Composites: Part A 40(2009)137-143 Contents lists available at ScienceDirect Composites: Part A ELSEVIER journalhomepagewww.elsevier.com/locate/compositesa Design and processing of a ceramic laminate with high toughness and strong interfaces S Bueno, C Baudin Instituto de Ceramica y Vidrio(CSIC. Campus de Cantoblanco, Kelsen 5, 28049 Madrid, Spain ARTICLE IN F O ABSTRACT Article history: An alumina-aluminium titanate laminate designed to Received 7 July 2008 trong interfaces is proposed. It is constituted by relative l stine ng b ite exsternct av eraspwith y c Received in revised form 13 October 2008 Accepted 28 October 2008 cracked internal layers to produce multiple crack deflection at the microstructural scale. The mo tant difference of the laminated structure proposed here and that of other laminates with high for crack deflection is that the crack deflection process occurs at local level, thus, delamination lengths are imited and delamination does not lead to the lost of structural integrity B Mechanical properties A symmetric structure formed by five layers has been design to minimise residual stresses taking into D Mechanical testing account the strain on cooling and the Youngs modulus of monolithic materials of the same compositio E Slip casting as those of the layers and fabricated using the same processing procedure as that of the laminate pecial care was given to adjust the processing variables that permitted the fabrication of the designed minated by sequential slip casting and sintering. ms of strength(4 modulus, work of fracture and apparent toughness. The two latter parameters have been determined by 3-points bending of Single-Edge-V-Notch-Beams (SEVNB) and fractographic analysis has been per- formed on the tested samples. The apparent toughness value at the point of failure(12 MPa m"/2)wa omparable to values reported for the stationary state of transformation-toughened ceramics. Work of fracture(62+3Jm-)was significantly higher(26%)than that obtained by calculation from the values corresponding to monolithic materials of the same composition as that of the layers, revealing the syn- ergic effect of the laminated structure on the mechanical behaviour of the material e 2008 Elsevier Ltd. All rights reserved. 1 Introduction and brittle aragonite platelets held together by a easily to deform and tough proteinaceous matrix make nacre a rigid material in for metals in structural applications that involve high temperature those of aragonite, which constitutes the 95 vol% of nacre. Several in severe erosive and corrosive environments and or compressive mechanisms leading to energy dissipation have been identified to loads. The major problem for the structural use of ceramics is re- occur during the fracture of nacre [2, 3 ] sliding of the aragonite lated with their brittle fracture mode, which implies the variation layers, stretching of the filaments in the proteinaceous matrix f strength of different components within the same batch as a and crack deflection around the aragonite plates. function of the distribution of strength limiting flaws. Although On the basis of the toughening mechanisms proposed for nad particularly weak components can be removed from the batch by two groups of materials have been developed. One of them com- proof testing, once a component enters service, subcritical growth bines relatively thick rigid external layers with thin internal layers of pre-existing flaws or the formation of new cracks, for instance capable of deformation and energy absorption during fracture a lead to unpredicted failure of the components [1. number of ceramic-metal and ceramic-polymer laminates have One of the most promising new approaches to avoid the lack of been developed on this basis [ 4-8 which main drawback is the mechanical reliability of ceramics is that of layered materials. Nat- lack of stability at high temperature due to the characteristics of ure offers a number of simple layered structures, such as shells or the metal and polymeric layers On the other hand, since the sem- teeth, which present improved failure behaviour as compared to inal work by Clegg et al. [9 ceramic-ceramic layered composites that of the individual components. For example, layers of stiff, hard have been designed and processed on the basis of weak interfaces between layers to originate crack deflection [1, 10-11 An alterna- Corresponding author. Tel. +34 91 7355840: fax: +34 91 7355843 tive way to produce crack deflection is to incorporate porous layers of the same composition between dense ceramic layers [12-14: natter 2008 Elsevier Ltd. All rights reserved
Design and processing of a ceramic laminate with high toughness and strong interfaces S. Bueno, C. Baudín * Instituto de Cerámica y Vidrio (CSIC), Campus de Cantoblanco, Kelsen 5, 28049 Madrid, Spain article info Article history: Received 7 July 2008 Received in revised form 13 October 2008 Accepted 28 October 2008 Keywords: A. Layered structures B. Mechanical properties D Mechanical testing E. Slip casting abstract An alumina–aluminium titanate laminate designed to combine high crack deflection capability with strong interfaces is proposed. It is constituted by relatively stiff and brittle external layers with microcracked internal layers to produce multiple crack deflection at the microstructural scale. The most important difference of the laminated structure proposed here and that of other laminates with high capability for crack deflection is that the crack deflection process occurs at local level, thus, delamination lengths are limited and delamination does not lead to the lost of structural integrity. A symmetric structure formed by five layers has been design to minimise residual stresses taking into account the strain on cooling and the Young’s modulus of monolithic materials of the same compositions as those of the layers and fabricated using the same processing procedure as that of the laminate. Special care was given to adjust the processing variables that permitted the fabrication of the designed laminated by sequential slip casting and sintering. Mechanical characterisation has been done in terms of strength (4-points bending), dynamic Young’s modulus, work of fracture and apparent toughness. The two latter parameters have been determined by 3-points bending of Single-Edge-V-Notch-Beams (SEVNB) and fractographic analysis has been performed on the tested samples. The apparent toughness value at the point of failure (12 MPa m1/2) was comparable to values reported for the stationary state of transformation-toughened ceramics. Work of fracture (62 ± 3 Jm2 ) was significantly higher (26%) than that obtained by calculation from the values corresponding to monolithic materials of the same composition as that of the layers, revealing the synergic effect of the laminated structure on the mechanical behaviour of the material. 2008 Elsevier Ltd. All rights reserved. 1. Introduction Ceramic materials are being proposed and used as substitutes for metals in structural applications that involve high temperature in severe erosive and corrosive environments and/or compressive loads. The major problem for the structural use of ceramics is related with their brittle fracture mode, which implies the variation of strength of different components within the same batch as a function of the distribution of strength limiting flaws. Although particularly weak components can be removed from the batch by proof testing, once a component enters service, subcritical growth of pre-existing flaws or the formation of new cracks, for instance by erosion, can lead to unpredicted failure of the components [1]. One of the most promising new approaches to avoid the lack of mechanical reliability of ceramics is that of layered materials. Nature offers a number of simple layered structures, such as shells or teeth, which present improved failure behaviour as compared to that of the individual components. For example, layers of stiff, hard and brittle aragonite platelets held together by a easily to deform and tough proteinaceous matrix make nacre a rigid material in which both toughness and strength are significantly higher than those of aragonite, which constitutes the 95 vol.% of nacre. Several mechanisms leading to energy dissipation have been identified to occur during the fracture of nacre [2,3]: sliding of the aragonite layers, stretching of the filaments in the proteinaceous matrix and crack deflection around the aragonite plates. On the basis of the toughening mechanisms proposed for nacre, two groups of materials have been developed. One of them combines relatively thick rigid external layers with thin internal layers capable of deformation and energy absorption during fracture. A number of ceramic-metal and ceramic-polymer laminates have been developed on this basis [4–8] which main drawback is the lack of stability at high temperature due to the characteristics of the metal and polymeric layers. On the other hand, since the seminal work by Clegg et al. [9] ceramic-ceramic layered composites have been designed and processed on the basis of weak interfaces between layers to originate crack deflection [1,10–11]. An alternative way to produce crack deflection is to incorporate porous layers of the same composition between dense ceramic layers [12–14]; 1359-835X/$ - see front matter 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.compositesa.2008.10.012 * Corresponding author. Tel.: +34 91 7355840; fax:+34 91 7355843. E-mail address: cbaudin@icv.csic.es (C. Baudín). Composites: Part A 40 (2009) 137–143 Contents lists available at ScienceDirect Composites: Part A journal homepage: www.elsevier.com/locate/compositesa
138 S Bueno, C Baudin/Composites: 2009)137-143 for deflection to be efficient very high values of the porosity are Al203 jar and balls during 4 h. These conditions were selected from needed [12]. Main drawback of the second approach is that crack a previous work [29] deflection limits the wear resistance of the materials Solid discs with 20 mm in diameter were slip cast in plaster of From all ceramics, alumina(Al2O3)presents the highest thermal Paris moulds in order to determine the casting rate of each suspen- stability together with high hardness sustained up to temperatures sion by determination of the dry wall thickness(Mitutoyo, JDU25 over 1200C. therefore is the natural ceramic for wear. Japan)after different casting times(1-16 min). For Young s modu- n this work, we present an alumina-based layered ceramic- lus and strength, plates with 70 x 70 x 6 mm dimensions were ceramic laminate designed on the basis of a combination of both obtained by slip casting of the a10 and A30 slips. The cast bodies discussed approaches. It is constituted by relatively stiff and brittle were carefully removed from the moulds and dried in air at room external layers and microcracked internal layers to produce multi- temperature for at least 24 h ple crack deflection at the microstructural scale, thus, limiting The reaction sintering behaviour was studied with a differential delamination lengths in order not to loose structural integrity. Spe- dilatometer(Adamel Lhomargy, DI24, France)to 1550C using cial care was given to adjust the processing variables that permit- small(5 x 5 x 5 mm)green samples of the monoliths with alt ed the fabrication of the designed laminated by sequential slip mina detector and the sintering schedule was selected from tl casting and sintering. tained results. To obtain the monolithic composites, the dried As internal material an alumina -aluminium titanate(Al2TiO5) plates were sintered in air in an electrical box furnace(termiber composite containing 30 vol% of aluiminium titanate previously Spain)at heating and cooling rates of 2C min, with 4 h dwell studied was chosen [15. In this material, extensive deflection and at 1200C and 2 h dwell at the maximum temperature, 1450C branching of the main crack along the pre-existing microcracks oc. Densities of the sintered compacts were determined by the archi- curred during fracture. An alumina -aluminium titanate composite medes method in water(European Standard En 1389: 2003)and with relatively high strength (10 voL% aluminium titanate, relative densities were calculated from these values and those of or= 261+6 MPa, 3-points bending, samples 2x 2.5 x 30 mm, theoretical densities calculated taking values of 3.99 g cm- for span 20 mm [16) was chosen to constitute the external layers. alumina(a-Al2O3, ASTM 42-1468)and 3.70 g cm- for aluminium External composite layers instead of pure alumina ones were chosen titanate(B-Al2TiOs, ASTM 26-0040) at the expense of strength(for monophase alumina or=456+ Additional sintering experiments(samples 12 x 5 mm) 29 MPa, 3-points bending, samples 2 x 2.5 x 30 mm, span 20 mm were performed in a differential dilatometer (Setaram, Setsys 16/ [16) in order to assure compatible sintering between the layers. 18, France)with alumina detector reproducing the same thermal On the one hand it is well-known that titania accelerates the initial treatment schedule as that used to obtain the final materials in or sintering of alumina[17-19] and, on the other, the formation of alu- der to determine strain during cooling. minium titanate is expansive(Ava 11%, calculated using density The sintered blocks of the monoliths were machined into bars of values of 3.99 g cm-for a-Al203 ASTM File 42-1468, 3.70 g cm-3 50 x 3 x 4 mm for dynamic Youngs modulus determinations. for B-Al, TiOs, ASTM File 26-0040, and 3. 89 g cm- for TiO2-anatasa, from the resonance frequency of the bars in flexure(Grindosonic, ASTM Files 21-1272)thus, an arrest of the shrinkage rate occurs at J.W. Lemmens, Belgium) and bend strength tests( 4-points bend- the temperature of aluminium titanate formation (1390.C) ing, 40-20 mm span, 0.5 mm min": Microtest, Spain) Reported Youngs modulus and bend strength values are the average of five The laminate design was done taking into account the total measurements and errors are strain of monolithic materials of both compositions when cooling A symmetric laminated structure of five layers was fabricated from the sintering temperature and their Youngs modulus in order by casting each suspension alternately Casting times were fixed to limit tensile residual stresses in the external layers to reach the desired layer thickness considering the casting kinet The work of fracture has been chosen as mechanical parameter ics and sintering shrinkage of each composition. The laminate had to establish the relationships between the mechanical behaviour of the central (1200 um)and outer layers(e2100 um )made of A10 the laminate and that of the constituent layers. The advantage of and the two inner layers(e300 um)of A30 this energy parameter is that it does not require any assumpt Microstructure of polished cross sections was characterised in a about the constitutive equation of the body with the crack to dis- field emission gun scanning electron microscopy(FEG-SEM, Hit- cuss its propagation [22]. Thus, it can be used to describe behav- achi, S-4700, Japan) Thermally etched (1440C-1min) specimens iours which separate from linearity and it is an additive were analysed. Additional observations were performed on chem- parameter that makes it possible to quantify the different contri- ically etched(hF 10 vol%-1 min) specimens in order to assure that butions to energy dissipation during fracture [23-27 the thermal etching did not produce further microcracking in the The apparent toughness as proposed by Clegg et al. 9. 28 for sintered specimens. laminates with weak interfaces has been used to compare the Single-Edge-V-Notch-Beams(SEvNB) of 4 x 6 x 50 mm,ma- mechanical performance of the proposed laminate with that of chined from the laminated sintered blocks, were tested in a other structural ceramics 3-points bending device using a span of 40 mm and a cross-head speed of 0.005 mm min(Microtest, Spain). The notches were ini tially cut with a 150 um wide diamond wheel. Using this slot as a guide, the remaining part of the notch was done with a razor blade sprinkled with diamond pastes of successively 6 and 1 um Speci The starting materials were commercial Alumina(a-Al2O3, Con- mens with relative notch depths of about 0.8 of the thickness of dea, HPAO5, USA)and titania(anatase-TiO2, Merck, 808, Germany) the first external layer, corresponding to 0. 26 of the specimen owders. Al 2O3/TiO2 mixtures with relative TiO2 contents of 5 and thickness( W), were tested. The tip radii of all notches were deter 15 wt% were prepared to obtain alumina/aluminium titanate com- mined from optical observations and they were always found to be posites with second phase contents of 10 and 30 vol% after reac- below 20 um. The curves load-displacement of the cross-head of tion sintering, named A10 and A30, respectively he load frame were recorded and corrected by subtracting the The mixtures were dispersed in deionised water by adding compliance of the testing set up(machine, supports, load cell 0.5 wt%(on a dry solids basis) of a carbonic acid based polyelectro- and fixtures, 1.5 x 10-m/N)determined by testing a thick lyte(Dolapix CE64, Zschimmer-Schwarz, Germany). Suspensions (25 x 25 x 100 mm )unnotched alumina bar. Three specimens were prepared to a solids loading of 50 vol% and ball milled with were tested and the curves were found to be practically coincident
for deflection to be efficient very high values of the porosity are needed [12]. Main drawback of the second approach is that crack deflection limits the wear resistance of the materials. From all ceramics, alumina (Al2O3) presents the highest thermal stability together with high hardness sustained up to temperatures over 1200 C, therefore, is the natural ceramic for wear. In this work, we present an alumina-based layered ceramicceramic laminate designed on the basis of a combination of both discussed approaches. It is constituted by relatively stiff and brittle external layers and microcracked internal layers to produce multiple crack deflection at the microstructural scale, thus, limiting delamination lengths in order not to loose structural integrity. Special care was given to adjust the processing variables that permitted the fabrication of the designed laminated by sequential slip casting and sintering. As internal material an alumina–aluminium titanate (Al2TiO5) composite containing 30 vol.% of aluiminium titanate previously studied was chosen [15]. In this material, extensive deflection and branching of the main crack along the pre-existing microcracks occurred during fracture. An alumina–aluminium titanate composite with relatively high strength (10 vol.% aluminium titanate, rf = 261 ± 6 MPa, 3-points bending, samples 2 2.5 30 mm3 , span 20 mm [16]) was chosen to constitute the external layers. External composite layers instead of pure alumina ones were chosen at the expense of strength (for monophase alumina rf = 456 ± 29 MPa, 3-points bending, samples 2 2.5 30 mm3 , span 20 mm [16]) in order to assure compatible sintering between the layers. On the one hand, it is well-known that titania accelerates the initial sintering of alumina [17–19] and, on the other, the formation of aluminium titanate is expansive (DV 11%, calculated using density values of 3.99 g cm3 for a-Al2O3, ASTM File 42-1468, 3.70 g cm3 for b-Al2TiO5, ASTM File 26-0040, and 3.89 g cm3 for TiO2-anatasa, ASTM Files 21-1272) thus, an arrest of the shrinkage rate occurs at the temperature of aluminium titanate formation (1390 C) [20,21]. The laminate design was done taking into account the total strain of monolithic materials of both compositions when cooling from the sintering temperature and their Young’s modulus in order to limit tensile residual stresses in the external layers. The work of fracture has been chosen as mechanical parameter to establish the relationships between the mechanical behaviour of the laminate and that of the constituent layers. The advantage of this energy parameter is that it does not require any assumptions about the constitutive equation of the body with the crack to discuss its propagation [22]. Thus, it can be used to describe behaviours which separate from linearity and it is an additive parameter that makes it possible to quantify the different contributions to energy dissipation during fracture [23–27]. The apparent toughness as proposed by Clegg et al. [9,28] for laminates with weak interfaces has been used to compare the mechanical performance of the proposed laminate with that of other structural ceramics. 2. Experimental The starting materials were commercial Alumina (a-Al2O3, Condea, HPA05, USA) and titania (anatase-TiO2, Merck, 808, Germany) powders. Al2O3/TiO2 mixtures with relative TiO2 contents of 5 and 15 wt.% were prepared to obtain alumina/aluminium titanate composites with second phase contents of 10 and 30 vol.% after reaction sintering, named A10 and A30, respectively. The mixtures were dispersed in deionised water by adding 0.5 wt.% (on a dry solids basis) of a carbonic acid based polyelectrolyte (Dolapix CE64, Zschimmer-Schwarz, Germany). Suspensions were prepared to a solids loading of 50 vol.% and ball milled with Al2O3 jar and balls during 4 h. These conditions were selected from a previous work [29]. Solid discs with 20 mm in diameter were slip cast in plaster of Paris moulds in order to determine the casting rate of each suspension by determination of the dry wall thickness (Mitutoyo, JDU25, Japan) after different casting times (1–16 min). For Young’s modulus and strength, plates with 70 70 6 mm3 dimensions were obtained by slip casting of the A10 and A30 slips. The cast bodies were carefully removed from the moulds and dried in air at room temperature for at least 24 h. The reaction sintering behaviour was studied with a differential dilatometer (Adamel Lhomargy, DI24, France) to 1550 C using small (5 5 5 mm3 ) green samples of the monoliths with alumina detector and the sintering schedule was selected from the obtained results. To obtain the monolithic composites, the dried plates were sintered in air in an electrical box furnace (Termiber, Spain) at heating and cooling rates of 2 C min1 , with 4 h dwell at 1200 C and 2 h dwell at the maximum temperature, 1450 C. Densities of the sintered compacts were determined by the Archimedes method in water (European Standard EN 1389:2003) and relative densities were calculated from these values and those of theoretical densities calculated taking values of 3.99 g cm3 for alumina (a-Al2O3, ASTM 42-1468) and 3.70 g cm3 for aluminium titanate (b -Al2TiO5, ASTM 26-0040). Additional sintering experiments (samples 12 5 5 mm3 ) were performed in a differential dilatometer (Setaram, Setsys 16/ 18, France) with alumina detector reproducing the same thermal treatment schedule as that used to obtain the final materials in order to determine strain during cooling. The sintered blocks of the monoliths were machined into bars of 50 3 4 mm3 for dynamic Young’s modulus determinations, from the resonance frequency of the bars in flexure (Grindosonic, J.W. Lemmens, Belgium) and bend strength tests (4-points bending, 40–20 mm span, 0.5 mm min1 ; Microtest, Spain). Reported Young’s modulus and bend strength values are the average of five measurements and errors are the standard deviations. A symmetric laminated structure of five layers was fabricated by casting each suspension alternately. Casting times were fixed to reach the desired layer thickness considering the casting kinetics and sintering shrinkage of each composition. The laminate had the central (1200 lm) and outer layers (ffi2100 lm) made of A10 and the two inner layers (ffi300 lm) of A30. Microstructure of polished cross sections was characterised in a field emission gun scanning electron microscopy (FEG-SEM, Hitachi, S-4700, Japan). Thermally etched (1440 C–1min) specimens were analysed. Additional observations were performed on chemically etched (HF 10 vol.%–1 min) specimens in order to assure that the thermal etching did not produce further microcracking in the sintered specimens. Single-Edge-V-Notch-Beams (SEVNB) of 4 6 50 mm3 , machined from the laminated sintered blocks, were tested in a 3-points bending device using a span of 40 mm and a cross-head speed of 0.005 mm min1 (Microtest, Spain). The notches were initially cut with a 150 lm wide diamond wheel. Using this slot as a guide, the remaining part of the notch was done with a razor blade sprinkled with diamond pastes of successively 6 and 1 lm. Specimens with relative notch depths of about 0.8 of the thickness of the first external layer, corresponding to 0.26 of the specimen thickness (W), were tested. The tip radii of all notches were determined from optical observations and they were always found to be below 20 lm. The curves load–displacement of the cross-head of the load frame were recorded and corrected by subtracting the compliance of the testing set up (machine, supports, load cell and fixtures, 1.5 107 m/N) determined by testing a thick (25 25 100 mm3 ) unnotched alumina bar. Three specimens were tested and the curves were found to be practically coincident. 138 S. Bueno, C. Baudín / Composites: Part A 40 (2009) 137–143
S. Bueno, C Baudin/Composites: Part A 40(2009)137-143 The fracture toughness parameters, i. e critical stress intensity fac- heating were coincident at 1240C (Fig. la)and the sintering rates tor, KIc, and work of fracture, wwo, were calculated from the curves were coincident at 1150C(Fig. 1b). A significant slope change obtained during the SEvNb tests. Optical and scanning electron occurred at about 1380C in agreement with the reported temper microscopy observations were performed on the fracture surfaces. ature for the expansive reaction between Al2O3 and TiO to form Additional laminated plates of 70 x 70 x 0.6 mm were fabri- aluminium titanate[20]. cated to get bars(50 x 3x 4 mm) for strength testing. In these From these curves(fig. 1). a two-step sintering treatment, with samples, the relationship between the width of the a10 and A30 a rather low heating rate 2C min-, was designed to favour co- layers (tAiotA30 5) was slightly higher than that between the sintering of A10 and A30 layers in the laminated structure. An ini- internal layers in the thick specimens(talo tA302 4) tial dwell of 4 h at 1200C was chosen for homogeneous shrinkage before reaction, this temperature being a compromise between 3. Results and discussion those for coincident levels of shrinkage, 1240C, and of shrinkage ate, 1150C.a2 h dwell at 1450C was selected for final sinte 3. 1. Laminate design and processin ing, as this was the temperature at which shrinkage was almost ar- rested in both composites. This relatively low temperature for final During slip casting of suspensions of both compositions, the sintering assured the control of grain growth well-known proportionality between the square of wall growth The properties of the monolithic materials of compositions A10 and the slip time was found. Casting kinetics of 0.9 and 1.2 mm2/ and A30 are summarised in Table 1. Strength values for composite min for the A10 and A30 slips were determined As described else- A10 are lower than those previously reported for small specimens where [29]. for a fixed wall thickness casting time shortens with tested in 3-point bending (of=261 6 MPa, 3-points bending. titania content samples 2x 2.5 x 30 mm, span 20 mm [16])as expected for the In Fig. 1, dynamic sintering curves for alumina and both studied experimental method used here. In the same sense, the dynamic omposites are plotted. Shrinkage was initially retarded for the Youngs modulus value determined in work is higher than tha composites with respect to that of pure alumina due to the pres- for the static one previously reported(201 t 1 GPa [16). ence of a second phase and then it was accelerated between laminated structure, the composition with the highest st 1200 and 1400C. For the composites, the shrinkage levels during was chosen for the external layers whereas that of the work of fracture was chosen for the internal ones To design the symmetric structure with five layers the level a sign of the expected residual stresses was analysed and the thick nesses of the layers were chosen in order to minimise tensile resid- 0.00 ual stresses. The residual stresses at the centre of the layers of a A10 symmetrical laminate can be evaluated using the simplified model A30 of a symmetric plate constituted by alternate layers of the same thickness having a uniform biaxial distribution of stresses across -0.04 each layer [30 Using this approach, the residual stresses at the -0.06 centre of the composite A10 and A30 layers, are given by -0.08 (1) -0.10 where AE is the thermal expansion mismatch between the layers, E -0.14 is the Youngs modulus and nA10, A30, tA10, A30 are the number of lay ers and their thickness The actual difference between the dimensional variations of the layers during cooling(AE, Eq.(1),(2)) was determined from the b0.0001 cooling part of the sintering curves of the monoliths shown in Fig. 2, as described previously for alumina-aluminium titanate A10 composites 31. In alumina-based materials, deformation mis- match at temperatures higher than 1200C can be accommodated by diffusion [32], whereas, from 1200C to room temperature, this mismatch originates stresses. Therefore, in order to evaluate the 右0002 stress level and sign in the laminate, the differences between the dimensional variations of 0.0003 "stress free"temperature, 1200C, were analysed In Fig 2b, these variations are plotted together for both monoliths, after correction for coincident dimensions at 1200oC. in the considered interval of Table 1 -0.0006 Properties of the monolithic materials used to design the laminated structure 1400 the laminate material Work of fracture (m-2) Young,'s modulus(GPa) Strength(MPa) 35±3I6 359±5 210±1 Fig. of the The curve of A30 pure alumina(A)is shown for comparative purposes. (a) 53±4I15 225±5 temperature.(b) Linear shrinkage rate, d(alLo)ldT, versus temperatur rinkage,△|lA10A3062±3
The fracture toughness parameters, i.e. critical stress intensity factor, KIC, and work of fracture, cWOF, were calculated from the curves obtained during the SEVNB tests. Optical and scanning electron microscopy observations were performed on the fracture surfaces. Additional laminated plates of 70 70 0.6 mm3 were fabricated to get bars (50 34 mm3 ) for strength testing. In these samples, the relationship between the width of the A10 and A30 layers (tA10/tA30 5) was slightly higher than that between the internal layers in the thick specimens (tA10/tA30 4). 3. Results and discussion 3.1. Laminate design and processing During slip casting of suspensions of both compositions, the well-known proportionality between the square of wall growth and the slip time was found. Casting kinetics of 0.9 and 1.2 mm2 / min for the A10 and A30 slips were determined. As described elsewhere [29], for a fixed wall thickness casting time shortens with titania content. In Fig. 1, dynamic sintering curves for alumina and both studied composites are plotted. Shrinkage was initially retarded for the composites with respect to that of pure alumina due to the presence of a second phase and then it was accelerated between 1200 and 1400 C. For the composites, the shrinkage levels during heating were coincident at 1240 C (Fig. 1a) and the sintering rates were coincident at 1150 C (Fig. 1b). A significant slope change occurred at about 1380 C in agreement with the reported temperature for the expansive reaction between Al2O3 and TiO2 to form aluminium titanate [20]. From these curves (Fig. 1), a two-step sintering treatment, with a rather low heating rate 2 C min1 , was designed to favour cosintering of A10 and A30 layers in the laminated structure. An initial dwell of 4 h at 1200 C was chosen for homogeneous shrinkage before reaction, this temperature being a compromise between those for coincident levels of shrinkage, 1240 C, and of shrinkage rate, 1150 C. A 2 h dwell at 1450 C was selected for final sintering, as this was the temperature at which shrinkage was almost arrested in both composites. This relatively low temperature for final sintering assured the control of grain growth. The properties of the monolithic materials of compositions A10 and A30 are summarised in Table 1. Strength values for composite A10 are lower than those previously reported for small specimens tested in 3-point bending (rf = 261 ± 6 MPa, 3-points bending, samples 2 2.5 30 mm3 , span 20 mm [16]) as expected for the experimental method used here. In the same sense, the dynamic Young’s modulus value determined in this work is higher than that for the static one previously reported (201 ± 1 GPa [16]). For the laminated structure, the composition with the highest strength was chosen for the external layers whereas that of the highest work of fracture was chosen for the internal ones. To design the symmetric structure with five layers the level and sign of the expected residual stresses was analysed and the thicknesses of the layers were chosen in order to minimise tensile residual stresses. The residual stresses at the centre of the layers of a symmetrical laminate can be evaluated using the simplified model of a symmetric plate constituted by alternate layers of the same thickness having a uniform biaxial distribution of stresses across each layer [30]. Using this approach, the residual stresses at the centre of the composite A10 and A30 layers, are given by: rA10 ¼ DeEA10 1 þ E0 A10nA10tA10 E0 A30nA30tA30 ð1Þ rA30 ¼ rA10 nA10 nA30 tA10 tA30 ð2Þ where De is the thermal expansion mismatch between the layers, E is the Young’s modulus and nA10, A30, tA10, A30 are the number of layers and their thickness. The actual difference between the dimensional variations of the layers during cooling (De, Eq. (1), (2)) was determined from the cooling part of the sintering curves of the monoliths shown in Fig. 2, as described previously for alumina–aluminium titanate composites [31]. In alumina-based materials, deformation mismatch at temperatures higher than 1200 C can be accommodated by diffusion [32], whereas, from 1200 C to room temperature, this mismatch originates stresses. Therefore, in order to evaluate the stress level and sign in the laminate, the differences between the dimensional variations of specimens during cooling from the ‘‘stress free” temperature, 1200 C, were analysed. In Fig. 2b, these variations are plotted together for both monoliths, after correction for coincident dimensions at 1200 C. In the considered interval of 800 1000 1200 1400 1600 -0.0006 -0.0005 -0.0004 -0.0003 -0.0002 -0.0001 0.0000 0.0001 d( ΔL/Lo)/dT [ºC-1] T [ºC] A A10 A30 800 1000 1200 1400 1600 -0.14 -0.12 -0.10 -0.08 -0.06 -0.04 -0.02 0.00 0.02 A A10 A30 ΔL/Lo T [ºC] Fig. 1. Dynamic sintering curves of the composites (A10 and A30). The curve of pure alumina (A) is shown for comparative purposes. (a) Linear shrinkage, DL/L0, versus temperature. (b) Linear shrinkage rate, d(DL/L0)/dT, versus temperature. Table 1 Properties of the monolithic materials used to design the laminated structure and of the laminate material. Work of fracture (Jm2 ) Young’s modulus (GPa) Strength (MPa) A10 35 ± 3 [16] 359 ± 5 210 ± 10 A30 53 ± 4 [15] 225 ± 5 136 ± 10 A10A30 62 ± 3 – 160 ± 5 S. Bueno, C. Baudín / Composites: Part A 40 (2009) 137–143 139
140 S Bueno, C Baudin/Composites: Part A 40(2009)137-143 0.00 liths Dynamic Young,'s modulus values from Table 1 were used for calculations which, being always higher than the static ones, would also lead to conservative values of the residual stresses Tensile -A30 stresses of about 15 MPa and e stresses of about -0.06 90 MPa were expected in the A10 and A30 layers, respectively In the case of the small bend strength specimens(50 x 3 x 4 mm). vith a relationship between the width of the central Al0 layer and A30 layers tAlo/tA30 5, the expected residual stresses would 0.12 be similar; tensile stresses of about 12 MPa and compressive stres ses of about 90 MPa were calculated -0.15 3.2. Microstructure In Table 2 the desired thicknesses according to Fig. 3 and the 1000 casting times required for the processing of the laminate plates by sequential slip casting are summarised. The obtained thick nesses, determined directly in the SEM, were systematically larger b0.0000 than the desired ones which could be explained by the relatively higher efficiency for filtration of the large plaster moulds used to -0.0005 cast the plates as compared to that of the small ones used to fab- ricate the disc specimens. Fig 4 shows characteristic microstructures of the specimens. At 000 ow magnification( Fig 4a). long microcracks(200 um)with large A10 crack opening displacement(Fig 4 b)were observed. These micI A30 cracks were not randomly oriented as occurred in the monolithic -0.0015 A30 material taken as a reference [15 but preferentially oriented parallel to the interfaces between the layers. This orientation can 00020 be explained taking into account that the internal layers in the laminate will be subjected to compression during cooling from sin- tering. The Poisson effect will lead to the development of tensile stresses perpendicular to the interfaces inside the internal layers 020040060080010001200 which, added to the local stresses developed at the grain bound- TrC] aries due to thermal expansion mismatch, will produce the forma- tion of the observed microcracks. The fact that edge cracks [33, 34 of the ty were not observed in this laminate in which internal compressive schedule: 2C/min, 4 h at residual stresses develop during cooling indicates that there was a b) Cooling from1200° significant release of residual stresses due to microcrack formation. 33 fracture behaviour temperatures (1200-50C)the shrinkage for the Al0 composit (e19.7 x 10)was slightly larger than that of A30 Fig. 5 shows a characteristic corrected load-displacement (e16.3 x 10-), which would imply tensile residual stresses in curve; the load increased linearly with the displacement of the the central and external high strength A10 layers and compressive testing frame up to a maximum value(150 N)after which a sud residual stresses in the internal A30 layers, respectively den load drop occurred. The initial linearity indicates that the ini- Taking into account that, according to Eq (1),(2), the minimisa- tial deformation behaviour was controlled by the linear and stiff tion of the tensile residual stresses of the external A10 layers external and central layers until the initiation of crack propagation would require relatively thin A30 internal layers and wider A10 at Koapp of e4.1 MPa m 2. The most interesting feature of the yers, the laminated structure shown in Fig 3 was selected a rela- curves was that the load drop was arrested at about 55 N and then, onship between the width of the central A10 layer and A30 layers the specimens admitted further deformation without tAlo/tA30 4 In order to simplify, a conservative model with the the load up to their complete failure. three a10 layers having the same width as the central one in the In Fig. 6 the fracture modes of the different layers are shown. All laminate(e1200 um, Fig 3)can be considered to calculate a the layers of composition A30 presented different fracture planes residual stresses using Eq (1).(2)and the properties of the mono- forming steps that followed the orientation of the pre-existing 2100um 1200um 4300um 300um Fig 3. Schematic illustration of the designed five layered laminated sti showing a bend bar and the notch orientation with respect to the layers. The thin A30 internal composite layers are ented with white colour and have a thickness of e300 um. The thic
temperatures (1200–50 C) the shrinkage for the A10 composite (ffi19.7 104 ) was slightly larger than that of A30 (ffi16.3 104 ), which would imply tensile residual stresses in the central and external high strength A10 layers and compressive residual stresses in the internal A30 layers, respectively. Taking into account that, according to Eq. (1), (2), the minimisation of the tensile residual stresses of the external A10 layers would require relatively thin A30 internal layers and wider A10 layers, the laminated structure shown in Fig. 3 was selected a relationship between the width of the central A10 layer and A30 layers tA10/tA30 ffi 4. In order to simplify, a conservative model with the three A10 layers having the same width as the central one in the laminate (ffi1200 lm, Fig. 3) can be considered to calculate a the residual stresses using Eq. (1), (2) and the properties of the monoliths. Dynamic Young’s modulus values from Table 1 were used for calculations which, being always higher than the static ones, would also lead to conservative values of the residual stresses. Tensile stresses of about 15 MPa and compressive stresses of about 90 MPa were expected in the A10 and A30 layers, respectively. In the case of the small bend strength specimens (50 3 4 mm3 ), with a relationship between the width of the central A10 layer and A30 layers tA10/tA30 ffi 5, the expected residual stresses would be similar; tensile stresses of about 12 MPa and compressive stresses of about 90 MPa were calculated. 3.2. Microstructure In Table 2 the desired thicknesses according to Fig. 3 and the casting times required for the processing of the laminate plates by sequential slip casting are summarised. The obtained thicknesses, determined directly in the SEM, were systematically larger than the desired ones which could be explained by the relatively higher efficiency for filtration of the large plaster moulds used to cast the plates as compared to that of the small ones used to fabricate the disc specimens. Fig. 4 shows characteristic microstructures of the specimens. At low magnification (Fig. 4a), long microcracks (200 lm) with large crack opening displacement (Fig. 4 b) were observed. These microcracks were not randomly oriented as occurred in the monolithic A30 material taken as a reference [15] but preferentially oriented parallel to the interfaces between the layers. This orientation can be explained taking into account that the internal layers in the laminate will be subjected to compression during cooling from sintering. The Poisson effect will lead to the development of tensile stresses perpendicular to the interfaces inside the internal layers which, added to the local stresses developed at the grain boundaries due to thermal expansion mismatch, will produce the formation of the observed microcracks. The fact that edge cracks [33,34] were not observed in this laminate in which internal compressive residual stresses develop during cooling indicates that there was a significant release of residual stresses due to microcrack formation. 3.3. Fracture behaviour Fig. 5 shows a characteristic corrected load–displacement curve; the load increased linearly with the displacement of the testing frame up to a maximum value (150 N) after which a sudden load drop occurred. The initial linearity indicates that the initial deformation behaviour was controlled by the linear and stiff external and central layers until the initiation of crack propagation at K0app of ffi4.1 MPa m1/2. The most interesting feature of the curves was that the load drop was arrested at about 55 N and then, the specimens admitted further deformation without increasing the load up to their complete failure. In Fig. 6 the fracture modes of the different layers are shown. All layers of composition A30 presented different fracture planes forming steps that followed the orientation of the pre-existing 800 -0.18 -0.15 -0.12 -0.09 -0.06 -0.03 0.00 ΔL/L0 T [ºC] 0 -0.0025 -0.0020 -0.0015 -0.0010 -0.0005 0.0000 ΔL/L0 T [ºC] A10 A30 1000 1200 1400 1600 200 400 600 800 1000 1200 A10 A30 Fig. 2. Dilatometric curves (DL/L0 = length variation) for green compacts of the two studied monoliths heat treated following the sintering schedule: 2 C/min, 4 h at 1200 C and 2h at 1450 C. (a) Complete thermal cycle. (b) Cooling from 1200 C, after correction for coincident dimensions at 1200 C. 1200 μm 2100 μm 2100 μm 300 μm 300 μm Fig. 3. Schematic illustration of the designed five layered laminated structure showing a bend bar and the notch orientation with respect to the layers. The thin A30 internal composite layers are represented with white colour and have a thickness of ffi300 lm. The thick external (ffi2100 lm) and central (ffi1200 lm) A10 layers are represented with grey colour. 140 S. Bueno, C. Baudín / Composites: Part A 40 (2009) 137–143
S. Bueno, C Baudin/Composites: Part A 40(2009)137-143 2 Casting times corresponding to the designed layer thicknesses of the laminated Measured thicknesses of layers 1 and 5 are not shown because the specimens were machined before polishing. igned layer Casting time s Measured layer 3 2100 1320-1360 360-375 A10-layer 5 A10a 0.01 0.02 0.03 0.04 d [mm] Fig. 5. Characteristic load-displacement curve of specimens notched with a relative notch length (a/w) of 0. 26. Semistable behaviour is shown with a sudden load decrease down to approximately 54 N. A30 A10 A10 心 A10 Fig. 6. Characteristic fracture features of the laminated specimens. FE-SEM micrograph of a fracture surface. The internal A30 layer shows different fracture planes forming steps that followed the orientation of the pre-existing microcrcaks the point of impingement into the crack to traverse the remaining part of the first A30 and the second Al0 layers. When the crack cemal A30 layers in the laminated reaches the second A 30 layer(Fig. 7b) it impinges also a pre-exist thermally etched surface (a)Low ing microcrack but in this case multiple cracks at different the interfaces between ks preferentially oriented parallel to distances of the point of impingement. Moreover, the same process of a characteristic microcrack. occurs with the emerging cracks. Thus, the microcracks present in the laminate play a role similar to that of the weak interfaces microcracks, which indicates that the main crack was deflected between layers in the crack-deflecting laminates [ 9, 28, 35-39 nd or branched along them; fracture of the A10 layers was flat while the delaminaton distances are limited to the microcrack relatively to that of the A30 ones. In Fig. 7 the crack paths in the lengths. The fact that no deflection along the central part of the lateral surfaces are shown. The propagation of the main cracks internal layers, were the residual compressive stresses would be as never completely straight due to the above mentioned crack the largest, was observed further support their low level. deflection and or branching processes, nevertheless, significant The aspect of the load-displacement curves can be inferred differences were found depending on the particular layer of the from the fracture features just discussed( Figs. 6 and 7). The initial laminate. The deflection distances through the first A10 layers ad drop observed, corresponding to the kinetic pr and part of the first A30 ones were very small and multiple crack- the main crack, will be associated to its propagation to reach the ing was not observed whereas significant deflection distances and second A30 layer. The observed crack deflection in the first A30 multiple cracking were observed in the other layers. Fig. yer(Fig. 7a) would not be enough as to produce significant strain how a main crack reaches a pre-existing microcrack, an in the specimens. Conversely, the multiple deflection and forma be arrested by it emerging at a certain distance(80- tion of secondary cracks perpendicular to the layers in the second
microcracks, which indicates that the main crack was deflected and/or branched along them; fracture of the A10 layers was flat relatively to that of the A30 ones. In Fig. 7 the crack paths in the lateral surfaces are shown. The propagation of the main cracks was never completely straight due to the above mentioned crack deflection and/or branching processes, nevertheless, significant differences were found depending on the particular layer of the laminate. The deflection distances through the first A10 layers and part of the first A30 ones were very small and multiple cracking was not observed whereas significant deflection distances and multiple cracking were observed in the other layers. Fig. 7a shows how a main crack reaches a pre-existing microcrack, and seems to be arrested by it emerging at a certain distance (80–100 lm) of the point of impingement into the crack to traverse the remaining part of the first A30 and the second A10 layers. When the crack reaches the second A30 layer (Fig. 7b) it impinges also a pre-existing microcrack but in this case multiple cracks emerge at different distances of the point of impingement. Moreover, the same process occurs with the emerging cracks. Thus, the microcracks present in the laminate play a role similar to that of the weak interfaces between layers in the crack-deflecting laminates [9,28,35–39] while the delaminaton distances are limited to the microcrack lengths. The fact that no deflection along the central part of the internal layers, were the residual compressive stresses would be the largest, was observed further support their low level. The aspect of the load–displacement curves can be inferred from the fracture features just discussed (Figs. 6 and 7). The initial load drop observed, corresponding to the kinetic propagation of the main crack, will be associated to its propagation to reach the second A30 layer. The observed crack deflection in the first A30 layer (Fig. 7a) would not be enough as to produce significant strain in the specimens. Conversely, the multiple deflection and formation of secondary cracks perpendicular to the layers in the second Table 2 Casting times corresponding to the designed layer thicknesses of the laminated structure described in Fig. 3 and measured thicknesses in the sintered materials. Measured thicknesses of layers 1 and 5 are not shown because the specimens were machined before polishing. Designed layer thickness, lm Casting time, s Measured layer thickness, lm A10–layer 1 2100 415 – A30–layer 2 300 90 350–370 A10–layer 3 1200 640 1320–1360 A30–layer 4 300 150 360–375 A10–layer 5 2100 1930 – Fig. 4. Characteristic microstructure of the internal A30 layers in the laminated structure. FE-SEM micrograph of a polished and thermally etched surface. (a) Low magnification observation showing microcracks preferentially oriented parallel to the interfaces between the layers. (b) Detail of a characteristic microcrack. 0.00 0 25 50 75 100 125 150 175 200 P [N] d [mm] 0.01 0.02 0.03 0.04 Fig. 5. Characteristic load–displacement curve of specimens notched with a relative notch length (a/W) of 0.26. Semistable behaviour is shown with a sudden load decrease down to approximately 54 N. Fig. 6. Characteristic fracture features of the laminated specimens. FE-SEM micrograph of a fracture surface. The internal A30 layer shows different fracture planes forming steps that followed the orientation of the pre-existing microcrcaks (Fig. 4a). S. Bueno, C. Baudín / Composites: Part A 40 (2009) 137–143 141
S Bueno, C. Baudin/Composites: Part A 40(2009)137-143 ondary cracks perpendicular to the layers in the second A30 layer a (Fig. 7b). From such load value(55 N, Fig. 5), the failure stress for A10题 the discerned formation of secondary cracks was calculated from the stress distribution under 3-points bending on a prismatic bar formed by layers with different elastic properties 40 by assuming a propagating crack with a length equal to the location of the cen- tre of the second A30 layer(a/w=0.76: Fig. 7b). From the failure stress, and taking into account a general stress intensity formula- This value is slightly lower than that reported for the most per forming crack-deflecting laminates(15-18 MPa m"/[9, 28, 35-39) constituted by Sic and graphite. Nevertheless, it is comparable to values reported for the stationary state of transformation-tough- ened ceramics(=6 MPa m"/[42]. 9-12 MPa m/2[43)), which are A10 the toughest oxides. 4100um The work of fracture value calculated from the area under the semistable fracture curves(Fig. 5. 62 +3 Jm higher(=26%) than that obtained by calculation, taking into ac count the additive character of the work of fracture from the work aob of fracture values of monolithic materials of composition Al0 same way, and the surface fraction of the crack corresponding to each layer(S10≈0.83ands3o≈0.17,ands10≈0.6,530≈0.4for specimens with notch lengths of 0. 4 and 0.8, respectively). This fact A30 reveals a synergic effect of the laminated structure on the mechan ical behaviour of the material The most important difference of the laminate posed here and that of other laminates with high capability for crack deflection is that the crack deflection and branching pro- cesses occur at local level, as demonstrated by the micrograph of fractured samples in Figs. 5 and 6. As a consequence, the new de- sign for ceramic laminates proposed allows reaching high apparent A10 toughness and work of fracture while maintaining the structu integrity of the piece after the initiation of crack propagation under 100 shear stresses as those that develop in wear applications. Further improvements of the proposed structure will be reached by lami- nated designs with larger numbers of thinner layers to originate linate. Polished later. of the samples. The sense of crack propagation was from the bottom to graceful fracture he micrographs (a) Crack branching and deflection of the main crack due to its interaction with the pre-existing microcracks ar Acknowledgments rrow shows the point at which the new crack is started. (b)Crack branching and multiple cracking are shown. The authors acknowledge the support of the Project MEC MAT2006-13480C02-01, Spain. A30 layer(Fig. 7b)will allow the specimens to deform without References increasing the load, as observed in the load-displacement curves before the complete failure. [1 Clegg W] Design of ceramics laminates for structural applications. Mater Sci Strength values are summarised in Table 1. The strength values for the layered material were lower than those of the A10 monolith [21 Gunnison KE, Sarikaya M, Liu J, Aksay IA Structure-mechanical property which can not be explained at this point, as the expected residual Baer e Sariaya M, editors. Hie ally structured materials materials stresses were very low and, moreover, they should have been par research society. Pittsburgh, Penn. USA: DA Tirrel Editors: 1992. tially released as discussed above. Special care should be taken in [3 Aksay IA, Sarikaya M. Bioinspired proces composite materials. In: Soga order to improve the quality of the microstructure of the external amics toward the 21st century. Tokyo: The Ceramic layers in this kind of structures. Nevertheless, strength values of [4] Hiltner A, Sung K, Shin E, Bazhenov S, Im J. Baer E Polymer microlayer the laminate were higher than those of the monolith of the same composites. In: Aksay IA, Baer E, Sarikaya M, editors. Hierarchically structured composition as that of the internal A30 layers while presenting a laterals. Materials research society. Pittsburgh, Penn. USA: DA Tirrel rong interfaces. J Am Ceram Soc 1995: 78(4): 1125-7. 4. Concluding remarks [61 Sakai M, Bradt RC Graphical methods for determining the nonlinear fracture From the load-displacement curves obtained and the observed fracture paths it is possible to calculate the apparent [7 Lutz EH, Brunings St E, Steinbrech RW. Steel-re ma ceramics. A the point of failure of the studied laminate [9, 28 was arrested at about 55 N, and then, the specimens 如~邮 [81 Rao MP. Sanchez-Herencia [91 Clegg W]. Kendall K, Alford NMcN, Birchall JD, Button Tw. A simple way to ther deformation due to multiple deflection and forn make tough ceramics. Nature 1990: 347: 45-57
A30 layer (Fig. 7b) will allow the specimens to deform without increasing the load, as observed in the load–displacement curves before the complete failure. Strength values are summarised in Table 1. The strength values for the layered material were lower than those of the A10 monolith which can not be explained at this point, as the expected residual stresses were very low and, moreover, they should have been partially released as discussed above. Special care should be taken in order to improve the quality of the microstructure of the external layers in this kind of structures. Nevertheless, strength values of the laminate were higher than those of the monolith of the same composition as that of the internal A30 layers while presenting a significantly higher toughness. 4. Concluding remarks From the load–displacement curves obtained and the observed fracture paths it is possible to calculate the apparent toughness at the point of failure of the studied laminate [9,28]. As previously discussed for the load–displacement curve in Fig. 5, the load drop was arrested at about 55 N, and then, the specimens admitted further deformation due to multiple deflection and formation of secondary cracks perpendicular to the layers in the second A30 layer (Fig. 7b). From such load value (55 N, Fig. 5), the failure stress for the discerned formation of secondary cracks was calculated from the stress distribution under 3-points bending on a prismatic bar formed by layers with different elastic properties [40] by assuming a propagating crack with a length equal to the location of the centre of the second A30 layer (a/W = 0.76; Fig. 7b). From the failure stress, and taking into account a general stress intensity formulation [41], an apparent toughness of ffi12 MPa m1/2 was calculated. This value is slightly lower than that reported for the most performing crack-deflecting laminates (15–18 MPa m1/2 [9,28,35–39]) constituted by SiC and graphite. Nevertheless, it is comparable to values reported for the stationary state of transformation-toughened ceramics (6 MPa m1/2 [42], 9–12 MPa m1/2 [43]), which are the toughest oxides. The work of fracture value calculated from the area under the semistable fracture curves (Fig. 5, 62 ± 3 Jm2 ) was significantly higher (26%) than that obtained by calculation, taking into account the additive character of the work of fracture, from the work of fracture values of monolithic materials of composition A10 (35 ± 3 Jm2 [16]) and A30 (53 ± 4 Jm2 [15]) processed in the same way, and the surface fraction of the crack corresponding to each layer (s10 0.83 and s30 0.17, and s10 0.86, s30 0.14 for specimens with notch lengths of 0.4 and 0.8, respectively). This fact reveals a synergic effect of the laminated structure on the mechanical behaviour of the material. The most important difference of the laminated structure proposed here and that of other laminates with high capability for crack deflection is that the crack deflection and branching processes occur at local level, as demonstrated by the micrographs of fractured samples in Figs. 5 and 6. As a consequence, the new design for ceramic laminates proposed allows reaching high apparent toughness and work of fracture while maintaining the structural integrity of the piece after the initiation of crack propagation under shear stresses as those that develop in wear applications. Further improvements of the proposed structure will be reached by laminated designs with larger numbers of thinner layers to originate graceful fracture. Acknowledgments The authors acknowledge the support of the Project MEC MAT2006-13480 C02-01, Spain. References [1] Clegg WJ. Design of ceramics laminates for structural applications. Mater Sci Technol 1998;14(6):483–95. [2] Gunnison KE, Sarikaya M, Liu J, Aksay IA. Structure-mechanical property relationships in a biological ceramic-polymer composite: nacre. In: Aksay IA, Baer E, Sarikaya M, editors. Hierarchically structured materials materials research society. Pittsburgh, Penn. USA: DA Tirrel Editors; 1992. [3] Aksay IA, Sarikaya M. Bioinspired processing of composite materials. In: Soga N, Kato A, editors. Ceramics toward the 21st century. Tokyo: The Ceramic Society of Japan; 1991. p. 136–49. [4] Hiltner A, Sung K, Shin E, Bazhenov S, Im J, Baer E. Polymer microlayer composites. In: Aksay IA, Baer E, Sarikaya M, editors. Hierarchically structured materials. Materials research society. Pittsburgh, Penn. USA: DA Tirrel Editores; 1992. p. 141–50. [5] Prakash O, Sarkar P, Nicholson P. Crack deflection in ceramic/ceramic laminates with strong interfaces. J Am Ceram Soc 1995;78(4):1125–7. [6] Sakai M, Bradt RC. Graphical methods for determining the nonlinear fracture parameters of silica and graphite refractory composites. In: Bradt RC, Evans AG, Hasselman DPH, Lange FF, editors. Fracture mechanics of ceramics, vol. 7. New York: Plenum; 1986. p. 127–42. [7] Lutz EH, Brunings St E, Steinbrech RW. Steel-reinforced plasma ceramics. A new multilayer design. Ceram Eng Sci Proc 1998;19(3):457–65. [8] Rao MP, Sánchez-Herencia AJ, Beltz GE, McMeeking RM, Lange FF. Laminar ceramics that exhibit a threshold strength. Science 1999;286(5437):102–5. [9] Clegg WJ, Kendall K, Alford NMcN, Birchall JD, Button TW. A simple way to make tough ceramics. Nature 1990;347:45–57. Fig. 7. Characteristic crack paths in the studied laminate. Polished lateral surfaces of the samples. The sense of crack propagation was from the bottom to the top of the micrographs. (a) Crack branching and deflection of the main crack in the first A30 layer due to its interaction with the pre-existing microcracks are shown. The arrow shows the point at which the new crack is started. (b) Crack branching and multiple cracking are shown. 142 S. 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