Mechanical properties and anisotropy in hydroxyapatite single crystals

Mechanical properties and anisotropy in hydroxyapatite single crystals

Scripta Materialia 57 (2007) 361–364 www.elsevier.com/locate/scriptamat Mechanical properties and anisotropy in hydroxyapatite single crystals B. Vis...

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Scripta Materialia 57 (2007) 361–364 www.elsevier.com/locate/scriptamat

Mechanical properties and anisotropy in hydroxyapatite single crystals B. Viswanath,a R. Raghavan,b U. Ramamurtyb and N. Ravishankara,* a

b

Materials Research Centre, Indian Institute of Science, Bangalore 560 012, India Department of Materials Engineering, Indian Institute of Science, Bangalore 560 012, India Received 27 March 2007; revised 5 April 2007; accepted 7 April 2007 Available online 24 May 2007

The mechanical behavior of single crystalline hydroxyapatite is examined through instrumented nano- and microindentation experiments on prism and basal planes. The results indicate that the c-axis orientation is stiff and hard, and significant anisotropy in properties was found. Nanoscale plasticity is observed in the form of pile-up along the nanoindentation edges and manifests as discrete displacement bursts in the load vs. depth of penetration curves. Toughness also shows significant anisotropy and is lower compared with the values reported for the polycrystalline hydroxyapatite. Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Anisotropy; Hardness; Elastic behavior; Toughness; Slip

The basic structural building units of all mineralized tissues, such as bone, are an organic matrix, which is primarily collagen, and hydroxyapatite (HA), an inorganic mineral, in an approximately 1:1 volumetric ratio. These units are organized into a hierarchy of structures, ranging from the nanometer scale to millimeters, which imparts high strength and toughness to the bone [1–5]. At the finer end of the hierarchical spectrum, the selfassembled collagen fibrils are impregnated with plateshaped apatite nanocrystals of HA that are 2–3 nm in thickness and tens of nanometers in length and width. Thus, in the context of understanding the mechanical behavior of the bone, it is imperative to have a good measure of the mechanical properties of the single crystal form of HA. However, conventional mechanical testing procedures require crystals of the order of centimeters, which are very difficult to grow. In the absence of direct experimental data, the deformation models of bone either assume the mechanical properties of single crystal HA or infer from measurements on polycrystalline samples or crystals with related structures [6,7]. In the recent past, instrumented indentation methods have proven to be an excellent means to probe mechanical properties of small volumes of materials and have been applied previously to study, for example, the spatial gradation in the mechanical properties of human tooth den* Corresponding author. E-mail: [email protected]

tine [8]. In this paper, we employ both nano- and microindentation techniques to probe the mechanical properties and the attendant anisotropy of flux-grown HA single crystals. HA single crystals were grown by the molten salt synthesis method [9] using NaCl as the flux. Calcined HA powder prepared by wet-chemical means was mixed with NaCl and the mixture was heated to 900 °C for 2 h followed by furnace cooling. Single crystals in the form of hexagonal prisms (60–70 lm long and 20– 30 lm in diameter) were separated from the flux by repeated washing with hot distilled water and drying at 100 °C for 12 h. The as-grown crystals, characterized by X-ray diffraction and Fourier transform infrared spectroscopy, are of phase-pure HA, predominantly bounded by the prism and the basal planes. By controlling the growth conditions, it was possible to change the aspect ratio of the crystals. For indentation experiments, the crystals were embedded in a matrix of tricalcium phosphate (TCP), another member of the apatite family, and sintered at 1200 °C for 3 h [10]. Larger aspect ratio crystals always presented the prism plane on the surface and hence lower aspect ratio crystals were used for the indentation experiments so that both the prism and basal planes could be indented. Nanoindentation experiments on both the basal and prism facets of the HA single crystals were conducted using a Berkovich diamond indenter (Hysitron Triboscope). A maximum load, Pmax, of up to 10 mN was used. In all cases, the

1359-6462/$ - see front matter Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2007.04.027

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rates of loading and unloading used were such that the loading and unloading time was 5 s with a pause at Pmax of 5 s. Load, P, vs. depth of penetration, h, were recorded and analyzed to extract the elastic modulus, E, and the hardness, H, using the Oliver and Pharr method [11]. Images of the indents were captured using the same indenter in the scanning probe microscope (SPM) mode. The microindentation experiments were performed using an instrumented microindenter (MHT, CSM Instruments) with a Vickers diamond tip. The loading and unloading rates were maintained at 1.67 mN s1 and a maximum load of 100 mN was used with a hold time at peak load of 10 s. Invariably, this load led to cracking of the indentation corners. The crack lengths were measured using a scanning electron microscope (SEM) and the fracture toughness, Kc, of both the prism and basal facets was estimated using the Evans and Charles method [12]. Figure 1a shows typical P–h curves obtained through nanoindentation on prism and basal planes. In all cases,

Figure 1. (a) P–h curves obtained from nanoindentation on the basal and prism facets, Pmax = 2 mN. The inset is a SPM image of an indent showing the pile-up around the residual impression, Pmax = 10 mN. (b) Line profile across an indent showing the formation of pile-up at the edges (across line shown in the SPM image).

discrete displacement bursts (pop-ins) each with a magnitude of 6–8 nm were seen. In both orientations, the first pop-ins were always observed to occur at a P of 500 ± 100 lN. Further pop-ins were observed upon continued increase of P. Initially, these were thought be a result of corner cracking. However, the SPM images (typical image shown in the inset of Fig. 1a) did not show any cracks. Instead, flow of the material along the edges of indentation impressions, referred to as pile-up in indentation literature, was noted. Figure 1b shows a line profile across an indent (dashed line in the inset of Fig. 1a) showing pile-up along the edge of the indent. Kooi et al. [13], who examined the nanoindentation response of the basal planes of Ti3SiC2, observed similar pop-ins in the P–h curves. With the aid of detailed transmission electron microscopy, they showed that these pop-ins are due to delamination of the basal plane – due to the relatively weaker bonding between these planes in this particular ceramic – and subsequent outof-plane kinking of the planes. However, in the present context, the possibility that the pop-ins are a result of such delaminations is unlikely as the crystal structure of HA is not of the layered type, as in Ti3SiC2. Further, chipping of the edges is also ruled out since it usually occurs during unloading, driven by the indentation induced residual stress, and no displacement bursts or serrations were observed in the unloading part of the P–h curves. Discrete displacement bursts were reported earlier during the nanoindentation of single crystal Al films, which were considered as intermittent microplastic events as a result of dislocation slip [14], and in amorphous alloys due to the nucleation and propagation of shear bands underneath the indenter [15]. Polycrystalline HA is known to exhibit plasticity at higher temperatures [16], but no plasticity has been reported at room temperature for single crystal HA. Plastic deformation is observed in fluorapatite single crystals above 1100 °C, with f1 0 1 0gh0 0 0 1i being the operative slip system [17]. Despite the high resistance to dislocation motion in ceramics, similar pop-ins in the P–h curves have been observed in sapphire, aragonite, GaAs and ZnO single crystals at room temperatures [18–20]. Therefore, the faceted pile-up region in the SPM image suggests that the observed pop-ins are indeed due to dislocation plasticity. Values of the E and H obtained through the nanoindentation experiments are listed in Table 1. Here, data obtained from two different experiments, one with a Pmax of 2 mN and the other where the sample is unloaded before the first pop-in, are listed. The latter’s unloading P–h curve overlaps that obtained during loading, again supporting that the pop-ins are associated with some plastic events. The following observations can be made from the data listed in Table 1. (i) Both H and E measured with the indentation normal to the basal plane are higher (by 36% and 22%, respectively) vis-a`-vis the prism plane. (ii) Both the E and H values are smaller for those experiments wherein the Pmax exceeds the first pop-in load. This is because the plastic flow of the material being indented leads to the pile-up of material against the indenter faces, thus increasing the effective contact area and in turn reducing

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Table 1. Mechanical properties of HA whiskers obtained from nanoindentation Pmax

Prism E (GPa)

2 mN Below pop-in Kc (MPa m0.5)

Basal H (GPa)

125.9 ± 1.6 8.8 ± 0.4 127.2 ± 10.4 9.5 ± 1.5 0.65

E (GPa)

H (GPa)

135.1 ± 1.3 9.7 ± 0.1 162 ± 1.9 14.8 ± 0.1 0.40

the contact pressure. This can be ascertained by examining the H/E2 values as reported in the literature [21]. The differences in H/E2 values obtained with Pmax below and above pop-ups are small (<6%). This further reinforces the implication that the pop-ins are due to plastic flow; otherwise H/E2 would not be the same, as cracking makes the material significantly more compliant locally. Due to the complexity of the stress state underneath a sharp indenter, it is difficult to pinpoint the origins of anisotropy in E and H. The higher stiffness of the basal plane is possibly due to the close-packed nature of these planes. Hardness, which is a measure of the resistance offered by the material to plastic deformation, of the basal planes is high because the maximum shear stresses will be acting on non-basal planes on which it is difficult to initiate plastic deformation because they require higher critical resolved shear stress. In this context, it is worth noting that Nowak et al. [22] examined the anisotropy in hardness of hexagonal crystals and show that the amplitude of the anisotropy in hardness is 30% of the average hardness value. Pei et al. [23] recently demonstrated that the nanoindentation technique, in combination with orientation imaging microscopy, can be successfully utilized to evaluate the elastic anisotropy of Zn – a metal with a hexagonal close-packed structure and hence anisotropic. In contrast to the nanoindentation experiments, microindentation experiments always show cracked corners. Figure 2a and b shows the microindents (generated with Pmax = 100 mN) with corner cracks on basal and prism facets, whereas Figure 2c is a higher-magnification image showing the slip steps within the indent. The steps are possibly due to the crushing of the material piled up due to plastic flow prior to cracking. As seen in Figure 2a (first indent from right) and b, for indentation close to an edge of the single crystals, the cracks run all the way to the edges of the crystal, i.e. cleave the crystal completely. For the calculation of Kc, we have considered only those cracks which end well within the crystal. Also, the crack lengths are highly anisotropic, depending on the orientation of the single crystal with respect to the indenter. We have taken the shortest crack lengths and hence the values of Kc obtained are an overestimate. The values of Kc of HA in prism and basal facets are found to be 0.65 and 0.40 MPa m0.5, respectively, indicating a strong anisotropy, with 40% difference. (For reference, the respective Kc values computed using the longest crack lengths are 0.37 and 0.13 MPa m0.5, whereas the average Kc values are 0.48 ± 014 and 0.28 ± 0.14 MPa m0.5, respectively.) Note that the cracks are produced normal to the basal and prism planes, hence the values of Kc obtained are not of the basal and the prism facets.

Figure 2. Microindents on (a) the prism facet and (b) the basal facet, and (c) high-magnification SEM image of the microindent on the prism facet showing crushed slip steps within the indent.

Reported values of Kc for the polycrystalline HA range between 0.6 and 1.5 MPa m0.5, depending upon several processing parameters [24]. Typically, the Kc of single crystals is much less than that of the polycrystalline aggregate, as the grain boundaries in the latter improves toughness by promoting crack deflection. Anisotropy in toughness is observed primarily due to the anisotropy in the surface energies of facets getting cracked. The basal facet of HA has a higher surface energy compared with the prism facet [25]. Thus, the Kc of HA is higher along a facet perpendicular to the c-axis than parallel to it. Since the HA nanocrystals are oriented in such a way that the c-axis coincides with

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the long axis of the bone, higher toughness perpendicular to it can be rationalized in terms of natural design against bone fracture under flexural loading [26]. In summary, mechanical properties of the HA single crystal has been determined by micro- and nanoindentation studies, which show that these crystals are mechanically anisotropic. It appears that the single crystals of HA align such that the c-axis coincides with the long axis of the bone, imparting higher compressive stiffness and hardness to the bone. Further, nanoscale plastic deformation was observed, suggesting possible connections to the size scale of the HA crystals and their ability to withstand some plastic deformation. Finally, the observed toughness anisotropy suggests that the composite structure is designed by nature in such a way that the flexural toughness is enhanced. This work is supported by the Council of Scientific and Industrial Research (CSIR) and the Department of Science and Technology (DST), Government of India. [1] H.S. Gupta, J. Seto, W. Wagermaier, P. Zaslansky, P. Boesecke, P. Fratzl, Proc. Natl. Acad. Sci. USA 103 (2006) 17741. [2] K. Tai, F.-J. Ulm, C. Ortiz, Nano Lett. 6 (2006) 2520. [3] D.B. Burr, Bone 31 (2002) 8. [4] H.S. Gupta, W. Wagermaier, G.A. Zickler, D. Raz-Ben Aroush, S.S. Funari, P. Roschger, H.D. Wagner, P. Fratz, Nano Lett. 5 (2005) 2108. [5] S. Weiner, H.D. Wagner, Annu. Rev. Mater. Sci. 28 (1998) 271. [6] J.L. Katz, K. Ukraincik, J. Biomech. 4 (1971) 221. [7] K. Teraoka, A. Ito, K. Onuma, T. Tateishi, S. Tsutsumi, J. Biomed. Mater. Res. 34 (1997) 269.

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