theoretical and ne ionIcs ELSEVIER Theoretical and applied fracture Mechanics 24(1995)13-19 High toughness silicon carbide/ graphite laminar composite by slip casting hang V. Krstic Department of Materials and A slip casting method has been developed to manufacture high toughness laminated SiC ceramics. The samples are produced by slip-casting alternate layers of Sic and graphite of various thicknesses ranging from 120 um to 600 m. After casting, the samples were dried and then pressureless sintered to high density. The graphite layer serves o deflect the crack and raise the apparent fracture toughness from 4 MPa m/ to 14 MPa m/2. a strong effect of SiC and graphite layer thickness on apparent fracture toughness and flexural strength was observed. 1. Introduction workers [4-6] have developed silicon carbide lam inated ceramics with carbon serving as weak in The concept of enhancing fracture toughness terfaces. The technique consists of mixing Sic of materials by introducing weak interfaces per- powder with 30-40 vol% of polyvinyl alcohol pendicular to the direction of crack propagation solution, and rolling the mass into flexible sheets has been most widely used in fibre composites [1]. 200 um thick. The sheets are then coated with The success of this approach has led to the devel- thin layer of carbon which, upon sintering, serves opment of ceramic matrix-fibre composites with as week interface. The coated sheets are then resistance to crack propagation an order of mag- debinderized to eliminate organic component and nitude higher than that of the monolithic ceramic subsequently sintered to achieve high density material [2,3]. However, due to the relatively high Graphite is particularly suitable for making ost of fibres and difficulties associated with the weak interfaces because it happens to be a suc incorporation of long fibres into ceramic matri- cessful sintering aid for SiC and at high tempera- ces, new approaches are being sought and new ture it does not react with Sic to form non- techniques for toughening are being developed stoichiometric compounds [7]. Although this sys- involving no fibres. Recently, Clegg and co- tem is relatively stable under neutral and reduc- ing environments, it begins to oxidize at 600C [7] In Corresponding author Tel. +1 613 545 2754, fax +1 613 5456610 The present paper reports the results on me- SSD0167-8442(95)000275
ELSEVIER Theoretical and Applied Fracture Mechanics 24 (1995) 13-19 theore6cal and applied fradure mechanics High toughness silicon carbide/graphite laminar composite by slip casting L. Zhang, V.D. Krstic * Department of Materials and Metallurgical Engineering, Queen's University, Kingston, Ontario, Canada Abstract A slip casting method has been developed to manufacture high toughness laminated SiC ceramics. The samples are produced by slip-casting alternate layers of SiC and graphite of various thicknesses ranging from 120 I~m to 600 I~m. After casting, the samples were dried and then pressureless sintered to high density. The graphite layer serves to deflect the crack and raise the apparent fracture toughness from 4 MPa. m 1/2 to 14 MPa. m 1/2. A strong effect of SiC and graphite layer thickness on apparent fracture toughness and flexural strength was observed. 1. Introduction The concept of enhancing fracture toughness of materials by introducing weak interfaces perpendicular to the direction of crack propagation has been most widely used in fibre composites [1]. The success of this approach has led to the development of ceramic matrix-fibre composites with resistance to crack propagation an order of magnitude higher than that of the monolithic ceramic material [2,3]. However, due to the relatively high cost of fibres and difficulties associated with the incorporation of long fibres into ceramic matrices, new approaches are being sought and new techniques for toughening are being developed involving no fibres. Recently, Clegg and co- * Corresponding author. Tel. + 1 613 545 2754, fax + 1 613 545 6610. workers [4-6] have developed silicon carbide laminated ceramics with carbon serving as weak interfaces. The technique consists of mixing SiC powder with 30-40 vol% of polyvinyl alcohol solution, and rolling the mass into flexible sheets 200 i~m thick. The sheets are then coated with thin layer of carbon which, upon sintering, serves as week interface. The coated sheets are then debinderized to eliminate organic component and subsequently sintered to achieve high density. Graphite is particularly suitable for making weak interfaces because it happens to be a suecessful sintering aid for SiC and at high temperature it does not react with SiC to form nonstoichiometric compounds [7]. Although this system is relatively stable under neutral and reducing environments, it begins to oxidize at 600"C [7] in air, and this may give rise to reduction in properties. The present paper reports the results on me- 0167-8442/95/$09.50 © 1995 Elsevier Science B.V. All rights reserved SSDI 0167-8442(95)00027-5
L. Zhang, V.D. Krstic/ Theoretical and Applied Fracture Mechanics 24(1995)13-19 chanical properties of laminated SiC ceramics powder-slurry, respectively. The fracture tough produced by a slip casting technique involving no ness was calculated using the following equation organic binder. Another novel feature of this [8] echnique is that graphite layer forming wea nterface between the Sic layers is converted into P(S-s)rvA SiC in the course of sintering BYw w 2(1-a) where P is the load at the fracture, s is the outer span, s is the inner span, B is the thickness, w 2. Materials fabrication the width, and Laminated structures were produced by slip F=1.9887-1.326a casting alternating layers of SiC and graphite in 349-0.68a+135a2)a(1-a) desirable thicknesses. Submicron size Sic pow der, produced by carbothermal reaction of silica and carbon [8), was mixed with 8 wt% AL, O, and where a=(w-a)/w,and(w-a)is the notch 4 wt%Y, 0, to produce a slurry with solid to depth liquid ratio of 30/70 by volume. The concentra At least five samples were tested for tion of carbon in a graphite slurry was between point. The microstructural analysis was per 2.5 and 5.0% by volume, depending on the de- formed using an optical and SEM microscopes sired thickness of the carbon layer in the struc- and the phase analyses was conducted using an ture. Slip casting was carried out in the plaster of X-ray diffraction technique paris mould in such a way to ensure that th interpenetration of Sic and carbon layer does occur [10]. This is done by adjusting the viscosity 3. Results and discussion of the slurry and increasing the rate of drying The thickness of the graphite layer was varied To compare the fracture behaviour of the lam from 3 to 20 um and that of SiC from 100 to 600 inated structures and a monolithic silicon car. um. The number of SiC layers in the structure bide, a series of four-point bend tests were per- was varied from 1 to 30. After slip casting, the formed. Fig. I shows a typical load-displacement reen samples were dried in air for 24 hours and curves for monolithic block and for laminated then sintered at 1870 C for 30 min The average materials of equivalent density. While the mono- density of the sintered samples was between 96 lithic SiC exhibited linear -elastic behaviour until and 98% of its theoretical density. The theoreti- cal density of the sintered samples was calculated on the basis of the initial concentration of SiC, The sintered samples were machined into rect ngular bars with dimensions 3 mm thick, 4mm wide, and 35 mm long and fractured using four point bend jig having an inner and outer span of 5 and 25 mm, respectively Three different notch widths were introduced into the samples in order to measure the effect of notch width on four point bend toughness. The 0.5 mm wide notch was introduced using a dia- mond wheel, whereas 0.225 mm and 0. 125 mm I'plocnten: mIry) wide notches were cut using diamond impreg Fig. I. Typical load-displacement curve for monolithic(a)and nated tungsten wire and copper wire with BC aminated (b) SiC ceramics
14 L. Zhang, V.D. Krstic / Theoretical and Applied Fracture Mechanics 24 (1995) 13-19 chanical properties of laminated SiC ceramics produced by a slip casting technique involving no organic binder. Another novel feature of this technique is that graphite layer forming weak interface between the SiC layers is converted into SiC in the course of sintering. 2. Materials fabrication Laminated structures were produced by slip casting alternating layers of SiC and graphite in desirable thicknesses. Submicron size SiC powder, produced by carbothermal reaction of silica and carbon [8], was mixed with 8 wt.% A120 3 and 4 wt.% Y20 3 to produce a slurry with solid to liquid ratio of 30/70 by volume. The concentration of carbon in a graphite slurry was between 2.5 and 5.0% by volume, depending on the desired thickness of the carbon layer in the structure. Slip casting was carried out in the plaster of paris mould in such a way to ensure that the interpenetration of SiC and carbon layer does not occur [10]. This is done by adjusting the viscosity of the slurry and increasing the rate of drying. The thickness of the graphite layer was varied from 3 to 20 ~m and that of SiC from 100 to 600 ~m. The number of SiC layers in the structure was varied from 1 to 30. After slip casting, the green samples were dried in air for 24 hours and then sintered at 1870°C for 30 min. The average density of the sintered samples was between 96 and 98% of its theoretical density. The theoretical density of the sintered samples was calculated on the basis of the initial concentration of SiC, A1203, Y203 and C. The sintered samples were machined into rectangular bars with dimensions 3 mm thick, 4 mm wide, and 35 mm long and fractured using four point bend jig having an inner and outer span of 5 and 25 mm, respectively. Three different notch widths were introduced into the samples in order to measure the effect of notch width on four point bend toughness. The 0.5 mm wide notch was introduced using a diamond wheel, whereas 0.225 mm and 0.125 mm wide notches were cut using diamond impregnated tungsten wire and copper wire with B4 C powder-slurry, respectively. The fracture toughness was calculated using the following equation [81: e (S-s) 3rc~d K'c B~/-W W 2(1-a) ' (1) where P is the load at the fracture, S is the outer span, s is the inner span, B is the thickness, W is the width, and F = 1.9887 - 1.326a (3.49 - 0.68a + 1.35a2)a(l - a) (1 + a) z , (2) where a = (W-a)/W, and (W-a) is the notch depth. At least five samples were tested for each point. The microstructural analysis was performed using an optical and SEM microscopes, and the phase analyses was conducted using an X-ray diffraction technique. 3. Results and discussion To compare the fracture behaviour of the laminated structures and a monolithic silicon carbide, a series of four-point bend tests were performed. Fig. 1 shows a typical load-displacement curves for monolithic block and for laminated materials of equivalent density. While the monolithic SiC exhibited linear-elastic behaviour until -4 t -\ _ • • ~,_, -, • "' : 7 rT7 Fig. 1. Typical load-displacement curve for monolithic (a) and laminated (b) SiC ceramics
L. Zhang, V.D. Krstic/ Theoretical and Applied fracture Mechanics 24(1995)13-19 o3focge Average Graphite Loyer Thickness (um) Fig. 2. The sequence of fracture of SiC laminates(four layers thout notch). The first drop in load (plot 1)is due to Fig. 3. Variation of strength with average graphite layer hickness in laminated sic ormation of first interfacial crack between first and second yers(from bottom). The second drop (plot 2)in load is due to fracture of the second layer, etc. the reduction of interfacial strength and the ease with which the cracks are initiated in the individ fracture, the laminated material showed a slow ual layers. Fig. 5 shows the path of a crack decrease of load after reaching a maximum level, through the laminated specimen during four-point along with large increase in roller displacement bend testing. Clearly, the position of crack initia prior to fracture. Another important difference tion at the surface of each individual Sic layer between the two materials is the shape of the varied from layer to layer, indicating that the load-displacement plots in the linear-elastic re- strength of these layers controls the strength of gion. An abrupt drop in load, followed by an the entire sample ncrease in load before the load reach a maxi- The relationship between the four-point bend mum value, was observed in all samples possess- fracture toughness and Sic and graphite layer ing high apparent fracture toughness(see Fig. 1). thickness is shown in Figs 6 and 7. In the case of A series of tests was conducted such that at the thinner graphite layers(5 um) there is a continu- moment when the first sign of drop in load was ous drop in toughness with the increase of Sic noticed the load was quickly relaxed to zero and layer thickness, whereas for thicker graphite lay he samples were examined microscopically(Fig. ers(10 um)a maximum in toughness was reached 2). It was found that all these samples developed t approximately 300 to 350 um thick SiC layers cracks along the interface between the first and the next adjacent layer. The ease at which crack- ing of the interface occurs appears to be con- trolled by the initial thickness of the carbon lay., soto- ers. Furthermore it was observed that both the carbon and sic layer thickness strongly influence he fracture response of the sample. Figs. 3 and 4 show the change of fracture strength with initial carbon and SiC layer thickness A maximum four- point bend strength was 2 Graphite Layer achieved with graphite layer thicknesses between 3 and 5 um and with Sic layer thicknesses be. 8 tween 300 and 600 um. a sharp drop in strength with samples having graphite layer thicknesses SiC Layer Thickness (urm) above 3 to 5 um is believed to be associated with g. 4. Variation of strength with average SiC layer thickness
L. Zhang, V.D. Krstic / Theoretical and Applied Fracture Mechanics 24 (1995) 13-19 15 4O 52 ' /i~ /' 12+1 i / i cl ~L _~_~._~_ F ~, _~ ~q,', I 2 3 4 5 h, elotive Disp,acernent (XO 1 rc/r~ ) Fig. 2. The sequence of fracture of SiC laminates (four layers without notch). The first drop in load (plot 1) is due to formation of first interracial crack between first and second layers (from bottom). The second drop (plot 2) in load is due to fracture of the second layer, etc. fracture, the laminated material showed a slow decrease of load after reaching a maximum level, along with large increase in roller displacement prior to fracture. Another important difference between the two materials is the shape of the load-displacement plots in the linear-elastic region. An abrupt drop in load, followed by an increase in load before the load reach a maximum value, was observed in all samples possessing high apparent fracture toughness (see Fig. 1). A series of tests was conducted such that at the moment when the first sign of drop in load was noticed, the load was quickly relaxed to zero and the samples were examined microscopically (Fig. 2). It was found that all these samples developed cracks along the interface between the first and the next adjacent layer. The ease at which cracking of the interface occurs appears to be controlled by the initial thickness of the carbon layers. Furthermore, it was observed that both the carbon and SiC layer thickness strongly influence the fracture response of the sample. Figs. 3 and 4 show the change of fracture strength with initial carbon and SiC layer thickness. A maximum four-point bend strength was achieved with graphite layer thicknesses between 3 and 5 Ixm and with SiC layer thicknesses between 300 and 600 p~m. A sharp drop in strength with samples having graphite layer thicknesses above 3 to 5 p,m is believed to be associated with z= E. F 2 D o L 550 46O 370 SiC Lay --0-- 500~,m 0 -- £-- 120/zm lOG I I I I 0 5 6 9 li2 lt5 Average Graphite Layer Thickness {/~m) Fig. 3. Variation of strength with average graphite layer thickness in laminated SiC. the reduction of interfacial strength and the ease with which the cracks are initiated in the individual layers. Fig. 5 shows the path of a crack through the laminated specimen during four-point bend testing. Clearly, the position of crack initiation at the surface of each individual SiC layer varied from layer to layer, indicating that the strength of these layers controls the strength of the entire sample. The relationship between the four-point bend fracture toughness and SiC and graphite layer thickness is shown in Figs. 6 and 7. In the ease of thinner graphite layers (5 p,m) there is a continuous drop in toughness with the increase of SiC layer thickness, whereas for thicker graphite layers (10 I~m) a maximum in toughness was reached at approximately 300 to 350 Ixm thick SiC layers. o [L £ ~2 7 [r o 550 1 • 4so ,5j-//~- 9- I r / // / 250 ~// I 150- 50 -- q ----+- 100 200 Graphite Layer --0-- 5~,m • -- 10/~m 300 400 500 600 700 SiC Layer Thickness (,u,m) Fig. 4. Variation of strength with average SiC layer thickness
L. Zhang, V.D. Krstic/Theor Applied Fracture Mechanics 24(1995)13-19 Fig. 7. Change of apparent fracture toughness with graphite layer thickness. Fig. 5. Polished section of the laminated SiC ceramics showing the graphite layer to form the Sic according to the crack path This maximum in fracture toughness, at particu SiO+2C→SiC+CO (4) lar SiC layer thickness, occurs in all sample Clearly, the extent of graphite conversion to where the ratio of SiC layer thickness to graphite Sic depends on the relative amount of carbon in layer thickness is around 30. It is believed that the graphite and the amount of Sio liberated this behaviour is associated with the extent of from the Sic layer. The amount of Sio formed graphite conversion to SiC which in turn controls depends, in turn, on the SiC layer thickness the strength of the interface. Because of the Thus, there will be some optimum SiC layer afinity between the carbon from the silicon car- thickness capable of producing sufficient amount bide and oxygen from the alumina sintering aid of Sio to completely convert graphite into SiC volatile Al,O component is formed [10] and form an optimum interfacial strength. Based SiC Al2O3+ Al20+ Sio + CO on this argument, one can predict that the higher As it forms, the Sio component diffuses the graphite layer thickness, the thicker the SiC through the sample and reacts with carbon from layer thickness is required to achieve an optimum interfacial strength. Typical micrograph of graphite layers before and after conversion is shown in Fig. 8. Inspection of Fig. 8. shows that before sintering, the graphite layers are in the form of network while after sintering the inter face is full of large and uniform size hexagona a-SiC crystals. X-ray analysis confirmed that these were indeed a-SiC crystals( Fig 9). Another important characteristic of this new class of SiC materials is the notch width insensi- tivity seldom found in brittle materials. Fig. 10 shows the variation of apparent fracture tough ness with notch width for monolithic and lami- nated SiC. As expected, a continuous decrease in Fig. 6. Measured variation of apparent fracture toughness acture toughness with decrease in notch width with average Sic layer thickness. was observed in monolithic SiC, whereas in the
16 L. Zhang, V.D. Krstic / Theoretical and Applied Fracture Mechanics 24 (1995) 13-19 ] • t : l,r.! ] ,~,l i i :: t •1 ] Fig. 7. Change of apparent fracture toughness with graphite layer thickness. Fig. 5. Polished section of the laminated SiC ceramics showing the crack path. This maximum in fracture toughness, at particular SiC layer thickness, occurs in all samples where the ratio of SiC layer thickness to graphite layer thickness is around 30. It is believed that this behaviour is associated with the extent of graphite conversion to SiC which in turn controls the strength of the interface. Because of the afinity between the carbon from the silicon carbide and oxygen from the alumina sintering aid, a volatile Al20 component is formed [10]: SiC + A1203 "-~ A]20 -t- SiO + CO. (3) As it forms, the SiO component diffuses through the sample and reacts with carbon from L ;? i: : i i : 1 i S ; Fig. 6. Measured variation of apparent fracture toughness with average SiC layer thickness. the graphite layer to form the SiC according to the reaction: SiO + 2C ~ SiC + CO. (4) Clearly, the extent of graphite conversion to SiC depends on the relative amount of carbon in the graphite and the amount of SiO liberated from the SiC layer. The amount of SiO formed depends, in turn, on the SiC layer thickness, Thus, there will be some optimum SiC layer thickness capable of producing sufficient amount of SiO to completely convert graphite into SiC and form an optimum interracial strength. Based on this argument, one can predict that the higher the graphite layer thickness, the thicker the SiC layer thickness is required to achieve an optimum interracial strength. Typical micrograph of graphite layers before and after conversion is shown in Fig. 8. Inspection of Fig. 8. shows that, before sintering, the graphite layers are in the form of network while after sintering the interface is full of large and uniform size hexagonal a-SiC crystals. X-ray analysis confirmed that these were indeed a-SiC crystals (Fig. 9). Another important characteristic of this new class of SiC materials is the notch width insensitivity seldom found in brittle materials. Fig. 10, shows the variation of apparent fracture toughness with notch width for monolithic and laminated SiC. As expected, a continuous decrease in fracture toughness with decrease in notch width was observed in monolithic SiC, whereas in the
L. Zhang, V.D. Krstic/ Theoretical and Applied Fracture Mechanics 24(1995)13-19 a Fig. 8. SEM micro h of graphite layer(a)before and( b)after sintering showing SiC crystals
L. Zhang, V.D. Krstic / Theoretical and Applied Fracture Mechanics 24 (1995) 13-19 17 (a) (b) Fig. 8. SEM micrograph of graphite layer (a) before and (b) after sintering showing SiC crystals
L. Zhang, V.D. Krstic/Theoretical and Applied Fracture Mechanics 24(1995)13-19 3000 conceivable that the strength of the laminated be controlled by the probability of crack during loading at the surface of the layers. These cracking and delamination processes determine th For example, when the as produced individual Sic layers are stacked together and loaded four-point bending, the strength of the structure dropped to 70 M o 8 the individual layers were polished before stack ing and then separated by graphite powder, the 9. X-ray diffraction pattem for the interfacial layer of strength of the sample did not exceed 200 MPa The reason for this large drop in strength is thought to be associated with the formation of cracks at the interface between the sic and case of laminated slabs the apparent fracture graphite layers due to frictional forces and rub- toughness remained unchanged. This behaviour is bing effects caused by the deflection of the indi- not surprising considering that the primary crack vidual layers. This processes was found to influ- extending from the notch will be deflected along ence critically the strength response of individual he weak interface and the fracture of the next layers in the laminated structure. If, for example Sic layer will be controlled by the initiation of a the sliding between the individual layers does not new crack on its surface It is well known that this create cracks, then the strength of the laminated process requires higher stress than that involving structure should be the same as that of a single extension of the existing cracks. It follows then layer (i.e. 800 MPa). The fact that the strength of that the apparent toughness of the laminated laminated structure was lower than that of the structure will be governed by the interfacial con- single isolated layer suggests that cracks are dition and the strength of the individual Sic formed at the surface of the individual layers layers. A series of strength tests were co when they slide relative to each other Thus, one on a single SiC layer (with thicknesses of the prerequisites for achieving high strength of from 0.3 to 1 mm) and found that the laminated structure consists of lowering the inter strength was in excess of 800 MPa. It is now quite facial strength between the layers and minimizing the likelihood of creating cracks on the surface of Based on these findings, it is clear that the apparent fracture toughness of the laminated is governed by the thickness of SiC and graphite layer, and on the strength of individual SiC lay- ers. Thus, using this argument, the apparent frac- ture toughness may be calculated using the ex CI Monolithic KIc=oYc(1-c/d) (5) where o is the fracture strength, C is the notch length, d is the thickness of the beam and y is a otch width (umi constant (normally 1. 12). Substituting values for Fig. 10. Change of measured fracture toughness with notch 0r-350 MPa, c=0.002 m, and d=500 X 10-6 m in Eq(12)gives KIc=6.3 MPa m/.This
18 L. Zhang, V.D. Krstic / Theoretical and Applied Fracture Mechanics 24 (1995) 13-19 i ,~;' r ,:~ i il ::; 2) 30 40 50 60 7( 80 2 [het@, deq Fig. 9. X-ray diffraction pattern for the interracial layer of SiC-graphite laminates after sintering. All major peaks are due to SiC. case of laminated slabs the apparent fracture toughness remained unchanged. This behaviour is not surprising considering that the primary crack extending from the notch will be deflected along the weak interface and the fracture of the next SiC layer will be controlled by the initiation of a new crack on its surface. It is well known that this process requires higher stress than that involving extension of the existing cracks. It follows then that the apparent toughness of the laminated structure will be governed by the interracial condition and the strength of the individual SiC layers. A series of strength tests were conducted on a single SiC layer (with thicknesses ranging from 0.3 to 1 mm) and found that the flexural strength was in excess of 800 MPa. It is now quite .? 1J 1 'i i( 9 ,1 2 7 I )( • Lominoted (I Monolithic 700 300 400 500 Notch Width (/~m) Fig. 10. Change of measured fracture toughness with notch width for monolithic and laminated SiC. conceivable that the strength of the laminated sample consisting of large number of layers will be controlled by the probability of crack initiation during loading at the surface of the layers. These cracking and delamination processes determine the ease with which the crack may be initiated. For example, when the as produced individual SiC layers are stacked together and loaded in four-point bending, the strength of the structure dropped to 70 MPa. Even when the surfaces of the individual layers were polished before stacking and then separated by graphite powder, the strength of the sample did not exceed 200 MPa. The reason for this large drop in strength is thought to be associated with the formation of cracks at the interface between the SiC and graphite layers due to frictional forces and rubbing effects caused by the deflection of the individual layers. This processes was found to influence critically the strength response of individual layers in the laminated structure. If, for example, the sliding between the individual layers does not create cracks, then the strength of the laminated structure should be the same as that of a single layer (i.e. 800 MPa). The fact that the strength of laminated structure was lower than that of the single isolated layer suggests that cracks are formed at the surface of the individual layers when they slide relative to each other. Thus, one of the prerequisites for achieving high strength of laminated structure consists of lowering the interfacial strength between the layers and minimizing the likelihood of creating cracks on the surface of the SiC layers. Based on these findings, it is clear that the apparent fracture toughness of the laminated slab is governed by the thickness of SiC and graphite layer, and on the strength of individual SiC layers. Thus, using this argument, the apparent fracture toughness may be calculated using the expression [6]: Ktc = fv¢? (1 - c/d) 2, (5) where ~f is the fracture strength, C is the notch length, d is the thickness of the beam and Y is a constant (normally 1.12). Substituting values for trf = 350 MPa, c = 0.002 m, and d = 500 × 10 -6 m in Eq. (12) gives Klc = 6.3 MPa" m 1/2. This
L. Zhang V.D. Krstic/Theoretical and Applied fracture Mechanics 24(1995)13-19 value for apparent fracture toughness is signifi- much larger strain to failure and that after the cantly smaller than the measured values. How- peak load is reached failure is still not catas- ever, when the strength of the individual layer is trophic used(i.e. 800 MPa. ) the predicted fracture tough- ness is 14.4 MPam/2, which is very close to experimentally measured Kic. It should be stressed however, that one of the problems asso- References ciated with Eq. (3)is that it does not include the effect of graphite layer thickness. As already dis- .E. Gordon, Mechanisam for the control gation in all-brittle system, Proc. R Soc. cussed, the strength and fracture toughness of 32,508520(1964) partially converted graphite layer depends criti [2]RA.. Sambell, D H. Bowen and D. C. Phillips, Carbon cally upon the thickness of the SiC layers. These fibre composites with ceramic and glass matrices, J. in turn, controls the fracture response of the 3] K.M. Prewo and J.J. Brennan, High-strength carbide entire structure fibre-reinforced glass-matrix composites. J. Mater. Sci. 15,463468(1980) [4]WJ. Clegg, K. Kendall, N M. Alford, D. Birchall and 4. Conclusion T.W. Button, A simple way to make tough ceramics Nature347,455-457(1990) Using a simple slip casting technique it has [5]wJ. Clegg and J D. Birchall, in: Proc. 4th int. Conf. on Fibre Reinforced Composites, edited by G. Gibson, Inst proven possible to produce laminated Sic ceram- Mechan Eng. Liv ics with apparent fracture toughness one order of [6] wJ. Clegg, The fracture and failure of laminar ceramic magnitude higher than that of a monolithic mate- composites, Acta Metall. Mater. 40, 3085-3093(1992). rial. Due to the presence of weak interfaces and [7] w.J. Clegg, K Kendall, N MCN. Alford, T w. Button and repeated crack initiation across each SiC layer, J D. Birchall, in: Brit. Ceram. ics, Vol 9, Institute of Ceramics. Stoke-on-Trent, UK, p. the structure exhibits no notch width sensitivity 263(1980) and is capable of supporting the same load as the [8] V D. Ktrstic, Prodaction of fine, high-purity beta silicon sample without the notch but of the same height bide powders, J. Am. Ceram. Soc. 75(1)170-174 It is shown that the toughening is controlled by (1992) both the Sic and graphite layer thickness. A [9] D.J. Munz, J. L. Shannon and R.t. Fracture calculation from Maximum four poir maximum in apparent toughness is achieved at of chevron notch specimens, fract. 16 Sic to graphite layer thickness of 30. This ratio provides the condition under which complete [10] L Zhang and V.D. Krstic, High Toughness Carbide and Method Thereof, U.s occurs leading to neration of an optimum interfacial strength Patent No.5443770(1995) [11]M.A. Mulla and V D. Krstic, Pressureless sintering of Another important characteristic of this new class B-SiC with Al2O, additions, J. Mater. Sci. 29, 934-938 of sic ceramics is that the material fails with (1994
L. Zhang, V.D. Krstic /Theoretical and Applied Fracture Mechanics 24 (1995) 13-19 19 value for apparent fracture toughness is significantly smaller than the measured values. However, when the strength of the individual layer is used (i.e. 800 MPa.) the predicted fracture toughness is 14.4 MPa. m 1/2, which is very close to experimentally measured K1c. It should be stressed, however, that one of the problems associated with Eq. (3) is that it does not include the effect of graphite layer thickness. As already discussed, the strength and fracture toughness of partially converted graphite layer depends critically upon the thickness of the SiC layers. These in turn, controls the fracture response of the entire structure. 4. Conclusions Using a simple slip casting technique it has proven possible to produce laminated SiC ceramics with apparent fracture toughness one order of magnitude higher than that of a monolithic material. Due to the presence of weak interfaces and repeated crack initiation across each SiC layer, the structure exhibits no notch width sensitivity and is capable of supporting the same load as the sample without the notch but of the same height. It is shown that the toughening is controlled by both the SiC and graphite layer thickness. A maximum in apparent toughness is achieved at SiC to graphite layer thickness of 30. This ratio provides the condition under which complete conversion of graphite to SiC occurs leading to generation of an optimum interracial strength. Another important characteristic of this new class of SiC ceramics is that the material fails with much larger strain to failure and that after the peak load is reached failure is still not catastrophic. References [1l J. Cook and J.E. Gordon, Mechanisam for the control of crack propagation in all-brittle system, Proc. K Soc. London A282, 508-520 (1964). [2l R.A.J. Sambell, D.H. Bowen and D.C. Phillips, Carbon fibre composites with ceramic and glass matrices, J. Mater. Sci. 7, 663-675 (1972). [3] K.M. Prewo and J.J. Brennan, High-strength carbide fibre-reinforced glass-matrix composites. J. Mater. Sci. 15, 463-468 (1980). [4] W.J. Clegg, K. Kendall, N.M. Alford, D. Birchall and T.W. Button, A simple way to make tough ceramics, Nature 347, 455-457 (1990). [5l W.J. Clegg and J.D. Birchall, in: Proc. 4th Int. Conf. on Fibre Reinforced Composites, edited by G. Gibson, Inst. of Mechan. Eng., Liverpool, pp. 179-184 (1990). [6] W.J. Clegg, The fracture and failure of laminar ceramic composites, Acta Metall. Mater. 40, 3085-3093 (1992). [7] W.J. Clegg, K. Kendall, N.McN. Alford, T.W. Button and J.D. Birchall, in: Brit. Ceram. Soc. Conf. Special Ceramics, Vol. 9, Institute of Ceramics. Stoke-on-Trent, UK, p. 263 (1980). [8] V.D. Ktrstic, Prodaction of fine, high-purity beta silicon carbide powders, J. Am. Ceram. Soc. 75(1) 170-174 (1992). [9] D.J. Munz, J.L Shannon and R.T. Bubsey, Fracture toughness calculation from Maximum load in four point bend tests of chevron notch specimens, Int. J. Fract. 16, 137 (1980). [10] L. Zhang and V.D. Krstic, High Toughness Carbide Ceramics by Slip Casting and Method Thereof, U.S. Patent No. 5,443,770 (1995). [11] M.A. Mulla and V.D. Krstic, Pressureless sintering of B-SiC with AI203 additions, J. Mater. ScL 29, 934-938 (1994)