Development of calcium stabilized nitrogen rich alpha-sialon ceramics along the Si3N4:1/2Ca3N2:3AlN line using spark plasma sintering

Calcium stabilized nitrogen rich sialon ceramics having a general formula of Ca x Si 12-2x Al 2x N 16 with x value in the range of 0.2-2.2 for compositions lying along the Si 3 N 4 :1/2Ca 3 N 2 :3AlN line were synthesized using nano/submicron size starting powder precursors and spark plasma sintering (SPS) technique. The development of calcium stabilized nitrogen rich sialon ceramics at a signicantly low sintering temperature of 1500°C (typically reported a temperature of 1700°C or greater) remains to be the highlight of the present study. The SPS processed sialons were characterized for their microstructure, phase and compositional analysis, physical and mechanical properties. Furthermore, a correlation was developed between the lattice parameters and the content (x) of the alkaline metal cation in the alpha-sialon phase. Well densied single-phase nitrogen rich alpha-sialon ceramics were achieved in the range of 0.53(3) ≤ x ≤ 1.27(3). A nitrogen rich alpha-sialon sample possessing a maximum hardness of 22.4 GPa and fracture toughness of 6.1 MPa.m 1/2 was developed.


Introduction
Sialon materials is a class of ceramics that have been thoroughly studied with respect to their phase stability regimes, thermo-mechanical properties, photoluminescence and oxidation behavior [1][2][3][4][5][6][7][8][9][10][11]. Most of the work on sialon ceramics has been concentrated around the two commonly known phases, namely alpha-sialon and beta-sialon. A compound of silicon nitride, beta-sialon having a general formula of Si 6z Al z O z N 8-z , is formed as a result of the chemical replacement of silicon-nitrogen bonds with aluminumoxygen bonds [12].
Compositions belonging to alpha-sialon phase are generally de ned by M x v Si 12-(m+n) Al m+n O n N 16-n (x and v represent solubility and valence of the cation M in alpha-sialon structure, respectively) where x < 2, x = mv, and m (Al-N) and n (Al-O) replace (m+n) (Si-N) bonds [13]. The substitution of silicon-nitrogen bonds with aluminum-nitrogen and aluminum-oxygen bonds results in a charge imbalance which is neutralized by the incorporation of M v+ ions such as Li + , Ca 2+ , Mg 2+ , Y 3+ and lanthanide ions [14][15][16][17]. Studies have been reported on synthesis and phase stability regime of single phase alpha sialons [14][15][16][17][18][19]. Several compositions of yttrium-stabilized alpha-sialons along the nitrogen rich line of Si 3 N 4 -YN:3AlN were synthesized by Sun et al. where the solubility limit of yttrium cation in alpha sialon was found out to be 0.43<x<0.8 [1]. Zhen-Kun et al. studied the formation of oxygen rich yttrium stabilized alpha-sialons along the Si 3 N 4 -Y 2 O 3 :9AlN and the solubility limit of yttrium cations was reported as 0.33<x<0.67 [14]. It is well known that synthesis of alpha-sialons occurs via a solution-reprecipitation mechanism, which involves the formation of an intermediary oxy-nitride liquid phase as a result of the reaction between the starting oxide and nitride precursors [20,21].
In contrast to alpha-sialons synthesized using yttrium or rare earth stabilizing cations, calcium stabilized alpha-sialons have gained considerable attention due to the higher solubility of calcium cation as compared to the former (maximum solubility value x max of 1.6 for Ca +2 as compared to x max value of 1.0 for Yb +2 ) [15,17]. Furthermore, Hassan Mandal, in his study on post sintering heat treatment of alpha sialons, reported that in contrast to rare-earth stabilized alpha-sialons, calcium stabilized alpha-sialons depicted complete resistance towards alpha to beta phase transformation in the temperature range of 1450-1500°C [22].
The higher stability of calcium stabilized alpha-sialons has been attributed to the lower valency of calcium cation, since the high temperature stability of alpha-sialon increases with increase in the solubility of charge stabilizing cation (x), which in turn increases with decrease in cation valency [23]. Figure 1 shows the Janecke prism and the calcium alpha-sialon plane highlighting the single-phase alphasialon region (oxygen rich as well as nitrogen rich) along with the neighboring phases at 1800°C [24]. Several studies conducted on the synthesis of oxygen rich alpha-sialons have reported the formation of elongated morphology of alpha-sialon as compared to the generally known equiaxed morphology, thus resulting in considerable improvement in the toughness of these materials [25][26][27]. However, there are very few studies that have been reported on the synthesis of nitrogen-rich alpha-sialons, mainly due to the fact that they are di cult to densify [23,28]. Studies on the synthesis of calcium stabilized nitrogen rich alphasialons using all nitride precursors is even more scarce in the literature due to the di culty in handling and storage of calcium nitride [23].
Xie et al. worked on the synthesis of calcium stabilized nitrogen rich alpha-sialons along the Si 3 N 4 -1/2Ca 3 N 2 :3AlN line at 1700°C using hot pressing technique [23]. The limits of single phase calcium alphasialon was reported to be 0.5<x<1.7. Later, Yanbing Cai et al. studied the synthesis of calcium stabilized nitrogen rich alpha-sialons using CaH 2 as a starting precursor instead of Ca 3 N 2 at 1800°C using hot pressing (HP) technique [28]. Single phase alpha sialon ceramics were obtained in the range of 0.50<x<1.38. Densi cation using conventional sintering processes, such as hot press (HP), is governed by the heat supplied via heating elements (comparatively much slower heating rate than SPS) and the externally applied pressure [29].
On the other hand, for the non-conventional and modern sintering process spark plasma sintering (SPS), in addition to the uniaxial applied pressure, the pulsed nature of the current creates spark discharge and gives rise to joule heating effect between the powder particles. The high rate of plasma discharge transmits and distributes the Joule heat throughout the material. This phenomenon results in a fast and uniform heat distribution within the sample, which ultimately leads to highly homogeneous samples with consistent densities. In addition, the high heating/cooling rate is bene cial as it restricts the grain growth as well as the formation of intermediate phase(s). Samples can be densi ed in very short time by SPS [29,30]. Raja et al. has reported the synthesis of oxygen rich calcium alpha-sialon ceramics at relatively low temperature of 1500°C using nano-sized precursors and SPS technique [31]. However, to our knowledge, synthesis of nitrogen rich sialons at low temperatures (below 1700°C) has not been reported in the literature.
The aim of the present study was to develop calcium stabilized nitrogen rich alpha-sialons along the Si 3 N 4 :1/2Ca 3 N 2 :3AlN line at a relatively low temperature of 1500°C using SPS and nano-/submicron-size starting powder precursors. The developed sialons were characterized for their microstructure, phase and compositional analysis, physical and mechanical properties. Furthermore, a correlation was developed between the lattice parameters and the content of alkaline metal cation (calcium) in the alpha-sialon.

Experimental Procedure
Calcium stabilized nitrogen rich alpha-sialon compositions having the general formula of Ca x Si 12-m Al m N 16 (where x = m/2), along the nitrogen rich line, were selected in this study. The nominal x value varied in the range of 0.2-2.2. Actually achievable compositions of alpha-sialon samples (by taking into account the oxide content on the powder surface) as listed in Table 1 were synthesized using starting powder precursors of alpha-Si 3 N 4 with an average particle size of ~300 nm (Ube Industries SN-10, Japan), AlN with the particle size of ~50nm (ChemPUR, Germany) and Ca 3 N 2 of ~74µm (~200 mesh, Sigma Aldrich, Germany very sensitive to moisture and air and quickly decomposes into calcium hydroxide and ammonia, powders of silicon nitride and aluminum nitride (required to synthesize 5 gm of a speci c composition) were initially mixed via ultrasonic-probe sonicator in ethanol medium. The sonicated powder mixture was then dried in oven at 95°C for a period of 24h to evaporate ethanol. The dried powder mixture of silicon nitride and aluminum nitride was then mixed with the required amount of calcium nitride using mortar and pestle in NEXUS II glove box (Vacuum Atmosphere Company, USA) under the presence of high purity argon gas. The powder mixture thoroughly mixed for about 15-20 minutes was then poured into a 20mm graphite die lined with BN, and the die was sealed with para lm. The sealed die was taken out of the glove box and placed within the SPS chamber. The powder mixtures were synthesized (compacted and sintered) using an SPS apparatus (FCT system, model HP D5, Germany), calibrated and maintained by the manufacturer, on a yearly basis. The schematic of the SPS machine is shown in gure 2. During the synthesis process, the temperature was monitored and controlled using a pyrometer, xed just above the upper punch, measuring the temperature only 3 mm away from the top surface of powder mixture inside the graphite die. Moreover, as shown in the schematic, the temperature of the die was also measured using a K-type thermocouple, placed in the center of the die, only 3 mm away from the powder sample in the die. The maximum temperature difference of less than ±20°C, measured between pyrometer and thermocouple, assured the indepth calibration.
The powder mixture was consolidated into 20mm diameter pellets at a sintering temperature of 1500°C with 30 minutes holding time in a vacuum environment and 50MPa of constant uniaxial pressure. The heating rate of 100°C/min was used in the synthesis, followed by constant temperature holding process, then samples were cooled down to room temperature within 5 minutes. Each sample was prepared twice and densi cation data was obtained from the SPS equipment. The shrinkage curves were generated using the experimental data (relative piston displacement, time and temperature values) collected and stored automatically on the FCT SPS software. The synthesized samples were cleaned of graphite using SiC grinding paper and were further ground and polished using Automet 300 Buehler grinding machine to a surface nish of 1μm. Phase analysis was performed using Bruker D8 X-ray powder diffractometer with Cu K α1 radiation (γ = 0.15416 nm). The tube current of 40mA, accelerating voltage of 40 kV, step size of 0.02 degrees and step time of 3sec were maintained during the entire scan. In order to calculate the lattice parameters of alpha-sialon accurately, silicon powder was used as an internal standard. The amount of alpha and beta phases was calculated using the following equation [32]: (see Equation 1 in the Supplemental Files) Where I α represent observed intensities of (102) and (210) peaks belonging to alpha-phase while I β represents observed intensities of (101) and (210) peaks belonging to beta-phase. W β represents fraction of beta-phase while K is the proportionality constant (0.518 for β(101) and α(102) peaks while 0.544 for β(210) and α(210) peaks [32]. The amount of minor phases was determined based on the relative intensity method. Highly polished and fractured surfaces of the synthesized samples were observed using a eld emission scanning electron microscope (FESEM, Lyra3, Tescan, Czech Republic) with an accelerating voltage of 20 kV. Phoenix focused ion beam (Helios G4 UX DualBeam, FEI) was used to prepare the lamellas for observation under a transmission electron microscope (TEM). High resolution imaging and selected area diffraction patterns were acquired using FEI's Titan transmission electron microscope equipped with EDX detector. On average, ve area EDX analyses were performed to calculate the cation composition. The software UnitCellWin and XRD peak positions were used to calculate the lattice parameters. The density of the synthesized disk shaped samples (after removal of graphite) was calculated using Archimedes' principle. Vickers hardness (HV 10 ) was measured on the surface of the polished samples using a universal hardness tester (Zwick-Roell, ZHU250, Germany). The indentation load of 10kg was adopted for the hardness test. Indentation method was used to calculate the fracture toughness (K 1c ) of the synthesized samples [33][34][35][36]. Evan's equation shown below was used to calculate the indentation fracture toughness. In the equation 'MCL' stands for the maximum crack length initiated from the indentation and 'd' is the average diagonal: (see Equation 2 in the Supplemental Files)  precursors at 1700°C and 1h holding time [23]. However, the reported shrinkage rate was much slower due to the slower reaction kinetics as a result of the conventional heating technique (slower heating rate) as well as due to larger size (>1μm) of the starting powder precursors. insu cient amount of liquid phase was provided to be the main reason [23]. This suggests that along with nano-size starting precursors and novel pulsed-based SPS technique, low melting temperature (1195°C) of calcium nitride plays a pivotal role in achieving well densi ed samples at 1500°C. This is in line with the literature where it has been reported that a small particle size of starting powders helps to synthesize dense sialon materials at low temperatures [39,40].   using HP technique [23]. Furthermore, Yabaing et al. in their work on the synthesis of nitrogen rich sialons at 1800°C, have also communicated the formation of AlN and CaSiAlN 3 phases along with alpha-phase for higher values of x (x>1.6) and the single phase alpha-sialon phase is reported to be in the range of 0.6 ≤ x ≤ 1.6 [24].

Phase Analysis
X-ray diffraction results of the samples synthesized at 1500°C are summarized in Table 3. Lattice parameters of alpha sialon are seen to increase with the increase in x value attributed to the fact that more amount of Si-N bonds are replaced by Al-N bonds as well as more amount of calcium ion is being incorporated in the structure. In calcium stabilized alpha-sialons, the charge imbalance caused as a result of the substitution of Si +4 for Al +3 is brought into balance by the addition of Ca +2 ions at the interstitial positions in the network. It is established that alpha-sialon unit cell contains two interstices per unit cell,   (4) 19.4 (7) 5.0 (4) α (87), β  Figure 6 shows the variation of alpha sialon lattice parameters with respect to change in x value. Similar behavior is reported by Yabning for samples synthesized at 1800 C ( gure 6) [24]. The dependence of alpha-sialon lattice parameters with the amount of Ca+2 can be well de ned by the following empirical relationships: Formation of alpha-sialon phase for a sample having x value as low as 0.2 along with the observation suggesting a linear dependence of alpha-sialon unit cell parameters with x value, provide a good reason to believe that calcium stabilized nitrogen rich alpha-sialon forms continuously in the range of 0<x<1.83.
Therefore, the previously reported low solubility limit of Ca +2 in alpha-sialons shall not be attributed to a structural limitation but should rather be ascribed to the fact that formation of single phase alpha-sialons near the nitrogen rich corner has not been attained due to the presence of the oxide layer contamination leading to the formation of beta-sialon phase. Figure 7 (a-f) depicts the secondary electron FESEM micrographs of the fractured surface of samples from the nitrogen rich sialon ceramics sintered at 1500°C. Primarily, a classical intergranular fracture was observed in all samples, along with some amount of grain pullout. Sample Ca-0.2 containing alpha-and beta-sialon phases exhibit elongated morphology of beta-sialon along with smaller equiaxed grains of alpha-sialon ( gure 7a). Samples with x values greater than 0.4 (Ca-0.6 to Ca-2.2) composed of single phase or almost single-phase nitrogen rich calcium alpha-sialon depicted ne equiaxed grains (a morphology typical of alpha-sialon phase). However, with the increase in x value (more Ca 3 N 2 ), alphasialon grains with elongated morphology are also observed. Additionally, it is observed that an increase in

Microstructure Analysis
x value yields an increase in the average grain size of alpha-sialon. The grain size distribution for the three samples (Ca-0.6, Ca-1.2 and Ca-2.0) is depicted in gure 8 where the average grain size is measured to increase from 286 +98nm to 385 +96nm to 468+169 nm respectively. having an elongated morphology as a result of preferential grain growth during the densi cation process has also been reported in the synthesis of hot-pressed oxygen rich Ca-alpha-sialons [11,38,41]. It is known that the grain growth of alpha-sialon takes place via solution, re-precipitation mechanism and that the presence of a liquid phase is necessary to promote grain growth. Therefore, the presence of elongated morphology of alpha-sialon grains as well as the grain growth for samples having higher x value (i.e. more have reported that intergranular phase (if present) is always amorphous and can exist as reasonably large pockets having a size in the range of 0.1 to 1μm or interfacial glassy lm with a minimum reported thickness of about 60nm [42]. In our study, the high-resolution TEM images of nitrogen rich sialons ( gure 9 b, d, and e) reveal that there is no sign of a glassy phase present at the grain boundaries nor at the junction of multiple grains. These results suggest that the synthesized nitrogen rich sialons would be more stable at high temperature (1400-1600°C) compared to the oxygen rich sialons containing an amorphous phase on the grain boundaries.

Mechanical Properties
Vickers hardness (HV 10 ) and fracture toughness (K 1c ) of the calcium stabilized nitrogen rich alpha-sialons synthesized at 1500°C are summarized in

Conclusion
Calcium stabilized nitrogen rich sialon having compositions along the Si 3 N 4 :1/2Ca 3 N 2 :3AlN line were synthesized. Enhanced reaction kinetics due to nano/submicron-sized precursors and high heating rates achieved with the aid of non-conventional SPS process allowed the preparation of well densi ed nitrogen rich sialon ceramics at relatively low temperature (1500°C as compared to 1700°C or 1800°C). Alpha sialon-phase for samples having Ca content (c) in the range of 0.15 < x < 1.83 was obtained. Formation of single phase nitrogen rich alpha-sialons was achieved for samples having nominal x value in the range of 0.4 < x < 1.6. Increase in the lattice parameters of alpha-sialon unit cell was observed to have a linear relationship with the amount of Ca +2 ions incorporated in the alpha structure. Increasing x value resulted in an increase in the average grain size from 286nm to 468nm.
Moreover, higher amount of Ca 3 N 2 starting powder was observed to facilitate the formation of elongated alpha-sialon grains. Calcium stabilized nitrogen rich sialon ceramics having very promising hardness and fracture toughness of 22.4 GPa and 6.1 MPa.m 1/2 respectively were developed. Furthermore, the absence of a glassy grain boundary phase in nitrogen rich sialons make them potentially viable candidates for high temperature applications. It is believed that such engineering ceramic materials would appeal to the industry owing to their unique properties and relatively low processing temperature.
Declarations Figure 2 Schematic of spark plasma sintering equipment showing the die and temperature measurement con guration.

Figure 3
Densi cation curves of several compositions of calcium stabilized nitrogen rich sialons sintered at 1500 C using spark plasma sintering technique, where Ca content varies between 0.2-2.2. The insert shows the complete densi cation curves while the main plot represents the selected high temperature region (heating from 1050°C to 1500°C followed by 30 min holding time).     Grain size distribution of nitrogen rich samples synthesized at 1500 C having the sample ID's of Ca-0.6, Ca-1.2 and Ca-2.0.