DOI: 10.1021/cg900868t
Nucleation and Early Stage Crystallization of Fluorapatite in Apatite-Mullite Glass-Ceramics
2010, Vol. 10 1111–1117
Kevin P. O’Flynn and Kenneth T. Stanton* School of Electrical, Electronic and Mechanical Engineering, University College Dublin, Belfield, Dublin 4, Ireland Received July 24, 2009; Revised Manuscript Received November 25, 2009
ABSTRACT: The crystallization morphology and activation energy of fluorapatite formation in apatite-mullite glass-ceramics has been examined. A systematic change in the amount of network modifier in the base glass has demonstrated that increasing levels of network modifier cause a corresponding increase in activation energy for fluorapatite crystallization and a decrease in fluorapatite crystal size. Apatite-mullite glass-ceramics are known to phase separate and nucleate prior to crystallization. To separate these effects during activation energy determination, a nucleation hold was used to ensure the activation energy is purely for crystal growth and is not skewed by any phase separation or nucleation occurring. It is postulated that greater amounts of network modifier in the system cause the base glass to phase separate and nucleate to a greater extent during heating thus limiting the crystal growth.
Introduction Apatite-mullite glass-ceramics (AMGCs) are under investigation as a possible alternative to hydroxyapatite for use in vivo as dental implants1 and as coatings for orthopedic metallic implants.2-4 When heat-treated, they possess a biphasic crystal structure of fluorapatite (FAp, Ca10(PO4)6F2) and mullite (Al6Si2O13), an inert phase with excellent mechanical properties.1 FAp is a bioactive, stable chemical analogue of hydroxyapatite (HA, Ca10(PO4)6(OH)2) and has a lower bioresorption rate than HA.5,6 The use of glass-ceramics over conventional ceramics is advantageous as they can be cast and molded to near net shape with no porosity as a glass. Unlike apatite-wollastonite, another commonly used glass-ceramic which must be crystallized as a powder prior to pressing and sintering,7 apatite-mullite glasses can be cast as a glass into the desired shape then crystallized as a bulk material. This allows for little or no postprocessing of the material after it has been formed. Rafferty et al.8 have shown that AMGCs undergo amorphous phase separation prior to crystallization via a nucleation and growth mechanism in the same glass system as used here. It has also been shown that fluorapatite in the SiO2Al2O3-CaO-P2O5-K2O-F-system forms in phase separated droplets and the growth of the apatite crystals seems to be limited by the size of these droplets which are on average 500 nm in diameter.9 Energy dispersive X-ray (EDX) shows these regions to be rich in calcium, fluorine, and phosphorus with diminished levels of aluminum and silicon. In SiO2Al2O3-MgO-K2O-CaO-P2O5-F glasses, the growth of the apatite phase is similarly limited by droplet size.10 There, the apatite crystallized to a hexagonal morphology approximately 2 μm in diameter but never grew beyond the droplet boundary.10 Within these droplets in some glass compositions, inclusions have been seen that transmission electron microscopy (TEM)-EDX indicate to be secondary glass-inglass phase separation occurring within the droplet itself and are high in aluminum.9 *To whom correspondence should be addressed. r 2010 American Chemical Society
Hill et al.11 have also shown that apatite-based glassceramics can phase separate not only by a nucleation and growth mechanism, but also via spinodal decomposition. They used an AMGC of composition 1.5SiO2-Al2O30.5P2O5-CaO-0.67CaF2. This glass was shown to undergo both spinodal and binodal decomposition during cooling following casting and subsequent reheating. The effect of amorphous phase separation on nucleation crystallization is very important when considering glassceramics. The phase separated regions are rich in the certain elements bringing the composition closer to that of the resulting crystals. Since the atomic species required for the crystal to grow are in abundance it is therefore easier for the crystal to grow; thermodynamically, we may say that there is a lower activation energy for crystal growth. In addition, the surfaces of the phase separated regions can act as heterogeneous nucleation sites which further reduces the barrier to nucleation. Where a glass exhibits an optimum nucleation temperature just above the glass transition temperature, this indicates that the glass undergoes amorphous phase separation prior to nucleation.11 Previous research has investigated the activation energy (Ea) for the formation of FAp in apatite-mullite glassceramics without the use of a nucleation hold and, while recognizing that the activation energy was dependent on the glass composition, has reported that “the activation energy plots were never very linear”.12 This may be because the material was nucleating and crystallizing concurrently. In such a case, the method will give an activation energy indicative of both the barrier to nucleation and the energy required to crystallize the glass together. Although the Matusita method,13 a commonly used technique for determining Ea, can be modified using Avrami parameters to account for an increasing number of nuclei, for example n=3 and m=4 for bulk nucleation and increasing numbers of nuclei, it is likely that a more accurate measurement can be made by having a constant number of nuclei, giving n =3 and m =3. This can be achieved by fully nucleating the sample at the optimum nucleation temperature, when the glass will have reached an equilbrium composition. Previous studies have looked at the effect of varying the Ca/P ratio of AMGC on the activation energy and showed Published on Web 02/02/2010
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O’Flynn and Stanton
Table 1. Summary of Results from DSC study of ONTa glass code
G250
G220
G200
G187
G175
G150
G125
G100
x value Tg,onset Tg,mid Tg,end Tp1 Tp1,end,ONT Tp2 ONT Tp1,ONT Tp2,ONT Eact Tp1 Tp1 error
2.5 609 629 649 682.94 883 833.69 610 681.81 850.98 -336145 21479
2.2 612.77 634.04 655.31 743.58 891 908.62 655 725.82 842.84 -423040 30112
2 631.26 650.21 669.17 755.32 905 961 670 729.87 877.83 -470726 21323
1.87 642.8 663.08 683.37 791.97 1011 991.38 680 744.72 972.69 -410798 30996
1.75 658.73 676.28 693.84 836 1073 1032 685 765.52 1048.93 -470810 67671
1.5 680.68 695.91 711.15 917 1085 1073 700 865.22 1058.16 -379721 44221
1.25 695.89 709.08 722.27 1008 1124 1101 715 958.03 1101.02 -330395 5550
1 708.35 726.46 744.57 1036 1152 1121 740 1030.1 1124.47 -291849 17310
a
All results are from runs done at 10 °C 3 min-1.
that it reduces with decreasing Ca/P ratios.14 It has been proposed that for AMGCs the activation energy may correspond to the activation energy for viscous flow as the glass structure is rearranged to form fluorapatite. In this case, the activation energy should decrease as the network connectivity decreases. The work presented here directly tests this hypothesis; glasses of different network connectivity are held at the optimum nucleation temperature for an hour to produce well-nucleated samples. In this case, it can be assumed that the activation energy is now solely representative of crystallization of FAp. Materials and Method Glass Synthesis. The particular apatite-mullite glass-ceramic system used here has generic composition of 1.5(5-x)SiO2-(5-x)Al2O3-1.5P2O5-(5-x)CaO-xCaF2 where x can vary between 1.0 and 2.5. The compositions used are listed in Table 1. This system has previously been examined with detailed thermal and X-ray analysis and is designed to maintain an apatitic Ca/P ratio of 1.67.1,15 The addition of fluorine to the glass acts as a network modifier, increasing the atomic mobility, and lowering the glass transition and crystallization temperatures. The reagents were mixed for 40 min in a ball mill without P2O5 due to the hygroscopic nature of the latter component. The P2O5 was then added and the mixture ball-milled for a further 15 min. The powder was then placed into a high-density lidded mullite crucible. Using an electric furnace, the crucible and charge were heated to 1440 °C and held at this temperature for 1 h. The glass melt was then poured directly into demineralized water to shock quench to room temperature, thereby avoiding crystallization and preserving the amorphous nature of the glass. The molten glass shatters upon cooling and is called glass frit. The frit was oven-dried in air for one day. During heating, low-fluorine glasses can lose fluorine at the surface in the form of silicon tetrafluoride and form anorthite (CaAl2Si2O8).12 To minimize unwanted surface effects, the frit was used for all thermal and phase analysis; there was therefore a minimal surface area for anorthite to form and results are representative of the bulk of the material. To create samples for microscopy, the glass was cast by remelting it in an alumina crucible and pouring into a graphite mold preheated to Tg þ 40 °C. The glass was then annealed by transferring both the graphite mold and glass to a furnace preheated to the same temperature. It was left at this temperature for 2 h before the furnace was turned off and the samples allowed to cool in the furnace. This produced well-annealed glass rods approximately 50 mm in length and 20 mm in diameter. Samples were sectioned using a Buehler Isomet 2000 PrecisionSaw into 2 mm thick slivers, using a diamond wafering blade. To examine the microstructural evolution of the glasses, samples were heated in a Carbolite furnace in air. Differential Scanning Calorimetry (DSC). Thermal analysis was performed using a Rheometric Scientific STA 1600 (Surrey, UK). Samples were placed in platinum-rhodium crucibles in a flowing dry nitrogen atmosphere. Each of the glasses exhibited a glass transition temperature onset Tg,onset, mid point Tg,mid, and end Tg,end, as well as two exothermic crytallization peaks, the first and
Figure 1. DSC traces of G200, indicating Tg,onset, Tg,mid, Tg,end, Tp1, Tp1,end, and Tp2. second of which are labeled Tp1 and Tp2, respectively, as shown in Figure 1. Also given in the latter diagram is the method for defining the end point of the crystallization exotherm, Tp1,end. Optimum Nucleation Temperature. The optimum nucleation temperature (ONT) is defined as the temperature at which the most number of stable nuclei form per volume element. The relative number of stable nuclei that form can be inferred from the crystallization temperature (Tp): a greater number of nuclei reduces the barrier to crystallization and lowers Tp. Therefore at the ONT, Tp will be a minimum. In practice, the ONT is determined by heating the glass to various temperatures between Tg,onset and Tp1, then holding at that temperatre for one hour before heating the sample to full crystallization. This is repeated for various temperatures before determining which hold temperature produces a minimum Tp. In all cases here, a heating rate of 10 °C.min-1 was used. A simple description of the rate of homogeneous nucleation, I, according to classical nucleation theory can be given as ΔG þ Q I ¼ A exp ð1Þ RT where ΔG* is the maximum free energy of activation for the formation of a stable nucleus, Q is the activation energy for the diffusion of molecules across the phase boundary, R is the universal gas constant, A is a constant, and T is absolute temperature.16 It may be seen that the nucleation rate is highly sensitive to the temperature. Matusita notes that “usually, the rate of crystal nucleation in glass reaches the maximum at a temperature somewhat higher than the glass transition temperature and then decreases rapidly with increasing temperature”.17 In this case it will be seen that the ONT occurs within the glass transition for all of the glasses observed. Activation Energy. The activitation energy, Ea, for crystallization can be determined in a number of ways. In particular, the Matusita17 or Kissinger18 methods are widely used. Kissinger18 states that for reactions observed using DSC at different heating rates, d ln
φ Tp2
1 d Tp
! ¼ -
Ea R
ð2Þ
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where φ is dT/dt. Ea can be easily determined by finding the slope of ln(φ/Tp2) against 1/Tp. The Kissinger method is useful as it is independent of nucleation and can be used for reactions of any rate order.18 The Matusita method is a refinement of Kissinger which requires knowledge of the nucleation mechanism. It proposes that φn d ln 2 Tp m d Tp
! ¼ -
Ea R
ð3Þ
where m is the number of dimensions in which the crystal grows and n is a numerical factor depending on the nucleation process.17 Matusita17 observes that n = m þ 1 for a quenched glass being heated at a constant rate containing no nuclei and n=m for a glass containing a sufficiently large number of nuclei, that is, after a nucleation hold. Three-dimensional growth (bulk crystallization) is given as m = 3, 2 for growth in thin film glass and 1 for surface crystallization. In the current experiment, heating rates of 5, 10, 20, 30, and 40 °C 3 min-1 were used for activation energy analysis. In all cases, each glass was held at its respective ONT for one hour then heated to full crystallization. Fully nucleating the samples ensures that the activation energy is purely for crystal growth and is not skewed by any phase separation or nucleation occurring. X-ray Diffraction (XRD). Powder XRD was performed to confirm the phases present in the samples. Prior to XRD, the sample was crystallized by heating in the DSC furnace to ensure accurate temperature and heating rate control. Analysis was carried out in a Huber 642 Guinier Diffractometer (Rimsting, Germany) with a quartz Johansson monochromator and copper target operating in subtractive transmission mode at 40 kV and 30 mA. X-rays were pure monochromatic Cu KR1 with λ=1.54056 A˚. Si powder (pure element grade, 99.5% pure, Johnson Matthey Alfa Products, Karlsruhe, Germany) was added at 10 wt % to all crystallized samples as an internal standard to allow correction for nonlinear peak shift. The Scherrer equation was used to determine particle size along a given reflection plane. Here, the (211) plane was measured as it corresponds to the 100% intensity peak for fluorapatite. The Scherrer equation states that the crystal size τ is given by: Kλ B cosðθÞ
ð4Þ
b2 ¼ B2 þ β2
ð5Þ
τ ¼
where b is the experimentally determined full width at half-maximum of the peak broadened due to to crystal size and instrumental factors, B is the broadening due to crystal size, β is the instrument broadening, K is the Scherrer constant chosen as 0.9 here, λ is the wavelength of the incident X-rays, and θ is the Bragg angle in radians. To make a more accurate measurement of the fwhm and peak θ value, a Gaussian curve19 was fitted to the data using a Levenberg-Marquardt algorithm which also provides an estimation of the error. The instrument broadening profile was determined by obtaining the diffraction pattern for lanthanum hexaboride (LaB6). Scanning Electron Microscopy (SEM). For microscopy, samples were cold mounted in epoxy, ground and polished to a
