Modulus- and Surface-Energy-Tunable Thiolâene for UV

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Modulus- and Surface Energy-Tunable Thiol-ene for UV Micromolding of Coatings Yuyang Du, Jun Xu, John D. Sakizadeh, Donovan G. Weiblen, Alon V. McCormick, and Lorraine F. Francis ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b06339 • Publication Date (Web): 29 Jun 2017 Downloaded from http://pubs.acs.org on July 2, 2017

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ACS Applied Materials & Interfaces

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Modulus- and Surface Energy-Tunable Thiol-ene for UV

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Micromolding of Coatings

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Yuyang Du, Jun Xu, John D. Sakizadeh, Donovan G. Weiblen, Alon V.

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McCormick* and Lorraine F. Francis*

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Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis,

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Minnesota 55455, United States

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ABSTRACT

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Micromolding of UV curable materials is a patterning method to fabricate microstructured

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surfaces that is an additive manufacturing process fully compatible with roll-to-roll systems. The

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development of micromolding for mass production remains a challenge because of the

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multifaceted demands of UV curable materials and the risk of demolding-related defects -

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particularly when patterning high aspect-ratio features. In this research, a robust micromolding

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approach is demonstrated that integrates thiol-ene polymerization and UV LED curing. The

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moduli of cured thiol-ene coatings were tuned over two orders of magnitude by simply adjusting

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the acrylate concentration of a coating formulation, the curing completed in all cases within 10

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seconds of LED exposure. Densely packed 50-µm-wide gratings were faithfully replicated in

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coatings ranging from soft materials to stiff highly crosslinked networks. Further, surface energy

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was modified with a fluorinated polymer, achieving a surface energy reduction of more than a

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half at a loading of 1 wt%, and enabling tall (100-µm) defect-free patterns to be attained. The

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demolding strengths of microstructured coatings were compared using quantitative peel testing,

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showing its decrease with decreasing surface energy, coating modulus, and grating height. This

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micromolding process, combining tunability in thermomechanical and surface properties, makes

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thiol-ene microstructured coatings attractive candidates for roll-to-roll manufacture. As a

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demonstration of the utility of the process, superhydrophobic surfaces are prepared using the

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system modified by the fluorinated polymer.

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KEYWORDS: replica molding, surface microstructures, high throughput, thiol-ene, mechanical

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properties, surface modification, peel test

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ACS Paragon Plus Environment

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INTRODUCTION The fabrication of nano/micro scale surface structures is essential to a variety of

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applications, including textured coatings1, 2, optical films3, superhydrophobic surfaces4, and

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flexible electronics5, 6. Photolithography, a conventional nanofabrication technique, involves

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curing through a photomask to achieve the spatially-selective reaction of photoresists and a

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subsequent developing step to remove unreacted materials. The development of

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photolithography for the mass production of structured surfaces is limited due both to the waste

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involved with etching and to its stringent requirements on processing conditions, so nanoimprint

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lithography has received growing research interest as an alternative lithographic technique.7, 8

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Unlike photolithography, nanoimprinting creates a surface pattern based on the mechanical

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deformation of materials using a stamp with surface relief structures, followed by pattern

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solidification assisted with either heat or ultraviolet (UV) exposure.

Figure 1. Schematic diagram of the UV micromolding process: (a) a liquid coating forms conformal contact with a patterned mold, (b) the coating is exposed to UV light through the mold, and (c) the cured coating obtains surface microstructures after the mold is peeled off.

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While nanoimprint lithography focuses on the fabrication of nanostructures and involves

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etching to transfer the pattern from the photoresist to the underlying (typically rigid) substrate,

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micromolding of UV curable materials (UV micromolding) is promising for micro-scale

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applications in which the final structured coating is the aim of the process. In micromolding, a

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liquid UV curable coating is deposited onto a substrate, a flexible mold with micro-scale features

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is then pressed into conformal contact with the coating, UV exposure solidifies the coating with

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the mold in place, and the mold is peeled off to obtain a coating with surface structure (Figure

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1).9 With the application of roller molds, UV micromolding is compatible with continuous roll-

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to-roll (R2R) processes and large area substrates.10, 11 Although processing rates on the order of

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meters-per-minute have been reported for continuous versions of imprint processes12,


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improvements in throughput are still needed for UV micromolding processes to be viable for

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mass production. Key challenges in boosting the throughput include the complete filling of mold

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cavities, the fast curing of coatings, and the clean release of surface microstructures; all of these

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are closely tied to the development of new UV-curable materials with low viscosity, high

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reactivity with no danger of oxygen inhibition, and prospects for low surface energy when

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needed.8

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Acrylates and epoxies are widely used UV curable materials for coating applications, so

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are potential candidates for UV micromolding. Acrylates have high curing speeds and desirable

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physical properties, such as a range of achievable modulus, using the large variety of

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commercially available monomers and oligomers.13-15 For UV micromolding, however, acrylates

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have two main disadvantages. First, due to their free-radical, chain-growth photopolymerization

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mechanism, the curing of acrylates is inhibited by oxygen and thus may lead to uncured edges

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under ambient curing conditions. Second, high levels of volume shrinkage during curing limit

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the accuracy of replicated patterns and lead to stress, which can cause cracking or distortion of

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the final product. The photoreaction of epoxy, proceeding via cationic, step-growth

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photopolymerization, is insensitive to oxygen and develops less shrinkage during curing.16

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However, the curing speeds of epoxy resins are generally lower, and their viscosities are higher

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compared to those of acrylates.

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As an alternative class of UV curable materials, thiol-ene chemistry is mediated by free-

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radical, step-growth photopolymerization, and it displays many properties desirable for high-

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throughput UV micromolding applications.17, 18 Benefitting from the free-radical reactions, thiol-

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ene systems undergo a rapid photopolymerization process at ambient temperatures; due to the

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step-growth character, the reaction is insensitive to oxygen, and it generates polymer networks

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with high homogeneity and little shrinkage. Furthermore, the range of commercially available

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acrylates can be incorporated homogeneously into the thiol-ene reactions, offering prospects for

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changing thermomechanical properties (e.g., modulus, Tg).19-21

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Previous work has shown the promise of thiol-ene coatings in imprint lithography of

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nano-scale features. Carter and coworkers22-24 reported the pattern replication capability of thiol-

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ene coatings, while Lin et al.25, 26 extended the coating chemistry to siloxane containing

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monomers to fabricate microstructures with enhanced etch resistance and thermal stability. Quite

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recently, Stadlober and coworkers27 used coatings based on polyurethane-acrylate oligomers and


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thiol monomers in R2R nano-imprinting processes and showed successful patterning at large

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areas. However, these studies, being focused on nanostructures, only addressed the fabrication

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of shallow features (less than 1 µm in height). The replication of micro-scale features with high

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aspect ratios still presented a substantial processing challenge. Such high aspect ratio microscale

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features are crucial to practical applications, e.g. enhancing the efficiency of electronic devices

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with tall conductive channels, or providing robust superhydrophobic microtextures. With such

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large features as those needed for micromolding, patterned coatings can be deformed due to the

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strong adhesion between the coating and the mold in the presence of interfacial microstructures;

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demolding may even fail, resulting in permanent contamination of the mold. It has been shown

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at the nano-scale that coatings with low surface energies will reduce demolding-related defects26,

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processes (where the whole mold is separated from a nanostructured coating in a normal

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separation direction).28-30

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, and investigations of the demolding force have shed light on this aspect of the nanoimprinting

In micromolding, especially for roll-to-roll processes, the peeling separation is preferred over the normal separation because of a much smaller peeling force and the roller geometries.

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Peel tests have been widely used in determining the interfacial strength between polymer-

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polymer surfaces and a recent work successfully applied it to study the geometrical effects in the

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peeling of a commercial adhesive from patterned PDMS surfaces.31 Therefore, conventional

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peel tests should provide an effective method to study the demolding strength in UV

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micromolding processes.

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In the present work, a pattern replication approach utilizing the thiol-ene

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photopolymerization and the UV micromolding process is reported that, within 10 seconds of

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UV LED exposure, replicates micro-scale gratings with a high pattern density at high accuracy.

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Starting from a well-established thiol-ene monomer system, the available thermomechanical and

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surface properties of cured thiol-ene networks were expanded through simple and robust

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materials design; further, performance relevant to demands of roll-to-roll micromolding was

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tested. The thermomechanical properties of cured coatings were adjusted over a wide range by

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changing the acrylate content in the formulation, so the coating can range from stiff polymer

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networks to soft materials. The surface energy of the thiol-ene coating was significantly reduced

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through the introduction of a fluorinated polymer at a 1 wt% loading, and this will allow

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microstructures with high aspect ratios to be attained. Finally, the demolding behavior of micro-


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scale gratings is evaluated with 90° peel tests, independently examining changes in the coating

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properties and in the pattern geometry, to find quantitative relations between the demolding

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strength and the coating surface energy, the Young’s modulus, and the pattern height. As one

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example of the promise of our approach using UV LED micromolding process combined with

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fluorinated thiol-ene coatings, we illustrate the fabrication of superhydrophobic surfaces with

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potential for high throughput.

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EXPERIMENTAL SECTION Materials. Figure 2 shows the chemical structures of materials. Pentaerythritol

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tetrakis(3-mercaptopropionate) (PTMP), 1,3,5-triallyl-1,3,5-triazine-2,4,6(1H,3H,5H)-trione

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(TTT), and 2,2-dimethoxy-2-phenylacetophenone (DMPA) were purchased from Sigma-Aldrich.

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1,6-Hexanediol ethoxylate diacrylate (HEDA) was provided by Sartomer. All the materials were

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used as received. Fluorinated acrylate copolymers were synthesized by free radical

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polymerization. See Supporting Information for more details.

Figure 2. Chemical structures of materials employed in coating formulations: (a) pentaerythritol tetrakis(3-mercaptopropionate) (PTMP), (b) 1,3,5-triallyl-1,3,5-triazine-2,4,6(1H,3H,5H)-trione (TTT), (c) 2,2-dimethoxy-2-phenylacetophenone (DMPA) (d) 1,6-hexanediol ethoxylate diacrylate (HEDA), and (e) fluorinated acrylate copolymer (F-PMMA).

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Coating Formulation Preparation. The coatings were formulated based on a

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stoichiometric balance of thiol functional groups (Tf) and ene functional groups (Ef). In the

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preparation of thiol-ene(-acrylate) coatings, acrylate functionalities (Af) were added to the base


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thiol-ene mixture, with the molar percentages of all three functional groups (Tf, Ef, and Af)

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totaling 100%.32 The fluorinated additive (Ff) was added at 1 wt%, and DMPA was added at

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0.25 wt%, of the coating mixture. Compositions of all the UV-curable coating formulations

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employed in this study are summarized in Table 1. All the components of a formulation were

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added into a glass vial and stirred at room temperature for 4 h. The glass vials were wrapped

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with aluminum foil to avoid any prepolymerization. In the preparation of TEF, F-PMMA was

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dissolved in 2 ml tetrahydrofuran (THF) and then added to the mixture of thiol (11.72 g) and ene

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(7.96 g) monomers. No phase separation was observed for the TEF mixture. Table 1. Composition of coating formulations Ef Formulation Tf (mol%) TE 50 TEA5 47.5 TEA10 45 TEA25 37.5 TEF 50 [a] All formulations include 0.25 wt% of DMPA.

(mol%) 50 47.5 45 37.5 50

Af (mol%) 5 10 25 -

Ff (wt%) 1

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Unpatterned Coating Preparation. Unpatterned coatings were prepared by casting

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coating formulations on glass substrates using a wire-wound rod. Especially, the applied

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fluorinated formulation was allowed to dry under ambient conditions for 5 min to allow removal

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of THF. Liquid coatings were exposed to UV light (at a wavelength of 365 nm and an intensity

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of 100 mW/cm2) from a LED source (Tangent Industries Inc.) at room temperature. The

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thickness of the obtained coating was approximately 100 µm. Cured coatings were directly used

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for subsequent characterizations unless otherwise noted.

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Kinetic Analysis. Fourier transform infrared spectroscopy (FTIR) was performed to

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monitor the photopolymerization kinetics of prepared coatings on a Nicolet 6700 FTIR

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spectrometer in the transmission mode.19, 33 Samples were prepared by sandwiching a thin layer

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of the uncured formulation (~25 µm) between two sodium chloride plates separated by a round

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steel spacer (0.025 × 25 mm). Samples were partially cured with the LED lamp at 100 mW/cm2

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using different curing times, and each sample was transferred to the FTIR spectrometer within 30

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seconds; with such a short dark time before measurement, dark reaction conversion was assumed

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to be negligible compared to the reaction during illumination. The FTIR spectrometer operated

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at 64 scans per spectrum and a 2 cm-1 resolution. Resulting conversions of functional groups

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were analyzed using the Thermo Scientific OMNIC software by integrating the area under


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characteristic IR absorption peaks (the thiol peak at 2570 cm-1, the ene peak at 3083 cm-1, and

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the acrylate double bond peak at 810 cm-1). Integrated values of the peak area were proportional

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to the corresponding concentrations of the functional group in the sample. The functional group

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conversions (x) at various curing times were calculated from the peak area after curing for time t

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(At) and that of uncured samples (A0): =

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−

× 100%

(1)

Thermal Analysis. Differential scanning calorimetry (DSC) was used to characterize

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thermal properties of cured coatings on a TA Q1000 calorimeter (TA instrument). Samples (5 -

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10 mg) were sealed in aluminum pans and heated from -50°C to 100°C at a heating rate of

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10°C/min. The second heating scan was used to determine the glass transition temperatures.

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Thermogravimetric (TGA) analysis was carried out using a Shimadzu TGA-50

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thermogravimetric analyzer and heating at 10°C/min up to a temperature of 500°C. The polymer

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degradation temperature was determined as the value at 10% weight loss. All samples were

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tested under a nitrogen atmosphere.

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Mechanical Properties. Standard tensile tests were conducted to measure the moduli of

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UV cured freestanding films on an RSA G2 solid analyzer (TA Instruments). Samples were

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prepared from coatings on silicone release films and cut into 2.5 cm × 1 cm × 100 µm specimens.

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Stress-strain behaviors were recorded at a crosshead speed of 1 mm/min under ambient

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conditions.

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Surface Analysis. Contact angle measurements were performed to measure the static

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contact angles of cured coatings using a drop shape analyzer (Krüss). The equipment was

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enclosed in a transparent plastic box to reduce exposure to organic solvents. Water and

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diiodomethane (Sigma-Aldrich) were used as solvents at a droplet size of 2 µL. Two specimens

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were prepared for each coating formulation, and three drops were made on each specimen. The

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contact angle of a droplet was measured at room temperature using the Advance software. The

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surface energy of a cured coating was analyzed according to ASTM D7490 based on the Owens-

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Wendt-Kaelble equation (Equation 2): 1 + cos / / = + 2

(2)

= +

(3)


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where θ is the contact angle of the testing liquid on the tested coating surface, and are the

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surface tension of the testing liquid and the surface energy of the solid coating, respectively, and

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and are the dispersion component and the polar component, respectively. The surface

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tension values for water and diiodomethane are known (water: = 72.8 dyn/cm, = 21.8

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dyn/cm, and = 51.0 dyn/cm; diiodomethane: = 50.8 dyn/cm, = 49.5 dyn/cm, and =

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1.3 dyn/cm). X-ray photoelectron spectroscopy (XPS) measurements were performed on an

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SSX-100 XPS system (Surface Science Laboratories) equipped with a monochromatic Al Kα X-

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ray source operating at 200 W under 10-9 Torr. All spectra were recorded with the Esca Capture

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software. Binding energies were calibrated with respect to C 1s at 285 eV.

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PDMS Mold Fabrication. The PDMS prepolymer and its curing agent (Sylgard-184,

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Dow Corning) were mixed at a 10:1 weight ratio and then vacuum degassed for approximately

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20 min until the mixture was clear. The mixture was poured onto an SU-8 master mold and

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degassed for another 20 min to remove any entrapped air bubbles. See Supporting Information

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for the photolithography process for the fabrication of SU-8 master molds. The PDMS layer was

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cured in an atmospheric oven at 90°C for 2 h. After cooling down to room temperature, the

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cured PDMS was peeled off from the SU-8 master to obtain the PDMS mold.

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Micromolding Experiments. Figure 1 presents the schematics of the procedures of a

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micromolding process. The prepared coating formulation was applied on a100-µm-thick PET

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substrate (3M) using a wire-wound bar with a wet film thickness of approximately 100 µm.

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Before the coating application, the substrates were roughened with silicone carbide powders,

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rinsed with acetone and water, and blow-dried with air. The PDMS mold was placed on the

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liquid coating from one end of the mold to the other with the patterned side facing down. The

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wetting front of the liquid coating moved along the grating direction to avoid air entrapment

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between the liquid coating and the mold. The coating was cured through the PDMS mold by UV

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LED light (365 nm and 100 mW/cm2). The PDMS mold has >95% transmittance to UV with a

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wavelength of 365 ± 5 nm, as shown in the Supporting Information Figure S7. After UV curing,

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the mold was peeled off by hand to obtain the microstructured coating.

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Patterned Coating Imaging. A Hitachi S4700 scanning electron microscopy (SEM)

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was used for obtaining images of fabricated microstructures in an oblique view and a cross-

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sectional view. The coating cross-sections were cryomicrotomed at -120°C using a Leica UC6


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microtome instrument equipped with a diamond knife. Samples were sputter-coated with a thin

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layer (5 nm) of Iridium before imaging.

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Peel Strength Measurement. The peel strength was measured with 90° peel tests on an

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RSA G2 solid analyzer, which operated in the spring mode using an upper tension fixture and a

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3D-printed test stage. The assembly of a cured coating and a PDMS mold was used for peel

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testing without any post-cure treatment. The coating on a PET substrate was attached to a piece

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of aluminum foil, which connected to the force sensor. The PDMS mold was mounted on a glass

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slide that provided rigid support and fixed on the test stage. The cured coating on substrate was

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peeled off from the mold on the solid analyzer, during which the peel force was recorded as a

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function of the displacement of the crosshead. A photo of the peeling apparatus midway through

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the demolding step is presented in Figure 3. For our experimental set-up, the peeling rate equals

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the crosshead speed, 30 mm/min, and the peeling angle is approximately 90°. The peel force (F)

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was normalized with respect to the sample width (w), yielding the peel strength (P): " = #/$

(4) 34

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In the present work, the peeling strength was calculated using the area method . For each

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measured curve of peeling strength versus displacement, the area under the plateau region was

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integrated and then divided by the corresponding displacement. The middle part of the plateau

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region (D = 5 mm) was used for calculation to remove any influences of the instabilities at the

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edge of the patterned area. The long and flexible aluminum foil could reduce the systematic

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error in the peeling angle by increasing the distance between the force sensor and the sample (d

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= 120 mm). The theoretical change of the peeling angle during a test was ± 1.2° [tan) + ,].

*

Figure 3. Schematic representation and picture of the apparatus for the peel strength measurement. Samples were peeled off at 90° and 30 mm/min.


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RESULTS AND DISCUSSION

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Photopolymerization Kinetics

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The FTIR results show that, for all the thiol-ene(-acrylate) formulations, the acrylate

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oligomers attained full conversions immediately upon UV exposure, while the conversions of

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thiol and ene functional groups increased rapidly initially and then both reached a limiting

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conversion value. (See Supporting Information Figure S1.) Bowman et al. observed similar

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results in thiol-ally ether-acrylate ternary systems and attributed it to the reactivity difference of

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between acrylates and ally ethers.32 Because the electron density on the acrylate double bond is

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higher than that of the ene double bond, the chain polymerization of acrylates dominates the

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initial stage of the curing process, generating acrylate oligomers. Following this stage, thiol and

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ene functional groups react to form crosslinked networks. Unfortunately, estimating the

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statistically expected crosslink density from this data is challenging due to the complex nature of

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the monomers. The limiting curing extent of thiol groups is slightly higher than that of ene

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groups because a proportion of thiols undergo chain transfer reaction with the acrylate double

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bonds.

Figure 4. Thiol group conversion as a function of the irradiation time in thiol-ene(-acrylate) photopolymerizations monitored by FTIR. The curing kinetics of TEF was identical to that of TE and thus not shown in the plot. Samples contained 0.25 wt% DMPA and were exposed to a 365 nm UV LED lamp at 100 mW/cm2.

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We investigated the effects of acrylate addition on the curing kinetics of coatings,

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represented by the conversion of thiol functional groups. As shown in Figure 4, all the

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formulations exhibited high photopolymerization rates and achieved plateau conversions after

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approximately 1.5 seconds of exposure. Curing was carried out for 10 seconds to ensure the


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limiting conversions were reached for all the samples. According to the light intensity and the

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exposure time, the curing dosage of the thiol-ene(-acrylate) coating system is approximately 150

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mJ/cm2, which is among the lowest dosage values for UV coatings used in micromolding

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applications. The curing dosage of a urethane acrylate coating35 is reported to be 300 mJ/cm2

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and the lowest reported value is 80 mJ/cm2 for an epoxy silicone coating16. Therefore, the thiol-

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ene(-acrylate) coating system shows great promise for use in high-throughput micromolding

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applications.

263

The limiting conversion increased with the acrylate concentration. Limiting thiol

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conversions for the coatings based on TE, TEA5, TEA10, and TEA25 were 77%, 81%, 86%, and

265

96%, respectively. The relatively low limiting conversion of TE results from the fact that highly

266

crosslinked networks restrict the mobility of the reactive species and thus further reactions of

267

functional groups. The addition of acrylate oligomers leads to higher limiting conversions,

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which can be explained by an enhanced mobility afforded by the flexible ether linkage along

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with a lower functionality of acrylate oligomers. In a previous kinetic analysis by Cramer and

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Bowman33 of TE coatings with a low intensity light source (0.8 mW/cm2), cured coatings

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reached 60% conversions after 2 minutes. With the higher intensity LED in this study, curing

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leads much faster to a higher limiting conversion. The light intensity effect might be attributed

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to either or both of two effects: free volume36 and temperature37. As the photopolymerization is

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faster at higher intensities, the actual free volume of the cured polymer networks exceeds the

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equilibrium free volume, resulting in increased motilities of reactive species and thus enhanced

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conversions. Moreover, high reaction rates may lead to more heat accumulation from the rapid

277

reaction, which may increase the conversion by enhancing the mobiliity of polymer networks.

278 279 280

Tunability of Thermomechanical Properties It has been well documented that thermomechanical properties of thiol-ene(-acrylate)

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coatings depend on the chemical structures of monomers, the crosslinking density, and the curing

282

extent. For instance, flexible ether groups lead to low glass transition temperatures and low

283

moduli, and ring structures result in stiff materials with glass transitions at higher temperatures.38

284

Moreover, the moduli of cured coatings increase with the crosslinking densities of polymer

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networks, which can be achieved by decreasing the length of spacer between functional groups,

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increasing the monomer functionality, or increasing the concentration of crosslinkers.23, 39


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Additionally, the curing extent of functional groups also has an effect on mechanical properties

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due to the plasticization effect of unreacted dangling chains.

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Based on our current knowledge of the thermomechanical properties of thiol-ene(-

290

acrylate) coatings38, 40, 41, we selected a tetrafunctional thiol monomer (PTMP) and a trifunctional

291

ene monomer (TTT), which photopolymerize into materials presenting a high Tg and a high

292

modulus, along with a diacrylate oligomer (HEDA) with a long and flexible spacer between

293

functional groups. In the curing of samples, a LED light source was used considering that its

294

high intensity (100 mW/cm2) could increase the curing speed and that the narrow wavelength

295

distribution (365 ± 5 nm) could reduce heat generation during curing.42 Samples were exposed

296

to a sufficient dosage of UV light (10 seconds) to ensure limiting conversions were attained and

297

cured samples were not subject to any thermal annealing process after UV curing.

298

The effects of the acrylate addition on the thermal properties of thiol-ene(-acrylate)

299

coatings were investigated using DSC. To complement the DSC, TGA analysis was carried out;

300

weight loss began at ~ 350°C, well above the maximum temperature used in DSC (see Table 2

301

and Supporting Information). For TE, an exothermic peak was observed during the first heating

302

cycle and disappeared in the second (See Supporting Information Figure S2), which could be due

303

to the continued polymerization of residual monomers at elevated temperatures, further

304

confirming the limited conversion value of TE from the FTIR experiment. In cases of coatings

305

containing acrylates, the first and the second heating cycles overlapped since almost complete

306

conversions were attained in these coatings. The second heating cycles were used for the

307

comparison of the glass transition behaviors of the coating system. Shown in Figure 5a, the DSC

308

curves of all the coatings exhibited a distinct and narrow glass transition region (~15˚C in width),

309

indicating that the polymer networks were homogenous upon the acrylate addition. The obtained

310

Tg of the TE (45˚C) was in agreement with the literature value (45˚C)38. Furthermore, the glass

311

transition temperatures decreased with increasing acrylate concentrations as summarized in

312

Table 2. The observed trend in Tg could be attributed to a combined effect of the chemical

313

structures and the network crosslink densities. It is inferred that acrylate oligomers introduce

314

flexible ether linkage into polymer networks and thus enhance the segmental mobility; at the

315

same time, the lower functionality of acrylates reduces the crosslinking density, resulting in a

316

less restricted segment movement.


12

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Figure 5. (a) DSC curves of thiol-ene(-acrylate) networks at a 10°C/min heating rate. Results from the second heating cycle were presented. (b) Stress-strain relationship of UV-cured thiol-ene(-acrylate) coatings from tensile tests under ambient conditions. Inset: low strain data. The crosshead speed was 1 mm/min.

317 Table 2. Thermomechanical properties of UV-cured thiol-ene(-acrylate) films Formulation Tg (˚C) E (MPa) Tdegradation (˚C) TE 45 1740 ± 153 368 TEA5 34 994 ± 190 361 TEA10 27 440 ± 101 357 TEA25 5 13 ± 2 352 TEF 41 522 ± 38 368 [a] The standard deviations of tensile moduli were based on testing results of five samples for each formulation.

318 319

The mechanical properties of cured thiol-ene(-acrylate) films were investigated by tensile

320

testing under ambient temperature, providing information on coating performances under typical

321

processing conditions and application environments. Samples were prepared using conditions

322

consistent with those for micromolding experiments (including curing conditions and thickness)

323

and were not subject to thermal annealing after curing. The elastic moduli of cured films, as

324

calculated from the slope of the stress-strain curve at small strains (0 - 2%), decreased by two

325

orders of magnitude with increasing the acrylate concentration. Consistent with previous

326

results20, 38, TE based samples showed a high modulus of 1.74 GPa, which is comparable to those

327

of commonnn thermoplastics43 (PMMA and PS) and epoxy resins44. With the addition of

328

acrylate oligomers, the significantly decreased modulus was due to the flexibility of acrylate

329

oligomers and a lower crosslink density.


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330

Page 14 of 26

Although the main purpose of mechanical testing in this study is to evaluate the elastic

331

deformation of thiol-ene(-acrylate), an interesting observation of plastic deformation was found.

332

In previous work by McNair et al.40, UV cured and thermally annealed thiol-ene based polymers

333

were found to be brittle with stress increasing monotonically with the strain until fracture; in

334

contrast, yielding was observed for TE networks containing thiourethane linkages. This yielding

335

was attributed to the sliding of polymer chains in the linear thiourethane domain due to the

336

disruption of hydrogen bonds. We suspect that yielding observed in our study is the result of

337

partial curing of monomers (as revealed in FTIR measurements), where unreacted functional

338

groups can cause sliding of polymer segments while under tension. In addition, the lower strain

339

rate, lack of thermal annealing, and smaller sample thickness used in our study might contribute

340

to the observed yielding. Additional experiments revealed that the yielding is less significant for

341

thicker samples tested at a higher strain rate (see Supporting Information Figure S4). More work

342

is needed to fully understand the plastic deformation of unannealed samples, and this would be

343

relevant and needed for future development of high throughput, continuous roll-to-roll

344

fabrication.

345 346 347

Surface Energy Reduction The surface energy of thiol-ene coatings was modified using a fluorinated acrylic

348

copolymer (F-PMMA), synthesized from methyl methacrylate and 1H, 1H, 2H, 2H-

349

perfluorodecyl acrylate by free radical polymerization (details in the Supporting Information).45

350

The chemical structure of F-PMMA is composed of a long fluorinated side chain to reduce the

351

surface energy and an acrylate backbone to ensure good miscibility in the liquid coating. THF

352

was chosen as the solvent for the fluorinated coating (TEF), because it is a good solvent for both

353

F-PMMA and the thiol-ene monomer mixture. Before curing, the TEF formulation is a clear

354

liquid without any phase separation. After coating application, the excess solvent was allowed to

355

evaporate for 5 min and during which F-PMMA is expected to migrate to the free surface, which

356

was experimentally confirmed, as discussed below.

357

The static contact angles of the thiol-ene coating (TE) and the fluorinated coating (TEF)

358

were examined using water and diiodomethane (oil with a low polarity). Both the water contact

359

angle (θw) and the diiodomethane contact angle (θi) significantly increased upon only 1 wt%

360

addition of F-PMMA (Figure 6a-d). Further increasing the F-PMMA loading resulted in


14

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361

marginal changes of contact angles, suggesting that the surface coverage of fluorinated side

362

groups saturates at 1 wt%. According to the surface elemental analysis by XPS (Figure 6e), TEF

363

exhibited strong signals related to the binding energy of F 1s at 688.6 eV46, which did not exist in

364

the spectrum of TE coatings. The results support the mechanism of surface energy reduction

365

proposed in previous studies on fluorinated surfaces47, 48: fluorinated side groups tend to

366

assemble at the air and liquid interface to reduce the surface tension of the coating system,

367

forming a fluorocarbon-rich layer at the surface. The surface segregation of fluorinated groups is

368

enthalpy driven due to the difference in surface tension between monomers and the fluorinated

369

polymers.49 In addition, the high molecular weight of F-PMMA facilitates the migration of

370

polymer chains to the liquid-air interface due to less entropic gain in the mixing process

371

comparing to the monomeric counterpart.

Figure 6. Cured TE surfaces with (a) a water droplet and (b) a diiodomethane droplet compared with cured TEF surfaces with (c) a water droplet and (d) a diiodomethane droplet. The TEF coating contains 1 wt% F-PMMA. The droplet size was 2 µL. (e) XPS survey spectra for cured TE and TEF coating surfaces.

372 Table 3. Surface properties of cured TE and TEF coatings. Formulation θw (˚) θi (˚) TE 63.4 ± 2.3 19.5 ± 1.3 TEF 98.6 ± 0.8 76.5 ± 2.8 [a] Standard deviations of contact angles were calculated from 12 measurements.

γ (mJ/m2) 52.2 20.2

373 374

The surface free energies (γ) of coatings were determined based on the Owens-Wendt-

375

Kaelble method (the calculation is described in the experimental section). According to Table 3,

376

the surface energy of a TE coating is similar to that of PET films (49.8 mJ/m2)50, a relatively


15


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377

high value among polymeric surfaces. With fluorinated additives, TEF exhibits a remarkable

378

low surface energy comparable to that of PDMS (~19.6 mJ/m2)51. Compared to a previous

379

study52 on the surface modification of thiol-ene coatings using a fluorinated monomer (at ~20

380

wt%), our approach based on the polymeric fluorinated additive is more environmentally

381

friendly and works at much lower loadings.

382

Page 16 of 26

Thermomechanical properties of films prepared from TEF were evaluated with DSC and

383

tensile testing. Comparing the DSC curves of TE and TEF in Figure 5a, the Tg value slightly

384

decreased upon the addition of F-PMMA, which could be explained by the plasticization effect

385

that the chain ends of F-PMMA enhance the mobility of the segments on polymer networks.

386

Moreover, tensile testing results show that the fluorinated additives resulted in a lower elastic

387

modulus of the cured coating. Since F-PMMA is not expected to appreciably affect the

388

crosslinking density of thiol-ene networks at such a low loading, it is inferred that the observed

389

decrease of modulus comes from a higher level of heterogeneity due to the segregation of

390

fluorinated polymers.

391

Before discussing coating performances in the micromolding process, we summarize the

392

materials properties of the prepared thiol-ene coating system including the modulus and the

393

surface energy (Figure 7). The elastic moduli vary across two orders of magnitude with the

394

change of the acrylate content. With 1 wt% of fluorinated acrylate copolymers, the surface

395

energy of the thiol-ene coating is reduced by more than a half. The ability to tune the modulus

396

and the surface energy independently is critical to a better understanding of the roles of material

397

properties in micromolding.

Figure 7. Summary of critical material properties of the thiol-ene(acrylate) coating system: (a) modulus tunability and (b) surface energy tunability.


16

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398


Pattern Fidelity

399

To examine the pattern replication capability of the prepared thiol-ene(-acrylate)

400

formulations, we tested micro-scale grating patterns with a high pattern-density (a line width of

401

50 µm and a spacing of 50 µm). Typical UV micromolding procedures (Figure 1) were

402

performed under ambient conditions using a PDMS stamp, which was fabricated from an SU-8

403

master. For the master fabrication, microstructures made of SU-8 photoresists on a silicon wafer

404

were directly used and the feature height was determined by the coating thickness of SU-8. As

405

opposed to the commonly used silicon masters (microstructures are fabricated by etching of

406

wafer and subsequent removal of photoresists), SU-8 masters possess smoother sidewall profiles.

407

It should be noted that the master and the mold were not subject to any surface treatment, e.g.

408

silanization. The PET substrates were roughened to enhance the adhesion between coating and

409

substrate.

Figure 8. SEM images of the surface microstructures of thiol-ene(-acrylate) coatings with different moduli. (a) TE coating with a modulus of 1.74 GPa, (b) TEA5 coating with a modulus of 994 MPa, (c) TEA10 coating with a modulus of 440 MPa, and (d) TEA25 coating with a modulus of 13 MPa. All pattern dimensions were 50 µm width by 50 µm spacing by 50 µm height. Images were taken at 45°.

410 411

Coating formulations with various acrylate concentrations were used to prepare gratings

412

of 20-µm-height. The patterns were successfully replicated from the PDMS mold in all the

413

formulations after 10 seconds of UV exposure at 100 mW/cm2. The fast replication of patterns

414

could be attributed to the efficiency of the photoinitiator (DMPA), the high reactivity of the

415

monomers, and the high intensity of the LED source. Moreover, the replicated surface

416

microstructures showed straight lines, sharp corners, and smooth sidewalls; no significant

417

difference was observed among the patterns in spite of the vastly different moduli of cured

418

coatings (Figure 8). Therefore, combining the LED curing technique and thiol-ene(-acrylate)


17


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419

coatings, the fabrication of micro-scale patterns was achieved within seconds at high accuracy

420

and the attainable material properties in UV micromolding processes were expanded.

421

Page 18 of 26

The performance of TEF in the UV micromolding process was evaluated, with the goal

422

of eliminating demolding-related defects by lowering the surface energy of cured coatings. With

423

10 seconds of UV illumination, cured TEF coatings could be easily peeled off from the mold.

424

Gratings of the same pattern density but higher aspect ratios (with 50 µm and 100 µm heights)

425

were successfully replicated using TEF. The resulting microstructures showed well-defined

426

shapes as presented in typical SEM images (Figure 9).

Figure 9. SEM images of microstructured TEF coatings with various grating heights: (a) 20 µm, (b) 50 µm, and (c) 100 µm. All the gratings were 50 µm wide and 50 µm spaced. Inset images are the cross sections of the microstructures.

427 Table 4. Modulus and surface energy values of elastomeric mold materials used in UV micromolding processes Material Modulus (MPa) Surface energy (mN/m2) TEF 522 20.2 PDMS (Sylgard 184)53 2.0 19.6 Polyurethane acrylate (PUA)15 20 - 320 20 - 60 Ethylene tetrafluoroethylene (ETFE)54, 55 1200 15.6 155 23.6 Perfluoropolyether (PFPE)51, 56

428 429

In addition to the direct application as UV curable resins, the TEF-based formulation

430

represents an exciting alternative to the traditional mold material PDMS. It is known that the

431

attainable pattern resolution of a mold material is determined by the modulus and the surface

432

energy; typically, a high modulus helps maintain the integrity of surface features while a low

433

surface energy is required for clean mold release. Table 4 compares the modulus and the surface

434

energy of TEF to those of mold materials currently used in UV micromolding processes.

435

Accordingly, TEF outperforms PDMS by a much higher modulus, and it is comparable to other

436

newly developed mold materials in consideration of the theoretically achievable resolution.

437


18

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438 439


Demolding Analysis In the replication of 100 µm tall gratings, demolding failed in all the formulations except

440

for TEF: TEF patterned coatings were peeled off from the PDMS mold without any

441

contamination of the mold; on the contrary, coatings without the fluorine addition were

442

completely removed from the PET substrate, adhering to the mold surface. (See Supporting

443

Information Figure S8.) In the thiol-ene(-acrylate) coatings, the modulus and the surface energy

444

of cured coatings can be adjusted independently, allowing for the investigation of their roles

445

during demolding. In the subsequent discussion, the demolding behaviors of different coating

446

formulations were quantitatively compared using peel tests. Because the peel strength is known

447

to correlate with the peeling angle and the peeling rate57, the reproducibility of tests was ensured

448

by controlling the peeling rate at 30 mm/min and the peeling angle at 90° (with an approximated

449

variance of ±1.2° in each measurement).

Figure 10. (a) Representative peel strength versus displacement of the peel front for patterned TEA10 and TEF, which have different surface energies but similar moduli. Grating height: 50 µm. (b) Averaged peel strengths of TEA10 and TEF at grating heights of 20 µm, 50 µm, and 100 µm. No peel strength is reported for TEA10 with 100 µm features because demolding failed. (c) Representative peel strength versus displacement for TEAx coatings with various moduli but similar surface energies. Grating height: 20 µm. (d) Averaged peel strength plotted against TEAx coatings at grating heights of 20 µm and 50 µm. All the patterns investigated are 50 µm wide and 50 µm spaced and the peel direction was perpendicular to the gratings. Errors bars in (b) and (d) were based on repeated tests of 5 samples.


19


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Page 20 of 26

450

The effects of surface energy on the peel strength were investigated using formulations of

451

TEF (γ = 20.2 mJ/m2, E = 522 MPa) and TEA10 (γ = 50.1 mJ/m2, E = 440 MPa) with the peeling

452

direction perpendicular to gratings. Typical curves for the peel strength as a function of the

453

displacement of the peel front is shown in Figure 10a. For both coatings, the peel strength rises

454

sharply when the peel front reaches the patterned area and remains at a plateau value across the

455

patterned region. The demolding behaviors of TEF (1% Ff) and TEA10 (no Ff) were compared

456

at different grating heights and summarized in Figure 10b. The peel strength is lower with the

457

fluorine addition because of the reduced surface energy of a coating. Moreover, patterns with

458

higher aspect ratios result in increased peel strength, which could be attributed to a larger

459

interfacial area between the microstructured coating and the mold. Especially, in the case of

460

TEA10 coating with a 100 µm grating height, the peel strength needed to separate the coating

461

from the mold exceeds the adhesion strength between the coating and the substrate. As a result,

462

fracture occurs at the coating/substrate interface, which explains the demolding failure observed

463

in previous micromolding experiments.

464

We further studied the effects of elastic modulus on the peeling strength of

465

microstructured coatings. With an increase of the acrylate content, the elastic modulus of the

466

cured coating decreases without a significant change in the surface energy. Figure 10c shows

467

that peeling strengths decrease with increasing the acrylate content of coating formulations. In

468

Figure 10d, similar trends were observed in 20 µm and 50 µm gratings (data for 100-µm-high

469

gratings were not available because demolding failed in coatings without any fluorine addition).

470

Therefore, in addition to the surface energy, decreasing the modulus of a coating also

471

significantly reduces the peel strength (demolding work). We propose the following explanation:

472

the work done by the external peel force is mainly consumed by the work of adhesion and the

473

elastic energy stored in the system during demolding. Gratings on the coating surface need to

474

deform elastically to be released from the mold. As the coating gets softer, less energy is

475

consumed during the elastic deformation of surface features, and thus the apparent work done by

476

the external force is smaller.

477 478

One illustration of the promise of this approach: fabricating superhydrophobic

479

microstructured TEF coatings


20

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480


Superhydrophobic surfaces have attracted growing interests due to their great potential in

481

applications such as self-cleaning surfaces and catalyst supports. A low-surface energy TEF

482

coating was patterned with 10 µm wide and 10 µm spaced microgratings using the micromolding

483

process demonstrated in Figure 11a. The fabricated microstructured TEF coating shows

484

superhydrophobic wetting, as evidenced by a blue-dyed water droplet sitting on the surface; the

485

surface contact angle of the patterned coating reaches 153° (Figure 11b). The

486

superhydrophobicity can be attributed to the inherent low surface energy of TEF and to the

487

densely packed microscale surface topography. Moreover, the patterned TEF surface is

488

mechanically robust with an elastic modulus of 500 MPa and thermally stable up to ca. 370°C.

489

The combination of the UV LED micromolding and the modified thiol-ene chemistry provides a

490

powerful approach with high throughput potential for the fabrication of superhydrophobic

491

surfaces.

Figure 11. (a) SEM image of the surface patterns of TEF. Gratings width is 10 µm, grating spacing is10 µm, and grating height is 20 µm. (b) Static water contact angle for TEF coatings patterned with micrograting arrays. Inset: photograph of a blue-dyed water droplet on the surface of a patterned TEF coating.

492 493

CONCLUSION

494

In summary, a UV curable coating system based on the thiol-ene chemistry was

495

developed for high-throughput micromolding processes compatible with R2R manufacturing.

496

Starting from commercially available monomers, thiol-ene(-acrylate) coatings cured into

497

materials with drastically different mechanical properties, with the Young’s moduli tunable over

498

two orders of magnitude through the variation of the acrylate concentration in a formation. A

499

fluorinated acrylate copolymer was designed for the surface modification of thiol-ene coatings,

500

which has not been realized in previous studies of thiol-ene systems. The thiol-ene coating


21


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Page 22 of 26

501

attained excellent hydrophobicity and oleophobicity at only 1 wt% loading of fluorinated

502

additives, achieving a surface energy as low as that of PDMS. With 10 seconds of UV LED

503

exposure, densely packed microstructures were accurately replicated in thiol-ene(-acrylate)

504

coatings of vastly different elastic moduli. Using the fluorinated coating formulation, we

505

replicated surface microstructures of high aspect ratios without any defect formation, which was

506

typically observed during the demolding of tall features. At controlled peeling angle and rate,

507

the demolding strengths of micro gratings were measured with 90° peel tests. It was suggested

508

that failures to release the structured coating from the mold would occur when the demolding

509

strength between the coating and the mold exceeded the adhesion strength at the

510

coating/substrate interface. Coatings with high moduli or/and surface energies and tall features

511

resulted in increased demolding strengths and thus demolding defects.

512 513

ASSOCIATED CONTENT

514

Supporting Information

515

The Supporting Information is available free of charge on the ACS Publications website.

516

Functional group conversions of TEA10 measured by real-time FTIR, DSC curves of TE

517

from the first and second heating cycles, synthesis of F-PMMA, size exclusion

518

chromatography and nuclear magnetic resonance spectroscopy of F-PMMA, fabrication

519

of the master mold by photolithography, and representative images of the demolding

520

failure of TEA10 and the clean release of TEF for 100-µm-tall features.

521 522

AUTHOR INFORMATION

523

Corresponding Authors

524

*E-mail: [email protected], [email protected]

525

Notes

526

The authors declare no competing financial interest.

527 528

ACKNOWLEDGMENTS

529

We thank Dr. Heonjoo Ha and Dr. Qiang Guo for the help with UV-Vis and TGA measurements

530

respectively. We are grateful to Dr. David Giles for the helpful discussion on peel testing. This

531

work was supported by the Industrial Partnership for Research in Interfacial and Materials


22

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532

Engineering (IPRIME), University of Minnesota. Parts of this work were performed at the

533

Nano-Fabrication Center and the Characterization Facility, University of Minnesota, which

534

receives partial support from NSF through the MRSEC program.

535 536

REFERENCES

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for Table of Contents use only

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Modulus- and Surface Energy-Tunable Thiol-ene for UV

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Micromolding of Coatings

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Yuyang Du, Jun Xu, John D. Sakizadeh, Donovan G. Weiblen, Alon V.

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McCormick* and Lorraine F. Francis*

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Department of Chemical Engineering and Materials Science, University of Minnesota,

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Minneapolis, Minnesota 55455, United States

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Modulus- and Surface-Energy-Tunable Thiolâene for UV

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