- Email: [email protected]
nHAp composite porous film-based coating of magnesium alloy
Accepted Manuscript Biocorrosion and Osteoconductivity of PCL/nHAp Composite Porous Film-Based Coating of Magnesium Alloy Abdalla Abdal-hay, Touseef Amna, Jae Kyoo Lim PII:
S1293-2558(12)00380-9
DOI:
10.1016/j.solidstatesciences.2012.11.017
Reference:
SSSCIE 4635
To appear in:
Solid State Sciences
Received Date: 25 May 2012 Accepted Date: 7 November 2012
Please cite this article as: A. Abdal-hay, T. Amna, J.K. Lim, Biocorrosion and Osteoconductivity of PCL/ nHAp Composite Porous Film-Based Coating of Magnesium Alloy, Solid State Sciences (2012), doi: 10.1016/j.solidstatesciences.2012.11.017. 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 proof before it is published in its final 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.
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
ACCEPTED MANUSCRIPT
RI PT
Biocorrosion and Osteoconductivity of PCL/nHAp Composite Porous Film-Based Coating of Magnesium Alloy
Abdalla Abdal-hay1,2, Touseef Amna3, Jae Kyoo Lim1,* 1
Dept. of Bionano System Engineering, College of Engineering, Chonbuk National University,
Chonbuk National University, Jeonju 561-756, Republic of Korea, Department of Animal Resources and Biotechnology, Chonbuk NationalUniversity, Jeonju 561756, Republicof Korea
AC C
EP
TE D
3
Dept. of Mechanical Design Engineering, Advanced Wind Power System Research Institute,
M AN U
2
SC
Jeonju 561-756, Republic of Korea,
*Corresponding author: Tel: +82-63-20-2321 Fax: +82-63-270-4439 E-mail: [email protected] (Jae Kyoo Lim)
1
ACCEPTED MANUSCRIPT
Abstract The present study was aimed at designing a novel porous hydroxyapatite/poly(εcaprolactone) (nHAp/PCL) hybrid nano-composite matrix on a magnesium substrate with high
RI PT
and low porosity. The coated samples were prepared using a dip-coating technique in order to enhance the bioactivity and biocompatibility of the implant and to control the degradation rate of magnesium alloys. The mechanical and biocompatible properties of the coated and uncoated
SC
samples were investigated and an in vitro test for corrosion was conducted by electrochemical polarization and measurement of weight loss. The corrosion test results demonstrated that both
M AN U
the pristine PCL and nHAp/PCL composites showed good corrosion resistance in SBF. However, during the extended incubation time, the composite coatings exhibited more uniform and superior resistance to corrosion attack than pristine PCL, and were able to survive severe localized corrosion in physiological solution. Furthermore, the bioactivity of the composite film was determined by the rapid formation of uniform CaP nanoparticles on the sample surfaces
TE D
during immersion in SBF. The mechanical integrity of the composite coatings displayed better performance (~34% higher) than the uncoated samples. Finally, our results suggest that the
EP
nHAp incorporated with novel PCL composite membranes on magnesium substrates may serve as an excellent 3-D platform for cell attachment, proliferation, migration, and growth in bone
AC C
tissue. This novel as-synthesized nHAp/PCL membrane on magnesium implants could be used as a potential material for orthopedic applications in the future. Keywords: Hydroxyapatite, PCL, Bio-corrosion; Dip-coating; Magnesium alloys; Tissue engineering scaffold
1. Introduction 2
ACCEPTED MANUSCRIPT
Magnesium is an amazing material with outstanding properties that make it the material of choice for a wide range of applications such as automotive, aircraft, aerospace, architectural, packaging, and biomaterials. On the other hand, magnesium and its alloys are susceptible to
RI PT
dissolution in an aqueous environment, especially those containing chloride ion electrolytes, mainly due to their very low corrosion potential [1]. However, the biodegradable nature of magnesium is a desirable advantage as compared [2, 3] with conventional implants made from
SC
materials such as cobalt/chromium-based alloys, titanium and its alloys, stainless steels, and so on. All of these materials are non-degradable [4-6], though it has been established that the
M AN U
desirable characteristics of implant materials include an ability to degrade after the bone has healed, as many problems with implants arise if the implants are not degradable [7]. However, considerable work is still needed to control the degradation rate and improve the corrosion resistance of magnesium before it can be accepted into wide use in orthopedic and cardiovascular
TE D
implants. Immense effort has been devoted to the improvement of the general corrosion resistance of magnesium-based alloys, such as by decreasing the impurity level [8], adding superior corrosion-resistant materials, and producing a more homogeneous microstructure [1].
EP
The reactivity of the surface of magnesium is high, and therefore surface modification could be an ideal way to improve biological functions, control the degradation rate, avoid severe localized
AC C
surface corrosion, and improve the mechanical interlocking between the implant and bones, particularly through the use of organic coatings. Organic coatings can act as barrier layers to separate the substrate metal from its environment during the initial healing period. Earlier studies [9],[10], [11], [12], have been conducted on the properties of pure polymeric coatings of magnesium-based alloys, but the results obtained were not encouraging, as they were unable to meet the biological function requirements. The main drawback of plain organic coatings is that 3
ACCEPTED MANUSCRIPT
they do not show any intrinsic bioactivity, or in other words, they do not induce any positive regenerative response in bone cells when implanted. More recently, organic/inorganic nanocomposites have attracted significant attention
RI PT
worldwide from both academic and industrial viewpoints [13], especially in terms of bone tissue regeneration [14]. Such nanocomposites have the potential for greater function and performance than pure organic or inorganic materials. Moreover, these nanocomposites mimic the
SC
multifunctional performance of natural bone and play an important role in tissue engineering by directing cell growth either seeded from within the 3-D membrane structure of the scaffold or
M AN U
migrating from the surrounding tissues [15]. Bone is a biomineral material, one-third of which is inorganic (mainly partially carbonated HAp on the nanometer scale) and two-thirds of which is organic (mainly collagen). The coating layer should be osteoconductive so that osteoprogenitor cells can adhere to and migrate on the scaffolds, differentiate, and finally form new bone after
TE D
complete degradation of the temporary scaffold. Nanohydroxyapatite (nHAp) is considered to be an ideal bioactive material whose composition and crystal structure is very similar to that of bone [7, 16], and has bone bonding ability in vivo. In addition, this material possesses excellent
EP
osteoconductivity, osteoinductivity, corrosion resistance, and high chemical stability, and has therefore, been used in many forms as a substitute for regeneration material. Despite these
AC C
advantages, hydroxyapatite has very poor mechanical stability, is inherently brittle, and difficult to process, thus limiting its uses in the regeneration of load-bearing bone defects [16]. Thus, coatings on the magnesium surface could help to improve the properties of implants such as cell affinity, osteoconductivity, and enhancement of corrosion resistance, leading to enhanced bone formation. Hence, synthesized polymeric/inorganic nanocomposites are promising materials for filling this need. Polycaprolactone is a biodegradable polymer, has been used in bone tissue 4
ACCEPTED MANUSCRIPT
engineering, and gives excellent results when mixed with an inorganic material [17] such as hydroxyapatite. Thus, keeping in mind the unique properties of PCL and hydroxyapatite, we attempted to improve magnesium implants via deposition of PCL/nHAp using a dip-coating
RI PT
method. We report here the possibility of these polymeric membranes for controlling the degradation of magnesium alloys under in vitro conditions, and address the cytocompatibility and mechanical integrity of the deposited samples during degradation. We introduce here for the
SC
first time a porous PCL/nHAp nanocomposite film with a broad range of properties which may
M AN U
potentially be applied to future clinical uses.
2. Experimental 2.1. Materials
TE D
The following chemicals were used without further purification or modification: calcium nitrate tetrahydrate (Ca(NO3)2.4H2O), diammonium hydrogen phosphate (NH4)2HPO4 (Showa Chemical, Japan), aqueous ammonia solution, ethanol (Samchun Pure Chemical, South Korea),
EP
dichloromethane (DCM) (Junse Chemical, Japan), and PCL (Mw: 70,000 to 90,000 by GPC, Sigma Aldrich, Korea). Hank’s balanced solution was utilized as a simulated body fluid (SBF) (
AC C
H2387, Sigma Aldrich, Korea). Die-cast magnesium alloy AM50 (Mg-4.98Al-0.29Mn in mass pct) was used as a substrate material(South-Korean Manufacturers, Korea).
2.2. Synthesis of HAp nanopowder The synthesis of HAp nanocrystals was carried out via wet chemical precipitation without any aid of surfactants or co-surfactants since they do not have adequate biological properties and 5
ACCEPTED MANUSCRIPT
are suspected to produce allergic reactions due to their non-biodegradability [18, 19]. HAp nanopowder was synthesized using a wet chemical procedure. 0.6 M (NH4)2HPO4 aqueous solution was added at a rate of 0.4 mL min-1 to 1 M Ca(NO3)2.4H2O aqueous solution with
RI PT
strong stirring. The Ca/P ratio was 1.67 when these solutions were mixed to produce stoichiometric HAp, and ammonium solution was used to adjust the pH of the solution to 11. After stirring the solution at room temperature (RT) for 3 h, it was heated to 90oC for 1.5 h under
SC
continuous stirring. The resultant precipitate was kept for 24 h at RT under vigorous stirring to produce a homogeneous solution. Finally, the solution was filtered, and the precipitate was
M AN U
washed several times by a mixture of distilled water and alcohol (volume ratio=1:1) until the pH was 7. The final precipitated HAp was dispersed in alcohol to modify the agglomeration of the obtained powder. The powder was dried under vacuum for 24 h and then calcined at 650οC for 4 h in ambient air with a heating rate of 2οC min-1. The obtained nanoscale HAp powder (Fig.1)
EP
2.3. Sample preparation
TE D
was used as a filler onto PCL matrices.
Die-cast magnesium alloy AM50 was chosen as a substrate material to validate the
AC C
degradation rate and the corrosion behavior under test conditions. AM50 was cut into square pieces (12 x 12 mm2 and 2 mm in thickness). Prior to the coating process, the samples were mechanically polished with 600-1200 grit waterproof adhesive paper and then polished by 1 µm diamond grinding. Thereafter, the samples were sonicated in acetone for 5 min to remove residual grease and distilled water and were finally dried.
6
ACCEPTED MANUSCRIPT
2.2. Preparation and coating of PCL/nHAp hybrid nanocomposite layer The deposited porous layer (PCL/nHAp) was prepared by a dip-coating method with
RI PT
suitable modifications. Two different (5% and 10% by wt) PCL solutions were prepared. Briefly, PCL granules were dissolved in dichloromethane (DCM) solvent under magnetic stirring for 5 h. Colloids were prepared by mixing the synthesized HAp nanoparticles with polymer solution. The
SC
obtained colloidal (PCL solution with HAp) solution was stirred continuously for ~18 h followed by sonication for 30 min at room temperature. Two suspension solutions were prepared using
M AN U
one concentration of nHAp (20% by wt) [20] with each PCL concentration. In total, four different formulations (2 pristine PCL solutions and 2 colloids) were prepared and investigated as coating layers for a Mg-alloy substrate. All the specimens (before dip-coatings) were kept over a hot plate at 160ºC for 10 min to remove moisture and entrapped air from the substrate surfaces. The prepared substrates were immersed into the prepared solutions for 30 s to allow
TE D
wetting of the substrate. In order to obtain a stain-free surface, the specimens were slowly, mechanically pulled out of the solution at a speed of 2 mm/s. The motor lifting the substrate
2.3.
AC C
drier (10 mbar)
EP
worked continuously and was vibration-free. All the coated samples were dried with a vacuum
Characterization
Internal characterization was performed using transmission electron microscopy (TEM, CM 200, Philips, USA) with an accelerating voltage of 200 kV. The surface morphology of the 7
ACCEPTED MANUSCRIPT
deposited layers was characterized using a scanning electron microscope (JEOL JSM 820) coupled with an energy dispersive spectrometer (EDS). The samples were uniformly sprayed onto carbon tape, a Pt coating was applied for 10 s onto the synthesized layers, and images were
RI PT
acquired at various magnifications. These images were processed and analyzed with ImageJ software (National Institute of Mental Health, Bethesda, Maryland, USA) to gather information on mean pore length, porosity, and uniformity of the coated layers. Prior to ImageJ processing,
SC
the scale was set to the scale bar recorded in the image to obtain lengths in micrometers instead of pixels. The specimen thicknesses were measured using ultrasonic measurement (coating
M AN U
thickness gauge OMEGA instrument, OM179-745) with a precision of 1 µm. The surface roughness was determined using a Mitutoyo Surftest SV−402 profilometer. The mean roughness (Ra) of 7-10 different tests was calculated for each sample. The travelling distance and the speed of the stylus were 1.775 mm and 0.1 mm s-1, respectively.
TE D
The properties of wettability and surface free energy of the treated samples were measured with de-ionized water contact angle measurements using a contact angle meter (GBX, Digidrop, France). De-ionized water was automatically dropped (drop diameter: 6 µm) onto the sample.
EP
Measurement was done after 1 s and 10 s from the application of the water droplet. Three specimens for each sample were selected and three different locations on each were recorded. All
2.4.
AC C
measurements were carried out at room temperature.
Electrochemical corrosion test
Prior to the polarization test, the samples were immersed in 1000 mL SBF for 20 min to establish a relatively stable open circuit potential. The electrochemical behavior of both the 8
ACCEPTED MANUSCRIPT
uncoated and coated samples was determined in standard simulated body fluids (SBF) at pH 7.4 using the potentiodynamic polarization test (263A, EG&G PAR, USA). Briefly, a three-electrode cell with the sample as the working electrode, Ag/AgCl/1 M KCl (Satd.) as the reference
RI PT
electrode, and platinum as the counter electrode was used. The area of the working electrode (control and deposited layer samples) exposed to the solution was 1 cm2. A 1 mVs-1 scanning rate was applied during the potentiodynamic polarization test. The changes in the free corrosion
SC
potential (Ecorr) were monitored as a function of time. The temperature was maintained at 37 ±0.5ºC during the test. The results of the corrosion tests performed on the samples were
2.5.
Mechanical properties test
M AN U
calculated by extrapolating the polarization curve according to ASTM-G102-89.
TE D
The bending strength of the samples was evaluated as the maximum point of the stressdisplacement curve using a three-point bending test, which was carried out in a standard laboratory atmosphere with a mechanical tester (Computerized Instron Universal Testing
EP
Instrument model 4206). The tests were conducted at an across-head speed of 1 mm/min and
AC C
with a span length of 25 mm. Samples were prepared with dimensions of 33×5×5 mm, which is the required size for testing. At least 5 specimens were tested for each sample. Average values and standard deviation were calculated. 2.6.
Immersion tests
An immersion test was performed in a standard simulated body fluid (Hank’s balanced salts), with 5 specimens for each sample being immersed in 100 mL of solution. The electrolyte 9
ACCEPTED MANUSCRIPT
was saturated with atmospheric oxygen without stirring during the experiments. The temperature was maintained at 37.5 ±0.2ºC during the test. The immersed samples were extruded at different interval points of time, gently rinsed with distilled water, and dried at room temperature to avoid
RI PT
cracking. Before weight loss measurement, all samples were immersed in chromic acid solution (200 g/L Cr2O3 + 10 g/L AgNO3) for 10 minutes to remove the residual coated layer and corrosion products. Samples were then rinsed in distilled water, then acetone, and then dried. The
SC
surface appearance of the coated and uncoated immersed samples before and after removing the
2.7.
M AN U
residual coatings and corrosion products were observed by SEM.
Osteoblast Cell Culture and MTT Assay
The MTT assay was performed to measure cell viability in the presence of pristine and
TE D
composite samples. MC3T3 cell lines were used and were cultured in Dulbecco’s modified Eagle’s medium (DMEM) (Gibco Co, USA), supplemented with 10% fetal bovine serum (FBS)
EP
in a humidified incubator with 5% CO2 and 95% relative humidity at 37℃. The test for the coated samples and uncoated substrates was carried out by using an indirect extraction method,
AC C
where the immersion extracts collected from the immersion test were used for culturing cells. The extraction media was used at full concentration after one day of incubation in a humidified atmosphere with 5% CO2 at 37℃. Cell morphology and cell proliferation were observed under optical microscopy (LEICA Microsystems EZ4D). To assess cell proliferation, an MTT assay was carried out as a function of incubation time (1 -3 days) by calculating the RGR percentage (RGR = optical density for coated specimen/optical density for negative control group × 100%). 10
ACCEPTED MANUSCRIPT
Cell viability was obtained by comparing the absorbance of cells cultured with different samples to that of control wells containing only cells. The results were expressed as the means ± standard error of the mean. All the quantitative data were statistically analyzed to be expressed as mean ±
RI PT
standard deviation. Statistical analysis was determined by single factor ANOVA. P-values of less
SC
than 0.05 were considered significant.
3. Results and discussion
M AN U
In a dip-coating process, a substrate is dipped into a liquid coating solution, left for a reasonable time to allow wettability, and then withdrawn from the solution at a controlled speed and allowed to dry. Coating thickness generally increases with faster withdrawal speed. The thickness is controlled by the balance of forces at the stagnation point on the liquid surface. A
TE D
faster withdrawal speed pulls more fluid up onto the surface of the substrate before it has time to flow back down into the solution. The thickness is primarily affected by fluid viscosity, fluid density, temperature, and surface tension. Besides the possibility of increasing thickness and
EP
controlling pore size when using high volatile solvents, this method allows for the simultaneous coating of both sides of a sheet as well as the coating of non-flat substrates. In this way dip-
AC C
coating provides a level of adhesion needed in the process of organic coating of metal. This method, while excellent for producing high-quality coatings, requires precise control and a clean environment. In general, dip-coating is a reasonable coating technique to use whenever the substrate size, weight, or geometry makes spin coating difficult or impossible [12, 21]. For these reasons, many studies in chemical and nanomaterials engineering in both academia and industry make use of the dip-coating process. 11
ACCEPTED MANUSCRIPT
3.1. Surface properties of the coated samples
RI PT
The surface properties of the as-synthesized protective coatings (pristine PCL and nHAp/PCL) were described on the basis of surface morphology, surface roughness, wettability, and the layer thickness of the surface film deposited on the magnesium bars. The surface
SC
morphology observed by scanning electron microscope of the coated samples is shown in Fig. 2. The SEM images show that the plain PCL film produced fairly interconnected pore networks at
M AN U
both the concentrations used (5% and 10%). The surface microstructure modifications may have an intrinsic influence on biological functions such as cell attachment and ingrowth and, ultimately, tissue formation when the implant comes into contact with tissues at early stages [22]. The 3-D pore structure had a relatively uniform distribution and was free of defects on the
TE D
primary coating layer (the closest layer on the substrate), which is the layer responsible for providing adhesion between the organic coating and metal substrate [23]. The size of the pore phase at 5 and 10 wt% plain PCL was approximately 10 ±2 µm and 1.5 ±0.72 µm, respectively.
EP
These features are typical of porous polymer prepared by polymer-poor phase and polymer-rich phase separation using highly volatile solvents such as dichloromethane in the presence of
AC C
humidity. With the incorporation of nHAp (20 wt%) in polymer (5 and 10 wt%) matrices, the pore structure size decreased, especially at 10/90 PCL/DCM, but the macropore structure was still clearly observed at 5/95 PCL/DCM (Fig.2c). Although the HAp particles were embedded inside the matrix, the pore walls of the polymer were perfectly smooth and no HAp was observed on the polymer pore surfaces. The good interfacial bonding strength of nHAp with the polymer matrix can be explained by the hydrogen bonding formed between the -OH groups of HA and the 12
ACCEPTED MANUSCRIPT
=O sites of the PCL polymer [20]. The major peaks of the deposited composite layer on magnesium substrate are shown in the spectrum (Fig. 2). From the EDX peaks, it can be
RI PT
concluded that the deposited film is relatively free of impurities.
Nano-scaled composite coatings on magnesium bars may offer a promising therapy for natural bone regeneration since the surface characteristics, such as surface roughness, have a
SC
direct impact on the formation of bone cells. The surface roughness of the composite coatings may show an appropriate value. The composite film which contained nHAp in a PCL film (5
M AN U
wt%) showed a surface roughness of about 0.543 µm, whereas the plain PCL was 0.325 µm (Table 1). It has already been reported that rougher surfaces can adsorb more fibronectin and improve cell attachment over smoother surfaces [24]. Collectively, our results showed that the surface roughness of magnesium coatings increases with an increase in HAp content within the
TE D
matrix, as shown in Table 1.
PCL is a biodegradable and biocompatible polymer with a slow degradation rate which
EP
shows promise as a magnesium coating in bone tissue-engineering applications [25-27]. The surface of PCL is hydrophobic and does not exhibit any physiological activity, which makes it
AC C
unfavorable for cell growth when it comes into contact with the living body [28]. Therefore, the cytocompatibility of these synthesized materials should be improved before they can be used in clinical settings. Variation in the surface topologies as well as material compositions can be used to control the wetting process or the adsorption performance of PCL. The existence of hydrophilic domains on a material’s surface will have great influence on surface properties. It is well known that the water contact angle (WCA) reflects superficial properties of a material. The 13
ACCEPTED MANUSCRIPT
WCA of plain PCL and nHAp/PCL deposited layers on magnesium substrates are listed in Table 1. The nHAp/PCL composite membranes had lower contact angles than that of pure PCL membranes. It was observed that the WCA gradually decreased with increasing amounts of
RI PT
nHAp in the PCL layers deposited on Mg bars (inset of Fig. 2a, c-d). The WCA was decreased from 77.6ο (plain PCL) to 71.3ο (composite layer), as shown in Table 1. This phenomenon might be because of the wettable nHAp particles incorporated into the membranes, thus making the
SC
composite membranes more hydrophilic than pristine PCL. Another possible explanation is the increase in surface energy caused by the surface roughness of nHAp on the PCL surface. In
M AN U
conclusion, after nHAp incorporation, the surface property of plain PCL was altered. The introduction of nHAp provides an opportunity to further modify the surface biocompatibility. As mentioned previously, layer thickness can be controlled by various parameters. Our results showed that the polymer layers exhibited different thicknesses under the same dip-coating
TE D
conditions, as shown in Table.1. The composite films were much thicker than the corresponding plain films. The average thickness of pristine PCL/DCM (5/95) was 12.5 µm while that of
EP
PCL/nHAp film was 15.3 µm, Table.1.
AC C
3.2. Electrochemical corrosion analysis Fig. 3 illustrates the typical overall potentiodynamic behaviors of magnesium bars and dipcoated magnesium samples tested in a standard simulated body fluid at 37οC. The analysis data of these curves are shown in Table 2. The corrosion potentials (Ecorr) of the coated samples with 5 and 10 wt% plain PCL and 5 wt% PCL/nHAp were 200, 631, and 130 mV more positive, respectively, than an unprotected magnesium bar. At the same time, the values of corrosion 14
ACCEPTED MANUSCRIPT
current (Icorr) of protected samples were lower than those of uncoated samples. Therefore, values for both Icorr and Ecorr illustrated that plain PCL and nHAp/PCL composite would be able to
3.3.
RI PT
enhance the corrosion resistance and thereby control the degradation rate of magnesium alloy.
Immersion tests
SC
Fig. 4 shows the representative SEM images for the uncoated and coated samples before
M AN U
removing the corrosion products and after immersion in SBF at 37οC for 20 days. It is clear from the SEM images that the samples coated with plain PCL experienced cleavage and a high level of attack by water molecules. Water diffusion through tiny micro or nano defects on the coating layer deteriorated the coatings and destroyed the adhesion of the organic layer on the substrate. Earlier reports [29] have indicated that when the water molecules come into contact with
TE D
magnesium surfaces, hydrogen gas is produced by the electrochemical reaction (Eq. 1). However, the porous layer of the pure polymer fabricated via dip-coating cannot provide a fully porous structure mainly due to the organic layer, which contains three films: the primary,
EP
intermediate, and top-coating films. The primary film is responsible for providing good adhesion
AC C
between the whole organic layer and the metal substrate. Hence, the primary film is a dense, non-porous structure. As a result, the production of hydrogen is hindered and the osmotic pressure is increased. The tiny micro pores or/and in the range of nanometers pores contained in the primer film are not adequate for the speed of hydrogen release required from the magnesium substrate, i.e. hydrogen gas cannot be removed at an adequate velocity. The hydrogen that is produced results in an increase in gas pressure beneath the polymer. The hydrogen permeability of the polymer is insufficient to release the gas at the rate it is beging produced, which leads to 15
ACCEPTED MANUSCRIPT
excess pressure that results in the bursting of the polymer film. However, the reaction itself is initiated by failures in the polymer coating, Fig. 4a. While composite coatings may contain fully porous structures, could establishe space between the nano-particles and the polymer matrix
RI PT
sufficient to allow evacuation of hydrogen gas via the polymer matrix layer. To inform our hypothesis, we found that the composite coating layer was still intact and resisted rupturing and/or delamination from their substrates as shown in Fig. 4ab-d. From these findings, it is
M AN U
stability of the coating layers on magnesium substrates.
SC
assumed that the incorporation of nHAp into the polymer matrix could provide a sufficient
Mg(s)+ 2H2O→Mg2++ 2OH-(aq)+H2(g) ↑ (overall reaction)
(1)
PCL surfaces are inherently hydrophilic, which leads to a decrease in the stimulation and
TE D
precipitation of apatite on their surfaces. However, it is known that polymer surfaces gradually change towards a more hydrophilic structure in polar aqueous environments. In a comparison of coated samples with pristine PCL, the nHAp/PCL coatings showed a uniform distribution of
EP
precipitated Ca/P particles of a uniform size of around 230 nm. The overall pore structure of the composite layer was remarkably unchanged and free of defects, and the characteristic
AC C
distribution of macropores in the microporous matrix was still observed, Fig. 4. It is speculated that over a longer incubation time, the precipitation of Ca-P increases, resulting in a uniformly dense layer of Ca-P on the polymer layer, and that this amorphous inorganic layer finally crystallizes into apatite. In support of our hypothesis, we found that some areas (shown in the dashed box in Fig. 4d) show a dense coating with a characteristic cauliflower structure typically associated with apatite. Again, if a dense CaP layer forms, this layer will be able to add further 16
ACCEPTED MANUSCRIPT
protection to the magnesium substrates further the high formation of bone-like. This finding confirmed that the introduction of hydroxyapatite into the PCL matrix improved the bioactivity of the scaffold. These results indicate that composite films stimulate CaP compound in a regular
RI PT
phase more than plain PCL does. A reasonable explanation for this phenomenon could be that there is a lower release of Mg+2 from the polymer matrix of composite-film-coated magnesium samples during the whole immersion period in comparison with samples coated with a plain PCL
SC
film. The slight release of magnesium ions could lead to a slightly increased pH, which in turn could assist in the nucleation of CaP in a uniform manner [30]. Another possible reason for the
M AN U
better protection seen with polymer matrix composite film-coated magnesium samples is the increase in pH, leading to the dissolution of ions from nHAp particles [31], as shown in Fig. 5. It should be mentioned that increased degradation of magnesium substrate without sufficient apatite formation can lead to deleterious results. Several reports have suggested that reduction of
TE D
the degradation rate of magnesium implants lowers the accumulation of hydrogen gas bubbles and halts the lowering of the local pH, which could otherwise lead to necrosis of tissues,
EP
infections, and alkaline poisoning [8, 29, 32].
SEM analysis showed remarkable changes in the morphology of the composite surface
AC C
during the degradation period in SBF, as shown in Fig. 6. After removing the corrosion products from the magnesium surface, it was apparent that the uncoated samples experienced a large amount of severe localized corrosion as well as pitting corrosion on the whole magnesium surface as shown in Fig. 6a. Continued localized corrosion at the valuable sites gives rise to nonuniformly corroded sites [33]. The corrosion morphology of the samples protected with composite films showed uniform corrosion on the magnesium surfaces, Fig. 6b-d. Quantitatively, 17
ACCEPTED MANUSCRIPT
the results of the immersion experiment clearly showed that PCL/nHAp composite films have better stability than those of plain PCL films, a property which could contribute to the slower degradation rate of the composite films compared with plain PCL films during the length of the
RI PT
incubation.
Fig.7 shows a graph of the total mass of magnesium dissolved into the SBF solution as a
SC
function of immersion time in the bath. In comparison with uncoated samples, the plain PCL and PCL matrix composite coatings decreased the corrosion rate of the magnesium alloy by factors
M AN U
of 3 and 4, respectively. These degradation rates of both uncoated and coated samples are in good agreement with the above pre-mentioned section. It is clear that both the coated samples showed almost no weight loss at the initial immersion time, especially the samples coated with composite films. This observation confirms that these coatings can provide a sufficient
several days [34, 35].
In vitro loss of mechanical integrity
EP
3.4.
TE D
protection at the first implantation time and avoid the inflammation stage which can last for
AC C
A material used for bone implant application should have a mechanical integrity sufficient for bone cells to attach, proliferate, and differentiate in a manner similar to native ECM until the healing process is complete. Hence, mechanical properties are an important factor for bioimplant materials. Fig. 8 illustrates the influence of in vitro degradation over 20 days on the bending strength of uncoated and nHAp/PCL-coated magnesium samples. From the selected stressdisplacement curves it can be seen that the samples protected with composite film showed better 18
ACCEPTED MANUSCRIPT
mechanical properties than uncoated ones. Furthermore, the tensile strength of the coated and uncoated samples decreased by around 13 and 34%, respectively, in comparison to preimmersion magnesium alloy (i.e. time point 0). This difference is likely due to surface defects
RI PT
such as severe localized corrosion and pitting corrosion formed during degradation. Conclusively, the mechanical degradation of magnesium-based polymer nanocomposites coating
Cell culture
M AN U
3.5.
SC
was significantly decreased by the incorporation of nano inorganic particles.
To avoid undesired effects when materials are implanted in vivo, candidate materials have to be assayed for their compatibility using cell cultures [36]. This step is essential to ensure that the optimal material for a specific biomedical application is used in patients. As mentioned
TE D
above, biocompatibility, cell attachment, migration, proliferation, and differentiation are significantly influenced by surface conditions and chemical compositions of the substrates [37]. Use of a biocompatible material such as HAp [17] in combination with a biodegradable synthetic
EP
polymer gives a ‘bioartificial advanced composite’ with enhanced biocompatibility and other biological functions. Furthermore, the ionic dissolution from the inorganic compounds,
AC C
particularly in the case of nHAp (Fig. 5), plays an essential role in such cellular processes as bone cell differentiation and the genes expression. Based on this consideration, MTT assay data showed that the incorporation of nHAp within the polymer matrix could improve the cell viability and survival after a 24 h incubation in comparison with plain PCL, Fig. 9. After a longer incubation time of three days, it was clear that the composite coating on the magnesium substrate resulted in more proliferation and continuous cell growth. These results confirmed that nHAp 19
ACCEPTED MANUSCRIPT
plays an important role in supporting cell proliferation and bone cell differentiation. In contrast, uncoated samples showed worse behavior in comparison with both of the coated samples. We predicted that the number of cells would be statistically increased for the coated Mg at all time
RI PT
intervals compared with uncoated samples, meaning that all of the coated samples have a significantly better initial biocompatibility than the uncoated sample. To verify this hypothesis, the respective optical microscopy morphologies of osteoblasts of the MC3T3 cell line after one
SC
day of culture in extracts from the immersion test were observed, Fig. 10. From these images it can be seen that the cell morphologies of cells treated with extracts taken from coated samples
M AN U
were normal and healthy and showed good survival behavior, especially the samples with the composite coating, while no cells were observed in cultures grown with extracts of the uncoated samples, indicating that uncoated Mg does not promote cell growth. However, it has been reported that Mg is inherently biocompatible [29] and that the aluminum content in the AM50
TE D
alloy does not have negative effects on the viability and proliferation of osteoblasts, as reported by Ciapetti et. al [36]. This observation may be due to the steep increase in pH or the release of ions into the medium from the uncoated samples [12, 38]. Based on the results obtained from the
EP
in vitro tests, it is believed that the incorporation of nHAp particles into the polymer matrix improves the biocompatibility and enhances the development of the connection between the
AC C
magnesium implants and natural bone. Our results are in agreement with all studies considered to have shown improved biocompatibility and higher expression of osteogenic markers when HAp is incorporated into the polymers [39]. This behavior may be partially explained by improved adsorption of extracellular matrix (ECM) proteins, such as fibronectin or vibronectin, to composites as compared with plain polymers. Hence, the finding of good compatibility in vitro
20
ACCEPTED MANUSCRIPT
makes nHAp/PCL advanced composite coating of magnesium worthy of further studies for
RI PT
development of surgical implants.
4. Conclusion
In summary, novel nHAp/PCL advanced composite film coatings on magnesium substrates
SC
were developed using a simple and cost effective dip-coating technique. In addition, various physicochemical properties such as surface properties and biocorrosion performance were
M AN U
investigated. The matrices were evaluated for in vitro cytocompatibility under simulated physiological conditions using an MTT assay. The surface modifications in the magnesium substrate induced promising changes in such characteristics as surface roughness and wettability. Our results demonstrated that the incorporation of nHAp particles into the polymer matrix layer
TE D
improved both the degradation resistance and biological function of the magnesium substrates. The bone implant showed good cytocompatibility to better mimic the surface characteristics of natural bone. In addition, it is speculated that the level of osteoblastic differentiation activity will
EP
also increase significantly with the incorporation of nHAp into the PCL polymer matrix composite coatings on magnesium implants. Further in-depth studies using an in vivo animal
AC C
model will be required to evaluate the degradation performance and biological functions of assynthesized hybrid polymer matrix nanocomposite coated on magnesium based implants.
Acknowledgments
21
ACCEPTED MANUSCRIPT
The authors gratefully acknowledge support for this research by the Ministry of Education, Science Technology (MEST) and National Research Foundation of Korea (NRF) through the
RI PT
Human Resource Training Project for Regional Innovation. References
[1] X. Zhang, G. Yuan, L. Mao, J. Niu, P. Fu, W. Ding, Journal of the Mechanical Behavior of
SC
Biomedical Materials. [2] J.A. Cowan, BioMetals, 15 (2002) 225-235.
M AN U
[3] M.E. Maguire, J.A. Cowan, BioMetals, 15 (2002) 203-210.
[4] M.P. Staiger, A.M. Pietak, J. Huadmai, G. Dias, Biomaterials, 27 (2006) 1728-1734. [5] T. Jinno, V.M. Goldberg, D. Davy, S. Stevenson, Journal of Biomedical Materials Research, 42 (1998) 20-29.
TE D
[6] R. Van Noort, Journal of Materials Science, 22 (1987) 3801-3811. [7] W. TIANSHI, Z. RENJI, A.Y. YONGNIAN, Journal of BIOACTIVE AND COMPATIBLE POLYMERS, 24 (2009) 169-182.
EP
[8] S. Zhang, X. Zhang, C. Zhao, J. Li, Y. Song, C. Xie, H. Tao, Y. Zhang, Y. He, Y. Jiang, Y. Bian, Acta Biomaterialia, 6 (2010) 626-640.
AC C
[9] X.N. Gu, Y.F. Zheng, Q.X. Lan, Y. Cheng, Z.X. Zhang, T.F. Xi, D.Y. Zhang4, Biomedical Materials, 4 (2009) 044109.
[10] J. Li, P. Cao, X. Zhang, S. Zhang, Y. He, Journal of Materials Science, 45 (2010) 60386045.
[11] K. Bai, Y. Zhang, Z. Fu, C. Zhang, X. Cui, E. Meng, S. Guan, J. Hu, Materials Letters, 73 (2012) 59-61. 22
ACCEPTED MANUSCRIPT
[12] L. Xu, A. Yamamoto, Colloids and Surfaces B: Biointerfaces. [13] J.K. Pandey, K. Raghunatha Reddy, A. Pratheep Kumar, R.P. Singh, Polymer Degradation and Stability, 88 (2005) 234-250.
RI PT
[14] W.-E. Teo, S. Ramakrishna, Composites Science and Technology, 69 (2009) 1804-1817. [15] M. M, N.-E. M, International Journal of Nanomedicine, 6 (2011) 1651-1659
372.
SC
[16] K.A. Khalil, S.W. Kim, H.Y. Kim, Materials Science and Engineering: A, 456 (2007) 368-
[17] H.-H. Lee, H.-S. Yu, J.-H. Jang, H.-W. Kim, Acta Biomaterialia, 4 (2008) 622-629.
M AN U
[18] H. Zhou, J. Lee, Acta Biomaterialia, (2011).
[19] B. Nasiri-Tabrizi, P. Honarmandi, R. Ebrahimi-Kahrizsangi, P. Honarmandi, Materials Letters, 63 (2009) 543-546.
[20] S.Z. Fu, X.H. Wang, G. Guo, S. Shi, M. Fan, H. Liang, F. Luo, Z.Y. Qian, Journal of
TE D
Biomedical Materials Research Part B: Applied Biomaterials, 97B (2011) 74-83. [21] T.F. Conceicao, N. Scharnagl, C. Blawert, W. Dietzel, K.U. Kainer, Corrosion Science, 52 (2010) 2066-2079.
EP
[22] Y. Zhu, C. Gao, J. Shen, Biomaterials, 23 (2002) 4889-4895. [23] P. Sørensen, S. Kiil, K. Dam-Johansen, C. Weinell, Journal of Coatings Technology and
AC C
Research, 6 (2009) 135-176.
[24] K.M. Woo, J. Seo, R. Zhang, P.X. Ma, Biomaterials, 28 (2007) 2622-2630. [25] E.-J. Choi, J.-K. Park, Polymer Degradation and Stability, 52 (1996) 321-326. [26] D. Darwis, H. Mitomo, F. Yoshii, Polymer Degradation and Stability, 65 (1999) 279-285. [27] X. Zheng, C.A. Wilkie, Polymer Degradation and Stability, 82 (2003) 441-450.
23
ACCEPTED MANUSCRIPT
[28] Q. Zhao, J. Tao, R.C.M. Yam, A.C.K. Mok, R.K.Y. Li, C. Song, Polymer Degradation and Stability, 93 (2008) 1571-1576. [29] B. Zberg, P.J. Uggowitzer, J.F. Loffler, Nat Mater, 8 (2009) 887-891.
RI PT
[30] L. Li, J. Gao, Y. Wang, Surface and Coatings Technology, 185 (2004) 92-98.
[31] Y. Ito, H. Hasuda, M. Kamitakahara, C. Ohtsuki, M. Tanihara, I.-K. Kang, O.H. Kwon, Journal of Bioscience and Bioengineering, 100 (2005) 43-49.
SC
[32] J.-H. Jo, J.-Y. Hong, K.-S. Shin, H.-E. Kim, Y.-H. Koh, Journal of Biomaterials Applications, 5 (2011) 8.
M AN U
[33] Y. Xin, T. Hu, P.K. Chu, Acta Biomaterialia, 7 (2011) 1452-1459.
[34] J. Yang, F. Cui, I. Lee, Annals of Biomedical Engineering, 39 (2011) 1857-1871. [35] S. Chen, S. Guan, S. Hou, L. Wang, S. Zhu, J. Wang, W. Li, Surface and Interface Analysis, 43 (2011) 1575-1580.
TE D
[36] G. Ciapetti, J.L. González-Carrasco, L. Savarino, M.A. Montealegre, S. Pagani, N. Baldini, Biomaterials, 26 (2005) 849-859.
[37] F. Yang, R. Murugan, S. Ramakrishna, X. Wang, Y.X. Ma, S. Wang, Biomaterials, 25
EP
(2004) 1891-1900.
[38] F. Witte, F. Feyerabend, P. Maier, J. Fischer, M. Störmer, C. Blawert, W. Dietzel, N. Hort,
AC C
Biomaterials, 28 (2007) 2163-2174.
[39] L. Shor, S. Güçeri, X. Wen, M. Gandhi, W. Sun, Biomaterials, 28 (2007) 5291-5297.
Figure captions and Tables
Fig. 1. TEM morphology of synthesized HAp nano-scale 24
ACCEPTED MANUSCRIPT
Fig. 2. Surface appearance in SEM of magnesium samples coated with, 5 and 10 wt% plain PCL (a and b) and nHAp at 5 and 10% by wt, PCL (c and d); SEM-EDX of the composite
RI PT
coated samples. Insets are the water contact angles Fig. 3. Potentiodynamic polarization curves for coated and uncoated samples
Fig. 4. Surface appearance in SEM after immersion in SBF for 20 days for the samples coated
SC
with (a) plain PCL, (b-d) composite coating layers. Parts c and d are respective circular
M AN U
areas in b.
Fig. 5. Schematic diagram illustrating the apatite-like nucleation and cell behavior on magnesium substrate coated with composite films
Fig. 6. Scanning electron micrographs of the uncoated (a) and composite film-coated (b) samples
TE D
after removing the corrosion products. Parts c and d are respective circular areas in b. Fig. 7. Total weight lost by magnesium alloys uncoated or coated with plain PCL and
EP
PCL/nHAp composite coated samples Fig. 8. Selected stress-displacement curves
AC C
Fig. 9. MTT cytotoxicity assay of the coated and uncoated magnesium samples after 1 and 3 days of incubation. The data are reported as the mean ± standard deviation (n=3 and p‹0.05)
25
ACCEPTED MANUSCRIPT
Fig. 10. Optical micrographs of osteoblast cell (MC3T3) morphology after one day of culture, including (a) uncoated bars of Mg, (b) bars coated with pristine PC, and (c) PCL/nHAp
AC C
EP
TE D
M AN U
SC
RI PT
composite-coated Mg bars
26
ACCEPTED MANUSCRIPT
Table 1. Layer thickness, surface roughness, and water contact angle of coated samples
Layer thickness (µm)
12.5±1.564
0.325±0.09
nHAp in 5% wt PCL
15.3±2.21
0.543±0.057
nHAp in 10% wt PCL
34.12±2.45
0.952±0.066
M AN U
*
10 S
77.6±0.55
77.9.45±0.45
75.7±2
73.3±2.15
71.3±1.23
70.5±1
SC
Plain PCL at 5% wt
1S
RI PT
Sample
Water contact angle ± STD (o)
Surface roughness*± STD (µm)
AC C
EP
TE D
The roughness values correspond to the average surface roughness (Ra)
27
ACCEPTED MANUSCRIPT
Table 2. Summary of corrosion potential voltage and current density values of uncoated and coated samples
5% wt PC
10% wt PCL
Ecorr (mV)
-1340
-1140
-709
-1210
Current density (µA cm-2)
103
7.723
1.721
3.6
SC M AN U TE D EP
AC C
Figure
28
PCL/nHAp
RI PT
Uncoated
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
S1293-2558(12)00380-9
DOI:
10.1016/j.solidstatesciences.2012.11.017
Reference:
SSSCIE 4635
To appear in:
Solid State Sciences
Received Date: 25 May 2012 Accepted Date: 7 November 2012
Please cite this article as: A. Abdal-hay, T. Amna, J.K. Lim, Biocorrosion and Osteoconductivity of PCL/ nHAp Composite Porous Film-Based Coating of Magnesium Alloy, Solid State Sciences (2012), doi: 10.1016/j.solidstatesciences.2012.11.017. 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 proof before it is published in its final 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.
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
ACCEPTED MANUSCRIPT
RI PT
Biocorrosion and Osteoconductivity of PCL/nHAp Composite Porous Film-Based Coating of Magnesium Alloy
Abdalla Abdal-hay1,2, Touseef Amna3, Jae Kyoo Lim1,* 1
Dept. of Bionano System Engineering, College of Engineering, Chonbuk National University,
Chonbuk National University, Jeonju 561-756, Republic of Korea, Department of Animal Resources and Biotechnology, Chonbuk NationalUniversity, Jeonju 561756, Republicof Korea
AC C
EP
TE D
3
Dept. of Mechanical Design Engineering, Advanced Wind Power System Research Institute,
M AN U
2
SC
Jeonju 561-756, Republic of Korea,
*Corresponding author: Tel: +82-63-20-2321 Fax: +82-63-270-4439 E-mail: [email protected] (Jae Kyoo Lim)
1
ACCEPTED MANUSCRIPT
Abstract The present study was aimed at designing a novel porous hydroxyapatite/poly(εcaprolactone) (nHAp/PCL) hybrid nano-composite matrix on a magnesium substrate with high
RI PT
and low porosity. The coated samples were prepared using a dip-coating technique in order to enhance the bioactivity and biocompatibility of the implant and to control the degradation rate of magnesium alloys. The mechanical and biocompatible properties of the coated and uncoated
SC
samples were investigated and an in vitro test for corrosion was conducted by electrochemical polarization and measurement of weight loss. The corrosion test results demonstrated that both
M AN U
the pristine PCL and nHAp/PCL composites showed good corrosion resistance in SBF. However, during the extended incubation time, the composite coatings exhibited more uniform and superior resistance to corrosion attack than pristine PCL, and were able to survive severe localized corrosion in physiological solution. Furthermore, the bioactivity of the composite film was determined by the rapid formation of uniform CaP nanoparticles on the sample surfaces
TE D
during immersion in SBF. The mechanical integrity of the composite coatings displayed better performance (~34% higher) than the uncoated samples. Finally, our results suggest that the
EP
nHAp incorporated with novel PCL composite membranes on magnesium substrates may serve as an excellent 3-D platform for cell attachment, proliferation, migration, and growth in bone
AC C
tissue. This novel as-synthesized nHAp/PCL membrane on magnesium implants could be used as a potential material for orthopedic applications in the future. Keywords: Hydroxyapatite, PCL, Bio-corrosion; Dip-coating; Magnesium alloys; Tissue engineering scaffold
1. Introduction 2
ACCEPTED MANUSCRIPT
Magnesium is an amazing material with outstanding properties that make it the material of choice for a wide range of applications such as automotive, aircraft, aerospace, architectural, packaging, and biomaterials. On the other hand, magnesium and its alloys are susceptible to
RI PT
dissolution in an aqueous environment, especially those containing chloride ion electrolytes, mainly due to their very low corrosion potential [1]. However, the biodegradable nature of magnesium is a desirable advantage as compared [2, 3] with conventional implants made from
SC
materials such as cobalt/chromium-based alloys, titanium and its alloys, stainless steels, and so on. All of these materials are non-degradable [4-6], though it has been established that the
M AN U
desirable characteristics of implant materials include an ability to degrade after the bone has healed, as many problems with implants arise if the implants are not degradable [7]. However, considerable work is still needed to control the degradation rate and improve the corrosion resistance of magnesium before it can be accepted into wide use in orthopedic and cardiovascular
TE D
implants. Immense effort has been devoted to the improvement of the general corrosion resistance of magnesium-based alloys, such as by decreasing the impurity level [8], adding superior corrosion-resistant materials, and producing a more homogeneous microstructure [1].
EP
The reactivity of the surface of magnesium is high, and therefore surface modification could be an ideal way to improve biological functions, control the degradation rate, avoid severe localized
AC C
surface corrosion, and improve the mechanical interlocking between the implant and bones, particularly through the use of organic coatings. Organic coatings can act as barrier layers to separate the substrate metal from its environment during the initial healing period. Earlier studies [9],[10], [11], [12], have been conducted on the properties of pure polymeric coatings of magnesium-based alloys, but the results obtained were not encouraging, as they were unable to meet the biological function requirements. The main drawback of plain organic coatings is that 3
ACCEPTED MANUSCRIPT
they do not show any intrinsic bioactivity, or in other words, they do not induce any positive regenerative response in bone cells when implanted. More recently, organic/inorganic nanocomposites have attracted significant attention
RI PT
worldwide from both academic and industrial viewpoints [13], especially in terms of bone tissue regeneration [14]. Such nanocomposites have the potential for greater function and performance than pure organic or inorganic materials. Moreover, these nanocomposites mimic the
SC
multifunctional performance of natural bone and play an important role in tissue engineering by directing cell growth either seeded from within the 3-D membrane structure of the scaffold or
M AN U
migrating from the surrounding tissues [15]. Bone is a biomineral material, one-third of which is inorganic (mainly partially carbonated HAp on the nanometer scale) and two-thirds of which is organic (mainly collagen). The coating layer should be osteoconductive so that osteoprogenitor cells can adhere to and migrate on the scaffolds, differentiate, and finally form new bone after
TE D
complete degradation of the temporary scaffold. Nanohydroxyapatite (nHAp) is considered to be an ideal bioactive material whose composition and crystal structure is very similar to that of bone [7, 16], and has bone bonding ability in vivo. In addition, this material possesses excellent
EP
osteoconductivity, osteoinductivity, corrosion resistance, and high chemical stability, and has therefore, been used in many forms as a substitute for regeneration material. Despite these
AC C
advantages, hydroxyapatite has very poor mechanical stability, is inherently brittle, and difficult to process, thus limiting its uses in the regeneration of load-bearing bone defects [16]. Thus, coatings on the magnesium surface could help to improve the properties of implants such as cell affinity, osteoconductivity, and enhancement of corrosion resistance, leading to enhanced bone formation. Hence, synthesized polymeric/inorganic nanocomposites are promising materials for filling this need. Polycaprolactone is a biodegradable polymer, has been used in bone tissue 4
ACCEPTED MANUSCRIPT
engineering, and gives excellent results when mixed with an inorganic material [17] such as hydroxyapatite. Thus, keeping in mind the unique properties of PCL and hydroxyapatite, we attempted to improve magnesium implants via deposition of PCL/nHAp using a dip-coating
RI PT
method. We report here the possibility of these polymeric membranes for controlling the degradation of magnesium alloys under in vitro conditions, and address the cytocompatibility and mechanical integrity of the deposited samples during degradation. We introduce here for the
SC
first time a porous PCL/nHAp nanocomposite film with a broad range of properties which may
M AN U
potentially be applied to future clinical uses.
2. Experimental 2.1. Materials
TE D
The following chemicals were used without further purification or modification: calcium nitrate tetrahydrate (Ca(NO3)2.4H2O), diammonium hydrogen phosphate (NH4)2HPO4 (Showa Chemical, Japan), aqueous ammonia solution, ethanol (Samchun Pure Chemical, South Korea),
EP
dichloromethane (DCM) (Junse Chemical, Japan), and PCL (Mw: 70,000 to 90,000 by GPC, Sigma Aldrich, Korea). Hank’s balanced solution was utilized as a simulated body fluid (SBF) (
AC C
H2387, Sigma Aldrich, Korea). Die-cast magnesium alloy AM50 (Mg-4.98Al-0.29Mn in mass pct) was used as a substrate material(South-Korean Manufacturers, Korea).
2.2. Synthesis of HAp nanopowder The synthesis of HAp nanocrystals was carried out via wet chemical precipitation without any aid of surfactants or co-surfactants since they do not have adequate biological properties and 5
ACCEPTED MANUSCRIPT
are suspected to produce allergic reactions due to their non-biodegradability [18, 19]. HAp nanopowder was synthesized using a wet chemical procedure. 0.6 M (NH4)2HPO4 aqueous solution was added at a rate of 0.4 mL min-1 to 1 M Ca(NO3)2.4H2O aqueous solution with
RI PT
strong stirring. The Ca/P ratio was 1.67 when these solutions were mixed to produce stoichiometric HAp, and ammonium solution was used to adjust the pH of the solution to 11. After stirring the solution at room temperature (RT) for 3 h, it was heated to 90oC for 1.5 h under
SC
continuous stirring. The resultant precipitate was kept for 24 h at RT under vigorous stirring to produce a homogeneous solution. Finally, the solution was filtered, and the precipitate was
M AN U
washed several times by a mixture of distilled water and alcohol (volume ratio=1:1) until the pH was 7. The final precipitated HAp was dispersed in alcohol to modify the agglomeration of the obtained powder. The powder was dried under vacuum for 24 h and then calcined at 650οC for 4 h in ambient air with a heating rate of 2οC min-1. The obtained nanoscale HAp powder (Fig.1)
EP
2.3. Sample preparation
TE D
was used as a filler onto PCL matrices.
Die-cast magnesium alloy AM50 was chosen as a substrate material to validate the
AC C
degradation rate and the corrosion behavior under test conditions. AM50 was cut into square pieces (12 x 12 mm2 and 2 mm in thickness). Prior to the coating process, the samples were mechanically polished with 600-1200 grit waterproof adhesive paper and then polished by 1 µm diamond grinding. Thereafter, the samples were sonicated in acetone for 5 min to remove residual grease and distilled water and were finally dried.
6
ACCEPTED MANUSCRIPT
2.2. Preparation and coating of PCL/nHAp hybrid nanocomposite layer The deposited porous layer (PCL/nHAp) was prepared by a dip-coating method with
RI PT
suitable modifications. Two different (5% and 10% by wt) PCL solutions were prepared. Briefly, PCL granules were dissolved in dichloromethane (DCM) solvent under magnetic stirring for 5 h. Colloids were prepared by mixing the synthesized HAp nanoparticles with polymer solution. The
SC
obtained colloidal (PCL solution with HAp) solution was stirred continuously for ~18 h followed by sonication for 30 min at room temperature. Two suspension solutions were prepared using
M AN U
one concentration of nHAp (20% by wt) [20] with each PCL concentration. In total, four different formulations (2 pristine PCL solutions and 2 colloids) were prepared and investigated as coating layers for a Mg-alloy substrate. All the specimens (before dip-coatings) were kept over a hot plate at 160ºC for 10 min to remove moisture and entrapped air from the substrate surfaces. The prepared substrates were immersed into the prepared solutions for 30 s to allow
TE D
wetting of the substrate. In order to obtain a stain-free surface, the specimens were slowly, mechanically pulled out of the solution at a speed of 2 mm/s. The motor lifting the substrate
2.3.
AC C
drier (10 mbar)
EP
worked continuously and was vibration-free. All the coated samples were dried with a vacuum
Characterization
Internal characterization was performed using transmission electron microscopy (TEM, CM 200, Philips, USA) with an accelerating voltage of 200 kV. The surface morphology of the 7
ACCEPTED MANUSCRIPT
deposited layers was characterized using a scanning electron microscope (JEOL JSM 820) coupled with an energy dispersive spectrometer (EDS). The samples were uniformly sprayed onto carbon tape, a Pt coating was applied for 10 s onto the synthesized layers, and images were
RI PT
acquired at various magnifications. These images were processed and analyzed with ImageJ software (National Institute of Mental Health, Bethesda, Maryland, USA) to gather information on mean pore length, porosity, and uniformity of the coated layers. Prior to ImageJ processing,
SC
the scale was set to the scale bar recorded in the image to obtain lengths in micrometers instead of pixels. The specimen thicknesses were measured using ultrasonic measurement (coating
M AN U
thickness gauge OMEGA instrument, OM179-745) with a precision of 1 µm. The surface roughness was determined using a Mitutoyo Surftest SV−402 profilometer. The mean roughness (Ra) of 7-10 different tests was calculated for each sample. The travelling distance and the speed of the stylus were 1.775 mm and 0.1 mm s-1, respectively.
TE D
The properties of wettability and surface free energy of the treated samples were measured with de-ionized water contact angle measurements using a contact angle meter (GBX, Digidrop, France). De-ionized water was automatically dropped (drop diameter: 6 µm) onto the sample.
EP
Measurement was done after 1 s and 10 s from the application of the water droplet. Three specimens for each sample were selected and three different locations on each were recorded. All
2.4.
AC C
measurements were carried out at room temperature.
Electrochemical corrosion test
Prior to the polarization test, the samples were immersed in 1000 mL SBF for 20 min to establish a relatively stable open circuit potential. The electrochemical behavior of both the 8
ACCEPTED MANUSCRIPT
uncoated and coated samples was determined in standard simulated body fluids (SBF) at pH 7.4 using the potentiodynamic polarization test (263A, EG&G PAR, USA). Briefly, a three-electrode cell with the sample as the working electrode, Ag/AgCl/1 M KCl (Satd.) as the reference
RI PT
electrode, and platinum as the counter electrode was used. The area of the working electrode (control and deposited layer samples) exposed to the solution was 1 cm2. A 1 mVs-1 scanning rate was applied during the potentiodynamic polarization test. The changes in the free corrosion
SC
potential (Ecorr) were monitored as a function of time. The temperature was maintained at 37 ±0.5ºC during the test. The results of the corrosion tests performed on the samples were
2.5.
Mechanical properties test
M AN U
calculated by extrapolating the polarization curve according to ASTM-G102-89.
TE D
The bending strength of the samples was evaluated as the maximum point of the stressdisplacement curve using a three-point bending test, which was carried out in a standard laboratory atmosphere with a mechanical tester (Computerized Instron Universal Testing
EP
Instrument model 4206). The tests were conducted at an across-head speed of 1 mm/min and
AC C
with a span length of 25 mm. Samples were prepared with dimensions of 33×5×5 mm, which is the required size for testing. At least 5 specimens were tested for each sample. Average values and standard deviation were calculated. 2.6.
Immersion tests
An immersion test was performed in a standard simulated body fluid (Hank’s balanced salts), with 5 specimens for each sample being immersed in 100 mL of solution. The electrolyte 9
ACCEPTED MANUSCRIPT
was saturated with atmospheric oxygen without stirring during the experiments. The temperature was maintained at 37.5 ±0.2ºC during the test. The immersed samples were extruded at different interval points of time, gently rinsed with distilled water, and dried at room temperature to avoid
RI PT
cracking. Before weight loss measurement, all samples were immersed in chromic acid solution (200 g/L Cr2O3 + 10 g/L AgNO3) for 10 minutes to remove the residual coated layer and corrosion products. Samples were then rinsed in distilled water, then acetone, and then dried. The
SC
surface appearance of the coated and uncoated immersed samples before and after removing the
2.7.
M AN U
residual coatings and corrosion products were observed by SEM.
Osteoblast Cell Culture and MTT Assay
The MTT assay was performed to measure cell viability in the presence of pristine and
TE D
composite samples. MC3T3 cell lines were used and were cultured in Dulbecco’s modified Eagle’s medium (DMEM) (Gibco Co, USA), supplemented with 10% fetal bovine serum (FBS)
EP
in a humidified incubator with 5% CO2 and 95% relative humidity at 37℃. The test for the coated samples and uncoated substrates was carried out by using an indirect extraction method,
AC C
where the immersion extracts collected from the immersion test were used for culturing cells. The extraction media was used at full concentration after one day of incubation in a humidified atmosphere with 5% CO2 at 37℃. Cell morphology and cell proliferation were observed under optical microscopy (LEICA Microsystems EZ4D). To assess cell proliferation, an MTT assay was carried out as a function of incubation time (1 -3 days) by calculating the RGR percentage (RGR = optical density for coated specimen/optical density for negative control group × 100%). 10
ACCEPTED MANUSCRIPT
Cell viability was obtained by comparing the absorbance of cells cultured with different samples to that of control wells containing only cells. The results were expressed as the means ± standard error of the mean. All the quantitative data were statistically analyzed to be expressed as mean ±
RI PT
standard deviation. Statistical analysis was determined by single factor ANOVA. P-values of less
SC
than 0.05 were considered significant.
3. Results and discussion
M AN U
In a dip-coating process, a substrate is dipped into a liquid coating solution, left for a reasonable time to allow wettability, and then withdrawn from the solution at a controlled speed and allowed to dry. Coating thickness generally increases with faster withdrawal speed. The thickness is controlled by the balance of forces at the stagnation point on the liquid surface. A
TE D
faster withdrawal speed pulls more fluid up onto the surface of the substrate before it has time to flow back down into the solution. The thickness is primarily affected by fluid viscosity, fluid density, temperature, and surface tension. Besides the possibility of increasing thickness and
EP
controlling pore size when using high volatile solvents, this method allows for the simultaneous coating of both sides of a sheet as well as the coating of non-flat substrates. In this way dip-
AC C
coating provides a level of adhesion needed in the process of organic coating of metal. This method, while excellent for producing high-quality coatings, requires precise control and a clean environment. In general, dip-coating is a reasonable coating technique to use whenever the substrate size, weight, or geometry makes spin coating difficult or impossible [12, 21]. For these reasons, many studies in chemical and nanomaterials engineering in both academia and industry make use of the dip-coating process. 11
ACCEPTED MANUSCRIPT
3.1. Surface properties of the coated samples
RI PT
The surface properties of the as-synthesized protective coatings (pristine PCL and nHAp/PCL) were described on the basis of surface morphology, surface roughness, wettability, and the layer thickness of the surface film deposited on the magnesium bars. The surface
SC
morphology observed by scanning electron microscope of the coated samples is shown in Fig. 2. The SEM images show that the plain PCL film produced fairly interconnected pore networks at
M AN U
both the concentrations used (5% and 10%). The surface microstructure modifications may have an intrinsic influence on biological functions such as cell attachment and ingrowth and, ultimately, tissue formation when the implant comes into contact with tissues at early stages [22]. The 3-D pore structure had a relatively uniform distribution and was free of defects on the
TE D
primary coating layer (the closest layer on the substrate), which is the layer responsible for providing adhesion between the organic coating and metal substrate [23]. The size of the pore phase at 5 and 10 wt% plain PCL was approximately 10 ±2 µm and 1.5 ±0.72 µm, respectively.
EP
These features are typical of porous polymer prepared by polymer-poor phase and polymer-rich phase separation using highly volatile solvents such as dichloromethane in the presence of
AC C
humidity. With the incorporation of nHAp (20 wt%) in polymer (5 and 10 wt%) matrices, the pore structure size decreased, especially at 10/90 PCL/DCM, but the macropore structure was still clearly observed at 5/95 PCL/DCM (Fig.2c). Although the HAp particles were embedded inside the matrix, the pore walls of the polymer were perfectly smooth and no HAp was observed on the polymer pore surfaces. The good interfacial bonding strength of nHAp with the polymer matrix can be explained by the hydrogen bonding formed between the -OH groups of HA and the 12
ACCEPTED MANUSCRIPT
=O sites of the PCL polymer [20]. The major peaks of the deposited composite layer on magnesium substrate are shown in the spectrum (Fig. 2). From the EDX peaks, it can be
RI PT
concluded that the deposited film is relatively free of impurities.
Nano-scaled composite coatings on magnesium bars may offer a promising therapy for natural bone regeneration since the surface characteristics, such as surface roughness, have a
SC
direct impact on the formation of bone cells. The surface roughness of the composite coatings may show an appropriate value. The composite film which contained nHAp in a PCL film (5
M AN U
wt%) showed a surface roughness of about 0.543 µm, whereas the plain PCL was 0.325 µm (Table 1). It has already been reported that rougher surfaces can adsorb more fibronectin and improve cell attachment over smoother surfaces [24]. Collectively, our results showed that the surface roughness of magnesium coatings increases with an increase in HAp content within the
TE D
matrix, as shown in Table 1.
PCL is a biodegradable and biocompatible polymer with a slow degradation rate which
EP
shows promise as a magnesium coating in bone tissue-engineering applications [25-27]. The surface of PCL is hydrophobic and does not exhibit any physiological activity, which makes it
AC C
unfavorable for cell growth when it comes into contact with the living body [28]. Therefore, the cytocompatibility of these synthesized materials should be improved before they can be used in clinical settings. Variation in the surface topologies as well as material compositions can be used to control the wetting process or the adsorption performance of PCL. The existence of hydrophilic domains on a material’s surface will have great influence on surface properties. It is well known that the water contact angle (WCA) reflects superficial properties of a material. The 13
ACCEPTED MANUSCRIPT
WCA of plain PCL and nHAp/PCL deposited layers on magnesium substrates are listed in Table 1. The nHAp/PCL composite membranes had lower contact angles than that of pure PCL membranes. It was observed that the WCA gradually decreased with increasing amounts of
RI PT
nHAp in the PCL layers deposited on Mg bars (inset of Fig. 2a, c-d). The WCA was decreased from 77.6ο (plain PCL) to 71.3ο (composite layer), as shown in Table 1. This phenomenon might be because of the wettable nHAp particles incorporated into the membranes, thus making the
SC
composite membranes more hydrophilic than pristine PCL. Another possible explanation is the increase in surface energy caused by the surface roughness of nHAp on the PCL surface. In
M AN U
conclusion, after nHAp incorporation, the surface property of plain PCL was altered. The introduction of nHAp provides an opportunity to further modify the surface biocompatibility. As mentioned previously, layer thickness can be controlled by various parameters. Our results showed that the polymer layers exhibited different thicknesses under the same dip-coating
TE D
conditions, as shown in Table.1. The composite films were much thicker than the corresponding plain films. The average thickness of pristine PCL/DCM (5/95) was 12.5 µm while that of
EP
PCL/nHAp film was 15.3 µm, Table.1.
AC C
3.2. Electrochemical corrosion analysis Fig. 3 illustrates the typical overall potentiodynamic behaviors of magnesium bars and dipcoated magnesium samples tested in a standard simulated body fluid at 37οC. The analysis data of these curves are shown in Table 2. The corrosion potentials (Ecorr) of the coated samples with 5 and 10 wt% plain PCL and 5 wt% PCL/nHAp were 200, 631, and 130 mV more positive, respectively, than an unprotected magnesium bar. At the same time, the values of corrosion 14
ACCEPTED MANUSCRIPT
current (Icorr) of protected samples were lower than those of uncoated samples. Therefore, values for both Icorr and Ecorr illustrated that plain PCL and nHAp/PCL composite would be able to
3.3.
RI PT
enhance the corrosion resistance and thereby control the degradation rate of magnesium alloy.
Immersion tests
SC
Fig. 4 shows the representative SEM images for the uncoated and coated samples before
M AN U
removing the corrosion products and after immersion in SBF at 37οC for 20 days. It is clear from the SEM images that the samples coated with plain PCL experienced cleavage and a high level of attack by water molecules. Water diffusion through tiny micro or nano defects on the coating layer deteriorated the coatings and destroyed the adhesion of the organic layer on the substrate. Earlier reports [29] have indicated that when the water molecules come into contact with
TE D
magnesium surfaces, hydrogen gas is produced by the electrochemical reaction (Eq. 1). However, the porous layer of the pure polymer fabricated via dip-coating cannot provide a fully porous structure mainly due to the organic layer, which contains three films: the primary,
EP
intermediate, and top-coating films. The primary film is responsible for providing good adhesion
AC C
between the whole organic layer and the metal substrate. Hence, the primary film is a dense, non-porous structure. As a result, the production of hydrogen is hindered and the osmotic pressure is increased. The tiny micro pores or/and in the range of nanometers pores contained in the primer film are not adequate for the speed of hydrogen release required from the magnesium substrate, i.e. hydrogen gas cannot be removed at an adequate velocity. The hydrogen that is produced results in an increase in gas pressure beneath the polymer. The hydrogen permeability of the polymer is insufficient to release the gas at the rate it is beging produced, which leads to 15
ACCEPTED MANUSCRIPT
excess pressure that results in the bursting of the polymer film. However, the reaction itself is initiated by failures in the polymer coating, Fig. 4a. While composite coatings may contain fully porous structures, could establishe space between the nano-particles and the polymer matrix
RI PT
sufficient to allow evacuation of hydrogen gas via the polymer matrix layer. To inform our hypothesis, we found that the composite coating layer was still intact and resisted rupturing and/or delamination from their substrates as shown in Fig. 4ab-d. From these findings, it is
M AN U
stability of the coating layers on magnesium substrates.
SC
assumed that the incorporation of nHAp into the polymer matrix could provide a sufficient
Mg(s)+ 2H2O→Mg2++ 2OH-(aq)+H2(g) ↑ (overall reaction)
(1)
PCL surfaces are inherently hydrophilic, which leads to a decrease in the stimulation and
TE D
precipitation of apatite on their surfaces. However, it is known that polymer surfaces gradually change towards a more hydrophilic structure in polar aqueous environments. In a comparison of coated samples with pristine PCL, the nHAp/PCL coatings showed a uniform distribution of
EP
precipitated Ca/P particles of a uniform size of around 230 nm. The overall pore structure of the composite layer was remarkably unchanged and free of defects, and the characteristic
AC C
distribution of macropores in the microporous matrix was still observed, Fig. 4. It is speculated that over a longer incubation time, the precipitation of Ca-P increases, resulting in a uniformly dense layer of Ca-P on the polymer layer, and that this amorphous inorganic layer finally crystallizes into apatite. In support of our hypothesis, we found that some areas (shown in the dashed box in Fig. 4d) show a dense coating with a characteristic cauliflower structure typically associated with apatite. Again, if a dense CaP layer forms, this layer will be able to add further 16
ACCEPTED MANUSCRIPT
protection to the magnesium substrates further the high formation of bone-like. This finding confirmed that the introduction of hydroxyapatite into the PCL matrix improved the bioactivity of the scaffold. These results indicate that composite films stimulate CaP compound in a regular
RI PT
phase more than plain PCL does. A reasonable explanation for this phenomenon could be that there is a lower release of Mg+2 from the polymer matrix of composite-film-coated magnesium samples during the whole immersion period in comparison with samples coated with a plain PCL
SC
film. The slight release of magnesium ions could lead to a slightly increased pH, which in turn could assist in the nucleation of CaP in a uniform manner [30]. Another possible reason for the
M AN U
better protection seen with polymer matrix composite film-coated magnesium samples is the increase in pH, leading to the dissolution of ions from nHAp particles [31], as shown in Fig. 5. It should be mentioned that increased degradation of magnesium substrate without sufficient apatite formation can lead to deleterious results. Several reports have suggested that reduction of
TE D
the degradation rate of magnesium implants lowers the accumulation of hydrogen gas bubbles and halts the lowering of the local pH, which could otherwise lead to necrosis of tissues,
EP
infections, and alkaline poisoning [8, 29, 32].
SEM analysis showed remarkable changes in the morphology of the composite surface
AC C
during the degradation period in SBF, as shown in Fig. 6. After removing the corrosion products from the magnesium surface, it was apparent that the uncoated samples experienced a large amount of severe localized corrosion as well as pitting corrosion on the whole magnesium surface as shown in Fig. 6a. Continued localized corrosion at the valuable sites gives rise to nonuniformly corroded sites [33]. The corrosion morphology of the samples protected with composite films showed uniform corrosion on the magnesium surfaces, Fig. 6b-d. Quantitatively, 17
ACCEPTED MANUSCRIPT
the results of the immersion experiment clearly showed that PCL/nHAp composite films have better stability than those of plain PCL films, a property which could contribute to the slower degradation rate of the composite films compared with plain PCL films during the length of the
RI PT
incubation.
Fig.7 shows a graph of the total mass of magnesium dissolved into the SBF solution as a
SC
function of immersion time in the bath. In comparison with uncoated samples, the plain PCL and PCL matrix composite coatings decreased the corrosion rate of the magnesium alloy by factors
M AN U
of 3 and 4, respectively. These degradation rates of both uncoated and coated samples are in good agreement with the above pre-mentioned section. It is clear that both the coated samples showed almost no weight loss at the initial immersion time, especially the samples coated with composite films. This observation confirms that these coatings can provide a sufficient
several days [34, 35].
In vitro loss of mechanical integrity
EP
3.4.
TE D
protection at the first implantation time and avoid the inflammation stage which can last for
AC C
A material used for bone implant application should have a mechanical integrity sufficient for bone cells to attach, proliferate, and differentiate in a manner similar to native ECM until the healing process is complete. Hence, mechanical properties are an important factor for bioimplant materials. Fig. 8 illustrates the influence of in vitro degradation over 20 days on the bending strength of uncoated and nHAp/PCL-coated magnesium samples. From the selected stressdisplacement curves it can be seen that the samples protected with composite film showed better 18
ACCEPTED MANUSCRIPT
mechanical properties than uncoated ones. Furthermore, the tensile strength of the coated and uncoated samples decreased by around 13 and 34%, respectively, in comparison to preimmersion magnesium alloy (i.e. time point 0). This difference is likely due to surface defects
RI PT
such as severe localized corrosion and pitting corrosion formed during degradation. Conclusively, the mechanical degradation of magnesium-based polymer nanocomposites coating
Cell culture
M AN U
3.5.
SC
was significantly decreased by the incorporation of nano inorganic particles.
To avoid undesired effects when materials are implanted in vivo, candidate materials have to be assayed for their compatibility using cell cultures [36]. This step is essential to ensure that the optimal material for a specific biomedical application is used in patients. As mentioned
TE D
above, biocompatibility, cell attachment, migration, proliferation, and differentiation are significantly influenced by surface conditions and chemical compositions of the substrates [37]. Use of a biocompatible material such as HAp [17] in combination with a biodegradable synthetic
EP
polymer gives a ‘bioartificial advanced composite’ with enhanced biocompatibility and other biological functions. Furthermore, the ionic dissolution from the inorganic compounds,
AC C
particularly in the case of nHAp (Fig. 5), plays an essential role in such cellular processes as bone cell differentiation and the genes expression. Based on this consideration, MTT assay data showed that the incorporation of nHAp within the polymer matrix could improve the cell viability and survival after a 24 h incubation in comparison with plain PCL, Fig. 9. After a longer incubation time of three days, it was clear that the composite coating on the magnesium substrate resulted in more proliferation and continuous cell growth. These results confirmed that nHAp 19
ACCEPTED MANUSCRIPT
plays an important role in supporting cell proliferation and bone cell differentiation. In contrast, uncoated samples showed worse behavior in comparison with both of the coated samples. We predicted that the number of cells would be statistically increased for the coated Mg at all time
RI PT
intervals compared with uncoated samples, meaning that all of the coated samples have a significantly better initial biocompatibility than the uncoated sample. To verify this hypothesis, the respective optical microscopy morphologies of osteoblasts of the MC3T3 cell line after one
SC
day of culture in extracts from the immersion test were observed, Fig. 10. From these images it can be seen that the cell morphologies of cells treated with extracts taken from coated samples
M AN U
were normal and healthy and showed good survival behavior, especially the samples with the composite coating, while no cells were observed in cultures grown with extracts of the uncoated samples, indicating that uncoated Mg does not promote cell growth. However, it has been reported that Mg is inherently biocompatible [29] and that the aluminum content in the AM50
TE D
alloy does not have negative effects on the viability and proliferation of osteoblasts, as reported by Ciapetti et. al [36]. This observation may be due to the steep increase in pH or the release of ions into the medium from the uncoated samples [12, 38]. Based on the results obtained from the
EP
in vitro tests, it is believed that the incorporation of nHAp particles into the polymer matrix improves the biocompatibility and enhances the development of the connection between the
AC C
magnesium implants and natural bone. Our results are in agreement with all studies considered to have shown improved biocompatibility and higher expression of osteogenic markers when HAp is incorporated into the polymers [39]. This behavior may be partially explained by improved adsorption of extracellular matrix (ECM) proteins, such as fibronectin or vibronectin, to composites as compared with plain polymers. Hence, the finding of good compatibility in vitro
20
ACCEPTED MANUSCRIPT
makes nHAp/PCL advanced composite coating of magnesium worthy of further studies for
RI PT
development of surgical implants.
4. Conclusion
In summary, novel nHAp/PCL advanced composite film coatings on magnesium substrates
SC
were developed using a simple and cost effective dip-coating technique. In addition, various physicochemical properties such as surface properties and biocorrosion performance were
M AN U
investigated. The matrices were evaluated for in vitro cytocompatibility under simulated physiological conditions using an MTT assay. The surface modifications in the magnesium substrate induced promising changes in such characteristics as surface roughness and wettability. Our results demonstrated that the incorporation of nHAp particles into the polymer matrix layer
TE D
improved both the degradation resistance and biological function of the magnesium substrates. The bone implant showed good cytocompatibility to better mimic the surface characteristics of natural bone. In addition, it is speculated that the level of osteoblastic differentiation activity will
EP
also increase significantly with the incorporation of nHAp into the PCL polymer matrix composite coatings on magnesium implants. Further in-depth studies using an in vivo animal
AC C
model will be required to evaluate the degradation performance and biological functions of assynthesized hybrid polymer matrix nanocomposite coated on magnesium based implants.
Acknowledgments
21
ACCEPTED MANUSCRIPT
The authors gratefully acknowledge support for this research by the Ministry of Education, Science Technology (MEST) and National Research Foundation of Korea (NRF) through the
RI PT
Human Resource Training Project for Regional Innovation. References
[1] X. Zhang, G. Yuan, L. Mao, J. Niu, P. Fu, W. Ding, Journal of the Mechanical Behavior of
SC
Biomedical Materials. [2] J.A. Cowan, BioMetals, 15 (2002) 225-235.
M AN U
[3] M.E. Maguire, J.A. Cowan, BioMetals, 15 (2002) 203-210.
[4] M.P. Staiger, A.M. Pietak, J. Huadmai, G. Dias, Biomaterials, 27 (2006) 1728-1734. [5] T. Jinno, V.M. Goldberg, D. Davy, S. Stevenson, Journal of Biomedical Materials Research, 42 (1998) 20-29.
TE D
[6] R. Van Noort, Journal of Materials Science, 22 (1987) 3801-3811. [7] W. TIANSHI, Z. RENJI, A.Y. YONGNIAN, Journal of BIOACTIVE AND COMPATIBLE POLYMERS, 24 (2009) 169-182.
EP
[8] S. Zhang, X. Zhang, C. Zhao, J. Li, Y. Song, C. Xie, H. Tao, Y. Zhang, Y. He, Y. Jiang, Y. Bian, Acta Biomaterialia, 6 (2010) 626-640.
AC C
[9] X.N. Gu, Y.F. Zheng, Q.X. Lan, Y. Cheng, Z.X. Zhang, T.F. Xi, D.Y. Zhang4, Biomedical Materials, 4 (2009) 044109.
[10] J. Li, P. Cao, X. Zhang, S. Zhang, Y. He, Journal of Materials Science, 45 (2010) 60386045.
[11] K. Bai, Y. Zhang, Z. Fu, C. Zhang, X. Cui, E. Meng, S. Guan, J. Hu, Materials Letters, 73 (2012) 59-61. 22
ACCEPTED MANUSCRIPT
[12] L. Xu, A. Yamamoto, Colloids and Surfaces B: Biointerfaces. [13] J.K. Pandey, K. Raghunatha Reddy, A. Pratheep Kumar, R.P. Singh, Polymer Degradation and Stability, 88 (2005) 234-250.
RI PT
[14] W.-E. Teo, S. Ramakrishna, Composites Science and Technology, 69 (2009) 1804-1817. [15] M. M, N.-E. M, International Journal of Nanomedicine, 6 (2011) 1651-1659
372.
SC
[16] K.A. Khalil, S.W. Kim, H.Y. Kim, Materials Science and Engineering: A, 456 (2007) 368-
[17] H.-H. Lee, H.-S. Yu, J.-H. Jang, H.-W. Kim, Acta Biomaterialia, 4 (2008) 622-629.
M AN U
[18] H. Zhou, J. Lee, Acta Biomaterialia, (2011).
[19] B. Nasiri-Tabrizi, P. Honarmandi, R. Ebrahimi-Kahrizsangi, P. Honarmandi, Materials Letters, 63 (2009) 543-546.
[20] S.Z. Fu, X.H. Wang, G. Guo, S. Shi, M. Fan, H. Liang, F. Luo, Z.Y. Qian, Journal of
TE D
Biomedical Materials Research Part B: Applied Biomaterials, 97B (2011) 74-83. [21] T.F. Conceicao, N. Scharnagl, C. Blawert, W. Dietzel, K.U. Kainer, Corrosion Science, 52 (2010) 2066-2079.
EP
[22] Y. Zhu, C. Gao, J. Shen, Biomaterials, 23 (2002) 4889-4895. [23] P. Sørensen, S. Kiil, K. Dam-Johansen, C. Weinell, Journal of Coatings Technology and
AC C
Research, 6 (2009) 135-176.
[24] K.M. Woo, J. Seo, R. Zhang, P.X. Ma, Biomaterials, 28 (2007) 2622-2630. [25] E.-J. Choi, J.-K. Park, Polymer Degradation and Stability, 52 (1996) 321-326. [26] D. Darwis, H. Mitomo, F. Yoshii, Polymer Degradation and Stability, 65 (1999) 279-285. [27] X. Zheng, C.A. Wilkie, Polymer Degradation and Stability, 82 (2003) 441-450.
23
ACCEPTED MANUSCRIPT
[28] Q. Zhao, J. Tao, R.C.M. Yam, A.C.K. Mok, R.K.Y. Li, C. Song, Polymer Degradation and Stability, 93 (2008) 1571-1576. [29] B. Zberg, P.J. Uggowitzer, J.F. Loffler, Nat Mater, 8 (2009) 887-891.
RI PT
[30] L. Li, J. Gao, Y. Wang, Surface and Coatings Technology, 185 (2004) 92-98.
[31] Y. Ito, H. Hasuda, M. Kamitakahara, C. Ohtsuki, M. Tanihara, I.-K. Kang, O.H. Kwon, Journal of Bioscience and Bioengineering, 100 (2005) 43-49.
SC
[32] J.-H. Jo, J.-Y. Hong, K.-S. Shin, H.-E. Kim, Y.-H. Koh, Journal of Biomaterials Applications, 5 (2011) 8.
M AN U
[33] Y. Xin, T. Hu, P.K. Chu, Acta Biomaterialia, 7 (2011) 1452-1459.
[34] J. Yang, F. Cui, I. Lee, Annals of Biomedical Engineering, 39 (2011) 1857-1871. [35] S. Chen, S. Guan, S. Hou, L. Wang, S. Zhu, J. Wang, W. Li, Surface and Interface Analysis, 43 (2011) 1575-1580.
TE D
[36] G. Ciapetti, J.L. González-Carrasco, L. Savarino, M.A. Montealegre, S. Pagani, N. Baldini, Biomaterials, 26 (2005) 849-859.
[37] F. Yang, R. Murugan, S. Ramakrishna, X. Wang, Y.X. Ma, S. Wang, Biomaterials, 25
EP
(2004) 1891-1900.
[38] F. Witte, F. Feyerabend, P. Maier, J. Fischer, M. Störmer, C. Blawert, W. Dietzel, N. Hort,
AC C
Biomaterials, 28 (2007) 2163-2174.
[39] L. Shor, S. Güçeri, X. Wen, M. Gandhi, W. Sun, Biomaterials, 28 (2007) 5291-5297.
Figure captions and Tables
Fig. 1. TEM morphology of synthesized HAp nano-scale 24
ACCEPTED MANUSCRIPT
Fig. 2. Surface appearance in SEM of magnesium samples coated with, 5 and 10 wt% plain PCL (a and b) and nHAp at 5 and 10% by wt, PCL (c and d); SEM-EDX of the composite
RI PT
coated samples. Insets are the water contact angles Fig. 3. Potentiodynamic polarization curves for coated and uncoated samples
Fig. 4. Surface appearance in SEM after immersion in SBF for 20 days for the samples coated
SC
with (a) plain PCL, (b-d) composite coating layers. Parts c and d are respective circular
M AN U
areas in b.
Fig. 5. Schematic diagram illustrating the apatite-like nucleation and cell behavior on magnesium substrate coated with composite films
Fig. 6. Scanning electron micrographs of the uncoated (a) and composite film-coated (b) samples
TE D
after removing the corrosion products. Parts c and d are respective circular areas in b. Fig. 7. Total weight lost by magnesium alloys uncoated or coated with plain PCL and
EP
PCL/nHAp composite coated samples Fig. 8. Selected stress-displacement curves
AC C
Fig. 9. MTT cytotoxicity assay of the coated and uncoated magnesium samples after 1 and 3 days of incubation. The data are reported as the mean ± standard deviation (n=3 and p‹0.05)
25
ACCEPTED MANUSCRIPT
Fig. 10. Optical micrographs of osteoblast cell (MC3T3) morphology after one day of culture, including (a) uncoated bars of Mg, (b) bars coated with pristine PC, and (c) PCL/nHAp
AC C
EP
TE D
M AN U
SC
RI PT
composite-coated Mg bars
26
ACCEPTED MANUSCRIPT
Table 1. Layer thickness, surface roughness, and water contact angle of coated samples
Layer thickness (µm)
12.5±1.564
0.325±0.09
nHAp in 5% wt PCL
15.3±2.21
0.543±0.057
nHAp in 10% wt PCL
34.12±2.45
0.952±0.066
M AN U
*
10 S
77.6±0.55
77.9.45±0.45
75.7±2
73.3±2.15
71.3±1.23
70.5±1
SC
Plain PCL at 5% wt
1S
RI PT
Sample
Water contact angle ± STD (o)
Surface roughness*± STD (µm)
AC C
EP
TE D
The roughness values correspond to the average surface roughness (Ra)
27
ACCEPTED MANUSCRIPT
Table 2. Summary of corrosion potential voltage and current density values of uncoated and coated samples
5% wt PC
10% wt PCL
Ecorr (mV)
-1340
-1140
-709
-1210
Current density (µA cm-2)
103
7.723
1.721
3.6
SC M AN U TE D EP
AC C
Figure
28
PCL/nHAp
RI PT
Uncoated
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT