Developing a Macro-scale SiC-cladding Behavior Model Based on Localized Mechanical and Thermal Property Evaluation on Pre- and Post-Irradiation SiC-SiC Composites Fuel Cycle Research and Development Peter Hosemann University of California, Berkeley Collaborators Oregon State University University of Illinois, Urbana Champaign Frank Goldner, Federal POC Yutai Katoh, Technical POC Project No. 15-8439
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Developing a Macro-scale SiC-cladding Behavior Model Based on Localized
Mechanical and Thermal Property Evaluation on Pre- and Post-Irradiation
SiC-SiC Composites
Fuel Cycle Research and Development Peter Hosemann
University of California, Berkeley
CollaboratorsOregon State University
University of Illinois, Urbana Champaign
Frank Goldner, Federal POCYutai Katoh, Technical POC
Project No. 15-8439
Project Title: Developing a macro-scale SiC-cladding behavior model based on localized
mechanical and thermal property evaluation on pre-and post-irradiation SiC-SiC composites
Reporting Frequency: Final Report December 2018
Recipient: UC-Berkeley
Award number: CFA-15-8439
Awarding Agency: DOE
Working Partners: University of California, Berkeley
Oregon State University
University of Illinois Urbana Champaign
General Atomics
University of Oxford
Principal Investigator:
P. Hosemann, Dep. of Nuclear Engineering; UC-Berkeley
Milestone deliverables outlined by DOE Work Package ...................................................................................... 3
Collaboration roles and responsibilities ............................................................................................................... 3
Project status .................................................................................................................................................. 4
Executive summary of achievements ................................................................................................................... 4
Budget status ....................................................................................................................................................... 4
Communication and reporting status .................................................................................................................. 4 Journal publications ...................................................................................................................................... 5 Conference publications & presentations .................................................................................................... 5
FEA model development............................................................................................................................. 60 Microscale Models ...................................................................................................................................... 60
2.4.7.1 Micropillar Compression Model .......................................................................................................... 60 2.4.7.2 Fiber Pushout Model ........................................................................................................................... 61
Mini-composite model ................................................................................................................................ 62 2.4.8.1 Modelling Approaches ......................................................................................................................... 63
Summary and Conclusions .............................................................................................................................. 66
List of Figures ..................................................................................................................................................... 71
Proposal overview
Proposal abstract
Silicon carbide is being investigated for accident tolerant fuel cladding applications due to its good neutronic
performance, high temperature strength, exceptional stability under irradiation, and reduced oxidation
compared to Zircaloy under accident conditions. The development and investigation of these materials is
particularly important in the light of the Fukushima event and subsequent emphasis in DOE on accident
tolerant fuel (ATF) concepts. In this work, we propose to develop and improve upon existing small-scale
mechanical and thermal characterization methods, including micro-cantilever bend tests and fiber push-out
tests, and time-domain thermo-reflectance measurements. These techniques will be applied to evaluate
micro-scale properties of SiC-SiC composite constituents (matrix, fibers, interphase). The results will be
coupled with meso-scale fiber and void structural information and be used as input to develop a
comprehensive, finite element-based model based on constituent properties. The goal of this multi-scale
model is to be used to predict anisotropic mechanical and thermal cladding behaviors. The model will be
benchmarked against measurements of bulk properties and validated for radiation-damaged materials by
utilizing irradiated SiC-SiC composites available from General Atomics. The characterization approach
developed under this program will have particular application in the property prediction of reactor irradiated
materials, as only small volumes need to be tested, minimizing associated hazards and costs. The micro-scale
test methods developed here will be used along with meso-scale structural data to inform a finite element
model of the full SiC-based cladding system. This model will capture fabrication and irradiation effects and
importantly include for the first-time separate material data for each system component. The incorporation
of these micro-scale effects will lead to a more accurate model of SiC-SiC composite behavior, enabling
further advancements in the development and design of SiC-based materials for improved fuel cladding
performance.
The aim of this work is to develop advanced localized material characterization techniques to directly
measure mechanical and thermal properties of the individual constituents of SiC-based claddings at the
relevant micro-scale. SiC-SiC composites will be evaluated before and after irradiation and these results will
be coupled with macro-scale properties and microstructural information in order to provide the input
parameters for a comprehensive finite element model, which will be developed in this program.
Milestone deliverables outlined by DOE Work Package
1) Developing micro bend bar testing, fiber push out testing, and other small-scale mechanical testing
techniques: micro-pillar compression and nano-indentation.
2) Finish the manufacturing of all SiC samples
3) Receiving neutron irradiated SiC/SiC from GA
4) Spatially resolved thermal conductivity measurements on the as produced materials.
5) Determination of thermal conductance of interfacial layers of as produced materials.
6) Tomography measurements of the SiC materials
7) Macroscopic property measurements
8) Simulation of macro-scale component testing for comparison with experimental data.
9) Delivering a FEA model that incorporates all data generated
10) Localized mechanical property measurements on the irradiated material
11) Final Report
Collaboration roles and responsibilities
PI university (UC Berkeley) was to conduct small-scale testing on the composite constituents with specific
focus on debond and friction parameters as a function of interface characteristics. In situ SEM/FIB micro-
mechanical testing and TEM microstructural investigation were performed. The in situ synchrotron X-ray
based tomography work was also led by this institution.
Industrial Partner I (General Atomics) was responsible for sample and component manufacturing and
fabricated SiC-SiC composites with a range of composite parameters. Industry partner I supported the bulk
mechanical property characterization, fracture analysis, and X-ray tomography work to investigate porosity
size and distribution effects.
University Partner I (Oregon State University) was to develop a FEA model for the SiC-SiC composite.
Mechanical, thermal, microstructural properties were compiled from the literature, and by conducting first
principles calculations when needed. The FEA model was pursued to predict macro-scale component
behavior. University Partner I also developed sub-models to describe interface failure mechanisms as a
function of irradiation and fabrication.
University Partner II (University of Illinois UC) was responsible for high spatial resolution measurements
of the thermal conductivity of fibers, matrix, and interfacial phases by time-domain thermoreflectance
(TDTR) between room temperature and 300°C. Ion beam irradiation experiments will also be carried out to
complement the studies of materials damaged in reactor environment.
Foreign Institution I (University of Oxford) was responsible for development of high temperature
micromechanical testing techniques. Focused ion beam machining was used to manufacture micro-cantilever
samples in SiC- SiC composite materials supplied by Industrial Partner I. These cantilevers sampled the key
microstructural features of the materials, including, fiber, matrix and fiber matrix interphase. Failure
mechanisms were studied using high resolution SEM and correlated with observed local microstructure.
Project status
Executive summary of achievements
• All milestones and deadlines have been met.
• 1D, 2D, and layered composites were successfully fabricated with variable parameters including fiber
type and PyC interface thickness. This showed versatility in manufacture capability and enabled
investigation of relationship between constituent properties and bulk behavior. Additional unique micro-
composites and specialized laminates were fabricated for enhanced fundamental studies.
• Microstructure characterization was successfully applied to evaluate macroscopic composite density and
constituent grain structure. This was achieved through XCT, HRTEM, EBSD and Transmission
Kickuchi Diffraction (TKD).
• Constituent thermo-mechanical properties were successfully evaluated using micro-cantilever testing,
fiber pushout, micro-pillar compression, nano-indentation, and thermoreflectance measurements. This
Additional neutron irradiated samples were received from ORNL under a RTE NSUF proposal. Again, many
of these samples underwent similar characterization developed in this NEUP, and the results are interesting
and stimulating to discuss in this context. These data have been included to support the data of this project
as to provide a more comprehensive view of the impact of irradiation damage. NEUP GA irradiated samples
are all SA3 fiber composites w/ ~100nm PyC first layer thickness. NSUF samples are HNLS composites with
~50-100nm first layer PyC25. The applied experimental techniques include micro-cantilever testing for
fracture toughness, nano-indentation for elastic modulus and hardness, micro-pillar compression for interface
strength and friction properties, and TEM to understand degradation/evolution of the constituent
microstructure.
2.4.4.2 Neutron irradiation:
UCB Berkeley carried out nano-indentation and micro-pillar compression on irradiated composites to
evaluate the hardness, elastic modulus, and interfacial debond properties as a function of dose and irradiation
temperature. Details regarding experimental analysis were already described in previous sections, therefore
this section quickly present the results and summarizes the impact of neutron irradiation effects.
Nano-indentation:
Diamond tip Berkovich indentations were carried out using a MicroMaterials Nanoindenter on control and
neutron irradiated SA3 composites across a temperature range RT to 500°C. Below in Figure 45 show the
change in hardness and modulus for the fibers and matrix as a function of temperature. As described in section
2.4.2.1, indentation data for thin <1μm PyC interfaces has always proved challenging to gather statistically
satisfying properties. As a result, the nano-indentation campaign presents data for the fiber and matrix. The
neutron irradiated sample received 4.5dpa @630°C, and is from the exact samples used for micro-pillar
compression of the PyC interface. This data is also tabulated and compared to ion irradiation in Table 5.
Figure 45 – (Top) Plot of the hardness and (bottom) Elastic Modulus as a function of constituent,
irradiation, and temperature.
These plots show that hardness and modulus decrease for both fiber and matrix as a function of temperature.
Also, it shows fundamentally that hardness is increased and modulus is slightly decreased for the irradiated
samples. This would suggest the neutron irradiation point defect damage increases resistance to deformation
while the local disorder decreases the elastic modulus. This is consistent with literature and expected as point
defects reduce ability for plastic deformation. Though, the decrease in the modulus over temperature is larger
than expected for this temperature range, dropping nearly 100 and 50 GPa for the matrix and fiber
respectively over 500°C. It can be seen that the fibers still have a large variation in their mechanical properties
even over temperature and irradiation. This is likely due to the large concentration and variability of carbon
precipitates formed during the manufacturing process.
Looking at the RT indentations from Oxford and UC Berkeley for the unirradiated matrix, we find very
consistent hardness equal to ~37GPa, but a large difference in modulus from 460GPa to 540GPa respectively.
460GPa is consistent with literature, while 540 is likely an artifact of the data processing or system calibration
on that run. This is currently undergoing investigation, and we advise the application of 460GPa as the model
input for modulus, but still follow the slope of the hardness vs temperature curve established by Berkeley.
Micro-pillar compression:
Compression testing was carried out in-situ SEM using a Hysitron PI88 Picoindenter to evaluate the influence
of irradiation of interface debond strength and friction characteristics. Samples that were tested include ~50-
140nm PyC interface of HNLS and SA3 fiber composites respectively. The HNLS sample was exposed to
1dpa at ~350°, while the SA3 sample experienced 4.5dpa at ~630°C. An additional 12dpa HNLS fiber sample
with 180nm PyC was tested. Figure 46 compares the control and irradiated samples by applying the Mohr-
Coulomb criterion where the intercept is the cohesive shear strength and the slope the internal friction
coefficient. This linear regression, and thereby property values are displayed on Figure 46 as well as within
Table 5 in the property values section of this document. Graphics of representative pillars of each condition
are overlaid in Figure 46. The analysis revealed a significant decrease in cohesive shear strength for irradiated
samples, however without much impact on friction coefficient.
Figure 46 - Mohr-Coulomb criterion applied to the unirradiated and irradiated interfaces. A fundamental
decrease in cohesive shear strength between unirradiated and irradiated interfaces was observed.
The group averaged pillars at ~60° for each condition reveal a fundamental decrease in shear strength of
~200MPa and ~75MPa for SA3_140nm and HNLS_50nm respectively. This suggests that degradation in
shear strength scales with increase in neutron damage. With respect to irradiation effects, it can be observed
that while the chemical debonding strength decreases with irradiation, while the friction coefficient appears
remain the same. As a whole, the interface was weakened, however the relative contribution of chemical
bonding and frictional resistance has shifted. This is likely a result of irradiation induced microstructural
y =0.24x + 160
y =0.25x + 266
y =0.27x + 280
changes within the turbostatic graphite-like structures 39,40 . For thought, consider a perfect graphite lattice
for which shear failure occurs at the weakly bound basal planes. As irradiation displaces atoms from the
hexagonal frame work, it is conceived that there is an effective reduction in interacting basal plane surface
area, resulting in reduced strength. In addition, the displacements may result in new stacking and basal plane
arrangement, some of which may have less strength than the others.
The HNLS_180nm 12dpa showed very poor consistency. The group averaged interface inclines for
HNLS_180 12dpa, shown in Figure 46, reveal that the data is not applicable for the MC criterion. This is
likely a result of the relatively large porosity that evolved at this irradiation dose and temperature. Ultimately,
this porosity is expected to be inducing stress concentrations during testing, thereby introducing significant
data scatter and straying from a representative debond strength. To better understand the evolved properties
and characteristics, Figure 47 compares side-by side close up SEM images of the control and irradiated
samples. For HNLS_50nm 1dpa, major defects are not observed. HNLS_180nm 12pda shows significant
defect evolution, and SA3_140nm 4.5dpa does not show show major defects. There has been extensive
research into defect evolution and crystallite reconstruction for graphite and graphite-like materials under
irradiation 40–42. It has been found that swelling/shrinkage, dissolution and restructuring of nano-crystallites
is strongly dependent on the fabrication parameters, as well as the irradiation dose and temperature. The
porosity defects observed for HNLS_180nm 12dpa are expected to be a result of relatively low irradiation
temperature of 280°C. It has been discussed that the activation energy for graphite-like materials to undergo
significant point defect recombination during irradiation is around 300°C. Therefore, the porosity is likely an
artifact of reduced interstitial defect mobility, allowing for biased vacancy cluster formation. However, most
literature data is associated with bulk nuclear grade graphite which can exhibit substantially different
microstructure compared to the semi-oriented nano-crystallite nature of thin deposited PyC. Therefore,
drawing major conclusions and characteristics of historical nuclear graphite may be misleading.’
Figure 47 –Before and after irradiation SEM images comparing the PyC interface structure. (Top)
HNLS_50nm control vs 1dpa at 350°C. (Middle) HNLS_180nm control vs ~12dpa at 280°. (Bottom)
SA3_140nm control vs ~4.5dpa at 630°.
2.4.4.3 Ion Irradiation
The ion irradiation campaign was performed at University of Surrey Ion Beam Center. It included Si-ion
irradiation of SA3 SiC-SiC samples from GA of two different grades – one with single-layered, and one with
multi-layered interphase. In order to create a quasi-uniform damage profile, consecutive irradiations at three
different energies were accomplished – 500 keV, 1 MeV and 2 MeV. These conditions Irradiations were
performed with two peak damage levels – 0.26 and 2.6 dpa. Using SRIM software, the ion doses necessary
to achieve a specified damage level were calculated, and set as follows:
For peak damage of 0.26 dpa:
2 MeV – 6e14 ions/cm2;
1MeV – 4e14 ions/cm2;
0.5 MeV – 2e14 ions/cm2.
For peak damage of 2.6 dpa:
2 MeV – 6e15 ions/cm2;
1MeV – 4e15 ions/cm2;
0.5 MeV – 2e15 ions/cm2
Sample temperature during irradiation was set to 300°C for 0.26 dpa irradiation, and to 750°C for 2.6 dpa
irradiations. Irradiated samples were characterized using nanoindentation and cantilever fracture testing.
Nanoindentation was done using sharp Berkovich indenter and continuous stiffness measurement mode
(CSM). Two measurement modes were implemented:
Shallow indents:
In this measurement mode, shallow indents of 400 nm deep were made at 1.5 µm spacing. Lines of indents
were then going from matrix into a fiber, oriented normal to the surface. This way, hardness and modulus
could be measured both in the matrix and fiber materials; several indents at different radial position within a
fiber allowed monitoring the changes of already radially non-uniform properties. A disadvantage of this mode
is that only part of the irradiated depth can be probed. Figure 48 below presents an example of such a linescan,
for a sample of a single-layer interphase grade at 2.6dpa. In order to produce a single point corresponding to
a specific location, values of hardness or modulus from a corresponding depth dependence, as measured by
CSM method, were averaged in the depth range of 300 – 380 nm.
Figure 48 - Comparison of typical linescans of hardness and modulus across matrix and fiber for 2.6dpa
ion irradiated and reference unirradiated samples.
Deep indents:
Here deeper indents of 1000 nm deep were made, forming rectangular arrays in the matrix and irregular
arrays to probe the fibers, at 15 – 20 µm spacing. This way, most of the irradiated depth is probed, but indents
are necessarily located farther from each other, not allowing the probing of several locations within a fiber.
An example of such a measurement of depth dependence of hardness and modulus is presented in Figure 49
below:
-20 -10 0 1010
20
30
40
50
Hard
ness, G
Pa
Position relative to fiber center, m
300 nm deep indents
Unirradiated
Irraidated
-20 -10 0 10
200
300
400
500
Modulu
s, G
Pa
Position relative to fiber center, m
300 nm deep indents
Unirradiated
Irradiated
Figure 49 - Typical comparison of depth dependence of hardness and modulus in the matrix for samples
irradiated to different damage levels.
Nanoindentation measurements indicate that overall effect of irradiation on hardness and elastic modulus is
small, both in matrix and in the fibers. Hardness slightly increases, and modulus slightly decreases, similar
to that observed for neutron irradiation. Further increase of dose doesn’t make a significant effect on hardness
and modulus.
On the other hand, microscopic examination of the indents reveals that irradiation leads to the noticeable
modification of the crack patterns surrounding the indents, shown in Figure 50 below. Prominent radial cracks
appear at the corners of the indents in unirradiated matrix material. However, following the irradiation
cracking is completely suppressed.
Figure 50 - Comparison of crack patterns around the indents in the matrix of unirradiated an irradiated
samples.
This suppression of cracking can be attributed to the stress field induced in near-surface region sue to ion
irradiation. This stress would counteract the crack propagation, suppressing the cracking.
Micro-Cantilever testing:
Following the same procedures presented in section 2.4.2.2, testing was performed using straight-notched
triangular cantilevers, manufactured using FIB at the interphases, in the matrix and in the fibers. Figure 51
below presents the comparison of fracture toughness between irradiated and unirradiated samples, separately
for each composite constituent.
0 200 400 600 800 10000
10
20
30
40
Ha
rdn
ess, G
Pa
Depth, nm
1000 nm deep indents in the matrix
0.26 dpa
2.6 dpa
0 200 400 600 800 10000
100
200
300
400
500
Modulu
s, G
Pa
Depth, nm
1000 nm deep indents in the matrix
0.26 dpa
2.6 dpa
2 µm
Unirradiated
2 µm
Irradiated – 2.6 dpa
Figure 51 - Comparison of fracture toughness as measured in different constituents, for unirradiated and
irradiated samples.
It can be seen that there is a clear trend towards an increase of toughness for the interphases and fibers.
Toughness of the interphases increases after irradiation, both for single- and multi-layered ones. There is a
noticeable difference (about factor of 2) between the interphase of the irradiated and unirradiated samples;
on the other hand, toughness of the irradiated interphases is similar, regardless of the irradiation conditions.
Toughness of the fibers progressively increases with the increase of irradiation dose and temperature. On the
other hand, there doesn’t seem to be a definitive trend in the properties of matrices. Although different trends
arise with respect to constituent material, it is still the case that fracture toughness of the interface is ~<1/4
the toughness of the matrix. This is important for matrix crack deflection, and maintaining pseudo-ductility.
This supports observation of graceful failure in irradiated composites25.
To compare the effects of irradiation type, we can compare the hardness and modulus of the 760°C 2.6dpa
Si ion irradiated to the 630°C 4.5dpa neutron irradiated sample. The hardness for the matrix was 42GPa and
39GPa respectively. The modulus for the matrix is 425 GPa and 500GPa respectively. This reveals a relative
increase of ~14% and 5% for ion irradiated and neutron irradiated hardness respectively. For the modulus,
we observed a relative reduction of 5% and 7% for ion irradiated and neutron irradiated respectively. It is
unexpected that matrix properties change much beyond 1dpa as SiC reaches point defect swelling plateau.
Although relatively similar, one explanation for ion irradiation and neutron irradiation discrepancy is that
irradiation was not at the same temperature relevant dose rate, so point defect evolution and recombination
may have affected the damage state. With that in mind, the data may suggest that hardness is more strongly
dependent on the type irradiation, and that modulus is much less sensitive. The mechanical properties for
control, ion, and neutron irradiated samples are tabulated and compared below in Table 5.
Table 5 - Comparison of irradiated constituent properties tested at ambient temperature. Constituent H (GPa) E(GPa) KI
(MPa-m1/2)
τo
(MPa)
μi
Unirr.
SA3
M 37, 37* 460, 540* 4.1 - -
F Gradient with
center = 14,
25* (averaged)
Gradient with
center 190,
360* (averaged)
2.05 - -
I (140nm PyC) - - 0.8 280* 0.27*
Neutron:
4.5 dpa, 716°C
SA3
M 39* 500* - ~100* ~0.27*^
F 27* 330* - - -
I (140nm PyC) - - - - -
Unirr (NSUF):
HNLS
M - - - - -
F - - - - -
I (50nm PyC) - - - 266* 0.25*
Neutron
(NSUF):
1 dpa, 330°C
HNLS
M - - - - -
F - - - - -
I (50nm PyC) - - - 160* 0.24*
Si ion
irradiation:
0.24dpa 300°C
SA3
M 40 445 5.2 - -
F - - 2.5 - -
I (140nm PyC) - - 1.5 - -
Si ion
Irradiation:
2.6dpa 760°C
SA3
M 42 425 3.8 - -
F Gradient with
center 15
Gradient center
200
2.75 - -
I (140nm PyC) - - 1.3 - -
*UCB data ^ Estimated assuming limited change in μi observed for HNLS
TDTR on He-ion irradiated single crystal SiC:
A 6H crystal structure SiC wafer was implanted with He ions at 300°C at fluences of 0.016, 0.08, and 0.16
nC-μm-2. Following implantation, the samples were cleaned by spraying them with IPA, and were not heated
in order to not disturb the irradiated damage. Following procedures for TDTR outlined in section 2.4.3, the
He implanted areas were TDTR mapped with 1 μm resolution. In trial TDTR measurements, the aluminum
interfacial thermal conductance was calculated to be 80 MW m-2 K-1, which was slightly lower than
anticipated. However, the picosecond acoustic echo of the Al transducer had a sharp peak, indicating that the
interface was clean. It is suspected that the lower interfacial thermal conductance is only as a result of the
swelling from the He implantation, and that no foreign debris was present between the transducer and sample.
Though the interfacial thermal conductance is low compared to its values for un-implanted SiC crystals,
because the implanted area thermal conductivity is so low, the measurement is much more sensitive to the
implanted SiC area than the interface’s thermal conductance. Averaging the mapping data together, it was
found that the thermal conductivity of the sample at 0.016 nC-μm-2 was 13 W m-1 K-1 , at 0.08 nC-μm-2 it
was 2.8 W m-1 K-1 , and at 0.16 nC-μm-2 it was 2.2 W m-1 K-1 . Figure 52 shows the TLDR map of the 0.16 nC-μm-2 implantation and corresponding trend in thermal conductivity as a function of dose. This map also shows that the unirradiated thermal conductivity was approximately 100 W m-1 K-1, which is comparable to the unirradiated SiC composite matrix material (which is 3C cubic crystal structure and polycrystalline) at room temperature of 80-120 W m-1 K-1 presented earlier. Figure 52d shows the thermal conductivity trend as a function of He-ion implantation. It is clear that increasing dose rapidly reduces this conductivity, which is anticipated because of the sensitivity of thermal conductivity and phonon energy dispersion to point defect concentration. Although He implantation is not expected in fission reactor environments, this data supports finding for rapid deterioration of thermal conductivity found for neutron irradiation of monolithic polycrystalline SiC presented by Snead et al23, and may prove useful for fusion applications down the road.
Figure 52 - Thermal conductivity mapping and profile of a 20x20 μm He implanted region of a 6H SiC
wafer. The fluence of the implanted region is 0.016 nC-μm-2. (a) Photo of the He implanted region of the
SiC wafer. (b) A 20x20 μm thermal conductivity map of the He implanted region, taken at 1 μm steps. Time
delay was set to 𝑡𝑡 = 150 ps, heat capacity was C = 2.21 J cm-3 K-1, and the interfacial Al thermal
conductance for this region was 80 MW m-2 K-1. (c) Thermal conductivity profile of the implanted region.
(d) Plotted trend of thermal conductivity across fluences of 0.016 nC-μm-2, 0.08 nC-μm-2, and 0.16 nC-μm-2.
Macroscopic composite characterization In order to develop a suitable model that is able to predict composite behavior based on constituent level
properties, the macroscopic characteristics and mechanical properties must be evaluated for validation. The
first step of this effort was to capture the inherent porosity of the bulk composite that initiate micro-cracks in
the matrix. The following steps were to characterize the mechanical response and deformation of these
composites. The sections below describe and summarize these efforts in detail.
2.4.5.1 X-ray Tomography Composite materials fabricated via chemical vapor infiltration (CVI) method are prone to have some internal
porosity due to the nature of the process. The amount of internal porosity is a function of CVI process
conditions, component geometry and fiber architecture. The amount and the geometrical shape of internal
voids have a direct effect on mechanical properties of the bulk material and therefore should be considered
when evaluating mechanical performance of a composite. In this work, x-ray computed tomography (XCT)
was performed on a planar and tubular specimens with the goal to measure internal porosity. Samples
evaluated were SA3 and HNLS fiber composites produced at GA under the funding of this NEUP
All specimens were scanned at 100 kV and 100 μA x-ray beam power. Each specimen was reconstructed into
a 3D volume from 720 individual projection scans. The specimens were positioned as close to the beam
source as geometrically allowable in order to obtained the highest resolution of the scans. Volume Graphics
VGStudio Max 2.2 software was used to analyze scanned specimens and to determine internal and external
surfaces. The enclosed internal porosity was found by selecting a three dimensional boundary that was within
the sample surface and that also contained all of the internal porosity (no open porosity from the surface was
included and no internal porosity is excluded from the bounds). Once the boundary was established a simple
porosity search was conducted that finds voxels (three dimensional resolution units) within the preset bounds
that were then identified as voids. The total void voxel count, total void voxel volume, volume of contiguous
voids and their locations were all calculated within the Volume Graphics software. The resulting porosity is
represented as a void percent fraction of the total specimen volume. It should be noted that the porosity
measurement using XCT technique is ultimately limited by the resolution of a scan.
The XCT scan results for planar and tubular specimens are shown in Table 6. A three dimensional
reconstruction with the porosity maps is shown in Figure 53. It should be noted that beam hardening feature
was used during volume reconstruction of the planar composite because of higher beam scattering near flat
edges of planar specimens. Beam hardening corrections were not used for the tubular specimen. The resulting
porosity of planar specimens was measured to be less than 0.5% of total volume. The tubular specimen
showed a network of internal voids located within the composite braid, closely matching the fiber
architecture. Measured porosity for this specimen was 6.5% of the total volume.
D
Table 6 – Planar and Tubular specimen XCT scan results
Figure 53 – (left) XCT scan of a planar specimen with identified porosity map. (Right) XCT image showing
internal porosity of tubular specimen.
XCT mini-composite testing:
In situ tensile tests were performed on mini-composites to evaluate micro-crack evolution. We conducted a
24-hour beam time experiment at the Advanced Light Source (Lawrence-Berkeley Lab, University of
California Berkeley) to aid in the development and advancement of experimental methods for in situ tensile
testing at elevated temperature with simultaneous x-ray tomography imaging. A detailed description of the
beam-line and test set up was published by Bale et al43. Our work sought out to run these tests at room
temperature, 700°C and 1000°C to evaluate temperature dependence on micro-crack evolution and
deformation.
Testing was carried out on Tyranno ZMI fiber minicomposites with ~70nm PyC monolayer interfaces
fabricated by GA. The tensile test set up was comprised of ball in socket joints to allow for mini composite
self-alignment. The load train was surrounded by an inert gas chamber and directional halogen heat lamps
capable of bringing the system to 1600°C. Figure 54 is a schematic of the test chamber44. The mini-composite
samples were then mounted using OMEGABOND “700” high temperature cement in a bored and externally
threaded insert. Each insert was subsequently threaded into the high temperature brass balls for socket
mounting in the tensile fixture. Specimen temperature was regulated using a mounted thermocouple and
feedback system to adjust the current to the heat lamps. Examples of the mounted specimens are show at the
top of Figure 55. During testing, X-rays enter through a 0.5cm alumina window as the entire chamber rotates
180 degrees to capture the projection scans need for reconstruction. Because of the limited viewing window,
a notch was introduced at the center of the gauge length to promote microcrack evolution in the scan
projections. It has been discussed that the stress concentration around a notch in ceramic mini-composites
can dissipate quickly once localize micro cracking take place33. This means that micro-crack spacing away
from the notch can be representative of an unnotched specimen. Micro-crack spacing is directly related to
the interface shear properties that determine how large of a debond length will evolve along the fiber axis
once the matrix crack has deflected34. Typically, a weaker interface results in a larger debond length and
thereby larger micro-crack spacing.
Planar composite Tubular composite
Resolution, μm 5.4 5.6
Average Width, mm 7.0 8.20
Average Thickness, mm 1.0 9.68
Beam Hardening Minimal None
Porosity, % <0.5% 6.5
Figure 54 – (Left)Schematic illustrations showing a specimen mounted between upper and lower grips.
(Right) External image of chamber with halogen lamps at 700°C
Testing was successfully carried out for three ambient and three 700°C mini-composite specimens.
Unfortunately, the high temperature cement consistently failed at 1000°C. Testing was carried out by
applying incremental loads via displacement-controlled steps of 5μm. A typical load vs time curve is shown
in Figure 55 for sample T5 that was tested at 700°C. The peak load for this test was 24lbs.
Figure 55 - (Top) Sequential stage of test set up; high temperature cement casting, in situ high temperature
loading, and post fracture evaluation. Note that a thermo couple was attached to this specimen to
appropriately track the thermal state. (Bottom) Typical load vs time plot following 5um incremental lading
steps.
Alumina window for x-ray penetration
5μm disp.-
controlled steps
Peak load ~ 24lbs
At each incremental load, an x-ray projection scan was taken. This enabled real-time visualization of
microcrack evolution, shown in Figure 56, and the ability to evaluate 3D reconstructions. These
reconstructions are made by stacking individual z-slices that represent the material cross section. Figure 56
shows a reconstructed cross section. The resolution of the tomography system offered 1.6um/pixel. Using
imageJ, the cross-sectional area and porosity could be extracted. The cross-sectional area of 0.53*10-6m2
reveal failure stress equal to ~200MPa, which is within the typical range for SiC/SiC mini-composites.
Porosity was found to be ~5%, which is also typical of CVI production as will be discussed in the following
section. Looking at the normal projection of Figure 56, micro-crack spacing is clearly visible. The observed
average micro-crack spacing, taken by dividing the observable micro-cracks into the height of the observable
window, is on the order of 0.35mm for both ambient and 700°C. This translates to an interfacial sliding
strength of ~45MPa for these mini-composite configurations. Details on this value is discussed in detail in
the next section regarding hysteresis testing. Similar micro-crack spacing for both ambient and 700°C
suggests minimal change in interface characteristics across this temperature change. This is expected
considering the system is in an intern environment and both carbon and SiC material is not expected to
degrade mechanically until much greater temperatures.
Figure 56 – (top) Tomographic reconstruction showing the cross-sectional geometry of a single tow tensile
test sample. (bottom left) Normal x-ray projection of ambient temperature test. Numbers denote visible
micro-cracks. (Bottom right) Normal x-ray projection of 700°C test, numbers denote visible micro-cracks.
We see this behavior as key to understanding fiber-matrix interaction. By creating simplified representations
of fractures bridged by fibers, we can investigate the matrix, fiber, and interaction parameters that support
development of subsequent fractures and control fracture spacing. If this fracture behavior does not appear
with a particular interaction condition (e.g. simple friction) then more complex models will be investigated.
Individual parameters (e.g. friction coefficient) that underpin working models can then be cycled back
through the research group for reconciliation with existing data.
2.4.5.2 Hysteresis testing of unidirectional mini-composites Hystersis testing explores a myriad of composite behavior and performance characteristics from residual
stresses to damage tolerance22,33,45–48. The specific goal of this effort was to link interface properties from
micro- to macro-scale and to illuminate fundamental differences betweent the two. There are two analytical
relationships that have been developed to evaluate interfacial sliding strength based on macroscopic
performance. The first is related to the hystersis loop width of the stress-strain data, and the second is related
to the micro-crack evolution and specifcally the distance between these micro-cracks. Hystersis loop width
measurments requires high fidelity strain measurement. Unfortunately our experimental set up had size and
spacing limitations that eliminated the option for high precision strain gauges. Additionally the SEM field of
view was such that digital image correlation (DIC) methods were not applicable. However, the SEM enabled
unique insight to the micro-crack evolution and spacing. The analytical model for interfacial sliding strength
is related to matrix crack spacing by the following equation48
𝜏 = 𝜎𝑠𝑅𝑓𝑉𝑚
2𝑉𝑓𝐿𝑠
where σs is the crack saturation stress (usually taken to be the peak stress before failure for PyC interface
containing composites46), Rf, is the fiber radius, V is the volume fraction for matrix and fiber respectively,
and Ls is the average measured crack spacing across the gauge length. This equation is derived from the
debond length of the defelcted crack, and more specifically, the resitantce to sliding behind that crack tip.
This resistance to sliding, often refered to as the sliding strength or shear “stress” (as opposed to shear
“strength” that is the resistance to initial debond), governs the fiber-matrix load sharing as the composite is
put in tension. If the sliding strength is large, the matrix carries more load, and is therefore more likely to
crack, with dependence on defect distribution in the matrix of course. This is what gives rise to relatively
uniform crack spacing observed in experimental hysteresis testing. Increasing the sliding strength decreases
the observed matrix crack spacing. Figure 57 is a schematic that describes load sharing as function of debond
length for a given sliding strength. Where the sliding strength is dependent on the residual stress, friction
coefficient and initial debond strength. Ls represents the average matrix crack spacing that occurs based on
the stress and defect distributions.
Figure 57 – Schematic of load sharing between fiber and matrix as a function of debond length. Debond
strength as well as sliding strength are dependent on residual stress and friction characteristics.24
Thereby it is important to note that the value obtained via hysteresis testing is fundamentally diffrenet from
the τdebond described earlier by the micro-pillar and pushout testing. Both are important parameters that
influence the debond length and energy abosorbtion (micro-cracking) of the composite as a whole. However,
they are not the same property and therefore cannot be directly compared. With that, the values presented
below are lower than the values found for either micro-pillar comporession or fiber pushout. This is
reasonable as the stress required to slide an already debonded interface is expected to be less than the measure
the true bond strength.
UC Berkeley received two sets of unidirectional, single tow (~500 fiber bundle), mini-composites. One set
contained HNLS fibers with 50nm PyC monolayer fiber-matrix interphase. The other set contained SA3
fibers and the same interphase. The unidirectional fiber architecture was fully infiltrated via CVI SiC
processing at GA, then sent to UC Berkeley. Images of the mini-composites were shown earlier in Figure 1.
Figure 58 shows typical cross-sections for the SA3 and HNLS composites. The area, porosity, fiber and
matrix volume fraction were extracted for each mini-composite. Porosity for the SA3 composites was ~8%
while HNLS showed ~4%, and fiber volume fraction was around ~25% for both.
Figure 58 – Fracture surface cross-sections of SA3 (left) and HNLS (right) mini-composites. The total area
and porosity are outlined in yellow. ImageJ measurement results, in μm2, are overlaid. Small circular black
dots in the cross-section are locations of fiber pullout.
Testing was performed using the Kammrath & Weiss Tensile and Compression Module in-situ XL30 Phillips
SEM, Figure 59a. Testing was performed using a 500N load cell and displacement control at 1μm/s. A
baseline hysteresis test schedule called for 5N load steps, returing back to the first infleciton point of the load
curve (~30N) associated with flexing and final alignment of custom grips. The instrument software was very
versitle and allowed for test pausing and redefinition of load step size and lower bounds when necessary, see
Figure 60. This enabled SEM imaging of the entire gauge length mid test, as well as insight to hystersis
behavior with microcracks fully open. Special attention was paid to the gripping configuration and alignment
during mounting. A custom designed ball-and-socket gripping system was machined in house at UC
Berkeley, shown in Figure 59a&c. The mini-composites were cut to ~5cm, 2cm was used for epoxy gripping
on both ends, leaving a 1cm gauge length. The ends were mounted in bored and externally threaded studs
and set up in an alignment ficture to cure, Figure 59b.
Figure 59 – A) Kammrath and Weiss module on SEM stage with loaded sample. B) Alignment fixture for
epoxy mounting and curing. C) SA3 mini-composite test specimen with SS balls on threaded studs. Ball and
socket joints were lubricated with colloidal graphite
In total, four SA3 and four HNLS minicomposites were tested. Figure 60 shows typical stress-strain curves
for SA3 (left) and HNLS (right). Stress was evaluated via the cross-sections of each sample tested, shown
in Figure 58. As alluded to earlier, strain data was only collected via cross-head output readings. Thereby,
Total area μm2
porosity μm2
only general characteristics of the stress-strain curve are worth discussing. It is noted that the averge slope of
each loop reduces as loading and unloading takes place. This is a measure of damage tolerance and is both
expected and desired in CMCs. This damage tolerance is a result of micro-crack evolution. Secondly we
observe that loop width is increasing as the test progresses. This is a sign of increased micro-cracks and
thereby increased debonded interface. The sliding of the debonded interface, or energy absorption, is why
hystersis is observed. Finally we can point out that the average onset of matrix cracking, or proportional limit
stress, was around ~175MPa for HNLS and ~125MPa for SA3. The ultimate tensile strength was on average
about 300MPa for both. These values align reasonably with those in literature5, though are on the high end.
This may be because we accounted for and subtracted out porosity from our cross-section for stress
calculation. The lower PLS value for SA3 is attributed to the slightly less homogenous cross-section and
increased porosity, leading to more defects and therefore crack initiation sites.
Figure 60 – Stress-strain curves for SA3 (left) and HNLS (right). The SA3 curve show the versatility of the
KW software to pause, and redefine load/unload regime to explore effects of interface degradation at
matrix crack saturation.
During testing, loading was paused and images were taken along the entire gauge length to capture the matrix
crack spacing. Figure 61 shows the resulting SEM images of gauge length and micro-cracking.
Figure 61 – a) Stitched SEM image of the gauge length prior to failure for SA3 mini-composite #4. The
gauge length is saturated with matrix cracks. B) Zoomed in image from the same SA3 sample. Showing
typical matrix crack spacing. Overlaid value for Ls is the average value across the entire gauge length. C)
HNLS minicompoiste after failure with matrix crack spacing still visible. Dotted lines attributed to flaking
following abrupt crack closure of the thin conductive coating that was deposited.
Applying the observed micro-cack spacing to the equation above provided the interfacial sliding strength
values shown in the table below.
τsliding
for mini-composites with 50nm monolayer PyC interphase
τ SA3
= 62 MPa
τ HNLS
= 17 MPa
These values follow the same trend observed for micro-pillar compression. The rougher fiber results in
increased resistance to failure. In this case, it is likely that this roughness is contributing significantly to the
frictional resistance to sliding. It has also been shown that roughness can increase residual clamping stress
during fabrication49. Both the residual stress and friction can explain the increased sliding strength for the
SA3 composite. These constituent level values are critically important from a modelling perspective. The
next generation of modelling needs both the constituent properties as well as information regarding the type
of damage that is evolving. This hysteresis testing in conjunction with the micro-mechanical testing has
provided a unique platform to identify these characterstics and will prove useful for the continued modelling
efforts.
2.4.5.3 Mechanical testing of woven composites General Atomics performed a series of mechanical testing of SiC/SiC composites. Planar and tubular shaped
composite samples were tested and were all fabricated at GA using the Chemical Vapor Infiltration method
(CVI). Planar composites were reinforced with Tyranno-SA3, and tubular composites were made with Hi-
Nicalon Type S fiber. Composite tubes with a fiber ratio of hoop to axial directions at 2:1 and a pyrolytic
carbon interphase layer ~100nm were used for elastomeric expending plug test that is based on ASTM C1819,
and c-ring testing at room and 800°C test temperatures based on ASTM C1323 standard, and 4-point bed
tests. The test method and results are described below. Provided the load versus displacement data for bulk
composite test specimens provides a standard to validate code against.
Flexural 4-point bend test results for 8 specimens are shown in Table 7. The composites in planar shape were
fabricated with Tyranno SA3 fiber. All specimens with the nominal size of 52 mm by 6.8 mm by 1.1 mm
were cut from a single plate using a waterjet technique. After the initial cut, the longest sides were polished
with a 30 micron diamond disc to remove large defects. The final width of the test specimens contained about
two repeating fiber unit cells. The loading pins of the test fixture were spaced by 1/4 of the loading span
which was 42 mm. A typical flexural stress versus extension is shown in Figure 62a, small load drops are
observed upon loading, representative of microcrack formation.
Table 7 - Flexural 4-point bend test results
For the flexural testing, while the specimens underwent a significant amount of bending, and showed signs
of micro-crack evolution, the fractured fragments indicated a more brittle like behavior with minimal fiber
pull out, shown in Figure 62b. In 7 out of 8 specimens fracture occurred at one or both loading pins, which
are typically considered non-valid tests in 4-point bend testing.
For 8 specimens: Extension, mm Flexural Stress, MPa
45. Curtin, W. A. Theory of Mechanical Properties of Ceramic‐Matrix Composites. J. Am. Ceram. Soc. 74, 2837–
2845 (1991).
46. Domergue, J. M., Vagaggini, E. & Evans, A. G. Relationships between Hysteresis Measurements and the
Constituent Properties of Ceramic Matrix Composites: II, Experimental Studies on Unidirectional Materials.
Journal of the American Ceramic Society 78, 2721–2731 (1995).
47. Kerans, R. J., Parthasarathy, T. A., Rebillat, F. & Lamon, J. Interface properties in high-strength nicalon/C/SiC
composites, as determined by rough surface analysis of fiber push-out tests. J. Am. Ceram. Soc. 81, 1881–
1887 (1998).
48. Lamon, J. & Bansal, N. Ceramic Matrix Composites: Materials, Modeling and Technology. (2015).
49. Buet, E. et al. Influence of surface fibre properties and textural organization of a pyrocarbon interphase on the
interfacial shear stress of SiC / SiC minicomposites reinforced with Hi-Nicalon S and Tyranno SA3 fibres. J.
Eur. Ceram. Soc. 34, 179–188 (2014).
50. Hashin, Z. Failure Criteria for Unidirectional Fiber Composites 1. J. Appl. Mech. 47, 329–334 (1980).
51. Hashin, Z. & Rotem, A. A Fatigue Failure Criterion for Fiber Reinforced Materials*. J. Compos. Mater. 7,
448–463 (1973).
52. Tabatabaei, S. A., Lomov, S. V. & Verpoest, I. Assessment of embedded element technique in meso-FE
modelling of fibre reinforced composites. Compos. Struct. 107, 436–446 (2014).
53. Lubliner, J., Oliver, J., Oller, S. & Oñate, E. A plastic-damage model for concrete. Int. J. Solids Struct. 25,
299–326 (1989).
List of Figures
Figure 1 - Single tow mini-composites with ZMI SiC fibers (left) and SA3 fibers (right), fabricated by GA
that were provided for characterization. .......................................................................................................... 7 Figure 2 - Six planar SiC-SiC panels fabricated for this work (top row are baseline panels with a single PyC
layer at the fiber matrix interface; bottom row are panels with multi-layer interphase). ................................. 7 Figure 3 – Examples of single-fiber mico-composites. Initial fabrication produced a non-continuous SiC
coating (left) while subsequent fabrication produced individual fibers coated with a uniform CVD SiC later
(right). .............................................................................................................................................................. 8 Figure 4 - SEM images of the polished surface of the composite: (a) general appearance, (b) close-up of the
vicinity of fiber bundle, (c) close-up with individual fibers, denoted by arrows. ............................................ 8 Figure 5 - (Left) SEM image of an edge polish required for pillar fabrication. (Right) Height image of a SiC
fiber in the SiC matrix. It can be seen that the graphite containing regions were removed preferentially during
polishing. ......................................................................................................................................................... 9 Figure 6 - STEM image of a typical region containing fibers and surrounding matrix, arrow indicates
submicron-sized pores. .................................................................................................................................... 9 Figure 7 – (Left) Image quality map of a large area, including several fibers and a surrounding region of the
matrix, artifact CVI rings are visible. (Right) STEM image of the matrix material. ..................................... 10 Figure 8 - A) SA3 TEM foil. B) SA3 first layer PyC/fiber interface with oriented graphitic structure. C SEM-
STEM image of HNLS foil. D) HNLS first layer PyC/fiber interface with oriented graphitic structure. E)
Diffraction ring pattern at HNLS first layer PyC with characteristic graphite rings (002).8 ......................... 11 Figure 9 - EELS spectra comparison of the pyrolytic carbon interphase and secondary carbon phase of the
SA3 fiber showing graphitic structure ........................................................................................................... 11 Figure 10 - TEM images: (a) cross-section of a fiber, dashed line corresponds to the central axis of it; (b)
fiber material with the location of EDX line scan denoted, arrows indicate inclusion at the grain boundaries;
(c) radial distribution of inclusion within the fiber; (d) EDS linescane – signals of C and Si. ...................... 12 Figure 11 -(a) TEM image of the microstructure and the corresponding IPFX map. The arrow represents the
direction of grain growth. The vertical dashed line in the IPFX map represents the location of the interlayer;
(b) pole figure of the matrix (top row) and fiber (bottom row). .................................................................... 13 Figure 12 - (a) Close-up of a typical matrix microstructural, with fringes within grains visible; (b) comparison
of a TKD orientation map and a TEM image of the same area. Red rectangles denote the twin boundaries,
and it is evident that the fringes do not directly correspond to these. ............................................................ 14 Figure 13 - High-resolution TEM image of an area within the matrix. The line follows is parallel to atomic
rows, and it is evident that areas of varying contrast (fringes) are not associated with different crystal
structure. ........................................................................................................................................................ 14 Figure 14 - (a) Example of a line of indents, crossing a matrix region and one of the fibers; (b) typical depth
dependence of hardness (filled symbols) and modulus (hollow symbols) as a function of depth for the indents
placed within the matrix (solid lines) and close to the center of a fiber (dashed lines); vertical lines denote
the depth range used for averaging. ............................................................................................................... 15 Figure 15 - Hardness (filled symbols) and modulus (hollow symbols) as a function of distance from the fiber
center averaged over multiple line scans crossing the fibers; vertical dashed lines denote the typical
dimensions of a fiber. .................................................................................................................................... 16 Figure 16 - Results of high-temperature nanoindentation – temperature dependences of hardness and
modulus. ........................................................................................................................................................ 17 Figure 17 - Typical cantilever in the (a) matrix;(b) fiber; (c) interphases; standard triangular cross-section is
visible. ........................................................................................................................................................... 17 Figure 18 - Typical notched cantilever in the matrix SiC.............................................................................. 18 Figure 19 - Results of fracture tests from cantilevers in different components of a composite; two values for
the matrix are from the longitudinal (hollow symbol) and transverse (filled symbol) cantilevers. Overlay – a
stress-strain curve of a typical test. ................................................................................................................ 19 Figure 20 - Fracture toughness of different components of a composite. Two values for the matrix are from
the longitudinal (hollow symbol) and transverse (filled symbol) cantilevers. ............................................... 19 Figure 21 - Typical fracture surfaces of the notched cantilevers in different constituents: (a) interphase
(remaining base of the cantilever, visible in the image, is in the fiber); (b) fiber; (c) matrix – longitudinal
direction; (d) matrix – transverse direction. Note the geometry of the visible straight notch. ...................... 20
Figure 22 - TEM images of the fractured cantilevers: (a) at the matrix-fiber interphase; insert – EDX
elemental map of the highlighted area; (b) in the bulk fiber, arrow denote the C precipitates at grain
boundaries, X denotes the instance of intergranular crack propagation, Y the instance of transgranular crack
propagation instances of crack propagation; (c) in the bulk matrix, transverse orientation; here and in (d) the
arrows are to guide the eye along the crack; (d) located in the bulk matrix, longitudinal orientation. Cantilever
in (a) is unnotched, all the others are notched. .............................................................................................. 22 Figure 23 - SEM image of the impression left by a flat-punch tip on the fiber (indicated by an arrow),
indicating minor plastic deformation during loading. ................................................................................... 24 Figure 24 - Comparison of (a) highly porous central part of a typical fiber tow, and (b) relatively monolithic
periphery ....................................................................................................................................................... 24 Figure 25 - Results of a number of push-out tests performed at the fibers (a) in the center and (b) at the
periphery of the tows. .................................................................................................................................... 25 Figure 26 - Cross-section of the near-interphase region following a push-out, indicating the crack path. .... 26 Figure 27 - (Top) Typical micro-pillar fabrication process using FIB milling techniques, overlays show the
extraction of and interface incline and resulting interfacial area. (Bottom) Comprehensive set of micro-pillar
interface conditions that were tested. Notably there is a a wide range of PyC interface thickness and fiber
type. ............................................................................................................................................................... 28 Figure 28 - Snapshots from live mico-pillar compression testing of SA3_700nm PyC. (Left) Diamond flat
punch applying load. (Right) Post failure with fracture surface maintained. ................................................ 28 Figure 29 - (a) Representative schematic of pillar fabrication. (b) Resolved stress state and force balance of
a representative micro-pillar structure. .......................................................................................................... 29 Figure 30 - Comparison of raw data for control sample HNLS_50nm, HNLS_1300nm, and
HNLS_50nm_1dpa to show fundamental reduction in strength with irradiation and thickness. .................. 30 Figure 31 - MC criterion applied to SA3 and HNLS micro pillars with varying PyC thickness. It is observed
that for the same thickness, SA3 shows fundamentally stronger debond shear strength. .............................. 30 Figure 32 - (Left) NEUP micro-pillar property values plotted versus PyC thickness. An empirical relationship
was fitted for application in modelling efforts. (Right) Literature fiber pushout debond strength values as a
function of PyC. ............................................................................................................................................ 31 Figure 33 – (Top) example fracture micro-pillar suitable TEM foil fabrication and TEM dark field image of
foil with slipped HNLS pillar. (Bottom) HRTEM image of deformed PyC layer as a result of failure. 002
graphite-like planes are visible. And suggest fracture is occurring along those weak basal planes. ............. 32 Figure 34 - A) TEM image of HNLS interphase pre-failure. B) TEM image of PyC interface post failure
suggesting cohesive failure in the PyC layer. ................................................................................................ 32 Figure 35 - (Left) SEM fractography evaluation of fracture HNLS and SA3 micro-pillars. SEM images of
fiber surfaces are reproduced from Sauder et al6. (Right) AFM scan of SA3_700nm fracture surface, and
corresponding tabulated data. ........................................................................................................................ 33 Figure 36 - Load vs displacement curve of pillar compression showing failure load and shaded area under
the curve (work of fracture) for extraction of the fracture release rate energy. HNLS_B micro-pillar with
thick PyC interphase prior to compression. ................................................................................................... 34 Figure 37 - Energy release rate as a function of interface angle grouped in three and interface layer thickness
(data call-outs in nanometers). The mode I energy release rate was calculated using Eq.3 Cedric’s relationship
ΓII = 3.5 *ΓI .................................................................................................................................................. 34 Figure 38- - Isolation of chemical bonding contribution to fracture energy release rate. Post interface failure
of micro-pillar from HNLS_A with pillar cap still intact. ............................................................................. 35 Figure 39 - Detail of structure of a SiC composite and its interphase (a) SEM picture of the Hi-Nicalon type
S fibers from a SiC composite. The interphase is clearly present. (b) Schematic of structure of the interphase
material structure, consisting of 40nm pyrolytic carbon (PyC), followed by four repeating units of 50nm
SiC/10 nm PyC .............................................................................................................................................. 36 Figure 40 - This shows the TDTR system layout. The Ti:Sapphire produces laser pulses that are split by a
PBS (Polarizing Beam Splitter) and two-tint wavelength filters into pump (red line) and probe (purple line)
pulses. The pump beam is modulated by an EOM (electro-optic modulator), and the probe beam goes through
a mechanical chopper. The pump beam heats the sample, while the probe beam measures the reflectance of
the transducer layer (e.g., Al). ....................................................................................................................... 37 Figure 41 - TDTR data fits for the thermal conductivity of the SiC/SiC composite components and the ratio
of the in-phase and out-of-phase voltage as a function of thermal conductivity. (a) Data and TDTR model fits
for all mapped areas of the SiC/SiC composite, using a spot size of ωo = 2.9 µm, (b) Calculated ratio values
for t = 150 ps, a spot size of ωo = 2.9 µm, a heat capacity of C = 2.21 J cm-3 K-1, and varying interface thermal
conductance ................................................................................................................................................... 38 Figure 42 - Diagram and a thermal conductivity map of an interphase of the fiber. (a) A diagram showing a
cross section of the SiC/SiC composite where a fiber has been puleed out of the matrix. According to the
geometry of these regions, the average angle of the fibers to the surface has been determined to be 5 degrees,
which creates an interphase area of ~4 µm. (b) A micrograph of a pulled fiber. (c) A thermal conductivity
map of the region at the end of a fiber that has been pulled out of the matrix. Circled in the area where full
TDTR measurements were taken. Spot size (ωo) is 2.9 µm, time delay = 150 ps, and heat capacity (C) is 2.21
J cm-3 K-1 ....................................................................................................................................................... 39 Figure 43 - Summary of thermal conductivity measurements of fibers and matrix as a function of temperature.
(a) A 50µm x 50 µm thermal conductivity map, taken at 1 µm steps, at room temperature (lighter region
indicates higher thermal conductivity), (b) A cross section through the map in figure5a (dashed white line)
at room temperature, (c) A cross section of the same region at different temperatures, (d) Summary of thermal
conductivity dependencies on temperatures for matrix and fiber, compared with ........................................ 40 Figure 44 - Radioactive materials storage cabinet. Source drawers are complete with lead shielding and
locking mechanisms ...................................................................................................................................... 41 Figure 45 – (Top) Plot of the hardness and (bottom) Elastic Modulus as a function of constituent, irradiation,
and temperature. ............................................................................................................................................ 43 Figure 46 - Mohr-Coulomb criterion applied to the unirradiated and irradiated interfaces. A fundamental
decrease in cohesive shear strength between unirradiated and irradiated interfaces was observed. .............. 44 Figure 47 –Before and after irradiation SEM images comparing the PyC interface structure. (Top)
HNLS_50nm control vs 1dpa at 350°C. (Middle) HNLS_180nm control vs ~12dpa at 280°. (Bottom)
SA3_140nm control vs ~4.5dpa at 630°. ....................................................................................................... 45 Figure 48 - Comparison of typical linescans of hardness and modulus across matrix and fiber for 2.6dpa ion
irradiated and reference unirradiated samples. .............................................................................................. 46 Figure 49 - Typical comparison of depth dependence of hardness and modulus in the matrix for samples
irradiated to different damage levels. ............................................................................................................ 47 Figure 50 - Comparison of crack patterns around the indents in the matrix of unirradiated an irradiated
samples. ......................................................................................................................................................... 47 Figure 51 - Comparison of fracture toughness as measured in different constituents, for unirradiated and
irradiated samples. ......................................................................................................................................... 48 Figure 52 - Thermal conductivity mapping and profile of a 20x20 μm He implanted region of a 6H SiC wafer.
The fluence of the implanted region is 0.016 nC-μm-2. (a) Photo of the He implanted region of the SiC wafer.
(b) A 20x20 μm thermal conductivity map of the He implanted region, taken at 1 μm steps. Time delay was
set to 𝑡𝑡 = 150 ps, heat capacity was C = 2.21 J cm-3 K-1, and the interfacial Al thermal conductance for this
region was 80 MW m-2 K-1. (c) Thermal conductivity profile of the implanted region. (d) Plotted trend of
thermal conductivity across fluences of 0.016 nC-μm-2, 0.08 nC-μm-2, and 0.16 nC-μm-2. .......................... 50
Figure 53 – (left) XCT scan of a planar specimen with identified porosity map. (Right) XCT image showing
internal porosity of tubular specimen. ........................................................................................................... 51 Figure 54 – (Left)Schematic illustrations showing a specimen mounted between upper and lower grips.
(Right) External image of chamber with halogen lamps at 700°C ................................................................ 52 Figure 55 - (Top) Sequential stage of test set up; high temperature cement casting, in situ high temperature
loading, and post fracture evaluation. Note that a thermo couple was attached to this specimen to
appropriately track the thermal state. (Bottom) Typical load vs time plot following 5um incremental lading
steps. .............................................................................................................................................................. 52 Figure 56 - (top) Tomographic reconstruction showing the cross-sectional geometry of a single tow tensile
test sample. (bottom left) Normal x-ray projection of ambient temperature test. Numbers denote visible
micro-cracks. (Bottom right) Normal x-ray projection of 700°C test, numbers denote visible micro-cracks.
pullout is observed. (c) Macroscopic view of fractured test specimen. ......................................................... 58 Figure 58 – (Left) The dog bone specimen shape. (Center) Image of an axial tension test setup. (Right) Stress-
strain curve for a dog bone shaped specimen from a planar tension test ....................................................... 58 Figure 59 -. (Right) Example of digital image correlation strain map during test. (Right) Typical fracture
surface observed after mechanical testing of planar SiC/Sic samples, showing limited fiber pull-out.) ....... 59 Figure 60 - A representative stress-strain curve from an elastomeric expanding plug test. .......................... 59
Figure 61 - Load versus crosshead extension for a typical RT and High Temperature c-ring tests .............. 60 Figure 62 – The leftmost image shows a finite element model using cohesive surfaces after decohesion has
occurred. Results from this model are shown as the solid green line on the plot at right. The center picture
shows a model with Mohr-Coulomb plasticity in a thin layer (which appears in red) just after the layer fails.
Results from this type of model are shown as circular red dots in. The remaining data points in the plot are
experimental results. ...................................................................................................................................... 61 Figure 63 - Fiber pushout model (left) and comparison with experiments (right). The 2D axisymmetric model
was revolved 180° for visualization purposes. Model predictions are plotted at right in the orange dashed
line. Note that the load at which the PyC layer fractures and the load begins to drop is determined in part by
residual stress between the fiber and matrix, which is unknown. .................................................................. 62 Figure 64 – Raw tomography projection scan of partially broken mini-composite with pre-machine notch on
right. Note the array of cracks through the composite matrix. ...................................................................... 62 Figure 65 - Simple two-dimensional single fiber model (left) and three-dimensional multi fiber model used
to explore the feasibility of a larger-scale mini-composite (right). These models used predefined cracks with
cohesive contact as well as cohesive contact between the matrix and fibers. While results from the two-
dimensional model were encouraging, convergence in the three-dimensional model was very poor and even
the simple model shown at right was computational expensive. ................................................................... 63 Figure 66 - Two-dimensional homogenized model of fiber "tow" with Hashin damage model. The red
elements show the damaged material and the predicted crack path .............................................................. 64 Figure 67 - Results from Embedded Element Technique with Concrete Damaged Plasticity material models.
This technique includes separate finite elements for the matrix and fibers. The matrix elements are shown
with contours for material degradation (a) and equivalent plastic strain (b). These figures show the
progression of failure in the matrix. Fiber elements are shown in (c). Here, contours have the same scale as
(a) and show degradation of the fiber material. Note that this model predicts that fracture propagates across
the specimen and does not capture the radial cracking pattern shown in testing. .......................................... 65