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hydroxyapatite interface
Diamond & Related Materials 100 (2019) 107561
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Diamond & Related Materials journal homepage: www.elsevier.com/locate/diamond
Preparation of reduced graphene oxide/hydroxyapatite nanocomposite and evaluation of graphene sheets/hydroxyapatite interface
T
Hassan Nosratia, Rasoul Sarraf Mamoorya, , Dang Quang Svend Leb, Cody Eric Büngerb ⁎
a b
Department of Materials Engineering, Tarbiat Modares University, Tehran, Iran Department of Clinical Medicine, Aarhus University, Denmark
ARTICLE INFO
ABSTRACT
Keywords: Hydroxyapatite Graphene Hydrothermal Gas injection Interface
In this investigation, reduced graphene oxide/hydroxyapatite (rGO/HA) hybrid powders have been synthesized using hydrogen gas injection into hydrothermal autoclave. The powders were then consolidated with spark plasma sintering. The results showed that this method of synthesis caused the obtained powders to have high crystallinity. Microscopic analysis confirmed the presence of rGO sheets with folding and wrinkling in the nanocomposite and indicated that various crystalline planes such as (002) and (300) played a role in the growth of hydroxyapatite crystals. The results of interface analysis (HA||rGO) showed that the HA is coherently connected by its (300)-planes with the surface of the rGO sheets and this coherency is accomplished in the rGO crosssection with the (002) planes of HA. The hardness and the Young's modulus of the composite samples were 5.9–6.6 GPa and 132–146 GPa respectively.
1. Introduction Synthesis of hydroxyapatite (HA) with nanotechnological approaches in nanosized dimensions and in various forms such as nanotubes and nanorods have improved the mechanical properties of this ceramic material [1–3]. The chemical composition (Ca10(PO4)6(OH)2), the ratio of calcium to phosphate (C/P ≈ 1.67), and the crystalline structure of HA are very similar to the mineral part of the human skeletal system. Also, having properties such as bioactivity and osteoconductivity make HA one of the most valuable materials for orthopedic applications [4–6]. The applications of this bioceramic are very broad include orthopedics, biosensors, catalysts and drug delivery [7–12]. However, the poor mechanical properties of HA such as fracture toughness, intrinsic brittleness, and poor wear resistance still limit the applications of HA despite its excellent biomaterial properties [13–15]. To make HA usable as an implant, one of the strategies is to enhance its mechanical properties with a second, reinforcing, phase. Various materials have been investigated including titanium oxide, aluminum oxide, and carbon nanostructures [16–20]. But among these materials, graphene, with its unique mechanical properties (Young's modulus ≈ 1 TPa and fracture strength ≈ 130 GPa), has attracted a great deal of attention and research [21,22]. Graphene has a honeycomb-like structure of a single carbon atom with the SP2 hybrid [23–25]. Its high
⁎
specific surface area (2630 m2 g−1) has made its reinforcing properties more effective. Also, its good biological properties have led to much research for medical applications including orthopedic surgery, bioimaging, and drug delivery [26–30]. Several studies have been published on the use of graphene and its derivatives to improve the mechanical and biological properties of HA [31–34]. The results of these investigations showed that the mechanical and biological properties of HA were improved by the addition of graphene. Graphene has been used as filler that blocks the crack growth with its two-dimensional structure [35–37]. One of the methods to synthesis of graphene is a hydrothermal method that the graphene oxide (GO) is used as precursor. Hydrogen under the high pressure and high temperature of the hydrothermal autoclave is the main factor in the reduction of GO, assembly as a gel, and forming three-dimensional rGOs. Using hydrothermal method is an in situ mechanism for the synthesis of HA-rGO hybrid powders because the presence of oxygen agents on GO sheets causes nucleation and growth of HA particles on the graphene surface [38–41]. Due to the high mechanical properties of graphene recovered from GO, the degree of reduction is very important. One of the most attractive methods is utilizing of hydrogen gas as a reducing agent. In one published report, a mixture of argon and hydrogen gases was used alongside palladium and the process temperature was 200 °C. Palladium caused hydrogen gas to automatically reduce the oxygen present on the graphene oxide surface [42]. In another
Corresponding author. E-mail address: [email protected] (R.S. Mamoory).
https://doi.org/10.1016/j.diamond.2019.107561 Received 18 July 2019; Received in revised form 21 September 2019; Accepted 27 September 2019 0925-9635/ © 2019 Elsevier B.V. All rights reserved.
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electron microscopy (HRTEM), X-ray photoelectron spectroscopy, and nanoindentation were used to investigate the type (coherent or incoherent), chemical bonds, and mechanical behavior of the interface.
Table 1 The primary chemicals used in the powder synthesis phase. No.
Chem.
Co.
Purity
Formulation
1
Calcium nitrate tetrahydrate Diammonium hydrogenphosphate DMF Ammonium solution GO
Merck
> 99%
Ca(NO3)2·4H2O
Merck
> 99%
(NH4)2HPO4
Sigma-Aldrich Merck Abalonyx
> 99.8% 25% 25 g/l DMF
(CH3)2NC(O)H NH4OH COxHy
2 3 4 5
2. Materials and methods According to the chemical formula of HA, the Ca/P ratio should be 1.67. Therefore, for the synthesis of primary powders, 7.84 g of calcium nitrate tetrahydrate was dissolved in 200 ml deionized water and 1.32 g of diammonium hydrogenphosphate dissolved in 150 ml deionized water to obtain two solutions containing calcium and phosphate ions as the basic HA ions.
report, hydrogen-rich water was used to reduce GO [43]. Further researches have been published in which hydrogen gas is injected into a hydrothermal autoclave, and the results showed that the presence of hydrogen gas in hydrothermal conditions would also reduce GO [44,45]. The mechanical behavior of composites made of HA and rGO depends on interface type between the two phases. Therefore, studying the details and behavior of the interface is very important. In this study, the hybrid nanostructured powders were synthesized using a hydrothermal method utilizing hydrogen gas injection to increase the reduction rate of GO. However, to have a greater effect on reduction of GO, hydrogen gas needs to be decomposed by the catalysts to the hydrogen atoms, because the hydrogen gas has little solubility in solution. So, hydrothermal temperature and time will probably have a stronger role in the synthesis of powders [42–46]. In order to investigate the effect of this process on the final nanocomposite properties, the spark plasma sintering (SPS) method has been used to consolidate these powders and to fabricate bulk samples. High-resolution transmission
2.1. Powders preparation The following steps were used to synthesize the powders, respectively. a) The solution containing calcium ions was added dropwise to a 20 ml stirred suspension of GO (3.13 mg/ml) (HA/1.5% rGO) with stirring continued for 1 h. b) The solution containing phosphate ions was dropwise added to the solution. c) The pH of the solutions was adjusted to 11 with ammonium solution. d) The resulting solution was poured into the Teflon (PTFE) vessel and transferred to the hydrothermal autoclave. The hydrothermal process was carried out for 5 h at 180 °C by injection of hydrogen gas at 10 bar. e) The powders were dried at oven for 24 h at 60 °C. The volume of the PTFE container was 340 ml and the total pressure was 25 bar when using gas injection. The primary chemicals used in the powder synthesis phase with the specifications are listed in Table 1.
Fig. 1. Dilatometric curves of the sintered samples (a, b) and the nanoindentation loading curves (c). 2
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Fig. 2. HAADF image (a), elemental analysis of calcium/phosphorus (b), phosphorus (c), oxygen (e), and calcium (f), for synthesized powders, EDS analysis (d), FESEM image (g), and TEM image (h) of powders.
STEM mode. Exposure times of 5 min were used to create elemental distribution maps, with satisfactory counting statistics, while minimizing potential problems such as beam damage and specimen drift. STEM images were obtained using a high angle annular dark field detector (HAADF). RG overlays of the STEM EDX elemental maps were made using the FIJI. Phase constituents of the samples were identified by X-ray diffraction (XRD, X'Pert Pro, Panalytical Co.) with a detector using Cu Kα radiation (λ = 1.5406 Å) generated at 40 kV and 40 mA and a 2θ scanning range from 10° up to 70° in steps of 0.02°. Fourier transform infrared spectroscopy (FTIR, VERTEX 70, Bruker Corp.) was carried out to identify the functional groups of the composites with a resolution of 4 cm−1 and a scan number of 8, with a spectral region from 400 to 4000 cm−1 using 2 cm−1 steps. Micro-Raman spectra were carried out using a Renishaw inVia spectrometer in the range of 300–3500 cm−1 (recording 5 times for 10 s of each accumulation) with a wavelength of 532 nm (green laser line in a backscattering configuration using a microscope with a 100× objective, 100% power) and an acquisition time of 10 s, which had been excited from an argon ion laser. An optical microscope was used with the Raman spectrometer. Instrumented nanoindentation experiments were conducted on the polished surfaces of samples using Grindosonic tester with a Berkovich tip (Fig. 1c). Elastic modulus and hardness were calculated from the load-displacement curve using Olive-Pharr method [47]. The other instruments used to characterize the samples include inductively coupled plasma (ICP) (DV7300, Optima Co.), Contact AFM (SPM, FemtoScan Co.), and X-ray photoelectron spectroscopy (XPS, Thermo ESCALAB 250XI). Imagej 1.52d software was used in this study.
2.2. Spark plasma sintering The SPS method was chosen for sintering the powders. To control the volume change during the sintering process, the heating rate was set at 100 °C/min to 700 °C and From 700 to 950 °C, the heating rate was 50 °C/min. The samples were stored at 950 °C for 7 min and then cooled down to ambient temperature in the furnace. The pressure applied during sintering was considered to be 50 MPa. The samples obtained from the sintering stage were polished after the density measurement and were evaluated mechanically. Fig. 1 shows the dilatometric curves of the sintered samples and the nanoindentation loading curves. Fig. 1a and b show that the volumetric variation of the sample lasts up to around 900 °C, so the temperature of 950 °C is chosen for sintering and therefore, the slower heating rate is chosen from 700 to 950 °C. The slope of the dilatometric diagram varies in two parts, one about 100 °C that is related to water evaporation and one between 600 and 900 °C that is related to powders shrinkage. The shrinkage has been completed at 950 °C. 2.3. Characterization The morphology of the powders and surfaces was observed by a Field Emission Scanning Electron Microscope (FESEM, Hitachi S4700 equipped with energy dispersive X-ray spectroscopy) and a portable Scanning Electron Microscope (SEM, TM-1000). The samples were mounted in an adhesive carbon film and Au coated by sputtering for its observation. HRTEM images were obtained on a TALOS F200A with a twin lens system, X-FEG electron source, Ceta 16 M camera and a super-X EDS detector. Spatially resolved elemental analysis, with a spatial resolution higher than 2 nm, was obtained using the same TALOS microscope in 3
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Fig. 3. XRD (a), FTIR analysis (b), and Raman spectroscopy (c) of the samples before and after sintering.
converted into powders and a very small fraction of them remained in the residual solution that could be discarded. ICP results and homogeneous element map Show that the ratio of calcium to phosphate in the final powders is about 1.67 [48]. The FESEM images (Fig. 2g) show that the wrinkled rGO sheets are assembled and the HA particles are placed on them and between the layers [32,33,49]. Also, the edges of the rGO sheets have a higher HA density than the intermediate surfaces. There is some agglomeration after synthesis of powders. Fig. 2h shows the TEM image of the synthesized powders. The edges of rGO sheets have a high density of HA particles. The HA particles are very fine with rod and prism morphology and indicate that there are preferential growth directions. In this case, there is a distribution of particle size. After the sample preparation operations, smaller particles remain on the surface of graphene sheets due to more contact. Also, these particles have pores that may be due to polycrystalline structure of HA [31,50]. Fig. 3 shows the XRD, FTIR analysis, and Raman spectroscopy of the samples before and after sintering. The XRD patterns of powders (Fig. 3a) are in perfect agreement with the standard pattern of HA (JPCDS 09-0432). According to this adaptation, HA has a high purity hexagonal crystalline structure. In other words, the XRD pattern of the HA/rGO powders is quite similar to pure HA. According to studies, GO has a peak in the range of 2θ = 10. After reduction, the peak disappears and rGO peak appears from reduced GO that has a marked peak in the range of 2θ = 26. Likely, the amorphous structure of rGO reduces its XRD peaks intensity compared to the HA. The peak in 2θ = 26 for HA associated with the (002) plane is more intense than the peak for HA/ rGO, and covers the graphene peak. These peaks are sharp and high
Table 2 The characteristics of the FTIR analysis for HA. Wavenumber (cm−1)
Bond
Mode
3400–3500 1095 1035 925 565
OeH PeO(H) PeO(H) PeO(H) PeO
Stretching Stretching Stretching Stretching Bending
vibration vibration vibration vibration
Table 3 The characteristics of the FTIR analysis for GO. Wavenumber (cm−1)
Bond
Mode
3400–3500 1055 1230 1395 1620
OeH CeO CeOH CeOeH C]C
Stretching vibration Stretching vibration Stretching vibration Deformation vibration Stretching vibration
3. Results and discussion Fig. 2 shows HAADF image, elemental analysis of calcium, phosphorus and oxygen for synthesized powders, EDS analysis, FESEM image and TEM image of powders. Elemental analysis shows that all available elements are homogenously distributed that is aligned with the EDS analysis (Fig. 2d). After the hydrothermal process, some of the residual solution was subjected to ICP analysis. Due to the initial ratio of calcium to phosphate (1.67), almost all the input elements were 4
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Fig. 4. Elemental analysis of carbon, calcium, phosphorus and oxygen for sintered samples, EDS analysis (a, b), and the FESEM images of fractured surface (c–e).
graphene sheets and is located at the end of the spectrum (2700 cm−1) [34]. The G bond is related to the shaking of the E2g of carbon atoms phonon with the SP2 hybrid, While D bond is related to the symmetric oscillations of the A1g of carbon atoms with the SP3 hybrid [52]. v1PO43− (PeO) symmetric stretching peaks at 962 cm−1 and v3PO43− (PeO) symmetric stretching peaks at 1049 cm−1 can be seen in the Raman spectrum which confirms the formation of the HA phase. Despite the low amount of GO, Raman signals of this substance are more visible than HA. In the Raman spectrum of composite powders, several peaks at 428 and 576 cm−1 were found that are related to phosphate groups in HA. The D and G bonds in the rGO have not had any displacement at the Raman spectra, indicating that the composite powders have been successfully synthesized. Regarding the Raman spectrum of rGO, the peaks intensity ratio (ID/IG) is significantly increased compared to primary GO which shows the chemical and thermal reduction that has caused structural disorder in the graphene network [49,50,54]. Low-intensity, broad and sharp peaks at 428 and 962 cm−1, due to the OePeO bending mode and the PeO stretching mode of PO4 groups respectively (symmetric stretching of tetrahedral oxygen atoms around phosphorus atoms), only are revealed in the HA phase with high crystallinity. The peak at 1049 cm−1 refers to apatite phosphate groups only visible in high-quality, stoichiometric HA with high crystallinity. As it is known in Raman spectra, graphene sheets are survived and there is no chemical reaction during sintering process. The ratio of D to G bond in sintered samples has decreased compared to the powdery state, which is probably related to graphene sheets degradation and defects caused by high pressure and temperature. Also, the intensity of the D bond has changed in comparison with the powder state, which shows that the structural change occurred in sintered samples. In this case, the 2D peak is more intense and thinner than the powdery state, due to the lower number of layers after heat treatment. D and G peaks have been shifted slightly to the right in the post-sinter mode due to
intensity due to high crystallinity [49,51,52]. The homogeneity of the chemical composition and high purity makes the HA nanoparticles more stable at high temperatures. The presence of rGO does not affect the stability of HA. In addition, the absence of rGO peaks is related to the rGO layer structure and the arrangement of atoms in three dimensions [33]. As it is known, all peaks in the XRD spectrum after sintering are in full compliance with HA (Fig. 3a). According to JCPDS PDF No 29-359 that is related to tricalcium phosphate, the highest peak of αTCP is at 2θ = 30.7, the second highest peak at 22.9, and the third highest peak at 34.2. All three peaks are absent in this spectrum, which is indicative of the non-degradation of HA during sintering process. Due to the low temperature and high speed of the SPS process, no phase separation or chemical decomposition has occurred. Peak intensities have increased compared to the powders that are obviously due to the increase in crystallinity of HA [33,34]. Tables 2 and 3 show the characteristics of the FTIR analysis as shown in Fig. 3b. After reduction of GO, the bonds associated with functional groups have been significantly reduced and some of the bonds have disappeared. The peaks related to CeO and CeOH bonds in the powders have been changed to higher absorption that is related to the GO reduction. These findings indicate that the composite powders contain rGO and HA [32,33,53]. After sintering, all of the HA related bonds have retained. Also, the FTIR spectrum shows that the rGOs have not been oxidized during sintering process and if that happens, it is negligible. In terms of absorbance, there are changes in the shape of the peaks that are likely due to changes in HA particles size after sintering process and SPS pressure [33]. Fig. 3c shows Raman spectroscopy of the samples before and after SPS. The G peak determines the CeC stretching in the graphene sheets which is located at the wave number of 1350 cm−1, the peak D is related to structural defects which is located at the wave number of 1600 cm−1, and the 2D peak is related to the number of layers of the 5
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Fig. 5. AFM images of the polished fracture surface with a grit of 2500.
Fig. 6 shows the force-displacement diagrams of the sintered samples along with the mechanical properties extracted from these graphs. Considering the conditions for the samples preparation, it is likely that two mechanisms including a higher degree of GO reduction or high crystallinity of HA are responsible for the mechanical properties of the final composites. In Fig. 6b, small porosity is observed due to the HA bar shaped morphology. The use of smaller particles and proper distribution reduces the number of pores. The force-displacement curve shows that the Berkovich indenter has hit a hole in its path. The part shown with the arrow shows the contact depth where the cavity is located. These changes are more evident in samples with more porosity. In some curves, these changes appear several times. These cases involve some errors in the calculations. The results show that the hardness and the Young's modulus of the composite samples are 5.9–6.6 GPa and 132–146 GPa, respectively, because of the rGO/HA interface behavior under load-bearing conditions [32–34]. Fig. 7 shows the XPS analysis of thin layer obtained from the sintered sample. XPS is very useful to analyze the chemical composition of carbon materials [55–57]. Fig. 7a shows the signals of Ca 2p and P 2p emerge in the XPS analysis of composites and confirms that HA have been survived after sintering. Fig. 7b shows that the carboxyl group remained similar to its GO counterpart [58–61] because of strong electrostatic integration between carboxyl groups of GO and Ca2+ that prevents from reducing carboxyl groups of GO. Fig. 8 shows the TEM image, and the FESEM image of sintered sample. Fig. 8a confirms that graphene sheets have been survived after the sintering process. Also, Fig. 8b confirms the presence of rGO decorated with HA nano particles after consolidation. Fig. 9 shows the HRTEM analysis of HA||HA interface in sintered sample. FFT and IFFT analysis of the area A and B show that the planes
high pressure and high temperature. Fig. 4 shows the elemental analysis of carbon, calcium, phosphorus and oxygen for sintered samples, the EDS analysis, and the FESEM images of fracture surface. The analysis of the elements in two modes (in beam and backscatter) confirms the homogeneous distribution of these elements, which is aligned with the EDS analysis (Fig. 4a, b). Fig. 4c shows the presence of rGO in three dimensions. The graphene layers are assembled together so that the HA particles are placed between them and on their surface. As it is known, the presence of this three-dimensional structure causes incomplete compression during sintering and increases porosity and it is expected that by increasing the amount of graphene the porosity will increase equally as previously. The presence of these porosities may in part reduce mechanical properties because they are localized to crack nucleation and to focus stresses, but according to previous studies, the presence of these porosities can increase osteoconductivity for these materials. Fig. 4d and e show that the fracture type is brittle because of the inherent properties of HA and the dents that appear on the surface relate to the morphology of the primary powders. Fig. 5 shows the AFM images of the polished fracture surface with a grit of 2500 that is prepared for mechanical analysis. Spherical areas formed on the surface were created by pull out of rGO 3D structures during polishing and may result in errors. The roughness chart, corresponding to the AB line on the two-dimensional image shows that the surface average roughness (Ra) is about 7.23 nm. The RMS roughness (Rq) is about 9.39 nm and Rt is about 40.76 nm. The length of the swept line is 3.25 μm, the maximum of Z is 7.8 μm and the minimum is 7.7 μm. The findings show that the synthesized composite surface has a good quality for nanoindentation. Therefore, the Berkovich indenter has been used to obtain mechanical properties. 6
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Fig. 6. The force-displacement diagrams of the sintered samples (a, b) along with the mechanical properties extracted from these graphs (c).
of HA that are the preferred directions for crystal growth are the same and the mismatch between these phases is much less than the incoherent interface (0.25). So, the interface of these phases is coherent. Fig. 10 shows the HRTEM analysis of rGO||HA interface in crosssection of rGO after SPS, SAED, and schematic illustration of interface parameters. In Fig. 10a, HA and graphene sheets identified for analysis. IFFT analysis show the (002) planes of HA that is the preferred direction for crystal growth. It is clear that in the cross-section of the graphene sheets, (002) planes of HA are fit with rGO. According to the schematic diagram in Fig. 10c, the d-spacing of the (002) planes of HA crystals is 0.344 nm, and the d-spacing of the rGO sheets is 0.34 nm. Therefore, the mismatch between these two phases is much less than the incoherent interface (0.25). The findings indicate that in the crosssection of rGO, (002) planes of HA are in line with rGO and the interface of these two phases is likely coherent. Fig. 10b shows the electron diffraction analysis of this region. The points shown in the image correspond to (002) plane of HA. The points far away from the center, which appear dispersed and disintegrating, are related to the graphene sheets which are characterized by the overlapping of the sheets and the wrinkles in the structure [62–65]. Fig. 11 shows the HRTEM analysis of rGO||HA interface on the surface of rGO and schematic illustration of interface parameters. rGO sheets are perpendicular to (002) planes of HA (Fig. 11a). As a result, (300) planes of HA are tangent to the surface of rGO. According to the values indicated in the schematic images (Fig. 11b), the atomic alignment of the crystalline planes with rGO sheets is less than the incoherent interface (0.25). Therefore, the interface between the two phases on the rGO surface is coherent. In this research, during the synthesis of HA, its (300) planes are prior to the (100) planes [31,51].
Fig. 7. XPS analysis of the thin layer obtained from the sintered sample (a), high resolution and fitted (b).
4. Conclusion In this study, the analysis of the powders showed that the hybrid 7
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Fig. 8. TEM image (a), and the FESEM image of sintered sample (b).
Fig. 9. HRTEM analysis of HA/HA interface in sintered sample.
powders, including rGO coated with HA particles were well synthesized. Microscopic analysis confirmed the presence of graphene sheets with folding and wrinkling in the powders and consolidated samples and indicated that various preferential directions played a role in the growth of HA crystals. Sintered samples analysis showed that rGO sheets and HA have been survived after sintering process. Mechanical analysis results showed that the hardness and the Young's modulus of
the composite samples are 5.9–6.6 GPa and 132–146 GPa respectively, because of the rGO/HA interface behavior under load-bearing conditions. The results of interface analysis (HA||rGO) showed that the HA is coherently connected by its (300)-planes with the surface of the rGO sheets and this coherency is accomplished in the rGO cross-section with the (002) planes of HA. The results of this study could increase the use of this composite type in orthopedic applications.
8
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Fig. 10. HRTEM analysis of rGO/HA interface in cross-section of rGO (a), SAED (b), and schematic illustration of interface parameters.
Fig. 11. HRTEM analysis of rGO/HA interface on the surface of rGO and schematic illustration of interface parameters (b). 9
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Preparation of reduced graphene oxide/hydroxyapatite nanocomposite and evaluation of graphene sheets/hydroxyapatite interface
T
Hassan Nosratia, Rasoul Sarraf Mamoorya, , Dang Quang Svend Leb, Cody Eric Büngerb ⁎
a b
Department of Materials Engineering, Tarbiat Modares University, Tehran, Iran Department of Clinical Medicine, Aarhus University, Denmark
ARTICLE INFO
ABSTRACT
Keywords: Hydroxyapatite Graphene Hydrothermal Gas injection Interface
In this investigation, reduced graphene oxide/hydroxyapatite (rGO/HA) hybrid powders have been synthesized using hydrogen gas injection into hydrothermal autoclave. The powders were then consolidated with spark plasma sintering. The results showed that this method of synthesis caused the obtained powders to have high crystallinity. Microscopic analysis confirmed the presence of rGO sheets with folding and wrinkling in the nanocomposite and indicated that various crystalline planes such as (002) and (300) played a role in the growth of hydroxyapatite crystals. The results of interface analysis (HA||rGO) showed that the HA is coherently connected by its (300)-planes with the surface of the rGO sheets and this coherency is accomplished in the rGO crosssection with the (002) planes of HA. The hardness and the Young's modulus of the composite samples were 5.9–6.6 GPa and 132–146 GPa respectively.
1. Introduction Synthesis of hydroxyapatite (HA) with nanotechnological approaches in nanosized dimensions and in various forms such as nanotubes and nanorods have improved the mechanical properties of this ceramic material [1–3]. The chemical composition (Ca10(PO4)6(OH)2), the ratio of calcium to phosphate (C/P ≈ 1.67), and the crystalline structure of HA are very similar to the mineral part of the human skeletal system. Also, having properties such as bioactivity and osteoconductivity make HA one of the most valuable materials for orthopedic applications [4–6]. The applications of this bioceramic are very broad include orthopedics, biosensors, catalysts and drug delivery [7–12]. However, the poor mechanical properties of HA such as fracture toughness, intrinsic brittleness, and poor wear resistance still limit the applications of HA despite its excellent biomaterial properties [13–15]. To make HA usable as an implant, one of the strategies is to enhance its mechanical properties with a second, reinforcing, phase. Various materials have been investigated including titanium oxide, aluminum oxide, and carbon nanostructures [16–20]. But among these materials, graphene, with its unique mechanical properties (Young's modulus ≈ 1 TPa and fracture strength ≈ 130 GPa), has attracted a great deal of attention and research [21,22]. Graphene has a honeycomb-like structure of a single carbon atom with the SP2 hybrid [23–25]. Its high
⁎
specific surface area (2630 m2 g−1) has made its reinforcing properties more effective. Also, its good biological properties have led to much research for medical applications including orthopedic surgery, bioimaging, and drug delivery [26–30]. Several studies have been published on the use of graphene and its derivatives to improve the mechanical and biological properties of HA [31–34]. The results of these investigations showed that the mechanical and biological properties of HA were improved by the addition of graphene. Graphene has been used as filler that blocks the crack growth with its two-dimensional structure [35–37]. One of the methods to synthesis of graphene is a hydrothermal method that the graphene oxide (GO) is used as precursor. Hydrogen under the high pressure and high temperature of the hydrothermal autoclave is the main factor in the reduction of GO, assembly as a gel, and forming three-dimensional rGOs. Using hydrothermal method is an in situ mechanism for the synthesis of HA-rGO hybrid powders because the presence of oxygen agents on GO sheets causes nucleation and growth of HA particles on the graphene surface [38–41]. Due to the high mechanical properties of graphene recovered from GO, the degree of reduction is very important. One of the most attractive methods is utilizing of hydrogen gas as a reducing agent. In one published report, a mixture of argon and hydrogen gases was used alongside palladium and the process temperature was 200 °C. Palladium caused hydrogen gas to automatically reduce the oxygen present on the graphene oxide surface [42]. In another
Corresponding author. E-mail address: [email protected] (R.S. Mamoory).
https://doi.org/10.1016/j.diamond.2019.107561 Received 18 July 2019; Received in revised form 21 September 2019; Accepted 27 September 2019 0925-9635/ © 2019 Elsevier B.V. All rights reserved.
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electron microscopy (HRTEM), X-ray photoelectron spectroscopy, and nanoindentation were used to investigate the type (coherent or incoherent), chemical bonds, and mechanical behavior of the interface.
Table 1 The primary chemicals used in the powder synthesis phase. No.
Chem.
Co.
Purity
Formulation
1
Calcium nitrate tetrahydrate Diammonium hydrogenphosphate DMF Ammonium solution GO
Merck
> 99%
Ca(NO3)2·4H2O
Merck
> 99%
(NH4)2HPO4
Sigma-Aldrich Merck Abalonyx
> 99.8% 25% 25 g/l DMF
(CH3)2NC(O)H NH4OH COxHy
2 3 4 5
2. Materials and methods According to the chemical formula of HA, the Ca/P ratio should be 1.67. Therefore, for the synthesis of primary powders, 7.84 g of calcium nitrate tetrahydrate was dissolved in 200 ml deionized water and 1.32 g of diammonium hydrogenphosphate dissolved in 150 ml deionized water to obtain two solutions containing calcium and phosphate ions as the basic HA ions.
report, hydrogen-rich water was used to reduce GO [43]. Further researches have been published in which hydrogen gas is injected into a hydrothermal autoclave, and the results showed that the presence of hydrogen gas in hydrothermal conditions would also reduce GO [44,45]. The mechanical behavior of composites made of HA and rGO depends on interface type between the two phases. Therefore, studying the details and behavior of the interface is very important. In this study, the hybrid nanostructured powders were synthesized using a hydrothermal method utilizing hydrogen gas injection to increase the reduction rate of GO. However, to have a greater effect on reduction of GO, hydrogen gas needs to be decomposed by the catalysts to the hydrogen atoms, because the hydrogen gas has little solubility in solution. So, hydrothermal temperature and time will probably have a stronger role in the synthesis of powders [42–46]. In order to investigate the effect of this process on the final nanocomposite properties, the spark plasma sintering (SPS) method has been used to consolidate these powders and to fabricate bulk samples. High-resolution transmission
2.1. Powders preparation The following steps were used to synthesize the powders, respectively. a) The solution containing calcium ions was added dropwise to a 20 ml stirred suspension of GO (3.13 mg/ml) (HA/1.5% rGO) with stirring continued for 1 h. b) The solution containing phosphate ions was dropwise added to the solution. c) The pH of the solutions was adjusted to 11 with ammonium solution. d) The resulting solution was poured into the Teflon (PTFE) vessel and transferred to the hydrothermal autoclave. The hydrothermal process was carried out for 5 h at 180 °C by injection of hydrogen gas at 10 bar. e) The powders were dried at oven for 24 h at 60 °C. The volume of the PTFE container was 340 ml and the total pressure was 25 bar when using gas injection. The primary chemicals used in the powder synthesis phase with the specifications are listed in Table 1.
Fig. 1. Dilatometric curves of the sintered samples (a, b) and the nanoindentation loading curves (c). 2
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Fig. 2. HAADF image (a), elemental analysis of calcium/phosphorus (b), phosphorus (c), oxygen (e), and calcium (f), for synthesized powders, EDS analysis (d), FESEM image (g), and TEM image (h) of powders.
STEM mode. Exposure times of 5 min were used to create elemental distribution maps, with satisfactory counting statistics, while minimizing potential problems such as beam damage and specimen drift. STEM images were obtained using a high angle annular dark field detector (HAADF). RG overlays of the STEM EDX elemental maps were made using the FIJI. Phase constituents of the samples were identified by X-ray diffraction (XRD, X'Pert Pro, Panalytical Co.) with a detector using Cu Kα radiation (λ = 1.5406 Å) generated at 40 kV and 40 mA and a 2θ scanning range from 10° up to 70° in steps of 0.02°. Fourier transform infrared spectroscopy (FTIR, VERTEX 70, Bruker Corp.) was carried out to identify the functional groups of the composites with a resolution of 4 cm−1 and a scan number of 8, with a spectral region from 400 to 4000 cm−1 using 2 cm−1 steps. Micro-Raman spectra were carried out using a Renishaw inVia spectrometer in the range of 300–3500 cm−1 (recording 5 times for 10 s of each accumulation) with a wavelength of 532 nm (green laser line in a backscattering configuration using a microscope with a 100× objective, 100% power) and an acquisition time of 10 s, which had been excited from an argon ion laser. An optical microscope was used with the Raman spectrometer. Instrumented nanoindentation experiments were conducted on the polished surfaces of samples using Grindosonic tester with a Berkovich tip (Fig. 1c). Elastic modulus and hardness were calculated from the load-displacement curve using Olive-Pharr method [47]. The other instruments used to characterize the samples include inductively coupled plasma (ICP) (DV7300, Optima Co.), Contact AFM (SPM, FemtoScan Co.), and X-ray photoelectron spectroscopy (XPS, Thermo ESCALAB 250XI). Imagej 1.52d software was used in this study.
2.2. Spark plasma sintering The SPS method was chosen for sintering the powders. To control the volume change during the sintering process, the heating rate was set at 100 °C/min to 700 °C and From 700 to 950 °C, the heating rate was 50 °C/min. The samples were stored at 950 °C for 7 min and then cooled down to ambient temperature in the furnace. The pressure applied during sintering was considered to be 50 MPa. The samples obtained from the sintering stage were polished after the density measurement and were evaluated mechanically. Fig. 1 shows the dilatometric curves of the sintered samples and the nanoindentation loading curves. Fig. 1a and b show that the volumetric variation of the sample lasts up to around 900 °C, so the temperature of 950 °C is chosen for sintering and therefore, the slower heating rate is chosen from 700 to 950 °C. The slope of the dilatometric diagram varies in two parts, one about 100 °C that is related to water evaporation and one between 600 and 900 °C that is related to powders shrinkage. The shrinkage has been completed at 950 °C. 2.3. Characterization The morphology of the powders and surfaces was observed by a Field Emission Scanning Electron Microscope (FESEM, Hitachi S4700 equipped with energy dispersive X-ray spectroscopy) and a portable Scanning Electron Microscope (SEM, TM-1000). The samples were mounted in an adhesive carbon film and Au coated by sputtering for its observation. HRTEM images were obtained on a TALOS F200A with a twin lens system, X-FEG electron source, Ceta 16 M camera and a super-X EDS detector. Spatially resolved elemental analysis, with a spatial resolution higher than 2 nm, was obtained using the same TALOS microscope in 3
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Fig. 3. XRD (a), FTIR analysis (b), and Raman spectroscopy (c) of the samples before and after sintering.
converted into powders and a very small fraction of them remained in the residual solution that could be discarded. ICP results and homogeneous element map Show that the ratio of calcium to phosphate in the final powders is about 1.67 [48]. The FESEM images (Fig. 2g) show that the wrinkled rGO sheets are assembled and the HA particles are placed on them and between the layers [32,33,49]. Also, the edges of the rGO sheets have a higher HA density than the intermediate surfaces. There is some agglomeration after synthesis of powders. Fig. 2h shows the TEM image of the synthesized powders. The edges of rGO sheets have a high density of HA particles. The HA particles are very fine with rod and prism morphology and indicate that there are preferential growth directions. In this case, there is a distribution of particle size. After the sample preparation operations, smaller particles remain on the surface of graphene sheets due to more contact. Also, these particles have pores that may be due to polycrystalline structure of HA [31,50]. Fig. 3 shows the XRD, FTIR analysis, and Raman spectroscopy of the samples before and after sintering. The XRD patterns of powders (Fig. 3a) are in perfect agreement with the standard pattern of HA (JPCDS 09-0432). According to this adaptation, HA has a high purity hexagonal crystalline structure. In other words, the XRD pattern of the HA/rGO powders is quite similar to pure HA. According to studies, GO has a peak in the range of 2θ = 10. After reduction, the peak disappears and rGO peak appears from reduced GO that has a marked peak in the range of 2θ = 26. Likely, the amorphous structure of rGO reduces its XRD peaks intensity compared to the HA. The peak in 2θ = 26 for HA associated with the (002) plane is more intense than the peak for HA/ rGO, and covers the graphene peak. These peaks are sharp and high
Table 2 The characteristics of the FTIR analysis for HA. Wavenumber (cm−1)
Bond
Mode
3400–3500 1095 1035 925 565
OeH PeO(H) PeO(H) PeO(H) PeO
Stretching Stretching Stretching Stretching Bending
vibration vibration vibration vibration
Table 3 The characteristics of the FTIR analysis for GO. Wavenumber (cm−1)
Bond
Mode
3400–3500 1055 1230 1395 1620
OeH CeO CeOH CeOeH C]C
Stretching vibration Stretching vibration Stretching vibration Deformation vibration Stretching vibration
3. Results and discussion Fig. 2 shows HAADF image, elemental analysis of calcium, phosphorus and oxygen for synthesized powders, EDS analysis, FESEM image and TEM image of powders. Elemental analysis shows that all available elements are homogenously distributed that is aligned with the EDS analysis (Fig. 2d). After the hydrothermal process, some of the residual solution was subjected to ICP analysis. Due to the initial ratio of calcium to phosphate (1.67), almost all the input elements were 4
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Fig. 4. Elemental analysis of carbon, calcium, phosphorus and oxygen for sintered samples, EDS analysis (a, b), and the FESEM images of fractured surface (c–e).
graphene sheets and is located at the end of the spectrum (2700 cm−1) [34]. The G bond is related to the shaking of the E2g of carbon atoms phonon with the SP2 hybrid, While D bond is related to the symmetric oscillations of the A1g of carbon atoms with the SP3 hybrid [52]. v1PO43− (PeO) symmetric stretching peaks at 962 cm−1 and v3PO43− (PeO) symmetric stretching peaks at 1049 cm−1 can be seen in the Raman spectrum which confirms the formation of the HA phase. Despite the low amount of GO, Raman signals of this substance are more visible than HA. In the Raman spectrum of composite powders, several peaks at 428 and 576 cm−1 were found that are related to phosphate groups in HA. The D and G bonds in the rGO have not had any displacement at the Raman spectra, indicating that the composite powders have been successfully synthesized. Regarding the Raman spectrum of rGO, the peaks intensity ratio (ID/IG) is significantly increased compared to primary GO which shows the chemical and thermal reduction that has caused structural disorder in the graphene network [49,50,54]. Low-intensity, broad and sharp peaks at 428 and 962 cm−1, due to the OePeO bending mode and the PeO stretching mode of PO4 groups respectively (symmetric stretching of tetrahedral oxygen atoms around phosphorus atoms), only are revealed in the HA phase with high crystallinity. The peak at 1049 cm−1 refers to apatite phosphate groups only visible in high-quality, stoichiometric HA with high crystallinity. As it is known in Raman spectra, graphene sheets are survived and there is no chemical reaction during sintering process. The ratio of D to G bond in sintered samples has decreased compared to the powdery state, which is probably related to graphene sheets degradation and defects caused by high pressure and temperature. Also, the intensity of the D bond has changed in comparison with the powder state, which shows that the structural change occurred in sintered samples. In this case, the 2D peak is more intense and thinner than the powdery state, due to the lower number of layers after heat treatment. D and G peaks have been shifted slightly to the right in the post-sinter mode due to
intensity due to high crystallinity [49,51,52]. The homogeneity of the chemical composition and high purity makes the HA nanoparticles more stable at high temperatures. The presence of rGO does not affect the stability of HA. In addition, the absence of rGO peaks is related to the rGO layer structure and the arrangement of atoms in three dimensions [33]. As it is known, all peaks in the XRD spectrum after sintering are in full compliance with HA (Fig. 3a). According to JCPDS PDF No 29-359 that is related to tricalcium phosphate, the highest peak of αTCP is at 2θ = 30.7, the second highest peak at 22.9, and the third highest peak at 34.2. All three peaks are absent in this spectrum, which is indicative of the non-degradation of HA during sintering process. Due to the low temperature and high speed of the SPS process, no phase separation or chemical decomposition has occurred. Peak intensities have increased compared to the powders that are obviously due to the increase in crystallinity of HA [33,34]. Tables 2 and 3 show the characteristics of the FTIR analysis as shown in Fig. 3b. After reduction of GO, the bonds associated with functional groups have been significantly reduced and some of the bonds have disappeared. The peaks related to CeO and CeOH bonds in the powders have been changed to higher absorption that is related to the GO reduction. These findings indicate that the composite powders contain rGO and HA [32,33,53]. After sintering, all of the HA related bonds have retained. Also, the FTIR spectrum shows that the rGOs have not been oxidized during sintering process and if that happens, it is negligible. In terms of absorbance, there are changes in the shape of the peaks that are likely due to changes in HA particles size after sintering process and SPS pressure [33]. Fig. 3c shows Raman spectroscopy of the samples before and after SPS. The G peak determines the CeC stretching in the graphene sheets which is located at the wave number of 1350 cm−1, the peak D is related to structural defects which is located at the wave number of 1600 cm−1, and the 2D peak is related to the number of layers of the 5
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Fig. 5. AFM images of the polished fracture surface with a grit of 2500.
Fig. 6 shows the force-displacement diagrams of the sintered samples along with the mechanical properties extracted from these graphs. Considering the conditions for the samples preparation, it is likely that two mechanisms including a higher degree of GO reduction or high crystallinity of HA are responsible for the mechanical properties of the final composites. In Fig. 6b, small porosity is observed due to the HA bar shaped morphology. The use of smaller particles and proper distribution reduces the number of pores. The force-displacement curve shows that the Berkovich indenter has hit a hole in its path. The part shown with the arrow shows the contact depth where the cavity is located. These changes are more evident in samples with more porosity. In some curves, these changes appear several times. These cases involve some errors in the calculations. The results show that the hardness and the Young's modulus of the composite samples are 5.9–6.6 GPa and 132–146 GPa, respectively, because of the rGO/HA interface behavior under load-bearing conditions [32–34]. Fig. 7 shows the XPS analysis of thin layer obtained from the sintered sample. XPS is very useful to analyze the chemical composition of carbon materials [55–57]. Fig. 7a shows the signals of Ca 2p and P 2p emerge in the XPS analysis of composites and confirms that HA have been survived after sintering. Fig. 7b shows that the carboxyl group remained similar to its GO counterpart [58–61] because of strong electrostatic integration between carboxyl groups of GO and Ca2+ that prevents from reducing carboxyl groups of GO. Fig. 8 shows the TEM image, and the FESEM image of sintered sample. Fig. 8a confirms that graphene sheets have been survived after the sintering process. Also, Fig. 8b confirms the presence of rGO decorated with HA nano particles after consolidation. Fig. 9 shows the HRTEM analysis of HA||HA interface in sintered sample. FFT and IFFT analysis of the area A and B show that the planes
high pressure and high temperature. Fig. 4 shows the elemental analysis of carbon, calcium, phosphorus and oxygen for sintered samples, the EDS analysis, and the FESEM images of fracture surface. The analysis of the elements in two modes (in beam and backscatter) confirms the homogeneous distribution of these elements, which is aligned with the EDS analysis (Fig. 4a, b). Fig. 4c shows the presence of rGO in three dimensions. The graphene layers are assembled together so that the HA particles are placed between them and on their surface. As it is known, the presence of this three-dimensional structure causes incomplete compression during sintering and increases porosity and it is expected that by increasing the amount of graphene the porosity will increase equally as previously. The presence of these porosities may in part reduce mechanical properties because they are localized to crack nucleation and to focus stresses, but according to previous studies, the presence of these porosities can increase osteoconductivity for these materials. Fig. 4d and e show that the fracture type is brittle because of the inherent properties of HA and the dents that appear on the surface relate to the morphology of the primary powders. Fig. 5 shows the AFM images of the polished fracture surface with a grit of 2500 that is prepared for mechanical analysis. Spherical areas formed on the surface were created by pull out of rGO 3D structures during polishing and may result in errors. The roughness chart, corresponding to the AB line on the two-dimensional image shows that the surface average roughness (Ra) is about 7.23 nm. The RMS roughness (Rq) is about 9.39 nm and Rt is about 40.76 nm. The length of the swept line is 3.25 μm, the maximum of Z is 7.8 μm and the minimum is 7.7 μm. The findings show that the synthesized composite surface has a good quality for nanoindentation. Therefore, the Berkovich indenter has been used to obtain mechanical properties. 6
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Fig. 6. The force-displacement diagrams of the sintered samples (a, b) along with the mechanical properties extracted from these graphs (c).
of HA that are the preferred directions for crystal growth are the same and the mismatch between these phases is much less than the incoherent interface (0.25). So, the interface of these phases is coherent. Fig. 10 shows the HRTEM analysis of rGO||HA interface in crosssection of rGO after SPS, SAED, and schematic illustration of interface parameters. In Fig. 10a, HA and graphene sheets identified for analysis. IFFT analysis show the (002) planes of HA that is the preferred direction for crystal growth. It is clear that in the cross-section of the graphene sheets, (002) planes of HA are fit with rGO. According to the schematic diagram in Fig. 10c, the d-spacing of the (002) planes of HA crystals is 0.344 nm, and the d-spacing of the rGO sheets is 0.34 nm. Therefore, the mismatch between these two phases is much less than the incoherent interface (0.25). The findings indicate that in the crosssection of rGO, (002) planes of HA are in line with rGO and the interface of these two phases is likely coherent. Fig. 10b shows the electron diffraction analysis of this region. The points shown in the image correspond to (002) plane of HA. The points far away from the center, which appear dispersed and disintegrating, are related to the graphene sheets which are characterized by the overlapping of the sheets and the wrinkles in the structure [62–65]. Fig. 11 shows the HRTEM analysis of rGO||HA interface on the surface of rGO and schematic illustration of interface parameters. rGO sheets are perpendicular to (002) planes of HA (Fig. 11a). As a result, (300) planes of HA are tangent to the surface of rGO. According to the values indicated in the schematic images (Fig. 11b), the atomic alignment of the crystalline planes with rGO sheets is less than the incoherent interface (0.25). Therefore, the interface between the two phases on the rGO surface is coherent. In this research, during the synthesis of HA, its (300) planes are prior to the (100) planes [31,51].
Fig. 7. XPS analysis of the thin layer obtained from the sintered sample (a), high resolution and fitted (b).
4. Conclusion In this study, the analysis of the powders showed that the hybrid 7
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Fig. 8. TEM image (a), and the FESEM image of sintered sample (b).
Fig. 9. HRTEM analysis of HA/HA interface in sintered sample.
powders, including rGO coated with HA particles were well synthesized. Microscopic analysis confirmed the presence of graphene sheets with folding and wrinkling in the powders and consolidated samples and indicated that various preferential directions played a role in the growth of HA crystals. Sintered samples analysis showed that rGO sheets and HA have been survived after sintering process. Mechanical analysis results showed that the hardness and the Young's modulus of
the composite samples are 5.9–6.6 GPa and 132–146 GPa respectively, because of the rGO/HA interface behavior under load-bearing conditions. The results of interface analysis (HA||rGO) showed that the HA is coherently connected by its (300)-planes with the surface of the rGO sheets and this coherency is accomplished in the rGO cross-section with the (002) planes of HA. The results of this study could increase the use of this composite type in orthopedic applications.
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Fig. 10. HRTEM analysis of rGO/HA interface in cross-section of rGO (a), SAED (b), and schematic illustration of interface parameters.
Fig. 11. HRTEM analysis of rGO/HA interface on the surface of rGO and schematic illustration of interface parameters (b). 9
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