《复合材料 Composites》课程教学资源（学习资料）第五章陶瓷基复合材料_Interphase effects on the bend strength and toughness of an oxide.pdf_大学文库

CERAMICS INTERNATIONAL ELSEⅤIER Ceramics International 28(2002)565-573 Interphase effects on the bend strength and toughness of an oxide fibre/oxide matrix composite Ramanan venkatesh School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia, 14300 Nibong Tebal, Seberang Perai Setatan Pulau Pinang, Malaysia Received 30 April 2001; received in revised form 12 October 2001; accepted 3 December 2001 Abstract The effect of alumina-20 wt ZrO2(PRD-166) fibre and Sno2 coating properties on the bending strength and toughness of alumina fibre/SnO2/glass matrix composites have been investigated. The mean strength of as-received alumina-20 wt ZrO2 fibres was 1380 MPa for a gage length of 17 mm and decreased with increase in heat treatment temperatures. It was also observed that as the Sno, coating thickness increased, roughness of the coating increased and this decreased the strength of the fibres. This rough ness effect had serious implications on the fracture characteristics of PRD-166/SnO2/glass and Saphikon/SnO2/glass matrix com- posites. PRD-166/SnO,glass matrix composites exhibited non-planar failure with fiber bridging and fibre debonding as major toughening mechanisms. Saphikon/SnO2/glass matrix composites failed in a tough manner with extensive fibre pullout. The differ- nce the failure mode between PRD-166/ SnO2/glass and Saphikon/SnO2/glass matrix composites was attributed to the clamping stress associated with fiber roughness at the PrD-166/SnO, interface as compared to the smoother Saphikon/ Sno, interface C 2002 Elsevier Science Ltd and Techna S.r. l. All rights reserved Keywords: Interface effects; Bend strength; Toughness; Oxide fibre composite Introduction irregularities at the fibre /matrix interface. In composites with strong bonding at the interface, cracks originating Fibre reinfor great potential for in the brittle matrix tensile cut through the fibres. improving strength and toughness of ceramic materials resulting in a planar brittle failure of the composite In [1-5]. Parameters that influence the properties of fibres composite with a weak bonding at the interface, when in ceramic matrix composites(CMCs) include: tensile matrix tensile strength is exceeded, multiple cracking of strength, strain to failure, Weibull modulus, aspect ratio the matrix takes place with the fibres having enough and surface roughness. Alumina, mullite and zirconia strength to bridge the cracks. Further increase in stress are the principal polycrystalline oxide fibres developed causes fibre debonding due to interfacial stress and [6-21]. Oxide fibre/oxide matrix composites are con- Poisson's effect. Continued stressing of the composite sidered for potential use at extremely high temperatures beyond fibre debonding causes the failure of the fibre (1400-1600C)and in severe environments [22-31]. along its length and then depending on residual stress, Failure strength and toughness of CMCs depend on a Poissons ratio of fibre and matrix and interfacial fric multitude of mechanisms involving matrix microcrack- tional stress, fibre pullout occurs. Fibre pullout is the ing, matrix prestressing, fibre debonding and fibre pull- main toughening mechanism in CMCs. Hence for a out. Strength and toughness of CMCs are greatly tough composite, the interface bonding should be influenced by the interfacial bonding at the fibre/matrix strong enough to allow load transfer but weak enough interface. Interfacial strength is a strong function of the to aid crack deflection, fibre debonding and fibre pull- degree of bonding(chemical or mechanical) between out Interfacial strength can be controlled by modifying fibre and matrix and the thermal mismatch between fibre the fibre/matrix reactions at the processing and service and matrix. Mechanical bonding is primarily due to temperatures either through proper selection of materi als or by means of interface gs. For a operate successfully the following conditions should be 0272-8842/02/S22.00C 2002 Elsevier Science Ltd and Techna S.r. l. All rights reserved. PII:S0272-8842(02)00011-1

Interphase effects on the bend strength and toughness of an oxide fibre/oxide matrix composite Ramanan Venkatesh School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia, 14300 Nibong Tebal, Seberang Perai Setatan, Pulau Pinang, Malaysia Received 30 April 2001; received in revised form 12 October 2001; accepted 3 December 2001 Abstract The effect of alumina–20 wt.% ZrO2 (PRD-166) fibre and SnO2 coating properties on the bending strength and toughness of alumina fibre/SnO2/glass matrix composites have been investigated. The mean strength of as-received alumina–20 wt.% ZrO2 fibres was 1380 MPa for a gage length of 17 mm and decreased with increase in heat treatment temperatures. It was also observed that as the SnO2 coating thickness increased, roughness of the coating increased and this decreased the strength of the fibres. This roughness effect had serious implications on the fracture characteristics of PRD-166/SnO2/glass and Saphikon/SnO2/glass matrix composites. PRD-166/SnO2/glass matrix composites exhibited non-planar failure with fiber bridging and fibre debonding as major toughening mechanisms. Saphikon/SnO2/glass matrix composites failed in a tough manner with extensive fibre pullout. The difference in the failure mode between PRD-166/SnO2/glass and Saphikon/SnO2/glass matrix composites was attributed to the clamping stress associated with fiber roughness at the PRD-166/SnO2 interface as compared to the smoother Saphikon/SnO2 interface. # 2002 Elsevier Science Ltd and Techna S.r.l. All rights reserved. Keywords: Interface effects; Bend strength; Toughness; Oxide fibre composite Introduction Fibre reinforcement offers a great potential for improving strength and toughness of ceramic materials [1–5]. Parameters that influence the properties of fibres in ceramic matrix composites (CMCs) include: tensile strength, strain to failure, Weibull modulus, aspect ratio and surface roughness. Alumina, mullite and zirconia are the principal polycrystalline oxide fibres developed [6–21]. Oxide fibre/oxide matrix composites are considered for potential use at extremely high temperatures (1400–1600 C) and in severe environments [22–31]. Failure strength and toughness of CMCs depend on a multitude of mechanisms involving matrix microcracking, matrix prestressing, fibre debonding and fibre pullout. Strength and toughness of CMCs are greatly influenced by the interfacial bonding at the fibre/matrix interface. Interfacial strength is a strong function of the degree of bonding (chemical or mechanical) between fibre and matrix and the thermal mismatch between fibre and matrix. Mechanical bonding is primarily due to irregularities at the fibre/matrix interface. In composites with strong bonding at the interface, cracks originating in the brittle matrix tensile cut through the fibres, resulting in a planar brittle failure of the composite. In a composite with a weak bonding at the interface, when matrix tensile strength is exceeded, multiple cracking of the matrix takes place with the fibres having enough strength to bridge the cracks. Further increase in stress causes fibre debonding due to interfacial stress and Poisson’s effect. Continued stressing of the composite beyond fibre debonding causes the failure of the fibre along its length and then depending on residual stress, Poisson’s ratio of fibre and matrix and interfacial frictional stress, fibre pullout occurs. Fibre pullout is the main toughening mechanism in CMCs. Hence for a tough composite, the interface bonding should be strong enough to allow load transfer but weak enough to aid crack deflection, fibre debonding and fibre pullout. Interfacial strength can be controlled by modifying the fibre/matrix reactions at the processing and service temperatures either through proper selection of materials or by means of interfacial coatings. For a coating to operate successfully the following conditions should be met. 0272-8842/02/$22.00 # 2002 Elsevier Science Ltd and Techna S.r.l. All rights reserved. PII: S0272-8842(02)00011-1 Ceramics International 28 (2002) 565–573 www.elsevier.com/locate/ceramint E-mail address: ram5nan@tm.net.mu (R. Venkatesh)

R Venkatesh/Ceramics International 28 (2002)565-573 1. Chemical compatibility with the matrix The alumina(Prd-166) fibres were coated with SnO 2. Refractoriness by a chemical vapor deposition technique. The alumina 3. Stability in oxidising, water vapor and corrosive fibre tows were placed in the central hot zone of the environments reactor and heated to the deposition temperature of 4. Providing a relatively weak fibre-coating inter- 500C. Dry nitrogen was the carrier gas for SnCl4. The phase allowing for fibre debonding and pullout. flow rate of nitrogen was I I/min. A second bubbler contained water heated to 80C through which oxygen In oxide fibre/oxide matrix composite systems, reac- was passed at a rate of 0.6 I/min. The deposition occur tions at the fibre/matrix interface, e.g. Al_O3/SiO2, red via the chemical reaction [38]. mullite/mullite leads to a strong interfacial bonding with the consequence of brittle behaviour of the composite. SnCl4(g)+ 2H20(g)+ SnO2(S)+4HCI(g) Various interphase coatings have been applied in oxide fibre/oxide matrix composites including, SnO2, TiO2, The microstructures of as-received fibres heat treated ZrO2, HfO2, monazite, magnetoplumbite, perouskite at 500, 600 and 900C for 90 minutes and SnO2 coated structures like BaTiO3, etc. [26-33]. In the present work, fibres were characterised using SEM and XRD. SEM the effect of fibre and coating properties on the bending was used to determine uniformity, morphology and strength and toughness of alumina fibre/glass matrix thickness of the coating. The fracture surfaces of the as- composites have been investigated SnO2 was chosen as received and SnO2 coated fibres were also characterised a coating since it has no reaction with alumina up to by SEM 1400 oC in a partial pressure of oxygen >10-7atm Single fibre tensile tests were carried out on as- [34, 35]. The strength of alumina-zirconia fibres as- received, heat treated and SnO, coated PRD-166 fibres received, heat treated and SnO2 coated were deter- A random selection of single fibres was made from the mined. Glass matrix composites fabricated using slurry impregnation technique and reinforced with two ypes of fibres, namely PrD-166(alumina-20 wt zir- conia) fibres and relatively smooth saphikon fibres were tested to investigate how the properties of the fibres can nfluence the bending strength and toughness of CMCs Experimental procedure The PRD-166 fiber used in the present work is a polycrystalline a-Al2O3 fiber, 20 um in diameter and containing 15-20 wt %Y2O3 partially stabilized zirco- nia particles. The properties of PRD-166 fiber are given in Table 1 [36]. The zirconia particles are dispersed throughout the fiber but primarily along the grain boundaries. Saphikon is a single crystal alumina fila- ment. The c-axis of the filament is oriented parallel to the fibre surface. The mechanical and physical proper ies of the saphikon filaments are given in Table 1 [37] N 5lA, a borosilicate glass, obtained from Owens Illinois Inc, was used as a matrix in the present study. Table Room-temperature properties of PRD-166 fiber and Saphikon fila- ment [ 36,37 Fibre Melting Density Tensile Tensile Thermal ao point(C)(g/cm)strength modulus expansion Mpa)(GPa)(×10-°/°C Saphikon 2053 3931503809.12∥/toc-axis) PRD-1662045 2070380 Fig. I.(a)Chevron notch specimen;(b) geometry of chevron notch

1. Chemical compatibility with the matrix. 2. Refractoriness. 3. Stability in oxidising, water vapor and corrosive environments. 4. Providing a relatively weak fibre-coating interphase allowing for fibre debonding and pullout. In oxide fibre/oxide matrix composite systems, reactions at the fibre/matrix interface, e.g. Al2O3/SiO2, mullite/mullite leads to a strong interfacial bonding with the consequence of brittle behaviour of the composite. Various interphase coatings have been applied in oxide fibre/oxide matrix composites including, SnO2, TiO2, ZrO2, HfO2, monazite, magnetoplumbite, perouskite structures like BaTiO3, etc. [26–33]. In the present work, the effect of fibre and coating properties on the bending strength and toughness of alumina fibre/glass matrix composites have been investigated. SnO2 was chosen as a coating since it has no reaction with alumina up to 1400 C in a partial pressure of oxygen >107 atm. [34,35]. The strength of alumina-zirconia fibres asreceived, heat treated and SnO2 coated were determined. Glass matrix composites fabricated using a slurry impregnation technique and reinforced with two types of fibres, namely PRD-166 (alumina–20 wt.% zirconia) fibres and relatively smooth saphikon fibres were tested to investigate how the properties of the fibres can influence the bending strength and toughness of CMCs. Experimental procedure The PRD-166 fiber used in the present work is a polycrystalline a-Al2O3 fiber, 20 mm in diameter and containing 1520 wt.% Y2O3 partially stabilized zirconia particles. The properties of PRD-166 fiber are given in Table 1[36]. The zirconia particles are dispersed throughout the fiber but primarily along the grain boundaries. Saphikon is a single crystal alumina filament. The c-axis of the filament is oriented parallel to the fibre surface. The mechanical and physical properties of the saphikon filaments are given in Table 1[37]. N 51A, a borosilicate glass, obtained from Owens Illinois Inc., was used as a matrix in the present study. 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 1l/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 via the chemical reaction [38], SnCl4ðgÞ þ 2H2OðgÞ ! SnO2ðsÞ þ 4HClðgÞ ð1Þ The microstructures of as-received fibres heat treated at 500, 600 and 900 C for 90 minutes 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 characterised by SEM. Single fibre tensile tests were carried out on asreceived, heat treated and SnO2 coated PRD-166 fibres. A random selection of single fibres was made from the Table 1 Room-temperature properties of PRD-166 fiber and Saphikon filament [36,37] Fibre Melting point (C) Density (g/cm3 ) Tensile strength (Mpa) Tensile modulus (GPa) Thermal expansion (106 / C) Saphikon 2053 3.9 3150 380 9.12 (// to c-axis) 7.95 (to c-axis) PRD-166 2045 4.2 2070 380 9.0 Fig. 1. (a) Chevron notch specimen; (b) geometry of chevron notch. 566 R. Venkatesh / Ceramics International 28 (2002) 565–573

R Venkatesh/Ceramics International 28(2002)565-573 material to be tested. The fibres were center-line moun- coeffecient of variation were then evaluated by Weibull ted on a paper frame. The fibres were centered over the analysis [39] frame and lightly stretched. A small amount of adhesive Alumina fiber reinforced glass matrix composites was then carefully placed at each end of the fibre. The were fabricated by a slurry impregnation technique [40] specimen gage length for all the fibres tested was 17 m The slurry consisted of glass frit, 2-propanol and an An Instron tensile testing machine (model 1 120)was organic binder to impart green strength to the tapes and used with a 5 N load cell. Before fibres were loaded onto facilitate their handling. For fabrication of alumina/ the machine, the diameter of the individual fibres was glass composites, a continuous process was employed to measured with an optical microscope. The frame was make unidirectional tapes. For fabrication of alumina, then gripped in the jaws of the testing machine and the SnO2/glass composites, the coated fibers were dipped in mounting frame was burned on the sides. The fibres the slurry and laid on mylar tapes to form prepeg tape were successively stressed to failure with a crosshead These unidirectional tapes were heated to 500oC in air peed of 0.25 mm/min. An average of 80 fibres in each to remove the binder and then hot pressed. The hot group, i.e. as-received, heat treated at 500, 600 and pressing was performed in a graphite lined die in argon 900oC and Sno, coated were tested. The mean tensile atmosphere at 925C and 3 MPa. strength, Weibull modulus, standard deviation and a Optical microscopy was used to evaluate the volume action and fiber distributions in the composites. The fracture surfaces of the composites were characterized using SEM. Three point bending tests on the glass matrix composites were conducted in the longitudinal direction. Bending tests were carried out on specimens having a span length(S) to thickness(W) ratio >8 and thickness(W) to breadth(B)ratio of 0.75. The three point bending tests were conducted on an Instron machine(model 1102)with a crosshead speed of 0.05 10 slope=B 7 slope= 0.10 Tensile Strength, a;( MPa) Tensile Strength, O:(MPa) Fig 3. Weibull plots (a) as-received PRD-166 fibres;(b)SnO2 coated PRD.166 fibres Table 2 Estimated Weibull parameters of as-received and heat treated alumina (PRD. 166)fibres Standard strength(MPa) deviation(MPa) variation(%) As-received 1375 418 1313 Fig. 2. Microstructure of as-received fibres (a)Rough longitudinal 600°C 1283 urface of the fibres:(b) zirconia particles dispersed throughout the 900°C 29

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 successively stressed to failure with a crosshead speed of 0.25 mm/min. An average of 80 fibres in each group, i.e. as-received, heat treated at 500, 600 and 900 C and SnO2 coated were tested. The mean tensile strength, Weibull modulus, standard deviation and coeffecient of variation were then evaluated by Weibull analysis [39]. Alumina fiber reinforced glass matrix composites were fabricated by a slurry impregnation technique [40]. The slurry consisted of glass frit, 2-propanol and an organic binder to impart green strength to the tapes and facilitate their handling. For fabrication of alumina/ glass composites, a continuous process was employed to make unidirectional tapes. For fabrication of alumina/ SnO2/glass composites, the coated fibers were dipped in the slurry and laid on mylar tapes to form prepeg tapes. These unidirectional tapes were heated to 500 C in air to remove the binder and then hot pressed. The hot pressing was performed in a graphite lined die in argon atmosphere at 925 C and 3 MPa. Optical microscopy was used to evaluate the volume fraction and fiber distributions in the composites. The fracture surfaces of the composites were characterized using SEM. Three point bending tests on the glass matrix composites were conducted in the longitudinal direction. Bending tests were carried out on specimens having a span length (S) to thickness (W) ratio > 8 and thickness (W) to breadth (B) ratio of 0.75. The three point bending tests were conducted on an Instron machine (model 1102) with a crosshead speed of 0.05 Fig. 2. Microstructure of as-received fibres. (a) Rough longitudinal surface of the fibres; (b) zirconia particles dispersed throughout the fibre. Fig. 3. Weibull plots (a) as-received PRD-166 fibres; (b) SnO2 coated PRD-166 fibres. Table 2 Estimated Weibull parameters of as-received and heat treated alumina (PRD- 166) fibres Fibre Mean tensile strength (MPa) Standard deviation (MPa) Coefficient of variation (%) As-received 1375 418 30 500 C 1313 386 29 600 C 1283 440 34 900 C 1083 320 29 R. Venkatesh / Ceramics International 28 (2002) 565–573 567

R Venkatesh/Ceramics International 28 (2002)565-573 parameter(a), standard deviation and coefficient of gives 1375(17/50) 3.5=1010 MPa. Both values are variation of as-received and heat treated alumina fibres. lower than measured by the cited authors. The A room temperature strength of 1375 MPa was difference in results obtained could be explained by pre obtained for a gage length of 17 mm. This value is lower sence of voids. It can also be seen from Table 2 that the than that reported by romine [36] who obtained a strength decreases with increase in heat treatment tem strength of 2100 MPa for a gage length of 6.9 mm and perature. Pysher and Tressler [44, 45] have reported the Yang et al. [43] who obtained a strength of 1180 MPa presence of minor elements (0.01-0.1%)of Si, and P for a gage length of 50 mm. There are two effects which could form a SiO2-P2Os glassy grain boundary explaining such results: one is the different gage lengths, phase in PRD-166 fibres at high temperatures. These the other the presence of voids( Fig 4)in the as-received glassy phase could reduce the strength of the fibres. It the Weibull strength at a given has been shown in earlier works that the alumina fibres probability depends on the stressed volume as v-1, for undergo transgranular failure up to 800 oC beyond fibres of the same diameter tested on different gage which the failure mode is intergranular [44, 45 lengths L, L', the strength ratio should be a/o'=L/L Table 3 shows variation of tensile strength. Weibull the result of the present research to be compared with modulus and scale paremeters of SnO2 coated fibres the one by Romine is, accordingly, 1375(17/6.9)/3.5 Again, the tensile strength decreases with increase in 1779 MPa. The value to be compared with Yang et al coating thickness. Some loss in strength can be attrib uted to exposures at 500C during SnO2 deposition Another source of strength reduction could be thermal stresses generated after deposition and subsequent cool ing of the alumina/tin dioxide composite fibre. Thermal stresses at the fibre/coating interface were calculated using a two-element cylinderical model [46, 47, and the results are shown in Table 4. The radial stress is tensile while the axial and circumferential stresses are tensile in the fibre and compressive in the coating. The axial tensile stress and radial stress increases with coating thickness. This state of stress could reduce the strength of the fibres with increase in coating thickness. In order to obtain a measure of the effect of rough ne properties of the fibres, amplitude of 50μ roughness of as-received and SnO2 coated fibres was evaluated. The pea evaluated in the present study using SEM micrographs and tracings of roughness profiles. These results are shown in Table 5. It can be seen that the amplitude of roughness, A increases with increase in coating thick- ness. Under an axial load this roughness could act as a notch and decrease the strength of the fibres with increase in coating thickness The roughness of fibre nd coating hav matrix composites. In order to further study the effect of fibre roughne of fibre alumina fibre(PRD-166) and Saphikon fibre on the fracture characteristics of alumina fibre (PRD-166)/ glass(AG), Saphikon fibre/glass (SG), alumina fibre (PRD-166)/SnOz/glass (ASG) and Saphikon fibre SnO2/glass (SSG)composites were investigated. Evalu- 10 m ating compressive roughness strain, A/r(amplitude of roughness/radius of fibre) using tracings of interphase micrographs Fig. 5(a)and(b)it was observed that the A/r value of PRD-166/SnO, interface was about nine Fig. 7.(a and b) Fracture surface of Saphikon alumina fibre/Sno times that of the Saphikon/SnO interface(Table 6). I glass matrix composites. The SnO2 coating on a relatively smooth was also found that, in PRD-166/ SnO2/glass matrix saphikon fibre results in a neat a long fibre pull out composites, the compressive radial strain induced due

parameter (a), standard deviation and coefficient of variation of as-received and heat treated alumina fibres. 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 [36] who obtained a strength of 2100 MPa for a gage length of 6.9 mm and Yang et al. [43] who obtained a strength of 1180 MPa for a gage length of 50 mm. There are two effects explaining such results: one is the different gage lengths, the other the presence of voids (Fig. 4) in the as-received alumina fibres. As the Weibull strength at a given probability depends on the stressed volume as V-1/b, for fibres of the same diameter tested on different gage lengths L, L0 , the strength ratio should be /0 =L/L0 ; the result of the present research to be compared with the one by Romine is, accordingly, 1375 (17/6.9)1/3.5 =1779 MPa. The value to be compared with Yang et al gives 1375 (17/50)1/3.5=1010 MPa. Both values are lower than measured by the cited authors. The difference in results obtained could be explained by presence of voids. It can also be seen from Table 2 that the strength decreases with increase in heat treatment temperature. Pysher and Tressler [44,45] 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 at high temperatures. These glassy phase could reduce the strength of the fibres. It has been shown in earlier works that the alumina fibres undergo transgranular failure up to 800 C beyond which the failure mode is intergranular [44,45]. Table 3 shows variation of tensile strength, Weibull modulus and scale paremeters of SnO2 coated fibres. Again, 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 after deposition and subsequent cooling of the alumina/tin dioxide composite fibre. Thermal stresses at the fibre/coating interface were calculated using a two-element cylinderical model [46,47], and the results are shown in Table 4. The radial stress is tensile while the axial and circumferential stresses are tensile in the fibre and compressive in the coating. The axial tensile stress and radial stress increases with coating thickness. This state of stress could reduce the strength of the fibres with increase in coating thickness. In order to obtain a measure of the effect of roughness on the properties of the fibres, amplitude of roughness of as-received and SnO2 coated fibres was evaluated. The peak-valley roughness amplitude, A, was evaluated in the present study using SEM micrographs and tracings of roughness profiles. These results are shown in Table 5. It can be seen that the amplitude of roughness, A increases with increase in coating thickness. Under an axial load, this roughness could act as a notch and decrease the strength of the fibres with increase in coating thickness. The roughness of fibre surface and coating have serious implications in the development of tough ceramic matrix composites. In order to further study the effect of fibre roughness, two different types of fibres namely, alumina fibre (PRD-166) and Saphikon fibre on the fracture characteristics of alumina fibre (PRD-166)/ glass (AG), Saphikon fibre/glass (SG), alumina fibre (PRD-166)/SnO2/glass (ASG) and Saphikon fibre/ SnO2/glass (SSG) composites were investigated. Evaluating compressive roughness strain, A/r (amplitude of roughness/radius of fibre) using tracings of interphase micrographs Fig. 5(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 6). It was also found that, in PRD-166/SnO2/glass matrix composites, the compressive radial strain induced due Fig. 7. (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. 570 R. Venkatesh / Ceramics International 28 (2002) 565–573

R Venkatesh/Ceramics International 28(2002)565-573 omposites was only 3%. It should be pointed out that roughness plays a major role in the fracture resistance SSG composites were fabricated using a small quantity of ceramic matrix composites Interfacial roughness in of fibers only to verify the importance of fiber CMCs can be controlled through the use of smooth fibres roughness. The bend strength increased with the volume or reducing the coating thickness fraction of the fibers. The strength of AG was slightl reater than AsG, possibly due to the strong chemical bonding at the fiber/matrix interface leading to better Conclusion load transfer in AG as compared to ASG composites The work of fracture evaluated from the area under the 1. The strength of alumina-ZrO, fibres(PRD-166 load-displacement curve for both coated and uncoated decreased with increase in heat treatment tem- composites is shown in Table 7. This parameter perature of fibres. This could be because of the increased with volume fraction of fibers in both uncoa- presence of a glassy second and or due to pro ted and coated fiber composites. The work of fracture of cessing defects in the fibre ASG composites was larger than that of AG composites 2. The strength of the fibres decreased with increase due to contributions from modulus mismatch, crack in SnOz coating thickness. It was observed that deflection, fiber bridging and fiber pullout. Fracture as the coating thickness increased, roughness of toughness as a function of volume fraction of fibers for the coating increased. This acted as a notch in both AG and AsG composites is shown in Table 7. The order to decrease the strength of the Al,O3/SnO toughness of AG composites obtained in this study is in composite fibre close agreement with that obtained by Michalske and 3. In CMCs, as rough ness Increa sed. fibre debond Hellmann [48]. The toughness of AG and ASG compo- ing decreased and toughness of composite sites increased with volume fraction of fibers. The decreased. PRD-166/SnO2/glass matrix compo- toughness of ASG composites was larger than that of sites exhibited non-planar failure with fiber brid- AG composites. The main contributors to the increase ging and debonding as major toughening in toughness of ASG as compared to AG composites mechanisms. Saphikon/ SnO/glass matrix com- are crack deflection, partial debonding, fiber bridging posites failed in a tough manner with extensive and partial fiber pullout fiber pullout Fracture surfaces of ASG composites are non planar, 4. The difference in the failure mode between prd. Fig. 6. Fig. 6 also shows that the predominant mechan- 6/SnO/glass and Saphikon/SnO2/glass matrix ism of toughening is fiber bridging and fibre debonding composites could be attributed to roughness of Partial pullout of fibers can also be seen. A higher the fibres magnification micrograph of the fracture surface of 5. It is important to control the roughness at the ASG composites along the fiber, [Fig. (6c)), shows fibre/coating and coating/matrix interfaces in clearly the partial removal of the coating and the rough order to develop tough ceramic matrix compo- fiber surface. Hence the primary toughening mechan sites. methods g roughness at the isms in ASG composites are crack deflection, fiber interfaces consist in the incorporation of smooth bridging, partial fiber debonding and pullout. It has fibres and in decreasing coating thickness been shown that as the roughness of interphase increa- es, the compressive clamping stress increases thereby affecting the fracture resistance of ceramic matrix com- posites [49-58]. This increase in the compressive clamp ing stress due to fibre roughness causes an increase in References the shear stress transfer at the interface beyond matrix cracking from fiber to matrix. This causes a reduction in [1 I.w. Donald, P w. McMillan, J Mater. Sci. 11(1976)146 2R. W. Rice, Ceramic matrix composite toughening mechanisms: the debond length, i.e., fibers break rather than debond an update, in Ceramic Science and Engineering Series, Vol. 6, as the matrix crack grows, resulting in a composite fracture surface in ASG with little or no fiber pullout on 3A.G. Evans, Perspective on the development of high-toughness the fracture surface Fracture surface of SSG compo- ceramics, J. Am. Ceram Soc. 73(1990)187 sites, Fig. 7(a and b), showed neat fiber debonding and 4W.B. Hillig, Strength and toughness of ceramic matrix compo- fiber pullout at the fiber/SnO2 interface, as confirmed [D.B. Marshall, J E. Ritter, Reliability of advan also by EDs on the pulled out saphikon fiber in Fig 8 ceramics and ceramic matrix composites-a review, In the region marked as a, the EDs analysis showed 66(1987)309 only Al and no Sn since SnO, and Al, O, have no [6 H.G. Sowman, DD. Johnson, in: K.s. Mazdiyani (Ed. ), Fibre mutual solubility. Hence fibre debonding and fibre pullout with lengths over 100 um take place in saphikon Technology, Noyes Publications, USA, 1990, pp l-/sand [7 G. Das, Ceram. Eng. Sci. Proc. 16(5)(1995)977-986 fibre/SnO2/glass matrix composites. Thus interfacial [8 G. Das, Ceram. Eng. Sci. Proc. 18(6)(1997)25-33

composites was only 3%. It should be pointed out that SSG composites were fabricated using a small quantity of fibers only to verify the importance of fiber roughness. The bend strength increased with the volume fraction of the fibers. The strength of AG was slightly greater than ASG, possibly due to the strong chemical bonding at the fiber/matrix interface leading to better load transfer in AG as compared to ASG composites. The work of fracture evaluated from the area under the load-displacement curve for both coated and uncoated composites is shown in Table 7. This parameter increased with volume fraction of fibers in both uncoated and coated fiber composites. The work of fracture of ASG composites was larger than that of AG composites due to contributions from modulus mismatch, crack deflection, fiber bridging and fiber pullout. Fracture toughness as a function of volume fraction of fibers for both AG and ASG composites is shown in Table 7. The toughness of AG composites obtained in this study is in close agreement with that obtained by Michalske and Hellmann [48]. The toughness of AG and ASG composites increased with volume fraction of fibers. The toughness of ASG composites was larger than that of AG composites. The main contributors to the increase in toughness of ASG as compared to AG composites are crack deflection, partial debonding, fiber bridging and partial fiber pullout. Fracture surfaces of ASG composites are non planar, Fig. 6. Fig. 6 also shows that the predominant mechanism of toughening is fiber bridging and fibre debonding. Partial pullout of fibers can also be seen. A higher magnification micrograph of the fracture surface of ASG composites along the fiber, [Fig. (6c)], 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. It has been shown that as the roughness of interphase increases, the compressive clamping stress increases thereby affecting the fracture resistance of ceramic matrix composites [49–58]. This increase in the compressive clamping stress due to fibre roughness causes an increase in the shear stress transfer at the interface beyond matrix cracking from fiber to matrix. This causes a reduction in the debond length, i.e., fibers break rather than debond as the matrix crack grows, resulting in a composite fracture surface in ASG with little or no fiber pullout on the fracture surface Fracture surface of SSG composites, Fig. 7 (a and b), showed neat fiber debonding and fiber pullout at the fiber/SnO2 interface, as confirmed also by EDS on the pulled out saphikon fiber in Fig. 8. In the region marked as a, the EDS analysis showed only Al and no Sn since SnO2 and Al2O3 have no mutual solubility. Hence fibre debonding and fibre pullout with lengths over 100 mm take place in saphikon fibre/SnO2/glass matrix composites. Thus interfacial roughness plays a major role in the fracture resistance of ceramic matrix composites. Interfacial roughness in CMCs can be controlled through the use of smooth fibres or reducing the coating thickness. Conclusion 1. 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 and/or due to processing defects in the fibres. 2. The strength of the fibres decreased with increase in SnO2 coating thickness. 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 Al2O3/SnO2 composite fibres. 3. In CMCs, as roughness increased, fibre debonding decreased and toughness of composite decreased. PRD-166/SnO2/glass matrix composites exhibited non-planar failure with fiber bridging and debonding as major toughening mechanisms. Saphikon/SnO2/glass matrix composites failed in a tough manner with extensive fiber pullout. 4. The difference in the failure mode between PRD- 166/SnO2/glass and Saphikon/SnO2/glass matrix composites could be attributed to roughness of the fibres. 5. 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 consist in the incorporation of smooth fibres and in decreasing coating thickness. References [1] I.W. Donald, P.W. McMillan, J. Mater. Sci. 11 (1976) 146. [2] R.W. Rice, Ceramic matrix composite toughening mechanisms: an update, in Ceramic Science and Engineering Series, Vol. 6, American Ceramic Society, Columbus, OH, p.589. [3] A.G. Evans, Perspective on the development of high-toughness ceramics, J. Am. Ceram. Soc. 73 (1990) 187. [4] W.B. Hillig, Strength and toughness of ceramic matrix composites, Ann. Rev. Mater. Sci. 17 (1987) 341. [5] D.B. Marshall, J.E. Ritter, Reliability of advanced structural ceramics and ceramic matrix composites–a review, Ceram. Bull. 66 (1987) 309. [6] H.G. Sowman, D.D. Johnson, in: K.S. Mazdiyani (Ed.), Fibre Reinforced Ceramic Composites: Materials Processing and Technology, Noyes Publications, USA, 1990, pp. 122–138. [7] G. Das, Ceram. Eng. Sci. Proc. 16 (5) (1995) 977–986. [8] G. Das, Ceram. Eng. Sci. Proc. 18 (6) (1997) 25–33. 572 R. Venkatesh / Ceramics International 28 (2002) 565–573

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[9] D.M. Wilson, D.C. Luenberg, S.L. Lieder, Ceram. Eng. Sci. Proc. 14 (1993) 7–8. [10] D.M. Wilson, D.C. Luenberg, S.L. Lieder, Ceram. Eng. Sci. Proc., 16(5) (1995) 1005–1014. [11] J.A. DiCarlo, S. Dutta, in: R. Lehmanm, S. El-Rahaiby, J. Wachatman Jr. (Eds.), Handbook on Continuous Fibre Reinforced Ceramic Matrix Composites, CIAC, Purdue University, IN, 1995, 137–183. [12] R. Venkatesh, Mater. Sci. Eng. A268 (1999) 47–54. [13] M.H. Stacey, Br. Ceram. Trans. 87 (1988) 168–172. [14] E. Rasterrer, W.K. Symes, Interceram (1983) 215–220. [15] J.D. Birchall, Br. Ceram. Soc. 82 (1983) 143–145. [16] T. Yogo, H. Iwahara, J. Mater Sci, 26 (1991) 5292–5296. [17] W. Gaubitt, W. Watzke, Scholz, D. Sporn, J. Sol–Gel Sci. Technol. 8 (1997) 29–33. [18] J. Dinwoodie, Canadian Ceram. Quarterly 2 (1996) 23–29. [19] A.R. Bunsell, H. Berger, Compos Sci. Technol. 51 (1994) 127– 133. [20] R. Venkatesh, S.R. Ramanan, J. Eur. Ceram. Soc. 20 (2000) 2543–2549. [21] R. Venkatesh, L.C. Teck, Z.A. Ahmad, A. Fauzi, J. Mater. Tech. 16(2) (2001) 103. [22] C.G. Levi, J.Y. Yang, B.J. Dalgeish, F.W. Zok, A.E. Evans, J. Am. Ceram. Soc. 81(8) (1998) 2077–2086. [23] R. Lundberg, L. Eckerbar, in: R. Naslain (Ed.), High Tech. CMCs II, Woodhead Publ., NY, 1994, pp. 95–103. [24] R.N. Singh, M.K. Brun, Ceram. Eng. Sci. Proc. 8 (1987) 636– 643. [25] O. Yeheskel, M.L. Balmer, D.C. Cranmer, Ceram. Eng. Sci. Proc. 9 (1988) 687–694. [26] R. Lundberg, L. Pejryd, E. Butler, M. Kelund, M. Nygreen, in: R. Naslain, L. Doumeingh (Eds.), Proc. Int. Conf. on HTCMC-I, Woodhead Publ., UK, 1993, pp. 167–174. [27] P.E.D. Morgan, D.B. Marshall, J. Am. Ceram. Soc. 78 (1995) 1553–1563. [28] M.K. Cinnibulk, Ceram. Eng. Sci. Proc. 16(5) (1995) 721–728. [29] M.K. Cinnibulk, Ceram. Eng. Sci. Proc. 15(5) (1994) 721–728. [30] K.K. Chawla, Z.R. Xu, R. Venkatesh, J.S. Ha, 9th Int. Conf. on Composites, Spain, 1993, pp. 789–795. [31] R. Venkatesh, P.K. Chakrabarthy, B. Siladitya, M. Chatterjee, D. Ganguli, Ceram. Int. 25 (1999) 539–543. [32] J.B. Davis, J.B. Lofrander, A.G. Evans, E. Bischoff, M.C. Emiliano, J. Am. Ceram. Soc. 76(5) (1993) 249–257. [33] R. Venkatesh, K.K. Chawla, J. Mater. Sci. Lett. 11 (1992) 650– 653. [34] P.E.D. Morgan, D.B. Marshall, J. Am. Ceram. Soc. 78 (6) (1995) 1553–1556. [35] K.K. Chawla, A. Choudhary, R. Venkatesh, J.R. Hellmann, Mater. Char. 31 (3) (1993) 167–172. [36] J.C. Romine, Ceram. Eng. Sci. Proc. 8 (1987) 1755–1760. [37] Saphikon data sheet, Saphikon, NH. [38] G.N. Advani, A.G. Jordan, C.H.P. Lupis, R.L. Longini, Thin Solid Films 62 (1979) 361–369. [39] W. Weibull, Appl. Mech. Rev. 5 (1952) 449–454. [40] K. Prewo, J.J. Brennan, G.K. Layden, Am. Ceram. Soc. Bull. 65 (1980) 305–312. [41] J.I. Bluhm, Eng. Fract. Mech. (1984) 221–225. [42] Wu-Shang Xian, Eng. Frac. Mech. 19 (1984) 271–276. [43] X. Yang, X. Hu, R.J. Day, R.J. Young, J. Mater. Sci. 27 (1992) 1409–1416. [44] J. Pysher, K.C. Goretta, R.S. Hodder, R.E. Tressler, J. Am. Ceram. Soc. 72 (2) (1989) 284–288. [45] D.J. Pysher, R.E. Tressler, J. Mater. Sci. 27 (1992) 423–428. 65. [46] H. Poritsky, Phy. 5 (1934) 406–410. [47] A.W. Hull, E.E. Burger, Phy. 5 (1934) 384–389. [48] T.A. Michalske, J.R. Hellmann, J. Am. Ceram. Soc. 71(1998) 725–732. [49] T.A. Parthasarathy, R.J. Kerans, J. Am. Ceram. Soc. 80(5) (1992) 2043–2055. [50] C.H. Hseuh, J. Mater. Sci. Lett. 16 (5) (1997) 350–358. [51] R.J. Kerans, T.A. Parthasarathy, J. Am. Ceram. Soc. 74 (1991) 1585–1600. [52] T.J. Mackin, J. Yang, P.D. Warren, J. Am. Ceram. Soc. 75(12) (1992) 3358–3365. [53] T.A. Parthasarathy, R.J. Kerans, J. Am. Ceram. Soc. 80 (8) (1992) 2043–2501. [54] H.P. Wang, C.L. Nelson, C.L. Lim, W.W. Gerberch, Interfacial stability and mechanicxal properties of alumina fibre reinforced Ti matrix composites, J. Mater. Res. 9 (2) (1994) 498. [55] P.D. Jero, T.A. Parthasarathy, Interface properties their measurement with fibre pushout tests, in R. Naslain (Ed.), High temperature Ceramic Matrix Composites, Goodhead Publishers, Cambridge, 1993, p. 401. [56] B.F. Sorensen, Effect of fibre roughness on the overall stress– transverse strain response of ceramic composites, Scripta Metall. Et Mater. 28 (1993) 435. [57] W.C. Carter, E.P. Butler, E.R. Fuller Jr., Scripta Metall. Mater. 25 (1991) 579–585. [58] T.J. Mackin, P.D. Warren, A.G. Evans, Acta Metall. 40 (1992) 1251–1259. R. Venkatesh / Ceramics International 28 (2002) 565–573 573

《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Interphase effects on the bend strength and toughness of an oxide

《复合材料 Composites》课程教学资源（学习资料）第五章陶瓷基复合材料_Interphase effects on the bend strength and toughness of an oxide