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Temperature and concentration dependence of SiC deposition on Nicalon fibers
Surface and Coatings Technology, 43/44 (1990) 167—175
167
TEMPERATURE AND CONCENTRATION DEPENDENCE OF SiC DEPOSITION ON NICALON FIBERS T. M. BESMANN, B. W. SHELDON and M. D. KASTER Oak Ridge National Laboratory, P.O. Box 2008, Oak Ridge, TN 37831-6063 (U.S.A.)
Abstract Chemical vapor infiltration of NicalonTM fiber bundles (approximately 500 filaments) was studied using methyltrichiorosilane to form the SiC matrix. The operating conditions were chosen to simulate those used in composite fabrication. It was determined that the deposition reaction has an activation energy of 66 ±20 kJ mol1 and is first order with respect to reactant concentration.
1. Introduction Fiber-reinforced composites are being fabricated using chemical vapor deposition to form the matrix. This process, which is referred to as chemical vapor infiltration (CVI), is one of only a few processes capable of incorporating continuous ceramic fibers in a ceramic matrix without chemically, ther. mally or mechanically damaging the reinforcing fibers. The high strength (approximately 400 MPa) and exceptional (for ceramics) fracture toughness (greater than 10 MPa m112) of these composites, combined with their refractory nature and resistance to erosion, corrosion and wear, make them ideal candidates for numerous advanced high temperature structural applications. Commercial interest is focused on composites fabricated from continuous carbon, silicon—carbon--oxygen or aluminosilicate fibers, all with a matrix of SiC. Silicon carbide is typically deposited from methyltrichiorosilane (MTS) in hydrogen. Deposition of stoichiometric SiC is facilitated by the 1:1 gram-atomic ratio of silicon to carbon in MTS. The current commercial CVI processes are isothermal—isobaric and require reactant transport into the free-standing fibrous preform by diffusion [1]. A different process for the fabrication of fiber-reinforced ceramic composites has been under development at Oak Ridge National Laboratory (ORNL) for several years [2, 3]. This process simultaneously utilizes thermal and pressure gradients and can reduce the infiltration time from weeks to less than 24 h. The efficient extension of the thermal-gradient, forced-flow technique to large and complex shapes and the further optimization of the isothermal—iso0257-8972/90/$3.50
c7 Elsevier Sequoia/Printed in The Netherlands
168
baric process will require accurate process models. Such analytical models are being developed by a number of researchers, who are focusing on linking mass and heat transfer with the reaction kinetics [4—il]. However, any model for infiltration requires accurate data for the inherent chemical and physical processes. One of the most important of these is the rate-limiting chemical reaction for the deposition of the matrix material. It is the objective of the current effort to examine this reaction by measuring the SiC deposition rates (from MTS in hydrogen) as a function of temperature and reactant concentration.
2. Experimental details 2.7. Materials The hydrogen used is of ultrahigh purity and the MTS (Alfa Products, Danvers, MA) is of commercial grade and has a reported purity of 98%. The fibers are polycarbosilane-derived Si--C—O (Nicalon, ceramic grade, Nippon Carbon Co.. Tokyo, Japan); these are washed in acetone to remove sizing and are then coated with approximately 0.2 ~im of graphitic carbon via the pyrolysis of propylene. The carbon precoat accurately represents a fibrous preform prepared for infiltration because such precoats are necessary in composites to protect the fibers from the infiltration gases and to obtain optimum mechanical properties. The fibers used are in the form of a tow (approximately 500 fibers), which is removed from precoated, plain-weave cloth. 2.2. Apparatus The fiber-coating apparatus is shown in Fig. 1. The fiber tow is carefully spread apart to provide adequate space between filaments for the reactant gases to flow. The fibers are then attached to a graphite ring using a graphite glue, with the ring situated, as shown in Fig. 1, in a graphite tube through which the reactant gases flow. Perforated graphite disks are shown on both the upstream and downstream sides of the filaments. The upstream disk serves to help preheat the reactant gas and to distribute the gas over the diameter of the tube. The downstream disk helps to heat the fibers radiatively. The graphite tube is contained within a fused silica tube sealed with stainless steel end fittings. The graphite tube and disks are heated via coupling to a 455 kHz r.f. field. The hydrogen and MTS are metered into the system by mass flow controllers (model 1259B with controller model 247B; MKS Instruments, Burlington, MA) and a vapor source (Tylan Source 5; Tylan, Carson, CA) respectively. The MTS vapor source is checked for calibration using the difference in the weight of MTS in the reservoir before and after a run. The calibration of the mass flow controllers is periodically checked with a wet-test meter. Both the reactant hydrogen and argon purge gases are gettered of oxygen by flowing through a bed of titanium sponge held at
169
RETAINING RING rRF COIL
///I/////////////II,,I//////I/I//I//II~
//I/II/fI/fII/////
I//////I////_
NICALON FIBERS SUPPLEMENTAL RADIANT HEATER
GRAPHITE SUSCEPTOR ///Il/I//////JJIfI
oil
Il//I/I/I/I//I/I
I/Ill
FUSED SILICA TUBE GAS DISTRIBUTOR/PREHEATER
Fig. 1. Apparatus for chemical vapor deposition onto separated fiber tows.
approximately 900 K. Temperature is measured with an optical pyrometer which can directly view the fibers through a window in the downstream stainless steel fitting and the holes in the downstream graphite disk. Corrections are applied for absorbance by the window, and the pyrometer is calibrated regularly and determined to be accurate at the melting point of gold (1338 K) to within 3 K. Procedure The reactor is initially evacuated with a rotary pump to approximately 15 Pa to test the integrity of the system. Gettered argon is then passed through the system bringing it to local atmospheric pressure (approximately 98 kPa), and the fiber sample is brought to the required temperature. Once a stable, desired temperature is reached, the gas stream is switched to hydrogen and the MTS is allowed to flow through a bypass around the reactor for 10 mm. The MTS—hydrogen stream is then switched to the reactor for the required deposition time. The hydrogen flow is kept constant at 500 cm3 min’ throughout all the runs. Small temperature adjustments are occasionally required, particularly when the MTS flow is initiated, because of the more efficient cooling by the MTS. After deposition, the MTS—hydro. gen stream is shut off, gettered argon is again allowed to flow through the reactor, and the power to the induction coil is shut off. The sample then cools to room temperature in the flowing argon. 2.3.
170
2.4. Analysis The fibers are removed from the graphite ring by cutting them at least 2--3 mm from the point at which they are glued. X-ray diffraction is used to confirm the nature of the deposited phase(s). Interference from the substrate fibers is negligible because of their amorphous character. Coating thickness is determined via scanning electron microscopy or optical metallography of the fiber cross-section.
3. Results Table 1 summarizes the results of the deposition runs. With the exception of sample 30, the average flow rates of MTS, determined by weight loss from the MTS reservoir, are within 10% of the set point. The low deposition temperature of runs 31 and 32 promotes the formation of elemental silicon as the major phase. The deposition times are varied such that sufficient coating is formed to yield an accurate measurement and to overcome any initial surface effects, but the times are limited so that the fiber coatings do not grow together. A typical metallographic cross-section is shown in Fig. 2. There is no apparent thickness dependence on the position within the fiber bundle as shown by comparisons of cross-sections at the ends and center.
TABLE 1 Data for the deposition of coatings on Nicalon fibersa Sample
Temperature (corrected) (K)
Time (mm)
MTS flow rate 3 min1) (cm
Linear growth rate (nmmin1)
25 26 27 30 31 32 33 35 36 37 39 40 42 44 45 47 49
1393 1292 1497 1436 1186 1230 1292 1393 1393 1393 1393 1393 1393 1393 1393 1393 1310
30 45 20 10 60 45 45 30 25 60 20 20 20 35 20 22 45
48.3 51.5 48.0 59.2 54.8 52.6 55.9 77.2 97.2 39.2 86.4 85.9 92.4 41.0 68.4 97.9 48.5
277 222 543 339 235 208 148 314 526 206 416 239 445 108 328 514 166
~Fibers precoated with approximately 0.2 pm of pyrolytic carbon.
171
S 20Mm
I000x
Fig. 2. Metallographic cross-section of coated fibers (sample 30).
The data for MTS flow rates of approximately 50 cm3 mm ‘ are shown in the Arrhenius plot in Fig. 3. The five points for which fl-SiC, as indicated by powder X-ray diffraction, is the predominantly deposited phase can be linearly least-squares fitted to the Arrhenius relation p
=
A exp(—E/RT)
(1)
where p is the deposition rate, A is the pre-exponential factor, E is the activation energy and R is the ideal gas law constant. The computed values are E = 72 ±13 kJ moI~ and ln A = 12 ±1.4 nm min* The data points 31 and 32 (shown as filled circles) display disproportionately high deposition rates and as noted previously, the coatings are predominantly elemental silicon. Such a temperature dependence of the deposited phases has also been observed by Langlais et al. [12]. Figure 4 shows a plot of the deposition rate as a function of MTS concentration in the reactant gas. From the slope, it appears that the rate has a non.zeroth-order dependence on the MTS concentration. Thus the rate expression for SiC deposition can be written
p =A’C’1 exp(—E/RT)
(2)
where A’ is the pre-exponential factor when the MTS concentration C is included and n is the reaction order. Although MTS is not stable under these conditions, but decomposes to carbon- and silicon-containing species, its initial concentration may still be representative of the actual reactant species concentrations [12, 13]. A linear least-squares fit of all the data in Table 1, with the exception of points 31 and 32, to the natural logarithm
172
7.0
‘~
6.5
—
6.0
—
P=1O1 kPa MTS FLOW = 5Ocm3/min H 3/min 2 FLOW = 500 cm
0 ~2(±13~kJ/moIOo~
(nm/mm)
[~
• Si ~-SiC~
4.5 4.0 6.5
7.0
7.5 l/T (K1)
8.0
(x1O4)
Fig. 3. Arrhenius plot of deposition rate is, reciprocal temperature.
I
600 500 400
P=lOlkPa -
~:
1393K
o
(nm/mm)300 Ii 200— 100 0 0.050
0
0
I
I
0.075 0.100 0.125 0.150 0.175 REACTANT GAS MOLE FRACTION MTS
Fig. 4. Plot of MTS concentration in the reactant gas vs. SiC growth rate.
0.200
173
of eqn. (2) yields ln A’ = 14 ±1.7 nm min’, E = 66 ±20 kJ mol~ and n = 1.1 ±0.2. Thus the rate-limiting deposition mechanism is apparently first order with respect to reactant concentration.
4. Discussion The measured activation energy and first-order concentration depen. dence indicate that a chemical kinetic mechanism is rate controlling rather than gas-phase mass transport. This is further confirmed by the cylindrical symmetry of the coatings deposited on the fibers, as illustrated in Fig. 2. If mass transport were rate limiting, there would be obvious differences in the coating thicknesses on the upstream and downstream surfaces of the fibers. The activation energy determined in this work differs from that (120 kJ mo1~)determined by Brennfleck et al. [14], who measured the deposi. tion rate of SiC on a carbon fiber at atmospheric pressure with an H2 to MTS ratio of 3.6. However, it agrees with the value of 68 kJ mol obtained by van Kemenade and Stemfoort [15] for deposition on graphite. As discussed by Besmann and Johnson [13], others who have measured rate dependences at atmospheric pressure have observed a first-order relationship to MTS concentration. Besmann and Johnson determined, at reduced pressure (3.3 kPa), a zeroth.order dependence on MTS concentration and an activation energy of 188 kJ mol’ for deposition on graphite. The different activation energies and reaction orders which have been measured by various investigators indicate that it may be difficult to describe SiC deposition accurately for MTS over a range of conditions with only one simple rate expression. The deposition of elemental silicon instead of SiC at relatively low temperatures has been observed previously [12, 161. The formation of silicon is not thermodynamically favored under the conditions used here; therefore, the presence of silicon must be caused by kinetic factors, probably those related to nucleation. By extrapolating the observed SiC rates in Fig. 3 to the lower temperatures where silicon formation is observed, it is apparent that silicon deposition is significantly faster than SiC deposition at the same temperature. This conclusion is still accurate when 2the linearthe growth rates s’~).If deposition arethe converted to silicon deposition (mol cm of silicon atoms in this systemrates occurs at Sicomparable rates during the formation of both silicon and SiC, then the lower SiC deposition rates that are observed can be attributed to the low reactivity of carbon. This assumption appears to agree with the current supposition that SiC forms by separate silicon and carbon deposition mechanisms, where the carbon reaction is rate limiting [17]. If the phase composition of the deposit is strictly determined by the growth rates for the competing silicon and SiC formation reactions, then the measured growth rates (mol Si cm2 s’) should be about equal at the temperature at which silicon deposition ceases to occur and SiC deposition begins. This is not observed in Fig. 3. At most of the temperatures at which -‘
174
SiC is deposited, the extrapolated silicon deposition rate is actually faster. In these cases the fact that the slower SiC deposition reaction is preferred suggests that SiC nucleation is favored over silicon nucleation at these temperatures. A factor that probably contributes to the absence of silicon at higher temperatures is that the degree of supersaturation with respect to silicon formation decreases with increasing temperature as the equilibrium concentration of vapor-phase silicon chlorides and chlorosilanes increases [13).
5. Conclusions The reaction order and activation energy were determined for the atmospheric pressure deposition of SiC on pyrolytic carbon-coated Nicalon from MTS carried in hydrogen. The observed preferential deposition of silicon at reduced temperatures qualitatively agrees with an SiC deposition mechanism in which carbon formation is rate limiting. SiC formation is thermodynamically favored at all of the conditions investigated. The formation of silicon at lower temperatures demonstrates the importance of kinetic factors in determining the phase composition. The observation that a slower SiC deposition reaction is dominant at higher temperatures indicates that the nucleation kinetics are important. The reactivity of carbon-containing species and the equilibrium concentrations of silicon-containing species appear to be important factors.
Acknowledgments The authors wish to thank G. C. Marsh for the metallography and L. Riester for the scanning electron microscopy and X-ray diffraction. C. A. Valentine prepared the manuscript for publication and H. R. Livesey provided the graphics. The research was sponsored by the U.S. Department of the Air Force, Air Force Wright Research and Development Center, WrightPatterson Air Force Base, OH, AFWAL MIPR FY1457-89-N5001, U.S. Department of Energy Interagency Agreement 1692-1692-Al under contract DE-ACO5-840R21400 with Martin Marietta Energy Systems, Inc.
References 1 P. J. Lamicq, G. A. Bernhart, M. M. Dauchier and J. G. Mace, Am. Ceram. Soc. Bull., 64 (2) (1986) 336. 2 D. P. Stinton, T. M. Besmann and R. A. Lowden, Am. Ceram. Soc. Bull., 67(2) (1988) 350. 3 T. M. Besmann, R. A. Lowden, D. P. Stinton and T. L. Starr, J. Phys. (Paris) Colloq., 50(1989) 229. 4 J. V. Rossignol, F. Langlais and R. Naslain, in MeD. Robinson, C. H. J. van den Brekel, G. W. Cullen and J. M. Blocher, Jr. (eds.), Proc. 9th mt. Conf. on Chemical Vapor Deposition. 1984, Electrochemical Society, Pennington, NJ, 1984, p. 596.
175 5 T. L. Starr, Ceram. Eng. Sci. Proc., 9 (7—8) (1988a) 803. 6 T. L. Starr, Proc. mt. Congr. on Whisker- and Fiber-Toughened Ceramics, ASM International, Materials Park, OH, 1988, pp. 243—252. 7 T. L. Starr and A. W. Smith, in T. M Besmann and B. M. Gallois (eds.), Chemical Vapor Deposition of Refractory Metals and Ceramics, Materials Research Society, Pittsburgh, PA, 1990, p. 55. 8 S. Middleman, J. Mater. Res. Soc., 4 (6) (1989) 1515. 9 R. R. Melkote and K. F. Jensen, in T. M. Besmann and B. M. Gallois (eds.), Chemical Vapor Deposition of Refractory Metals and Ceramics, Materials Research Society, Pittsburgh, PA, 1990, p. 67. 10 N.H. Tai and T.-W. Chou, in T. M. Besmann and B. M. Gallois (eds.), Chemical Vapor Deposition of Refractory Metals and Ceramics, Materials Research Society, Pittsburgh, PA, 1990, p. 61. 11 5. V. Sotirchos and M. M. Tomadakis, in T. M. Besmann and B. M. Gallois (eds.), Chemical Vapor Deposition of Refractory Metals and Ceramics, Materials Research Society, Pittsburgh, PA, 1990, p. 73. 12 F. Langlais, C. Prebende, B. Tarride and R. Naslain, in M. Ducarroir, C. Bernard and L. Vandenbulcke (eds.), Proc. 7th European Conf. on Chemical Vapor Deposition, Colloq. Phys., 650 (5) (1989) 93. 13 T. M. Besmann and M. L. Johnson, in V. J. Tennery (ed), Proc. 3rd mt. Symp. on Ceramic Materials and Components for Engines, American Ceramic Society, Westerville, OH, 1989, p. 443. 14 K. Brennfieck, E. Fitzer, G. Schoch and M. Dietrich, in McD. Robinson, C. H. J. van den Brekel, G. W. Cullen and J. M. Blocher, Jr. (eds.), Proc. 9th mt. Conf. on Chemical Vapor Deposition, 1984, Electrochemical Society, Pennington, NJ, 1984, p. 649. 15 A. W. C. van Kemenade and C. F. Stemfoort, J. Cryst. Growth, 12 (1972) 13. 16 J. Schlichting, Powder Metall. mt., 12 (3,4) (1980) 141. 17 G. S. Fischman and W. T. Petuskey, J. Am. Ceram. Soc., 91 (1985) 185.
167
TEMPERATURE AND CONCENTRATION DEPENDENCE OF SiC DEPOSITION ON NICALON FIBERS T. M. BESMANN, B. W. SHELDON and M. D. KASTER Oak Ridge National Laboratory, P.O. Box 2008, Oak Ridge, TN 37831-6063 (U.S.A.)
Abstract Chemical vapor infiltration of NicalonTM fiber bundles (approximately 500 filaments) was studied using methyltrichiorosilane to form the SiC matrix. The operating conditions were chosen to simulate those used in composite fabrication. It was determined that the deposition reaction has an activation energy of 66 ±20 kJ mol1 and is first order with respect to reactant concentration.
1. Introduction Fiber-reinforced composites are being fabricated using chemical vapor deposition to form the matrix. This process, which is referred to as chemical vapor infiltration (CVI), is one of only a few processes capable of incorporating continuous ceramic fibers in a ceramic matrix without chemically, ther. mally or mechanically damaging the reinforcing fibers. The high strength (approximately 400 MPa) and exceptional (for ceramics) fracture toughness (greater than 10 MPa m112) of these composites, combined with their refractory nature and resistance to erosion, corrosion and wear, make them ideal candidates for numerous advanced high temperature structural applications. Commercial interest is focused on composites fabricated from continuous carbon, silicon—carbon--oxygen or aluminosilicate fibers, all with a matrix of SiC. Silicon carbide is typically deposited from methyltrichiorosilane (MTS) in hydrogen. Deposition of stoichiometric SiC is facilitated by the 1:1 gram-atomic ratio of silicon to carbon in MTS. The current commercial CVI processes are isothermal—isobaric and require reactant transport into the free-standing fibrous preform by diffusion [1]. A different process for the fabrication of fiber-reinforced ceramic composites has been under development at Oak Ridge National Laboratory (ORNL) for several years [2, 3]. This process simultaneously utilizes thermal and pressure gradients and can reduce the infiltration time from weeks to less than 24 h. The efficient extension of the thermal-gradient, forced-flow technique to large and complex shapes and the further optimization of the isothermal—iso0257-8972/90/$3.50
c7 Elsevier Sequoia/Printed in The Netherlands
168
baric process will require accurate process models. Such analytical models are being developed by a number of researchers, who are focusing on linking mass and heat transfer with the reaction kinetics [4—il]. However, any model for infiltration requires accurate data for the inherent chemical and physical processes. One of the most important of these is the rate-limiting chemical reaction for the deposition of the matrix material. It is the objective of the current effort to examine this reaction by measuring the SiC deposition rates (from MTS in hydrogen) as a function of temperature and reactant concentration.
2. Experimental details 2.7. Materials The hydrogen used is of ultrahigh purity and the MTS (Alfa Products, Danvers, MA) is of commercial grade and has a reported purity of 98%. The fibers are polycarbosilane-derived Si--C—O (Nicalon, ceramic grade, Nippon Carbon Co.. Tokyo, Japan); these are washed in acetone to remove sizing and are then coated with approximately 0.2 ~im of graphitic carbon via the pyrolysis of propylene. The carbon precoat accurately represents a fibrous preform prepared for infiltration because such precoats are necessary in composites to protect the fibers from the infiltration gases and to obtain optimum mechanical properties. The fibers used are in the form of a tow (approximately 500 fibers), which is removed from precoated, plain-weave cloth. 2.2. Apparatus The fiber-coating apparatus is shown in Fig. 1. The fiber tow is carefully spread apart to provide adequate space between filaments for the reactant gases to flow. The fibers are then attached to a graphite ring using a graphite glue, with the ring situated, as shown in Fig. 1, in a graphite tube through which the reactant gases flow. Perforated graphite disks are shown on both the upstream and downstream sides of the filaments. The upstream disk serves to help preheat the reactant gas and to distribute the gas over the diameter of the tube. The downstream disk helps to heat the fibers radiatively. The graphite tube is contained within a fused silica tube sealed with stainless steel end fittings. The graphite tube and disks are heated via coupling to a 455 kHz r.f. field. The hydrogen and MTS are metered into the system by mass flow controllers (model 1259B with controller model 247B; MKS Instruments, Burlington, MA) and a vapor source (Tylan Source 5; Tylan, Carson, CA) respectively. The MTS vapor source is checked for calibration using the difference in the weight of MTS in the reservoir before and after a run. The calibration of the mass flow controllers is periodically checked with a wet-test meter. Both the reactant hydrogen and argon purge gases are gettered of oxygen by flowing through a bed of titanium sponge held at
169
RETAINING RING rRF COIL
///I/////////////II,,I//////I/I//I//II~
//I/II/fI/fII/////
I//////I////_
NICALON FIBERS SUPPLEMENTAL RADIANT HEATER
GRAPHITE SUSCEPTOR ///Il/I//////JJIfI
oil
Il//I/I/I/I//I/I
I/Ill
FUSED SILICA TUBE GAS DISTRIBUTOR/PREHEATER
Fig. 1. Apparatus for chemical vapor deposition onto separated fiber tows.
approximately 900 K. Temperature is measured with an optical pyrometer which can directly view the fibers through a window in the downstream stainless steel fitting and the holes in the downstream graphite disk. Corrections are applied for absorbance by the window, and the pyrometer is calibrated regularly and determined to be accurate at the melting point of gold (1338 K) to within 3 K. Procedure The reactor is initially evacuated with a rotary pump to approximately 15 Pa to test the integrity of the system. Gettered argon is then passed through the system bringing it to local atmospheric pressure (approximately 98 kPa), and the fiber sample is brought to the required temperature. Once a stable, desired temperature is reached, the gas stream is switched to hydrogen and the MTS is allowed to flow through a bypass around the reactor for 10 mm. The MTS—hydrogen stream is then switched to the reactor for the required deposition time. The hydrogen flow is kept constant at 500 cm3 min’ throughout all the runs. Small temperature adjustments are occasionally required, particularly when the MTS flow is initiated, because of the more efficient cooling by the MTS. After deposition, the MTS—hydro. gen stream is shut off, gettered argon is again allowed to flow through the reactor, and the power to the induction coil is shut off. The sample then cools to room temperature in the flowing argon. 2.3.
170
2.4. Analysis The fibers are removed from the graphite ring by cutting them at least 2--3 mm from the point at which they are glued. X-ray diffraction is used to confirm the nature of the deposited phase(s). Interference from the substrate fibers is negligible because of their amorphous character. Coating thickness is determined via scanning electron microscopy or optical metallography of the fiber cross-section.
3. Results Table 1 summarizes the results of the deposition runs. With the exception of sample 30, the average flow rates of MTS, determined by weight loss from the MTS reservoir, are within 10% of the set point. The low deposition temperature of runs 31 and 32 promotes the formation of elemental silicon as the major phase. The deposition times are varied such that sufficient coating is formed to yield an accurate measurement and to overcome any initial surface effects, but the times are limited so that the fiber coatings do not grow together. A typical metallographic cross-section is shown in Fig. 2. There is no apparent thickness dependence on the position within the fiber bundle as shown by comparisons of cross-sections at the ends and center.
TABLE 1 Data for the deposition of coatings on Nicalon fibersa Sample
Temperature (corrected) (K)
Time (mm)
MTS flow rate 3 min1) (cm
Linear growth rate (nmmin1)
25 26 27 30 31 32 33 35 36 37 39 40 42 44 45 47 49
1393 1292 1497 1436 1186 1230 1292 1393 1393 1393 1393 1393 1393 1393 1393 1393 1310
30 45 20 10 60 45 45 30 25 60 20 20 20 35 20 22 45
48.3 51.5 48.0 59.2 54.8 52.6 55.9 77.2 97.2 39.2 86.4 85.9 92.4 41.0 68.4 97.9 48.5
277 222 543 339 235 208 148 314 526 206 416 239 445 108 328 514 166
~Fibers precoated with approximately 0.2 pm of pyrolytic carbon.
171
S 20Mm
I000x
Fig. 2. Metallographic cross-section of coated fibers (sample 30).
The data for MTS flow rates of approximately 50 cm3 mm ‘ are shown in the Arrhenius plot in Fig. 3. The five points for which fl-SiC, as indicated by powder X-ray diffraction, is the predominantly deposited phase can be linearly least-squares fitted to the Arrhenius relation p
=
A exp(—E/RT)
(1)
where p is the deposition rate, A is the pre-exponential factor, E is the activation energy and R is the ideal gas law constant. The computed values are E = 72 ±13 kJ moI~ and ln A = 12 ±1.4 nm min* The data points 31 and 32 (shown as filled circles) display disproportionately high deposition rates and as noted previously, the coatings are predominantly elemental silicon. Such a temperature dependence of the deposited phases has also been observed by Langlais et al. [12]. Figure 4 shows a plot of the deposition rate as a function of MTS concentration in the reactant gas. From the slope, it appears that the rate has a non.zeroth-order dependence on the MTS concentration. Thus the rate expression for SiC deposition can be written
p =A’C’1 exp(—E/RT)
(2)
where A’ is the pre-exponential factor when the MTS concentration C is included and n is the reaction order. Although MTS is not stable under these conditions, but decomposes to carbon- and silicon-containing species, its initial concentration may still be representative of the actual reactant species concentrations [12, 13]. A linear least-squares fit of all the data in Table 1, with the exception of points 31 and 32, to the natural logarithm
172
7.0
‘~
6.5
—
6.0
—
P=1O1 kPa MTS FLOW = 5Ocm3/min H 3/min 2 FLOW = 500 cm
0 ~2(±13~kJ/moIOo~
(nm/mm)
[~
• Si ~-SiC~
4.5 4.0 6.5
7.0
7.5 l/T (K1)
8.0
(x1O4)
Fig. 3. Arrhenius plot of deposition rate is, reciprocal temperature.
I
600 500 400
P=lOlkPa -
~:
1393K
o
(nm/mm)300 Ii 200— 100 0 0.050
0
0
I
I
0.075 0.100 0.125 0.150 0.175 REACTANT GAS MOLE FRACTION MTS
Fig. 4. Plot of MTS concentration in the reactant gas vs. SiC growth rate.
0.200
173
of eqn. (2) yields ln A’ = 14 ±1.7 nm min’, E = 66 ±20 kJ mol~ and n = 1.1 ±0.2. Thus the rate-limiting deposition mechanism is apparently first order with respect to reactant concentration.
4. Discussion The measured activation energy and first-order concentration depen. dence indicate that a chemical kinetic mechanism is rate controlling rather than gas-phase mass transport. This is further confirmed by the cylindrical symmetry of the coatings deposited on the fibers, as illustrated in Fig. 2. If mass transport were rate limiting, there would be obvious differences in the coating thicknesses on the upstream and downstream surfaces of the fibers. The activation energy determined in this work differs from that (120 kJ mo1~)determined by Brennfleck et al. [14], who measured the deposi. tion rate of SiC on a carbon fiber at atmospheric pressure with an H2 to MTS ratio of 3.6. However, it agrees with the value of 68 kJ mol obtained by van Kemenade and Stemfoort [15] for deposition on graphite. As discussed by Besmann and Johnson [13], others who have measured rate dependences at atmospheric pressure have observed a first-order relationship to MTS concentration. Besmann and Johnson determined, at reduced pressure (3.3 kPa), a zeroth.order dependence on MTS concentration and an activation energy of 188 kJ mol’ for deposition on graphite. The different activation energies and reaction orders which have been measured by various investigators indicate that it may be difficult to describe SiC deposition accurately for MTS over a range of conditions with only one simple rate expression. The deposition of elemental silicon instead of SiC at relatively low temperatures has been observed previously [12, 161. The formation of silicon is not thermodynamically favored under the conditions used here; therefore, the presence of silicon must be caused by kinetic factors, probably those related to nucleation. By extrapolating the observed SiC rates in Fig. 3 to the lower temperatures where silicon formation is observed, it is apparent that silicon deposition is significantly faster than SiC deposition at the same temperature. This conclusion is still accurate when 2the linearthe growth rates s’~).If deposition arethe converted to silicon deposition (mol cm of silicon atoms in this systemrates occurs at Sicomparable rates during the formation of both silicon and SiC, then the lower SiC deposition rates that are observed can be attributed to the low reactivity of carbon. This assumption appears to agree with the current supposition that SiC forms by separate silicon and carbon deposition mechanisms, where the carbon reaction is rate limiting [17]. If the phase composition of the deposit is strictly determined by the growth rates for the competing silicon and SiC formation reactions, then the measured growth rates (mol Si cm2 s’) should be about equal at the temperature at which silicon deposition ceases to occur and SiC deposition begins. This is not observed in Fig. 3. At most of the temperatures at which -‘
174
SiC is deposited, the extrapolated silicon deposition rate is actually faster. In these cases the fact that the slower SiC deposition reaction is preferred suggests that SiC nucleation is favored over silicon nucleation at these temperatures. A factor that probably contributes to the absence of silicon at higher temperatures is that the degree of supersaturation with respect to silicon formation decreases with increasing temperature as the equilibrium concentration of vapor-phase silicon chlorides and chlorosilanes increases [13).
5. Conclusions The reaction order and activation energy were determined for the atmospheric pressure deposition of SiC on pyrolytic carbon-coated Nicalon from MTS carried in hydrogen. The observed preferential deposition of silicon at reduced temperatures qualitatively agrees with an SiC deposition mechanism in which carbon formation is rate limiting. SiC formation is thermodynamically favored at all of the conditions investigated. The formation of silicon at lower temperatures demonstrates the importance of kinetic factors in determining the phase composition. The observation that a slower SiC deposition reaction is dominant at higher temperatures indicates that the nucleation kinetics are important. The reactivity of carbon-containing species and the equilibrium concentrations of silicon-containing species appear to be important factors.
Acknowledgments The authors wish to thank G. C. Marsh for the metallography and L. Riester for the scanning electron microscopy and X-ray diffraction. C. A. Valentine prepared the manuscript for publication and H. R. Livesey provided the graphics. The research was sponsored by the U.S. Department of the Air Force, Air Force Wright Research and Development Center, WrightPatterson Air Force Base, OH, AFWAL MIPR FY1457-89-N5001, U.S. Department of Energy Interagency Agreement 1692-1692-Al under contract DE-ACO5-840R21400 with Martin Marietta Energy Systems, Inc.
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