100 110

Time [min] Fig. 1. Temperature and pressing profiles during various hot-pressing runs. 10 MPa of pressure on the die was applied at 25 min that resulted in rapid consolidation of the metal and fuel mixture.

1. The geometry of the BISO and TRISO-S particles was extensively characterized using a well-developed industrial procedure [5,6] to provide accurate input parameters for the kinetic study. Mixtures of zirconium metal powder and various coated fuel particles were poured into a graphite die for hot-pressing; no cold-pressing was performed prior to hot-pressing. After hotpressing at temperatures ranging from 900 to 1300 °C a fully dense metal matrix was formed around the coated fuel particles. Photomicrographs of various coated fuel particles embedded in the zirconium metal matrix have been included in a prior publication and are omitted here [2]. The pressure during the hot-pressing was 10 MPa and was only applied when the furnace was at the maximum temperature for each processing trial. The furnace was controlled using the readings from K-type thermocouples for processing trials at and below 1100 °C, while the high-temperature runs were controlled using the reading from a C-type thermocouple. In addition to the K- and C-type thermocouples, a pyrometer was also used to collect additional temperature information during the processing trails. Fig. 1 shows the temperature and the press profile during various hot-pressing runs. Note the remarkably fast densification rate for zirconium metal powder once the 10 MPa pressure is applied over the wide temperature range between 900 and 1300 °C.

Si(1 1 1) monochromator is used to select the energy of X-rays used in this study with a relative bandwidth of 104. The HXRM employs a gold zone plate of size 160 lm and an outmost zone width of 150 nm. The HXRM can achieve a focal spot of 200 nm and a focal flux higher than 109 photons/s/0.01% bandwidth, corresponding to a photon density gain of 105 [8]. The selection of the horizontal focal spot size is realized by use of a white-beam slit located 46 m upstream from the HXRM. The detector for diffraction studies used in this study was a Princeton Quad-RO CCD with 2084  2084 pixels and a pixel size of 24 lm. The samples prepared for the transmission experiment were thin slices of the microencapsulated particles (TRISO-S, TRISO-Z, and TRISO-F) embedded in the zirconium metal matrix by hotpressing at 1050 °C. The slices had been cut from the hot-pressed pellet and were then polished on both sides to produce flat and parallel 150 lm thick cross sections. The thin cross sections were necessary for the transmission experiment. An X-ray energy of 10.1 keV (k = 0.12276 nm) was used during the l-XRD experiment. At this energy the attenuation distances for PyC, SiC, and Zr are 2.8  103, 1.4  102, and 2.2  101 lm, respectively [9]. The thin slices were also meant to alleviate the complexities associated with distinguishing various interfaces due to the spherical geometry of the shells. Fig. 2 shows the front view and the side view of the slice and the direction of X-ray beam for the transmission experiments. The shape of the incoming X-ray beam was that of a 0.2 lm  0.2 lm square. Diffraction ring patterns were collected at 0.2 lm intervals using a Si CCD detector. The line profiles were taken perpendicular to the tangential line at the surface of the particle. The scan line lengths were 80 lm in order to span the various coating layers and the Zr metal matrix. Given the position of the CCD detector, a specific range of diffraction angles (2h) was mapped. The CCD detector was placed at various distances away from the sample or positioned at different angles with respect to the incoming X-ray beam in order to map a wide range of diffraction angles. The data from the various scans were then bundled together during the analysis. The Debye ring patterns were azimuthally integrated to tally peak intensity as a function of angle for each spatial location. Fig. 3 shows the contour map of peak intensity as a function of diffraction angle and spatial location for the interface between Zr metal and TRISO-S, TRISO-F, and TRISOZ particles. All the peaks in the TRISO-F/TRISO-Z–zirconium metal interfaces have been identified, whereas some of the peaks have been left unidentified for the TRISO-S–Zr case to be analyzed later. While clear evidence for the presence of reaction products at the interface is present in the cases of the SiC-Zr and PyC–Zr interfaces, no new phases are present at the ZrC–Zr interface, though this

Zr Matrix

OPyC SiC IPyC Buffer ZrO2

Beam Direction

3. Characterization results 3.1. Micro X-ray diffractometry (l-XRD)

l-XRD was performed at the Argonne National Laboratory Advanced Photon Source (APS) using the 2-ID-D micro-diffraction beamline. A hard X-ray microprobe (HXRM) focuses the radiation generated from APS undulator A [7] to a spot size limited by the outermost zone width of the zone-plate optics. A double-crystal

~800µm

(a)

~150µm

(b)

Fig. 2. Front view (a) and side view (b) of the experimental slices used during the lXRD experiment.

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SiC 311

PyC-Zr

2.5

ZrC 222

Measured Calculated

ZrC 311

SiC Zr ZrC Zr2Si

2.0

Zr 10-13

1.5

4

ZrC 220

Sqrt (Intensity) [ ×10 a.u.]

2θ [o]

SiC 220

Zr 10-12

SiC

Zr SiC 200 ZrC 200 Zr 10-11

SiC 111

Zr 0002 ZrC 111

1.0 0.5 2.5

SiC-Zr

2.0 1.5

x [µm] 1.0 ZrC 400

0.5 25

SiC 311

30

35

40

45

50

55

60

65

2θ [°]

ZrC 222 Zr 20-21 Zr 11-22 ZrC 311 Zr 10-13

2θ [o]

SiC 220 Zr 11-20 ZrC 220

Fig. 4. Integrated peak intensity as a function of the diffraction angle over the reaction volume at the interface between Zr metal and TRISO-F (top) and TRISO-S (bottom) particles.

No Data

SiC

OPyC

ZrC

Zr

SiC 200 ZrC 200 Zr 10-11 Zr 0002 ZrC 111 Zr 10-10

SiC 111

x [µm] ZrC 400

ZrC 222

2θ [o]

ZrC 311

ZrC 220

No Data

IPyC

ZrC

Zr

ZrC 200

of the intensity to magnify low-intensity peaks. At the PyC–Zr interface, formation of crystalline ZrC is clearly apparent, whereas no diffraction peaks are expected from the amorphous PyC. The chemical interaction at the SiC–Zr interface proved more complex. A large number of studies on the ternary Zr–Si–C system have been performed with results reported as early as 1956 [12,13]. Chen et al. provide a complete thermodynamic phase stability overview of this body of studies up to recent years [14]. Upon investigation of the ternary Zr–Si–C phase diagrams, a number of reaction products are expected to form at the SiC–Zr interface. Calculated SiC–Zr isoplethal sections based on data from Chen et al. are shown in Fig. 5. Reaction of SiC–Zr has been specifically studied and at 900–1100 °C and a reaction path identified [15]. Specifically, formation of ZrC, Zr5Si3Cx, and Zr2Si at the interface is expected. In Fig. 4, formation of ZrC and Zr2Si was confirmed whereas Zr5Si3Cx was not included in the refinement fit. Although the Rietveld refinement takes into account the overlapping peaks, the small amount of the ternary phase present prevents it from being clearly distinguished in the pattern.

ZrC 111

3.2. Scanning electron microscopy x [µm] Fig. 3. Peak intensity as a function of the diffraction angle and spatial location for the interface between Zr metal and TRISO-S (top), TRISO-F (middle) and TRISO-Z (bottom) particles.

might be expected through inspection of the binary Zr–C phase diagram [10]. However, it is expected that the ZrC layer on the surface of the particle will undergo an increase in thickness at the expense of reduction in C/Zr ratio. An alternative representation of the diffraction patterns is provided in Fig. 4, where the peak intensity was integrated over the distance encompassing the reaction region at the PyC–Zr and SiC–Zr interfaces. Rietveld refinement was performed using the General Structure Analysis System [11] to fit the experimentally obtained patterns with structure factor calculations. The calculated intensities were compared to the experimental data in this manner to provide a clearer analysis. The results are plotted as square root

The SEM study was performed with a JEOL 6500 FEG-type microscope operating at 15 kV accelerating potential. Metallographic specimens were prepared to examine the cross section of BISO and TRISO-S particles embedded in Zr metal after fabrication at the different processing temperatures. As discussed in the previous section, the interaction between PyC outer layer in the BISO particle and the Zr matrix results in formation of ZrC. In the case of the SiC–Zr interface, a series of binary and ternary compounds form. Fig. 6 shows the interface between the matrix and the particles at various processing temperatures. The backscattered electron images at the PyC–Zr interface show contrast among various ZrC grains due to their varying crystallographic orientation (a channeling effect). At the SiC–Zr interface the contrast is due to both Z-contrast among the various phases as well as crystallographic orientation of the grains. In both cases, the reaction layer thickness along with the grain size increases with processing temperature. A small fraction of pores can be observed in the ceramic layers grown at 1300 °C.

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4000

2500 Liquid +SiC+ZrC

2000 1500

bcc+L +ZrC

Zr5Si3Cx +ZrC

SiC+ZrC+ZrSi2 bcc(Zr)+Zr3Si +ZrC

1000

0 0.2

0.4

0.6

0.8

ZrC

SiC +ZrC +α-ZrSi

hcp(Zr)+Zr3Si +ZrC

0

900

600

SiC+ZrC+α-ZrSi

500

1200

300 0.45

1.0

0.50

Zr5Si3C

Liquid +Zr5Si3Cx+ZrC

Zr2Si+Zr Si 5 3Cx+ZrC

Liquid +ZrC

Zr5Si3Cx +ZrC +α-ZrSi

Zr5Si3Cx+α-Zr5Si4+ZrC

Temperature, C

3000

ZrC +ZrSi2 +α-ZrSi

α-Zr5Si4+ZrC+α-ZrSi

Liquid

Liquid +graphite

Temperature, C

3500

x+ZrC

1500

Zr2Si +Zr3Si +ZrC

Zr3Si +ZrC

0.55

Mole Fraction, Zr

Zr3Si +Zr5Si3Cx bcc(Zr) +ZrC +Zr3Si +ZrC

0.60

0.65

hcp(Zr) +Zr3Si +ZrC

0.70

0.75

Mole Fraction, Zr

Fig. 5. Calculated SiC–Zr isoplethal sections. The Si/C ratio is fixed at unity.

900°C

1000°C

1100°C

1200°C

1300°C Zr

ZrC

IPyC

Zr Zr-SiC interaction layer

SiC

10 µm

Fig. 6. Backscattered electron image of PyC–Zr (top) and SiC–Zr (bottom) interfaces as function of temperature after 1 h of hot-pressing.

x

Cross section

h r1

t

θ

r2

midplane

ginal particle and the particle after interaction layer growth are denoted as r1 and r2, respectively. x is the diameter of the circular cross section of the particle, and t is the apparent thickness of the interaction layer given any cross section at height h away from the midplane. The true thickness of the interaction layer is the difference between the two radii r1 and r2 that is given by:

0qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 a2 sin2 ðhÞ þ 1 r2  r1 ¼ r1 @  1A;

a

ð1Þ

where a is defined as Fig. 7. Schematic of a spherical particle cross sectioned at height h away from the midplane.

The true interaction layer thickness cannot be easily discerned from the micrographs in Fig. 6 owing to the spherical nature of the particles. Effectively the interaction layer thickness is always exaggerated unless the imaged cross section is exactly at the particle mid-plane. Fig. 7 depicts a spherical particle with a cross section made at a position in the upper half. The radius of the ori-

a¼

r1 : x=2 þ t

ð2Þ

The true interaction layer thickness is calculated using Eq. (1) given r1 that is determined initially by examining the particles prior to their insertion in the metal matrix (Table 1) and x and t that are measured during the post-processing microscopy. Fig. 8 shows the variation in the true interaction layer thickness at the surface of BISO and TRISO-S particles after hot-pressing at various processing temperatures for 1 h.

K.A. Terrani et al. / Journal of Nuclear Materials 437 (2013) 171–177

the Bernoulli equation-nonlinear but with a known exact solution. Substituting back for the concentration gradient into the solution the carbon concentration gradient at any point inside the ZrC shell is given by

Interaction Layer Thickness [μm]

12 PyC-Zr SiC-Zr

10

dC 1 ¼ ; dr rðb þ k1 rÞ

8 6

ð7Þ

4

where k1 is a constant. After integrating Eq. (7) one arrives at the solution for the steady state concentration for any thickness of the ZrC layer:

2

C¼

0 900

1000

1100

1200

  1 b þ k1 r ln þ k2 ; b r

bð1  nÞ ; nr 1  r 2

Fig. 8. Interaction layer thickness at the surface of BISO and TRISO-S particles vs. 1 h hot-pressing temperature.

k1 ¼

4. Growth kinetics

k2 ¼ C 1 

4.1. ZrC layer growth

where n is given by

Formulation of the ZrC layer growth can be approached here as a diffusion-limited reaction. The diffusing species is carbon that travels through the ZrC lattice to arrive at the ZrC–Zr interface and react to contribute to layer growth. Formation and the effects of any porosity are ignored in this analysis. The transient diffusion equation for carbon in the ZrC layer in spherical coordinates is as follows:

n¼

  @C 1 @ @C ¼ 2 r2 D ; @t r @r @r

ð3Þ

where C is the C/Zr ratio in the ZrC lattice and D is the diffusivity of carbon in the ZrC lattice. Note that both sides of Eq. (3) can be multiplied by the number density of zirconium in ZrC if one wishes to calculate the number density of carbon. By examining the Zr–C binary phase diagram [10] it is noted that a wide range of hypostoichiometry for carbon in the ZrC lattice exists; however, any positive deviation from the stoichiometric composition will result in precipitation of carbon. Assuming layer growth is not interface reaction limited, it is expected that the C/Zr ratio at the ZrC–PyC interface will be always fixed at 1 (at least for T < 2900 °C, above which maximum C/Zr ratio decreases even further) and C/Zr ratio at the Zr–ZrC interface will be given by the minimum hypostoichiometric composition at any given temperature. A review of diffusivity of carbon in ZrC is provided by Van Loo et al. [16], and it is reported to be dependent on C/Zr ratio with the following form:

D ¼ D ebC eE=RT ;

ð4Þ

where D, b, and E are reported to be 82.1 cm2/s, 9.2, and 304 kJ/ mole, respectively. To solve the diffusion–reaction problem one can assume steady state conditions whereby Eq. (3) yields the following:

  D eE=RT d 2 bC dC r ¼ 0: e dr dr r2

ð5Þ

By differentiating the term in parenthesis, Eq. (5) can be rewritten as



dC dr

2

2

þr

d C dr

2

þ2

dC ¼ 0: dr

ð8Þ

where k2 is also a constant. Given the boundary conditions discussed earlier, the constants k1 and k2 for any ZrC layer shell with an inner and outer radii of r1 and r2 are defined as

1300

Temperature [ C]

br

175

ð6Þ

Fortunately, by substituting any arbitrary parameter for the concentration gradient (dC/dr), Eq. (6) reduces to a special manifestation of

ð9Þ

  1 b þ r 1 k1 ; ln b r1

r 2 bðC 2 C 1 Þ e r1

ð10Þ

ð11Þ

and C1 and C2 are the C/Zr ratio at the PyC–ZrC and ZrC–Zr interfaces, respectively. Once the steady state diffusion profile for any layer thickness is determined, the growth kinetics can be determined given the diffusive flux of carbon atoms into the interface. Again, note that the interface reaction kinetics are assumed to be rapid and are not considered as the rate-limiting step. The ZrC shell thickness increase is then given by

 dr J r2 dC  ¼ ¼ Dr2  ; dt dr r2 N

ð12Þ

where Jr2 is the diffusive flux of carbon atoms into the ZrC–Zr interface and N is the number density of zirconium atoms in the ZrC lattice. The concentration gradient at the interface is given by Eq. (7). The following integration can then be formulated to correlate time with layer thickness:

Z

r2

rðb þ k1 rÞdr ¼

r1

Z

t

0

Dr2 dt:

ð13Þ

By performing the integration, the relationship between layer thickness and time is given by

2r 32  3r 1 r 22 þ r 31 ¼ k3 t:

ð14Þ

In this assumed that the ZrC layer grows into the Zr metal matrix. Applying initial condition r1 = r2 at t = 0, k3 is determined to be

k3 ¼

6Dr2 r 1 bðC1 C2 Þ  1Þ: ðe b

ð15Þ

Given the methodology developed above, the diffusion-limited growth model is compared to the experimental layer thickness data, as shown in Fig. 9. The model drastically under predicts the experimental data. It has been observed that the interdiffusion coefficient in the NaCl-type structure of ZrC is essentially equal to the intrinsic diffusion coefficient of carbon in the lattice [16]. Therefore faster zirconium atom diffusion through the grains in the ZrC layer is not expected. Better fits with the experimental results can be achieved if the pre-exponential term in the diffusion coefficient in Eq. (4) is increased arbitrarily by two orders of magnitude, as shown in Fig. 9. Effectively, it is assumed that the diffusion coefficient of carbon during the thin layer growth with small grains is signifi-

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ZrC Layer Thickness [μm]

10

take 200 years at 700 °C. Assuming irradiation does not significantly enhance the transport processes across the layer, the particle and coating geometries are expected to remain stable under irradiation during the fuel’s lifetime.

BISO Particles Hot-Pressed for 1 Hour

1

4.2. SiC–Zr interaction layer growth 0.1 Experimental Data Diffusion Limited Model Do × 100

0.01

Grain Boundary Diffusion [17]

900

1000

1100

1200

1300

Temperature [ C] Fig. 9. Comparison between the experimental ZrC layer growth data and prediction results from Eq. (14) with diffusion coefficient from Ref. [16] and 1-D Cartesian parabolic growth with grain boundary diffusion coefficient from Ref [17].

5. Conclusions

Interaction Layer Thickness [μm]

25 20 15

As discussed earlier, SiC–Zr interaction results in formation of various phases at the interface. In SiC–M couples where M is a metal that forms carbide, it has been shown that both M and carbon are the diffusing species in the reaction zone [15]. The growth kinetics is not treated here in the level of detail presented for ZrC layer growth. Instead, a review of the few reported data is performed and compared with the results here [15,18]. Specifically, Fig. 10 examines the growth kinetics at 1100 °C. Sub-parabolic, even sub-cubic, kinetics appear to be the case at this temperature. The complex layered aggregate structure with altering grain size throughout the thickness could explain the observed kinetics by reducing the effective diffusion path for the reactant species.

1100 C growth data This study Vishnyakov et al. Bhanumurthy et al.

10

m=0.3

5

1

10

100

Time [h] Fig. 10. SiC–Zr interaction layer thickness as a function of time at 1100 °C.

cantly higher if other mechanisms aside from the bulk variant are considered. Another way to achieve a better fit with the experimental data is to increase the magnitude of the concentration coefficient b in Eq. (4) approximately twofold to account for possibly more rapid carbon diffusion at very low C/Zr ratios that have not been captured in Ref. [16]. Note that the growth during the temperature ramp and cool down periods is ignored in the model, though it is expected to be of little contribution to the total layer thickness. Sarian and Criscione reported the concentration–independent bulk and grain boundary diffusivities during their early studies on carbon diffusion in ZrC [17]. If one assumes one-dimensional Cartesian parabolic growth where the thickness square is essentially equal to twice the product of time and diffusivity (l2 = 2Dt), the calculated thickness using grain boundary diffusivities from Ref. [17] is even smaller than what is predicted by the model here as shown in Fig. 9. Extrapolation of the experimental data to predict lower temperature growth, pertaining to nominal fuel operating temperatures, is useful. The maximum fuel temperature is deemed to be on the order of 700 °C during normal operating conditions in the LWR fuel pin geometry [1]. Extrapolation of the experimental data points using the developed model with rapid diffusivity shows that for the ZrC layer thickness to increase to 20 lm from 10 lm would

During production of M3 fuel specimens, the zirconium metal matrix chemically interacts with the outermost coating layer of the fuel particle. The overall extent and nature of this interaction, given the material at the interface present alongside zirconium were evaluated for temperatures ranging between 900 and 1300 °C under a 1 h processing window. The interaction is the least noticeable for the case of a ZrC outermost coating layer. For the case where PyC is the outermost layer, this interaction results in formation of a ZrC coating layer at the interface in an in situ manner. The thickness of the coating layer is likely governed by carbon diffusion through the ZrC layer and can be controlled during processing by adjusting time and temperature. The effective diffusivity of carbon across the ZrC layer during layer growth is roughly two orders of magnitude larger than what has been reported in literature. Extrapolation of the fast growth kinetics observed here to lower temperatures where the fuel is expected to operate under nominal conditions shows that layer thickness growth is negligible during the fuel’s lifetime. Finally, if the outermost coating layer is SiC, a complex chemical interaction with zirconium matrix is expected that will result in formation of various binary and even ternary compounds. The kinetics of interaction layer growth is well below what would be expected from simple parabolic kinetics. Acknowledgments The aid and technical insight of Eliot Specht and Theodore Besmann at ORNL is gratefully acknowledged. Use of the Advanced Photon Source was supported by the U.S. Department of Energy, Office of Basic Energy Sciences. The JEOL6500 FEG SEM was supported by ORNL’s Shared Research Equipment (ShaRE) User Facility, which is sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, US Department of Energy. The work presented in this paper was also partially supported by the Advanced Fuels Campaign of the Fuel Cycle R&D program in the Office of Nuclear Energy, US Department of Energy and Laboratory Directed R&D funds at ORNL. The United States Government retains, and by accepting the article for publication, the publisher acknowledges that the United States Government retains, a non-exclusive, paid-up, irrevocable, worldwide license to publish or reproduce the published form of this work, or allow others to do so, for United States Government purposes.

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References [1] K.A. Terrani, L.L. Snead, J.C. Gehin, J. Nucl. Mater. 427 (2012) 209–224. [2] K.A. Terrani, J.O. Kiggans, L.L. Snead, J. Nucl. Mater. 427 (2012) 79–86. [3] D.L. Hanson, A review of radionuclide release from HTGR cores during normal operation, EPRI Report 1009382, Electric Power Research Institute, 2004. [4] D. Petti, J. Maki, J. Hunn, P. Pappano, C. Barnes, J. Saurwein, S. Nagely, J. Kendall, R. Hobbins, JOM 62 (9) (2010) 62. [5] J.R. Price, D.B. Aykac, J.D. Hunn, A.K. Kercher, R.N. Morris, in: Proceedings of Machine Vision Applications in Industrial Inspection XIV, SPIE, vol. 6070, 2006, p. 60700H. [6] J.R. Price, D.B. Aykac, J.D. Hunn, A.K. Kercher, in: Proceedings of Machine Vision Applications in Industrial Inspection XV, SPIE, vol. 6503, 2007, p. 650302. [7] R.J. Dejus, I.B. Vasserman, S. Sasaki, E.R. Moog, Report ANL/APS/TB-45, Argonne National Laboratory, Argonne, IL, USA, 2002.

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[8] Z. Cai, B. Lai, W. Yun, P. Ilinski, D. Legnini, J. Maser, W. Rodrigues, AIP Proc. 507 (2000) 472–477. [9] LBNL, Center for X-ray Optics, . [10] A.F. Guillermet, J. Alloys Compd. 217 (1995) 69–89. [11] A.C. Larson, R.B. Von Dreele, General Structure Analysis System (GSAS), Los Alamos National Laboratory Report LAUR 86-748, 2000. [12] L. Brewer, O. Krikorian, J. Electrochem. Soc. 103 (1956) 38–51. [13] H. Nowotny, B. Lux, H. Kudielka, Monatsh. Chem. 87 (1956) 447–470. [14] H.M. Chen et al., J. Alloys Compd. 474 (2009) 76–80. [15] K. Bhanumurthy, R. Schmid-Fetzer, Scripta Mater. 45 (2001) 547–553. [16] F.J.J. Van Loo, W. Wakelkamp, G.F. Bastin, R. Metselaar, Solid State lonics 32 (33) (1989) 824–832. [17] S. Sarian, J.M. Criscione, J. Appl. Phys. 38 (1967) 1794. [18] L.P. Vishnyakov, V.P. Moroz, V.A. Pisarenko, A.V. Samelyuk, Powder Metall. Met. Ceram. 46 (1–2) (2007).