4.07

Radiation Effects in SiC and SiC–SiC

L. L. Snead and Y. Katoh
Oak Ridge National Laboratory, Oak Ridge, TN, USA

T. Nozawa
Japan Atomic Energy Agency, Rokkasho, Aomori, Japan

ß 2012 Elsevier Ltd. All rights reserved.

4.07.1
4.07.2
4.07.3
4.07.4
4.07.4.1
4.07.4.2
4.07.4.3
4.07.4.4
4.07.5
4.07.6
References

Introduction
Irradiation-Induced Swelling and Microstructure of Pure SiC
Irradiation-Induced Thermal Conductivity Degradation of Monolithic SiC
Effect of Irradiation on the Mechanical Properties of Monolithic SiC
Elastic Modulus of Monolithic SiC
Hardness of Monolithic SiC
Fracture Toughness of Monolithic SiC

Strength and Statistical Variation in Strength for Monolithic SiC
Irradiation Creep of SiC
Silicon Carbide Composites Under Irradiation

Abbreviations
ATR
BSD
BSR
CVD
CVI
dpa
DuET
ETR
FB
fcc
HFBR
HFIR
HFIR-METS
HNLS
HP
JMTR
Kd
Kgb
Kirr
Knonirr
Krd
Ku
PLS
PS


Advanced test reactor
Black spot dot
Bend stress relaxation
Chemical vapor deposition
Chemical vapor infiltration
Displacement per atom
Dual-beam facility for energy
science and technology
Engineering test reactor
Fluidized bed
Face-centered cubic
High-flux beam reactor
High Flux Isotope Reactor
High-flux isotope reactor – mapping
elevated temperature swelling
Hi Nicalon Type S
Hot-pressing
Japan materials testing reactor
Thermal conductivity by defect
scattering
Thermal conductivity by grain
boundary scattering
Irradiated thermal conductivity
Nonirradiated thermal conductivity
Thermal conductivity by radiation
Thermal conductivity by Umklapp
scattering
Proportional limit stress
Pressureless sintering


PW
PyC
SEM
SiC/SiC
composite
SW
TEM
Tirr
TRISO
TySA
UTS

215
216
221
224
224
226
226
227
233
234
239

Plain weave
Pyrolytic carbon
Scanning electron microscopy
Silicon carbide fiber reinforced
silicon carbide matrix composite
Satin weave

Transmission electron microscopy
Irradiation temperature
TRIstructural ISOtropic
Tyranno SA
Ultimate tensile stress

4.07.1 Introduction
Silicon carbide (SiC) has been studied and utilized
in nuclear systems for decades. Its primary use was,
and still is, as the micro pressure vessel for hightemperature gas-cooled reactor fuels. For these
so-called TRI-ISOtropic (TRISO) fuels, the SiC is
deposited via a gas-phase decomposition process over
two layers of pyrolytic graphite surrounding the fuel
kernel. In addition to being strong enough to withstand the pressure buildup from the fission product
gas liberated, this SiC layer must also withstand
chemical attack from metallic fission products such
as palladium and the mechanical loads derived from
irradiation-induced dimensional changes occurring
in the pyrolytic graphite. More recent nuclear applications of SiC include its use as structural composites
215


216

Radiation Effects in SiC and SiC–SiC

support of nuclear fuel coating1–9 and more recently,
for various nuclear applications such as structural SiC
composites.10 Before proceeding, it is important to
distinguish neutron-induced effects on high-purity

materials, such as single crystal and most forms of
chemical vapor deposited (CVD) SiC, from those on
lower purity forms such as sintered with additives,
reaction-bonded, or polymer-derived SiC. It is well
understood that the presence of significant second
phases and/or poorly crystallized phases in these materials leads to unstable behavior under neutron irradiation,11–14 as compared to stoichiometric materials,
which exhibit remarkable radiation tolerance. Discussion and data for this section refer only to high purity,
stoichiometric, near-theoretical density SiC, unless
otherwise specified. Rohm and Haas (currently Dow
Chemicals) CVD SiC is an example of such material.
The irradiation-induced microstructural evolution of CVD SiC is roughly understood and has
been reviewed recently by Katoh et al.15 An updated
version of the microstructural evolution map is shown
in Figure 1. However, the contribution of the defects
themselves to the swelling in SiC is less understood.
Below several hundred Kelvin, the observable

(i.e., SiC/SiC) for high-temperature gas-cooled reactors and for fusion power systems. The possibility of
using composite and monolithic SiC thermal insulators for both fusion and fission systems is also being
investigated. Moreover, both monolithic and composite forms of SiC are being investigated for use in
advanced sodium fast, advanced liquid salt-cooled,
and advanced light water reactors.
In this chapter, the effects of neutron irradiation
on relatively pure, radiation resistant forms of SiC
are discussed. This chapter has been limited to the
effects of irradiation on the microstructure, and the
mechanical and thermal properties of SiC, although it
is recognized that environment aspects such as oxidation and corrosion will also be factors in eventual
nuclear application. These areas are not discussed here.


4.07.2 Irradiation-Induced Swelling
and Microstructure of Pure SiC
The neutron-induced swelling of SiC has been well
studied for low and intermediate temperatures ($293–
1273 K). Originally, this material was investigated in
1600

1400

6

7

7

6
6

6
7

Irradiation temperature (ЊC)

1200

6

1000

800


7

3

Black spot
defects(BSD)
and/or
unidentified
small loops

6

6

BSD and/or
unidentified
small loops

7

Frank
loops

Frank loops

6
1

Unfaulted loops

and/or network
Voids

1

1

6

6

4

600

6

6
1

1

6
3

Larger loops
dislocation network
voids

1


1

7

5

6

5

2

1
6

1

6

2
2

(1973)4

400

200
0.1


1. Price
2. Yano (1998)17
3. Senor (2003)18
4. Iseki (1990)19
5. Katoh, neutron (2006)15
6. Katoh, ion (2006)15
7. Snead (2007)16

1

2

5

10
Fluence (dpa)

100

1000

Figure 1 Updated summary of the microstructural development in cubic SiC during neutron and self-ion irradiation.
Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.


Radiation Effects in SiC and SiC–SiC

Tirr = 300 °C, 6 dpa

Tirr = 800 °C, 7.7 dpa

Loop number density » 3.3e + 23 m-3
Mean loop diameter = 3.0 nm

Dot number density » 2.2e + 24 m-3
Mean dot diameter = 1 nm

g = 200

217

40 nm
g = 200

50 nm

Figure 2 Microstructure for CVD neutron irradiated at 573 and 1073 K.

microstructure of neutron-irradiated SiC is
described as containing ‘black spots, which are most
likely tiny clusters of self-interstitial atoms in various
indeterminate configurations. For irradiation temperatures less than about 423 K, accumulation of
strain due to the irradiation-produced defects can
exceed a critical level above which the crystal
becomes amorphous. This has been shown in the
case of both self-ion irradiation and fast neutron irradiation.20–22 As shown by Katoh et al.,23 the swelling at
323 K under self-ion irradiation increases logarithmically with dose until amorphization occurs. The
swelling of neutron- and ion-amorphized SiC has
been reported to be 10.8% for 343 K irradiation.22
However, there is evidence that the density of amorphous SiC will depend on the conditions of irradiation
(dose, temperature, etc.)24

For temperatures above the critical amorphization temperature (423 K), the swelling increases
logarithmically with the dose until it approaches
saturation, with a steady decrease in the saturation
swelling level with increasing irradiation temperature. The dose exponents of swelling during the
logarithmical period are in many cases close to twothirds, as predicted by a kinetic model assuming
planar geometry for interstitial clusters.25 This temperature regime is generally referred to as the pointdefect swelling regime and can be roughly set
between 423 and 1273 K. As an example of how
these ‘black spot’ defects mature in the point-defect
swelling regime, Figure 2 shows neutron-irradiated
microstructures at 573 and 1073 K for doses

consistent with a saturation in density. While these
microstructural features are generically classified
as ‘black spots,’ the defects formed at 1073 K are
clearly coarser compared to those formed under
573 K irradiation.
The approach to saturation swelling is shown for
High Flux Isotope Reactor (HFIR) neutron irradiated Rohm and Haas CVD SiC in Figure 3. In
this figure, the swelling is depicted in both logarithmic (Figure 3(a)) and linear (Figure 3(b)) plots. In
addition to the approach to saturation, this figure
highlights two other characteristics of neutroninduced swelling of SiC. First, the swelling of SiC is
highly temperature dependent. For the data given in
Figure 3, the 1 dpa and saturation values of swelling
at 473 K are approximately five times that for
1073 K irradiation. This reduced swelling with
increasing irradiation temperature is primarily attributed to enhanced recombination of cascadeproduced Frenkel defects due to lower interstitial
clustering density at higher temperatures. The second characteristic swelling behavior to note is that
the swelling saturates at a relatively low dose. For
damage levels of a few dpa (typically months in a
fission power core), the swelling in the point-defect

recombination range has found its saturation value.
At higher temperatures such as 1173–1673 K,4,18,26
Frank faulted loops of the interstitial type become
the dominant defects observed by transmission electron microscopy (TEM). It has also been reported
that Frank faulted loops appear for lower temperature neutron irradiation at extremely high doses.27


218

Radiation Effects in SiC and SiC–SiC

3
CVD SiC

CVD SiC
2

200 °C

1

400 °C

2.5

200 °C
600 °C

Swelling (%)


Swelling (%)

2

650 °C

300 °C

1.5
400 °C

800 °C

1
600 °C

800 °C

0.1

0.5

0
0
(a)

1

2


3
4
5
6
Neutron dose (dpa)

7

8

0.01
(b)

0.1
1
Neutron dose (dpa)

Tirr = 200 °C

Tirr = 300 °C

Tirr = 400 °C

Tirr = 500 °C

Tirr = 600 °C

Tirr = 800 °C

10


30

Figure 3 Swelling of SiC in the intermediate temperature point defect swelling regime. Reproduced from Snead, L. L.;
Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

Under silicon ion irradiation at 1673 K, the development of Frank loops into dislocation networks through
unfaulting reactions at high doses is reported.26 The
volume associated with dislocation loops in irradiated
SiC has been estimated to be on the order of 0.1%.28
At temperatures where vacancies are sufficiently
mobile, vacancy clusters can be formed. Threedimensional (3D) cavities (or voids) are the only
vacancy clusters known to commonly develop to
large sizes in irradiated SiC. The lowest temperature
at which void formation was previously reported
under neutron irradiation is 1523 K.4 Senor reported
the lack of void production after neutron irradiation
to 0.9 dpa at 1373 K, although voids were observed
after subsequent annealing at 1773 K for 1 h.18 Under
silicon ion irradiation, voids start to form at 1273 K
at very low density and become major contributors
to swelling at irradiation conditions of 1673 K at
>10 dpa.29 Positron annihilation and electron paramagnetic resonance studies have shown that the
silicon vacancy in cubic SiC becomes mobile at
1073–1173 K.30,31 Therefore, it would not be surprising for void swelling to take place at as low
as $1273 K at high doses, particularly for low damage
rate irradiations.

As previously mentioned, data on swelling of
SiC in the high-temperature ‘void swelling’ regime

has been somewhat limited. Recently, however,
work has been carried out in the $1173–1773 K
range for Rohm and Haas CVD SiC irradiated in
HFIR. Of particular significance to that experiment
is the confidence in irradiation temperature owing to
the melt-wire passive thermometry.32 Recent TEM
imaging by Kondo28 clearly shows the evolution of
complex defects. As an example, Figure 4 indicates
sparse void formation on stacking faults for material
irradiated at 1403 K. Significant growth of voids
commences at 1723 K. The well-faceted voids
appeared to be tetrahedrally bounded by planes,
which likely provide the lowest surface energy in
cubic SiC. In many cases, voids appeared to be aligned
on stacking faults at all temperatures. However, intragranular voids unattached to stacking faults were also
commonly observed at 1723 K. The evolution of dislocation microstructures at 1403–1723 K is shown
in Figure 5. In this temperature range, dislocation
loops are identified to be Frank faulted loops of interstitial type. Evolution of the dislocation loops into
dislocation networks was confirmed for irradiation
at 1723 K.


Radiation Effects in SiC and SiC–SiC

20 nm

(a)

1280 ЊC, 5.0 dpa


1130 ЊC, 8.5 dpa

1130 ЊC, 1.8 dpa

(b)

(c)
1450 ЊC, 8.5 dpa

1450 ЊC, 5.0 dpa

1450 ЊC, 1.8 dpa

(d)

219

(e)

(f)

Figure 4 Evolution of voids in high-temperature irradiated CVD SiC.

1130 ЊC, 8.5 dpa

1130 ЊC, 1.8 dpa

30 nm

(a)


(b)

(c)
1450 ЊC, 5.0 dpa

1450 ЊC, 1.8 dpa

(d)

1280 ЊC, 5.0 dpa

g

(e)

1450 ЊC, 8.5 dpa

g

(f)

g

Figure 5 Evolution in dislocation networks for high-temperature irradiated CVD SiC.

Figure 6 plots both historical data, recently published, and unpublished data on the swelling behavior
of SiC over a wider range of temperature.16,33 This
plot is limited to literature data on high-purity CVD
SiC. A divergence from point-defect ‘saturated’

swelling to unsaturated swelling is observed in the
1273–1473 K range, although additional data in this
temperature range as a function of fluence would be
required to precisely define such behavior. Above
1373 K, there exists a clear unsaturated swelling
behavior for CVD SiC. The three divergent curves

drawn in Figure 6 represent data taken at nominally $1.75, 5.0, and 8.5 Â 1025 n mÀ2 (E > 0.1 MeV)
(assumed 1.75, 5.0, and 8.5 dpa). In the 1373–1473 K
temperature range, volumetric swelling is apparently
at a minimum, although it increases from $0.2%
to $0.4% to $0.7% for $1.75, 5.0, and 8.5 dpa,
respectively. Clearly, the swelling in this temperature
range has not saturated by 10 dpa. Above this minimum
in swelling, the data indicates a continual swelling
increase to the highest irradiation temperature
of $1773–1873 K. At $1773 K, measured swelling


220

Radiation Effects in SiC and SiC–SiC

Snead 2006
Snead 2006
Snead, unpub.

Price 1973
Blackstone 1971
Price 1969


Saturable regime
point defect swelling

Amorphization
regime

Price 1973,#2
Price 1973
Senor 2003
Nonsaturable regime
void swelling

20

10
7
5

Swelling (%)

3

8.5 dpa

2

5 dpa

1

0.7

1.75 dpa

0.5
0.3
0.2

0.1
0

200

400

600
800
1000
1200
Irradiation temperature (°C)

1400

1600

Figure 6 Irradiation-induced swelling of SiC to high irradiation temperatures. Reproduced from Snead, L. L.; Nozawa, T.;
Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

was $0.4, 1.0, and 2.0% for $1.75, 5.0, and 8.5 dpa,
respectively. It was also noted in the study by Snead

et al.33 that at $1773 K, surface reaction between SiC
and the graphite holder had taken place. However, a
loss of silicon from the surface cannot be ruled out.
Figure 6 includes historical data for swelling
above 1273 K.3,4,18,22,34,35 Specifically, Senor et al.18
report swelling for the same type of CVD SiC irradiated in this study when irradiated in a watermoderated fission reactor (the ATR) as well. Their
maximum dose, irradiation temperature, and swelling
data were $1 dpa, $1373 Æ 30 K, and 0.36 Æ 0.02%,
respectively. The irradiation temperature quoted
in Senor et al.’s work was a best estimate, although
the authors also provide an absolute bound of
1073–1473 K for their experiment. The maximum
swelling in their work (0.36 Æ 0.02% at $1 dpa) is
somewhat higher than the $0.25% swelling at
2 dpa, $1373 K, of the trend data in Figure 6. This
is seen from the rightmost figure of Figure 6. Also
seen in the figure is the high-temperature swelling of
Price.3,4,34 The Price data, which are in the dose range
of about 4–8 dpa, are in fair agreement with the

measured swelling of the Snead data16,33 of Figure 6.
The highest swelling material ($1523 K, $6 and
10 dpa) shows the largest discrepancy, although if the
temperature error bands quoted by the various authors
are taken into account, the data are conceivable more
in alignment. It is also noted that the Price material
may have had some excess silicon leading to higher
swelling as compared to stoichiometric material.
As mentioned earlier, the microstructural evolution of irradiated SiC is roughly understood, at least
for temperatures up to $1373 K. The swelling near

the critical amorphization temperature ($423 K) is
classically described as the differential strain between
the single interstitial, or tiny interstitial clusters,
immobile vacancies, and antisite defects. As the temperature increases above the critical amorphization
temperature, the number of defects surviving the
postcascade thermally activated recombination is
reduced and the mobility of both silicon and carbon
interstitials becomes significant. For temperatures
exceeding $1273 K, microstructural studies have
noted the presence of both Frank loops and tiny
voids, indicating limited mobility of vacancies.


Radiation Effects in SiC and SiC–SiC

The apparent increase in swelling with dose in the
1373–1873 K range seen in Figure 6 and the observed
production of voids are interesting considering that
the maximum irradiation temperature ($1773 K)
in Figure 6 is $0.65 of the melting (dissociation)
temperature (Tm) for SiC. Here, we have assumed
Olesinski and Abbaschian’s36 value of 2818 K where
stoichiometric SiC transforms into C þ liquid phase.
This value of 0.65Tm is high when viewed in comparison to fcc metal systems where void swelling typically begins at $0.35Tm, goes through a maximum
value, and decreases to nil swelling by $0.55Tm. (It is
noted that the melting and dissociation temperatures
of SiC are somewhat variable in the literature. However, even considering this variability, the previous
statement is accurate). If, as the swelling data seems
to indicate, the voids in SiC are continuing to grow in
SiC irradiated to 1773 K, the energies for diffusion of

either the Si or C vacancy or both must be quite high,
as are the binding energies for clustered vacancies.
This has been shown through theoretical work in
the literature.37–40 However, it is to be noted that
the defect energetics obtained from this body of
work, and in particular those of the Si and C vacancies
within SiC, vary widely. Perhaps, the work of
Bockstedte et al.,39 which follows an ab initio approach,
is the most accurate, yielding a ground state migration energy of 3.5 and 3.4 eV for Si and C vacancies,
respectively. It was also noted by Bockstedte et al.41
that the assumed charge state of the vacancy affects
the calculated migration energy. Specifically, the carbon vacancy in the þ1 and þ2 charge state increases
from 3.5 to 4.1 and 5.2 eV, respectively, and that
of silicon in the –1 and –2 charge state decreases
from 3.4 to 3.2 and 2.4 eV, respectively. Several
papers discuss the vacancy and vacancy cluster mobility measured experimentally. The silicon monovacancy has been shown to be mobile below 1273 K.
Using electron spin resonance, Itoh et al.30 found the
irradiation-produced T1 center in 3C–SiC disappearing above 1023 K. The T1 center was later confirmed to
be an Si vacancy.31 Using electron spin resonance, the
carbon vacancy in 6H–SiC is shown to anneal above
1673 K.42 Using isochronal annealing and positron
lifetime analysis, Lam et al.40 have shown a carbon–
silicon vacancy complex to dissociate above $1773 K
for the same 6H single crystal materials studied here.
From a practical nuclear application point of
view, the swelling data for CVD SiC can be broken
down into the amorphization regime (<423 K), the
saturable point-defect swelling regime (423–1073 K)
range, and the unsaturated void swelling regime,


221

which occurs for irradiation temperature >1273 K.
From the data of Figure 6, it is still unclear where
the actual transition into the unsaturated swelling
begins. Furthermore, while there is an increase in
swelling in the 1273–1773 K range, as the dose is
increased from $1.75, 5.0, and 8.5 Â 1025 n mÀ2
(E > 0.1 MeV), swelling is close to linear with neutron
doses, and it is unclear how swelling will increase
as a function of dose above 10 dpa. For example,
swelling by voids estimated from the TEM examination accounts for only relatively small fractions of the
total swelling even in the void swelling regime. Analogous to the typical swelling behavior in metals, void
growth may cause steady-state swelling after a certain
transition dose regime. However, dose dependence of
the swelling due to the nonvoid contribution remains
to be understood. Extrapolation of swelling outside of
the dose range of Figure 6 is therefore problematic.

4.07.3 Irradiation-Induced Thermal
Conductivity Degradation of
Monolithic SiC
According to Lee et al.,43 the effect of neutron irradiation on the specific heat of SiC was negligibly small.
The specific heat of SiC is therefore assumed to be
unchanged by neutron irradiation, although this has
not been verified at high dose. A single study5 also
indicated that stored energy (Wigner energy) occurs
in SiC irradiated in the point defect regime. The
relative amount of stored energy appears to be less
than that of graphite.44

Because of a low density of valence band electrons,
thermal conductivity of most ceramic materials, SiC
in particular, is based on phonon transport. For a
ceramic at the relatively high temperature associated
with nuclear applications, the conduction heat can be
generally described by the strength of the individual
contributors to phonon scattering: grain boundary
scattering (1/Kgb), phonon–phonon interaction (or
Umklapp scattering 1/Ku), and defect scattering
(1/Kd). Because scattering of each of these types
occurs at differing phonon frequencies and can be
considered separable, the summed thermal resistance for a material can be simply described as the
summation of the individual components; that is,
1/K ¼ 1/Kgb þ 1/Ku þ 1/Kd. As seen in Figure 7, the
unirradiated thermal conductivity of SiC is highly
dependent on the nature of the material (grain size,
impurities, etc.) and the temperature. While materials
can be optimized for low intrinsic defect, impurity,


222

Radiation Effects in SiC and SiC–SiC

Legend

500

Reference


N/R

Material

Note
Note
Single Crystal

Rohm and Haas Co.

CVD

Grain size ~5µm

Senor et al. (1996)10

CVD

Morton CVD

Graebner et al. (1998)48

CVD

Morton CVD

Pickering et al. (1990)49

CVD


Grain size ~10µm

Rohde (1991)45

Taylor et al. (1993)46
47

Highly pure and dense
single-/poly-crystals

400

CVD

Grain size ~3 µm

50

CVD

Grain size >10 µm

50

CVD

Grain size <10 µm

Price (1973)


FB

Grain size <5 µm

Price (1969)3

FB

Grain size <5 µm

Shaffer (1991)51

PS

Li et al. (1998)37

-

Thermal conductivity (W m–1 K–1)

Collins et al. (1990)
Collins et al. (1990)
34

300

Calculated value

PS: pressureless sintering, HP: hot-pressing
CVD: chemically vapor deposition, FB: fluidized bed

N/R: not reported
Poly-crystal, large grains

200

Kth = [–0.0003 + 1.05ϫ10–5 T]–1
Poly-crystal, small grains
Porous poly-crystal,
small grains

100

0
0

500

1000

1500

2000

Temperature (K)
Figure 7 Thermal conductivity of various forms of SiC as a function of temperature. Reproduced from Snead, L. L.;
Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

and grain boundary scattering, the temperaturedependent phonon scattering cannot be altered and
tends to dominate at high temperature (above about
673 K for SiC).

The effect of irradiation on SiC in the temperature range of $423–1073 K (the point defect regime)
is to produce simple defects and defect clusters that
very effectively scatter phonons. For ceramics possessing high thermal conductivity, the irradiationinduced defect scattering quickly dominates, with
saturation thermal conductivity typically achieved
by a few dpa. Moreover, as the irradiation-induced
defect scattering exceeds the phonon–phonon scattering, the temperature dependence of thermal conductivity is much reduced or effectively eliminated.
The rapid decrease as well as saturation in thermal
conductivity of CVD SiC upon irradiation in the
point-defect regime has been reported by several
authors.8,34,45,52,53 Figure 8 shows this rapid decrease
in thermal conductivity for fully dense CVD SiC,
including new data, previous data from the
authors,52,53 and that of Rohde.45 It is noted that the
data of Thorne is omitted as the material was of
exceptionally low density for a CVD SiC material.
Moreover, the data of Price34 is published with a
range of fluence that is not valuable in the figure.

In recent papers by Snead on the effects of neutron irradiation on the thermal conductivity of ceramics,53 and specifically on SiC,16,52 the degradation
in thermal conductivity has been analyzed in terms of
the added thermal resistance caused by the neutron
irradiation. The thermal defect resistance is defined
as the difference between the reciprocals of the
irradiated and nonirradiated thermal conductivity
(1/Krd ¼ 1/Kirr – 1/Knonirr). This term can be related
directly to the defect type and concentration present
in irradiated ceramics.53 Moreover, this term can be
used as a tool to compare the thermal conductivity
degradation under irradiation of various ceramics or,
for example, various forms of SiC. It has been shown

that, for certain high purity forms of alumina, the
accumulation of thermal defect resistance is very
similar even though the starting thermal conductivities of the materials are substantially different. Similarly, CVD SiC was shown to have a
similar accumulation of thermal defect resistance as
a hot-pressed form of SiC with substantially lower
($90 W mÀ1 KÀ1) unirradiated thermal conductivity.
The utility of this finding is that if the thermal defect
resistance is mapped as a function of irradiation temperature and dose for a form of high-purity CVD
SiC, it can be applied to determine the thermal


Radiation Effects in SiC and SiC–SiC

Room temperature thermal conductivity (W m−1 K–1)

300

223

Nonirradiated thermal conductivity 381 ± 26 W m−1 K−1
(except Rohde, 108 W m−1 K−1)

Tirr = 200 °C
Tirr = 800 °C

100

50
Tirr = 80 °C


30
Tirr = 750-850 °C
Tirr = 600 °C
Tirr = 480-550 °C
Tirr = 400-480 °C

10

Tirr = 375 °C
Tirr = 200 °C

CVD SiC
5
0.001

0.01

0.1

1

10

Fast neutron dose (´ 1025 n m-2 E > 0.1 MeV)
Figure 8 Degradation in room-temperature thermal conductivity for CVD SiC. (Rohde data designated as Â.) Reproduced
from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

conductivity of any high-purity CVD SiC, independent of the starting thermal conductivity. The accumulation in thermal defect resistance generated from
the data of Figure 8 is shown in Figure 9.
Another result of the previously reported analysis

on irradiated CVD SiC16,52 is that the thermal defect
resistance appears to be directly proportional to the
irradiation-induced swelling, although the data-set
for making the previous assertion was somewhat limited. A compilation plot including the previous dataset as well as the new data of Figure 9 is shown in
Figure 10. It is clear from this plot that a linear
relationship exists between swelling and thermal
defect resistance. Moreover, there does not appear
to be any effect of irradiation temperature on this
result. The fact that the thermal defect resistance
is proportional to the irradiation-induced swelling
allows a rough estimate of thermal conductivity.
As measurement of thermal conductivity for the
SiC TRISO shell is not practical, while measurement
of density is routine, this finding allows an indirect
determination of thermal conductivity by measurement of the density change in the TRISO SiC shell
by means of a density gradient column or some other
technique.

The thermal conductivity degradation discussed
up to this point has been for irradiation temperature
associated with the point defect regime. For irradiation above this temperature (the nonsaturable
void swelling regime), the thermal properties are
not expected to saturate (at least at low dpa). The
primary reason for this is that the formation of voids
and other complex defects in the high-temperature
regime (which contributes to the unsaturated swelling
as seen in Figure 6) contributes to phonon scattering,
and these defects will not saturate. Moreover, it has
been shown that the linear relationship that existed
between swelling and thermal defect resistance (as

seen in Figure 10) does not exist in this elevated
temperature irradiation regime.16,52 This underlines
the fact that the phonon scattering and swelling
are not controlled by the same defects in the lower
temperature ‘saturable,’ and elevated temperature
‘nonsaturable’ irradiation regimes. A compilation
plot of room-temperature thermal conductivity as a
function of irradiation temperature for the saturable
and nonsaturable temperature regimes is given in
Figure 11.
By comparison to the unirradiated roomtemperature conductivity value of $280 W mÀ1 KÀ1,


224

Radiation Effects in SiC and SiC–SiC

Thermal defect resistance (m K W–1)

Tirr = 80 °C

Tirr = 200 °C
Tirr = 375 °C

0.1

Tirr = 400-480 °C
Tirr = 480-550 °C
Tirr = 600 °C


Rhode, Tirr = 80 °C
Tirr = 750-850 °C

0.01

Tirr = 80 °C
Tirr = 200 °C
Tirr = 375 °C
Tirr = 400-480 °C
Tirr = 480-550 °C
Tirr = 600 °C
Tirr = 750-800 °C

0.001
0.001

0.01

0.1

1

10

Fast neutron dose (´ 1025 n m-2 E > 0.1 MeV)
Figure 9 Thermal defect resistance for stoichiometric CVD SiC as a function of neutron dose. (Rohde data designated
as Â.) Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371,
329–377.

it is clear that the thermal conductivity degradation

in the highest temperature regimes is less dramatic,
even though the swelling is rapidly increasing (see
Figure 6). This is opposite to the behavior in the
lower temperature, saturable regime, where high
swelling corresponds to extreme reduction in thermal
conductivity. Unfortunately, the data on thermal conductivity reduction in the nonsaturable regime is
limited, and given the lack of knowledge of the specific defects governing the phonon scattering, it is not
possible to accurately predict behavior outside of the
data-set of Figure 6.
Data presented thus far has been limited to measurement of thermal conductivity at room temperature. As described in Figure 7, there is a dramatic
dependence of thermal conductivity on measurement temperature. The temperature dependence
of irradiated materials can be found by applying
the temperature dependence of unirradiated SiC
(the Umklapp thermal resistance term) to the
as-neutron-degraded room-temperature values. This
approximation (dashed lines) is compared to actual

data (solid lines) in Figure 12 and shows fair correspondence. However, it is clear that such a treatment
systematically underestimated the thermal conductivity degradation. This implies temperature dependence on the defect scattering that is not presently
understood.

4.07.4 Effect of Irradiation on
the Mechanical Properties of
Monolithic SiC
4.07.4.1

Elastic Modulus of Monolithic SiC

Figure 13 summarizes the irradiation temperature
dependence of the elastic modulus change. Irradiation generally reduces modulus to a greater extent for

lower temperature irradiation. The modulus reduction
becomes negligible when irradiation temperature
reaches or exceeds $1273 K. There seems to be an
indistinct stage between 1073 and 1273 K. As expected,
the elastic modulus trends with ‘point defect swelling’
of SiC. Point defect swelling is an isotropic volume


Radiation Effects in SiC and SiC–SiC

Tirr = 200 °C

Tirr = 400-480 °C

Tirr = 300 °C

Tirr = 480-550 °C X Tirr = 80 °C

Tirr = 375 °C

Tirr = 600 °C

225

Tirr = 750-850 °C

0.12

X


100

0.1
Open symbols
Right
0.08

0.06

0.04

Thermal defect resistance (m K W-1)

Room temperature thermal conductivity (W m-1 K–1)

0.14

10
0.02

Closed symbols
Left
X

0
0

0.5

1


1.5

2

2.5

Swelling (%)
Figure 10 The room-temperature thermal conductivity and thermal defect resistance as a function of irradiation-induced
density change. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007,
371, 329–377.

expansion that is believed to occur by lattice relaxation
due to accumulated isolated point defects and small
point defect clusters during irradiation at temperatures where vacancies are not readily mobile. In SiC
irradiated in the point defect swelling regime, a fairly
good agreement between dimensional expansion and
lattice spacing has been confirmed by X-ray diffractometry studies. In contrast, the data in the nonsaturable swelling regime is somewhat limited, although the
data suggest that there is little reduction in elastic
modulus in spite of the swelling being relatively
large. However, in this regime, the defects responsible
for swelling are voids and other relatively large defects,
which would have less of an effect on elastic modulus
as compared to point defects.
An estimation of the influence of lattice relaxation on elastic modulus was attempted using the
Tersoff potential.54 The result predicted a linear lattice

swelling of 1% causing approximately 10% reduction
in elastic modulus (Figure 14). The predicted elastic
modulus change could be varied by more than 10%

depending on a selection of interatomic potential, with
the Tersoff potential giving a relatively high sensitivity
of modulus to the interatomic distance among various
empirical interatomic potential functions. Therefore,
the measured elastic modulus changes observed in
this experiment are generally greater than the theoretical prediction. It is noted that the methods applied
for generating the data of Figure 14 are various and
of differing quality. Typically, elastic modulus as
measured by nanoindentation, which sometimes is the
only alternative for miniature specimens, tends to give
widely scattered and less reliable data than the mechanical or sonic modulus methods. Nonetheless, it is
clear that the lattice expansion is a major cause of the
irradiation-induced elastic modulus reduction in SiC.


226

Radiation Effects in SiC and SiC–SiC

Amorphization
regime

Saturable regime

Unsaturable regime

120
Nonirradiated conductivity ~280 W m-1 K-1
Rohm and Haas CVD SiC


Room temperature thermal
conductivity (W m-1 K-1)

100
~1.75 dpa
80

~5 dpa

60

~8.5 dpa
40

20

0
0

200

400

600
800
1000
Irradiation temperature (°C)

1200


1400

1600

Figure 11 Room-temperature thermal conductivity in the saturable and nonsaturable regime. Reproduced from Snead, L.
L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

4.07.4.2

Hardness of Monolithic SiC

The irradiation effect on nanoindentation hardness
of Rohm and Haas CVD SiC in a fluence range of
0.1–18.7 dpa is summarized in Figure 15. It is interesting to note that the nanoindentation hardness exhibits
relatively small scatter for the individual experiments,
and the trend in data as a function of temperature is
uniform. This observation is in contrast to both the
flexural strength and the indentation fracture toughness data, which indicate a broad peak at an intermediate temperature and a relatively large scatter. It is
worth noting that nanoindentation hardness of brittle
ceramics is, in general, determined primarily by the
dynamic crack extension resistance in the near surface
bulk material, and therefore should be more relevant to
fracture toughness than to plastic deformation resistance. However, surface effects of the original sample
affect the nanoindentation hardness less, as the samples
are generally polished prior to testing.
4.07.4.3 Fracture Toughness of
Monolithic SiC
The effect of irradiation on the fracture toughness of
Rohm and Haas CVD SiC is summarized in Figure 16.
This compilation plots data using the Chevron

notched beam technique, although the bulk of the

data sets report Vicker’s or nanoindentation generated data.55–57 The general trend is that the
irradiation-induced toughening seems to be significant at 573–1273 K for the indentation fracture toughness data, in spite of the decrease in elastic modulus,
which confirms the increase in fracture energy caused
by irradiation. The scatter of the indentation fracture
toughness data among different experiments is likely
caused by both the condition of the sample surface and
the lack of standardized experimental procedures. Typically, indentation should be applied on the polished
surfaces, but conditions of polishing are not always
provided in literature. Moreover, the crack length measurements are done using optical microscopy, conventional scanning electron microscopy (SEM), or field
emission SEM, all of which may give very different
crack visibility. In addition, a few different models
have been used for derivation of the fracture toughness.
In conclusion, indentation fracture toughness techniques can be used only for qualitative comparison within
a consistent set of experiments. It is noted that the
experiment employing the Chevron notched beam
technique also indicates the irradiation-induced toughening, although scatters of toughness values were even
greater. These results lead to the conclusion that,
in the intermediate irradiation temperature range, the
increase of the fracture toughness of SiC exists.


Radiation Effects in SiC and SiC–SiC

227

120
500


Rohm and Haas CVD SiC irradiated
in HFIR

Rohm and Haas CVD SiC irradiated in HFIR at 800 ˚C

400
300

100
Tirr = 1500 ЊC

100

Thermal conductivity (W m-1 K–1)

Thermal conductivity (W m-1 K–1)

200

Nonirradiated

80
60
50

0.05 dpa
++

40


+ +
+

+

30

+

0.5 dpa
+

1.94 dpa
4.3 dpa

20
Calculated based on
Umklapp scattering

1.75 dpa
80

60

Tirr = 1500 ЊC
7.5 dpa

40

20


Calculation based on
Umklapp scattering

Fit to data

Fit to data

10
0

(a)

200
400
600
800
1000
Irradiation and measurement temperature ( ˚C)

0
0

400
500
600
700
100
200
300

Irradiation and measurement temperature ( ЊC)

(b)

800

40
Rohm and Haas CVD SiC irradiated in HFIR
35

Thermal conductivity (W m-1 K–1)

T = 1020 ЊC
irr
1.75 dpa

30

+
Tirr = 1050 ЊC

25

5 dpa

+

+
+ +


+

1.75 dpa

+

+

5 dpa

+ + +

Tirr = 1060 ЊC

20

+

7.5 dpa

8.5 dpa

15

10

Calculation based on
Umklapp scattering
Fit to data


5

0
(c)

0

100 200 300 400 500 600 700 800
Irradiation and measurement temperature ( ˚C)

Figure 12 Effect of temperature on the conductivity of irradiated SiC. (a) Tirr ¼ 1073 K, (b) Tirr ¼ 1773 K, and
(c) Tirr ¼ 1293–1333 K. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater.
2007, 371, 329–377.

4.07.4.4 Strength and Statistical Variation
in Strength for Monolithic SiC
There have been several studies on the effect of neutron irradiation on the strength of various types of SiC
forms including reaction-bonded, sintered, pressureless sintered, and CVD SiC materials.1,11,13,14,58,60–65

The strength of SiC depends significantly on stoichiometry under neutron irradiation. Both the sintered
SiC and the reaction-bonded SiC forms exhibit
significant deterioration in strength by neutron irradiation (Figure 17).13 The presence of impurities
such as sintering additives for sintered SiC and excess


228

Radiation Effects in SiC and SiC–SiC

1.20


Relative Young’s modulus

1.10

1.00

0.90

0.80

0.70

0.60
200

Osborne (1999)55,HFIR, 2 dpa
Nogami (2002)56,HFIR + HFBR, 0.15–7.7 dpa
Park (2003)57,DuET 5.1 MeV Si, 3 dpa
Katoh (2005)58,HFIR – 14J, 6.0–7.7 dpa
Snead (2007)16, HFIR – METS, 0.7–8.6 dpa
Price(1977)59,ETR, 2.8–12.2 dpa
Snead (2007)16, HFIR, 0.7–4.2 dpa

400

600

Nanoindentation


4pt. bend
Sonic resonance

800 1000 1200 1400
Irradiation temperature (K)

1600

1800

2000

Figure 13 Irradiation temperature dependence of irradiated elastic modulus of CVD SiC, at ambient temperature,
normalized to unirradiated values. The error bars are showing standard deviations for all the neutron data points and ranges
of data scatter for the ion data points. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A.
J. Nucl. Mater. 2007, 371, 329–377.

1.20
Nanoindentation

Relative Young’s modulus

1.10

4pt. bend
Sonic resonance

Osborne (1999)55,HFIR, 2 dpa
Nogami (2002)56,HFIR + HFBR, 0.15–7.7 dpa
Park (2003)57,DuET 5.1MeV Si, 3 dpa

Katoh (2005)58,HFIR - 14J, 6.0–7.7 dpa
Snead (2007)16, HFIR - METS, 1.7–8.6 dpa
Price (1977)59, ETR, 2.8–12.2 dpa
Snead (2007)16, HFIR, 0.7–4.2 dpa
Model (Tersoff potential)

1.00

0.90

0.80

0.70
0.0

0.5

1.0

1.5
2.0
2.5
Volumetric swelling (%)

3.0

3.5

Figure 14 Irradiation-induced change of elastic modulus versus swelling of CVD SiC. An estimation of the influence of
lattice relaxation on elastic modulus is calculated using Tersoff potential. Reproduced from Snead, L. L.; Nozawa, T.;

Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.


Radiation Effects in SiC and SiC–SiC

229

1.6

Irradiated nano-indentation hardness
normalized to unirradiated values

1.4
1.2
1
0.8
0.6
0.4
0.2
0
200

Katoh (2005)58, HFIR-14J, 6.0–7.7 dpa
Nogami (2002)56, HFIR+HFBR, 0.2–7.7 dpa
Osborne (1999)55, HFIR, 0.1–5.0 dpa
Park (2003)57, DuET 5.1 MeV Si, 3 dpa

400

600


800

1000

1200

1400

1600

Irradiation temperature (K)
Figure 15 Nanoindentation hardness of CVD SiC at ambient temperature as a function of irradiation temperature.
Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

2
1.8

Irradiated fracture toughness
normalized to unirradiated values

1.6
1.4
1.2
1
0.8
0.6
0.4
0.2
0

200

Nogami (2002)56, HFIR+HFBR, 0.2–7.7 dpa
Osborne (1999)55, HFIR, 0.1–5.0 dpa
Park (2003)57, DuET 5.1 MeVSi, 3 dpa
Snead (2007)16, HFIR, 0.7–4.2 dpa

400

600

800

1000

1200

1400

1600

1800

Irradiation temperature (K)
Figure 16 Indentation fracture toughness of CVD SiC at ambient temperature as a function of irradiation temperature.
Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

Si for reaction-bonded SiC, which typically segregate
to grain boundaries during sintering, tends to have
a significant influence on strength under neutron

irradiation. For the case of sintered SiC with boron

compounds as sintering additives, the reaction of
B(n, a)7Li causes the accumulation of helium bubbles at and near the grain boundary phases under
neutron irradiation.60–63 In contrast, unmatched

10


230

Radiation Effects in SiC and SiC–SiC

1.5

Normalized strength

Hot pressed and sintered SiC forms

1.0

Matheney (1979)64
Matthews (1974)11
Iseki (1990)19 RB(reaction bonded)
Iseki (1990)19 PLS(pressureless sintered)

0.50

Iseki (1990)19 HP(hot pressed)
Price (1982)63 NC-430

Price (1982)63 Carborundum
Dienst (1991)62 0.75% B
Dienst (1991)62 1.9% B
Corelli (1983)60 NC430

0.0

0

0.1

1
Dose (dpa)

10

100

Figure 17 Fluence dependence of irradiated flexural strength of hot-pressed and sintered SiC normalized to unirradiated
strength. Irradiation is variable but in the saturable swelling regime. Reproduced from Newsome, G. A.; Snead, L. L.;
Hinoki, T.; Katoh, Y.; Peters, D. J. Nucl. Mater. 2007, 371, 76–89.

2.0
CVD silicon carbide

Normalized strength

1.5

1.0


Snead
Price
Dienst
Newsome 300 ЊC
Newsome 500 ЊC
Newsome 800 ЊC

0.50

0.0
0.1

1

10

100

Dose (dpa)
Figure 18 Flexural strength of CVD SiC at ambient temperature as a function of irradiation dose. Reproduced from
Newsome, G. A.; Snead, L. L.; Hinoki, T.; Katoh, Y.; Peters, D. J. Nucl. Mater. 2007, 371, 76–89.

swelling between Si and SiC for reaction-bonded SiC
causes disruption at the grain boundary, severely
reducing the strength.11,19,60–64 Meanwhile, the highpurity materials such as CVD SiC exhibit superior
irradiation resistance.
The irradiation effect on flexural strength of Rohm
and Haas CVD SiC in a fluence range of 0.15–30 dpa
is summarized in Figure 18. In comparing Figure 18


with Figure 17, it is clear that CVD SiC retains
stability in strength to a much higher dose than the
sintered and reaction-bonded forms of SiC. It is to be
noted that in Figure 18, the data of Dienst does
indicate a significant as-irradiated degradation in
strength around 15 dpa. However, such degradation
is not seen for the $30 dpa irradiation of Snead. It is
speculated that the degradation in the Dienst data may


Radiation Effects in SiC and SiC–SiC

have been due to statistical limitations of the study
and/or due to issues with sample handling postirradiation. This issue is discussed in the Dienst reference.65
A compilation of strength data as a function of irradiation temperature is given in Figure 18, indicating no
apparent correlation for the dose and temperature
ranges studied. However, as with the fracture toughness data, irradiation-induced strengthening seems to
be significant at 573–1273 K. The large scatter in flexural strength of brittle ceramics is inevitable, as the
fracture strength is determined by the effective fracture toughness, morphology, and characteristics of the
flaw that caused the fracture. Irradiation possibly
modifies both the flaw characteristics and the fracture
toughness through potential surface modification,
relaxation of the machining-induced local stress,
modifications of elastic properties, and fracture
energy.
A typical means of describing the failure of ceramics is through the use of Weibull statistics, which is
a departure from the analysis of data that is assumed
to follow a normal Gaussian distribution. In the twoparameter Weibull formalism, sometimes referred to
as a weakest-link treatment, the failure probability F

is described as
F ðxÞ ¼ 1 À e Àðx

m

=x0 Þ

where m is the Weibull modulus and xo is the
distribution size parameter. A change in the Weibull
statistics, indicating a higher scatter in as-irradiated
flexural strength has been observed by previous
authors, although the point could not be made
convincingly because of limitations in the number
of tests observed. In the earliest work known to the
authors, Sheldon66 noted a 14% decrease in crushing
strength of highly irradiated CVD SiC shells with
an increase of the coefficient of variation from 8%
to 14%. Price63 went on to a 4-point bend test
using relatively thin ($0.6 mm) strips of CVD SiC
deposited onto a graphite substrate. In his work,
the flexural strength following a $9.4 Â 1025 n mÀ2
(E > 0.1 MeV) irradiation was unchanged within the
statistical scatter, but the scatter itself increased from
about 10 to 30% of the mean flexural strength as
described assuming a normal distribution. Unfortunately, there were not sufficient samples in Price’s
work to infer Weibull parameters. In more recent
work by Dienst,65 the Weibull modulus was reported
to decrease from about 10 for irradiation of $1 Â 1026
n mÀ2 (E > 0.1 MeV). However, it is worth noting
that the Dienst work used a very limited sample

population (about 10 bars.)

231

Statistically meaningful data sets on the effect
of flexural strength of CVD SiC have been reported by Newsome and coworkers14 and Katoh and
coworkers.58,67 Figure 19 shows a compilation Weibull
plot of the flexural strength of unirradiated and irradiated Rohm and Haas CVD SiC taken from the two
separate irradiation experiments carried out by Newsome and cowokers14 and Katoh and coworkers.58,67
The sample population was in excess of 30 for each
case. In Figure 19(a), the data was arranged by irradiation temperature, including data for unirradiated
samples and 1.5–4.6 Â 1026 n mÀ2 (E > 0.1 MeV) dose
range. It is likely that the Weibull modulus decreased
by irradiation, appearing to be dependent on irradiation temperature. This is not easily visualized through
inspection of Figure 19(a) unless one notes that there
are significantly more low stress fractures populating
the 573 K population. The scale parameters of flexural
strength of unirradiated materials and materials irradiated at 573, 773, and 1073 K were 450, 618, 578, and
592 MPa, respectively. The Weibull modulus of the
flexural strength of unirradiated materials and materials irradiated at 573, 773, and 1073 K were 9.6, 6.2, 5.5,
and 8.7, respectively, with significant uncertainty.
The work of Katoh, on identical material irradiated at the same temperature as in the Newsome
work, is at a slightly higher irradiation dose than the
data of Newsome. As seen in Figure 19(b), the effect
on the Weibull modulus undergoes a trend similar
to that of Newsome, although the modulus for the
773 K and 1073 K irradiation of Katoh remained
almost unchanged. Given the data discussed on the
effect of irradiation on the Weibull modulus and scale
parameter of CVD SiC bend bars, it is clear that

the material is somewhat strengthened and that the
Weibull modulus may undergo irradiation-induced
change, with the greatest decrease occurring for the
lowest temperature irradiation.
The fracture strength and failure statistics of
tubular SiC ‘TRISO surrogates’ have been determined by the internal pressurization test and the
results are plotted in Figure 20. Thin-walled tubular
SiC specimens of 1.22 mm outer diameter, 0.1 mm
wall thickness, and 5.8 mm length were produced by
the fluidized-bed technique alongside TRISO fuels.68
The specimens were irradiated in the HFIR to 1.9 and
4.2 Â 1025 n mÀ2 (E > 0.1 MeV) at 1293 and 1553 K.
In the internal pressurization test, tensile hoop stress
was induced in the wall of the tubular specimens by
compressively loading a polyurethane insert.68,69
In Figure 20, Weibull plots of the flexural
strength and internal pressurization fracture strength


232

Radiation Effects in SiC and SiC–SiC
si (MPa)
200

3

300

400


500

600

800

1000

500 ºC, 2.0 dpa
m = 5.5

2

Nonirrad.
m = 9.6

1

ln(ln(1/(1–Fi)))

0
-1
-2
800 ºC, 2.0 dpa
m = 8.7

-3
-4
-5


300 ºC, 2.0 dpa
m = 6.2

-6
5.0

5.5

6.0

(a)

6.5

7.0

ln(si)
si (MPa)
3

200

300

400

600

800


1000

300 ºC, 6.0 dpa
m = 5.5

2
1

500

Nonirrad.
m = 9.9

ln(ln(1/(1–Fi)))

0

-1
500 ºC, 6.0 dpa
m = 10.8

-2
-3
-4

800 ºC, 7.7 dpa
m = 7.9

-5

-6
5.0

5.5

(b)

6.0

6.5

7.0

ln(si)

Figure 19 Weibull plots of flexural strength of unirradiated and irradiated CVD SiC in the dose range of
(a) 1.5–4.6 Â 1025 n mÀ2 (E > 0.1 MeV) from Newsome14 and (b) 7.7 Â 1025 n cmÀ2 (E > 0.1 MeV) from Katoh.58

of unirradiated and irradiated samples are presented.
As with the Newsome and Katoh data, the sample
population is large enough to be considered statistically
meaningful. From this data, the mean fracture stress
of tubular specimens is seen to increase to 337 MPa
(from 297 MPa) and the Weibull modulus slightly
decreased to 3.9 (from 6.9) after irradiation to
1.9 Â 1025 n mÀ2 (E > 0.1 MeV) dpa at 1293 K. The
mean fracture stresses and Weibull moduli at
4.2 Â 1025 n mÀ2 (E > 0.1 MeV) were similar to those
at 1.9 dpa. The results for 4.2 dpa irradiation indicate


that by increasing the irradiation temperature from
1293 to 1553 K, no discernible change in fracture
stress distribution occurred. The horizontal shift
indicates a simple toughening or an increase in fracture toughness alone. While the data for these surrogate TRISO samples, irradiated through internal
compression, are somewhat limited, the findings indicate that the trend in strength and statistics of failure
are consistent with those found for the bend bars.
Therefore, the general findings of the bend bar irradiation on strength and Weibull modulus appear


Radiation Effects in SiC and SiC–SiC

3
2

200

si (MPa)
400
500

300

800

1000

Nonirrad.
m = 7.6

1


ln(ln(1/(1–Fi)))

600

233

1280 ЊC, 4.2 dpa
m = 3.8

0
-1
-2

1020 ЊC, 4.2 dpa
m = 5.4

-3
1020 ЊC, 1.9 dpa
m = 4.4

-4
-5
5.0

5.5

6.0
ln(si)


6.5

7.0

Figure 20 Weibull statistical fracture strength of CVD SiC measured by the internal pressurization test. Reproduced
from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.

appropriate for application to TRISO fuel modeling. Specifically, a slight increase in the mean
strength is expected (although it may be less significant at higher temperatures), and the statistical
spread of the fracture data as described by the
Weibull modulus may broaden. Unfortunately, a
precise description of how the Weibull modulus
trends with irradiation dose and temperature is not
yet possible, although within the dose range and
temperature covered by the data in Figures 19 and
20, a modest reduction is possible.

4.07.5 Irradiation Creep of SiC
Irradiation creep is defined as the difference in
dimensional changes between a stressed and an
unstressed sample irradiated under identical conditions. Irradiation creep is important for structural
materials for nuclear services as it is a major
contributor to the dimensional instability of irradiated materials at temperatures where thermal
creep is negligible. However, studies on irradiation
creep of SiC(-based materials) have so far been very
limited, although it is of high importance for the
behavior of the SiC TRISO shell.

Price published the result of the irradiation
creep study on CVD SiC in 1977.59 In this work,

elastically bent strip samples of CVD SiC were irradiated in a fission reactor, and the steady-state
creep compliance was estimated to be in the order
of 10–38 (Pa dpa mÀ2 (E > 0.18 MeV))À1 at 1053–
1403 K. Scholz and coworkers measured the transient
creep deformation of SCS-6 CVD SiC-based fiber,
which was torsionally loaded under penetrating proton or deuteron beam irradiation.70–73 They reported
several important observations including the linear
stress and flux dependency of the tangential primary
creep rate at 873 K, and the negative temperature
dependence of primary creep strain at the same
dose. Recently, Katoh et al. determined the bend stress
relaxation (BSR) creep in Rohm and Haas CVD SiC
and Hoya monocrystalline 3C–SiC during irradiation in HFIR and JMTR at 673–1353 K.74 The results
reported for CVD SiC are summarized in Table 1.
In the BSR irradiation creep experiment by Katoh
et al., the creep strain for CVD SiC exhibited a weak
temperature dependence at <0.7 dpa in the $673–$
1303 K temperature range, whereas a major transition
at higher doses likely exists between $1223 and
$1353 K. Below $1223 K, the creep strain appeared
highly nonlinear with neutron fluence because of the


234
Table 1
Tirr ( C)

CVD SiC
400
600

600
640
700
750
1030
1080
3C–SiC
640
700
1030
1080

Radiation Effects in SiC and SiC–SiC

Irradiation creep data for CVD SiC from bend stress relaxation experiments
Fluence
(dpa)

Reactor

Initial/final
bend stress
(MPa)

Initial/final bend
strain (Â10À4)

Creep strain
(Â10À4)


BSR ratio m

Average creep
compliance Â10À6
(MPa dpa)À1

0.6
0.2
0.6
3.7
0.7
0.6
0.7
4.2

JMTR
JMTR
JMTR
HFIR
HFIR
JMTR
HFIR
HFIR

82/60
81/57
81/46
87/36
102/72
80/55

85/61
101/8

1.80/1.39
1.80/1.31
1.80/1.05
1.95/0.83
2.27/1.64
1.80/1.27
1.94/1.42
2.29/0.19

0.41
0.49
0.75
1.12
0.63
0.53
0.52
2.10

0.77
0.73
0.58
0.42
0.72
0.71
0.73
0.08


0.97
3.5
2.0
0.50
1.1
1.3
0.97
0.91

3.7
0.7
0.7
4.2

HFIR
HFIR
HFIR
HFIR

87/30
102/90
86/57
101/1

1.94/0.68
2.27/2.06
1.94/1.31
2.29/0.02

1.26

0.21
0.63
2.27

0.35
0.87
0.67
0.01

0.59
0.34
1.2
1.1

early domination of the transient irradiation creep.
The transient creep is speculatively caused by the
rapid development of defect clusters and the structural
relaxation of as-grown defects during early stages of
irradiation damage accumulation. At $1353 K, irradiation creep mechanisms, which are common to metals,
are likely operating.
In metals, steady-state irradiation creep rates
are generally proportional to the applied stress and
neutron (or other projectiles) flux, f,75,76 and therefore, irradiation creep compliance, B, has been conveniently introduced75:
eic ¼ sðBf þ DSÞ
where S is void swelling and D is a coefficient of
swelling–creep coupling. Ignoring the swelling–
creep coupling term (valid in the saturable swelling
regime), preliminary estimations of the steady-state
irradiation creep compliance of CVD SiC were
given to be 2.7 Æ 2.6 Â 10–7 and 1.5 Æ 0.8 Â 10–6 (MPa

dpa)À1 at $873–$1223 K and $1353 K, respectively.
If linear-averaged, creep compliances of 1–2 Â 10–6
(MPa dpa)À1 were obtained for doses of 0.6–0.7 dpa at
all temperatures. Monocrystalline 3C–SiC samples
exhibited a significantly smaller transient creep strain
by 0.7 dpa and a greater subsequent deformation
when loaded along <011> direction.
To better define the irradiation creep behavior
of SiC and the underlying physical mechanisms,
it will be essential to further examine the stress dependence, dose dependence, effect of crystallographic
orientation, microstructures of the crept samples, and
the coupling between irradiation creep and swelling.

4.07.6 Silicon Carbide Composites
Under Irradiation
SiC composites are a family of materials of varied
constituents and architectures. Up to the point of
writing this chapter, nuclear-grade SiC composites
(those specifically developed for application in fast
neutron environments and exhibiting neutron irradiation damage resistance) are more precisely defined
as continuous fiber-reinforced ceramic composites.
The history of development for these materials has
been reviewed in a number of publications.29,77–79
The primary constituents of these nuclear-grade composites are the continuous SiC fiber, a fiber/matrix
interphase material that can be SiC or pyrolytic graphite or a combination of the two, and a matrix of SiC
infiltrated into the woven fiber preform. The most
common matrix material is derived from chemical
vapor infiltration (CVI), and is essentially identical in
structure, properties, and irradiation response to the
CVD SiC discussed in previous sections. While there

has been little direct study on the effects of irradiation
on the material properties of the SiC interphase, it can
be assumed that it would also behave in a similar
manner to the SiC matrix. However, the effect of
neutron irradiation on pyrolytic graphite interphase
(if used) will be substantially different from that on
both matrix and fiber. While the effect of irradiation
on the underlying properties of graphite interphase
has not been well studied, it can be assumed that the
interphase will behave in a similar manner to nuclear
graphite (discussed in Chapter 4.05, Radiation Damage of Reactor Pressure Vessel Steels).


Radiation Effects in SiC and SiC–SiC

Y

(a)

235

(b)
X

Z
Y
(c)

300 mm


SiC matrix

SiC/PyC
multilayer

PyC
1 mm

Hi-Nicalon type-S fiber

Figure 21 Example of braided nuclear-grade SiC/SiC composite. Fiber: Hi-Nicalon™ Type-S; Interphase: Multilayer
SiC with pyrolytic carbon; Matrix: CVI SiC deposited through an isothermal process. Reproduced from Nozawa, T.;
Lara-Curzio, E.; Katoh, Y.; Shinavski, R. J. Tensile properties of advanced SiC/SiC composites for nuclear control rod
applications. Wiley: 2007; pp 223–234.

An example of an SiC/SiC composite that has been
developed for high-temperature gas-cooled reactor
control rod applications is shown in Figure 21. The
basic textile weaving of the composite is evident on
inspection of Figure 21(a). In this case, a Æ55 weave
is depicted. For the polished section of Figure 21(b),
large voids, which are an unavoidable characteristic
of chemical vapor infiltrated materials and also the
primary reason why it is difficult to produce gasimpermeable SiC/SiC composite, are clearly observed.
In Figure 21(c), the complicated structure of the
interphase is seen. In this case, alternating layers of
SiC and pyrolytic graphite have been applied. The
pyrolytic graphite layer between the SiC layers is
quite thin (tens of nanometer), with a relatively
thick graphite layer in contact with the fiber itself.

From the earliest study of SiC/SiC composites
under irradiation, it was clear that the fiber was the
key to performance. As with the impure forms of
SiC monolithic ceramics (cf. Figure 17), the impure
and oxygen-rich early grades of SiC fiber (trade name
Nicalon™) were quite unstable under neutron irradiation.12,80,81 Researchers were able to directly link
an irradiation-induced shrinkage of the SiC-based

fibers with debonding of the fiber–matrix interface
that severely compromised the ability to load the
high-strength fibers.80 Composite mechanical properties such as strength suffered appreciably.
With continued evolution of the fiber systems
to increasingly pure, stoichiometric materials, the
irradiation stability improved significantly. Presently,
there are two commercial fiber systems used
in nuclear-grade composites, both of which have relatively low impurity contents and are approaching a
1:1 stoichiometry. Specifically, the $11 micron
Hi-Nicalon™ Type-S fiber has the nominal chemistry of SiC1.05, 0.2%-O, while the $7.5 and $10 mm
Tyranno™ SA-3 fibers have the nominal chemistry
of SiC1.07, 0.5% Al. Study has revealed that these
‘near stoichiometric’ fibers exhibit irradiationinduced swelling similar to that of CVD,82 thus
avoiding the debonding phenomenon mentioned in
the previous paragraph. For this reason, composites
fabricated from these materials are superior under
irradiation to their predecessors. Consistent with
the discussion of properties of irradiated monolithic
SiC, the following discussion will be limited to the
more pure, near stoichiometric fiber materials.



236

Radiation Effects in SiC and SiC–SiC

Weibull mean strength (GPa)

4

280

3

300

470

800

500

470
770

2
280

950
950

1


Offset for nonirradiated
Hi-Nicalon™ type-S, Katoh (2010)67
Tyranno™-SA3, Katoh (2010)67
Hi-Nicalon™ type-S, Nozawa (2004)83

0
0

2

4
6
Neutron dose (dpa)

8

10

Figure 22 Effect of neutron irradiation on fiber strength. Data labels indicate the nominal irradiation temperature (in  C) for
Hi-Nicalon™ Type-S (upright) and Tyranno™ SA-3 (oblique) fibers. Reproduced from Katoh, Y.; Snead, L. L.; Nozawa, T.;
Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48–61.

The effect of neutron irradiation on the Weibull
mean strength of individual ‘near stoichiometric’
fibers is given in Figure 22.83,84 Within inherent statistical scatter, no change in strength is observed for
either the Hi-Nicalon™ Type-S or the Tyranno™
SA-3 bare fibers. The numbers inset to the figure
indicate the irradiation temperature of the SiC fibers,
with no apparent function of irradiation temperature

on strength observed. From the same study, the effect
of irradiation on composite properties is also
observed. Figure 2367 gives the proportional limit
stress for which the load departs from elastic behavior
and the ultimate tensile strength. As with the fiber
data, and the data for monolithic CVD SiC (Figure 18),
the composite strength does not exhibit any statistically
meaningful change. Supporting studies14,82,83,85–87 on
the strength in tension or bending of neutron-irradiated
stoichiometric fiber composites support the fact
that at least up to $40 dpa, composite strength is
not significantly affected by irradiation. A recent
study88 on the fracture toughness of irradiated and
unirradiated Hi-Nicalon™ Type-S composites also
reports no appreciable change. However, a minor difference in the fracture surface (length of fiber pull out)
and a trend in the fiber–matrix interphase properties
are reported,89 suggesting that mechanical property
evolution may occur at higher doses.

In the unirradiated state, the thermal conductivity
of SiC composites is dependent on variables including the fibers and matrix constituents, processing,
and the level of porosity. For the nuclear composite
considered here, there is considerable thermal conductivity anisotropy and temperature dependence typical
of all ceramics. This is demonstrated in Figure 24,
which gives the measured and calculated thermal
conductivity for the two nuclear-grade SiC composites.90 Presented are Hi-Nicalon™ Type-S fiber
and Tyranno™ SA fiber composites, each matrix
infiltrated through CVI.58 Architectures included
balanced (1:1:1 for x:y:z) and unbalanced (1:1:4) 3D
forms and 2D laminates (SW: satin weave, PW:

plain weave.) In each case, a pyrolytic graphite
interphase was applied. The conductivity for all
materials is presented in the through thickness direction (perpendicular to the plate and the fabric for the
2D composite.) This typically represents the lowconductivity direction.
As evident from Figure 24 and the supporting
analysis by Katoh,90 the fiber makes a significant contribution to the thermal conductivity of these highly
stoichiometric fiber composites, and this conductivity
is a fairly strong function of temperature. However, the
absolute conductivity is only a fraction of that for the
highest thermal conductivity CVD SiC (cf. Figure 7.)


Radiation Effects in SiC and SiC–SiC

237

400
Hi-Nicalon-S CVI UTS
Tyranno-SA3 CVI UTS
Hi-Nicalon-S CVI PLS

Nonirradiated

420

300
400

Tensile stress (MPa)


Tyranno-SA3 CVI PLS

460
480
690

760

780

610

UTS

530
380

200

220

1000
480

570 610

PLS

100


350

0
0

1

2

4
3
Neutron dose (dpa)

5

6

Figure 23 Effect of neutron dose on tensile proportional limit and ultimate tensile stresses for composites. Data labels
indicate the nominal irradiation temperature in  C for Hi-Nicalon™ Type-S (upright) and Tyranno™ SA-3 (oblique) composites.
Reproduced from Katoh, Y.; Snead, L. L.; Nozawa, T.; Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48–61.

80
3D 1:1:4 TySA/PyC, through-thickness, model
3D 1:1:1 TySA/PyC, through-thickness, model
2D-PW TySA/PyC, through-thickness, model
5HSW HNLS/PyC, through-thickness, model
3D 1:1:4 TySA/PyC through-thickness, experiment
3D 1:1:1 TySA/PyC through-thickness, experiment
2D-PW TySA/PyC through-thickness, experiment
5HSW HNLS/PyC through-thickness, experiment


Thermal conductivity (W m-1 K–1)

70
60
50
40
30
20
10
0
0

200

400

800
600
Temperature (ЊC)

1000

1200

Figure 24 Thermal conductivity of representative nuclear-grade SiC/SiC composite in unirradiated condition. Reproduced
from Katoh, Y.; Nozawa, T.; Snead, L. L.; Hinoki, T.; Kohyama, A. Fus. Eng. Des. 2006, 81, 937–944.

As with the CVD SiC discussed in section 4.07.3,
silicon carbide composite also undergoes significant

degradation in thermal conductivity because of
neutron irradiation. The data is somewhat limited;

however, Figure 25 gives the ambient throughthickness thermal conductivity for a plain weave
Hi-Nicalon™ Type-S, multilayer SiC interphase,
and CVI SiC matrix composite. It is noted that, in


238

Radiation Effects in SiC and SiC–SiC

Knonirr = 10.1 ± 2.2

Thermal conductivity at ambient (W m-1 K–1)

10

8

6
Tirr ~ 200 ЊC
Tirr ~ 800 ЊC
4
~ 600 ЊC
~ 400 ЊC

2
Nicalon Type-S fiber composite
0


0.001

0.01

0.1
Dose (dpa)

1

10

Figure 25 Effect of neutron irradiation on the through-thickness thermal conductivity of Hi-Nicalon™ Type-S, CVI matrix
composite.

1

Thermal defect resistance (W m-1 K-1)-1

~ 800 ЊC composite

0.1
Tirr ~ 200 ЊC
composite

0.01

0.001

Tirr ~ 800 ЊC

CVD SiC

Tirr ~ 200 ЊC
CVD SiC

0.001

0.01

0.1
Dose (dpa)

1

10

Figure 26 Comparison of the thermal defect resistance for neutron irradiated CVD SiC and Hi-Nicalon™ Type-S,
CVI matrix composite.


Radiation Effects in SiC and SiC–SiC

comparison to the conductivity shown in Figure 24
(second from lowest curve), the ambient throughthickness thermal conductivity for the material of
Figure 25 is relatively low (10.2 Æ 2.2 W mÀ1 KÀ1).
This is mostly ascribed to the large porosity for that
composite. Nevertheless, the figure clearly shows a
significant, irradiation temperature-dependent reduction in thermal conductivity as a function of irradiation
dose. The fact that this is temperature dependent
suggests that the degradation is due to the production of stable point defects and clusters, as discussed

in Section 4.07.3, although this may not be the
sole factor determining the degradation. Figure 26
provides the accumulated thermal defect resistance
at the lowest and highest irradiation temperature for
the composite materials of Figure 25, compared with
high-conductivity CVD SiC. It is interesting to note
that the thermal defect resistance for the composite,
while accumulating in the same manner as that of
the CVD SiC, is about an order of magnitude greater
than that of CVD SiC at a given dose (at least prior
to saturation.) This greater accumulation of thermal
defect resistance has been recently observed by
Katoh.67 The reason for this is unclear, although it is
plausible that, in addition to defect production, propagation of internal interfaces (e.g., cracks) in the composite is occurring under irradiation. It is also possible
that the defects population responsible for phonon
scattering for the composite material is stabilized at a
higher level than that of the highly pure CVD SiC.90

13.

14.
15.
16.

17.
18.
19.
20.
21.
22.

23.
24.
25.

26.

27.
28.
29.
30.
31.
32.

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