Refractory transition-metal oxides such as zirconia are an industrially important class of ceramic materials ^{[1] and [2]}. Zirconia and zirconia-based ceramics are of both scientific and technological interest as structural and functional materials due to the superior properties of zirconia. Zirconia (ZrO₂) is hard with good mechanical properties and has a high fracture toughness and hardness, possesses a high melting point (2710 °C), and is resistant to chemical attack by strong acids ^[3] along with excellent refractance. Zirconia is also stable in oxidizing environments, allows oxygen diffusion through its bulk ^[4], exhibits low thermal conductivity at high temperature, good thermal stability and resistance to thermal shock, large ionic conductivity, and is electrically insulating. These properties enable zirconia to be widely used, e.g. as an abrasive, as a hard resistant coating for cutting tools, in oxygen electrodes and sensors, and in high temperature engine components ^{[1], [2], [5], [6] and [7]}. Few materials show as much potential as ZrO₂ and ZrO₂-based ceramics, because of such a wide range of advanced engineering ceramic applications. To produce these ceramics, high-purity powders with uniform particles of submicron size and good sintering properties are required ^{[8] and [9]}.

Zirconia has four well-defined polymorphic forms at atmospheric pressure: the monoclinic, tetragonal, orthorhombic and cubic phases ^[10]. Only the monoclinic form is stable at room temperature. During the last decades an active research has been undertaken in different laboratories, in order to obtain zirconia and zirconia-based ceramic powders with the required characteristics of size, purity, uniformity, crystallinity and etc. As a result of the intensive work done, a wide variety of physical and chemical preparation techniques were found ^[11].

Conventionally, ZrO₂ is synthesized from the high temperature decomposition of zirconium compounds such as nitrates, carbonates, or silicates ^{[2], [3], [12], [13] and [14]}. In recent years low temperature synthetic methods have been developed. Nowadays due to the need for high performance products in ceramic industry there has been an increasing interest in the synthesis of nanopowders. The nanostructured zirconia-based ceramic powders are gaining increasing importance and various routes have been employed for the production of nanocrystalline zirconia-based particles, such as co-precipitation ^[15], sol–gel preparation method ^[16], hydrothermal synthesis ^{[17] and [18]}, polymerized complex processes ^[19] and gel-combustion process ^{[15], [20], [21], [22], [23], [24] and [25]}.

One of the most powerful techniques for the synthesis of a wide range of materials such as zirconia is a solid-state process named mechanochemical treatment. This technique is a non-equilibrium solid-state process in which the final product retains a very fine amorphous or nanocrystalline structure. This process is versatile, simple and cost effective and can be used to synthesize a wide variety of materials with a high yield rate on a commercial scale. This process can also be designed in such a way to synthesize nanocrystalline particles dispersed within a soluble matrix. In other words, chemical precursors react, either during milling or at the subsequent heat treatment stage, to form a nanocrystalline powder embedded within a soluble salt matrix phase. The ultrafine powder is subsequently recovered by selective removal of the matrix phase through washing with an appropriate solvent. The ability to decrease time and temperature of the chemical reaction at low energy consumption is another advantage of this technique ^{[26], [27] and [28]}.

The aim of the present work was to investigate on the mechanochemical treatment of anhydrous ZrCl₄ and CaO to obtain more information on the nature of this process and to determine the effect of thermal treatment on the structural evolutions and morphological characteristics of the nanopowder and also to obtain the activation energies of crystallite growth for the different polymorphs of the zirconia phase. The mechanism of the nanocrystallite growth is also discussed.

Zirconia was prepared according to the previously published work ^[29]. The starting materials were anhydrous ZrCl₄ (Acros) and CaO (Alfa Aesar). Stoichiometric amounts of ZrCl₄ and CaO were used according to reaction without any process control agent. The starting materials were weighted and prepared in a glovebox under inert atmosphere and after sealing the milling jar, it was transferred to the milling machine. Mechanical treatment of reactant mixtures was carried out in a FRITSCH planetary mill (planetary micro mill pulverisette7) at 5000 r/min under inert atmosphere using 12 and 4 alumina balls of 10 and 15 mm in diameter respectively and an alumina jar of 80 ml. The ball to powder mass ratio of 10:1 and 5 h milling time was used in this study. All of the samples were prepared in a glovebox before the milling procedure to avoid moisture adsorption.

The heat treatment of the prepared powders was carried out at a rate of 100 °C/h, under air atmosphere with 1 h holding time at the maximum temperature. The milled powders were annealed at 250 °C, where all the starting powders had reacted. The reaction produced CaCl₂(2H₂O) and Zr(OH)₄ phases ^[29]. The removal of the salt-by product (CaCl₂(2H₂O)) was performed by washing the powder with distilled water several times using a sonifier and centrifuge. The washed powder was dried in an oven and all of the subsequent heat treatments were carried out in air at temperatures between 400 and 1700 °C. The phase composition and crystal structures of the prepared powders were investigated at room temperature using X-ray powder diffraction measurements. X-ray powder diffraction data were recorded under ambient conditions on a high resolution laboratory X-ray powder diffractometer (Bruker D8 ADVANCE, Mo Kα₁ radiation from primary Ge(220)-Johannson type monochromator, Lynx-Eye position sensitive detector (PSD) in Debye-Scherrer geometry) at 50 kV and 40 mA. All the X-ray samples were filled in low absorbing glass capillaries of 0.3 mm diameter (Hilgenberg glass No. 14) and sealed in a glovebox under argon atmosphere using a hot wire. Data were taken in steps of 0.007° 2 θ from 2° to 55° of 2 θ at a rate of 0.06°/min. The samples were spun during measurement for better particle statistics. A quantitative phase analysis of the measured powder diffraction patterns using the PDF-2 database (ICCD, 2007) in combination with the program MATCH ^[30] was used to identify the phases. For Rietveld refinement the Topas4-2 program ^[31] was used. To overcome difficulties and limitations of other methods of microstructural analysis and to consider all the benefits of the whole profile fitting methodology, Rietveld's powder structure refinement procedure based on pseudo-Voigt (pV) profile fitting function has been adopted in the present study. The thermal behavior of the washed calcined as-milled powder was investigated using a DSC analyzer under air atmosphere at a heating rate of 10 °C min^-1 and a sample size of about 10 mg per run. The microstructures of the samples were analyzed using a transmission electron microscope (JEOL 4000 FX) and a scanning electron microscope. Raman spectroscopy was used for further verification of the presence of tetragonal zirconia.

The starting powder mixture of ZrCl₄ and CaO in stoichiometric molar ratio was milled with the aim of inducing the mechanical reaction according to reaction (1).

ZrCl₄ + 2CaO = ZrO₂ + 2CaCl₂ (1)

According to our previous work ^[29], the synthesis mechanisms maybe described by the following reactions. After milling the ZrCl₄ and CaO mixture for 5 h at ball to powder ratio (BPR) = 10, ZrCl₄ is partially amorphous while the crystallite size of CaO is slightly decreased (according to the XRD results not shown here). The formation of zirconium hydroxide occurred after the heat treatment of the milled powder which completed at 250 °C ^[29]. Based on thermodynamic calculations the amount of H₂O present in air is about 285 times more than needed. Finally the washed heat treated powder containing amorphous Zr(OH)₄ converts into ZrO₂ with removing the hydroxyl groups during the second heat treatment stage as shown in reaction .

Zr(OH)₄ = ZrO₂ + 2H₂O (2)

Fig. 1 presents the differential scanning analysis (DSC) curve of the dried zirconium hydroxide powder. After heating the sample from room temperature up to 1100 °C, two endothermic peaks at around 87 °C and 177 °C are present due to the release of adhesive water and structural water respectively. Furthermore, two major exothermic peaks due to crystallization of tetragonal ZrO₂ from amorphous Zr(OH)₄ and the conversion of the tetragonal to the monoclinic ZrO₂ polymorph were observed at 500 and 998 °C respectively. These experimental data are in agreement with the results reported by Dodd and McCormick ^[32].

Fig. 1.

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	Fig. 1. DSC curve of the zirconium hydroxide powder heat treated up to around 1200 °C (heating rate 10 K/min).

The zirconium hydroxide powder was calcined for 1 h in the temperature range of 400–1700 °C. The whole pattern Rietveld refinement of the powders calcined at 500 °C and 1300 °C are shown in Fig. 2. It is clear that the heat treatment not only results in the transformation of the amorphous zirconium hydroxide into crystalline phases but also increases the crystallite size of the crystalline phases. Furthermore, increasing the temperature results in the phase transformations as follows:

Amorphous zirconium hydroxide → t-ZrO₂ → m-ZrO₂ (3)

Fig. 2.

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	Fig. 2. XRD patterns of the heat treated samples for 1 h at 500 °C (a) and 1300 °C (b).

The results of the Rietveld refinement of the X-ray diffraction patterns along with the summarized results from our previous work ^[29] for the temperature range of 400–1700 °C are presented in . Fig. 3 presents the weight change of the amorphous, tetragonal and monoclinic phases with temperature according to the Rietveld refinements. According to Fig. 3 the amount of amorphous zirconium hydroxide phase decreases with increasing the temperature and finally converts to t-ZrO₂ at temperatures higher than 600 °C. In the case of t-ZrO₂ it can be observed that the weight percent of this phase increases to a maximum value at around 450 °C and then suddenly decreases up to around 700 °C. The t-ZrO₂ still exists in slight amounts up to around 1100 °C when the transition from t-ZrO₂ into m-ZrO₂ is finished. For the m-ZrO₂ it can be observed that its weight percent increases from 400 °C to 1700 °C. This increase suddenly happens in the range of 500 °C–700 °C which is related to the sudden decrease of t-ZrO₂ due to its phase transformation. Furthermore it should be mentioned that the diffraction pattern of the tetragonal phase is almost identical to that of the cubic phase except for the “splitting” of a number of peaks to form doublets, caused by the slight elongation of the c-axis. When the peaks are broad, for instance due to small crystal size, it can be difficult to distinguish the diffraction patterns of the two phases unless the high-angle peaks are examined ^{[33] and [34]} or vibrational spectroscopy is used ^[35]. There are a number of reports for crystalline zirconia that have a cubic-like diffraction pattern, but are shown to have the tetragonal structure when non-diffraction techniques are used, such as Raman spectroscopy ^{[36] and [37]}, infra-red spectroscopy ^[38], EXAFS ^[39], and Perturbed Angular Correlations method ^{[40] and [41]}, or a high resolution diffraction technique is applied ^[42]. Although our Rietveld refinements were applied on patterns obtained from a high resolution diffractometer giving good R_wp and goodness of fit (GOF) values which confirms the existence of tetragonal zirconia instead of cubic zirconia but Raman spectroscopy was also applied to completely ensure that tetragonal zirconia is the present phase.

Table 1.

Table 1.

Table 1. Crystallite size (nm) and weight percent (wt%) of zirconium hydroxide samples calcined at different temperatures for the corresponding phases along with the Rwp and GOF of the Rietvield refinement Calcination temp. (°C) Amorphous phase (Zr(OH)4) Tetragonal ZrO2 Monoclinic ZrO2 Rwp GOF Crystallite size wt% Crystallite size wt% Crystallite size wt% 400 – 66.45 37.14 26.41 7.8 7.14 0.06574 1.279 420 – 45.50 40.59 43.10 9 11.40 0.06780 1.337 450 – 23.41 44.29 60.17 12.7 16.42 0.07788 1.401 500 – 19.50 45.63 58.42 15.4 22.08 0.08268 1.531 600 – 4.54 33.74 18.52 20.3 76.94 0.07077 1.315 700 – – 32.80 6.63 26.9 93.37 0.07109 1.314 800 – – 38.24 5.62 39.1 94.38 0.07060 1.392 900 – – 30.69 2.36 52.2 97.64 0.06809 1.393 1000 – – 41.05 2.09 70.9 97.91 0.07196 1.330 1100 – – 41.69 1.26 96.1 98.74 0.07771 1.408 1200 – – – – 115.7 100 0.07759 1.590 1300 – – – – 134.49 100 0.07662 1.502 1500 – – – – 175.35 100 0.07182 1.229 1700 – – – – 198.59 100 0.07499 1.401 Note:$R_{\text{wp}}=\sqrt{\left(\sum w_m{\left(Y_{o,m}-Y_{c,m}\right)}^2\right)/\left(\sum w_m{Y}_{o,m}^2\right)}$ and$\text{GOF}=\sqrt{\left(\sum w_m{\left(Y_{o,m}-Y_{c,m}\right)}^2\right)/\left(M-P\right)}$, where Yo,m and Yc,m are the observed and calculated data respectively at data point m, M the number of data points, P the number of parameters and w m the weighting given to data point m.

Table 1. Crystallite size (nm) and weight percent (wt%) of zirconium hydroxide samples calcined at different temperatures for the corresponding phases along with the Rwp and GOF of the Rietvield refinement

Fig. 3.

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	Fig. 3. Weight percent of different phases vs temperature.

The Raman spectra of the monoclinic and tetragonal phase are well-known and can be very clearly distinguished ^{[43] and [44]}. Most of the reported spectra for tetragonal zirconia are obtained from different stabilized t-ZrO₂ (“doped” with different kinds of oxides) ^{[44], [45] and [46]}, or from pure zirconia at high temperatures ^{[43] and [47]}. However due to the very similar crystal structure of “metastable” and stabilized t-ZrO₂ the main features of the spectrum are expected to be the same. A number of Raman spectra of the metastable tetragonal phase have been published, mostly with poor signal-to-noise ^{[35], [36], [37], [48] and [49]}. The Raman spectrum of the stabilized-cubic phase is known ^[44], and is quite distinguishable from that of the tetragonal phase, but no spectrum has been reported for the pure cubic phase.

According to our Rietveld refinements t-ZrO₂ is present without any dopants such as Ca, Cl etc. According to Southon pure metastable t-ZrO₂ has a unit cell tetragonality ( c/ a ratio) in the range of 1.017–1.020 depending on the calcinations temperature ^[50]. The presence of any dopant in the structure of t-ZrO₂ will change the c/ a ratio to higher or lower values. To assure that the Cl^- ions were completely removed from the washed powder they were checked by 0.1 N AgNO₃ solution after washing them several times with DI-water. However the energy dispersive X-ray spectroscopy (EDAX) of transmission electron microscopy (TEM) and SEM experiments (not shown here) and also the Raman spectroscopy verified the absence of any doping elements in the t-ZrO₂ phase. According to results reported by Stefanic et al. the Raman spectroscopy of pure metastable tetragonal zirconia exhibits all or some of the following bands: 642, 463, 314, 267 and 147 cm^{-1 [49]}. The existence of any dopants results in a shift in the mentioned bands. Furthermore it has been shown by Yashima et al. ^[45] that the intensity ratio of pure and stabilized phases is different. They observed that the intensity ratio of the band at approximately 460 and 640 cm^-1 increases from zero at 20 mol% dopant (cubic phase) to approximately 0.35 at 6 mol%. This change is attributed by the authors to be a function of oxygen displacement along the c-axis in the unit cell, which increases with decreasing dopant. Thus it can be stated that the oxygen displacement in the metastable tetragonal phase is large in comparison to the stabilized phase, and is consistent with the trend observed by Yashima et al. ^[45]. The Raman spectroscopy of the sample calcined at 1100 °C is shown in Fig. 4 with the corresponding bands of t-ZrO₂ marked as “t” while the other bands correspond to the m-ZrO₂ phase which are indicated as “m”.

Fig. 4.

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	Fig. 4. Raman spectroscopy of the sample calcined at 1100 °C.

The evolution of crystallite size of the ZrO₂ powder during calcination has been investigated and the results obtained from Rietveld analysis of the X-ray patterns are summarized in Fig. 5 and Fig. 6 for the tetragonal and monoclinic phase, respectively. In Fig. 5(a) it can be seen that the crystallite size of t-ZrO₂ does not have a continuous increasing trend. The behavior of t-ZrO₂ is such that, after a critical average size is reached, it converts to m-ZrO₂. It should be mentioned that the t-ZrO₂ is a metastable phase at this temperature while m-ZrO₂ is the thermodynamically stable phase. This non-continuous growth behavior is due to the presence of the un-stabilized pure t-ZrO₂ phase. For the m-ZrO₂ phase according to Fig. 6(a) it can be observed that by increasing the temperature the crystallite size increases. The increase in crystallite size is attributed to the typical effect of temperature on the crystal growth. Based on Fig. 6(a) it can be mentioned that the growth rate of the nanocrystallites of m-ZrO₂ is nearly linear. However the growth rate of crystallites with sizes out of the nano-range changes, this is due to the fact that in the nanosize range the porosity is quite high and the pores are interconnected to maintain smaller crystal sizes ^[51]. In the case of crystallites with sizes out of the nano-range the bridging of fine particles form continuous grain boundary networks which increases the crystal sizes with a higher rate.

Fig. 5.

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	Fig. 5. (a) Effect of calcination temperatures on the crystallite size of t-ZrO2 and (b) plot of ln D against 1/ T.

Fig. 6.

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	Fig. 6. (a) Effect of calcination temperatures on the crystallite size of m-ZrO2 and (b) plot of ln D against 1/ T.

As shown in Fig. 5(b) and Fig. 6(b) the nanocrystallite growth during annealing can be obtained by the Scott equation by plotting the straight line of ln D against 1/ T, assuming that the nanocrystallite growth is homogeneous ^{[52], [53], [54] and [55]} as follows:

$D=A\exp \left(\frac{-E}{RT}\right)$ (4)

where D is the crystallite size obtained from the XRD pattern (in this research the whole pattern Rietveld refinement method is used to obtain more accurate results), A is the constant, E is the activation energy for nanocrystallite growth, R is the ideal gas constant and T is the absolute temperature of heat treatment. The activation energy for the t-ZrO₂ and m-ZrO₂ phases is calculated to be around 13,715 and 27,333 J/mol, respectively. In the case of t-ZrO₂, increasing trends are considered and the presented value is the mean value of the activation energy obtained from the increasing trends. By increasing the temperature above 1100 °C, surface diffusion can be activated and it can cause sintering and formation of necks between particles. It can be considered that the crystallite grows by means of an interfacial reaction ^[54].

Fig. 7 shows the TEM micrographs of the powders calcined at 400 and 420 °C. As can be seen from Fig. 7(a) and (b) the size of the as formed particles from the amorphous zirconium hydroxide are about 5 nm. From the TEM results and the Rietvield refinement it can be concluded that we have a distribution of crystallite and particle size. As the crystallite size of t-ZrO₂ passes the critical size it transforms to m-ZrO₂ leaving the smaller sized t-ZrO₂ behind, which will convert to m-ZrO₂ upon further heat treatment after passing the critical crystallite size. In Fig. 7(a) a crystalline zirconia particle is shown which has been crystallized at 400 °C from the amorphous zirconium hydroxide phase (the particle is shown with high magnification in the inset) while in Fig. 7(b) agglomerated freshly formed particles calcined at 420 °C are shown. In addition the TEM micrographs show that most of the particles are spherical in shape.

Fig. 7.

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	Fig. 7. TEM image of the nanopowders heat treated at (a) 400 °C and (b) 420 °C.

Fig. 8 presents the SEM images of the amorphous zirconium hydroxide powder heat treated at 400, 700, 800, 900, 1100 and 1200 °C. As seen in Fig. 8(a) individual spherical particles (probably crystalline zirconia particles) are present in a porous network (most likely due to the phase transformation of amorphous zirconium hydroxide to t-ZrO₂). In Fig. 8(b) the powder calcined at 700 °C is shown consisting of nanoparticles. Increasing the temperature up to 800 °C, results in the clinging of particles forming larger agglomerated structures as shown in Fig. 8(c). After calcination at 900 °C it can be seen that individual particles are gathered together forming big spherical agglomerates ( Fig. 8(d)). Increasing the temperature to 1100 °C results in slightly larger particles while at 1200 °C the particles grow significantly as shown in Fig. 8(e) and (f), respectively. Generally, it can be concluded that with the increase in temperature, the activated mechanism is surface diffusion and interface reaction which result in increased particle size, neck formation and more agglomeration at higher calcination temperatures. In the case of low calcination temperature the porosity is quite high and most pores are interconnected to maintain smaller crystal sizes. For the higher calcination temperatures continuous grain boundary networks have been formed due to the bridging of fine particles to increase the crystal sizes. The agglomeration and densification processes continue with the high calcination temperatures as seen in the SEM images ( Fig. 8).

Fig. 8.

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	Fig. 8. SEM images of the powders heat treated at different temperatures: (a) 400 °C, (b) 700 °C, (c) 800 °C, (d) 900 °C, (e) 1100 °C and (f) 1200 °C.

Nanocrystalline ZrO₂ particles have been successfully prepared by heat treatment of as-milled powder obtained by mechanochemical reaction of ZrCl₄ and CaO as raw materials. It has been found that after milling, heat treatment and washing, the formed zirconium hydroxide transforms to tetragonal zirconia, and monoclinic zirconia upon further heat treatment. The calculation of the activation energy for the tetragonal and monoclinic polymorphs is 13,715 and 27,333 J/mol, respectively, indicating that the nanocrystalline of these polymorphs grows by means of an interfacial reaction.