Carboxy Anhydride Polymerization in Nanoporous Anodic Alumina

Feb 19, 2009 - the size of the O-ring (Scheme 1B)sinner diameter ) 7.6 mm and thickness ... Solution Reservoir. 3180 J. Phys. Chem. B, Vol. 113, No. 1...
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J. Phys. Chem. B 2009, 113, 3179–3189

3179

In situ Characterization of N-Carboxy Anhydride Polymerization in Nanoporous Anodic Alumina K. H. Aaron Lau, Hatice Duran,* and Wolfgang Knoll Max Planck Institute for Polymer Research, Ackermannweg 10, D-55128 Mainz, Germany ReceiVed: October 30, 2008; ReVised Manuscript ReceiVed: December 15, 2008

Poly(γ-benzyl-L-glutamate) (PBLG) has been a popular model polypeptide for a range of physicochemical studies, and its modifiable ester side chains make it an attractive platform for various potential applications. Thin films of Poly(γ-benzyl-L-glutamate) PBLG were surface grafted within nanoporous anodic alumina (AAO) by surface-initiated polymerization of the N-carboxy anhydride of benzyl-L-glutamate (BLG-NCA). The grafting process was characterized by optical waveguide spectroscopy (OWS), infrared spectroscopy (FT-IR), and scanning electron microscopy (SEM). OWS was able to track the PBLG layer thickness increase in situ, and ex situ FT-IR gave complementary information on the PBLG chain’s secondary structure. Transitions in the PBLG growth rate could be correlated with transitions in the polypeptide secondary structure. The emergence of a three-dimensional, anisotropic PBLG morphology within the cylindrical pores of the AAO membrane was also identified as the grafted PBLG average layer thickness increased. Comparison of the PBLG/AAO results with those on a planar silicon dioxide surface indicated that both the conformational transitions and the PBLG nanostructure development could be attributed to the confining geometry within the pores of the nanoporous AAO matrix. The use of a nanoporous AAO matrix, combined with the surface grafting of a thin film of PBLG chains with multiple modifiable side chains, could potentially offer a nanoporous platform with a very high density of functional sites. I. Introduction The tethering of rodlike, synthetic polypeptides on solid surfaces with a controlled orientation and molecular conformation is desired for a range of applications, including chiral separation membranes,1 optoelectronic devices,2-5 and biosensing.6-8 Polypeptides are characterized by a tunable secondary structure with a range of unique properties: they can exist in R-helical, β-sheet, and random coil conformations with R-helices and β-sheets constituting the secondary structures of proteins,9 and helical peptides are also considered to be promising materials for the construction of supramolecular systems.10 In addition, the transition from the R-helix to the random coil state in solution shares key characteristics with protein denaturation.11 Poly(γ-benzyl-L-glutamate) (PBLG) has been a popular model synthetic polypeptide for the study of a range of physicochemical phenomena, including conformational transitions,12 anisotropic dielectric properties of macromolecules,13-16 and surface grafting behavior of polypeptides.17-23 Among its attractive properties, PBLG has the most stable R-helical structure among the polyglutamates, and its chirality also makes it a suitable platform for applications in chiral separation.1,24 Moreover, as first proposed by Flory, its flexible ester side chains favor solubility of the R-helical conformers in most organic solvents25 and may be easily modified.1 Surface grafting of PBLG on planar surfaces has been extensively investigated for creating high density polypeptide layers with an aligned helical orientation. In order to control the orientation and grafting density, two strategies have generally been applied: (i) ring opening polymerization of the N-carboxy anhydride (NCA) monomer via surface attached initiator groups (grafting from)18-21 and (ii) coupling of preformed R-helical * Corresponding author. E-mail: [email protected]. Tel.: 496131-379-545. Fax: 49-6131-379-100.

PBLG to solid surfaces (grafting to).19,22,23,26 In the “grafting to” approach, due to the strong interaction between polyglutamate molecules and many surfaces, the grafted polypeptide molecules tend to lie parallel to the surface, thus blocking other surface binding sites and hindering subsequent chemisorption.18 Also, due to strong dipole interactions, PBLG molecules already grafted to the surface have a tendency to repel other helices having the same directional sense, resulting in lower grafting surface densities.19 At the same time, the accessibility of a monomer to surface-reactive groups in the “grafting from” approach is larger than that of a “grafting to” preformed polymer, therefore many grafting from approaches have been suggested,18,21 although the degree of polymerization achievable for surface-initiated polymerization can be highly dependent on solvent and monomer loading techniques.18 Experimentally, Chang and Frank19 reported higher PBLG grafting densities for the “grafting to” than the “grafting from” approach, but the resulting alignment of grafted to siloxane-PBLG was relatively disordered on silicon oxide surfaces. Heise et al.27 used long chain alkyl spacers to adjust the density of initiating groups at the surface and obtained uniform PBLG films, but surface chain densities were not reported. Taken together, these studies show that the chain density and the helical alignment are strongly dependent on the preparation conditions. For biosensing applications, it is desirable to maximize the number of surface binding sites in order to enhance detection sensitivities. The use of a nanoporous matrix to provide an extra internal surface for binding is one step toward that goal.28,29 Surface grafting of functional polymer brushes to create more binding sites per unit area represents another strategy.7,30 To achieve this goal, we investigated the surface grafting of PBLG on the internal pore surfaces of nanoporous anodic alumina oxide (AAO) membranes via the grafting from approach. AAO

10.1021/jp809593d CCC: $40.75  2009 American Chemical Society Published on Web 02/19/2009

3180 J. Phys. Chem. B, Vol. 113, No. 10, 2009 membranes are formed by anodization of Al films in an acidic electrolyte.31,32 They are characterized by close-packed, cylindrical nanopores that run straight through the thickness of the membrane and their tunable nanostructure has attracted great interest in nanotechnology33-35 and optical applications.29,36 For PBLG surface grafting, the AAO surface was first functionalized using a silane terminated with a primary amine via gas phase deposition37 to provide a high density initiator surface groups for NCA polymerization. In our previous studies,6,7 we demonstrated that we can sensitively monitor silane functionalization of nanoporous AAO by optical waveguide spectroscopy (OWS) and effective medium theory (EMT) analysis. OWS is a simple and versatile technique for characterizing the thickness and optical properties of thin films,38,39 and it provides a nondestructive tool for monitoring the growth of organic layers inside nanoporous thin films.7,28,29 The optical response of such nanostructured materials can then be analyzed by EMT, based on the nanostructure morphology and the bulk dielectric responses of the matrix material and pore-filling medium.29,40,41 Here, we applied the technique to characterize, in situ, the surface-initiated polymerization of PBLG within a nanoporous AAO thin film with pores 60-70 nm in diameter and around 1 µm in thickness. The present investigation of the PBLG film growth process within a nanoporous matrix is unprecedented. Furthermore, the OWS data was complemented by ex situ measurements of the polypeptide’s secondary structure using Fourier transform infrared spectroscopy (FT-IR), and the final film morphology, using scanning electron microscopy (SEM). Different stages of the PBLG film growth in the AAO pores with different growth kinetics and secondary structures were observed. In addition, in situ OWS analysis was able to discern the emergence of an anisotropic PBLG morphology at high degrees of polymerization, which was corroborated by SEM microscopy. II. Experimental Section 1. Materials and Characterization. The N-carboxy anhydride of the benzyl-L-glutamate (BLG-NCA) monomer of PBLG was synthesized by phosgenation of L-glutamic acid (Merck Biosciences) with stoichiometric quantities of triphosgene (Sigma-Aldrich) in dry ethyl acetate (Fluka). The monomer was purified two times by recrystallization in n-hexane. 1H NMR spectra were recorded on a Bruker Dpx250 spectrometer, using the residual proton resonance signal of the deuterated dichloro methane (CD2Cl2) as the internal standard. 1H NMR (CD2Cl2 250 MHz): δ 2.17 (m, (ppm): β-CH2), 2.52 (m, γ-CH2), 4.32 (t, C-H), 5.14 (s, -CH2-benzylic), 6.31 (s, N-H), 7.41 (s, Ar-H). (Caution! Triphosgene decomposes to highly toxic phosgene on heating and upon reaction with any nucleophile. EVen a trace of moisture leads to the formation of phosgene.42 Therefore, work with this substance should be conducted in a highly ventilated fume hood to prevent exposure, and reaction byproduct should be neutralized immediately.) Dichloroacetic acid (99+%, Acros Organics), chloroform (Merck), and anhydrous dimethylformamide (DMF; Merck) were used as received. 3-Aminopropyltriethoxysilane (APTES; Sigma-Aldrich) was dried and distilled prior to use. Anhydrous tetrahydrofuran (THF) was prepared by the following procedure: 1 mL 1-1-Diphenylethylene was added to THF, and then, n-BuLi was added dropwise until the THF became red. The anhydrous THF was finally obtained by distillation and used immediately. 2. BLG-NCA Polymerization. (See Scheme 1.) APTES was used as the surface attached initiator. Substrates (AAO and SiO2; see below) were exposed to silane vapors at 100 °C under

Lau et al. SCHEME 1: (A) PBLG Surface Grafting on the APTES Functionalized Pore Surfaces of AAO Thin Films via Surface Initiated BLG-NCA Polymerization. (B) OWS Setup with Liquid Flow Cell Assembly and Peristaltic Pump for Cycling Solution between the Liquid Cell and Solution Reservoir

vacuum (2-3 mbar) for 3 h, then rinsed with copious amounts of acetone.37 The polymerization was carried out in anhydrous THF at room temperature, and the concentration of BLG-NCA in THF was 100 mM. A Teflon-PTFE liquid flow cell fitted with Teflon tubings and sealed with a Chemraz O-ring (Greene, Tweed & Co., Germany), was used together with a peristaltic pump (Reglo Digital, Ismatec, Germany; operated at 2-10 mL/ min) for in situ OWS measurements (see section 5 below). The PTFE/Chemraz liquid cell setup was used to prevent THF evaporation from the liquid cell (to maintain a constant BLGNCA concentration) and to minimize the introduction of water impurities (which could cleave an anhydride ring and lead to formation of unbound polymer in the solution). In addition, the peristaltic pump cycled BLG-NCA solution from the flow cell to an Ar-filled solution container in order to prevent the accumulation of carbon dioxide produced during the reaction. The volume of the flow cell was 0.045 mL and was defined by the size of the O-ring (Scheme 1B)sinner diameter ) 7.6 mm and thickness ∼ 1 mm. The Teflon tubing also introduced a volume of ∼1 mL to the flow cell setup. A 15-20 mL reservoir of the monomer solution was used for each experiment. The polymerization was initiated by flowing in the BLG-NCA solution (at 2 mL/min) and terminated by exchanging the NCA solution with pure THF, followed by continuous flushing of the flow cell with fresh THF for 1 h (at 10 mL/min for the first 10 min, then at 2 mL/min). For ex situ IR and GPC measurements

PBLG Surface Grafting in Nanoporous Alumina (see below), a cleaning post-treatment reported in the literature18 was also applied: the substrates were (i) sonicated in dichloroacetic (DCA)/chloroform 1/9 (v/v) solution (pH ) 5.0) for 30 min, (ii) washed with copious of chloroform and dried with Ar, (iii) immersed in anhydrous DMF for 24 h at room temperature, and (iv) finally washed with THF and dried with Ar. 3. Nanoporous AAO Waveguide Layers. AAO thin film waveguides were prepared based on a previously described method.29 A ∼1 µm Al thin film was deposited onto glass substrates and subsequently anodized31,32 to form nanoporous AAO. In this report, a thin layer of Al was used as the optical coupling layer, and either a 1- or 2-step43 anodization process was employed. High index LaSFN9 glass slides (εLaSFN9 ) 3.406 at λ ) 632.8 nm, Hellma Optik, Halle, Germany) were used as substrates, and Al was deposited by DC magnetron sputtering in an Ar plasma (provided by Institut fu¨r Mikrotechnik Mainz, Germany). The Al films, connected as the anode opposite a Pt counter electrode, were anodized at a constant potential of 40 V in a beaker of 0.3 M oxalic acid in deionized water at 2 °C. For 1-step anodized AAO, the anodization was allowed to proceed until only 25 nm Al remained (∼40 min), at which point the anodization current was switched off. For 2-step anodized AAO, the Al films were first anodized for 10 min such that a ∼300 nm sacrificial layer of AAO was formed. This was removed by 3 h of immersion in 5 wt % H3PO4. Anodization was then applied a second time until, as for the 1-step samples, only 25 nm Al remained. The 2-step process slightly improved the spatial distribution of the pore openings but otherwise did not alter the AAO pore structure. The average pore diameter (Dpore) for both types of samples was then widened to 60-70 nm by immersion in 5 wt % H3PO4 for 50-60 min. The final layer structure was the following: LaSFN9/Al(25 nm)/ AAO(600-1000 nm). 4. SiO2 Planar Waveguides. SiO2 waveguides were prepared on LaSFN9 glass substrates, and the layer structure was the following: LaSFN9/Cr(1 nm)/Ag(30 nm)/Cr(1 nm)/ SiO2(sputtered, ∼1 µm)/SiO2(sol-gel, 45 nm). The Cr layers were used to promote adhesion between the Ag and the glass/ SiO2 layers. The ∼1 µm thick SiO2 layer was deposited by RFplasma sputtering directly from a SiO2 target (400 W, 8 × 10-2 mbar Ar, using an Edwards Autolab 500). An additional SiO2 sol-gel layer was prepared on the sputtered SiO2 in order to optimize the solvent stability of the waveguide.44 The SiO2 layers had a dielectric constant identical to that of bulk amorphous SiO2 (εSiO2 ) 2.1, measured by OWS at λ ) 632.8 nm). The top sol-gel SiO2 layer also exhibited a smooth surface with a residual roughness Rrms ) 0.5 nm (see Supporting Information Figure S1). 5. Optical Waveguide Spectroscopy (OWS) and Effective Medium Theory (EMT). A customized instrument setup in the Kretschmann configuration38 was employed to measure the waveguide responses of the AAO and SiO2 thin films deposited on glass substrates (Scheme 1B). The backside of the substrate was attached to the base of a symmetric glass prism (with a dielectric constant identical to that of the glass substrate) by optical immersion oil. A laser (λ ) 632.8 nm) was incident through the prism-substrate assembly and reflected off the thin metal coupling layer in between the oxide thin film and the substrate (i.e., the Al layer in the AAO samples, and the Ag layer in the SiO2 samples) as the incidence angle (θ) was varied. At specific angles determined by the thickness and the dielectric constant(s) of the oxide film (εfilm), the laser was coupled into the film and such waveguide modes were measured as sharp

J. Phys. Chem. B, Vol. 113, No. 10, 2009 3181 minima in a reflectivity (R) vs θ scan. Our instrument was able to obtain an angle resolution better than ∆θ < 0.005°. Different sets of waveguide modes were coupled under different incidence polarizations: transverse electric (TE) modes were excited for light with an electric field component parallel to the waveguiding film surface; transverse electric (TM) modes were excited for light with a magnetic field component parallel to the film surface. These TE and TM modes were indexed according to the number of nodes in their electromagnetic field distributions.38 From the angles of the waveguide mode reflectivity minima, εfilm could be derived by Fresnel calculations, which exactly describe the layered geometry of the thin film waveguide in the Kretschmann configuration;38,39 tracking the coupling angle of a mode enabled real time, in situ monitoring of changes in the thickness or the dielectric constant of the film.38 The EMT approximation was used to analyze εfilm of the nanoporous AAO (εAAO) with respect to the film’s nanostructure.29,45 Due to the anisotropic nature of the long, cylindrical AAO pores, the dielectric response of the AAO thin film was also anisotropic and was described by the infinite, prolate ellipsoid approximation within Maxwell-Garnett theory:40,41 ⊥ εAAO ) εalumina + fpore(εpore - εalumina)

| εAAO ) εalumina

(1A)

εalumina + 1/2(1 + fpore)(εpore - εalumina) εalumina + 1/2(1 - fpore)(εpore - εalumina) (1B)

⊥ | where εAAO and εAAO are, respectively, the dielectric constant components normal and parallel to the AAO film surface, fpore is the pore volume fraction within the AAO, εalumina () 2.68)6 is the dielectric constant of bulk anodic alumina at λ ) 633 nm, and εpore is the (effective) dielectric constant within the pores. For a blank AAO film in air, εpore ) εair ) 1; in THF, εpore ) εTHF ) 1.980. With the addition of an organic film on the internal pore surfaces (e.g., APTES or PBLG), the volume within the pores would be occupied by a combination of the organic material and the pore filling medium (i.e., air or solvent). In the simplest approximation, εpore was estimated by the volume-weighted-averageoftheconstituents’dielectricconstants:29

εpore ) εmedium + forganic(εorganic - εmedium)

(2)

where forganic is the volume fraction of the organic layer within the cylindrical pores. The organic layer thickness (dorganic) could then be calculated from forganic and the pore diameter by assuming that the organic film formed a conformal layer on the pore surfaces:

dorganic )

Dpore (1 - √1 - forganic) 2

(3)

A more sophisticated analysis of εpore is presented in section III.4. 6. Scanning Electron Microscopy (SEM). SEM measurements were performed with a LEO Gemini 1530 SEM. The electron acceleration voltage used was between 1 and 6 kV. 7. Fourier Transform Infrared (FT-IR) Spectroscopy. FTIR spectra were recorded in reflectance mode on a Nicolet Magna-IR 850 spectrometer equipped with a Nic-Plan Microscope. Omnic series software was utilized for data acquisition.

3182 J. Phys. Chem. B, Vol. 113, No. 10, 2009

Figure 1. OWS R vs θ measurements of a nanoporous AAO thin film before and after surface functionalization with a 1.7 nm APTES layer. The waveguide mode minima are labeled by their polarizations and the number of nodes in their optical fields. The solid symbols represent data for the blank AAO film, the open symbols are data for the APTES functionalized film, and the curves are the corresponding Fresnel calculations.

For each sample, 1000 scans were taken at a resolution of 4 cm-1 with 1 h induction time for N2 exposure (to eliminate noise from atmospheric water). Further, the AAO thin film samples for FT-IR had a 200 nm Al layer, compared with the 25 nm layers for OWS measurements, to maximize the reflected signal. The samples were placed in the spectrometer so that the external, top surface of the AAO membrane was perpendicularly to the plane of the incident beam. In order to achieve the highest signal-to-noise ratio, the beam was polarized parallel to the external AAO top surface. The characteristics peptide peaks were calculated by deconvulation of the amide I and amide II regions using Gauss-Lorentzian fitting. 8. Molecular Weight and Polydispersity Characterization by Gel Permeation Chromatography (GPC). A GPC instrument equipped with three PVXL-TSK columns of 103, 104, and 106 Å pore sizes was employed for the molecular weight and molecular weight distribution determination of the PBLG films. A polystyrene solution in DMF was used as the eluent. The measurements were performed at 60 °C with 1 mL/min flow rate. A pump (Waters 515) was used with refractive index (ERC RI-101) and ultraviolet at 270 nm (SOMA, UV S-3702) detectors. 9. Etching of AAO Templates for GPC Measurements. The grafted PBLG material was released from AAO substrates by selective dissolution of the alumina using an aqueous HF solution (45%) at 0 °C for 1.5 h and subsequent neutralization of the suspension by repeated filtration steps. III. Results and Discussion 1. Characterization of the AAO Thin Film and APTES Surface Functionalization. OWS38,39 is a powerful yet experimentally simple technique for independently characterizing the thickness and the (anisotropic) dielectric constant(s) of optically transparent thin films. In OWS measurements, a planar thin film acts as a slab waveguide, in which optical waves are confined within the thickness of the film but could otherwise freely propagate in directions parallel to the film surface. Figure 1 shows the reflectivity (R) vs incidence angle (θ) measurements of a 2-step anodized nanoporous AAO thin film in air, before and after surface functionalization with an APTES layer. For the blank film, the angles of the TE and TM waveguide mode reflectivity minima represented38,39 an AAO film thickness of 0.64 µm and components of the film dielectric constants normal ⊥ | , εAAO } of {1.757, 1.619}. and parallel to the film surface {εAAO

Lau et al.

Figure 2. SEM images of the nanoporous AAO thin film prepared on a glass substrate. The inset shows the cylindrical pores that run straight through the film thickness.

These dielectric constant values could be best-fit by an EMT approximation (eq 1A) corresponding to a pore fraction (fpore) of 55%. Assuming a close-packed pore center-to-center distance ∼100 nm typical of AAO films anodized in oxalic acid (see below), the value of fpore obtained implied an average pore diameter (Dpore) of ∼76 nm. The OWS/EMT analysis was corroborated by SEM characterization. Figure 2 shows the top surface of the AAO thin film, and pore openings with diameters in the range of 60-80 nm could be observed. Although the pores exhibited only local ordering, an average pore center-to-center distance ∼100 nm was measured. The lack of long-range pore ordering is typical of AAO prepared from vacuum deposited Al films, for which the initial Al film is too thin to enable the growth of a thick, sacrificial AAO layer for hexagonal pore ordering to develop by self-organization.29,34 The cross-section of the AAO film is displayed in the inset and shows the typical AAO nanostructure that is characterized by cylindrical pores that run straight through the film thickness. Moreover, the film thickness shown in the cross-sectional image is consistent with the OWS measurement of 0.64 µm. After APTES surface functionalization by vapor deposition, Figure 1 shows that the waveguide mode minima had shifted ⊥ , to higher angles, which corresponded to increases in {εAAO | εAAO} to {1.809, 1.669}. In the blank AAO film, the pores were completely filled with air. After APTES functionalization, a certain fraction of the pore volume was occupied by APTES. Since the dielectric constant of APTES (εAPTES ) 2.019)46 is higher than air (εair ) 1.000); therefore, the dielectric constant within the pores (εpore), hence also εAAO, increased (see eq 1A). The average APTES layer thickness estimated by eq 3 was 1.9 ( 0.12 nm, compared with the theoretical APTES monolayer thickness of ∼1 nm.46,47 FT-IR measurements also confirmed the addition of the amine surface-initiator group after APTES deposition (data not shown). 2. Complementary OWS, FT-IR, and SEM Characterization of BLG-NCA Surface-Initiated Polymerization. PBLG surface grafting by surface-initiated, ring-opening polymerization of ΒLG-NCA monomers on APTES modified nanoporous AAO thin film matrix was characterized in situ by OWS (Scheme 1A). The monomer concentration used was 100 mM in anhydrous THF, and the reaction was carried out at room temperature. On the OWS setup, the AAO sample was mounted with a liquid flow cell and a peristaltic pump system (Scheme 1B). The flow cell was initially filled with pure anhydrous THF, and the polymerization commenced as the THF was exchanged with the BLG-NCA solution by the peristaltic pump. To minimize undesirable side reactions and premature termination, the monomer source bottle was filled with Ar.

PBLG Surface Grafting in Nanoporous Alumina

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Figure 4. Ex situ FT-IR scans for BLG-NCA polymerized on nanoporous AAO substrate with 100 mM BLG-NCA concentration in THF after various time intervals.

Figure 3. (A) In situ OWS R vs θ measurements in THF of the nanoporous AAO film before and at the end of PBLG grafting. The symbols represent data, and the curves are the matching Fresnel calculations. (B) Corresponding real-time TM1 angle shift and dielectric constant measurements. The average PBLG layer thickness estimated from the measured anisotropic dielectric constants is also indicated on the outer left axis. The inset SEM image shows the top surface of the AAO film with decreased pore diameters (compare with Figure 2) due to addition of the grafted PBLG layer.

The APTES functionalized AAO sample used for the in situ BLG-NCA polymerization experiment was the same as that already shown in Figure 1 for the R vs θ measurements in air. Comparison between the R vs θ measurements of the APTES functionalized sample in air (Figure 1) with the measurement in the THF-filled flow cell (Figure 3A) shows that the waveguide mode minima have all shifted to higher angles with the introduction of THF. Since the dielectric constant of THF, εTHF ) 1.980, was larger than air, εair ) 1.0005, infiltration of the pores with THF led to an increase of the dielectric constant ⊥ | , εAAO } correspondingly increased within the pores and {εAAO to {2.303, 2.282}. That is, the finding that these angles increased verified that the pores became filled with THF.29 Figure 3A also shows the R vs θ measurements after a 25 h polymerization step. It is seen that surface grafting of a PBLG layer on the internal pore surfaces of the AAO further increased the waveguide modes to higher θ. Analogous to the APTES deposition, since the dielectric constant of PBLG (εPBLG ) 2.403)16 is larger than εTHF, the addition of a PBLG layer on the pore surfaces resulted in the increase of {ε⊥AAO, ε|AAO}. Indeed, ⊥ , the extent of the surface grafting and the increases in {εAAO | εAAO} were of such a magnitude that, by the end of the experiment, an additional waveguide mode could be coupled into the AAO thin film at s-polarization (TE1). To monitor the polymerization process in real time, the angle position of the TM1 waveguide mode, was tracked in situ, and R vs θ scans were taken periodically to measure the increases ⊥ | in {εAAO , εAAO } (Figure 3B). The corresponding increase in

dPBLG, as estimated from eq 3 and related calculations, is also indicated in the figure. The total TM1 angle shift after 25 h surface-initiated polymerization was 4.6° and indicated the addition of a PBLG layer with an average thickness dPBLG ∼ 17 nm. This increase in dPBLG was corroborated by SEM characterization of the average pore size, which decreased to ∼40 nm at the end of the polymerization step (Figure 3B inset). The chemical nature and secondary structures of the surface grafted PBLG layer were also confirmed by IR spectroscopy measurements (see below). In Figure 3B, a rapid initial increase in the angle shift (∆θ ∼ 0.2°) was observed at the beginning of the surface-initiated polymerization process with the introduction of the 100 mM BLG-NCA solution. This indicated the adsorption of a layer of monomers about 1 nm thick. ∆θ and dPBLG then continued to increase slowly for about 2-3 h, followed by a time trace that increased in slope to a quasilinear region within 5-17 h. This increase in slope might have origins in both an increase in PBLG layer dielectric constant and an increase in the BLG-NCA polymerization rate, both of which could be accompanied by a PBLG conformational change from β-sheets to R-helices as the degree of polymerization (DPn) increases beyond the oligomer length.11 Such a DPn dependent conformational transition has also been demonstrated by several research groups,48-50 which might be attributed to favorable intramolecular hydrogen bonding in ordered R-helical structures. Independent of the PBLG preparation conditions (solution, vapor phase, or film cast), the coexistence of an oligomeric fraction (random coil or β-sheet) has been observed for 3 < DPn < 10, and for chains above DPn ) 27 (a thickness of ∼4 nm) PBLG was shown to adopt a mostly stable R-helix conformation.18,19,27,48 The anticipated conformational transition in the secondary structure was identified by ex situ IR spectroscopy characterization of the surface grafted PBLG. Figure 4 shows the IR scans for PBLG growth as a function of time (2, 5, 12, 17, 23, and 25 h). Separate samples were used for this series of ex situ IR measurementssBLG-NCA polymerization was allowed to proceed uninterrupted until the specified time, and the samples were rinsed in THF and dried in Ar before IR measurements. The backbone amide I and amide II bands of the three most frequently occurring conformations of polyglutamates, i.e. R-helix (amide I ∼ 1650 cm-1, amide II ∼ 1546-1550 cm-1), β-sheet (amide I ∼1625 or 1630 cm-1, amide II ∼ 1530 cm-1), and random coil (amide I ∼ 1658 or 1670 cm-1, amide II ∼

3184 J. Phys. Chem. B, Vol. 113, No. 10, 2009 1535 cm-1), appear at distinct positions in the IR spectrum.11,18,19,47 Absorbances and positions of the amide I and amide II peaks corresponding to the different conformers were obtained by deconvolution of the spectra around the amide I and amide II regions using Gauss-Lorentzian fitting. Note that contributions corresponding to the aromatic ring (assumed to be centered at 1498 and 1516 cm-1)10 were discounted when calculating absorbances from the amide II regions in the deconvolution procedure. Furthermore, the spectra shown in Figure 4 were normalized according to the peak intensity of the side-chain carbonyl stretch at 1730-1738 cm-1, which is considered to be invariant to the secondary structure.19 Thus Figure 4 only shows the relative contributions of the R-helix, β-sheet, and random coil peaks and not their absolute amounts. The IR spectrum after 2 h was relatively noisy, and it was difficult to quantify the different components of the secondary structure. However, it is clear from the rough positions of the peaks, corresponding to the amide I and II positions, that peptide material had been deposited. After 5 h, the characteristic amide I and amide II peaks became more visible and were centered around 1625 and 1530 cm-1, respectively, corresponding to mainly β-sheet conformation. There was also an additional shoulder at 1669 cm-1 around the amide I region, which might have indicated contributions from the disordered random coil conformation. The ratios between the random coil, β-sheet, and R-helix components of both the amide I and amide II peaks gave roughly 22% random coil, 50% β-sheet, and 28% R-helix content. (See Supporting Information Figure S2 for detailed views of the deconvoluted peaks.) Referring back to the OWS measurement (Figure 3B), dPBLG was approximately 4 nm at this stage of the polymerization, which corresponded to a degree of polymerization (DPn) of ∼20. The IR spectrum at 12 h was very similar to the spectrum at 5 h, except that the peak of the shoulder at 1669 cm-1 at 5 h appeared to have shifted to 1654 cm-1 at 12 h, which is indicative of the R-helical conformation. Indeed peak deconvulation showed that the percentage of secondary structures were 14% random coil, 44% β-sheet, and 42% R-helix content, which shows that random coil content decreased while the amount of R-helical chains increased. After 12 h dPBLG was approximately 10 nm, as measured by OWS, which corresponded to DPn ∼ 80 if a helix tilt angle of 37° was assumed. Previous observations indicated that a DPn > 10 is necessary for the onset of R-helix formation.11,51 Thus the mixture of conformations detected after 5 h and the increased R-helical fraction after 12 h are consistent with the reported relationship between a transition to the R-helical conformation and higher DPn. Further polymerization until 17 h (PBLG layer ∼14 nm, Figure 3B) resulted in a double maxima (1650 and 1626 cm-1) at the amide I position, representing clearly both the β-sheet and the R-helix peaks. At the same time, the amide II region appeared wider and indicated a higher intensity of the R-helix component. Peak deconvolution indicated that the β-sheet portion of PBLG decreased to roughly 40%, while the R-helical content correspondingly increased to 60%. Thus the IR measurements for 5-17 h, concurrent with the approximately linear increase in the average PBLG layer thickness (Figure 3B), indicated a continuing increase in the R-helical content as DPn increased and the R-helix conformation became more favorable. As the polymerization proceeded beyond 17 h, the rate of polymerization slowed and appeared to approach an asymptotic value, as indicated by a decrease in the slope of the time trace in Figure 3B. In Figure 4, the IR spectrum taken at 23 h, when the PBLG layer thickness was around 16 nm, shows two strong

Lau et al. peaks at the amide I positionsat around 1650 cm-1 (R-helix, 37%) and 1623 cm-1 (β-sheet, 63%)swhile the amide II position between 1590 and 1490 cm-1 showed one main peak at 1520 cm-1 (β-sheet) and two shoulder peaks at 1539 cm-1 (random coil) and 1554 cm-1 (R-helix). However, the IR spectrum at the end of the 25 h polymerization step, taken after rinsing the sample with copious amounts of anhydrous THF, shows that the majority of the PBLG material had the R-helix (∼70%) secondary structure (amide I at 1650 cm-1 and amide II at 1548 cm-1) with a minority amount of only β-sheet (∼30%) material (amide I at 1625 cm-1 and amide II at 1525 cm-1). That is, the 25 h spectrum was more similar to the 17 h spectrum than the 23 h data. Thus it appears that the higher β-sheet content and a small amount of random coil material observed for the 23 h sample was likely due to an accumulation of physisorbed β-sheet material deposited from the monomer solution that was then removed by THF rinsing at the end of the 25 h polymerization step. The minor β-sheet fractions still remaining at 25 h after THF rinsing might also be physisorbed material that were bound too strongly to be removed by THF rinsing alone (see below). Possible sources of the physisorbed components include side reactions mediated by solution impurities or self-polymerization. Note that, as mentioned above, the normalized IR spectra in Figure 4 only show the relative contributions from the different conformations and not their absolute amounts. As discussed above, the increase of the PBLG layer thickness/ polymerization rate slowed appreciably between polymerization times of 17 and 23 h and 14-16 nm of PBLG layer growth, which coincided with the high physisorbed content observed in the IR spectrum at 23 h. The exact mechanism leading to the decrease in polymerization rate is not entirely clear. At the end of the 25 h polymerization step, the diameter of the original ∼76 nm wide pores have decreased to only ∼40 nm (Figure 3B inset) and the cross-sectional area of the pores has correspondingly decreased by ∼70%. It may appear plausible that the decrease in the polymerization rate was correlated with a decrease in total monomer transport into the pores. However, the pore length of the AAO templates were chosen to be only ∼1 µm, which is much shorter than the diffusion length for the present AAO pore system (see the Supporting Information). On the other hand, previous investigations have found that physisorbed chains52 and accumulation of oligomer units from the monomer solution (side reactions in solution11) may physically block access to initiator and propagation sites on the surface and decrease the polymerization rate at later stages of PBLG surface grafting.18 To investigate whether the β-sheet content still remaining after in situ THF rinsing at 25 h consisted of physisorbed or grafted material, we performed an additional ex situ cleaning posttreatment on the 25 h sample. The post-treatment consisted of sonicating the PBLG-AAO sample in DCA/chloroform for 30 min and keeping it in anhydrous DMF for 24 h (see Experimental).18,19,47 Indeed, after the post-treatment, the amide I and amide II peaks attributed to the R-helix conformation were retained whereas the peaks attributed to β-sheet was removed (Figure 4). Therefore, purely surface grafted PBLG layers may be prepared within nanoporous AAO thin film templates may be prepared under the present reaction conditions by application of a cleaning post-treatment. Furthermore, we speculate that a significant portion of the β-sheet and random coil content detected earlier, after polymerization times of 5 h when the PBLG thickness had already increased substantially beyond the oligomer length to >4 nm, could also be attributed to the presence of physisorbed material.

PBLG Surface Grafting in Nanoporous Alumina

Figure 5. Average thickness of the PBLG layer grown within the pores of AAO films (samples with ∼60 and ∼76 nm pores) compared with the thickness grown on top of a flat SiO2 film. The corresponding curves are only for guiding the eye.

j w and M j n) and polydispersity index The molecular weight (M (PDI) of the grafted PBLG at the end of the 25 h polymerization process was also characterized by GPC, both before and after the cleaning post-treatment. Larger AAO templates prepared from bulk aluminum disks 5 cm in diameter, with pore diameter identical to the in situ AAO samples and 100 µm in thickness, were used for GPC analysis because large amounts of material was required for the chromatography measurements. PBLG from AAO templates that were not subjected to the additional postj n ∼ 30 900, and PDI ∼ 6.75, j w ∼ 208 700, M treatment had M j n ∼ 31 300, j w ∼ 183 600, M while post-treated PBLG films had M j w by and PDI ∼ 5.86. Thus, the post-treatment reduced M j ∼12% while Mn was essentially unchanged. Although the reduction in Mw appears to run contrary to the FT-IR analysis above, which showed the removal of the low DPn β-sheet and random coil components by the post-treatment, the modest j w could also have been attributed to sample difference in M variability, especially during preparation for GPC measurements j w measured (see the Experimental Section). Due to the smaller M upon post-treatment, the PDI of the PBLG was also reduced, although it was still much higher than PBLG polymerized in bulk solutions.8,11 The high PDI for the surface-initiated polymerization might be a result of reduced transport of monomers to the shorter chains within the PBLG brush. Furthermore, a j n ∼ 31 000 corresponds to a DPn ∼ 150. Given a theoretical M tilt angle of 37° with respect to the grafting surface23 (see discussion in the following section) and a monomer length of 0.15 nm for ordered PBLG R-helices,47 this DPn corresponded to a PBLG film thickness of ∼14 nm. The film thickness estimated from Mn compares well with the OWS measured final film thickness of ∼17 nm at 25 h. In order to further confirm the trend of the polymerization process within the nanoporous AAO film matrix, the BLG-NCA polymerization experiment was repeated on another AAO sample. To gain information on the extent to which the nanopore geometry had an influence on the rate of polymerization, the average pore diameter of the second AAO sample was chosen to be smaller at ∼60 nm. The experiment was also repeated on a flat substratesa 1060 nm thick SiO2 waveguide. The SiO2 film surface was functionalized with APTES using a procedure identical to that applied to the AAO films, and IR measurements confirmed that the APTES density on the SiO2 was similar to that on AAO films (see section III.3). Figure 5 compares the PBLG layer thickness increase with time on the two AAO samples and on the flat surface. It is seen

J. Phys. Chem. B, Vol. 113, No. 10, 2009 3185 that the basic growth trends for both AAO samples were similar, i.e. (i) an increase in the slope of the dPBLG vs time curve between 2 and 5 h, followed by (ii) a quasi-linear increase in dPBLG, and finally (iii) a decrease in the slope toward longer polymerization times. However, the time at which this final slow down in dPBLG appeared to occurred earlier for the AAO sample with ∼60 nm pores than for the sample with ∼76 nm pores. Also accompanying the earlier decrease in PBLG growth rate was a smaller final dPBLG ) 14 nm after ∼24 h polymerization. The effect of the nanoporous geometry was more clearly displayed by differences in the basic growth trend between polymerization on the AAO and planar SiO2 samples. Figure 5 shows that the transition to a higher slope in dPBLG vs time on the planar SiO2 appeared much more gradually and at a later time than on the AAO samples. In OWS, the sensitivity to changes in the ad-layer thickness on top of a waveguide is much lower than the sensitivity to changes in εAAO caused by pore processes within a nanoporous AAO matrix.29 Thus the transition point in dPBLG vs time on the planar SiO2 could only be specified broadly, around the period between 6 and 20 h. However, in the absence of pore confinement on the planar surface, there was no decrease in the rate of dPBLG increase throughout the polymerization process. Thus, dPBLG at 24 h on the planar sample actually exceeded that within the ∼76 nm pores. Together with the comparison between AAO samples in which smaller pores led to smaller final dPBLG and earlier saturation in the PBLG growth rate, the differences in dPBLG vs time trends between the AAO samples and planar SiO2 film suggested that the slow down in PBLG growth rate in the AAO samples was an effect of the nanoporous geometry of the AAO. Although the AAO and flat SiO2 substrates exhibited different kinetic regimes during the polymerization process, the overall PBLG layer growth rates averaged over 25 h were similar for all samples at ∼0.8 nm/h (Figure 5). Molecular transport in and out of the AAO pores was enabled by diffusion. Thus, the fact that there was essentially no difference in the overall grafting rates between the nanoporous and flat substrates implied that the polymerization process within the nanopores was not limited by diffusion of BLG-NCA monomers into pores in the present experiments. In any case, as discussed earlier, the diffusion length of the system is much larger than the ∼1 µm pore length of the AAO template (see the Supporting Information). The overall PBLG growth rate was also quite low. Therefore diffusion was an efficient transport mechanism for the BLGNCA monomers under the present reaction conditions. At the end of the polymerization step on the nanoporous AAO film matrix (25 h), we also performed an additional experiment in which the PBLG/AAO film was annealed at 120 °C for 2 h under vacuum. Consistent with previous reports,18,19,52 a high quality R-helical PBLG layer was obtained, as indicated by IR measurements showing high definition peaks at 1650 and 1546-1550 cm-1 assigned to the R-helix conformation (Figure 4: 25 h + ∆). This is further evidence that PBLG layers exhibiting properties consistent with previous reports, of bulk polymerized PBLG or PBLG grafted on flat surfaces, could be grafted by our present method within a nanoporous AAO matrix. 3. Orientation and Chain Density of PBLG. In addition to the conformation analysis, IR spectroscopy can be used to study the orientation and chain density of R-helical polypeptides.10,17,18,47 In an R-helix structure, the amide I transition moment is oriented roughly parallel to the backbone axis owing to intramolecular hydrogen bonding, while the amide II transition moment is roughly perpendicular to the backbone axis. In contrast, the orientations are reversed in a β-sheet

3186 J. Phys. Chem. B, Vol. 113, No. 10, 2009 structure, i.e. the amide I band is perpendicular to the backbone axis while the amide II band is parallel. Thus, the PBLG orientation in the film, for either the R-helix or the β-sheet content, can be characterized by comparison of the relative absorbances of the corresponding amide I and amide II bands. (The absorbances were normalized to the side chain carbonyl stretch at 1730-1738 cm-1, which is considered to be invariant to the backbone orientation.)19 Here we concentrate on the calculation of the average tilt angle of the R-helices from the surface normal (θ):10,23

(2SHStI + 1) AI )C ) AII (2SHStII + 1) 1 2 (3cos2 θ - 1) 2 C 1 2 (3cos2 θ - 1) 2

[ [

][ 21 (3cos γ - 1)] + 1 (4) ][ 21 (3cos γ - 1)] + 1 2

I

2

II

where A is the observed absorbance calculated from integrating each amide peak, SH and St are respectively the orientational order parameters of the helical axis and the transition moments, and the subscripts refer to the amide I and amide II peaks of the R-helix conformation. C is a proportionality constant that relates the intrinsic oscillator strengths of the amide I and amide II vibrational modes, and its reported value at 1.5 ( 0.2 has been calculated from the amide I/amide II absorbance ratio measured for a KBr pellet prepared with randomly oriented R-helical PBLG.47,53,54 Lastly, γi represent the angles between the helical axis and the transition moments of the amide I and amide II vibrations and were taken to be γΙ ) 39° and γΙΙ ) 75°, respectively.47,53 As an example, for polymerization within an AAO film with ∼76 nm pores in 100 mM BLG-NCA in THF, the average R-helix tilt angle after 25 h was found to be 36° with respect to the AAO pore surfaces. On the flat SiO2 substrate using the identical polymerization conditions, the tilt was 38° with respect to the flat surface. These results are in good agreement with reported values.18,23,47,55 For example, Wieringa47 obtained 32 ( 5° with respect to the substrate, while Chang and Frank18 reported values of less than 45° for vapor phase deposited PBLG films. Moreover, Enriquez and Samulski23 calculated a theoretical orientation of 37° for perfectly ordered helices arrayed on a surface. A rough estimate of the chain density of PBLG R-helices may also be determined from both the average R-helix tilt angle calculated above and the PBLG film thickness measured by OWS. It was assumed that the helices formed a close-packed array of uniform, tilted cylinders. Using a polymer density of 1.32 g/cm3,47 we calculated 4-5 chains/nm2. Here it is worthwhile to mention that the PBLG grafting density should be intimately related to the density of APTES surface-initiator molecules. Although we were not able to measure the molecular density of the 1-2 nm APTES layers, our measured PBLG grafting density was consistent with literature values of the chain density of gas phase deposited APTES, which is between 5 and 6 molecules/nm2 on a flat aluminum oxide surface.46,56 4. Anisotropic PBLG Nanostructures. The PBLG layer thickness increase shown in Figure 3B was based on fitting the OWS measurements of the anisotropic dielectric response of ⊥ | and εAAO , to a simple effective medium the AAO film, i.e. εAAO theory (EMT) approximation which assumed that grafting of the PBLG was uniform and conformal on the internal pore surfaces of the AAO (eqs 2 and 3). Moreover, the PBLG layer

Lau et al. was assumed to have an isotropic dielectric constant (εPBLG ) 2.403).16 For these assumptions, the modeled ε⊥AAO and ε|AAO have a persistent anisotropy based on the alignment of the cylindrical ⊥ | > εAAO for any AAO pores such that, according to eq 1A, εAAO volume fraction of PBLG inside the pores. However, close ⊥ | and εAAO (Figure examination of the measured values of εAAO ⊥ | 3B) showed that as polymerization proceeded, εAAO and εAAO | approached each other, and by ∼20 h, εAAO was even slightly larger than ε⊥AAO. Note however that the PBLG thickness (dPBLG) plotted in Figure 3B represents the best fit value when using the simple EMT approximation (eqs 2 and 3) to model the OWS measurements. In effect, the procedure simply fitted values for dPBLG that were most consistent with the average of the measured ⊥ | and εAAO . Hence the estimated dPBLG shown in Figure 3B εAAO was approximately correct, although the discrepancy in the ⊥ | and εAAO indicated an inconsistency with anisotropy of εAAO the assumption of a conformal, optically isotropic PBLG layer. Quantitatively, the R vs θ waveguide measurements (Figure 3B) indicated an increase of {ε⊥AAO, ε|AAO} from {2.303, 2.282} (i.e., ⊥ | ⊥ | , εAAO ) to {2.458, 2.460} (i.e., εAAO < εAAO ) at the end of εAAO polymerization at 25 h. For the other sample shown in Figure 5 with smaller ∼60 nm pores, ε|AAO also became closer in value ⊥ ⊥ | but remained slightly smaller ({εAAO , εAAO } increased to εAAO from {2.447, 2.422} to {2.545, 2.541}). This is expected from the lower absolute PBLG content within the AAO film (smaller Dpore and smaller dPBLG). The AAO oxide matrix is stable in THF at room temperature and pressure, thus its optical anisotropy arising from the oriented cylindrical pore structure is expected to remain constant throughout the experiment. Therefore the PBLG structure within the pores must have developed an anisotropy opposite in orientation, and large enough to balance (or even overcompensate) the optical anisotropy of the nanoporous AAO structure. At the same time, the intrinsic birefringence of R-helical PBLG is not expected to impart an optical anisotropy to the PBLG/ ⊥ | ∼ εAAO . The AAO composite film that could bring εAAO difference in dielectric responses of R-helical PBLG between components parallel and normal to the helical axis is ∼0.03 at visible wavelengths,16,57,58 while the anisotropy of the AAO film that was compensated for by the PBLG layer was ∼0.01 (Figure 3B). However, IR measurements (section III.2) indicated an average R-helix tilt angle of ∼36° against the pore surfaces of the cylindrical AAO pores running straight through the thickness of the AAO film. That is, the R-helices were oriented slightly closer to the internal pore surface than to the pore surface normal, which would lead to a slightly higher contribution of the intrinsic PBLG birefringence for the component of the dielectric response along pore surface (normal to the top film surface). Therefore the intrinsic PBLG birefringence, if it had any effect at all, would only be expected to enhance the optical ⊥ | > εAAO ) rather than reverse anisotropy of the AAO film (εAAO that anisotropy. On the other hand, a nanostructured rearrangement of PBLG material within the pores may generate a dielectric response of the magnitude required to explain the OWS measurement. In particular, a PBLG morphology within the pores that has elements aligned normal to the internal pore surfaces and parallel to the AAO top film surface, could impart an optical anisotropy within the pores such that the component of the pore dielectric response parallel to the AAO film surface | ) is larger than the component normal to the film surface (εpore ⊥ ). This anisotropy within the pores would be generated in (εpore a manner analogous to how the aligned cylindrical pores of the ⊥ | > εAAO . Consenative AAO nanostructure give rise to εAAO ⊥ | ) could quently, an anisotropy within the pores (εpore < εpore

PBLG Surface Grafting in Nanoporous Alumina

Figure 6. SEM cross-sections of the PBLG structures protruding normal off the pore surfaces after 24-25 h polymerization, for the samples with ∼76 (A) and ∼60 nm (B) average pore diameters.

balance the anisotropy of the AAO film to give the roughly isotropic dielectric response measured of the PBLG/AAO hybrid membranes after 25 h polymerization. As anticipated by the measured changes in optical anisotropy, SEM characterization of both the PBLG/AAO samples after 25 h polymerization (Figure 6) showed the existence of PBLG nanostructures that protruded perpendicular from the AAO pore walls. Identical nanostructures were observed for AAO samples regardless of whether they were rinsed only in THF or subjected to the DCA post-treatment discussed in section III.2 (see Supporting Information Figure S3). In Figure 6, the anisotropic PBLG structures were observed as ribbed undulations on the surfaces of the ∼76 nm pores (Figure 6A), and as filamentary/ lamellae features in the smaller, ∼60 nm wide pores (Figure 6B). Thus the filaments in the smaller pores appeared to be further developed structures of the same basic morphology as the structures in the larger pores. The morphology of the structures is also suggestive of solvent dewetting or bubble formation within the pores. However, Figure 6 shows that the protruding structures were distributed through the entire lengths of the AAO pores, and such an extensive accumulation of voids or a gaseous phase within the pores would ⊥ | , εAAO } that are drastically smaller than the have implied {εAAO observed values because εTHF is significantly larger than εgas ∼ ⊥ , εair ∼ εvacuum ) 1. For example, according to eqs 1A-3, {εAAO | εAAO} of the AAO sample with 76 nm pores with a PBLG layer 16 nm thick (Figure 3B) but otherwise with pores filled with air would be {2.358, 2.267}, which are significantly smaller than the measured values of {2.458, 2.460}. Thus the observed protruding structures were unlikely to be artifacts arising from dewetting effects. Other morphologies reminiscent of the protruding PBLG structures include those caused by Rayleigh instability observed when thin polymer pore coatings within nanoporous AAO are annealed above their glass transition temperatures (Tg),59 and PBLG fibrils observed on thick PBLG films grafted on flat substrates.18 We carried out PBLG surface grafting at room temperature, which is higher than Tg(PBLG) ∼12 °C.14 However, the dominant wavelength of the aforementioned Rayleigh instability should be on the order of 21/2πDpore ∼ 300 nm,59 which is much larger than for the present structures observed by SEM (∼50 nm, Figure 6). On the other hand, Chang and Frank observed fibril-like structures as thick films of PBLG (∼42 nm) were grafted, depending on the solvent chosen, on flat SiO2/Si and mica surfaces.18 The formation of such

J. Phys. Chem. B, Vol. 113, No. 10, 2009 3187

Figure 7. (A) Idealized 3-dimensional PBLG structure, with both a uniform pore coating and nanostructures protruding normal off the pore surfaces, on which the EMT model is based. (B) Best fit amounts of PBLG material contained in the uniform coating (open squares), as protruding anisotropic structures (open circles), and the summed amount (solid squares), expressed in terms of AAO pore volume fractions filled for the sample with ∼76 nm pores (Figure 3 and 6A). The corresponding curves are only for guiding the eye.

nanofibrils on solid substrates were attributed to three dominant effects: hydrophobic interaction between substrate and adsorbing (oligo)peptides; electrostatic/dipole attractions between substrate and PBLG molecules; and chain associations along the length of the fibers.18 Larger PBLG fibers with diameters ranging from 100 to 280 nm have also been prepared by the precipitation of PBLG from DMF solution.60,61 In our case, we also obtained fibrils from films of PBLG surface-grafted on flat SiO2/Si substrates, which had dPBLG thicker than the sample shown in Figure 5 (see Supporting Information Figure S4). Thus the ribbed/fibrillar PBLG material extending normal from the AAO pore surfaces observed for films with average thickness 5 h, the grafted material had a predominantly R-helical conformation while significant random coil and β-sheet content were physisorbed onto the PBLG layer. If the deposition of the physisorbed material had indeed induced a decrease in the polymerization rate at 17-23 h (see the discussion of section III.2), the temporally correlated normal slowing in the decrease of f PBLG might indicate that the protruding nanostructures observed in Figure 6 had consisted of surface grafted R-helical PBLG. Nonetheless, caution should be exercised in the interpreting the results of the anisotropic optical model, due the approximations involved and the limited data available in this first report of PBLG surface grafting within nanoporous AAO. IV. Conclusions Using complementary OWS, FT-IR, and SEM characterization, we were able to verify that surface-initiated BLG-NCA polymerization could be used to graft PBLG layers at a high density within the pores of AAO to obtain PBLG/AAO hybrid membranes. Thin films of PBLG were grafted on the AAO pore by surface-initiated polymerization of BLG-NCA monomer in THF solution. Complementary OWS and FT-IR measurements

showed that BLG-NCA polymerization within the nanoporous AAO films proceeded according to 3 stages that correlated with transitions in the proportions of R-helical, β-sheet, and random coil secondary structures. In particular, a higher rate of PBLG film growth appeared to be accompanied by a higher R-helical content. The average R-helix tilt angle and the average PBLG chain density were also characterized by IR spectroscopy. In addition, in situ OWS analysis was able to discern the emergence of a three-dimensional PBLG morphology toward high degrees of polymerization, and complementary SEM microscopy showed fibril-like PBLG nanostructures protruding perpendicular from the AAO pore surfaces. Comparison of the PBLG/AAO results with a parallel experiment performed on a planar SiO2 surface suggested that both the staged process of BLG-NCA polymerization and the development of PBLG three-dimensional nanostructures were intimately linked to the confined geometry of the nanoporous AAO matrix. The present investigation of the PBLG film growth process within a nanoporous matrix, and with the aid of an in situ methodology such as OWS, is unprecedented. The ability of the nondestructive OWS technique to anticipate nanostructured PBLG morphologies within the nanoporous AAO also demonstrated the technique’s utility and versatility. Furthermore, FTIR measurements verified that R-helical PBLG layers on the AAO pore walls with high orientational order could be obtained by thorough cleaning of the PBLG/AAO sample to remove any oligomeric peptide material physisorbed on the as-polymerized high-density PBLG layers. Thus the use of a nanoporous AAO matrix, combined with the surface grafting of PBLG chains with multiple modifiable side chains, may offer a substrate platform with a potentially very high density of functional sites for, e.g. high sensitivity biosensing applications. Moreover, the development of fibril-like PBLG nanostructures within the PBLG/AAO nanoporous membranes could potentially lead to important pharmaceutical applications in the extraction and chiral separation of drug analytes by membrane filtration or sequestration. Acknowledgment. H.D. thanks the European Union, Mari Curie Intra-European Fellowship (MEIF-CT-2005-024731), for financial support. We thank Gunnar Glasser and Juergen Thiel at the Max Planck Institute for Polymer Research, for SEM analysis and the preparation of anhydrous THF, respectively. We also gratefully acknowledge Stefan Schmidt and Dr. Peter Detemple at the Institut fu¨r Mikrotechnik Mainz for providing Al thin film deposition services and Dr. M. Steinhart at the Max Planck Institute of Microstructure Physics for providing large area porous alumina templates. Supporting Information Available: Characterizations of the SiO2 waveguide surface roughness, the deconvulated IR spectrum at 5 h, PBLG fibrillar structures on SiO2, and a discussion on the diffusion length. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Lee, N. H.; Frank, C. W. Polymer 2002, 43, 6255–6262. (2) Naijo, Y.; Ikeda, K.; Tsunoda, M.; Tanaka, T.; Kobayashi, N. Mol. Cryst. Liq. Cryst. 2003, 406, 1–1. (3) Ciardelli, F.; Pieroni, O. Photoswitchable Polypeptides. In Molecular Switches; Ben, L. F., Ed.; Wiley-VCH Verlag: Berlin, 2001; pp 399-441. (4) Machida, S.; Urano, T. I.; Sano, K.; Kato, T. Langmuir 1997, 13, 576–580. (5) Whitesell, J. K.; Chang, H. K. Mol. Cryst. Liq. Cryst. 1994, 240, 251–258. (6) Duran, H.; Lau, K. H. A.; Lu¨bbert, A.; Jonas, U.; Steinhart, M.; Knoll, W. Biopolymers for Biosensors: Polypeptide Nanotubes for Optical Biosensing. In Polymers for Biomedical Applications; Mahapatro, A.,

PBLG Surface Grafting in Nanoporous Alumina Kulshrestha, A. S., Eds.; Oxford University Press: New York, 2008; Vol. 977, pp 371-390. (7) Duran, H.; Lau, K. H. A.; Jonas, U.; Knoll, W. PMSE Prepr. 2006, 51, 345–347. (8) Deming, T. J. AdV. Mater. 1997, 9, 299–311. (9) Lesk, A. Introduction to Protein Science: Architecture, Function, and Genomics; Oxford University Press: New York, 2004. (10) Miura, Y.; Kimura, S.; Imanishi, Y.; Umemura, J. Langmuir 1998, 14, 6935–6940. (11) Block, H. Poly (gamma-Benzyl-L-Glutamate) and Other Glutamic Acid Containing Polymers; Gordon & Breach Science Publishers: New York, 1983. (12) Lin, J.; Liu, N.; Chen, J.; Zhou, D. Polymer 2000, 41, 6189–6194. (13) Gitsas, A.; Floudas, G.; Dietz, M.; Mondeshki, M.; Spiess, H. W.; Wegner, G. Macromolecules 2007, 40, 8311–8322. (14) Papadopoulos, P.; Floudas, G.; Klok, H. A.; Schnell, I.; Pakula, T. Biomacromolecules 2004, 5, 81–91. (15) Hartmann, L.; Kratzmu¨ller, T.; Braun, H.-G.; Kremer, F. Macromol. Rapid Commun. 2000, 21, 814–819. (16) Cassim, J. Y.; Taylor, E. W. Biophys. J. 1965, 5, 531–552. (17) Wang, Y.; Chang, Y. C. AdV. Mater. 2003, 15, 290–293. (18) Chang, Y. C.; Frank, C. W. Langmuir 1998, 14, 326–334. (19) Chang, Y. C.; Frank, C. W. Langmuir 1996, 12, 5824–5829. (20) Wieringa, R. H.; Schouten, A. J. Macromolecules 1996, 29, 3032– 3034. (21) Whitesell, J. K.; Chang, H. K. Science 1993, 261, 73–76. (22) Machida, S.; Urano, T. I.; Sano, K.; Kawata, Y.; Sunohara, K.; Sasaki, H.; Yoshiki, M.; Mori, Y. Langmuir 1995, 11, 4838–4843. (23) Enriquez, E. P.; Samulski, E. T. Self-assembled R-helical polypeptide films. Mater. Res. Soc. Symp. Proc. 1992, 255, 423-434. (24) Smuleac, V.; Butterfield, D. A.; Bhattacharyya, D. Chem. Mater. 2004, 16, 2762–2771. (25) Flory, P. J.; Leonard, W. J. J. Am. Chem. Soc. 1965, 87, 2102– 2108. (26) Enriquez, E. P.; Gray, K. H.; Guarisco, V. F.; Linton, R. W.; Mar, K. D.; Samulski, E. T. Behavior of rigid macromolecules in self-assembly at an interface. In 38th National Symposium of the American Vacuum Society, 4th ed.; AVS: Seattle, WA, 1992; Vol. 10; pp 2775-2782. (27) Heise, A.; Menzel, H.; Yim, H.; Foster, M. D.; Wieringa, R. H.; Schouten, A. J.; Erb, V.; Stamm, M. Langmuir 1997, 13, 723–728. (28) Cameron, P. J.; Jenkins, A. T. A.; Knoll, W.; Marken, F.; Milsom, E. V.; Williams, T. L. J. Mater. Chem. 2008, 18, 4304–4310. (29) Lau, K. H. A.; Tan, L. S.; Tamada, K.; Sander, M. S.; Knoll, W. J. Phys. Chem. B 2004, 108, 10812–10818. (30) Yu, F.; Persson, B.; Lofas, S.; Knoll, W. J. Am. Chem. Soc. 2004, 126, 8902–8903. (31) O’Sullivan, J. P.; Wood, G. C. Proc. R. Soc. London, Series A 1970, 317, 511–543.

J. Phys. Chem. B, Vol. 113, No. 10, 2009 3189 (32) Li, F.; Zhang, L.; Metzger, R. M. Chem. Mater. 1998, 10, 2470– 2480. (33) Martin, C. R. Science 1994, 266, 1961–1966. (34) Lei, Y.; Cai, W.; Wilde, G. Prog. Mater. Sci. 2007, 52, 465–539. (35) Kriha, O.; Zhao, L.; Pippel, E.; Go¨sele, U.; Wehrspohn, R. B.; Wendorff, J. H.; Steinhart, M.; Greiner, A. AdV. Funct. Mater. 2007, 17, 1327–1332. (36) Mikulskas, I.; Juodkazis, S.; Tomasiunas, R.; Dumas, J. G. AdV. Mater. 2001, 13, 1574–1577. (37) Jonas, U.; Krueger, C. J. Supramol. Chem. 2002, 2, 255–270. (38) Knoll, W. Annu. ReV. Phys. Chem. 1998, 49, 569–638. (39) Raether, H. Surface-Plasmons on Smooth and Rough Surfaces and on Gratings; Springer: Berlin/Heidelberg, 1988; Vol. 111. (40) Choy, T. C. EffectiVe Medium Theory: Principles and Applications, 1st ed.; Oxford University Press: New York, 1999; Vol. 102. (41) Aspnes, D. E. Thin Solid Films 1982, 89, 249–262. (42) Damle, S. B. Chem. Eng. News 1993, 71, 4. (43) Masuda, H.; Satoh, M. Jpn. J. Appl. Phys 1996, 35, L126-L129. (44) Kambhampati, D.; Nielsen, P. E.; Knoll, W. Biosens. Bioelectron. 2001, 16, 1109–1118. (45) Foss, C. A.; Hornyak, G. L.; Stockert, J. A.; Martin, C. R. J. Phys. Chem. 1992, 96, 7497–7499. (46) Kurth, D. G.; Bein, T. Langmuir 1995, 11, 3061–3067. (47) Wieringa, R. H.; Siesling, E. A.; Werkman, P. J.; Angerman, H. J.; Vorenkamp, E. J.; Schouten, A. J. Langmuir 2001, 17, 6485–6490. (48) Rinaudo, M.; Domard, A. Biopolymers 1976, 15, 2185–2199. (49) Sisido, M.; Imanishi, Y.; Okamura, S. Biopolymers 1969, 7, 937– 947. (50) Lundberg, R. D.; Doty, P. J. Am. Chem. Soc. 1957, 79, 3961– 3972. (51) Bradbury, E. M.; Crane-Robinson, C.; Goldman, H.; Rattle, H. W. E. Nature 1968, 217, 812–816. (52) Chang, Y. C.; Frank, C. W.; Forstmann, G. G.; Johannsmann, D. J. Chem. Phys. 1999, 111, 6136–6143. (53) Niwa, M.; Morikawa, M.-a.; Higashi, N. Angew. Chem., Int. Ed. 2000, 39, 960–963. (54) Tsubio, M. J. Polym. Sci. 1962, 59, 139–153. (55) Jaworek, T.; Neher, D.; Wegner, G.; Wieringa, R. H.; Schouten, A. J. Science 1998, 279, 57–60. (56) Kurth, D. G.; Bein, T. Langmuir 1993, 9, 2965–2973. (57) de Campos Vidal, B. Cell. Mol. Biol. 1986, 32, 109–112. (58) References 16 and 56 report the anisotropy in terms of refractive indices. The corresponding value in dielectric constants is ∆ε ) 2∆n〈n〉, where 〈n〉 is the average refractive index. (59) Chen, J. T.; Zhang, M.; Russell, T. P. Nano Lett. 2007, 7, 183– 187. (60) Ishikawa, S.; Kurita, T. Biopolymers 1964, 2, 381–393. (61) Rybnikar, F.; Geil, P. H. Biopolymers 1972, 11, 271–278.

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