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ArticleJuly 10, 2014

Conformational Evolution of Elongated Polymer Solutions Tailors the Polarization of Light-Emission from Organic Nanofibers
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Cite this: Macromolecules 2014, 47, 14, 4704–4710

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https://doi.org/10.1021/ma500390v

Published July 10, 2014

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Abstract

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Polymer fibers are currently exploited in tremendously important technologies. Their innovative properties are mainly determined by the behavior of the polymer macromolecules under the elongation induced by external mechanical or electrostatic forces, characterizing the fiber drawing process. Although enhanced physical properties were observed in polymer fibers produced under strong stretching conditions, studies of the process-induced nanoscale organization of the polymer molecules are not available, and most of fiber properties are still obtained on an empirical basis. Here we reveal the orientational properties of semiflexible polymers in electrospun nanofibers, which allow the polarization properties of active fibers to be finely controlled. Modeling and simulations of the conformational evolution of the polymer chains during electrostatic elongation of semidilute solutions demonstrate that the molecules stretch almost fully within less than 1 mm from jet start, increasing polymer axial orientation at the jet center. The nanoscale mapping of the local dichroism of individual fibers by polarized near-field optical microscopy unveils for the first time the presence of an internal spatial variation of the molecular order, namely the presence of a core with axially aligned molecules and a sheath with almost radially oriented molecules. These results allow important and specific fiber properties to be manipulated and tailored, as here demonstrated for the polarization of emitted light.

Subjects

Introduction

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Fibers (1-4) are typically formed upon the solidification of a tiny filament drawn from a viscous polymer solution or melt, (5-7) whose thinning follows a very complex dynamics. (8-10) Understanding how polymer chains modify their conformation at nanoscale, and to what extent they keep their configuration in solid nanostructures, is fundamental for many applications and for controlling the resulting physical properties of fibers. (11, 12) For example, polymers are typically considered bad thermal conductors, but aligning their chains in 1-dimensional (1D) nanostructures allows their thermal conductivity to be improved approaching the single-molecule limit (about 350 W m^–1 K^–1 for polyethylene). (5) Similarly, charge mobilities (μ) in organic semiconductor films are typically low (most often <10^–1 V cm^–2 s^–1), although in single π-conjugated polymer chains μ can be of the order of hundreds of V cm^–2 s^–1. (13) In the bulk, the disordered supramolecular assembly limits the charge mobility, whereas 1D nanostructures show an increase of 1–3 orders of magnitude of μ. (14, 15) The alignment of π-conjugated molecules is also effective to improve the amplification of light by stimulated emission, (16) the macroscopic quantum spatial coherence of the exciton state, (17) and the polarization of emitted light. In general, stretching a semidilute polymer solution by an electrostatic field is very effective to prime the formation of fibers, potentially resulting in a structure mostly composed of ordered and aligned chains. (18-20) Little is known however about the nanoscale features induced by elongational dynamics and about how these features can be exploited to tailor and control macroscale properties of solid nanostructures.

In this paper, we employ the unique features of scanning near-field optical microscopy (SNOM) (21-23) to investigate at nanoscale polymer fibers produced by electrospinning. Absorption measurements with nm-spatial resolution and polarization modulation provide insight into the nanoscale variation of molecular alignment, evidencing an unexpected change from axial to radial molecular orientation upon moving from the fiber axis to its surface. The formation of such complex structures occurs close to the polymer jet start, as demonstrated by modeling the evolution of the conformation of the polymer chains network.

Experimental Section

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Conjugated Polymer Nanofibers

Fibers are produced by electrospinning a solution (70–200 μM polymer) of poly[2-methoxy-5-(2-ethylhexyloxy)-1,4-phenylenevinylene] (MEH-PPV) (molecular weight 380 000 g/mol, American Dye Source Inc.). Sprayed films of microbeads and microfibers are obtained at concentrations >200 μM. The polymer is dissolved in a 1:4 (weight:weight) mixture of dimethyl sulfoxide (DMSO) and tetrahydrofuran (THF). The electrospinning system consists of a microprocessor dual drive syringe pump (33 Dual Syringe Pump, Harvard Apparatus Inc.), feeding the polymer solution through the metallic needle at constant rate (10 μL/min). A 11 kV bias is applied between the needle and a metallic collector (needle-collector distance 6 cm), made of two Al stripes positioned at a mutual distance of 2 cm. The MEH-PPV nanofibers are collected on a 1 × 1 cm² quartz substrate for optical investigation. Arrays of uniaxially aligned nanofibers are also produced by using a rotating collector (4000 rpm, corresponding to a linear velocity of 30 m/s at the disk edge) for emission polarization measurements, featuring similar morphology and optical properties as samples deposited on the Al stripes. The fiber morphology is investigated by scanning electron microscopy (SEM) using a Nova NanoSEM 450 system (FEI), with an acceleration voltage of 5–10 kV.

Polarized Emission

Optical images of the fibers are obtained by confocal microscopy, using an inverted microscope (Eclipse Ti, Nikon) equipped with a confocal laser scanning head (A1R MP, Nikon). An Ar⁺ ion laser (λ_exc = 488 nm) excites the fibers through an oil immersion objective with numerical aperture, N.A. = 1.4. The intensity of the light transmitted through the sample, measured by a photomultiplier, is recorded synchronously to the confocal acquisition of the laser-excited fluorescence. The polarization of the emission of individual nanofibers at different polymer concentrations is characterized by a microphotoluminescence system, composed by a diode laser excitation source (λ = 405 nm) coupled to an inverted microscope (IX71, Olympus). The laser beam propagates perpendicular to the substrate on which the fiber is deposited, and it is focused on the sample through a 20× objective (N.A. = 0.5, spot size 30 μm). Furthermore, the fiber is positioned with its longitudinal axis almost parallel to the incident laser polarization. The PL emitted by individual nanofibers is collected along the direction perpendicular to the substrate, by means of an optical fiber, and dispersed in a monochromator (USB 4000, Ocean Optics). The polarization of the emission is analyzed by a polarization filter mounted on a rotating stage and positioned between the emitting MEH-PPV nanofiber and the collecting optics. The system response is precisely analyzed in order to avoid artifacts due to the collection and measurement apparatus.

SNOM

A polarization-modulation near field microscopy system is used to analyze the linear dichroism of tens of individual MEH-PPV fibers. The SNOM system, operating in emission mode, excites samples in the optical near-field of a tapered optical fiber probe (Nanonics), with a nominal aperture of 50 nm, delivering a near-field power up to the tens of nW range (λ = 473 nm). The signal transmitted by the samples is collected by an aspherical lens (N.A. = 0.55, diameter of 13 mm) and sent onto a photomultiplier. The polarization modulation relies on a photoelastic modulator (PEM-100, Hinds Instruments), behaving as a waveplate with periodically modulated retardation. The modulator is followed by a λ/4 waveplate and the whole system is conceived in order to send into the optical fiber probe radiation linearly polarized along a direction periodically oscillating in the transverse plane. The photomultiplier signal is split and sent into two different digital dual lock-in amplifiers (Stanford Research SR830DSP). The first one, referenced to the polarization modulator frequency, f, provides with an output (hereafter called AC) representative of the sample response to polarized radiation, whereas the second lock-in, referenced to a slow modulation frequency f′ (f/f′ >10) of the laser amplitude, is used to determine the optical transmission averaged over all polarization states (DC output).

The dichroic ratio of sample, γ = ((I_∥ – I_⊥)/(I_∥ + I_⊥)), where I_∥ and I_⊥ are the transmitted intensity for polarization aligned along two mutually orthogonal directions, respectively, is quantitatively evaluated from the ratio AC/DC. This requires to model the behavior of the whole optical chain and to account for the residual optical activity of its components, including the optical fiber probe (see Supporting Information). Indeed, reference measurements performed on bare substrates provide a dichroic ratio around zero as expected (see also Supporting Information). Moreover, the polarization state of the light incident on the sample is also characterized, by rotating the linear polarization of the light coupled into the SNOM fiber using a λ/2 waveplate, and measuring the intensity transmitted by a linear polarizer used as sample for each position of the λ/2 waveplate (see also Supporting Information). We have measured a ratio between the maximum and the minimum intensity transmitted by the polarizer in the range 10¹–10². Overall, calibration experiments allow any contribution of the measurement setup to the obtained results to be ruled out.

Results and Discussion

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Figure 1 shows SEM pictures of MEH-PPV fibers produced by electrospinning solutions with different polymer concentrations. The fibers generally feature a ribbon shape, with average width in the range 500–600 nm and width:height ratio of about 10:1. The average fiber width increases by roughly 30% upon increasing the polymer concentration in the 70–200 μM range. In addition, fibers emit bright light, allowing the chain order to be investigated by optical methods (Figure 2a). Figure 2b shows the confocal transmission micrograph of excitation laser light, collected by crossed polarizers (analyzer axis perpendicular to the incident laser polarization), for nanofibers positioned at 0°, 65°, and 90° with respect to the incident laser polarization. A significant transmitted signal can be measured only for fibers positioned at 65°, indicating optical anisotropy which is expectedly the result of a preferential molecular alignment along the fiber length.

Figure 1. SEM images of electrospun MEH-PPV fibers realized by varying the solution polymer concentration in the range 70–200 μM. The corresponding polymer volume fraction, ϕ, is 0.025 (a), 0.036 (b), 0.054 (c), and 0.064 (d), respectively. Scale bar: 20 μm. Inset in part c: Zoomed micrograph of an individual fiber highlighting its ribbon shape. Scale bar: 2 μm.

Figure 2. (a) Fluorescence confocal micrograph of conjugated polymer fibers. Scale bar: 10 μm. (b) Confocal map of the exciting laser intensity transmitted by the fibers, collected simultaneously to the emission map in part (a). The polarization of the excitation laser (highlighted by the horizontal arrow) is aligned parallel to the longitudinal axis of the horizontal fiber, whereas the axis of the analyzer (highlighted by the vertical arrow) is positioned perpendicularly to the incident laser polarization.

Polarized near-field absorption microscopy (Figure 3a) provides a direct measurement of the spatial variation of polymer alignment, through the map of the local dichroism, γ, i.e. the normalized difference between the transmission of radiation polarized along two mutually orthogonal directions. The map (Figure 3b) is determined by the distribution and anisotropy of absorbing chromophores (see Supporting Information). Here, the most important finding is the spatial variation of dichroism and, consequently, of molecular alignment (Figure 3c). Unexpectedly, the sign of the dichroic ratio, γ, is not constant throughout the fiber, because of regions showing preferential absorption of light polarized along or across the fiber axis (for the scan shown, they correspond to negative or positive γ, respectively).

Figure 3. (a) Schematics of the polarization-modulation SNOM measurement. PEM: photoelastic modulator, PMT: photomultiplier. (b) Map of the dichroic ratio of a single MEH-PPV fiber. The dichroic ratio is zero for nonoptically active regions (background contribution subtracted, see Supporting Information). (c) Line profile analysis displaying the cross sections, along the dashed segment in part b, of γ (continuous line) and of topography got simultaneously with the optical data (dotted line). The change in sign of γ when crossing the fiber (dashed horizontal line corresponding to γ = 0) indicates different alignments of the polymer with respect to the fiber axis.

These results suggest the presence of a core, with width ∼40% of fiber diameter, where chromophore dipoles preferentially align along the fiber length, whereas molecules closer to the fiber border show a preferential radial orientation. The decrease of the dichroic ratio from positive to null values nearby the fiber edges (i.e., for positions roughly ≤0.5 μm and ≥1.5 μm according to the horizontal axis of Figure 3c) may be, instead, affected by the local curvature of the ribbon-shaped fibers and will not be considered in the following analysis. To learn more about the origin of the found spatial variation of the molecular alignment, we use a model of the polymer network and perform simulations of its dynamics, as previously developed for fully flexible and semiflexible polymer chains. (24-26) Simulations are here aimed at better rationalizing the observed chain orientation in the core, and at assessing the relevant process variables determining such orientation, thus ultimately allowing the macroscopic physical properties of the electrospun fibers to be tailored and controlled. In this approach, a semiflexible conjugated polymer chain is modeled as a series of N rigid segments, each of length b = nd (n spherical beads of diameter d ≅ 1.2 nm, each bead consisting of two chemical monomers). The segment length b represents the average distance between two neighboring bonding defects along the chain backbone, where a bonding defect introduces local flexibility in the chain. (27) The corresponding defects concentration, using two chemical monomers per bead, is (2n)⁻¹ of monomers. The chain conformational correlation is lost above the scale of a segment due to the bonding defects, and therefore the rigid segment b is regarded as a Kuhn segment, and a freely jointed chain model is assumed. Fully flexible polymers are a particular case of the model (n = 1), (24) and generality is retained by using the segmental aspect ratio parameter, n, to specify the degree of chain flexibility.

In general, the high entanglement of chains creating a connective network determines the viscoelastic property of semidilute solutions. An entanglement can be simply defined as a topological constraint that inhibits intercrossing of two chains. The conformation of the entangled polymer network in the semidilute solution and the interactions relevant to the solvent type are described by scaling laws. When the segmental aspect ratio is high, an entanglement strand (i.e., a chain section between two adjacent entanglements) has the same length scale as the network correlation length (mesh size), ξ, the end-to-end distance of an unperturbed subchain containing N_s rigid segments. Given the aspect ratio n, polymer volume fraction ϕ, and solution properties expressed by Flory’s exponent, (28) ν and Flory’s interaction parameter, χ, the number of rigid segments in a subchain for good solvents is (Supporting Information):

(1)and the corresponding correlation length is ξ ≈ b[(1–2χ)/n]^2v−1N_s^v. (29) The exponent v is 0.5 for ideal chains, corresponding to θ-solvents, and ∼0.6 for real chains, corresponding to good and athermal solvents.

The mapping of N_s as a function of n²ϕ, for different solvent qualities, is depicted in Figure 4a. The effect of various solvents is discussed in detail in the Supporting Information. When the calculated N_s is above the N limit (upper dotted line, designating the overlap concentration ϕ*), the polymer network is not sufficiently entangled for elastic stretching. The limit n = 1 (lower dotted line) designates the minimal selectable n value. For our solvent mixture (THF:DMSO 4:1 weight:weight), the interaction parameter can be estimated as χ ≅ 0.38, (30) and for our polymer volume fraction, ϕ = 0.025, the transition from ideal to real chain conformation occurs at n ≅ 2.7 beads (point B′ in Figure 4a). The corresponding defects concentration (19% of monomers) is much higher than typical values (<10%, equivalent to n > 5), (27) and therefore the conformation of subchains is close to ideal (θ-solvent line in Figure 4a, right to point B′). At the high temperature limit (athermal solvent), the transition from ideal to real conformation occurs at n ≅ ϕ^–1/3 ≅ 3.4 beads (point B in Figure 4a), equivalent to 15% defects. Thus, as a chain is stiffer (higher n) it is more likely to be practically ideal, regardless of the solvent quality, provided that sufficient entanglement exists. At low concentration, when n < ϕ^–1/2 (left to point A in Figure 4a), the subchain consists of many segments and does not interact with other chains [Figure 4b(i)]. When n ≈ ϕ^–1/2 (point A in Figure 4a), the correlation length ξ is of the same scale as the segment length b [Figure 4b(ii)].

Figure 4. (a) Plot of N_s vs n²ϕ and solvent quality. The θ-solvent curve marks the crossover between good and poor solvents. The dotted lines constitute the upper and lower limits for ϕ = 0.025. Polymer molecular weight = 380,000 g/mol, equivalent to N_beads = 730. Points B and B′, plotted for ϕ = 0.025 for Flory’s interaction parameter χ ≅ 0 and χ ≅ 0.38, respectively, mark the transition from ideal subchains (right) to real subchains (left). Prefactors are omitted for sake of simplicity. (b) Crossover (point A in part a) of the polymer network conformation with respect to the scale of the correlation length, ξ (circles) and the segment length, b: (i) regular semidilute, ξ > b, (ii) crossover, ξ ≈ b, and (iii) different chains intermix within a single correlation volume, ξ ≈ b. (c) Simulation of subchains during electrospinning. The axial mesh size ξ_∥, radial mesh size ξ_⊥, and orientation parameter O are plotted vs the axial position, z, along the jet. ξ_∥ is compared to the theoretical model (dotted line). The position close to full subchain extension is designated by z_s. Parameters used: ideal chain, ϕ = 0.025, n = 5 beads, d = 1.2 nm, ξ₀ ≅ 20 nm, N_s = 14 segments. Jet dynamics is from Figure S4 (see Supporting Information).

However, when chains are not fully flexible, the correlation volume is not completely occupied by a single segment, and further increase of n and/or ϕ is possible. The network then crosses over to a state where different chains intermix within a single correlation volume [Figure 4b(iii)], increasing the probability of interchain overlap. The increased interaction between neighboring rigid segments may lead to nematic ordering and enhanced orientation, according to Onsager theory. For the volume fraction used in the experiment, ϕ = 0.025, this crossover occurs at n ≅ 8.6 beads, corresponding to defects concentration of ∼6% of monomers.

On these bases, the evolution of the polymer conformation under dynamic tension can be described by a beads-and-spring lattice model and a 3D random walk simulation. (24) To this aim, the calculated number of segments per subchain, N_s, the initial correlation length, ξ₀, and the pertaining experimental conditions are used as input to the simulations. The jet velocity is derived from the measured jet radius a, subjecting each subchain to a hydrodynamic force induced by the solvent, as well as to entropic elastic forces applied by its neighboring subchains.

The simulation results are presented in Figure 4c. It is seen that subchains fully extend within less than 1 mm from the jet start (position z_s), while contracting laterally, and their segments become fully oriented along the jet axis. The theoretical expression for the axial stretching (dotted line in Figure 4c), derived for linear elasticity, is given by (24) ξ_II ≈ (v/v₀)ξ₀ = (a₀/a)²ξ₀, (a₀ and v₀ are the jet initial radius and velocity, respectively), whereas the orientational parameter (dashed line in Figure 4c) is defined as O = (3/2)⟨cos²δ⟩ – (1/2), δ being the angle between a rigid segment of the polymer molecule and the longitudinal axis. An example of the conformational evolution of a single subchain under the same conditions is shown in Figure 5.

The polymer chain is entangled with other chains in the solution (Figure 5a,b). Each subchain (an entanglement strand) starts from an equilibrium conformation at the jet start (Figure 5c), proceeds through intermediate stretching, and approaches full extension and lateral contraction (Figure 5d). The subchain conformation is sensitive to the average concentration of defects. A change in the defects concentration from 10% of monomers to 15% raises the subchain size from 14 to 67 segments and increases the correlation length. The lateral contraction of individual subchains affects the conformation of the whole polymer network, narrowing its radius a_p faster than the narrowing of the jet radius a. The dominant effect is of axial stretching and lateral contraction, resulting in compacting of the network toward the jet core. Thus, the model and simulation predict axial alignment at the jet center, while closer to the jet boundary the stretching effect is not dominant and other mechanisms prevail. These may include the influence of surface charges as recently investigated in polyamide 6 nanofibers. (31) These charges can be tailored by changing the polarity of the applied voltage, (32) and can lead to distinct properties such as enhanced surface energy compared to solution-processed films.

Figure 5. (a) Polymer network at rest. (b) Single chain with N = 146. (c) Examples of single subchains, left N_s = 14 (n = 5, 10% defects), right N_s = 67 (n = 3.4, 15% defects). (d) Stretched subchains, N_s = 14, z = 0.08 mm (top) and z = *z_s* = 0.16 mm (bottom).

In our model, full extension is approached when ξ_II ≈ bN_s, i.e. at the axial position z_s and a corresponding jet radius a_s. The axial position of full stretching, omitting the effect of n, scales as (Supporting Information):

(2)

The estimated z_s for various solvent qualities, allowing the axial position of full stretching and consequently the resulting fiber properties for each particular nanofabrication experiment to be predicted, is shown in Figure 6a. Typically z_s < 1 mm and a₀/a_s = 2–10, close to the jet start. Considering that the final radius reduction ratios in electrospinning are typically 10²–10⁴, substantial stretching occurs quite early in the process. For given polymer concentration and molecular weight, when n is larger (i.e., longer segments, equivalent to lower defects concentration), full stretching is approached at a higher jet radius and lower z_s; however, at the same time, the number of entanglements per chain N/N_s is higher and therefore the solution viscosity will be larger, increasing z_s. In contrast to the radius reduction ratio a₀/a_s, the axial position z_s is strongly affected by the jet rheology, resulting in a concentration dependence with a large positive exponent, as well as added dependence on the molecular weight (Figure 6a). In addition to its dependence on the molecular weight and concentration as expressed in eq 2, z_s strongly depends on the intensity of the electrostatic field E and the jet initial velocity v₀. It can be shown that this dependence may be approximated by z_s ∼ v₀^1/2E^-1, meaning that a high strain rate, caused by high E and low v₀, should result in earlier stretching.

Figure 6. (a) Plot showing the axial position where subchains approach full extension, z_s/a₀, normalized by N^3/2, vs the polymer volume fraction ϕ and solvent quality, for n = 1. The dotted line constitutes the lower limit imposed by N_s < N. Prefactors are omitted for sake of simplicity. Points B and B′ are explained in Figure 4a. Insets: Plot of the normalized nanofiber emission intensity vs the angle between the fiber and the analyzer axis, measured on fibers electrospun from a solution with ϕ = 0.03 (left inset) and on a sprayed film for comparison (right inset). (b) Polarization ratio, r_pol, vs solution volume fraction, ϕ. The dashed line is a guide for the eyes. An unpolarized sample (sprayed film) has r_pol = 1. (c, d) Confocal images of nanofiber polarized emission. The laser-excited emission is filtered through an analyzer with axis (highlighted by arrows) parallel and perpendicular to the fiber axis, respectively. (e) Experimental distributions of the nanofiber polarization ratio, r_pol, at different polymer concentrations.

A substantial axial stretching of chains is therefore predicted during the initial stage of the elongational flow, causing lateral contraction of the polymer network toward the center of the jet, as well as orientation of chain segments along the jet axis. Similar results obtained for fully flexible chains (particular case with n = 1) have been confirmed by X-ray imaging of high strain rate electrified jets. (24, 25) In electrospinning, the electric field provides the flow of the semidilute solution with a characteristic increasing velocity along the jet axis, with a strain rate that continuously increases the elastic stretching of the polymer network and reduces network relaxation. Our model shows that full chain stretching is approached at a higher jet radius as the chain is stiffer, at a region where the mass loss rate due to evaporation is still low.

When stretching is less dominant (e.g., at low electric field and high flow rate), the rapid solvent evaporation can adversely affect the polymer matrix, creating a porous nanofiber structure. Dominant evaporation can also lead to a rapid solidification of the jet surface, limiting further solvent loss from the core. (33-35) The presence of residual solvent content in the jet core would allow for chain relaxation, thus disfavoring the retention of alignment in the core as recently reported for electrospun polyvinyl-alcohol fibers. (36) Instead, when stretching is dominant as in the present case, the polymer network compacts toward the center, producing an increase of the density close to the jet axis. (26) The here observed ribbons (Figure 1) are hence likely affected by concurring effects rather than jet skin collapse, such as flattening and relaxation processes occurring at the impact onto the substrate. This would be consistent with the presence of a slowly evaporating solvent component, i.e. with a jet time-of-flight which is comparable with the drying time scale, (35, 37) and supported by the joints observed in SEM micrographs of intersecting deposited fibers (e.g., Figure 1b,c).

This description is in agreement with polarization modulation measurements (Figure 3), showing a change in the sign of the dichroic ratio along the fiber radius, and indicating a preferred axial alignment of molecules at the fiber core, whereas molecules closer to the fiber boundary possess a preferred radial alignment. Thus, at the jet center axial stretching is dominant, and one can anticipate a propensity for interchain interaction and π–π stacking and consequently high extent of local crystallinity. (38) At the boundary region of the jet, where the polymer concentration is reduced, (25, 26) entanglement may be low or nonexistent, allowing partial or full relaxation of chains back to their coil-shaped equilibrium state. This mechanism is also supported by a recent study on entanglement loss in extensional flow, (39) showing that the electrospinning process causes partial untangling of the polymer network when stretching is faster than the chains relaxation time.

Overall, full extension of the network occurs at an earlier stage of the jet (lower z_s) if the solution initial concentration, the polymer molecular weight, and the solvent quality are lower, accounting for lower network entanglement. Under such conditions, the likelihood that the extended conformation, and the associated axial molecular alignment, will partially remain in the polymer structure after solidification is higher. This enables tailoring the physical properties of fibers, such as the far-field, macroscale emission from fibers. The z_s(ϕ) diagram of Figure 6a clearly relates the chain alignment, and hence the resulting polarization, to the polymer volume fraction (i.e., solution concentration). Indeed, by our approach we obtain a fine-tuning of the polarization ratio of the fiber emission (r_pol = I_∥/I_⊥, given by the ratio between the photoluminescence intensity parallel, I_∥, and perpendicular, I_⊥, to the fiber axis, respectively), increasing up to about 5 by gradually decreasing the solution concentration down to a volume fraction ϕ = 0.025, as shown in Figure 6b–e. This demonstrates the possibility of tailoring specific fiber properties by the relevant process parameters. The ultimate r_pol values may also benefit from the higher density in the core, (26) as well as from electronic energy-transfer mechanisms, which strongly affect the emission properties of conjugated polymers. (40, 41) Conjugated polymer fibers frequently show red-shifted absorption compared to spincast films (Figure S1a and ref 42), a property consistent with a longer effective conjugation length consequence of the stretched conformation. In fact, the elongational dynamics of solutions leads to extended structures having interchain alignment. The bonding defects concentration, (2n)⁻¹ of monomers, determines chain flexibility, and appears as a possible key factor in controlling the desired morphology.

Conclusions

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In summary, anisotropy at nanoscale is investigated in polymer fibers by polarization modulation SNOM absorption measurements, evidencing a variation of molecular orientation from axial to radial upon moving from the fiber axis to its surface. Modeling the evolution of the conformation of the chains network allows us to identify key parameters for controlling molecular alignment, as demonstrated by the fine control of the emission polarization. The found complex internal structure and assessment of the key influencing process variables open new perspectives for tailoring the molecular morphology and resulting fiber properties.

Supporting Information

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Theoretical model rational and details, fibers optical properties, and technical details about SNOM measurements. This material is available free of charge via the Internet at http://pubs.acs.org.

ma500390v_si_001.pdf (526.79 kb)

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Author Information

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Corresponding Authors
- Andrea Camposeo - National Nanotechnology Laboratory of Istituto Nanoscienze-CNR, via Arnesano, I-73100 Lecce, Italy; Email: [email protected]
- Israel Greenfeld - Department of Mechanical Engineering, Technion−Israel Institute of Technology, Haifa 32000, Israel; Email: [email protected]
- Dario Pisignano - National Nanotechnology Laboratory of Istituto Nanoscienze-CNR, via Arnesano, I-73100 Lecce, Italy; Dipartimento di Matematica e Fisica “Ennio De Giorgi”, Università del Salento, via Arnesano I-73100 Lecce, Italy; Email: [email protected]
Authors
- Francesco Tantussi - Dipartimento di Fisica “Enrico Fermi” and CNISM, Università di Pisa, Largo Bruno Pontecorvo 3, I-56127 Pisa, Italy; Istituto Nazionale di Ottica INO−CNR, Sezione di Pisa, I-56127 Pisa (Italy)
- Maria Moffa - National Nanotechnology Laboratory of Istituto Nanoscienze-CNR, via Arnesano, I-73100 Lecce, Italy
- Francesco Fuso - Dipartimento di Fisica “Enrico Fermi” and CNISM, Università di Pisa, Largo Bruno Pontecorvo 3, I-56127 Pisa, Italy; Istituto Nazionale di Ottica INO−CNR, Sezione di Pisa, I-56127 Pisa (Italy)
- Maria Allegrini - Dipartimento di Fisica “Enrico Fermi” and CNISM, Università di Pisa, Largo Bruno Pontecorvo 3, I-56127 Pisa, Italy; Istituto Nazionale di Ottica INO−CNR, Sezione di Pisa, I-56127 Pisa (Italy)
- Eyal Zussman - Department of Mechanical Engineering, Technion−Israel Institute of Technology, Haifa 32000, Israel
Author Contributions
These authors contributed equally.
Notes
The authors declare no competing financial interest.

Acknowledgment

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We gratefully acknowledge the financial support of the United States-Israel Binational Science Foundation (BSF Grant 2006061), the RBNI-Russell Berrie Nanotechnology Institute, and the Israel Science Foundation (ISF Grant 770/11). The research leading to these results has received funding from the European Research Council under the European Union’s Seventh Framework Programme (FP/2007-2013)/ERC Grant Agreement No. 306357 (ERC Starting Grant “NANO-JETS”). The authors also gratefully thank S. Pagliara for sample preparation, E. Caldi for assistance in the SNOM measurements, S. Girardo for imaging of the polymer jet, and V. Fasano for confocal images.

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Shen, Sheng; Henry, Asegun; Tong, Jonathan; Zheng, Ruiting; Chen, Gang
Nature Nanotechnology (2010), 5 (4), 251-255CODEN: NNAABX; ISSN:1748-3387. (Nature Publishing Group)
Bulk polymers are generally regarded as thermal insulators, and typically have thermal conductivities on the order of 0.1 W m-1 K-1 (ref.). However, recent work suggests that individual chains of polyethylene-the simplest and most widely used polymer-can have extremely high thermal cond. Practical applications of these polymers may also require that the individual chains form fibers or films. Here, we report the fabrication of high-quality ultra-drawn polyethylene nanofibers with diams. of 50-500 nm and lengths up to tens of millimeters. The thermal cond. of the nanofibers was found to be as high as ∼104 W m-1 K-1, which is larger than the conductivities of about half of the pure metals. The high thermal cond. is attributed to the restructuring of the polymer chains by stretching, which improves the fiber quality toward an ideal' single cryst. fiber. Such thermally conductive polymers are potentially useful as heat spreaders and could supplement conventional metallic heat-transfer materials, which are used in applications such as solar hot-water collectors, heat exchangers and electronic packaging.
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Reneker, Darrell H.; Chun, Iksoo
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The electrospinning process for the prepn. of polymer fibers with diams. of 40-2000 nm is reviewed with 41 refs. Electrospinning uses elec. forces to produce polymer fibers with nanometer-scale diams. Electrospinning occurs when the elec. forces at the surface of a polymer soln. or melt overcome the surface tension and cause an elec. charged jet to be ejected. When the jet dries or solidifies, an elec. charged fiber remains. This charged fiber can be directed or accelerated by elec. forces and then collected in sheets or other useful geometrical forms. More than 20 polymers, including poly(ethylene oxide), nylon, polyimide, DNA, polyaramid, and polyaniline, have been electrospun in the lab. Most were spun from soln., although spinning from the melt in vacuum and air was also demonstrated. Electrospinning from polymer melts in a vacuum is advantageous because higher fields and higher temps. can be used than in air.
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Arinstein, Arkadii; Burman, Michael; Gendelman, Oleg; Zussman, Eyal
Nature Nanotechnology (2007), 2 (1), 59-62CODEN: NNAABX; ISSN:1748-3387. (Nature Publishing Group)
Polymer materials of reduced size and dimensionality, such as thin films, polymer nanofibers and nanotubes, exhibit exceptional mech. properties compared with those of their macroscopic counterparts. We discuss here the abrupt increase in Young's modulus in polymer nanofibers. Using scaling estn. we show that this effect occurs when, in the amorphous (non-cryst.) part of the nanofibers, the transversal size of regions consisting of orientation-correlated macromols. is comparable to the nanofibers diam., thereby resulting in confinement of the supramol. structure. We suggest that in polymer nanofibers the resulting supramol. microstructure plays a more dominant role in the deformation process than previously thought, challenging the commonly held view that surface effects are most significant. The concept we develop also provides a way to interpret the obsd. - but not yet understood - temp. dependence of Young's modulus in nanofibers of different diams.
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Nature Physics (2008), 4 (2), 149-154CODEN: NPAHAX; ISSN:1745-2473. (Nature Publishing Group)
When a liq. is subject to a sufficiently strong elec. field, it can be induced to emit thin fluid jets from conical tip structures that form at its surface. Such behavior has both fundamental and practical implications, from rain drops in thunder clouds to pendant drops in electrospray mass spectrometry. But the large difference in length scales between these microscopic/nanoscopic jets and the macroscopic drops and films from which they emerge has made it difficult to model the electrohydrodynamic (EHD) processes that govern such phenomena. Here, simulations and expts. are reported that enable a comprehensive picture of the mechanisms of cone formation, jet emission and break-up that occur during EHD tip streaming from a liq. film of finite cond. Simulations show that EHD tip streaming does not occur if the liq. is perfectly conducting or perfectly insulating, and enable to develop a scaling law to predict the size of the drops produced from jet break-up.
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Nature Physics (2010), 6 (8), 625-631CODEN: NPAHAX; ISSN:1745-2473. (Nature Publishing Group)
Break-up of viscoelastic filaments is pervasive in both nature and technol. If a filament is formed by placing a drop of saliva between a thumb and forefinger and is stretched, the filament's morphol. close to break-up corresponds to beads of several sizes interconnected by slender threads. Although there is general agreement that formation of such beads-on-a-string (BOAS) structures occurs only for viscoelastic fluids, the underlying physics remains unclear and controversial. The physics leading to the formation of BOAS structures is probed by numerical simulation. Computations reveal that viscoelasticity alone does not give rise to a small, satellite bead between two much larger main beads but that inertia is required for its formation. Viscoelasticity, however, enhances the growth of the bead and delays pinch-off, which leads to a relatively long-lived beaded structure. We also show for the first time theor. that yet smaller, sub-satellite beads can also form as seen in expts.
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Sessile and pendant droplets of polymer solns. acquire stable shapes when they are elec. charged by applying an elec. p.d. between the droplet and a flat plate, if the potential is not too large. These stable shapes result only from equil. of the elec. forces and surface tension in the cases of inviscid, Newtonian, and viscoelastic liqs. In liqs. with a nonrelaxing elastic force, that force also affects the shapes. It is widely assumed that when the crit. potential [Jgr]0* has been reached and any further increase will destroy the equil., the liq. body acquires a conical shape referred to as the Taylor cone, having a half angle of 49.3°. The Taylor cone corresponds essentially to a specific self-similar soln., whereas there exist nonself-similar solns. which do not tend toward a Taylor cone. Thus, the Taylor cone does not represent a unique crit. shape: there exists another shape, which is not self-similar. The obsd. half angles are much closer to the new shape. In this article a theory of stable shapes of droplets affected by an elec. field is proposed and compared with data acquired in our exptl. work on electrospinning of nanofibers from polymer solns. and melts.
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Nature Nanotechnology (2007), 2 (10), 647-652CODEN: NNAABX; ISSN:1748-3387. (Nature Publishing Group)
We control the chain conformation of a semiconducting polymer by encapsulating it within the aligned nanopores of a silica host. The confinement leads to polarized, low-threshold amplified spontaneous emission from the polymer chains. The polymer enters the porous silica film from only one face and the filling of the pores is therefore graded. As a result, the profile of the index of refraction in the film is also graded, in the direction normal to the pores, so that the composite film forms a low-loss, graded-index waveguide. The aligned polymer chains plus naturally formed waveguide are ideally configured for optical gain, with a threshold for amplified spontaneous emission that is twenty times lower than in comparable unoriented polymer films. Moreover, the optimal conditions for ASE are met in only one spatial orientation and with one polarization. The results show that nanometer-scale control of semiconducting polymer chain orientation and position leads to novel and desirable optical properties.
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Nature Physics (2006), 2 (1), 32-35CODEN: NPAHAX; ISSN:1745-2473. (Nature Publishing Group)
Macroscopic quantum coherence has been obsd. in some many-body systems including superconductors, quantum liqs. and cold atom condensates, but never for a single quasi-particle state. In an ideal semiconductor, excitons (electron-hole pairs bound by the Coulomb interaction) can, in principle, exist as delocalized plane waves extending over the entire vol. However, any kind of disorder prevents long-range spatial coherence from emerging. There has been evidence for the formation of macroscopic coherent states only in condensate phases such as in the case of microcavity polaritons condensation or in a dense quasi-two-dimensional exciton gas. It is unclear however, whether in this latter case the observations are really related to macroscopic coherence. Here, we show that a single exciton state in an individual ordered conjugated polymer chain, shows macroscopic quantum spatial coherence reaching tens of micrometres, limited by the chain length. The spatial coherence of the k = 0 exciton state is demonstrated by selecting two spatially sepd. emitting regions of the chain and observing their interference.
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Electrospinning gives rise to polymer nano-fibers. The spinning process is characterized by strong deformations of the polymer material taking place during the spinning process and a very rapid structure formation process happening within milliseconds. We were interested in the influence of the peculiar spinning process on the structures of nanofibers. For this purpose, we analyzed the internal structures of nano-fibers spun from polyamide-6 and poly lactide with an av. diam. of about 50 nm. The fibers were partially cryst., with degrees of crystallinity not significantly smaller than those found for less rapidly quenched and much thicker melt-extruded fibers. The annealing of polyamide fibers at elevated temps. resulted in a transformation from the disordered γ modification to the more highly ordered α modification, and this again was in close agreement with the response of melt-extruded fibers. The orientation of the crystals along the fiber axis was strongly inhomogeneous: it was, on av., very weak, yet it could be quite pronounced locally. Small elongations of approx. 10% resulted in well-developed homogeneous crystal orientations.
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Macromolecules

Cite this: Macromolecules 2014, 47, 14, 4704–4710

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Abstract
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Figure 1
Figure 1. SEM images of electrospun MEH-PPV fibers realized by varying the solution polymer concentration in the range 70–200 μM. The corresponding polymer volume fraction, ϕ, is 0.025 (a), 0.036 (b), 0.054 (c), and 0.064 (d), respectively. Scale bar: 20 μm. Inset in part c: Zoomed micrograph of an individual fiber highlighting its ribbon shape. Scale bar: 2 μm.
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Figure 2
Figure 2. (a) Fluorescence confocal micrograph of conjugated polymer fibers. Scale bar: 10 μm. (b) Confocal map of the exciting laser intensity transmitted by the fibers, collected simultaneously to the emission map in part (a). The polarization of the excitation laser (highlighted by the horizontal arrow) is aligned parallel to the longitudinal axis of the horizontal fiber, whereas the axis of the analyzer (highlighted by the vertical arrow) is positioned perpendicularly to the incident laser polarization.
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Figure 3
Figure 3. (a) Schematics of the polarization-modulation SNOM measurement. PEM: photoelastic modulator, PMT: photomultiplier. (b) Map of the dichroic ratio of a single MEH-PPV fiber. The dichroic ratio is zero for nonoptically active regions (background contribution subtracted, see Supporting Information). (c) Line profile analysis displaying the cross sections, along the dashed segment in part b, of γ (continuous line) and of topography got simultaneously with the optical data (dotted line). The change in sign of γ when crossing the fiber (dashed horizontal line corresponding to γ = 0) indicates different alignments of the polymer with respect to the fiber axis.
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Figure 4
Figure 4. (a) Plot of N_s vs n²ϕ and solvent quality. The θ-solvent curve marks the crossover between good and poor solvents. The dotted lines constitute the upper and lower limits for ϕ = 0.025. Polymer molecular weight = 380,000 g/mol, equivalent to N_beads = 730. Points B and B′, plotted for ϕ = 0.025 for Flory’s interaction parameter χ ≅ 0 and χ ≅ 0.38, respectively, mark the transition from ideal subchains (right) to real subchains (left). Prefactors are omitted for sake of simplicity. (b) Crossover (point A in part a) of the polymer network conformation with respect to the scale of the correlation length, ξ (circles) and the segment length, b: (i) regular semidilute, ξ > b, (ii) crossover, ξ ≈ b, and (iii) different chains intermix within a single correlation volume, ξ ≈ b. (c) Simulation of subchains during electrospinning. The axial mesh size ξ_∥, radial mesh size ξ_⊥, and orientation parameter O are plotted vs the axial position, z, along the jet. ξ_∥ is compared to the theoretical model (dotted line). The position close to full subchain extension is designated by z_s. Parameters used: ideal chain, ϕ = 0.025, n = 5 beads, d = 1.2 nm, ξ₀ ≅ 20 nm, N_s = 14 segments. Jet dynamics is from Figure S4 (see Supporting Information).
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Figure 5
Figure 5. (a) Polymer network at rest. (b) Single chain with N = 146. (c) Examples of single subchains, left N_s = 14 (n = 5, 10% defects), right N_s = 67 (n = 3.4, 15% defects). (d) Stretched subchains, N_s = 14, z = 0.08 mm (top) and z = z_s = 0.16 mm (bottom).
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Figure 6
Figure 6. (a) Plot showing the axial position where subchains approach full extension, z_s/a₀, normalized by N^3/2, vs the polymer volume fraction ϕ and solvent quality, for n = 1. The dotted line constitutes the lower limit imposed by N_s < N. Prefactors are omitted for sake of simplicity. Points B and B′ are explained in Figure 4a. Insets: Plot of the normalized nanofiber emission intensity vs the angle between the fiber and the analyzer axis, measured on fibers electrospun from a solution with ϕ = 0.03 (left inset) and on a sprayed film for comparison (right inset). (b) Polarization ratio, r_pol, vs solution volume fraction, ϕ. The dashed line is a guide for the eyes. An unpolarized sample (sprayed film) has r_pol = 1. (c, d) Confocal images of nanofiber polarized emission. The laser-excited emission is filtered through an analyzer with axis (highlighted by arrows) parallel and perpendicular to the fiber axis, respectively. (e) Experimental distributions of the nanofiber polarization ratio, r_pol, at different polymer concentrations.
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  High strain rate extensional flow of a semidilute polymer soln. can result in fragmentation caused by polymer entanglement loss, evidenced by appearance of short nanofibers during electrospinning. The typically desired outcome of electrospinning is long continuous fibers or beads, but, under certain material and process conditions, short nanofibers can be obtained, a morphol. that has scarcely been studied. Here we study the conditions that lead to the creation of short nanofibers, and find a distinct parametric space in which they are likely to appear, requiring a combination of low entanglement of the polymer chains and high strain rate of the electrospinning jet. Measurements of the length and diam. of short nanofibers, electrospun from PMMA dissolved in a blend of CHCl3 and DMF, confirm the theor. prediction that the fragmentation of the jet into short fibers is brought about by elastic stretching and loss of entanglement of the polymer network. The ability to tune nanofiber length, diam. and nanostructure, by modifying variables such as the molar mass, concn., solvent quality, elec. field intensity, and flow rate, can be exploited for improving their mech. and thermodn. properties, leading to novel applications in engineering and life sciences. © 2013 Wiley Periodicals, Inc. J. Polym. Sci., Part B: Polym. Phys., 2013.
40. 40
  Hwang, I.; Scholes, G. D. Chem. Mater. 2011, 23, 610
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  Nguyen, T.-Q.; Wu, J.; Doan, V.; Schwartz, B. J.; Tolbert, S. H. Science 2000, 288, 652
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  Li, D.; Babel, A.; Jenekhe, S. A.; Xia, Y. Adv. Mater. 2004, 16, 2062– 2066
  42
  Nanofibers of conjugated polymers prepared by electrospinning with a two-capillary spinneret
  Li, Dan; Babel, Amit; Jenekhe, Samson A.; Xia, Younan
  Advanced Materials (Weinheim, Germany) (2004), 16 (22), 2062-2066CODEN: ADVMEW; ISSN:0935-9648. (Wiley-VCH Verlag GmbH & Co. KGaA)
  Uniform nanofibers composed of poly[2-methoxy-5-(2-ethylhexyloxy)-1,4-phenylenevinylene] and its blends with poly(3-hexylthiophene) are formed by co-electrospinning their solns. with poly(vinylpyrrolidone) (PVP) through a coaxial capillary system, followed by extn. of the PVP phase (see Figure). Compared with spin-cast films, the fibers show reduced phase sepn. and improved energy-transfer efficiency.
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Supporting Information
Theoretical model rational and details, fibers optical properties, and technical details about SNOM measurements. This material is available free of charge via the Internet at http://pubs.acs.org.
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Conformational Evolution of Elongated Polymer Solutions Tailors the Polarization of Light-Emission from Organic NanofibersClick to copy article linkArticle link copied!

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Abstract

Introduction

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Conjugated Polymer Nanofibers

Polarized Emission

SNOM

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Conformational Evolution of Elongated Polymer Solutions Tailors the Polarization of Light-Emission from Organic Nanofibers
Click to copy article linkArticle link copied!