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carbon composites in seawater conditions
Composites Science and Technology 87 (2013) 50–57
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Tribological behavior of three-dimensional needled ceramic modified carbon/carbon composites in seawater conditions Yanzhi Cai a,b,⇑, Xiaowei Yin a, Shangwu Fan a, Litong Zhang a, Laifei Cheng a a b
National Key Laboratory of Thermostructure Composite Materials, Northwestern Polytechnical University, 127#, Youyi Road, Xi’an, Shaanxi 710072, PR China College of Materials and Mineral Resources, Xi’an University of Architecture and Technology, 13#, Yanta Road, Xi’an, Shaanxi 710055, PR China
a r t i c l e
i n f o
Article history: Received 17 November 2012 Accepted 18 July 2013 Available online 13 August 2013 Keywords: A. Carbon fibers A. Ceramics B. Friction/wear Slurry infiltration
a b s t r a c t Silicon carbide (SiC) or boron carbide (B4C) ceramic filler was introduced into a three-dimensional needled carbon fiber integrated felt by unidirectional pressure slurry infiltration–filtration to prepare ceramic modified carbon/carbon (C/C) composites. The contents of SiC and B4C introduced by a one-shot infiltration were about 18 wt% and 16 wt% respectively. The fade of coefficient of friction (COF) in seawater conditions for C/C composites was significantly reduced by ceramic modification. SiC modified C/C composite was superior to C/SiC for its excellent stable friction without any fade in seawater conditions. The weak hygroscopicity and the ability to produce fresh debris continuously in seawater conditions were the primary factors for ceramic modified C/C composites to reduce and even avoid the fade. The double lubrication action from water film and silicon oxide film increased the COF fade for C/SiC. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Aircraft brakes represent one of the most significant and mature markets for carbon/carbon (C/C) composites [1]. Three-dimensional (3D) needled C/C composites have an obvious advantage in braking applications. C/C composites not only have prominent structural properties of high specific strength and specific modulus, but also excellent functional characteristics such as good thermal conductivity, high specific heat, low density, good wearability, and self-lubricating capability. 3D needled C/C introduces fiber bundles perpendicular to the lamina direction which further improves the bonding strength and thermal conductivity between laminas. However, C/C brakes suffer from their insufficient stability in coefficient of friction (COF) caused by humidity [2,3]. What is particularly important, high corrosion resistance is required other than the good water resistance for C/C composite aeroplane brakes used in the marine environment. The extremely low COF is a bottle-neck for C/C composite brakes being used in marine environments. Due to the limitations of C/C composites, it is necessary to introduce some materials acting as friction modifiers. By adjusting the compositions and microstructures of matrices, a wide variety of properties can be produced in the C/C composites. Park et al. [4,5] prepared C/C composite laminates with molybdenum disili⇑ Corresponding author at: College of Materials and Mineral Resources, Xi’an University of Architecture and Technology, 13#, Yanta Road, Xi’an, Shaanxi 710055, PR China. Tel./fax: +86 29 82205245. E-mail address: [email protected] (Y. Cai). 0266-3538/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.compscitech.2013.07.026
cide filler, the bulk density, graphitization degree and mechanical properties being effectively improved. Seghi et al. [6,7] fabricated chopped or 3D needled C fibers reinforced composites containing a hybrid matrix (C and boron nitride) which provided outstanding wear resistance, incurring nearly zero wear across the entire testing spectrum. The authors prepared 3D needled carbon/silicon carbide (C/SiC) composites with graphite filler, and the friction/wear properties [8,9] and mechanical properties [10] were improved greatly. However, the reports are few in the studies of the wet tribological behavior of 3D ceramic modified C/C composites, and also few in the tribological behavior of C/C composites in seawater conditions. In this study, SiC or boron carbide (B4C) ceramic was introduced to modify pyrolytic carbon (PyC) matrix, and the composites with double matrixes of C and ceramic reinforced by 3D C fibers were prepared. SiC and B4C both have good corrosion resistance, abrasion resistance and chemical inertness. In addition, B4C has high specific heat which can reduce the temperature rise in the braking process. The ceramic filler was introduced into a 3D needled C fiber integrated felt by a new process, unidirectional pressure slurry infiltration–filtration (UPSIF). This new process has an advantage over liquid silicon (Si) infiltration (LSI) process in avoiding C fibers damage by molten Si and avoiding residual Si. The optimization combination of double matrixes was hoped to improve the adaptability to harsher environments of brake materials. The comparisons of tribological behavior in seawater conditions between ceramic modified C/C, traditional C/C and C/SiC composites had been drawn. The friction stabilization mechanism in seawater conditions for ceramic modified C/C composites was discussed.
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2. Experimental methods 2.1. Samples preparation The 3D needled C fiber integrated felt with a density of 0.55 g/ cm3 was used as the preform. It was fabricated by repeatedly overlapping the layers of 0° weftless ply, short-cut fiber web and 90° weftless ply with needle-punching step by step. All the C fiber types were Polyacrylonitrile (PAN) based C fiber (T 700, 12 K tow, Toray, Japan). The fiber volume contents of the weftless plies, the webs and the needled fibers were 24.0%, 4.5% and 1.5%, respectively. Aimed to overcome the drawbacks of C/C brakes, the ceramic constituent was introduced into the fiber preform by UPSIF, which was followed by chemical vapor infiltration (CVI) to prepare ceramic modified C/C composites. The fiber preform (/ 100 mm 15 mm) was placed in a unidirectional pressure infiltration–filtration device (Northwestern Polytechnical University, Xi’an, PR China). The bottom side of the preform clung to a millipore membrane. The fiber preform was infiltrated with the slurry including SiC or B4C filler powder in water medium under an inert gas pressure-driving from top to bottom. Then the slurry was separated by the millipore membrane at the bottom of the fiber preform. The fiber preform itself is a kind of deep bed filtration medium. Most of the filler grains were retained within the fiber preform and the surface of the millipore membrane. Water and a small quantity of filler grains were flowed out. The infiltration pressure was about 0.6–1.0 MPa and time was 10–20 min. The solid volume fraction in the slurry was 8–12%. SiC and B4C filler powders with average grain sizes of 1.0 and 1.5 lm respectively. All samples were performed only a one-shot infiltration–filtration process and then removed from the device and dried. The preform containing SiC or B4C filler was then obtained. The CVI process was used to transform the preform containing the ceramic filler to the composite of C fiber reinforced double matrix of C and ceramic. The temperature for CVI was about 800– 1000 °C and time was 300–400 h. Propylene was used as a precursor and argon as a carrier and diluting gas. The composites prepared with the above method were named C/C-SiCf and C/C-B4Cf respectively according to the filler types, the subscript ‘‘f’’ said filler. The traditional 3D needled C/C and C/SiC composites were used for comparison. The temperature and time for CVI process to prepare the C/C composite were about 800–1000 °C and 400–700 h. The C/SiC composite [11] was prepared by CVI combined with LSI steps. The CVI process to manufacture the porous C/C preform was performed at about 800–1000 °C for 400–700 h.
2.2. Braking tests The tests of braking properties were operated on MM-1000 testing machine which was described previously in detail [11]. The tests were carried out with one rotating disk of / 76 mm / 52 mm 15 mm (76 mm in outer diameter, 52 mm in inner diameter and 15 mm in thickness) pressed against one stationary disk of / 90 mm / 55 mm 15 mm. The given braking speed, braking pressure, and inertia were 7200 r/min (25 m/s), 0.85 MPa, and 0.235 kg m2, respectively. During braking tests, some parameters such as rotating speed n (r/min), braking torque M (N m), surface temperature of brake disks T (°C), and braking time t (s) were automatically recorded. There were two different seawater conditions for wet braking. Firstly, the couple friction surfaces were poured with enough seawater before braking, and a brake test was performed immediately then. After the wet braking, braking was performed continuously
at the same condition (the brake discs being no longer watered) until the brake curve restore to the dry braking curve. Secondly, in order to further compare the adaptability to adverse environments, the couple friction surfaces of the brake discs were poured with sufficient seawater, standing for 48 h to dry and performing the brake test then. The seawater used in this test was artificial seawater prepared according to the standard of substitute ocean water in American Society of Testing Materials (ASTMs) [12]. Its chemical compositions are shown in Table 1. The tests were repeated 20 times for dry condition and 5 times for seawater conditions. The COF could be calculated from the following equation.
M ¼ lðr1 þ r2 ÞP=2
ð1Þ
where M is braking torque, l COF, P braking pressure, r1 inner radius, and r2 outer radius. The fade of COF in seawater condition was determined by the following equation.
Dð%Þ ¼
lwet =ldry 1 100
ð2Þ
where lwet is the average COF in seawater condition and ldry in dry condition. 2.3. Microstructure and phase analysis The open porosities and bulk densities of samples were measured by Archimedes’ method. The microstructures and friction surface morphology were examined by scanning electron microscope/energy dispersive X-ray spectroscopy (SEM-EDX, S-4200, Hitachi, Japan) and optic microscope (OM, OLYMPUS PM-T3). The phases were analyzed using X-ray diffractometry (XRD, PANALYTIAL, X’PERT, Nether-lands). The mass contents of C and SiC in C/ C-SiCf were determined according to the gravimetric analysis [13]. The content of C was measured by burning it off at 700 °C for 20 h in air, and the SiC-content could be calculated. 3. Results and discussion 3.1. Microstructure characteristics The physical properties of ceramic modified and unmodified C/ C composites are listed in Table 2. Compared with the unmodified C/C composite, the densities of ceramic modified C/C composites were slightly lower, and their open porosities were slightly higher, being due to a shorter densification process. Compared with C/SiC composite (with a density of 2.10 g/cm3 and an open porosity of 3% [11]), their densities were greatly reduced, which is beneficial for reducing weight of brake materials.
Table 1 Chemical compositions of seawater. Constituents Concentration (g/L)
NaCl 24.53
MgCl26H2O 5.20
Na2SO4 4.09
CaCl2 1.17
KCl 0.69
NaHCO3 0.20
KBr 0.10
Table 2 Physical properties of C/C composites with or without ceramic filler.
a
Composites
C/C-SiCf
C/C-B4Cf
C/C
Density (g/cm3) Open porosity (%) Volume content of filler (vol%) Mass content of fillera (wt%)
1.66 15.8 9.5 18.5/18.0
1.68 16.1 10.4 15.8/–
1.76 11.1 – –
Mass content of filler was calculation value/test value.
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Other than the mass percent of SiC in C/C-SiCf obtained by the gravimetric analysis, the volume content and mass content of the ceramic filler in the modified C/C composites were also calculated in this study. The calculation formulas were as follows:
vf
¼ 100% ðm2 m1 Þ=ðpr 2 dqf Þ
ð3Þ
wf ¼ 100% ðv f qf Þ=qc
ð4Þ
where vf and wf are the volume content and mass content of the filler respectively, qf and qc the densities of the filler and the corresponding modified C/C composite respectively, m1 and m2 the dry weight of C fiber preform before and after infiltration respectively, r and d the radius and thickness of the modified C/C composite respectively. It is clear that the calculated value was close to the test value for SiC mass content in C/C-SiCf. So it can be extrapolated that it is feasible and effective to obtain the filler contents in the modified C/C composites using the above formulas. The C/SiC composite was composed of 65 wt% C, 27 wt% SiC, and 8 wt% Si [11]. The SiC content introduced by UPSIF was less than that introduced by LSI. But the SiC content in C/C-SiCf was close to 20 wt% which still indicates a high infiltration efficiency of this new process. Fig. 1 shows the XRD analysis results of ceramic modified and unmodified C/C composites. The (0 0 2) peak of each material was sharp, indicating a high crystallinity was obtained in C matrices. The intensity of (0 0 2) peak for C/C-SiCf being lower than that for C/C because the C content in C/C-SiCf was lower than that in C/C, but the (0 0 2) diffraction angle for C/C-SiCf was very close to that for C/C. So, SiC filler had no obvious effect on the graphitization of C matrix. The intensity of (0 0 2) peak for C/C-B4Cf was very close to that for C/C despite the lower C content in C/C-B4Cf than in C/C, moreover the shift of the (0 0 2) peak to the right occurred for C/CB4Cf. The calculated value of the interlayer spacing d002 of C matrix in these composites from Bragg equation was also listed in Fig. 1. It is clear that the interlayer spacing of C matrix in C/C-B4Cf approached nearest to that of graphite crystal (0.3354 nm). These results indicate a structure of higher crystallite graphite in C/C-B4Cf than in C/C and C/C-SiCf. As an excellent graphitization catalyst, B4C could accelerate the process of C matrix graphitization. The whole morphologies of ceramic modified and unmodified C/C composites are shown in Fig. 2. For the 3D needled C fiber preform, in the short-cut webs and inter-bundles fibers are sparse, where pores are large; within fiber bundles fibers are dense, where pores are much smaller. There were few pores inside the fiber bundles but some large pores within the short-cut webs and inter-bundles in the unmodified C/C composite (Fig. 2(a)). Introducing ceramic filler decreased greatly the pore size in the short-cut webs
8000
(002) +
+ C; Δ SiC; * B4C
7000 d=0.3403nm
Intensity / a.u.
6000 d=0.3367nm
5000 4000
(c)
3000 ∗ ∗
2000 1000
∗
d=0.3399nm
Δ
0 10
20
(b)
∗
Δ
30
Δ
Δ
40
(a) 50
2 /° Fig. 1. XRD pattern of C/C composites with or without ceramic filler: (a) C/C-SiCf; (b) C/C-B4Cf; and (c) C/C.
and inter-bundles, and the pore size distribution in the whole samples adding ceramic filler was more uniform (Fig. 2(b) and (c)). The pore size distribution test result (Fig. 2(d)) conformed to the microstructure characteristics. Compared with the unmodified C/C composite, a unimodal pore size distribution (about 40– 150 lm) became a bimodal pore size distribution (about 0–6 lm and 10–100 lm) for ceramic modified C/C composites. A prior deposition for PyC in small pore regions, ceramic filler, on the contrary, can fill large pores more efficiently. The large pores in shortcut webs and inter-bundles were divided into smaller pores by filler grains entered, PyC was then deposited in these smaller pores efficiently. Fig. 3 shows the backscattered photos of the distribution in PyC matrix of the ceramic phase introduced by UPSIF or LSI technique. The ceramic phases in all composites were mainly distributed in the sparse fiber areas. But the ceramic phase had better dispersion between fiber tows in ceramic modified C/C composites than in C/ SiC. For the former, the ceramic grains was introduced into the fiber preform in advance by UPSIF, and then PyC was deposited in the surfaces of fibers and ceramic grains by CVI, so ceramic grains were separated from each other by C matrix (Fig. 3(a) and (b)). For the latter, on the contrary, the PyC was deposited within the fiber preform through the CVI process, and then liquid Si was infiltrated into the C/C porous preform and transformed to SiC by in-situ synthesis. Because of a prior deposition of PyC in dense fiber regions, there were large residual pores in sparse fiber regions which became liquid Si infiltration channel, resulting in a continuous SiC layer in these regions (Fig. 3(c)). Moreover, Si gathered between fiber tows in C/SiC (Fig. 3(c)). Si reacted firstly with C around it to form SiC. Si and C were gradually separated by the increased SiC phase, requiring diffusion through SiC layer to contact each other for further reaction. More the SiC produced, longer the diffusion distance, and silicification reaction was controlled by diffusion, resulting in large tracts of residual Si.
3.2. Friction behavior in seawater conditions The typical braking curves and recovery curves under the first seawater condition are shown in Fig. 4. The corresponding COF and its fade are listed in Table 3. The frictional performance of ceramic modified C/C discs was less influenced by seawater conditions compared with the unmodified C/C discs. Especially C/C-SiCf discs exhibited excellent stable friction without any fade. There was a small fade in the COF of C/C-B4Cf discs. The wet COF and its fade of C/C-B4Cf discs were close to those of C/SiC discs, and they both required only one braking to restore to the dry braking curve. For the unmodified C/C disc, however, the COF under seawater condition had a much more serious fade and recovered slowly. The first braking curves under the second seawater condition are shown in Fig. 5. The corresponding COF and its fade are also listed in Table 3. There was still no fade for C/C-SiCf discs. C/CB4Cf and C/SiC discs both had a severer fade by comparison with the test result in the first seawater condition. C/C-B4Cf discs restored to the dry braking curve at the 4th braking and C/SiC discs restored at the 3rd braking when braking was performed continuously after the first braking. The unmodified C/C discs showed failure and braking impossible after corrosion by seawater for 48 h. No braking to a stop occured when performing braking eight times in succession on C/C discs. So the adaptability to corrosion and moisture environment of ceramic modified C/C composites was much better than that of the unmodified C/C composite. C/C-SiCf outperformed C/SiC in resistance to corrosion and moisture, despite the lower SiC content in C/C-SiCf than in C/SiC; C/C-B4Cf had a worse adaptability to harsh environment than C/SiC.
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(b)
(c)
(d) 0.005 Specific volume / ml g-1
(a)
0.004
C/C C/C-SiCf C/C-B4Cf
0.003 0.002 0.001 0.000 1
10
100
Pore diameter / µm Fig. 2. SEM photo of whole morphology and pore size distribution of C/C composites with or without filler: (a) C/C; (b) C/C-SiCf; (c) C/C-B4Cf; and (d) pore size distribution.
(a)
(b)
(c) SiC
Si
fiber PyC
Fig. 3. Ceramic phase distribution in C matrix: (a) disperse SiC grains in short-cut web of C/C-SiCf; (b) disperse B4C grains in short-cut web of C/C-B4Cf; and (c) gathered SiC and residual Si in short-cut web of C/SiC.
3.3. Wet friction surfaces and debris The COF fade in wet conditions is generally ascribed to the lubrication of the wet layer on the friction surface. So the wet fric-
tion behavior could be related to two factors. One was the hydrophilicity and chemical reactivity with water of material constituents. Another was damaging effects on the lubrication film from the abrasive behavior of material constituents. The authors
Y. Cai et al. / Composites Science and Technology 87 (2013) 50–57
µ
54
0.8 C/C 0.7 0.6 0.5 0.4 (a) 0.3 0.2 0.1 0.0 0 2
4
6
8
10
12
14
16
(g)
(f)
(e)
(d)
(c)
(b)
18
20
22
24
26
28
30
32
µ
Time / s 0.8 C/C-SiCf 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 2
(a)
4
(g)
6
8
10
12
Time / s
µ
0.8 C/C-B4Cf 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 2
(a)
4
6
(b)
8
10
(g)
12
14
16
µ
Time / s 0.8 C/SiC 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 2
(a)
(g)
(b)
4
6
8
10
12
14
Time / s Fig. 4. COF-time curves for braking under the first seawater condition and recovery curves: (a) braking curve in seawater condition; (b) the 1st recovery curve; (c) the 2nd recovery curve; (d) the 3rd recovery curve; (e) the 4th recovery curve; (f) the 5th recovery curve; and (g) braking curve in dry condition.
Table 3 COF and their fade values in seawater conditions.
Braking in the first seawater condition First braking in the second seawater condition Braking in dry condition
l D (%)
l D (%)
l
C/CSiCf
C/CB4Cf
C/C
C/ SiC
0.36 2.9
0.34 15
0.25 36
0.35 13
0.36 2.9
0.29 28
Braking failure
0.33 18
0.35
0.4
0.39
0.40
µ
Samples
0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0
C/SiC C/C-B4Cf
C/C-SiCf
0
2
4
6
8
10
12
14
16
18
Time / s Fig. 5. COF-time curves for first braking under the second seawater condition.
found in the experiment the seawater flowed along the friction surfaces of the ceramic modified C/C discs but was sucked into the unmodified C/C discs quickly after braking under dry condition. The results listed in Table 2 show the open porosities of modified C/C composites were higher than that of the unmodified C/C
composite. So the hydrophilic property was greatly influenced by the hygroscopicity of material constituents besides the open porosity. PyC matrix crystal structure affected the hydrophilic properties of brake discs. Despite the lamellar crystalline structure of
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Y. Cai et al. / Composites Science and Technology 87 (2013) 50–57
graphite, its lubricating property is not intrinsic. Its COF increase is often ascribed to the presence of dangling bonds [14]. Hence, its low friction also results from the complete deactivation of all the dangling bonds. A large number of dangling bonds produced in the wear process were able to be rapidly deactivated by chemisorptions of water and oxygen from the environment. No dangling bond is produced by a preferential cleavage along the basal planes. A graphite particle is broken transversally to its basal planes, new prismatic or edge surfaces are created, which normally display dangling bonds with high chemical reactivity, since their formation implies the breakage of covalent bonds [15]. C matrix in modified and unmodified C/C composites had different graphitization degree. Lower the graphitization degree of PyC, farer from perfect crystal and more defective sites in this PyC structure, brittle fracture is more likely to occur. PyC changed into debris with smaller grain and thereby more dangling bonds were produced, PyC was thus more susceptible to the water. Both SiC and B4C has the weak capacity to absorb water. There was no obvious difference in the graphitization of C matrix in C/C-SiCf and C/C, and a structure of higher crystallite graphite in C/C-B4Cf than in C/C. So these ceramic modified C/C discs had a less tendency for water adsorption. Fig. 6 shows the friction surfaces of different discs after braking under the second seawater condition. The friction surfaces of C/C discs were covered with a layer of uniform, compact and polished film (Fig. 6(a)). Wear should be produced in the surface film. But in terms of braking in seawater conditions for unmodified C/C discs – both braking immediately after the couple friction surfaces being poured with seawater and the continuous braking process after corrosion by seawater for 48 h, the worn debris was too few to be collected. The wear process not only destroys the lubricating film, but also results in the dangling bond produced in PyC. Wear debris rarely produced indicates the surface film was difficult to be destroyed. Water and NaCl were main constituents in the chemical compositions of seawater in this experiment. Chloride was believed to have real lubrication effect [16]. The lubricating effect from both salt and water and their interactive effect on the friction surface resulted better lubricating. If the surface shear stress was not big en-
ough to remove the film, less wear and even non-wear could be expected. So the COF for unmodified C/C discs in seawater conditions had a serious fade and was difficult to be restored. Literature [16] pointed out the surface film could be destroyed only by a braking pressure exceeding a threshold value. But this increased pressure would raise a more rigorous requirement for the mechanical strength and toughness of materials. For braking test under the second seawater condition, water molecules was firstly attracted and combined with dangling bonds in PyC within the surfaces and subsurfaces of C/C discs. Then the seawater kept infiltrating into the deeper part of discs as they evaporated slowly during the dry process. Salt molecules continued to go deeply through diffusion after moisture being evaporated completely. This whole process was long, thus water and salt molecules entering the inside of discs were more, so the COF for C/C discs faded so severely as to be failure in braking. Similar to the unmodified C/C discs, the friction surfaces of ceramic modified C/C discs were covered with a thin layer of uniform and compact film (Fig. 6(b) and (c)) in seawater condition. But different from the unmodified C/C discs, the ceramic filler uniformly mixed with PyC in the film in the modified discs (white dots in the Optical micrographs). Fresh debris produced constantly by the effects of Ploughing and grinding on the PyC matrix by the filler grains, ensuring enough dangling bonds in PyC to maintain a high COF. The ploughing effect itself led to higher COF. Ceramic filler mixed with PyC could improve the ‘grip’ of the couple friction surfaces. What was more, B4C and SiC both have good chemical inertness and corrosion resistance. The hard ceramic grains slid and turned with the rotate disk since the beginning of braking, and the fresh debris could be produced immediately. The film formation and destruction was in a dynamic equilibrium. The cycle of the disruption of debris film and the forming of subsequent new debris film was done repeatedly throughout the friction process. The fresh worn debris was produced during each cycle. So there were two factors to decide the friction behavior in seawater conditions. One was the absorption of seawater in the friction surfaces, the other was the ploughing and grinding action of fillers in the frictional surfaces.
(b)
(a)
20µm
20µm
(d)
(c)
SiC
20µm
20µm
Fig. 6. Optical micrograph of friction surface for braking under the second seawater condition: (a) C/C; (b) C/C-SiCf; (c) C/C-B4Cf; and (d) C/SiC.
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Ceramic grains repartitioned within the PyC in the friction surfaces during the grinding process. Ceramic phase was mainly distributed in the sparse fibers areas before braking, but ceramic grains slid and turned in the process of grinding, leading to ceramic fillers well-distributed in the whole PyC matrix in the friction surfaces. Fig. 6(c) shows there were a lot of B4C grains in the dense fiber area. For C/SiC discs, SiC tended to accumulate in friction surfaces (Fig. 6(d)). SiC accumulated area flaking off from the friction surface due to brittle rupture was possible. So SiC content being higher in C/SiC than in C/C-SiCf did not mean SiC abrasive action on the lubricant film being more effective in C/SiC than in C/CSiCf. Fig. 7 shows the morphology and EDX analysis results of debris for braking under the first seawater condition. For two kinds of ceramic modified C/C composites (Fig. 7(a)–(d)), worn debris was mainly consisted of flaky and granular debris with a small amount of particles aggregate, showing a less seawater absorption. The friction stability of C/C-B4Cf was worse than that of C/C-SiCf,
(a)
severer the wet conditions, wider their performance gap became. Literature [17] mentioned the COF of B4C reduced when rubbed against itself with the increasing air humidity. B4C has good chemical inertness, but in the moist environment a slow oxidation as tribochemical reaction was possible. However, the EDX analysis result for worn debris of C/C-B4Cf discs shows the oxygen content was lower. This should be attributed to the B4C grains being encompassed by PyC matrix which was highly graphitized. The PyC matrix had relative weak hydrophilicity and good oxidation resistance, so the B4C grains were protected. Debris of C/SiC discs was mainly in the shape of strip, flake and granule (Fig. 7(e)), which was similar to the debris morphology of ceramic modified C/C discs. Oxygen content in debris of C/SiC discs was much higher than those of ceramic modified C/C discs (Fig. 7(f)). SiC has better oxidation resistance than B4C, but SiC was accumulative in C/SiC. The aggregate residual Si was also contained in C/SiC. The tribochemical reaction between Si and wet environment formed silicon oxide (SiO2) film [18].
(b) 3200
C/C-SiCf
2800
Element At% C 91.98 Si 3.55 O 3.61 Na 0.51 Cl 0.30 Mg 0.05
Counts
2400 2000
C
1600 1200 800 Si
400 O
Na Mg
Cl
0 0
50
100
150
200
250
300
Energy / KeV
(c)
(d) 3200 2800
C/C-B4Cf Element At% C 80.50 B 16.08 O 2.73 Na 0.35 Cl 0.31 Mg 0.03
C
Counts
2400 2000 1600 1200 800 400
B
O
Na Mg
Cl
0 0
50
100
150
200
250
300
Energy / KeV
(e)
(f)
3200
C/SiC Element At% C 73.56 Si 8.36 O 17.24 Na 0.41 Cl 0.41 Mg 0.02
2800
Counts
2400
C
2000
Si
1600 1200 O
800 400
Cl
Na Mg
0 0
50
100
150
200
250
300
Energy / KeV Fig. 7. Debris for braking under the first seawater condition: (a) debris of C/C-SiCf discs; (b) EDX analysis of debris in (a); (c) debris of C/C-B4Cf discs; (d) EDX analysis of debris in (c); (e) debris of C/SiC discs; and (f) EDX analysis of debris in (e).
Y. Cai et al. / Composites Science and Technology 87 (2013) 50–57
SiðsÞ þ 2H2 OðlÞ ! SiO2 ðsÞ þ 2H2 ðgÞ
ð5Þ
The double lubrication action from water film and SiO2 film caused a lower COF in seawater condition for C/SiC discs than for C/C-SiCf discs. But the lubricating film was easily disrupted due to the high SiC content in C/SiC, so the COF was restored fast. 4. Conclusions (1) 3D needled ceramic modified C/C composites: C/C-B4Cf and C/C-SiCf were prepared using UPSIF combined with CVI. The contents of SiC and B4C introduced by a one-shot infiltration were about 18 wt% and 16 wt% respectively. (2) The pore size distribution in C/C composites became more uniform due to ceramic filler addition. A structure of higher crystallite graphite was in C/C-B4Cf than in C/C and C/C-SiCf. Different from the gathered SiC and residual Si in C/SiC, ceramic grains were dispersed and uniform in the modified C/C composites. (3) The COF fade in seawater conditions for C/C composites was significantly reduced by ceramic modification. C/C-SiCf was superior to C/SiC for its excellent stable friction without any fade in seawater conditions. (4) The weak hygroscopicity and the ability to produce fresh debris continuously in seawater conditions were the primary factors for ceramic modified C/C composites to relieve and even avoid the fade. The double lubrication action from water film and SiO2 film increased the COF fade for C/SiC.
Acknowledgements The authors acknowledge the financial supports of Natural Science Foundation of China (Contract No. NSFC50672076) and Shaanxi Province Department of Education Fund (2013JK0925). The authors would like to give their special thanks to Prof. Dr. Yongdong Xu for his kind help on many aspects of scientific research and his continual support in the time of difficulties. Without these, this work would have been not possible.
57
References [1] Hutton TJ, McEnaney B, Crelling JC. Structural studies of wear debris from carbon–carbon composite aircraft brakes. Carbon 1999;37:907–16. [2] Chen JD, Chern Lin JH, Ju CP. Effect of humidity on the tribological behavior of carbon–carbon composites. Wear 1996;193(1):38–47. [3] Blanco C, Bermejo J, Marsh H, Menendez R. Chemical and physical properties of carbon as related to brake performance. Wear 1997;213:1–12. [4] Park S-J, Cho M-S, Lee J-R, Pak P-K. Influence of molybdenum disilicide filler on carbon–carbon composites. Carbon 1999;37:1685–9. [5] Park S-J, Cho M-S. Effect of anti-oxidative filler on the interfacial mechanical properties of carbon–carbon composites measured at high temperature. Carbon 2000;38:1053–8. [6] Seghi S, Fabio B, Economy J. Carbon/carbon-boron nitride composites with improved wear resistance compared to carbon/carbon. Carbon 2004;42: 3043–8. [7] Seghi S, Lee J, Economy J. High density carbon fiber/boron nitride matrix composites: fabrication of composites with exceptional wear resistance. Carbon 2005;43:2035–43. [8] Cai Y, Fan S, Liu H, Zhang L, Cheng L, Dong B, et al. Microstructures and improved wear resistance of 3D needled C/SiC composites with graphite filler. Compos Sci Technol 2009;69:2447–53. [9] Cai Y, Fan S, Yin X, Zhang L, Cheng L, Jiang J, et al. Effects of graphitization degree in 3D needled C/SiC composites on tribological properties. Int J Appl Ceram Technol 2011;8(2):317–28. [10] Cai Y, Fan S, Liu H, Zhang L, Cheng L, Jiang J, et al. Mechanical properties of a 3D needled C/SiC composite with graphite filler. Mater Sci Eng A 2010;527: 539–43. [11] Fang S, Zhang L, Xu Y, Cheng L, Lou J, Zhang J. Microstructure and properties of 3D needle-punched carbon/silicon carbide brake materials. Compos Sci Technol 2007;67:2390–8. [12] ASTM International. ASTM D1141-98. Standard practice for the preparation of substitute ocean water. USA: West Conshohocken, PA; 1998. [13] Frieb M, Krenkel W, Brandt R, Neuer G. Influence of process parameters on the thermophysical properties of C/C-SiC. In: Weinheim S, editor. High temperature ceramic matrix composites, vol. 4. Wiley-VCH Press; 2001. p. 328–33. [14] Brendlé M, Stempflé P. Triboreactions of graphite with moisture – a new model of triboreactor for integrating friction and wear. Wear 2003;254: 818–26. [15] Gardos MN, Soriano BL. The effect of environment on the tribological properties of polycrystalline diamond films. J Mater 1990;5:2599–609. [16] Hanawa T, Asami K, Asaoka K. Repassivation of titanium and surface oxide film regeneration in simulated bioliquid. J Biomed Mater Res 1998;40:530–8. [17] Umeda K, Enomoto Y, Mitsui A, Mannami K. Friction and wear of boride ceramies in air and water. Wear 1993;169:63–8. [18] Fouquet S, Rollin M, Pailler R, Bourrat X. Tribological behaviour of composites made of carbon fibres and ceramic matrix in the Si–C system. Wear 2008;264(9–10):850–6.
Contents lists available at ScienceDirect
Composites Science and Technology journal homepage: www.elsevier.com/locate/compscitech
Tribological behavior of three-dimensional needled ceramic modified carbon/carbon composites in seawater conditions Yanzhi Cai a,b,⇑, Xiaowei Yin a, Shangwu Fan a, Litong Zhang a, Laifei Cheng a a b
National Key Laboratory of Thermostructure Composite Materials, Northwestern Polytechnical University, 127#, Youyi Road, Xi’an, Shaanxi 710072, PR China College of Materials and Mineral Resources, Xi’an University of Architecture and Technology, 13#, Yanta Road, Xi’an, Shaanxi 710055, PR China
a r t i c l e
i n f o
Article history: Received 17 November 2012 Accepted 18 July 2013 Available online 13 August 2013 Keywords: A. Carbon fibers A. Ceramics B. Friction/wear Slurry infiltration
a b s t r a c t Silicon carbide (SiC) or boron carbide (B4C) ceramic filler was introduced into a three-dimensional needled carbon fiber integrated felt by unidirectional pressure slurry infiltration–filtration to prepare ceramic modified carbon/carbon (C/C) composites. The contents of SiC and B4C introduced by a one-shot infiltration were about 18 wt% and 16 wt% respectively. The fade of coefficient of friction (COF) in seawater conditions for C/C composites was significantly reduced by ceramic modification. SiC modified C/C composite was superior to C/SiC for its excellent stable friction without any fade in seawater conditions. The weak hygroscopicity and the ability to produce fresh debris continuously in seawater conditions were the primary factors for ceramic modified C/C composites to reduce and even avoid the fade. The double lubrication action from water film and silicon oxide film increased the COF fade for C/SiC. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Aircraft brakes represent one of the most significant and mature markets for carbon/carbon (C/C) composites [1]. Three-dimensional (3D) needled C/C composites have an obvious advantage in braking applications. C/C composites not only have prominent structural properties of high specific strength and specific modulus, but also excellent functional characteristics such as good thermal conductivity, high specific heat, low density, good wearability, and self-lubricating capability. 3D needled C/C introduces fiber bundles perpendicular to the lamina direction which further improves the bonding strength and thermal conductivity between laminas. However, C/C brakes suffer from their insufficient stability in coefficient of friction (COF) caused by humidity [2,3]. What is particularly important, high corrosion resistance is required other than the good water resistance for C/C composite aeroplane brakes used in the marine environment. The extremely low COF is a bottle-neck for C/C composite brakes being used in marine environments. Due to the limitations of C/C composites, it is necessary to introduce some materials acting as friction modifiers. By adjusting the compositions and microstructures of matrices, a wide variety of properties can be produced in the C/C composites. Park et al. [4,5] prepared C/C composite laminates with molybdenum disili⇑ Corresponding author at: College of Materials and Mineral Resources, Xi’an University of Architecture and Technology, 13#, Yanta Road, Xi’an, Shaanxi 710055, PR China. Tel./fax: +86 29 82205245. E-mail address: [email protected] (Y. Cai). 0266-3538/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.compscitech.2013.07.026
cide filler, the bulk density, graphitization degree and mechanical properties being effectively improved. Seghi et al. [6,7] fabricated chopped or 3D needled C fibers reinforced composites containing a hybrid matrix (C and boron nitride) which provided outstanding wear resistance, incurring nearly zero wear across the entire testing spectrum. The authors prepared 3D needled carbon/silicon carbide (C/SiC) composites with graphite filler, and the friction/wear properties [8,9] and mechanical properties [10] were improved greatly. However, the reports are few in the studies of the wet tribological behavior of 3D ceramic modified C/C composites, and also few in the tribological behavior of C/C composites in seawater conditions. In this study, SiC or boron carbide (B4C) ceramic was introduced to modify pyrolytic carbon (PyC) matrix, and the composites with double matrixes of C and ceramic reinforced by 3D C fibers were prepared. SiC and B4C both have good corrosion resistance, abrasion resistance and chemical inertness. In addition, B4C has high specific heat which can reduce the temperature rise in the braking process. The ceramic filler was introduced into a 3D needled C fiber integrated felt by a new process, unidirectional pressure slurry infiltration–filtration (UPSIF). This new process has an advantage over liquid silicon (Si) infiltration (LSI) process in avoiding C fibers damage by molten Si and avoiding residual Si. The optimization combination of double matrixes was hoped to improve the adaptability to harsher environments of brake materials. The comparisons of tribological behavior in seawater conditions between ceramic modified C/C, traditional C/C and C/SiC composites had been drawn. The friction stabilization mechanism in seawater conditions for ceramic modified C/C composites was discussed.
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2. Experimental methods 2.1. Samples preparation The 3D needled C fiber integrated felt with a density of 0.55 g/ cm3 was used as the preform. It was fabricated by repeatedly overlapping the layers of 0° weftless ply, short-cut fiber web and 90° weftless ply with needle-punching step by step. All the C fiber types were Polyacrylonitrile (PAN) based C fiber (T 700, 12 K tow, Toray, Japan). The fiber volume contents of the weftless plies, the webs and the needled fibers were 24.0%, 4.5% and 1.5%, respectively. Aimed to overcome the drawbacks of C/C brakes, the ceramic constituent was introduced into the fiber preform by UPSIF, which was followed by chemical vapor infiltration (CVI) to prepare ceramic modified C/C composites. The fiber preform (/ 100 mm 15 mm) was placed in a unidirectional pressure infiltration–filtration device (Northwestern Polytechnical University, Xi’an, PR China). The bottom side of the preform clung to a millipore membrane. The fiber preform was infiltrated with the slurry including SiC or B4C filler powder in water medium under an inert gas pressure-driving from top to bottom. Then the slurry was separated by the millipore membrane at the bottom of the fiber preform. The fiber preform itself is a kind of deep bed filtration medium. Most of the filler grains were retained within the fiber preform and the surface of the millipore membrane. Water and a small quantity of filler grains were flowed out. The infiltration pressure was about 0.6–1.0 MPa and time was 10–20 min. The solid volume fraction in the slurry was 8–12%. SiC and B4C filler powders with average grain sizes of 1.0 and 1.5 lm respectively. All samples were performed only a one-shot infiltration–filtration process and then removed from the device and dried. The preform containing SiC or B4C filler was then obtained. The CVI process was used to transform the preform containing the ceramic filler to the composite of C fiber reinforced double matrix of C and ceramic. The temperature for CVI was about 800– 1000 °C and time was 300–400 h. Propylene was used as a precursor and argon as a carrier and diluting gas. The composites prepared with the above method were named C/C-SiCf and C/C-B4Cf respectively according to the filler types, the subscript ‘‘f’’ said filler. The traditional 3D needled C/C and C/SiC composites were used for comparison. The temperature and time for CVI process to prepare the C/C composite were about 800–1000 °C and 400–700 h. The C/SiC composite [11] was prepared by CVI combined with LSI steps. The CVI process to manufacture the porous C/C preform was performed at about 800–1000 °C for 400–700 h.
2.2. Braking tests The tests of braking properties were operated on MM-1000 testing machine which was described previously in detail [11]. The tests were carried out with one rotating disk of / 76 mm / 52 mm 15 mm (76 mm in outer diameter, 52 mm in inner diameter and 15 mm in thickness) pressed against one stationary disk of / 90 mm / 55 mm 15 mm. The given braking speed, braking pressure, and inertia were 7200 r/min (25 m/s), 0.85 MPa, and 0.235 kg m2, respectively. During braking tests, some parameters such as rotating speed n (r/min), braking torque M (N m), surface temperature of brake disks T (°C), and braking time t (s) were automatically recorded. There were two different seawater conditions for wet braking. Firstly, the couple friction surfaces were poured with enough seawater before braking, and a brake test was performed immediately then. After the wet braking, braking was performed continuously
at the same condition (the brake discs being no longer watered) until the brake curve restore to the dry braking curve. Secondly, in order to further compare the adaptability to adverse environments, the couple friction surfaces of the brake discs were poured with sufficient seawater, standing for 48 h to dry and performing the brake test then. The seawater used in this test was artificial seawater prepared according to the standard of substitute ocean water in American Society of Testing Materials (ASTMs) [12]. Its chemical compositions are shown in Table 1. The tests were repeated 20 times for dry condition and 5 times for seawater conditions. The COF could be calculated from the following equation.
M ¼ lðr1 þ r2 ÞP=2
ð1Þ
where M is braking torque, l COF, P braking pressure, r1 inner radius, and r2 outer radius. The fade of COF in seawater condition was determined by the following equation.
Dð%Þ ¼
lwet =ldry 1 100
ð2Þ
where lwet is the average COF in seawater condition and ldry in dry condition. 2.3. Microstructure and phase analysis The open porosities and bulk densities of samples were measured by Archimedes’ method. The microstructures and friction surface morphology were examined by scanning electron microscope/energy dispersive X-ray spectroscopy (SEM-EDX, S-4200, Hitachi, Japan) and optic microscope (OM, OLYMPUS PM-T3). The phases were analyzed using X-ray diffractometry (XRD, PANALYTIAL, X’PERT, Nether-lands). The mass contents of C and SiC in C/ C-SiCf were determined according to the gravimetric analysis [13]. The content of C was measured by burning it off at 700 °C for 20 h in air, and the SiC-content could be calculated. 3. Results and discussion 3.1. Microstructure characteristics The physical properties of ceramic modified and unmodified C/ C composites are listed in Table 2. Compared with the unmodified C/C composite, the densities of ceramic modified C/C composites were slightly lower, and their open porosities were slightly higher, being due to a shorter densification process. Compared with C/SiC composite (with a density of 2.10 g/cm3 and an open porosity of 3% [11]), their densities were greatly reduced, which is beneficial for reducing weight of brake materials.
Table 1 Chemical compositions of seawater. Constituents Concentration (g/L)
NaCl 24.53
MgCl26H2O 5.20
Na2SO4 4.09
CaCl2 1.17
KCl 0.69
NaHCO3 0.20
KBr 0.10
Table 2 Physical properties of C/C composites with or without ceramic filler.
a
Composites
C/C-SiCf
C/C-B4Cf
C/C
Density (g/cm3) Open porosity (%) Volume content of filler (vol%) Mass content of fillera (wt%)
1.66 15.8 9.5 18.5/18.0
1.68 16.1 10.4 15.8/–
1.76 11.1 – –
Mass content of filler was calculation value/test value.
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Other than the mass percent of SiC in C/C-SiCf obtained by the gravimetric analysis, the volume content and mass content of the ceramic filler in the modified C/C composites were also calculated in this study. The calculation formulas were as follows:
vf
¼ 100% ðm2 m1 Þ=ðpr 2 dqf Þ
ð3Þ
wf ¼ 100% ðv f qf Þ=qc
ð4Þ
where vf and wf are the volume content and mass content of the filler respectively, qf and qc the densities of the filler and the corresponding modified C/C composite respectively, m1 and m2 the dry weight of C fiber preform before and after infiltration respectively, r and d the radius and thickness of the modified C/C composite respectively. It is clear that the calculated value was close to the test value for SiC mass content in C/C-SiCf. So it can be extrapolated that it is feasible and effective to obtain the filler contents in the modified C/C composites using the above formulas. The C/SiC composite was composed of 65 wt% C, 27 wt% SiC, and 8 wt% Si [11]. The SiC content introduced by UPSIF was less than that introduced by LSI. But the SiC content in C/C-SiCf was close to 20 wt% which still indicates a high infiltration efficiency of this new process. Fig. 1 shows the XRD analysis results of ceramic modified and unmodified C/C composites. The (0 0 2) peak of each material was sharp, indicating a high crystallinity was obtained in C matrices. The intensity of (0 0 2) peak for C/C-SiCf being lower than that for C/C because the C content in C/C-SiCf was lower than that in C/C, but the (0 0 2) diffraction angle for C/C-SiCf was very close to that for C/C. So, SiC filler had no obvious effect on the graphitization of C matrix. The intensity of (0 0 2) peak for C/C-B4Cf was very close to that for C/C despite the lower C content in C/C-B4Cf than in C/C, moreover the shift of the (0 0 2) peak to the right occurred for C/CB4Cf. The calculated value of the interlayer spacing d002 of C matrix in these composites from Bragg equation was also listed in Fig. 1. It is clear that the interlayer spacing of C matrix in C/C-B4Cf approached nearest to that of graphite crystal (0.3354 nm). These results indicate a structure of higher crystallite graphite in C/C-B4Cf than in C/C and C/C-SiCf. As an excellent graphitization catalyst, B4C could accelerate the process of C matrix graphitization. The whole morphologies of ceramic modified and unmodified C/C composites are shown in Fig. 2. For the 3D needled C fiber preform, in the short-cut webs and inter-bundles fibers are sparse, where pores are large; within fiber bundles fibers are dense, where pores are much smaller. There were few pores inside the fiber bundles but some large pores within the short-cut webs and inter-bundles in the unmodified C/C composite (Fig. 2(a)). Introducing ceramic filler decreased greatly the pore size in the short-cut webs
8000
(002) +
+ C; Δ SiC; * B4C
7000 d=0.3403nm
Intensity / a.u.
6000 d=0.3367nm
5000 4000
(c)
3000 ∗ ∗
2000 1000
∗
d=0.3399nm
Δ
0 10
20
(b)
∗
Δ
30
Δ
Δ
40
(a) 50
2 /° Fig. 1. XRD pattern of C/C composites with or without ceramic filler: (a) C/C-SiCf; (b) C/C-B4Cf; and (c) C/C.
and inter-bundles, and the pore size distribution in the whole samples adding ceramic filler was more uniform (Fig. 2(b) and (c)). The pore size distribution test result (Fig. 2(d)) conformed to the microstructure characteristics. Compared with the unmodified C/C composite, a unimodal pore size distribution (about 40– 150 lm) became a bimodal pore size distribution (about 0–6 lm and 10–100 lm) for ceramic modified C/C composites. A prior deposition for PyC in small pore regions, ceramic filler, on the contrary, can fill large pores more efficiently. The large pores in shortcut webs and inter-bundles were divided into smaller pores by filler grains entered, PyC was then deposited in these smaller pores efficiently. Fig. 3 shows the backscattered photos of the distribution in PyC matrix of the ceramic phase introduced by UPSIF or LSI technique. The ceramic phases in all composites were mainly distributed in the sparse fiber areas. But the ceramic phase had better dispersion between fiber tows in ceramic modified C/C composites than in C/ SiC. For the former, the ceramic grains was introduced into the fiber preform in advance by UPSIF, and then PyC was deposited in the surfaces of fibers and ceramic grains by CVI, so ceramic grains were separated from each other by C matrix (Fig. 3(a) and (b)). For the latter, on the contrary, the PyC was deposited within the fiber preform through the CVI process, and then liquid Si was infiltrated into the C/C porous preform and transformed to SiC by in-situ synthesis. Because of a prior deposition of PyC in dense fiber regions, there were large residual pores in sparse fiber regions which became liquid Si infiltration channel, resulting in a continuous SiC layer in these regions (Fig. 3(c)). Moreover, Si gathered between fiber tows in C/SiC (Fig. 3(c)). Si reacted firstly with C around it to form SiC. Si and C were gradually separated by the increased SiC phase, requiring diffusion through SiC layer to contact each other for further reaction. More the SiC produced, longer the diffusion distance, and silicification reaction was controlled by diffusion, resulting in large tracts of residual Si.
3.2. Friction behavior in seawater conditions The typical braking curves and recovery curves under the first seawater condition are shown in Fig. 4. The corresponding COF and its fade are listed in Table 3. The frictional performance of ceramic modified C/C discs was less influenced by seawater conditions compared with the unmodified C/C discs. Especially C/C-SiCf discs exhibited excellent stable friction without any fade. There was a small fade in the COF of C/C-B4Cf discs. The wet COF and its fade of C/C-B4Cf discs were close to those of C/SiC discs, and they both required only one braking to restore to the dry braking curve. For the unmodified C/C disc, however, the COF under seawater condition had a much more serious fade and recovered slowly. The first braking curves under the second seawater condition are shown in Fig. 5. The corresponding COF and its fade are also listed in Table 3. There was still no fade for C/C-SiCf discs. C/CB4Cf and C/SiC discs both had a severer fade by comparison with the test result in the first seawater condition. C/C-B4Cf discs restored to the dry braking curve at the 4th braking and C/SiC discs restored at the 3rd braking when braking was performed continuously after the first braking. The unmodified C/C discs showed failure and braking impossible after corrosion by seawater for 48 h. No braking to a stop occured when performing braking eight times in succession on C/C discs. So the adaptability to corrosion and moisture environment of ceramic modified C/C composites was much better than that of the unmodified C/C composite. C/C-SiCf outperformed C/SiC in resistance to corrosion and moisture, despite the lower SiC content in C/C-SiCf than in C/SiC; C/C-B4Cf had a worse adaptability to harsh environment than C/SiC.
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Y. Cai et al. / Composites Science and Technology 87 (2013) 50–57
(b)
(c)
(d) 0.005 Specific volume / ml g-1
(a)
0.004
C/C C/C-SiCf C/C-B4Cf
0.003 0.002 0.001 0.000 1
10
100
Pore diameter / µm Fig. 2. SEM photo of whole morphology and pore size distribution of C/C composites with or without filler: (a) C/C; (b) C/C-SiCf; (c) C/C-B4Cf; and (d) pore size distribution.
(a)
(b)
(c) SiC
Si
fiber PyC
Fig. 3. Ceramic phase distribution in C matrix: (a) disperse SiC grains in short-cut web of C/C-SiCf; (b) disperse B4C grains in short-cut web of C/C-B4Cf; and (c) gathered SiC and residual Si in short-cut web of C/SiC.
3.3. Wet friction surfaces and debris The COF fade in wet conditions is generally ascribed to the lubrication of the wet layer on the friction surface. So the wet fric-
tion behavior could be related to two factors. One was the hydrophilicity and chemical reactivity with water of material constituents. Another was damaging effects on the lubrication film from the abrasive behavior of material constituents. The authors
Y. Cai et al. / Composites Science and Technology 87 (2013) 50–57
µ
54
0.8 C/C 0.7 0.6 0.5 0.4 (a) 0.3 0.2 0.1 0.0 0 2
4
6
8
10
12
14
16
(g)
(f)
(e)
(d)
(c)
(b)
18
20
22
24
26
28
30
32
µ
Time / s 0.8 C/C-SiCf 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 2
(a)
4
(g)
6
8
10
12
Time / s
µ
0.8 C/C-B4Cf 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 2
(a)
4
6
(b)
8
10
(g)
12
14
16
µ
Time / s 0.8 C/SiC 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 2
(a)
(g)
(b)
4
6
8
10
12
14
Time / s Fig. 4. COF-time curves for braking under the first seawater condition and recovery curves: (a) braking curve in seawater condition; (b) the 1st recovery curve; (c) the 2nd recovery curve; (d) the 3rd recovery curve; (e) the 4th recovery curve; (f) the 5th recovery curve; and (g) braking curve in dry condition.
Table 3 COF and their fade values in seawater conditions.
Braking in the first seawater condition First braking in the second seawater condition Braking in dry condition
l D (%)
l D (%)
l
C/CSiCf
C/CB4Cf
C/C
C/ SiC
0.36 2.9
0.34 15
0.25 36
0.35 13
0.36 2.9
0.29 28
Braking failure
0.33 18
0.35
0.4
0.39
0.40
µ
Samples
0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0
C/SiC C/C-B4Cf
C/C-SiCf
0
2
4
6
8
10
12
14
16
18
Time / s Fig. 5. COF-time curves for first braking under the second seawater condition.
found in the experiment the seawater flowed along the friction surfaces of the ceramic modified C/C discs but was sucked into the unmodified C/C discs quickly after braking under dry condition. The results listed in Table 2 show the open porosities of modified C/C composites were higher than that of the unmodified C/C
composite. So the hydrophilic property was greatly influenced by the hygroscopicity of material constituents besides the open porosity. PyC matrix crystal structure affected the hydrophilic properties of brake discs. Despite the lamellar crystalline structure of
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graphite, its lubricating property is not intrinsic. Its COF increase is often ascribed to the presence of dangling bonds [14]. Hence, its low friction also results from the complete deactivation of all the dangling bonds. A large number of dangling bonds produced in the wear process were able to be rapidly deactivated by chemisorptions of water and oxygen from the environment. No dangling bond is produced by a preferential cleavage along the basal planes. A graphite particle is broken transversally to its basal planes, new prismatic or edge surfaces are created, which normally display dangling bonds with high chemical reactivity, since their formation implies the breakage of covalent bonds [15]. C matrix in modified and unmodified C/C composites had different graphitization degree. Lower the graphitization degree of PyC, farer from perfect crystal and more defective sites in this PyC structure, brittle fracture is more likely to occur. PyC changed into debris with smaller grain and thereby more dangling bonds were produced, PyC was thus more susceptible to the water. Both SiC and B4C has the weak capacity to absorb water. There was no obvious difference in the graphitization of C matrix in C/C-SiCf and C/C, and a structure of higher crystallite graphite in C/C-B4Cf than in C/C. So these ceramic modified C/C discs had a less tendency for water adsorption. Fig. 6 shows the friction surfaces of different discs after braking under the second seawater condition. The friction surfaces of C/C discs were covered with a layer of uniform, compact and polished film (Fig. 6(a)). Wear should be produced in the surface film. But in terms of braking in seawater conditions for unmodified C/C discs – both braking immediately after the couple friction surfaces being poured with seawater and the continuous braking process after corrosion by seawater for 48 h, the worn debris was too few to be collected. The wear process not only destroys the lubricating film, but also results in the dangling bond produced in PyC. Wear debris rarely produced indicates the surface film was difficult to be destroyed. Water and NaCl were main constituents in the chemical compositions of seawater in this experiment. Chloride was believed to have real lubrication effect [16]. The lubricating effect from both salt and water and their interactive effect on the friction surface resulted better lubricating. If the surface shear stress was not big en-
ough to remove the film, less wear and even non-wear could be expected. So the COF for unmodified C/C discs in seawater conditions had a serious fade and was difficult to be restored. Literature [16] pointed out the surface film could be destroyed only by a braking pressure exceeding a threshold value. But this increased pressure would raise a more rigorous requirement for the mechanical strength and toughness of materials. For braking test under the second seawater condition, water molecules was firstly attracted and combined with dangling bonds in PyC within the surfaces and subsurfaces of C/C discs. Then the seawater kept infiltrating into the deeper part of discs as they evaporated slowly during the dry process. Salt molecules continued to go deeply through diffusion after moisture being evaporated completely. This whole process was long, thus water and salt molecules entering the inside of discs were more, so the COF for C/C discs faded so severely as to be failure in braking. Similar to the unmodified C/C discs, the friction surfaces of ceramic modified C/C discs were covered with a thin layer of uniform and compact film (Fig. 6(b) and (c)) in seawater condition. But different from the unmodified C/C discs, the ceramic filler uniformly mixed with PyC in the film in the modified discs (white dots in the Optical micrographs). Fresh debris produced constantly by the effects of Ploughing and grinding on the PyC matrix by the filler grains, ensuring enough dangling bonds in PyC to maintain a high COF. The ploughing effect itself led to higher COF. Ceramic filler mixed with PyC could improve the ‘grip’ of the couple friction surfaces. What was more, B4C and SiC both have good chemical inertness and corrosion resistance. The hard ceramic grains slid and turned with the rotate disk since the beginning of braking, and the fresh debris could be produced immediately. The film formation and destruction was in a dynamic equilibrium. The cycle of the disruption of debris film and the forming of subsequent new debris film was done repeatedly throughout the friction process. The fresh worn debris was produced during each cycle. So there were two factors to decide the friction behavior in seawater conditions. One was the absorption of seawater in the friction surfaces, the other was the ploughing and grinding action of fillers in the frictional surfaces.
(b)
(a)
20µm
20µm
(d)
(c)
SiC
20µm
20µm
Fig. 6. Optical micrograph of friction surface for braking under the second seawater condition: (a) C/C; (b) C/C-SiCf; (c) C/C-B4Cf; and (d) C/SiC.
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Ceramic grains repartitioned within the PyC in the friction surfaces during the grinding process. Ceramic phase was mainly distributed in the sparse fibers areas before braking, but ceramic grains slid and turned in the process of grinding, leading to ceramic fillers well-distributed in the whole PyC matrix in the friction surfaces. Fig. 6(c) shows there were a lot of B4C grains in the dense fiber area. For C/SiC discs, SiC tended to accumulate in friction surfaces (Fig. 6(d)). SiC accumulated area flaking off from the friction surface due to brittle rupture was possible. So SiC content being higher in C/SiC than in C/C-SiCf did not mean SiC abrasive action on the lubricant film being more effective in C/SiC than in C/CSiCf. Fig. 7 shows the morphology and EDX analysis results of debris for braking under the first seawater condition. For two kinds of ceramic modified C/C composites (Fig. 7(a)–(d)), worn debris was mainly consisted of flaky and granular debris with a small amount of particles aggregate, showing a less seawater absorption. The friction stability of C/C-B4Cf was worse than that of C/C-SiCf,
(a)
severer the wet conditions, wider their performance gap became. Literature [17] mentioned the COF of B4C reduced when rubbed against itself with the increasing air humidity. B4C has good chemical inertness, but in the moist environment a slow oxidation as tribochemical reaction was possible. However, the EDX analysis result for worn debris of C/C-B4Cf discs shows the oxygen content was lower. This should be attributed to the B4C grains being encompassed by PyC matrix which was highly graphitized. The PyC matrix had relative weak hydrophilicity and good oxidation resistance, so the B4C grains were protected. Debris of C/SiC discs was mainly in the shape of strip, flake and granule (Fig. 7(e)), which was similar to the debris morphology of ceramic modified C/C discs. Oxygen content in debris of C/SiC discs was much higher than those of ceramic modified C/C discs (Fig. 7(f)). SiC has better oxidation resistance than B4C, but SiC was accumulative in C/SiC. The aggregate residual Si was also contained in C/SiC. The tribochemical reaction between Si and wet environment formed silicon oxide (SiO2) film [18].
(b) 3200
C/C-SiCf
2800
Element At% C 91.98 Si 3.55 O 3.61 Na 0.51 Cl 0.30 Mg 0.05
Counts
2400 2000
C
1600 1200 800 Si
400 O
Na Mg
Cl
0 0
50
100
150
200
250
300
Energy / KeV
(c)
(d) 3200 2800
C/C-B4Cf Element At% C 80.50 B 16.08 O 2.73 Na 0.35 Cl 0.31 Mg 0.03
C
Counts
2400 2000 1600 1200 800 400
B
O
Na Mg
Cl
0 0
50
100
150
200
250
300
Energy / KeV
(e)
(f)
3200
C/SiC Element At% C 73.56 Si 8.36 O 17.24 Na 0.41 Cl 0.41 Mg 0.02
2800
Counts
2400
C
2000
Si
1600 1200 O
800 400
Cl
Na Mg
0 0
50
100
150
200
250
300
Energy / KeV Fig. 7. Debris for braking under the first seawater condition: (a) debris of C/C-SiCf discs; (b) EDX analysis of debris in (a); (c) debris of C/C-B4Cf discs; (d) EDX analysis of debris in (c); (e) debris of C/SiC discs; and (f) EDX analysis of debris in (e).
Y. Cai et al. / Composites Science and Technology 87 (2013) 50–57
SiðsÞ þ 2H2 OðlÞ ! SiO2 ðsÞ þ 2H2 ðgÞ
ð5Þ
The double lubrication action from water film and SiO2 film caused a lower COF in seawater condition for C/SiC discs than for C/C-SiCf discs. But the lubricating film was easily disrupted due to the high SiC content in C/SiC, so the COF was restored fast. 4. Conclusions (1) 3D needled ceramic modified C/C composites: C/C-B4Cf and C/C-SiCf were prepared using UPSIF combined with CVI. The contents of SiC and B4C introduced by a one-shot infiltration were about 18 wt% and 16 wt% respectively. (2) The pore size distribution in C/C composites became more uniform due to ceramic filler addition. A structure of higher crystallite graphite was in C/C-B4Cf than in C/C and C/C-SiCf. Different from the gathered SiC and residual Si in C/SiC, ceramic grains were dispersed and uniform in the modified C/C composites. (3) The COF fade in seawater conditions for C/C composites was significantly reduced by ceramic modification. C/C-SiCf was superior to C/SiC for its excellent stable friction without any fade in seawater conditions. (4) The weak hygroscopicity and the ability to produce fresh debris continuously in seawater conditions were the primary factors for ceramic modified C/C composites to relieve and even avoid the fade. The double lubrication action from water film and SiO2 film increased the COF fade for C/SiC.
Acknowledgements The authors acknowledge the financial supports of Natural Science Foundation of China (Contract No. NSFC50672076) and Shaanxi Province Department of Education Fund (2013JK0925). The authors would like to give their special thanks to Prof. Dr. Yongdong Xu for his kind help on many aspects of scientific research and his continual support in the time of difficulties. Without these, this work would have been not possible.
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