Excavating the Role of Nanofiller Aspect Ratio, Amorphic - American

Article pubs.acs.org/JPCC

Tailor-Made Fibrous Nanohydroxyapatite/Polydimethylsiloxane Composites: Excavating the Role of Nanofiller Aspect Ratio, Amorphicity, and Noncovalent Surface Interaction Nabarun Roy† and Anil K. Bhowmick*,†,‡ †

Indian Institute of Technology, Kharagpur-721302, India Department of Chemistry, Indian Institute of Technology, Patna-800013, India

‡

S Supporting Information *

ABSTRACT: This work holds significance in designing biomimetic implants where biocompatibility and strength are two critical concerns. In the work mentioned herein, hydroxyapatite nanofibers of varying aspect ratios have been synthesized by the soft template assisted method. These nanoparticulates have been successfully introduced into the polydimethylsiloxane (PDMS) matrix through a solution casting technique, and this is probably the first successful attempt made. Nanocomposites prepared with longer nanofibers (300−500 nm length) showed more improvement in various physicomechanical and thermal properties compared with those prepared with lower aspect ratio nanofibers (10.3 versus 4.5). In-depth analysis revealed that a combined effect of several factors leads to these differences. Besides the most important factor, viz., aspect ratio of the nanofibers, the amorphicity of nanofibers and adherence of polymer on the nanofiller surface, i.e., noncovalent surface modification of nanofiller and restriction of microcrystalline domain formation by the nanofiller, also owe their responsibility.

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toughness.12,13 This amalgamates properties of both the constituents without sacrificing any of them in a much more efficient manner compared with the conventional methods. Hence, scientists endeavor to mimic bones and joints by generating hybrid materials with biodegradable and biocompatible polymers and integrate hydroxyapatite nanoparticles in them.14−20 Silicone rubber, owing to its chemical inertness, good flexibility, and excellent biocompatibility, has been used in constructing body part implants and other medical devices since the 1960s (first successful breast implant).21,22 Nowadays, arthritic knees and joints are effectively replaced by silicone rubber prosthesis. However, the efficiency of the implants is a function of weight-bearing load and often ends up in implant failure. The sole reason for this is poor mechanical properties of silicone rubber. Carbonaceous fillers, although they impart very

INTRODUCTION A recent trend in the field of polymer science and nanotechnology involves crafting elegant hybrids suitable for smart applications such as drug delivery,1−3 biomimetic implants,4−6 etc. The unconventional properties of biocompatibility and biodegradability of various novel polymers foster the researchers of this field to generate materials which can readily impersonate natural bones and muscles through biomedical improvisations. Hydroxyapatite, Ca10(PO4)6(OH)2, is a natural constituent of human bone where it possesses a nanometric lamellar morphology intertwined with collagen fibers.7−9 The inorganic constituent is basically responsible for imparting stiffness to the bone. In recent years, hydroxyapatite has widely been utilized in the field of tissue engineering owing to its outstanding biocompatibility and incredible mechanical properties.10,11 It has been observed that applicability of various biocompatible polymers in biomimetic implants is restricted due to their poor mechanical properties. Hence, nanoparticulates are integrated into the polymer matrices to improve their mechanical properties such as tensile strength and © 2012 American Chemical Society

Received: November 10, 2011 Revised: March 10, 2012 Published: March 28, 2012 8763

dx.doi.org/10.1021/jp210835a | J. Phys. Chem. C 2012, 116, 8763−8772

The Journal of Physical Chemistry C

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good strength to the silicone matrix, are cytotoxic in nature.23,24 On the other hand, hydroxyapatite (HA) is biocompatible as well as possesses high modulus and might serve as a superior candidate in implant formulation. An important facet of property improvement in polymer nanocomposites is the aspect ratio of the nanofiller used. The nanofillers with greater aspect ratio provide more sites for physical cross-linking and thereby may result in better polymer−filler interaction. Though a literature survey yields innumerous reports on synthesis of hydroxyapatite nanorods of different aspect ratios,25−31 utilization of these in improving properties of polymer matrices is rather scanty.32−34 There are only very few reports based on spherical hydroxyapatite (HA)/ polydimethylsiloxane (PDMS) nanocomposites.35,36 However, lack of any literature warrants a detailed and systematic study on HA nanofibers/PDMS nanocomposites. The present study is basically comprised of two parts. In the first part, soft template induced synthesis of the HA nanofibers was executed. Although the method of preparation is well established in the literature,25−31 it was manipulated in terms of nature of reagents and reaction conditions. Two different sets of variations were introduced in the reaction scheme (designated as method A and method B), and this procedure resulted in synthesis of nanofibers of varying mean aspect ratios. The nanofibers prepared were characterized extensively. Morphological analysis was done by High-Resolution Transmission Electron Microscopy (HRTEM) and Field Emission Scanning Electron Microscopy (FESEM) analysis. Chemical composition was examined through Energy-Dispersive X-ray Spectroscopy (EDX) and Fourier Transform Infrared Spectroscopy (FTIR) studies. The crystallinity of the nanoparticles was determined through Selected-Area Electron Diffraction (SAED) and Wide-Angle X-ray Diffraction (WAXD) analysis in the powdered form. In addition, the effect of calcination of the nanoparticles on morphology and crystallinity was studied and discussed in depth. The second part of the study involves successful incorporation of the nanofiller into the PDMS matrix. The thermodynamic feasibility of the nanocomposite formation in terms of free energy change of the system was determined by thorough Attenuated Total Reflection−Fourier Transform Infrared Spectroscopy (ATR−FTIR) studies. This was followed by a systematic structure−property study of the nanocomposites prepared with different forms of HA nanofibers. This involved measurements of mechanical, dynamic mechanical, and thermal properties. The difference in dispersion which governs the nanocomposite properties was explained through morphological analyses by HRTEM and FESEM studies. This paper highlights several factors which affect property improvement in nanocomposites. Apart from the prominent role of aspect ratio of the nanofiller, its amorphicity is expected to show its effect in polymer reinforcement. In addition, noncovalent surface modification of the nanofiller has been intentionally carried out to facilitate better dispersion of the otherwise difficult to disperse hydroxyapatite. However, the presence of nanofillers may restrict the process of microcrystalline domain formation in the polymer. This phenomenon will also affect the properties of the nanocomposites. Such a comprehensive investigation of various factors has not been reported on such a system to the best knowledge of the authors.

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EXPERIMENTAL SECTION

Materials. Calcium nitrate tetrahydrate (Ca(NO3)2·4H2O), cetrimide (CH3(CH2)13N+(CH3)3 Br−), and diammoniumhydrogenphosphate ((NH4)2HPO4) were procured from Merck, India. Acetic acid and polyethylene glycol (PEG 400) were supplied by Merck, India, and used as supplied. Octamethylcyclotetrasiloxane [(CH3)2SiO]4 (D4), required for the synthesis of PDMS, has a boiling point of 175 °C, viscosity of 1.396, density of 0.955, and purity >99% (GC) and was obtained from Momentive Performance Materials, Bangalore, India. D4 was freshly distilled before use. 1,1,3,3-Tetramethyl-1,3-divinyldisiloxane (purity 97%) with a boiling point of 139 °C and density of 0.809, which was used for functionalization of the PDMS chains with vinyl end group, was supplied by Sigma-Aldrich. A curing system consisted of platinum catalyst (Pt catalyst in U10, where U-10 is a vinyl PDMS system with molecular weight 74 400 and viscosity 10 Pa·s having a hydride content of 0.05 mmol/g) and the hydride cross-linker polymethylhydrogenosiloxane (V430) with the chemical formula Me 3 Si(OSiMe2)x(OSiMeH)yOSiMe3, where x and y are 10, having a hydride content of 4.3 mmol/g. These were supplied by Momentive Performance Materials, Bangalore, India. Potassium hydroxide used as the initiator in the anionic ring-opening polymerization of D4 was purchased from Merck, Mumbai, India. Synthesis of Hydroxyapatite (HA) Nanofibers. HA nanofibers were synthesized using the soft template assisted method using (NH4)2HPO4 and Ca(NO3)2·4H2O. A template was generated using cetrimide and PEG 400. Nanofibers of varying aspect ratio and crystallinity were synthesized by varying pH of the reaction medium and the time for calcination of the as-prepared nanofibers. The uncalcined and calcined samples prepared by methods A (normal pH) and B (acidic pH) are designated as HA(U), HA(C) and HB(U), HB(C), respectively. Hydroxyapatite is synthesized from calcium nitrate and diammoniumhydrogenphosphate as follows 10Ca(NO3)2 + 6(NH4)2 HPO4 + 2H 2O → Ca10(PO4 )6 (OH)2 + 12NH4NO3 + 8HNO3

Details of both the methods are provided in the Supporting Information section S1. Synthesis of HA Nanofiber/PDMS Nanocomposites. Vinyl end-capped PDMS was synthesized by anionic ringopening polymerization of octamethylcyclotetrasiloxane (D4). Nanocomposites were prepared by a solution casting method. In this method, the synthesized vinyl-terminated PDMS was dissolved in 50 mL of toluene. In a separate beaker, 2 wt % HA was dispersed in 10 mL of toluene by stirring followed by ultrasonication. Different masterbatches were prepared with HA (uncalcined and calcined) as mentioned earlier. This nanofiller suspension was added to the homogeneous solution of the polymer and stirred thoroughly for 2 h followed by addition of curing agents (Pt catalyst and Si−H cross-linker). The resultant mixture was degassed and cast in a Teflon petri dish and left undisturbed overnight. The transparent sheet thus obtained upon solvent evaporation was subjected to vacuumdrying at 80 °C. The sample designation and the compositions are summarized in Table 1. The amount of polymer (13 g), catalyst (0.05 g), and Si−H cross-linker (0.32 g) is the same in all the samples. 8764



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glass transition temperature (Tg). The error in the measurement was ±1 °C Reinforcement of the nanofiller in the nanocomposite was evidenced by examining the mechanical properties of the samples. The tests were carried out with the tensile specimens punched out from the vacuum-dried samples with an ASTM Die-C on a Zwick UTM, model − Z010 (Zwick GmbH and Co., Ulm, Germany), at a cross-head speed of 500 mm/min at 25 °C. The thermal stability of the nanocomposites was determined by TGA analysis using a Perkin-Elmer Instrument, Diamond TG-DTA. The heating rate during the run was set at 20 °C/ min where around 5 mg of the sample was programmed to be heated under an air atmosphere up to 800 °C. The data were analyzed by Perkin-Elmer analysis software. The Tmax (temperature at which maximum degradation takes place) and Ti (onset temperature of degradation) were obtained from the PerkinElmer data analysis software. The error in the measurement was ±1 °C.

Table 1. Composition of HA Nanofiber/PDMS Nanocomposites with Sample Designation sample

filler (2 wt %)

VPHAU2 VPHAC2 VPHBU2 VPHBC2

HA(U) HA(C) HB(U) HB(C)

Characterization. Morphological studies of the nanofillers were executed through High-Resolution Transmission Electron Microscope (HRTEM) analysis. A minute quantity of the sample was added to deionized water and was dispersed by ultrasonication at 45 kHz for 1 h. A drop of this dispersion was added through a microliter syringe on a carbon-coated copper grid. Microscopy was performed with JEOL 2100, Japan. A transmission electron microscope was operated at an accelerating voltage of 200 kV. The morphology of the nanocomposites was determined through HRTEM analysis of the ultramicrotomed samples. This provided a sound knowledge regarding the state of dispersion of HA in the polymer matrix. Besides this, Field Emission Scanning Electron Microscopy (FESEM) was also pursued with an FESEM S4800 Hitachi microscope at an acceleration voltage 10.0 kV with a working distance of 10.1 mm. The polycrystalline nature of the prepared nanofibers was also examined through Wide-Angle X-ray Diffraction (WAXD) analysis. WAXD patterns of the HA were obtained through a Philips X-ray diffractometer (model PW-1710) with the aid of crystal monochromated Cu Kα radiation in the Bragg’s angle range of 2−40°. The operating voltage and current were, respectively, 40 kV and 20 mA. WAXD studies of the nanocomposite samples were also carried out to examine the effect of the nanofiller in the restriction of the crystallization process during nanocomposite formation. To authenticate the procedure of the nanofiller synthesis, the absorption bands in the FTIR spectra of the synthesized nanofillers were investigated. FTIR studies using a PerkinElmer FTIR spectrophotometer (model spectrum RX I), within a range of 400−4400 cm−1 using a resolution of 4 cm−1, were carried out by preparing KBr pellets of the powdered samples. An average of 16 scans was acquired for each sample. Surface analysis of the nanocomposites was done through Attenuated Total Reflection (ATR)−Fourier Transform Infrared Spectra (FTIR) studies. ATR−FTIR spectra of the samples were recorded through an infrared spectrophotometer (Nicolet Nexus, Madison, WI, USA) within a range of 650−4000 cm−1, taking a resolution of 4 cm−1. A ZnSe prism was used for ATR−FTIR spectroscopy measurements. An average of 120 scans was reported for each spectrum recorded. Various applications of PDMS require determination of the mechanical properties under dynamic conditions. Dynamic mechanical analysis of the nanocomposites with sample specification (12.6 mm × 6.65 mm × 1.2 mm) was carried out using a DMA of TA Instruments (model Q800). The conditions applied for testing the samples in a tensile mode were: a constant frequency of 1 Hz, a strain of 0.05%, and a temperature range from −125 to 50 °C at a heating rate of 3 °C/min. Storage modulus (E′) and loss tangent (tan δ) were measured as a function of temperature for all the samples under the same conditions. The temperature corresponding to the glass−rubber transition showed a prominent peak in tan δ versus temperature plot. This temperature was designated as

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RESULTS AND DISCUSSION Synthesis and Morphology of HA Nanofibers. The HA nanofibers were prepared using the soft template approach. The samples prepared at pH 7 resulted in longer nanofibers with an aspect ratio 10.3, while those prepared at pH 4 yielded shorter nanofibers (aspect ratio 4.5). The crystalline nature of the nanofibers was confirmed from SAED analysis, while elemental analysis was done through EDX measurement. Details of synthesis and HRTEM, SAED, and EDX analysis of the HA nanofibers are included in Supporting Information sections S1, S2.1, and S2.2.1 and Figures S1, S2, and S3, respectively. The HA nanofibers prepared by the above methods were used in the synthesis of PDMS nanocomposites. Synthesis and Characterization of Nanocomposites. Morphology of the Nanocomposites. The nanocomposites prepared with H(A) and H(B) showed marked difference in the state of dispersion of the filler in the polymer matrix. The highest degree of dispersion was attained in the case of the nanocomposites prepared with HA(U) nanofibers. This is evident from Figure 1(a). Here individual nanofibers are found to be dispersed in the polymer matrix.

Figure 1. Representative HRTEM images of (a) VPHAU2 and (b) VPHBU2.

Nanofiller dispersion was poor for the samples prepared with H(B). Figure 1(b) which represents the HRTEM micrograph of VPHBU2 speaks out clearly about the inferior quality of dispersion of the nanofibers. The small fibers undergo significant agglomeration as is evident from the figure. 8765



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VPHBU2 exhibited a D0.1 value of just 1.36%, it increased to 6.30% for VPHAU2. Nanocomposites prepared with calcined HA nanofibers resulted in deterioration in the extent of dispersion. The nanocomposites VPHAC2 and VPHBC2 showed, respectively, much lower D0.1 values of 2.09% and 1.23%. Hence, it is well-explained from this method of quantification that nanofiller distribution was better and more homogeneous for VPHAU2. Comparison of the XRD pattern of the as-prepared and calcined HA nanofibers prepared by methods A and B is depicted in Figure 4(a). In both cases, distinct and sharp peaks

FESEM analysis of the samples also revealed vividly the difference in dispersion of the nanofibers in various samples. While VPHAU2 (Figure 2(a)) showed much better dispersion

Figure 2. Representative FESEM images of (a) VPHAU2 and (b) VPHBU2.

with individual HA nanofibers well-dispersed in the polymer matrix, significant agglomeration was observable in the case of VPHBU2 (Figure 2(b)). It is also understandable from the FESEM images that HA nanofibers in VPHAU2 possess a higher aspect ratio compared with those in VPHBU2. A quantitative approach was adopted to differentiate the extent of dispersion of the HA nanofibers in various nanocomposites. A dispersion degree parameter D0.1 was employed to determine the extent of nanofiber dispersion in the nanocomposites.37−39 The parameter lies in the range of 0.9x̅ to 1.1x,̅ where x̅ is the mean spacing between the HA nanofibers. The higher the D0.1 value, the larger the amount of spacing data in proximity to the value of mean spacing. This is the consequence of homogeneous and better dispersion. Free path distribution follows a log-normal distribution model,37 and D0.1 is expressed as

Figure 4. WAXD profiles of (a) powdered HA nanofiber samples and (b) unfilled and HA nanofiber filled PDMS vulcanizates.

were observed though slight narrowing of the peaks that took place for the calcined samples. There is also a coherent increase in the peak intensity upon calcination which is consistent with the fact that complete crystal formation takes place upon calcination. All the WAXD profiles exhibit well-resolved diffraction peaks in the 2θ range of 10−40°. The reflections in the WAXD profiles of the nanofibers were well in accordance with JCPDS card 09-0432 (hexagonal phase of HA).40 Crystallite size was calculated using the Scherrer method41 and Williamson−Hall isotropic strain model (W−H−ISM),42 and the results are compiled in Table 2. Details of the calculations have been provided in Supporting Information section S2.2.2 and Figure S5. The crystallite size determined by both the methods showed a higher value for the sample HA(U). The size also increased drastically on calcination.

D0.1 = 1.1539 × 10−2 + 7.5933 × 10−2(x ̅ /s) + 6.6838 × 10−4(x ̅ /s)2 − 1.9169 × 10−4(x ̅ /s)3 + 3.9201 × 10−6(x ̅ /s)4

(1)

where s is the standard deviation. A histogram was constructed with the nanofiber spacing data as shown in Figure 3, and a log-normal model was imposed upon it. The x̅ and s values obtained from it were used to calculate the D0.1 values for the nanocomposites. While

Table 2. Crystalline Size Determination by Various Methods for Powder Samples and Nanocomposites crystallite size determination Scherrer method

Figure 3. Plot of frequency versus distance between nanofibers for VPHAU2. 8766

sample

(hkl)

2θ

Lc (nm)

W−H method (nm)

HA(U) HA(C) HB(U) HB(C) VPH0 VPHAU2 VPHBU2

(200) (200) (210) (210) (110) (110) (110)

21.2 21.2 28.9 28.9 12.5 11.9 11.4

36 65 17 22 1.9 1.6 1.5

34 66 18 22 -

lattice strain (W−H method) 5.35 4.90 7.27 5.57 -

× × × ×

10−4 10−4 10−4 10−4



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promising tool in determining the change ΔHS during nanocomposite formation. According to Maiti and Bhowmick,46 a shift of the IR peaks toward lower wavenumber in the ATR−FTIR spectra of nanocomposites corresponds to a negative value of ΔHS. This principle can be subtly used to enunciate the thermodynamic feasibility of the nanocomposite formation. ΔHS can thus be calculated from the Fowkes equation47 relating shift in peak position (Δν) and enthalpy change of the system. This is shown as follows

Lattice strain was found to undergo a significant decrease upon calcination.43 Crystallographic analysis clarifies that HA acquires a hexagonal unit cell where a equals to b and c is along the nanofiber axis. Crystallite size determination by peak broadening in WAXD studies confirms the fact that crystallite size of the nanoparticles increased upon calcination. This observation has consistency with that of Klinkaewnarong et al.44 According to Teng et al.,28 the driving force behind the increase in crystal size of the HA nanofibers on calcination is the lowering of high surface energy by fusion of smaller nanocrystals. Hence, random orientation of the HA nanofibers in the PDMS matrix results in marked reduction in the intensities of the peaks with an l index. In fact, peaks corresponding to diffraction in the planes (111), (002), and (211) showed a prominent collapse in intensity or even permanent disappearance in the WAXD profiles of the nanocomposites (Figure 4(b)). A comparison of the WAXD profiles of the nanocomposites is shown in Figure 4(b). In these WAXD profiles, the unfilled as well as the HA filled samples exhibit a prominent peak around 12.5° (2θ) along with a hump around 22°. These are due to the centered tetragonal unit cell of the polymer.45 Nanocomposites with both types of HA nanofibers demonstrate this phenomenon: a prominent shift of the peak at 12.5° toward lower 2θ value is observed along with peak broadening. The profiles reflect the restriction of the microcrystalline domain formation in the polymer matrix upon nanofiller incorporation. A comparison of microcrystalline domain size is shown in Table 2. However, additional peaks are visible in the XRD traces of both the nanocomposites. The nanocomposite with H(A) shows a cluster of peak in the 2θ range of 25−30°, while a similar array of peaks appears from 30 to 34° for the nanocomposite prepared with H(B). These additional peaks are in good coherence with those observed in the WAXD profiles of the nanofiller (Figure 4(a)). Synthesis of Nanocomposites: Thermodynamic Aspect. Nanocomposite formation does not merely mean mixing of the two components. Successful nanocomposite preparation implies proper mixing of the two components in a homogeneous fashion. Dispersion of the nanofiller in the polymer matrix is good only when the polymer has the ability to wet the nanofiller surface. Basically thermodynamic parameters owe their responsibility to the successful nanocomposite formation which is reflected in the property improvement of the nanocomposites. The entire process is governed by the free energy of mixing which is defined as ΔG E = ΔHE − T ΔSE

For the polymer

(2)

ΔGC = ΔHC − T ΔSC

For the nanofiller

(3)

ΔHS = 0.236 × Δν

Figure 5(a) shows a comparison of the ATR−FTIR spectra of the filled and the unfilled PDMS vulcanizates. The peak at

Figure 5. (a) Comparison of FTIR spectra of unfilled and HA nanofiber filled PDMS vulcanizates. (b) Comparison of the FTIR spectra of the uncalcined and calcined H(A) and H(B) samples.

1039 cm−1, which corresponds to asymmetric Si−O−Si stretching, was shifted by ∼20 cm−1 toward the lower wavenumber side. A similar shift was observed for the peak corresponding to Si−O−Si skeletal stretching around 800 cm−1. However, no shift in position was observed for the peaks around 2964 (asymmetric CH stretching), 1408, and 1259 cm−1 (CH3 asymmetric and symmetric deformation, respectively). This gives a comprehensive idea of some sort of interaction between the polymer and the filler which is responsible for improvement in various properties of the nanocomposites, as discussed in the subsequent sections. The values of enthalpy change of the systems are compiled in Table 3. The enthalpy change, ΔHS, is lowest with VPHBU2 and VPHBC2. This shows that nanocomposites were successfully

Here, ΔHE, ΔHC, ΔSE, and ΔSC are the respective changes in enthalpy and entropy of the polymer and the nanofiller during the process of nanocomposite formation. From the above equation, the free energy of the system can be formulated as follows ΔGS = ΔG E + ΔGC = (ΔHE + ΔHC) − T (ΔSE + ΔSC) = ΔHS − T ΔSS

(5)

(4)

The process of nanocomposite formation will be thermodynamically facilitated only when ΔGS is negative. This is possible only when ΔHS is negative and ΔSS is positive. FTIR is a 8767



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Table 3. Comparison of the Enthalpy and Energy Change Values for the Nanocomposite Preparation sample VPH0 VPHAU2 VPHAC2 VPHBU2 VPHBC2

peak position (cm−1) 1039 1019 1020 1023 1023

± ± ± ± ±

1 1 2 2 2

ΔHS (kcal/mol)

ΔGS (kcal/mol)

−4.720 ± −4.484 ± −3.776 ± −3.776 ±

−4.721 −4.485 −3.777 −3.777

0.472 0.944 0.944 0.944

± ± ± ±

0.472 0.944 0.944 0.944

prepared with PDMS as the continuous phase and HA nanofiber as the dispersed phase. An attempt was made in estimating the free energy change in mixing of the components during nanocomposite preparation. To execute this, entropy change was calculated using the Flory−Huggins theory derived from statistical mechanics.48−50 The entropy change of the system is formulated as ΔSS = −k[NC ln(NC/N ) + NE ln(xNE/N )]

(6)

ΔSS = −k[NC ln(ϕC) + NE ln(ϕE)]

(7)

ϕC = NC/N and ϕE = xNE/N

(8)

Figure 6. Comparison of the stress−strain plots of unfilled PDMS and hydroxyapatite nanofiber filled PDMS nanocomposites.

order, the tensile modulus increased by 211%, 113%, 43%, and 10%. The modulus at 50% and 100% elongation also followed the same trend. Elongation at break decreased in the order: VPHAU2 > VPHAC2 > VPHBU2 > VPHBC2. The factors which owe significant contribution to the property enhancement are nanofiber aspect ratio and nanofiber crystallinity and noncovalent surface modification of the nanofiller surface with polymer molecules. Nanofiller aspect ratio has a significant role in determining the physicomechanical properties of the nanocomposites. This was observed by Peng et al.54 who investigated the effect of aspect ratio of HA nanofiller in strengthening an electrospun polylactic acid scaffold. Improvement in mechanical properties is proportional to the aspect ratio of the filler, the latter being governed by several factors such as geometry and dispersion of the filler. This, in turn, is dependent on certain crucial parameters, the most important being surface energy of the components. Summarization of the results and calculations show that HA(U) with the largest value of aspect ratio (10.3) showed the greatest improvement in mechanical properties. However, the nanofiber prepared by method B (see Supporting Information) with lower aspect ratio 4.5 exhibited lesser improvement in mechanical properties. Determination of the aspect ratio of the nanofibers prepared by the two methods is discussed in detail in the Supporting Information (section S2.2.1 and Figure S4). Although the nanofiber aspect ratio does not change much with increasing crystallinity, a huge difference in magnitude of various properties was observed for the nanocomposites prepared with uncalcined and calcined nanofibers. This is because the high crystallinity of the nanofillers facilitates regularity in arrangement of the atoms and attainment of regular shape with smooth surfaces. Crystallinity was calculated from the XRD profiles of the HA nanofibers using the Scherrer and Williamson−Hall isotropic strain model (W−H−ISM). The results obtained as compiled in Table 2 show that crystallinity of the nanofiller increased upon calcination. As reported by Pang and Bao,55 HA nanoparticles with low crystallinity are irregular with poorly defined contour. However, upon calcination, nanoparticles with smooth surface are obtained. Irregular surfaces of the HA nanofiller are efficient in inducing enhanced polymer−filler adhesion through mechanical interlocking phenomena.56 This is evident from the different chain dynamics in the polymer molecules attached

Or

Here, ϕC and ϕE designate the volume fractions of the nanofiller and polymer, respectively, while NC and NE are the number of molecules of the respective components and k is the Boltzmann constant. Typical calculations reveal that ΔSS is ∼22.76 × 10−7 kcal/mol during nanocomposite formation. Hence, enthalpic control is significant in favoring the fabrication process, and for this system ΔGS ≅ ΔHS. ΔGS values are recorded in Table 3. These follow exactly the same trend as ΔHS values, as expected. From theoretical models, it is well illuminated that net entropy change during nanocomposite preparation is negligible or tends to zero.51 The probable underlying reason is that entropy gain due to debundling or dispersion of HA nanofibers is counterbalanced by the entropy loss due to polymer chain interpenetration between the fibers. Hence, the enthalpic term has a significant role in the free energy change during the hybrid material formation. As reviewed by Kumar and Krishnamoorti,52 a nonbonding force of interaction (as in this case hydrogen bonding) between the components leads to increased enthalpic contributions, thereby favoring the process. In addition, the concentration of the nanoparticles is also very low, thereby excluding the probability of nematic or columnar phase formation as explained by Onsager.53 Properties of the Nanocomposites. Mechanical Properties. Figure 6 shows the representative stress−strain curves of the nanocomposites prepared with unfilled and filled PDMS vulcanizates, and the results of the mechanical properties are summarized in Table 4. The stress−strain curves maintain their rationality and are well justified when compared with the characteristic stress−strain curve of a conventional vulcanized elastomer. Tensile strength and tensile modulus of the nanocomposites prepared with H(A) showed a dramatic increase in magnitude compared with those prepared with H(B). Moreover, uncalcined nanofibers were found to impart better reinforcement compared with the calcined ones. For instance, tensile strength increased by 109%, 72%, 64%, and 23% for VPHAU2, VPHAC2, VPHBU2, and VPHBC2, respectively. In the same 8768



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Table 4. Comparison of Various Properties of Neat Rubber Vulcanizate and Its Nanocomposites with Different Types of HA Nanofibers storage modulus (MPa)

sample VPH0 VPHAU2 VPHAC2 VPHBU2 VPHBC2

Emod (kPa) 140 436 299 201 154

± ± ± ± ±

20 14 23 5 12

tensile strength (TS) (kPa) 167 349 287 274 206

± ± ± ± ±

12 17 8 16 9

elongation at break (EAB) (%)

temperature of maximum degradation Tmax (°C)

rate of degradation (%/min)

± ± ± ± ±

350 549 539 535 523

35 17 18 11 15

506 340 242 291 135

59 25 19 25 17

at −120 °C at −75 °C 1210 720 630 1640 1080

376 170 210 530 363

at 25 °C

tan δ height at Tg

0.164 0.335 0.281 0.267 0.187

0.195 0.094 0.109 0.111 0.121

Figure 7. Comparison of (a) tan δ versus temperature and (b) storage modulus versus temperature plots for unfilled and HA nanofiber filled PDMS vulcanizates. (c) Schematic representation of the restriction of microcrystalline domain formation in the nanocomposites.

are set free to interact with the PDMS matrix. The molecular weight of PEG 400 being low, the density of the hydroxyl group per unit area is appreciably high. This possibly enables the pendent hydroxyl functionalities of PEG to interact with the O atom in the PDMS backbone. This is probably the reason for the unique morphology of the nanocomposite and a plausible explanation for the unprecedented increase in modulus values. Thermal Properties. Oxidative thermal stability of the nanocomposites increased markedly upon incorporation of HA. Improvement in temperature of maximum degradation (Tmax) and temperature of onset of degradation (Ti) was observed for all the nanocomposites. However, maximum improvement was observed in the case of VPHAU2. This is shown in the results of Table 4. Tmax increased by 199, 189, 185, and 173 °C for VPHAU2, VPHAC2, VPHBU2, and VPHAU2, respectively. Rate of degradation was also found to undergo a drastic decrease upon nanofiller incorporation. A significant improvement of barrier properties in heat transmission for nanofillers with higher aspect ratio also

to or in the vicinity of the geometrically constricted nanoparticulate surface.57,58 Hence, with uncalcined nanofibers, the polymer molecules find more anchoring sites due to the structural defects in the nanofiber. This facilitates stronger polymer−filler interface formation. As obvious from the FTIR spectrum of HA(U) powder in Figure 5(b), there is a peak assigned to the PEG layer around 1115 cm−1 which is adsorbed on the HA nanofiber surface apart from the absorptions for the PO43− unit (564 and 602 cm−1 for triply degenerated bending vibration (ν4) of PO4−3, P−O symmetric stretching vibration at 963 cm−1 (ν1)) and for the OH unit (strong bands at 3570 and 632 cm−1). The driving force for this polymer wrapping is the electrostatic force of attraction between the Ca2+ ion of the nanofiller and the O atom of the PEG molecules. Noncovalent surface wrapping of PEG molecules facilitates better dispersion of the nanofillers and stronger interface formation which also account for the increase. The PEG layer acts as a form of noncovalent surface functionalization of the nanofiller. However, the −OH groups 8769



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plot where both the peaks at −57 °C and −30 °C are replaced by humps or are considerably minimized. In fact, the microcrystalline domains in the polymer matrix act as crosslinks, thereby imparting a greater strength to the matrix. These domains owe a good contribution to the modulus value. The nanofillers with higher aspect ratio are more proficient in restricting the microcrystalline domain formation. Hence, consequence of reduction of the number as well as size of the crystalline domains is a marked reduction in lowtemperature storage moduli.12,64 This is consistent with the WAXD results of the nanocomposites where nanocomposites showed significantly reduced size of the microcrystalline domains (according to crystallite size determined through the Scherrer method in Table 2). From room temperature onward, the modulus is predominantly governed by the reinforcing effect of the filler and restriction of the macromolecular chains.65 This is well reflected from the storage moduli values for the nanocomposites. The loss in crystalline domains also took place in the nanocomposites with H(B). However, due to a lower aspect ratio, the effect was not very marked. A hightemperature modulus clearly showed a better reinforcement effect of the fillers with higher aspect ratios. Modeling of Modulus of the Nanocomposites. To quantify the extent of contribution of the aspect ratio of the filler in property improvement, the moduli values (both static and dynamic) were compared with predictions from various models. The Halpin−Tsai model66−68 shows good agreement with the experimental values of the modulus of the nanocomposites with reinforcing fibrous filler. This is shown as follows

improves the oxidative thermal stability of the respective nanocomposite. Thermal stability is related to restriction in mass transport of the volatiles generated upon degradation through the barrier mechanism of the nanofiller along with insulation of the polymer molecules by the latter.59 Nanofillers with higher aspect ratio are more competent in improving the thermal stability of the nanocomposites. Hence, in coherence with the observation of Liu and Yang,60 nanocomposites prepared with H(A) nanofillers showed improved thermal stability. Dynamic Mechanical Properties. Incorporation of the HA nanofiber into the PDMS matrix resulted in a marked improvement in the room temperature storage modulus for all the nanocomposites. This is depicted in Figure 7(b). This marked improvement is due to enhanced polymer−filler interaction which was reflected in the lowering of the peak at Tg in the tan δ versus temperature plot (Figure 7(a)). However, there was, as such, no marked improvement in low-temperature storage modulus of the nanocomposites. Interestingly, for VPHAU2 and VPHAC2, the low-temperature modulus dropped below that of the unfilled PDMS vulcanizate. Besides the peak corresponding to Tg at ∼−117 °C,61 the tan δ versus temperature plot (Figure 7(a)) for the unfilled system showed a prominent peak around −57 °C which is characteristic of the crystalline domains present in the PDMS matrix62 and a peak around −30 °C due to melting of the crystalline domains.61 However, in the nanocomposites the crystalline peak was replaced by a broad hump spanning a temperature range of 100 °C in the same region. The crystalline melting peak also showed a prominent reduction in height. An in-depth analysis of the storage modulus plot (Figure 7(b)) also leads to a similar observation. While the unfilled polymer showed a three-step descent corresponding to glass transition and melting of crystalline domains, these transitions were not very well-defined in the storage modulus plots for the nanocomposites. Moreover, the nanocomposites failed to retain the high modulus, which is evident from the modulus values at −33 °C. However, all the nanocomposites irrespective of the nature of the HA used showed a pretty high room temperature modulus. The nanocomposites VPHA showed higher room temperature storage modulus values due to the higher nanofiber aspect ratio compared with VPHB nanocomposites. The nanocomposites prepared with calcined nanofibers showed a decrease in storage modulus value over the entire temperature span. For instance, at 25 °C, VPH0 showed a storage modulus value of 0.164 MPa, whereas the modulus value increased to 0.335 MPa VPHAU2. A more detailed compilation of the data related to the magnitude of storage moduli and height of tan δ around Tg is shown in Table 4. The dynamic modulus also shows some unique trends in the low-temperature region. A significantly reduced low-temperature moduli value for the nanocomposites is the consequence of two independent phenomena. One of them is the plasticizing effect of the PEG molecules adhered to the nanofiber surface as discussed in the previous section. This phenomenon is more prominent for the uncalcined sample VPHAU2 which has nanofibers coated with PEG molecules. The plasticizing effect of the PEG molecules is so prominent that the low-temperature modulus of the nanocomposite falls below that of VPH0. Besides this, incorporation of the filler in the polymer matrix restricts the growth of microcrystalline domains63 in the polymer as shown in the pictorial representation in Figure 7(c). This is also well reflected from the tan δ versus temperature

(1 + ξηϕ) E = Em (1 − ηϕ)

(9)

where η depends on the shape and orientation of the filler in the polymer matrix and is formulated as η=

Ef Em Ef Em

−1 +ξ

(10)

Ef being the modulus of the filler (180 GPa in this case).54Here, E and Em are the modulus for the nanocomposite and pure polymer, and α and ϕ are, respectively, the aspect ratio and filler volume fraction. ξ is the shape factor and is equivalent to twice the aspect ratio.The Halpin−Tsai model is applicable to both static and dynamic moduli of the nanocomposites. In addition, assuming a strong interface formation, the Hui and Shia model for the tensile modulus was also used for the system.69,70 According to the model: Longitudinal modulus Ec 1 = Em 1−

ϕ ξ

(11)

and Transverse modulus Ec 1 = ϕ⎡1 Em 1 − 4 ⎣ξ +

3 ⎤ ξ+ Λ⎦

(12)

where 8770


The Journal of Physical Chemistry C ξ=ϕ+

⎡ (1 − g )α 2 − Em + 3(1 − ϕ)⎢ ⎢⎣ E f − Em α2 − 1

Article g 2

⎤ ⎥ ⎥⎦

The phenomena of polymer adsorption and hydrogen bonding owe their responsibility partly to this property improvement. Enhancement in properties is evident from the shift in the position of the absorptions in the ATR−FTIR spectra of the nanocomposites. The filler loading and cross-linker amount being the same in every sample, the nanofiller aspect ratio is the key factor in determining the extent of improvement in various properties. It was found that maximum improvement in tensile strength and modulus, storage modulus, and oxidative thermal stability follow the order VPHAU2 > VPHAC2 > VPHBU2 > VPHBC2. Compared with the lower mean aspect ratio nanofibers prepared by method B, long nanofibers prepared by method A resulted in better reinforcement of the polymer matrix. An increase in 109%, 72%, 64%, and 23% tensile strength was observed for VPHAU2, VPHAC2, VPHBU2, and VPHBC2, respectively. In addition, nanofiller amorphicity is also responsible partly for this improvement in properties. Stronger scaffold formation and adherence of the PDMS matrix to the PEG-coated HA nanofibers have prominent contribution to the improved properties. The property improvements were correlated with the morphology of the nanocomposites through HRTEM and FESEM analysis of the nanocomposites. Dispersion was at its best for VPHAU2 with the highest dispersion degree parameter D0.1 value of 6.30%. This was reflected in maximum improvement in properties of the nanocomposite. In a similar way, morphology analysis was utilized in elucidating property improvement for the other nanocomposites. However, incorporation of HA into the polymer matrix resulted in a reduction in crystallinity of the polymer as observed from the XRD studies. This phenomenon also affected the properties of the nanocomposites.

(13)

and ⎡ 3(α 2 + 0.25)g − 2α 2 ⎤ ⎥ Λ = (1 − ϕ)⎢ α2 − 1 ⎣ ⎦

(14)

Here ⎡ ln(2α) − 1 ⎤ −4 g=1−⎢ ⎥ + O(α ) ⎣ ⎦ α2

(15)

where O symbolizes the error of the approximations. Figure 8 is a comparison of the experimental values of the modulus with various modulus models. The inset is a schematic

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ASSOCIATED CONTENT

S Supporting Information *

Figure 8. Comparison of experimental values and predicted values for the modulus ratio (E/Em). Inset shows a model of the interface between the hydroxyapatite nanofiber (HANF) (coated with PEG) and PDMS.

Full description of the synthesis of HA nanofibers, morphology analysis through HRTEM, chemical composition through EDX, and crystallinity variation through SAED analysis. In addition, particle size determination through Image Analysis and details of crystal size determination through Scherrer method and Williamson−Hall isotropic strain model (W−H−ISM) are also incorporated here. This material is available free of charge via the Internet at http://pubs.acs.org.

representation of the interface between the polymer and the HA nanofibers for the H(A) sample. It is found that there is a significant deviation between the experimental data and predictions, particularly for VPHAU2 and VPHAC2, although the basic trend of variation is similar. The huge deviation in the experimental values and the predicted values is the outcome of the several factors discussed in the preceding sections. The aspect ratio of HA(U) is highest with an additional advantage of noncovalent surface modification. This facilitates maximum improvement in the modulus for the VPHAU2 nanocomposite.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]/[email protected]. Phone: 91-3222-283180/91-612-2277380. Fax: 91-3222-220312/91612-2277384.

■

CONCLUSIONS HA nanofibers were synthesized by two methods and were characterized extensively through morphological analyses. HRTEM and WAXD studies were used to determine the dimension of the as-prepared and calcined nanofillers. The mean aspect ratio of the nanofillers prepared by the two methods varied significantly. These were successfully incorporated into the polymer matrix. Enthalpy and hence free energy change for the whole process implied a thermodynamically favored phenomenon. Various physicomechanical and thermal properties showed a dramatic improvement upon a low level of filler incorporation.

Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The financial grant for this work by the Council of Scientific and Industrial Research (CSIR), New Delhi, is gratefully acknowledged.

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REFERENCES

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