Accepted Manuscript

Title: Electrophoretic Deposition of NanostructuredHydroxyapatite Coating on AZ91 Magnesium Alloy Implantswith Different Surface Treatments

Author: Ramin Rojaee Mohammadhossein Fathi KeyvanRaeissi

PII: S0169-4332(13)01599-7DOI: http://dx.doi.org/doi:10.1016/j.apsusc.2013.08.108Reference: APSUSC 26251

To appear in: APSUSC

Received date: 14-5-2013Revised date: 19-8-2013Accepted date: 23-8-2013

Please cite this article as: R. Rojaee, M. Fathi, K. Raeissi, ElectrophoreticDeposition of Nanostructured Hydroxyapatite Coating on AZ91 Magnesium AlloyImplants with Different Surface Treatments, Applied Surface Science (2013),http://dx.doi.org/10.1016/j.apsusc.2013.08.108

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 proofbefore it is published in its final form. Please note that during the production processerrors may be discovered which could affect the content, and all legal disclaimers thatapply to the journal pertain.

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Electrophoretic Deposition of Nanostructured Hydroxyapatite Coating on AZ91 Magnesium 1

Alloy Implants with Different Surface Treatments2

Ramin Rojaee1*, Mohammadhossein Fathi1, 2, Keyvan Raeissi33

1Biomaterials Research Group, Department of Materials Engineering, Isfahan University 4of Technology, Isfahan, 84156-83111, Iran5

2Dental Materials Research Center, Isfahan University of Medical Sciences, Isfahan, Iran.6

3Department of Materials Engineering, Isfahan University of Technology, Isfahan, 84156-83111, 7

Iran8

9

*Corresponding author, Tel.: +989119901024, Fax: +983113912752, E-mail: raminrojaee@aim.com

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Abstract2

Bio-absorbable magnesium (Mg) based alloys have been introduced as innovative orthopedic3

implants during recent years. It has been specified that rapid degradation of Mg based alloys in 4

physiological environment should be restrained in order to be utilized in orthopedic trauma 5

fixation and vascular intervention. In this developing field of healthcare materials, micro-arc 6

oxidation (MAO), and MgF2 conversion coating were exploited as surface pre-treatment of 7

AZ91 magnesium alloy to generate a nanostructured hydroxyapatite (n-HAp) coating via 8

electrophoretic deposition (EPD) method. X-ray diffraction (XRD), Scanning Electron 9

Microscopy (SEM), Fourier Transform Infrared spectroscopy (FT-IR), and Transmission 10

Electron Microscopy (TEM) techniques were used to characterize the obtained powder and 11

coatings. The potentiodynamic polarization tests were carried out to evaluate the corrosion 12

behavior of the coated and uncoated specimens, and in vitro bioactivity evaluation were 13

performed in simulated body fluid. Results revealed that the MAO/n-HAp coated AZ91 Mg alloy 14

samples with a rough topography and lower corrosion current density leads to a lower Mg 15

degradation rate accompanied by high bioactivity.16

17

Keywords: Magnesium, Bio-degradable implant, Nanostructured hydroxyapatite, Electrophoretic18

deposition, Micro arc oxidation, Fluoride conversion coating.19

20

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1. Introduction2

Metallic biomaterials compose major quota of orthopedic implants, nowadays. Prolonged 3

therapy and possible release of undesirable ions of conventional implants such as stainless steels 4

and cobalt–chromium alloys in body environment has demanded new bio-degradable/ bio-5

absorbable materials [1]. In this developing field of healthcare materials, magnesium alloys are 6

posed as potential bio-degradable/ bio-absorbable metallic implants. Although, an injured tissue 7

can consume Mg alloys implants during the healing period, the implants suffer from high 8

corrosion rate (i.e. high mass loss and abundant rate of hydrogen evolution) in body environment9

[2]. Therefore, further treatments should be fulfilled to increase the corrosion resistance of Mg 10

alloys, and consequently improve the hurt bone response to Mg alloys during post implantation11

[3]. Producing protective bioceramics coatings on Mg alloys would enhance their 12

biocompatibility and decelerate their degradation rate in physiological environments.13

Hydroxyapatite (HAp-Ca10(PO4)6OH2) is the main inorganic component of bones, which its 14

application as a coating provides advantages intimately attributed to its nanoscale structure15

[4,5]. Moreover, former investigations revealed that the nano-sized HAp can promote 16

mechanical properties, bone cell adhesion and proliferation within the injured area in comparison 17

with the micro-sized HAp [3]. Thus, it seems that applying the nanostructured HAp (n-HAp) on 18

Mg alloys give rise to a good bioactivity and superior corrosion resistance. Many methods are 19

utilized to coat bioceramics on metallic substrates. Among these, electrophoretic deposition 20

(EPD) seems to be more suitable to produce a homogeneous and dense ceramic, polymer and 21

composite coatings for biomedical applications [6]. In EPD technique, direct current (DC) is 22

applied to suspended powder particles in a liquid medium to charge the particles and deposit 23

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them onto a conductive substrate of opposite charge [7]. Farrokhi-Rad and Ghorbani applied1

EPD technique to coat TiO2 nanoparticles on the 304 stainless steel and studied the 2

electrophoretic mobility of titania nanoparticles in different suspension mediums [8]. Mehdipour3

et al. deposited a kind of bioactive glass-contained Mg on 316L stainless steel via EPD to 4

improve the bioactivity of the substrate [9]. Besides, the biocompatibility of Mg-based alloy 5

surfaces could be improved via controlling their degradation rate utilizing fluoride conversion 6

coating and surface modification through micro arc oxidation (MAO) procedure [10,11].7

Fluoride conversion coatings are MgF2 layers produced by treating Mg substrate in hydrofluoric 8

acid (HF) [12]. Moreover, MAO technique is able to create a thick and stable porous ceramic 9

coating on Mg alloys [3]. The aim of this study was introduction of fluoride conversion coating 10

and MAO-produced layer as intermediate layers for n-HAp coating; also their effect on bio-11

corrosion behavior of these coating systems was investigated.12

2. Materials and methods13

2.1. Substrate surface preparation14

AZ91 magnesium alloy with chemical composition (wt. %) of 8.63 % Al, 0.59 % Zn, 0.17 %15

Mn, <0.05 % Fe, <0.05 % Cu and balance Mg was used as substrate. The as-cast AZ91 alloy was 16

heat-treated according to ASTM B661. AZ91 alloy samples were cut into dimensions of 20 mm17

× 10 mm × 2 mm. The surface roughness of both sides of the specimens was adjusted at Ra = 0.20 18

± 0.03 µm, using successive finer silicon carbide abrasive paper up to grit # 600. Then, the 19

specimens were ultrasonically cleaned (WUC-D10H, Wisd Laboratory Equipments, Germany) in 20

acetone for 30 min and dried quickly in warm stream of air.21

2.2. Fluoride conversion coating treatment22

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For fluoride conversion treatment, the samples were immersed vertically in a plastic bottle 1

containing hydrofluoric acid 48 wt.% (HF, GR, Merck) at ambient temperature for 24 h [12]. 2

Then, the treated samples were rinsed by deionized water and dried in warm air.3

2.3. MAO coating procedure4

An aqueous saline solution electrolyte containing 100 g/l sodium hydroxide (NaOH) - 100 g/l5

sodium silicate (Na2SiO3) - 20 g/l potassium hydroxide (KOH) was provided as an anodizing6

solution. The composition of an anodizing bath was opted to induce Na, Si and K accompanied 7

with Mg and O elements into the surface of the AZ91 alloy samples. These elements are known 8

to be essential for metabolic processes and incorporating within the healing period [13].9

Plasma arc oxidation was performed under a constant applied voltage of 30V.cm-2 for 30 min. 10

Initial temperature of the anodizing solution was set at 25 ˚C. AZ91 magnesium alloy samples 11

were put as anodes (working) electrode and a stainless steel 316L sample was introduced as a 12

cathode (counter) electrode with the surface ratio of 1:1. The distance between the working and 13

counter electrodes was set 20 mm. The cell was connected to a DC power supply (IPC-14

SL20200J, 20 A, 200 V, Iran). After formation of anodic films, the samples were ultrasonically 15

rinsed with acetone and distilled water, and then dried in air.16

2.4. n-HAp powder preparation and coating17

2.4.1. n-HAp powder preparation18

n-HAp powders were synthesized using a simple sol-gel method [14]. For this purpose, calcium 19

nitrate tetrahydrate (Ca(NO3)2.4H2O, GR, Merck) and phosphorous pentoxide (P2O5, GR, 20

Merck) were dissolved separately in ethanol (C2H5OH, GR, Merck) as calcium (Ca) and 21

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phosphorous (P) precursors, respectively. Ca-precursor was added drop-wise into the P-1

precursor. The mixture was continuously stirred for about 24 h at ambient temperature. The gel 2

was dried at 80 °C for 24 h in an oven (FN 120, NÜVE, Turkey). Then, the powders were heated 3

at a rate of 5 °C/min up to 600 °C in a muffle furnace (FHP 05, WiseTherm, Germany) and were 4

stored for 1 h at this temperature [14]. The synthesized powders were cooled to ambient 5

temperature. The heat-treated n-HAp powders were milled using a high-energy planetary ball 6

mill (PM 100, Retsch, Germany) with a zirconia vial and zirconia balls to produce more 7

homogenous n-HAp powders in the particle sizes. The ball milling parameters were selected as 8

follows: ball-to-powder weight ratio: 25:1, rotational speed: 250 rpm, and time: 5 h. Finally, this9

obtained powder was rinsed 3 times with ethanol and double distilled water in a centrifuge10

(EBA20,Hettich, Germany) at 6000 rpm each for 5 min to remove any impurities [15]. Finally, 11

powder was dried in oven at 80 ˚C for 5 h.12

2.4.2. Determination of particle dispersion stability13

In order to find the best pH for colloidal suspension used in EPD technique, the stability of 14

particles in a suspension as a function of pH was determined using a zeta potential-sizer 15

(ZEN3600, Malvern, England). The absolute values of zeta potential indicate the strength of 16

interactions among between colloidal particles and determines the degree of repulsion forces 17

among particles in a suspension [16,17]. Another important factor in determining the mobility of 18

the particles in a suspension is its electrical conductivity. It has been reported that too conductive 19

suspensions sustain very low particles motion and too resistive colloidal liquids would lost their 20

particles stability [7]. In the current study, the pH value has shown to impress the conductivity of 21

n-HAp suspension. Based on the zeta potential measurements, pH of the suspension was adjusted 22

using 0.1 mol/l nitric acid 65 wt. % (HNO3, GR, Merck) in order to obtain an stable suspension 23

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condition for n-HAp powder deposition on AZ91 alloy samples. The pH values were monitored 1

by a pH meter (GLP 22+, Crison, Spain).2

2.4.3. n-HAp coating3

In order to provide an electrophoretic suspension medium, 50 g/l n-HAp powders were dispersed 4

in methanol (CH3OH, GR, Merck). It is worth noting that the purity degree of suspension is a 5

vital point. Incorporating humidity and water also causes hydrolyzing which produces H2 and O26

gases during coating procedure and provides a porous and non-homogenous deposition on the 7

substrate. The prepared suspension stirred at 400 rpm for 30 min at 25˚C in a beaker. The 8

suspension was ultrasonicated for 1 h to achieve a homogeneous dispersion of n-HAp particles. 9

Prior to EPD, the suspension was settled for 1 h to avoid deposition of flocculated particles. 10

Pristine AZ91 magnesium alloy, HF-treated and MAO-treated samples were used as the cathodes11

(negative charge). A graphite plate with 20mm×10mm×3mm dimensions was used as an anode 12

(positive charge). The distance between the cathode and the counter electrode was set to 20 mm. 13

n-HAp powders deposition was performed in two steps in order to achieve a more homogenous14

n-HAp coating. The stepwise deposition avoids any cracking during drying shrinkage as reported 15

before [17]. First deposition step was performed at a constant voltage of 50V.cm-2 for 3min. The 16

coated samples were kept at room temperature for 30 min for evaporating solvent of the coated 17

surface. Then, second coating step was deposited at constant voltage of 30 V.cm-2 for 5 min on 18

the previous layer to provide a dense n-HAp coating. Finally, the samples were put at ambient 19

temperature for drying.20

2.5. Characterization21

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A monochromatic copper Kα radiation was selected for X-ray diffraction (XRD) studies, using a1

diffractometer (XRD, Philips Xpert, The Netherlands) operated at 40 kV and 30 mA at 2θ range 2

of 20˚– 60˚ at 0.02˚/s step size.3

Fourier transform-infrared spectroscopy (FTIR, Jasco-680, Japan) was utilized to identify and 4

confirm the functional groups of the coatings in the range of 2000–400 cm−1.5

Transmission electron microscopy (TEM, EMlOC Zeiss, Germany) was used to characterize the 6

synthesized hydroxyapatite nanoparticles. For this purpose, the samples were prepared by 7

dispersing the powders in ethanol. Then, the suspension was put in an ultrasonic bath for 30 min. 8

The prepared suspension was dropped on an amorphous carbon-coated copper grid for further 9

analysis.10

Scanning electron microscopy (SEM, Philips XL 30: Eindhoven, The Netherlands) equipped 11

with an energy dispersive X-ray spectroscopy (EDS, Seron Technology AIS 2300C, Korea) was12

used to study the microstructure and morphology of the coatings, distinguish the formation of 13

apatite sedimentation on the samples surface after immersion in biological testing environment, 14

and to determine the elemental distribution of the deposited coatings, respectively.15

2.6. Surface roughness measurement16

The surface roughness (Ra) of the treated samples were measured using a roughness tester 17

(Taylor-Hobson Surtronic Duo, Leicester, England) with a cut-off 0.8 mm. Five different 18

measurements were collected and the mean values and standard deviation were reported, 19

accordingly.20

2.7. In vitro bioactivity evaluation21

It is known that the bioactive materials can form bone-like apatite on their surfaces in an 22

acellular and protein-free simulated body fluid (SBF) having ion concentrations nearly equal to 23

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those of human blood as produced in the living body [18]. The majority of bioactive implants1

bond to living bone through this apatite layer. It should be pointed out that the apatite 2

precipitates can bond to living bone through the apatite layer in the living body, as long as the 3

material does not contain any components that induces cytotoxic or antibody reactions. Despite 4

this constraints, investigation of apatite formation on the surface of a material in SBF is useful 5

for predicting the in vivo bone bioactivity behavior [19,20]. Several SBF compositions have 6

been proposed by Kokubo et al. [20]. However, these SBF procedures involve the risk of 7

premature precipitation of calcium phosphates [21]. Thus, in the following study, SBF was 8

prepared according to Bohner and Lemaitre [21] recipe in order to investigate the bioactivity 9

behavior of the coatings. This SBF preparation script is the revised procedure of Kokubo and 10

Takadama’s [20]. Table 1 indicates the SBF ions concentrations.11

Various groups of specimens with the effective surface area of 1 cm2 were incubated in SBF 12

solution and maintained in a water bath (NB9-NÜVE, Turkey) at 37±1 ˚C within 28 days in 13

order to investigate the precipitation of the apatite deposits on the samples.14

2.8. Evaluation of electrochemical corrosion behavior15

A tri-electrode cell set-up was used with a standard platinum electrode as a counter and an 16

Ag/AgCl (saturated with potassium chloride [KCl]) as a reference electrode. Each sample was17

immersed in 500 ml SBF solution at 37±1 ˚C and the atmosphere was open to air. The 18

potentiodynamic polarization tests were carried out after the samples were immersed in the test 19

solution for 1 h, so the electrode/solution interface reached a quasi-steady state in the electrolyte20

[22,23]. All the electrochemical measurements were carried out using an AMETEK21

potentiostat/galvanostat (PARSTAT 2273). The potentiodynamic polarization measurements 22

were monitored from -250 mV vs. OCP to 800 mV vs. Ag/AgCl with a scan rate of 1 mV/s. The 23

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corrosion potentials (Ecorr.), corrosion current densities (iCorr.), anodic Tafel slopes (βa), and 1

cathodic Tafel slopes (βc) were extracted from the potentiodynamic polarization plots using Tafel 2

extrapolation and linear polarization methods [24].3

The polarization resistance (Rp) was calculated using the Stern–Geary equation according to 4

ASTM G59 (Eq. 1).5

2 303a c Corr . a cp iR / . ( ) (1)6

All the experiments were carried out at least three times and the mean values and standard 7

deviations were reported.8

3. Results and discussion9

3.1. Zeta potential/conductivity measurements10

The movement of ceramic particles in a suspension fluid is defined as the electrophoretic 11

mobility (µ) of the charged particles in a suspension [7]. This parameter is supported by the pH, 12

ionic strength and the viscosity of a suspension [17]. The electrophoretic mobility (µ) is defined 13

as Eq. 2 [17]. 14

µ = ζ.ε/4π.η (2)15

Where ζ, ε, η are the zeta potential, dielectric constant, and viscosity of the medium, respectively 16

[17]. As the methanol is chosen as a suspension medium for n-HAp deposition, ε and η could be 17

assumed almost constant. Thus, zeta potential is the main parameter in the particles’ mobility [7]. 18

It is worth noting that in the presence of DC electric field, the charged particles move in a 19

suspension toward the oppositely charged electrode [25]. The electrophoretic velocity (ν) is 20

given by the Eq. 3 [25].21

ν = E.µ (3)22

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Where E is the potential gradient (V/cm), and µ is the electrophoretic mobility [25]. Eq. 3 shows 1

that a more electrophoretic velocity is attained in higher values of applied voltages.2

Fig. 1 shows the zeta potential and conductivity values of the suspension as a function of pH. It 3

can be seen that a higher zeta potential occurs at 4.5 <pH<6.5. Researchers reported that the 4

ionic strength of the suspension is closely dependent on the suspension conductivity [26]. As the 5

conductivity of a suspension increases, a larger amount of free ions will form in the suspension. 6

This phenomenon causes aggregation of colloidal particles and consequently, reduces the 7

mobility of particles [27]. Nevertheless, it has been found that a small amount of free ions in a 8

suspension acts as a electromotive force to move the particles towards an oppositely charged 9

electrode [27]. Thus, the optimum pH value of the present n-HAp suspension was fitted at 6.0, 10

where a higher zeta potential and also a suitable value of conductivity were monitored.11

3.2. Coating characterization12

3.2.1. X- ray diffraction and TEM studies13

Fig. 2 illustrates the X-ray diffraction pattern of pristine AZ91, MgF2 and MAO-treated AZ91 14

samples. As shown, pristine AZ91 microstructure is composed of α-phase (pure Mg), and β-15

phase (Al12Mg17). XRD pattern of MAO-coated sample indicates that it is mainly composed of 16

MgO, MgSiO3 and Mg2SiO4. Formation of these phases confirms that Mg, Si and O elements 17

were engaged in chemical reactions during MAO process [28]. Fig. 2 also confirms that the 18

fluoride treated sample is covered by MgF2 layer. The peaks of the substrate alloy are also 19

detected in the XRD pattern of coated samples.20

A typical X-ray diffraction pattern of n-HAp powder has been shown in Fig. 3. The crystallite 21

size of n-HAp powder was calculated using Scherrer’s equation [29,30]. The average crystallite 22

size was about 30 nm estimated from (002) diffraction peak. Fig. 4 shows TEM photomicrograph 23

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of the n-HAp powder following the heat treatment at 600 ˚C. Morphology and size of the n-HAp 1

particles could be observed which confirms the nanocrystalline nature of the powder. The 2

average particle size of n-HAp powder was estimated to be about 45 nm in range.3

3.2.2. FT-IR spectroscopy analysis4

It is known that the peak positions in an infrared spectrum correlate with molecular composition 5

and structure [31]. Fig. 5 shows chemical composition of the n-HAp, MAO and MgF2-treated 6

samples obtained by FTIR spectra. A widened band relating to the vibration of adsorbed water is 7

observed between 1630-1640 cm-1 in n-HAp and MAO-treated samples (bending mode) [32]. 8

The IR signals in the range of 860-880 and 1400-1500 cm-1 indicate characteristic peaks of 9

carbonate phases and arise from the anti-symmetrical C-O stretching vibration and bending 10

resonance, respectively [32–34]. The peaks at 473 and 631 cm-1 are assigned to the stretching 11

modes of structural hydroxyl groups (O-H) in n-HAp lattice structure [29]. The bands at 572, 12

601, 1042 and 1089 cm-1 correspond to P-O bending vibration [30]. A peak around 962 cm-1 is 13

assigned to the PO43- tetrahedron stretching resonance [30]. On the other hand, a weak resonance 14

peak is detected at 424 cm-1in MAO-treated sample. This peak is attributed to Mg-O longitudinal 15

acoustic mode [35]. The characteristic stretching band related to SiO6 group is appeared at 68516

cm-1 indicating the presence of MgSiO3 in MAO-treated surface [36]. Also, the typical resonance 17

peaks of forsterite is observed at 903 and 1022 cm-1 indicating SiO4 stretching vibrations [37,38].18

With regard to MgF2-treated sample, the peak around 431 cm-1 is ascribed to Mg-F stretching 19

vibration [39]. It has been reported that MgF2 acts as a transparent layer against infrared 20

radiation in the range of 600–4000 cm-1, thus, it does not show further characteristic frequencies 21

[40]. Table 2 summarizes the FTIR results.22

3.2.3. SEM evaluation, surface and cross section studies23

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Figs. 6- 8 show SEM micrographs and elemental analyses of the surface and cross section of the 1

MgF2, MAO, n-HAp layers formed on the substrate, respectively. Figs. 6a, b reveal formation of 2

a dense and uniform layer of MgF2 with a mean thickness of 1.31 ± 0.02 μm. The average surface 3

roughness of this layer was 0.23 ± 0.01 μm. Fig. 6c shows the elemental analysis and maps of 4

elemental distribution for fluoride-treated surface. O, F, Mg and Al elements were detected in 5

this structure. It is worth noting that Al and Mg are mainly originated from the substrate.6

Fig.7 shows the typical SEM micrograph of surface and cross section of MAO layer on AZ91 7

sample. As is seen, the surface is consisted of several circular pores arise due to the 8

inhomogeneous growing pattern of the coating and trapping of molten oxide and gas bubbles 9

during the coating growth process [1,41]. This event caused a porous outer layer with 2–5 μm in 10

diameter accompanied with a dense inner layer. The average surface roughness of MAO-treated 11

surface was assessed 2.22 ± 0.05 μm and the mean thickness of MAO layer was 21.11 ± 2.51 μm.12

The corresponding EDS surficial elemental analysis and scanning map of the MAO layer13

confirm the distribution of O, Na, Mg, Si, and K elements within the structure. The presence of 14

these elements in MAO layer suggests that they were incorporated into the surface structure 15

during anodizing process [3].16

Fig. 8 shows the SEM surface morphology and the cross sections of n-HAp coated samples at the 17

opted conditions as described in section 2.5. The n-HAp coatings were consisted of an outer 18

course layer with the mean thicknesses of 87.31 ± 4.52 μm on the MAO (Fig. 8b) and 82.67 ± 19

1.33 μm on the MgF2 (Fig. 8c) intermediate layers. The n-HAp coating showed an average 20

roughness of 5.05 ± 0.72 µm in the both cases. Previously, Deligianni et al. [42] stated that the 21

roughness of hydroxyapatite surface affects cellular response, and has a direct effect on 22

enhancing cell adhesion and proliferation. EDS pattern of n-HAp coating is shown in Fig. 8d. 23

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Quantitative estimation of atomic percentages in EDS analysis and pre-calculated Ca/P ratio1

(“34.1:18.88” vs. “1.67:1”) was in the same range, showing that Ca, P-precursors were 2

incorporated completely during n-HAp production reaction.3

3.3. Corrosion behavior evaluation4

Fig. 9 illustrates typical potentiodynamic polarization curves of MgF2 and MAO-treated AZ915

alloy (with and without n-HAp coating) in SBF solution, and pristine AZ91 as control specimen.6

The average values of electrochemical parameters were extracted and could be observed in Table 7

3. It can be seen that (i) all types of the coatings (MgF2, MgF2/n-HAp, MAO, MAO/n-HAp)8

ennobled the corrosion potential. This could be a good indication of the coating stability [43]; (ii) 9

all types of the coatings decreased the substrate corrosion current density, significantly. There is 10

no significant difference between the corrosion current densities of MgF2 and MgF2/n-HAp or 11

MAO and MAO/n-HAp samples. However, it should be remarked that the major role of the n-12

HAp coating was enhancing the surface bioactivity. While, it could be seen that it also improved13

the corrosion resistance in some extent; (iii) the slope of anodic branch is much steeper for all 14

types of the coated samples comparing AZ91 substrate. This could be a response of the higher 15

activation energy in releasing Mg2+ ions, causing lower corrosion rates [44]; (iv) Rp is one of the 16

most important variables in direct proportion to the corrosion resistance [45]. Rp measurements 17

demonstrate that the interlayers lead to significant increase in the corrosion resistance, i.e. the 18

MgF2 and MAO layers can provide good barriers against the substrate rapid degradation.19

Consequently, it is considerable that the interlayers also play an important role in improvement 20

of the biocompatibility of AZ91 alloy [46].21

3.4. In vitro bioactivity responses22

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One of the most efficient parameters in governing the bioactivity of materials is to investigate the 1

precipitation behavior of bone-like apatites on their surface [47]. The bioactivity process will 2

take place at the material periphery via releasing and adsorbing ions [48]. Thus, different 3

coatings show different morphology of the precipitated apatites [49]. The surface morphology of 4

the coatings following immersion within 4 weeks in SBF solution was observed with SEM, and 5

are represented in Fig. 10. Pristine AZ91 sample was also immersed in the same solution for 6

comparison. It can be seen that the semi-spherical crystals of apatite-like precipitations had been 7

deposited densely on the pristine AZ91 sample (Fig. 10a). Since the precipitations removed from 8

the surface, micro-pits were observed on the surface (red circles in the Fig. 10b). Fig. 10c shows9

the fluoride-treated surface after soaking in SBF solution. The surface presented a crackled 10

appearance due to the incorporation of the interlayer’s ions (Mg2+, F-) in mineralization process11

[49]. It is worth noting that a low density of bone-like apatite crystals formed on MgF2 surface.12

This could be due to the presence of fluorine in the coating which decreases the bio-absorbability 13

of calcium phosphates in SBF solution [50]. Previously, Yan et al. [45] and Zahrani et al. [50]14

have reported a similar behavior of apatite sedimentations on fluoridated coatings. Fig. 10d15

represents the apatite precipitations on MgF2/n-HAp surface. By comparing Fig. 10c and Fig. 16

10d, it could be deduced that n-HAp coating on MgF2 layer was able to speed up the apatite 17

sedimentation reaction compared to that happens on MgF2 layer. Thus, more precipitation 18

occurs.19

Figs. 10e and 10f demonstrate the behavior of the surface bioactivity on MAO-treated surface 20

and MAO/n-HAp coating, respectively. Both MAO and MAO/n-HAp coated surfaces were 21

covered with dense apatite precipitations. This indicates that they are highly prone to induce 22

precipitation of bone-like apatites and stimulate matrix mineralization. Presence of several ions 23

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within the surface would enhance the bone-forming ability of implant materials [13]. Moreover, 1

by the higher the surface roughness (n-HAp>MAO>MgF2), greater bioactivity occurs [51]. The2

morphology and condition of the precipitated bone-like apatites can be observed in Fig. 10.3

Combining the results of immersion test and potentiodynamic polarization measurements4

conducted that the best bioactivity and corrosion resistance for MAO/n-HAp coating in SBF5

solution. It can be concluded that this type of coating system (n-HAp layer as a top-coat and 6

MAO layer as an inter-layer) can provide a suitable framework for cell adhesion, encouraging 7

growth and stable healing due to its surface roughness which also has the benefits of controlled 8

Mg2+ ion release from the substrate.9

4. Conclusion10

The present literature introduced the MgF2 and MAO coatings as the inter-layers between AZ91 11

alloy substrate and n-HAp bioceramic top-coat. Presence of several ions within the coat layer 12

provides promising agents necessary for enhancing the bone-forming ability of implant 13

materials. Higher corrosion resistance of the MAO/n-HAp coated sample led to the tailored 14

degradation kinetics of AZ91 alloy substrate as biodegradable implants. Moreover, its rough 15

surface also plays an effective role in the healing process. The presented data support the 16

research that the MAO/n-HAp coating is highly prone to be utilized on Mg-based bio-absorbable 17

metallic implants that offer improved patient outcomes.18

Acknowledgements19

The authors are grateful for support of this research by Biomaterials Research Group of Isfahan 20

University of Technology.21

22

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1

Legend of Figures2

Fig. 1. Zeta potential/conductivity of a n-HAp/methanol suspension as a function of pH.3

Fig. 2. XRD patterns of pristine AZ91alloy, MgF2 and MAO coatings.4

Fig. 3. XRD pattern of sol–gel derived n-HAp powder after heating at 600 ˚C for 1 h.5

Fig. 4. TEM micrograph of sol–gel derived n-HAp powder after heating at 600 ˚C for 1 h.6

Fig. 5. FTIR spectra of the MgF2, MAO and n-HAp coatings.7

Fig. 6. Morphologies of; (a) the surface, (b) cross section, (c) elemental compositions at the 8

selected area quantified by EDS analysis and corresponding element distribution images of 9

fluoride-treated AZ91 alloy.10

Fig. 7. Morphologies of; (a) surface, (b) cross section, (c) elemental compositions at the selected 11

area quantified by EDS analysis and corresponding element distribution images of MAO-treated 12

AZ91 alloy.13

Fig. 8. Morphologies of; (a) the surface of the n-HAp coating, (b) cross section of the MAO/n-14

HAp coating, (c) cross section of the MgF2/n-HAp coating, (d) and elemental compositions at the 15

selected area quantified by EDS analysis of the n-HAp coating.16

Fig. 9. Electrochemical corrosion polarization curves of (a) the pristine AZ91, and the samples 17

treated with; (b) MgF2 coating, (c) MgF2/n-HAp coating, (d) MAO coating, and (e) MAO/n-HAp18

coating.19

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Fig. 10. SEM photomicrographs of the surface of (a) the pristine AZ91, (b) the pristine AZ91 1

after removing precipitations and corrosion products, (c) MgF2 coating, (d) MgF2/n-HAp2

coating, (e) MAO coating, and (f) MAO/n-HAp coating, after 28 days soaking in the SBF.3

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Table 1. Nominal ion concentrations of SBF in comparison with those in human blood plasma [21].2

3

4

5

6

7

8

9

10

11

12

Ion concentrations (mM)

Blood plasma SBF

Na+ 142.0 142.0

K+ 5.0 5.0

Mg2+ 1.5 1.5

Ca2+ 2.5 2.5

HCO3- 27.0 4.2

HPO42- 1.0 1.0

SO42- 0.5 0.5

Cl- 103.0 147.96

pH 7.2-7.4 ≈7.4

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Table 2. Infrared band positions of the studied samples.2

3

Wavenumber (cm-1)n-HApcoating

MAOcoating

MgF2

coatingAssignment References

--- 424 --- Mg-O (longitudinal acoustic) [35]--- --- 431 Mg-F (stretching) [39,40]

473 --- --- O-H (stretching) [29]572 --- --- P-O (bending) [30]601 --- --- P-O (bending) [30]631 --- --- O-H (stretching) [29]--- 685 --- Si-O (stretching) [36]

876 865 --- C-O (bending) [32,34]--- 903 --- Si-O (stretching) [37,52]

962 --- --- P-O (stretching) [30]--- 1022 --- Si-O (stretching) [38]

1042 --- --- P-O (bending) [30]1089 --- --- P-O (bending) [30]1432 1411 --- C-O (anti-symmetrical stretching) [33]1457 1453 --- C-O (anti-symmetrical stretching) [33]1633 1631 --- H-O-H (bending) [32]

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Table 3. Electrochemical data related to potentiodynamic polarization curves.2

Samples Ecorr., mV icorr., µA/cm2 Βa, mV/dec Βc, mV/dec Rp, Ω/cm2

Pristine AZ91 -1448 (51) 22.14 (3.21) 30 (4) 120 (6) 471

MgF2 coated AZ91 -1371 (50) 7.36 (0.63) 104 (5) 486 (31) 5 × 103

MgF2/n-HAp coated AZ91 -1400 (24) 2.92 (0.71) 75 (8) 506 (34) 9.7 × 103

MAO coated AZ91 -1280 (30) 0.58 (0.01) 83 (6) 125 (4) 3.7 × 104

MAO/n-HAp coated AZ91 -1323 (28) 0.45 (0.05) 96 (6) 141 (8) 5.5 × 104

3

4

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Research highlights1

Nano-hydroxyapatite (n-HAp) particles synthesized using a sol-gel process.2

Stable suspension of n-HAp powders was prepared in methanol medium.3

Fluoride and MAO pre-treatments were applied on AZ91 samples as intermediate layers.4

Electrophoretic deposition process was used for coating n-HAp particles.5

The MAO/n-HAp/AZ91 system is highly prone to be utilized as bio-absorbable implants.6

7

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Graphical Abstract (for review)

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Figure1

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Figure2

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Figure3

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Figure4

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Figure5

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Figure6

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Figure7

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Figure8

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Figure9

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Figure10