Novel niobium and silver toughened hydroxyapatite nanocomposites with enhanced mechanical and biological properties for load-bearing bone implants

Applied Materials Today 15 (2019) 531–542

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Novel niobium and silver toughened hydroxyapatite nanocomposites with enhanced mechanical and biological properties for load-bearing bone implants Pengbo Wei a,b , Ju Fang a , Liming Fang c , Kefeng Wang d , Xiong Lu e , Fuzeng Ren a,∗ a

Department of Materials Science and Engineering, Southern University of Science and Technology, Shenzhen, Guangdong 518055, China Department of Mechanical and Aerospace Engineering, The Hong Kong University of Science and Technology, Clear Water Bay, Kowloon, Hong Kong, China School of Materials Science and Engineering, South China University of Technology, Guangzhou, Guangdong 510641, China d National Engineering Research Center for Biomaterials, Sichuan University, Chengdu, Sichuan 610064, China e Key Laboratory of Advanced Technologies of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Chengdu, Sichuan 621000, China b c

a r t i c l e

i n f o

Article history: Received 4 April 2019 Accepted 13 April 2019 Keywords: Hydroxyapatite Nanocomposite Interface Antibacterial Implants

a b s t r a c t An ideal load-bearing implant material for hard tissue replacement requires a combination of excellent mechanical properties and biological functions. However, such combination can hardly be achieved by monolithic bioceramics or metals. To this end, by combining the individual characteristic advantages of hydroxyapatite (HA), niobium (Nb) and silver (Ag), we have developed novel Nb and Ag co-reinforced HA nanocomposites by high energy ball milling (HEBM) and spark plasma sintering (SPS). The fabricated HA–20Nb and HA–15Nb–5Ag (wt%) show nanocomposite microstructure with metallic reinforcements uniformly dispersed in the HA matrix. The formation of a nanothick Ca4 Nb2 O9 transition layer at the interfaces of HA/Nb enables high interface strength between the reinforcements and the HA matrix. Addition of Nb or Nb–Ag could significantly increase the compressive strength and fracture toughness of HA. In vitro and in vivo evaluations further show that addition of Nb could promote osteoblast proliferation, increase the osteogenic differentiation and enhance the osteointegration ability. Incorporation of Ag could remarkably increase the antibacterial activity both against Gram-positive organism Staphylococcus aureus and Gram-negative organism Escherichia coli. Thus, the developed new nanocomposites can bridge the gap between the mechanical and biofunctional requirements for load-bearing bone implants. © 2019 Elsevier Ltd. All rights reserved.

1. Introduction Calcium phosphate (Ca-P) bioceramics have been widely used as medical implants for the repair and reconstruction of diseased or damaged hard tissues [1–6]. Among the synthetic Ca-P bioceramics, hydroxyapatite (HA) is one of the most attractive biomaterial for clinical use, due to the chemical and structural resemblance with the inorganic constituent of natural bone/teeth [7,8] and excellent bioactivity, biocompatibility and osteoconductivity [9–11]. However, the intrinsic brittleness (KIC ≈ 1 MPa m1/2 ), low mechanical strength and poor fatigue resistance of HA have limited its application for load-bearing bone implants in comparison with other metallic biomaterials [12–14].

∗ Corresponding author. E-mail address: [email protected] (F. Ren). https://doi.org/10.1016/j.apmt.2019.04.009 2352-9407/© 2019 Elsevier Ltd. All rights reserved.

To enhance the mechanical properties of HA, in particular, fracture toughness, considerable efforts have been devoted to developing HA-based composites with various second phase reinforcements such as bioglass [14–16], mullite [17,18], carbon nanotube [19–23], various fibers [24–27] and ductile metals [28–35]. Among such toughening process, the reinforcement of HA with a ductile metallic phase is a promising strategy due to the possibility of combining the mechanical strength and toughness of biocompatible metals with the osteoconductive property of HA [36]. Of these ductile metallic reinforcements, despite with certain success in titanium [28–35], iron [36–39] and silver [40–44] reinforced bioceramics from previous dedicated efforts, the challenging issues leading to low toughening efficacy remain not fully solved, which include: (1) low relative density from conventional sintering; (2) rarely retained metallic phase after densification due to oxidation or reaction with the HA decomposition products; and (3) the weak interface bonding between the HA matrix and the metallic reinforcements. Moreover, from biological aspects, the slower

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healing, low osteointegration ability and lack of antimicrobial infection of HA-based composite are also unsolved [34]. Among the 70 metals in the periodic table, niobium (Nb) is characterized by excellent biocompatibility, superior corrosion resistance, pronounced ductility and low Young’s modulus [45–51]. Previous investigations also show that Nb supports fibroblast growth and osteointegration [52,53]. In addition, Nb has low magnetic susceptibility and is magnetic resonance imaging compatible [54–56]. Recently, hot press sintered zirconia/Nb biocermet has demonstrated excellent biocompatibility and osteoconductivity [53]. Silver not only provides broad spectrum of antimicrobial activity [57], but also its high thermal expansion coefficient enables the sintered composite under residual compressive stresses and therefore, improves the crack growth resistance [40–44]. However, the reinforcing effect of only silver is very limited. Our recent studies have shown that severe plastic deformation (SPD) in the highly immiscible Nb and Ag can result in the formation of supersaturated nanocrystalline solid solutions, while during the subsequent sintering, Ag precipitated out into the Nb matrix and such Nb–Ag binary nanostructured alloys demonstrate excellent wear resistance [58,59]. In this context, by combining the abovementioned individual characteristic advantages of Nb, Ag and HA, in the present study, we propose to develop Nb and Ag co-reinforced HA nanocomposites. To achieve this, we have first fabricated the nanocomposites by high energy ball milling (HEBM) and spark plasma sintering (SPS). On the basis of a detailed characterization on the microstructure, especially the interfaces among the Nb, Ag and HA of the cermets, we have then evaluated the mechanical properties and illustrated the toughening mechanism. Finally, in vitro and in vivo evaluations on the novel nanocomposites have been performed to reveal the effects of addition of Nb and Ag on the biological performance. Such new nanocomposites can bridge the gap between the mechanical and biofunctional requirements for load-bearing bone implants.

2. Experimental 2.1. Materials fabrication 2.1.1. Powder preparation The HA powder was provided by the National Engineering Research Center for Biomaterials in Sichuan University, China. The HA nanoparticles have an average size of ∼60 nm in length and ∼30 nm in width and the phase have been confirmed by X-ray diffraction (XRD) previously [60]. Commercially pure niobium (Alfa Aesar, 1–5 ␮m, 99.99%) and silver (Alfa Aesar, 1–3 ␮m, 99.99%) were used as starting powders. Two cermet compositions of HA–20 wt% Nb (HA–20Nb) and HA–15 wt% Nb–5 wt% Ag (HA–15Nb–5Ag) along with pure HA as control were selected in this study. To avoid possible chemical reactions among Nb, Ag and HA during HEBM, the metallic powder (Nb or Nb–Ag) and HA powder were ball milled separately. First, the asreceived HA powder was ball milled for 3 h and used as the matrix. Then, niobium and silver powders with nominal weight ratio of 3:1 (Nb–25 wt% Ag) were ball milled (SPEX 8000D) for 9 h to form a supersaturated solid solution in an argon glove box. To investigate the forced atomic mixing process, the phase formation of Nb–Ag powders ball milled for 1, 3, 6 and 9 h was analyzed by XRD, respectively. In comparison to Nb–Ag, pure Nb was also ball milled for 9 h. The size and morphology of the 3 h-ball milled HA and the 9 h-ball milled Nb and Nb–25 wt% Ag powders were examined by scanning electron microscopy (SEM; TESCAN, MIRA3, Czech). Then, the ball milled Nb and Nb–Ag powders were added to HA with nominal composition of HA–20Nb and HA–15Nb–5Ag, respectively. To

mix the metallic phase and HA phase homogeneously, the powder mixtures were subjected to vibrational mixing for 3 h.

2.1.2. Consolidation of the ball milled powders into bulk by SPS After ball milling and vibrational mixing, pure HA, HA–20Nb and HA–15Nb–5Ag powders were packed into the high-purity graphite (99.99%) dies with 10 mm inner diameter and then compacted by SPS (SPS-211Lx, Fujidempa Kogyo. Co., Ltd.) at the temperature of 900 ◦ C and pressure of 60 MPa in a vacuum environment. The temperature profile during SPS is shown in Figure S1 in the Supporting Information. The obtained bulk cylinders were then ground by silicon carbide emery papers down to 1200-grit size and sequentially polished with diamond suspension of 3, 1 and 0.05 ␮m. The density of as-compacted cylinders was measured by Archimedes method.

2.2. Phase identification and microstructural characterization The phase formation of the ball-milled powders and the ascompacted bulk samples were analyzed by XRD recorded on a diffractometer in the 2 range from 10◦ to 90◦ using Cu–K␣ radi˚ 45 kV, 200 mA) with a ation (Rigaku Smartlab-9KW,  = 1.54056 A, step size of 0.02◦ and speed duration time of 5◦ min−1 . The Nb/Nb–Ag distribution in the HA matrix was revealed by SEM with energy dispersive X-ray (EDX; Oxford X-MaxN ) elemental maps. The microstructure of the bulk samples was first characterized by secondary electron (SE) imaging of focus ion beam (FIB; FEI Helios NanoLabTM 600i)-milled cross sections. Then, finer microstructure with particular focus on the interfaces between metallic phase and HA of the samples was characterized by transmission electron microscopy (TEM; FEI Tecnai F30). High angle annular dark field scanning TEM (HAADF-STEM) with attached EDX was also employed to distinguish the Nb, Ag and HA phase in the cermets. The TEM samples were prepared by FIB milling using the standard lift-out technique. To minimize surface damage by Ga+ ion implantation, the acceleration voltage and beam current were reduced to 2 kV and 23 pA, respectively, during the final step of polishing.

2.3. Mechanical properties An ideal load-bearing implant material for hard tissue replacement should possess excellent mechanical properties, including adequate modulus, high strength, fracture toughness and hardness to provide reliable stress support. The compressive strength and modulus were measured using uniaxial compressive tests on the cylinders of 3 mm in diameter and 5 mm in length with loading rate of 0.1 mm/min by universal testing machine (UTM; MTS, Sintech 10/D) at ambient temperature. The fracture toughness (KIC ) of three groups of samples was measured by a single-edge notched beam (SENB) method under the three-point bending condition [61]. Prior to SENB testing, the assintered cylinders were cut into 1.5 mm × 2 mm × 9 mm bars with a notch of 0.5 mm in depth and 0.1 mm in thickness. The span was 8 mm and the cross-head speed was set at 0.05 mm/min. The fracture toughness was calculated using the following equation: KIC =

3 PS √ aF 2 W 2S

a W

(1)

where P, S, W, B, and a are the applied load, span, height, thickness and crack length of a specimen, respectively (Supplementary Figure

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S2). The geometry correction factor function F proposed by Fett [62] is given as following, F(˛)

1 (1 − ˛)3/2

[A0 ˛ + (1 − ˛)(A1 + A2 ˛ + A3 ˛2 + A4 ˛3 + A5 ˛4 )] (2)

where A0 = 0.3738, A1 = 1.0426, A2 = −2.2336, A3 = 3.9625, A4 = −3.3469, and A5 = 1.0121. During the above strength and fracture toughness measurements, at least five independent samples were measured per composition and the mean was reported with standard deviations. The fracture surface of the specimens was examined by SEM. To check the possible anisotropy of fracture toughness induced by SPS, the SENB tests were performed both parallel (||) and perpendicular to (⊥) the pressing direction of SPS. Schematic diagrams of the geometrical details and orientation of the samples for SENB tests are presented in Supplementary Figure S2. The hardness was measured by Vickers diamond pyramidal indenter (HXD-1000TMC/LCD, Shanghai Taiming Optical Instrument Co., Ltd.) under a load of 300 gf for 10 s. For each sample, a 10 × 10 matrix was indented with 500 ␮m intervals between adjacent indents. The average hardness with standard deviations was reported. In order to understand the interaction of a crack path with microstructure and further to reveal the toughening mechanism, the indented regions on monolithic HA, HA–20Nb and HA–15Nb–5Ag were observed. 2.4. Antibacterial activity To investigate the effect of addition of Nb/Nb–Ag metallic reinforcements on the antibacterial activity, we have evaluated the antibacterial activity of the two cermets both against the Gram-negative bacteria, Escherichia coli (ATCC 25922), and the Gram-positive bacteria, Staphylococcus aureus (ATCC 29213) and made a comparison with monolithic HA. The three groups of samples (n = 3 for each group) with size of ∅10 mm × 2 mm were autoclave-sterilized and rinsed in 1 mL of phosphate buffer saline (PBS) for 1 min before being challenged with 400 ␮L of the diluted overnight culture of E. Coli or S. aureus, within the wells of a 24-well microtiter plate. Overnight bacterial suspensions were then diluted in PBS to 1 × 107 colony forming units (CFUs) mL−1 of E. Coli and S. aureus, as confirmed by colony counts on Mueller-Hinton agar (MHA). Plates were incubated at 37 ◦C for 24 h. Bacterial viability was determined after antimicrobial exposure by plating 0.1 mL of test culture onto MHA plates in triplicate. 2.5. In vitro study 2.5.1. Osteoblast-like MC3T3-E1 cell culture Mouse osteoblast-like MC3T3-E1 cell culture was performed to investigate the cellular response to the developed HA–20Nb and HA–15Nb–5Ag cermets along with monolithic HA as control. After sterilization by an autoclave for 1 h, the three groups of sample disks were placed in 48-well plates (Corning, NY, USA). 1 mL of ␣-minimum essential medium (␣-MEM) (Minimum Essential Medium alpha-Medium, Gibco, Invitrogen Inc.) containing 10% fetal bovine serum (FBS) (Hyclone, USA), 100 I.U./mL penicillin and 100 ␮g/mL streptomycin (Antibiotic-Antimycotic, Hyclone, USA) was added per well. Mouse pre-osteoblast MC3T3-E1 cells were cultured in a humidified atmosphere of 95% air and 5% CO2 at 37 ◦ C. The cells were cultured at a seeding density of 3 × 104 cells/disk. 2.5.2. Cell proliferation by MTT assays Cell proliferation on the three group of sample discs (n = 3 for each group) were analyzed by MTT (3-(4,5-dimethylthiazol2-yl)-2,5-diphenyltetrazolium bromide) assay kit following the manufacturer’s protocol (Sigma–Aldrich). After incubation for 1,

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4 and 7 days in 48-well plates, respectively, 20 ␮L of 0.5 mg/mL MTT solution was added to each well. The plates were incubated for an additional 4 h at 37 ◦ C, and then the medium in each well was replaced by 500 ␮L of DMSO to dissolve the formazan crystals. Absorbance at a wavelength of 570 nm of the solution in each well was measured using a microplate reader (Biotek-Cytation 3). 2.5.3. Cell viability Cell viability was analyzed by live/dead staining. Cells after 1 day and 4 days incubation on cermet discs were stained with 4 mM calcein-AM and 2 mM propidium iodide for 10 min. Then, the cells were washed with PBS and air dried. The viability of cells was visualized by confocal laser scanning microscopy (CLSM; TCS SP8, Leica, Germany). The living cells were stained green while dead cells were stained red. 2.5.4. Cell morphology Cell morphology on the two cermets and HA discs was observed by SEM. After incubation for 1 day and 4 days, culture media was removed and cells were rinsed with PBS followed immediately by fixation using 2.5% glutaraldehyde for 10 min. Once cells were fixed, gradually dehydration process was carried out using serial dilutions of ethanol (30%, 50%, 75%, 95%, and 100%). The discs were freeze dried and coated with gold prior to SEM observation. 2.5.5. In vitro osteogenic differentiation The osteogenic differentiation of cells on the three groups of samples was determined by the alkaline phosphatase (ALP) activity assay. 1 × 106 cells were seeded on each sample in a 48-well plate. After 7 days incubation, the medium was removed, and the cells were washed twice with PBS. Afterward, 200 ␮L of 1% TritonX-100 was added to lyse the cells. The cell lysis was performed by freeze–thaw treatment between −80 ◦ C and 4 ◦ C for three times. The total protein concentration of the cell lysate of each sample was detected by a BCA Kit (Jiancheng Biotech, China). ALP activity was measured by an ALP Assay Kit (Jiancheng Biotech, China). The ALP activity was calculated by normalized with total protein concentration of the same cell lysate. The experiment was performed with five independent samples per group. 2.6. In vivo study 2.6.1. Animal model Female New Zealand white rabbits (6-month-old, weight 2.8–3.2 kg) were used as animal model for surgery. All procedures of animal experiments were carried out under the permission and regulations of the Animal Care and Use Committee of Southern University of Science and Technology. Briefly, rabbits were anesthetized with pentobarbital (2%, 30 mg/kg) via ear vein injection. The skin over anterior femur was shaved, and then sterilized with 10% povidone-iodine solution. A skin incision was created and the femur was exposed by separating the muscles and sarcolemma. Two medical defects of 5 mm in depth and 1 mm in diameter were created using a pilot drill at the femoral cortical bone under constant rinse of saline for cooling. The two defects were more than 10 mm apart from each other. Sample implants (HA, HA–20Nb and HA–20Nb–5Ag) with the same size of the defect were inserted. The periosteum and skin were carefully sutured, respectively. The surgical wound on skin was rinsed with povidone-iodine solution. After 3 months of post-surgery, rabbits were sacrificed by the injection of lethal dose of pentobarbital. The femurs were harvested, fixed in 4% buffered formaldehyde solution for subsequent micro-computer tomography (micro-CT) and histological examinations.

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Fig. 1. XRD patterns of ball milled powders and as sintered bulk samples: (a) ball milled HA powder and as sintered HA; (b) Nb–25 wt% Ag powders after ball milling for varying time intervals; (c) ball milled HA–15Nb–5Ag powder and as sintered HA–15Nb–5Ag; and (d) ball milled HA–20Nb powder (powder mixture of 80 wt% of ball-milled HA and 20 wt% of ball-milled Nb) and as sintered HA–20Nb. Note that HA was ball milled for 3 h in (a, c, and d), both Nb and Nb–25 wt% Ag was ball milled for 9 h in (c and d).

2.6.2. Ex vivo micro-CT analysis High-resolution images of all the specimens were obtained from micro-CT scanner (Skyscan 1176, Bruker, Germany) with the source voltage and source current of 90 kV and 270 ␮A, respectively, and an exposure time of 550 ms. The sagittal, coronal and axial planes of each implant were reconstructed. The evaluated area was considered as a ring of 470 ␮m radius from implant surface in order to analyze the bone volume per total volume (BV/TV) and bone-toimplant contact ratio (BC). 2.6.3. Histological analysis Hard tissue histological analysis was conducted to evaluate the healing of defects and the osteointegration between the implant and new bone. After the ex vivo micro-CT scans, the bone specimens containing the implants were dehydrated using 100% ethanol and embedded in methyl methacrylate for 1 month. The position of the implants was located via radiologic examinations. After photopolymerization of the infiltrated blocks, sections of 50 ␮m were cut by a microtome (SYNTEK Co., Ltd., Japan). The prepared slides were stained with 5% acid fuchsin and 1% toluidine blue and observed by using a transmitted light microscope (Eclipse E100, Nikon, Japan) with an imaging system (DS-U3, Nikon, Japan). 2.7. Statistical analysis The data were presented as mean ± standard deviation. Oneway analysis of variance (ANOVA) was used to analyze the significant differences between samples. A value of p < 0.05 was considered statistically significant.

3. Results 3.1. Powder morphology and phase compositions The SEM images of the ball-milled HA, Nb and Nb–Ag particles are presented in Supplementary Figure S3. The ball-milled HA shows agglomerates with an average size of 2.9 ␮m, while the ballmilled Nb and Nb-25 wt% Ag powders show big flakes, with an average size of 8.5 ␮m and 10.5 ␮m, respectively. Fig. 1 presents the XRD patterns of the powders and the sintered bulk HA, HA–20Nb and HA–15Nb–5Ag samples. Fig. 1a confirms that the HA phase is retained after ball milling for 3 h and the subsequent densification at 900 ◦ C. XRD patterns of Nb and Ag powder mixtures after ball milling for different time intervals are presented in Fig. 1b. After 1 h, we can still see the Ag reflections despite with weak intensity, but after 3 h, the Ag reflections are completely gone and the Nb diffraction peaks shift to smaller 2 angles, which suggest that SPD during ball milling induced the formation of a supersaturated Nb-rich solid solution. Extending the ball milling time from 3 h to 6 h broadens the diffraction peaks of the Nb-rich phase, which suggests further grain size reduction and/or micro-strain increase in the supersaturated phase. The XRD patterns show almost no difference between 6 and 9 h. Thus, we stopped the ball milling of Nb–Ag at 9 h and mixed them with HA according to the nominal compositions. The HA–15Nb–5Ag powder mixture only contains Nb-rich phase and HA, but after SPS into bulk at 900 ◦ C, Ag reflections are present (Fig. 1c). This means that Ag precipitated out from the Nb-rich solid solution during sintering. As for the HA–20Nb, no extra peaks are present after sintering,

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Table 1 A comparison of relative density and mechanical properties of spark plasma sintered HA–20Nb and HA–15Nb–5Ag nanocomposites with the monolithic HA. The fracture toughness data we choose is from SENB test with the loading direction parallel to the pressing direction of SPS. Samples

Relative density (%)

Young’s modulus (GPa)

Compressive strength (MPa)

Vickers hardness (HV)

Fracture toughness (MPa m1/2 ) (||)

HA HA–20Nb HA–15Nb–5Ag

99.7 ± 0.1 99.2 ± 0.2 99.3 ± 0.2

64.1 ± 4.9 72.3 ± 6.1 74.4 ± 3.8

245.7 ± 27.8 311.4 ± 30.4 340.0 ± 26.8

429.6 ± 13.6 464.3 ± 23.8 456.3 ± 43.7

3.32 ± 0.23 4.17 ± 0.17 4.40 ± 0.25

Fig. 3. Secondary electron images of the HA/metallic phase interfaces from HA–20Nb (a) and HA–15Nb–5Ag (b) nanocomposites milled by FIB. Fig. 2. SEM images and corresponding elemental mappings of HA (a–c), HA–20Nb (d–g) and HA–15Nb–5Ag (h–l).

however, both Nb and HA diffraction peaks become much sharper, indicating grain growth and release of residual strain (Fig. 1d). In agreement with our expectation, no detectable new phase formation and decomposition of HA are present in the bulk samples. 3.2. Microstructure and interfaces Density measurements show that all the spark plasma sintered (SPSed) bulk samples at the present processing parameters reach nearly full density, exceeding 99% of their corresponding theoretical density, as shown in Table 1. Addition of metallic phase (Nb/Nb–Ag) only slightly (≤0.5%) reduces the relative density compared with monolithic HA, unlike iron-containing HA/Ti composites via pressureless sintering where incorporation of 15 wt% Ti–Fe particles leads to up to ∼15% reduction in relative density [33]. This should be attributed to the nanostructured powders via HEBM, and the fast heating rate and pulsed electric current assistance via SPS. SEM images along with elemental mappings can reveal the metallic phase distribution in the HA matrix, as shown in Fig. 2. In contrast to monolithic HA with Ca and P full coverage on the surface (Fig. 2a), the composite shows homogeneous distribution of Nb particles/Nb–Ag agglomerates in the HA matrix (Fig. 2d–l). Over 90% of the metallic particles in the HA matrix have aspect ratios smaller than 1.5 (Fig. 2d and h). EDX elemental mappings

confirm that the particles in bright contrast are Nb/Nb–Ag. For HA–20Nb composite, the Nb particle size scatters from 0.8 ␮m to 30 ␮m, while for HA–15Nb–5Ag, since Ag precipitated from Nb-rich supersaturated solid solution matrix, Nb particles and Ag precipitates agglomerate together which can hardly be distinguished by EDX elemental mapping. The Nb–Ag agglomerate size scatters from 0.5 ␮m to 50 ␮m. SEM images of the polished surfaces also confirm the high relative density of bulk samples. Almost no pores were observed in monolithic HA and only several small pores with size of several microns are found in the composites. In order to investigate the interfaces between the HA and the metallic reinforcements (Nb/Nb–Ag), we use FIB to mill cross-sections near the interface regions in both HA–20Nb and HA–15Nb–5Ag cermets. Both cermets show jagged interfaces between the HA matrix and the metallic reinforcements, as shown in Fig. 3. Compared with HA–20Nb, the HA–15Nb–5Ag cermet shows more intimate contact with segregated Ag along the interfaces (Fig. 3b). This is probably due to that local temperature during SPS may exceed the melting point of silver (961 ◦ C) [63], which causes silver melting and precipitation in areas of lower chemical potential where particles are in not close contact like liquid-phase sintering [64]. Fig. 3b also reveals that except the Ag segregated at the HA/metallic reinforcements interface regions, Ag precipitates scatter from 10 to 110 nm (with average size of 40 nm) are uniformly dispersed in the Nb matrix. Detailed TEM/STEM based characterization has also been performed on the two cermets along with monolithic HA to reveal

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Fig. 4. Microstructure and elemental compositions of HA–20Nb and HA–15Nb–5Ag. (a) and (b) are HAADF-STEM and corresponding bright field TEM images of HA–20Nb, respectively; (c) shows the EDX spectra obtained from Nb, Nb/HA interface and HA; (d) is SAED pattern of the interface transition layer; (e) is HAADF-STEM image of HA–15Nb–5Ag nanocomposite; and (f) shows the EDX spectra from three contrasts in (e).

more microstructural information at the nanoscale. The representative bright field TEM images of HA obtained from the three bulk samples are shown in Supplementary Figure S4. The HA in the three samples have similar average grain size of ∼60 nm. Selected area electron diffraction (SAED) patterns further confirm the phase composition of HA. Addition of the Nb/Nb–Ag reinforcements does not change the microstructure of HA matrix. Then, we turn to characterize the microstructure of the two cermets. HAADF-STEM imaging which provides atomic number (Z−) contrast was used to distinguish different phases present in the composites. Fig. 4a and b shows the HAADF-STEM and corresponding bright field TEM images of the HA/Nb interface in HA–20Nb cermet. The Nb in HA–20Nb shows elongated grains with mean long axis of 390 nm and aspect ratio of 3. A transition layer with thickness of ∼150 nm was observed at the interface between HA and Nb. EDX spectrum (Fig. 4c) reveals that this nanothick transition layer contains Ca, O and Nb. The phase of the transition layer is identified to be Ca4 Nb2 O9 by SAED (Fig. 4d), which suggests reaction bonding between the HA matrix and Nb reinforcement. Such reaction bonding should yield higher interface strength than mechanical or physical bonding. More importantly, such nanothick transition layers were only localized at the Nb/HA interfaces which make most of the Nb phase retained, as confirmed by XRD in Fig. 1c and d. This should be attributed to the advantage of SPS where the micro-plasma discharges result in localized and momentary heating of the particles surfaces and thus, formation of ‘nano-necks’ between particles [65]. Fig. 4e shows the HAADF-STEM image of HA–15Nb–5Ag cermet. The bright, gray and dark contrasts are Ag, Nb and HA respectively, as also confirmed by EDX spectra (Fig. 4f). The Nb in HA–15Nb–5Ag shows much smaller grain size than that in HA–20Nb, with average long axis of 250 nm and aspect ratio of 2.5. The reduced grain size of Nb should be due to the Ag addition which inhibits the grain growth during processing [59].

3.3. Mechanical properties Fig. 5 shows the compressive stress-strain curves of HA–20Nb and HA–15Nb–5Ag cermets with monolithic HA. The modulus and compressive strength of SPSed monolithic HA are 64.1 ± 4.9 GPa and 245.7 ± 27.8 MPa, respectively. Addition of the Nb/Nb–Ag metallic reinforcements slightly increases the modulus (12–16%),

Fig. 5. The compressive stress–strain curves of the HA, HA–20Nb, and HA–15Nb–5Ag.

as revealed in similar slope in the elastic regions of stress-strain curves. However, the compressive strength increases 27% and 38% for HA–20Nb and HA–15Nb–5Ag cermets, respectively. The fracture toughness of three groups of samples measured by SENB method (the loading direction of SENB parallel to the pressing direction of SPS) under the three-point bending condition show that as compared with KIC of monolithic HA of 3.32 ± 0.23 MPa m1/2 , the HA–20Nb and HA–15Nb–5Ag show 25% and ∼33% increase in the 25% KIC value, respectively. The HA–15Nb–5Ag has the highest KIC of 4.40 ± 0.25, respectively. A comparison of the fracture toughness values measured along (||) and perpendicular to (⊥) the pressing direction of SPS is presented in Table S1. No apparent difference was found between the fracture toughness values measured from the two directions, indicating that no anisotropy was present in the SPSed bulk samples. The fracture surfaces of HA–20Nb and HA–15Nb–5Ag are distinct from that of HA (Supplementary Figure S5). The fractograph of HA shows characteristic feature of brittle fracture, while microviods and pullout of the metallic particles were observed in fracture surface of the HA–Nb and HA–15Nb–5Ag nanocomposites, indicating interface debonding between the metallic particles and the HA matrix. It is noticed that the fracture toughness in the present study shows very close value to previous SPSed monolithic HA measured by four-point flexural SENB [34], which confirms the reliability our toughness measurements. It is also found that addition of

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Fig. 6. The representative optical photographs (a) and statistical analysis of total colony forming units of E. coli (b) and S. aureus (c) bacterial colonies after incubation with HA, HA–20Nb and HA–15Nb–5Ag at 37 ◦ C for 24 h with the same dilution. n = 3, *p < 0.05, **p < 0.01, ***p < 0.001.

the Nb/Nb–Ag reinforcements into the HA matrix slightly (6–8%) increases microhardness. This finding differs from previous study on iron-containing HA/Ti composites via pressureless sintering, which shows that increasing the amount of Ti–Fe particles causes linear reduction in microhardness and elastic modulus [33]. This is probably attributed to nanostructured characteristics of our metallic reinforcements. The measured modulus, compressive strength, fracture toughness and hardness are summarized in Table 1. 3.4. Antibacterial activity The antibacterial activity of the HA, HA–20Nb and HA–15Nb–5Ag against Gram-positive organism S. aureus and Gram-negative organism E. coli was evaluated by the platecounting method. The representative images of viable bacterial colonies after 24 h co-incubation with the sample discs are shown in Fig. 6a. Statistical analysis on the numbers of CFUs is included in Fig. 6b and c. It clearly shows that upon against E. coli, the HA–20Nb (E. coli, 175 ± 3) has very close numbers of CFUs with the HA control group HA (E. coli, 182 ± 10), but upon against S. aureus, the HA–20Nb (S. aureus, 162 ± 4) group has a little smaller number of CFUs than that of HA (S. aureus, 197 ± 1). This suggests that niobium also has weak antibacterial activity against S. aureus. Remarkably, compared with HA and HA–20Nb, the CFUs of both E. coli and S. aureus are dramatically reduced upon incubation with HA–15Nb–5Ag discs (E. coli, 62 ± 6, reduced ∼66%; S. aureus, 24 ± 5, reduced ∼88%), which suggests the HA–15Nb–5Ag cermet has strong antibacterial activity. Such strong antibacterial activity may effectively prevent microbial infection induced inflammation after implantation. 3.5. In vitro cytocompatibility To evaluate the effect of addition of Nb/Nb–Ag reinforcements on the cytocompatibility of HA, the MC3T3-E1 cell proliferation of

the three groups of discs were assessed by MTT assay. As shown in Fig. 7a, steady increase of cell proliferation is observed on all discs during 7 days’ culturing, indicating that all samples show no cytotoxicity to MC3T3-E1 cells. Among the three groups, the HA–20Nb shows the highest cell proliferation rate at all the time intervals, which means that addition of Nb can improve the cytocompatibility. Compared with HA, the HA–15Nb–5Ag group has almost the same cell proliferation rate at day 1, slightly reduced at day 4 but then increased at day 7. These results suggest that despite introducing Ag into the HA matrix, the HA–15Nb–5Ag nanocomposite has no cytotoxicity but even enhances the long term cell proliferation. Live/dead staining results (Fig. 7b) observed by CLSM show no significant difference among HA, HA–20Nb and HA–15Nb–5Ag. All sample discs have a dense concentration of living cells attached on the surface after 4-day incubation and no dead cells were found, suggesting good cytocompatibility even though the antimicrobial agent Ag was introduced to the nanocomposite. This may be attributed to the low relative amount as well as the uniform distribution of Ag in Nb phase (as confirmed by SEM in Fig. 2). Hence, HA–15Nb–5Ag shows almost no difference in cytocompatibility when compared with pure HA and HA–20Nb. SEM observation of cell morphology shows that all the discs support osteoblast attachment and spreading (Supplementary Figure S6). No impaired cell morphology was observed. Almost the whole surface is covered with cells at day 4. The enlarged images of a single cell at 4-day incubation clearly reveal the well spreading of osteoblasts on all discs (Supplementary Figure S6g–i, denoted by white arrows). The extended cellular body and the stretching of filopodia of cells are beneficial to cell communications and extracellular matrix (ECM) deposition. The SEM observations confirm good cell affinity of our fabricated HA–20Nb and HA–15Nb–5Ag cermets. ALP assay was performed to evaluate the osteogenic differentiation of osteoblasts seeded on the samples (Fig. 7c). The cells on HA–20Nb and HA–15Nb–5Ag show ∼70% and ∼20% increase in the ALP activity, respectively, when compared with cells seeded on

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Fig. 7. The osteoblastic MC3T3-E1 cells proliferation (a), cell viability (b), and alkaline phosphatase (ALP) activity after cultivation on the HA, HA–20Nb, and HA–15Nb–5Ag bulk samples. *p < 0.05, ***p < 0.001.

Table 2 Bone contact ratio and bone volume/tissue volume (BV/TV) of HA, HA–20Nb and HA–15Nb–5Ag rod implants at 3 months after implantation.

HA HA–20Nb HA–15Nb–5Ag

Bone contact ratio (%)

BV/TV (%)

98.5 97.4 95.2

96.1 98.5 93.0

HA. Since ALP expression is closely related to cell mineralization during osteogenic differentiation, the ALP assay results clearly show that Nb substantially promotes osteoblast maturation, which finally promotes osteogenic differentiation [66].

3.6. In vivo biocompatibility 3.6.1. Micro-CT analysis of bone regeneration To investigate the defect recovery after surgery, micro-CT analysis was performed. The bone contact ratio and bone to tissue volume (Table 2) are calculated by a ring-like model with a 470 ␮m radius from the implant surface. The represented micro-CT slice images and reconstructed bone contact models are shown in Fig. 8. The results show that all implants have high bone contact ratios and BV/TV over 90%, indicating the new bone grows in close contact with the implant surface and the new bone density is almost the same with the host cortical bone of femur. The data also demonstrate that the two cermets have very close bone contact ratios and BV/TV to that of HA.

Fig. 8. Micro-CT analysis on the HA (a), HA–20Nb (b) and HA–15Nb–5Ag (c) rod implants at 3 months after implantation. The bone-implant contact ratio and bone volume/tissue volume (BV/TV) were measured by a ring-like model. The reconstructed contact region of HA, HA–20Nb and HA–15Nb–5Ag implants are shown in (d)–(f), respectively.

3.6.2. Histological analysis of bone regeneration To examine the in vivo biocompatibility and osteointegration of the nanocomposites, histological analysis was performed in comparison with HA. The new bone formation around the implants and the interfacial morphology between the implants and newly formed bone are shown in Fig. 9. After 3 months’ implantation, the implants of HA, HA–20Nb and HA–15Nb–5Ag are all associated directly with the compact bone of femur. The new bone tissue that

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Fig. 9. Optical images of stained slices of histological specimens (rod implant + bone), extracted after 3 month’s implantation: (a) HA, (b) HA–20Nb, and (c) HA–15Nb–5Ag. The enlarged interface regions between bone and implants of HA, HA–20Nb and HA–15Nb–5Ag rod implants are shown in (d)–(f), respectively. Notes: NB, new bone; V, vascular space; OC, osteocyte.

stained red by fuchsin is found to surround the implants, covering both the surface end and the apical end of the rods. Osteocytes are found to embedded within the new bone, which is a key evidence of bone maturation. The newly formed bone is also strongly vascularized. For the HA and HA–20Nb groups, visible bone fusion can be seen at the boundary of implant and new bone (Fig. 9d and e). In contrast, slightly less bone fusion is present in the group of HA–20Nb-5Ag (Fig. 9f). No dissociation layer or fibrous tissue is observed at the interfaces between the three groups of implants and new bone. No inflammatory infiltration and atypical cells are observed in all groups of slices. The histological results indicate the HA–20Nb and HA–15Nb–5Ag implants show similar osteointegration ability as well as biocompatibility to HA.

4. Discussion The primary purpose of the present work is to enhance the mechanical strength, fracture toughness of HA while maintain its excellent bioactivity, biocompatibility and osteointegration ability, which will extend its application for load-bearing bone implant material. Moreover, to prevent implant-associated infections which frequently cause implant failure, we also aim to increase the antimicrobial activity of the developed nanocomposite. To achieve these goals, we have developed a Nb–Ag co-reinforced nanocomposite by HEBM and SPS. As expected, the HEBM process forces the immiscible Nb–Ag to form a nanocrystalline supersaturated solid solution and during densification at elevated temperature, the silver precipitates out into the niobium matrix. The fast heating rate and short sintering time of SPS enable the fabricated cermet high relative density and nanocomposite microstructure. One important finding is that addition of 20 wt% Nb or 15 wt% Nb together with 5 wt% Ag significantly increases the compressive strength and fracture toughness of monolithic HA. The SPSed HA–15Nb–5Ag nanocomposite shows the long crack toughness of up to 4.40 ± 0.25 MPa m1/2 , at the same level of the best achieved toughness of the previously developed HA–10Ti biocomposites [34]. Such toughness is also at least three times higher than the HA biocomposites developed by conventional sintering techniques. In order to reveal the toughening mechanism, the crack-microstructure interactions were investigated via

Vickers indentation and SEM observations. Since HA–20Nb and HA–15Nb–5Ag nanocomposites show similar toughening mechanism, here, we only show the representative post-indentation SEM images of HA–15Nb–5Ag along with monolithic HA. For monolithic HA, as shown in Fig. 10a and b, the crack paths emanating from the indent are straight with average radial length up to 29.5 ± 3.5 ␮m, which is typical for brittle ceramics. Fig. 10c-f show typical crack paths for HA–15Nb–5Ag nanocomposites. The average radial crack length is sharply reduced to 12.8 ± 4.4 ␮m. Fig. 10c and d demonstrates that the Nb–Ag agglomerates/HA interfaces deflect the propagating crack away from the principal direction, resulting in a tortuous path to release stress. This provides direct evidence for crack deflection toughening mechanism. The interface debonding leading to crack deflection is basically energy-dissipating process that can result in an increase in toughness or work of fracture. The crack-wake interfacial debonding may be attributed to the presence of Nb, Ag and/or Ca4 Nb2 O9 phase, as confirmed in Figs. 3 and 4. Fig. 10e and f shows the typical example of crack impeding (crack tip shielding). Since fracture toughness of the metallic phases is remarkably higher than HA matrix locally, when the crack tips encounter the Nb–Ag agglomerates, the resistance to crack propagation is increased and more energy will be consumed, and thus, the cracks could be arrested by such metallic agglomerates/particles. These results suggest that niobium particles or Nb–Ag agglomerates are highly effective in preventing crack propagation in HA matrix. The mechanisms of crack deflection and crack impeding are responsible for the toughness enhancement in HA–Nb and HA–Nb–Ag nanocomposites. The other key finding is that incorporation of Nb into the HA matrix could promote osteoblast proliferation, increase the ALP activity, and enhance the osteointegration ability, as confirmed by in vitro and in vivo evaluations on HA and HA–20Nb. The excellent bioactivity of niobium has also been observed in previous studies [48,67,53,68] although the exact mechanism remains unclear. A widely accepted view lies in the formation of Nb-OH upon exposure to air or any electrolyte through hydroxylating the surface oxide [69]. The abundant amphoteric Nb–OH groups can not only facilitate the apatite nucleation, but also can bond the terminal carboxyl and amino groups of amino acids and proteins [53]. When the 5 wt% antimicrobial agent silver was introduced to the nanocomposite while maintaining the same weight percentage of the total metallic reinforcements, the antibacterial activity both against E. coli and

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Fig. 10. SEM images of the indents (Vickers hardness at 300 gf for 10 s) and radial cracks of HA (a and b) and HA–15Nb–5Ag (c–f). (b), (d) and (f) show high magnification images of the selected area in (a), (c) and (e), respectively. (c and d) show the evidence of crack deflection; (e and f) show the evidence of crack impeding (crack tip shielding).

S. aureus are remarkably increased, as shown for Nb–15Nb–5Ag nanocomposite. The accurate antibacterial mechanism of silver is also not fully understood and still in debate [70,71], it is generally accepted that silver ions could damage the external bacterial membranes, bind to essential intracellular components like DNA and thus prevent the bacteria from performing even their most basic functions [71–73]. The negative effects of incorporation of silver on cytocompatibility, ALP activity and osteointegration ability of the nanocomposite is negligible, especially from the long-term view, as evidenced in Figs. 7–9. Therefore, the developed new composite implants can bridge the gap between the mechanical and biofunctional requirements for load-bearing bone implants. Despite that reinforcing bioceramics using second metallic phase has been reported previously for Ti, Fe and Ag, but incorporation of Nb or binary immiscible nanostructured Nb–Ag phases to HA is novel. Our present study also systematically explored the correlation among the processing, microstructure, mechanical properties and in vitro and in vivo biocompatibility of the fabricated nanocomposites, which should provide deep insights into developing novel nanocomposites for load bearing bone implants and such Nb/Nb–Ag reinforcements should also be easily extended to toughen other nanocomposites. Finally, it also should be mentioned that we have only investigated two compositions besides the control sample HA. Future work should be devoted to the effect of volume fraction of the reinforcements on the mechanical properties and biological performance. There is still much composition space yet to be explored and optimized. Furthermore, the present nanocomposites were fabricated in the dense form; developing porous nanocomposite scaffold is another promising direction. 5. Conclusions We have systematically explored the processing, microstructure, mechanical properties and in vitro and in vivo biocompatibility of HA–20Nb and HA–15Nb–5Ag nanocomposites. The following major conclusion can be drawn: (a) Highly dense HA–20Nb and HA–15Nb–Ag nanocomposites were fabricated by HEBM and SPS. The obtained nanocomposites have the microstructure with metallic phase uniformly dispersed in the HA matrix. The HA matrix has the average grain size of 60 nm and the Ag has the average precipitate size of 40 nm. The Nb phase is in elongated shape, with long axis of

(b) (c)

(d)

(e)

390 nm in HA–20Nb and 250 nm in HA–15Nb–5Ag, with aspect ratio of ∼3. A nanothick transition layer formed at the HA/Nb interfaces with the phase identified to be Ca4 Nb2 O9 . In comparison to monolithic HA with compressive strength of 245.7 ± 27.8 MPa and fracture toughness of 3.32 ± 0.23 MPa m1/2 , the compressive strength was increased 27% and 38%, and the fracture toughness was increased 25% and 33% for the HA–20Nb and HA–15Nb–5Ag nanocomposites, respectively. The HA–15Nb–5Ag nanocomposite shows strong antibacterial activity both against Gram-positive organism S. aureus and Gram-negative organism E. coli but still maintain the similar biocompatibility. Compared with HA, the HA–15Nb–5Ag nanocomposite shows 66% and 88% reduction in the numbers of CFUs incubated under the same condition. In vitro and in vivo evaluations show that addition of Nb into the HA matrix could promote osteoblast proliferation, increase the ALP activity, and enhance the osteointegration ability.

To sum up, the fabricated nanocomposites with combined excellent mechanical properties and biofunctions should have great potential in load-bearing bone implant. Acknowledgements This work was financially supported by the National Key Research and Development Program of China (2016YFB0700803) and the Fundamental Research Program of Shenzhen (Grant Nos. JCYJ20170307110418960 and JCYJ20170412153039309). This work was also supported by the Pico Center at SUSTech that receives support from Presidential Fund and Development and Reform Commission of Shenzhen Municipality. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.apmt.2019.04.009. References [1] L.L. Hench, Bioceramics: from concept to clinic, J. Am. Ceram. Soc. 74 (1991) 1487–1510.

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