calcium phosphate coating by pulse electrodeposition on anodized titanium: Biocorrosion and mechanical behavior

Author’s Accepted Manuscript Development of graphene oxide/calcium phosphate coating by pulse electrodeposition on anodized titanium: biocorrosion and mechanical behavior Leila Fathyunes, Jafar Khalil-Allafi, Maryam Moosavifar www.elsevier.com/locate/jmbbm

PII: DOI: Reference:

S1751-6161(18)30825-7 https://doi.org/10.1016/j.jmbbm.2018.11.011 JMBBM3067

To appear in: Journal of the Mechanical Behavior of Biomedical Materials Received date: 27 May 2018 Revised date: 12 November 2018 Accepted date: 12 November 2018 Cite this article as: Leila Fathyunes, Jafar Khalil-Allafi and Maryam Moosavifar, Development of graphene oxide/calcium phosphate coating by pulse electrodeposition on anodized titanium: biocorrosion and mechanical behavior, Journal of the Mechanical Behavior of Biomedical Materials, https://doi.org/10.1016/j.jmbbm.2018.11.011 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Development of graphene oxide/calcium phosphate coating by pulse electrodeposition on anodized titanium: biocorrosion and mechanical behavior Leila Fathyunesa*, Jafar Khalil-Allafib, Maryam Moosavifarc a

Department of Materials Science Engineering, University of Bonab, 5551761167, Bonab, Iran Research Center for Advanced Materials, Faculty of Materials Engineering, Sahand University of Technology, 5133511996, Tabriz, Iran. c Department of Chemistry, Faculty of Science, University of Maragheh, 55181-83111, Maragheh, Iran b

* Corresponding author: E-mail: [email protected]

Abstract In this work, graphene oxide (GO) reinforcement was used to improve the strength and fracture toughness of the calcium phosphate (CaP) coating applied on the anodized titanium using pulse electrodeposition. The results showed that the CaP coating consisted of mixed phases of octacalcium phosphate (OCP), dicalcium phosphate dehydrate (DCPD) and hydroxyapatite (HAp); however, compositing of this coating with GO caused deposition of the pure HAp phase. Moreover, the nanohardness and elastic modulus for the CaP-GO coating increased over 52% and 41%, respectively, as compared to those measured for the GO-free coating. An improvement of about 16% in the adhesion strength of the CaP coating composited with GO to the anodized titanium was also arisen from improving integrity, crystallinity and decreasing the elastic modulus mismatch of this coating with titanium substrate. Finally, uniformity in the microstructure and more biostability of the CaP-GO coating led to its better protection against the corrosion of anodized titanium. Keywords: calcium phosphate, graphene oxide, pulse electrodeposition, biocorrosion resistance, mechanical properties.

1. Introduction During the last decades there is a remarkably increasing demand for using tissue-engineered devices in order to repair or replace damaged human parts [1,2]. Substantial percentage of the implants used in the orthopedic surgery are made of metallic materials such as 316L stainless

steel, cobalt-chromium, titanium and its alloys [1,3]. Among them, titanium and its alloys are widely recognized as very promising biomaterials because of providing the appropriate mechanical properties required to replace the hard tissues such as the high strength-to-weight ratio, fracture toughness and wear resistance [3–5]. The elastic modulus of titanium (98-105 GPa) has less difference with that of bone tissues (20–30 GPa) when it is compared to the other metals intended for biomedical applications such as Ti–6Al–4V alloy (105-114 GPa), stainless steel (193-205 Gpa) and cobalt-chromium alloy (200-240 Gpa) [6,7]. According to Wolff’s law of bone remodeling, the stronger mismatch between the Young’s modulus of bone and implant is associated with the negative effect of stress-shielding, which in prolonged times causes the bone resorption, implant loosening and its failure [1,8]. The most challenging issue on the application of bio-inert titanium is its low osseointegration and bone-bonding ability [3,9,10]. In order to overcome this problem, various bioactive materials have been applied as a coating on titanium implants [3]. The bioactive materials could form a strong bonding with bone tissue as well as guide the growth of bone along the interface [11,12]. Among these materials, the calcium phosphate (CaP) ceramics such as hydroxyapatite (HAp) because of having the chemical compositions similar to that of natural bone mostly use in dental implants, orthopedics and drug delivery systems [12]. The main drawback of HAp is its low wear resistance, fracture strength (less than 120 MPa), fracture toughness (0.8–1.2 MPam1/2) as compared to the human bone with fracture strength of ~150 MPa and fracture toughness of 2–12 MPam1/2 [13–15]. Consequently, CaP ceramics such as HAp are an unsuitable candidate under biomechanical loading conditions [16,17] and, therefore, to combine their bioactivity with the mechanical properties of metals, these ceramics apply as a coating on the load-bearing metallic implants [18,19]. However, low mechanical and tribological characteristics on the implant surface leads to its failure under friction and wear conditions [20]. Recently, a scientific attention has been devoted to the carbon nanomaterials such as carbon nanotubes (CNTs), graphene or its derivatives for improving strength and fracture toughness of the brittle HAp [21–24]. Graphene and its derivatives possess a series of merits including large specific surface area (up to 2630 m2g–1), excellent mechanical properties (e.g. tensile strength of 130 GPa and Young’s modulus of 0.5–1 TPa), biocompatibility, biostability, antibacterial properties and the ability to initialize the apatite mineralization [24–

26]. Besides, graphene and its derivatives synthesis in a large scale with significantly low cost and less toxicity compared to CNTs [24,27]. The energy required to pull out a sheet-like graphene reinforcement is also greater than that of a single-walled or multi-walled CNTs, mainly owing to its higher contact area with the matrix [28]. However, the poor dispersion of graphene reinforcement as well as existence a weak interfacial adhesion between graphene and matrix hinder to achieve the expected mechanical properties in the graphene-ceramic composites [25]. In this regard, the oxygen-containing functional groups on the surface of graphene oxide (GO) could improve its dispersion [12] and biocompatibility [29]. Up to date, various methods have been used to apply HAp-GO coatings such as electrophoretic [30], biomimetic [29], vacuum cold spraying [31], and electrodeposition [32,33]. The need for low-cost equipments, good control of the coating composition and the possibility to coat complex shapes have led to an increased research in the electrodeposition method. Li et al. [34] reported that adding of GO reinforcement exhibited the HAp-based coating with less cracks. Zeng et al [32] applied the HAp-GO coating using electrodeposition at low current density of 0.6 mA/cm2 and reported improving the uniformity and crystallinity of the coating in the presence of GO. They further post-treated the electrodeposited coatings in 0.1 mol/L NaOH for 2 h. It is clear that this step was carried out due to the conversion of all other CaP phases like octa calcium phosphate (OCP) and dicalcium phosphate dehydrate (DCPD) to the most stable HAp phase [19]. In the above perspective, we aimed to directly apply the CaP-GO composite coating containing the predominant phase of HAp by electrodeposition at higher current density. It should be mentioned that producing a considerable amount of H2 gas through the water reduction at high current density could decrease the adhesion strength of the coating to the substrate [33]. To overcome this problem, we used the anodizing to modify the surface of titanium before applying the CaP-based coatings as well as the pulse current to provide an off time for the escape of H2 gases from the substrate surface during electrodeposition. In previous papers, we studied the biocompatibility and biomineralization ability of the CaP-GO composite coating applied by pulse electrodeposition in the absence [35] and presence of ultrasonic waves [33]. To the best of our knowledge, there is lack of information about the effect of GO reinforcement on the CaP bioceramic being appropriate for use as a coating on the load-

bearing metallic implants. Thus, in the present work, the surface morphology, phase composition and mechanical properties of the CaP-GO coating applied by pulse electrodeposition on the anodized titanium were studied. Moreover, the role of composition and microstructure of the coating on its ability to protect the anodized titanium substrate against the corrosion was evaluated.

2. Materials and methods 2. 1. Synthesis of graphene oxide (GO) According to the modified Hummers’ method, GO was synthesized through the harsh oxidation of graphite (G) purchased from Sigma-Aldrich in a mixture of 69 mL sulfuric acid (H2SO4, 98%, merk), 1.5 g sodium nitrate (NaNO3, Sigma Aldrich) and 6 g potassium permanganate (KMnO4, 99.9%, Sigma Aldrich). The obtained graphite oxide was washed with 10% HCl aqueous solution, deionized water and alcohol, respectively, until the pH of the solution became neutral. At the end, the exfoliation of graphite oxide into GO was carried out by an ultrasound treatment of the aqueous solution of graphite oxide with the power of 250 W for 1h. 2.2. Anodizing of titanium substrate Before applying the CaP-based coatings, the surface modification of pure titanium (99.9%, ASTM F67 Grade 1) with dimensions of 10×10 ×1 mm was carried out using anodizing. First, the titanium substrate was mechanically abraded with SiC papers from 120 to 5000 grit and was chemically polished in HNO3/HF (3:1 in volume ratio) solution for 30s. Then, this substrate was anodized in the ethylene glycol-based electrolyte containing 4wt. % H3PO4 and 0.25 wt. % HF at a constant voltage of 60 V and temperature of 35 °C for 50 min. At the next step, the anodized titanium was immersed in 0.5 M NaOH for 2 min at 50 °C and was heat treated for 1h at 500 °C to crystallize the amorphous TiO2. 2.3. Applying CaP and CaP-GO coatings Preparation of the electrolyte used for electrodeposition of the CaP-based coatings was initially started with separately dissolution of 0.008 M Ca(NO3)2.4H2O and 0.005 M NH4H2PO4 in double distilled water to maintain the Ca2+ to PO43- molar ratio of 1.67 for stoichiometric HAp. Then, 1 M NaNO3 and 6mL/L H2O2 were added to above electrolyte and its pH was adjusted to 6 using

dilute tris-hydroxymethyl aminomethane buffer. For applying the CaP-GO coating, 30 mg/L of GO was also added to the electrolyte. A magnetic stirring of the electrolyte was controlled at a speed of 120 rpm. To decrease the carbonate contamination, the electrolyte was purged during electrodeposition by argon gas. The anodized titanium as the cathode and graphite as the anode were set with a fixed distance of 2 cm inside the electrolyte. The CaP and CaP-GO coatings were applied by the pulse current electrodeposition at a constant current density of 15 mA/cm2, duty cycle of 0.1 [ton: 1s , toff: 9s] and temperature of 65 °C. 2.4. Characterization The surface morphology of the anodized titanium and coatings were observed by a field emission scanning electron microscope (FESEM, Tescan MIRA3 FEG-SEM) operated at an accelerating voltage of 15 kV, equipped with an energy dispersive spectrometer (EDS). The chemical bands of the synthesized GO and the powdery samples scrapped from the coatings were studied using a Fourier transform infrared absorption spectra (FTIR, BRUKER Tensor 27) over a frequency range of 400– 4000 cm-1. Transmission electron microscopy (TEM, CM30 Philips) was used to analyze the morphology of GO and coatings. In preparing the TEM samples, GO and the scraped coatings were dispersed in ethanol and then a few drops of suspension was deposited on copper-coated carbon grid. After evaporation of the ethanol solvent, samples were observed using TEM at an acceleration voltage of 150 kV. The crystal structure of the coatings was also characterized by Selected Area Electron Diffraction (SAED) pattern. The surface composition and elemental state of the CaP-GO coating, and the CaP coating as the reference, were determined by X-ray photoelectron spectroscopy (XPS, Bestec GmbH). This test was carried out under ultra-high vacuum of 10-10 mbar using a monochromated Al Kα radiation (X-ray source, hν= 1486.6 eV) with a 45° electron collection angle. The survey spectra in the range of 0–1200 eV was recorded with an energy resolution of 1.5 eV. Moreover, high-resolution spectra over C1s, Ca2p, P2p and O1s peaks were acquired in the pass energy of 20 eV and 0.03 eV steps. The electrochemical measurements of open circuit potential (OCP) and potentiodynamic polarization were also carried out to evaluate the biocorrosion behavior of bare, anodized, and

coated samples in Ringer's solution at 37 °C using Potentiostat/Galvanostat Autolab (model PGSTAT302N). The saturated calomel electrode and platinized platinum electrode were considered as the reference and auxiliary electrodes, respectively. Before the beginning of each test, the samples were immersed for 30 min in the Ringer's solution to establish the open circuit potential (OCP). The potential was increased at a scan rate of 1 mV/s by sweeping from -800 mV to +500 mV. The corrosion potential (Ecorr) and corrosion current density (Icorr) were determined by Tafel extrapolation method from the polarization curve. Moreover, the protection efficiency (η%) of titanium substrate resulting from anodizing as well as applying the CaP and CaP-GO coatings was estimated using the Eq. (1) [36]:

( )

Eq. (7)

Where, i0corr and icorr are the corrosion current density of the bare titanium and the anodized or anodized/coated titanium, respectively. The topography of the CaP and CaP-GO coated samples was scanned using atomic force microscopy (AFM, Flex, Nonosurf Company) in noncontact mode over the surface area of 10 × 10 µm2. The surface roughness parameters including arithmetic mean deviation of the roughness profile (Sa), root mean square roughness (Sq), maximum height of peaks (Sp), maximum depth of valleys (Sv) and maximum height within the evaluation length (Sy=|Sp|+|Sv|) were calculated using the Nanosurf C3000 software. The mechanical properties of the coatings were assessed using a nanoindentation machine (TriboScope®, Hysitron Inc.) equipped with a Berkovich diamond indenter. The coated samples were mounted in resin and the nanoindentation test was carried out on the randomly selected sites of their polished cross-section. The load–penetration depth curves were recorded in a controlled constant maximum load of 6000μN, loading time of 30 s, holding time of 10 s and unloading time of 30 s. Fig. 1 shows the measurable parameters of the nanoindentation technique from the typical load-penetration depth curve. In this figure, hmax is the maximum penetration depth corresponded to the maximum load of Pmax and hf is the final penetration depth. The nanohardness (H) and Young’s modulus (E) of the coatings were measured from the

obtained curves by Oliver–Pharr (O-P) method. According to this method, nanohardness can be calculated through dividing the maximum load (Pmax) by the actual contact area between nanoindenter and materials (A) [Eq. (2) and (3)].

Eq. (2) Eq. (3)

The C1 to C8 are constants determined using standard calibration method, and hc is the penetration depth obtained from Eq. (4):

(

Eq. (4)

)

Where, ε is ≈0.75 for the Berkovich tip and hmax is the maximum penetration depth. Moreover, the contact stiffness (S) is described as the slope of the first one-third unloading part of the load-penetration depth curve (Eq. (5)).

(

)

√

Eq. (5)

In the above equation, β= 1.034 and CA= 2/√π are considered for the Berkovich tip, and E* is the effective Young’s modulus of the indenter and the sample set. Considering Oliver–Pharr model, E* can be expressed as Eq. (6): (

)

(

)

Eq. (6)

In this expression, the subscripts of i and s denote indenter and sample and the values of Ei (Young’s modulus) and νi (Poisson’s ratio) for Berkovich diamond indenter are 1140 GPa and 0.07, respectively [13,37,38].

Besides, the area involved by the loading–unloading curve represents the amount of plastic deformation work (Wp) done during the indentation test [37]. At last, the adhesion strength between the coatings and the anodized titanium was measured according to ASTM F1044-05 standard. For this purpose, a coated sample was bonded to the bare sample with epoxy resin. The shear load was applied using a universal testing machine (Santam STM50, Engineering Design Company, Iran). This machine was programmed at a crosshead speed of 1 mm/min until the failure of coating. The adhesion strength (σ) was calculated using Eq. (7) [17]:

Eq. (7)

Where P is the ultimate pull-out load and A is the surface area of the sample.

3. Results and discussion The SEM images in Figs. 2a and b display the top and cross-sectional views of TiO2 nanotubular structure on the anodized titanium. The EDS spectrum of this oxide layer in Fig. 2c shows the presence of Na, Ti and O elements. The existence of Na is owing to the immersion of the anodized titanium in 0.5 M NaOH. In fact, alkaline treatment of the anodized titanium was carried out because of improving the adhesion strength [19,33] and crystallinity [19] of the coating electrodeposited on the titanium substrate. Additionally, Fig. 2d shows the current density–time (I–t) curve recorded with a digital voltmeter during the anodizing of titanium. Within the first few seconds, the current density shows a sharp decrease to a minimum value, which is attributed to the formation of a compact oxide layer with poor electrical conductivity (stage I). Afterward, tiny pores on the TiO2 layer are formed due to the attack of fluoride ions as well as the electric field-assisted dissolution of the oxide layer. These pores, by providing more pathways for ions, lead to an increase in the current (stage II). In the next stage, a slight decrease in the current density could be observed as a consequence of achieving a higher insulation by the growth of TiO2 around the pores (stage III). At last, variations in the current density stay almost constant (stage IV). During this stage, there are both chemical oxidation and dissolution

processes and the individual pores compete for the available current to form a self-assembled nanotubular structure. The surface morphology and the corresponding EDS spectra of the CaP and CaP-GO coatings are also displayed in Fig. 3. According to results of the EDS elemental analysis, the Ca/P atomic ratio in the CaP and CaP-GO coatings is 1.57 and 1.63, respectively, and the reasons for the appearance of C peak are the presence of GO sheets and/or carbonate contaminations in the coating. Moreover, as it is evident from Fig. 3a, microstructure of the CaP coating is non-uniform so that the larger size crystals can be found in some regions marked with yellow circles. While, keeping the same condition for electrodeposition, the CaP-GO coating in Fig. 3b represents a more uniform microstructure. This difference could be related to the presence of GO sheets in the coating. As a matter of fact, the water reduction during electrodeposition produces hydroxyl ions, causing an increase in the pH near the cathodically polarized substrate [39]. In this condition, the CaP ceramics become insoluble and precipitate on the surface of the anodized titanium substrate. Moreover, Ti-OH groups can be formed on the anodized titanium due to the bonding between hydroxyl and titanium ions of TiO2 layer [33]. Therefore, in the absence of GO, the hydroxyls and Ti-OH groups are responsible for deposition of the CaP coating. In addition to these, the occurrence of a large percentage of oxygen-containing functional groups such as hydroxyl and carboxyl on the basal plane and edges of GO sheets also plays a pivotal role in the nucleation of CaP crystals during electrodeposition of the CaP-GO coating. In this context, the rising of pH at the substrate results in deprotonation of carboxyl groups (COOH) and, thereby, GO sheets become more negatively charged [27]. These deprotonated carboxyl groups (COO-) as well as hydroxyls (OH-) facilitate the electrostatic interaction of Ca2+ ions with GO sheets [27,28]. Moreover, Ca2+ ions could be absorbed onto GO sheets through ion exchange with H+ of carboxyl groups [32]. In the next stage, PO43- ions easily combine with Ca2+ ions through electrovalent bonds to from apatite [27,32]. Therefore, GO sheets serve as effective nucleation sites for the CaP crystals, causing the formation of a refined and relatively uniform microstructure for the CaP-GO composite coating in comparison with the CaP one. Fig. 4 reveals FTIR spectra of GO and powdery samples scrapped from the CaP and CaP-GO coatings. For the synthesized GO in Fig. 4a, FTIR spectrum shows the bands related to oxygen

functional groups of alkoxy, epoxy and hydroxyl at the wave numbers of 1053 cm-1, 1252 cm-1 and 1408 cm-1, respectively [40,41]. The band found at 1721 cm-1 is assigned to carboxyl groups on the edge of GO basal planes or conjugated carbonyl groups [28]. The sp2 hybridized carbon atoms in the graphitic domains of GO or absorbed hydroxyl groups reveal a band at 1625 cm -1 [42,43]. A broad band centered at 3431 cm-1 is also associated with hydroxyl groups [40] and the weak bands at 2853 and 2926 cm-1 are corresponded to methylene groups [28,32,44]. Moreover, the FTIR spectra of coatings in Figs. 4b and c show infrared bands belong to the phosphate groups (PO43-) of HAp at 474, 563, 603, 973, 1031, and 1097 cm-1 [35,45]. The weak or shoulder bands at 526, 667, 923, and 1103 cm-1 are assigned to the acidic phosphate groups (HPO42-), indicating the presence of OCP and DCPD phases in the coating [28,46,47]. The broad and high-intensity band centered at 3431 cm-1 [48] and the band at 1651 cm-1 [23,48] are related to the physically adsorbed water molecules. The absorbed hydroxyl band of HAp is also appeared at 636 cm-1 [23,48]. The rest bands are corresponded to carbonates that might be originated from adsorbing the atmospheric CO2 by the electrolyte during the electrodeposition [35]. According to the obtained results, although we applied a higher current density than that used by Zeng et al. [32] to electrodeposit the HAp phase, but some metastable OCP and DCPD phases besides the carbonated HAp (CHAp) are formed in the CaP coating. In this regard, adding of GO seems to be helpful for achieving the only CHAp phase in the CaP-GO coating. This can be due to the effective role of oxygen-containing functional groups of GO sheets on the nucleation of HAp crystals [28,49]. Moreover, the results of our previous study [35] showed a decrease in the rate of electrodeposition when GO sheets were present in the electrolyte. This means that during the electrodeposition of the CaP-GO composite coating there is more time to produce the adequate hydroxyl ions required for electrocrystallization of the HAp phase. Further, in the spectrum of the CaP-GO coating, there are two weak bands related to methylene groups at 2853 and 2926 cm-1, most likely owing to the presence of GO sheets in this composite coating. The XPS analysis was used to study the chemical composition of the upmost layer of coatings. The Ca, P, O and C peaks in survey spectra of both coatings further confirm the formation of CHAp phase (see Figs. 5a and b). The peaks of C and O could also be related to the GO sheets. Moreover, Na peak in both spectra is originated from sodium nitrate used to improve the ionic

strength of the electrolyte. As can be seen, the presence of GO sheets in the CaP-GO composite coating increases the intensity of C peak. The high-resolution XPS spectra of the Ca2p, P2p, O1s and C1s recorded for the CaP-GO coating are also depicted in Figs. 5c-f. The high-resolution XPS spectrum of Ca2p shows two prominent peaks of Ca2p3/2 (347.3 eV) and Ca2p1/2 (350.8 eV) [50,51]. The P2p at binding energy of 133.1 eV is assigned to the phosphate groups (PO43−) in HAP [50,51]. The high-resolution O1s represents two intense peaks centering at 529.84 and 531.77 eV. The first peak is associated with oxygen in the hydroxyl [50,51] and the second peak is related to the oxygen functional groups in GO sheets [51] or the oxygen atoms of phosphate groups [51]. Furthermore, as it is evident from Fig. 5f, the C1s XPS spectrum consists of three carbon peaks at 284.4, 286.1 and 287.8 eV belonging to C-C (sp2 hybridized carbon atoms), C-O (alkoxy and epoxy) and C=O (carbonyl) groups of GO, respectively [32,51]. Overall, the XPS results demonstrate the successful incorporation of GO reinforcement into the CaP-GO composite coating applied by pulse electrodeposition. Fig. 6a shows the bright-field TEM image of the synthesized GO. The transparency of GO sheet indicates its successful exfoliation from graphite oxide using ultrasound treatment. Moreover, the oxygenated groups on the edges and surface of GO cause some ripples and wrinkles accompanied with this sheet. Bright-field TEM images and the corresponding SAED patterns of the CaP and CaP-GO coatings collected along the [100] zone axis are also represented in Figs. 6 (b, c) and (d, f), respectively. Bright-field images in Fig. 6b and d show that the CaP coating consists of some larger size platelike crystals marked with yellow arrows, whereas more uniform CaP crystals in size are formed in the CaP-GO coating. These results are in good agreement with SEM observations. Moreover, the fully covering of wrinkled GO sheets with the CaP crystals in the CaP-GO coating (see Fig. 6d) can be due to the fact that the chemically induced defects on GO act as effective sites for attachment of these crystals. The oxygen functional groups on the surface of GO also adsorb the CaP crystals through electrostatic interactions [16]. According to the simulated SAED patterns of hexagonal HAp and triclinic OCP reported by Mokabber and et al. [52], the CaP coating is composed of both HAp and OCP phases, although, the much brighter spots of (002)HAp and (022)HAp in Fig. 6c indicate that HAp with the atomic structure of hexagonal is dominant phase in

this coating. Moreover, the SAED pattern of the CaP-GO coating in Fig. 6e matches well with that of the HAp phase, showing the formation of only HAp crystals orientated along the c-axis in the composite coating. Consequently, TEM patterns obtained for both coatings are in good agreement with the FTIR results. It should be mentioned that the XRD investigation of the CaP and CaP-GO coated samples in the our previous paper [35] yielded similar results. This could be justified by the fact that the nucleation of OCP is energetically more favorable than that of HAp mainly owing to the higher surface energy of HAp [53]. Therefore, in the CaP-GO coating, the nucleation sites on GO sheets provide an appropriate surface for the HAp formation. The spots of (002)HAp and (030)OCP in Fig. 6c and (002)HAp in Fig. 6e also show the plan spacing of about 0.352, 0.304 and 0.348 nm, respectively, which are close to those calculated for these plans of carbonated HAp (CHAp) and OCP using the lattice parameters. Generally, metal ions released into the body owing to the corrosion of metallic implants can result in an allergic reactions when they come in contact with the surrounding tissues [1]. Moreover, Corrosion can adversely affect the mechanical properties of the metallic implants [23]. Knowing the necessity of improving the corrosion resistance of metallic implants for their successful application, the current work has been studied the effects of anodizing as well as CaP and CaP-GO coatings on the biocorrosion of titanium substrate in Ringer’s solution. One of the ways to do it is monitoring the open circuit potential (OCP) of samples in corrosive environment as a function of time. Fig. 7 shows OCP plots of the bare, anodized and anodized/coated titanium samples in Ringer’s solution. The steady state potentials extracted from these plots as OCP values are given in Table 1. According to the obtained results, the OCP of bare sample shifts a slight to the negative direction and then reaches a steady state potential of -466 mV, whereas OCP values of the anodized, anodized/CaP coated and anodized/CaP-GO coated samples after a shift in the noble direction become steady at -358, -283 and -237 mV, respectively. The noblest value of OCP for the anodized/ CaP-GO coated sample reveals the better corrosion protection of the composite coating as compared to the CaP one. The potentiodynamic polarization measurements for bare, anodized and anodized/coated samples are also performed in Ringer’s solution. Fig. 8 shows the obtained polarization curves. A passive region is detectable for all samples in the anodic branch of polarization curves, in which

the current density remains almost constant. It is observable from Fig. 8 that the anodized and anodized/coated samples begin to display the passive behavior very soon as compared to the bare titanium. The electrochemical parameters including the corrosion potential (Ecorr), corrosion current density (Icorr) and protection efficiency (η%) calculated from the polarization curves by Tafel extrapolation are tabulated in Table 1. Anodizing causes a shift in Ecorr of the bare titanium towards the noble direction from -559 to -423 mV and a decrease in Icorr from 1.06 to 0.32 µA/cm2. In addition, the passive current density sharply drops after anodizing. Indeed, the barrier role of the TiO2 layer against the corrosive physiological fluids leads to a decrease in the biocorrosion rate and achieving the protection efficiency of about 69% for the anodized sample. Additionally, the anodized/CaP coated sample displays a better corrosion resistance (E corr of -362 mV, Icorr of 0.21 µA/cm2) and; in this case, the protection efficiency increases to 80%. This can be due to the fact that in the coated sample both TiO2 (bottom layer) and the poor electrical conductive CaP coating (surface layer) provide corrosion protection of the titanium substrate. Finally, although incorporation of GO into the CaP coating does not significantly shift the Ecorr value (-369 mV), but there is a decrease in the Icorr value to 0.13 µA/cm2 and thereby the anodized/CaP-GO coated sample displays the highest protection efficiency of 88%. Comparing to the CaP/anodized coated sample, the passive current density is also lower for the sample coated with the CaP-GO. The difference in the protection properties of the CaP and CaP-GO coatings can be related to their microstructure. In fact, the CaP crystals with plate-like structure grow outwardly during electrodeposition to form the porous coating on the titanium surface (see Fig. 3), in which the pores are the preferred paths for corrosive mediums. Existence the strong and coherence interconnection between the GO reinforcement and CaP crystals in the CaP-GO coating effectively prevent the body fluids from penetrating through the coating reacting with the titanium substrate to corrode it. As a result, reinforcing of CaP coating with GO sheets is effective to improve its corrosion protection property. Besides, Liang et al. reported that the surface topography of the coating affects its corrosion resistance and the coating with smoother surface exhibits more protection [54]. In contrast, based on the findings of Srinivasan et al., the higher corrosion resistance is accessible for the rougher surface [8]. Hence, in the current work, we investigated how the roughness of the CaP-

based coating could influence the biocorrosion of the titanium substrate. For this purpose, 2D AFM images (color map) and line graphs of the coated samples on Z scale are displayed in Fig. 9. The surface roughness parameters extracted from these images are also tabulated in Table 2. As it is evident from the results, the CaP coating shows an irregular surface topography, so that the higher values of maximum height of peaks (Sp), maximum depth of valleys (Sv), and maximum height within the evaluation length (Sy) are obtained for this coating in comparison with the CaPGO one. Indeed, such surface topography is caused by the non-uniformity in the microstructure of the CaP coating (see Fig. 3a). Knowing that the valleys and pores in the coating are places for the accumulation of corrosive fluids, the larger and deeper valleys due to having the high capacity for the fluids result in the penetration of a larger volume of the corrosive mediums to the substrate. This means that the smoother surface topography of the CaP-GO coating is responsible for its higher biocorrosion inhibition when compared to the CaP coating. The corrosion behavior of the coating is also affected by its composition so that the coating with more chemical stability could offer the better corrosion protection [55]. In this regard, Figs. 10 (a, b) and (c, d) show the SEM images of the anodized/CaP coated and anodized/CaP-GO coated samples before and after the polarization test, respectively. As can be seen from Fig. 10b, tiny pores are detectable on the surface of CaP coating after polarization because of its dissolution. The possibility of formation of the less stable phases like OCP and DCPD in the CaP coating is a good reason for appearance these tiny dissolved areas. Besides, the surface roughness is challenging issue in association with increasing the dissolution [56]. It is worth mentioning that these dissolved areas in the coating facilitate the reaching of corrosive fluids to the titanium substrate, resulting in the substrate corrosion. The mechanical stability is important aspect in increasing the lifespan of biomaterials [5]. Hence, development of new composites with both biocompatibility and enhanced mechanical properties has attracted much interest in recent years. In the present work, we investigated the effect of GO reinforcement on the mechanical behavior of the CaP-based coatings using nanoindentation test. Fig. 11 shows the load-penetration depth curves for these coatings and the values of nanohardness (H), Young's modulus (E), (H/E) and (H3/E2) ratios calculated according to O–P model with integrated software are summarized in Table 3. The H/E ratio can

be used to evaluate the fracture toughness [15] and also is a suitable parameter for predicting wear resistance of the materials [57,58]. The H3/E2 ratio proposed by Johnson is also associated with the coating resistance to plastic deformation [57,58]. The results in Table 3 show that applying the CaP-based coatings is desirable to decrease the great Young's modulus mismatch at the interface of titanium implant (98-105 GPa) and bone tissues (20–30 GPa). As mentioned in the introduction, the less mechanical mismatch at the interface is helpful to minimize the stress shielding effect. Moreover, it can be clearly observed from Fig. 11 that the CaP and CaP-GO coatings reveal the maximum penetration depth of about 192 and 151 nm, respectively, at the same indentation load. The difference in the penetration depth arises from the fact that the softer materials exhibit less resistant to indentation. Incorporation of GO sheets also increases the nanohardness and elastic modulus of the CaPbased coatings about 52%, 41%, respectively. Indeed, the high surface area of GO sheets makes it possible to create a large ratio of CaP/GO interfaces-to-volume in the CaP-GO coating. Under this condition, the efficient transfer of GO strengthening effect to the CaP matrix is responsible for improvement the mechanical properties of this composite coating. Increasing the elastic modulus of the CaP coating from 35.7 to 50.3 GPa by adding GO sheets also reduces the mechanical mismatch at the coating/titanium interface, resulting in a decreased possibility of coating delamination and stress-shielding. Moreover, the CaP-GO coating represents the higher H/E ratio (0.064) as compared to that measured for the CaP one (0.058), which indicates the improvement of wear resistance and fracture toughness of the CaP coating after its compositing with GO. The toughening mechanisms proposed in the ceramic composites reinforced with GO are crack branching, crack bridging, pull-out and crack deflection by GO sheets [15,28,30]. In fact, when a crack interacts with sheet-like GO, it is deflected and/or branched by this reinforcement to create a tortuous propagation path. Under this condition, the crack needs a more energy for propagation. Besides, the rough and wrinkled surface of GO sheets enables a large contact area and mechanical interlocking with the CaP matrix, causing a more required energy for pull-out of GO from the matrix. Our results also show a higher H3/E2 ratio for the CaP-GO coating (0.013) than that of the CaP one (0.007). This means that incorporation of GO sheets into the CaP coating leads to an

increase in its resistance to plastic deformation. In this case, based on the area encompassed by the loading–unloading curve, the CaP coating undergoes greater plastic deformation work of about 0.55 nJ as compared to 0.43 nJ measured for the CaP-GO coating. Decreasing the plastic work of the indentation is a consequence of obtaining the higher hardness in the composite coating. Therefore, the brittle nature of the CaP bioceramic could improve significantly by GO, and hence, the CaP-GO composite is more suitable candidate to apply as a coating on the titanium implants used especially under tensions. The adhesion strength of the coating is the crucial factor affecting the lifetime [31] and mechanical properties of the coated implant in the human body [2]. In this study, applying the high current density of 15 mA/cm2 results in detachment of the CaP coating on the bare titanium. In fact, both gaseous bubbles produced owing to the water reduction during electrodeposition and the insufficient interlocking between the coating and substrate could deteriorate the adhesion strength [59,60]. By introducing the nanotubular TiO2 layer on the titanium surface, the bonding in the coating/substrate interface improves. In this case, the adhesion strength of the CaP coating to the anodized titanium is measured to be about 13.3 ±1.4 MPa. The enhancement of the adhesion arises from the mechanical interlocking between the nantubular structure of TiO2 and the coating [61]. Besides, incorporation of GO into the CaP coating leads to an improvement of about 16% in the adhesion strength to 15.4±1.7 MPa. This value is around the least adhesive strength of 15 MPa required for biomedical applications of the HAp coatings in the ISO standard 13779-2 [17]. Indeed, the adhesion strength comprises the combination of adhesive and cohesive strength [23,62]. The adhesive strength is affected by the interlocking at the coating-substrate interface and the residual stress in the coating [62]. Existence a large difference between the elastic modulus of the coating and the substrate can weaken the adhesion strength [63]. Considering the results of nanoindentation test, the closer elastic modulus of titanium substrate to the CaP-GO composite coating causes a decrease in the residual stress, thereby improving the adhesion of this composite coating to the substrate. The cohesive strength of the coated implant is also dependent on the porosity, crystallinity, crack propagation through the coating, etc. The presence of pores can deteriorate the mechanical integrity and adhesion strength of the coating to the substrate [17]. Therefore, reducing the

porosity of the CaP coating through its compositing with GO reinforcement play an effective role in enhancing its cohesion strength. Moreover, as mentioned before, there is a strong interconnection between GO sheets and the CaP matrix because of the large specific surface area of GO sheets as well as existence a sufficient bonding between the functional groups of GO and CaP crystals. Therefore, GO sheets could effectively inhibit the further propagation and opening of the crack tip, leading to a strong bonding strength at the CaP-GO coating/substrate interface. Increasing the crystallinity of the coating also leads to its higher adhesion to the substrate [23]. According to the obtained results from our previous study [35], the crystallinity of the CaP coating applied by electrodeposition improves in the presence of GO sheets. Hence, the other reason for the higher adhesion strength of the CaP-GO coating to the substrate could be its more crystallinity in comparison with the CaP one. The stability of the CaP-based coating is probably another factor influencing the failure or success of the coated implant, so that dissolution of the coating can deteriorate its mechanical stability and adhesion strength in the early lifetime of implantation [32]. In this regard, the CaPGO coating, which is free of the less stable OCP and DCPD phases, could possess the more stability, and so better adhesion to the substrate over time. At last, Saber-Samandari et al. reported that missing the hydroxyl groups in the HAp coating causes a decrease in its mechanical properties and adhesion strength [64]. In this regard, the presence of additional hydroxyl groups at the wavenumber of 636 cm-1, as observed in the FTIR spectrum of the CaP-GO coating (Fig. 4c), could be the another reason for improving the adhesion strength of this composite coating to the substrate. 4. Conclusion The CaP coating reinforced with GO sheets was fabricated using pulse electrodeposition on the nanotubular structure of TiO2. FESEM images showed that the CaP-GO coating had a refined and more uniform microstructure in comparison with the CaP one. In fact, oxygen-containing functional groups on GO sheets by providing the significant number of nucleation sites for the CaP crystals led to a structure refining. According to the results of FTIR and SAED patterns, only HAp phase with a hexagonal atomic structure was crystallized in the CaP-GO coating, whereas,

both hexagonal HAp and triclinic OCP phases were formed in the CaP coating. Moreover, based on the polarization results, the protection efficiency for the anodized, anodized/CaP coated and anodized/CaP-GO coated samples was calculated to be 69%, 80% and 88%, respectively. The formation of TiO2 barrier layer on the anodized sample is responsible for the protection of titanium substrate against corrosion. The higher biocorrosion protection of the titanium substrate by the CaP-GO coating, as compared to the CaP one, could be also related to the inhibitory role of GO sheets in penetration of corrosive body fluids through the insulation coating to titanium. Moreover, the more biostable coating containing only HAp phase formed in the presence of GO which in turn increased the dissolution resistance of the protective coating. The surface topography of the coating also affected its corrosion protection so that the rougher CaP coating provided the deeper valleys for accumulation of the corrosive fluids in the coating. Furthermore, the nanohardness and elastic modulus of the CaP coating increased from 2.1 to 3.2 GPa and 35.7 to 50.3 GPa, respectively, after its compositing with GO sheets. The H/E ratio of the CaP coating also increased in the presence GO reinforcement from 0.058 to 0.064, revealing the raised toughness fracture and wear resistance of the composite coating. The reasons for improving the mechanical properties of the CaP-GO coating were efficient transfer of GO strengthening effect to the CaP matrix, branching and deflection of the crack in its propagation path by GO sheets. The higher elastic modulus of the CaP-GO coating could decrease the mechanical mismatching at the coating/ titanium substrate interface, and therefore, the possibility of stress-shielding occurrence. At last, existence the additional hydroxyl groups in the structure of the CaP-GO composite coating and the higher crystallinity and compactness of this coating resulted in improvement of it adhesion strength to the anodized titanium (15.4 MPa) as compared to that of the CaP coating (13.3 MPa).

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Figures captions Fig. 1: Typical force-penetration depth curve showing the measurable parameters of the nanoindentation test.

Fig. 2:, FESEM images taken from the top (a) and cross-sectional (b) views of TiO2 nanotubes, the EDS spectrum of the anodized surface (c), current density versus time curve recorded at the anodizing voltage of 60 V and temperature of 35 °C for 50 min (d). Fig. 3: FESEM images and EDS spectra of the CaP (a) and CaP-GO (b) coating applied using pulse electrodeposition at a constant current density of 15 mA/cm2 and duty cycle of 0.1. Fig. 4: FTIR spectra of the synthesized GO (a) and the CaP (b) and GO-CaP (c) coatings. Fig. 5: XPS survey spectra of the CaP and CaP-GO coatings (a); high-resolution XPS spectra for the selected Ca2p (b), P2p (c), O1s (d), and C1s (e) of the CaP-GO coating. Fig. 6: Bright-field image of the synthesized GO sheet (a). Bright-field image (b) and SAED pattern (c) of the CaP coating. Bright-field image (d) and SAED pattern (e) of the CaP-GO coating. Fig. 7: Variation of open circuit potential (OCP) with time for bare (a), anodized (b), anodized/CaP coated, and anodized/CaP-GO coated titanium in Ringer’s solution at 37 °C. Fig. 8: Potentiodynamic polarization curves for bare (a), anodized (b), anodized/CaP coated, and anodized/CaP-GO coated titanium in Ringer’s solution at 37 °C. Fig. 9: 2D AFM image and line graph of the CaP (a, b) and CaP-GO (c, d) coatings, respectively. Fig. 10: SEM images of the CaP (a, b) and CaP-GO (c, d) coated samples before and after the polarization test, respectively. Fig. 11: Load-penetration depth curves obtained from nanoindentation test for the CaP and CaPGO coatings.

Table. 1: The electrochemical parameters and protection efficiency calculated from polarization curves for the bare, anodized and coated samples. Samples Bare Ti Anodized Ti Anodized/CaP coated Ti Anodized/CaPGO coated Ti

OCP (mV vs. SCE) -466 -358

Ecorr (mV vs. SCE) -559 -423

Icorr (µA/cm2) 1.06 0.32

Protection efficiency (η%) 69

-283

-362

0.21

80

-237

-369

0.13

88

Table. 2: The surface roughness parameters extracted from AFM images for the CaP and CaP-GO coatings. Coatings CaP CaP-GO

Sa(nm) 450.68 234.42

Sq(nm) 561.62 297.24

Sp(nm) 1576.3 943.5

Sv(nm) -1467.2 -1121.4

Sy (nm) 3043.5 2064.9

Table. 3: Nanoindentation results for the CaP and CaP-GO coatings electrodeposited on the anodized titanium Coatings CaP CaP-GO

Nano-hardness (H, GPa) 2.1±0.6 3.2±0.4

Young's modulus (E, GPa) 35.7±4.2 50.3±3.5

(H/E)

H3/E2

0.058 0.064

0.007 0.013

Highlights 

The CaP-GO bioceramic coating was electrochemically deposited on the anodized titanium.



The presence of GO reinforcement caused an increase in the nanohardness and elastic modulus of the CaP coating over 52% and 41%, respectively, and the elastic modulus mismatch of the CaP-GO coating with the titanium substrate reduced. Therefore, in comparison with the CaP; the CaP-GO composite is more suitable candidate to apply as a coating on the titanium implant used especially under tension.



The bonding strength of this biocomposite coating to the anodized titanium also improved about 16% compared to the graphene-free CaP coating.



The barrier role of GO sheets and the more uniformity and crystal refinement in the microstructure of the CaP-GO coating caused its better protection against the corrosion of the anodized titanium, whereas, some larger crystals in the CaP coating provided the deeper valleys in the coating for the rapid contact of the corrosive medium with the substrate.