Availableonlineatwww.sciencedirect.com Science Direct E噩≈RS ELSEVIER Joumal of the European Ceramic Society 28(2008)1961-1971 www.elsevier.comlocate/jeurceramsoc fracture behaviour of microcrack -free alumina-aluminium titanate ceramics with second phase nanoparticles at alumina grain boundaries S Bueno a m.H. Berger b.R. Moreno a c.Baudin a a Instituto de Ceramica y Vidrio(CSIC). C. Kelsen 5, 28049 Madrid, Spain Ecole des Mines de Paris, Centre des Materiaur, 91003 Evry Cedex, france Received 12 October 2007: received in revised form 19 December 2007; accepted 4 January 2008 Available online 4 March 2008 Alumina+10 vol %o aluminium titanate composites were obtained by colloidal filtration and reaction sintering of alumina and titania. The materials were dense with aluminium titanate grains of average sizes 2.2-2. 4 um located mainly at alumina triple points. The reaction sintering schedule promoted the formation of additional nanometric grains, identified as aluminium titanate using STEM-EDX analysis between the alumina grains. This special microstructure led to a change of the toughening mechanism from the typical crack bridging reported for microcrack-free composites fabricated from alumina and aluminium titanate powders to microcracking The identification of microcracking as the main toughening mechanism was done from the analysis of stable fracture tests of SENVB samples in three points bending and fractographic observations. Monophase alumina materials with similar grain sizes were used as referer Different fracture toughness parameters were derived from the load-displacement curves: the critical stress intensity factor, KiC, the critical energy release rate, Gic, the J-Integral and the work of fracture, ywoF, and the R curves were also built. The comparison between the linear elastic fracture arameters and the non-linear ones revealed significant toughening and faw tolerance o 2008 Elsevier Ltd. all rights reserved. Keywords: D. Al2O3; D. Al2 TiOs; C Mechanical properties: C. Toughening: B. Nanocomposites Introduction 10-6°C-1,a25-100c=-2:7×10-6°C-)13and mina shows limited anisotropy(aa25-10000C=8.4 x 10-6 6oc-I The use of ceramic materials in structural applications is ac25-10000C=9.2 x 10-6oC-), 4 thus, high tensile or com- limited by the"ilaw sensitive"fracture, occurring sponta- pressive stresses, depending on the particular crystallographic neously from natural flaws, inherent to the brittle behaviour. orientation of the grains, would develop during cooling from The"flaw tolerance"approach deals with the development of the sintering temperature at the grain-matrix interfaces due to microstructures that originate toughening mechanisms to reduce thermal expansion mismatch. Depending on grain size and the the sensitivity of the strength to the size of any processing or characteristics of the grain boundaries, microcracking might induced flaw, thus improving the reliability of the materials. occur during cooling from sintering and/or during fracture. Such mechanisms originate an increasing resistance with con- In the early 90s alumina-aluminium titanate composites tinued crack extension, rising R-curve behaviour, and most of with aluminium titanate contents 20-30 vol% obtained from them are caused by localized internal residual stresses in the alumina and aluminium titanate mixtures, were studied by materials other authors. Crack bridging by second phase agglomerates Alumina(Al2O3-aluminium titanate(Al2TiO5) materials nd by lar was e n offer improved flaw tolerance and toughness 4-12 Ther- ing mechanism leading to R-curve behaviour, assessed by hal expansion of aluminium titanate is highly anisotr- the indentation-strength method; no toughness values were opic(a25-100c010.9×10-60C-1,ab25-100c20.5× Corresponding author. Tel: +3491 7355840: fax: +3491 7355843 In this work. B-Al2TiOs orthorhombic crystal is described by a b-face cen- E-mail address: cbaudin @icv csices(C. Baudin) tered unit cell, space group Bbmm, a=9.439A, b=9.647A c=3.593A 0955-2219/S-see front matter o 2008 Elsevier Ltd. All rights reserved. doi: 10.1016/j-jeurceramsoc 2008.01.01
Available online at www.sciencedirect.com Journal of the European Ceramic Society 28 (2008) 1961–1971 Fracture behaviour of microcrack-free alumina–aluminium titanate ceramics with second phase nanoparticles at alumina grain boundaries S. Bueno a, M.H. Berger b, R. Moreno a, C. Baud´ın a,∗ a Instituto de Cer ´amica y Vidrio (CSIC). C. Kelsen 5, 28049 Madrid, Spain b Ecole des Mines de Paris, Centre des Mat´eriaux, 91003 Evry Cedex, France Received 12 October 2007; received in revised form 19 December 2007; accepted 4 January 2008 Available online 4 March 2008 Abstract Alumina + 10 vol.% aluminium titanate composites were obtained by colloidal filtration and reaction sintering of alumina and titania. The materials were dense with aluminium titanate grains of average sizes 2.2–2.4 m located mainly at alumina triple points. The reaction sintering schedule promoted the formation of additional nanometric grains, identified as aluminium titanate using STEM–EDX analysis between the alumina grains. This special microstructure led to a change of the toughening mechanism from the typical crack bridging reported for microcrack-free composites fabricated from alumina and aluminium titanate powders to microcracking. The identification of microcracking as the main toughening mechanism was done from the analysis of stable fracture tests of SENVB samples in three points bending and fractographic observations. Monophase alumina materials with similar grain sizes were used as reference. Different fracture toughness parameters were derived from the load–displacement curves: the critical stress intensity factor, KIC, the critical energy release rate, GIC, the J-Integral and the work of fracture, γWOF, and the R curves were also built. The comparison between the linear elastic fracture parameters and the non-linear ones revealed significant toughening and flaw tolerance. © 2008 Elsevier Ltd. All rights reserved. Keywords: D. Al2O3; D. Al2TiO5; C. Mechanical properties; C. Toughening; B. Nanocomposites 1. Introduction The use of ceramic materials in structural applications is limited by the “flaw sensitive” fracture, occurring spontaneously from natural flaws, inherent to the brittle behaviour. The “flaw tolerance” approach deals with the development of microstructures that originate toughening mechanisms to reduce the sensitivity of the strength to the size of any processing or induced flaw, thus improving the reliability of the materials.1–3 Such mechanisms originate an increasing resistance with continued crack extension, rising R-curve behaviour, and most of them are caused by localized internal residual stresses in the materials. Alumina (Al2O3)–aluminium titanate (Al2TiO5) materials can offer improved flaw tolerance and toughness.4–12 Thermal expansion of aluminium titanate is highly anisotropic (αa25–1000 ◦C = 10.9 × 10−6 ◦C−1, αb25–1000 ◦C = 20.5 × ∗ Corresponding author. Tel.: +34 91 7355840; fax: +34 91 7355843. E-mail address: cbaudin@icv.csic.es (C. Baud´ın). 10−6 ◦C−1, αc25–1000 ◦C = −2.7 × 10−6 ◦C−1) 1 13 and alumina shows limited anisotropy (αa25–1000 ◦C = 8.4 × 10−6 ◦C−1, αc25–1000 ◦C = 9.2 × 10−6 ◦C−1),14 thus, high tensile or compressive stresses, depending on the particular crystallographic orientation of the grains, would develop during cooling from the sintering temperature at the grain–matrix interfaces due to thermal expansion mismatch. Depending on grain size and the characteristics of the grain boundaries, microcracking might occur during cooling from sintering and/or during fracture. In the early 90s alumina–aluminium titanate composites with aluminium titanate contents 20–30 vol.% obtained from alumina and aluminium titanate mixtures, were studied by other authors.4–7 Crack bridging by second phase agglomerates and by large alumina grains was identified as the toughening mechanism leading to R-curve behaviour, assessed by the indentation–strength method; no toughness values were 1 In this work, -Al2TiO5 orthorhombic crystal is described by a b-face centered unit cell, space group Bbmm, a = 9.439 A, ˚ b = 9.647 A, ˚ c = 3.593 A. ˚ 0955-2219/$ – see front matter © 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.jeurceramsoc.2008.01.017
S. Bueno et al. /Journal of the European Ceramic Society 28 (2008)1961-1971 provided for the fine-grained materials with homogeneous where E=E/(1-v)is the generalized Youngs modulus for microstructures. All materials presented different levels of plane strain(E is the Youngs modulus and v is the Poisson's microcracks in the"as sintered"state. A latter work on ratio) microcrack-free and fine-grained alumina+ 10 vol %o aluminium The activation of toughening mechanisms during the frac- titanate fabricated from alumina and aluminium titanate mix- ture of ceramic materials gives rise to inelastic strain processes tures showed that second phase grains as well as matrix grains that produce additional release of the elastic energy accumulated could act as bridges in the wake of the propagating crack. 8 in the material at the moment of fracture initiation and/or con- This material presented increased thermal shock resistance than tributing to the retardation of crack growth.The inelastic strain a monophase alumina of similar grain size while aining levels achieved in ceramic materials can be enough to restrict the strength direct utilization of linear elastic fracture toughness parameters The initial objective of this work was to investigate the since they become dependent on testing and specimen geometry possibilities of crack bridging in fine-grained, homogeneous for non-linear materials. 18-20 and microcrack -free alumina-10 vol %o aluminium titanate com The rising R-curve behaviour, increasing Kic or GIc posites for flaw tolerance. Reaction sintering of alumina and with crack extension(Aa), has traditionally been the most titania was used as processing route. I 5 The microstruc- utilized approach to analyze deviations from the linear tures of the reaction sintered materials were different than behaviour induced by toughening in dense and fine-grained nat of the previously studied material, with a bimodal dis- ceramics. 9, 21-22 In equilibrium conditions, the applied stress tribution of aluminium titanate grains with nanoparticles intensity factor, KI, is balanced by the crack growth resistance, located at the alumina grain boundaries. The characteriza- K, and maximum values of this latter, Ko, are reached when tion of the fracture process in the composites and monophase the process zone is completely developed. alumina materials, combining different fracture parameters In order to build the R curve of the materials, crack growth gether with fractographic observations, has allowed determin- resistance and crack length values during crack extension are ing the extreme effect of the grain boundary characteristics needed. The"in situ" measurement of crack length can be a prob- in the fracture process. The major toughening mechanism lem especially for materials such as alumina-aluminiumtitanate identified in the composite studied here has been microcrack- composites, constituted by phases with large differences in hard- ness and in which residual stresses are present. The low quality of polished surfaces of relatively large specimens(e.g: bending 2. Quantification of fracture toughness bars with lateral face dimension 50mm x 6 mm)of such mate- rials makes the identification and monitoring of the propagating In general, the linear elastic fracture behaviour of ceramic crack enormously difficult Alternatively, the R curves can be determined by the indirect critical stress intensity factor in mode I, KiC, and critical strain method that defines an equivalent crack length as a function energy release rate, GIC. For three-point-bend beams, the val- of the elastic compliance of the specimen, C 23-25 For par- ues of Kic can be determined from the notch depths and the allelepiped bars with straight through notches tested in three maximum loads reached in the tests according to the general points bending, the expression provided by Guinea et al. can stress intensity formulation, valid for any notch depth, a, in line be utilized(eq (4)) elastic materials(Eq(1)6 (CEB) 3PL 2BW3/2 X Y(a) [CEB+qI(CEB)2+q2(CEB)+g3 where P is the maximum load, L is the span, B and w are the where E, a and B have the same meaning as before(eq.(1) width and the thickness of the bars, a is the normalized notch and qi(i=1, 2, 3) are parameters that depend on the L/W ratio length(a=alW) and Y(a)is a shape function depending on the (2.5 <(L/W)< 16) span to thickness ratio(L/W, Eq(2)) In a lesser extent, the non-linear fracture toughness paran J-integral and work of fracture, ywoF, have been used br (199+0.83a-0.31a2+0.14a3+4W/L) and co-workers. 26 Bradt and co-workers 27 and Sakai et al. 18 Y(a)=x(-0.09-0.42a+0.82a2-0.31a2) (1-a)32×(1+3a) (2) to characterize ceramic materials with coarser microstructures and higher levels of non-linearity such as refractories and fiber From Kic and Young s modulus, Gic can be calculated according The J-integral is an energy term that generalizes the energy to the analysis of Irwin that relates the stress-derived fracture release rate, G, to include non-linear elastic and inelastic toughness(Kic)and the energy-derived fracture toughness( Gic) behaviours and that describes the total energy of the crack-tip for plane strain conditions(Eq ( 3): stress-strain field. The critical value, JIC, constitutes a fracture criterion for materials where the toughening occurs along lim- ited crack propagation such as those that present small bridging
1962 S. Bueno et al. / Journal of the European Ceramic Society 28 (2008) 1961–1971 provided for the fine-grained materials with homogeneous microstructures. All materials presented different levels of microcracks in the “as sintered” state. A latter work on microcrack-free and fine-grained alumina + 10 vol.% aluminium titanate fabricated from alumina and aluminium titanate mixtures showed that second phase grains as well as matrix grains could act as bridges in the wake of the propagating crack.8 This material presented increased thermal shock resistance than a monophase alumina of similar grain size while maintaining strength. The initial objective of this work was to investigate the possibilities of crack bridging in fine-grained, homogeneous and microcrack-free alumina–10 vol.% aluminium titanate composites for flaw tolerance. Reaction sintering of alumina and titania was used as processing route. 15 The microstructures of the reaction sintered materials were different than that of the previously studied material, with a bimodal distribution of aluminium titanate grains with nanoparticles located at the alumina grain boundaries. The characterization of the fracture process in the composites and monophase alumina materials, combining different fracture parameters together with fractographic observations, has allowed determining the extreme effect of the grain boundary characteristics in the fracture process. The major toughening mechanism identified in the composite studied here has been microcracking. 2. Quantification of fracture toughness In general, the linear elastic fracture behaviour of ceramic materials is quantified by the following toughness parameters: critical stress intensity factor in mode I, KIC, and critical strain energy release rate, GIC. For three-point-bend beams, the values of KIC can be determined from the notch depths and the maximum loads reached in the tests according to the general stress intensity formulation, valid for any notch depth, a, in linear elastic materials (Eq. (1)) 16: KI = 3PL 2BW3/2 × Y(α) (1) where P is the maximum load, L is the span, B and W are the width and the thickness of the bars, α is the normalized notch length (α = a/W) and Y(α) is a shape function depending on the span to thickness ratio (L/W, Eq. (2)). Y(α)= √α(1.99 + 0.83α − 0.31α2 + 0.14α3 + 4(W/L) ×(−0.09 − 0.42α + 0.82α2 − 0.31α3)) (1 − α) 3/2 × (1 + 3α) (2) FromKIC and Young’s modulus,GIC can be calculated according to the analysis of Irwin that relates the stress-derived fracture toughness (KIC) and the energy-derived fracture toughness (GIC) for plane strain conditions (Eq. (3)): GIC = K2 IC E (3) where E = E/(1 − ν2) is the generalized Young’s modulus for plane strain (E is the Young’s modulus and ν is the Poisson’s ratio). The activation of toughening mechanisms during the fracture of ceramic materials gives rise to inelastic strain processes that produce additional release of the elastic energy accumulated in the material at the moment of fracture initiation and/or contributing to the retardation of crack growth.17 The inelastic strain levels achieved in ceramic materials can be enough to restrict the direct utilization of linear elastic fracture toughness parameters since they become dependent on testing and specimen geometry for non-linear materials.18–20 The rising R-curve behaviour, increasing KIC or GIC with crack extension (a), has traditionally been the most utilized approach to analyze deviations from the linear behaviour induced by toughening in dense and fine-grained ceramics.19,21–22 In equilibrium conditions, the applied stress intensity factor, KI, is balanced by the crack growth resistance, KR, and maximum values of this latter, K∞, are reached when the process zone is completely developed. In order to build the R curve of the materials, crack growth resistance and crack length values during crack extension are needed. The “in situ” measurement of crack length can be a problem especially for materials such as alumina–aluminium titanate composites, constituted by phases with large differences in hardness and in which residual stresses are present. The low quality of polished surfaces of relatively large specimens (e.g.: bending bars with lateral face dimension 50 mm × 6 mm) of such materials makes the identification and monitoring of the propagating crack enormously difficult. Alternatively, the R curves can be determined by the indirect method that defines an equivalent crack length as a function of the elastic compliance of the specimen, C. 23–25 For parallelepiped bars with straight through notches tested in three points bending, the expression provided by Guinea et al.16 can be utilized (Eq. (4)): α = (CE B) 1/2 [CE B + q1(CE B) 1/2 + q2(CE B) 1/3 + q3] 1/2 (4) where E , α and B have the same meaning as before (Eq. (1)) and qi (i = 1, 2, 3) are parameters that depend on the L/W ratio (2.5 ≤ (L/W) ≤ 16). In a lesser extent, the non-linear fracture toughness parameter J-integral and work of fracture, γWOF, have been used by Li and co-workers,26 Bradt and co-workers27 and Sakai et al.18 to characterize ceramic materials with coarser microstructures and higher levels of non-linearity such as refractories and fiber reinforced ceramic matrix composites. The J-integral is an energy term that generalizes the energy release rate, G, to include non-linear elastic and inelastic behaviours and that describes the total energy of the crack-tip stress–strain field.28 The critical value, JIC, constitutes a fracture criterion for materials where the toughening occurs along limited crack propagation such as those that present small bridging zones.29
S. Bueno et al. / Journal of the European Ceramic Society 28(2008)1961-1971 1963 There are two main different procedures to determine 1389: 2003)and relative densities were calculated from these JiC,26-27,30 either based on the determination of the energy values and those of theoretical densities calculated taking values absorbed by the specimen, given by the area under the corre- of 3.99 gcm-3for alumina(ASTM 42-1468)and 3.70 g cm-3 sponding load-crack opening displacement curve, 30 or from for aluminium titanate(ASTM 26-0040 load-displacement curves by conducting tests on two spec Microstructural characterization on polished and thermally imens with different crack lengths. Both methods require etched(20C below the sintering temperature during 1 min) the identification of the propagating crack. In this work, a surfaces was performed by field emission gun scanning elec- graphical procedure was used- in which Jic is calculated tron microscopy(FEG-SEM, Hitachi, S-4700, Japan). The (Eq (5))from the difference between the areas under the load true average grain size was determined by the linear inter- (P)displacement(8)curves of the notched non-linear speci- cept method considering at least 200 grains for each phase and mens(Al) and an unnotched linear elastic specimen of the same g the correction factor 4.Chemical profiles across grain material(AE) for equal maximum loads(Pmax) boundaries were achieved by STEM-EDX (energy dispersed x- ray spectroscopy, coupled with scanning transmission electron X(Al-Ae) (5) microscopy, Tecnai F20-ST, The Netherlands) at 200 kV. Thin (W-a)B (W一a)B foils were prepared by mechanical polishing of a 3 mm diame ere a, W and B have the same meaning as before(Eq (1)). ter disk up to 15 um in thickness followed by Ar milling(PlPs The work of fracture is defined as the mean energy con- Gatan, USA, operating at 5 kV with a beam incidence of 6%0) sumption required forming the unit fracture surface area and Bars of 25 mm x 2 mm x 2.5 mm were diamond machined the additional process zone. It accounts an average value of from the sintered blocks for bend strength tests(three points, the whole fracture process that does not require any as 20 mm span, 0.5 mm min; Microtest, Spain). Engineering tion on the constitutive equation of the cracked body to deal stress-strain curves were calculated from the load values and with crack growth problems as discussed by Sakai et al. 8 The the displacement of the central part of the samples recorded work of fracture is obtained by dividing the work done on the during the bending tests and static Youngs modulus was deter specimen to propagate the crack, given by the area under the mined from the initial linear part of the curves. Given results load-displacement curves, by the area of the newly created for strength and static Youngs modulus are the average of five surfaces. For parallelepiped bars with straight through notches determinations and the standard deviation tested in flexure, this area is twice the area of the unnotched part Strength was also determined for specimens of a previously of the cross-section of the specimens studied A10 composite(named A10AT)fabricated from pow ders of Al2O3(90 vol %)and Al2TiO5 (10 vol %o)obtained by 3. Experimental procedure reaction of Al2O3 and TiO2 powders and sintered at 1500C, the starting Al2O3 and TiO2 powders used were the same as in Monoliths of monophase alumina (A) and alu- this work. nina+10 vol 9 aluminium titanate(AlO)composites were Single-Edge-V-Notch-Beams(SEVNB )of 4 mm x 6mm x manufactured by colloidal filtration from aqueous alumina, 50 mm were tested in a three points bending device using a span Al2O3, and titania, TiO2, suspensions using the optimum green of 40 mm and a cross-head speed of 0.005 mm min(Microtest, processing conditions previously established. 5.31 A mixture Spain). The compliance of the whole testing system(machine, of alumina ( =95 wt % )and titania (5 wt %)was used to supports, load cell and fixtures) was determined by testing a obtain the sintered composition with 10 vol. of aluminium thick(25 mm x 25 mm x 100 mm)unnotched alumina bar. The titanate, Al2TiO5. The starting materials were commercial obtained value was 1. 5x 10-m/N. The notches were initially a-Al2O3( Condea, HPAO5, USA) and TiO2-anatase(Merck, cut with a 150 um wide diamond wheel. Using this slot as a 808, Germany) powders. The powders were dispersed in guide, the remaining part of the notch was done with a razor deionised water by adding 0.5 wt %(on a dry solids basis) blade sprinkled with diamond pastes of successively 6 and 1 um of a carbonic acid-based polyelectrolyte (Dolapix CE64, Three relative notch depths, a, with approximately 0.4, 0.5 and Zschimmer-Schwarz, Germany). Suspensions were prepared to 0.6 of the thickness of the samples(w) were tested. The tip a solids loading of 50 vol. and ball milled with alumina radii of all notches were determined from optical observations and balls during 4 h. and they were always found to be below 20 um. The curves Plates of the materials with 70 mm x 70 mm x 10 mm dimen- load-displacement of the cross-head of the load frame were sions were obtained by slip casting, removed from the moulds recorded. All curves were corrected by subtracting the com- and dried in air at room temperature for at least 24 h. Sinter- pliance of the testing set up. ing of the green plates was performed in air in an electrical Additional tests were performed with unnotched specimens box furnace(Termiber, Spain) at heating and cooling rates of up to loads(=20N) well below the starting of the non-linear 2Cmin-I,with 4h, dwell at 1200C during heating and two behaviour and the obtained values of stiffness were used to different treatments at the maximum temperature: 2h, dwell at calculate JIc following the procedure described above(Eq (5)) 1450C and 3 h, dwell at 1550C. For all tests, samples were The fracture toughness parameters, i.e., critical stress inten- diamond machined from the sintered blocks sity factor, KiC, critical strain energy release rate, GIC, critical Densities of the sintered compacts were determined by J-integral, JiC, and work of fracture, ywoF, were calculated the Archimedes's method in water(European Standard en from the curves obtained during the sEvnb tests for the three
S. Bueno et al. / Journal of the European Ceramic Society 28 (2008) 1961–1971 1963 There are two main different procedures to determine JIC, 26–27,30 either based on the determination of the energy absorbed by the specimen, given by the area under the corresponding load–crack opening displacement curve,30 or from load–displacement curves by conducting tests on two specimens with different crack lengths.26 Both methods require the identification of the propagating crack. In this work, a graphical procedure was used27 in which JIC is calculated (Eq. (5)) from the difference between the areas under the load (P)–displacement (δ) curves of the notched non-linear specimens (AI) and an unnotched linear elastic specimen of the same material (AE) for equal maximum loads (Pmax): JIC= 2 (W − a)B × δmax 0 Pdδ= 2 (W − a)B × (AI − AE) (5) where a, W and B have the same meaning as before (Eq. (1)). The work of fracture is defined as the mean energy consumption required forming the unit fracture surface area and the additional process zone. It accounts an average value of the whole fracture process that does not require any assumption on the constitutive equation of the cracked body to deal with crack growth problems as discussed by Sakai et al.18 The work of fracture is obtained by dividing the work done on the specimen to propagate the crack, given by the area under the load–displacement curves, by the area of the newly created surfaces. For parallelepiped bars with straight through notches tested in flexure, this area is twice the area of the unnotched part of the cross-section of the specimens. 3. Experimental procedure Monoliths of monophase alumina (A) and alumina + 10 vol.% aluminium titanate (A10) composites were manufactured by colloidal filtration from aqueous alumina, Al2O3, and titania, TiO2, suspensions using the optimum green processing conditions previously established.15,31 A mixture of alumina (∼=95 wt.%) and titania (∼=5 wt.%) was used to obtain the sintered composition with 10 vol.% of aluminium titanate, Al2TiO5. The starting materials were commercial α-Al2O3 (Condea, HPA05, USA) and TiO2-anatase (Merck, 808, Germany) powders. The powders 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 alumina jar and balls during 4 h. Plates of the materials with 70 mm × 70 mm × 10 mm dimensions were obtained by slip casting, removed from the moulds and dried in air at room temperature for at least 24 h. Sintering of the green plates was performed in air in an electrical box furnace (Termiber, Spain) at heating and cooling rates of 2 ◦C min−1, with 4 h, dwell at 1200 ◦C during heating and two different treatments at the maximum temperature: 2 h, dwell at 1450 ◦C and 3 h, dwell at 1550 ◦C. For all tests, samples were diamond machined from the sintered blocks. Densities of the sintered compacts were determined by the Archimedes’s 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 cm−3 for alumina (ASTM 42-1468) and 3.70 g cm−3 for aluminium titanate (ASTM 26-0040). Microstructural characterization on polished and thermally etched (20 ◦C below the sintering temperature during 1 min) surfaces was performed by field emission gun scanning electron microscopy (FEG-SEM, Hitachi, S-4700, Japan). The true average grain size was determined by the linear intercept method considering at least 200 grains for each phase and using the correction factor 4/π. 32 Chemical profiles across grain boundaries were achieved by STEM–EDX (energy dispersed Xray spectroscopy, coupled with scanning transmission electron microscopy, Tecnai F20-ST, The Netherlands) at 200 kV. Thin foils were prepared by mechanical polishing of a 3 mm diameter disk up to 15m in thickness followed by Ar+ milling (PIPS Gatan, USA, operating at 5 kV with a beam incidence of 6%). Bars of 25 mm × 2 mm × 2.5 mm were diamond machined from the sintered blocks for bend strength tests (three points, 20 mm span, 0.5 mm min−1; Microtest, Spain). Engineering stress–strain curves were calculated from the load values and the displacement of the central part of the samples recorded during the bending tests and static Young’s modulus was determined from the initial linear part of the curves. Given results for strength and static Young’s modulus are the average of five determinations and the standard deviation. Strength was also determined for specimens of a previously studied A10 composite8 (named A10AT) fabricated from powders of Al2O3 (90 vol.%) and Al2TiO5 (10 vol.%) obtained by reaction of Al2O3 and TiO2 powders33 and sintered at 1500 ◦C; the starting Al2O3 and TiO2 powders used were the same as in this work. Single-Edge-V-Notch-Beams (SEVNB) of 4 mm × 6 mm × 50 mm were tested in a three points bending device using a span of 40 mm and a cross-head speed of 0.005 mm min−1 (Microtest, Spain). The compliance of the whole testing system (machine, supports, load cell and fixtures) was determined by testing a thick (25 mm × 25 mm × 100 mm) unnotched alumina bar. The obtained value was 1.5 × 10−7 m/N. The notches were initially cut with a 150 m 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 m. Three relative notch depths, α, with approximately 0.4, 0.5 and 0.6 of the thickness of the samples (W) were tested. The tip radii of all notches were determined from optical observations and they were always found to be below 20 m. The curves load–displacement of the cross-head of the load frame were recorded. All curves were corrected by subtracting the compliance of the testing set up. Additional tests were performed with unnotched specimens up to loads (∼=20 N) well below the starting of the non-linear behaviour and the obtained values of stiffness were used to calculate JIC following the procedure described above (Eq. (5)). The fracture toughness parameters, i.e., critical stress intensity factor, KIC, critical strain energy release rate, GIC, critical J-integral, JIC, and work of fracture, γWOF, were calculated from the curves obtained during the SEVNB tests for the three
S. Bueno et al. /Journal of the European Ceramic Society 28 (2008)1961-1971 Table I Properties of the materials: average grain size(G), relative density (p), static Young s modulus(E) and three points bending strength(of) GA(S.D. )(um) GAT(.D)(um) p(S D )( theoretical) E(S D )(GPa) ar(SD )(MPa) 3.5(0.3) 981(0.3) 379(8) 456(29) 5.5(0.6) 981(0.5) 349(31) A10-1450 3.2(0.4 2.2(0.1) 973(0.5) 301(4) 261(6 10-1550 3.9(0.3 24(0.2) 97200.3) 272(10) 230(1) AlOAT 985 367(5) 360(31) A: alumina, AT: aluminum titanate: S D. standard deviation notch depths utilized. Reported values are the average of three(0.85-0.87)was slightly lower than the stoichiometric(0.89) determinations and errors are the standard deviations. R curves Nevertheless, results of these semi quantitative analyses are valid were determined from the load versus displacement curves cor- for comparative purposes. No Ti was detected inside the alumina responding to tests performed with a relative notch depth of 0.6 grains of the composites(Fig 2a), whose analyses were simi- of the thickness of The fracture surfaces of tested strength and SEVNB speci mens were characterized by FEG-SEM. Also small samples of the lateral faces(face dimension 50 mm x 6 mm)containing the notches and the cracks were polished and chemically etched (HF-10 vol %0-3 min) in order to observe the zones surround ing the propagating cracks to characterize the process zones. In order to complement the fractographic observations, polished surfaces of composite samples indented with a Vickers point using 50N during 10s, were also observed. 4. Results and discussion 4. Microstructure The microstructures of both aluminas were typical of mate- rials fabricated from high-purity submicron alumina powders The material sintered at 1450C was constituted by equidimen sional grains with a narrow distribution of relatively small sizes whereas that sintered at 1550C presented a coarser microstruc ture with a wide distribution of sizes and pore trapping associated 04m with exaggerated grain growth. The microstructural parameters together with the density, static Youngs modulus and strength values are reported in Table 1 The composites presented micrometer sized(2.2-2. 4 um, Fig. la and b, Table 1)aluminium titanate grains homoge neously distributed and located mainly at alumina triple points and grain boundaries and alumina grains of sizes similar to those of the monophase alumina sintered at 1450C(3. 2-3.9 um, Fig. la, Table 1). Submicrometric second phase grains were also observed inside the alumina grains and occasionally at grai boundaries(Fig. la). Additional nanometer sized grains were bserved at grain boundaries by SEM(Fig. lb) In Fig. 2 characteristic STEM observations for the ites sintered at 1550C together with EDX chemical analysis are shown. The ratios(wt %)Al/O(E1.4)and T1/O(E1.3)in the than those corresponding to the stoichiometric, 0.68 and 0.60, electron micrographs of polishedand s of the stud grains of aluminium titanate( Fig. 2a)were always well higher Fig. 1. Characteristic micr AlO composites. Scanning ermally etched surfaces. Alumina grains respectively. The Ka. B radiations emitted by light elements have appear with dark grey colour whereas micrometer sized aluminium titanat lower energies and are preferentially absorbed by carbon con- cromenc. ghter gray shade.(a)Composite A10 sintered at 1450CSubmi- tamination formed during the spot analyses. This induces an (b)Composite A10 sintered at 1550C. Detail of nanosized(arrows)aluminium underestimation of oxygen concentration. Also, the ratio Ti/al titanate grains located at the boundaries between the alumina grains
1964 S. Bueno et al. / Journal of the European Ceramic Society 28 (2008) 1961–1971 Table 1 Properties of the materials: average grain size (G), relative density (ρ), static Young’s modulus (E) and three points bending strength (σf) GA (S.D.) (m) GAT (S.D.) (m) ρ (S.D.) (% theoretical) E (S.D.) (GPa) σf (S.D.) (MPa) A-1450 3.5 (0.3) – 98.1 (0.3) 379 (8) 456 (29) A-1550 5.5 (0.6) – 98.1 (0.5) 376 (6) 349 (31) A10-1450 3.2 (0.4) 2.2 (0.1) 97.3 (0.5) 301 (4) 261 (6) A10-1550 3.9 (0.3) 2.4 (0.2) 97.2 (0.3) 272 (10) 230 (1) A10AT 98.5 (0.1) 367 (5) 360 (31) A: alumina, AT: aluminum titanate; S.D.: standard deviation. notch depths utilized. Reported values are the average of three determinations and errors are the standard deviations. R curves were determined from the load versus displacement curves corresponding to tests performed with a relative notch depth of 0.6 of the thickness of the samples. The fracture surfaces of tested strength and SEVNB specimens were characterized by FEG-SEM. Also small samples of the lateral faces (face dimension 50 mm × 6 mm) containing the notches and the cracks were polished and chemically etched (HF-10 vol.%–3 min) in order to observe the zones surrounding the propagating cracks to characterize the process zones. In order to complement the fractographic observations, polished surfaces of composite samples indented with a Vickers point using 50 N during 10 s, were also observed. 4. Results and discussion 4.1. Microstructure The microstructures of both aluminas were typical of materials fabricated from high-purity submicron alumina powders. The material sintered at 1450 ◦C was constituted by equidimensional grains with a narrow distribution of relatively small sizes whereas that sintered at 1550 ◦C presented a coarser microstructure with a wide distribution of sizes and pore trapping associated with exaggerated grain growth. The microstructural parameters together with the density, static Young’s modulus and strength values are reported in Table 1. The composites presented micrometer sized (2.2–2.4m, Fig. 1a and b, Table 1) aluminium titanate grains homogeneously distributed and located mainly at alumina triple points and grain boundaries and alumina grains of sizes similar to those of the monophase alumina sintered at 1450 ◦C (3.2–3.9m, Fig. 1a, Table 1). Submicrometric second phase grains were also observed inside the alumina grains and occasionally at grain boundaries (Fig. 1a). Additional nanometer sized grains were observed at grain boundaries by SEM (Fig. 1b). In Fig. 2 characteristic STEM observations for the composites sintered at 1550 ◦C together with EDX chemical analysis are shown. The ratios (wt.%) Al/O (∼=1.4) and Ti/O (∼=1.3) in the grains of aluminium titanate (Fig. 2a) were always well higher than those corresponding to the stoichiometric, 0.68 and 0.60, respectively. The K, radiations emitted by light elements have lower energies and are preferentially absorbed by carbon contamination formed during the spot analyses. This induces an underestimation of oxygen concentration. Also, the ratio Ti/Al (0.85–0.87) was slightly lower than the stoichiometric (0.89). Nevertheless, results of these semi quantitative analyses are valid for comparative purposes. No Ti was detected inside the alumina grains of the composites (Fig. 2a), whose analyses were simiFig. 1. Characteristic microstructures of the studied A10 composites. Scanning electron micrographs of polished and thermally etched surfaces. Alumina grains appear with dark grey colour whereas micrometer sized aluminium titanate grains have lighter gray shade. (a) Composite A10 sintered at 1450 ◦C. Submicrometric second phase grains inside the alumina matrix are pointed by arrows. (b) Composite A10 sintered at 1550 ◦C. Detail of nanosized (arrows) aluminium titanate grains located at the boundaries between the alumina grains.
S. Bueno et al. /Journal of the European Ceramic Sociery 28(2008)1961-1971 1965 3 0100020003000400050006000 Energy (ev C 1000200030 400050006000 100nm 95 Fig. 2. Characteristic scanning transmission electron microscopy(STEM)observations for the A10 composites sintered at 1550C together with EDX chemical nalysis(au -arbitrary units). (a) Alumina and aluminium titanate grains. No Ti was detected inside the alumina grains (b)Chemical profile along a line traversing an alumina/alumina grain boundary showing enrichment in Ti Negligible Si contents are detected. lar to those of the monophase specimens. However, the eDx should be aluminium titanate, formed by reaction of the thermo- line profiles across alumina grain boundaries in the composites dynamically incompatible compounds alumina and titania. The (Fig. 2b) showed a systematic evidence of Ti segregation at the fact that such particles were not observed by STEM should be alumina/alumina grain boundaries. Values from 0.5 to 2.5 Ti due to the relatively small portions of material characterized by wt% were detected with no systematic variation with alumina this method(two samples were observed) graIn size. The presence of the major impurity in the starting pow- 4.2. Toughness parameters ders, Si, was also investigated and only no Si or negligible Si contents were found in the grain boundaries(Fig. 2b). More- The load-displacement curves for both composites and for over, STEM-EDX analysis evidenced diffusion of titanium ions the three relative notch sizes showed stable fracture In Fig. 3 across the alumina grain boundaries during sintering. Thus, the characteristic curves for specimens with a relative notch length composition of the nanosized particles found by SEM(Fig. 1b) a/W=0.5 are shown. Controlled fracture was difficult to achieve
S. Bueno et al. / Journal of the European Ceramic Society 28 (2008) 1961–1971 1965 Fig. 2. Characteristic scanning transmission electron microscopy (STEM) observations for the A10 composites sintered at 1550 ◦C together with EDX chemical analysis (a.u. = arbitrary units). (a) Alumina and aluminium titanate grains. No Ti was detected inside the alumina grains. (b) Chemical profile along a line traversing an alumina/alumina grain boundary showing enrichment in Ti. Negligible Si contents are detected. lar to those of the monophase specimens. However, the EDX line profiles across alumina grain boundaries in the composites (Fig. 2b) showed a systematic evidence of Ti segregation at the alumina/alumina grain boundaries. Values from 0.5 to 2.5 Ti wt.% were detected with no systematic variation with alumina grain size. The presence of the major impurity in the starting powders, Si, was also investigated and only no Si or negligible Si contents were found in the grain boundaries (Fig. 2b). Moreover, STEM–EDX analysis evidenced diffusion of titanium ions across the alumina grain boundaries during sintering. Thus, the composition of the nanosized particles found by SEM (Fig. 1b) should be aluminium titanate, formed by reaction of the thermodynamically incompatible compounds alumina and titania. The fact that such particles were not observed by STEM should be due to the relatively small portions of material characterized by this method (two samples were observed). 4.2. Toughness parameters The load–displacement curves for both composites and for the three relative notch sizes showed stable fracture. In Fig. 3 characteristic curves for specimens with a relative notch length a/W = 0.5 are shown. Controlled fracture was difficult to achieve