Fabrication and creep properties of eutectic-composition Al2O3/YAG/YSZ sintered composites

1 Fabrication and creep properties of eutectic-composition Al2O3/YAG/YSZ sintered composites F.A. Huamán-Mamani*, C. Jiménez-Holgado** and M. Jiménez-Melendo*** * Departamento de Ciencias Naturales, Universidad Católica San Pablo. Arequipa, Peru ** Department of Chemistry, School of Science, University of Ioannina. 45110 Ioannina, Greece ***Departamento de Física de la Materia Condensada, Universidad de Sevilla. 41080 Sevilla, Spain ***Corresponding author: mel[email protected] ABSTRACT Three-phase alumina/YAG/yttria-stabilized cubic zirconia (YSZ) composites were fabricated by a solid-state reaction route starting from commercial powders of Al2O3, Y2O3 and monoclinic ZrO2. The final phases Al2O3, YAG and YSZ were obtained after calcination of the powder mixtures at 1400 °C. Dense bulk composites were obtained after sintering, with a homogeneous microstructure of fine and equiaxed grains with sizes of 1 m. Compressive mechanical tests were performed at 1300 – 1450 °C in air at constant load and at constant initial strain rate. A brittle-to-ductile transition was found with increasing temperature. Grain boundary sliding is the main deformation mechanism in the ductile regime, characterized by a stress exponent of 2 and by the absence of dislocation activity and changes in grain morphology. Alumina seems to be the rate-controlling phase owing to the improvement in creep resistance by the presence of yttrium and zirconium of the other two phases. 2 Keywords: Alumina; Composite material; Creep; Microstructure; Grain boundary sliding. 3 1. Introduction Alumina-based composites are probably the most used engineering oxide ceramics nowadays in structural applications, particularly at high temperatures, because of their excellent chemical stability and mechanical properties. Although pure monolithic alumina is inherently brittle, with a low fracture toughness at low and high temperatures [1,2], its mechanical properties can be considerably enhanced by the incorporation of other phases in composite structures, such as fibers, particulates, layers, etc. In particular, particulate composites formed by equiaxed and approximately similar-sized particles exhibit improved thermal and chemical stability relative to their single-phase constituents, thus maintaining optimal mechanical properties of strength, toughness, and thermal shock resistance at elevated temperatures. In this way, dual phase Al2O3/ZrO2 [3–8] and Al2O3/YAG [4,9–11] composites have been produced with higher strength, fracture toughness, and creep resistance than their single-phase counterparts. Remarkably, these fine-grained composites exhibit metal-like superplasticity, which contrasts with the premature failure of undoped monolithic alumina owing to the extensive grain growth and cavitation during high-temperature deformation [2,12]. Directionally-solidified alumina-based eutectic ceramics grown from the melt have received much attention in the last years because of their excellent mechanical properties even at temperatures close to the eutectic temperature [13–19], which derive from the special lamellar microstructures obtained during solidification. In particular, directionally-solidified ternary Al2O3/YAG/ZrO2 eutectic melts have been shown to have a fracture toughness as high as 8 MPa.m1/2 at room temperature [16], and the retention of large flexural, tensile and compressive strengths up to temperatures close to the eutectic temperature [13,14]. The fabrication routes necessary to produce directionally-solidified eutectic ceramics are, however, economically 4 expensive, no suitable for mass production and impose serious restrictions for obtaining bulk components with custom shapes and sizes. Furthermore, the physical properties of these eutectic ceramics are very anisotropic, depending strongly on the solidification direction. These drawbacks can be mitigated, at least partially, by using sintered particulate composites produced by conventional solid-state reaction routes. Very few studies are, however, concerned with Al2O3/YAG/ZrO2 (thereinafter AYZ) sintered composites, which have been fabricated with different phase contents: equivolumetric (33 vol%) [20], eutectic-composition [21] and 5 vol% YAG + 5 vol% ZrO2 particulate-reinforced Al2O3 [22]. Precisely on this last material, fabricated by surface modification of the alumina powder with inorganic precursors of the second phases, Palmero et al. [23] have reported the only study addressing the high-temperature mechanical behavior of AYZ sintered ceramics; by means of four-point bending tests from room temperature up to 1500 °C, the authors found an enhanced deformability at temperatures above 1400 °C (though the maximum strain was limited to 4% due to restrictions of the experimental setup). The aim of the present study is therefore twofold: first, to fabricate and characterize the crystalline phases and microstructure of three-phase AYZ composites with the eutectic composition produced by a conventional ceramic processing route; and second, to investigate their creep response at high temperatures by constant crosshead speed and constant load tests in correlation with microstructural observations, in order to assess the microscopic mechanisms involved in the plastic deformation. The creep of the threephase composite is compared to that of the single-phase constituents. 2. Experimental procedure 2.1. Starting materials 5 Samples of polycrystalline AYZ composites with the eutectic composition (65.8 mol% Al2O3, 15.6 mol% Y2O3, 18.6 mol% monoclinic ZrO2 [24]) were produced via a conventional solid-state reaction route. Appropriate amounts of high-purity commercial powders (  99.99%, SigmaAldrich) of Al2O3, Y2O3 and monoclinic ZrO2 were dry ball-milled in agate media for 1 h at 150 rpm using a planetary ball mill (Pulverisette 6, Fritsch, Germany). The resulting powders were calcined at different temperatures between 1200 and 1600 °C in air for 10 h and reground again using the same milling conditions as for the initial mixtures to eliminate possible agglomerates. A calcination temperature as high as 1600 °C was selected because it was reported [25] that the formation of the YAG phase started at 1300 °C in Al2O3/Y2O3 powder mixtures processed by solidstate reaction, obtaining the single-phase YAG only at temperatures above 1600 °C; below this temperature, the intermediate phases YAlO3 (YAP, orthorhombic perovskite) and Y4Al2O9 (YAM, monoclinic) were also present in the mixtures from 1100 to 1400 °C and from 1200 to 1600 °C, respectively. The particle size distribution of the various powder mixtures was analyzed by laser light scattering (Malvern Mastersizer 2000 with distilled water as carrier medium), resulting a number-based median particle size d50 of 1.20 m for the as-received powder, 0.92 m after the first milling and finally 0.76 m after calcination and second milling. The calcined powders were analyzed by X-ray diffraction in order to determine the nature and content of the crystalline phases (see below). The powder mixtures were uniaxially pressed at 150 MPa into 20-mm diameter pellets and then isostatically cold pressed at 210 MPa. The resulting green pellets were sintered in air at 1500 °C for 10 h with low heating and cooling rates of 5 °C/min. The bulk density of the composites was determined from weight/dimensions measurements. 2.2. Structural and microstructural characterization 6 X-ray structure analyses were systematically performed on the calcined AYZ powders as well as in the sintered composites in order to ensure the presence of the required phases. X-ray powder diffractograms were obtained using a Bruker D8 Advance A25 X-ray diffractometer with Cu K radiation and Ni filter, equipped with a scintillation detector in  – 2 Bragg-Bentano configuration (X-ray Laboratory, CITIUS, University of Sevilla, Spain). A continuous scan mode was used to collect 2 data in the 10 – 120° range in steps of 0.015° and a scanning speed of 1.8 °/min. Collected X-ray spectra were first processed by the Le Bail method followed by Rietveld refinement using the TOPAS 4.2 Bruker AXS software package to quantify the nature and amount of the different crystalline phases in the powders. The microstructural characterization of as-sintered and deformed composites was carried out using high-resolution scanning electron microscopy (HRSEM), particularly in back-scattered electron image mode to discern the different crystalline phases, and transmission electron microscopy (TEM) (Microscopy Service, CITIUS, University of Sevilla, Spain). In order to reveal the grain boundaries for SEM observations, sections were cut from the samples and mechanically polished using up to 0.25 m-grade diamond paste, and then thermally etched at 1300 °C for 2 h in air. The relevant morphological parameters, grain size d (taken as the equivalent planar diameter d = (4 x grain area/)1/2), form factor F (defined as F = 4 x grain area/(grain perimeter)2) and preferential orientation, were measure by using a semiautomatic image analyzer; over 300 grains were measured for each phase. Thin films for TEM observations were obtained from the as-fabricated and deformed samples following a classical procedure of grinding and ion-thinning until electron transparency of sliced sections. Elemental composition analysis was performed by energy dispersive X-ray spectroscopy (EDS) in both SEM and TEM to characterize the various phases present in the composites. Wavelength dispersive X-ray fluorescence spectroscopy (WDXRF Malvern Pananalytical Axios, Germany, X-Ray Laboratory, 7 CITIUS, University of Sevilla, Spain) was also used to determine the chemical composition of the sintered composites. 2.3. Mechanical tests Prismatic specimens of 5 x 3 x 3 mm in size were cut from the sintered pellets and used for mechanical experiments. Two types of compressive tests were carried out in air at temperatures T between 1300 and 1450 °C: (i) under constant load in a creep machine, at nominal stresses  between 50 and 120 MPa; the raw data, instantaneous specimen length l(t) vs. time t, were plotted as log −  curves, where  is the true strain (= ln(lo/l(t)), with lo the initial length) and  = d /dt is the instantaneous strain rate. And (ii) at constant cross-head speed in an universal testing machine at different initial strain rates o ; the recorded data, load vs. time, were analyzed in  -  curves, where  and  are the true stress and the true strain, respectively. After the completion of tests, the specimens were cooled under load to preserve the final microstructure. The mechanical data were analyzed using the standard high-temperature power law for steadystate deformation [26]: np A d exp( Q /RT) −  =  − (1) where A is a parameter depending on the deformation mechanism, n is the stress exponent, p is the grain size exponent, Q is the activation energy for creep and R is the gas constant. The parameters n and Q, characteristics of the deformation mechanism, were measured from stress and temperature changes, respectively. The specimens were typically deformed to total strains of 50-60 % (unless premature failure occurred) for subsequent microstructural observations. 8 3. Results and discussion 3.1. Phase composition of the calcined powders Figure 1 shows the X-ray diffractograms of the as-milled and calcined powders at temperatures between 1200 and 1600 °C. The patterns are rather complex because up to seven different phases could be identified at intermediate temperatures. Every peak in the diffractograms was indexed according to one of the following PDF-2002 database reference patterns (International Centre for Diffraction Data, ICDD): α-Al2O3 (No. 05-0712), Y2O3 (No. 89-5591), monoclinic ZrO2 (No. 36-0420), Y4Al2O9 (YAM, No. 34-0368), YAlO3 (YAP, No. 74-1334), Y3Al5O12 (YAG, No. 330040) and Zr0.72Y0.28O1.862 (YSZ, No. 77-2112). The volume fractions of the crystalline phases were quantified by Rietveld refinement, and are gathered in Table 1. In the as-ground powder (Figure 1 and Table 1), the three original phases Al2O3, Y2O3 and monoclinic ZrO2 were identified, indicating that no additional phases formed during the milling process. At 1200 °C, seven different phases were detected: Al2O3, Y2O3, monoclinic and cubic ZrO2, and the three phases of the Al2O3/Y2O3 system, YAM, YAP and YAG. YAM and YAP are undesired phases because they exhibit poorer mechanical strengths than YAG. Regarding the zirconia phase, the Zr0.72Y0.28O1.862 reference pattern (No. 77-2112) provided the best match to the observed XRD profiles. After Rietveld refinement, a lattice parameter of a = 5.156 ± 0.003 Å was obtained, corresponding to an yttrium composition of 0.27 ± 0.02 at% (equivalent to 15.6 ± 2.0 mol% Y2O3 fully-stabilized cubic ZrO2) according to the Vegard’s law [27]. This cubic phase derives from the original low-symmetry monoclinic phase through alloying with substitutional yttrium cations provided by the Y2O3 phase. The yttria molar content agrees with the 9 compositions of 15.5 and 16 mol% Y2O3 reported in directionally-solidified AYZ eutectic composites [17] and in 5 vol% ZrO2 + 5 vol% YAG-reinforced Al2O3 sintered composite [22], respectively. Regarding the evolution of the yttrium aluminate phases, the YAM and YAP phases gradually disappear with increasing the calcination temperature (Figure 1 and Table 1). Single-phase YAG is completely formed at 1400 °C, which is 200 °C lower than the critical temperature reported by Ikesue et al. [25] in the Al2O3/Y2O3 system. This result suggests that the presence of the zirconia phase inhibits the coarsening of the alumina particles, resulting in faster diffusion processes. Higher calcination temperatures practically do not introduce further phase changes. The sintered compacts were finely crushed and powdered for further XRD analyses, which were found to be virtually identical to the calcined powders. From Rietveld refinement, the sintered composites consist of 39.1 vol% Al2O3, 40.1 vol% YAG and 20.8 vol% YSZ (Table 1), indicating the absence of significant phase evolution during the sintering step. Table 1 also collects the phase contents expected by assuming a complete reaction of the oxide precursor powders as well as the values reported in directionally-solidified AYZ eutectic melts [14] and in eutecticcomposition AYZ sintered polycrystals [21] (in this case, the zirconia phase was stabilized in the tetragonal structure due to the use of 3 mol% yttria-stabilized tetragonal zirconia 3YTZP powder as precursor), showing a fair agreement with the present values estimated from XRD spectra. 3.2. Microstructural characterization of the sintered composites Figure 2 depicts representative SEM and TEM micrographs of the as-sintered composites. Figure 2(a) corresponds to a secondary electron SEM image showing a homogeneous distribution of 16 recovery-controlled creep mechanism, which needs the establishment or the modification of a dislocation substructure prior to the attainment of the steady state. These features highlight the large deformability of the AYZ composites, based on the superior microstructural stability of the phases. As noted above, a value of n = 2 has been systematically reported in fine-grained metals and ceramics [7,8,28–30,32,33], where grain boundary sliding is the dominant deformation mechanism. This is also the value predicted by most theoretical and semiphenomenological models developed to explain the superplastic behavior of fine-grained materials [29]. The set of experimental results found in the present work: a stress exponent close to 2, extended steady states of deformation, absence of changes in grain size and morphology after large strains and the absence of dislocation substructures and creep transients, allow to conclude that the creep of the three-phase AYZ sintered composites takes place by grain boundary sliding. 3.5.2. Activation energy Q For grain boundary sliding, the steady-state deformation is usually controlled by the diffusion necessary to accommodate the local stresses generated during sliding. The creep activation energy Q (Eq. (1)) can be thus identified with the diffusion energy of the slowest moving species in the material along the faster path. Q has been measured by upand down-temperature changes during steady-state creep, as illustrated in Figure 6(b) for changes between 1400 and 1450 °C at a fixed initial stress of 103 MPa. An average value of Q = 680 ± 50 kJ/mol was found, which is substantially higher than the creep energies reported for the single-phase constituents: 400-500 kJ/mol in Al2O3 [4,33,37,38], 430-520 kJ/mol in YSZ [4,28,39] and 530-580 kJ/mol in YAG [4,40]. These creep energies are consistent, though with some scatter, with the values reported 17 for cation diffusion (the slowest moving ionic species) in cubic zirconia [28,41] and YAG [40]. The mechanisms of diffusion in alumina are however still uncertain [42–44] due to its highly extrinsic character and also because the diffusion energies along grain boundaries are substantially higher than in volume, contrary to what is generally found in other materials. Despite the scatter, it must be noted that oxygen diffuses several orders of magnitude faster than aluminum both in bulk and along (sub)grain boundaries when measured on the same suite of crystals [45– 47], thus becoming the rate-limiting species in alumina. For the sake of clarity, Table 3 compiles the creep and the corresponding diffusion energies for the single-phase monoliths, along with creep energies for binary alumina/zirconia [4] and alumina/YAG [4,48] composites. 3.6. Strain rate-controlling phase in the AYZ composite This discrepancy between the creep energies of the AYZ composite Q = 680 kJ/mol and the single-phase constituents is also encountered when comparing their strain rates. Figure 7 shows the variation of the grain size-compensated strain rate with stress at 1400 °C for the AYZ composite (black solid line) and the monolithic counterparts (slashed areas). Strain rates have been normalized with a grain size exponent p = 2 in Eq. (1), intermediate between the values of 1 and 3 reported for grain boundary sliding [7,10,28,29,33,37] (in fact, the discussion is not affected by the choice of p due to the small range of grain sizes involved). It can be seen that the strain rates of the single-phase monoliths are similar among them, but higher than that of the AYZ composite. This unexpected result has been also reported in binary Al2O3/YAG [4,48] and Al2O3/ZrO2 composites [4] (AY and AZ, respectively, in Figure 7). The discrepancy can be understood based on the highly extrinsic character exhibited by alumina noted previously; even a few ppm of impurities/dopants can change drastically the strain rate 18 because of the hindering of the diffusion processes. It has been reported a decrease by one order of magnitude in the strain rate of Zr-doped alumina with respect to undoped alumina [33,37], and by two orders of magnitude when doped with Y [4,31,37] (Figure 7); the reduction was even more drastic, by a factor of 400, when doubly doped with yttrium and zirconium [31]. Such a behavior has been related to the segregation of dopant cations along the alumina grain boundaries, which restrains the diffusion processes along this path and then retards the creep rate. Cation segregation at grain boundaries has been reported in different ceramic oxides such as yttrium-stabilized zirconia [28], yttrium-doped barium zirconate [30] and other perovskite oxides for solid oxide fuel cells [49]. The intergranular fracture mode of the alumina grains noted above in the brittle region is most likely to occur by the same effect. Furthermore, the enhanced creep resistance exhibited by cation-doped monolithic Al2O3 relative to undoped alumina was systematically accompanied by a significant increase in the creep energy, raising from Q = 400 – 500 kJ/mol in nominally pure alumina up to 700 – 800 kJ/mol in Yand Zr-doped alumina [4,33,37,38] (Table 3). A similar increase in Q was reported in binary alumina/zirconia and alumina/YAG composites relative to the monolithic constituents (Table 3). The self-diffusion activation energy of oxygen measured in pristine and yttrium-doped alumina also follow the same trend observed in creep (Table 3). Based on the previous results, it is reasonable to assume that the creep behavior of the alumina phase in the three-phase AYZ composite is modified by the presence of the yttrium and zirconium ions provided by the other two phases, which most likely segregated to the alumina grain boundaries during the sintering step. The load transfer between grains can be now modelled by the isostrain condition (the harder phase controls the overall strain rate) [4,9,12] if the three phases are coupled in the composite, as expected for particulate interpenetrated 19 phases. In this case, the strain and strain rate are the same for each phase, and the composite (applied) stress  is given by  = ∑𝑓𝑖σ𝑖 (i = Al2O3, YAG, YSZ), where fi and i are the volume fraction and stress supported for each phase. Assuming that alumina is the rate-controlling phase in the composite, which is present with a volume fraction of f = 39.1% (Table 1), the model predicts that this phase would carry a stress of /f ≅ 2.6, equivalent to a creep rate of (1/f)n ≅ 7 (with n = 2, the stress exponent) times lower than that of the composite (shown by a black dotted lined in Figure 7). The model predicts thus a reduction of the alumina strain rate by at least a factor of 200 compared to undoped alumina (Figure 7), in close agreement with the behavior reported for yttriumand zirconium-doped alumina. This result, along with the change in fracture mode and the restrained grain growth of the alumina phase in the composite with respect to undoped monolithic alumina [50], suggests that alumina is strongly influenced by the presence of the cations of the other two phases, changing the nature and properties of their grain boundaries. Furthermore, the present results indicate that Al2O3/Y3Al5O12/YSZ sintered composites could compete with melt-growth eutectic analogs by tailoring the microstructure through the control of the processing parameters to achieve the desired properties. 4. Conclusions Al2O3/YAG/cubic ZrO2 sintered composites with the ternary eutectic composition have been successfully fabricated by solid-state reaction. Starting with Al2O3, Y2O3 and monoclinic zirconia powders as precursors, X-ray diffraction analysis indicates that the final phases of Al2O3, YAG and Y2O3 fully-stabilized cubic zirconia, without other intermediate phases, are obtained at a calcination temperature of 1400 °C. Bulk composites with a relative density of 98.5% were obtained after sintering at 1500 °C in air for 10 h. The volume fractions measured for each phase are consistent with the values found in melt-growth AYZ eutectic composites. The 20 microstructure is formed by a homogeneous distribution of equiaxed and fine grains with average sizes of 1 m for the three phases. Compressive mechanical tests were performed at constant cross-head speed and at a constant load in air at 1300 – 1450 °C. A gradual brittle-ductile transition occurs when the temperature increases. In the ductile region, extended steady states of deformation are attained without signals of macroscopic failure. In this regime, the shape and size of the grains remain unaffected by the deformation process, even after strains larger than 50%. The stress exponent n measured in this region is nearly 2, as found in fine-grained superplastic metals and ceramics. These results, together with the absence of dislocation activity in the grains and creep transitories, suggest that grain boundary sliding is the main deformation mechanism in steady-state conditions. As the temperature decreases, the local stresses generated during grain sliding are no longer fully accommodated, and microcavitation takes place along grain boundaries; however, the combination of elevated plasticity and great ability to withstand such flaws resulted in a creep damage-tolerant regime. Eventually, the microcavities coalesce into microcracks leading to the final failure of the composite. 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X-ray diffractograms of Al2O3, Y2O3 and monoclinic ZrO2 powder mixtures after milling and calcination at temperatures between 1200 and 1600 °C for 10 h in air. Bragg reflections were indexed according to the PDF-2002 reference patterns (International Centre for Diffraction Data, ICDD) as shown in the top of the figure. Figure 2. Microstructure of AYZ sintered composites: (a) secondary electron SEM micrograph; (b) backscattered electron SEM micrograph of the same area showing the three phases: Al2O3 (dark), YAG (gray) and cubic ZrO2 (white); (c) and (d) TEM micrographs showing the absence of secondary phases along grain boundaries and the presence of occasional dislocation arrays in the larger alumina grains. For the sake of comparison, the microstructure of a melt-growth AYZ eutectic composite with an interphase spacing similar to the grain size in the sintered composite is also shown: (e) transverse and (f) longitudinal cross-sections relative to the solidification direction (same color contrasts than in (b)). Figure 3. Grain size distributions of the Al2O3, YAG and yttria fully-stabilized cubic ZrO2 (YSZ) phases in eutectic-composition AYZ sintered composites. Figure 4. True stress - true strain curves for AYZ sintered composites deformed at an initial strain rate of o = 2x10-5 s-1 as a function of temperature. The  –  curve for the directionally-solidified AYZ composite (DSE, dashed curve) of Figure 2 deformed at 1400 °C is also shown for comparison. 33 Figure 5. Microstructure of deformation of AYZ sintered composites: (a) secondary and (b) backscattered electron SEM micrographs of the same area of a specimen 50%-strained in the ductile region at 1400 °C; (c) and (d) same as before for a specimen 50%-strained in the semibrittle region at 1350 °C; (e) fracture surface of a specimen failed at 1300 °C; and (f) TEM image of a sample 50%-deformed in steady-state conditions at 1400 °C. Loading direction is indicated by arrows. Figure 6. Creep curves plotted as log  −  for AYZ sintered composites deformed in steady state. Several determinations of (a) the stress exponent n and (b) the activation energy for creep Q (Eq. 1) by stress and temperature changes, respectively, are shown. Figure 7. Grain size-compensated strain rate against stress at 1400 °C for the three-phase AYZ sintered composite (solid black line). The average behavior of alumina (d = 3.0 m [4], 1.1 m [33], 0.5 m [52]), YAG (1.7 m [4], 1.7 – 5.3 m [40]) and cubic zirconia (1.8 m [4], 3.0 m [53]) monoliths (dashed areas) is shown, along with data for alumina doped with 103 ppm Y3+ (2.6 m [4], 2.2 m [37]) and with 103 ppm Zr4+ (0.7 m [33], 1.7 m [37]), and for binary alumina/zirconia (AZ) (2.3 m [4]) and alumina/YAG (AY) (2.0 m [4], 4.0 m [48]) composites. The strain rate of the alumina phase (volume fraction of 39.1 vol%) predicted by the isostrain model to explain the overall creep rate of the AYZ composite is also shown (dotted black line). 34 Figure 1 35 Figure 2 10 m (c) (e) (f) (a) 10 m (b) 5 m 1 m (d) 1 m (d) 5 m 36 Figure 3 37 Figure 4 38 Figure 5 (a) (c) (d) (e) (f) (b) 10 m 5 m 5 m 5 m 5 m 2 m 5 m 39 Figure 6 40 Figure 7