Al2O3/Y3Al5O12 (YAG)/ZrO2 Composites by Single-step Powder Synthesis and Spark Plasma Sintering

Journal of the European Ceramic Society 44 (2024) 7180–7188 Available online 10 May 2024 0955-2219/© 2024 The Author(s). Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/bync-nd/4.0/). Al 2 O 3 /Y 3 Al 5 O 12 (YAG)/ZrO 2 composites by single-step powder synthesis and spark plasma sintering Maryam Vakhshouri a , * , Aliasghar Najafzadehkhoee b , Ali Talimian a , Cristina L´ opez Pernia c , Rosalía Poyato d , ´ Angela Gallardo-L´ opez c , d , Felipe Guti´ errez-Mora c , d , Anna Prnova b , Duˇ san Galusek a , b a FunGlass, Alexander Dubˇ cek University of Trenˇ cín, ˇ Studentsk´ a 2, Trenˇ cín 91150, Slovak Republic b Joint Glass Centre of the IIC SAS, TnUAD and FChFT STU, ˇ Studentsk´ a 2, Trenˇ cín 91150, Slovak Republic c Dpto. de Física de la Materia Condensada, Universidad de Sevilla, Apdo. 1065, Sevilla 41080, Spain d Instituto de Ciencia de Materiales de Sevilla (CSIC-US), Am´ erico Vespucio 49, Sevilla 41092, Spain ARTICLE INFO Keywords: Al 2 O 3 -YAG ZrO 2 Composites Spark plasma sintering ABSTRACT Alumina-yttrium aluminum garnet (YAG)-zirconia composites are often produced by the melt solidification method. In the present study, we investigated the fabrication of α -Al 2 O 3 /Y 3 Al 5 O 12 (YAG)/ZrO 2 composite by Spark Plasma Sintering (SPS) of powders synthesized by Pechini’s sol-gel method. The ternary composites with homogenous microstructure and high density were produced by SPS at 1300 ◦C for 15 min. The addition of ZrO 2 promoted the sintering of composites, resulting in a higher density and, in turn, higher hardness. A change in the indentation fracture behavior as the result of ZrO 2 addition was observed. 1. Introduction Ceramic composites based on alumina are of the most widely used oxide materials in engineering applications because of their high mechanical performance both at room and elevated temperatures [1,2]. Producing Al 2 O 3 -based composites with superior mechanical performance has thus been the holy grail of the past few decades [2–8]. Among Al 2 O 3 -based ceramics, Al 2 O 3 /Y 3 Al 5 O 12 (YAG) composites are very promising because of their high thermal and chemical stability, mechanical strength, and creep resistance [9–13]. The physical properties of Al 2 O 3 /Y 3 Al 5 O 12 (YAG) composites, e.g. elastic modulus, meet the requirements for inducing zirconia phase transformation upon fracture [14–19]. Adding ZrO 2 to Al 2 O 3 /YAG composites is thus expected to improve the mechanical performance of the composites even further. For example, Al 2 O 3 /YAG/ZrO 2 composites produced by unidirectional solidification exhibited a fracture toughness as high as 8 MPa m 0.5 [1,5,6,20,21] however, the melt-solidification processes are not economically viable for mass production. Therefore, to efficiently produce Al 2 O 3 /YAG/ZrO 2 composites, alternative production routes have been developed, including rapid solidification [21–24], laser floating zone (LFZ) [12], edge-defined film-fed growth (EFG) [25], micro-pulling down (µ-PD)[26], laser-engineered net shaping (LENS) [16], selective laser melting (SLM) [27], laser directed energy deposition (LDED) [28,29], induction heating zone melting (IHZM) [10] and plasma spray deposition (PSD) [30]. However, the mentioned fabrication methods require melting or high-temperature sintering, which produces stabilized cubic ZrO 2 , preventing mehanisms related to transformation toughening [5,16,31]. Decreasing the sintering temperature of Al 2 O 3 /YAG/ZrO 2 composites is of great importance in preventing the formation of stabilized ZrO 2 ; as such, several attempts have been made to produce Al 2 O 3 /YAG/ZrO 2 composites by sintering Al 2 O 3 , yttria, and zirconia precursor powders [11,31–34]. However, due to the partial densification and, in turn, the presence of microstructural inhomogeneities, any influence of ZrO 2 addition on the mechanical properties is ambiguous. Therefore, it is unclear whether adding zirconia to Al 2 O 3 /YAG composites improves their mechanical performance, and fracture toughness in particular. This gap in knowledge is the primary motivation of the present study. Pechini’s sol-gel synthesis method enables effective homogenous mixing of precursor materials, thereby potentially improving the uniformity of the synthesized powders [35]. In the present study, Al 2 O 3 /YAG/ZrO 2 composite powders were synthesized using Pechini’s * Correspondence to: FunGlass – Centre for Functional and Surface Functionalized Glass, Alexander Dubˇ cek University of Trenˇ cín, ˇ Studentsk´ a 2, 911 50 Trenˇ cín, Slovakia. E-mail address: [email protected] (M. Vakhshouri). Contents lists available at ScienceDirect Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc https://doi.org/10.1016/j.jeurceramsoc.2024.05.004 Received 7 February 2024; Received in revised form 1 May 2024; Accepted 3 May 2024 Journal of the European Ceramic Society 44 (2024) 7180–7188 7181 methods; this way, homogenous ceramic composites were obtained. Here, we focus on the fabrication of homogenous Al 2 O 3 /YAG/ZrO 2 composite powder and the influences of ZrO 2 on the phase evolution, sintering behavior, and mechanical properties of the composites. 2. Experimental procedures Al(NO 3 ) 3 .9 H 2 O (99.9%, Sigma Aldrich, Germany), Y 2 O 3 (99.9%, Treibacher Industry, Austria), and ZrOClO 2 .xH 2 O (99.9%, Alfa Aesar) were used as raw materials in this study. Al 2 O 3 -Y 3 Al 5 O 12 composite powders with the eutectic composition of 76.8Al 2 O 3 -23.2 mol% Y 2 O 3 were synthesized by Pechini’s method. The details of the synthesis procedure are described elsewhere [36]. Briefly, yttrium nitrate solution, prepared by dissolving Y 2 O 3 in diluted HNO 3 , was mixed with an aqueous solution of Al(NO 3 ) 3 ; then, an aqueous solution of citric acid (C 3 H 5 O.(COOH) 3 , CentralChem, Bratislava, SK) and ethylene glycol (C 2 H 4 (OH), CentralChem, Bratislava, SK), with a molar ratio of 1:1, was added to the prepared solution. Afterward, the mixture was transferred to a reflux reactor and stirred at 85◦C for 2 h. The temperature was then increased to 150◦C to promote the polymerization process and remove the solvent to obtain a dried powder. The dried powder was then crushed and calcined at 1200◦C for 5 h in an ambient atmosphere. Ready-to-Press (RTP) powder was prepared by the addition of a few drops of deionized (DI) water, pan granulation, and passing through a sieve with a 50 mesh screen. ZrO 2 -containing powders containing 5 and 10 mol% of ZrO 2 were synthesized by introducing appropriate amounts of aqueous solution of zirconium source (ZrOClO 2 .xH 2 O) to the Y 3+ /Al 3+ solution. ZrO 2 substituted both Al 2 O 3 and YAG, and the molar ratio of Al 2 O 3 :Y 2 O 3 was kept constant (76.8:23.2) for all compositions. Pellets of the prepared powders were produced using spark plasma sintering (515 S, Dr. Sinter, Inc., Kanagawa, Japan, CITIUS 1 ). For sample preparation, 1.0 g of the powder was poured into a graphite die with an inner diameter of 10.0 mm and lined with graphite paper. The samples were heated to 700◦C in 3 min; then, the temperature was increased at a constant heating rate of 100◦C min −1 to 1300◦C and then kept for 15 min. Sintering was performed under a vacuum (5–9 Pa). A constant uniaxial pressure of 75 MPa was applied above 700◦C. Graphite debris was mechanically removed from the surface of sintered samples, and the samples were subsequently heat treated in a furnace at 800◦C for 1 h in an ambient atmosphere to remove carbon residue. The density of samples was measured using Archimedes’ principle in deionized water. The theoretical density of samples was estimated following the rule of mixtures and considering the theoretical density of α -Al 2 O 3 =3.98 g cm −3 , YAG=4.57 g cm −3 , and ZrO 2 =6.00 g cm −3 . Vickers’ indentation was performed on the polished surface of sintered samples using a micro-hardness tester (Duramin-1, Struers, Germany); the measurements were carried out in the ambient air (RH≈45%) and at room temperature (20◦C). At least 5 indentations were produced at each load (2.0, 3.0, and 5.0 kgf) with a dwell time of 15 s. The Vickers hardness H V was calculated from the applied normal load, P, and the size of residual hardness imprint: HV=1.8544 P d2(1) where d represents the average length of projected diagonals of the residual imprint. The indentation fracture toughness, K IFT , of the samples, was calculated as [37]: KIFT =0.016  E HV √P c3 2 (2) where E is Young’s modulus and c is the average of half crack length. The Young’s modulus was estimated using the pulse-echo method. The development of median cracks was confirmed post-indentation by removing a layer of material (ca. 2 µm) by ceramographic polishing [38]. The thermal decomposition and chemical reactions of synthesized powders upon heating were investigated using a Simultaneous Thermal Analyzer (STA, 449 F1 Jupiter®, Netzsch, Germany) in Differential Scanning Calorimetry-Thermogravimetry (DSC-TG) configuration upon heating to 1350◦C at a constant heating rate of 10◦C min −1 . The measurements were carried out under the flow of synthetic air (N 2 -20%O 2 ). Prior to the measurements, the powders were dried at 100◦C overnight. Small amounts of synthesized powders were heated at a constant heating rate of 10◦C min −1 up to 1400◦C with one hour dwell time at 700◦C to remove residual carbon; then, the samples were quenched to room temperature by removing them from the furnace at different temperatures between 1000◦C and 1400◦C to investigate phase evolution of the synthesized powders. The phase composition was investigated by X-ray powder diffraction (XRD) using an X-ray diffractometer (Empyrean, Malvern Pan- Analytical, Netherlands). The XRD measurements were carried out using nickel-filtered Cu K α radiation (λ 1 =1.5406 Å and λ 2 =1.5444 Å) in a θ:2θ scan mode (Bragg-Brentano focusing method), and 2θ between 10 and 90◦at a step size of 0.026◦and a scan step time of 96.390 s. The samples were rotated along ϕ-axis at the speed of 120 rpm. Quantitative phase analyses were performed by fitting the measured data using the Rietveld refinement method in MAUD software [39]. Small pieces were cut from samples, mounted in a conductive resin, and polished carefully using diamond abrasive disks down to 0.5 µm grit size. To prevent charging, the polished cross-sections were coated with a 10 nm thick platinum layer using a sputter coater (Q150-ES, Quorum, Chiddingly, UK). Microstructural examinations were performed using a scanning electron microscope, SEM, (JSM-7600 F, JEOL, Japan) equipped with a Low-Angle Backscattered Electron detector (LABE, JEOL, Japan). 3. Results The DSC and TG thermograms of Al 2 O 3 -Y 3 Al 5 O 12 samples (Al 2 O 3 - YAG) and of the samples containing 10 mol% of ZrO 2 , A-YAG-10Z, are shown in Fig. 1. The Al 2 O 3 -YAG and A-YAG-10Z samples showed similar DSC-TG results; the TG curve contained three main weight-loss events: first, the weight of the sample gradually decreased up to ca. 300◦C; then, there was a rapid weight-loss between 300◦C and 600◦C. Afterward, the weight of the sample remained constant until 1000◦C, at which point it started decreasing slowly again. While the weight of Al 2 O 3 -YAG samples reached a plateau above 1250◦C, the weight loss of A-YAG-10Z samples proceeded until the end of the measurement at 1300◦C. While a major exothermic peak was observed between 300 and 600◦C, there was no significant change in the DSC signal below 300◦C. Such behavior indicates that the weight loss below 300◦C was a result of the evaporation of volatile residues from the synthesis process. The burnout of residual carbon at higher temperatures was responsible for the intense exothermic peak centered at ca. 480◦C, corresponding to a significant weight loss [40,41]. A small exothermic peak was observed at about 935◦C for both Al 2 O 3 -YAG and A-YAG-10Z samples, which is attributed to the crystallization of yttrium aluminum garnet (YAG) [40,42,43]. Although a broad asymmetric peak between 950 and 1150◦C appeared next to the YAG crystallization peak in Al 2 O 3 -YAG samples, the A-YAG-10Z sample showed only a small deviation from the baseline in the DSC curves. Understanding the formation mechanism and crystallization kinetics of phases of the synthesized powders, particularly at low temperatures below 600◦C, requires further studies. However, considering the main events in the DSC curves (Fig. 1) and the densification of heat-treated powders (Fig. 3), the thermal behavior of synthesized powders was further studied in the temperature interval between 700◦C and 1400◦C. The synthesized powders were subjected to 1 CITIUS: Centro de Investigaci´ on, Tecnología e Innovaci´ on de la Universidad de Sevilla M. Vakhshouri et al. Journal of the European Ceramic Society 44 (2024) 7180–7188 7182 a temperature slightly (ca 50◦C) higher than the maxima of the peaks in the DSC records (Fig. 1). Then, the samples were removed from the furnace and quenched at different temperatures to minimize possible phase transformation. Fig. 2 shows the X-ray diffraction patterns of Al 2 O 3 -YAG, and of the samples containing 10 mol% ZrO 2 (A-YAG-10Z) quenched from 700, 1000, and 1200◦C. While the samples quenched from 700◦C lacked longrange order, the samples heated up to 1000◦C contained crystalline phases. The Al 2 O 3 -YAG samples exhibited diffraction maxima corresponding to yttrium aluminate garnet and transition aluminum oxides Fig. 1. DSC/TG curves of as-synthesized powders upon heating at the constant rate of 10 ◦C.min −1 under synthetic air flow: (a) Al 2 O 3 -Y 3 Al 5 O 12 , (Al 2 O 3 -YAG), (b) Al 2 O 3 -Y 3 Al 5 O 12 +10 mol% ZrO 2 , (A-YAG-10Z). Fig. 2. XRD patterns of synthesized powders: (a) Al 2 O 3 -Y 3 Al 5 O 12 , (Al 2 O 3 -YAG), (b) Al 2 O 3 -Y 3 Al 5 O 12 +10 mol% ZrO 2 , (A-YAG-10Z) after subjecting to heat-treatment at 700, 1000 and 1200◦C. ( α -Al 2 O 3 (●), Y 3 Al 5 O 12 (◼), c/t-ZrO 2 (◆), m-ZrO 2 (*), transition-Al 2 O 3 phases (▾)) Inset: zoom in on the peak associated with α -Al 2 O 3 (Fig. 2-a) and c-ZrO 2 (Fig. 2-b)). Note: the insets in Fig. 2-a are depicted using a square root scale for the vertical axis to augment the diffraction peaks. M. Vakhshouri et al. Journal of the European Ceramic Society 44 (2024) 7180–7188 7183 such as η -Al 2 O 3 and θ-Al 2 O 3 ; for ZrO 2 -containing samples, the main diffraction lines were associated with the ZrO 2 phases (tetragonal or cubic). Although the incorporation of yttrium cations into the zirconia stabilizes the tetragonal or cubic structures [44], the diffraction peak at 28◦, associated with monoclinic zirconia, indicates that yttrium ions might be inhomogeneously distributed, which resulted in insufficient stabilization and crystallization of monoclinic phase. The diffraction lines corresponding to α -Al 2 O 3 (corundum) and YAG were well-developed in the samples quenched from 1200◦C; however, diffraction maxima associated with transition Al 2 O 3 phases were still present. The sample containing 10 mol% ZrO 2 exhibited low intensity diffraction lines of zirconium oxide in addition to Al 2 O 3 and YAG phases. The introduction of ZrO 2 to the composite composition affected the thermal decomposition behavior and the crystallization of Al 2 O 3 and YAG phases; heat treatment at 1200◦C resulted in the formation of phases desired in this study, i.e., corundum, yttrium aluminum garnet, and zirconia. Therefore, the synthesized powders were heat-treated at 1200◦C for 5 h to ensure complete phase evolution and residual carbon removal. Fig. 3 shows the normalized displacement of the punch during spark plasma sintering of the composite powders upon heating to 1300◦C at the constant heating rate of 100◦C min −1 ; the punch displacement data were normalized with respect to the mass of the powder in the die. The changes in SPS pressure as a function of temperature are presented at a secondary horizontal axis. The sample composed of only Al 2 O 3 -Y 3 Al 5 O 12 (Al 2 O 3 -YAG) was characterized by a higher densification threshold temperature and a smaller final normalized densification than those containing 5.0 and 10.0 mol% ZrO 2 . While Al 2 O 3 -YAG samples exhibited a densification threshold at about 1060◦C, the samples containing 5.0 and 10.0 mol% of ZrO 2 started to densify at 860◦C and 900◦C, respectively. The densification of Al 2 O 3 -YAG samples slowed down above 1200◦C and continued until the end of the SPS process. The ZrO 2 - containing samples showed a moderate punch displacement vs. temperature above 1160◦C; A-YAG-10Z samples were characterized by zero punch displacement (plateau) above 1250◦C. Fig. 4 shows the XRD patterns measured on the pellets prepared by SPS. The Al 2 O 3 -YAG is composed of α -Al 2 O 3 and YAG; cubic-ZrO 2 was detected in the samples with added ZrO 2 . The quantification of the phase composition was carried out by Rietveld refinement, assuming that the substitution of Zr 4+ and Y 3+ is reflected only in the change of unit-cell parameters. Table 1 summarizes the phase composition, grain size, relative density, and mechanical properties of produced pellets. Given the phase composition of produced samples (Fig. 4), it was assumed that the samples contained only α -Al 2 O 3 , YAG, and cubic-zirconia (depending on the composition): Poisson’s ratio of the composite samples was estimated following the inverse rule of mixtures [45]. Poisson’s ratios are 0.23 for Al 2 O 3 , 0.25 for YAG, and 0.31 for cubic zirconia [46,47]. The addition of ZrO 2 to the samples resulted in an increase in the final relative density: ca. 97% for ZrO 2 -containing samples against ca. 95% for Al 2 O 3 -YAG samples; the result is in line with the densification behavior during SPS (Fig. 3). While ZrO 2 -containing samples are characterized by a higher hardness than Al 2 O 3 -YAG samples (~ 15–16 GPa vs. ~13 GPa), the indentation fracture toughness of samples is almost the same, about 3.8 MPa m 0.5 . It is worth mentioning that the measured hardness values of Al 2 O 3 -based ceramics is affected both by the indentation load and the formation of cracks upon indentation [48,49]. Vickers hardness of alumina and alumina-based ceramics is very sensitive to the used load due to the Indentation Size Effect (ISE) phenomenon and the effects of grain boundaries. This is more significant in fine-grained (below 3 µm) ceramics [49]. Although using larger loads produces cracks surrounding indentation, the measured hardness is not load-dependent; this allows for comparison of the hardness of samples. The hardness of Al 2 O 3 -YAG samples lower than the hardness of the ZrO 2 -containing samples is attributed to the lower density and to the presence of residual porosity in the microstructure. Fig. 5 shows SEM micrographs taken from the cross-sections of samples produced by SPS. While large pores remained in the Al 2 O 3 -YAG samples, no residual porosity was observed in the samples with added ZrO 2 . The microstructure is composed of a continuous matrix of yttrium aluminum garnet (YAG) with randomly distributed α -Al 2 O 3 and cubic- ZrO 2 grains. While A-YAG-5Z sample contained relatively coarse grains, the microstructures of Al 2 O 3 -YAG and A-YAG-10Z were more finegrained (Table 1). The secondary particles, e.g. c-ZrO 2 , can act as Zenner’s sites, preventing grain growth by grain boundary pinning. However, a minimum concentration of the Zenner particles is required to suppress grain growth [50]. In the present case, cubic-ZrO 2 is concentrated in the form of large grains within the YAG phase in A-YAG-5Z samples, which does not provide enough Zenner particles to suppress grain growth. In contrast, the larger concentration of c-ZrO 2 in A-YAG-10Z prevents growth of Al 2 O 3 and Y 3 Al 5 O 12 , resulting in smaller Fig. 3. Punch displacement (normalized by weight) during SPS for: Al 2 O 3 - Y 3 Al 5 O 12 , (Al 2 O 3 -YAG), Al 2 O 3 -Y 3 Al 5 O 12 +5 mol% ZrO 2 , (A-YAG-5Z), and Al 2 O 3 -Y 3 Al 5 O 12 +10 mol% ZrO 2 , (A-YAG-10Z) samples. Fig. 4. XRD patterns of pellets produced from Al 2 O 3 -Y 3 Al 5 O 12 , (Al 2 O 3 -YAG), Al 2 O 3 -Y 3 Al 5 O 12 +5 mol% ZrO 2 , (A-YAG-5Z), and Al 2 O 3 -Y 3 Al 5 O 12 +10 mol% ZrO 2 , (A-YAG-10Z) SPS (1300◦C, 15 min, 75 MPa). ( α -Al 2 O 3 (●), Y 3 Al 5 O 12 (◼), c-ZrO 2 (◆)). M. Vakhshouri et al. Journal of the European Ceramic Society 44 (2024) 7180–7188 7184 grains. However, providing a definite explanation for this result is challenging and requires detailed studies on the critical concentration of ZrO 2 in Al 2 O 3 /Y 3 Al 5 O 12 /ZrO 2 composites. Fig. 6 shows exemplary Vickers’ indentations produced on the polished surface of SPSed samples using an indentation load of 29.4 N; radial cracks were visible surrounding the hardness imprints in all samples. Moreover, the ZrO 2 -containing samples exhibited light birefringence patterns around the indentation, which indicated that lateral cracks were also produced and extended beneath the surface. It should be noted that in oxide ceramics, the lateral and radial cracks might initiate and grow during various stages of indentation, such as loading and unloading [51].The formation of secondary cracks upon sharp indentation is influenced by different parameters and is, thus, complex. Therefore, further detailed examinations are required to investigate the cracking behavior of samples. The indentation results indicated that the fracture mechanisms of ZrO 2 -containing samples are different from those of the binary Al 2 O 3 -YAG composite. The addition of ZrO 2 to the Al 2 O 3 -YAG composite has changed the indentation cracking behavior of the ceramic composite. 4. Discussion Although solid-state synthesis is a common method to produce Al 2 O 3 -YAG composites, it requires high-temperature treatment to eliminate the chemical inhomogeneities or intermediate phases such as YAM and YAP [52,53]. Pechini’s method enabled the synthesis of a homogenous mixture of precursors, i.e., Al 2 O 3 and Y 2 O 3 , that directly yields yttrium aluminate garnet (YAG) and corundum ( α -Al 2 O 3 ) during the thermal treatment (Figs. 1 and 2). While YAG crystallized directly from the synthesized powder, corundum was produced through a series of transformations of Al 2 O 3 phases [54]. Although adding ZrO 2 affected the crystallization behavior of α -Al 2 O 3 and YAG (Fig. 2), composite powders containing the main desired phases, α -Al 2 O 3 and YAG, were obtained for all prepared compositions. The synthesized powders were successfully densified using SPS. The process that combines Pechini’s method and SPS thus shows promise as a fabrication method of Al 2 O 3 -based ceramics. Al 2 O 3 -YAG samples exhibited diffraction peaks corresponding to YAG and transition Al 2 O 3 phases such as θ-Al 2 O 3 ; in contrast, the diffraction maxima associated with various (monoclinic or tetragonal) ZrO 2 phases are the major features of diffraction patterns of ZrO 2 -con- taining samples. It should be noted that various ZrO 2 phases cannot be clearly distinguished due to the peak broadening related to their poor crystallinity and low concentration of some phases. Although Zr 4+ ions might impact the reaction steps of the powder synthesis and intermediate species at low temperatures [55–57], heat treatments above 1000◦C are required to remove residual carbon and produce crystalline phases [36,58,59]. The phase evolution of the sample was determined using XRD (Figure S1). Some transition Al 2 O 3 phases, i.e. γ-Al 2 O 3 , 300–500◦C and δ-Al 2 O 3 , 700–800◦C, [54] have been reported to crystallize at temperatures lower than the initial calcination of powder, 700◦C. However, no distinguishable diffraction peaks can be observed in the diffraction. The small concentration of transition Al 2 O 3 phases compared to the amorphous matrix and their small crystallite size are considered responsible for the absence of the diffraction peaks. Nevertheless, the diffraction peaks associated with γ-Al 2 O 3 and δ-Al 2 O 3 indicate that the mentioned transition Al 2 O 3 phases have crystallized upon heating. Moreover, the diffraction peaks corresponding to θ-Al 2 O 3 in the XRD pattern collected from the powder quenched from 1000◦C show that the produced phases, γ-Al 2 O 3 and δ-Al 2 O 3 , were replaced by θ-Al 2 O 3 . Eventually, all Al 2 O 3 was transformed into α -Al 2 O 3 . Such phase development resembles the thermal decomposition of boehmite (AlOOH) and its transformation into α -Al 2 O 3 [54,60]. A more detailed consideration of the crystallization of Al 2 O 3 phases upon heating, however, is outside the scope of this study. While the crystallization of Al 2 O 3 phases was only slightly affected when ZrO 2 was added to the composite powder, the crystallization temperature of YAG was affected by the addition of ZrO 2 significantly and shifted to a higher temperature (from 1050◦C to 1100◦C-Fig S1-a vs. c). After heat treatment at 1000◦C, the powder with 5 wt% ZrO 2 contained both tetragonal zirconia and YAG; the diffraction peaks associated with tetragonal zirconia were slightly stronger in the XRD pattern of the sample with 10 wt% of ZrO 2 . In samples subjected to the heat treatment above 1100◦C, a shift of the tetragonal zirconia diffraction lines was observed, and eventually, at >1300◦C, the samples showed only the diffraction lines corresponding to cubic yttria-doped zirconia. The gel network structure determines the crystallization behavior of sol-gel powders, the formation of individual phases, and, in turn, the properties of the final product. The cross-linking of citric acid and ethylene glycol produces a gel network with tangled chains that entrap the metal ions, Al 3+ , Y 3+ , and Zr 4+ , as well as the solvent [35,57,61]. The produced gel preserves the random distribution of metal ions and exhibits crystallization behavior similar to Al 2 O 3 -YAG composites prepared by the glass-ceramic route [36,59,62,63]. The metal ions are adjacent to the carboxyl groups of the gel network [19,36]. Therefore, it is intuitively expected that the metal cations maintain the structure of the first neighboring anions, such as Al-OOH, even after burnout and dissociation of carboxyl groups [57,63,64]. Then, oxide phases are formed as the temperature increases; this agrees well with the sequence of the formation of transition Al 2 O 3 upon heating (γ→δ→θ→ α ; Figure S1). It is also interesting to point out the impact of ZrO 2 on the crystallization of phases using the field strength (FS) concept (ratio of ionic charge to its radius) for the metal cations in the system. While the field strength for Zr 4+ ions is almost identical to the field strength of Al 3+ , it is significantly larger than that of Y 3+ . Thus, it is expected that Zr 4+ ions affect the yttrium-oxygen bonds and change the crystallization behavior of yttrium aluminate garnet. Given the field strength sequence, ZrO 2 crystallizes before YAG during the heat treatment. Despite the sequential formation of the phases, α -Al 2 O 3 and cubic-ZrO 2 are randomly distributed in a continuous network of YAG after spark plasma sintering. Fig. 5c shows a representative microstructure of the sample produced by SPS from the powder containing 10 mol% ZrO 2 ; the microstructure consists of a continuous network of YAG grains in which ZrO 2 grains and islands of Al 2 O 3 grains are embedded. The addition of zirconia also affects the densification of Al 2 O 3 -YAG composites and increases the final density of samples after SPS; Table 1 Phase composition (vol%), relative density, and mechanical properties of the SPSed samples. Sample composition Composition (vol%) Grain size (µm) Relative Density (%) Elastic Modulus (GPa) Poisson’s Ratio Hardness (GPa) Indentation Fracture Toughness (MPa m 0.5 ) YAG α -Al 2 O 3 c- ZrO 2 YAG α -Al 2 O 3 c- ZrO 2 Al 2 O 3 -Y 3 Al 5 O 12 (YAG) 55.6 44.4 ND 1.1 ±0.2 1.1 ± 0.8 – 94.7 ±0.2 342 ±5 0.24 12.7 ±0.6 3.7 ±0.2 A-YAG-5Z 55.6 41.8 3.5 1.7 ±0.4 2.1 ±0.8 1.2 ±0.6 97.3 ±0.5 350 ±5 0.24 15.1 ±0.9 4.1 ±0.1 A-YAG-10Z 51.0 42.0 8.0 1.2 ±0.2 1.3 ±0.5 0.7 ±0.2 97.8 ±0.1 352 ±6 0.25 16.1 ±0.8 3.8 ±0.2 M. Vakhshouri et al. Journal of the European Ceramic Society 44 (2024) 7180–7188 7185 moreover, the samples containing ZrO 2 are characterized by a lower densification threshold than the samples without ZrO 2 . YAG is the continuous phase in the produced composites (Fig. 5) and is expected to be the main contributor to the shrinkage. However, despite the known benefits of ZrO 2 addition to Al 2 O 3 bodies, ZrO 2 is also known to slow down the densification of Al 2 O 3 [18,65]. Therefore, it is plausible to assume that only the interactions of ZrO 2 and YAG account for the improved densification of the composites. Given the ionic radii of cations (Zr 4+ : 86 pm, Y 3+ : 104 pm vs Al 3+ 67.5 pm), Y 3+ and Zr 4+ can partially substitute each other in YAG and ZrO 2 structures [66,67]. When zirconium ions take the place of yttrium ions in the YAG structure, one mole of Zr 4+ takes the place of two moles of Y 3+ and produces one mole of oxygen vacancies to compensate for the charge caused by unfilled sites of the structure. Since the incorporation of Zr 4+ into Al 3+ sites is limited [68,69], if any, the interactions of Zr 4+ and aluminum sites can be neglected. Thus, Y 2 O 3 sites can be considered available for Zr 4+ doping and the stoichiometric reaction described in Kroger-Vink Notation: ZrO2→ Y2O3Zr• Y+V/// Y+2O× O+V•• O(3) The production of vacancies in YAG facilitates the diffusion of Y 3+ and oxygen and, in turn, promotes the densification of samples. Tetragonal-ZrO 2 was transformed to cubic-ZrO 2 when the samples are subjected to a temperature above 1200◦C (Figure S1), which is corroborated by the reported temperature for phase transformation of yttrium-dopped ZrO 2 [46]. Moreover, in this work, cubic-ZrO 2 was present in the SPS samples (Fig. 4). The small particle size or limited incorporation of Y 3+ into ZrO 2 might be responsible for the formation of tetragonal-ZrO 2 at lower temperatures. When the composite powders were subjected to higher temperatures, Y 3+ diffused from YAG particles to ZrO 2 particles and stabilized cubic ZrO 2 ; this resulted in an increase of the unit-cell parameter of cubic zirconia from 5.09 Å to ca. 5.16 Å, as determined by Rietveld refinement of powder diffraction data, Fig. 4. The depletion of YAG from yttrium due to the diffusion of Y 3+ into zirconia particles can change the defect chemistry of YAG and, in turn, promote the densification of YAG. The faster diffusion of Y 3+ may have also contributed to the grain growth observed for A-YAG-5Z. In contrast, the grain size of A-YAG-5Z was larger than the grain size of the composite containing 10 mol% of ZrO 2 ; the Zenner pinning of the YAG matrix grain boundaries by smaller ZrO 2 and Al 2 O 3 grains is assumed to be responsible for the observed microstructure refinement. The A-YAG-10Z contains a sufficient number of ZrO 2 particles, which enable suppression of grain growth via the Zenner pinning mechanism [70–72]. The most significant impact of adding ZrO 2 to Al 2 O 3 -YAG composites is the increase of Vickers hardness; it increased from ~ 13 GPa for samples without ZrO 2 to 16 GPa for ZrO 2 -containing composites. The higher density and smaller grain size of A-YAG-10Z are responsible for the improved mechanical properties. Although the ZrO 2 -containing samples exhibited a slightly higher fracture toughness than those without ZrO 2 , tetragonal to monoclinic phase transformation of zirconia is unlikely to contribute to the fracture resistance increase since a limited amount, if any, of tetragonal phase is present in the samples after SPS (Fig. 4). Nevertheless, zirconia has influenced the indentation fracture resistance of samples: either by improving the densification or changing the fracture mechanism. ZrO 2 particles are dispersed within the YAG matrix of the composite (Fig. 5). Considering the differences between the coefficient of thermal expansions of cubic ZrO 2 and YAG (9.7 ×10 −6 K −1 vs. 6.9 ×10 −6 K −1 ), a radial tension is expected to be generated surrounding the secondary particles during the cooling of samples, resulting in crack deflection and improving the fracture toughness [73]. In addition to the immediate role of c-ZrO 2 particles in crack deflection, they might improve the fracture toughness of the composites through the crack shielding process. Detailed studies on the roles of c-ZrO 2 particles in increasing indentation fracture resistance are, however, required to reveal the mechanisms at play in the increase of fracture toughness. Figure S2 shows the indentation imprints produced in samples at two different loads: 19.6 and 49.1 N. Lateral cracks can be seen in zirconiacontaining samples where the radius of lateral cracks is comparable to or larger than the radius of radial/median cracks. Although the formation of lateral and radial cracks is complex and could occur during various stages of indentation, their formation is governed by the mechanical Fig. 5. Microstructure of samples by SPS at 1300 ◦C for 15 min from (a) Al 2 O 3 - Y 3 Al 5 O 12 , (Al 2 O 3 -YAG), (b) Al 2 O 3 -Y 3 Al 5 O 12 +5 mol% ZrO 2 , (A-YAG-5Z), and (c) Al 2 O 3 -Y 3 Al 5 O 12 +10 mol% ZrO 2 , (A-YAG-10Z). The images were taken in backscattered (Z-contrast) mode; from the darkest to the brightest are α -Al 2 O 3 , Y 3 Al 5 O 12 , and c-ZrO 2 .The white arrows point at insolated pores. M. Vakhshouri et al. Journal of the European Ceramic Society 44 (2024) 7180–7188 7186 characteristics of the material, for example, Young’s modulus and Poission’s ratio. The radius ratio of radial/median cracks to the lateral cracks (2 C/2 R) was therefore used to quantify possible changes in the fracture mechanism as a function of zirconia concentration. Fig. 7 shows the 2 C/2 R ratio as a function of the indentation load of ZrO 2 -con- taining samples. While a trend emerges from the ratio of crack lengths in samples with 5 mol% of ZrO 2 , the 2 C/2 R ratio for samples containing 10 wt% ZrO 2 decreased when the indentation load was increased to 49.1 N indicating the increased contribution of lateral cracking (2 C) in the indentation fracture. The flaws produced from lateral cracks play a significant role in surface wear and are responsible for visible scratches and surface wear [51]. The produced Al 2 O 3 -YAG-ZrO 2 composites lack transformation toughening mechanisms; however, their response to sharp contacts, i.e., indentations, can be tailored by changing the concentration of ZrO 2 . Detailed studies are required to understand the impacts of zirconia on the wear behavior of this composite family and the possible energy dissipation mechanisms during scratch or upon sharp contact. 5. Conclusions Alumina-yttrium aluminum garnet (YAG)-zirconia ceramics were produced by synthesizing the composite powder via Pechini’s sol-gel method, followed by spark plasma sintering at 1300◦C. Pechini’s method enables the production of composite bodies with homogenous microstructure containing YAG, α -Al 2 O 3 , and cubic ZrO 2 ,. The addition of ZrO 2 to the composite increases the sintered density, which, in turn, results in an increase in the Vickers hardness of produced bodies. The addition of ZrO 2 and its influence on the densification of composite bodies result in a change in the indentation fracture mechanism of Al 2 O 3 -YAG composites, which is reflected in the distinct formation of lateral cracks during Vickers indentation. Contributions The manuscript was written through the contribution of all authors. The idea was conceived by D.G. M.V produced the samples, characterized them, and prepared the first draft. CLP, RP, AGL, and FGM performed Spark Plasma Sintering and mechanical testing. Thermal analyses were carried out by A.P. A.N. and A.T. analyzed results and edited the text. All authors contributed to the text and gave their approval to the final version of the manuscript. CRediT authorship contribution statement Anna Prnova: Formal analysis. Duˇ san Galusek: Writing – review & editing, Conceptualization, Funding acquisition. ´ Angela Gallardo- L´ opez: Data curation. Felipe Guti´ errez-Mora: Validation. Cristina L´ opez Pernia: Formal analysis. Rosalía Poyato: Methodology. Aliasghar Najafzadehkhoee: Formal analysis. Ali Talimian: Writing – review & editing. Maryam Vakhshouri: Writing – original draft, Investigation. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments This project has received funding from the European Union’s Horizon 2020 research and innovation programme under grant agreement No 739566. Financial support of this work by the grants APVV 19–0010, and VEGA 2/0028/21 is gratefully acknowledged. We appreciate the support provided by JECS TRUST under contract 2021286. The authors would like to thank the aid of CITIUS at the University of Seville for using the SPS facility. We appreciate the helpful discussions and useful comments of Prof. Dr. Raúl BERMEJO on the mechanical properties and indentation behavior of composite samples. Fig. 6. Indentation cracks produced using 29.4 N on SPSed pellets from (a) Al 2 O 3 -Y 3 Al 5 O 12 , (Al 2 O 3 -YAG), (b) Al 2 O 3 -Y 3 Al 5 O 12 +5 mol% ZrO 2 , (A-YAG-5Z), and (c) Al 2 O 3 -Y 3 Al 5 O 12 +10 mol% ZrO 2 , (A-YAG-10Z). Birefringence patterns due to lateral cracks are marked by dashed arcs. Fig. 7. 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