- Email: [email protected]
SiC composites fabricated by PIP
Journal Pre-proof Effect of pyrolysis temperature on the mechanical evolution of SiCf/SiC composites fabricated by PIP Lian-Yi Wang, Rui-Ying Luo, Guang-Yuan Cui PII:
S0272-8842(19)32616-1
DOI:
https://doi.org/10.1016/j.ceramint.2019.09.087
Reference:
CERI 22854
To appear in:
Ceramics International
Received Date: 30 June 2019 Revised Date:
23 August 2019
Accepted Date: 9 September 2019
Please cite this article as: L.-Y. Wang, R.-Y. Luo, G.-Y. Cui, Effect of pyrolysis temperature on the mechanical evolution of SiCf/SiC composites fabricated by PIP, Ceramics International (2019), doi: https://doi.org/10.1016/j.ceramint.2019.09.087. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Ltd.
Effect of pyrolysis temperature on the mechanical evolution of SiCf/SiC composites fabricated by PIP Lian-Yi Wang1, Rui-Ying Luo*,1, Guang-Yuan Cui1 (1 School of Physics and Nuclear Energy Engineering, Beihang University, Beijing, 100191, China * Corresponding author. E-mail address: [email protected] (R.-Y. Luo). Tel. & Fax: +86-010-82338267)
Abstract Three types of SiCf/SiC composites with a four-step three-dimensional SiC fibre preform and pyrocarbon interface fabricated via precursor infiltration and pyrolysis at 1100
, 1300
, and 1500
were heat-treated at 1300
under argon atmosphere for
50 h. The effects of the pyrolysis temperature on the microstructural and mechanical properties of the SiCf/SiC composites were studied. With an increase in the pyrolysis temperature, the SiC crystallite size of the as-fabricated composites increased from 3.4 to 6.4 nm, and the flexural strength decreased from 742 ± 45 to 467 ± 38 MPa. After heat treatment, all the samples exhibited lower mechanical properties, accompanied by grain growth, mass loss, and the formation of open pores. The degree of mechanical degradation decreased with an increase in the pyrolysis temperature. The composites fabricated at 1500
exhibited the highest property retention rates
with 90% flexural strength and 98% flexural modulus retained. The mechanism of the mechanical evolution after heat treatment was revealed, which suggested that the thermal stability of the mechanical properties is enhanced by the high crystallinity of the SiC matrix after pyrolysis at higher temperatures. Key words: SiCf/SiC composite; pyrolysis temperature; microstructure; mechanical properties; heat treatment.
1 Introduction Silicon-carbide-fibre-reinforced silicon carbide matrix (SiCf/SiC) composites are
receiving increasing attention as potential candidates for high-temperature components (e.g. in turbine engines) owing to their excellent properties, including low density, excellent high-temperature mechanical performance with good mechanical damage tolerance (which enables them to overcome the inherent brittleness of monolithic SiC), exceptional oxidation and corrosion resistance, and low thermal expansion coefficient [1–6]. Various processing routes are used to fabricate SiCf/SiC composites, including chemical vapour infiltration (CVI) [7,8], precursor infiltration and pyrolysis (PIP) [9–11], liquid silicon infiltration [12,13], and the nano-infiltration and transient eutectic phase process [14,15]. Among these routes, the PIP process, in which a fibre preform is densified using the precursor-derived ceramic (PDC) SiC matrix, has been extensively studied owing to the designability of the polymer precursor molecules, lower pyrolysis temperature, low cost, and ease of processing [16–19]. However, the thermal stability of PDC SiC significantly influences the application of PIP SiCf/SiC composites at elevated temperatures [20–22]. Studies have demonstrated that the changes in the microstructure of the PDC SiC bulk material, which consists of nanocrystalline domains of SiC, amorphous SiOxCy, and carbon impurities, depend mainly on the processing and heat-treatment conditions (specifically, pyrolysis temperature, and/or exposure time). The SiC crystallite size tended to increase, accompanied by a decomposition of amorphous SiOxCy into SiC nanocrystalline domains and gaseous SiO and CO. Further, numerous cracks and pores were formed owing to volume shrinkage [23–26]. Moreover, the microstructural changes exhibited similar trends with an increase in the pyrolysis temperature and/or time. Rahman et al. [27] studied the structural changes in PDC SiC as a function of the pyrolysis temperature. The results revealed that the crystallite size of the SiC nanocrystalline domain increased from 6 to 22 nm as the pyrolysis temperature was increased from 900 to 1400
. Furthermore, the degree of crystallinity increased with
an increase in the pyrolysis temperature. Poerschke et al. [28] observed an increase in the crystallite size and crystallinity during the pyrolysis of precursor-derived SiC ceramics. The SiC samples fabricated at 800
were initially amorphous. The
crystallite size of the nanocrystalline SiC and the degree of crystallinity increased to 6 nm and 68%, respectively, after 100 h of annealing at 1300
, whereas the crystallite
size and degree of crystallinity increased to 124 nm and 80%, respectively, after 100 h of annealing at 1500
. When PDC SiC is used as the matrix in composites, both the
pyrolysis-temperature-dependent microstructure and the microstructural evolution at elevated temperatures significantly affect the properties of the composites. Therefore, to determine the applicability of SiCf/SiC composites fabricated via PIP under thermal conditions, extensive characterisation of the microstructural and mechanical evolution during treatment at high temperatures, especially as a function of the processing conditions, is required. However, only limited data are available in the literature. In this study, SiCf/SiC composites with a pyrocarbon (PyC) interface fabricated via PIP at various temperatures between 1100 °C and 1500
were heat-treated at 1300
under Ar atmosphere. The pyrolysis-temperature-dependent microstructural and mechanical evolution of the as-fabricated and heat-treated composites were studied. The characteristics of the PDC SiC bulk material fabricated under various temperatures and of SiC fibres heat-treated under different conditions were studied as controls. The performance evolution of the SiC fibres during preparation and heat treatment was evaluated. Accordingly, the mechanism of the mechanical evolution after heat treatment was revealed.
2. Materials and experimental procedure 2.1 Sample fabrication The SiCf/SiC composites examined in this study were fabricated via PIP, reinforced with Amosic SiC fibres provided by Xiamen University, China, and precoated with a PyC interphase. The fabrication process is shown in Fig. 1. SiC fibres were used to prepare the four-step three-dimensional fibre preform with a fibre volume fraction of ~0.4. Before the PIP process, a PyC layer was deposited on the preform via CVI. Propylene (C3H6) was applied for the PyC deposition at 850
and 2 kPa for 10 h. After interphase deposition, the SiC matrix was prepared to densify the preform via PIP using a 50 wt% polycarbosilane (PCS, provided by the National University of Defense Technology, China)/xylene solution as the precursor. The coated fibre preforms were dipped into the solution under vacuum for more than 2 h. Subsequently, the preforms were pyrolysed at 1100 °C, 1300 °C, or 1500
for 1
h under a protective Ar atmosphere. The PIP procedure was repeated until the weight increase of the composite was less than 1%. The samples are denoted as CMC-1100, CMC-1300, and CMC-1500, where CMC indicates the SiCf/SiC composite and the numbers indicate the corresponding pyrolysis temperatures. The composites were then machined into rectangular bars measuring approximately 40 × 4 × 3 mm3 for tests. A PDC SiC bulk material was prepared via the same pyrolysis process used for the composites, where PCS was pyrolysed at 1100 °C, 1300 °C, or 1500
for 1 h under a
protective Ar atmosphere. The mass of the PCS before and after pyrolysis was measured
five
times
using
an
electronic
balance
(resolution:
0.1
mg).
The ceramic yields of PCS pyrolysed under various temperatures were calculated as
ω% =
m - m0 × 100% m0
,
(1)
where m is the average mass of the sample after pyrolysis and m0 is the average initial mass of the sample. Subsequently, the product was crushed to obtain particles of sizes less than 100 µm. The samples are denoted as PDC-1100, PDC-1300, and PDC-1500, where PDC indicates the PDC SiC powder and the numbers indicate the corresponding pyrolysis temperatures.
2.2 Heat treatment The SiC fibres were annealed at 1100 °C, 1300 °C, or 1500
under a protective
Ar atmosphere for the entire PIP holding time; these samples are denoted as F-1100, F-1300, and F-1500, respectively. The objective was to obtain fibres in a similar state to that of the as-fabricated composites. The as-received fibres and composites were moved to a high-temperature horizontal tube furnace (GSL-1700X, Ke Jing Materials Technology Co. Ltd., China) and heated to 1300
under flowing Ar (0.1 L/min). The
fibres and composites were characterised after 25 h and 50 h of annealing.
2.3 Mechanical property tests A monofilament tensile test of the SiC fibres was conducted. A schematic diagram of the sample shape and size is shown in Fig. 2. A single SiC fibre of length ~150 mm was separated from the fibre bundle and glued to a supporting paper for ~30 mm of the length at both ends. Subsequently, the fibre was loaded at the centre of the unsupported length of 90 mm at a loading speed of 0.1 mm/min. More than 10 samples were analysed. The flexural strength and modulus of the composites before and after 50 h of heat treatment were characterised using a three-point bending test. The dimensions of the samples were 40 × 4 × 3 mm3, and they were tested at a span of 30 mm and a loading speed of 0.5 mm/min. Five samples of each type were analysed, and the average values of the five samples are reported.
2.4 Microstructural characterisation Archimedes’ principle was employed to measure the open porosity of the composites before and after heat treatment. The mass change of the composites during the heat-treatment processing was measured using an electronic balance. Five samples of each type were analysed for both tests, and the average values of the five samples are reported. The phases of the SiC fibre, PDC SiC powder, and SiCf/SiC composites treated under various conditions were analysed using X-ray diffraction (XRD, D/Max 2500PC, Rigaku, Japan) with Cu Kα radiation (λ = 0.154 nm) and 2θ = 10°–90°. The scan rate was 4°/s. Further, the crystallite size of the (111) plane for each sample was calculated using Scherrer’s formula as follows: D hkl =
kλ βcos θ
,
(2)
where Dhkl is the crystallite size, k is the Scherrer constant, β is the peak width at half-height, and θ is the Bragg diffraction angle. The transverse surfaces of the as-fabricated and heat-treated samples and the
microstructures of the as-fractured samples were observed using scanning electron microscopy (SEM, S4800, Hitachi, Japan). The elemental analyse of the transverse surface was performed with an energy-dispersive X-ray spectroscopy (EDS, equipped with SEM).
3 Results and discussion 3.1 Characterisation of the PDC SiC powder fabricated at various pyrolysis temperatures Fig. 3 shows the XRD patterns of the PDC SiC powder prepared at various pyrolysis temperatures. β-SiC was the only crystalline phase detected using XRD, and the peaks at 35.7°, 60.0°, and 71.8° correspond to the (111), (220), and (311) planes of
β-SiC crystalline grains, respectively. All these samples exhibited characteristic broad halo peaks with diffuse intensity, which indicated an amorphous/nanocrystalline structure. These peaks became sharper and more intense when the pyrolysis temperature was increased to 1500
. Two new peaks at 41.4° and 75.5°, which
correspond to the (002) and (022) planes, respectively, were observed for PDC-1500. These findings suggest that the amorphous structure was converted to nanocrystalline domains with an increase in the pyrolysis temperature. Furthermore, the crystallite sizes of the (111) plane for the three samples, as analysed using Scherrer's formula, are shown in Table 1. The sizes also increased with an increase in the pyrolysis temperature, rising from 2.3 to 6.8 nm, further indicating the development of the nanocrystalline structure in the samples. The above results are validated by a comparison with test data from the literature in Table 1. There was also a negative correlation between the ceramic yield and pyrolysis temperature as shown in Table 1. The relatively high ceramic yield decreased from 68 ± 5% to 57 ± 6% when the temperature was increased from 1100 to 1500
. The mass
loss is attributed to the formation of gaseous SiO and/or CO as the SiOxCy decomposed.
3.2 Microstructural and mechanical evolution of SiC fibres The SiC fibres were annealed at 1100 °C, 1300 °C, or 1500
for the entire PIP
holding time to evaluate the effects of the preparation process on the microstructure and mechanical properties of the SiC fibres, which are denoted as F-1100, F-1300, and F-1500, respectively. Subsequently, the as-received fibres were heat-treated at 1300
under Ar atmosphere. The microstructure and mechanical properties of the
fibres without heat treatment were measured as a control. The XRD patterns of the SiC fibres annealed under each preparation condition are shown in Fig. 4(a), and the crystallite sizes of the (111) plane for these samples are listed in Table 2 (samples with 0 h of heat treatment). All these samples exhibited characteristic broad halo peaks. Further, there were three main peaks at 35.7°, 60.0°, and 71.8°, which correspond to the (111), (220), and (311) planes of β-SiC crystalline grains, and a weak peak at 41.4° that corresponds to the (002) plane of β-SiC crystalline grain. The average grain size of the fibres treated with various preparation processes was 5.7–6.1 nm for the untreated fibre, F-1100, and F-1300, and it increased to 7.1 nm for F-1500, indicating that grain growth occurred in the SiC fibres when the pyrolysis temperature was increased to 1500 After heat treatment at 1300
.
for various times [the XRD patterns shown in Fig.
4(b)–(d)], the grain size was almost constant for each type of sample as listed in Table 2. The average grain size was 5.7–6.8 nm for the F-1100 and F-1300 fibres and was higher for the F-1500 fibres (7.1–8.6 nm). These observations suggested that the microstructure of the SiC fibres was insensitive to temperature during heat treatment at 1300
.
The monofilament tensile properties of the SiC fibres annealed under each preparation condition were determined and are shown in Table 2 and Fig. 5. The monofilament tensile strength of the untreated fibres was also determined for reference. The tensile strength decreased slightly, from 2.74 ± 0.32 GPa (untreated fibres) to 2.69 ± 0.39 and 2.46 ± 0.34 GPa after the fibres were annealed during preparation at 1100 °C and 1300
, respectively. When the pyrolysis temperature was
increased to 1500
, the tensile strength decreased significantly to 1.51 ± 0.31 GPa,
indicating that the properties of the SiC fibres were degraded at the relatively high pyrolysis temperature of ~1500 After heat treatment at 1300
. for 50 h, the monofilament tensile properties of the
SiC fibres were tested and are shown in Table 2 and Fig. 5. The strength of the heat-treated F-1100, F-1300, and F-1500 fibres decreased to 2.25 ± 0.68, 2.04 ± 0.47, and 1.33 ± 0.51 GPa (decreases of ~16%, ~17%, and ~12%), respectively. The relatively weak effect of temperature on the decrease in tensile strength indicated that the SiC fibres were thermally stable at the examined temperatures.
3.3 Microstructural evolution of the SiCf/SiC composites 3.3.1 Dependence of microstructural evolution of as-fabricated SiCf/SiC composites on pyrolysis temperature The microstructural evolution of the as-fabricated composites (the samples with 0 h of heat treatment) is shown in Table 3. With an increase in the pyrolysis temperature, the density of the composites increased slightly from 2.21 ± 0.03 to 2.30 ± 0.02 g/cm3. The porosity of the composites exhibited a decreasing trend, from 5.6 ± 0.4% (CMC-1100) and 5.7 ± 0.3% (CMC-1300) to 5.1 ± 0.5% (CMC-1500). The XRD patterns of the as-fabricated SiCf/SiC composites are shown in Fig. 6(a), which are similar to those of the PDC SiC samples fabricated at each temperature. Three broad halo peaks appeared at 35.7°, 60.0°, and 71.8° for the CMC-1100 and CMC-1300 samples, and two new peaks appeared at 41.4° and 75.5° when the pyrolysis temperature was increased to 1500
. The crystallite sizes of the (111) plane
for the three as-fabricated composites, as analysed using Scherrer’s formula, are shown in Table 3. The crystallite sizes of the composites increased with an increase in the pyrolysis temperature, i.e. they were 3.4, 4.8, and 6.4 nm for CMC-1100, CMC-1300, and CMC-1500, respectively. Owing to the insignificant change in the grain sizes of the SiC fibres during preparation, the grain growth of the composites is mostly attributed to the PDC SiC matrix. The differences in the crystallite sizes of the composites compared with those of both the PDC SiC powder and the SiC fibres
might be associated with the strong dependence of the crystallisation behaviour of PDC ceramics on the local microstructure [28].
3.3.2 Dependence of microstructural evolution of heat-treated SiCf/SiC composites on pyrolysis temperature The microstructures of the SiCf/SiC composites after heat treatment at 1300 were characterised. The mass loss, open porosity, and crystallite sizes of the composites after heat treatment under various conditions are listed in Table 3. The phases of the composites after heat treatment under each condition were analysed using XRD as shown in Fig. 6(b)–(d). With an increase in the exposure time at 1300
, the aforementioned peaks became sharper and more intense. Furthermore,
the peaks at 41.4° and 75.5° appeared after 25 h of heat treatment for CMC-1100 and CMC-1300. In addition, the crystallite sizes after various exposure times are shown in Fig. 7, where the as-fabricated composite (0 h of treatment) serves as a control. The crystallite sizes increased monotonically with an increase in the exposure time. After heat treatment for 50 h, the crystallite size increased from 3.4, 4.8, and 6.4 nm to 38.8, 25.9, and 17.7 nm for CMC-1100, CMC-1300, and CMC-1500, respectively. These findings suggest that the SiC crystallites became coarser and the amorphous structure was converted to nanocrystalline domains with an increase in the exposure time. The evolution of mass loss rates with an increase in the exposure time is plotted in Fig. 8. Each sample lost mass. The mass loss rates for CMC-1100, CMC-1300, and CMC-1500 after heat treatment for 50 h were 1.37 ± 0.12%, 0.61 ± 0.12%, and 0.28 ± 0.06%, respectively. Those phenomena are attributed to the decomposition of SiOxCy accompanied by the formation of gas (SiO and/or CO). As shown in Fig. 9, the open porosity for each sample, measured using Archimedes’ principle, exhibited an increasing trend, and it rose to 9.8 ± 0.4%, 8.2 ± 0.5%, and 6.4 ± 0.4% for CMC-1100, CMC-1300, and CMC-1500, respectively, after heat treatment for 50 h. The effect of heat treatment on the open pores was observed using SEM, with the transverse surfaces of the three heat-treated samples (the heat-treated samples of CMC-1100, CMC-1300, and CMC-1500) shown in Fig. 10
and the ones of the corresponding untreated samples considered as the control. The low-magnification images of the transverse surfaces for each sample were similar. The structures of the as-fabricated and heat-treated CMC-1300 samples are shown in Fig. 10 (a) and (b), respectively, as representative structures. The open pores of the samples consist of large pores between bundles and small pores between the fibres in a fibre bundle. Comparing the structures of the transverse surfaces before and after heat treatment, it can be observed that the heat-treatment process had a more significant effect on the pores between the fibres in a fibre bundle. The three untreated samples exhibited a high density in the fibre bundles with the pores between fibres filled by the matrix (presented in Fig. 10 (c-1), (d-1), and (e-1)). In contrast, a large number of micro-cracks were formed in the fibre bundles after the heat treatment, and penetrating cracks were observed in the heat-treated CMC-1100 samples. These cracks, which are attributed to volume shrinkage caused by grain growth and/or decomposition, resulted in an increase in the open porosity. Comparing the structures of the three heat-treated samples presented in Fig. 10 (c-2), (d-2), and (f-2), it can be observed that the structures of CMC-1100 were different from those of the other samples, indicating that SiC particles (as confirmed using EDS in Fig. 11) with a regular geometric shape were attached to the sample surface. These structures indicated that CMC-1100 could easily form SiC with higher crystallinity after the heat treatment. As the SiC fibres maintained a stable microstructure during heat treatment, the microstructural evolution of the PDC SiC matrix resulted in the dramatic grain growth of the composites. Note that the CMC-1100 samples exhibited greater grain growth and showed both the highest mass loss and the highest open porosity after 50 h of heat treatment. These findings confirm that the microstructural evolution of CMC-1100 was more intense than that of the other samples and indicate that the PDC SiC matrix fabricated under lower pyrolysis temperatures was less thermally stable.
3.4 Mechanical evolution of the SiCf/SiC composites 3.4.1 Dependence of mechanical evolution of as-fabricated SiCf/SiC composites on pyrolysis temperature The mechanical properties of the as-fabricated SiCf/SiC composites prepared at various pyrolysis temperatures are listed in Table 4. The typical stress–strain curves and the fracture morphology of the as-fabricated SiCf/SiC composites are shown in Fig. 12 and Fig. 13 (a)–(c), respectively. For the as-fabricated samples prepared at different pyrolysis temperatures, the mechanical behaviours indicate the pseudo-plastic fracture mode, where three regions appear in the typical stress–strain curves. The stress first increased almost linearly with an increase in the strain in the elastic region, where only elastic elongation of the constituents and marginal interfacial sliding occurred. The mechanical behaviour then entered the pseudo-plastic region, where the slope decreased with an increase in the loads, and matrix crack initiation and interfacial debonding occurred. Finally, the stress decreased after the initial failure in the breakdown region, where fibre breakdown accompanied by fibre pull-out occurred, as shown on the fracture surfaces in Fig. 13 (a)–(c). As the pyrolysis temperature was increased from 1100 to 1300
, the flexural
strength of the as-fabricated samples decreased slightly, from 742 ± 45 to 649 ± 51 MPa. However, when the pyrolysis temperature was increased to 1500
, the flexural
strength declined sharply to 467 ± 38 MPa. Conversely, the modulus increased with an increase in the pyrolysis temperature, rising from 99 ± 7 GPa(CMC-1100) to 108 ± 4 GPa (CMC-1500). The decrease in strength is attributed to the degradation of the properties of the SiC fibres at relatively high pyrolysis temperatures, especially at ~1500
. By contrast,
the increase in modulus could be understood according to the model of the modulus for composites, as follows:
Ec = EfVf + EmVm
(3)
where Ec, Ef, and Em are the moduli of the composite, fibres, and matrix, respectively;
Vf and Vm are the volume contents of the fibres and matrix, respectively. According to previous studies [18,23,27], crystalline PDC SiC growth accompanied by increasing modulus of PDC SiC occurs with an increase in the pyrolysis temperature. Therefore, the modulus of the composites increases, according to the model in Eq. (3).
3.4.2 Dependence of mechanical evolution of heat-treated SiCf/SiC composites on pyrolysis temperature The mechanical properties of the heat-treated SiCf/SiC composites are listed in Table 4, and the fracture morphology and typical stress–strain curves are shown in Fig. 12 and Fig. 13 (d)–(f), respectively. For the heat-treated samples, the mechanical parameters still indicate pseudo-ductile fracture mode, as shown in Fig. 13 (d)–(f). However, a large yield before fracture appears in the stress–strain curves of the heat-treated CMC-1100 samples, which differ from the curves of the heat-treated CMC-1300 and CMC-1500 samples. This finding indicates lower load-bearing and load-transfer capacity of the matrix in CMC-1100. All the samples exhibited a degradation of the mechanical properties after heat treatment. The strength and modulus decreased to 371 ± 21, 481 ± 61, and 421 ± 39 MPa and 62 ± 6, 88 ± 5, and 106 ± 6 GPa for the CMC-1100, CMC-1300, and CMC-1500 samples, respectively, after heat treatment for 50 h. The heat-treated CMC-1100 samples exhibited the most severe property degradation, retaining only 51% of their flexural intensity and 63% of their flexural modulus. In contrast, the mechanical properties of the heat-treated CMC-1500 samples were less degraded, as they retained 90% of their flexural strength and 98% of their flexural modulus. Further, the highest flexural strength after pyrolysis among the three types of samples was 481 ± 61 MPa for CMC-1300 (retention of 74%). The change in the properties is attributed to the changes in the microstructure. As both the microstructure and mechanical properties of the SiC fibres remained stable during heat treatment at 1300
, the lower mechanical properties of the composites
are attributed to matrix degradation. The volume shrinkage and cracking in the matrix caused by PDC SiC matrix crystallisation and decomposition yield a more porous
structure, which weakens both the load-bearing and load-transfer capacities. Moreover, the microstructural transformation is greater at lower pyrolysis temperatures. These dual functions would first generate mechanical failure of the matrix and then limit the bearing capacity of the fibres in the composites. Consequently, the properties of the composites are degraded. Furthermore, the thermal stability is enhanced by the high crystallinity of the SiC matrix fabricated at higher pyrolysis temperatures.
4 Conclusion SiCf/SiC composites with a four-step three-dimensional SiC fibre preform and PyC interface fabricated via PIP at various pyrolysis temperatures (1100 °C, 1300 °C, and 1500
) were heat-treated at 1300
under Ar atmosphere. The dependence of the
microstructural and mechanical evolution of the as-fabricated and heat-treated composites on the pyrolysis temperature was studied. As the pyrolysis temperature was increased from 1100 to 1500
, the flexural strength of the as-fabricated
composites decreased from 742 ± 45 to 467 ± 38 MPa, whereas the modulus and SiC crystallite size increased from 99 ± 7 to 108 ± 4 GPa and from 3.4 to 6.4 nm, respectively. The decrease in strength is attributed to degradation of the properties of the SiC fibres during preparation. Further, the enhanced grain growth in the PDC SiC matrix at higher pyrolysis temperatures caused the increase in the modulus and crystallite size of the composites. After heat treatment for 50 h at 1300
, all the
samples exhibited a degradation of the mechanical properties, accompanied by grain growth, mass loss, and open pore formation. The flexural strength of the CMC-1100, CMC-1300, and CMC-1500 samples after 50 h of heat treatment decreased to 371 ± 21, 481 ± 61, and 421 ± 39 MPa, and the SiC crystallite size increased to 38.8, 25.9, and 17.7 nm, respectively. The heat-treated CMC-1100 samples exhibited the most dramatic microstructural transformation and property degradation, retaining only 51% of their flexural strength and 63% of their flexural modulus. By contrast, the CMC-1500 composites maintained the smallest SiC crystallite size and exhibited the least mechanical property degradation, retaining 90% of their flexural strength and
98% of their flexural modulus. As both the microstructure and mechanical properties of the SiC fibres remained stable during heat treatment at 1300
, the microstructural
evolution and mechanical property degradation of the composites are attributed to the degradation of the PDC SiC matrix (including SiC grain growth and decomposition of amorphous SiOxCy into nanocrystalline SiC domains and gaseous SiO and CO). Further, the thermal stability was enhanced by the high crystallinity of the SiC matrix fabricated at higher pyrolysis temperatures.
Acknowledgements This work was supported by the National Natural Science Foundation of China (No. 21071011).
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2017;43(10):7387-92. [18] Moraes KV, Interrante LV. Processing, Fracture Toughness, and Vickers Hardness of Allylhydridopolycarbosilane-Derived Silicon Carbide. Journal of the American Ceramic Society. 2003;86(2):342-6. [19] Dong SM, Katoh Y, Kohyama A, Schwab ST, Snead LL. Microstructural evolution and mechanical performances of SiC/SiC composites by polymer impregnation/microwave
pyrolysis
(PIMP)
process.
Ceramics
International.
2002;28(8):899-905. [20] Larson NM, Zok FW. In-situ 3D visualization of composite microstructure during polymer-to-ceramic conversion. Acta Materialia. 2018;144:579-89. [21] Colombo P, Mera G, Riedel R, Soraru GD. Polymer-Derived Ceramics: 40 Years of Research and Innovation in Advanced Ceramics. Journal Of the American Ceramic Society. 2010;93(7):1805-37. [22] Nechanicky MA, Chew KW, Sellinger A, Laine RM. α-Silicon carbide/β-silicon carbide particulate composites via polymer infiltration and pyrolysis (PIP) processing using polymethylsilane. J Eur Ceram Soc. 2000;20(4):441-51. [23] Zunjarrao SC, Rahman A, Singh RP. Characterization of the Evolution and Properties of Silicon Carbide Derived From a Preceramic Polymer Precursor. Journal Of the American Ceramic Society. 2013;96(6):1869-76.
[24] Gou Y, Wang H, Jian K. Formation of carbon-rich layer on the surface of SiC fiber by sintering under vacuum for superior mechanical and thermal properties. J Eur Ceram Soc. 2017;37(3):907-14. [25] Cao S, Wang J, Wang H. Effect of heat treatment on the microstructure and tensile strength of KD-II SiC fibers. Materials Science and Engineering: A. 2016;673:55-62. [26] Kaur S, Riedel R, Ionescu E. Pressureless fabrication of dense monolithic SiC ceramics from a polycarbosilane. J Eur Ceram Soc. 2014;34(15):3571-8. [27] Rahman A, Zunjarrao SC, Singh RP. Effect of degree of crystallinity on elastic properties of silicon carbide fabricated using polymer pyrolysis. J Eur Ceram Soc. 2016;36(14):3285-92. [28] Poerschke DL, Braithwaite A, Park D, Lauten F. Crystallization behavior of polymer-derived Si-O-C for ceramic matrix composite processing. Acta Materialia. 2018;147:329-41.
Tables Table 1. Evolution of crystallite size and ceramic yield with an increase in the pyrolysis temperature
Sample
Pyrolysis temperature ( )
Crystallite size (nm)
Ceramic yield(%)
PDC-1100
1100
2.3
68±5
PDC-1300
1300
4.1
66±7
PDC-1500
1500
6.8
57±6
1100
2
-
1300
3
-
1500
7
-
[28]
Table 2. Crystallite sizes and tensile strength of SiC fibres for various exposure times
Exposure time
Crystallite size
Tensile strength
(h)
(nm)
(GPa)
-
5.8
2.74±0.32
0
5.7
2.69±0.39
25
6.0
-
50
6.2
2.25±0.68
0
6.1
2.46±0.34
25
6.3
-
50
6.8
2.04±0.47
0
7.1
1.51±0.31
25
8.5
-
50
8.6
1.33±0.51
Sample Untreated fibre
F-1100
F-1300
F-1500
Table 3. Microstructural evolution of the SiCf/SiC composites for various pyrolysis temperatures and exposure times
Pyrolysis temperature
Exposure time
Density
Open porosity
Crystallite size
Mass loss
( )
(h)
(g/cm-3)
(%)
(nm)
(%)
0
2.21±0.03
5.6±0.4
3.4
0
25
-
8.7±0.3
33.8
1.13±0.21
50
-
9.8±0.4
38.8
1.37±0.12
0
2.25±0.04
5.7±0.3
4.8
0
25
-
7.6±0.2
21.6
0.53±0.08
50
-
8.2±0.5
25.9
0.61±0.12
0
2.30±0.02
5.1±0.5
6.4
0
25
-
6.1±0.4
17.3
0.23±0.02
50
-
6.4±0.4
17.7
0.28±0.06
Sample
CMC-1100
CMC-1300
CMC-1500
1100
1300
1500
Table 4. Mechanical properties of as-fabricated and the heat-treated SiCf/SiC composites
As-fabricated samples
Heat-treated samples (50h)
Flexural
Flexural
Flexural
Residual
Flexural
Residual
strength
modulus
strength
strength ratio
modulus
modulus ratio
(MPa)
(GPa)
(MPa)
(%)*
(GPa)
(%)*
CMC-1100
742±45
99±7
371±21
51
62±6
63
CMC-1300
649±51
107±7
481±61
74
88±5
82
CMC-1500
467±38
108±4
421±39
90
106±6
98
Sample
*Calculated from the average data for the samples before heat treatment
Figures
Fig. 1. Flowchart of SiCf/SiC composite fabrication process
Fig. 2. Schematic of monofilament tensile test sample
Intensity
(111) (002)
(022) (311) (222)
PDC-1500
PDC-1300 PDC-1100
10
20
30
40 50 60 2θ (degrees)
70
80
90
Fig. 3. XRD patterns of the PDC SiC powder prepared at various pyrolysis temperatures
(111) (002)
(022)
(111) (002)
(311)
a
b
0 h HT
F-1100 HT
F-1300
25 h
F-1100
0h
untreated fibre
10
20
(311)
50 h
Intensity
Intensity
F-1500
(022)
30
untreated fibre
40 50 60 2θ (degrees)
(111) (002)
(022)
70
80
90
10
20
30
40 50 60 2θ (degrees)
(111) (002)
(311)
c
d
F-1300 HT
F-1500 HT
(022)
70
80
90
80
90
(311)
50 h
Intensity
Intensity
50 h 25 h
25 h
0h
0h
untreated fibre
10
20
30
untreated fibre
40 50 60 2θ (degrees)
70
80
90
10
20
30
40 50 60 2θ (degrees)
70
Fig. 4. XRD patterns of SiC fibres treated under various conditions: simulated preparation process condition (a) and heat-treated F-1100 (b), F-1300 (c), and F-1500 (d). HT: heat treatment.
Monofilament tensile strength (GPa)
3.0
Untreated fibre
State 1 State 2
2.5 2.0 1.5 1.0 F-1100
F-1300
F-1500
Fig. 5. Monofilament tensile strength of the SiC fibres annealed under each preparation condition (state 1) and heat-treated for 50 h (state 2)
(022) (311) (222)
(111) (002)
c
0 h HT
CMC-1300
CMC-1500
Intensity
Intensity
(111) (002)
a
CMC-1300
(022) (311) (222)
50 h
25 h CMC-1100
10
20
0h
30
40 50 60 2θ (degrees)
80
90
10
20
30
(022) (311) (222)
40 50 60 2θ (degrees) (111) (002)
c
d
CMC-1300
CMC-1500 HT
Intensity
Intensity
(111) (002)
70
50 h
25 h
0h
0h
20
30
40 50 60 2θ (degrees)
70
80
90
80
90
(022) (311) (222)
50 h
25 h
10
70
10
20
30
40 50 60 2θ (degrees)
70
80
90
Fig. 6. XRD patterns of the SiCf/SiC composites fabricated at various temperatures (a) and heat-treated CMC-1100 (b), CMC-1300 (c), and CMC-1500 (d) samples. HT: heat-treated.
Crystallite size (nm)
40
CMC-1100 CMC-1300 CMC-1500
30
20
10
0 0
25 Exposure time (h)
50
Fig. 7. Evolution of SiC crystallite size in the SiCf/SiC composites as a function of exposure time during the heat treatment at 1300 ℃
Mass loss rate (%)
1.5 a
1.0
0.5
CMC-1100 CMC-1300 CMC-1500
0.0 0
25 Exposure time (h)
50
Fig. 8. Evolution of mass loss rate for the SiCf/SiC composites as a function of exposure time during the heat treatment at 1300 ℃
11
Open porosity (%)
10 9 8 7 6 CMC-1100 CMC-1300 CMC-1500
5 4 0
25 Exposure time (h)
50
Fig. 9. Evolution of open porosity for the SiCf/SiC composites as a function of exposure time during the heat treatment at 1300 ℃
Fig. 10. Transverse surfaces of SiCf/SiC composites before and after heat treatment for 50 h: the low-magnification images for the as-fabricated CMC-1300 (a) and the heat-treated CMC-1300 (b);
the high-magnification images in the fibre bundle for CMC-1100 (c), CMC-1300 (d), and CMC-1500 (e), and (1, 2) for the as-fabricated sample and the heat-treated sample, respectively.
Si
Au
Counts
C
0.5
1.0
1.5
2.0 2.5 keV
3.0
3.5
4.0
Fig. 11. EDS curve of the particles shown in Fig.10 (c-2)
800 500 Stress (MPa)
Stress (MPa)
600
400
As-fabricated CMC-1100 CMC-1300 CMC-1500
200
0 0.000
0.005
0.010
400 300 200 Heat-treated CMC-1100 CMC-1300 CMC-1500
100
0.015
Strain
0 0.000
0.005
0.010 Strain
0.015
0.020
Fig. 12. Typical stress-strain curves of the as-fabricated and heat-treated (50 h) SiCf/SiC composites
Fig. 13. Fracture morphology of as-fabricated samples: CMC-1100 (a), CMC-1300 (b), and CMC-1500 (c), and the heat-treated (50 h) samples: CMC-1100 (d), CMC-1300 (e), and CMC-1500 (f)
S0272-8842(19)32616-1
DOI:
https://doi.org/10.1016/j.ceramint.2019.09.087
Reference:
CERI 22854
To appear in:
Ceramics International
Received Date: 30 June 2019 Revised Date:
23 August 2019
Accepted Date: 9 September 2019
Please cite this article as: L.-Y. Wang, R.-Y. Luo, G.-Y. Cui, Effect of pyrolysis temperature on the mechanical evolution of SiCf/SiC composites fabricated by PIP, Ceramics International (2019), doi: https://doi.org/10.1016/j.ceramint.2019.09.087. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Ltd.
Effect of pyrolysis temperature on the mechanical evolution of SiCf/SiC composites fabricated by PIP Lian-Yi Wang1, Rui-Ying Luo*,1, Guang-Yuan Cui1 (1 School of Physics and Nuclear Energy Engineering, Beihang University, Beijing, 100191, China * Corresponding author. E-mail address: [email protected] (R.-Y. Luo). Tel. & Fax: +86-010-82338267)
Abstract Three types of SiCf/SiC composites with a four-step three-dimensional SiC fibre preform and pyrocarbon interface fabricated via precursor infiltration and pyrolysis at 1100
, 1300
, and 1500
were heat-treated at 1300
under argon atmosphere for
50 h. The effects of the pyrolysis temperature on the microstructural and mechanical properties of the SiCf/SiC composites were studied. With an increase in the pyrolysis temperature, the SiC crystallite size of the as-fabricated composites increased from 3.4 to 6.4 nm, and the flexural strength decreased from 742 ± 45 to 467 ± 38 MPa. After heat treatment, all the samples exhibited lower mechanical properties, accompanied by grain growth, mass loss, and the formation of open pores. The degree of mechanical degradation decreased with an increase in the pyrolysis temperature. The composites fabricated at 1500
exhibited the highest property retention rates
with 90% flexural strength and 98% flexural modulus retained. The mechanism of the mechanical evolution after heat treatment was revealed, which suggested that the thermal stability of the mechanical properties is enhanced by the high crystallinity of the SiC matrix after pyrolysis at higher temperatures. Key words: SiCf/SiC composite; pyrolysis temperature; microstructure; mechanical properties; heat treatment.
1 Introduction Silicon-carbide-fibre-reinforced silicon carbide matrix (SiCf/SiC) composites are
receiving increasing attention as potential candidates for high-temperature components (e.g. in turbine engines) owing to their excellent properties, including low density, excellent high-temperature mechanical performance with good mechanical damage tolerance (which enables them to overcome the inherent brittleness of monolithic SiC), exceptional oxidation and corrosion resistance, and low thermal expansion coefficient [1–6]. Various processing routes are used to fabricate SiCf/SiC composites, including chemical vapour infiltration (CVI) [7,8], precursor infiltration and pyrolysis (PIP) [9–11], liquid silicon infiltration [12,13], and the nano-infiltration and transient eutectic phase process [14,15]. Among these routes, the PIP process, in which a fibre preform is densified using the precursor-derived ceramic (PDC) SiC matrix, has been extensively studied owing to the designability of the polymer precursor molecules, lower pyrolysis temperature, low cost, and ease of processing [16–19]. However, the thermal stability of PDC SiC significantly influences the application of PIP SiCf/SiC composites at elevated temperatures [20–22]. Studies have demonstrated that the changes in the microstructure of the PDC SiC bulk material, which consists of nanocrystalline domains of SiC, amorphous SiOxCy, and carbon impurities, depend mainly on the processing and heat-treatment conditions (specifically, pyrolysis temperature, and/or exposure time). The SiC crystallite size tended to increase, accompanied by a decomposition of amorphous SiOxCy into SiC nanocrystalline domains and gaseous SiO and CO. Further, numerous cracks and pores were formed owing to volume shrinkage [23–26]. Moreover, the microstructural changes exhibited similar trends with an increase in the pyrolysis temperature and/or time. Rahman et al. [27] studied the structural changes in PDC SiC as a function of the pyrolysis temperature. The results revealed that the crystallite size of the SiC nanocrystalline domain increased from 6 to 22 nm as the pyrolysis temperature was increased from 900 to 1400
. Furthermore, the degree of crystallinity increased with
an increase in the pyrolysis temperature. Poerschke et al. [28] observed an increase in the crystallite size and crystallinity during the pyrolysis of precursor-derived SiC ceramics. The SiC samples fabricated at 800
were initially amorphous. The
crystallite size of the nanocrystalline SiC and the degree of crystallinity increased to 6 nm and 68%, respectively, after 100 h of annealing at 1300
, whereas the crystallite
size and degree of crystallinity increased to 124 nm and 80%, respectively, after 100 h of annealing at 1500
. When PDC SiC is used as the matrix in composites, both the
pyrolysis-temperature-dependent microstructure and the microstructural evolution at elevated temperatures significantly affect the properties of the composites. Therefore, to determine the applicability of SiCf/SiC composites fabricated via PIP under thermal conditions, extensive characterisation of the microstructural and mechanical evolution during treatment at high temperatures, especially as a function of the processing conditions, is required. However, only limited data are available in the literature. In this study, SiCf/SiC composites with a pyrocarbon (PyC) interface fabricated via PIP at various temperatures between 1100 °C and 1500
were heat-treated at 1300
under Ar atmosphere. The pyrolysis-temperature-dependent microstructural and mechanical evolution of the as-fabricated and heat-treated composites were studied. The characteristics of the PDC SiC bulk material fabricated under various temperatures and of SiC fibres heat-treated under different conditions were studied as controls. The performance evolution of the SiC fibres during preparation and heat treatment was evaluated. Accordingly, the mechanism of the mechanical evolution after heat treatment was revealed.
2. Materials and experimental procedure 2.1 Sample fabrication The SiCf/SiC composites examined in this study were fabricated via PIP, reinforced with Amosic SiC fibres provided by Xiamen University, China, and precoated with a PyC interphase. The fabrication process is shown in Fig. 1. SiC fibres were used to prepare the four-step three-dimensional fibre preform with a fibre volume fraction of ~0.4. Before the PIP process, a PyC layer was deposited on the preform via CVI. Propylene (C3H6) was applied for the PyC deposition at 850
and 2 kPa for 10 h. After interphase deposition, the SiC matrix was prepared to densify the preform via PIP using a 50 wt% polycarbosilane (PCS, provided by the National University of Defense Technology, China)/xylene solution as the precursor. The coated fibre preforms were dipped into the solution under vacuum for more than 2 h. Subsequently, the preforms were pyrolysed at 1100 °C, 1300 °C, or 1500
for 1
h under a protective Ar atmosphere. The PIP procedure was repeated until the weight increase of the composite was less than 1%. The samples are denoted as CMC-1100, CMC-1300, and CMC-1500, where CMC indicates the SiCf/SiC composite and the numbers indicate the corresponding pyrolysis temperatures. The composites were then machined into rectangular bars measuring approximately 40 × 4 × 3 mm3 for tests. A PDC SiC bulk material was prepared via the same pyrolysis process used for the composites, where PCS was pyrolysed at 1100 °C, 1300 °C, or 1500
for 1 h under a
protective Ar atmosphere. The mass of the PCS before and after pyrolysis was measured
five
times
using
an
electronic
balance
(resolution:
0.1
mg).
The ceramic yields of PCS pyrolysed under various temperatures were calculated as
ω% =
m - m0 × 100% m0
,
(1)
where m is the average mass of the sample after pyrolysis and m0 is the average initial mass of the sample. Subsequently, the product was crushed to obtain particles of sizes less than 100 µm. The samples are denoted as PDC-1100, PDC-1300, and PDC-1500, where PDC indicates the PDC SiC powder and the numbers indicate the corresponding pyrolysis temperatures.
2.2 Heat treatment The SiC fibres were annealed at 1100 °C, 1300 °C, or 1500
under a protective
Ar atmosphere for the entire PIP holding time; these samples are denoted as F-1100, F-1300, and F-1500, respectively. The objective was to obtain fibres in a similar state to that of the as-fabricated composites. The as-received fibres and composites were moved to a high-temperature horizontal tube furnace (GSL-1700X, Ke Jing Materials Technology Co. Ltd., China) and heated to 1300
under flowing Ar (0.1 L/min). The
fibres and composites were characterised after 25 h and 50 h of annealing.
2.3 Mechanical property tests A monofilament tensile test of the SiC fibres was conducted. A schematic diagram of the sample shape and size is shown in Fig. 2. A single SiC fibre of length ~150 mm was separated from the fibre bundle and glued to a supporting paper for ~30 mm of the length at both ends. Subsequently, the fibre was loaded at the centre of the unsupported length of 90 mm at a loading speed of 0.1 mm/min. More than 10 samples were analysed. The flexural strength and modulus of the composites before and after 50 h of heat treatment were characterised using a three-point bending test. The dimensions of the samples were 40 × 4 × 3 mm3, and they were tested at a span of 30 mm and a loading speed of 0.5 mm/min. Five samples of each type were analysed, and the average values of the five samples are reported.
2.4 Microstructural characterisation Archimedes’ principle was employed to measure the open porosity of the composites before and after heat treatment. The mass change of the composites during the heat-treatment processing was measured using an electronic balance. Five samples of each type were analysed for both tests, and the average values of the five samples are reported. The phases of the SiC fibre, PDC SiC powder, and SiCf/SiC composites treated under various conditions were analysed using X-ray diffraction (XRD, D/Max 2500PC, Rigaku, Japan) with Cu Kα radiation (λ = 0.154 nm) and 2θ = 10°–90°. The scan rate was 4°/s. Further, the crystallite size of the (111) plane for each sample was calculated using Scherrer’s formula as follows: D hkl =
kλ βcos θ
,
(2)
where Dhkl is the crystallite size, k is the Scherrer constant, β is the peak width at half-height, and θ is the Bragg diffraction angle. The transverse surfaces of the as-fabricated and heat-treated samples and the
microstructures of the as-fractured samples were observed using scanning electron microscopy (SEM, S4800, Hitachi, Japan). The elemental analyse of the transverse surface was performed with an energy-dispersive X-ray spectroscopy (EDS, equipped with SEM).
3 Results and discussion 3.1 Characterisation of the PDC SiC powder fabricated at various pyrolysis temperatures Fig. 3 shows the XRD patterns of the PDC SiC powder prepared at various pyrolysis temperatures. β-SiC was the only crystalline phase detected using XRD, and the peaks at 35.7°, 60.0°, and 71.8° correspond to the (111), (220), and (311) planes of
β-SiC crystalline grains, respectively. All these samples exhibited characteristic broad halo peaks with diffuse intensity, which indicated an amorphous/nanocrystalline structure. These peaks became sharper and more intense when the pyrolysis temperature was increased to 1500
. Two new peaks at 41.4° and 75.5°, which
correspond to the (002) and (022) planes, respectively, were observed for PDC-1500. These findings suggest that the amorphous structure was converted to nanocrystalline domains with an increase in the pyrolysis temperature. Furthermore, the crystallite sizes of the (111) plane for the three samples, as analysed using Scherrer's formula, are shown in Table 1. The sizes also increased with an increase in the pyrolysis temperature, rising from 2.3 to 6.8 nm, further indicating the development of the nanocrystalline structure in the samples. The above results are validated by a comparison with test data from the literature in Table 1. There was also a negative correlation between the ceramic yield and pyrolysis temperature as shown in Table 1. The relatively high ceramic yield decreased from 68 ± 5% to 57 ± 6% when the temperature was increased from 1100 to 1500
. The mass
loss is attributed to the formation of gaseous SiO and/or CO as the SiOxCy decomposed.
3.2 Microstructural and mechanical evolution of SiC fibres The SiC fibres were annealed at 1100 °C, 1300 °C, or 1500
for the entire PIP
holding time to evaluate the effects of the preparation process on the microstructure and mechanical properties of the SiC fibres, which are denoted as F-1100, F-1300, and F-1500, respectively. Subsequently, the as-received fibres were heat-treated at 1300
under Ar atmosphere. The microstructure and mechanical properties of the
fibres without heat treatment were measured as a control. The XRD patterns of the SiC fibres annealed under each preparation condition are shown in Fig. 4(a), and the crystallite sizes of the (111) plane for these samples are listed in Table 2 (samples with 0 h of heat treatment). All these samples exhibited characteristic broad halo peaks. Further, there were three main peaks at 35.7°, 60.0°, and 71.8°, which correspond to the (111), (220), and (311) planes of β-SiC crystalline grains, and a weak peak at 41.4° that corresponds to the (002) plane of β-SiC crystalline grain. The average grain size of the fibres treated with various preparation processes was 5.7–6.1 nm for the untreated fibre, F-1100, and F-1300, and it increased to 7.1 nm for F-1500, indicating that grain growth occurred in the SiC fibres when the pyrolysis temperature was increased to 1500 After heat treatment at 1300
.
for various times [the XRD patterns shown in Fig.
4(b)–(d)], the grain size was almost constant for each type of sample as listed in Table 2. The average grain size was 5.7–6.8 nm for the F-1100 and F-1300 fibres and was higher for the F-1500 fibres (7.1–8.6 nm). These observations suggested that the microstructure of the SiC fibres was insensitive to temperature during heat treatment at 1300
.
The monofilament tensile properties of the SiC fibres annealed under each preparation condition were determined and are shown in Table 2 and Fig. 5. The monofilament tensile strength of the untreated fibres was also determined for reference. The tensile strength decreased slightly, from 2.74 ± 0.32 GPa (untreated fibres) to 2.69 ± 0.39 and 2.46 ± 0.34 GPa after the fibres were annealed during preparation at 1100 °C and 1300
, respectively. When the pyrolysis temperature was
increased to 1500
, the tensile strength decreased significantly to 1.51 ± 0.31 GPa,
indicating that the properties of the SiC fibres were degraded at the relatively high pyrolysis temperature of ~1500 After heat treatment at 1300
. for 50 h, the monofilament tensile properties of the
SiC fibres were tested and are shown in Table 2 and Fig. 5. The strength of the heat-treated F-1100, F-1300, and F-1500 fibres decreased to 2.25 ± 0.68, 2.04 ± 0.47, and 1.33 ± 0.51 GPa (decreases of ~16%, ~17%, and ~12%), respectively. The relatively weak effect of temperature on the decrease in tensile strength indicated that the SiC fibres were thermally stable at the examined temperatures.
3.3 Microstructural evolution of the SiCf/SiC composites 3.3.1 Dependence of microstructural evolution of as-fabricated SiCf/SiC composites on pyrolysis temperature The microstructural evolution of the as-fabricated composites (the samples with 0 h of heat treatment) is shown in Table 3. With an increase in the pyrolysis temperature, the density of the composites increased slightly from 2.21 ± 0.03 to 2.30 ± 0.02 g/cm3. The porosity of the composites exhibited a decreasing trend, from 5.6 ± 0.4% (CMC-1100) and 5.7 ± 0.3% (CMC-1300) to 5.1 ± 0.5% (CMC-1500). The XRD patterns of the as-fabricated SiCf/SiC composites are shown in Fig. 6(a), which are similar to those of the PDC SiC samples fabricated at each temperature. Three broad halo peaks appeared at 35.7°, 60.0°, and 71.8° for the CMC-1100 and CMC-1300 samples, and two new peaks appeared at 41.4° and 75.5° when the pyrolysis temperature was increased to 1500
. The crystallite sizes of the (111) plane
for the three as-fabricated composites, as analysed using Scherrer’s formula, are shown in Table 3. The crystallite sizes of the composites increased with an increase in the pyrolysis temperature, i.e. they were 3.4, 4.8, and 6.4 nm for CMC-1100, CMC-1300, and CMC-1500, respectively. Owing to the insignificant change in the grain sizes of the SiC fibres during preparation, the grain growth of the composites is mostly attributed to the PDC SiC matrix. The differences in the crystallite sizes of the composites compared with those of both the PDC SiC powder and the SiC fibres
might be associated with the strong dependence of the crystallisation behaviour of PDC ceramics on the local microstructure [28].
3.3.2 Dependence of microstructural evolution of heat-treated SiCf/SiC composites on pyrolysis temperature The microstructures of the SiCf/SiC composites after heat treatment at 1300 were characterised. The mass loss, open porosity, and crystallite sizes of the composites after heat treatment under various conditions are listed in Table 3. The phases of the composites after heat treatment under each condition were analysed using XRD as shown in Fig. 6(b)–(d). With an increase in the exposure time at 1300
, the aforementioned peaks became sharper and more intense. Furthermore,
the peaks at 41.4° and 75.5° appeared after 25 h of heat treatment for CMC-1100 and CMC-1300. In addition, the crystallite sizes after various exposure times are shown in Fig. 7, where the as-fabricated composite (0 h of treatment) serves as a control. The crystallite sizes increased monotonically with an increase in the exposure time. After heat treatment for 50 h, the crystallite size increased from 3.4, 4.8, and 6.4 nm to 38.8, 25.9, and 17.7 nm for CMC-1100, CMC-1300, and CMC-1500, respectively. These findings suggest that the SiC crystallites became coarser and the amorphous structure was converted to nanocrystalline domains with an increase in the exposure time. The evolution of mass loss rates with an increase in the exposure time is plotted in Fig. 8. Each sample lost mass. The mass loss rates for CMC-1100, CMC-1300, and CMC-1500 after heat treatment for 50 h were 1.37 ± 0.12%, 0.61 ± 0.12%, and 0.28 ± 0.06%, respectively. Those phenomena are attributed to the decomposition of SiOxCy accompanied by the formation of gas (SiO and/or CO). As shown in Fig. 9, the open porosity for each sample, measured using Archimedes’ principle, exhibited an increasing trend, and it rose to 9.8 ± 0.4%, 8.2 ± 0.5%, and 6.4 ± 0.4% for CMC-1100, CMC-1300, and CMC-1500, respectively, after heat treatment for 50 h. The effect of heat treatment on the open pores was observed using SEM, with the transverse surfaces of the three heat-treated samples (the heat-treated samples of CMC-1100, CMC-1300, and CMC-1500) shown in Fig. 10
and the ones of the corresponding untreated samples considered as the control. The low-magnification images of the transverse surfaces for each sample were similar. The structures of the as-fabricated and heat-treated CMC-1300 samples are shown in Fig. 10 (a) and (b), respectively, as representative structures. The open pores of the samples consist of large pores between bundles and small pores between the fibres in a fibre bundle. Comparing the structures of the transverse surfaces before and after heat treatment, it can be observed that the heat-treatment process had a more significant effect on the pores between the fibres in a fibre bundle. The three untreated samples exhibited a high density in the fibre bundles with the pores between fibres filled by the matrix (presented in Fig. 10 (c-1), (d-1), and (e-1)). In contrast, a large number of micro-cracks were formed in the fibre bundles after the heat treatment, and penetrating cracks were observed in the heat-treated CMC-1100 samples. These cracks, which are attributed to volume shrinkage caused by grain growth and/or decomposition, resulted in an increase in the open porosity. Comparing the structures of the three heat-treated samples presented in Fig. 10 (c-2), (d-2), and (f-2), it can be observed that the structures of CMC-1100 were different from those of the other samples, indicating that SiC particles (as confirmed using EDS in Fig. 11) with a regular geometric shape were attached to the sample surface. These structures indicated that CMC-1100 could easily form SiC with higher crystallinity after the heat treatment. As the SiC fibres maintained a stable microstructure during heat treatment, the microstructural evolution of the PDC SiC matrix resulted in the dramatic grain growth of the composites. Note that the CMC-1100 samples exhibited greater grain growth and showed both the highest mass loss and the highest open porosity after 50 h of heat treatment. These findings confirm that the microstructural evolution of CMC-1100 was more intense than that of the other samples and indicate that the PDC SiC matrix fabricated under lower pyrolysis temperatures was less thermally stable.
3.4 Mechanical evolution of the SiCf/SiC composites 3.4.1 Dependence of mechanical evolution of as-fabricated SiCf/SiC composites on pyrolysis temperature The mechanical properties of the as-fabricated SiCf/SiC composites prepared at various pyrolysis temperatures are listed in Table 4. The typical stress–strain curves and the fracture morphology of the as-fabricated SiCf/SiC composites are shown in Fig. 12 and Fig. 13 (a)–(c), respectively. For the as-fabricated samples prepared at different pyrolysis temperatures, the mechanical behaviours indicate the pseudo-plastic fracture mode, where three regions appear in the typical stress–strain curves. The stress first increased almost linearly with an increase in the strain in the elastic region, where only elastic elongation of the constituents and marginal interfacial sliding occurred. The mechanical behaviour then entered the pseudo-plastic region, where the slope decreased with an increase in the loads, and matrix crack initiation and interfacial debonding occurred. Finally, the stress decreased after the initial failure in the breakdown region, where fibre breakdown accompanied by fibre pull-out occurred, as shown on the fracture surfaces in Fig. 13 (a)–(c). As the pyrolysis temperature was increased from 1100 to 1300
, the flexural
strength of the as-fabricated samples decreased slightly, from 742 ± 45 to 649 ± 51 MPa. However, when the pyrolysis temperature was increased to 1500
, the flexural
strength declined sharply to 467 ± 38 MPa. Conversely, the modulus increased with an increase in the pyrolysis temperature, rising from 99 ± 7 GPa(CMC-1100) to 108 ± 4 GPa (CMC-1500). The decrease in strength is attributed to the degradation of the properties of the SiC fibres at relatively high pyrolysis temperatures, especially at ~1500
. By contrast,
the increase in modulus could be understood according to the model of the modulus for composites, as follows:
Ec = EfVf + EmVm
(3)
where Ec, Ef, and Em are the moduli of the composite, fibres, and matrix, respectively;
Vf and Vm are the volume contents of the fibres and matrix, respectively. According to previous studies [18,23,27], crystalline PDC SiC growth accompanied by increasing modulus of PDC SiC occurs with an increase in the pyrolysis temperature. Therefore, the modulus of the composites increases, according to the model in Eq. (3).
3.4.2 Dependence of mechanical evolution of heat-treated SiCf/SiC composites on pyrolysis temperature The mechanical properties of the heat-treated SiCf/SiC composites are listed in Table 4, and the fracture morphology and typical stress–strain curves are shown in Fig. 12 and Fig. 13 (d)–(f), respectively. For the heat-treated samples, the mechanical parameters still indicate pseudo-ductile fracture mode, as shown in Fig. 13 (d)–(f). However, a large yield before fracture appears in the stress–strain curves of the heat-treated CMC-1100 samples, which differ from the curves of the heat-treated CMC-1300 and CMC-1500 samples. This finding indicates lower load-bearing and load-transfer capacity of the matrix in CMC-1100. All the samples exhibited a degradation of the mechanical properties after heat treatment. The strength and modulus decreased to 371 ± 21, 481 ± 61, and 421 ± 39 MPa and 62 ± 6, 88 ± 5, and 106 ± 6 GPa for the CMC-1100, CMC-1300, and CMC-1500 samples, respectively, after heat treatment for 50 h. The heat-treated CMC-1100 samples exhibited the most severe property degradation, retaining only 51% of their flexural intensity and 63% of their flexural modulus. In contrast, the mechanical properties of the heat-treated CMC-1500 samples were less degraded, as they retained 90% of their flexural strength and 98% of their flexural modulus. Further, the highest flexural strength after pyrolysis among the three types of samples was 481 ± 61 MPa for CMC-1300 (retention of 74%). The change in the properties is attributed to the changes in the microstructure. As both the microstructure and mechanical properties of the SiC fibres remained stable during heat treatment at 1300
, the lower mechanical properties of the composites
are attributed to matrix degradation. The volume shrinkage and cracking in the matrix caused by PDC SiC matrix crystallisation and decomposition yield a more porous
structure, which weakens both the load-bearing and load-transfer capacities. Moreover, the microstructural transformation is greater at lower pyrolysis temperatures. These dual functions would first generate mechanical failure of the matrix and then limit the bearing capacity of the fibres in the composites. Consequently, the properties of the composites are degraded. Furthermore, the thermal stability is enhanced by the high crystallinity of the SiC matrix fabricated at higher pyrolysis temperatures.
4 Conclusion SiCf/SiC composites with a four-step three-dimensional SiC fibre preform and PyC interface fabricated via PIP at various pyrolysis temperatures (1100 °C, 1300 °C, and 1500
) were heat-treated at 1300
under Ar atmosphere. The dependence of the
microstructural and mechanical evolution of the as-fabricated and heat-treated composites on the pyrolysis temperature was studied. As the pyrolysis temperature was increased from 1100 to 1500
, the flexural strength of the as-fabricated
composites decreased from 742 ± 45 to 467 ± 38 MPa, whereas the modulus and SiC crystallite size increased from 99 ± 7 to 108 ± 4 GPa and from 3.4 to 6.4 nm, respectively. The decrease in strength is attributed to degradation of the properties of the SiC fibres during preparation. Further, the enhanced grain growth in the PDC SiC matrix at higher pyrolysis temperatures caused the increase in the modulus and crystallite size of the composites. After heat treatment for 50 h at 1300
, all the
samples exhibited a degradation of the mechanical properties, accompanied by grain growth, mass loss, and open pore formation. The flexural strength of the CMC-1100, CMC-1300, and CMC-1500 samples after 50 h of heat treatment decreased to 371 ± 21, 481 ± 61, and 421 ± 39 MPa, and the SiC crystallite size increased to 38.8, 25.9, and 17.7 nm, respectively. The heat-treated CMC-1100 samples exhibited the most dramatic microstructural transformation and property degradation, retaining only 51% of their flexural strength and 63% of their flexural modulus. By contrast, the CMC-1500 composites maintained the smallest SiC crystallite size and exhibited the least mechanical property degradation, retaining 90% of their flexural strength and
98% of their flexural modulus. As both the microstructure and mechanical properties of the SiC fibres remained stable during heat treatment at 1300
, the microstructural
evolution and mechanical property degradation of the composites are attributed to the degradation of the PDC SiC matrix (including SiC grain growth and decomposition of amorphous SiOxCy into nanocrystalline SiC domains and gaseous SiO and CO). Further, the thermal stability was enhanced by the high crystallinity of the SiC matrix fabricated at higher pyrolysis temperatures.
Acknowledgements This work was supported by the National Natural Science Foundation of China (No. 21071011).
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Tables Table 1. Evolution of crystallite size and ceramic yield with an increase in the pyrolysis temperature
Sample
Pyrolysis temperature ( )
Crystallite size (nm)
Ceramic yield(%)
PDC-1100
1100
2.3
68±5
PDC-1300
1300
4.1
66±7
PDC-1500
1500
6.8
57±6
1100
2
-
1300
3
-
1500
7
-
[28]
Table 2. Crystallite sizes and tensile strength of SiC fibres for various exposure times
Exposure time
Crystallite size
Tensile strength
(h)
(nm)
(GPa)
-
5.8
2.74±0.32
0
5.7
2.69±0.39
25
6.0
-
50
6.2
2.25±0.68
0
6.1
2.46±0.34
25
6.3
-
50
6.8
2.04±0.47
0
7.1
1.51±0.31
25
8.5
-
50
8.6
1.33±0.51
Sample Untreated fibre
F-1100
F-1300
F-1500
Table 3. Microstructural evolution of the SiCf/SiC composites for various pyrolysis temperatures and exposure times
Pyrolysis temperature
Exposure time
Density
Open porosity
Crystallite size
Mass loss
( )
(h)
(g/cm-3)
(%)
(nm)
(%)
0
2.21±0.03
5.6±0.4
3.4
0
25
-
8.7±0.3
33.8
1.13±0.21
50
-
9.8±0.4
38.8
1.37±0.12
0
2.25±0.04
5.7±0.3
4.8
0
25
-
7.6±0.2
21.6
0.53±0.08
50
-
8.2±0.5
25.9
0.61±0.12
0
2.30±0.02
5.1±0.5
6.4
0
25
-
6.1±0.4
17.3
0.23±0.02
50
-
6.4±0.4
17.7
0.28±0.06
Sample
CMC-1100
CMC-1300
CMC-1500
1100
1300
1500
Table 4. Mechanical properties of as-fabricated and the heat-treated SiCf/SiC composites
As-fabricated samples
Heat-treated samples (50h)
Flexural
Flexural
Flexural
Residual
Flexural
Residual
strength
modulus
strength
strength ratio
modulus
modulus ratio
(MPa)
(GPa)
(MPa)
(%)*
(GPa)
(%)*
CMC-1100
742±45
99±7
371±21
51
62±6
63
CMC-1300
649±51
107±7
481±61
74
88±5
82
CMC-1500
467±38
108±4
421±39
90
106±6
98
Sample
*Calculated from the average data for the samples before heat treatment
Figures
Fig. 1. Flowchart of SiCf/SiC composite fabrication process
Fig. 2. Schematic of monofilament tensile test sample
Intensity
(111) (002)
(022) (311) (222)
PDC-1500
PDC-1300 PDC-1100
10
20
30
40 50 60 2θ (degrees)
70
80
90
Fig. 3. XRD patterns of the PDC SiC powder prepared at various pyrolysis temperatures
(111) (002)
(022)
(111) (002)
(311)
a
b
0 h HT
F-1100 HT
F-1300
25 h
F-1100
0h
untreated fibre
10
20
(311)
50 h
Intensity
Intensity
F-1500
(022)
30
untreated fibre
40 50 60 2θ (degrees)
(111) (002)
(022)
70
80
90
10
20
30
40 50 60 2θ (degrees)
(111) (002)
(311)
c
d
F-1300 HT
F-1500 HT
(022)
70
80
90
80
90
(311)
50 h
Intensity
Intensity
50 h 25 h
25 h
0h
0h
untreated fibre
10
20
30
untreated fibre
40 50 60 2θ (degrees)
70
80
90
10
20
30
40 50 60 2θ (degrees)
70
Fig. 4. XRD patterns of SiC fibres treated under various conditions: simulated preparation process condition (a) and heat-treated F-1100 (b), F-1300 (c), and F-1500 (d). HT: heat treatment.
Monofilament tensile strength (GPa)
3.0
Untreated fibre
State 1 State 2
2.5 2.0 1.5 1.0 F-1100
F-1300
F-1500
Fig. 5. Monofilament tensile strength of the SiC fibres annealed under each preparation condition (state 1) and heat-treated for 50 h (state 2)
(022) (311) (222)
(111) (002)
c
0 h HT
CMC-1300
CMC-1500
Intensity
Intensity
(111) (002)
a
CMC-1300
(022) (311) (222)
50 h
25 h CMC-1100
10
20
0h
30
40 50 60 2θ (degrees)
80
90
10
20
30
(022) (311) (222)
40 50 60 2θ (degrees) (111) (002)
c
d
CMC-1300
CMC-1500 HT
Intensity
Intensity
(111) (002)
70
50 h
25 h
0h
0h
20
30
40 50 60 2θ (degrees)
70
80
90
80
90
(022) (311) (222)
50 h
25 h
10
70
10
20
30
40 50 60 2θ (degrees)
70
80
90
Fig. 6. XRD patterns of the SiCf/SiC composites fabricated at various temperatures (a) and heat-treated CMC-1100 (b), CMC-1300 (c), and CMC-1500 (d) samples. HT: heat-treated.
Crystallite size (nm)
40
CMC-1100 CMC-1300 CMC-1500
30
20
10
0 0
25 Exposure time (h)
50
Fig. 7. Evolution of SiC crystallite size in the SiCf/SiC composites as a function of exposure time during the heat treatment at 1300 ℃
Mass loss rate (%)
1.5 a
1.0
0.5
CMC-1100 CMC-1300 CMC-1500
0.0 0
25 Exposure time (h)
50
Fig. 8. Evolution of mass loss rate for the SiCf/SiC composites as a function of exposure time during the heat treatment at 1300 ℃
11
Open porosity (%)
10 9 8 7 6 CMC-1100 CMC-1300 CMC-1500
5 4 0
25 Exposure time (h)
50
Fig. 9. Evolution of open porosity for the SiCf/SiC composites as a function of exposure time during the heat treatment at 1300 ℃
Fig. 10. Transverse surfaces of SiCf/SiC composites before and after heat treatment for 50 h: the low-magnification images for the as-fabricated CMC-1300 (a) and the heat-treated CMC-1300 (b);
the high-magnification images in the fibre bundle for CMC-1100 (c), CMC-1300 (d), and CMC-1500 (e), and (1, 2) for the as-fabricated sample and the heat-treated sample, respectively.
Si
Au
Counts
C
0.5
1.0
1.5
2.0 2.5 keV
3.0
3.5
4.0
Fig. 11. EDS curve of the particles shown in Fig.10 (c-2)
800 500 Stress (MPa)
Stress (MPa)
600
400
As-fabricated CMC-1100 CMC-1300 CMC-1500
200
0 0.000
0.005
0.010
400 300 200 Heat-treated CMC-1100 CMC-1300 CMC-1500
100
0.015
Strain
0 0.000
0.005
0.010 Strain
0.015
0.020
Fig. 12. Typical stress-strain curves of the as-fabricated and heat-treated (50 h) SiCf/SiC composites
Fig. 13. Fracture morphology of as-fabricated samples: CMC-1100 (a), CMC-1300 (b), and CMC-1500 (c), and the heat-treated (50 h) samples: CMC-1100 (d), CMC-1300 (e), and CMC-1500 (f)