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C–ZrC–SiC composites
Journal of Alloys and Compounds 645 (2015) 206–212
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Effect of carbon nanotubes on the toughness, bonding strength and thermal shock resistance of SiC coating for C/C–ZrC–SiC composites Qian-gang Fu ⇑, Lei Zhuang, He-jun Li, Lei Feng, Jun-yi Jing, Bi-yi Tan State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China
a r t i c l e
i n f o
Article history: Received 11 April 2015 Received in revised form 28 April 2015 Accepted 30 April 2015 Available online 8 May 2015 Keywords: C/C–ZrC–SiC composites SiC coating Carbon nanotube Toughness Bonding strength Thermal shock
a b s t r a c t In order to improve the toughness, interface bonding strength and thermal shock resistance of SiC coating for C/C–ZrC–SiC composites, carbon nanotubes (CNTs) were prepared by injection chemical vapor deposition and attempted as the reinforcement materials in the SiC coating. After incorporating CNTs, the hardness and elastic modulus of the SiC coating increased by 26.37% and 28.23%, respectively. The interface bonding strength between SiC coating and C/C–ZrC–SiC composites was enhanced by 53.31%. The mass loss of the SiC coated C/C–ZrC–SiC composites after thermal shock between 1773 K and room temperature for 15 times decreased from 5.98% to 1.98%. The incorporation of CNTs can effectively improve the toughness, interface bonding strength and thermal shock resistance of SiC coating due to the nanoscale toughening mechanism of CNTs by pullout, bridging and crack deflection. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Introducing ultra-high temperature ceramics (UHTCs), such as ZrC and SiC [1,2] into carbon/carbon (C/C) composites for the purpose of improving the ablation resistance has been highlighted for several years [3–5]. Unfortunately, C/C–UHTCs composites are lack of stability in air at elevated temperature, mainly due to the rapid oxidation and degradation of carbon, carbon fibers and ceramic matrix [6]. Anti-oxidation coating, typically SiC [7,8], has thus been employed to protect C/C–UHTCs composites from burn-off. Up to now, numbers of techniques have been explored to prepare SiC coating, such as slurry [9], pack cementation [10] and chemical vapor deposition (CVD) [11]. Among these techniques, CVD is an effective technique to produce dense and homogenous SiC coating with various shape [12]. The interface bonding between the coating and substrate is regarded as a key factor to determine the oxidation resistant ability of the coated specimens [12], but currently, there is still a challenge in improving the adhesion ability of CVD-SiC coating obtained by the decomposition of Methyltrichlorosilane (MTS) [13]. Additionally, the mismatch of coefficients of thermal expansion (CTEs) between SiC coating and substrates will lead to cracking in the coating, leaving diffusion paths for oxygen when the coated specimens exposed to an ⇑ Corresponding author. Tel.: +86 29 88494197; fax: +86 29 88495764. E-mail address: [email protected] (Q.-g. Fu). http://dx.doi.org/10.1016/j.jallcom.2015.04.223 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
oxidizing environment at high temperature [14]. The above-mentioned poor adhesion ability and CTEs mismatch between CVD-SiC coating and C/C–UHTCs limit their potential applications in extremely environment. Carbon nanotube (CNT) is a relatively new type of nano-structured material [15], which has attracted much attention as a potential reinforcement element for structural composites, due to their high specific surface area, exotic mechanical, thermal and electrical properties [16]. Recently, Wang et al. have testified that spreading multi-walled carbon nanotubes (MWCNTs) into C/UHTCs contributed to a 29.7% increase in the flexural strength and a 27.9% increase in the fracture toughness [17]. The main problem of CNTs as the reinforcement materials is associated with their natural tendency to agglomerate considerably due to their large surface areas and aspect ratios [18]. Injection chemical vapor deposition (ICVD) method is a process for effectively and uniformly fabricating high-purified CNTs [19]. Up to now, little research has been reported about using ICVD-CNTs as the toughening materials in CVD-SiC coating for C/C–UHTCs composites. In order to improve the interface adhesion ability and alleviating the mismatch of the CTEs between SiC coating and C/C–ZrC–SiC substrates, it is worth attempting to deposit CNTs on C/C–ZrC–SiC composites by ICVD and construct a CNT/SiC buffer layer. In this work, an ICVD technique was employed to deposit CNTs on the surface of C/C–ZrC–SiC substrates, and then the porous CNTs layer was filled with CVD-SiC coating. The microstructures of CNTs and the CNTs–SiC coating were investigated. The effect of CNTs on
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the toughness, interface bonding strength and thermal shock resistance of the CNTs–SiC coating for C/C–ZrC–SiC composites were discussed. 2. Experimental procedure T300 carbon fiber felts (0.45 g/cm3) were densified by chemical vapor infiltration (CVI) to obtain porous C/C composites with a density of 1.2 g/cm3. Xylene used as solvent, zirconium-containing polymer (PZC) (Institute of Process Engineering, Chinese Academy of Science, China) and polycarbosilane (PCS) (National University of Defense Technology, Changsha, China) were dissolved and mixed with a mass radio of 3:1 to form a homogenous solution, which was used as ZrC–SiC precursor. After 30 min immersion in the precursor under vacuum, composites were took out and dried in a drying oven at 373 K for 30–36 h. And then, the dried composites were undergone a pyrolysis process at 1773–1873 K for 2 h in a flowing Ar atmosphere. Repeat the infiltration-pyrolysis process for several times until the mass increase percentage of C/C–ZrC–SiC composites was no more than 1%. Cubic specimens (10 10 10 mm3) used as substrates were cut from bulk C/C–ZrC–SiC composites with a density of 2.15 g/cm3, and hand-abraded with 300-grit SiC paper and cleaned in an ultrasonic device with ethanol. CNTs were synthesized on the C/C–ZrC–SiC substrates using ICVD method under ambient pressure in a horizontal quartz reactor equipped with an electrical furnace. A mixture solution of 1.2 wt.% ferrocene, 76.8 wt.% ethanol and 22 wt.% ethylenediamine was fed into the reaction tube at a rate of 10 mL/h by a syringe pump. The reaction temperature was 1000–1058 K and the growth time was 15 min. After preparing the porous network of the CNTs, CVD was employed to obtain SiC coating on the C/C–ZrC– SiC substrates. H2 and Ar were used as the carrier gas and the diluent gas, respectively. MTS as the precursor was brought into the reaction chamber by bubbling H2. The flow rates of H2, MTS and Ar were 800 mL/min, 100 mL/min and 100 mL/min, respectively. The deposition temperature was 1373 K and the deposition time was 6 h. The fabrication process of CNTs–SiC coated C/C–ZrC–SiC composites is depicted in Fig. 1. The nanoindentation test was carried out using Nano-Indenter™ XP (MTS System Corp., USA) system with a diamond indenter on the well-polished cross-section of the upper part of the coating to measure the hardness and elastic modulus of the coating. The maximum indentation load was set at a constant value (400 mN). Load and depth were recorded continuously during the indentation process. The hardness and elastic were calculated from the load–depth curves using the Oliver and Pharr method [20]. Scratch tester (WS-2005 multi-functional tester. China) equipped with a C diamond pinhead (cone apex angle 120, tip radius 0.2 mm) was subjected to estimate the interface bonding strength between SiC coating and C/C–ZrC–SiC composites. Scratch test was carried out by applying a constantly changing load which ranged from 0 to 16 N during sliding on the 5 mm path. Thermal cycling test was performed between 1773 K and room temperature in air to investigate the thermal shock resistance of the coating. After oxidation for 5 min at 1773 K in a furnace, the coated specimens were taken out and cooled down to room temperature in air. An electronic balance with a sensitivity of ±0.1 mg was used to weigh the coated specimens. And then, the coated specimens were put into the furnace for the next cycle. Cumulative mass changes of the coated specimens after each thermal cycle were reported as a function of the thermal cycling time. The morphologies of CNTs and SiC coating were analyzed by a scanning electron microscopy (SEM, JSM-6700). The structure of CNTs were carried out by transmission electron microscopy (TEM, Tecnai F30G2). The graphitization degree of the substrates before and after grafting CNTs was investigated by Raman spectroscopy (Renishaw) using an inVia micro-Raman spectrometer with an Ar ion laser of 514.5 nm wavelength. The crystalline structures of SiC coating was analyzed by X-ray diffraction with Cu Ka radiation. (XRD, X Pert PRO, PANalytical, Almelo, The Netherlands).
3. Results and discussion 3.1. Microstructure of CNTs and SiC coating on C/C–ZrC–SiC substrates Fig. 2 presents the typical surface morphology of CNTs on C/C– ZrC–SiC substrates by ICVD. From Fig. 2a, it can be observed that
Coated specimen with CNTs
Specimen with CNTs
C/C-ZrC-SiC specimen ICVD
CVD
Coating
Fig. 1. The schematic diagram of preparing CNTs–SiC coated C/C–ZrC–SiC composites.
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C/C–ZrC–SiC substrates were uniformly and homogeneously covered with a 3D porous network involving numerous of random-oriented CNTs. These CNTs are generally several tens to one hundred micrometers in diameter with a curved structure, as shown in Fig. 2b. Fig. 3a displays the TEM observation, indicating that the CNTs by ICVD have a bamboo-like structure since the periodical fringes can be seen across the nanotubes. In addition, the nanotubes appear high purity without isolated amorphous carbon. High-magnification TEM image (Fig. 3b) of CNTs shows that both the nanotube walls and the inner caps are well graphitized, implying their high crystallinities. Raman spectra were applied to characterize the degree of graphitization of CNTs. There are two main bands in Raman spectra, which are the disorder-induced D band (about 1348 cm1) and the tangential G band (about 1575 cm1). The intensity ratios of D band and G band (ID/IG) were calculated. The lower the ID/IG is, the higher degree of graphitization the specimens have. The substrate without CNTs holds an ID/IG of 1.081, while the substrate covering with CNTs has the ID/IG of 0.972, as displayed in Fig. 3c. The decrease of ID/IG value suggests a higher graphitization degree after depositing CNTs. Fig. 4a presents the cross-section image of the CNTs–SiC coating before polishing. The inner CNT/SiC layer shows a relatively rough morphology, and abundant CNTs can be observed. This inner layer is about 15 lm. Fig. 4b exhibits the cross-section image of the CNTs–SiC coating after polishing. The whole coating is of a dense and homogeneous structure with a thickness of about 60 lm, bonding well with C/C–ZrC–SiC substrates. Fig. 4c displays the XRD analysis of SiC coating. There are three obvious peaks of 71.838°, 60.088° and 35.708°, which are in good agreement with the diffractions of (3 1 1), (2 2 0) and (1 1 1) crystalline planes of b-SiC [21]. SiC coating was realized by the decomposition of MTS based on the following reaction:
CH3 SiCl3 ðgÞ ! SiC ðsÞ þ 3HCl ðgÞ
ð1Þ
Surface image of SiC coating without CNTs is shown in Fig. 4d. A dense surface with a microcrack can be clearly found. When cooling from the fabricating temperature of 1373 K to room temperature, the mismatch of the CTEs between the coating and substrates would lead to some microcracks emerged on the surface inevitably. The breadth of the widest microcrack for the coating without CNTs is about 1.21 lm, while the widest microcrack of coating with CNTs is about 0.74 lm (Fig. 4e). The incorporation of CNTs contributes to a decrease of microcrack size. 3.2. Toughness of CNTs–SiC coating Fig. 5 shows the representative load–depth curves of the coating from nanoindentation. The maximum indentation depth of CNTs–SiC coating is lower than that of SiC coating with the same maximum indentation load (400 mN). This indicates that the hardness of the coating with CNTs is higher than that of the coating without CNTs. Additionally, the unloading curves of coating without and with CNTs could be fitted as unary linear regression equations, as followed:
y ¼ 0:767x 413:201
ð2Þ
y ¼ 0:802x 377:579
ð3Þ
According to Eqs. (2) and (3), the slopes of unloading curves of the coating without and with CNTs are 0.767 and 0.802, respectively, suggesting the incorporation of CNTs contributes to an increase in elastic modulus of the coating. The corresponding results are shown in Table 1. Before incorporating the CNTs, the hardness and elastic modulus of the coating are 32.16 ± 0.38 and 317.74 ± 17.34 GPa, respectively. After the incorporation of CNTs,
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Fig. 2. Surface SEM images of the 3D porous network of the CNTs on C/C–ZrC–SiC substrates: (a) low-magnification; and (b) high-magnification.
Fig. 3. TEM images of the as-synthesized CNTs: (a) low-magnification; (b) high-magnification; and (c) Raman spectra for C/C–ZrC–SiC substrates before and after depositing CNTs.
the hardness and elastic modulus of the coating are 40.64 ± 0.45 and 407.43 ± 20.12 GPa, respectively. Hence, the hardness and elastic modulus of SiC coating are enhanced by 26.37% and 28.23% respectively due to the incorporation of CNTs. Generally, the toughness of the ceramic is correlated to their hardness and elastic modulus. Thus it can be deduced that the toughness of SiC coating could be improved by the incorporation of CNTs. High magnification SEM image of the fracture surface of the CNTs–SiC coating is displayed in Fig. 6. It is interesting to find that, the pulling-out and bridging of CNTs through the SiC matrix can be observed. The improvement of toughness of SiC coating would be responsible for the following reasons. On one hand,
when the CNTs form bridges across the gaps or coating grains, the nanoscale mechanical interlocking could be constructed and stress would be effectively transferred to the CNTs due to their exotic mechanical properties, restraining the propagating and widening of microcrack. On the other hand, during the process of crack propagation, CNTs pulled out from the SiC matrix could consume much energy, which is also attributed to improve the hardness and elastic modulus of SiC coating. Thus it can be seen that, CNTs play a positive role on the loading transfer and coating toughening not only by bearing a part of the externally applied load through bridging mechanism but also by CNTs pulled out from SiC matrix.
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(a)
CNT
Inner CNT/SiC
Outer SiC 1 µm (b)
(c)
Resin Coating Substrate
25 µm
(d)
(e)
4 µm
10 µm
4 µm
10 µm
Fig. 4. XRD pattern and SEM images of the coated specimens: (a) cross-section SEM image of the CNTs–SiC coating before polishing; (b) cross-section SEM image of the CNTs– SiC coating after polishing; (c) XRD pattern of the coating surface; (d) surface SEM image of the SiC coating without CNTs; and (e) surface SEM image of the SiC coating with CNTs.
3.3. Bonding strength of CNTs–SiC coating Acoustic emission (AE) scratch test was employed to estimate the interface bonding strength between the coating and substrates. The AE was recorded as a function of loading force. For the coated specimens without CNTs, the first obvious AE was detected at 2.12 N. After the incorporation of CNTs, there is an evident improvement of interface bonding strength and the first AE was detected at 3.25 N, as shown in Fig. 7a. Friction scratch test was applied to further investigate the interface bonding strength. It is well known that the friction is proportional to the constantly changing load exerted by the diamond
Table 1 Mechanical properties of the coating.
Fig. 5. Representative nanoindentation.
load–displacement
curves
of
the
coating
from
SiC coating
Hardness (GPa)
Elastic modulus (GPa)
Without CNTs With CNTs
32.16 ± 0.38 40.64 ± 0.45
317.74 ± 17.34 407.43 ± 20.12
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Bridging Debonding Pulling-out
CNTs
Bridging
100 nm
Fig. 6. SEM image of the coating with CNTs showing fracture surface. Fig. 8. Thermal shock curves of the coated specimens in air at 1773 K.
pinhead. Hence, the abrupt vanish of friction illustrated the pinhead did not put pressure on the specimens and the failure of the coating happened. For the coated specimens without and with CNTs, the first obvious vanishes of friction were detected at 2.04 and 3.22 N, respectively, as displayed in Fig. 7b, which were in approximation with the results of acoustic emission scratch test. The enhancement of interface bonding strength could be contributed to the CNTs firmly rooted in C/C–ZrC–SiC substrates. 3.4. Thermal shock resistance of CNTs–SiC coating The coated composites were subjected to thermal shock test between 1773 K and room temperature in air and the mass loss curves are shown in Fig. 8. All of the coated specimens display an initial slightly mass gain process and reach to about 0.26% and 0.02% after 4-time thermal cycling for the coated specimens with and without CNTs, respectively. During the thermal shock test at 1773 K in air, the coating would react with oxygen, as follows:
2SiC ðsÞ þ 3O2 ðgÞ ! 2SiO2 ðsÞ þ 2CO ðgÞ
ð4Þ
SiC ðsÞ þ 2O2 ðgÞ ! SiO2 ðsÞ þ CO2 ðgÞ
ð5Þ
According to reactions (4) and (5), the oxidation of SiC was a mass gain process. Combined with XRD pattern (Fig. 9a), there is a newly added phase of SiO2 on the coating surface. EDS analysis (Fig. 9b) shows that the coating consists of 14.08% oxygen element.
Thus it can be inferred that SiC was partly oxidized into SiO2 film which could act as an oxidation inhibitor due to its low oxygen diffusion coefficient at 1773 K. The mass loss curves of coated specimens for 5–15-time thermal shock cycling were fitted as linear equations. It can be seen from Fig. 8, the slopes of mass loss for the coated specimens without and with CNTs are 0.552 and 0.187, respectively, suggesting a lower oxidation rate of coated specimens with CNTs compared to that of the coated specimens without CNTs. After 15-time thermal cycling, the coated specimens without CNTs show a 5.98% mass loss, while the coated specimens with CNTs loss 1.98% of their mass, indicating the better thermal shock resistance for the coated specimens with CNTs. The oxidation rate of specimens is strongly affected by the diffusion of oxygen through the microcracks in the coating that were formed during the cooling from 1773 K. By assuming the gradient of oxygen is homogeneous and oxygen diffusion through the microcracks follows the Fick Law, the oxygen diffusion rate through the microcracks can be expressed as:
R ¼ ADO @P O2 =@x ¼ nAC DO PO2 =d
ð6Þ
where AC is the opening area of microcracks, DO the diffusion coefficient, PO2 the pressure of O2 in air, n the number of microcracks and d the thickness of coating. The oxidation rate of specimens is proportional to the opening area of microcracks, which play an
Fig. 7. Variation of acoustic emission signal (a) and friction (b) with sliding load for the coated specimens.
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(b)
(a)
(c)
(d) Spot 1
Fig. 9. XRD pattern (a), EDS analysis (b) and surface SEM images of the SiC coated specimens after thermal shock cycling: (c) without CNTs; (d) with CNTs.
important effect on the thermal shock resistance of the coated specimens. Surface SEM images of the coating with and without CNTs after 15-time thermal cycling are shown in Fig. 9c and d. After undergone the thermal shock between 1773 K and room temperature for 15 times, microcracks can be observed on the coating without CNTs and the maximum width of microcracks is about 4.17 lm. While for the coating with CNTs, the maximum width of microcrack is about 1.93 lm. That is, the opening area of microcracks is restrained due to the construction of CNT/SiC stress buffer layer.
For the coated specimens with CNTs, there is a decrease of microcrack size by about 53.71% after 15-time thermal cycling, as displayed in Fig. 10. The improvement of thermal shock resistance of the CNTs–SiC coated specimens could be attributed to the following reasons. Firstly, CNT/SiC inner coating as a stress buffer layer could effectively reduce the mismatch of the CTEs between C/C–ZrC–SiC substrates and SiC outer coating, resulting in alleviating the thermal stress emerged during the thermal shock between 1773 K and room temperature. Secondly, the coating would be in tensile stress when cooling from elevated temperature to room temperature, causing cracking in coating once the tensile stress exceeded the strength of coating. A schematic was made to further illustrate the reinforcement mechanism of CNTs in the coating, shown in Fig. 11. The propagation of microcrack in the CNTs–SiC coating resulted in the CNTs bridging and finally pulling out from SiC coating. This process could consume much energy and restrain the expansion and widening of microcrack. Additionally, when the microcracks encountered the CNT rods, their deflection would happen [22]. The microcrack firstly propagated along rod boundary which acted as a direction leader. Since the energy for the propagation of microcrack was supplied continuously by the
Crack deflection
Cracks
CNT Coating Substrate Fig. 10. The breadth of the widest microcrack of the coated specimens before and after thermal cycling.
(a)
(b)
Fig. 11. The schematic diagram of the coating: (a) without CNTs; and (b) with CNTs.
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external stress field, after propagating to the end of the rod, the microcrack would turn to another direction. The microcrack deflection led to a zigzag path in the coating and the resistance increase of the crack propagation due to an unfavorable stress state. Thirdly, incorporating CNTs into SiC coating has been confirmed to be advantageous to the interface bonding strength, which made the coating difficult to peeling off from the substrates, leading to an evident improvement of thermal shock resistance. 4. Conclusions The effect of ICVD-CNTs on the toughness, interface bonding strength and thermal shock resistance of SiC coating for C/C– ZrC–SiC composites was investigated. The incorporation of CNTs led to an increase of hardness and elastic modulus of the coating. The interface bonding strength was enhanced and the microcrack size after 15-time thermal shock cycling in air was restrained due to the introduction of CNTs. The CNT/SiC buffer layer could effectively alleviate the mismatch of the CTEs between SiC coating and C/C–ZrC–SiC composites, enhancing the toughness, interface bonding strength and thermal shock resistance of SiC coating due to the nanoscale reinforcement mechanism of CNTs by the mechanical interlocking of bridging, pullout and crack deflection. Acknowledgments This work has been supported by National Natural Science Foundation of China under Grant Nos. 51221001 and 51222207, Project supported by the Research Fund of the State Key Laboratory of Solidification Processing (NWPU), China (Grant No. 12-BZ-2014). References [1] H.J. Li, L. Liu, Y.D. Zhang, K.Z. Li, X.H. Shi, Y.L. Zhang, W. Feng, Effect of high temperature heat treatment on the ablation of SiC–ZrB2–ZrC particles modified C/C composites in two heat fluxes, J. Alloys. Comp. 621 (2015) 18–25. [2] K.Z. Li, J. Xie, Q.G. Fu, H.J. Li, L.J. Guo, Effects of porous C/C density on the densification behavior and ablation property of C/C–ZrC–SiC composites, Carbon 57 (2013) 161–168. [3] D.D. Jayaseelan, R.G. Sa, P. Brown, W.E. Lee, Reactive infiltration processing (RIP) of ultra high temperature ceramics (UHTC) into porous C/C composite tubes, Eur. Ceram. Soc. 31 (2011) 361–368. [4] X.T. Shen, K.Z. Li, H.J. Li, H.Y. Du, W.F. Cao, F.T. Lan, Microstructure and ablation properties of zirconium carbide doped carbon/carbon composites, Carbon 48 (2010) 344–351.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Effect of carbon nanotubes on the toughness, bonding strength and thermal shock resistance of SiC coating for C/C–ZrC–SiC composites Qian-gang Fu ⇑, Lei Zhuang, He-jun Li, Lei Feng, Jun-yi Jing, Bi-yi Tan State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, Shaanxi 710072, PR China
a r t i c l e
i n f o
Article history: Received 11 April 2015 Received in revised form 28 April 2015 Accepted 30 April 2015 Available online 8 May 2015 Keywords: C/C–ZrC–SiC composites SiC coating Carbon nanotube Toughness Bonding strength Thermal shock
a b s t r a c t In order to improve the toughness, interface bonding strength and thermal shock resistance of SiC coating for C/C–ZrC–SiC composites, carbon nanotubes (CNTs) were prepared by injection chemical vapor deposition and attempted as the reinforcement materials in the SiC coating. After incorporating CNTs, the hardness and elastic modulus of the SiC coating increased by 26.37% and 28.23%, respectively. The interface bonding strength between SiC coating and C/C–ZrC–SiC composites was enhanced by 53.31%. The mass loss of the SiC coated C/C–ZrC–SiC composites after thermal shock between 1773 K and room temperature for 15 times decreased from 5.98% to 1.98%. The incorporation of CNTs can effectively improve the toughness, interface bonding strength and thermal shock resistance of SiC coating due to the nanoscale toughening mechanism of CNTs by pullout, bridging and crack deflection. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Introducing ultra-high temperature ceramics (UHTCs), such as ZrC and SiC [1,2] into carbon/carbon (C/C) composites for the purpose of improving the ablation resistance has been highlighted for several years [3–5]. Unfortunately, C/C–UHTCs composites are lack of stability in air at elevated temperature, mainly due to the rapid oxidation and degradation of carbon, carbon fibers and ceramic matrix [6]. Anti-oxidation coating, typically SiC [7,8], has thus been employed to protect C/C–UHTCs composites from burn-off. Up to now, numbers of techniques have been explored to prepare SiC coating, such as slurry [9], pack cementation [10] and chemical vapor deposition (CVD) [11]. Among these techniques, CVD is an effective technique to produce dense and homogenous SiC coating with various shape [12]. The interface bonding between the coating and substrate is regarded as a key factor to determine the oxidation resistant ability of the coated specimens [12], but currently, there is still a challenge in improving the adhesion ability of CVD-SiC coating obtained by the decomposition of Methyltrichlorosilane (MTS) [13]. Additionally, the mismatch of coefficients of thermal expansion (CTEs) between SiC coating and substrates will lead to cracking in the coating, leaving diffusion paths for oxygen when the coated specimens exposed to an ⇑ Corresponding author. Tel.: +86 29 88494197; fax: +86 29 88495764. E-mail address: [email protected] (Q.-g. Fu). http://dx.doi.org/10.1016/j.jallcom.2015.04.223 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
oxidizing environment at high temperature [14]. The above-mentioned poor adhesion ability and CTEs mismatch between CVD-SiC coating and C/C–UHTCs limit their potential applications in extremely environment. Carbon nanotube (CNT) is a relatively new type of nano-structured material [15], which has attracted much attention as a potential reinforcement element for structural composites, due to their high specific surface area, exotic mechanical, thermal and electrical properties [16]. Recently, Wang et al. have testified that spreading multi-walled carbon nanotubes (MWCNTs) into C/UHTCs contributed to a 29.7% increase in the flexural strength and a 27.9% increase in the fracture toughness [17]. The main problem of CNTs as the reinforcement materials is associated with their natural tendency to agglomerate considerably due to their large surface areas and aspect ratios [18]. Injection chemical vapor deposition (ICVD) method is a process for effectively and uniformly fabricating high-purified CNTs [19]. Up to now, little research has been reported about using ICVD-CNTs as the toughening materials in CVD-SiC coating for C/C–UHTCs composites. In order to improve the interface adhesion ability and alleviating the mismatch of the CTEs between SiC coating and C/C–ZrC–SiC substrates, it is worth attempting to deposit CNTs on C/C–ZrC–SiC composites by ICVD and construct a CNT/SiC buffer layer. In this work, an ICVD technique was employed to deposit CNTs on the surface of C/C–ZrC–SiC substrates, and then the porous CNTs layer was filled with CVD-SiC coating. The microstructures of CNTs and the CNTs–SiC coating were investigated. The effect of CNTs on
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the toughness, interface bonding strength and thermal shock resistance of the CNTs–SiC coating for C/C–ZrC–SiC composites were discussed. 2. Experimental procedure T300 carbon fiber felts (0.45 g/cm3) were densified by chemical vapor infiltration (CVI) to obtain porous C/C composites with a density of 1.2 g/cm3. Xylene used as solvent, zirconium-containing polymer (PZC) (Institute of Process Engineering, Chinese Academy of Science, China) and polycarbosilane (PCS) (National University of Defense Technology, Changsha, China) were dissolved and mixed with a mass radio of 3:1 to form a homogenous solution, which was used as ZrC–SiC precursor. After 30 min immersion in the precursor under vacuum, composites were took out and dried in a drying oven at 373 K for 30–36 h. And then, the dried composites were undergone a pyrolysis process at 1773–1873 K for 2 h in a flowing Ar atmosphere. Repeat the infiltration-pyrolysis process for several times until the mass increase percentage of C/C–ZrC–SiC composites was no more than 1%. Cubic specimens (10 10 10 mm3) used as substrates were cut from bulk C/C–ZrC–SiC composites with a density of 2.15 g/cm3, and hand-abraded with 300-grit SiC paper and cleaned in an ultrasonic device with ethanol. CNTs were synthesized on the C/C–ZrC–SiC substrates using ICVD method under ambient pressure in a horizontal quartz reactor equipped with an electrical furnace. A mixture solution of 1.2 wt.% ferrocene, 76.8 wt.% ethanol and 22 wt.% ethylenediamine was fed into the reaction tube at a rate of 10 mL/h by a syringe pump. The reaction temperature was 1000–1058 K and the growth time was 15 min. After preparing the porous network of the CNTs, CVD was employed to obtain SiC coating on the C/C–ZrC– SiC substrates. H2 and Ar were used as the carrier gas and the diluent gas, respectively. MTS as the precursor was brought into the reaction chamber by bubbling H2. The flow rates of H2, MTS and Ar were 800 mL/min, 100 mL/min and 100 mL/min, respectively. The deposition temperature was 1373 K and the deposition time was 6 h. The fabrication process of CNTs–SiC coated C/C–ZrC–SiC composites is depicted in Fig. 1. The nanoindentation test was carried out using Nano-Indenter™ XP (MTS System Corp., USA) system with a diamond indenter on the well-polished cross-section of the upper part of the coating to measure the hardness and elastic modulus of the coating. The maximum indentation load was set at a constant value (400 mN). Load and depth were recorded continuously during the indentation process. The hardness and elastic were calculated from the load–depth curves using the Oliver and Pharr method [20]. Scratch tester (WS-2005 multi-functional tester. China) equipped with a C diamond pinhead (cone apex angle 120, tip radius 0.2 mm) was subjected to estimate the interface bonding strength between SiC coating and C/C–ZrC–SiC composites. Scratch test was carried out by applying a constantly changing load which ranged from 0 to 16 N during sliding on the 5 mm path. Thermal cycling test was performed between 1773 K and room temperature in air to investigate the thermal shock resistance of the coating. After oxidation for 5 min at 1773 K in a furnace, the coated specimens were taken out and cooled down to room temperature in air. An electronic balance with a sensitivity of ±0.1 mg was used to weigh the coated specimens. And then, the coated specimens were put into the furnace for the next cycle. Cumulative mass changes of the coated specimens after each thermal cycle were reported as a function of the thermal cycling time. The morphologies of CNTs and SiC coating were analyzed by a scanning electron microscopy (SEM, JSM-6700). The structure of CNTs were carried out by transmission electron microscopy (TEM, Tecnai F30G2). The graphitization degree of the substrates before and after grafting CNTs was investigated by Raman spectroscopy (Renishaw) using an inVia micro-Raman spectrometer with an Ar ion laser of 514.5 nm wavelength. The crystalline structures of SiC coating was analyzed by X-ray diffraction with Cu Ka radiation. (XRD, X Pert PRO, PANalytical, Almelo, The Netherlands).
3. Results and discussion 3.1. Microstructure of CNTs and SiC coating on C/C–ZrC–SiC substrates Fig. 2 presents the typical surface morphology of CNTs on C/C– ZrC–SiC substrates by ICVD. From Fig. 2a, it can be observed that
Coated specimen with CNTs
Specimen with CNTs
C/C-ZrC-SiC specimen ICVD
CVD
Coating
Fig. 1. The schematic diagram of preparing CNTs–SiC coated C/C–ZrC–SiC composites.
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C/C–ZrC–SiC substrates were uniformly and homogeneously covered with a 3D porous network involving numerous of random-oriented CNTs. These CNTs are generally several tens to one hundred micrometers in diameter with a curved structure, as shown in Fig. 2b. Fig. 3a displays the TEM observation, indicating that the CNTs by ICVD have a bamboo-like structure since the periodical fringes can be seen across the nanotubes. In addition, the nanotubes appear high purity without isolated amorphous carbon. High-magnification TEM image (Fig. 3b) of CNTs shows that both the nanotube walls and the inner caps are well graphitized, implying their high crystallinities. Raman spectra were applied to characterize the degree of graphitization of CNTs. There are two main bands in Raman spectra, which are the disorder-induced D band (about 1348 cm1) and the tangential G band (about 1575 cm1). The intensity ratios of D band and G band (ID/IG) were calculated. The lower the ID/IG is, the higher degree of graphitization the specimens have. The substrate without CNTs holds an ID/IG of 1.081, while the substrate covering with CNTs has the ID/IG of 0.972, as displayed in Fig. 3c. The decrease of ID/IG value suggests a higher graphitization degree after depositing CNTs. Fig. 4a presents the cross-section image of the CNTs–SiC coating before polishing. The inner CNT/SiC layer shows a relatively rough morphology, and abundant CNTs can be observed. This inner layer is about 15 lm. Fig. 4b exhibits the cross-section image of the CNTs–SiC coating after polishing. The whole coating is of a dense and homogeneous structure with a thickness of about 60 lm, bonding well with C/C–ZrC–SiC substrates. Fig. 4c displays the XRD analysis of SiC coating. There are three obvious peaks of 71.838°, 60.088° and 35.708°, which are in good agreement with the diffractions of (3 1 1), (2 2 0) and (1 1 1) crystalline planes of b-SiC [21]. SiC coating was realized by the decomposition of MTS based on the following reaction:
CH3 SiCl3 ðgÞ ! SiC ðsÞ þ 3HCl ðgÞ
ð1Þ
Surface image of SiC coating without CNTs is shown in Fig. 4d. A dense surface with a microcrack can be clearly found. When cooling from the fabricating temperature of 1373 K to room temperature, the mismatch of the CTEs between the coating and substrates would lead to some microcracks emerged on the surface inevitably. The breadth of the widest microcrack for the coating without CNTs is about 1.21 lm, while the widest microcrack of coating with CNTs is about 0.74 lm (Fig. 4e). The incorporation of CNTs contributes to a decrease of microcrack size. 3.2. Toughness of CNTs–SiC coating Fig. 5 shows the representative load–depth curves of the coating from nanoindentation. The maximum indentation depth of CNTs–SiC coating is lower than that of SiC coating with the same maximum indentation load (400 mN). This indicates that the hardness of the coating with CNTs is higher than that of the coating without CNTs. Additionally, the unloading curves of coating without and with CNTs could be fitted as unary linear regression equations, as followed:
y ¼ 0:767x 413:201
ð2Þ
y ¼ 0:802x 377:579
ð3Þ
According to Eqs. (2) and (3), the slopes of unloading curves of the coating without and with CNTs are 0.767 and 0.802, respectively, suggesting the incorporation of CNTs contributes to an increase in elastic modulus of the coating. The corresponding results are shown in Table 1. Before incorporating the CNTs, the hardness and elastic modulus of the coating are 32.16 ± 0.38 and 317.74 ± 17.34 GPa, respectively. After the incorporation of CNTs,
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Fig. 2. Surface SEM images of the 3D porous network of the CNTs on C/C–ZrC–SiC substrates: (a) low-magnification; and (b) high-magnification.
Fig. 3. TEM images of the as-synthesized CNTs: (a) low-magnification; (b) high-magnification; and (c) Raman spectra for C/C–ZrC–SiC substrates before and after depositing CNTs.
the hardness and elastic modulus of the coating are 40.64 ± 0.45 and 407.43 ± 20.12 GPa, respectively. Hence, the hardness and elastic modulus of SiC coating are enhanced by 26.37% and 28.23% respectively due to the incorporation of CNTs. Generally, the toughness of the ceramic is correlated to their hardness and elastic modulus. Thus it can be deduced that the toughness of SiC coating could be improved by the incorporation of CNTs. High magnification SEM image of the fracture surface of the CNTs–SiC coating is displayed in Fig. 6. It is interesting to find that, the pulling-out and bridging of CNTs through the SiC matrix can be observed. The improvement of toughness of SiC coating would be responsible for the following reasons. On one hand,
when the CNTs form bridges across the gaps or coating grains, the nanoscale mechanical interlocking could be constructed and stress would be effectively transferred to the CNTs due to their exotic mechanical properties, restraining the propagating and widening of microcrack. On the other hand, during the process of crack propagation, CNTs pulled out from the SiC matrix could consume much energy, which is also attributed to improve the hardness and elastic modulus of SiC coating. Thus it can be seen that, CNTs play a positive role on the loading transfer and coating toughening not only by bearing a part of the externally applied load through bridging mechanism but also by CNTs pulled out from SiC matrix.
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(a)
CNT
Inner CNT/SiC
Outer SiC 1 µm (b)
(c)
Resin Coating Substrate
25 µm
(d)
(e)
4 µm
10 µm
4 µm
10 µm
Fig. 4. XRD pattern and SEM images of the coated specimens: (a) cross-section SEM image of the CNTs–SiC coating before polishing; (b) cross-section SEM image of the CNTs– SiC coating after polishing; (c) XRD pattern of the coating surface; (d) surface SEM image of the SiC coating without CNTs; and (e) surface SEM image of the SiC coating with CNTs.
3.3. Bonding strength of CNTs–SiC coating Acoustic emission (AE) scratch test was employed to estimate the interface bonding strength between the coating and substrates. The AE was recorded as a function of loading force. For the coated specimens without CNTs, the first obvious AE was detected at 2.12 N. After the incorporation of CNTs, there is an evident improvement of interface bonding strength and the first AE was detected at 3.25 N, as shown in Fig. 7a. Friction scratch test was applied to further investigate the interface bonding strength. It is well known that the friction is proportional to the constantly changing load exerted by the diamond
Table 1 Mechanical properties of the coating.
Fig. 5. Representative nanoindentation.
load–displacement
curves
of
the
coating
from
SiC coating
Hardness (GPa)
Elastic modulus (GPa)
Without CNTs With CNTs
32.16 ± 0.38 40.64 ± 0.45
317.74 ± 17.34 407.43 ± 20.12
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Bridging Debonding Pulling-out
CNTs
Bridging
100 nm
Fig. 6. SEM image of the coating with CNTs showing fracture surface. Fig. 8. Thermal shock curves of the coated specimens in air at 1773 K.
pinhead. Hence, the abrupt vanish of friction illustrated the pinhead did not put pressure on the specimens and the failure of the coating happened. For the coated specimens without and with CNTs, the first obvious vanishes of friction were detected at 2.04 and 3.22 N, respectively, as displayed in Fig. 7b, which were in approximation with the results of acoustic emission scratch test. The enhancement of interface bonding strength could be contributed to the CNTs firmly rooted in C/C–ZrC–SiC substrates. 3.4. Thermal shock resistance of CNTs–SiC coating The coated composites were subjected to thermal shock test between 1773 K and room temperature in air and the mass loss curves are shown in Fig. 8. All of the coated specimens display an initial slightly mass gain process and reach to about 0.26% and 0.02% after 4-time thermal cycling for the coated specimens with and without CNTs, respectively. During the thermal shock test at 1773 K in air, the coating would react with oxygen, as follows:
2SiC ðsÞ þ 3O2 ðgÞ ! 2SiO2 ðsÞ þ 2CO ðgÞ
ð4Þ
SiC ðsÞ þ 2O2 ðgÞ ! SiO2 ðsÞ þ CO2 ðgÞ
ð5Þ
According to reactions (4) and (5), the oxidation of SiC was a mass gain process. Combined with XRD pattern (Fig. 9a), there is a newly added phase of SiO2 on the coating surface. EDS analysis (Fig. 9b) shows that the coating consists of 14.08% oxygen element.
Thus it can be inferred that SiC was partly oxidized into SiO2 film which could act as an oxidation inhibitor due to its low oxygen diffusion coefficient at 1773 K. The mass loss curves of coated specimens for 5–15-time thermal shock cycling were fitted as linear equations. It can be seen from Fig. 8, the slopes of mass loss for the coated specimens without and with CNTs are 0.552 and 0.187, respectively, suggesting a lower oxidation rate of coated specimens with CNTs compared to that of the coated specimens without CNTs. After 15-time thermal cycling, the coated specimens without CNTs show a 5.98% mass loss, while the coated specimens with CNTs loss 1.98% of their mass, indicating the better thermal shock resistance for the coated specimens with CNTs. The oxidation rate of specimens is strongly affected by the diffusion of oxygen through the microcracks in the coating that were formed during the cooling from 1773 K. By assuming the gradient of oxygen is homogeneous and oxygen diffusion through the microcracks follows the Fick Law, the oxygen diffusion rate through the microcracks can be expressed as:
R ¼ ADO @P O2 =@x ¼ nAC DO PO2 =d
ð6Þ
where AC is the opening area of microcracks, DO the diffusion coefficient, PO2 the pressure of O2 in air, n the number of microcracks and d the thickness of coating. The oxidation rate of specimens is proportional to the opening area of microcracks, which play an
Fig. 7. Variation of acoustic emission signal (a) and friction (b) with sliding load for the coated specimens.
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(b)
(a)
(c)
(d) Spot 1
Fig. 9. XRD pattern (a), EDS analysis (b) and surface SEM images of the SiC coated specimens after thermal shock cycling: (c) without CNTs; (d) with CNTs.
important effect on the thermal shock resistance of the coated specimens. Surface SEM images of the coating with and without CNTs after 15-time thermal cycling are shown in Fig. 9c and d. After undergone the thermal shock between 1773 K and room temperature for 15 times, microcracks can be observed on the coating without CNTs and the maximum width of microcracks is about 4.17 lm. While for the coating with CNTs, the maximum width of microcrack is about 1.93 lm. That is, the opening area of microcracks is restrained due to the construction of CNT/SiC stress buffer layer.
For the coated specimens with CNTs, there is a decrease of microcrack size by about 53.71% after 15-time thermal cycling, as displayed in Fig. 10. The improvement of thermal shock resistance of the CNTs–SiC coated specimens could be attributed to the following reasons. Firstly, CNT/SiC inner coating as a stress buffer layer could effectively reduce the mismatch of the CTEs between C/C–ZrC–SiC substrates and SiC outer coating, resulting in alleviating the thermal stress emerged during the thermal shock between 1773 K and room temperature. Secondly, the coating would be in tensile stress when cooling from elevated temperature to room temperature, causing cracking in coating once the tensile stress exceeded the strength of coating. A schematic was made to further illustrate the reinforcement mechanism of CNTs in the coating, shown in Fig. 11. The propagation of microcrack in the CNTs–SiC coating resulted in the CNTs bridging and finally pulling out from SiC coating. This process could consume much energy and restrain the expansion and widening of microcrack. Additionally, when the microcracks encountered the CNT rods, their deflection would happen [22]. The microcrack firstly propagated along rod boundary which acted as a direction leader. Since the energy for the propagation of microcrack was supplied continuously by the
Crack deflection
Cracks
CNT Coating Substrate Fig. 10. The breadth of the widest microcrack of the coated specimens before and after thermal cycling.
(a)
(b)
Fig. 11. The schematic diagram of the coating: (a) without CNTs; and (b) with CNTs.
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external stress field, after propagating to the end of the rod, the microcrack would turn to another direction. The microcrack deflection led to a zigzag path in the coating and the resistance increase of the crack propagation due to an unfavorable stress state. Thirdly, incorporating CNTs into SiC coating has been confirmed to be advantageous to the interface bonding strength, which made the coating difficult to peeling off from the substrates, leading to an evident improvement of thermal shock resistance. 4. Conclusions The effect of ICVD-CNTs on the toughness, interface bonding strength and thermal shock resistance of SiC coating for C/C– ZrC–SiC composites was investigated. The incorporation of CNTs led to an increase of hardness and elastic modulus of the coating. The interface bonding strength was enhanced and the microcrack size after 15-time thermal shock cycling in air was restrained due to the introduction of CNTs. The CNT/SiC buffer layer could effectively alleviate the mismatch of the CTEs between SiC coating and C/C–ZrC–SiC composites, enhancing the toughness, interface bonding strength and thermal shock resistance of SiC coating due to the nanoscale reinforcement mechanism of CNTs by the mechanical interlocking of bridging, pullout and crack deflection. Acknowledgments This work has been supported by National Natural Science Foundation of China under Grant Nos. 51221001 and 51222207, Project supported by the Research Fund of the State Key Laboratory of Solidification Processing (NWPU), China (Grant No. 12-BZ-2014). References [1] H.J. Li, L. Liu, Y.D. Zhang, K.Z. Li, X.H. Shi, Y.L. Zhang, W. Feng, Effect of high temperature heat treatment on the ablation of SiC–ZrB2–ZrC particles modified C/C composites in two heat fluxes, J. Alloys. Comp. 621 (2015) 18–25. [2] K.Z. Li, J. Xie, Q.G. Fu, H.J. Li, L.J. Guo, Effects of porous C/C density on the densification behavior and ablation property of C/C–ZrC–SiC composites, Carbon 57 (2013) 161–168. [3] D.D. Jayaseelan, R.G. Sa, P. Brown, W.E. Lee, Reactive infiltration processing (RIP) of ultra high temperature ceramics (UHTC) into porous C/C composite tubes, Eur. Ceram. Soc. 31 (2011) 361–368. [4] X.T. Shen, K.Z. Li, H.J. Li, H.Y. Du, W.F. Cao, F.T. Lan, Microstructure and ablation properties of zirconium carbide doped carbon/carbon composites, Carbon 48 (2010) 344–351.
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