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Mechanical properties of porous ceramics based on 30 nm hydroxyapatite powder for potential hard tissue engineering applications

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https://www.eduzhai.net American Journal of Biomedical En gineer in g 2012, 2(6): 278-286 DOI: 10.5923/j.ajbe.20120206.07 The Mechanical Properties of a Porous Ceramic Derived from a 30 nm Sized Particle Based Powder of Hydroxyapatite for Potential Hard Tissue Engineering Applications Gé rrard Eddy Jai Poine rn*, Ravi Krishna Brundavanam, Xuan Thi Le, De rek Fawce tt M urdoch Applied Nanotechnology Research Group. Department of Physics, Energy Studies and Nanotechnology School of Engineering and Energy, M urdoch University, M urdoch, Western Australia 6150, Australia Abstract In this paper, synthesised nanometre sized hydro xyapatite (nano-HAP) powders co mposed of spherical 30 ± 5 nm part icles were co mpacted and sintered at temperatures ranging from 650 to 1250℃ to form ceramics of varying porosity and mechanical strength. The size, c rystalline structure and morphology of both the synthesised nano-HAP particle powders and the compacted and sintered ceramics were investigated using both X-ray diffraction (XRD) and Field Emission Scanning Electron Microscopy (FESEM). Fu rther characterisation techniques such as Brunauer-Emmett-Teller (BET) particle surface area, porosity, bulk density, Vickers hardness and yield strength at various sintering temperatures were tested on sintered pellets and evaluated. Keywords Hydro xyapatite, Nano-Ceramics, Crystalline Structure, Porosity, Vickers Hardness, Yield Strength 1. Introduction Hydro xyapatite (HAP)[Ca10(OH)2 (PO4)6] is a hexagonal structured ceramic co mposed of calcium phosphate groups that is very similar to the mineral co mponent found in bone tissue. It is this close chemical similarity between synthetic HAP and the natural inorganic bone matrix co mponent that has resulted in the extensive research effort to emp loy synthetic HAP as a bone substitute and/or replacement in several clinical procedures[1, 2]. Synthetic HAP is one of the most attractive ceramics for many bio med ical engineering applications due to its four main inherent properties: 1) it’s biocompatibility to surrounding body tissues; 2) its biodegradability in situ is slow; 3) it provides good osteoconductivity; and 4) it has good osteoinductivity capabilit ies[3-5]. A recent investigation by Taniguchi et al. revealed that sintered HAP in contact with cellu lar material provided a good biocompatible response to soft tissues such as skin, muscle and gums[6]. It is this favourable tissue response that has made synthetic HAP an ideal candidate for a wide range of hard t issue engineering applications such as; bone repair, bone augmentation, coating metal-co mposite implants and as *Corresponding author: g.poinern@murdoch.edu.au (Gérrard Eddy Jai Poinern) Published online at https://www.eduzhai.net Copyright © 2012 Scientific & Academic Publishing. All Rights Reserved a filling material in both bone and teeth surgical procedures [7-9]. For many of these biomed ical engineering applicat ions, a dense HAP ceramic material, wh ich exh ibits the required mechanica l properties for load bearing, is needed to meet the requirements of an effect ive hard tissue scaffold. The scaffolds architecture has a significant ro le to play in determining how effect ive the surrounding bone tissue integrates with the implanted scaffold[10]. The rate and degree of bone in-growth into the scaffold structure or matrix is influenced by the pore size and the interconnecting properties of the pores within the scaffold matrix[11, 12]. The porosity, or storage capacity, of porous HAP can be defined as the percentage of space in a materia l not occupied by the ceramic matrix, usually called the volumetric porosity. In the case of cortical bone the pore size ranges fro m 1 to 100 µm with typical volu metric porosities ranging fro m 5 to 10%[13]. HAP ceramics with low volu metric porosities generally have high mechanical p roperties but provide only limited opportunities for cell and tissue in-growth. Porous HAP ceramics have the potential to provide a good biological environ ment to promote cell adhesion, cellular interactions, proliferation, and migration. The downside to an increase in porosity is the decrease in its mechanical properties, such as strength, stiffness and elastic modulus. Both the cellu lar response and the mechanical p roperties are dependent upon the pore size, porosity, interconnecting porosity and pore distribution[2]. The A merican Society for 279 American Journal of Biomedical Engineer ing 2012, 2(6): 278-286 Testing Materials (ASTM) has defined porous materials into three classifications, interconnecting (open pores), non-connecting (closed pores), or a combination of both[14]. The interconnecting porosity within the matrix scaffold is a very important property since it permits fluid flo w through the porous ceramic material, wh ich is crucial fo r cellular growth. The effectiveness of the flu id flow through this natural plu mbing system is called the permeability. Pore sizes smaller than 10 µm define a micro-porous structure in the ceramic. The size of these micro -pores usually prevents the influ x of cells into the pore structure, but flu ids are still ab le to penetrate into and flow through the scaffold structure. When the pore size is greater than 10 µm they are classified as macro-pores and are large enough for both cells and flu ids to enter the pore structure. The macro-pores have a large pore surface area to bulk vo lu me ratio which actively pro motes cell adhesion, cellular interactions, proliferat ion, and migration. Using techniques that can assist in forming a porous structure within the ceramic has many advantages, since the architecture of the porous structure can be controlled. The ab ility to select the pore size, pore geo metry and the interconnecting porosity allo ws cell and tissue growth throughout the porous scaffold structure. For example bone forming cells, osteoblasts grow far mo re efficiently when they are attached to a substrate rather than being suspended in a culture mediu m. In fact the ability of natural bone to bond with HAP is a major reason why many researchers have studied and continue to investigate this ceramic for b io medical applicat ions[15]. When the pores are large and open, the HAP matrix is usually formed into a strut like structure, which forms the pore walls. The resulting interconnecting open face pore structure produces a network of struts that form flow channels throughout the matrix. The pore channels are conducive for cell and tissue colonisation. In addition, the pore channels provide high permeab ility flows for the cells to be supplied with nutrients and for the removal of metabolic waste products from normal cellu lar activ ity[16]. This type of pore structure is ideal for cells and tissue growth, but it has poor mechanical properties. On the other hand, if the pores are closed, the HAP matrix forms a network of interconnecting plate like structures that produces a high density solid. This configuration prevents the passage of flu ids or cells to neighbouring pores or the rest of the scaffold. Since the architecture of the scaffold requires a significant amount of porosity to accommodate fluid transfer and tissue in growth, an effective balance between porosity and the mechanical properties such as strength must be achieved. It simp ly co mes down to the fact that when the porosity increases, the strength, along with other mechanical properties of the scaffold decrease rapidly. Because strength is an important property of a load-bearing scaffold, the internal architecture of the scaffold structure must be carefully considered since it strongly in fluences the ability of the scaffold to resist load[17]. Therefore, by adjusting the porosity, it is possible to fine-tune the strength of the scaffold for a site-specific bio med ical application[18]. Furthermore, as the porosity increases, there is an increase in matrix surface area which is exposed to the environmental effects of erosion, changing surface chemistry and cellular activity. For examp le, bioactive materials like HAP, exhib it surface modificat ion with time when exposed to bodily flu ids and cellular activ ity. So the application of high porosity, low strength biomaterials is restricted to non-load bearing applications. In this particular applicat ion, where the initial function of the scaffold is not load-bearing, then the porous scaffold can provide an effective functional imp lant or bio med ical device. Therefore, being able to fabricate and fine tune a particular strength/porosity HAP scaffold provides greater flexib ility in addressing the needs of the site specific application. To achieve this goal there is a need to use powder compaction-sintering techniques to fabricate a dense HAP ceramic with the appropriate degree of porosity to meet the requirements of a successful tissue scaffold platform. The process usually begins with the synthesis of the HAP powder v ia a wet chemical method and subsequent heat drying. The drying stage usually forms large agglomerates, which often form structurally weak HAP ceramics with inhomogeneous pores after compaction and sintering. To prevent particle agglo meration, the use of ultrasonic irradiat ion during wet milling is an efficient means of dispersing and de-agglomerating the sample part icles during the grinding process[19]. After the milling process, the resultant powders are composed of fine, uniformly sized particles in the nanometre range, and both homogeneous in phase and chemical co mposition with very little particle agglomerat ion[20, 21]. The strength of the resultant powder, co mpaction and sintered HAP ceramic mainly depends on grain size, grain distribution, porosity, and other micro-structural defects. Hardness testing is frequently used to characterise mechanical properties, such as strength, of bulk solid samples and thin films. The hardness technique involves pressing a hard indenter of well-defined geometry into the surface of the sample under a predetermined load[22]. Indentation is considered an attractive method for assessing the mechanical properties of materials since it is basically a non-destructive in contrast to other techniques such as bending, compression and extension of samples. The Vickers diamond pyramid indenter, whose opposite faces have the included angle of 136º is one of the mostly widely accepted indenter technique for determin ing the hardness of ceramic materials and is used in this research work[23]. The aim of this study was to compare the porosity and strength of HAP ceramic scaffo lds produced via a using powder compaction-sintering technique. The synthesised nano-HAP powders were compacted and sintered at temperatures ranging fro m 650 to 1250ºC to form ceramics of varying porosity and mechanical strength. The size, crystalline structure and morphology of the nano-HAP particle powders synthesised were investigated using both X-ray diffract ion (XRD) and Field Emission Scanning Electron Microscopy (FESEM ). In addit ion, the size, crystalline structure and morphology of the co mpacted and Gérrard Eddy Jai Poinern et al.: The M echanical Properties of a Porous Ceramic Derived from a 30 nm 280 Sized Particle Based Powder of Hydroxyapatite for Potential Hard Tissue EngineeringApplications sintered ceramics were also studied using XRD and FESEM. Further characterisation techniques such as particle surface area, porosity, relative density, hardness and yield strength at various sintering temperatures was investigated. 2. Materials and Methods 2.1. Materials The main reactants used to synthesis the n-HAP powders were calciu m nitrate tetrahydrate[Ca(NO3)2.4H2O] and potassium di-hydrogen phosphate[KH2PO4], while the pH control of the solutions was achieved by the addition of ammon iu m hydro xide[NH4OH]. All analytical grade reagents used in this work were supplied by Chem-Supply (Australia), wh ile Sig ma-A ldrich (Castle Hill, NSW, Australia) supplied the Stearic acid,[C18H36O2] which was used as a binding agent during the powder compaction procedure. The solutions containing the reactants were synthesised under the influence of ultrasound irradiat ion, which was provided by an UP50H Ultrasound Processor[50 W, 30 kHz, M S7 Sonotrode (7mm diameter, 80 mm length)] supplied by Hielscher Ultrasound Technology. All aqueous solutions used throughout this study were made using Milli-Q® water (18.3 MΩ cm-1) produced by an ultrapure water system (Barnstead Ultrapure Water System D11931; Thermo Scientific, Dubuque, IA). Figure 1. Schematic of the synthesis procedure used to produce nano-HAP powders and the sintering temperatures used in the processing of the test pellets 2.2. Synthesis of n-HAP P owders The procedure for producing the n-HAP powder begins with adding a 40 mL solution of 0.32M calciu m n itrate tetrahydrate into a small g lass beaker. The pH of the solution is then adjusted to 9.0 by slowly adding and mixing approximately 2.5 mL of ammoniu m hydro xide. The resulting solution was then exposed to ultrasonic irradiation for 1 h, with the processor set to 50 W and maximu m amp litude. At the end of the first hour a 60 mL solution of 0.19 M potassium d i-hydrogen phosphate was then slowly added drop-wise into the first solution while undergoing a second hour of ultrasonic irradiat ion. During the second hour, the Calciu m/Phosphate[Ca/P] rat io was maintained at 1.67, while the p H o f the solution was checked and maintained at 9.0. At the end of the second hour, the solution was then filtered using centrifugation (15,000 g) fo r 20 minutes at room temperature, the resultant white precipitate sample was then placed into a fused silica crucib le, which was supplied by Rojan Advanced Ceramics Pty Ltd, Western Australia. The crucible was then placed into a standard domestic household micro wave (1100W at 2450 MHz-LG® Australia) for a thermal treat ment period of 40 minutes at a power setting of 100 %. At the end of the thermal treat ment the sample ended up as a white agglomerated mass. Once cooled, the sample was ball milled to break up the agglo merations and produce an ultrafine n-HAP powder, see schematic procedure presented in Figure 1. Th is synthesis procedure is repeated until a suffic ient a mount of the n-HAP powder was available for advanced characterisation and for the manufacture of pellets needed for fu rther testing to determine properties such as porosity, density, hardness and elastic modulus. 2.3. XRD and FES EM Characterisati on Techni ques The size, crystalline structure and morphology of both the synthesised nano-HAP particle powders and the compacted sintered ceramics were investigated using both X-ray diffraction (XRD) and Field Emission Scanning Electron Microscopy (FESEM ) techniques. Powder XRD spectra were recorded at roo m temperature, using a Siemens D500 series diffracto meter[Cu Kα = 1.5406 Å radiation source] operating at 40 kV and 30 mA . The diffraction patterns were co llected over a 2θ range fro m 20° to 60° with an incre mental step size of 0.04° using flat plane geometry. The acquisition time was set at 2 seconds fro each scan. The powder XRD spectrum was used to identify the purity of the final nano-HAP powders and any other phases that were present, see Figure 2. The crystalline size of the particles in the powders was calculated using the Debye-Scherrer equation[Equation 5] fro m the respective XRD patterns and estimated fro m the corresponding FESEM micrographs, see Figures 3. In addition, the fract ion of crystalline phase (Xc) present in the HAP ceramic pellets sintered at various treatment temperatures was calculated using the crystallinity equation proposed by Landi et al.[Equation 6], see Figure 4[24]. The morphological and macro-structural features of the nano-HAP powders were investigated using FESEM. A ll micrographs were taken using a h igh resolution FESEM [Zeiss 1555 VP-FESEM ] at 3 kV with a 30 µm aperture operating under a pressure of 1×10-10 Torr. FT-IR spectroscopy investigations were carried out using a Bruker Optics IFS 66 series FT-IR spectrometer. The KBr pellet technique was used, in wh ich 2 g of nano-HAP powder was mixed with 5-10 g of spectroscopic grade KBr and then compressed at around 15 kPa to form a d isk. In addition, the FESEM micrographs were also used to estimate the nano-HAP particle size by graphically measuring the size of each particle. The part icle size of every particle in a 500 n m 281 American Journal of Biomedical Engineer ing 2012, 2(6): 278-286 square grid was measured and then the mean particle size and its weight is measured. The porosity of the open pores is was determined fro m the collected data. then calculated using Archimedes’ method and the following 2.4. Compacti on and Sintering of nano-HAP Samples After the init ial synthesis, nano-HAP powders underwent exhaustive ball milling to break up particle agglo merations and produce an ultrafine n-HAP powder that was suitable for for mula : Porosity = (m2-m1)/ (m2-m3) x 100 % (1) where m1 is the dry weight of the pellet measured in air, m2 is the weight of the water saturated pellet and m3 is the weight of the water saturated pellet suspended in Milli-Q® water. the compaction-sintering stage of the fabrication procedure. The measurements were carried out in t rip licate and the To improve densification and to improve the compaction mean results of both density and porosity are presented in properties of the nano-HAP powder, a binding agent was Table 1. added and thoroughly mixed with the powder[1% wt. of Stearic acid]. The binding agent is expelled fro m the powder 2.6. Vickers Hardness Measurements during the sintering process. The powder is then cold The hardness of a materia l is defined by the quotient of the compressed in a cy lindrical mould by a manually operated applied load (P) and the contact area (A ) between the single action axial hydraulic ram pressurized to 70 M Pa and indenter and the sample. All Vickers testing procedures use a maintained at this pressure for 1 h. The co mpaction 136º pyramidal diamond indenter that forms a square procedure was repeated to ensure co mplete co mpaction indentation when pressed into the surface of the test throughout the pellet; a typical pellet had a mean dia meter of specimen under an accurately controlled test load. The test 18.61 ± 0.05 mm and a length of 18.34 ± 0.05 mm. The equipment used in this study was a Zwick/Materialprufung compacted powder samples (green samples) or pellets were 3212 Hardness tester, Germany and the measurements were then sintered in a p rogrammable high temperature muffle carried out at Curtin University’s Material’s Science furnace[Model 60 SL, Kiln Manufacturers of Western laboratory Bentley Perth Western Australia. The test load of Australia] at treat ment temperatures of 650℃, 850 ºC, 1050℃ 10.2 kg was applied for a specific t ime period (dwell time: and 1250℃ for a period of 2 h. After sintering, the samples 10-15 s) before the indenter was removed leaving a square were then permitted to cool down to room temperature and shaped indentation in the surface of the test specimen. The then stored in airtight containers ready fo r further ch aracteris atio n . testing procedure was carried performed in accordance with the testing procedures outlined in ASTM E384[25]. The 2.5. Brunauer-Emmett-Teller (B ET) Surface Area, Density and Porosity 2.5.1. Brunauer-Emmett-Te lle r Surface Area Measurement The Brunauer-Emmett-Teller (BET) surface area measure ments of the nano-HAP powders were carried out by the Australian Co mmon wealth Scientific and Research hardness was then calculated using the standard formula: Hv = Ao P/d2 (2) where the constant Ao is contact area of the Vickers pyramid geometry defined by: Ao = d2/2 sin (θ/2) (3) Ao =1.8544 for[GPa] The size of the indentation diagonal d was determined Organisations, (CSIRO) Particle Analysis Servicesl optically fro m an FESEM image and a typical image is aboratory (PAS) in Perth, Western Australia. The adsorption presented in Figure 3 (b). Both diagonals were measured and technique used nitrogen gas to carry out the surface area the mean value was then used to calculate the Vickers measurement; the results are presented in Table 1 and Figure hardness of the sample using equation 3 and then the yield 5. strength of the samp le was calculated using equation 4 below. 2.5.2. Bu lk Density and Porosity Calcu lations The bulk density was determined by first weighing the mass of the pellet that had been previously dried under Both the Vickers hardness and yield strength (Y)[26] results are presented in Table 1. Y = Hv/ 3 (4) vacuum at 80 ℃ for 48 h. Th is was then followed by measurement of the dimensions of a pellet, which was used together with the mass data to calculate its density. The open pore volume was determined by immersing a prev iously vacuum dried pellet into a small beaker containing a s olution of Milli-Q® water for 2h by wh ich t ime no air bubbles left the pellet and indicated that the water saturation of the pellet was 100%. The pellet was then removed fro m the solution, weighed and then submerged back into the Milli-Q® water 3. Results and Discussions The XRD pattern of the init ial synthesised nano-HAP powder, along with the XRD patterns of the HAP pellets sintered at temperatures of 650℃, 850℃, 1050℃ and 1250℃ for a period of 2 h are presented in Figure 2. All samp le XRD patterns, except the 1250℃ one revealed the presence of crystalline nano-HAP phases. Gérrard Eddy Jai Poinern et al.: The M echanical Properties of a Porous Ceramic Derived from a 30 nm 282 Sized Particle Based Powder of Hydroxyapatite for Potential Hard Tissue EngineeringApplications Figure 2. XRD spectrum of synthesised nano-HAP powder and nano-HAP powders compressed at 70 MPa and thermally treated at various sintering t emp erat ures Table 1. Physical properties determined from the characterisation of test pellets thermally treated at various sintering temperatures Sin t erin g Temp (ºC) 650 850 1050 1250 BET Surface Area (m2/g) 10.977 5.438 1.137 0.039 Density P ellet (Dry) (g/cm3) 1.449 1.452 1.952 2.606 Po ro sit y (Open) (%) 53.02 46.66 36.78 16.90 Cry st allite Size (XRD) (± 5 nm) 323 491 911 1224 Cry st al Size (FESEM) (± 5 nm) 376 518 965 1318 Cry st allin ity (Xc) (%) 77.28 87.25 90.32 - Hardness (Hv) (MP a) 157 189 364 1727 Yield St ren gth (Y) (MP a) 52.3 63.0 121.3 575.7 These phases were found to be consistent with the phases listed in the ICDD database, with the main (h k l) indices for nano-HAP: (002), (211), (300), (202), (130), (002), (222) and (213) being indicated in Figure 2. The calciu m phosphate peaks found in the XRD pattern for 1250℃ where not HAP, but were identified as a potassium calciu m phosphate phase. The crystalline size, t(hkl), of the synthesized nano-HAP powder and the sintered HAP pellets was calculated fro m the XRD pattern using the Debye-Scherrer equation[27-30]. t(hkl ) = 0.9λ B cos θ(hkl) (5) where, λ is the wavelength of the monochromatic X-ray beam, B is the Full Width at Half Maximu m (FWHM ) of the peak at the maximu m intensity, θ(hkl) is the peak diffraction angle that satisfies Bragg’s law for the (h k l) plane and t(hkl) is the crystallite size. The (002) reflection peak fro m the XRD pattern was used to calculate the nano-HAP crystallite size in this study fro m the Debye-Scherrer equation and was estimated to have a mean value of 30 ± 5 n m. The calculation was done so that a comparison could be made between the crystallite size of the synthesised nano-HAP powder and the growth in the crystallite size with increasing sintering temperature. Since the degree of structural order or crystallinity of a material has a significant influence on properties such as density, hardness and strength, the percentage crystallin ity of the material was calcu lated using the equation proposed by Landi et.al[24], see equation 6. XC = 1-(V112-300/I300) (6) where XC is the percentage crystallinity of the material, V112-300 is the intensity of the trough between the (112) and (300) peaks and I300 is the intensity of (300) peak. The results of both crystallite size and the degree of crystallinity are presented in Table 1 and Figure 4. Initial XRD investigation revealed that the synthesis and microwave treated nano-HAP powder had a mean crystallite size of 30 ± 5 n m and with a crystallinity value of 48.17 %. The subsequent XRD spectra taken fro m powder samples taken fro m the various test pellets revealed that the sintering process promoted both crystallite growth and enhancement of crystallin ity. The init ial sintering temperature of 650℃ had produce an open porous HAP structure with a mean crystallite size o f 323 ± 5 n m and a crystallinity of 77.28 %. The higher sintering temperature of 1050℃ also formed an open porous HAP structure with a mean crystallite size of 911 ± 5 n m and a crystallin ity of 90.32 %. However, the higher sintering temperature of 1250℃ did not form a HAP 283 American Journal of Biomedical Engineer ing 2012, 2(6): 278-286 structure but instead produced a potassium calciu m phosphate phase with a mean crystallite size of 1224 n m (1.224 µm). The FESEM microscopy technique was used to investigate the size and morphology of the nano-HAP powders synthesised in this study. A typical image of the ultra-fine nano HAP powder synthesised in this work is presented in Figure 3(a). Inspection of the image presented in Figure 3(a) reveals the presence of a spherical particle morphology, wh ich is similar to the particle mo rphologies previously reported in the literature[25-30]. The mean particle size of the synthesised nano-HAP powder was determined fro m the FESEM images and found to have a mean value of 28 ± 5 n m, which was found to be favourably comparable with the calculated mean value of 30 ± 5 n m fro m the XRD spectra. The FESEM technique was also used to examine the size and morphology of the crystallite sizes formed in the sintered pellets. Figure 3 presents images of the 650℃, 850℃, 1050℃ and 1250℃ isotherms used to investigate the effects of sintering temperature on the growth and morphology of particles in each pellet. The overall trend seen in the FESEM micrographs (c) to (e) is an increasing particle size and a mo rphological structure becoming mo re angular, cubic and rectangular in nature with increasing sintering temperature. However, at the higher sintering temperature of 1250 ºC[Figure 3(f)], the character of the particles in the sample is quite different fro m those presented in earlier micrographs. The greater size diversity and shape change confirm the XRD results that indicated a phase change as verified by the presence of a potassium calciu m phosphate phase. Figure 3. (a) FESEM of synthesised nano-HAP ultrafine powders; (b) a typical FESEM micrographs of an indentation produced on the surface of a pellet duringthe Vickers hardnesstest; (c) to (f) are FESEM micrographs of samples thermally treated at sintering temperatures: (c) 650 ºC; (d) 850 ºC; (e) 1050 ºC; and (f) 1250℃ Further characterisation of the pellets involved determining their respective BET surface areas, porosities and Vickers hardness values which could then be used to calculate the yield strength of the compacted HAP pellet. After init ial co mpacting of the pellets, each pellet set was thermally t reated at various sintering temperatures. At the lowest sintering temperature of 650℃, both the BET surface area value of 10.977 m2/g and the open porosity value of 53.02% were at their highest levels see Figure 5. While at the higher sintering temperature 1050℃ both the BET surface area value of 1.137 m2/g and the open porosity value of 36.78 % were at their h ighest values for the HAP phase. The reduction in both surface area and porosity resulted fro m crystal growth during therma l treat ment. The lowest surface area and porosity occurred at the highest sintering temperature of 1250℃. However, at this temperature the HAP phase was transformed into a potassium calciu m phosphate phase. The much larger mean crystal size of 1.224 µm within the matrix of the test pellet significantly reduced the surface area down to 0.039 m2/g and also contributed to the reduced open porosity of 16.9 %. The increased sintering temperature also produced a significant increase in the Vickers hardness value of each sintered pellet; see Figure 6. At the lowest sintering temperature of 650 ℃, the mean Vickers hardness value was 157 MPa, wh ich slightly increased to 189 MPa at the higher sintering temperature of 850℃. The hardness value then increased to 364 MPa at 1050 ℃ before dramatically increasing to 1727 M Pa at 1250℃. The phase changed detected by the XRD results is also confirmed by the hardness values which clearly indicate that a different, much harder phase material is present after sintering at this higher temperature. This is also reflected in the estimated yield strength for the pellets at each sintering temperature, see Figure 5. At 650 ºC the yield strength is at its lowest mean value of 52.3 MPa, wh ich slightly increases up to 63.0 M Pa at 850℃. The higher sintering temperature of 1050℃ sees a significant increase in yield strength to 121.3 MPa, which is approximately double the 850℃ value. The dramat ic increase in yield strength to 575.7 M Pa at 1250℃ indicates that the potassium calciu m phosphate phase present at this temperature is much stronger than the earlier HAP phase that was present at lower sintering temperatures. The fabricated pellets at the lower end of the sintering temperature range were predo minantly characterised by high porosity and high permeability. The large open porosity of around 50 % has the potential to permit the mig ration of cells into the pellet matrix and colonise the large surface areas within the internal structure. The large open porosity architecture also has the potential to pro mote the flo w of nutrients to the cells deep with in the mat rix and subsequently, the removal of metabolic waste products resulting fro m cell activity. This open pore architecture is ideal for cell colonisation and tissue in-growth, but lacks the mechanical strength needed for load bearing applications. At the lowest sintering temperature of 650℃, the y ield strength is only 52.5 MPa. Since the ideal tissue scaffold requires a significant amount of porosity to accommodate fluid transfer and tissue in growth, there needs to be an effective balance between the open porosity and the strength required to meet Gérrard Eddy Jai Poinern et al.: The M echanical Properties of a Porous Ceramic Derived from a 30 nm 284 Sized Particle Based Powder of Hydroxyapatite for Potential Hard Tissue EngineeringApplications the load demands of the scaffold for a specific application. Figure 4. Crystal size growth and increasing crystallinity with increasing sintering temperature porosity of 16.9 %. Un fortunately, the increased temperature also induced the decomposition of the HAP, wh ich in turn formed a potassium calciu m phosphate phase, a similar HAP decomposition has been reported by Yeong et al[31]. This transformation is evident in the XRD spectrum presented in Figure 2 and the FESEM micrograph presented in Figure 3(f), which of a wide range of particles ranging in size fro m 400 nm to 3 µm. Figure 6 illustrates the dependence of Vickers hardness of the sintered HAP pellets on sintering temperature. The hardness increases almost linearly with rising sintering temperature fro m 650℃ up to 1050℃. Fro m 1050 ℃ onwards, the hardness rapidly increases with increasing sintering temperature and density as the open porosity decreases. Currently, the biocompatib ility of the 30 n m sized hydroxyapatite porous ceramic is undergoing in vivo investigation in an Ovine (Sheep) model. The results of which will be published in a future article. Preliminary results indicate that there is only a very minor inflammatory response, with the imp lanted pellets receiving a favourable cellu lar response. At this stage, fibroblast cells can be seen attaching to the surface of a typical ceramic pellet 4 weeks after the imp lantation date. Figure 5. The effect of increasing open porosity in the test pellets on Yield Strength and BET surface area Figure 7. Fibroblast cells attaching onto the surface of a hydroxyapatite ceramic pellet made from n-HAP powders 4. Conclusions Figure 6. The effect of sintering temperature on the open porosity and Vickers hardness of test pellets The increased sintering temperature of 850℃ produced a slight increase in the yield strength (63 M Pa), which corresponded to a reduction in the porosity of around 6.4% to a value of 46.7%. It is not until the higher sintering temperature of 1050℃ that a significant increase in strength is achieved (121.3 MPa). At this temperature the resulting open porosity was around 37%. However, the higher sintering temperature of 1250℃ did produce dramat ically higher yield strength of 575.7 MPa with a much lower Nano-crystalline nano-HAP powders co mposed of spherical 30±5 n m particles were formed under the influence ultrasound irradiation and then thermally treated using a conventional microwave oven. The powders were then compressed at 70 MPa to form pellets which were then sintered at s intered at temperatures ranging from 650 to 1250 ºC. The study confirmed that Vickers hardness significantly increased with increasing sintering temperature and density as the porosity decreased. Meanwhile, the yield strength of the ceramic pellets decreased as both the porosity and BET surface area increased. During the sintering process both the crystal size and crystallin ity of the ceramics steadily increased with sintering temperature. Despite being advantageous, the reduction in mat rix porosity to increase the yield strength of the ceramic for load bearing applications, the reduced porosity would significantly reduce 285 American Journal of Biomedical Engineer ing 2012, 2(6): 278-286 cell migrat ion and proliferat ion within the ceramic matrix. Hence, further work is needed to reach a compro mise between porosity, strength and cellular integration potential of this nano-HAP derived ceramic for various bio medical ap p licatio n s . ACKNOWLEDGEMENTS Dr Derek Fawcett would like to thank the Bill & Melinda Gates Foundation for their research fellowship. The authors would like to acknowledge the assistance of Prof. Arie Van Riessen and Mr. Zhenhua Luo in performing the hardness tes tin g . van Blitterwijk, K. De Groot. 3D microenvironment as essential element for osteoinduction by biomaterials. Biomaterials, Vol. 26, No. 17, (2005), p. 3565-3575. [12] U. 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