ATERIALS GENGE S ENGIEERING ELSEVIER Materials Science and Engineering A268(1999)47-54 Mechanical properties of alumina fiber/glass matrix composites with and without a tin dioxide interface Ramanan venkatesh School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia (KCP), Tronoh, 31750, Perak, Malaysia Received 15 November 1998: received in revised form 15 February 1999 Abstract Alumina+ zirconia(PRD-166)and Saphikon fibers reinforced glass matrix composites with and without a SnO, interphase were prepared by slurry infiltration and their mechanical characteristics were evaluated. Longitudinal bend strength increased with volume fraction of fibers in both PRD-166/glass and PRD-166/SnO,/glass matrix composites. PRD-166/glass matrix composites failed in a brittle manner whereas PRD-166/SnO, glass matrix composites exhibited non-planar failure with crack deflection and fiber bridging as major toughening mechanisms. Saphikon /SnO2/glass matrix composites failed in a tough manner with extensive fiber pullout. The difference in the failure mode between PRD-166/Sno /glass and Saphikon/SnO,glass matrix composites w lue to fiber roughness. The toughness of PRD-166/ SnO /glass matrix composites was due to crack deflection, fiber bridging, partial fiber debonding and some fiber pullout. Major toughening mechanisms in Saphikon /SnO glass matrix composites were iber debonding and fiber pullout with minor contributions due to crack deflection and fiber bridging o 1999 Elsevier Science S.A. All rights reserved Keywords: Glass matrix composites; Tin dioxide interface: Saphikon matrix microcracking, matrix prestressing, fiber debonding, crack deflection and fiber pullout [1-5]. For Ceramics have strong ionic/covalent bonds and very a tough CMC, the interface between the fiber and the few slip systems compared to metals. This makes ce- matrix should be strong enough for load transfer but ramic materials strong but brittle. The high thermal weak enough to allow crack deflection. The interface stability of these materials coupled with their low den- bonding can be controlled either by selecting fiber and sity and high elastic moduli make them very attractive matrix materials which are thermodynamically stable at for high temperature applications [1] structural the processing and service temperatures or by applying materials, monolithic ceramics suffer from two impor- coatings that act as diffusion barriers, thereby prevent tant reliability issues, namely high sensitivity to process- ing a strong bond between the fiber and the matrix [6 ing and service generated flaws, and the inability to Different coatings that have been tried include C, Bn, tolerate stress levels. To increase the work of fracture SiC, TiC, ZrO2, SnO2, Monazite, magnetoplumbite, etc or the toughness of monolithic ceramics, numerous [6-10 ughening methods have been employed. These Alumina type oxide fibers appear to have consider- clude toughening by dispersion of second phase parti ble potential in reinforcing ceramic and glass matrix cles, i.e. transformation toughening and toughening by composites due to its oxidative stability and moderately incorporation of whiskers, ductile fibers or ceramic high strength and stiffness. Glass matrix composites fibers. Fiber reinforcement offers a great potential for offer a great commercial potential regarding tempera- improving strength and toughness of ceramic materials ture and environmental stability due to their ease of Increase in strength and toughness of CMCs can be fabrication, tailoring of properties achieved through a multitude of mechanisms namel processing. and glass react to form a strong che E-mail address: ramSnan @tm net. my(R. Venkatesh) bond at the interface thereby resulting in a brittle 0921-5093/99/S- see front matter c 1999 Elsevier Science S.A. All rights reserved. PI:s0921-5093099)00115-X
Materials Science and Engineering A268 (1999) 47–54 Mechanical properties of alumina fiber/glass matrix composites with and without a tin dioxide interface Ramanan Venkatesh School of Materials and Mineral Resources Engineering, Uni6ersiti Sains Malaysia (KCP), Tronoh, 31750, Perak, Malaysia Received 15 November 1998; received in revised form 15 February 1999 Abstract Alumina+zirconia (PRD-166) and Saphikon fibers reinforced glass matrix composites with and without a SnO2 interphase were prepared by slurry infiltration and their mechanical characteristics were evaluated. Longitudinal bend strength increased with volume fraction of fibers in both PRD-166/glass and PRD-166/SnO2/glass matrix composites. PRD-166/glass matrix composites failed in a brittle manner whereas PRD-166/SnO2/glass matrix composites exhibited non-planar failure with crack deflection and fiber bridging as major toughening mechanisms. Saphikon/SnO2/glass matrix composites failed in a tough manner with extensive fiber pullout. The difference in the failure mode between PRD-166/SnO2/glass and Saphikon/SnO2/glass matrix composites was due to fiber roughness. The toughness of PRD-166/SnO2/glass matrix composites was due to crack deflection, fiber bridging, partial fiber debonding and some fiber pullout. Major toughening mechanisms in Saphikon/SnO2/glass matrix composites were fiber debonding and fiber pullout with minor contributions due to crack deflection and fiber bridging. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Glass matrix composites; Tin dioxide interface; Saphikon 1. Introduction Ceramics have strong ionic/covalent bonds and very few slip systems compared to metals. This makes ceramic materials strong but brittle. The high thermal stability of these materials coupled with their low density and high elastic moduli make them very attractive for high temperature applications [1]. As structural materials, monolithic ceramics suffer from two important reliability issues, namely high sensitivity to processing and service generated flaws, and the inability to tolerate stress levels. To increase the work of fracture or the toughness of monolithic ceramics, numerous toughening methods have been employed. These include toughening by dispersion of second phase particles, i.e. transformation toughening and toughening by incorporation of whiskers, ductile fibers or ceramic fibers. Fiber reinforcement offers a great potential for improving strength and toughness of ceramic materials. Increase in strength and toughness of CMCs can be achieved through a multitude of mechanisms namely matrix microcracking, matrix prestressing, fiber debonding, crack deflection and fiber pullout [1–5]. For a tough CMC, the interface between the fiber and the matrix should be strong enough for load transfer but weak enough to allow crack deflection. The interface bonding can be controlled either by selecting fiber and matrix materials which are thermodynamically stable at the processing and service temperatures or by applying coatings that act as diffusion barriers, thereby preventing a strong bond between the fiber and the matrix [6]. Different coatings that have been tried include C, BN, SiC, TiC, ZrO2, SnO2, Monazite, magnetoplumbite, etc. [6–10]. Alumina type oxide fibers appear to have considerable potential in reinforcing ceramic and glass matrix composites due to its oxidative stability and moderately high strength and stiffness. Glass matrix composites offer a great commercial potential regarding temperature and environmental stability due to their ease of fabrication, tailoring of properties and low cost of processing. Alumina and glass react to form a strong chemical E-mail address: ram5nan@tm.net.my (R. Venkatesh) bond at the interface thereby resulting in a brittle 0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S0921-5093(99)00115-X
tkatesh/Materials Science and Engineering 4268(1999)47 Table I Room-temperature properties of PRD-166 fiber and Saphikon filament Fiber Melting point Density (g/cm) Tensile strength Tensile modulus Thermal expansion(x 10-6/C) 9. 12(parallel to c-axis); 7.95(perp. to c-axis) PRD-1662045 9.0 composite. It has been shown that tin dioxide which saphikon /Sno2glass composites. The fracture has no diffusion in alumina up to 1400C in partia urfaces of AG. asc and ssg were characterized pressure of oxygen >10-/atmospheres [11, 12] and using SEM. Percent ers aligned parallel to the very little diffusion in glass [8] could be an ideal candi- fiber axis and fibers broken in as-fabricated AG and date for a barrier coating. The present work examines ASG composites were evaluated by lightly etching the the strength and toughness of alumina-fiber(PRD-166) composite surface with HF and observing under SEM lber reinforced glass matrix composites with and with- Three-point bend tests on AG, ASG, SG and SsG out a tin dioxide interphase. were conducted in the longitudinal direction, LI and transverse directions Ti and T2. The orientation of the fibers with espect to the applied load 2. Materials and experimental procedure samples Li, T and T2 is shown in Fig. 1. Bend tests were carried out on specimens having a span length(S) The PRD-166 fiber used in the present work is a breadth(B)ratio of 0.75. The three-point bed to to thickness () ratio >8 and thickness (w polycrystalline a-Al,O3 fiber, 20 um in diameter and contains 15-20 wt%Y,O were conducted on an Instron machine(model partially stabilized zirconia with a crosshead speed of 0.005 cm/min. The bend particles.The properties of PRD-166 fiber are given in strength, o is given as Table 1 [13. The zirconia particles are dispersed throughout the fiber but primarily along the grain 0,=3/2(PS/BW2) boundaries. The grain size of alumina, as determined by the linear intercept method was about 0.5 um, and where P is the maximum load the ize of zirconia particles was 0.33 Three-point bend tests were also conducted on AG Saphikon is a single crystal alumina filament. The and ASG composites at 200, 400 and 600C c-axis of the filament is oriented parallel to the fiber Fracture toughness of AG, asG, SG and SSg were surface. The mechanical and physical properties of the evaluated using chevron notch specimens. A chevron Saphikon filaments are given in Table 1 [14]. N51A, a notch specimen is as shown in Fig. 2a. A specimen borosilicate glass(Owens Illinois ) was used. Some me- geometry having a span-to-thickness ratio of 4 and chanical properties of N51A glass and SnO2 are given thickness-to-width ratio of 1.5 was used to evaluate the Table 2 fracture toughness. The three-point bend tests were Alumina fiber reinforced glass matrix composites conducted on an Instron machine(model 1102) with vere fabricated by slurry impregnation techniques[15]. crosshead speed of 0.005 cm/min. The fracture tough The slurry consisted of glass frit, 2-propanol and an ness(Kle) was evaluated by the following equation organic binder to impart green strength to the tapes Ke=(P/Bw 1/2)Yc(do 41) and facilitate their handling. For fabrication of alu (2) mina/glass composites, a continuous process was em ployed to make unidirectional tapes. For fabrication of Table 2 Properties of N51A glass and tin dioxide lumina/SnO2 glass composites, fiber tows of 5 cm in length were coated with SnO, using a CVD process by Tin dioxide Glass reaction of SnCL and H,0 at 500C. The coated fibers were then dipped in the slurry and laid on mylar tapes E(GPa) 72 to form prepeg tapes. These unidirectional tapes were VHN(GPa) 0.7-0.8 heated to 500C in air to remove the binder and then Density (g/cm) hot pressed. The hot pressing was performed in a Thermal expansion(10-C) graphite lined die in argon atmosphere at 925C and 3 Melting point(oC) MPa Annealing point (C) Optical microscopy was used to evaluate the volume Softening point (C) fraction and fiber distributions alumina/glass (AG) Boiling point (C) 1800-1900 Molecular weight alumina / SnO /glass (ASG), Saphikon/glass (SG)and
48 R. Venkatesh Venkatesh / Materials Science and Engineering A Materials Science and Engineering A268 (1999) 47–54 268 (1999) 47–54 Table 1 Room-temperature properties of PRD-166 fiber and Saphikon filament Fiber Density (g Melting point /cm3 ) Tensile strength Thermal expansion (×10−6 Tensile modulus /°C) (°C) (MPa) (GPa) Saphikon 3.9 2053 3150 380 9.12 (parallel to c-axis); 7.95 (perp. to c-axis) PRD-166 4.2 2045 2070 380 9.0 composite. It has been shown that tin dioxide which has no diffusion in alumina up to 1400°C in partial pressure of oxygen \10−7 atmospheres [11,12] and very little diffusion in glass [8] could be an ideal candidate for a barrier coating. The present work examines the strength and toughness of alumina-fiber (PRD-166) fiber reinforced glass matrix composites with and without a tin dioxide interphase. 2. Materials and experimental procedure The PRD-166 fiber used in the present work is a polycrystalline a-Al2O3 fiber, 20 mm in diameter and contains 15–20 wt.% Y2O3 partially stabilized zirconia particles. The properties of PRD-166 fiber are given in Table 1 [13]. The zirconia particles are dispersed throughout the fiber but primarily along the grain boundaries. The grain size of alumina, as determined by the linear intercept method was about 0.5 mm, and the average size of zirconia particles was 0.33 mm. Saphikon is a single crystal alumina filament. The c-axis of the filament is oriented parallel to the fiber surface. The mechanical and physical properties of the Saphikon filaments are given in Table 1 [14]. N51A, a borosilicate glass (Owens Illinois), was used. Some mechanical properties of N51A glass and SnO2 are given in Table 2. Alumina fiber reinforced glass matrix composites were fabricated by slurry impregnation techniques [15]. 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, fiber tows of 5 cm in length were coated with SnO2 using a CVD process by reaction of SnCl4 and H2O at 500°C. The coated fibers were then 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 alumina/glass (AG), alumina/SnO2/glass (ASG), Saphikon/glass (SG) and saphikon/SnO2/glass (SSG) composites. The fracture surfaces of AG, ASG, SG, and SSG were characterized using SEM. Percent of fibers aligned parallel to the fiber axis and fibers broken in as-fabricated AG and ASG composites were evaluated by lightly etching the composite surface with HF and observing under SEM. Three-point bend tests on AG, ASG, SG and SSG were conducted in the longitudinal direction, L1 and transverse directions T1 and T2. The orientation of the fibers with respect to the applied load in bend test samples L1, T1 and T2 is shown in Fig. 1. Bend 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 bend tests were conducted on an Instron machine (model 1102) with a crosshead speed of 0.005 cm/min. The bend strength, sb, is given as sb=3/2(PS/BW 2 ) (1) where P is the maximum load. Three-point bend tests were also conducted on AG and ASG composites at 200, 400 and 600°C. Fracture toughness of AG, ASG, SG and SSG were evaluated using chevron notch specimens. A chevron notch specimen is as shown in Fig. 2a. A specimen geometry having a span-to-thickness ratio of 4 and thickness–to-width ratio of 1.5 was used to evaluate the fracture toughness. The three-point bend tests were conducted on an Instron machine (model 1102) with a crosshead speed of 0.005 cm/min. The fracture toughness (K1c) was evaluated by the following equation Klc=(P/BW 1/2 )Yc(a0, a1) (2) Table 2 Properties of N51A glass and tin dioxide Tin dioxide Glass E (GPa) 72 233 VHN (GPa) 1.13 0.63 K1c (MPa m – 0.7–0.8 1 2 ) Density (g/cm 6.95 3 ) – Thermal expansion(10−6 /°C) 5.23 7 Melting point (°C) 1630 – Annealing point (°C) 570 – Softening point (°C) – 785 Boiling point (°C) – 1800–1900 Molecular weight – 150
The bend strength of AG and Asg matrix com posites as a function of volume fraction of fibers shown in Table 3. The bend strength increased with 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. Typical stress-displacement curves obtained during bend tests on AG and ASG composites are shown in Fig 4. AG composites showed a linear stress displacement curve typical for a brittle CMC with strong fiber /matrix bonding. The stress increased lin- early up to a maximum beyond which catastrophic failure of the composite occurred. ASG composites exhibited a tougher stress-displacement behavior. The relatively graceful characteristics of the stress-displace- ment curve in ASG composites can be attributed to crack deflection and fiber bridging Fracture surfaces of AG composites are planar(Fig 5)while the fracture surfaces of ASG composites are non-planar(Fig. 6). Fig. 6 also shows that the predom- Fig. 1. Orientation of fibers with respect to applied load in a inant mechanism of toughening is through crack deflec three-point test (a) Longitudinal. (b) Transverse I1(e)Transverse tion with some fiber bridging. Partial pullout of fibers can also be seen. A higher magnification micrograph of where P is maximum load, and Ycao, a, is a dimen- the fracture surface of ASG composites along the fiber sionless stress intensity factor. From a slice model [16] (Fig. 6c)shows clearly the partial removal of the coat for the specimen geometry used, Yc can be evaluated as ing and the rough fiber surface. Hence the primary (5.639+274x+18.93x-43.42x8+3389x where a=a/W, o=dow and a,=a,/w, a is the crack length, ao is the initial crack length(distance from line of load application to tip of chevron notch) and a, is the length of chevron notch at the specimen surface as shown in Fig. 2b 3. Results and discussion Fig. 3(a, b)shows the distribution of the fibers in AG and ASG composites. Volume fraction of the fibers in AG and ASG composites ranged from 12 to 40% and 20 to 36%, respectively. Volume fraction of SG and SSG composites was only 3%. It should be pointed out hat SSG composites were fabricated using a small quantity of fibers only to verify the importance of fiber roughness. Density measurements of AG and AS composites showed porosity to be less than 5%. It was found that as the volume fraction of fibers increased fiber breakage increased while fiber misorientation de creased [18] Fig. 2.(a) Chevron notch specimens. (b)Geometry of chevron notch
R. Venkatesh Venkatesh / Materials Science and Engineering A Materials Science and Engineering A268 (1999) 47–54 268 (1999) 47–54 49 Fig. 1. Orientation of fibers with respect to applied load in a three-point test. (a) Longitudinal. (b) Transverse T1. (c) Transverse T2. The bend strength of AG and ASG matrix composites as a function of volume fraction of fibers is shown in Table 3. The bend strength increased with 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. Typical stress-displacement curves obtained during bend tests on AG and ASG composites are shown in Fig. 4. AG composites showed a linear stressdisplacement curve typical for a brittle CMC with strong fiber/matrix bonding. The stress increased linearly up to a maximum beyond which catastrophic failure of the composite occurred. ASG composites exhibited a tougher stress-displacement behavior. The relatively graceful characteristics of the stress-displacement curve in ASG composites can be attributed to crack deflection and fiber bridging. Fracture surfaces of AG composites are planar (Fig. 5) while the fracture surfaces of ASG composites are non-planar (Fig. 6). Fig. 6 also shows that the predominant mechanism of 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. 6c) shows clearly the partial removal of the coating and the rough fiber surface. Hence the primary where P is maximum load, and Yc(a0, a1) is a dimensionless stress intensity factor. From a slice model [16] for the specimen geometry used, Yc can be evaluated as [17] Yc(a0, a1) =(5.639+27.44a0+18.93a0 2−43.42a0 3+338.9a0 4 ) (3) where a=a/W, a0=a0/W and a1=a1/W, a is the crack length, a0 is the initial crack length (distance from line of load application to tip of chevron notch) and a1 is the length of chevron notch at the specimen surface as shown in Fig. 2b. 3. Results and discussion Fig. 3(a,b) shows the distribution of the fibers in AG and ASG composites. Volume fraction of the fibers in AG and ASG composites ranged from 12 to 40% and 20 to 36%, respectively. Volume fraction of SG and SSG 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. Density measurements of AG and ASG composites showed porosity to be less than 5%. It was found that as the volume fraction of fibers increased, fiber breakage increased while fiber misorientation decreased [18]. Fig. 2. (a) Chevron notch specimens. (b) Geometry of chevron notch
400 Fig. 4. Stress-displacement curve for a uncoated and coated alumina/ of ASG composites was greater than AG composites due to contributions from modulus mismatch. crack 20p deflection, fiber bridging and some fiber pullout. Frac ture surface of SSG composites(Fig. 7a, b )showed neat Fig 3. As processed composites (a) Alumina fiber/glass.(b) Alumina fiber debonding and fiber pullout at the fiber/ Sno interface. Energy dispersive analysis on the pulled out Saphikon fiber surface showed only Al and no Sn [22] oughening mechanisms in ASG composites are crack Transverse strength of AG and AsG composites in deflection, fiber bridging, partial fiber debonding and direction Ti and T2 is shown in Table 3. The strengths pullout because of the relatively weak fiber/coating of these composites in the direction Ti and T2 are about interface [19-21 half those in the longitudinal direction. Aligned fiber The work of fracture evaluated from the area under composites are very weak when stressed perpendicular the load-displacement curve for both coated and un to the fiber axis because the fibers do not contribute to coated composites is shown in Table 4. Work of frac- strength. Table 3 shows that the composites behave ture increased with volume fraction of fibers in both uncoated and coated fiber composites. Work of fracture Table 3 Bend strength of AG and ASG composites in L, T and T> directions r(% Bend strength(MPa) TI T, 060 ASG Fig. 5. Fracture surface of PRD-166 fiber/glass composite showing
50 R. Venkatesh Venkatesh / Materials Science and Engineering A Materials Science and Engineering A268 (1999) 47–54 268 (1999) 47–54 Fig. 3. As processed composites. (a) Alumina fiber/glass. (b) Alumina fiber/SnO2/glass. Fig. 4. Stress-displacement curve for a uncoated and coated alumina/ glass composite during longitudinal bend test. of ASG composites was greater than AG composites due to contributions from modulus mismatch, crack deflection, fiber bridging and some fiber pullout. Fracture surface of SSG composites (Fig. 7a,b) showed neat fiber debonding and fiber pullout at the fiber/SnO2 interface. Energy dispersive analysis on the pulled out Saphikon fiber surface showed only Al and no Sn [22]. Transverse strength of AG and ASG composites in direction T1 and T2 is shown in Table 3. The strengths of these composites in the direction T1 and T2 are about half those in the longitudinal direction. Aligned fiber composites are very weak when stressed perpendicular to the fiber axis because the fibers do not contribute to strength. Table 3 shows that the composites behave toughening mechanisms in ASG composites are crack deflection, fiber bridging, partial fiber debonding and pullout because of the relatively weak fiber/coating interface [19–21]. The work of fracture evaluated from the area under the load-displacement curve for both coated and uncoated composites is shown in Table 4. Work of fracture increased with volume fraction of fibers in both uncoated and coated fiber composites. Work of fracture Fig. 5. Fracture surface of PRD-166 fiber/glass composite showing planar brittle failure. Table 3 Bend strength of AG and ASG composites in L, T1 and T2 directions V Bend strength (MPa) f (%) L T1 T2 AG 12 65 110 60 20 140 – – 26 100 205 75 30 – – 215 42 80 230 70 ASG 24 60 65 120 36 150 75 75 46 – – 190
shows the fracture surface of AG composites. Some breakage of the fibers can be noted. the decrease in high may be either due to porosity of the glass matrix or due to increased breakage of the fibers during processing with increase in volume fraction. The strength of ASG composites is seen to increase up to a volume fraction of 0.36. Fig. 8b shows the fracture surface of ASG Work of fracture of Ag and ASg tially showed the same behavior as the bend strength in directions Ti and T2(table 4). The work of fracture of 50 ASG composites was greater than AG composites be- cause of the nonplanar fracture mode in coated composites Fracture toughness as a function of volume fraction of fibers for both AG and ASG composites is shown in Table 5. The toughness of AG composites obtained in this study is in close agreement with that obtained by Michalske and Hellmann [23]. The toughness of AG and ASG composites increased with volume fraction of fibers. The toughness of ASG composites was greater 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 some fiber pullout Strength of AG and ASg composites decreased with increasing temperature (Table 6). The decrease in 10m strength was not very significant up to 400@C. At 600C softening of the glass matrix occurs. Fracture surfaces of AG and ASG composites at 400oC(Fig. 9a, b) showed that as temperature increased the effect of the various toughening mechanisms decreased namely, crack deflection, fiber bridging, and crack front debonding of the fiber. Post fabrication residual stresses arising due to ther mal expansion mismatch between fiber and the matrix govern the bonding at the fiber matrix interface. Using a two-element (fiber and matrix) and three-element model (fiber, coating and interface) thermal stresses evaluated on both AG and ASG composites showed Table 4 Work of fracture of AG and ASG composites in L. T1 and T2 directions 5 Work of fracture (/m2) L T Fig. 6. Fracture surface of PRD.166 fiber/SnO,/glass composite(a) Non-planar failure. (b) Partial pullout of PRD-166 fiber. (c)Parti removal of coating and rough fiber surface 770 strength of AG composites increased up to a volume SG essentially the same in directions T and T,. The 900 fraction of fibers of 0. 25 and then decreased. Fig. 8a
R. Venkatesh Venkatesh / Materials Science and Engineering A Materials Science and Engineering A268 (1999) 47–54 268 (1999) 47–54 51 Fig. 6. Fracture surface of PRD-166 fiber/SnO2/glass composite. (a) Non-planar failure. (b) Partial pullout of PRD-166 fiber. (c) Partial removal of coating and rough fiber surface. shows the fracture surface of AG composites. Some breakage of the fibers can be noted. The decrease in strength of AG composites at high volume fractions may be either due to porosity of the glass matrix or due to increased breakage of the fibers during processing with increase in volume fraction. The strength of ASG composites is seen to increase up to a volume fraction of 0.36. Fig. 8b shows the fracture surface of ASG composites. Work of fracture of AG and ASG composites essentially showed the same behavior as the bend strength in directions T1 and T2 (Table 4). The work of fracture of ASG composites was greater than AG composites because of the nonplanar fracture mode in coated composites. Fracture toughness as a function of volume fraction of fibers for both AG and ASG composites is shown in Table 5. The toughness of AG composites obtained in this study is in close agreement with that obtained by Michalske and Hellmann [23]. The toughness of AG and ASG composites increased with volume fraction of fibers. The toughness of ASG composites was greater 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 some fiber pullout. Strength of AG and ASG composites decreased with increasing temperature (Table 6). The decrease in strength was not very significant up to 400°C. At 600°C softening of the glass matrix occurs. Fracture surfaces of AG and ASG composites at 400°C (Fig. 9a,b) showed that as temperature increased, the effect of the various toughening mechanisms decreased namely, crack deflection, fiber bridging, and crack front debonding of the fiber. Post fabrication residual stresses arising due to thermal expansion mismatch between fiber and the matrix govern the bonding at the fiber matrix interface. Using a two-element (fiber and matrix) and three-element model (fiber, coating and interface) thermal stresses evaluated on both AG and ASG composites showed Table 4 Work of fracture of AG and ASG composites in L, T1 and T2 directions Vf (%) Work of fracture (J/m2 ) L T1 T2 AG 12 100 120 220 26 150 420 175 42 130 770 150 ASG 24 580 175 180 36 900 220 250 essentially the same in directions T1 and T2. The strength of AG composites increased up to a volume fraction of fibers of 0.25 and then decreased. Fig. 8a
100 po 50 10 50m Fig. 8. Fracture Fig. 7. Fracture surface of Saphikon fiber/SnO,/glass composite(a) (a) Note fiber breakage.(b)Porosity in high volume fraction com- Neat fiber pullout. (b) Debonding at the fiber/coating interface. posites. hat the radial thermal stress in both the composites Jero and Kerans and modeled the effect of fiber rough- were tensile in nature and had a maximum value of 109 ness on the fracture behavior of CMCs [26-31].Kerans and 50 MPa at the fiber/SnO2 interfaces in PRD-166/ and Parthasarathy (26, 31] have described the contribu- SnO2 glass and Saphikon/SnO2/glass composites,re- tion of interfacial roughness with relative fiber/matrix pectively [24]. This radial stress combined with weak displacement during debonding at the fiber-matrix in- interfacial bonding between alumina and SnO2 are de sirable attributes from a toughness point of view [ll Table 5 Once we eliminate interfacial chemical bonding in a Chevron notch fracture toughness of AG and ASG composites CMC, interfacial roughness assumes an important role Fiber volume (% Kl(MPa m) Work of fracture (/m) in controlling the fracture of the CMC. Recent studies on the importance of fiber roughness have evolved since AG the first paper by Jero and Kerans [25]. They reported the importance of fiber roughness using fiber pushout experiments on thin slices of Sic monofilaments in a borosilicate glass matrix composite in which residual 24 compressive stresses existed perpendicular to the inter- 36 3.3 face. Other researchers have confirmed the results of
52 R. Venkatesh Venkatesh / Materials Science and Engineering A Materials Science and Engineering A268 (1999) 47–54 268 (1999) 47–54 Fig. 7. Fracture surface of Saphikon fiber/SnO2/glass composite. (a) Neat fiber pullout. (b) Debonding at the fiber/coating interface. Fig. 8. Fracture surface of transverse tested alumina/glass composites. (a) Note fiber breakage. (b) Porosity in high volume fraction composites. Jero and Kerans and modeled the effect of fiber roughness on the fracture behavior of CMCs [26–31]. Kerans and Parthasarathy [26,31] have described the contribution of interfacial roughness with relative fiber/matrix displacement during debonding at the fiber–matrix inthat the radial thermal stress in both the composites were tensile in nature and had a maximum value of 109 and 50 MPa at the fiber/SnO2 interfaces in PRD-166/ SnO2/glass and Saphikon/SnO2/glass composites, respectively [24]. This radial stress combined with weak interfacial bonding between alumina and SnO2 are desirable attributes from a toughness point of view [11]. Once we eliminate interfacial chemical bonding in a CMC, interfacial roughness assumes an important role in controlling the fracture of the CMC. Recent studies on the importance of fiber roughness have evolved since the first paper by Jero and Kerans [25]. They reported the importance of fiber roughness using fiber pushout experiments on thin slices of SiC monofilaments in a borosilicate glass matrix composite in which residual compressive stresses existed perpendicular to the interface. Other researchers have confirmed the results of Table 5 Chevron notch fracture toughness of AG and ASG composites K1c (MPa m 1 2 ) Work of fracture (J/m2 Fiber volume (%) ) AG 16 2 50 26 2.3 120 36 2.6 150 ASG 24 160 2.8 36 3.3 220
Bend strength of AG and ASG composites versus temperature M Temperature(°Cv(%) ASG 222 Ⅴ iscous flow Viscous flow terface. They have described three regions namely an unslipped region ahead of matrix crack tip(I), a com- pletely debonded region (III)and a region (II) which m extends with increasing misfit strain from end of region I to the beginning of region II. The effective normal stress, on at the interface in region ill due to roughness and residual stresses has been given as [26]. Fig 10. Morphology of interfaces in(a)alumina/glass(b) Saphikon 20 on=i-qEmEJEdl+vm)+ Em(I-vollAoAT+ A/r] where Ao is the mismatch between the thermal expan- sion coefficients of the fiber and matrix. at is the temperature change during cooling, q is an adjustable parameter(equal to l for an infinite matrix), Em, Im, Er and vr are the Youngs modulus and Poisson ratio of the matrix and fiber respectively, A is the amplitude of f the fiber The amplitude of fiber roughness have been experi mentally determined from fiber surface profiles using atomic force microscopy [29-32], profilometry [33], and terference [34]. In the pre ase. sem of longitudinal section of the fibers has been used [24] 20 gm Roughness parameter and thermal stress contribution PRD-166/SnO, interfa Saphikon/SnO, interfac Fig 9. High temperature (400 C)tested composites (a)Alumina/ AxAT 0.0013 glass.(b) Alumina/SnO2/glass
R. Venkatesh Venkatesh / Materials Science and Engineering A Materials Science and Engineering A268 (1999) 47–54 268 (1999) 47–54 53 Table 6 Bend strength of AG and ASG composites versus temperature Temperature (°C) ASG Vf (%) AG 25 200 42 230 200 42 180 200 400 175 150 42 600 42 Viscous flow Viscous flow Fig. 10. Morphology of interfaces in (a) alumina/glass; (b) Saphikon/ SnO2/glass. terface. They have described three regions namely an unslipped region ahead of matrix crack tip (I), a completely debonded region (III) and a region (II) which extends with increasing misfit strain from end of region I to the beginning of region II. The effective normal stress, sn at the interface in region III due to roughness and residual stresses has been given as [26], Fig. 9. High temperature (400°C) tested composites. (a) Alumina/ glass. (b) Alumina/SnO2/glass. sn={−qEmEf /[Ef (1+nm)+Em(1−nf )]}[DaDT+A/r] (4) where Da is the mismatch between the thermal expansion coefficients of the fiber and matrix, DT is the temperature change during cooling, q is an adjustable parameter (equal to 1 for an infinite matrix), Em, nm, Ef and nf are the Young’s modulus and Poisson ratio of the matrix and fiber respectively, A is the amplitude of roughness and r is the radius of the fiber. The amplitude of fiber roughness have been experimentally determined from fiber surface profiles using atomic force microscopy [29–32], profilometry [33], and optical interference [34]. In the present case, SEM of longitudinal section of the fibers has been used [24]. Table 7 Roughness parameter and thermal stress contributions PRD-166/SnO2 interface Saphikon/SnO2 interface A/r 0.026 0.003 DaDT 0.0013 0.001
R. Venkatesh/ Materials Science and Engineering 4268(1999)47-54 Evaluating A/r and the residual thermal strain, AxAT Rockwell Science Center for providing the Saphikon at the PRD-166/ SnO2 and Saphikon /SnO2 interfac fibers (Fig. 10a, b)and using Eq(4), it was observed that the A/r value of PRD-166/SnO, interface was about nine times that of the Saphikon/SnO, interface, (Table 7). It References was also found that, in PRD-166/SnO2/glass matrix composites, the compressive radial strain induced due [K.K. Chawla, Ceramic Matrix Composites, Chapman Hall o fiber roughness was about 20 times greater than the London. 1993. tensile thermal radial strain. In Saphikon/SnO,/glass [2R.W.Rice, Ceramic matrix composite toughening mechanisms: An date. in: eramic Science and Engineering Series. Vol. 6 matrix composites, the roughness induced compressive strain was only three times greater than the tensile 3A.G. Evans, J. Am. Cer. Soc. 73(1990)187 thermal strain. This indicates the strong mechanical 4 W.B. Hillig, Annu. Rev. Mater. Sci. 17(1987)341 clamping due to fiber roughness at the fiber/SnO, inter- 5 D B. Marshall, J.E. Ritter, Cer. Bull. 66(1987)309 6K.K. Chawla, Z.R. Xu, R. Venkatesh, J.S. Ha, Interface Engineer face in ASG composites. As fiber roughness increased ing in some oxide/oxide composites, in: 9th International Confer increased shear stress transfer at the interface beyond ence on Composites (ICCM-9), July 12, 1993, Madrid, Spai matrix cracking from fiber to matrix causes a reduction [7 P.E. D Morgan, D B. Marshall,J. Am. Cer. Soc. 78(6)(1995)1553. he debond length. i.e. fibers broke rather than [8 A Maheshwari, KK. Chawla, T.A. Michalske, Mater. Sci. Eng debonded as the matrix crack grew, resulting in a A107(1989)269 9 M. Cinibulk, Cer. Eng. Sci. Proc. 15(5)(1995)72 composite fracture surface in ASG with little or no fiber (10 M. Cinibulk, Cer. Eng. Sci. Proc. 16(5)(1995) pullout on the fracture surface. Similar results have [l K.K. Chawla, A. Choudhary, R. Venkatesh, J.R. Hellmann been obtained by Chawla et al. using atomic force Mater.Char.31(3)(1993)167 microscopy [35 [12 P.E. D. Morgan, R. M. Housley, J. Am. Cer. Soc. 789(1)(1995) 13J.C. Romine, Cer. Eng. Sci. Proc. 8(1987)755. [14 Data Sheet, Saphikon, NH. 4. Conclusion [15 K.M. Prewo, J.J. Brennan, G.K. Layden, Am Cer. Soc. Bull. 65 It has now been realized that fiber surface roughness [16 J.I. Bluhm, Eng. Fracture Mech. 9(1975)593 [17 Wu Shang-Xian, Eng. Fracture Mech. 19(1984)221 and residual stresses in the composite after processing [18] S.N. Patankar, R Venkatesh, KK Chawla, Scripta Metall. Mate play a very important role in governing the strength nd toughness of CMCs As shown in the present work, [9J. Cook, J.E. Gordon, Proc. R. Soc. Lond. A282(1964)508. composites with tensile radial stress at the interface and 20A. G. Evans, M.Y. Ye, J W.Hutchinson, J Am. Cer. Soc. 72(1989) 2300 ery rough fiber surface exhibit very small debond 21]CC. Wu, S.W. Freiman, R.W. Rice, J.J. Mecholsky, J Mater. Sci. length and fail in a brittle manner. Alternatively in 3(1978)2659 composites with tensile radial stress at the interface and [22] R. Venkatesh, k.K. Chawla, J Mater. Sci. Lett. Il(1992) very smooth interface, the fiber/matrix bonding may be 23] T.A. Michalske, J.R. Hellmann, J. Am. Cer. Soc. 71(199 weak enough to compromise the strength of the 224 K.K. Chawla, M. Ferber, Z.R. Xu, R Venkatesh, Mater. Sci Eng A162(1993)35 225] P D. Jero, RJ. Kerans, Scripta Metall. Mater. 24(1991)2315 Another important effect of fiber roughness is abra- [26]RJ Kerans,TAParthasarathy, J.Am. Cer Soc. 74(1991)1585 sion at the fiber/matrix interface. Abrasion is especially w.C. Carter, E.P. Butler, E.R. Fuller Jr, Scripta Metall. Mater important in evaluating cyclic properties of If it occurs abrasion can decrease the interfacial fric [28] T.J. Mackin, P D. Warren, A.G. Evans, Acta Metall. 40(1992) 1251 tional stress and affect the interfacial properties of [29) P D Jero,RJ Kerans, T.A. Parthasarathy, J. Am. Cer. Soc. 7 CMCS. a judicious selection of material systems and (11)(1991)2793 processing can control the interfacial roughness and 30] T.J. Mackin, J. Yang, P D. Warren, J Am. Cer. Soc. 756(12)(1992) thermal stresses that can be utilized in developing tough CMCs B1]TA Parthasarathy, R. Kerans, J. Am. Cer. Soc. 80( 8)(1992) 32 H P. Wang, T.J. Nelson, C L Lim, w.w. Gerberch, J Mater Res 9(2)(1994)498 Acknowledgements 3]P D Jero, T.A. Parthasarathy, Interface properties: their ment with fiber pushout tests. in: R. Naslain(Ed. ) High This work was done at New Mexico Tech. Socorro NM and supported by the US Office of Naval Research [34] B.F. Sorensen, Scripta Metall. Mater. 28(1993)435 (contract no. N00014-89-J-1459). The author would like 35KK. Chawla, Z.R. Xu, A. Hlinak, Y.w. Chung Characteristics to thank Dr K.K. Chawla for his support. Thanks are of three alumina fibers by atomic force opy, in: N P Bansal also due to Dr S. Balsone of Wright-Patterson Air (Ed ) Advances in Ceramic Matrix Composites, Ceramic Trans- actions, vol 38, American Ceramic Society, Columbus, OH, 1993 Force Base and Drs a.h. muir. Jr.. and j. porter of the
54 R. Venkatesh Venkatesh / Materials Science and Engineering A Materials Science and Engineering A268 (1999) 47–54 268 (1999) 47–54 Evaluating A/r and the residual thermal strain, DaDT at the PRD-166/SnO2 and Saphikon/SnO2 interfaces, (Fig. 10a,b) and using Eq. (4), it was observed that the A/r value of PRD-166/SnO2 interface was about nine times that of the Saphikon/SnO2 interface, (Table 7). It was also found that, in PRD-166/SnO2/glass matrix composites, the compressive radial strain induced due to fiber roughness was about 20 times greater than the tensile thermal radial strain. In Saphikon/SnO2/glass matrix composites, the roughness induced compressive strain was only three times greater than the tensile thermal strain. This indicates the strong mechanical clamping due to fiber roughness at the fiber/SnO2 interface in ASG composites. As fiber roughness increased, increased shear stress transfer at the interface beyond matrix cracking from fiber to matrix causes a reduction in the debond length, i.e. fibers broke rather than debonded as the matrix crack grew, resulting in a composite fracture surface in ASG with little or no fiber pullout on the fracture surface. Similar results have been obtained by Chawla et al. using atomic force microscopy [35]. 4. Conclusion It has now been realized that fiber surface roughness and residual stresses in the composite after processing play a very important role in governing the strength and toughness of CMCs. As shown in the present work, composites with tensile radial stress at the interface and very rough fiber surface exhibit very small debond length and fail in a brittle manner. Alternatively in composites with tensile radial stress at the interface and very smooth interface, the fiber/matrix bonding may be weak enough to compromise the strength of the composite. Another important effect of fiber roughness is abrasion at the fiber/matrix interface. Abrasion is especially important in evaluating cyclic properties of composites. If it occurs, abrasion can decrease the interfacial frictional stress and affect the interfacial properties of CMCs. A judicious selection of material systems and processing can control the interfacial roughness and thermal stresses that can be utilized in developing tough CMCs. Acknowledgements This work was done at New Mexico Tech, Socorro, NM and supported by the US Office of Naval Research (contract no. N00014-89-J-1459). The author would like to thank Dr K.K. Chawla for his support. Thanks are also due to Dr S. Balsone of Wright-Patterson Air Force Base and Drs A.H. Muir, Jr., and J. Porter of the Rockwell Science Center for providing the Saphikon fibers. References [1] K.K. Chawla, Ceramic Matrix Composites, Chapman & Hall, London, 1993. [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, J. Am. Cer. Soc. 73 (1990) 187. [4] W.B. Hillig, Annu. Rev. Mater. Sci. 17 (1987) 341. [5] D.B. Marshall, J.E. Ritter, Cer. Bull. 66 (1987) 309. [6] K.K. Chawla, Z.R. Xu, R. Venkatesh, J.S. Ha, Interface Engineering in some oxide/oxide composites, in: 9th International Conference on Composites (ICCM-9), July 12, 1993, Madrid, Spain. [7] P.E.D. Morgan, D.B. Marshall, J. Am. Cer. Soc. 78 (6) (1995) 1553. [8] A. Maheshwari, K.K. Chawla, T.A. Michalske, Mater. Sci. Eng. A107 (1989) 269. [9] M. Cinibulk, Cer. Eng. Sci. Proc. 15 (5) (1995) 721. [10] M. Cinibulk, Cer. Eng. Sci. Proc. 16(5) (1995). [11] K.K. Chawla, A. Choudhary, R. Venkatesh, J.R. Hellmann, Mater. Char. 31 (3) (1993) 167. [12] P.E.D. Morgan, R.M. Housley, J. Am. Cer. Soc. 789 (1) (1995) 263. [13] J.C. Romine, Cer. Eng. Sci. Proc. 8 (1987) 755. [14] Data Sheet, Saphikon, NH. [15] K.M. Prewo, J.J. Brennan, G.K. Layden, Am Cer. Soc. Bull. 65 (1980) 305. [16] J.I. Bluhm, Eng. Fracture Mech. 9 (1975) 593. [17] Wu Shang-Xian, Eng. Fracture Mech. 19 (1984) 221. [18] S.N. Patankar, R. Venkatesh, K.K. Chawla, Scripta Metall. Mater. 25 (1991) 361. [19] J. Cook, J.E. Gordon, Proc. R. Soc. Lond. A282 (1964) 508. [20] A.G. Evans, M.Y. Ye, J.W. Hutchinson, J. Am. Cer. Soc. 72 (1989) 2300. [21] C.C. Wu, S.W. Freiman, R.W. Rice, J.J. Mecholsky, J. Mater. Sci. 13 (1978) 2659. [22] R. Venkatesh, K.K. Chawla, J. Mater. Sci. Lett. 11 (1992) 650. [23] T.A. Michalske, J.R. Hellmann, J. Am. Cer. Soc. 71 (1998) 725. [24] K.K. Chawla, M. Ferber, Z.R. Xu, R. Venkatesh, Mater. Sci. Eng. A162 (1993) 35. [25] P.D. Jero, R.J. Kerans, Scripta Metall. Mater. 24 (1991) 2315. [26] R.J. Kerans, T.A. Parthasarathy, J. Am. Cer. Soc. 74 (1991) 1585. [27] W.C. Carter, E.P. Butler, E.R. Fuller Jr., Scripta Metall. Mater. 25 (1991) 579. [28] T.J. Mackin, P.D. Warren, A.G. Evans, Acta. Metall. 40 (1992) 1251. [29] P.D. Jero, R.J. Kerans, T.A. Parthasarathy, J. Am. Cer. Soc. 74 (11) (1991) 2793. [30] T.J. Mackin, J. Yang, P.D. Warren, J. Am. Cer. Soc. 756 (12) (1992) 3358. [31] T.A. Parthasarathy, R.J. Kerans, J. Am. Cer. Soc. 80 (8) (1992) 2043. [32] H.P. Wang, T.J. Nelson, C.L. Lim, W.W. Gerberch, J. Mater. Res. 9 (2) (1994) 498. [33] P.D. Jero, T.A. Parthasarathy, Interface properties: their measurement with fiber pushout tests, in: R. Naslain (Ed.), High Temperature Ceramic Matrix Composites, Goodhead, Cambridge, 1993, p. 401. [34] B.F. Sorensen, Scripta Metall. Mater. 28 (1993) 435. [35] K.K. Chawla, Z.R. Xu, A. Hlinak, Y.W. Chung, Characteristics of three alumina fibers by atomic force microscopy, in: N.P. Bansal (Ed.), Advances in Ceramic Matrix Composites, Ceramic Transactions, vol. 38, American Ceramic Society, Columbus, OH, 1993, p. 725