Surface Orientation of Polystyrene Based Polymers: Steric Effects from

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Surface Orientation of Polystyrene Based Polymers: Steric Effects from Pendant Groups on the Phenyl Ring Joseph L. Lenhart,*,†,§ Daniel A. Fischer,‡ Tanya L. Chantawansri,† and Jan W. Andzelm† †

U.S. Army Research Laboratory, 4600 Deer Creek Loop, Aberdeen Proving Ground, Maryland 21005, United States National Institute of Standards and Technology, 100 Bureau Drive, Gaithersburg, Maryland 20899, United States § Sandia National Laboratories, Albuquerque, New Mexico 87185, United States ‡

S Supporting Information *

ABSTRACT: Near edge X-ray absorption fine structure (NEXAFS) coupled with molecular dynamics simulations were utilized to probe the orientation at the exposed surface of the polymer film for polystyrene type polymers with various pendant functional groups off the phenyl ring. For all the polymers, the surface was oriented so that the rings are nominally normal to the film surface and pointing outward from the surface. The magnitude of this orientation was small and dependent on the size of the pendant functional group. Bulky functional groups hindered the surface orientation, leading to nearly unoriented surfaces. Depth dependent NEXAFS measurements demonstrated that the surface orientation was localized near the interface. Molecular dynamics simulations showed that the phenyl rings were not oriented strongly around a particular “average tilt angle”. In contrast, simulations demonstrate that the phenyl rings exhibit a broad distribution of tilt angles, and that changes in the tilt angle distribution with pendant functionality give rise to the observed NEXAFS response. The more oriented samples exhibit a higher probability of phenyl ring orientation at angles greater than 60 degrees relative to the plane of the films surface. transform infrared spectroscopy,28,29 neutron and X-ray reflectivity, 30−33 fluorescence spectroscopy, 34,35 contact angle,36,37 and secondary ion and time-of-flight mass spectroscopy.38,39 In addition, techniques are continually emerging to probe organic material interfacial region, including second harmonic and sum frequency generation spectroscopy,40,41 angle resolved ultraviolet photoelectron spectroscopy,42 various X-ray scattering techniques,43,44 metastable-atom electron spectroscopy,45 and near edge X-ray absorption fine structure (NEXAFS).46,47 NEXAFS provides two unique measurement capabilities that complement these other techniques including (1) a high intensity synchrotron source allowing for monochromatic incident excitation and enabling enhanced peak resolution, and (2) polarized excitation enabling investigation of bond orientation. NEXAFS has proven useful at monitoring organization of self-assembled monolayers,48,49 interfacial composition in lithographic polymers,8,50−52 chemistry in polymeric materials,53,54 chemical modification of polymeric surfaces,55 as well as the near surface composition and orientation of block copolymers.56−59 In this study, NEXAFS is utilized to probe the surface chemistry and orientation for a series of spun cast polymeric films after annealing above the polymer glass transition temperature. The polymer films were polystyrene based, with pendant groups of gradually varying size

1. INTRODUCTION Polymeric thin films and coatings are utilized extensively in many areas including the microelectronics industry,1 automobile manufacturing, paints, and protective coatings,2 organic and molecular electronics,3 release coatings,4 biotechnology,5 sensors,6 and display devices,7 among others. With current trends in device miniaturization in addition to increasingly complex material formulations, the interfacial properties of polymers can play a critical role in device performance as well as the performance of many general materials systems such as adhesives, encapsulants, foams, filled polymers, and coatings. Polymer interfacial properties can be very different than their bulk properties, due to variations in chemical composition,8 molecular orientation,9 mobility,10 crystallinity, and microstructure11 near the interface. Specific examples include changes in the glass transition temperature and coefficients of thermal expansion for thin polymeric films,12−18 surface segregation of low surface tension species toward the interface,19−21 segregation of low molecular weight species and chain ends near interfaces,22−24 and wetting changes for polymers on surfaces that have various chemical functionality.25 To understand and control reliability and aging issues associated with polymer interfaces, it is important to characterize the properties, structure, and composition of polymers near the interface. Numerous techniques have been exploited to monitor the chemistry, structure, and properties of polymeric surfaces and films. Some of the most common techniques include X-ray photoelectron and Auger electron spectroscopy,26,27 Fourier This article not subject to U.S. Copyright. Published 2012 by the American Chemical Society

Received: August 30, 2012 Revised: October 18, 2012 Published: October 18, 2012 15713

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the edge jump at 315 eV, thereby removing the spectral dependence on the total number of atoms probed, as well as the incident X-ray beam footprint on the sample surface. Therefore, changes in the post edge jump normalized spectra represent real changes in chemical bonding, functional group density, or orientation. 2.3. Molecular Dynamics Simulations. Molecular dynamics (MD), as implemented in the Discover module of Materials Studio,65 was used to study the orientation of the phenyl ring at the surface of the thin film for numerous polystyrene-type polymers. Interactions between molecules are described using the COMPASS (Condensed-phase Optimized Molecular Potentials for Atomistic Simulation Studies) force field.66,67 The COMPASS force field was developed to predict structural and thermophysical properties for many inorganic and organic fluids and also polymers.68,69 The potential energy expressions of COMPASS contain valence and nonbonded terms. The nonbonded terms include a Coulombic term for electrostatic interactions and a Lennard−Jones function describing van der Waals interactions. The valence terms include diagonal and cross-coupling functions representing bond stretching, bond angle bending, and dihedral angle distortions. The nonbonded terms were parametrized to reproduce experimental data for polymers and liquids such as heats of vaporization, while the bonded, valence term parameters were obtained from fits to quantum mechanical results for molecular structures and vibrational frequencies. Initially, the bulk and slab models of polymers were constructed with the Amorphous Cell program of Materials Studio.65 This program uses a Monte Carlo technique to build an amorphous structure as a three- or twodimensional periodic cell.70 In order to avoid generation of high-energy configurations, the “lookahead” feature of the Amorphous Cell program was used with a standard value of 5. Every simulated system was optimized using eight amorphous cell structures, and the average values are reported. After each periodic cell was created, it was energy minimized followed by an extensive annealing procedure to equilibrate the system. Constant volume and temperature (NVT) simulations were performed on our simulation cells. The nonbonded energies were calculated using the group-based summation method instead of the popular Ewald summation. Although both methods are known to produce similar results, the group-based method is more computationally efficient.71 We utilized the predefined medium quality setting corresponding to a cutoff, spline width, and buffer width of 9.5, 1, 0.5 Å, respectively, where a tail correction was applied to the nonbonded interactions. The temperature was controlled using the Andersen thermostat algorithm72 with a collision ratio of 0.5. The following steps were taken to build the polymer system: (1) Use the homopolymer builder tool to build polymers of PS, and PTBS, each with 10 repeat monomers. 2) Build a confined simulation cell of the polymers at the experimental densities of 1.048 and 0.947 g/cm3 for PS and PTBS, respectively, using the amorphous cell builder where each cell is composed of 10 atactic chains. Since we are considering a thin film, we will set the x,y dimension larger than z, where x and y are at least 2 times the cutoff distance of 9.5 Å. Instead of using experimental densities, we could have also calculate the density using NPT (constant particle, pressure, and temperature) simulations, but our previous tests with PS have shown that COMPASS predicts density within 2% of experimental values.73 (3) Minimize the structure using the conjugate gradient method to remove unphysical overlap of the molecules for a maximum of 35 000 steps. (4) Anneal the structure by ramping the temperature from 300 to 800 to 300 K over 300 ps followed by additional minimization; repeat five times. A strong repulsive potential was used at both the top and bottom of the box to ensure that the polymer thin film remains confined during the annealing process and that periodic replicas in the z direction are not observed. (5) Remove the potential and increase the vacuum length in the z direction beyond the cutoff so that the polymer does not see its periodic image. Run NVT simulations at 25 °C for an additional 1 ns to alleviate any orientational bias of the phenyl ring due to the prescribed potential. (6) Calculate the orientation of the styrene ring relative to the x−y plane using a written Fortran90 program for a given film thickness using the final snapshot. We took care to determine the appropriate thickness of the film over which to analyze the MD simulations. Initially we choose several film thicknesses at the film surface (including 5, 10, and 15 Å) that were

and chemical functionality off the phenyl ring. The purpose of this work is to investigate the impact of pendant group size, position, and polarity on polymer surface composition and orientation at the exposed interface for these films.

2. EXPERIMENTAL SECTION 2.1. Materials and Methods. Polystyrene [Mw = 102 500; Mw/Mn = 1.05], poly(α-methylstyrene) [Mw = 100 300; Mw/Mn = 1.05], poly(4tert-butylstyrene) [Mw = 132 000; Mw/Mn = 1.04], poly(4-methylstyrene) [Mw = 82 500; Mw/Mn = 1.09], and poly(4-hydroxystyrene) [Mw = 8800; Mw/Mn = 1.07] were utilized as received from Polymer Source, Inc. (Montreal, Canada).60 All the polymers were dissolved in high pressure liquid chromatography grade toluene at 3 mass % except for poly(4-hydroxystyrene), which was dissolved in propylene glycol monomethyl ether acetate. The polymeric films were spun cast onto silicone wafers at 1500 rpm for 2 min, under a vapor pressure of the spincasting solvent. The samples were then annealed under vacuum at 20 °C above the polymer glass transition temperature for 20 h. Previous work has shown that the thickness of annealed PS films with similar spincasting solution concentrations, molecular weights, and spinning speeds is approximately 150 nm.55 Therefore, the thickness of the films in this study is substantially larger than the NEXAFS probing depth (1−10 nm). In addition, these films are thick enough to exhibit bulk behavior, as quantified by recent work examining the glass transition of polymeric films as a function of film thickness.12−18 For annealing purposes, the glass transition temperature of polystyrene was 102 °C,18 poly(4methylstyrene) was 108 °C,61 poly(4-tert-butylstyrene) was 147 °C,62 poly(α-methylstyrene) was 170 °C,63,64 and polyhydroxystyrene was 150 °C.17 2.2. Near-Edge X-ray Absorption Fine Structure (NEXAFS) Spectroscopy. NEXAFS measurements were conducted at the National Institute of Standards and Technology (NIST) U7A beamline of the National Synchrotron Light Source at Brookhaven National Laboratory. A monochromator with 600 line/mm grating, providing ±0.15 eV resolution, was used for all NEXAFS experiments. The monochromator energy scale was calibrated by the carbon K-edge π* transition of graphite at 285.35 eV. All the spectra were recorded at room temperature in the NIST − Dow material characterization chamber (beamline U7A) at 10−6 Pa. The spectra were normalized to the incident beam intensity, I0, by collecting the total electron yield intensity from a gold coated 90% transmitting grid placed in the incoming X-ray beam path. Surface sensitive partial electron yield measurements were made (probe depth of approximately 1−6 nm) by applying a negative bias on the entrance grid of the channeltron electron detector. The bias was changed to selectively probe different depths into the polymer film. In order to investigate orientation of the polymer films, the spectra were collected with the incident angle ranging from 20° (glancing) to 90° (normal) relative to the film surface. Since the polymer films are not strongly oriented like self-assembled monolayers, for example, it is important to recognize that the NEXAFS spectra do not exhibit dramatic changes. This effect coupled with surface roughness of the spun cast films (∼5 Å) and imperfect polarization of the excitation beam (∼90% polarize) means that the changes in the NEXAFS spectra as a function of angle will be subtle, typically ranging from 10 to 30%. For the NEXAFS spectra in this paper, the experimental standard uncertainty in the peak position is similar to the grating resolution of ±0.15 eV. The relative uncertainty in the NEXAFS intensity was less than ±3%, as determined by multiple scans on selected samples. More detailed assessment of data uncertainty is provided throughout the paper during the appropriate discussion points for each figure. Figure S1 (in the Supporting Information) presents a schematic of the principles of NEXAFS and shows a characteristic carbon edge partial electron yield NEXAFS spectrum for polystyrene. The NEXAFS data in this paper are normalized according to the description in the Supporting Information. Specifically, the NEXAFS spectra in this paper are pre-edge jump baselined, by subtracting background signal from the spectrum. In addition, most of the spectra are also post-edge jump normalized, by normalizing the NEXAFS signal (at each energy) by the edge jump intensity. The post edge jump normalization was done by quantifying 15714

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capable of encompassing an entire phenyl ring in the x−y plane, but were still an adequate distance away from the center of the simulation cell where the phenyl rings will be oriented as in the bulk. It should also be noted that the film is no longer flat in the z dimension after the system is allowed to reorient itself without the repulsive potential in step 5. It is instead modulated, so the thickness cannot simply be the diameter of the phenyl ring. We observed similar behavior for the MD simulations analyzed over these different film probing depths; therefore, we will only report results for a 10 Å film thickness. As will be demonstrated in the Results and Discussion section of the paper, at the electron yield detector bias utilized in these experiments, 50% of the NEXAFS electron yield signal comes from the top 10 Å of the film surface. Therefore, the analysis depth of 10 Å for the MD simulations is suitable for qualitative comparison with experimental results. With the MD simulations, we considered a total of 8 replicas for each polymer to obtain a statistical average. We have confirmed that enlarging the film thickness does not affect conclusions of this simulation. Representative surfaces for PS and PTBS are shown in Figure 1, where phenyl groups at surface are

Figure 2. Structures of the various polymers utilized in this study.

impact of polar functional groups on surface organization. The functional group for all the polymers was in the para-position off the ring, with the exception of poly(α-methylstyrene) [PAMS], which contains the methyl group off the chain backbone. Because the incident synchrotron radiation is polarized, NEXAFS can be utilized to probe molecular orientation. Figure 3 shows a schematic of a polystyrene monomer and highlights

Figure 3. Schematic of a polystyrene molecule and the various orbital transitions that can be probed with NEXAFS. θ indicates the angle of excitation relative to the plane of the film surface.

Figure 1. Representative surfaces for PS (a) and PTBS (b) films, where phenyl groups at the surface are highlighted.

the various orbital transitions that can be detected with NEXAFS. The C1s → σ*C−C transition (denoted by C−C σ*) is oriented with the polymer chain backbone. The C1s → σ*C−H transition (denoted by C−H σ*) is oriented perpendicular to the chain backbone. A pendant phenyl ring is normal to the chain backbone and contains a C1s → π*CC transition (denoted by CC π*) that is perpendicular to the phenyl group and parallel to the chain backbone. Since the incident X-ray radiation is polarized, with the electric field vector perpendicular to the direction of the light propagation, measuring the NEXAFS spectra at different angles can be utilized to investigate the average orientation of these absorption transitions. Figure 4 shows the NEXAFS electron yield spectra for a PS film, measured in both the normal and glancing mode. The PS sample exhibits a pronounced C1s → π*CC transition at 285.1

highlighted. To calculate the angle of the phenyl ring relative to the film surface, the angle was calculated between two vectors: the vector encompassing the minimum and maximum coordinates for the carbons in the phenyl group, and the z axis represented by the plane of the film surface.

3. RESULTS AND DISCUSSION Figure 2 shows the structures of the polymers utilized in this study. The size of the nonpolar pendant group was gradually varied from -hydrogen [polystyrene, PS], to -methyl [poly(4methylstyrene), P4MS], to -tert-butyl [poly(4-tertbutylstyrene), PTBS], in order to investigate the impact of pendant group size on the surface orientation of the rings. In addition, a -hydroxyl group [poly(4-hydroxystyrene), PHS] was utilized to probe the 15715

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report, we will provide clear evidence of the outward orientation of the rings by measuring the NEXAFS spectra at different probing depths. Consistent with the absorption at 285.1 eV, a second π* transition, C1s → π2*CC (denoted by C−C π2*), occurs at 288.9 eV and is slightly higher the normal mode than the glancing mode. Other transitions including the C1s → σ*C−H transition at 287.4 eV, the C1s → σ*C−C transition (labeled C−C σ*) at 293.6 eV, and the C1s → σ2*CC transition near 300 eV change in small amounts that are consistent with orientation of the phenyl ring (see Figure S2 of the Supporting Information for an expanded version of Figure 4 demonstrating these peak areas). These peaks that exhibit smaller changes are not adequate to quantify the polymer orientation within the measurement error. In general, π* transitions are sharp in the NEXAFS spectra, while σ* transitions are much broader, due to the impact of nuclear vibrations and the orientation of the transitions relative to the nuclear alignment in the chemical bonds.74 Since the dominant transition that exhibits a change between normal and glancing measurements, with these poly(styrene) based polymers, is the C1s → π*CC transition at 285.1 eV, this will be the focus for the remainder of the paper. In Figure 4, the electron yield spectra are measured at a high negative voltage bias (250 eV) on the electron yield detector entrance grid. This means that only electrons with energy greater than 250 eV will be detected. As discussed in the Experimental Section and in the Supporting Information, upon photon absorption, the molecules can release either an Auger electron or a photon to dissipate the excitation energy. Atoms with low atomic numbers, such as carbon, primarily release electrons. The electrons emitted from carbon are Auger electrons from the carbon KLL transition, which have an energy of 263 eV. As these Auger electrons travel through the film, they can lose energy due to inelastic scattering mechanism, which can also cause the emission of additional lower energy electrons. The electron yield NEXAFS measurements collect both these carbon Auger electrons as well as lower energy electrons from inelastic scattering processes, depending on a retarding voltage bias on the electron yield detector. Due to the retarding voltage bias of 250 eV, the predominant electrons collected in the spectra in Figure 4

Figure 4. Pre- and post-edge jump normalized NEXAFS partial electron yield spectra for a polystyrene film measured in both the normal and glancing mode at a electron yield entrance grid negative voltage bias of 250 eV over an energy range of 280 to 310 eV, with the inset showing an expansion of the C1s → π*CC transition at 285.1 eV. Figure S2 in the Supporting Information shows details of other absorption peaks.

eV and is more intense when measured in the normal mode than the glancing mode, indicating that the phenyl rings are nominally oriented perpendicular to the film surface, since the electric field vector is perpendicular to the light propagation. The inset in Figure 4 highlights this region. While the schematic in Figure 3 shows the phenyl rings perpendicular to the surface and pointed outward, the difference in the C1s → π*CC transition between normal and glancing mode NEXAFS spectra (shown in Figure 4) only indicates that the rings are nominally perpendicular to the surface, and cannot distinguish between inward or outward orientation of the rings relative to the film surface. Later in this

Figure 5. (a) Peak areas for the C1s → π*CC transition NEXAFS spectra (at 250 eV negative voltage bias) as a function of angle for polystyrene (PS), poly(4-methylstyrene) (P4MS), and poly(4-tert-butylstyrene) (PTBS). (b) Peak areas for the C1s → π*CC transition NEXAFS spectra (at 250 eV negative voltage bias) as a function of angle for poly(α-methyl styrene) (PAMS), poly(4-methylstyrene) (P4MS), and poly(4-hydroxystyrene) (PHS). (c) Peak areas for the C1s → π*CC transition NEXAFS spectra (at 250 eV negative voltage bias) normalized with the fraction of carbon atoms with aromatic character in PS, P4MS, and PTBS monomers. The repeat data points represent a second measurement from the sample. The uncertainty in the peak area is less than ±3% as determined by multiple measurements. The data scatter is an additional indicator of the data uncertainty. Lines are best fits for the data with slope and intercept for best fit line indicated in the figure. In addition, Table S1 in the Supporting Information presents a range of slope and intercept values that can adequately fit the data points. 15716

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group on the chain backbone hinders the surface orientation of the phenyl rings. The slope of the lines in Figure 5 can potentially be utilized to quantify an average tilt angle. However, we find the slope changes better suited for assessing qualitative changes in orientation for several reasons. First, spun-cast polymer films have a roughness of approximately 5 Å.77 Even with perfectly oriented phenyl rings, this surface roughness will decrease the angular dependence of the C1s → π*CC transition, leading to a skewed assessment of average tilt angle. Second, the incident X-ray beam is not perfectly polarized (∼90 polarization), leading to a further decrease in the observed slope of π* peak area verses angle. In addition, the calculation of a tilt angle implies strong orientation of the phenyl rings in a narrow distribution around the average tilt angle. As molecular dynamics simulations will clearly show later in the paper, the phenyl groups exhibit a broad distribution of tilt angles. It is these changes in tilt angle distribution that give rise to the observed NEXAFS angular dependent data. Given these factors that convolute the quantitative interpretation of an average tilt angle, the slope data in Figure 5 (and Table S1 in the Supporting Information) presents a relative representation of the extent of surface phenyl orientation in these polymers. The changes in the C1s → π*CC transition as a function of angle in from Figure 5 are subtle, and smaller than those for highly organized aromatic-based self-assembled monolayers deposited on smooth gold surfaces.78 However, the angle dependent signal changes from the C1s → π*CC transition in Figure 5 are comparable to the changes observed in rubbed polystyrene films.79 With the experiments on rubbed films performed by Liu et al.,79 the polystyrene films were rubbed with a velour cloth, which imparts orientation to the phenyl rings on the film surface. This orientation can be quantified by the dichroic ratio (DR) defined as (Ipl − Ipd)/(Ipl + Ipd), where Ipl and Ipd are the C1s → π*CC transition intensity for the polystyrene films measured parallel and perpendicular to the rubbing direction, respectively. The DRs observed by Liu et al.79 with rubbed PS films ranged from 0.1 to 0.15. If we define a DR from Figure 5 as (I90 − I0)/(I90 + I0), where I90 and I0 are the C1s → π*CC transition intensity for the polystyrene film extrapolated to a measurement angle of 90° and 0°, respectively, we get DR values of approximately 0.11, similar to the DR values observed by Liu et al.79 in rubbed polystyrene films. Therefore, the spectral changes observed in Figure 5 are reasonable given the potential for phenyl ring orientation in polystyrene polymers. Figure 5b demonstrates a complete lack of angular dependence in the C1s → π*CC transition for both PAMS and PHS, indicating that these polymers exhibit a random distribution of phenyl ring tilt angles. However, PS, P4MS, and PTBS (Figure 5a) clearly show an angular dependence. A key question remains about the changes in phenyl ring tilt angle distribution that can cause the slope changes observed for these polymers. While MD simulations will shed additional insight into this question, Figure 5c demonstrates that changes in orientation between PS, P4MS, and PTBS are subtle. In particular, Figure 5c plots the C1s → π*CC transition for these polymers after normalizing the peak area with the fraction of carbon atoms in the monomer that have aromatic character. This allows the data to be collapsed onto a common scale, enabling a closer examination of the data uncertainty. A slope difference still exist for the three polymers with PS > P4MS > PTBS. The difference between PS and P4MS is barely detectable. However, PTBS has a statistically significant slope change when compared to PS exhibiting a higher normalized π* peak area than PS at large cos2(θ) and a lower

are the KLL carbon Auger electrons, which have an inelastic mean free path of approximately 14 Å in organic materials, following the universal curve fitting calculation described by Seah and Dench.75 An effective mean free path can be utilized to estimate the probing depth of the NEXAFS experiment as a function of the retarding bias on the electron yield detector. Previous work by Klein et al.55 showed that an electron yield detector grid voltage bias of 250 eV gives an effective mean free path (effective probing depth) of approximately 1.6 nm for PSbased samples. Assuming the electron yield signal decays exponentially with film thickness, at an effective mean free path of 1.6 nm approximately 50% of the electron yield signal is coming from the top 1 nm of the film surface. Therefore, with a detector bias of 250 eV, the NEXAFS measurements are extremely surface sensitive. Figure 5a plots the area of the C1s → π*CC transition at 285.1 eV for three different polymers (PS, P4MS, and PTBS) as a function of the measurement angle, θ, which is varied from normal (90°) to glancing (20°) relative to the film surface. The increments are 10°. The peak areas were measured after the spectra were both pre- and post-edge jump normalized as in Figure 4. For each of the polymers, the peak area is highest in the normal mode (cos2θ = 0), and gradually decreases when the measurement angle decreases toward the glancing mode (cos2θ → 1). This indicates that, nominally, the styrene rings are oriented perpendicular to the film surface, where the C1s → π*CC transition would be more pronounced in the normal mode. The lines in Figure 5 represent the best linear fit to the data for each polymer. Table S1 in the Supporting Information shows a range of slope and intercept values that provide visually acceptable fits to the π* peak area as a function of cos2(θ) for the different polymers. The slope of these lines qualitatively represents the extent of phenyl ring orientational order. A slope of zero is indicative of random orientation. As the slope increases, the average orientation of the phenyl rings that are perpendicular to the films surface is increasing, or the probability of phenyl ring orientation at large angles relative to the film surface is increasing. As was mentioned previously, the NEXAFS data in Figure 5 does not decipher whether the phenyl rings are pointed outward or inward from the film surface. However, later experiments in this report will clearly indicate outward orientation of the rings. The slope in Figure 5a decreases in the order of PS > P4MS > PTBS. As the size of the chemical functionality in the para-position of the phenyl ring increases, the extent of the ring orientation decreases. Steric effects from a bulky pendant group on the ring hinder surface orientation. The extremely bulky tertbutyl group nearly eliminates the surface orientation of the phenyl rings, as only a 12% change is observed from normal to glancing mode based on the slope of the best fit line. Figure 5b shows similar data as Figure 5a with PHS, PAMS, and P4MS. The addition of polar hydroxyl functionality on PHS eliminates surface ordering of the rings. Hydrogen bonding between the monomers of PHS may lead to a more random surface orientation, and similar reasoning was used to explain a lack of PHS orientation at polar silica interfaces.76 It is also interesting to compare PAMS and P4MS. The difference between these polymers is the position of the methyl group. For P4MS, the methyl species is in the para position on the phenyl ring, whereas in PAMS the methyl group is attached to the same backbone carbon as the phenyl ring (see structures in Figure 2). The PAMS also has almost no angular dependence on the π* peak area, indicating that the presence of the bulky methyl 15717

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Figure 6. Area of the C1s → π*CC transition (π* peak area) at an entrance grid voltage bias on the electron yield detector of either 250 or 200 eV for (a) PS, (b) P4MS, (c) PTBS, (d) PAMS, and (e) PHS. The data scatter is an indicator of the data uncertainty. In addition, Table S1 in the Supporting Information presents a range of slope and intercept values that can adequately fit the data points.

π* peak area than PS at small cos2(θ). Still the difference in angle dependent NEXAFS spectra are small, indicating that changes in the probability distribution of phenyl ring tilt angle are also subtle. The data in Figures 4 and 5 qualitatively illustrates that the phenyl rings are oriented perpendicular to the surface. Before addressing the changes in phenyl ring tilt angle distribution that

occur, we first want to demonstrate the preference for ourtward ring orientation (rings pointing away from the film surface) for these polymers. By adjusting a negative voltage bias on the electron yield detector, different effective surface sampling depths can be probed, enabling us to determine if the rings are oriented outward. This depth profiling technique has been utilized to study self-assembled monolayers,80 relaxation of thin 15718

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polymer films,79 and surface composition of photoresist films.8,51 Effective probing depths can be varied from approximately (1 to 6) nm by adjusting the retarding voltage on the electron yield detector.55 Following the work of Klein et al.,55 we measured the NEXAFS spectra for these films at 200 eV retarding voltage bias on the electron yield detector, which provides an effective probing depth of 2.6 nm (as opposed to 1.6 nm at 250 eV detector grid bias). A comparison of PS films measured at 200 and 250 eV retarding bias is shown in the Supporting Information Figure S3, which clearly demonstrates a larger electron yield intensity with the 200 eV retarding bias, consistent with a deeper probing depth. Figure 6 compares the area of the C1s → π*CC transition (π* peak area) at 285.1 eV as a function of the measurement angle (cos2 θ) and at a detector bias of both 250 or 200 eV, for the various polymers in this study. As with Figure 5, the peak areas were integrated after both pre and post edge jump normalization. For each polymer, the π* peak area is larger at every angle with the 250 eV voltage bias when compared with the spectra measured at a voltage bias of 200 eV. Since the 250 eV spectra are more surface sensitive, this confirms that the phenyl rings are pointed outward from the film surface and away from the bulk polymer, for each polymer. While subtle orientation is still observed at 200 eV detector bias for these polymers, the difference between the normal and glancing measurements (line slopes in Figure 6) is lower for each of these polymers at 200 eV when compared to 250 eV. This indicates that the surface orientation of the phenyl rings is localized at the polymer surface, and does not propagate substantially into the film. Surface orientation of polystyrene has been studied utilizing vibrationally resonant sum frequency generation (VRSFG).40,76,81−84 In these studies, the phenyl rings on polystyrene were also observed to be oriented nominally outward and perpendicular to the film surface, similar to the results in this study. Initial work by Dhinojwala et al.40 showed that the phenyl rings were oriented perpendicular to the film surface at the exposed interface and parallel to the buried PS/sapphire interface. One study by Briggman et al. estimated a tilt angle for the phenyl rings at the exposed surface near 57° relative to the surface normal.81 Other studies on polystyrene films82,83 have confirmed well ordered phenyl rings oriented outward from the film surface and nominally perpendicular to the surface, and that the phenyl ring orientation can be altered upon solvent or liquid exposure. The NEXAFS studies in this report corroborate these studies and, in addition, demonstrate that the pendant functionality on the phenyl ring of polystyrene-like polymers can alter the surface orientation, with an increase in the bulkiness of the pendant functionality leading to a more random orientation of the phenyl rings. While only the film surface was probed with this study, it is reasonable to expect that a bulky pendant functionality can also hinder the packing and orientation of phenyl rings near a buried interface. The buried interface is difficult to probe with NEXAFS; however, VR-SFG studies on these polymers could be utilized to verify this hypothesis. For example, Wilson et al.76 showed that the PS phenyl ring orientation was away from the bulk polymer at both the film surface and buried interface of PS with spin-on glass. The results of this current paper may be generalizable to other surface organization phenomena, such as the packing of self-assembled monolayers, the organization of surface layers in organic devices, or the interfacial assembly of molecules with surfactant like character.

4. COMPARISON OF EXPERIMENTAL NEXAFS DATA AND MOLECULAR DYNAMICS SIMULATIONS The NEXAFS data demonstrates orientation in the polystyrene based polymers that is dependent on the size of the pendant groups on the phenyl ring. It is also important to understand, at least qualitatively, the extent of the orientation at the surface for these polymers, as well as changes in the phenyl ring tilt angle distribution that can cause changes in the NEXAFS data. To estimate how the angular dependence of π* peak area in the NEXAFS spectra relates to the probability of ring orientation, a simple geometric model was developed as illustrated in Figure 7

Figure 7. Schematic of the geometric model that describes the probability of ring orientation within a particular angular range, as well as the ability of the ring to absorb the polarized incident photons as a function of angle.

and eq 1. In particular, the incident radiation is polarized, and when measurements are made at a “normal angle”, the incident light is oriented 90° relative to the film surface and the electric field of the incident radiation is oriented parallel to the film surface. When the measurements are made in a “glancing angle”, the incident light is oriented 20° relative to the film surface and the electric field of the incident radiation is oriented at 70° relative to the film surface. As shown in qualitatively in Figure 7, the phenyl rings have a probability of being located in various angular regimes (0−30°, 30−60°, or 60−90°). In addition, due to the amorphous nature of the PS-based films, a symmetry in the probability of orientation will exist, where the fractions of phenyl rings oriented in 0−30°, 30−60°, or 60−90° will be equal to the fraction of ring orientation in the angular ranges of 150−180°, 120−150°, or 90−120°, respectively. The probability of X-ray absorption (and therefore electron yield signal, since the electron quantum yield is not dependent on angle and the electron yield detector, is positioned such that the takeoff angle is independent of the angle of the excitation light) depends on cos2(Δθ), where Δθ represents the angular difference between the electric field of the incident radiation and the absorption dipole of the C1s → π*CC transition. Therefore, the relative NEXAFS π* peak area can be assessed as a function excitation angle, assuming a profile for the probability of phenyl ring orientation as a function of tilt angle using the following equation. θr = 180 °

π *(area)α

∑ θr = 0 °

x θr⎡⎣cos2(Δθ )⎤⎦ (1)

Equation 1 basically states that the π* peak area at a particular incident excitation angle, θ, is proportional to the summation over all potential phenyl ring orientation angles (θr = 0−180°, 15719

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where θr is the phenyl ring orientation relative to the film surface) of the fraction of phenyl rings oriented at a particular angle, xθr, multiplied by the cos2(Δθ), where Δθ describes the angular difference between the electric field of the incident radiation and the absorption dipole of the C1s → π*CC transition, which is perpendicular to the plane of the phenyl ring. In order to calculate NEXAFS π* peak area at a particular incident excitation angle, a probability of phenyl ring orientation was assigned for every 10° range, and the midpoint of that range was utilized to as θr. For example, a probability of phenyl ring orientation in the angular range of 0−10°, 10−20°, 20−30°, ...170−180° relative to the film surface would assigned a specific value. The π* peak area would be calculated using eq 1, assuming that all the rings in that angular range were oriented at 5°, 10°, 15°, ...175° respectively. The fraction of phenyl rings totaled over all angles must equal 1. The orientation profiles presented in the following figures sum the phenyl rings oriented in 30° increments, where 0−30° represents the percent of phenyl rings oriented in the angular ranges of 0−30° and 150−180° (which are equal due to symmetry as described above); 30−60° represents the percent of phenyl rings oriented in the angular ranges of 30−60° and 120− 150°; and 60−90° represents the percent of phenyl rings oriented in the angular ranges of 60−90° and 90−120°. Figure 8 compares angular dependent π* peak areas using eq 1, and assuming an angular dependent profile for phenyl ring orientation. Figure 8a and b shows these calculations for PS and PTBS, respectively. The diamond shaped symbols represent angular dependent peak areas (computed using eq 1) for a completely unoriented polymer surface where the probability of phenyl ring orientation in an angular range of 0−30° (150− 180°), 30−60° (120−150°), 60−90° (90−120°), is 1/3 for each. No change in the π* peak area is observed as a function of angle with the nonoriented sample simulation. The square symbols in Figure 8a and b represent the π* peak areas (computed using eq 1) for a highly oriented polymer surface, where 100% of the phenyl rings are oriented in the angular range of 60−90° (90− 120°). The π* peak area exhibits a strong angular dependence. The open circles in Figure 8a and b represent the experimentally observed NEXAFS data for the polymer films measured at 250 eV detector bias. The NEXAFS data lies between the nonoriented and highly oriented calculations. However, the experimental data is closer to that calculated for a nonoriented sample, demonstrating that the extent of surface orientation in these spun cast films is small, even for the most oriented PS. The “×” and “+” shaped symbols in Figure 8a and b represent two different “fits” to the experimental NEXAFS data, using eq 1 and assuming a specific profile shown in Figure 8c for the probability of phenyl ring orientation as a function of the angular range. Figure 8c displays the percentage of phenyl rings oriented in the specific angular ranges of 0−30° (150−180°), 30−60° (120− 150°), 60−90° (90−120°) for “fit #1” and “fit #2” of both the PS and PTBS experimental data shown in Figure 8a and b, respectively. It is important to note that, despite these profile differences, these numerous angular profiles for the PS and PTBS provided adequate fits to the experimental data. Since a unique solution does not exist to fit the experimental data, an accurate measure of the probability of phenyl ring orientation as a function of angular range cannot be obtained solely from the NEXAFS angular dependent data. However, some generality can be made. For example, due to the symmetry assumed in the onedimensional simulation described in Figure 7, the “effective magic angle”, where a highly oriented sample looks similar to a

Figure 8. (a) π* peak area as a function of measurement angle for PS showing experimental data (○); data calculated using eq 1 for PS and assuming random surface phenyl ring order (◇); data calculated using eq 1 for PS and assuming strong phenyl ring orientation (□); and fits to the experimental data (× and + symbols) calculated using eq 1 and assuming the angular distribution profiles shown in Figure 8c for PS fit #1 and PS fit #2. (b) Similar data for PTBS. (c) Angular profiles for PS and PTBS that describe the percent of phenyl rings oriented in a specific angular range, where both fit #1 and fit #2 provide a prediction of the experimental data shown in (a) and (b). Dashed line demonstrates the angular profile for a nonoriented surface where the fraction of functional groups in each angular range is 1/3; (a) and (b) also illustrate fits to the experimental data assuming a single average tilt angle of 49° or 46°, respectively, for the PS or PTBS. 15720

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Figure 9. (a) MD simulations describing the probability of phenyl ring orientation within a particular angular range for PS and PTBS. (b) π* peak area as a function of measurement angle calculated using eq 1 and assuming the MD simulation profile describing the probability of phenyl ring orientation in a particular angular range for PS and PTBS films.

calculated from the MD simulation profiles of phenyl ring orientation probability as a function of angle and using eq 1. The MD simulation results support the conclusion that the PS films are more oriented than PTBS, and qualitatively predict the experimentally observed trends of π* peak area as a function of cos2(θ). The results also indicate that the percentage of phenyl rings in the range of 60−90° is larger than the fraction between 0 and 30° for both PS and PTBS, as concluded with the previous discussion. In addition, the simulations show that the fraction of phenyl rings oriented between 60 and 90° is larger for PS than for PTBS, which will cause the slope difference between the two polymers in the plots of the π* peak area as a function of cos2(θ). A particularly important point demonstrated by the MD simulations is that the phenyl rings for both PS and PTBS were not oriented strongly around a specific “average tilt angle”. In contrast the simulations demonstrate that the phenyl rings exhibit a broad distribution of tilt angles. Subtle changes in the distribution of the phenyl ring orientation with increasing pendant group size lead to the observed changes in angular dependent NEXAFS data. In particular, the more oriented polymers have a higher fraction of phenyl rings at tilt angles in the 60−90° range. The MD simulations predict a larger slope (approximately double) than is observed with the experimental data, as shown in Figure 9b. The smaller experimental slopes are likely due to the imperfect polarization of the synchrotron radiation (∼90% polarized) and surface roughness of the spun cast polymer films. In addition, differences between simulation and experimental conditions such as the chain length could be responsible for this discrepancy. For example, the simulation models the polymers as short oligomers, which are more mobile than their higher molecular weight polymer counterparts. Even so, the MD simulations demonstrate that the majority of phenyl rings for both polymers are oriented in a range of 30−60° relative to the surface plane, which agrees qualitatively with the results of Briggman et al.,81 as well as demonstrating that PS has a higher probability of phenyl ring orientation at angles from 60 to 90°. To elucidate the driving force for orientation of the phenyl group at the surface, we characterize the volume of the pendant

completely unoriented sample, is 45°. Calculations using angular distributions with an increasing percentage of phenyl rings that exhibit angles greater than 45° led to stronger slopes in the NEXAFS data fits, with more oriented polymers having a higher fraction of phenyl rings at tilt angles above 45°. Another fitting option that consistently explains the larger slope for PS π* peak area as a function of cos2(θ) is where a smaller fraction of phenyl rings were oriented in the 0−30° (150−180°) angular range, and a larger fraction of phenyl rings oriented in 60−90° (90−120°) angular range, when compared with the PTBS fit. While the precise phenyl ring orientation distribution cannot be determined from the NEXAFS data, the various fits all demonstrate a PS surface with improved phenyl ring organization, when compared to the PTBS surface. The differences in orientation, however, are due to subtle changes in phenyl ring tilt angle distribution, as indicated by the small slope changes in angle dependent NEXAFS data. The NEXAFS data in Figure 8a and b can also be adequately fit by assuming an average tilt angle of 49° and 46°, respectively, for the PS and PTBS polymers. The small changes in calculated average tilt angle further indicate that the changes in phenyl ring orientation or orientation probability as a function of angle are subtle between PS and PTBS. Interestingly, molecular dynamics simulations in the next section of the paper will show that the phenyl rings were not oriented strongly around a specific “average tilt angle”. In contrast, the MD simulations demonstrate that the phenyl rings exhibit a broad distribution of tilt angles. Subtle changes in the phenyl ring distribution with increasing pendant group size lead to the observed changes in angular dependent NEXAFS data. In addition to simple geometric calculations using eq 1 and an assumed ring orientation probability as a function of angle, molecular dynamics (MD) simulations were conducted for the PS and PTBS films as described in the Experimental Section. Figure 9 illustrates the MD simulation results, where Figure 9a represents the MD simulated profiles that describe phenyl ring orientation as a function of angular range. Figure 9b presents the calculated π* peak area as a function of cos2(θ), that would be 15721

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groups via a Connolly surface. The Connolly surface,85 which is often used to approximate the contour of the electron density for a prescribed molecule, was calculated for a two monomer chain of PS and PTSB using a module of Materials Studio and is presented in Figure 10. The occupied volume enclosed by the

1.6 nm) demonstrated orientation of the phenyl rings, as shown by the angular dependent NEXAFS data. However, NEXAFS measurements performed with a negative grid bias of 200 eV (probe depth approximately 2.6 nm) exhibited smaller angular dependence, demonstrating the surface organization was localized. (c) The surface orientation of the phenyl rings was weak for all polymers, and therefore, the dif ferences in the probability distribution of phenyl ring tilt angles as a f unction of pendant group size were subtle. Even with the most oriented PS film surface, the experimental slopes of π* peak area as a function of cos2(θ) were very small when compared to calculated values for strongly oriented films. (d) Molecular dynamics simulations showed that the phenyl rings were not oriented strongly around a particular “average tilt angle”. In contrast, the simulations demonstrate that the phenyl rings exhibit a broad distribution of tilt angles. Changes in the phenyl ring tilt angle distribution with increasing pendant group size lead to the observed changes in angular dependent NEXAFS data. In particular, the more oriented polymers have a higher fraction of phenyl rings at tilt angles above 60°.

Figure 10. Connolly surface for two monomer chains of PS (left) and PTBS (right).

Connolly surface for PS and PTBS for a probe radius of 1 Å is 234 Å3 and 363 Å3, respectively, which emphasizes the bulkiness of PTBS. Figure 10 also shows that the bulky tert-butyl group causes the rings to tilt away from each other due to steric hindrance and electron density cloud repulsion of the bulky group, which can hinder the packing and orientation of phenyl rings at the interface. To quantify this, the tilt angle between the two rings was calculated by taking the dot product between the vectors that runs through the carbons in the 1 and 4 position on the phenyl ring for both monomers in the dimer, where an angle of 0° denotes no tilt. For PS and PTBS, the tilt angle was calculated to be 2.9° and 17.3°, respectively. The relative orientation of the phenyl ring as revealed through Figure 10 qualitatively explains the lack of surface order of PTBS compared to PS at the surface.

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ASSOCIATED CONTENT

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AUTHOR INFORMATION

* Supporting Information S

Additional figures and table as described in the text. This material is available free of charge via the Internet at http://pubs.acs.org.

Corresponding Author

*E-mail: [email protected]. Telephone: 410-3061940. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We would also like to thank B. C. Rinderspacher for useful discussion. This work was initiated at Sandia National Laboratories by J.L.L. and is being continued at the United States Army Research Laboratory, Aberdeen Proving Ground, by J.L.L., T.L.C., and J.W.A. Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lockheed Martin Company, for the United States Department of Energy’s National Nuclear Security Administration under Contract No. DE-AC0494AL85000. Certain commercial equipment and materials are identified in this paper in order to specify adequately the experimental procedure. In no case does such identification imply recommendations by the Army Research Laboratory or does it imply that the material or equipment identified is necessarily the best available for this purpose.

5. CONCLUSIONS NEXAFS coupled with MD simulations were exploited to monitor the surface orientation of a series of polystyrene-based polymers with various chemical functionality pendant on the phenyl rings. The following general conclusions can be made based on these results: (a) Surface orientation of the phenyl rings for these polymers was dependent on the size and type of pendant f unctionality. Polystyrene exhibited the strongest surface order. This surface orientation of the rings decreased when pendant functional groups were attached to the phenyl ring. This decrease in orientation became more prominent as the size of the pendant group increases, as poly(4-methylstyrene) exhibited less order than polystyrene, and poly(4-tertbutylstyrene) exhibited even smaller surface orientation. The presence of a hydroxyl group with poly(4hydroxystyrene) or a methyl group on the polymer backbone with poly(α-methylstyrene) completely eliminated the orientation of the phenyl rings. Molecular dynamics simulations qualitatively confirmed the experimental results by comparing PS with PTBS. (b) For all the polymers, the rings were pointed outward and from the f ilm surface and away f rom the bulk polymer. The surface orientation was localized near the film surface. NEXAFS measurements performed with an electron yield detector negative grid bias at 250 eV (probe depth approximately

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REFERENCES

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