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alumina composites
Scripta Materialia, Vol. 41, No. 7, pp. 749 –754, 1999 Elsevier Science Ltd Copyright © 1999 Acta Metallurgica Inc. Printed in the USA. All rights reserved. 1359-6462/99/$–see front matter
Pergamon
PII S1359-6462(99)00212-2
FRACTURE BEHAVIOR OF CENTRIFUGALLY CAST MULTILAYER ALUMINA/ALUMINA COMPOSITES E. N. Drewry, R. J. Moon, K. J. Bowman, and K. P. Trumble, School of Materials Engineering, Purdue University, West Lafayette, Indiana 47907-1289 USA
J. Brehm* *Department of Materials Science, University of Technology, Darmstadt D-64287 Darmstadt, Germany (Received June 8, 1999) (Accepted June 20, 1999) Keywords: Brittle fracture and fracture toughness; Centrifugation, Composites, Multilayers Introduction Increased fracture toughness in brittle materials can be achieved by incorporating weak (porous) interfaces (1,2). These improvements have been attributed to energy absorption through crack deflection and frictional bridging along the weaker interfaces (1). While improved fracture properties are attractive for thermal and structure ceramic components, difficulties in producing a specific degree of interfacial porosity for a wide range of ceramics have limited application of this approach. Studies investigating the influence of weak interlayers in brittle composites (1–10) have examined systems ranging from tape cast multilayered composites of alternating fully dense and porous interlayers (7) to brittle plates bonded with weaker thermoplastic adhesives (8 –9). During flexural loading, cracks typically deflect along the weaker interfaces, giving rise to a step-wise drop in the loaddisplacement behavior rather than a catastrophic failure typical of non-deflecting brittle materials. Residual stress distributions within layered ceramic composites can also produce similar fracture behavior, where alternating layers of tensile and compressive stress states due to thermal expansion mismatch can promote crack bifurcation (10). Advances in centrifugal consolidation of colloidal suspensions (11–14) have made the technique an attractive alternative for producing layered composites. Tailored microstructures can be achieved by incorporating suspensions that are either flocculated (11) to inhibit particle segregation or dispersed (12–14) to allow segregation due to differences in particle size and density. This paper summarizes the initial results on the fracture behavior of alumina and alumina/agglomerated alumina multilayered composites produced by centrifugal consolidation. Experimental Procedure Four aqueous alumina suspensions (Table 1), each consisting of 15 vol% solids loading, were used to centrifugally cast the composite tiles (Table 2). The alumina powders were mixed in deionized water 749
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TABLE 1 Suspension Compositions Suspension A Suspension B
A16SG*——(0.4 m average grain size) A16SG* Tabular——(30–80 m diameter hard agglomerates) A16SG* Tabular A16SG* A10UNG*——(30–100 m diameter hard agglomerates)
15 vol% 14.7 vol% 0.3 vol% 14.4 vol% 0.6 vol% 14.4 vol% 0.6 vol%
Suspension C Suspension D
* ALCOA, Bauxite, AR
with 0.5 vol% Darvan 821A (R.T. Vanderbilt Company Inc., Norwalk, CT) and further dispersed by ultrasonicating (Model W-380, Heat Systems-Ultrasonics, Farmingdale, NY) for 20 minutes in pulse mode. A cobinder system (15) consisting of 2.6 vol% hydroxyethyl cellulose and 1.1 vol% poly(ethylene glycol) was then added. After the cobinder addition, 0.1 vol% anti-foaming agent (Dow 2210) was added, and the suspensions were simultaneously stirred and ultrasonicated for an additional 10 minutes. Sequential centrifugation of the suspensions was used to produce the multilayered tiles. Each layer was produced by adding a constant volume of suspension, either 5, 7.5, or 10 ml, to the 76 mm diameter casting containers. The containers were placed in the centrifuge rotor (Model RC 3C Plus, Sorvall, Newtown, CT) and centrifuged for ten minutes at 2000 RPM, which corresponds to an acceleration approximately 1000 times that of standard gravity. During this cycle, the alumina particles settle to the bottom, and the remaining water was then decanted. This process was repeated 15 to 30 times to produce tiles with the desired thickness. After final decanting, the casting containers were placed directly into a drying oven at 60°C for ⬃1 h or until the compact began to pull away from the walls of the container. The Plaster of Paris base and the alumina compact were then removed from the casting container and returned to the drying oven for at least 6 h. Casting integrity during drying and handling was greatly enhanced by the additions of the cobinder system (16). In experiments without cobinder additions, cracking typically began shortly after the final decanting, even when the tiles were dried over week-long periods in a humidity-controlled environment. TABLE 2 Apparent SENB Fracture Toughness of Multilayered Alumina Composites
Tile Tile Tile Tile Tile Tile Tile Tile Tile Tile
1 2 3 4 5 6 7 8 9 10
Suspension
Layer Thickness (m)
KIC Orientation 1
KIC Orientation 2
A B C D C C C D D D
500 500 500 500 400 500 800 400 500 800
5.6 (0.4)* 4.8 (0.4)* 8.0 (0.4)* 6.5 (0.3)* — — — — — —
5.1 (0.4)* 4.2 (0.2)* 9.0 (0.4)* 8.0 (0.8)* 4.3 (0.3)** 5.0 (0.6)** 7.6 (0.6)** 4.1 (0.3)** 4.6 (0.2)** 7.0 (1.0)**
KIC in MPa䡠m1/2; *3 specimens tested; **8 to 10 specimens tested.
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The tiles were fired at 1600°C for 4 h in air. The initial ramp rate was set at 200°C/h, which allowed binder pyrolysis, and the ramp down was set at 300°C/h. The percent of theoretical density for Tiles 1, 2, 3, and 4 were 98, 95, 93, and 91, respectively. The surfaces of the fired multilayered tiles were initially ground flat and parallel on a surface grinder with a 240 grit diamond wheel. Test specimens having the dimensions of 40 ⫻ 7.0 ⫻ 3.5 mm were then cut from the tiles with a high speed diamond saw. The test specimens were ground and polished using diamond abrasives on an automatic polisher (Spectrum 2000, Leco Corporation, St. Joseph, MI) down to 0.05 m colloidal silica finish, after which a notch was cut in the center of each test specimen with a 150 m wide diamond blade. The identical procedure was used to prepare control specimens of a commercial monolithic alumina (Friatec-Degussit A1-23, nominal grain size ⫽ 10 m). Fracture toughness was measured by the single edge notched beam (SENB) method following the DIN NMP 51 109 standard in four-point bending (17). Test specimens were loaded at a rate of 0.1 mm/min in an MTS Sintech 30/D machine (MTS System Corporation, Eden Prairie, MN) equipped with a fully articulating four-point bend fixture (Model CU-FL-34, Wyoming Test Fixtures Inc., Laramie, WY). Inner and outer loading spans were 18 mm and 36 mm, respectively. Load vs. displacement data was acquired using MTS TestPad v1.0 software. Fracture toughness calculations were based on the peak load required to propagate the machined notch. The increased sample cross-sections (40 ⫻ 7 ⫻ 3.5 mm from the specified 45 ⫻ 4 ⫻ 3 mm) for multilayered samples was used to increase the number of layers that a crack would have to propagate through before final fracture. The influence of the dimensional change on KIC was investigated by testing commercial monolithic alumina control samples for both geometries. Six specimens from each geometry had a resultant KIC of 4.3 (0.4) MPa䡠m1/2, and due to this constancy in KIC, the increased sample cross-section was incorporated for the remaining experiments. The KIC of the commercial alumina control samples was higher than the reported ⬃3 MPa䡠m1/2 of alumina previously tested by this method (18,19). The ⬃180 m notch-width used in this study likely resulted in artificially high KIC values (18 –22), where notch widths less than 100 m are necessary to minimize notch-width influence on KIC in alumina (18,19). The KIC values of the multilayered composites may also be artificially high, however, comparisons between the resulting KIC from different layer microstructures provides the relative KIC associated with a particular layer microstructure.
Results and Discussion Samples produced by the centrifugation process have two characteristic macroscopic features, layer curvature and tilt, which are inherent in the centrifugation process (23). Particles are deposited having curvature similar to the rotation radius based on radial force equilibrium and acceleration. Additionally, particle size segregation occurs within the individual layers as shown in Figure 1. The layer interfaces in Tile 3 (Fig. 1b) were more distinct, as compared to Tile 1 (Fig. 1a), due to the large fraction of porosity resulting from the segregation of the hard alumina agglomerates during centrifugation. The fracture results for the layered alumina specimens are given in Table 2. Tiles 1, 2, 5, 6, 8, and 9 failed catastrophically as is typical for an alumina monolith, however, large crack deflections and continued loading after the peak load were observed for Tiles 3, 4, 7, and 10. The differences in fracture behavior between Tiles 1, 2, and 3 resulted from the increased porosity concentration caused by the increased alumina agglomerate additions. The “weaker” bonding at layer-layer interfaces allowed crack deflection along the interfaces. Additionally, Tiles 3 and 4 (weaker interfaces) exhibited differences in fracture toughness depending on whether the crack propagated away (orientation 1) or towards (orientation 2) the center of curvature
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Figure 1. Microstructures of a given layer within (a) Tile 1, and (b) Tile 3. The increased porosity at the layer interface in Tile 3 was associated with the segregation of the large alumina agglomerates during centrifugation.
of the layers (Figure 2 and Table 2). Specimens tested in orientation 1 typically exhibited crack deflection less than 1 mm before kinking out of the interface to ultimate failure. Specimens loaded with notch orientation 2 exhibited crack deflection with continued crack growth along the first adjoining interface until reaching the limit of the inner loading span. Subsequently, the crack propagated normal to the interfaces until ultimate failure occurred. The increased fracture toughness for orientation 2 is believed to result from the crack propagating closer to Mode I where the layer curvature facilitates deviation of the crack into the interface. In contrast, for Tiles 1 and 2 (stronger interfaces), layer curvature had little effect on the fracture toughness results since large scale crack deflection did not occur. Tiles 5, 6, 7 and Tiles 8, 9, 10 were used to investigate the influence of layer thickness (400, 500, and 800 m) on KIC and the results are plotted in Figure 3. The increase in KIC with thicker layers corresponded to an increase in the degree of crack deflection at layer interfaces. The thicker layers obtained by injecting a larger slurry volume, produced thicker porous interlayers resulting from the larger amounts of alumina agglomerates which settled preferentially at the layer interface during centrifugation. The differences in KIC between Tiles 3 and 4 to that of Tiles 6 and 9, each having 500 m layer thickness, resulted from differences in porosity at the layer interface. The segregation of porosity at the layer interface for Tiles 3 and 4 was similar to Tiles 7 and 10 respectively and resulted in similar KIC values. Slight changes in the suspension preparation and insertion into the casting containers may account for the change in porosity and thus the variation of KIC.
Figure 2. (a) Crack propagation behavior observed in Tile 3 showing the influence of the notch orientation where, (b) shows the notch orientations 1 and 2 with respects to the layer curvature.
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Figure 3. The increasing KIC with thicker layers was associated with a larger segregation of porosity at the layer interface allowing for crack deflection to occur.
Summary Multilayered alumina and alumina/agglomerated alumina composites were produced by centrifugal consolidation. The extent of porosity at layer interfaces was varied by controlling the volume fraction of agglomerated additions to the suspension and by controlling layer thickness. Porosity at the layer interfaces decreases the interface bond strength, allowing cracks to kink and propagate along layer interfaces (parallel to layers). The likelihood of crack deflection along layer interfaces increased as the extent of porosity at the layer interfaces increased. The resulting KIC of multilayered composites increased as the amount of crack deflection increased where the apparent KIC transverse to the layers was nearly twice that of a commercial monolithic alumina. Additionally, multilayered composites without agglomerated alumina additions demonstrated fracture behavior and KIC values similar to commercial monolithic alumina. Acknowledgement This work has been supported by the Army Research Office: MURI grant no. DAAH04-96-1-0331. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.
A. J. Phillipps, W. J. Clegg, and T. W. Clyne, Composites. 25, 524 (1994). W. Lee, S. J. Howard, and W. J. Clegg, Acta Mater. 44, 3905 (1996). H. C. Cao and A. G. Evans, Mech. Mater. 7, 295 (1989). A. G. Evans, M. Ruehle, B. J. Dalgleish, and P. G. Charalambides, Metall. Trans. A. 21, 2419 (1990). A. A. Mammoli, A. L. Graham, I. E. Reimanis, and D. L. Tullock, Acta Metall. Mater. 43, 1149 (1995). C. M. Warwick and W. J. Clegg, Ceram. Eng. Sci. Proc. 72, 847 (1995). W. J. Clegg, K. S. Blanks, J. B. Davis, and F. Lanckmans, Key Eng. Mater. 132–136, 1866 (1997). C. Folsom, F. W. Zok, and F. Lange, J. Am. Ceram. Soc. 77, 689 (1994). C. Folsom, F. W. Zok, and F. Lange, J. Am. Ceram. Soc. 77, 2081 (1994). M. Oechsner, C. Hillman, and F. F. Lange, J. Am. Ceram. Soc. 79, 1834 (1996). J. C. Chang, B. V. Velamakanni, F. F. Lange, and D. S. Pearson, J. Am. Ceram. Soc. 74, 2201 (1991). W. Huisman, T. Graule, and L. J. Gauckler, J. Eur. Ceram. Soc. 13, 33 (1994). W. Huisman, T. Graule, and L. J. Gauckler, J. Eur. Ceram. Soc. 15, 811 (1995). R. K. Roeder, K. P. Trumble, and K. J. Bowman, J. Am. Ceram. Soc. 80, 27 (1997).
754
15. 16. 17. 18. 19. 20. 21. 22. 23.
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A. S. Barnes, J. S. Reed, and E. M. Anderson, Am. Ceram. Soc. Bull. 76, 77 (1997). E. Drewry, M. S. Thesis, Purdue University (August 1998). DIN NMP 51 109, German Standard for Fracture Toughness Analysis in Four-Point Flexure. R. J. Primas and R. Gstrein, Fatigue Fract. Eng. Mater. Struct. 20, 513 (1997). R. J. Damani, Key Eng. Mater. 132–136, 484 (1997). R. Damani, R. Gstrein, and R. Danzer, J. Eur. Ceram. Soc. 16, 695 (1996). J. Ku¨bler, Ceram. Eng. Sci. Proc. 18, 155 (1997). T. Nishida, Y. Hanaki, and G. Pezzotti, J. Am. Ceram. Soc. 77, 606 (1994). R. Moon, E. Drewry, K. Bowman, and K. Trumble, in 5th FGM Conference Proceedings, Dresden, Germany, Trans. Tech. Publications.
Pergamon
PII S1359-6462(99)00212-2
FRACTURE BEHAVIOR OF CENTRIFUGALLY CAST MULTILAYER ALUMINA/ALUMINA COMPOSITES E. N. Drewry, R. J. Moon, K. J. Bowman, and K. P. Trumble, School of Materials Engineering, Purdue University, West Lafayette, Indiana 47907-1289 USA
J. Brehm* *Department of Materials Science, University of Technology, Darmstadt D-64287 Darmstadt, Germany (Received June 8, 1999) (Accepted June 20, 1999) Keywords: Brittle fracture and fracture toughness; Centrifugation, Composites, Multilayers Introduction Increased fracture toughness in brittle materials can be achieved by incorporating weak (porous) interfaces (1,2). These improvements have been attributed to energy absorption through crack deflection and frictional bridging along the weaker interfaces (1). While improved fracture properties are attractive for thermal and structure ceramic components, difficulties in producing a specific degree of interfacial porosity for a wide range of ceramics have limited application of this approach. Studies investigating the influence of weak interlayers in brittle composites (1–10) have examined systems ranging from tape cast multilayered composites of alternating fully dense and porous interlayers (7) to brittle plates bonded with weaker thermoplastic adhesives (8 –9). During flexural loading, cracks typically deflect along the weaker interfaces, giving rise to a step-wise drop in the loaddisplacement behavior rather than a catastrophic failure typical of non-deflecting brittle materials. Residual stress distributions within layered ceramic composites can also produce similar fracture behavior, where alternating layers of tensile and compressive stress states due to thermal expansion mismatch can promote crack bifurcation (10). Advances in centrifugal consolidation of colloidal suspensions (11–14) have made the technique an attractive alternative for producing layered composites. Tailored microstructures can be achieved by incorporating suspensions that are either flocculated (11) to inhibit particle segregation or dispersed (12–14) to allow segregation due to differences in particle size and density. This paper summarizes the initial results on the fracture behavior of alumina and alumina/agglomerated alumina multilayered composites produced by centrifugal consolidation. Experimental Procedure Four aqueous alumina suspensions (Table 1), each consisting of 15 vol% solids loading, were used to centrifugally cast the composite tiles (Table 2). The alumina powders were mixed in deionized water 749
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TABLE 1 Suspension Compositions Suspension A Suspension B
A16SG*——(0.4 m average grain size) A16SG* Tabular——(30–80 m diameter hard agglomerates) A16SG* Tabular A16SG* A10UNG*——(30–100 m diameter hard agglomerates)
15 vol% 14.7 vol% 0.3 vol% 14.4 vol% 0.6 vol% 14.4 vol% 0.6 vol%
Suspension C Suspension D
* ALCOA, Bauxite, AR
with 0.5 vol% Darvan 821A (R.T. Vanderbilt Company Inc., Norwalk, CT) and further dispersed by ultrasonicating (Model W-380, Heat Systems-Ultrasonics, Farmingdale, NY) for 20 minutes in pulse mode. A cobinder system (15) consisting of 2.6 vol% hydroxyethyl cellulose and 1.1 vol% poly(ethylene glycol) was then added. After the cobinder addition, 0.1 vol% anti-foaming agent (Dow 2210) was added, and the suspensions were simultaneously stirred and ultrasonicated for an additional 10 minutes. Sequential centrifugation of the suspensions was used to produce the multilayered tiles. Each layer was produced by adding a constant volume of suspension, either 5, 7.5, or 10 ml, to the 76 mm diameter casting containers. The containers were placed in the centrifuge rotor (Model RC 3C Plus, Sorvall, Newtown, CT) and centrifuged for ten minutes at 2000 RPM, which corresponds to an acceleration approximately 1000 times that of standard gravity. During this cycle, the alumina particles settle to the bottom, and the remaining water was then decanted. This process was repeated 15 to 30 times to produce tiles with the desired thickness. After final decanting, the casting containers were placed directly into a drying oven at 60°C for ⬃1 h or until the compact began to pull away from the walls of the container. The Plaster of Paris base and the alumina compact were then removed from the casting container and returned to the drying oven for at least 6 h. Casting integrity during drying and handling was greatly enhanced by the additions of the cobinder system (16). In experiments without cobinder additions, cracking typically began shortly after the final decanting, even when the tiles were dried over week-long periods in a humidity-controlled environment. TABLE 2 Apparent SENB Fracture Toughness of Multilayered Alumina Composites
Tile Tile Tile Tile Tile Tile Tile Tile Tile Tile
1 2 3 4 5 6 7 8 9 10
Suspension
Layer Thickness (m)
KIC Orientation 1
KIC Orientation 2
A B C D C C C D D D
500 500 500 500 400 500 800 400 500 800
5.6 (0.4)* 4.8 (0.4)* 8.0 (0.4)* 6.5 (0.3)* — — — — — —
5.1 (0.4)* 4.2 (0.2)* 9.0 (0.4)* 8.0 (0.8)* 4.3 (0.3)** 5.0 (0.6)** 7.6 (0.6)** 4.1 (0.3)** 4.6 (0.2)** 7.0 (1.0)**
KIC in MPa䡠m1/2; *3 specimens tested; **8 to 10 specimens tested.
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The tiles were fired at 1600°C for 4 h in air. The initial ramp rate was set at 200°C/h, which allowed binder pyrolysis, and the ramp down was set at 300°C/h. The percent of theoretical density for Tiles 1, 2, 3, and 4 were 98, 95, 93, and 91, respectively. The surfaces of the fired multilayered tiles were initially ground flat and parallel on a surface grinder with a 240 grit diamond wheel. Test specimens having the dimensions of 40 ⫻ 7.0 ⫻ 3.5 mm were then cut from the tiles with a high speed diamond saw. The test specimens were ground and polished using diamond abrasives on an automatic polisher (Spectrum 2000, Leco Corporation, St. Joseph, MI) down to 0.05 m colloidal silica finish, after which a notch was cut in the center of each test specimen with a 150 m wide diamond blade. The identical procedure was used to prepare control specimens of a commercial monolithic alumina (Friatec-Degussit A1-23, nominal grain size ⫽ 10 m). Fracture toughness was measured by the single edge notched beam (SENB) method following the DIN NMP 51 109 standard in four-point bending (17). Test specimens were loaded at a rate of 0.1 mm/min in an MTS Sintech 30/D machine (MTS System Corporation, Eden Prairie, MN) equipped with a fully articulating four-point bend fixture (Model CU-FL-34, Wyoming Test Fixtures Inc., Laramie, WY). Inner and outer loading spans were 18 mm and 36 mm, respectively. Load vs. displacement data was acquired using MTS TestPad v1.0 software. Fracture toughness calculations were based on the peak load required to propagate the machined notch. The increased sample cross-sections (40 ⫻ 7 ⫻ 3.5 mm from the specified 45 ⫻ 4 ⫻ 3 mm) for multilayered samples was used to increase the number of layers that a crack would have to propagate through before final fracture. The influence of the dimensional change on KIC was investigated by testing commercial monolithic alumina control samples for both geometries. Six specimens from each geometry had a resultant KIC of 4.3 (0.4) MPa䡠m1/2, and due to this constancy in KIC, the increased sample cross-section was incorporated for the remaining experiments. The KIC of the commercial alumina control samples was higher than the reported ⬃3 MPa䡠m1/2 of alumina previously tested by this method (18,19). The ⬃180 m notch-width used in this study likely resulted in artificially high KIC values (18 –22), where notch widths less than 100 m are necessary to minimize notch-width influence on KIC in alumina (18,19). The KIC values of the multilayered composites may also be artificially high, however, comparisons between the resulting KIC from different layer microstructures provides the relative KIC associated with a particular layer microstructure.
Results and Discussion Samples produced by the centrifugation process have two characteristic macroscopic features, layer curvature and tilt, which are inherent in the centrifugation process (23). Particles are deposited having curvature similar to the rotation radius based on radial force equilibrium and acceleration. Additionally, particle size segregation occurs within the individual layers as shown in Figure 1. The layer interfaces in Tile 3 (Fig. 1b) were more distinct, as compared to Tile 1 (Fig. 1a), due to the large fraction of porosity resulting from the segregation of the hard alumina agglomerates during centrifugation. The fracture results for the layered alumina specimens are given in Table 2. Tiles 1, 2, 5, 6, 8, and 9 failed catastrophically as is typical for an alumina monolith, however, large crack deflections and continued loading after the peak load were observed for Tiles 3, 4, 7, and 10. The differences in fracture behavior between Tiles 1, 2, and 3 resulted from the increased porosity concentration caused by the increased alumina agglomerate additions. The “weaker” bonding at layer-layer interfaces allowed crack deflection along the interfaces. Additionally, Tiles 3 and 4 (weaker interfaces) exhibited differences in fracture toughness depending on whether the crack propagated away (orientation 1) or towards (orientation 2) the center of curvature
752
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Figure 1. Microstructures of a given layer within (a) Tile 1, and (b) Tile 3. The increased porosity at the layer interface in Tile 3 was associated with the segregation of the large alumina agglomerates during centrifugation.
of the layers (Figure 2 and Table 2). Specimens tested in orientation 1 typically exhibited crack deflection less than 1 mm before kinking out of the interface to ultimate failure. Specimens loaded with notch orientation 2 exhibited crack deflection with continued crack growth along the first adjoining interface until reaching the limit of the inner loading span. Subsequently, the crack propagated normal to the interfaces until ultimate failure occurred. The increased fracture toughness for orientation 2 is believed to result from the crack propagating closer to Mode I where the layer curvature facilitates deviation of the crack into the interface. In contrast, for Tiles 1 and 2 (stronger interfaces), layer curvature had little effect on the fracture toughness results since large scale crack deflection did not occur. Tiles 5, 6, 7 and Tiles 8, 9, 10 were used to investigate the influence of layer thickness (400, 500, and 800 m) on KIC and the results are plotted in Figure 3. The increase in KIC with thicker layers corresponded to an increase in the degree of crack deflection at layer interfaces. The thicker layers obtained by injecting a larger slurry volume, produced thicker porous interlayers resulting from the larger amounts of alumina agglomerates which settled preferentially at the layer interface during centrifugation. The differences in KIC between Tiles 3 and 4 to that of Tiles 6 and 9, each having 500 m layer thickness, resulted from differences in porosity at the layer interface. The segregation of porosity at the layer interface for Tiles 3 and 4 was similar to Tiles 7 and 10 respectively and resulted in similar KIC values. Slight changes in the suspension preparation and insertion into the casting containers may account for the change in porosity and thus the variation of KIC.
Figure 2. (a) Crack propagation behavior observed in Tile 3 showing the influence of the notch orientation where, (b) shows the notch orientations 1 and 2 with respects to the layer curvature.
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Figure 3. The increasing KIC with thicker layers was associated with a larger segregation of porosity at the layer interface allowing for crack deflection to occur.
Summary Multilayered alumina and alumina/agglomerated alumina composites were produced by centrifugal consolidation. The extent of porosity at layer interfaces was varied by controlling the volume fraction of agglomerated additions to the suspension and by controlling layer thickness. Porosity at the layer interfaces decreases the interface bond strength, allowing cracks to kink and propagate along layer interfaces (parallel to layers). The likelihood of crack deflection along layer interfaces increased as the extent of porosity at the layer interfaces increased. The resulting KIC of multilayered composites increased as the amount of crack deflection increased where the apparent KIC transverse to the layers was nearly twice that of a commercial monolithic alumina. Additionally, multilayered composites without agglomerated alumina additions demonstrated fracture behavior and KIC values similar to commercial monolithic alumina. Acknowledgement This work has been supported by the Army Research Office: MURI grant no. DAAH04-96-1-0331. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.
A. J. Phillipps, W. J. Clegg, and T. W. Clyne, Composites. 25, 524 (1994). W. Lee, S. J. Howard, and W. J. Clegg, Acta Mater. 44, 3905 (1996). H. C. Cao and A. G. Evans, Mech. Mater. 7, 295 (1989). A. G. Evans, M. Ruehle, B. J. Dalgleish, and P. G. Charalambides, Metall. Trans. A. 21, 2419 (1990). A. A. Mammoli, A. L. Graham, I. E. Reimanis, and D. L. Tullock, Acta Metall. Mater. 43, 1149 (1995). C. M. Warwick and W. J. Clegg, Ceram. Eng. Sci. Proc. 72, 847 (1995). W. J. Clegg, K. S. Blanks, J. B. Davis, and F. Lanckmans, Key Eng. Mater. 132–136, 1866 (1997). C. Folsom, F. W. Zok, and F. Lange, J. Am. Ceram. Soc. 77, 689 (1994). C. Folsom, F. W. Zok, and F. Lange, J. Am. Ceram. Soc. 77, 2081 (1994). M. Oechsner, C. Hillman, and F. F. Lange, J. Am. Ceram. Soc. 79, 1834 (1996). J. C. Chang, B. V. Velamakanni, F. F. Lange, and D. S. Pearson, J. Am. Ceram. Soc. 74, 2201 (1991). W. Huisman, T. Graule, and L. J. Gauckler, J. Eur. Ceram. Soc. 13, 33 (1994). W. Huisman, T. Graule, and L. J. Gauckler, J. Eur. Ceram. Soc. 15, 811 (1995). R. K. Roeder, K. P. Trumble, and K. J. Bowman, J. Am. Ceram. Soc. 80, 27 (1997).
754
15. 16. 17. 18. 19. 20. 21. 22. 23.
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Vol. 41, No. 7
A. S. Barnes, J. S. Reed, and E. M. Anderson, Am. Ceram. Soc. Bull. 76, 77 (1997). E. Drewry, M. S. Thesis, Purdue University (August 1998). DIN NMP 51 109, German Standard for Fracture Toughness Analysis in Four-Point Flexure. R. J. Primas and R. Gstrein, Fatigue Fract. Eng. Mater. Struct. 20, 513 (1997). R. J. Damani, Key Eng. Mater. 132–136, 484 (1997). R. Damani, R. Gstrein, and R. Danzer, J. Eur. Ceram. Soc. 16, 695 (1996). J. Ku¨bler, Ceram. Eng. Sci. Proc. 18, 155 (1997). T. Nishida, Y. Hanaki, and G. Pezzotti, J. Am. Ceram. Soc. 77, 606 (1994). R. Moon, E. Drewry, K. Bowman, and K. Trumble, in 5th FGM Conference Proceedings, Dresden, Germany, Trans. Tech. Publications.