Effect of ZrO2 addition on the mechanical properties of porous TiO2 bone scaffolds

Materials Science and Engineering C 32 (2012) 1386–1393

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Effect of ZrO2 addition on the mechanical properties of porous TiO2 bone scaffolds Hanna Tiainen a, Georg Eder a, b, Ola Nilsen c, Håvard J. Haugen a,⁎ a b c

Department of Biomaterials, Institute for Clinical Dentistry, University of Oslo, PO Box 1109 Blindern, NO-0317 Oslo, Norway Institute of Medical and Polymer Engineering, Chair of Medical Engineering, Technische Universität München, Boltzmannstrasse 15, 85748 Garching, Germany Department of Chemistry, University of Oslo, PO Box 1033 Blindern, NO-0315 Oslo, Norway

a r t i c l e

i n f o

Article history: Received 9 December 2011 Received in revised form 8 February 2012 Accepted 12 April 2012 Available online 20 April 2012 Keywords: Scaffold Mechanical strength TiO2 ZrO2 Porosity Bone tissue engineering

a b s t r a c t This study aimed at the investigation of the effect of zirconium dioxide (ZrO2) addition on the mechanical properties of titanium dioxide (TiO2) bone scaffolds. The highly biocompatible TiO2 has been identified as a promising material for bone scaffolds, whereas the more bioinert ZrO2 is known for its excellent mechanical properties. Ultra-porous TiO2 scaffolds (>89% porosity) were produced using polymer sponge replication with 0–40 wt.% of the TiO2 raw material substituted with ZrO2. Microstructure, chemical composition, and pore architectural features of the prepared ceramic foams were characterised and related to their mechanical strength. Addition of 1 wt.% of ZrO2 led to 16% increase in the mean compressive strength without significant changes in the pore architectural parameters of TiO2 scaffolds. Further ZrO2 additions resulted in reduction of compressive strength in comparison to containing no ZrO2. The appearance of zirconium titanate (ZrTiO4) phase was found to hinder the densification of the ceramic material during sintering resulting in poor intergranular connections and thus significantly reducing the compressive strength of the highly porous ceramic foam scaffolds. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Biocompatible ceramic materials have recently attracted increasing interest as porous scaffolds that stimulate and guide the natural bone regeneration in the repair of non-healing, or critical size, bone defects [1–3]. In order to provide optimal conditions for tissue regeneration, the scaffold structure must allow cell attachment onto its surface as well as sufficient space for cell proliferation and unobstructed tissue ingrowth [4]. Therefore, structural properties, such as porosity and pore morphology, of the 3D bone scaffold construct play a crucial role in the success of scaffold-based bone regeneration as has been emphasised in several studies done both in vitro and in vivo [5–8]. A bone scaffold is required to have a well-interconnected pore network with large pore volume and an average pore connection size exceeding 100 μm in order to ensure viable cell infiltration and attachment, nutrient and waste product transportation, vascularisation, and passage of newly formed bone tissue through the entire scaffold volume [9–13]. Together with the interconnectivity, appropriate pore morphology and average pore size larger 300 μm are also necessary to provide adequate space and permeability for viable bone formation in a nonresorbable scaffold structure [5]. In addition, the scaffold material itself must of course be highly biocompatible with favourable surface properties for bone cell attachment and differentiation in order to allow the formation of direct bone-to-scaffold interface [4,11].

⁎ Corresponding author. Tel.: + 47 22 85 23 50; fax: + 47 22 85 23 51. E-mail address: [email protected] (H.J. Haugen). 0928-4931/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.msec.2012.04.014

Due to the excellent biocompatibility of titanium dioxide (TiO2), porous three-dimensional TiO2 constructs have been proposed as promising scaffolding material for inducing bone formation from the surrounding tissue in the restoration of large bone defects [14]. Recent studies have shown the fabrication of highly porous ceramic TiO2 foams with pore architectural properties well-matched for those required from a bone scaffold and the capacity to promote adhesion and proliferation of osteoblasts and human mesenchymal stem cells (hMSC) on the entire scaffold surface in vitro [15–17]. However, increased porosity and pore size are known to have a detrimental effect on mechanical strength and consequently reduce the structural integrity of a scaffold. Compressive strength values of approximately 2.5 MPa were reported for the novel TiO2 scaffolds at overall porosity of ~85% [16], while the compressive strength of trabecular bone is typically 2–12 MPa [18]. Since the scaffold structure is also required to provide mechanical stability for the defect site, the use of the TiO2 scaffolds in load-bearing environment is somewhat limited due to the relatively low mechanical properties of the highly porous TiO2 foam structure. Thus, scaffolds with enhanced mechanical performance are needed in order to provide adequate structural support during the healing of large segmental bone defects in sites where heavy local loading can be expected. Zirconium dioxide (ZrO2) is a biocompatible ceramic material generally known for its excellent mechanical strength and toughness [19,20]. The production of porous ZrO2 and ZrO2 toughened aluminium oxide (Al2O3; ZTA) bone scaffolds with compressive strength values over 20 MPa at porosities above 70% has previously been reported [21,22]. However, both of these materials are considered bioinert at

H. Tiainen et al. / Materials Science and Engineering C 32 (2012) 1386–1393

physiological conditions, and therefore surface modification treatments, such as calcium phosphate coatings, are necessary to produce more bioactive scaffold surface that permits direct contact between the scaffold surface and bone tissue without fibrous encapsulation [21,23]. Crystalline TiO2, on the other hand, has been shown to have bioactive properties as its surface can support chemical bonding to bone via the formation of a bone-like apatite layer [24–26], but it lacks the good fracture strength and toughness of ZrO2. Therefore, combining the favourable osteogenic properties of TiO2 with the excellent mechanical properties of ZrO2 may allow the fabrication of highly osteoconductive bone scaffolds with sufficient mechanical strength for load-bearing environment, while still maintaining the highly reticulated pore structure required for viable bone regeneration. Hence, the aim of the present study was to investigate the effect of ZrO2 addition on the mechanical properties of porous TiO2 scaffolds. Ceramic TiO2 foams were fabricated by polymer sponge replication with 0–40 wt.% of the TiO2 substituted with ZrO2. The microstructure, chemical composition, and pore architectural features of the prepared TiO2–ZrO2 composite scaffolds were analysed and related to compressive strength in order to establish the relationship between the amount of added ZrO2 and the mechanical strength of the porous scaffold structure.

2. Materials and methods

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as for the ceramic foam samples, and the discs were used for EDX and XRD analysis. 2.2. Scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDX) The initial visualisation and optical observation of the microstructure of the prepared scaffolds was performed using a scanning electron microscope (TM-1000, Hitachi High-Technologies, Japan). The samples were mounted on aluminium stubs with conductive carbon tape and viewed with backscattered electrons at 15 kV accelerating voltage. Some high magnification images were taken using scanning electron microscope FEI Quanta™ 200 equipped with field emission gun (FEI Company, Hillsboro, Oregon, USA). Environmental SEM mode was used in combination with 15 kV accelerating voltage. In addition, EDX analysis (energy-dispersive X-ray spectroscopy) was performed using EDX detector (EDAX, Mahwah, New Jersey, USA) (take-off angle: 37°, voltage: 15 kV). 2.3. X-ray diffraction (XRD) The crystal structures of prepared sample discs were examined by X-ray diffraction (XRD), performed with D8 Discover powder diffractometer with a Ge (111) monochromator providing CuKα1 radiation and LynxEye detector (Bruker, Karlsruhe, Germany). The acquisition was done using θ–2θ configuration.

2.1. Sample preparation 2.4. Micro-computed tomography (micro-CT) Polymer sponge replication was used to produce the reticulated ceramic foam scaffolds. Ceramic slurry was prepared by gradual addition of 65 g of ceramic powder in total to 25 ml of sterilised water. In order to avoid coagulation and to control the viscosity, the pH of the slurry was kept at 1.5 for the entire duration of stirring with small additions of 1 M HCl. The dispersed powder consisted of either pure TiO2 (Kronos 1171, Kronos Titan GmbH, Leverkusen, Germany; cleaned with 1 M NaOH prior to use) or TiO2 with 1–40 wt.% of total amount of powder replaced by ZrO2 (Z-611, ABSCO Materials, Haverhill, Suffolk, UK). To some batches containing 30 and 40% ZrO2, 3 mol% of Y2O3 (Sigma-Aldrich Chemie GmbH, Steinheim, Germany) was added based on the amount of ZrO2. After dispersing the powder in water, stirring was continued for 2.5 h at 5000 rpm (Dispermat Ca-40, VMAGetzmann GmbH, Reichshof, Germany). The process parameters have been extensively studied and therefore only summarised [16]. Cylindrical polyurethane foam templates (60 ppi, Bulbren S, Eurofoam GmbH, Wiesbaden, Germany), 10 mm in both diameter and height, were coated with the prepared slurry. Excess slurry was squeezed out of the foam templates between two polymer foam sheets. The samples were then placed on a porous ceramic plate and allowed to dry at room temperature for at least 16 h before sintering. For the burnout of the polymer, the scaffolds were slowly heated to 450 °C at a heating rate of 0.5 °C/min. After 1 h holding time at 450 °C, the temperature was raised to 1500 °C at a rate of 3 °C/min and the sintering time at this temperature was set to 40 h (HTC-08/16, Nabertherm GmbH, Lilienthal, Germany). The sintered scaffolds were then cooled back to room temperature at the cooling rate of 5 °C/min. For some scaffolds batches, cooling rate was reduced to 0.5 °C/min between 1300 and 800 °C. Furthermore, some scaffolds have been sintered for additional 10 h at 1600 °C. In addition, discs with diameter of 16 mm and height of 5 mm were prepared by compressing the powder in a custom made tool using a vertical force of 40 kN (YK10BM, Sealy Power Tools, Suffolk, UK). The compressed discs consisted of TiO2 including 5, 30 and 40 wt.% ZrO2. Samples with 30 wt.% ZrO2 were also prepared with added Y2O3. The powders were mixed well together prior to compression in order to achieve even distribution of TiO2 and ZrO2 in the compressed green bodies. Similar sintering schedule was applied

Micro-computed tomography was used to determine the threedimensional microstructure of the scaffolds. The samples were mounted on a plastic sample holder and scanned with desktop 1172 micro-CT imaging system (SkyScan, Aartselaar, Belgium) at 6 μm voxel resolution using source voltage of 100 kV and current of 100 μA with 0.5 mm aluminium filter. The samples were rotated 180° around their vertical axis and three absorption images were recorded every 0.4° of rotation. These raw images of the samples were reconstructed with the standard SkyScan reconstruction software (NRecon) to serial coronal-oriented tomograms using 3D cone beam reconstruction algorithm. For the reconstruction, beam hardening was set to 20% and ring artefact reduction to 12. The image analysis of the reconstructed axial bitmap images was performed using the standard SkyScan software (CTan and CTvol) and included thresholding and despeckling (removing objects smaller than 500 voxels and not connected to the 3D body). In order to eliminate potential edge effects, a cylindrical volume of interest (VOI) with a diameter of 5 mm and a height of 2.5 mm was selected in the centre of the scaffold. The porosity was then calculated as 100%—vol.% of binarised object in the VOI. 2.5. Compressive strength The mechanical strength was investigated in a compressive test (Zwicki, ZwickRoell, Ulm, Germany). The compression tests were performed in accordance with DIN EN ISO 3386 at room temperature using a load cell of 1 kN with preloading force set to be 0.5 N. The scaffolds were compressed along their long axes at a compression speed of 100 mm/min until failure. The force and displacement were recorded throughout the compression and converted to stress and strain based on the initial scaffold dimensions. 2.6. Statistical analysis Normality and equal variance tests were performed prior to further statistical testing. When the datasets were found normally distributed, statistical comparison of different data groups was performed using Student's t-test or one-way analysis of variance (ANOVA) test followed

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by post hoc tests for pairwise comparisons performed using Holm–Sidak or Tukey test. The datasets that failed normality or equal variance test were analysed using non-parametric Mann–Whitney U test or Kruskal–Wallis one-way ANOVA with multiple comparisons performed using Dunn's test. Statistical significance was considered at a probability p b 0.05 and n = 10 unless otherwise specified. A correlation study was performed with a bivariate regression analysis, Spearman Rank Order correlation. The results were interpreted as follows: small correlation if 0.1 b |ρ| b 0.3; medium correlation if 0.3 b |ρ| b 0.5; strong correlation if 0.5 b |ρ| b 1 and p b 0.05.[27] A negative ρ indicated a negative correlation, whereas a positive ρ indicated a positive correlation (ρ = Spearman's rank correlation coefficient). All statistical analysis was performed using software SigmaPlot 11 (Systat Software Inc, San Jose, USA). 3. Results 3.1. Microstructure and chemical composition Typical appearances of the microstructure of the prepared ceramic foams are shown in Figs. 1 and 2. Pure TiO2 scaffolds did not differ markedly in appearance from scaffolds containing 1, 5, and 10 wt.% ZrO2 or pure TiO2 scaffolds with varied heating schedule. All scaffolds prepared with b 20 wt.% ZrO2 were characterised by single-phase microstructure consisting of rutile (Figs. 1 and 3), while the addition of ZrO2 was observed to shift the lattice parameters of rutile TiO2 towards a somewhat larger unit cell as the larger Zr 4 +-cations dissolved in the TiO2 lattice. The densification caused by the sintering process resulted in a relatively large overall grain size, and the grains were generally well-integrated via uniform grain boundary regions. However, as the amount of ZrO2 increased, the structure of some grain boundaries was visibly altered (Fig. 4A) and more porosity was observed at triple junctions. Furthermore, some transgranular cracks were also observed in some of the struts of scaffolds containing 5 and 10 wt.% ZrO2. In samples containing ≥20 wt.% ZrO2, a second phase became apparent as can be seen in Fig. 2. The two-phase grain structure consisted mainly of Zr-dissolved rutile and zirconium titanate (ZrTiO4), while the presence of srilankite (Ti2ZrO6) could not be ruled out (Fig. 3). EDX analysis on the microstructure visualised with SEM revealed

lower percentage of zirconium in the darkest grains in comparison to paler grains, and therefore this phase was identified as rutile. In addition, the XRD patterns indicated that a small amount of residual ZrO2 was present in these samples. The relic ZrO2 phases were characterised as monoclinic ZrO2 for samples containing no Y2O3, while its presence resulted in yttrium doped-ZrO2 phase which was clearly visible in samples containing Y2O3 and 40 wt.% ZrO2 (Fig. 2D). The small ZrTiO4 grains became visible on the surface of the struts of the scaffolds with 20 wt.% (Fig. 2A). These grains were randomly distributed over the whole scaffold material and several of the superficial ZrTiO4 grains were discovered protruding out from the roughly planar surface of the Zr-dissolved rutile grains. The grain size of the rutile phase decreased dramatically, while some of the ZrTiO4 grains grew larger as the ZrO2 concentration was increased to 30 wt.%. (Fig. 2B). In addition, majority of the grains had a slightly bulbous surface. Moreover, the grains were poorly interconnected as there were some gaps at the grain boundaries and porosity at triple junction increased in comparison the samples containing less ZrO2. When the cooling rate between 1300 °C and 800 °C was lowered, the ZrTiO4 grains became more irregularly shaped and even more disconnected from the rutile grains (Fig. 5A). In addition, more microcracks were discovered on the scaffold surface after the altered cooling rate was applied. As the amount of added ZrO2 was further increased to 40 wt.%, the ZrTiO4 became more prominent of the two phases. The overall structure can be described as similar to the structure of scaffolds containing 30 wt.% ZrO2, although the ZrTiO4 grains increased in size whereas the Zr-dissolved rutile grains were clearly somewhat smaller in diameter (Fig. 2C). The addition Y2O3 had a distinct effect on the microstructure of the prepared composite scaffolds (Figs. 2D and 4C). In samples containing ≥30 wt.% ZrO2 with Y2O3 addition, three separate phases were clearly visible. Grains of the third phase, which consisted of yttriumdoped ZrO2, became considerably more frequent and larger in size as the ZrO2 concentration was increased to 40 wt.%. Multitude of transgranular microcracks intersecting all types of phases present on the scaffold surface was also observed on the surfaces of both scaffold batches. Furthermore, the intergranular porosity was further increased from that of similar scaffold batches with no added Y2O3 as manifested by numerous large gaps at triple junctions and grain boundaries. A relatively large number of poorly reacted residual grains were also visible on the surface of scaffolds containing Y2O3, particularly at grain

Fig. 1. SEM images showing the microstructure of scaffolds containing 0 (A), 1 (B), 5 (C) and 10 wt.% (D) ZrO2. Only one single phase is visible.

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Fig. 2. SEM images showing the microstructure of scaffolds containing 20 (A), 30 (B) and 40 wt.% (C) ZrO2. Additionally, small amount of Y2O3 was added to scaffolds containing 40 wt.% ZrO2 (D). Two distinct phases can be visibly identified. In (D) also a third phase is present.

boundary regions, and these flat and irregular grains were frequently protruded out of the surface plane. Similar residual grains were also found on samples with no added Y2O3 but were considerably fewer in number (Fig. 4B). Scaffolds containing 0, 5, 20 and 40 wt.% ZrO2 were sintered for additional 10 h at 1600 °C in order to investigate the influence of higher sintering temperature on the densification of the ceramic structure (n = 5 for each batch). No significant changes were observed in the microstructure of samples that were composed of only rutile phase (0 and 5 wt.% ZrO2) while the microscopic appearance of scaffolds containing 20 and 40 wt.% ZrO2 was visibly altered as can be seen in shown in Fig. 5. In comparison to scaffolds manufactured by normal sintering parameters, the grain size of rutile phase increased markedly in samples containing 20 wt.% ZrO2, and large portions of the scaffold surface had no visible ZrTiO4 grains. Those ZrTiO4 grains that were observed displayed a very angular appearance, and the rutile grains in close contact with ZrTiO4 grains were furrowed at the phase boundaries and poor contact between the two phases was evident. Furthermore, the grain boundaries of rutile grains adjacent to the ZrTiO4 grains were highly irregular as shown in Fig. 5B and similar structures that were seen in Fig. 4A were observed at the grain boundaries in areas where no ZrTiO4 grains were visible. Regarding the scaffolds containing 40% ZrO2, the ZrTiO4 phase increased in amount and grain size while the overall appearance remained very similar to that of scaffolds that underwent normal heating cycle.

3.2. Pore structure Micro-CT analysis was performed in order to gather information on pore morphology and architecture of the scaffolds. The macroscopic structure of all produced scaffolds was a reticulated, fully open foam structure with open porosity constituting over 99.9% of the total porosity. The pore volume of the foams was characterised by spherical macropores of approximately 400 μm and struts with average diameter of ~ 60 μm which together formed a highly interconnected pore structure. The characteristic pore architectural features of several scaffold batches are given in Table 1.

The average porosity of the prepared scaffolds range from 89 to 92.8% with tendency towards higher porosity with increasing ZrO2 concentration and medium correlation was found the amount of ZrO2 and porosity (ρ = 0.359, n = 69). However, the ZrO2 addition was not found to affect the overall strut or pore size significantly (p > 0.05), whereas the addition of Y2O3 into the TiO2–ZrO2 system strongly correlated with decreasing porosity and pore size (ρ = −0.689 and ρ = −0.819 respectively, n = 39) while medium correlation was found for increasing strut size (ρ = 0.385). The added Y2O3 also caused visibly more shrinkage in the overall scaffold dimensions during sintering, particularly when the scaffolds contained 30 wt.% ZrO2. Lower cooling rate between 1300 and 800 °C resulted in reduced mean strut diameter and thus increased porosity in the case of ‘0% ZrO2’ samples, whereas such an effect was not seen for the twophase ‘30% ZrO2’ samples (Table 1). Additional sintering at 1600 °C was found to result in more shrinkage in all aspects apart from porosity of the porous scaffold structure in comparison to normal heating cycle. 3.3. Mechanical properties The compressive strength of the prepared scaffolds was found to range from a lowest value of 0.04 MPa to a highest value of 0.87 MPa, and the strength values of different scaffold batches were compared (Fig. 6). The compressive strength of the control group (0% ZrO2) was found to be significantly higher compared to all other groups apart from the batch containing 1 wt.% ZrO2 which had a significantly higher mean strength. The presence of Y2O3 significantly reduced the strength of scaffolds with 40 wt.% ZrO2 as the mean strength diminished from 0.336 MPa to 0.055 MPa when Y2O3 was introduced to the material composition. However, such a reduction in the compressive strength values was not detected when preparing scaffolds with 30 wt.% ZrO2 with and without added Y2O3 (0.176 vs. 0.157 MPa respectively, p = 0.077), although the power of the statistical test was below the desired power of 0.8. Applying lower cooling rate between 1300 °C and 800 °C led to significantly lower values of compressive strength for scaffolds containing 30 wt.% ZrO2 and therefore consisting of rutile and ZrTiO4 phases (p b 0.001), while the

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and ZrTiO4. No correlation was found between compressive strength and porosity for the two-phase scaffolds or when all batches were included in the analysis, but a medium correlation was detected for single-phase scaffolds with higher porosity resulting in lower strength. The compressive strength of scaffolds containing b20 wt.% ZrO2 is also plotted against porosity in Fig. 6B. 4. Discussion

Fig. 3. XRD pattern of the samples with various amounts of ZrO2 and Y2O3. The samples containing b 20 wt.% ZrO2 contained only rutile phase while ZrO2 addition led to multiphase composition with rutile and ZrTiO4 as well as some relic ZrO2 phases. + = ZrTiO4, × = ZrO2 (baddeleyite), * = Y0.15Zr 0.85O1.93, unlabelled peaks = rutile.

strength of the scaffolds containing no ZrO2 was not found significantly altered (p = 0.449). The additional high temperature sintering at 1600 °C significantly increased the strength of the scaffolds from the ‘20% ZrO2’ batch (p = 0.002), whereas no significant changes in the compressive strength values of the three other batches tested (0, 5, and 40 wt.% ZrO2) were detected after additional sintering at 1600 °C. Compressive strength was correlated with porosity, strut size, and the amount of ZrO2 and Y2O3 (Table 2). Overall, strong correlation was found between strut size and compressive strength with smaller strut sizes resulting in higher strength values. Strong correlation was also found between the compressive strength and the added amount of ZrO2 as well as between the compressive strength and the amount of Y2O3, with smaller amounts resulting in higher compressive strength in both cases. The correlation study was also performed separately for the single (b20 wt.% ZrO2) and two-phase (≥20 wt.% ZrO2) batches. The amount of ZrO2 strongly correlated with the compressive strength of the single-phase scaffolds but was not found to affect significantly the strength of the scaffolds composed of rutile

The purpose of this study was to evaluate the mechanical properties of TiO2–ZrO2 ceramic foams in attempt to produce highly porous bone scaffolds with improved mechanical strength, while still maintaining the desired pore architectural features necessary for unobstructed bone tissue ingrowth. Prior studies have reported the fabrication of ultra-porous TiO2 foams that display well-interconnected pore volume and appropriate pore size, but with an average compressive strength of 2.5 MPa, these scaffolds lack the structural integrity required in restoration of large segmental bone defects [16]. Since the porosity and accessible pore volume of the scaffold structure should not be significantly reduced, improvement in the mechanical properties of the scaffold material, particularly its fracture toughness, is necessary in order to produce bone scaffolds capable of sufficient load-bearing for orthopaedic applications. Al2O3 and ZrO2 are structural ceramics that are frequently used in dental and orthopaedic application, such as dental crowns and femoral heads in total hip prostheses, due to their chemical inertness, high wear resistance, and good mechanical strength [28]. Ceramic TiO2, on the other hand, has been shown to have excellent biocompatibility, especially in contact with bone tissue, but as a more brittle material would benefit from the higher strength and fracture toughness of Al2O3 and ZrO2, both of which have previously been used in reinforcement of biocompatible ceramics [22,29]. However, the formation of an aluminium titanate (Al2TiO5) phase at temperatures above 1280 °C restricts the use of Al2O3 in strengthening the porous TiO2 scaffold structure [30]. The anisotropic thermal expansion of Al2TiO5 leads to expansion in the a and b directions of the unit cell, while contraction occurs along the c axis resulting in severe cracking as the extremely high crystallographic anisotropy in thermal expansion causes thermal stresses in the ceramic structure upon cooling [31,32]. In a previous study, long holding time at sintering temperature of at least 1500 °C was shown to result in partial elimination of the residual triangular hollow space within the scaffold struts, caused by the burnout of the polymer template, and thus significantly improving the mechanical strength of the scaffold structure [15]. Therefore, the appearance of the Al2TiO5 phase could not be suppressed during the sintering of TiO2–Al2O3 composite scaffolds. Such extreme anisotropy in thermal expansion has not been reported for zirconium titanate (ZrTiO4) which is the intermediate compound in the binary TiO2–ZrO2 system. Zirconium titanate is wellknown compound in the field of electroceramics as a thermally stable dielectric material for ceramic capacitors and as an oscillator at microwave frequencies [33], and the phase evolution and equilibrium phases in the TiO2–ZrO2 system have been widely researched [34–37]. Furthermore, the influence of TiO2 addition on the stabilisation of tetragonal ZrO2 has also been the subject of much experimental research [38–40]. However, very little information regarding the effect of ZrO2 addition on the mechanical strength of TiO2 matrix has, to the authors' knowledge, been previously reported in the relevant literature [41]. As predicted by the phase diagram for the TiO2–ZrO2 system [34], ZrO2 additions b20 wt.% in ceramic slurry used for the production of the porous TiO2 scaffolds resulted in single-phase ceramic foams consisting of Zr-dissolved rutile TiO2, whereas scaffolds containing higher concentrations ZrO2 exhibited two-phase microstructure of ZrTiO4 and Zr-dissolved rutile (Figs. 1–3). At temperatures above 1400 °C, the rutile solid solution phase can dissolve up to 10–20 mol% of ZrO2 with increasing ZrO2 concentrations leading to progressive

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Fig. 4. ZrO2 addition resulted in reduced densification of the ceramic material during sintering. Poor intergranular connections seen on samples containing 10 wt.% (A) and 40 wt.% (B) ZrO2. The added Y2O3 reduced the quality of grain boundaries even further and the appearance of a Zr-rich phase resulted in several transgranular microcracks in the foam struts (C). Fracture surface shown in (D) showcases the local discontinuities at the phase boundaries.

increase in the length of both a and c axes of the tetragonal rutile lattice as the larger Zr 4 +-cations occupy some of the octahedral Ti 4 +-cation sites [34,42]. After the solubility limit was reached, an intermediate ZrTiO4 solid solution phase was formed, and the amount and grain size of the ZrTiO4 phase was observed to increase with increasing ZrO2 concentration as can be seen in Fig. 2. Incipient decomposition of the formed intermediate ZrTiO4 phase back to its constituent phases may have occurred during cooling. According to X-ray diffraction profile of the prepared samples, small amount of monoclinic ZrO2 was present in both the single- and two-phase ceramics. However, the samples used for the XRD analysis were prepared by compressing dry powders into discs and subsequently sintered at 1500 °C for 40 h, and therefore the ZrO2 and TiO2 particles were more agglomerated and less homogenously dispersed in the sample volume. Thus, some relic ZrO2 grains may remain in the compressed sample discs and give rise to the ZrO2 peaks detected in the X-ray diffraction profile (Fig. 3). Such residual ZrO2 grains are less likely to exist in the scaffold samples. Furthermore, the precipitation of Ti2ZrO6 and the high to low temperature transformation of ZrTiO4 is very sluggish [35], but this exsolution of ZrO2 from the formed equilibrium phases at temperatures lower than the sintering temperature of 1500 °C could explain the small number of poorly reacted grains detected on the grain boundary regions of the two-phase samples, particularly at high ZrO2 concentrations (Fig. 4B), as well as the curious microstructures formed at some

grain boundaries in scaffolds containing 10 wt.% ZrO2. Although, these microstructural features are more likely to consist of Ti-rich (Zr,Ti)2O4 phase than the monoclinic ZrO2 indicated by the XRD analysis, and also to contain a small amount of siliceous impurities from the amorphous grain boundary phase that has been shown to be present at the superficial grain boundaries in the prepared TiO2 scaffolds [43]. The compressive strength of the prepared scaffolds was found to correlate strongly with the amount of added ZrO2 with increasing ZrO2 concentration leading to significantly reduced mechanical strength. This relationship was particularly noticeable in the singlephase scaffolds consisting of Zr-dissolved rutile. Although there was only slight variation in the pore architectural features of the different scaffolds groups, medium correlation was also found between the porosity and compressive strength of the single-phase scaffolds. However, as can be seen in Fig. 6B, where the strength values of the scaffolds are plotted against their porosity, the reduction in strength for scaffolds containing 5 and 10 wt.% cannot be explained merely by the change in porosity, particularly as the compressive strength of the scaffolds with 1 wt.% ZrO2 increased in comparison to the control scaffolds with no added ZrO2 despite the higher mean porosity. Due to the larger ionic radii of the substitutional Zr 4 +-cations, the rutile lattice is dilated giving rise to an isotropic stress field around the solute ion. The altered lattice parameters may also induce residual thermal stresses in the scaffold struts upon cooling. These stress fields

Fig. 5. Grain structure of scaffolds with 30 wt.% ZrO2 and diminished cooling rate of 0.5 °C/min between 1300 and 800 °C (A), 20 wt.% ZrO2 and higher sintering temperature of 1600 °C (B) and 40 wt.% ZrO2 also sintered at 1600 °C (C).

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Table 1 Selected pore architectural parameters of the prepared scaffolds (mean ± SD). *significant difference in comparison to 0% ZrO2,^significant difference in comparison to corresponding batch without Y2O3 (p b 0.05), n = 10, h = lower cooling rate, Y = Y2O3, # per 50 mm3 of scaffold (solid volume= 4.2 ± 0.8 mm3). Batch

0%ZrO2 0%ZrO2 h 1%ZrO2 5%ZrO2 10%ZrO2 20%ZrO2 30%ZrO2 30%ZrO2 h 30%ZrO2 + Y 40%ZrO2 40%ZrO2 + Y

Porosity

Pore size

Strut size

Surface area#

%

μm

μm

mm2

90.9 ± 0.8 92.8 ± 0.9* 91.7 ± 0.9 91.9 ± 0.9 92.1 ± 1.0 92.5 ± 1.4* 92.2 ± 1.1 92.0 ± 1.1 91.4 ± 1.0 92.5 ± 1.7* 89.0 ± 1.6*,^

453 ± 17 448 ± 6 456 ± 7 413 ± 15* 441 ± 25 466 ± 8 438 ± 12 449 ± 5 369 ± 25*,^ 447 ± 29 367 ± 10*,^

66.9 ± 6.3 52.1 ± 2.3* 58.0 ± 4.7 56.5 ± 4.3 64.8 ± 11.7 64.9 ± 10.7 81.8 ± 16.8 85.8 ± 15.1 61.5 ± 5.3^ 62.7 ± 7.7 103.3 ± 18.7^

261 ± 14 245 ± 17 254 ± 14 262 ± 13 246 ± 15 225 ± 21* 230 ± 11* 229 ± 9* 282 ± 22^ 244 ± 36 286 ± 12^

may then interact with the compressive stress field of an advancing crack and add to the fracture toughness of the material. However, with a 16% increase in the average compressive strength, only limited improvement in the mechanical properties of the scaffold structure were achieved by dissolution of the small amount of Zr4 +-cations in the TiO2 lattice. When the ZrO2 concentration was increased to 5 and 10 wt.%, the compressive strength reduced considerably and was lower than the strength of scaffolds containing no ZrO2 (Table 2). It appears that the stresses induced by the substitutional Zr4 +-cations and the increased lattice parameters led to cracks in the material as an increasing amount of microcracks and some fractured struts were

Fig. 6. (A) The relationship between compressive strength and the amount of ZrO2 added to the TiO2 slurry. *statistical significance against 0 wt.% ZrO2 (p b 0.05), n = 10 (B) Compressive strength plotted against porosity for scaffolds containing b 20 wt.% ZrO2.

Table 2 Correlation study between compressive strength, porosity, strut size and amount of ZrO2 and Y2O3. Results shown on the first row include all experimental groups with normal heating cycle (n = 90), whereas second and third row show correlation coefficients for groups containing b20 wt.% (a, n = 40) and ≥20 wt.% ZrO2 but without Y2O3 (b, n = 30), respectively. Small correlation if 0.1 b |ρ| b 0.3, medium correlation if 0.3 b |ρ| b 0.5, strong correlation if 0.5 b |ρ| b 1 when *p b 0.05 or **p b 0.01. Spearman's rank correlation coefficient

Porosity

Strut size

Amount of ZrO2

Amount of Y2O

Compressive strength Compressive strengtha Compressive strengthb

0.036 − 0.354* − 0.323

− 0.527** 0.237 − 0.158*

− 0.810** − 0.628** − 0.070

− 0.522** – –

observed in these scaffold groups, which compromise the structural integrity of the scaffold structure. In addition, intergranular microporosity and exsolution of ZrO2 from the rutile lattice during cooling increased in scaffolds with more dissolved ZrO2, especially in scaffolds with 10 wt.% ZrO2, causing alterations in the microstructure of the scaffold material and further reducing the compressive strength. The mechanical strength was further reduced due to the appearance of ZrTiO4 solid solution phase. Typically, two-phase ceramics are prone to microcracking due to thermal stresses caused by mismatch of the thermal expansion coefficients of the constituent phases, and thus resulting in low mechanical strength. However, despite the slight anisotropy in the thermal expansion of ZrTiO4 (αa298-1073K =8.0× 10− 6 K− 1, αb298-1073K =10.0×10− 6 K− 1, αc298-1073K =6.2×10− 6 K− 1, Pbcn setting) [44], the overall thermal coefficient of this compound is almost identical to that of rutile TiO2. Thermal stresses are therefore not expected to cause significant cracking of the ceramic composite and the observed microcracks in the prepared two-phase TiO2–ZrO2 scaffolds were very few in number. Furthermore, the lack of thermally induced microcracking was also shown for pure ZrTiO4 samples in a recent study by López-López et al. [45]. Therefore, thermal stresses are not considered to contribute significantly to the reduction in compressive strength of the scaffolds. For scaffolds composed of rutile and ZrTiO4, no correlation was found between the scaffold strength and pore architectural features. Since the denser ZrTiO4 phase is expected to increase the strength of scaffold structure, the reduction in compressive strength arises from the poor densification of the two-phase microstructure. The lowest average strength was recorded for the scaffolds containing 30 wt.% ZrO2 which also had the most inhomogeneous microstructure, the highest degree of intergranular porosity, and the largest strut diameter indicating impeded grain growth and densification during sintering. As the ZrO2 concentration was increased to 40 wt.%, ZrTiO4 became the more prominent phase allowing better densification and more homogeneous distribution of the two phases. The size and distribution of flaws, such as micropores and cracks, have a detrimental effect on strength of highly porous ceramic foams under compressive loading. As the foam structure is compressed, some of the thin ceramic struts are bent causing failure by brittle fracture when a critical loading force is reached. Therefore, the pre-existing flaws in the scaffold material can significantly reduce the structural integrity of the ceramic scaffold structure due to the local stress amplification at the flaw site. The formation of ZrTiO4 occurs at temperatures lower than that required for initiation of densification, thus impeding the full densification of the ceramic compact during reaction sintering of TiO2 and ZrO2 [45]. It has been reported that ZrTiO4 ceramics sintered in the absence of additional fluxes produce polycrystalline ceramics with low density, typically 60 to 63% of the theoretical density, whereas the addition of a flux, such as ZnO or CuO, can considerably reduce the residual microporosity [33]. However, the added sintering aids may compromise biocompatibility, especially if the liquid phase sintering results in amorphous or poorly crystallised grain boundary phase which may be susceptible to dissolution in physiological conditions. Slowing the cooling rate from 5 to 0.5 °C/min between 1300 and 800 °C further diminished the mechanical strength of the scaffolds

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containing ZrTiO4. The slower cooling rate allowed ordering in the disordered ZrTiO4 polymorph, and the stresses induced by the volume change associated with this ordering transition caused the formation of microcracks in the scaffold structure [46]. In addition, the ordering transition typically goes hand in hand with change in composition towards higher Ti-contents and the observed changes in microstructure are related to this exsolution of ZrO2 [34]. Additional sintering in 1600 °C resulted in increased grain size of rutile phase, except in scaffolds containing 40 wt.% ZrO2 in which the ZrTiO4 grains increased in size, and reduction in the overall scaffold dimensions. In general, sintering in 1600 °C had no significant effect on the mechanical strength of the scaffolds. However, the compressive strength of the scaffolds containing 20 wt.% increased due to this additional sintering. Since the higher sintering temperature allowed higher ZrO2 concentration to dissolve in the rutile lattice, only a small number of faceted ZrTiO4 grains appeared on the sample surface. The reduction in the number of ZrTiO4 grains and their better integration in the rutile matrix improved the mechanical strength due to diminished size and distribution of flaws in the scaffold material. Adding a small amount of Y2O3 in the TiO2-ZrO2 scaffolds resulted in extensive microcracking in the scaffold material and led to dramatic decrease in the compressive strength. Zr-rich phase, which is considered to consist of Ti4 + and Y 3 +-dissolved ZrO2, was observed to appear in the scaffold microstructure after Y2O3 addition, although this phase was not particularly prominent in samples containing 30 wt.% ZrO2. Since the reduction in strength and the number of transgranular microcracks were considerably more pronounced in scaffolds with 40 wt.% of ZrO2, it seems evident that the appearance of the Zr-rich phase is responsible for the severe microcracking of the samples containing small amount Y2O3. Furthermore, the formation of very small quantities of Y2Ti2O7 cannot be completely ruled out [47], although such compound was not detected in the XRD analysis. The numerous protruding superficial grains are likely to have been formed due to reduced reaction rate between ZrO2 and TiO2 during sintering and exsolution of ZrO2 during cooling, indicating that the added Y2O3 had an adverse effect on the stability of the equilibrium phases at 1500 °C. It has also been suggested that Y2O3 increases the ordering rate of the disordered ZrTiO4 phase [46]. The volume change associated with the ordering transition may also contribute the appearance of microcracks in scaffolds containing yttrium. In addition, the significant increase in strut size indicates that the densification during sintering was further impeded, which was also observed as the strong correlation between strut size and compressive strength. As discussed above, the poor densification during sintering increased the amount and size of internal and superficial flaws in the scaffold material, which together with the numerous pre-existing microcracks have an exceedingly detrimental effect on the mechanical strength of the prepared scaffolds. 5. Conclusions The purpose of this study was to investigate effect of ZrO2 addition on the mechanical properties of ultra-porous TiO2 ceramic foams in attempt to produce highly porous bone scaffolds with improved mechanical strength, while still maintaining the desired pore architectural features necessary for unobstructed bone tissue ingrowth. Overall, the added ZrO2 did not significantly affect the desired pore architectural features of the TiO2 scaffolds. However, dissolution of ZrO2 in the TiO2 lattice and the formation of the intermediate ZrTiO4 phase prevent the use of ZrO2 particles to toughen the scaffold material by crack deflection or by transformation toughening mechanism, where the metastable tetragonal ZrO2 particles transform to monoclinic crystal structure. Moreover, the ZrO2 addition was generally found to have an adverse effect on the mechanical properties of the scaffold structure as only the lowest ZrO2 concentration (1 wt.%) resulted in improved strength with a 16% increase in the average

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