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phosphate glass composites
Biomaterials 19 (1998) 1735 — 1743
A quantitative study of the sintering and mechanical properties of hydroxyapatite/phosphate glass composites D.C. Tancred*, B.A.O. McCormack, A.J. Carr Bioengineering Research Centre, Department of Mechanical Engineering, University College Dublin, Dublin, Ireland Received 11 August 1997; accepted 8 March 1998
Abstract Previous work has shown that small additions of a phosphate glass (CaO—P O ) can significantly enhance the sinterability and 2 5 strength of hydroxyapatite. However, there are no quantitative phase analyses available for these materials which would provide indicators of biocompatibility and resorbability. Similarly, there is little information available about the mechanical properties, especially with high glass additions. In this study, the effects of sintering hydroxyapatite with phosphate glass additions of 2.5, 5, 10, 25, and 50 wt.% are quantified. Each composition was sintered over a range of temperatures, and quantitative phase analysis was carried out using XRD. In addition, the microstructures were studied using RLOM and SEM, and mechanical properties (Vickers hardness, K , and MOR) measured. These results may be used to indicate which compositions and processing conditions may IC provide materials suitable for use in hard tissue replacement. Composites containing up to 10 wt.% glass additions formed dense HA/TCP composite materials possessing flexural strength and fracture toughness values up to 200% those of pure HA. The HA/TCP ratio was strongly dependent on the percentage glass addition. Higher glass additions resulted in composites containing b-TCP together with large amounts of a- or b-calcium pyrophosphate, and having similar mechanical strengths to pure HA. ( 1998 Published by Elsevier Science Ltd. All rights reserved Keywords: Hydroxyapatite; Glass; Calcium phosphate; Composite; Bioceramic
1. Introduction Calcium phosphate ceramics, especially hydroxyapatite (HA) and the more resorbable b-tricalcium phosphate (b-TCP), are widely used for hard tissue replacement due to their biocompatibility and osteoconductive properties [1, 2]. HA/b-TCP composites (frequently termed biphasic calcium phosphates; BCPs) have been developed to allow variation in implant resorption rates whilst retaining useful bioactive properties. Other calcium phosphates such as a-TCP, tetracalcium phosphate, and b-calcium pyrophosphate (b-CPP), though bioactive, have proven less useful as bone replacement materials due to excessively high resorption rates. In general, applications of calcium phosphates in the body have been limited by low strength and low fracture toughness. Numerous techniques have been investigated in attempts to improve the mechanical properties for particular applications or implant configurations, through advanced powder processing techniques [3, 4] *Corresponding author. Tel.: #353 1 706 1752; fax: #353 1 706 1756.
or by formation of HA-containing composite materials. Such composites aim to retain their useful bioactive properties whilst providing more suitable mechanical properties for particular applications. These include fibre reinforcement of HA [5, 6], HA/polyethylene [7], and HA/Al O [8] composites. 2 3 In recent years increasing interest has been shown in sintering of HA with glass additions. There are two main motivations for sintering HA with a glassy phase—to enhance the densification and therefore the mechanical strength by acting as a sintering aid, and to enhance bioactivity through the combination of two bioactive phases [9, 10]. Sintering with low glass additions may promote densification through liquid-phase sintering, resulting in composite materials with enhanced mechanical properties. For example, small additions of phosphate glasses have been shown to significantly enhance composite flexural strength (by up to 400%), and fracture toughness (by up to 200%) [11—13]. However, sintering in the presence of a glass was found to cause degradation of HA due to reaction with the glass to yield either b-TCP or a-TCP. Glass composition and the level of
0142-9612/98/$ — See front matter ( 1998 Published by Elsevier Science Ltd. All rights reserved. PII S 0 1 4 2 - 9 6 1 2 ( 9 8 ) 0 0 0 8 2 - 9
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glass additions have a large effect on both the phase assembly and the mechanical properties of the resulting composites. In particular, the presence of sodium in a glass was found to promote greater degradation of HA [9, 11]. Mechanical properties of HA composites containing Na O, though better than for pure HA, were 2 poorer than for soda-free glass composites [11, 12]. Use of a phosphate glass (in the system CaO—P O ) 2 5 ensures that the sintered body contains only calcium phosphate phases, which are likely to be biocompatible and possibly bioactive. Previous work [11] showed that glasses containing a higher Ca/P ratio caused less degradation of HA during sintering. The Ca/P ratio is practically limited to about 50 : 50 mol% CaO : P O , due to 2 5 a sharp increase in the liquidus temperature at higher Ca/P ratios [14]. In this work, the effects of phosphate glass additions, from the system CaO—P O , on the sintering, phase 2 5 composition, and mechanical properties of HA are analysed. The objective is to provide a more comprehensive understanding of the mechanical and possible biochemical properties of such composites, which should help to determine the optimum composites for use as implant materials.
2. Materials and methods 2.1. Preparation of biomaterials Synthetic hydroxyapatite was prepared using a wet method based on addition of calcium hydroxide to orthophosphoric acid at 85°C according to the reaction [15]: 6H PO #10Ca(OH) PCa (PO ) (OH) #18H O. 3 4 2 10 46 2 2 Phosphate glass was prepared using conventional glassforming techniques. Reagent-grade CaCO and P O 3 2 5 were melted in a fine-grained mullite crucible (Zedmark Refractories, Dewsbury, UK) at 1070°C for 2 h, to produce a glass of 50 : 50 CaO : P O molar ratio. The result2 5 ing melt was quenched in water, and the glass frit collected and quickly dried in a conventional oven at 100°C to minimise corrosion due to contact with water. HA and glass frit powders were separately dry ballmilled for 1 h in an alumina pot and passed through a 45 lm sieve. HA and composite powders containing 2.5, 5, 10, 25 and 50 wt% glass were then prepared and wet ball-milled in 60 g batches in 500 ml polypropylene NalgeneTM pots together with 180 ml isopropyl alcohol (IPA) and approximately 650 g of 7 mm diameter cylpeb magnesia-stabilised zirconia milling media. Powders were wet milled on ball-mill rollers at 140 rpm for 24 h. Recovered slips were passed through a 38 lm sieve, dried at 70°C and then passed through a 106 lm sieve to break any agglomerates formed on drying.
The powders were mixed with 4 wt% organic binder (Bindemittel B11/V, Degussa, Germany) to improve green strength and uniaxially pressed at 120 MPa in a 25 mm diameter steel die using 2 g of powder per disk. Following curing of the binder, samples were heated to 350°C at 2°C/min to remove the binder. Sintering was carried out in a muffle furnace, in air, by heating at 4°C/min to sintering temperature which was held for 3 h, followed by cooling in the furnace at 4°C/min. All samples were surface ground on one side (R "1—1.5 lm) for ! mechanical testing. 2.2. Compositional analysis Phase analysis was performed by X-ray diffraction using a Huber 642 Guinier diffractometer, fitted with a quartz Johansson monochromator, operating in subtractive transmission mode at 50 kV and 40 mA. Radiation was pure monochromatic Cu Ka (j"1.54056 As ). 1 Samples were ground to a fine powder using an agate mortar and pestle and mixed with Si powder as an internal standard to allow peak correction, and for quantitative analysis. Samples were scanned from h"0—50° in increments of 0.02°, using a count time of 5 s, and a counter slit width of 0.2 mm. Peak positions were corrected using a calibration curve obtained from the diffraction peaks of the internal standard. Lattice parameters were calculated (with 95% confidence intervals) for HA and b-TCP phases of the sintered composites. Quantitative phase analysis was performed by preparing XRD calibration curves relating the integrated intensities of the strongest peak of individual component phases to that of the internal standard. Estimation of the percentage a-TCP present was performed using the I/I #03 values (taken from the JCPDS PDF or experimentally determined using standard mixtures where JCPDS data were unavailable) for a-TCP and b-TCP. Due to the high resolution of the Guinier diffractometer, as well as full crystallisation of the glass phase for all additions, difficulties in quantitative phase analysis experienced by other researchers [12] due to peak overlapping were absent. The quantitative analysis ignored any effects of deviations from stoichiometry of individual phases. Lattice parameters for the composite series showed appreciable deviations did occur for both HA and b-TCP. 2.3. Mechanical testing Flexural strength testing was performed using a disk bend test method [16]. The test applies an equi-biaxial loading to thin disc specimens by loading with concentric rings of steel ball bearings. Testing was performed using a Lloyd 6000S mechanical testing machine, fitted with a 500 N load cell, using a cross-head speed of
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0.5 mm min~1. Samples were tested with the ground surface in tension. Vickers hardness was measured using a Mitutoyo AVK-C2 hardness tester. Fracture toughness was determined using a Vickers indentation technique [17]. A minimum of six samples were used for each mechanical test. Ninty five per cent confidence intervals were calculated for each series. 2.4. Microstructures Microstructures for all sintered materials were examined using a JEOL 35C SEM operating at 25 kV. Polished surfaces were etched using 0.1 M lactic acid. Hydroxyapatite samples were etched to reveal the grain structure using an etching time of 10 min. Phosphate glass composite series were etched for 10 s to reveal the location of individual phases. Although longer etching times showed greater grain boundary detail, it became difficult to distinguish between individual phases and to determine the extent and type of porosity.
3. Results 3.1. Phase analysis Analysis by XRD showed the HA synthesised in this study was stable up to 1350°C in air. Lattice parameters for HA, calcined at 900°C for 15 h, were a" 9.4198$0.0007 As and c"6.8821$0.0008 As . No detectable second phase was observed at any sintering temperature. Sintered composites containing low phosphate glass additions (up to 10 wt%) consisted of HA/b-TCP or HA/b-TCP/a-TCP, depending on the percentage glass addition and sintering temperature. For 2.5 and 5 wt% glass additions, the HA and TCP contents were largely independent of the sintering temperature with the percentage HA decreasing only slightly over the sintering temperature range. At 1350°C, a small amount of b-TCP converted into a-TCP. There was no significant decrease in HA content at this sintering temperature. Glass additions of 10 wt% resulted in greater degradation of HA. The quantity of residual HA was lower (about 10—40 wt%) across the temperature range. Above 1250°C, there was an increase in the degradation of HA accompanied by both an increase in the b-TCP phase content and the formation of a-TCP (Fig. 1). The lattice parameters of HA and b-TCP changed significantly as the percentage glass was increased, indicating a change in stoichiometry due to either lattice vacancies or substitutions. With an increasing percentage glass addition, the quantity of the residual HA content decreased. In addition, the composition moved farther from stoichiometry as was evident from a decrease in
Fig. 1. Phase composition of HA/phosphate glass (PG) composites.
Table 1 Lattice parameters for HA/phosphate glass composites (95% confidence limit of final digit in parentheses) Composition
Temperature (°C)
a (A_ )
c (A_ )
HA
1250 1300 1350
9.4154 (4) 9.4126 (7) 9.4134 (8)
6.8865 (5) 6.8832 (8) 6.8839 (9)
2.5%
1250 1300 1350
9.4118 (7) 9.4081 (8) 9.4113 (8)
6.8877 (9) 6.8882 (9) 6.888 (1)
5.0%
1250 1300 1350
9.411 (1) 9.408 (1) 9.411 (1)
6.887 (1) 6.888 (2) 6.888 (2)
10.0%
1250 1300 1350
9.404 (5) 9.403 (3) 9.409 (3)
6.893 (6) 6.895 (4) 6.891 (4)
a and an increase in c, away from stoichiometric values (Table 1 and Fig. 2). Lattice parameters for b-TCP moved towards stoichiometry with increasing percentage glass incorporation (and with increasing percentage bTCP in the composite). For 25 wt% additions, sintering at 1100°C resulted in approximately a 1 : 1 ratio of HA : b-TCP. Further increase in temperature to 1150°C caused complete degradation of HA to form b-CPP. Higher temperatures resulted in the transformation of all b-CPP to a-CPP. Glass additions of 50 wt% resulted in elimination of all HA to form b-TCP/b-CPP in a ratio of about 1 : 25. The phase composition was constant up to a sintering temperature of 1150°C. Sintering at 1200°C was found to cause transformation of all b-CPP to a-CPP. 3.2. Densification and microstructures Densification of HA was almost complete at 1200°C, reaching a maximum of 97.8% theoretical density at
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Fig. 2. Variation in lattice parameters of HA for HA/phosphate glass (PG) composites.
1300°C. The percentage densification of each composite sample was adjusted to allow for the difference in density of component phases (Fig. 3). Composites based on 2.5 and 5 wt% glass additions attained maximum densification above 1250°C, similar to that of HA. Slightly lower densification was achieved for composites containing 10 wt% glass additions, whose density decreased rapidly above 1250°C, due to increased reaction of HA, the formation of significant quantities of a-TCP and the generation of large pores (Fig. 5d). Samples based on 25 wt% glass additions showed most densification occurred by 1150°C. At 1200°C, there was a slight decrease in the measured density due to transformation of b-CPP (o"3.129 g cm~3) to a-CPP (o"2.930 g cm~3), however densification continued up to 1250°C. Composites containing 50 wt% glass showed a monotonic increase in density with temperature. Pure HA showed rapid grain growth above 1200°C. Average grain size increased from 1.0$0.05 lm at 1200°C to 2.47$0.13 lm at 1350°C. Residual porosity above 1200°C consisted of submicron pores existing at grain boundaries (Fig. 4). Glass additions of up to 10 wt% resulted in an increasing degree of porosity with increasing percentage glass addition when sintered at 1200°C, probably due to reaction between the glass and HA. Between 1200 and 1250°C, there was a substantial decrease in porosity for each composite series as the materials reached almost maximum density. Residual porosity principally existed at phase boundaries. This
Fig. 3. Densification of HA/phosphate glass composites (% represents percentage of theoretical density).
Fig. 4. Microstructures of sintered HA (etched for 10 min, 0.1 M lactic acid): (a) 1250°C, (b) 1350°C.
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Fig. 5. Microstructures of HA/phosphate glass (PG) composites (etched for 10 s, 0.1 M lactic acid). (a) 5 wt% PG,1250°C; (b) 5 wt% PG, 1350°C; (c) 10 wt% PG, 1250°C; (d) 10 wt% PG, 1350°C. Note: A"HA, B"TCP, P"Pore.
was consistent with previous reports of pore formation associated with OH~ release from HA during sintering [9, 18]. Pores were typically less than 1 lm except where there was a further decrease in the amount of HA present due to continued reaction with the glass (which occurred for 10% additions, sintered above 1250°C). It was not possible to distinguish between a-TCP and b-TCP regions in microstructures. TCP did not exist along grain boundaries as might have been expected in a liquid-phase sintered composite, but instead as isolated areas dispersed throughout the matrix. As the percentage glass addition (and the percentage TCP) was increased, TCP areas grew and coalesced to form a continuous TCP matrix with interspersed HA areas. Glass additions of 25 wt% caused a phase change between 1100°C and 1150°C from about 50HA : 50bTCP to 50b-TCP : 50b-CPP. A further increase in temperature to 1200°C caused transformation of b-CPP to a-CPP. Although these phase changes did not prevent densification, the combined effects of densification and phase changes had a large influence on the strength of the composites. Composites based on 50 wt% additions showed little change in microstructure or composition over the sintering temperature range investigated. Resid-
ual porosity was associated primarily with the presence of b-TCP. 3.3. Mechanical testing 3.3.1. Flexural strengths Comparison of composite flexural strength data with that of HA (Table 2), showed maximum average strength was achieved for 5 wt% glass additions sintered at 1350°C (61.9$7.5 MPa) and was about twice that of HA (31.4$3.1 MPa) which was obtained for samples sintered at 1200°C. Maximum strengths for 10 wt% additions were about 160% of HA values and occurred at a sintering temperature of 1300°C. There was no significant improvement in maximum strength for 2.5 wt% additions, however, unlike HA there was no significant decrease in average strength when sintered beyond the optimum sintering temperature. The optimum strength for each series was achieved at higher temperatures than for HA alone. All composites based on low glass additions ()10 wt%) showed a significant decrease in residual porosity between 1200°C and 1250°C. This corresponded to a significant increase in strength over the same sintering temperature range (Table 2).
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Table 2 Flexural strength in MPa of HA and HA/phosphate glass composites (95% confidence limit in parentheses, n"6) Sintering temperature (°C) 1000 1050 1100 1150 1200 1250 1300 1350
HA
20.5 31.4 17.9 20.9
2.5% PG
(3.4) (3.1) (1.9) (1.9)
15.6 32.3 27.4 25.2
(2.8) (4.0) (4.4) (6.9)
Incorporation of 25 wt% glass resulted in similar strength values to HA at 1150°C. Between 1100 and 1150°C, there was a large decrease in porosity. This was accompanied by a reaction in which all HA was eliminated and was converted largely to b-TCP. The phase change does not seem to have disrupted sintering as there was a significant improvement in strength and densification by 1150°C. The transformation of such a large amount of HA over a narrow temperature range could have been expected to interrupt sintering due to volumetric change, as well as OH~ release. However, transformation may have occurred while there was still a substantial amount of open porosity which prevented entrapment of released water vapour. Sintering at between 1150 and 1200°C transformed all b-CPP to a-CPP. There was a significant decrease in strength over the same temperature range. Microstructural examination showed that the a-CPP component consisted of large grains which increased in size rapidly above 1200°C. The combined effect of volumetric change and coarse grain size may have been sufficient to weaken the material. There was no statistically significant change (at a 95% confidence level) in average flexural strengths with increasing sintering temperature for composites containing 50 wt% glass additions. Nor was there a significant difference in strength between the maximum strength of pure HA and these composites, for any sintering temperature.
3.3.2. Hardness and fracture toughness The fracture toughness of composites based on glass additions of up to 10 wt% increased with increasing percentage glass addition (Fig. 6). Although the increase was not always significant for adjacent compositions, the general trend applied consistently across the whole series. There was a corresponding decrease in hardness with increasing glass additions (Fig. 7). No K value was IC obtained for samples containing 10 wt% glass, sintered at 1300°C, due to extensive lateral vent crack formation on indentation.
5% PG
14.2 21.4 40.0 42.2 61.9
(3.0) (3.9) (3.5) (3.9) (7.5)
10% PG
19.6 21.8 45.9 50.5 19.9
(4.2) (4.9) (4.4) (2.6) (4.7)
25% PG
23.7 37.6 11.3 14.0
(5.0) (2.0) (1.0) (0.8)
50% PG 21.8 29.1 28.3 28.0
(7.9) (6.8) (5.0) (11.8)
Fig. 6. Fracture toughness of HA and HA/phosphate glass (PG) composite materials.
4. Discussion Composites containing phosphate glass additions showed two types of behaviour. Those containing up to 10 wt% glass additions formed dense, strong, HA/TCP composites. Higher additions led to the formation of more porous composites containing large amounts of b-CPP or a-CPP. For low glass additions, the HA/TCP ratio depended strongly on the percentage glass added. Thus, it is likely that implant resorption, which is strongly dependent on the HA/TCP ratio, may be controlled by altering the percentage glass addition. The osteoconductivity of HA, of b-TCP, and of HA/b-TCP composites (biphasic calcium phosphates), has been clearly demonstrated by numerous animal and clinical trials [19]. In addition, BCPs formed by sintering of HA with glass (of
D.C. Tancred et al. / Biomaterials 19 (1998) 1735—1743
Fig. 7. Vickers hardness of HA and HA/phosphate glass (PG) composite materials.
the system CaO—P O —Na O has been demonstrated as 2 5 2 bioactive and osteoconductive [20]. Previous work has found that BCP implants may achieve a better tissue response and stronger bone ingrowth, in macroporous form than either HA or b-TCP alone [21]. Although it may not be possible to define an optimal HA/TCP ratio applicable to every situation, due to variations in implant characteristics as well as surgical site and application, there is evidence from previous work [21] that compositions of about 60 : 40 to 80 : 20 may be advantageous. However, further research into the effects of variations of the HA/b-TCP ratio is required to determine what effect variations in glass additions may have on surface activity or bone ingrowth. Small amounts of a-TCP formed in composites sintered at the highest sintering temperatures. Despite its much higher resorption rate the quantity of a-TCP is likely to be sufficiently low not to significantly affect the overall degradation rate. However, it is possible that if the a-TCP formed occurred primarily at the grain boundaries, the degradation may be significantly increased due to preferential dissolution of the grain boundaries. The microstructural location of a-TCP was not determined but is assumed to be associated with b-TCP regions (Fig. 5). The maximal bending strength of HA (31.4$3.1 MPa at 1200°C) is lower than some values available in the literature. The disk bend test method employed in this work provides a more conservative estimate of strength due to the equibiaxial stress state. Higher values are attributable to testing by either 3 or 4-point bend test
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[22, 23], of isostatically pressed [24] or colloidally processed materials [25]. Santos et al. [12], who used a test method and material preparation techniques similar to those of this study, obtained slightly lower strength values for HA. The higher values in this study may be attributed to attainment of a higher density at a temperature sufficiently low to avoid significant grain growth. The compositions of the strongest composite materials in this study are close to those of BCPs previously found to have good osteoconductive properties [20, 21, 26]. Additions of 5 and 10 wt% glass resulted in strength increases over HA of about 160—200%. These strength increases were lower than measured in previous work [11, 12] for similar compositions using similar test methods, however values for the strongest materials were comparable and within the error ranges of the cited works. In this work, additions of 2.5 wt% glass did not produce a significant strength increase compared to HA. Thus, a minimum phosphate glass addition of 5—10 wt% is required for the reinforcement of HA. This differs from earlier work [12], which concluded that the addition of 2.5% phosphate glass was sufficient to achieve maximum, with no further increase in strength obtained from higher glass additions. Microstructural characterisation suggested that high composite strength is derived from almost full densification and the formation of fully interpenetrating matrices of HA and TCP (Fig. 5b). The strength of the series based on 10 wt% additions decreased when sintered above 1300°C. Previous work [27] suggested that composite strength may decrease due to microcracking arising from the combined expansions associated with phase transformation of HA to b-TCP and of b-TCP to a-TCP. In this work, expansion associated with these phase transformations did not seem to cause sufficient defects to result in deterioration of mechanical strength. This may be due to the more limited reaction of HA with added glass. However, where appreciable degradation of HA occurred following almost full densification, substantial microstructural deterioration in the form of large pore formation (possibly due to OH~ release), appears to have contributed to a loss of mechanical strength (Fig. 5d). Composites based on high glass additions contained large amounts of either a-CPP or b-CPP. Dense composites containing a-CPP were mechanically very weak. b-CPP, which is present as a precursor phase in the process of bone formation [28], has previously been used as a bone replacement material [28—30]; however, there is little detailed information available regarding its biological behaviour, particularly with respect to osteoconductivity and rate of biodegradation, though both a-CPP and b-CPP are thought to be more resorbable than b-TCP. Kitsugi et al. [30, 31] found that dense b-CPP appeared to have bone bonding ability similar to HA or
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b-TCP. Although strengths similar to HA were measured for composites with 25 wt% glass additions, sintered at 1200°C, its composition of about a 1 : 1 ratio of bTCP : b-CPP is likely to make it an unsuitable material, particularly if used in porous form, due to excessively high resorption rates. Consequently, applications for composites based on high phosphate glass additions are likely to be limited by high resorption rates but may prove useful where dense, fully resorbable materials or where resorbable carrier matrices are required.
5. Conclusions The compositional, mechanical, and microstructural properties of HA/phosphate glass composites, including some compositions not previously characterised, have been investigated. Composites based on low glass additions (up to 10 wt%) offer the greatest potential due to their close compositional similarity to established osteoconductive BCPs as well as their high flexural strengths and improved fracture toughness. Glass additions of up to 10 wt% resulted in HA/TCP composites. The HA/TCP ratio was strongly influenced by the percentage glass addition, possibly allowing some control of implant resorption by varying percentage glass addition. The degree of reaction between glass and HA was lower than in a previous study of composites containing up to 5 wt% glass additions. Glass additions of at least 5 wt% were required to achieve significant strength improvement. Flexural strength increases over HA of 160—200% were obtained for composites containing 5 and 10 wt% glass additions. Higher phosphate glass additions resulted in composites containing large amounts of a-CPP or b-CPP with similar mechanical strengths to pure HA. Acknowledgements The authors wish to thank Dr. Robert Hill, University of Limerick, for providng access to glass-making facilities. This work was financially supported by Forbart (the Government Agency for Enterprise, Innovation, Investment and Development). References [1] LeGeros RZ. Biodegradation and bioresorption of calcium phosphate ceramics. Clin Mater 1993;14:65—88. [2] Hench LL, Wilson J, editors. An Introduction to Bioceramics, vol. 1. Singapore, World Scientific, 1993. [3] Murray MGS, Wang J, Ponton CB, Marquis PM. An improvement in processing of hydroxyapatite. J Mater Sci 1995; 30:3061—74. [4] Lelie´vre F, Bernache-Assollant D, Chartier T. Influence of powder characteristics on the rheological behaviour of hydroxyapatite ceramics. J Mater Sci: Mater Med 1996;7:489—94.
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D.C. Tancred et al. / Biomaterials 19 (1998) 1735—1743 [27] Knowles JC. Development of hydroxyapatite with enhanced mechanical properties: effect of high glass additions on mechanical properties and phase stability of hydroxyapatite. Br Ceramic Trans 1994;93:100—3. [28] Lin FH, Lin CC, Lu CM, Liu HC, Sun JS, Wang CY. Mechanical properties and histological evaluation of sintered b-Ca P O 2 2 7 with Na P O ) 10H O addition. Biomaterials 1995;6:793—802. 4 2 7 2 [29] Hubbard WG, Hirthe WM, Mueller KH. Physiological calcium phosphate implants. In: Kahn A, editor. Proc. 26th Annual Conf.
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on Engineering in Medicine and Biology, vol. 15. Minneapolis, Minesota, 1973:198. [30] Kitsugi T, Yamamuro T, Nakamura T, Kotani S, Kokubo T, Takeuchi H. Four calcium phosphate ceramics as bone substitutes for non weight-bearing. Biomaterials 1993;14:216—23. [31] Kitsugi T, Yamamuro T, Nakamura T, Oka M. Transmission electron microscopy observations at the interface of bone and four types of calcium phosphate ceramics with different calcium/phosphate molar ratios. Biomaterials 1995;16:1101—7.
A quantitative study of the sintering and mechanical properties of hydroxyapatite/phosphate glass composites D.C. Tancred*, B.A.O. McCormack, A.J. Carr Bioengineering Research Centre, Department of Mechanical Engineering, University College Dublin, Dublin, Ireland Received 11 August 1997; accepted 8 March 1998
Abstract Previous work has shown that small additions of a phosphate glass (CaO—P O ) can significantly enhance the sinterability and 2 5 strength of hydroxyapatite. However, there are no quantitative phase analyses available for these materials which would provide indicators of biocompatibility and resorbability. Similarly, there is little information available about the mechanical properties, especially with high glass additions. In this study, the effects of sintering hydroxyapatite with phosphate glass additions of 2.5, 5, 10, 25, and 50 wt.% are quantified. Each composition was sintered over a range of temperatures, and quantitative phase analysis was carried out using XRD. In addition, the microstructures were studied using RLOM and SEM, and mechanical properties (Vickers hardness, K , and MOR) measured. These results may be used to indicate which compositions and processing conditions may IC provide materials suitable for use in hard tissue replacement. Composites containing up to 10 wt.% glass additions formed dense HA/TCP composite materials possessing flexural strength and fracture toughness values up to 200% those of pure HA. The HA/TCP ratio was strongly dependent on the percentage glass addition. Higher glass additions resulted in composites containing b-TCP together with large amounts of a- or b-calcium pyrophosphate, and having similar mechanical strengths to pure HA. ( 1998 Published by Elsevier Science Ltd. All rights reserved Keywords: Hydroxyapatite; Glass; Calcium phosphate; Composite; Bioceramic
1. Introduction Calcium phosphate ceramics, especially hydroxyapatite (HA) and the more resorbable b-tricalcium phosphate (b-TCP), are widely used for hard tissue replacement due to their biocompatibility and osteoconductive properties [1, 2]. HA/b-TCP composites (frequently termed biphasic calcium phosphates; BCPs) have been developed to allow variation in implant resorption rates whilst retaining useful bioactive properties. Other calcium phosphates such as a-TCP, tetracalcium phosphate, and b-calcium pyrophosphate (b-CPP), though bioactive, have proven less useful as bone replacement materials due to excessively high resorption rates. In general, applications of calcium phosphates in the body have been limited by low strength and low fracture toughness. Numerous techniques have been investigated in attempts to improve the mechanical properties for particular applications or implant configurations, through advanced powder processing techniques [3, 4] *Corresponding author. Tel.: #353 1 706 1752; fax: #353 1 706 1756.
or by formation of HA-containing composite materials. Such composites aim to retain their useful bioactive properties whilst providing more suitable mechanical properties for particular applications. These include fibre reinforcement of HA [5, 6], HA/polyethylene [7], and HA/Al O [8] composites. 2 3 In recent years increasing interest has been shown in sintering of HA with glass additions. There are two main motivations for sintering HA with a glassy phase—to enhance the densification and therefore the mechanical strength by acting as a sintering aid, and to enhance bioactivity through the combination of two bioactive phases [9, 10]. Sintering with low glass additions may promote densification through liquid-phase sintering, resulting in composite materials with enhanced mechanical properties. For example, small additions of phosphate glasses have been shown to significantly enhance composite flexural strength (by up to 400%), and fracture toughness (by up to 200%) [11—13]. However, sintering in the presence of a glass was found to cause degradation of HA due to reaction with the glass to yield either b-TCP or a-TCP. Glass composition and the level of
0142-9612/98/$ — See front matter ( 1998 Published by Elsevier Science Ltd. All rights reserved. PII S 0 1 4 2 - 9 6 1 2 ( 9 8 ) 0 0 0 8 2 - 9
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glass additions have a large effect on both the phase assembly and the mechanical properties of the resulting composites. In particular, the presence of sodium in a glass was found to promote greater degradation of HA [9, 11]. Mechanical properties of HA composites containing Na O, though better than for pure HA, were 2 poorer than for soda-free glass composites [11, 12]. Use of a phosphate glass (in the system CaO—P O ) 2 5 ensures that the sintered body contains only calcium phosphate phases, which are likely to be biocompatible and possibly bioactive. Previous work [11] showed that glasses containing a higher Ca/P ratio caused less degradation of HA during sintering. The Ca/P ratio is practically limited to about 50 : 50 mol% CaO : P O , due to 2 5 a sharp increase in the liquidus temperature at higher Ca/P ratios [14]. In this work, the effects of phosphate glass additions, from the system CaO—P O , on the sintering, phase 2 5 composition, and mechanical properties of HA are analysed. The objective is to provide a more comprehensive understanding of the mechanical and possible biochemical properties of such composites, which should help to determine the optimum composites for use as implant materials.
2. Materials and methods 2.1. Preparation of biomaterials Synthetic hydroxyapatite was prepared using a wet method based on addition of calcium hydroxide to orthophosphoric acid at 85°C according to the reaction [15]: 6H PO #10Ca(OH) PCa (PO ) (OH) #18H O. 3 4 2 10 46 2 2 Phosphate glass was prepared using conventional glassforming techniques. Reagent-grade CaCO and P O 3 2 5 were melted in a fine-grained mullite crucible (Zedmark Refractories, Dewsbury, UK) at 1070°C for 2 h, to produce a glass of 50 : 50 CaO : P O molar ratio. The result2 5 ing melt was quenched in water, and the glass frit collected and quickly dried in a conventional oven at 100°C to minimise corrosion due to contact with water. HA and glass frit powders were separately dry ballmilled for 1 h in an alumina pot and passed through a 45 lm sieve. HA and composite powders containing 2.5, 5, 10, 25 and 50 wt% glass were then prepared and wet ball-milled in 60 g batches in 500 ml polypropylene NalgeneTM pots together with 180 ml isopropyl alcohol (IPA) and approximately 650 g of 7 mm diameter cylpeb magnesia-stabilised zirconia milling media. Powders were wet milled on ball-mill rollers at 140 rpm for 24 h. Recovered slips were passed through a 38 lm sieve, dried at 70°C and then passed through a 106 lm sieve to break any agglomerates formed on drying.
The powders were mixed with 4 wt% organic binder (Bindemittel B11/V, Degussa, Germany) to improve green strength and uniaxially pressed at 120 MPa in a 25 mm diameter steel die using 2 g of powder per disk. Following curing of the binder, samples were heated to 350°C at 2°C/min to remove the binder. Sintering was carried out in a muffle furnace, in air, by heating at 4°C/min to sintering temperature which was held for 3 h, followed by cooling in the furnace at 4°C/min. All samples were surface ground on one side (R "1—1.5 lm) for ! mechanical testing. 2.2. Compositional analysis Phase analysis was performed by X-ray diffraction using a Huber 642 Guinier diffractometer, fitted with a quartz Johansson monochromator, operating in subtractive transmission mode at 50 kV and 40 mA. Radiation was pure monochromatic Cu Ka (j"1.54056 As ). 1 Samples were ground to a fine powder using an agate mortar and pestle and mixed with Si powder as an internal standard to allow peak correction, and for quantitative analysis. Samples were scanned from h"0—50° in increments of 0.02°, using a count time of 5 s, and a counter slit width of 0.2 mm. Peak positions were corrected using a calibration curve obtained from the diffraction peaks of the internal standard. Lattice parameters were calculated (with 95% confidence intervals) for HA and b-TCP phases of the sintered composites. Quantitative phase analysis was performed by preparing XRD calibration curves relating the integrated intensities of the strongest peak of individual component phases to that of the internal standard. Estimation of the percentage a-TCP present was performed using the I/I #03 values (taken from the JCPDS PDF or experimentally determined using standard mixtures where JCPDS data were unavailable) for a-TCP and b-TCP. Due to the high resolution of the Guinier diffractometer, as well as full crystallisation of the glass phase for all additions, difficulties in quantitative phase analysis experienced by other researchers [12] due to peak overlapping were absent. The quantitative analysis ignored any effects of deviations from stoichiometry of individual phases. Lattice parameters for the composite series showed appreciable deviations did occur for both HA and b-TCP. 2.3. Mechanical testing Flexural strength testing was performed using a disk bend test method [16]. The test applies an equi-biaxial loading to thin disc specimens by loading with concentric rings of steel ball bearings. Testing was performed using a Lloyd 6000S mechanical testing machine, fitted with a 500 N load cell, using a cross-head speed of
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0.5 mm min~1. Samples were tested with the ground surface in tension. Vickers hardness was measured using a Mitutoyo AVK-C2 hardness tester. Fracture toughness was determined using a Vickers indentation technique [17]. A minimum of six samples were used for each mechanical test. Ninty five per cent confidence intervals were calculated for each series. 2.4. Microstructures Microstructures for all sintered materials were examined using a JEOL 35C SEM operating at 25 kV. Polished surfaces were etched using 0.1 M lactic acid. Hydroxyapatite samples were etched to reveal the grain structure using an etching time of 10 min. Phosphate glass composite series were etched for 10 s to reveal the location of individual phases. Although longer etching times showed greater grain boundary detail, it became difficult to distinguish between individual phases and to determine the extent and type of porosity.
3. Results 3.1. Phase analysis Analysis by XRD showed the HA synthesised in this study was stable up to 1350°C in air. Lattice parameters for HA, calcined at 900°C for 15 h, were a" 9.4198$0.0007 As and c"6.8821$0.0008 As . No detectable second phase was observed at any sintering temperature. Sintered composites containing low phosphate glass additions (up to 10 wt%) consisted of HA/b-TCP or HA/b-TCP/a-TCP, depending on the percentage glass addition and sintering temperature. For 2.5 and 5 wt% glass additions, the HA and TCP contents were largely independent of the sintering temperature with the percentage HA decreasing only slightly over the sintering temperature range. At 1350°C, a small amount of b-TCP converted into a-TCP. There was no significant decrease in HA content at this sintering temperature. Glass additions of 10 wt% resulted in greater degradation of HA. The quantity of residual HA was lower (about 10—40 wt%) across the temperature range. Above 1250°C, there was an increase in the degradation of HA accompanied by both an increase in the b-TCP phase content and the formation of a-TCP (Fig. 1). The lattice parameters of HA and b-TCP changed significantly as the percentage glass was increased, indicating a change in stoichiometry due to either lattice vacancies or substitutions. With an increasing percentage glass addition, the quantity of the residual HA content decreased. In addition, the composition moved farther from stoichiometry as was evident from a decrease in
Fig. 1. Phase composition of HA/phosphate glass (PG) composites.
Table 1 Lattice parameters for HA/phosphate glass composites (95% confidence limit of final digit in parentheses) Composition
Temperature (°C)
a (A_ )
c (A_ )
HA
1250 1300 1350
9.4154 (4) 9.4126 (7) 9.4134 (8)
6.8865 (5) 6.8832 (8) 6.8839 (9)
2.5%
1250 1300 1350
9.4118 (7) 9.4081 (8) 9.4113 (8)
6.8877 (9) 6.8882 (9) 6.888 (1)
5.0%
1250 1300 1350
9.411 (1) 9.408 (1) 9.411 (1)
6.887 (1) 6.888 (2) 6.888 (2)
10.0%
1250 1300 1350
9.404 (5) 9.403 (3) 9.409 (3)
6.893 (6) 6.895 (4) 6.891 (4)
a and an increase in c, away from stoichiometric values (Table 1 and Fig. 2). Lattice parameters for b-TCP moved towards stoichiometry with increasing percentage glass incorporation (and with increasing percentage bTCP in the composite). For 25 wt% additions, sintering at 1100°C resulted in approximately a 1 : 1 ratio of HA : b-TCP. Further increase in temperature to 1150°C caused complete degradation of HA to form b-CPP. Higher temperatures resulted in the transformation of all b-CPP to a-CPP. Glass additions of 50 wt% resulted in elimination of all HA to form b-TCP/b-CPP in a ratio of about 1 : 25. The phase composition was constant up to a sintering temperature of 1150°C. Sintering at 1200°C was found to cause transformation of all b-CPP to a-CPP. 3.2. Densification and microstructures Densification of HA was almost complete at 1200°C, reaching a maximum of 97.8% theoretical density at
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Fig. 2. Variation in lattice parameters of HA for HA/phosphate glass (PG) composites.
1300°C. The percentage densification of each composite sample was adjusted to allow for the difference in density of component phases (Fig. 3). Composites based on 2.5 and 5 wt% glass additions attained maximum densification above 1250°C, similar to that of HA. Slightly lower densification was achieved for composites containing 10 wt% glass additions, whose density decreased rapidly above 1250°C, due to increased reaction of HA, the formation of significant quantities of a-TCP and the generation of large pores (Fig. 5d). Samples based on 25 wt% glass additions showed most densification occurred by 1150°C. At 1200°C, there was a slight decrease in the measured density due to transformation of b-CPP (o"3.129 g cm~3) to a-CPP (o"2.930 g cm~3), however densification continued up to 1250°C. Composites containing 50 wt% glass showed a monotonic increase in density with temperature. Pure HA showed rapid grain growth above 1200°C. Average grain size increased from 1.0$0.05 lm at 1200°C to 2.47$0.13 lm at 1350°C. Residual porosity above 1200°C consisted of submicron pores existing at grain boundaries (Fig. 4). Glass additions of up to 10 wt% resulted in an increasing degree of porosity with increasing percentage glass addition when sintered at 1200°C, probably due to reaction between the glass and HA. Between 1200 and 1250°C, there was a substantial decrease in porosity for each composite series as the materials reached almost maximum density. Residual porosity principally existed at phase boundaries. This
Fig. 3. Densification of HA/phosphate glass composites (% represents percentage of theoretical density).
Fig. 4. Microstructures of sintered HA (etched for 10 min, 0.1 M lactic acid): (a) 1250°C, (b) 1350°C.
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Fig. 5. Microstructures of HA/phosphate glass (PG) composites (etched for 10 s, 0.1 M lactic acid). (a) 5 wt% PG,1250°C; (b) 5 wt% PG, 1350°C; (c) 10 wt% PG, 1250°C; (d) 10 wt% PG, 1350°C. Note: A"HA, B"TCP, P"Pore.
was consistent with previous reports of pore formation associated with OH~ release from HA during sintering [9, 18]. Pores were typically less than 1 lm except where there was a further decrease in the amount of HA present due to continued reaction with the glass (which occurred for 10% additions, sintered above 1250°C). It was not possible to distinguish between a-TCP and b-TCP regions in microstructures. TCP did not exist along grain boundaries as might have been expected in a liquid-phase sintered composite, but instead as isolated areas dispersed throughout the matrix. As the percentage glass addition (and the percentage TCP) was increased, TCP areas grew and coalesced to form a continuous TCP matrix with interspersed HA areas. Glass additions of 25 wt% caused a phase change between 1100°C and 1150°C from about 50HA : 50bTCP to 50b-TCP : 50b-CPP. A further increase in temperature to 1200°C caused transformation of b-CPP to a-CPP. Although these phase changes did not prevent densification, the combined effects of densification and phase changes had a large influence on the strength of the composites. Composites based on 50 wt% additions showed little change in microstructure or composition over the sintering temperature range investigated. Resid-
ual porosity was associated primarily with the presence of b-TCP. 3.3. Mechanical testing 3.3.1. Flexural strengths Comparison of composite flexural strength data with that of HA (Table 2), showed maximum average strength was achieved for 5 wt% glass additions sintered at 1350°C (61.9$7.5 MPa) and was about twice that of HA (31.4$3.1 MPa) which was obtained for samples sintered at 1200°C. Maximum strengths for 10 wt% additions were about 160% of HA values and occurred at a sintering temperature of 1300°C. There was no significant improvement in maximum strength for 2.5 wt% additions, however, unlike HA there was no significant decrease in average strength when sintered beyond the optimum sintering temperature. The optimum strength for each series was achieved at higher temperatures than for HA alone. All composites based on low glass additions ()10 wt%) showed a significant decrease in residual porosity between 1200°C and 1250°C. This corresponded to a significant increase in strength over the same sintering temperature range (Table 2).
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Table 2 Flexural strength in MPa of HA and HA/phosphate glass composites (95% confidence limit in parentheses, n"6) Sintering temperature (°C) 1000 1050 1100 1150 1200 1250 1300 1350
HA
20.5 31.4 17.9 20.9
2.5% PG
(3.4) (3.1) (1.9) (1.9)
15.6 32.3 27.4 25.2
(2.8) (4.0) (4.4) (6.9)
Incorporation of 25 wt% glass resulted in similar strength values to HA at 1150°C. Between 1100 and 1150°C, there was a large decrease in porosity. This was accompanied by a reaction in which all HA was eliminated and was converted largely to b-TCP. The phase change does not seem to have disrupted sintering as there was a significant improvement in strength and densification by 1150°C. The transformation of such a large amount of HA over a narrow temperature range could have been expected to interrupt sintering due to volumetric change, as well as OH~ release. However, transformation may have occurred while there was still a substantial amount of open porosity which prevented entrapment of released water vapour. Sintering at between 1150 and 1200°C transformed all b-CPP to a-CPP. There was a significant decrease in strength over the same temperature range. Microstructural examination showed that the a-CPP component consisted of large grains which increased in size rapidly above 1200°C. The combined effect of volumetric change and coarse grain size may have been sufficient to weaken the material. There was no statistically significant change (at a 95% confidence level) in average flexural strengths with increasing sintering temperature for composites containing 50 wt% glass additions. Nor was there a significant difference in strength between the maximum strength of pure HA and these composites, for any sintering temperature.
3.3.2. Hardness and fracture toughness The fracture toughness of composites based on glass additions of up to 10 wt% increased with increasing percentage glass addition (Fig. 6). Although the increase was not always significant for adjacent compositions, the general trend applied consistently across the whole series. There was a corresponding decrease in hardness with increasing glass additions (Fig. 7). No K value was IC obtained for samples containing 10 wt% glass, sintered at 1300°C, due to extensive lateral vent crack formation on indentation.
5% PG
14.2 21.4 40.0 42.2 61.9
(3.0) (3.9) (3.5) (3.9) (7.5)
10% PG
19.6 21.8 45.9 50.5 19.9
(4.2) (4.9) (4.4) (2.6) (4.7)
25% PG
23.7 37.6 11.3 14.0
(5.0) (2.0) (1.0) (0.8)
50% PG 21.8 29.1 28.3 28.0
(7.9) (6.8) (5.0) (11.8)
Fig. 6. Fracture toughness of HA and HA/phosphate glass (PG) composite materials.
4. Discussion Composites containing phosphate glass additions showed two types of behaviour. Those containing up to 10 wt% glass additions formed dense, strong, HA/TCP composites. Higher additions led to the formation of more porous composites containing large amounts of b-CPP or a-CPP. For low glass additions, the HA/TCP ratio depended strongly on the percentage glass added. Thus, it is likely that implant resorption, which is strongly dependent on the HA/TCP ratio, may be controlled by altering the percentage glass addition. The osteoconductivity of HA, of b-TCP, and of HA/b-TCP composites (biphasic calcium phosphates), has been clearly demonstrated by numerous animal and clinical trials [19]. In addition, BCPs formed by sintering of HA with glass (of
D.C. Tancred et al. / Biomaterials 19 (1998) 1735—1743
Fig. 7. Vickers hardness of HA and HA/phosphate glass (PG) composite materials.
the system CaO—P O —Na O has been demonstrated as 2 5 2 bioactive and osteoconductive [20]. Previous work has found that BCP implants may achieve a better tissue response and stronger bone ingrowth, in macroporous form than either HA or b-TCP alone [21]. Although it may not be possible to define an optimal HA/TCP ratio applicable to every situation, due to variations in implant characteristics as well as surgical site and application, there is evidence from previous work [21] that compositions of about 60 : 40 to 80 : 20 may be advantageous. However, further research into the effects of variations of the HA/b-TCP ratio is required to determine what effect variations in glass additions may have on surface activity or bone ingrowth. Small amounts of a-TCP formed in composites sintered at the highest sintering temperatures. Despite its much higher resorption rate the quantity of a-TCP is likely to be sufficiently low not to significantly affect the overall degradation rate. However, it is possible that if the a-TCP formed occurred primarily at the grain boundaries, the degradation may be significantly increased due to preferential dissolution of the grain boundaries. The microstructural location of a-TCP was not determined but is assumed to be associated with b-TCP regions (Fig. 5). The maximal bending strength of HA (31.4$3.1 MPa at 1200°C) is lower than some values available in the literature. The disk bend test method employed in this work provides a more conservative estimate of strength due to the equibiaxial stress state. Higher values are attributable to testing by either 3 or 4-point bend test
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[22, 23], of isostatically pressed [24] or colloidally processed materials [25]. Santos et al. [12], who used a test method and material preparation techniques similar to those of this study, obtained slightly lower strength values for HA. The higher values in this study may be attributed to attainment of a higher density at a temperature sufficiently low to avoid significant grain growth. The compositions of the strongest composite materials in this study are close to those of BCPs previously found to have good osteoconductive properties [20, 21, 26]. Additions of 5 and 10 wt% glass resulted in strength increases over HA of about 160—200%. These strength increases were lower than measured in previous work [11, 12] for similar compositions using similar test methods, however values for the strongest materials were comparable and within the error ranges of the cited works. In this work, additions of 2.5 wt% glass did not produce a significant strength increase compared to HA. Thus, a minimum phosphate glass addition of 5—10 wt% is required for the reinforcement of HA. This differs from earlier work [12], which concluded that the addition of 2.5% phosphate glass was sufficient to achieve maximum, with no further increase in strength obtained from higher glass additions. Microstructural characterisation suggested that high composite strength is derived from almost full densification and the formation of fully interpenetrating matrices of HA and TCP (Fig. 5b). The strength of the series based on 10 wt% additions decreased when sintered above 1300°C. Previous work [27] suggested that composite strength may decrease due to microcracking arising from the combined expansions associated with phase transformation of HA to b-TCP and of b-TCP to a-TCP. In this work, expansion associated with these phase transformations did not seem to cause sufficient defects to result in deterioration of mechanical strength. This may be due to the more limited reaction of HA with added glass. However, where appreciable degradation of HA occurred following almost full densification, substantial microstructural deterioration in the form of large pore formation (possibly due to OH~ release), appears to have contributed to a loss of mechanical strength (Fig. 5d). Composites based on high glass additions contained large amounts of either a-CPP or b-CPP. Dense composites containing a-CPP were mechanically very weak. b-CPP, which is present as a precursor phase in the process of bone formation [28], has previously been used as a bone replacement material [28—30]; however, there is little detailed information available regarding its biological behaviour, particularly with respect to osteoconductivity and rate of biodegradation, though both a-CPP and b-CPP are thought to be more resorbable than b-TCP. Kitsugi et al. [30, 31] found that dense b-CPP appeared to have bone bonding ability similar to HA or
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b-TCP. Although strengths similar to HA were measured for composites with 25 wt% glass additions, sintered at 1200°C, its composition of about a 1 : 1 ratio of bTCP : b-CPP is likely to make it an unsuitable material, particularly if used in porous form, due to excessively high resorption rates. Consequently, applications for composites based on high phosphate glass additions are likely to be limited by high resorption rates but may prove useful where dense, fully resorbable materials or where resorbable carrier matrices are required.
5. Conclusions The compositional, mechanical, and microstructural properties of HA/phosphate glass composites, including some compositions not previously characterised, have been investigated. Composites based on low glass additions (up to 10 wt%) offer the greatest potential due to their close compositional similarity to established osteoconductive BCPs as well as their high flexural strengths and improved fracture toughness. Glass additions of up to 10 wt% resulted in HA/TCP composites. The HA/TCP ratio was strongly influenced by the percentage glass addition, possibly allowing some control of implant resorption by varying percentage glass addition. The degree of reaction between glass and HA was lower than in a previous study of composites containing up to 5 wt% glass additions. Glass additions of at least 5 wt% were required to achieve significant strength improvement. Flexural strength increases over HA of 160—200% were obtained for composites containing 5 and 10 wt% glass additions. Higher phosphate glass additions resulted in composites containing large amounts of a-CPP or b-CPP with similar mechanical strengths to pure HA. Acknowledgements The authors wish to thank Dr. Robert Hill, University of Limerick, for providng access to glass-making facilities. This work was financially supported by Forbart (the Government Agency for Enterprise, Innovation, Investment and Development). References [1] LeGeros RZ. Biodegradation and bioresorption of calcium phosphate ceramics. Clin Mater 1993;14:65—88. [2] Hench LL, Wilson J, editors. An Introduction to Bioceramics, vol. 1. Singapore, World Scientific, 1993. [3] Murray MGS, Wang J, Ponton CB, Marquis PM. An improvement in processing of hydroxyapatite. J Mater Sci 1995; 30:3061—74. [4] Lelie´vre F, Bernache-Assollant D, Chartier T. Influence of powder characteristics on the rheological behaviour of hydroxyapatite ceramics. J Mater Sci: Mater Med 1996;7:489—94.
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