C MATERIALIA Pergamon Acta mater.48(20004715-4720 www.elsevier.com/locate/actamat EFFECT OF LAYER INTERFACES ON THE HIGH TEMPERATURE MECHANICAL PROPERTIES OF ALUMINA/ZIRCONIA LAMINATE COMPOSITES M. JIMENEZ-MELENDO*, F. GUTIERREZ-MORA and A DOMINGUEZ-RODRIGUEZ Departamento de Fisica de la Materia Condensada, Universidad de Sevilla, Aptdo 1065, 41080 Sevilla, Abstract-High-temperature plastic deformation of laminar composites containing layers of Al2O3 and a mixture of 85 voL %Al2O3+15 vol. %3 mol%Y2O stabilized tetragonal ZrO2(ZTA) produced by sequenti slip casting is examined in uniaxial compression testing. The mechanical behavior is compared with that of monolithic Al]O, and ZTA synthesized by the same technique. The layered composites exhibit better cree properties at low strain rate than either of the two constituent materials. The good interfacial adhesion of he layers imparts creep resistance and ductility simultaneously to the laminates. 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. AlI rights reserved. Keywords: Structural ceramics; Layered structures; Creep 1. INTRODUCTION of pure Al2O3 and Al2O3+ 15 vol. ZrO doped with In the last few years, ceramic/ceramic lamellar struc- 3 mol%Y2O3 produced by sequential slip casting to tures have emerged as promising candidates to over- high-temperature mechanical properties,which is structural applications (1-3). These structures bring a essential for the design of materials and unique opportunity for tailoring the mechanical geometry for high-temperature applications properties by stacking layers of different thickness and composition in a suitable sequence. Laminar 2. EXPERIMENTAL PROCEDURES opposites are presently synthesized by different pro- cessing routes, including tape casting, sequential slip Green layered compacts of alternate layers of pure casting, electrophoretic deposition and colloidal tech- Al O3 and a mixture of 85 vol. Al O3+ 15 vol% niques [4-8]. Drastic increases in strength and ZrO2 doped with 3 mol%Y,O,(ZTA)were obtained especially in fracture toughness at room temperature by sequential slip casting; details of the starting pow- have been achieved in alumina/zirconia laminar com- ders and the casting process can be found elsewhere posites because of various crack-shielding phenom- [4, 12]. Monolithic Al2O3 and ZTA were also fabri- ena related to the presence of the layers cated by the same technique for the sake of compari- (delamination, crack deflection, etc. )[5,8-10. son. After sintering at 1550.C in air, the density of Beautiful examples of systems that exhibit crack all monolithic and layered compacts, measured by deflection at interfaces occur in nature and have a Archimede's method, was close to the theoretic plate-like architecture that include shells and teeth density. In order to study the effect of the layer thick [11]. It has also been shown that the laminar micro- ness, two multilayered configurations were produced structure influences the creep properties of multilay- composites"L', consisting of Al2 O3 layers 125 um ered composites with strong interfaces [6]. This work thick and ZTA layers 300 um thick; and composites extends a preliminary investigation 12] on laminates "F", where the thickness of both types of layers was roximately 150 Imens mn hom all correspondence should be addressed. Tel: from the laminated compacts with the largest dimen- +34-95-4550938;fax:+34-95-46120 ion(the loading axis)either parallel (composites lab- E-mail address: melendo(@cica.es (M. Jimenez. eled"PAR") or perpendicular(composites labeled Melendo) PER ) to the layer interfaces. Compression tests 1359-6454100/520.00@ 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved PI:S1359-6454(00)00262-7
Acta mater. 48 (2000) 4715–4720 www.elsevier.com/locate/actamat EFFECT OF LAYER INTERFACES ON THE HIGHTEMPERATURE MECHANICAL PROPERTIES OF ALUMINA/ZIRCONIA LAMINATE COMPOSITES M. JIME´ NEZ-MELENDO*, F. GUTIE´ RREZ-MORA and A. DOMI´NGUEZ-RODRI´GUEZ Departamento de Fı´sica de la Materia Condensada, Universidad de Sevilla, Aptdo 1065, 41080 Sevilla, Spain Abstract—High-temperature plastic deformation of laminar composites containing layers of Al2O3 and a mixture of 85 vol.% Al2O3+15 vol.% 3 mol% Y2O3-stabilized tetragonal ZrO2 (ZTA) produced by sequential slip casting is examined in uniaxial compression testing. The mechanical behavior is compared with that of monolithic Al2O3 and ZTA synthesized by the same technique. The layered composites exhibit better creep properties at low strain rate than either of the two constituent materials. The good interfacial adhesion of the layers imparts creep resistance and ductility simultaneously to the laminates. 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Structural ceramics; Layered structures; Creep 1. INTRODUCTION In the last few years, ceramic/ceramic lamellar structures have emerged as promising candidates to overcome the inherent brittleness of ceramics for use in structural applications [1–3]. These structures bring a unique opportunity for tailoring the mechanical properties by stacking layers of different thickness and composition in a suitable sequence. Laminar composites are presently synthesized by different processing routes, including tape casting, sequential slip casting, electrophoretic deposition and colloidal techniques [4–8]. Drastic increases in strength and especially in fracture toughness at room temperature have been achieved in alumina/zirconia laminar composites because of various crack-shielding phenomena related to the presence of the layers (delamination, crack deflection, etc.) [5, 8–10]. Beautiful examples of systems that exhibit crack deflection at interfaces occur in nature and have a plate-like architecture that include shells and teeth [11]. It has also been shown that the laminar microstructure influences the creep properties of multilayered composites with strong interfaces [6]. This work extends a preliminary investigation [12] on laminates * To whom all correspondence should be addressed. Tel.: +34-95-4550938; fax: +34-95-4612097. E-mail address: melendo@cica.es (M. Jime´nezMelendo) 1359-6454/00/$20.00 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S13 59-6454(00)00262-7 of pure Al2O3 and Al2O3 + 15 vol.% ZrO2 doped with 3 mol% Y2O3 produced by sequential slip casting to understand the effect of the layer interfaces on their high-temperature mechanical properties, which is essential for the design of materials and layer geometry for high-temperature applications. 2. EXPERIMENTAL PROCEDURES Green layered compacts of alternate layers of pure Al2O3 and a mixture of 85 vol.% Al2O3 + 15 vol.% ZrO2 doped with 3 mol% Y2O3 (ZTA) were obtained by sequential slip casting; details of the starting powders and the casting process can be found elsewhere [4, 12]. Monolithic Al2O3 and ZTA were also fabricated by the same technique for the sake of comparison. After sintering at 1550°C in air, the density of all monolithic and layered compacts, measured by Archimede’s method, was close to the theoretical density. In order to study the effect of the layer thickness, two multilayered configurations were produced: composites “L”, consisting of Al2O3 layers 125 µm thick and ZTA layers 300 µm thick; and composites “F”, where the thickness of both types of layers was approximately 150 µm. Rectangular specimens of 4×2×2 mm3 were cut from the laminated compacts with the largest dimension (the loading axis) either parallel (composites labeled “PAR”) or perpendicular (composites labeled “PER”) to the layer interfaces. Compression tests
4716 JIMENEZ. MELENDO et al- ALUMINA/ZIRCONIA LAMINATE COMPOSITES were carried out using an Instron universal testing machine at 1400C in air at constant cross-head speeds between I and 20 um/min, corresponding to nitial strain rates e between 4x10- and 8x10-/s The recorded data, load vs time, were plotted in o-E urves, were o and e are the engineering stress and true strain, respectively. Compressive creep tests at constant load were also performed in a creep machine made of alumina [13], using procedures described elsewhere [14]. These tests were carried out in air between 1400 and 1450%c under nominal stresses 82 between 50 and 100 MPa: both T and o were changed incrementally during creep. The recorded data, instan- aneous length vs time, were plotted as log E-E curves. Monolithic Al-O and ZTA were also deformed under identical experimental conditions The microstructural characterization of the as- received and deformed laminated composites and monoliths was performed using optical and scanning electron microscopy(Microscopy Service, University of Sevilla, Spain). To this end, parallel and perpen samples and mechanically polished. The sections were then thermally etched at 1350 C for 30 min in air. The morphological parameters of the various hases were characterized by using a semiautomatic Fig. 1. SEM micrographs of Al,O,/ZTA (85 vol %Al 03+15 F with layers of approximately the same thickness. (a) Low agnification(dark layers are Al2O bright layers are ZTA 3. EXPERIMENTAL RESULTS b)Detail of an interface between the Al,O, layer and the ZTa Figure 1(a) shows the cross-section of the as- of relatively small tetragonal ZrO, grains(bright phase)and received laminated composite F, with both Al2O3 arge alumina grains(dark phase) (dark phase)and ZTA layers having about the same thickness. The two types of layers are well defined vith straight interfaces, and no significant residua porosity was observed at the layer interfaces [F T=1400° I(b)). The grains in the Al,O, layer Fig. I(b)] exhibit well-faceted boundaries and sharp triple points, witl an average grain size d(the equivalent planar diam- a150 area) of 3 um. On the other hand, the ZTA layers [Fig. I(b)] show the typical duplex microstructure of alumina/zirconia composites [15], with relatively large grains of Al,O,(d=I um, dark phase) and small 3YTZP grains(d=0.3 um, bright phase)uni- formly dispersed throughout the layer, both having LPAR quiaxed shapes. The remarkable resistance to coars ning at elevated temperatures of such composites L [16, 17] has been exploited to achieve like-metal superplasticity [15, 181. Identical morphological STRAIN (O f laminated composite for the larger grain size of the Al,O, and 3YTZP L(AlzO, layers 125 m, Z. grains in monolithic ZTA (1.3 and 0.5 um, nd parallel(LPAR)to the respectively) with respect to those measured in the loading axis as a func initial strain rate Eo laminates Figure 2 shows the o-e curves for multilayered creep regime was attained for both layer geometries, composites"L with interface planes that are oriented characterized by a constant flow stress, the PAR con arallel(LPAR) and perpendicular(PeR)to the figuration shows a higher creep strength than that for compression axis, as a function of the initial strain PER. As the strain rate increases, the LPAR samples rate. For the lower strain rate, an extensive secondary exhibit a gradual failure after attaining a maximum
4716 JIME´ NEZ-MELENDO et al.: ALUMINA/ZIRCONIA LAMINATE COMPOSITES were carried out using an Instron universal testing machine at 1400°C in air at constant cross-head speeds between 1 and 20 µm/min, corresponding to initial strain rates e˙o between 4×1026 and 8×1025 /s. The recorded data, load vs time, were plotted in s–e curves, were s and e are the engineering stress and true strain, respectively. Compressive creep tests at constant load were also performed in a creep machine made of alumina [13], using procedures described elsewhere [14]. These tests were carried out in air between 1400 and 1450°C under nominal stresses between 50 and 100 MPa; both T and s were changed incrementally during creep. The recorded data, instantaneous length vs time, were plotted as log e˙–e curves. Monolithic Al2O3 and ZTA were also deformed under identical experimental conditions. The microstructural characterization of the asreceived and deformed laminated composites and monoliths was performed using optical and scanning electron microscopy (Microscopy Service, University of Sevilla, Spain). To this end, parallel and perpendicular sections to the interfaces were cut from the samples and mechanically polished. The sections were then thermally etched at 1350°C for 30 min in air. The morphological parameters of the various phases were characterized by using a semiautomatic image analyzer. 3. EXPERIMENTAL RESULTS Figure 1(a) shows the cross-section of the asreceived laminated composite F, with both Al2O3 (dark phase) and ZTA layers having about the same thickness. The two types of layers are well defined with straight interfaces, and no significant residual porosity was observed at the layer interfaces [Fig. 1(b)]. The grains in the Al2O3 layer [Fig. 1(b)] exhibit well-faceted boundaries and sharp triple points, with an average grain size d (the equivalent planar diameter, defined as d = (4A/p) 1/2, where A is the grain area) of 3 µm. On the other hand, the ZTA layers [Fig. 1(b)] show the typical duplex microstructure of alumina/zirconia composites [15], with relatively large grains of Al2O3 (d = 1 µm, dark phase) and small 3YTZP grains (d = 0.3 µm, bright phase) uniformly dispersed throughout the layer, both having equiaxed shapes. The remarkable resistance to coarsening at elevated temperatures of such composites [16, 17] has been exploited to achieve like-metal superplasticity [15, 18]. Identical morphological characteristics were found in the monoliths, except for the larger grain size of the Al2O3 and 3YTZP grains in monolithic ZTA (1.3 and 0.5 µm, respectively) with respect to those measured in the laminates. Figure 2 shows the s–e curves for multilayered composites “L” with interface planes that are oriented parallel (LPAR) and perpendicular (LPER) to the compression axis, as a function of the initial strain rate. For the lower strain rate, an extensive secondary Fig. 1. SEM micrographs of Al2O3/ZTA (85 vol.% Al2O3+15 vol.% ZrO2 doped with 3 mol% Y2O3) multilayer composites F with layers of approximately the same thickness. (a) Low magnification (dark layers are Al2O3, bright layers are ZTA). (b) Detail of an interface between the Al2O3 layer and the ZTA layer; this latter layer shows a duplex microstructure consisting of relatively small tetragonal ZrO2 grains (bright phase) and large alumina grains (dark phase). Fig. 2. Stress–strain curves at 1400°C of laminated composites L (Al2O3 layers 125 µm; ZTA layers 300 µm thick) with interface planes perpendicular (LPER) and parallel (LPAR) to the loading axis as a function of the initial strain rate e˙o. creep regime was attained for both layer geometries, characterized by a constant flow stress; the PAR con- figuration shows a higher creep strength than that for PER. As the strain rate increases, the LPAR samples exhibit a gradual failure after attaining a maximum
JIMENEZ-MELENDO et al- ALUMINA/ZIRCONIA LAMINATE COMPOSITES 4717 stress of about 200 MPa, independent of the initial strain rate. In contrast, the LPER samples still display a rather steady-state deformat ter a con- trolled stress dre E=2x10-5s-l igure 3 displays a constant load creep test for a LPER specimen showing several stress and tempera- ture changes. Despite the large final strain of the sam- 100 ple(E-60%), no evidence of macroscopic failure was found. The data have been analyzed using the stan- dard high-temperature power law for steady-state 5 deformation [14: E=4x106s-1 LPAR E=Ao c FPAR where A is a parameter depending on the deformation STRAIN(.) mechanism,R is the gas constant, n is the stress Fig 4 Stress-strain curves at 1400%C for laminated co exponent, p is the grain size exponent, and g is the L and F with interface planes parallel to the stress axis(LPAR activation energy. From the steady-state strain rate and FPAR)as a function of the initial strain rate e. E, before and after e2 the change in stress or tempera- ture(Fig. 3), the stress exponent and the activation energy is given, respectively, by [14] structural observations of deformed laminated ce posites(Fig. 5)[12] show a general creep damage of the Al2O, layers, with very little microcracking by (2) cavity coalescence. The ZTA layers are almost fea tuneless, indicating that the deformation of this phase was accommodated preferentially by grain boundary sliding. In all the cases the interfaces show an excel- @=RT -. log(ex/e)) (3) lent adhesion, maintaining their structural integrity The mechanical behavior of the laminates with PER and PaR geometry is compared with that of the Average values of n=1.9 and 0=710 kJ/mol were monolithic materials in Fig. 6 at selected initial strain determined from Fig. 3 by using equations(2) and rates. Al,O3 exhibits a brittle behavior-as a conse- quence of the high purity of the starting powder The effect of the layer thickness on the stress-(99.99%[12)- which contrasts with the large ain behavior is shown in Fig. 4 for layered com- ductility exhibited by doped Al,O, where an exten- posites with interface planes oriented parallel to the sive secondary creep regime is attained even at tem- stress axis(LPAR and FPAR), as a function of the peratures as low as 1200 C [19]. Two failure mode initial strain rate. Despite the different volume frac- can be discerned in pure Al2O3. At the higher strain tion of the Al2O, phase in the LPAR(VAL,0, =0. 24)rate[Fig. 6(a)], the sample fails very quickly with and FPAR (A0, =0.47)samples, both microarchi- no plastic deformation prior to failure; the fracture tectures display basically the same behavior. Micro- is dictated by the development of several cracks by LPER Q=710710 670750( kI/mol) 10° 1400°C →60MPa 73 MPa STRAIN(%) ig. 3. Creep curve plotted as log e vs e for a multi posite L with interface planes perpendicular to the applied Fig. 5. SEM micrograph of a laminated composite FPAR after tress(LPER) Several determinations of the stress exponent n deformation(T=1400C, E=40%), showing good interfacial and the activation energy o by stress and temperature changes adhesion between Al,, (left)and ZTA layers. The stress axis
JIME´ NEZ-MELENDO et al.: ALUMINA/ZIRCONIA LAMINATE COMPOSITES 4717 stress of about 200 MPa, independent of the initial strain rate. In contrast, the LPER samples still display a rather steady-state deformation regime after a controlled stress drop. Figure 3 displays a constant load creep test for a LPER specimen showing several stress and temperature changes. Despite the large final strain of the sample (e|60%), no evidence of macroscopic failure was found. The data have been analyzed using the standard high-temperature power law for steady-state deformation [14]: e˙ = Asn d2p exp(2Q/RT) (1) where A is a parameter depending on the deformation mechanism, R is the gas constant, n is the stress exponent, p is the grain size exponent, and Q is the activation energy. From the steady-state strain rate e˙1 before and after e˙2 the change in stress or temperature (Fig. 3), the stress exponent and the activation energy is given, respectively, by [14]: n = log(e˙2/e˙1) log(s2/s1) (2) Q = R T1T2 T22T1 log(e˙2/e˙1) (3) Average values of n = 1.9 and Q = 710 kJ/mol were determined from Fig. 3 by using equations (2) and (3). The effect of the layer thickness on the stress– strain behavior is shown in Fig. 4 for layered composites with interface planes oriented parallel to the stress axis (LPAR and FPAR), as a function of the initial strain rate. Despite the different volume fraction of the Al2O3 phase in the LPAR (VAl2O3 = 0.24) and FPAR (VAl2O3 = 0.47) samples, both microarchitectures display basically the same behavior. MicroFig. 3. Creep curve plotted as log e˙ vs e for a multilayer composite L with interface planes perpendicular to the applied stress (LPER). Several determinations of the stress exponent n and the activation energy Q by stress and temperature changes are shown. Fig. 4. Stress–strain curves at 1400°C for laminated composites L and F with interface planes parallel to the stress axis (LPAR and FPAR) as a function of the initial strain rate e˙o. structural observations of deformed laminated composites (Fig. 5) [12] show a general creep damage of the Al2O3 layers, with very little microcracking by cavity coalescence. The ZTA layers are almost featureless, indicating that the deformation of this phase was accommodated preferentially by grain boundary sliding. In all the cases the interfaces show an excellent adhesion, maintaining their structural integrity (Fig. 5). The mechanical behavior of the laminates with PER and PAR geometry is compared with that of the monolithic materials in Fig. 6 at selected initial strain rates. Al2O3 exhibits a brittle behavior — as a consequence of the high purity of the starting powder (>99.99% [12]) — which contrasts with the large ductility exhibited by doped Al2O3, where an extensive secondary creep regime is attained even at temperatures as low as 1200°C [19]. Two failure modes can be discerned in pure Al2O3. At the higher strain rate [Fig. 6(a)], the sample fails very quickly with no plastic deformation prior to failure; the fracture is dictated by the development of several cracks by Fig. 5. SEM micrograph of a laminated composite FPAR after deformation (T = 1400°C, e = 40%), showing good interfacial adhesion between Al2O3 (left) and ZTA layers. The stress axis is shown by arrows
4718 JIMENEZ-MELENDO et al. ALUMINA/ZIRCONIA LAMINATE COMPOSITES I=1400°C ZTA T=1400°C 160 9 10 2Al2O3[23] LPAR LPER STRAIN (% STRESS (MPa) Fig. 7. Variation in strain rate with stress at 1400 C for mor T=1400°C hic ZTA (squares); the dashed malized to the grain size in the ZTA layer of laminated com- 80 posites. The creep rate of monolithic Al2O3(circles)[23] with the same grain size as the Al2O, layers in the laminates is also 60 compare directly the e-o behavior of monolithic ZTA ZTA with that of the ZTA layers in the laminates, the strain ate has been normalized to the same d( dashed line E。=4x10s1 in Fig. 7) by using p=2 [15] in equation(1).Figure LPAR 7 also displays the steady-state creep rate at 1400oC PER of monolithic Al2O, with d=3.2 um [23], similar to (b) the grain size of the Al,O3 layers in the present lami- STRAIN () Fig. 6. Comparison of stress-strain behavior at 1400@C of lam 4. DISCUSSION nated composites LPAR and LPER wit Al2O3 and ZTA at(a)b。=2×10-3/s,(b)=4×10-6/s The most striking result of the stress-strain curves displayed in Fig. 6 is that the layered materials simul- taneously exhibit large ductility and high flow stress coalescence of two-grain boundary cavities. At the As shown below, such a behavior is not expected lower strain rate (b)1. the failure is delayed and from a simple composite creep model based on the eventually occurs by the coalescence of creep dam- bulk properties of the two constituent materials [24] age, consisting of a large density of two-grain bound- indicating clearly that the layer interfaces play a fun- ary cavities that are homogeneously distributed damental role in the overall mechanical behavior of throughout the sample; such a microstructure is very the layered composites. French et al. [24]have similar to that observed in the Al,O3 layers of lami duplex microstructures nated materials(Fig. 5)[12]. These two failure modes by assuming an isostress(alternated plates of each have been reported previously in pure Al, O, at high phase aligned perpendicularly to the applied stress) and low stresses, respectively [20, 21]. On the other or isostrain(plates are parallel to the applied stress) hand,monolithic ZTA shows a large ductility(Fig model, which coincide with the geometries PER and 6)because of the enhancement of grain boundary PAR, respectively, of the present laminates. In sliding by the addition of the tetragonal ZrO2 phase; PER configuration, each layer is submitted to the a tensile elongation to failure of 550% has been same stress(isostress condition), and the overall com- reported in a 10 vol. ZrO2-dispersed Al, O, com- posite strain rate Ec can be described by the relation posite[18]. Figure 7 illustrates the variation of ste- [24] lithicZTg eep rate with stress at 1400%C for mono- The fit of experimental data yields a stress Ee=VzTAEzTA valo, e, exponent n=2, the natural value for superplasticity in metals [22]. An activation energy (=720 kJ/mol was measured by temperature changes, in agreement where V is the volume fraction of each phase in the with previous results in ZTA composites with ZrO, laminated composite (zTA =0.76 in the LPER additions between 20 and 80 vol %[15]. In order to samples). As noted in the previous section, grain
4718 JIME´ NEZ-MELENDO et al.: ALUMINA/ZIRCONIA LAMINATE COMPOSITES Fig. 6. Comparison of stress–strain behavior at 1400°C of laminated composites LPAR and LPER with that of monolithic Al2O3 and ZTA at (a) e˙o = 2×1025 /s, (b) e˙o = 4×1026 /s. coalescence of two-grain boundary cavities. At the lower strain rate [Fig. 6(b)], the failure is delayed and eventually occurs by the coalescence of creep damage, consisting of a large density of two-grain boundary cavities that are homogeneously distributed throughout the sample; such a microstructure is very similar to that observed in the Al2O3 layers of laminated materials (Fig. 5) [12]. These two failure modes have been reported previously in pure Al2O3 at high and low stresses, respectively [20, 21]. On the other hand, monolithic ZTA shows a large ductility (Fig. 6) because of the enhancement of grain boundary sliding by the addition of the tetragonal ZrO2 phase; a tensile elongation to failure of 550% has been reported in a 10 vol.% ZrO2-dispersed Al2O3 composite [18]. Figure 7 illustrates the variation of steady-state creep rate with stress at 1400°C for monolithic ZTA. The fit of experimental data yields a stress exponent n>2, the natural value for superplasticity in metals [22]. An activation energy Q = 720 kJ/mol was measured by temperature changes, in agreement with previous results in ZTA composites with ZrO2 additions between 20 and 80 vol.% [15]. In order to Fig. 7. Variation in strain rate with stress at 1400°C for monolithic ZTA (squares); the dashed line is the strain rate normalized to the grain size in the ZTA layer of laminated composites. The creep rate of monolithic Al2O3 (circles) [23] with the same grain size as the Al2O3 layers in the laminates is also shown. compare directly the e˙–s behavior of monolithic ZTA with that of the ZTA layers in the laminates, the strain rate has been normalized to the same d (dashed line in Fig. 7) by using p = 2 [15] in equation (1). Figure 7 also displays the steady-state creep rate at 1400°C of monolithic Al2O3 with d = 3.2 µm [23], similar to the grain size of the Al2O3 layers in the present laminates. 4. DISCUSSION The most striking result of the stress–strain curves displayed in Fig. 6 is that the layered materials simultaneously exhibit large ductility and high flow stress. As shown below, such a behavior is not expected from a simple composite creep model based on the bulk properties of the two constituent materials [24], indicating clearly that the layer interfaces play a fundamental role in the overall mechanical behavior of the layered composites. French et al. [24] have developed a creep model for duplex microstructures by assuming an isostress (alternated plates of each phase aligned perpendicularly to the applied stress) or isostrain (plates are parallel to the applied stress) model, which coincide with the geometries PER and PAR, respectively, of the present laminates. In the PER configuration, each layer is submitted to the same stress (isostress condition), and the overall composite strain rate e˙c can be described by the relation [24]: e˙c = VZTAe˙ZTA + VAl2O3 e˙Al2O3 (4) where V is the volume fraction of each phase in the laminated composite (VZTA = 0.76 in the LPER samples). As noted in the previous section, grain
JIMENEZ-MELENDO et al. ALUMINA/ZIRCONIA LAMINATE COMPOSITES 4719 boundary sliding is the primary deformation mech- For LPER composites(d=2 mm, h=300 um), equ anism in monolithic ZTA, characterized by n=2(Fig. ation (6) gives o/bulk= 2.3, which is in good 7)and @=720 kJ/mol. Monolithic Al,O, exhibits a agreement with the experimental value noted above much lower activation energy (=400-500 kJ/mol This result corroborates that sequential slip casting is [23, 25-28] and a stress exponent between I and 2 an adequate technique to obtain strong junctions, depending on grain size, n is around 2 for dsl um which in turn increases the creep resistance of the [25-27 and decreases towards I with increasing laminar composites grain size(Fig. 7)[23, 25-28]. Since the creep rate For the PAR geometry, the strain and the strain of monolithic ZTA is much faster than that of mono- rate are the same for each layer(isostrain case). The lithic AlO, in the stress range studied(Fig. 7)and composite creep model [24] predicts in this case that VzTA3VALO,, the strain rate of the laminate [equ- the strain rate of the laminate will be controlled by ation(4)) can be simplified the phase with higher creep strength(AlO present case, Fig. 7). The composite stress o is then Ec= vztaezta (5)given by [24] The values of n=1.9 and 0=710 kJ/mol found in LPER composites(Fig 3)indicate that the reep rate Ec is controlled by the ZTA layers, in agree- where oAl0, is the stress supported by the Al O3 lay ment with equation (5). However, the corresponding ers. For the PAR composites strainedat flow stress o of the laminate is higher than that E.=4x10-6/s(EzTa =EAlo, for this configuration), expected from the model. For the LPER composite the composite flow stress is o=80 MPa [Fig. 6(b) deformed at Ec=4x10-6/s [Fig. 6(b)1, equation (5)and equation(7) predicts that the Al2O, layers would gives EzTA=5x10-Is, which corresponds to a flow be submitted to a stress of 330 MPa in the lpar stress azTA=22 MPa for the ZTa phase(Fig. 7). samples(VALo, =0.24)and 170 MPa in the FPAR composite is 62 MPa [Fig. 6(b)1, which is 2.8 times lithic Al O, at this strain rate [-100 MPa, Fig. 6(b)] higher than the expected value. According to the However, the composites exhibit extended steady model, the laminated composite is assumed to be sim- state regimes(Fig. 4), with no evidence of macro- ply a superposition of layers with no interaction scopic failure. Although this geometry is more com- between them [equation(4). The breakdown of the plex to analyze quantitatively than the PER configur model can thus be explained on the basis of the con- ation, the previous result, along with the trains imposed by the more creep resistant layers on microstructural observations(Fig. 5)[121, show that the deformation of the softer layers through interface the presence of the interfaces is again the key feature bonding. Assuming perfectly bonded interfaces, the that modifies the creep properties of the A1,O, layers in-plane strains must be the same at the layer inter- with respect to bulk Al203, imparting damage toler faces, resulting in an in situ flow stress of the Zta ance to the laminar composites layers higher than that exhibited by monolithic(bulk) ZTA. This effect has been previously reported in composite materials with ductile phases strongly 5 CONCLUSIONS bonded to brittle matrixes(WC-Co, WC-Ni, Al, O3 AL, etc. ) where the in situ flow stress of the ductile The mechanical properties of Al2 0, /(85 vol% binder phase is several times higher than its bulk flow Al20 +15 vol. ZrO, doped with 3 mol% Y,oj)lay- stress value(29-31). The present situation is equival. ered composites produced by sequential slip casting ent to the problem of deformation of a thin ductile have been investigated under compression at 1400C layer(ZTA) between rigid platens (Al,0, ) with a The composites were stressed both parallel and per- strongly bonded interface between the layer and the pendicular to the layer planes. After testing, the layer platens [29, 32]. In such a loading situation, the effec- interfaces maintain their structural integrity. The tive flow stress off of the ductile phase Is given as: comparison with monolithic Al,O, and ZTA pro- duced by the same processing technique shows that the laminated composites exhibit enhanced ductility (6)(characteristic of monolithic ZTA)and creep resist ance (characteristic of monolithic Al2O3)simul- by composite creep model based on dual where oulk is the bulk flow stress in the absence of properties of the two constituent ds. The y constraint, k is the maximum shear stress which improvement in mechanical properties is essentially is taken as 0.577 [29, 32], and d and h are respect- related to the presence of strong interfaces in the ively the width and the thickness of the ductile la
JIME´ NEZ-MELENDO et al.: ALUMINA/ZIRCONIA LAMINATE COMPOSITES 4719 boundary sliding is the primary deformation mechanism in monolithic ZTA, characterized by n>2 (Fig. 7) and Q = 720 kJ/mol. Monolithic Al2O3 exhibits a much lower activation energy Q = 400–500 kJ/mol [23, 25–28] and a stress exponent between 1 and 2 depending on grain size; n is around 2 for d#1 µm [25–27] and decreases towards 1 with increasing grain size (Fig. 7) [23, 25–28]. Since the creep rate of monolithic ZTA is much faster than that of monolithic Al2O3 in the stress range studied (Fig. 7) and VZTA>3VAl2O3 , the strain rate of the laminate [equation (4)] can be simplified to: e˙c = VZTAe˙ZTA (5) The values of n = 1.9 and Q = 710 kJ/mol found in LPER composites (Fig. 3) indicate that the composite creep rate e˙c is controlled by the ZTA layers, in agreement with equation (5). However, the corresponding flow stress sc of the laminate is higher than that expected from the model. For the LPER composite deformed at e˙c = 4×1026 /s [Fig. 6(b)], equation (5) gives e˙ZTA = 5×1026 /s, which corresponds to a flow stress sZTA = 22 MPa for the ZTA phase (Fig. 7). This value would be in turn equal to sc (isostress condition), but the actual flow stress of the laminated composite is 62 MPa [Fig. 6(b)], which is 2.8 times higher than the expected value. According to the model, the laminated composite is assumed to be simply a superposition of layers with no interaction between them [equation (4)]. The breakdown of the model can thus be explained on the basis of the constrains imposed by the more creep resistant layers on the deformation of the softer layers through interface bonding. Assuming perfectly bonded interfaces, the in-plane strains must be the same at the layer interfaces, resulting in an in situ flow stress of the ZTA layers higher than that exhibited by monolithic (bulk) ZTA. This effect has been previously reported in composite materials with ductile phases strongly bonded to brittle matrixes (WC–Co, WC–Ni, Al2O3– Al, etc.), where the in situ flow stress of the ductile binder phase is several times higher than its bulk flow stress value [29–31]. The present situation is equivalent to the problem of deformation of a thin ductile layer (ZTA) between rigid platens (Al2O3) with a strongly bonded interface between the layer and the platens [29, 32]. In such a loading situation, the effective flow stress seff of the ductile phase is given as: seff sbulk = 1 + k 3S d hD (6) where sbulk is the bulk flow stress in the absence of any constraint, k is the maximum shear stress which is taken as 0.577 [29, 32], and d and h are respectively the width and the thickness of the ductile layer. For LPER composites (d = 2 mm, h = 300 µm), equation (6) gives seff/sbulk = 2.3, which is in good agreement with the experimental value noted above. This result corroborates that sequential slip casting is an adequate technique to obtain strong junctions, which in turn increases the creep resistance of the laminar composites. For the PAR geometry, the strain and the strain rate are the same for each layer (isostrain case). The composite creep model [24] predicts in this case that the strain rate of the laminate will be controlled by the phase with higher creep strength (Al2O3 in the present case, Fig. 7). The composite stress sc is then given by [24]: sc = VAl2O3 sAl2O3 (7) where sAl2O3 is the stress supported by the Al2O3 layers. For the PAR composites strained at e˙c = 4×1026 /s (e˙ZTA = e˙Al2O3 for this configuration), the composite flow stress is sc>80 MPa [Fig. 6(b)], and equation (7) predicts that the Al2O3 layers would be submitted to a stress of 330 MPa in the LPAR samples (VAl2O3 = 0.24) and 170 MPa in the FPAR samples (VAl2O3 = 0.47). In both cases, these values are much higher than the fracture stress for monolithic Al2O3 at this strain rate [|100 MPa, Fig. 6(b)]. However, the composites exhibit extended steadystate regimes (Fig. 4), with no evidence of macroscopic failure. Although this geometry is more complex to analyze quantitatively than the PER configuration, the previous result, along with the microstructural observations (Fig. 5) [12], show that the presence of the interfaces is again the key feature that modifies the creep properties of the Al2O3 layers with respect to bulk Al2O3, imparting damage tolerance to the laminar composites. 5. CONCLUSIONS The mechanical properties of Al2O3/(85 vol.% Al2O3+15 vol.% ZrO2 doped with 3 mol% Y2O3) layered composites produced by sequential slip casting have been investigated under compression at 1400°C. The composites were stressed both parallel and perpendicular to the layer planes. After testing, the layer interfaces maintain their structural integrity. The comparison with monolithic Al2O3 and ZTA produced by the same processing technique shows that the laminated composites exhibit enhanced ductility (characteristic of monolithic ZTA) and creep resistance (characteristic of monolithic Al2O3) simultaneously. This behavior cannot be explained by a composite creep model based on the individual properties of the two constituent materials. The improvement in mechanical properties is essentially related to the presence of strong interfaces in the laminated composites
4720 JIMENEZ-MELENDO et al. ALUMINA/ZIRCONIA LAMINATE COMPOSITES Acknowledgements-This research has been supported 15. Clarisse. L. Baddi. R. Bataille. A J, Duclos, CICYT No MAT 97/0562-C02-01(Minist Educa R and Vicens, J. Acta mater. 1997, 45, 3843 Ciencia, Madrid, Spain). The authors would like to thank Dr 16. Lange, F. F. and Hirlinger, M dm Ceram. Soc Jose S Moya (ICM-CSIC, Madrid, Spain) for kindly supplying the samples 17. French, F D, Harmer, M. P, Chan, H. M. and Miller, G J. Am. Cera. Soc. 199 REFERENCES Scripta mater., 1998, 38, 33 19. Cannon, R. M, Rhodes, w.H. and Heuer, A.H.J.Amm. I. Harmer. M. P. Chan. H. M. and Miller. G. A.. dm. Ceram. Soc.. 1980 63. 4 Ceram.Soc,1992,75,1715. 20. Dalgleish D. J. Slamovich E. B and Evans. A G.JAm. 2. Moya, J. S. Adv. Mater. 1995, 7, 18 Ceram. Soc. 1985. 68. 575 3. Marshall, B. Am. Ceram. Soc. Bull., 1992, 71, 969. 21. Robertson, A G, Wilkinson, D. G. and Caceres, C. H.J. 4. Requena, J, Moreno, R. and Moya, J. S.J. Am Ceram. 4m. Ceram Soc., 1991, 74, 91 Soc.1989.72.1511 22. Nieh, T. G, Wadsworth, J. and Sherby, O. D. Superplas- 5. Chartier, T. and Rouxel, T. J. Eur. Cera. Soc., 1997 city in Metals and Ceramics, Cambridge University 17,299 Press. Cambridge. 1997 6. Jimenez-Melendo, M, Clauss, C, Dominguez-Rodriguez, 23. De Arellano-Lopez, A.R., Cumbrera, F. L, Dominguez. A, De Portu, G, Roncari, E. and Pinasco, P. Acta mater. Rodriguez, A, Coretta, K. C. and Routbort, J. L.J. Am. 998.46.3995 Ceram. Soc,1990,73,1297 7. Prakash, O, Sarkar, P and Nicholson, P S.J. m. Ceram. 24. French, F. D, Zhao, J, Harmer, M. P Chan, H. M. and Miller. G.A. Am Ceram. Soc. 1994.77 2137 8. Marshall, D. B. and Ratto, J. J.J. Am Ceram. Soc., 1991, 25. Chen, I.-w. and Xue, L. A.J.Am. Ceram. Soc., 1990 74.2979 9. Kuo, D.H. and Kriven, W. M.J. Am. Ceram Soc., 1997, 26. Wakai, F, Nagano, T. and lga, T.J. Am. Ceram. Soc., 1997.80.2361 10. Oeschner, M, Hillman, C and Lange, F. F. J./m. Ceram. 27. Yoshida, H, Ikuhara, Y and Sakuma, T.J. Mater:. Res c,1996,79,1834 998,13,2597 11. Lange,F.F.Si.ld,1998,63,33 28. Cannon, R. M, Rhodes, W. H. and Heuer, A. H.J. Am. 12. Jimenez-Melendo, M Clauss, C, Dominguez-Rodriguez, Ceram. Soc. 1980. 63.4 A, Sanchez-Herencia, A. J. and Moya, J. Am. Cere 29. Ravichandran, K. S. Acta metall. mater., 1994 Soc.1997.80.2126 30. Ravichandran, K. S tall. mater:. 1992 13. Gervais, H, Pellissier, B and Castaing, J. Rev. Int Htes. 31. Okuyama, T. Mater. Sci. En 家院 Temp. Refract., 1978, 15, 43 14. Bretheau, T, Castaing, J, Rabier, J and Veyssiere, P Adv. 32. Schroede Webster, D. A.J. Appl. Mech., 1949, Phs.,1979,28,835
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