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The effects of heat treatment and coating roughness on the strength of alumina-zirconia fibres
Composites Science and Technology 62 (2002) 205–212 www.elsevier.com/locate/compscitech
The effects of heat treatment and coating roughness on the strength of alumina-zirconia fibres Ramanan Venkatesh School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia, 14300 Nibong Tebal, Pulau Pinang, Malaysia Received 5 February 2001; received in revised form 18 April 2001; accepted 25 September 2001
Abstract The strengths of Al2O3–20 wt.% ZrO2 fibres as-received, heat treated and coated with SnO2 have been determined. The mean strength of as-received fibres was 1380 MPa for a gage length of 17 mm. The strengths of fibres heat treated at 500, 600 and 900 C decreased by 5, 7 and 21%, respectively, relative to the as-received fibres. This could be because of a secondary glassy phase formation at the grain boundaries. It was found that the accuracy of the tensile strength distribution increased with the number of samples tested and at a sample number of 80, the variation in the tensile strength was about 2%. The thickness and roughness of the SnO2 coating had a profound effect on the strength of the coated fibres. It was observed that as the coating thickness increased, roughness of the coating increased. This acted as a notch in order to decrease the strength of the fibres. Control of roughness at the fibre/coating and/or coating/matrix interphase is very important in the development of tough ceramic–matrix composites. As the interphase roughness increases, fibre debonding and pull-out decrease because of the increased compressive clamping stress. Hence for increased toughness in CMCs through fibre debonding and pull-out mechanisms, a relatively smooth interphase is necessary. The interphase roughness can be controlled either by selecting smooth fibres or by decreasing coating roughness. # 2002 Published by Elsevier Science Ltd. Keywords: A. Fibres; A. Coating; B. Interphase; Alumina
1. Introduction High-performance ceramic fibres are important in the development of ceramic–matrix composites for hightemperature applications. Characteristics of fibres that influence properties of CMCs include strength, strain to failure, statistical variation in strength, aspect ratio, microstructure and thermal stability. In applications where stability in air at high temperatures is a prime objective, oxide fibres have to be considered since they offer better oxidation resistance at high temperatures. Principal methods of forming oxide ceramic fibres include the spinning of aqueous alumina slurry containing spinning additives or spinning of polymeric solutions using sol-gel techniques [1–6]. Alumina, mullite and zirconia fibres are the principal oxide fibres developed for ceramic matrix reinforcement [7–15]. In order to be used for composite applications, these fibres should have good strengths of about 2 GPa and long-term application temperatures greater than 1200 C. E-mail address: [email protected]
Considerable work exists in the literature for the development of oxide fibre/oxide matrix composites [16–25]. In oxide fibre/oxide matrix composite systems a reaction at the fibre/matrix interface, e.g. Al2O3/SiO2, mullite/mullite could lead to a strong interfacial bonding and hence a brittle composite. Barrier layers that prevent chemical reaction between the fibre and the matrix, e.g. SnO2, BN, monazite, hibonite, etc. have to be introduced at the fibre/ matrix interface [20–25]. In the present study, SnO2 was chosen as a coating on the alumina fibre surface because it does not react with alumina up to 1400 C in partial pressure of oxygen > 107 atmospheres [26]. The present study is an investigation of the effect of heat treatment and SnO2 coating thickness on the strength of alumina (PRD-166) fibres and its implications towards the development of tough ceramic matrix composites.
2. Tensile strength of ceramic fibres The mean strength of ceramic fibre decreases with increase in volume of material due to increased probability
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of defects in the fibre. The effect of random distribution of a single defect on the strength of a solid has been described by Weibull [27]. According to the Weibull distribution, probability of failure of a fibre at a stress, , can be given as, FðÞ ¼ 1- exp ð1Þ where is a scale parameter and is known as Weibull modulus and is inversely related to scatter in the strength of fibres. The value of the Weibull modulus for most ceramic fibres is low, often around 5 which reflects the large variation in their strength. Eq. (1) can be written as, ln ln½ð1=1-Fð ÞÞ ¼ ln þ ln
ð2Þ
In an actual test, strengths of N fibres tested are arranged in an increasing order and the failure probability of the ith sample, Pfi, can be estimated from an estimator of the kind shown below, Pfi ¼ i=N þ 1
ð3Þ
Substituting this estimator into Eq. (2), we can write, ln ln½ðN þ 1=N þ 1-iÞ ¼ ln i þ ln
ð4Þ
A plot of ln ln [(N+1/N+1-i)] against ln i will then yield a straight line with slope and intercept ln . Once and are known, the Weibull mean tensile strength, f, is then given as, f ¼ 1= ð1 þ 1=Þ
ð5Þ
where (1+1/) is the gamma function of (1+1/) and can be obtained from standard mathematical tables. Standard deviation Sf is given as, 1=2 Sf ¼ 1= ð1 þ 2=Þ- 2 ð1 þ 1=Þ
ð6Þ
Coefficient of variation C.V., can then be given as, C:V: ¼ 100 ðSf =f Þ
ð7Þ
It has been shown that the standard deviation, Sf divided by mean modulus can be related to the number of fibres tested as [28], Sf = ¼ ð1=nÞ1=2
ð8Þ
where n is the number of fibres tested. The variation of mean tensile strength of the fibres, f for a 90% confidence limit can then be written as [28], f ¼ f 1:645 Sf =n1=2
ð9Þ
It can be seen from Eq. (9) that the variation in the mean tensile strength decreases with an increase in the number of fibres tested.
3. Experimental procedure The alumina (PRD-166) fibres were coated with SnO2 by a chemical vapor deposition technique. The alumina fibre tows were placed in the central hot zone of the reactor and heated to the deposition temperature of 500 C. Dry nitrogen was the carrier gas for SnCl4. The flow rate of nitrogen was 1 l/min. A second bubbler contained water heated to 80 C through which oxygen was passed at a rate of 0.6 l/min. The deposition occurred as per the chemical reaction given below [29], SnCl4 ðgÞ þ 2H2 O ðgÞ ! SnO2 ðsÞ þ 4HClðgÞ
ð10Þ
The microstructures of PRD-166 fibres as-received, heat treated at 500, 600 and 900 C for 90 min and SnO2 coated fibres were characterised using SEM and XRD. SEM was used to determine uniformity, morphology and thickness of the coating. The fracture surfaces of the asreceived and SnO2 coated fibres were also studied by SEM. Single fibre tensile tests were carried out on asreceived, heat treated and SnO2 coated fibres. A random selection of single fibres was made from the material to be tested. The fibres were center-line mounted on a paper frame. The fibres were centered over the frame and lightly stretched. A small amount of adhesive was then carefully placed at each end of the fibre. The specimen gage length for all the fibres tested was 17 mm. An Instron tensile testing machine (model 1120) was used with a 5 N load cell. Before fibres were loaded onto the machine, the diameter of the individual fibres was measured with an optical microscope. The frame was then gripped in the jaws of the testing machine and the mounting frame was burned on the sides. The fibres were then stressed to failure at a crosshead speed of 0.25 mm/min. An average of 80 fibres in each group, i.e. asreceived, heat treated at 500, 600 and 900 C and SnO2 coated were tested. The mean tensile strength, Weibull modulus, standard deviation and coefficient of variation were then evaluated by Weibull analysis as described in Section 2.
4. Results and discussion Fig. 1(a and b) shows the microstructure of alumina PRD-166 fibres. The surface of PRD-166 fibres has a rough cobblestone structure as is shown in Fig. 1(a). The grain size of alumina as determined by lineal intercept method was about 0.5 mm and that of zirconia particles was 0.33 mm. The zirconia particles are dispersed
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throughout the fibres but primarily along grain boundaries, Fig.1(b). The dispersion of 20 wt.% zirconia in PRD-166 fibre impedes grain growth and thereby improves the strength of these fibres [30]. XRD showed the zirconia particles to be primarily in tetragonal form. The transformation of the tetragonal zirconia particles to the monoclinic form at the crack tip during crack propagation could improve the toughness of the alumina-zirconia fibres [30]. Romine has shown that PRD-166 fibres maintain their strengths up to 1200 and even at 1400 C, 75% of their strength was maintained. He also observed that when PRD-166 fibres were exposed to high temperatures, the maximum strength and change in microstructure was observed in the first few hours. After several hours, the degradation of the fibres stabilised and, thereafter, the strength was retained for prolonged periods. This was
attributed to zirconia particles acting as grain growth inhibitors [30]. Table 1 shows the Weibull parameters of as-received and heat treated alumina fibres. It can be seen that the strength decreases with increase in heat treatment temperature. A room temperature strength of 1375 MPa was obtained for a gage length of 17 mm. This value is lower than that reported by Romine [30] who obtained a strength of 2100 MPa for a gage length of 6.9 mm. The decrease in strength as compared to Romine’s data could be due to increase in gage length. As the length of the fibres increased, the volume of material increased. Probability of finding flaws in a material increases with its volume. Processing voids (Fig. 2) could also decrease the strength of as-received fibres. These voids could have been formed because of the collapse of the internal structure of the fibres due to differential temperature distribution between its exterior and interior, known as core-sheath structure [31]. It has been shown in an earlier work that processing defects, e.g. sticking of fibres, shots, porous fibres, cracking and core-sheath structure can significantly reduce the strength of ceramic fibres [31]. Processing flaws have also been observed in PRD166 fibres by Pysher et al. [32]. Yang et al. [33] obtained Table 1 Mean strength, standard deviation and coeffecient of variation of asreceived and heat treated alumina (PRD-166) fibres
Fig. 1. Microstructure of as-received fibres . (a) Longitudinal surface of fibres showing the rough cobblestone structure, (b) zirconia particles dispersed throughout the fibre.
Fibre
Mean tensile strength (MPa)
Standard deviation (MPa)
Coefficient of variation (%)
As-received 500 C 600 C 900 C
1375 1313 1283 1083
4185 386 440 320
30 29 34 29
Fig. 2. Fracture surface of an as-received fibre showing processing voids.
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a strength of 1180 MPa for a gage length of 50 mm. The strength of the PRD-166 fibres were shown to decrease with increase in gage length which was again attributed to fibre roughness and presence of voids [33] The strength of fibres heat treated at 500, 600 and 900 C decreased by 5, 7 and 21% respectively as compared to as-received fibres, (Table 1). It has been shown that PRD-166 fibres retained most of their strength up to 800 but exhibited extensive deformation and no significant strengths above 1200 [32,34]. They have reported the presence of minor elements (0.01–0.1%) of Si, and P which could form a SiO2–P2O5 glassy grain boundary phase in PRD-166 fibres. These glassy phases could reduce the strength of these fibres at high temperatures. Fracture surfaces of as-received and fibres heat treated at 900 C are shown in Fig. 3. Fig. 3(a) shows the brittle transgranular fracture surface of the
alumina (PRD-166) as-received fibres. Fracture surface of fibres heat treated at 900 C, [Fig. 3(b)], showed brittle intergranular failure probably because of the glassy phase at the grain boundaries. Grain pullout can also be seen. It has been shown in earlier works that the alumina fibres fail transgranular up to 800 C beyond which their failure mode is intergranular [32]. Fibres coated with SnO2 showed brittle failure, (Fig. 4). Table 2 shows variation of Weibull parameters of SnO2 coated fibres. The tensile strength decreases with increase in coating thickness. Some loss in strength can be attributed to exposures at 500 C during SnO2 deposition. Another source of strength reduction could be thermal stresses generated during deposition. Thermal stress distribution using a two-element cylinderical model [35,36] showed radial tensile stress at fibre/coating interphase while the axial and circumferential stress was tensile in the fibre and compressive in the coating, (Table 3). Axial thermal stress in the fibre and radial stress at fibre/SnO2 interphase which are tensile in nature increased with coating thickness, (Table 3). It then is expected that the tensile stress in the fibre and the fact that SnO2 coating is
Fig. 4. Fracture surace of SnO2 coated fibre showing brittle failure. Table 2 Mean strength, standard deviation and coeffecient of variation of asreceived and SnO2 coated alumina (PRD-166) fibres
Fig. 3. Fracture surface of fibre (a) as-received (b) fibres calcined at 900 C showing brittle intergranular failure.
Fibre
Mean tensile standard (MPa)
Standard deviation (MPa)
Coefficient of variation (%)
As-received SnO2 coated (0.4 mm) SnO2 coated (0.5 mm) SnO2 coated (0.8 mm) SnO2 coated (2.0 mm) SnO2 coated (10 mm)
1375 1060 966 851 702 166
418 386 440 320 440 320
30 25 28 33 32 34
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R. Venkatesh / Composites Science and Technology 62 (2002) 205–212 Table 3 Radial ( r), circumferential ( ), and axial stresses ( z) at the alumina fibre/SnO2 interphase. Subscript f denotes the fibre and s the coating Thickness of SnO2
fr= f= sr (MPa)
fz (MPa)
sz (MPa)
s (MPa)
0.4 0.5 0.8 2.0
22 28 42 53
45 55 88 97
540 535 524 502
526 518 497 485
weaker than alumina could cause a reduction in the strength of alumina/SnO2 composite fibre with an increase in coating thickness. A plot of f/ f of as-received and SnO2 coated fibres versus number of samples tested is shown in Fig. 5. It can be seen that the deviation in the strength of the fibres decreases with increase in number of samples tested. The deviation in mean tensile strength of as-received fibres is about 1.6% and for SnO2 coated fibres is about 1.8%. In order to evaluate the strength of ceramic fibres, number of fibres tested have to be sufficiently high to have a low variation and therefore a high reliability. The rough Al2O3/SnO2 interphase and the coating surface can be seen in Fig. 6. In order to get a measure of the effect of roughness on the properties of the fibres, amplitude of roughness of as-received and SnO2 coated fibres were evaluated. The amplitude of fibre roughness
Fig. 6. SnO2 coated alumina (PRD-166) fibre showing rough PRD166/SnO2 interphase and coating surfaces.
have been experimentally determined from fibre surface profiles using atomic force microscopy [37–40], profilometry [41] and optical interference [42]. The peak-valley roughness amplitude, A, was evaluated in the present study using SEM micrographs and tracings of roughness profiles. These results are as shown in Table 4. It can be seen that the amplitude of roughness, A increased with increase in coating thickness. Under an axial load, this roughness could act as a notch and decrease the strength of the fibres. Hence it is very important to control the roughness in as-received and coated fibres. Roughness in alumina fibres can be controlled by improved processing through second phase additives, controlling the calcination temperatures and the amount of organics additives during the processing and also by decreasing the coating thickness [6,31,43]. The roughness of the fibre surface and coating have serious implications in the development of tough ceramic matrix composites. It has been shown that as the roughness of interphase increases, the compressive clamping stress increases thereby affecting the fracture characteristics of ceramic matrix composites [37–39, 43–46]. In the present study the effect of two different types of fibres namely, alumina fibre (PRD-166) and Saphikon fibre on the fracture characteristics of alumina fibre (PRD-166)/ SnO2/glass (ASG) and Saphikon fibre/ SnO2/glass (SSG) composites have been investigated. Evaluating Table 4 Amplitude of roughness, A with coating thickness of the fibres
Fig. 5. Deviation of tensile strength of PRD-166 and PRD-166/SnO2 coated fibres with number of samples for a 90% confidence limit.
Thickness of SnO2 (mm)
Amplitude, A (mm)
0.0 0.4 0.5 0.8 2.0 10.0
0.26 0.45 0.53 0.88 1.8 4.0
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Fig. 7. (a and b). Morphology of interfaces in alumina PRD-166 fibre/ SnO2 and saphikon fibre/SnO2.
amplitude of roughness, A and roughness strain A/r at the PRD–166/SnO2 and Saphikon/SnO2 interfaces, [Fig. 7(a and b)] it was observed that the A/r value of PRD–166/SnO2 interface was about nine times that of the Saphikon/SnO2 interface (Table 5). This indicates the strong mechanical clamping due to fiber roughness at the fiber/SnO2 interface in ASG composites. Fracture surfaces of ASG composites are non planar (Fig. 8). Fig. 8a shows that the predominant mechanism of Table 5 Amplitude of roughness, A, and roughness strain A/r of PRD-166/ SnO2 and Saphikon/SnO2 interphase
PRD-166/SnO2 Saphikon/SnO2
A
R (mm)
A/r
0.26 0.13
10 43.5
0.026 0.003
Fig. 8. (a and b). Fracture surface of PRD-166 alumina fibre/SnO2/ glass matrix composites showing partial debonding and fibre pullout. Note the extremely rough PRD-166 fibre.
toughening is through crack deflection with some fiber bridging. Partial pullout of fibers can also be seen. A higher magnification micrograph of the fracture surface of ASG composites along the fiber (Fig. 8b), shows clearly the partial removal of the coating and the rough fiber surface. Hence the primary toughening mechanisms in ASG composites are crack deflection, fiber bridging, partial fiber debonding and pullout. As fiber roughness increases, the compressive clamping stress causes an increase in the shear stress transfer at the interface beyond matrix cracking from matrix to fibre. This causes a reduction in the debond length, i.e. fibers break rather
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coated fibres decreased with the increase in number of samples tested and at a sample size of 80 the variation in the tensile strength was about 2%. The strength of the fibres decreased with increase in coating thickness. The thickness and roughness of SnO2 coating had a profound effect on the strength of the coated fibres. It was observed that as the coating thickness increased, roughness of the coating increased. This acted as a notch in order to decrease the strength of the composite fibres. This roughness induced effect has serious implications on the development of ceramic matrix composites. In CMCs, as roughness increased, compressive clamping stress at the fibre/matrix interface increased, fibre debonding decreased and toughness of the composite decreased. Hence, it is important to control the roughness at the fibre/ coating and coating/matrix interfaces in order to develop tough ceramic matrix composites. Methods of decreasing roughness at the interfaces include the incorporation of smooth fibres and decreasing coating roughness. References
Fig.9. (a and b). Fracture surface of saphikon alumina fibre/SnO2/ glass matrix composites. The SnO2 coating on a relatively smooth saphikon fibre results in a neat a long fibre pull out.
than debond as the matrix crack grows, resulting in a composite fracture surface in ASG with little or no fiber pullout. Fracture surface of SSG composites, [Fig. 9(a and b)], showed neat fiber debonding and fiber pullout at the fiber/SnO2 interface. This was because of the smooth saphikon/SnO2 interphase. Hence interfacial roughness plays a major role in the governing the fracture characteristics of ceramic matrix composites. Interfacial roughness in CMCs can be controlled through the use of smooth fibres or reducing the coating thickness and roughness.
5. Conclusion It was observed that the strength of alumina–ZrO2 fibres (PRD-166) decreased with increase in heat treatment temperature of fibres. This could be because of the presence of a glassy second phase in the fibres. It was found that the variation of the tensile strength of as-received and
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[39] Parthasarathy TA, Kerans RJ. J Am Ceram Soc 1992;80(8):2043– 501. [40] Wang HP, Nelson TJ, Lim CL, Gerberch WW. Interfacial stability and mechanical properties of alumina fiber reinforced ti matrix composites. J Mater Res 1994;9(2):498. [41] Chawla KK, Xu ZR, Hlinak A, Chung YW. Characteristics of three alumina fibers by atomic force microscopy. In: Bansal NP, editor. Advances in ceramic matrix composites, ceramic transaction, vol, 38. OH: Am Ceram Soc, 1993. p. 725. [42] Jero PD, Parthasarathy TA. Interface properties their measurement with fiber pushout tests. In: Naslain R, editor. High temperature ceramic matrix composites. Cambridge: Goodhead Publishers, 1993. p. 401. [43] Venkatesh, R., Teck, L.C., Ahmad, Z.A. and fauzi, A. J Mater Tech (in press). [44] Carter WC, Butler EP, Fuller Jr. ER. Scripta Metall Mater 1991; 25:579–85. [45] Mackin TJ, Warren PD, Evans AG. Acta Metall 1992;40:1251– 9. [46] Venkatesh R, Chawla KK. J Mater Sci Lett 1992;11:650–4.
The effects of heat treatment and coating roughness on the strength of alumina-zirconia fibres Ramanan Venkatesh School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia, 14300 Nibong Tebal, Pulau Pinang, Malaysia Received 5 February 2001; received in revised form 18 April 2001; accepted 25 September 2001
Abstract The strengths of Al2O3–20 wt.% ZrO2 fibres as-received, heat treated and coated with SnO2 have been determined. The mean strength of as-received fibres was 1380 MPa for a gage length of 17 mm. The strengths of fibres heat treated at 500, 600 and 900 C decreased by 5, 7 and 21%, respectively, relative to the as-received fibres. This could be because of a secondary glassy phase formation at the grain boundaries. It was found that the accuracy of the tensile strength distribution increased with the number of samples tested and at a sample number of 80, the variation in the tensile strength was about 2%. The thickness and roughness of the SnO2 coating had a profound effect on the strength of the coated fibres. It was observed that as the coating thickness increased, roughness of the coating increased. This acted as a notch in order to decrease the strength of the fibres. Control of roughness at the fibre/coating and/or coating/matrix interphase is very important in the development of tough ceramic–matrix composites. As the interphase roughness increases, fibre debonding and pull-out decrease because of the increased compressive clamping stress. Hence for increased toughness in CMCs through fibre debonding and pull-out mechanisms, a relatively smooth interphase is necessary. The interphase roughness can be controlled either by selecting smooth fibres or by decreasing coating roughness. # 2002 Published by Elsevier Science Ltd. Keywords: A. Fibres; A. Coating; B. Interphase; Alumina
1. Introduction High-performance ceramic fibres are important in the development of ceramic–matrix composites for hightemperature applications. Characteristics of fibres that influence properties of CMCs include strength, strain to failure, statistical variation in strength, aspect ratio, microstructure and thermal stability. In applications where stability in air at high temperatures is a prime objective, oxide fibres have to be considered since they offer better oxidation resistance at high temperatures. Principal methods of forming oxide ceramic fibres include the spinning of aqueous alumina slurry containing spinning additives or spinning of polymeric solutions using sol-gel techniques [1–6]. Alumina, mullite and zirconia fibres are the principal oxide fibres developed for ceramic matrix reinforcement [7–15]. In order to be used for composite applications, these fibres should have good strengths of about 2 GPa and long-term application temperatures greater than 1200 C. E-mail address: [email protected]
Considerable work exists in the literature for the development of oxide fibre/oxide matrix composites [16–25]. In oxide fibre/oxide matrix composite systems a reaction at the fibre/matrix interface, e.g. Al2O3/SiO2, mullite/mullite could lead to a strong interfacial bonding and hence a brittle composite. Barrier layers that prevent chemical reaction between the fibre and the matrix, e.g. SnO2, BN, monazite, hibonite, etc. have to be introduced at the fibre/ matrix interface [20–25]. In the present study, SnO2 was chosen as a coating on the alumina fibre surface because it does not react with alumina up to 1400 C in partial pressure of oxygen > 107 atmospheres [26]. The present study is an investigation of the effect of heat treatment and SnO2 coating thickness on the strength of alumina (PRD-166) fibres and its implications towards the development of tough ceramic matrix composites.
2. Tensile strength of ceramic fibres The mean strength of ceramic fibre decreases with increase in volume of material due to increased probability
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of defects in the fibre. The effect of random distribution of a single defect on the strength of a solid has been described by Weibull [27]. According to the Weibull distribution, probability of failure of a fibre at a stress, , can be given as, FðÞ ¼ 1- exp ð1Þ where is a scale parameter and is known as Weibull modulus and is inversely related to scatter in the strength of fibres. The value of the Weibull modulus for most ceramic fibres is low, often around 5 which reflects the large variation in their strength. Eq. (1) can be written as, ln ln½ð1=1-Fð ÞÞ ¼ ln þ ln
ð2Þ
In an actual test, strengths of N fibres tested are arranged in an increasing order and the failure probability of the ith sample, Pfi, can be estimated from an estimator of the kind shown below, Pfi ¼ i=N þ 1
ð3Þ
Substituting this estimator into Eq. (2), we can write, ln ln½ðN þ 1=N þ 1-iÞ ¼ ln i þ ln
ð4Þ
A plot of ln ln [(N+1/N+1-i)] against ln i will then yield a straight line with slope and intercept ln . Once and are known, the Weibull mean tensile strength, f, is then given as, f ¼ 1= ð1 þ 1=Þ
ð5Þ
where (1+1/) is the gamma function of (1+1/) and can be obtained from standard mathematical tables. Standard deviation Sf is given as, 1=2 Sf ¼ 1= ð1 þ 2=Þ- 2 ð1 þ 1=Þ
ð6Þ
Coefficient of variation C.V., can then be given as, C:V: ¼ 100 ðSf =f Þ
ð7Þ
It has been shown that the standard deviation, Sf divided by mean modulus can be related to the number of fibres tested as [28], Sf = ¼ ð1=nÞ1=2
ð8Þ
where n is the number of fibres tested. The variation of mean tensile strength of the fibres, f for a 90% confidence limit can then be written as [28], f ¼ f 1:645 Sf =n1=2
ð9Þ
It can be seen from Eq. (9) that the variation in the mean tensile strength decreases with an increase in the number of fibres tested.
3. Experimental procedure The alumina (PRD-166) fibres were coated with SnO2 by a chemical vapor deposition technique. The alumina fibre tows were placed in the central hot zone of the reactor and heated to the deposition temperature of 500 C. Dry nitrogen was the carrier gas for SnCl4. The flow rate of nitrogen was 1 l/min. A second bubbler contained water heated to 80 C through which oxygen was passed at a rate of 0.6 l/min. The deposition occurred as per the chemical reaction given below [29], SnCl4 ðgÞ þ 2H2 O ðgÞ ! SnO2 ðsÞ þ 4HClðgÞ
ð10Þ
The microstructures of PRD-166 fibres as-received, heat treated at 500, 600 and 900 C for 90 min and SnO2 coated fibres were characterised using SEM and XRD. SEM was used to determine uniformity, morphology and thickness of the coating. The fracture surfaces of the asreceived and SnO2 coated fibres were also studied by SEM. Single fibre tensile tests were carried out on asreceived, heat treated and SnO2 coated fibres. A random selection of single fibres was made from the material to be tested. The fibres were center-line mounted on a paper frame. The fibres were centered over the frame and lightly stretched. A small amount of adhesive was then carefully placed at each end of the fibre. The specimen gage length for all the fibres tested was 17 mm. An Instron tensile testing machine (model 1120) was used with a 5 N load cell. Before fibres were loaded onto the machine, the diameter of the individual fibres was measured with an optical microscope. The frame was then gripped in the jaws of the testing machine and the mounting frame was burned on the sides. The fibres were then stressed to failure at a crosshead speed of 0.25 mm/min. An average of 80 fibres in each group, i.e. asreceived, heat treated at 500, 600 and 900 C and SnO2 coated were tested. The mean tensile strength, Weibull modulus, standard deviation and coefficient of variation were then evaluated by Weibull analysis as described in Section 2.
4. Results and discussion Fig. 1(a and b) shows the microstructure of alumina PRD-166 fibres. The surface of PRD-166 fibres has a rough cobblestone structure as is shown in Fig. 1(a). The grain size of alumina as determined by lineal intercept method was about 0.5 mm and that of zirconia particles was 0.33 mm. The zirconia particles are dispersed
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throughout the fibres but primarily along grain boundaries, Fig.1(b). The dispersion of 20 wt.% zirconia in PRD-166 fibre impedes grain growth and thereby improves the strength of these fibres [30]. XRD showed the zirconia particles to be primarily in tetragonal form. The transformation of the tetragonal zirconia particles to the monoclinic form at the crack tip during crack propagation could improve the toughness of the alumina-zirconia fibres [30]. Romine has shown that PRD-166 fibres maintain their strengths up to 1200 and even at 1400 C, 75% of their strength was maintained. He also observed that when PRD-166 fibres were exposed to high temperatures, the maximum strength and change in microstructure was observed in the first few hours. After several hours, the degradation of the fibres stabilised and, thereafter, the strength was retained for prolonged periods. This was
attributed to zirconia particles acting as grain growth inhibitors [30]. Table 1 shows the Weibull parameters of as-received and heat treated alumina fibres. It can be seen that the strength decreases with increase in heat treatment temperature. A room temperature strength of 1375 MPa was obtained for a gage length of 17 mm. This value is lower than that reported by Romine [30] who obtained a strength of 2100 MPa for a gage length of 6.9 mm. The decrease in strength as compared to Romine’s data could be due to increase in gage length. As the length of the fibres increased, the volume of material increased. Probability of finding flaws in a material increases with its volume. Processing voids (Fig. 2) could also decrease the strength of as-received fibres. These voids could have been formed because of the collapse of the internal structure of the fibres due to differential temperature distribution between its exterior and interior, known as core-sheath structure [31]. It has been shown in an earlier work that processing defects, e.g. sticking of fibres, shots, porous fibres, cracking and core-sheath structure can significantly reduce the strength of ceramic fibres [31]. Processing flaws have also been observed in PRD166 fibres by Pysher et al. [32]. Yang et al. [33] obtained Table 1 Mean strength, standard deviation and coeffecient of variation of asreceived and heat treated alumina (PRD-166) fibres
Fig. 1. Microstructure of as-received fibres . (a) Longitudinal surface of fibres showing the rough cobblestone structure, (b) zirconia particles dispersed throughout the fibre.
Fibre
Mean tensile strength (MPa)
Standard deviation (MPa)
Coefficient of variation (%)
As-received 500 C 600 C 900 C
1375 1313 1283 1083
4185 386 440 320
30 29 34 29
Fig. 2. Fracture surface of an as-received fibre showing processing voids.
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a strength of 1180 MPa for a gage length of 50 mm. The strength of the PRD-166 fibres were shown to decrease with increase in gage length which was again attributed to fibre roughness and presence of voids [33] The strength of fibres heat treated at 500, 600 and 900 C decreased by 5, 7 and 21% respectively as compared to as-received fibres, (Table 1). It has been shown that PRD-166 fibres retained most of their strength up to 800 but exhibited extensive deformation and no significant strengths above 1200 [32,34]. They have reported the presence of minor elements (0.01–0.1%) of Si, and P which could form a SiO2–P2O5 glassy grain boundary phase in PRD-166 fibres. These glassy phases could reduce the strength of these fibres at high temperatures. Fracture surfaces of as-received and fibres heat treated at 900 C are shown in Fig. 3. Fig. 3(a) shows the brittle transgranular fracture surface of the
alumina (PRD-166) as-received fibres. Fracture surface of fibres heat treated at 900 C, [Fig. 3(b)], showed brittle intergranular failure probably because of the glassy phase at the grain boundaries. Grain pullout can also be seen. It has been shown in earlier works that the alumina fibres fail transgranular up to 800 C beyond which their failure mode is intergranular [32]. Fibres coated with SnO2 showed brittle failure, (Fig. 4). Table 2 shows variation of Weibull parameters of SnO2 coated fibres. The tensile strength decreases with increase in coating thickness. Some loss in strength can be attributed to exposures at 500 C during SnO2 deposition. Another source of strength reduction could be thermal stresses generated during deposition. Thermal stress distribution using a two-element cylinderical model [35,36] showed radial tensile stress at fibre/coating interphase while the axial and circumferential stress was tensile in the fibre and compressive in the coating, (Table 3). Axial thermal stress in the fibre and radial stress at fibre/SnO2 interphase which are tensile in nature increased with coating thickness, (Table 3). It then is expected that the tensile stress in the fibre and the fact that SnO2 coating is
Fig. 4. Fracture surace of SnO2 coated fibre showing brittle failure. Table 2 Mean strength, standard deviation and coeffecient of variation of asreceived and SnO2 coated alumina (PRD-166) fibres
Fig. 3. Fracture surface of fibre (a) as-received (b) fibres calcined at 900 C showing brittle intergranular failure.
Fibre
Mean tensile standard (MPa)
Standard deviation (MPa)
Coefficient of variation (%)
As-received SnO2 coated (0.4 mm) SnO2 coated (0.5 mm) SnO2 coated (0.8 mm) SnO2 coated (2.0 mm) SnO2 coated (10 mm)
1375 1060 966 851 702 166
418 386 440 320 440 320
30 25 28 33 32 34
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R. Venkatesh / Composites Science and Technology 62 (2002) 205–212 Table 3 Radial ( r), circumferential ( ), and axial stresses ( z) at the alumina fibre/SnO2 interphase. Subscript f denotes the fibre and s the coating Thickness of SnO2
fr= f= sr (MPa)
fz (MPa)
sz (MPa)
s (MPa)
0.4 0.5 0.8 2.0
22 28 42 53
45 55 88 97
540 535 524 502
526 518 497 485
weaker than alumina could cause a reduction in the strength of alumina/SnO2 composite fibre with an increase in coating thickness. A plot of f/ f of as-received and SnO2 coated fibres versus number of samples tested is shown in Fig. 5. It can be seen that the deviation in the strength of the fibres decreases with increase in number of samples tested. The deviation in mean tensile strength of as-received fibres is about 1.6% and for SnO2 coated fibres is about 1.8%. In order to evaluate the strength of ceramic fibres, number of fibres tested have to be sufficiently high to have a low variation and therefore a high reliability. The rough Al2O3/SnO2 interphase and the coating surface can be seen in Fig. 6. In order to get a measure of the effect of roughness on the properties of the fibres, amplitude of roughness of as-received and SnO2 coated fibres were evaluated. The amplitude of fibre roughness
Fig. 6. SnO2 coated alumina (PRD-166) fibre showing rough PRD166/SnO2 interphase and coating surfaces.
have been experimentally determined from fibre surface profiles using atomic force microscopy [37–40], profilometry [41] and optical interference [42]. The peak-valley roughness amplitude, A, was evaluated in the present study using SEM micrographs and tracings of roughness profiles. These results are as shown in Table 4. It can be seen that the amplitude of roughness, A increased with increase in coating thickness. Under an axial load, this roughness could act as a notch and decrease the strength of the fibres. Hence it is very important to control the roughness in as-received and coated fibres. Roughness in alumina fibres can be controlled by improved processing through second phase additives, controlling the calcination temperatures and the amount of organics additives during the processing and also by decreasing the coating thickness [6,31,43]. The roughness of the fibre surface and coating have serious implications in the development of tough ceramic matrix composites. It has been shown that as the roughness of interphase increases, the compressive clamping stress increases thereby affecting the fracture characteristics of ceramic matrix composites [37–39, 43–46]. In the present study the effect of two different types of fibres namely, alumina fibre (PRD-166) and Saphikon fibre on the fracture characteristics of alumina fibre (PRD-166)/ SnO2/glass (ASG) and Saphikon fibre/ SnO2/glass (SSG) composites have been investigated. Evaluating Table 4 Amplitude of roughness, A with coating thickness of the fibres
Fig. 5. Deviation of tensile strength of PRD-166 and PRD-166/SnO2 coated fibres with number of samples for a 90% confidence limit.
Thickness of SnO2 (mm)
Amplitude, A (mm)
0.0 0.4 0.5 0.8 2.0 10.0
0.26 0.45 0.53 0.88 1.8 4.0
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Fig. 7. (a and b). Morphology of interfaces in alumina PRD-166 fibre/ SnO2 and saphikon fibre/SnO2.
amplitude of roughness, A and roughness strain A/r at the PRD–166/SnO2 and Saphikon/SnO2 interfaces, [Fig. 7(a and b)] it was observed that the A/r value of PRD–166/SnO2 interface was about nine times that of the Saphikon/SnO2 interface (Table 5). This indicates the strong mechanical clamping due to fiber roughness at the fiber/SnO2 interface in ASG composites. Fracture surfaces of ASG composites are non planar (Fig. 8). Fig. 8a shows that the predominant mechanism of Table 5 Amplitude of roughness, A, and roughness strain A/r of PRD-166/ SnO2 and Saphikon/SnO2 interphase
PRD-166/SnO2 Saphikon/SnO2
A
R (mm)
A/r
0.26 0.13
10 43.5
0.026 0.003
Fig. 8. (a and b). Fracture surface of PRD-166 alumina fibre/SnO2/ glass matrix composites showing partial debonding and fibre pullout. Note the extremely rough PRD-166 fibre.
toughening is through crack deflection with some fiber bridging. Partial pullout of fibers can also be seen. A higher magnification micrograph of the fracture surface of ASG composites along the fiber (Fig. 8b), shows clearly the partial removal of the coating and the rough fiber surface. Hence the primary toughening mechanisms in ASG composites are crack deflection, fiber bridging, partial fiber debonding and pullout. As fiber roughness increases, the compressive clamping stress causes an increase in the shear stress transfer at the interface beyond matrix cracking from matrix to fibre. This causes a reduction in the debond length, i.e. fibers break rather
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coated fibres decreased with the increase in number of samples tested and at a sample size of 80 the variation in the tensile strength was about 2%. The strength of the fibres decreased with increase in coating thickness. The thickness and roughness of SnO2 coating had a profound effect on the strength of the coated fibres. It was observed that as the coating thickness increased, roughness of the coating increased. This acted as a notch in order to decrease the strength of the composite fibres. This roughness induced effect has serious implications on the development of ceramic matrix composites. In CMCs, as roughness increased, compressive clamping stress at the fibre/matrix interface increased, fibre debonding decreased and toughness of the composite decreased. Hence, it is important to control the roughness at the fibre/ coating and coating/matrix interfaces in order to develop tough ceramic matrix composites. Methods of decreasing roughness at the interfaces include the incorporation of smooth fibres and decreasing coating roughness. References
Fig.9. (a and b). Fracture surface of saphikon alumina fibre/SnO2/ glass matrix composites. The SnO2 coating on a relatively smooth saphikon fibre results in a neat a long fibre pull out.
than debond as the matrix crack grows, resulting in a composite fracture surface in ASG with little or no fiber pullout. Fracture surface of SSG composites, [Fig. 9(a and b)], showed neat fiber debonding and fiber pullout at the fiber/SnO2 interface. This was because of the smooth saphikon/SnO2 interphase. Hence interfacial roughness plays a major role in the governing the fracture characteristics of ceramic matrix composites. Interfacial roughness in CMCs can be controlled through the use of smooth fibres or reducing the coating thickness and roughness.
5. Conclusion It was observed that the strength of alumina–ZrO2 fibres (PRD-166) decreased with increase in heat treatment temperature of fibres. This could be because of the presence of a glassy second phase in the fibres. It was found that the variation of the tensile strength of as-received and
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