Quantitative analysis of the structural stability and degradation ability of hydroxyapatite and zirconia composites synthesized in situ

Abstract

In situ synthesis has been attempted in the present study to form hydroxyapatite [HAP, Ca₁₀(PO₄)₆(OH)₂] and zirconia [ZrO₂] composite powders. The synthesis has been attempted using an aqueous precipitation technique in which alterations have been made to the starting Ca/P molar ratio of calcium phosphate suspensions together with five different concentrations of ZrO₂ precursors. An additive in the form of yttria to stabilize tetragonal ZrO₂ (t-ZrO₂) has not been used in the present investigation. The precipitated powders resulting were heat treated at different temperatures before determination of the phase stability and degradation ability of the HAP–ZrO₂ composites. The powders were characterized by X-ray diffraction, Raman spectroscopy and quantitative analysis by Rietveld refinement. The results from the investigation confirmed the formation of HAP–ZrO₂ composites at 900 °C. The Ca/P precursors with molar ratios of 1.67 and 1.69 have yielded the formation of β-tricalcium phosphate [β-TCP, β-Ca₃(PO₄)₂] with a content of less than 5 wt% together with HAP and t-ZrO₂ at 900 °C. The Ca/P precursors with a molar ratio of 1.68 have yielded only HAP and t-ZrO₂ composite powders at 900 °C. The increased addition of ZrO₂ precursors had resulted in the enhanced formation of t-ZrO₂ with a simultaneous reduction in the amount of the HAP phase formed. However, heat treatment at 1000 °C and 1100 °C has initiated the degradation of HAP and t-ZrO₂ composite powders resulting in the formation of monoclinic ZrO₂ (m-ZrO₂). It is surmised from the present results that there is active involvement of the ZrO₂ content in the degradation behaviour of the HAP–ZrO₂ composite powders.

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1. Introduction

The stoichiometric hydroxyapatite [HAP, Ca₁₀(PO₄)₆(OH)₂] and several ionic substitutions in HAP has found many applications in the treatment of bone- and tooth-related defects for the past three decades because of their excellent biocompatibility and chemical similarities with natural bone and the minerals in teeth.^1–6 The inferior mechanical features of synthetic HAP in comparison with the biological bone have restricted its applications to the form of particulates or granules to fill either bone- or tooth-related defects. Several attempts have been made to recognize synthetic HAP as a suitable substitute in the replacement of hard tissues where the demands for high mechanical properties are necessary. Therefore, HAP has been used in the form of a reinforcement either with metals such as titanium alloy and stainless steels or with bioinert ceramics such as alumina (Al₂O₃), zirconia (ZrO₂) and titania (TiO₂). The HAP coatings on metallic implants have serious limitations such as metallic corrosion, and weak interfacial bonding between the metal and the ceramic, which is expected to result in the dissolution of the ceramic.^7–10 The plasma spray technique used to develop a HAP coating on metallic implants resulted in strong adhesion between the ceramic and the metallic alloy; however, such coatings had resulted in the phase decomposition of stoichiometric HAP to yield secondary phases. Al₂O₃ has been known for its features such as high mechanical stability, corrosion resistance and high wear resistance.¹¹ Generally the phase stabilization of Al₂O₃ occurs at a higher temperature whereas the phase decomposition of HAP starts to occur beyond 1200 °C and this leads to the dehydroxylation and conversion of HAP to tetracalcium phosphate [Ca₄(PO₄)₂O] and α-tricalcium phosphate [α-TCP, α-Ca₃(PO₄)₂]. This mismatch in the thermal stability of HAP and Al₂O₃ composites has led to difficulties in the preservation of their respective phases in the HAP/Al₂O₃ composite and some researchers have also reported on the chemical interaction between HAP and Al₂O₃ to form a CaAl₂O₄ phase.^12,13 The investigation of the reinforcement of TiO₂ with HAP has been mainly based on the fabrication of HAP/TiO₂ composite coatings on titanium metal. The reason for the development of such composite coatings on titanium metal is mainly based on the fact that the reinforcement of bioinert TiO₂ in the HAP matrix would substantially improve the mechanical reliability of HAP, thereby leading to the increased stability of the coatings.^14,15 The thermal decomposition of HAP/TiO₂ composites to form calcium oxide (CaO), calcium titanate and the combination of rutile TiO₂ and anatase TiO₂ have also been reported.¹⁶ ZrO₂ has been known to possess better mechanical features when compared to Al₂O₃ ¹⁷ and their comparative investigations have concluded that ZrO₂ has shown better resistance to fracture compared to Al₂O₃.¹⁸

It is well known that ZrO₂ exists in three polymorphs: monoclinic (stable from room temperature to 1170 °C with P2₁/c symmetry), tetragonal (stable between 1170 °C to 2370 °C with P4₂/nmc symmetry) and cubic (stable above 2370 °C with Fm3m symmetry).¹⁹ Among these polymorphs, tetragonal ZrO₂ (t-ZrO₂) has been found to be suitable for hard tissue replacements because it has favourable features such as wear resistance, chemical resistance, high fracture toughness and hardness. Many attempts have been made to combine the advantages of biocompatible HAP and mechanically stable ZrO₂ to form a composite which will be suitable for hard tissue replacements. The reported methods for the fabrication of HAP–ZrO₂ composites include the mechanical mixing of either a commercially procured or a synthesized individual HAP and ZrO₂ particles to form a composite.^20–25 The conventional heat treatment of the mechanically mixed individual powder particles of HAP and ZrO₂ give the formation secondary phases such as CaZrO₃, α-Ca₃(PO₄)₂, Ca₄(PO₄)₂O, and CaO, in addition to the presence of HAP and ZrO₂ phases. Thus, the presence of additional phases has led to difficulties in the precise determination of the quantitative content of HAP and ZrO₂ in the composites. Some of the studies have also reported on the application of a spark plasma sintering process for mechanically mixing the individual HAP and yttria (Y₂O₃)-stabilized ZrO₂ (y-ZrO₂) particles to yield a composite consisting only of the phases of HAP and ZrO₂ particles.^26–28 Such HAP–ZrO₂ composites have resulted in a variable mechanical strength that was solely dependent on the amount y-ZrO₂ added to the composites. Thus, much of the reported research on the development and fabrication of HAP–ZrO₂ composites has only dealt with the mechanical mixing of the separately synthesized HAP and ZrO₂ particles to form composites. The drawbacks in such studies are the difference in the grain growth of HAP and ZrO₂ particles during sintering, non-uniformity in the mixing level of the individual HAP and ZrO₂ particles throughout the composite matrix and lack of precise determination of the quantitative content of individual particles in the matrix. Thus, the present study has been focussed on the in situ synthesis of five different HAP–ZrO₂ composites starting from the solution precursors and followed by the application of the Rietveld refinement technique for the quantitative determination of the resulting individual HAP and ZrO₂ particles in the composites.

2. Materials and methods

2.1 Powder synthesis

An aqueous co-precipitation technique was used in the present investigation to synthesise different concentrations of HAP–ZrO₂ ceramic composite powders. Calcium nitrate tetrahydrate [Ca(NO₃)₂·4H₂O, HiMedia, India], ammonium di-hydrogen orthophosphate [NH₄H₂PO₄, HiMedia, India] and zirconium oxychloride octahydrate [ZrOCl₂·8H₂O, HiMedia, India] were used as precursors for Ca²⁺, PO₄³⁻ and ZrO₂. For the synthesis of the different HAP–ZrO₂ composite powders, stock solutions of appropriate molar concentrations of Ca(NO₃)₂ and NH₄H₂PO₄ were prepared separately in Millipore water. The molar concentrations of Ca(NO₃)₂ and NH₄H₂PO₄ precursor solutions were maintained for all the five different concentrations in order to ensure the stoichiometric Ca/P molar ratio of 1.67 for all the five different compositions. Five different concentrations of a ZrOCl₂·8H₂O stock solution ranging from 0.1 M to 0.5 M were prepared separately in Millipore water. In addition to the Ca/P molar ratio of 1.67, two more Ca/P ratios of 1.68 and 1.69 were also synthesized in situ with the ZrO₂ precursors to form a wide range of composite powders. A total of 15 different concentrations were prepared and their respective sample codes for all the 15 different compositions are indicated in Table 1.

Table 1 Molar concentrations of the precursors used in the synthesis of HAP–ZrO₂ composite powders

No.	Sample code	Molar concentrations of precursors			Ca/P molar ratio	Ca/Zr molar ratio
		Ca(NO3)2·4H2O	NH4H2PO4	ZrOCl2·8H2O
1.	1.67CaP_1xZrO2	0.5 M	0.2994 M	0.1	1.67	5.00
2.	1.67CaP_2xZrO2	0.5 M	0.2994 M	0.2	1.67	2.50
3.	1.67CaP_3xZrO2	0.5 M	0.2994 M	0.3	1.67	1.67
4.	1.67CaP_4xZrO2	0.5 M	0.2994 M	0.4	1.67	1.25
5.	1.67CaP_5xZrO2	0.5 M	0.2994 M	0.5	1.67	1.00
6.	1.68CaP_1xZrO2	0.5 M	0.2976 M	0.1	1.68	5.00
7.	1.68CaP_2xZrO2	0.5 M	0.2976 M	0.2	1.68	2.50
8.	1.68CaP_3xZrO2	0.5 M	0.2976 M	0.3	1.68	1.67
9.	1.68CaP_4xZrO2	0.5 M	0.2976 M	0.4	1.68	1.25
10.	1.68CaP_5xZrO2	0.5 M	0.2976 M	0.5	1.68	1.00
11.	1.69CaP_1xZrO2	0.5 M	0.2959 M	0.1	1.69	5.00
12.	1.69CaP_2xZrO2	0.5 M	0.2959 M	0.2	1.69	2.50
13.	1.69CaP_3xZrO2	0.5 M	0.2959 M	0.3	1.69	1.67
14.	1.69CaP_4xZrO2	0.5 M	0.2959 M	0.4	1.69	1.25
15.	1.69CaP_5xZrO2	0.5 M	0.2959 M	0.5	1.69	1.00

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As a preliminary step in the synthesis, the stock solution of NH₄H₂PO₄ was added dropwise to the continuously stirred Ca(NO₃)₂ solution. The temperature of the stirred suspension was maintained at 80 °C throughout the reaction. After the completion of the NH₄H₂PO₄ addition, the pH of the suspension was increased from the initial value of ∼4.00 to 10.00 by the addition of concentrated NH₄OH solution. The resulting mixture was allowed to stir for another 30 minutes to ensure the formation of a calcium phosphate precursor. Then, the separately prepared stock solution of ZrOCl₂ was added dropwise to the continuously stirred calcium phosphate mixture and the resulting mixtures containing the precursors of Ca²⁺, PO₄³⁻ and ZrO₂ were continuously stirred for two hours. The resulting mixture was then transferred to a hot air oven and dried at 120 °C for 48 hours. The dried precipitates obtained were ground into fine powder with a pestle and mortar and this powder was considered to be the as-prepared composite powder. Pure HAP and pure ZrO₂ have also been prepared simultaneously by adapting a similar aqueous precipitation technique for comparative purposes.

2.2 Characterization of powders

All the finely ground as prepared powder samples were heat treated at various temperatures to investigate the ability of the formation of a varied range of HAP–ZrO₂ composites. The powder samples were subjected to heat treatment at a rate of 5 °C per minute over various temperature ranges with a soaking time of two hours in a muffle furnace (MATRI-MC 2265 A, India) and subjected to X-ray diffraction (XRD) to observe any change in the phase behaviour and crystallinity. The phase purity and composition of all the composite powders were determined using a powder X-ray diffractometer (RIGAKU, Ultima IV, Japan). XRD studies for all the powders were carried out with Cu Kα radiation (λ = 1.5406 nm) produced at 40 kV and 30 mA to scan the diffraction angles (2θ) between 10° and 70° with a step size of 0.02° 2θ per second. Phase determinations were made using standard International Centre for Diffraction Data (ICDD) card no. 00-009-0432 for HAP, 01-079-1765 for t-ZrO₂, 01-071-4810 for cubic ZrO₂ (c-ZrO₂), 01-083-0944 for monoclinic ZrO₂ (m-ZrO₂) and 00-009-0169 for β-Ca₃(PO₄)₂.

The vibrational modes for HAP–ZrO₂ composite powders were determined by using Raman spectroscopy. All the vibrational modes of HAP–ZrO₂ composite powders after heat treatment at 900 °C, 1100 °C and 1100 °C were determined by using the back scattering geometry of a confocal Raman microscope (Renishaw, UK). All the powder samples were excited at a wavelength of 785 nm by a semiconductor diode laser (0.5% of power) with a data acquisition time of 30 seconds. The powder samples were analysed in the spectral range of 100 to 1200 cm⁻¹.

2.3 Quantitative analysis by Rietveld refinement

Rietveld analysis was used to provide quantitative information about the phases present in the sample. In the present study, quantitative phase analysis was performed for stoichiometric HAP, ZrO₂ and five different HAP–ZrO₂ composites samples by using the Rietveld method using a GSAS-EXPGUI software package.^29,30 For the analysis by the Rietveld refinement, an average of three scans were recorded for each powder samples with Cu Kα radiation (λ = 1.5406 nm) produced at 40 kV and 30 mA to scan the diffraction angles (2θ) between 10° and 70° with a step size of 0.02° 2θ per second using an X-ray powder diffractometer (RIGAKU, Ultima IV, Japan). Quantitative refinement was carried out for all the five different HAP–ZrO₂ composites, and the stoichiometric HAP and ZrO₂ were determined for powder samples heat treated at 900 °C, 1000 °C and 1100 °C. All the standard crystallographic information was obtained from the American mineralogist crystal structure database. The standard crystallographic data for the refinement of HAP, t-ZrO₂, β-Ca₃(PO₄)₂ and m-ZrO₂ were obtained from Wilson et al.,³¹ Howard et al.,³² Yashima et al.³³ and Smith and Newkirk,³⁴ respectively. In structural refinement, numerous cycles were run to achieve the quantitative analysis, weight fraction and bond length of the composite powders. In the first step of the refinement, all the structural parameters were fixed to values from the literature. Then, during the successive refinement cycles, numerous parameters were allowed to vary according to the relative weight of the observed phases. The following refinement sequence has been used as a standard for all the structures: scale factor, zero shift, background as Chebyshev polynomial of the fifth grade, peak profile, and lattice parameters. Fittings were performed using pseudo-Voigt peak profile functions and a preferred orientation along [001] was accounted for with the Marsh model. The fractional coordinates, isotropic temperature and atomic parameters were employed during refinement.

3. Results and discussion

3.1 Formation of HAP–ZrO₂ composite powders

The as-prepared powders were subjected to heat treatment at various temperatures in order to analyse the phase transformation behaviour. The XRD patterns of the as prepared powders (Fig. 1 and 2) were characterized by broad diffraction bands and no significant information could be obtained because of the excessive amorphous content present in the samples which had mainly arisen from the precursor impurities (NH₄⁺, Cl⁻, NO₃²⁻ and adsorbed H₂O) of the aqueous precipitation technique. Fig. 1 and 2 present the corresponding XRD patterns of the 1.68CaP_1xZrO₂ and 1.68CaP_3xZrO₂ samples after heat treatment at different temperatures. The heat treatment at 500 °C had resulted in the emergence of diffraction patterns characteristic of the HAP phase and no peaks were present for the ZrO₂ phase. The heat treatment at 700 °C had resulted in the formation of well-defined diffraction patterns that corresponded to the HAP phase and the formation of the ZrO₂ phase could be identified by the appearance of a small peak at the 2θ angle of 30°. Successive heat treatments at higher temperatures ranging from 800 °C to 1100 °C had resulted in the enhancement of diffraction patterns that corresponded to the presence of both HAP and ZrO₂ phases. However, the heat treatment at 1000 °C and 1100 °C had resulted in the formation of very small intensity diffraction patterns, which were characteristic of the β-Ca₃(PO₄)₂ phase, which has not been identified in the sample that was heat treated at 900 °C. The characteristic peaks for the presence of m-ZrO₂ phase were detected at 1100 °C for all the samples, despite the varied ZrO₂ precursor concentration added during the synthesis and also the varied Ca/P molar concentration maintained during the synthesis. It was observed from a previous investigation by the current authors³⁵ that pure ZrO₂ without the addition of any precursors had resulted in the formation of mixtures comprising c-ZrO₂ and m-ZrO₂ phase at 500 °C and heat treatment at elevated temperatures (700 °C, 900 °C and 1100 °C) had resulted in the enhanced presence of m-ZrO₂ with the corresponding presence of a c-ZrO₂ phase at the reduced level. It should be noted that the presence of the t-ZrO₂ phase was not detected at any of the heat treatment temperatures used in the investigation of the pure ZrO₂ system. The results from the present investigation have shown that the presence of the HAP phase had restricted the formation of the c-ZrO₂ and m-ZrO₂ phases at the temperatures studied until 1000 °C was reached and simultaneously it had stabilized the t-ZrO₂ phase.

Fig. 1 XRD patterns of the 1.68CaP_1xZrO₂ sample after heat treatment at different temperatures.

Fig. 2 XRD patterns of the 1.68CaP_3xZrO₂ sample after heat treatment at different temperatures.

The diffraction patterns presented in Fig. 1 and 2 show only the presence of pure HAP and t-ZrO₂ phases at 900 °C, which resulted in the formation of HAP–ZrO₂ composite powders when the Ca/P molar ratio was maintained at 1.68 during the synthesis. The diffraction patterns of the composite powders had initiated the formation of β-Ca₃(PO₄)₂and m-ZrO₂ after heat treatment at 1000 °C and 1100 °C. Both the HAP phase and the t-ZrO₂ had shown excellent matches with the standard ICDD diffraction patterns of card nos 00-009-0432 for HAP and 01-079-1765 for t-ZrO₂. The increased level of ZrO₂ precursor added during the synthesis had resulted in a corresponding enhancement in the intensity of the diffraction patterns observed for t-ZrO₂ phase with the simultaneous reduction in the intensity of diffraction patterns observed for the HAP phase. A similar trend could be noticed in the XRD patterns presented in Fig. 1 and 2, with, however, the formation of an additional phase in the form of β-Ca₃(PO₄)₂ and m-ZrO₂ phases, which had confirmed the decomposition of the HAP phase at 1000 °C. A previous study had illustrated the comparative decomposition mechanism of mechanically mixed HAP–ZrO₂ composite after heat treatment in air and hot isostatic pressing. According to work reported by Adolfsson et al.,²⁶ the heat treatment in air at less than 950 °C had resulted in the decomposition of the HAP–ZrO₂ composite by the reactions given next:

Ca₁₀(PO₄)₆(OH)₂ → Ca₁₀(PO₄)₆(OH)_2−2xO_x + xH₂OCa₁₀(PO₄)₆(OH)_2−2xO_x + yZrO₂(tetragonal) → Ca₃(PO₄)₂ + CaO(ZrO₂(cubic))_y + (1 − x)H₂O (1)

However, the pressure-assisted sintering, hot isostatic pressing or spark plasma sintering of the HAP–ZrO₂ composite samples have prevented the decomposition of HAP–ZrO₂ composites until 1200 °C is reached.^23,36,37 It should be noted that in the research reported above, the yttria precursor has been used as a stabilizing agent for the preservation of the t-ZrO₂ phase. In the present investigation, without the usage of a yttria precursor, the heat treatment at 900 °C did not cause the decomposition of the HAP–ZrO₂ composite with the maintained precursor molar ratio of Ca/P being 1.68, whereas at 1000 °C and 1100 °C, the formation of the m-ZrO₂ and the β-Ca₃(PO₄)₂ phase has been detected in all the composite powders. The decomposition of HAP in the HAP–ZrO₂ composite system during heat treatment in air has been attributed to vacancy formation during the release of structural water.²⁶ In the present investigation, HAP–ZrO₂ composite powders have been formed by an in situ precipitation technique and during the heat treatment process, the gradual elimination of volatile impurities (NH₄⁺, Cl⁻, NO₃²⁻ and adsorbed H₂O) that generally arise during the aqueous precipitation technique has to occur, followed by the release of structural water. This elimination of volatile impurities has been evident from the XRD patterns (Fig. 1 and 2) through the conversion of observed broad diffraction patterns at lower temperatures (500 °C) to the sharp diffraction patterns at higher temperatures (900 °C). The elimination of structural water is initiated only after the removal of volatile impurities has been completed. According to the present results, the decomposition of HAP has started to occur by the formation of m-ZrO₂ and β-Ca₃(PO₄)₂ at both 1000 °C and 1100 °C. The XRD patterns for all the powder samples after heat treatment at 1200 °C had resulted in the formation of α-TCP thus signifying the further degradation of the composites. However, the XRD results have not given the exact quantitative information for the yield of HAP, ZrO₂ and β-Ca₃(PO₄)₂ phases in all the compositions investigated.

3.2 Structural analysis of the HAP–ZrO₂ composite powders

The results from the XRD analysis have confirmed the formation of HAP–ZrO₂ composites at 900 °C and heat treatment at 1000 °C has resulted in the formation of m-ZrO₂ and β-Ca₃(PO₄)₂ phases. However, the structural information of the phases formed could not be specified from the XRD analysis and thus, a quantitative analysis in the form of a Rietveld refinement has been employed for a precise analysis of the structural and compositional behaviour of the HAP–ZrO₂ composite powders. It has been already reported by Vasanthavel et al.³⁵ and Nandha Kumar et al.³⁸ that the Rietveld refinement technique can be employed for the quantitative analysis of the bioceramics. Fig. 3 presents the Rietveld analysis pattern of the powder diffraction data of the HAP and HAP–ZrO₂ composite powders after heat treatment. Table 2 presents the data of R_wp, R_p, χ², and R_Bragg obtained from the refinement of different HAP–ZrO₂ composites after heat treatment at three different temperatures (900 °C, 1000 °C and 1100 °C) and these values showcase the refinement standards executed in the present investigation. The powder samples heat treated at 1200 °C have not been subjected to a refinement study because of the allotropic phase conversion behaviour of β-TCP to α-TCP that generally occurs at 1180 °C which is followed by the reconversion of α-TCP to β-TCP during cooling to the room temperature. As can be seen from the Table 2, the goodness of fit (GoF) values were relatively low, i.e., all were approximately less than 2%, which is considered as acceptable according to basic principle of a GoF less than 4% and a R_wp of less than 20%.³⁹ The refined lattice parameters and unit cell volume of the stoichiometric HAP and the HAP phase present in all the HAP–ZrO₂ composite powders confirms the hexagonal setting that crystallizes in P6₃/m (176) space group at all the heat treatment temperatures of 900 °C, 1000 °C and 1100 °C (Table 3). However, the heat treatment at 1000 °C had resulted in the additional formation of a β-Ca₃(PO₄)₂ phase, which possesses a rhombohedral structure that tends to crystallize in the R3c (167) space group. In a similar manner, the refined lattice data of the ZrO₂ phase in all the HAP–ZrO₂ composite powders confirms the tetragonal setting that crystallizes in the P4₂/nmc (137) space group at the investigated temperatures of 900 °C, 1000 °C and 1100 °C. The formation of an additional ZrO₂ phase at 1000 °C has confirmed its monoclinic setting in the P21/c¹⁴ space group. Despite the different amounts of ZrOCl₂ precursor that were added during the synthesis, the a-axis and c-axis lattice parameter values of the HAP phase have shown a uniform trend in all the HAP–ZrO₂ composite powders at 900 °C, 1000 °C and 1100 °C, whereas the lattice parameter values of t-ZrO₂ have also indicated a uniform trend in all the HAP–ZrO₂ composite powders at 900 °C, 1000 °C and 1100 °C. The lattice data obtained for the HAP phase and the t-ZrO₂ phase over various temperatures indicates the fact that the individual phases are unique in the composite powders.

Fig. 3 Rietveld analysis pattern of powder diffraction data of the HAP and HAP–ZrO₂ composite powders after heat treatment.(a) Stoichiometric HAP at 900 °C, (b) 1.68CaP_5xZrO₂ at 1000 °C, (c) 1.68CaP_5xZrO₂ at 1100 °C *The solid lines are calculated intensities and the marked ones are the intensities observed. The difference between the observed and calculated intensities is plotted below the profile.

Table 2 Rietveld agreement factors obtained for the different HAP–ZrO₂ composite powders after heat treatment at different temperatures

Sample code	900 °C				1000 °C				1100 °C
	cRp^a	cRwp^a	χ^2^b	RBragg	cRp^a	cRwp^a	χ^2^b	RBragg	cRp^a	cRwp^a	χ^2^b	RBragg
a cRp and cRwp represent the Rietveld agreement factors.b χ^2 is a goodness of fit (GoF) indicator defined by (Rwp/Rexp)^2 in which Rexp is the expected weighted profile factor. RBragg demonstrates the agreement between the observed and computed reflection during refinement.
1.67CaP_1xZrO2	09.80	7.42	1.542	05.45	09.57	07.30	1.497	05.82	09.63	7.42	1.616	09.98
1.67CaP_2xZrO2	09.30	6.76	1.576	09.10	08.38	06.45	1.319	05.80	14.03	10.5	3.809	11.99
1.67CaP_3xZrO2	09.21	6.81	1.681	11.73	09.56	07.27	1.836	08.19	07.25	5.65	1.063	03.82
1.67CaP_4xZrO2	09.44	6.94	1.843	10.53	07.99	06.07	1.349	07.53	07.92	6.13	1.325	04.94
1.67CaP_5xZrO2	08.80	6.55	1.683	08.15	08.08	06.03	1.442	04.89	07.56	5.83	1.284	04.96
1.68CaP_1xZrO2	11.03	8.22	1.970	05.75	09.71	07.43	1.475	05.84	11.64	9.03	2.166	09.52
1.68CaP_2xZrO2	10.92	8.22	2.145	08.95	09.92	07.84	1.828	06.74	10.85	8.38	2.258	10.15
1.68CaP_3xZrO2	11.13	8.23	2.510	06.95	09.95	07.79	1.893	08.48	09.27	7.20	1.710	07.41
1.68CaP_4xZrO2	10.53	7.80	2.345	06.13	08.62	06.79	1.509	07.09	08.96	6.94	1.703	08.57
1.68CaP_5xZrO2	11.05	7.97	2.717	08.76	06.86	05.19	1.025	04.22	07.49	5.70	1.264	03.62
1.69CaP_1xZrO2	10.36	7.97	1.753	06.90	11.20	08.57	1.945	07.74	10.71	8.28	1.863	08.36
1.69CaP_2xZrO2	09.43	7.08	1.621	09.39	11.55	09.13	2.418	09.26	12.31	9.51	2.727	14.11
1.69CaP_3xZrO2	09.62	7.14	1.873	08.51	08.96	07.02	1.556	06.61	10.57	8.11	2.266	10.85
1.69CaP_4xZrO2	11.62	8.23	2.853	11.04	09.76	07.46	1.989	09.97	09.62	7.15	1.993	10.89
1.69CaP_5xZrO2	11.46	8.06	2.897	09.70	08.10	06.13	1.435	05.43	08.53	6.44	1.578	07.06

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Table 3 Structural parameters of HAP–ZrO₂ composite powders determined from Rietveld refinement analysis after heat treatment at three different temperatures

Sample Code	Structural parameters at 900 °C				Structural parameters at 1000 °C				Structural parameters at 1100 °C
	HAP		t-ZrO2		HAP		t-ZrO2		HAP		t-ZrO2
	a-axis (Å)	c-axis (Å)	a-axis (Å)	c-axis (Å)	a-axis (Å)	c-axis (Å)	a-axis (Å)	c-axis (Å)	a-axis (Å)	c-axis (Å)	a-axis (Å)	c-axis (Å)
1.67CaP_1xZrO2	9.5517(6)	6.8342(4)	3.5978(3)	5.1782(7)	9.5333(6)	6.8422(4)	3.5994(3)	5.1794(6)	9.4942(4)	6.8517(3)	3.5964(4)	5.1792(4)
1.67CaP_2xZrO2	9.5603(7)	6.8253(5)	3.5984(2)	5.1789(5)	9.5512(6)	6.8336(4)	3.5983(2)	5.1823(4)	9.4843(6)	6.8625(5)	3.5971(2)	5.1816(5)
1.67CaP_3xZrO2	9.5556(8)	6.8289(6)	3.5994(3)	5.1787(6)	9.5441(7)	6.8353(5)	3.5997(2)	5.1797(5)	9.4842(4)	6.8568(3)	3.5957(1)	5.1792(2)
1.67CaP_4xZrO2	9.5733(7)	6.8204(6)	3.5984(2)	5.1814(4)	9.5539(6)	6.8296(4)	3.5989(2)	5.1815(4)	9.4867(4)	6.8586(4)	3.5961(1)	5.1807(2)
1.67CaP_5xZrO2	9.5792(8)	6.8128(7)	3.5977(2)	5.1801(4)	9.5337(6)	6.8363(5)	3.5972(2)	5.1805(5)	9.4862(5)	6.8594(4)	3.5958(1)	5.1821(2)
1.68CaP_1xZrO2	9.5509(6)	6.8359(4)	3.5996(3)	5.1786(8)	9.4995(4)	6.8560(3)	3.5998(2)	5.1784(5)	9.4902(4)	6.8549(3)	3.5977(2)	5.1784(4)
1.68CaP_2xZrO2	9.5663(8)	6.8307(6)	3.6029(3)	5.1776(7)	9.4829(7)	6.8482(5)	3.5959(3)	5.1719(5)	9.4766(6)	6.8491(5)	3.5942(2)	5.1703(4)
1.68CaP_3xZrO2	9.5733(8)	6.8203(6)	3.5984(3)	5.1808(6)	9.4895(6)	6.8545(5)	3.5979(2)	5.1775(4)	9.4841(4)	6.8563(4)	3.5960(1)	5.1783(3)
1.68CaP_4xZrO2	9.5624(9)	6.8275(7)	3.6002(3)	5.1787(6)	9.4879(5)	6.8554(4)	3.5975(2)	5.1775(3)	9.4795(5)	6.8578(4)	3.5971(1)	5.1769(3)
1.68CaP_5xZrO2	9.5548(1)	6.8155(9)	3.5946(3)	5.1748(7)	9.4764(5)	6.8588(5)	3.5958(2)	5.1815(2)	9.4865(5)	6.8560(5)	3.5954(1)	5.1814(3)
1.69CaP_1xZrO2	9.5708(7)	6.8256(5)	3.6041(6)	5.1753(9)	9.5244(4)	6.8402(3)	3.6002(2)	5.1750(6)	9.4957(3)	6.8580(2)	3.5967(1)	5.1840(4)
1.69CaP_2xZrO2	9.5650(7)	6.8263(5)	3.6006(3)	5.1790(6)	9.5114(6)	6.8499(4)	3.6015(2)	5.1771(5)	9.4955(5)	6.8546(4)	3.5990(2)	5.1783(4)
1.69CaP_3xZrO2	9.5762(9)	6.8203(7)	3.6001(3)	5.1808(7)	9.5138(5)	6.8466(4)	3.5970(2)	5.1808(3)	9.4929(5)	6.8548(4)	3.5970(1)	5.1793(3)
1.69CaP_4xZrO2	9.5671(9)	6.8257(9)	3.6008(4)	5.1816(7)	9.5001(6)	6.8528(5)	3.5982(2)	5.1808(4)	9.4910(5)	6.8555(4)	3.5967(1)	5.1785(3)
1.69CaP_5xZrO2	9.5773(8)	6.8175(9)	3.6005(4)	5.1800(8)	9.5150(5)	6.8428(5)	3.5966(2)	5.1818(3)	9.4953(5)	6.8510(4)	3.5955(1)	5.1812(3)

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3.3 Quantitative phase content

Table 4 represents the individual phase compositions of the powder samples obtained from Rietveld refinement at three different heat treatment temperatures, which have been synthesised by maintaining different Ca/P ratios and varying the ZrOCl₂ precursor concentrations. From Table 4, it is observed that HAP and t-ZrO₂ have occurred as the major phases, as expected, in all 15 samples after heat treatment at three different temperatures. At the heat treatment temperature of 900 °C, the phase content of t-ZrO₂ has been found to increase in line with the corresponding increase in the amount of added ZrOCl₂ precursor. In other words, the phase content of HAP has been found in decreasing order with the simultaneous increase in the content of t-ZrO₂. An interesting feature is that the increase in the Ca/P molar ratio from 1.67 to 1.69 has also resulted in the gradual enhancement of the HAP phase formed in the composite powders. The synthesis performed using a Ca/P molar ratio of 1.67 and 1.69 has resulted in the formation of a minor amount of additional phase in the form of β-TCP other than HAP and t-ZrO₂at 900 °C. All the previously reported work on the formation of HAP–ZrO₂ composite powders has used yttria as a precursor to stabilize the t-ZrO₂ phase.^40–43 The present study has seen the stabilization of t-ZrO₂ without the addition of any stabilizer during the synthesis. It has been reported by many authors that other than yttria, there are many additives in the form of CaO, MgO, La₂O₃ and rare earth elements that could be used to stabilize the t-ZrO₂.^44–47 It is also a well-known fact that depending on the extent of the calcium deficiency, the formation of β-TCP along with the HAP phase could not be avoided during heat treatment.^4–6 Thus, it could be stated from the present investigation that the excess Ca²⁺ ions present in the calcium phosphate suspension has played an active role in the stabilization of t-ZrO₂, thus creating a calcium deficiency and thereby leading to the formation of β-TCP during heat treatment. It should be noted that the weight fractions of all the β-TCP formed were found to be less than 5 wt%. In the case of the Ca/P precursors which had a maintained molar ratio of 1.68 during the synthesis, no β-TCP phase has been detected at the heat treatment temperature of 900 °C. Much research has been reported on the effect of the Ca/P molar ratio on the thermal stability and degradation behaviour of stoichiometric HAP. Osaka et al.⁴⁸ states that the HAP synthesized by maintaining a Ca/P molar ratio of 1.68 has been found to be thermally stable until 1300 °C without phase degradation. Some of the reports also state that synthesis using a Ca/P molar ratio of >1.70 results in the non-stoichiometry of the HAP formed, which then gets thermally degraded into β-TCP and CaO beyond 1300 °C.⁴⁹ Generally, the Ca/P molar ratio required to synthesize stoichiometric HAP is considered to be 1.67 and the resulting HAP starts to degrade beyond 1200 °C. Thus a minor alteration in the Ca/P molar ratio during the synthesis of stoichiometric HAP has resulted in the varied level of its thermal stability. However, the addition of ZrO₂ precursor during the synthesis have resulted in the earlier degradation of the HAP phase, thus yielding a minor amount of β-TCP phase (less than 5 wt%) at 900 °C for the Ca/P molar ratios of 1.67 and 1.69. The addition of a varied range of ZrO₂ precursors has not affected the HAP–ZrO₂ composite powders at 900 °C for the synthesis performed with a Ca/P molar ratio of 1.68.

Table 4 Quantitative phase compositions of HAP–ZrO₂ composite powders determined from Rietveld refinement analysis after heat treatment at three different temperatures

Sample Code	Phase composition in weight fractions at 900 °C			Phase composition in weight fractions at 1000 °C				Phase composition in weight fractions at 1100 °C
	HAP	t-ZrO2	β-TCP	HAP	t-ZrO2	β-TCP	m-ZrO2	HAP	t-ZrO2	β-TCP	m-ZrO2
1.67CaP_1xZrO2	73.00	22.87	04.13	62.77	21.77	13.66	01.80	60.25	18.68	19.36	01.71
1.67CaP_2xZrO2	58.17	36.91	04.92	48.84	34.70	14.82	01.64	52.71	21.10	20.19	06.00
1.67CaP_3xZrO2	53.47	44.39	02.14	44.81	40.53	08.26	06.40	38.96	25.65	16.29	19.10
1.67CaP_4xZrO2	50.29	48.18	01.53	37.19	45.61	08.70	08.50	34.97	26.96	12.11	25.96
1.67CaP_5xZrO2	44.14	54.60	01.26	32.90	51.20	06.90	09.00	26.12	29.53	13.25	31.10
1.68CaP_1xZrO2	75.81	24.19	—	64.52	20.13	12.53	02.82	63.66	19.55	15.86	00.93
1.68CaP_2xZrO2	62.68	37.32	—	50.84	35.87	10.56	02.73	53.40	25.22	15.73	05.65
1.68CaP_3xZrO2	58.99	41.01	—	46.00	39.66	07.94	06.40	39.00	31.26	14.98	14.76
1.68CaP_4xZrO2	52.72	47.28	—	38.15	42.55	09.13	10.17	35.49	33.59	13.81	17.11
1.68CaP_5xZrO2	48.64	51.36	—	34.75	46.00	08.25	11.00	28.99	24.01	15.37	31.63
1.69CaP_1xZrO2	77.32	20.40	02.28	67.46	20.85	09.49	02.20	66.51	16.69	11.96	04.84
1.69CaP_2xZrO2	64.20	33.72	02.08	56.46	33.34	07.74	02.46	55.28	26.85	11.23	06.64
1.69CaP_3xZrO2	59.64	38.82	01.54	49.12	40.18	05.23	05.47	41.93	32.77	11.58	13.72
1.69CaP_4xZrO2	53.35	44.80	01.85	39.18	41.67	07.78	11.37	36.80	35.05	10.63	17.52
1.69CaP_5xZrO2	50.52	46.81	02.67	36.16	43.84	08.00	12.00	29.13	36.38	10.66	23.83

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Heat treatment at 1000 °C and 1100 °C has revealed a significant change in the phase composition of all the powder samples. Despite the varied level of the Ca/P molar ratio maintained during the synthesis, the HAP–ZrO₂ composite powders have shown phase degradation at both the temperatures of 1000 °C and 1100 °C by the formation of a considerable amount of m-ZrO₂ (Table 4). The phase content of the β-TCP formed has been found to be comparably higher at 1000 °C and 1100 °C when compared to their content at 900 °C. The phase content of the newly formed m-ZrO₂ phase at 1100 °C has been seen at higher levels when compared to its content at 1000 °C, thus signifying that there is a phase degradation of HAP–ZrO₂ composite powders at 1000 °C. Furthermore, with the increased addition of ZrO₂ precursor during synthesis, the phase content of m-ZrO₂ has been found to increase. It is important to mention here that in any of the compositions investigated at three different heat treatment temperatures, neither c-ZrO₂ nor the formation of CaO has been detected in trace levels. The results obtained from the present investigation are a significant contradiction to the reported results shown in eqn (1) above, which shows that the degradation of HAP–ZrO₂ composites is initiated by the formation of the m-ZrO₂ and CaO phases after the heat treatment temperature of 950 °C. Reports in the literature have proposed a different kind of mechanism for the formation of m-ZrO₂ and c-ZrO₂ during the degradation of HAP–ZrO₂ composites. The formation of c-ZrO₂ has been attributed to the decomposition of HAP to yield CaO, which in turn is consumed by the t-ZrO₂ to be converted into c-ZrO₂.^50,51 The formation of m-ZrO₂ has been attributed to the presence of an insufficient amount of CaO from the degradation of HAP to stabilize the c-ZrO₂.⁵² The comparison of the phase content of the different compositions in Table 4 indicates that the enhanced content of m-ZrO₂ at 1100 °C causes the rapid degradation of the composite powders when compared with the degradation behaviour at 1000 °C. A greater amount of the phase content of m-ZrO₂ has been formed with the corresponding increase in the ZrO₂ content in the samples. Thus, the results from the present investigation infers that the ZrO₂ content has an active role in the degradation behaviour of HAP–ZrO₂ composite powders.

3.4 Raman spectroscopy

Fig. 4–6 represent the Raman spectra of the HAP–ZrO₂ composite powders after heat treatment at different temperatures, together with the spectra of pure HAP. From Fig. 4–6, it is observed that the bands representing the PO₄ groups of the HAP phase^53,54 have also been seen in pure HAP and also for all the HAP–ZrO₂ composite powders. The bands observed around the region of 426 cm⁻¹ and 458 cm⁻¹ correspond to the ν₂ bending vibrations of the PO₄³⁻ ion. The bands observed in the region of 578, 589 and 610 cm⁻¹ correspond with the ν₄ fundamental vibrational modes that arise from the triply degenerate bending vibrations. The high intensity band seen at 958 cm⁻¹ has arisen from the symmetric stretching of the ν₁ fundamental vibrational modes. The bands located at 1033, 1049 and 1073 cm⁻¹ have been attributed to the asymmetric stretching vibrations of the ν₃ fundamental vibrational modes of the PO₄ bonds. The Raman spectra for the powders after heat treatment at 900 °C (Fig. 4) confirm the bands corresponding to the t-ZrO₂ phase.

Fig. 4 Raman spectra for the HAP–ZrO₂ composite powders heat treated at 900 °C with the synthesis done at a Ca/P molar ratio of 1.68.

The bands observed in the region around 266, 146 and 312 cm⁻¹ correspond with the presence of t-ZrO₂.^55,56 At the heat treatment temperature of 900 °C, the intensity of the characteristic bands of the HAP phase has decreased with the corresponding increase in the intensity of bands of the t-ZrO₂ phase.

For Raman spectra obtained for the powder samples after heat treatment at 1000 °C and 1100 °C (Fig. 5 and 6), the bands corresponding to the HAP and t-ZrO₂ phases have been identified in the spectrum. Other than the bands of the HAP and t-ZrO₂ phases, the bands corresponding to the presence of m-ZrO₂ phase have been confirmed from the bands witnessed in the spectral region around 178, 189, 333, 347 and 380 cm⁻¹.^57,58 Since the overall quantitative yield of m-ZrO₂ has been found to a lesser extent at 1000 °C when compared with its content at 1100 °C, the bands of m-ZrO₂ at 1000 °C have been seen with less intensity towards 1100 °C. Additionally, the bands of m-ZrO₂ have been observed with increasing intensity, with their increasing content at 1100 °C. Thus, the results obtained from the Raman spectroscopy have been found to be in agreement with the quantitative analysis from the Rietveld refinement.

Fig. 5 Raman spectra for the HAP–ZrO₂ composite powders heat treated at 1000 °C with the synthesis done at a Ca/P molar ratio of 1.68.

Fig. 6 Raman spectra for the HAP–ZrO₂ composite powders heat treated at 1100 °C with the synthesis done at a Ca/P molar ratio of 1.68.

Summary

A wide range of HAP–ZrO₂ composite powders have been obtained in the present investigation by adapting the in situ aqueous precipitation technique. Minor alterations in the Ca/P molar ratio together with five different precursor concentrations of ZrO₂ have been varied during the synthesis. Heat treatment of the resulting powders has induced the formation of an HAP–ZrO₂ composite and the resulting phase content has been found to vary depending on the molar ratio of the Ca/P precursors and the concentration of the ZrOCl₂ precursor. A minor amount of β-TCP (less than 5 wt%) together with the HAP and t-ZrO₂ has been detected for the Ca/P precursors with molar ratios of 1.67 and 1.69. For the precursor with the Ca/P molar ratio of 1.68, however, the formation of HAP and t-ZrO₂ was only confirmed after heat treatment at 900 °C. Heat treatment at 1000 °C and 1100 °C has confirmed the phase degradation of the HAP–ZrO₂ composite powders by virtue of the formation of significant amounts of m-ZrO₂. The results from the present study show that the stabilization of the t-ZrO₂ phase had occurred without the addition of any stabilizing agent in the form of yttria, which has been a usual practice during the formation of HAP–ZrO₂ composites. In the absence of any stabilizing agent, the presence of Ca²⁺ ions in the calcium phosphate suspension had played a significant role in the t-ZrO₂ phase stabilization during heat treatment. The increased addition of ZrOCl₂ precursor had resulted in the corresponding enhancement in the phase content of ZrO₂ in the composites with the simultaneous reduction in the HAP content. The presence of a higher content of ZrO₂ in the composites had led to the enhanced formation of m-ZrO₂, thus signifying the rapid degradation of the composite powders.

Acknowledgements

The authors gratefully acknowledge the financial support received from the Council of Scientific and Industrial Research, India (CSIR Scheme no. 22(612)/12/EMR-II). The authors are also grateful for the use of the instrumentation at the Central Instrumentation Facility (CIF) of Pondicherry University.

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