Secondary-Structure-Mediated Hierarchy and ... - ACS Publications

Article Cite This: Biomacromolecules XXXX, XXX, XXX−XXX

pubs.acs.org/Biomac

Secondary-Structure-Mediated Hierarchy and Mechanics in Polyurea−Peptide Hybrids Lindsay E. Matolyak,† Chase B. Thompson,‡ Bingrui Li,†,# Jong K. Keum,∥ Jonathan E. Cowen,⊥ Richard S. Tomazin,⊥ and LaShanda T. J. Korley*,‡,§

Downloaded via DURHAM UNIV on July 13, 2018 at 04:04:50 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

†

Department of Macromolecular Science and Engineering, Case Western Reserve University, 2100 Adelbert Road, Cleveland, Ohio 44106-7202, United States ‡ Department of Materials Science and Engineering, University of Delaware, 127 The Green, Newark, Delaware 19716, United States § Department of Chemical and Biomolecular Engineering, University of Delaware, 150 Academy Street, Newark, Delaware 19716, United States ∥ Center for Nanophase Materials Sciences and Chemical and Engineering Materials Division, Oak Ridge National Laboratory, 1 Bethel Valley Road, Oak Ridge, Tennessee 37830, United States ⊥ Swagelok Center for Surface Analysis of Materials, Case Western Reserve University, 10900 Euclid Avenue, Cleveland, Ohio 44106-7202, United States S Supporting Information *

ABSTRACT: Peptide−polymer hybrids combine the hierarchy of biological species with synthetic concepts to achieve control over molecular design and material properties. By further incorporating covalent cross-links, the enhancement of molecular complexity is achieved, allowing for both a physical and covalent network. In this work, the structure and function of poly(ethylene glycol) (PEG)−network hybrids are tuned by varying peptide block length and overall peptide content. Here the impact of poly(ε-carbobenzyloxy-L-lysine) (PZLY) units on block interactions and mechanics is explored by probing secondary structure, PEG crystallinity, and hierarchical organization. The incorporation of PZLY reveals a mixture of α-helices and β-sheets at smaller repeat lengths (n = 5) and selective α-helix formation at a higher peptide molecular weight (n = 20). Secondary structure variations tailored the solid-state film hierarchy, whereby nanoscale fibers and microscale spherulites varied in size depending on the amount of α-helices and βsheets. This long-range ordering influenced mechanical properties, resulting in a decrease in elongation-at-break (from 400 to 20%) with increasing spherulite diameter. Furthermore, the reduction in soft segment crystallinity with the addition of PZLY resulted in a decrease in moduli. It was determined that, by controlling PZLY content, a balance of physical associations and self-assembly is obtained, leading to tunable PEG crystallinity, spherulite formation, and mechanics.

1. INTRODUCTION

driven by changes in chemical composition (e.g. cross-linking and peptide sequence). The ability to program self-assembly into materials as a means of mechanical reinforcement is an attractive prospect to researchers, allowing for the tuning of materials’ morphology and mechanics through discrete changes in chemical composition. Neat peptide sequences, however, suffer from thermal and pH sensitivity, often leading to unfolding of secondary structures and uncontrolled aggregation of the sequences.4 Conjugation of peptide segments to synthetic polymers can improve the stability of peptide assemblies by reinforcing secondary structure formation and preventing uncontrolled aggregation of unfolded domains.4 For example, the coupling of poly(ethylene glycol) (PEG) chains to an

As the demand for robust synthetic materials grows, researchers often turn to natural systems to gain insight into the precise mechanics afforded by their structures. The assembly of amino acid residues along a peptide or protein through an array of physical interactions affords secondary structures such as β-sheets and α-helices, which can contribute to a material’s stiffness or elasticity by unfolding under tension.1 For example, the muscular protein titin contains a series of α-helices and a single-stranded β-sheet that, under tension, unravel as a means of energy dissipation.2 Elastin, a major protein found in the extracellular matrix, achieves stability and reinforcement through the use of covalent crosslinks between short peptide segments, leading to high extensibility and significant elastic recovery.3 These natural materials display properties dictated by the hierarchical arrangement of their physical interactions, which are, in turn, © XXXX American Chemical Society

Received: May 12, 2018 Revised: June 20, 2018

A

DOI: 10.1021/acs.biomac.8b00762 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules

Figure 1. Reaction scheme for covalently cross-linked polyurea−PZLY network hybrids, where n = 5 or 20 and x = ∼44. In addition, the sample nomenclature is included along with the anticipated secondary structure.

of PEG can be exploited as a means to substantially influence the hierarchical assembly in PEG−peptide conjugated systems, leading to architectures that are tunable through changes in interactions between PEG and peptide segments. In our previous study, we incorporated a poly(εcarbobenzyloxy-L-lysine)-block-PEG-block-poly(ε-carbobenzyloxy-L-lysine) (PZLY−PEG−PZLY) triblock soft segment into a linear polyurea/polyurethane framework, demonstrating that changes in hard segment length through chain extension drove phase evolution and shifted mechanics by promoting competitive hydrogen-bonding interactions between the peptide, hard segment, and PEG.14 In this research, PZLY− PEG−PZLY triblock copolymers with different peptide lengths (5 or 20 repeat units) were introduced into a PEG-based polyurea cross-linked network; these elastomeric, hard segment cross-link junctions confine the PEG−peptide soft segments and isolate their hydrogen-bonding interactions to the soft phase. This approach allows for the tunability of peptide secondary structure, PEG crystallinity, and long-range ordering based on the physical interactions between PEG and the peptide segments, which is anticipated to modulate morphological or mechanical responsiveness to external stimuli by imparting changes to these physical associations through hydration, pH, or temperature.

alanine-rich polypeptide sequence improved the thermal stability of α-helical structures under acidic conditions, allowing for near-complete recovery of secondary structure of the polypeptide after heating.5 Furthermore, the application of polymer−peptide hybrid materials can also offer pathways to generate materials that mimic the morphological and mechanical behavior of biological systems such as spider silk,6,7 elastin,3,8 and collagen,9 closely reproducing assembly responses without having to express entire peptide sequences. Of the many synthetic polymers available for conjugation to peptidic sequences, PEG has seen the most exposure, due in part to its hydrophilicity and innate biocompatibility.10 Whereas PEG has been widely applied as a means to stabilize peptide structures, the potential for PEG to participate in hydrogen-bonding interactions and to crystallize offers avenues for further manipulation of peptide behavior to elicit morphological and mechanical changes. For instance, in a poly(γ-benzyl-L-glutamate)−PEG−poly(γ-benzyl-L-glutamate) (PBLG−PEG−PBLG) triblock copolymer, the presence of low volume fractions of PBLG segments led to the formation of microphase-separated domains, whereas higher fractions of peptide segments inhibited the crystallinity of PEG and induced phase mixing.11 PEG’s crystallinity can also be used to inhibit β-sheet assembly in certain systems while amplifying α-helical formations, attributed to the ability of PEG and poly(ethylene oxide) (PEO) to form 72 helices that closely mimic the pitch of α-helices.5,10,12 In one instance, amyloid β peptide fragments were conjugated to PEG, resulting in the disruption of fibrilization in peptide sequences with weaker propensity to form β-sheet fibrils.13 This antagonistic behavior

2. EXPERIMENTAL SECTION 2.1. Materials. Desmodur N 3300 A (HDI trimer) with 21.8% isocyanate content was used as received from Bayer MaterialScience. ε-carbobenzyloxy-L-lysine (ZLY or Z), extra-dry dimethylacetamide (DMAc), and extra-dry N-dimethylformamide (DMF) were used as B


Article

Biomacromolecules

under vacuum until a constant weight was achieved to remove any residual THF. The gel fraction was calculated using eq 123 É ÅÄÅ ij m − mf yzÑÑÑÑ ÅÅ zzÑÑ100 gel fraction (%) = ÅÅÅ1 − jjj i j mi zzÑÑÑ ÅÅ (1) ÅÇ k {ÑÖ

received from Sigma-Aldrich. Tetrahydrofuran (THF) was purified utilizing a Vacuum Atmosphere’s solvent purification system. α,ωbis(amine) poly(ethylene glycol) (PEG, 3400 g/mol) was purchased from Sinopeg Biotech (China) and dried at 130 °C under vacuum for a minimum of 16 h before use. ε-carbobenzyloxy-L-lysine Ncarboxyanhydride (ZLY-NCA) was prepared according to literature procedure.15,16 2.2. Synthesis. 2.2.1. Poly(ε-carbobenzyloxy-L-lysine)n-blockpoly(ethylene glycol)-block-poly(ε-carbobenzyloxy- L -lysine) n . (PZLY-b-PEG-b-PZLY) triblocks were synthesized via amine-initiated NCA polymerization utilizing a previously reported method.16,17 For example, the synthesis of Z20-b-PEG-b-Z20 was performed in a flamedried 100 mL round-bottomed flask with a magnetic stirrer. In a nitrogen (N2) atmosphere glovebox, 4.0 g (13 mmol) of ZLY−NCA was dissolved in 25 mL of 25/75 THF/DMAc mixture. To the solution, 1.0 g (0.3 mmol) of PEG was predissolved in 25 mL of THF and added via a drop funnel over 30 min to the reaction flask. The reaction was stirred at room temperature (RT) for 20 h. The final solution was precipitated into diethyl ether; the precipitate was filtered and dried under vacuum until constant weight was achieved (∼24 h). A sticky wax-like solid with yield of 82% was recovered. 1H NMR was utilized to characterize the triblock molecular weight (Figure S1). Additional molecular weight characterization techniques, such as gel permeation chromatography, were not conducted due to apparent fractionation of the sample solutions when passing through a filter during sample preparation.16−22 This solubility is the result of the lysine protecting group, which imparts hydrophobic character and structural changes that prevent full dissolution without the aid of heat, salts, or solvents.16−22 2.2.2. Synthesis of Network Polyurea Hybrids. To fabricate the cross-linked polyurea networks, a commercially available, trifunctional (f ≈ 3) HDI (Desmodur N 3300 A) was utilized. Because of the polyfunctionality of the isocyanate, the ratio of isocyanate groups to amine groups can be tuned to vary the degree of cross-linking. For this investigation, a calculated isocyanate/amine ([NCO]/[NH2]) ratio of 1.125 was used to achieve cross-linked networks, modeling the method previously utilized.18 A flame-dried, 100 mL round-bottomed flask with magnetic stirrer was placed inside a glovebox under a N2 atmosphere. As an example, 0.25 g (0.04 mmol) of PZLY5-b-PEG-bPZLY5 triblock and 0.74 g (0.22 mmol) of PEG were dissolved in a flask containing 50 mL of DMF. To dissolve the triblock copolymer and PEG, the solution was heated to 60 °C and agitated using a magnetic stir plate. Upon complete dissolution, the solution was cooled to room temperature, 0.1 g (0.2 mmol) of predissolved HDI trimer in 20 mL of DMF was added to the flask, and the solution was allowed to stir at room temperature for 1 min prior to pouring the mixture into a round glass Petri dish (∼9 cm diameter). The reaction was placed inside a desiccator under N2 for 1 h before heating in an oven at 80 °C under N2 purge. The reaction was held under these conditions for 1 week, after which the sample was tacky to the touch. The film was then placed under vacuum at 80 °C for 24 h to remove any residual DMF. After 24 h, the film was allowed to cool to room temperature while under vacuum. The resulting film (∼10 wt % peptide content) was clear, colorless, and flexible. Full synthetic quantities for each sample with corresponding amounts of each component (PEG, HDI, and triblock) are given in Figure S2. Figure 1 details the chemical synthesis of network polymer−peptide hybrids. It is noted that full dissolution is necessary for good film formation. Depending on PZLY length and content, it was necessary to heat the triblock and PEG solution to 80 °C under agitation to promote dissolution. Additional DMF was also added to the higher peptide content reactions to prevent premature gelation. Under these requirements, cooling to room temperature was an important step before reaction with HDI trimer. 2.3. Characterization. 2.3.1. Degree of Covalent Cross-Links via Gel Fraction. The gel fraction of each covalently cross-linked film was determined via THF Soxhlet extraction completed over a 24 h period. A poly(tetrafluoroethylene) (PTFE)-coated fiberglass fabric with 181/2 gauge holes was used as a thimble. Each film was weighed before and after extraction; the extracted film was collected and dried

where mi is the initial weight of the film and mf is the weight of the film after extraction. Resulting gel fractions are tabulated in Figure S2. 2.3.2. Attenuated Total Reflectance−Fourier Transform Infrared Spectroscopy (ATR-FTIR). The secondary structures of cross-linked PZLY−polyurea films were investigated using a Cary 680 Agilent ATR-FTIR system fitted with a diamond ATR accessory. Peaks in the 1600−1680 cm−1 region were fit with Gaussian functions; the location and width of the peaks were determined from the second derivative spectra, facilitating the assignment of PZLY secondary structures via calculation of the relative areas under the peaks.24,25 2.3.3. Atomic Force Microscopy (AFM). A Dimension 3100 Veeco Digital Instruments (Bruker equipped with a NanoScope IIIa controller and Quadrex signal processor) was used to observe the morphology of the network polyurea films. Phase images were collected in tapping mode with silicon tips from Bruker (333−380 kHz, 110−140 μm). 2.3.4. Scanning Electron Microscopy. Scanning electron microscopy (SEM) images were taken on a JEOL JSM-6510LV apparatus operating at a voltage of 5 kV. Prior to performing SEM, the samples were sputter-coated with palladium (∼7.5 nm thickness) using a Denton desk IV sputtering system operating at a 45° tilt. Freely available software (ImageJ) was utilized for size characterization by measuring the diameter of spherulites.26−28 Each spherulite was measured twice to provide an estimated diameter for the individual spherulite, and a series of 100 spherulites was measured for each sample to determine the average spherulite diameter. 2.3.5. Small-Angle X-ray Scattering. Small-angle X-ray scattering (SAXS) data were acquired at the Center for Nanophase Materials Sciences (CNMS) in Oak Ridge National Laboratory on an Anton Paar SAXSess mc2 apparatus. The scattered beam was recorded on a CCD detector (PI-SCX, Roper) with a pixel resolution of 2084 × 2084 and pixel dimensions of 24 × 24 μm2. The data collection time was 20 min. For the measurements, X-rays were generated at 40 kV/ 50 mA at a beam wavelength of λ = 1.541 Å (Cu Kα radiation). The generated X-ray beam was slit-collimated using a Kratky camera, giving rise to a beam size of 18 mm (length) × 0.6 mm (width), and the collected SAXS data were desmeared and expressed as intensity versus q, where q = (4π sin θ)/λ after subtraction of detector dark current and background scattering. Data were processed using Origin software version 2016. 2.3.6. Differential Scanning Calorimetry. Differential scanning calorimetry (DSC) was performed on a TA Instrument Q100 at a heating rate of 10 °C/min over the temperature range of −80 to 150 °C under a nitrogen atmosphere. The degree of crystallinity was calculated by dividing the melting enthalpy (ΔH) by the enthalpy of melting required for a 100% crystalline PEG sample (ΔHf = 196.8 J/ g) of the second heating cycle.29 2.3.7. Polarized Optical Microscopy. To analyze the birefringence of the films, polarized optical microscopy (POM) images were obtained under cross polarizers on a Leica DM 2500 M microscope equipped with a quartz halogen lamp. All samples were examined in the solid-state at 10× magnification with 0.30 aperture. For in situ temperature measurements, a mounted heat stage from Instec (STC200) was utilized to vary the temperature from 20 to 150 °C at a rate of 5 °C/min. Initial images were taken at 5 °C intervals after isothermal holds for 2 min until 60 °C (±0.3 °C), at which point the samples no longer exhibited birefringence. Samples were then heated to 150 °C and allowed to equilibrate for 5 min before rapid cooling to room temperature. After a 5 min hold at room temperature, the samples were re-examined to detect morphological changes. 2.3.8. X-ray Diffraction. X-ray diffraction (XRD) patterns were acquired using a Bruker D8 Discover with VÅNTEC 500 2D detector with Co Kα as the X-ray source (1.789 Å). The sample-to-detector distance was 21.35 cm, as determined via calibration with a silver C


Article

Biomacromolecules behenate standard. Two frames were collected beginning at 17° and moved in 20° increments for an angular range of 2 to 52°. Each frame was collected for 600 s. 2.3.9. Tensile Testing. A Zwick/Roell mechanical testing instrument equipped with a 500 N load cell was used to examine tensile behavior. Films for tensile testing were cut using a dogbone steel die according to a modified ASTM D1708 with the dimensions scaled down by a factor of 2. During the cutting process, the sample was placed between two sheets of Mylar to minimize stress concentrations at the edges of the material. All samples were elongated to failure at room temperature at a constant strain rate of 100% the initial gauge length per minute (∼10 mm/min). The modulus was determined using the 1% secant method due to the nonlinearity of the tensile curves. The reported mechanical properties were an average of a minimum of three samples.

block. The HDI trimer was utilized as the hard segment for comparison with the 1,6-hexamethylene diisocyanate (HDI; f = 2) linear analogs previously reported and its success as a crosslinker for other PU systems.18,31,32 Prior to network formation, the triblock PZLYn-b-PEG-bPZLYn was synthesized via ring-opening polymerization of εcarbobenzyloxy-L-lysine N-carboxyanhydride (NCA) initiated by amine-terminated PEG.15,16 PEG was chosen to enhance solubility, participate in hydrogen bonding, and aid in biointerfacing.10,33,34 The PZLY repeat length is defined by the ratio of NCA to PEG. The block lengths of PZLY5-b-PEGb-PZLY5 and PZLY20-b-PEG-b-PZLY20 were confirmed by 1H NMR (Figure S1) with calculated molecular weights of 6 and 13.9 kg/mol, respectively. To form the network hybrid, PZLYnb-PEG-b-PZLYn (n = 5 or 20) and excess PEG−bisamine are reacted with the HDI trimer via step-growth polymerization. The PZLY content is controlled by the addition of excess PEG. The synthetic pathway from triblock synthesis to PZLY−PU network formation is shown in Figure 1. To obtain solid-state films for mechanical analysis, network formation must occur simultaneously with film fabrication. Utilizing previously explored conditions for film fabrication, an amine to isocyanate [NH2]/[NCO] ratio of 1.125 was utilized.18 The gel fractions obtained via Soxhlet extraction are reported in Figure S2. The gel fraction of the network control, PEG−HDI is 90%, which is higher than the 55−65% of PZLY−PUs. This difference in gel fraction is attributed to a decrease in solubility of triblocks in DMF compared with PEG−bisamine. By keeping the gel fraction consistent, the focus of this work is attributing the confined peptide/PEG interactions to the resulting hierarchial formation and mechanical properties. In the next sections, PZLY content and repeat length are used to tune PZLY/PEG associations that are analyzed via secondary structure changes, the appearance of lamellar morphologies, and the presence of spherulite formation. In this way, modulating physical associations via synthetic motifs can be related to resulting material properties. 3.2. Secondary Structure Characterization. To probe the secondary structure of PZLY at varying peptide repeat lengths and weight percentages, ATR-FTIR analysis was conducted on the PZLY−PU hybrids. ATR-FTIR of the carbonyl stretching region 1600−1680 cm−1 was examined for relative secondary structures due to the high sensitivity to peptide structures in this region. Specifically, amide I bands give rise to various stretching frequencies due to the unique molecular geometry and hydrogen bonding pattern of peptide secondary structures. β-sheets are expressed between 1620 to 1630 cm−1, and bands of 1645 to 1655 cm−1 were assigned to α-helical formation.17,35 With PZLY repeat lengths of n = 5, primarily β-sheet secondary structures are expected to dominate. In contrast, when n = 20, an α-helical intramolecular organization is anticipated.16,17 Interestingly, the β-sheet content in PZLY5-bPEG-b-PZLY5 was limited to ∼40% (Figure S3). It is proposed that this unanticipated secondary structure preference is due to the intermolecular interactions between PZLY and PEG, which reduces β-sheet self-assembly.11 Additionally, the reaction solvent (DMF) is helicogenic, further facilitating intramolecular hydrogen bonding in the PZLY block.36 Thus primarily α-helical structures were displayed in both PZLY5-bPEG-b-PZLY5 and PZLY20-b-PEG-b-PZLY20.

3. RESULTS AND DISCUSSION A previous investigation of linear PEG-based, PZLY polyurea hybrids revealed an increase in α-helical formation and amine/ ether hydrogen bonding with increasing peptide content, suggesting enhanced intermolecular hydrogen bonding between the PZLY segments and the soft blocks in nonchain extended films.30 PZLY in these samples acted as the hard block, increasing the physical cross-links within the solid-state films and enhancing the mechanical extensibility up to ∼900%. Whereas the peptide associations with hard and soft segments drove the phase behavior and mechanics in linear PEG−PUs, the addition of a covalent network may impart further opportunities to control architecture and material properties. The overlay of covalent cross-linking within the hard block will locate the peptide physical associations within the soft block between the permanent network points. This synthetic design allows for probing of solid-state material properties with respect to ordered structures selectively in the soft domain. In contrast with our previous study of covalently cross-linked, PDMS−peptide hybrids, the PEG block is expected to associate with the PZLY unit within the soft domain, providing an additional handle to control material response via a balance between inter- and intramolecular hydrogen bonding and hierarchical organization. 3.1. Synthetic Strategy. Poly(ε-carbobenzyloxy-L-lysine) (PZLY or Z) was chosen as the peptide component due to the ability to control secondary structure as a function of block length.16,19 A block length of either 5 or 20 repeats was investigated to promote either β-sheets (intermolecular) or αhelical (intramolecular) character within the peptide block, respectively.7,25 Additionally, potential side-chain functionality can be achieved via further deprotection of the lysine block, expanding potential applications of these materials toward responsive systems in future studies. Peptide content (10 and 20 wt %) was varied within the network hybrids to observe the influence of peptide hydrogen bonding on nanostructure formation and mechanics. The following nomenclature was utilized for PZLY (Z) hybrids: ZX−Y, where X is the peptide repeat length of either 5 or 20 and Y is the weight percentage within the sample (10 or 20 wt %) (Figure 1). For example, a sample containing the triblock poly(ε-carbobenzyloxy-Llysine)20-block-poly(ethylene glycol)-block-poly(ε-carbobenzyloxy-L-lysine)20 (PZLY20-b-PEG-b-PZLY20) and an overall peptide content of 10 wt % is labeled Z20−10. Control network films without PZLY were also fabricated using the nomenclature PEG−HDI. The hexamethylene polyisocyanate (HDI trimer) was the cross-linker for the covalent PZLY−PU networks, placing the permanent network within the hard D


Article

Biomacromolecules

sheet formation. In contrast, Z20−10 and Z20−20 PU network hybrids selectively exhibit α-helices, which, while favorable for a PZLY block length of 20, may also be due to the aforementioned reasons for the promotion of α-helical conformations. Specifically, with the use of a nonhelicogenic solvent, such as THF, selectivity toward α-helices was no longer observed and a mixture of secondary structures became apparent (Figure S4). Thus, several physical contributions cooperatively aid in the assembly of secondary structures in PZLY−PU network hybrids, resulting in an overall favorability for intramolecular hydrogen bonding. Figure 3 represents network PEG−PU hybrids with increasing α-helical content ranging from 60% (Z5−20) to 100% (Z20−10 and Z20−20). The features in red indicate peptide secondary structure, whereby β-sheets are shown as an arrow and α-helices are shown as a ribbon. The variation in peptide structure can influence the organization of PEG (blue), affecting the percent crystallinity as well as the type of crystal packing.11,12,37,38 While Figure 3 represents the impact of secondary structure on soft phase ordering, it is acknowledged that additional hydrogen bonding between and within the hard block as well as peptide side-chain interactions may also exist, although not represented here. It is anticipated that the enhanced interaction between PZLY α-helices and PEG soft segment will result in varying soft phase organization, which is further explored through microscopy and X-ray analysis. 3.3. Morphological Development in PZLY−PU Network Hybrids. It is expected that the variation in inter- and intramolecular physical associations between the Z5−Y and Z20−Y hybrids plays a role in the long-range ordering and resulting morphology of the network hybrids. To probe hierarchy within the samples, several microscopy techniques were employed. Initially, POM was used to examine the birefringence and spherulitic structures of the PEG block.39−42

Analysis of the network PZLY−PU secondary structure showed a strong preference for α-helix formation (Figure 2).

Figure 2. ATR-FTIR spectra of network PUs in the amide I region to analyze secondary structure. The dashed lines indicate the location of secondary structure peaks. The calculated relative β-sheet content is also reported.

Z5−10 and Z5−20 polyurea networks exhibit an increase in βsheet structures with increasing PZLY content, attaining a maximum β-sheet content of 40%. The mixture of α-helices and β-sheets in the Z5 network series is a result of several factors. In a prior investigation, network formation in peptide− PU hybrids has shown that geometric constraints lead to weaker hydrogen bonding between neighboring peptides.18 Additionally, the film fabrication conditions, such as use of a helicogenic solvent and prolonged heat exposure, aid in helical structure formation. Finally, competitive intermolecular associations between the PZLY block and PEG disrupt β-

Figure 3. Proposed secondary structure influence on polyurea−PZLY hybrids. As the relative α-helix content increases (from Z5−20 to Z5−10 and further Z20−10 and Z20−20) the soft segment ordering (BLUE) was tuned. Regions in red indicate peptide, while objects in green are hard segments. E


Article

Biomacromolecules

Figure 4. Microscopy images of network PU hybrids. In the top row, images from POM show maltese cross patterns. POM images are at 10× magnification (scale bar = 50 μm). In the center row, SEM images highlight the spherulite structures with a scale bar of 50 μm. In the bottom row, AFM phase images of 1 × 1 μm show a nanoscale fibrous morphology.

is utilized to probe the impact on average spherulite size and potential PEG ordering. To better visualize the spherulitic morphology and measure the spherulite size of network films, SEM was utilized (Figure 4).26−28 The spherulite diameters of PEG−HDI and Z5−10 are similar (∼115 μm). However, at the higher PZLY content (20 wt %) in the Z5 series, the spherulitic structures are faint, with only diffuse interfaces vaguely apparent, indicating a more phase-mixed structure. Upon increasing the peptide block length (Z20−10), the spherulite diameter increases to ∼215 μm, and a sharp interface is observed. Whereas spherulites are still formed in Z20−20, their diameter is reduced to ∼160 μm with only a diffuse interface displayed. However, the Z20−20 spherulite diameter is notably higher than the control films, suggesting that competitive interactions of α-helices promote increased intermolecular associations between PEG and PZLY. This higher spherulite size for the Z20 series may be the result of an increase in α-helical content that aids in crystal formation due to the similarity in crystal sizes of the PEG 72 helix and an α-helix, which will be explored further via wide-angle X-ray scattering studies.12,44,46,47 Thus, the incorporation of peptides into network PEG−PU films was harnessed to control PEG crystal growth and hierarchy based on secondary structure and peptide content. To further analyze the lamellar assembly of semicrystalline PEG, atomic force microscopy was used to visualize the hard and soft segments on the surface of PU films. Phase images of PUs display hard segments as lighter regions and the soft domain as darker regions due to the higher phase offset of the higher modulus component. As observed in Figure 4, PEG forms well-defined lamellar fibers with evidence of hard segment aggregation noted as bright spots in PEG−HDI films. PEG lamellar organization is present in all samples except for Z5−20, which appears to be fully phase-mixed, consistent with POM and SEM results. With the aid of microscopy methods, network PZLY−PU hybrids exhibit a hierarchical organization of structure at three levels: molecular

In control PEG−HDI polyurea networks, a well-defined maltese cross morphology is exhibited, whereby the dark regions correspond to where the axis of the polarizer is oriented parallel to the principle axis of a folded crystal chain (Figure 4). Optical anisotropy, shown as a maltese cross, is the result of PEG crystals folding into a lamellar structure.43−45 This microscopic evidence indicates that the PEG soft segment maintains its crystalline character upon the incorporation of a covalent network and suggests microphase segregation of the PU hybrid. With the incorporation of PZLY, variations in spherulite size and interdomain sharpness occur. While Z5−10 exhibited an overall spherulite size and domain sharpness similar to that of the control film, increasing the peptide content (Z5−20) results in highly disrupted spherulite formation. This disruption is believed to be an effect of intermolecular interactions between PZLY and PEG that inhibit crystal growth. With increasing PZLY block length (Z20−10 and Z20−20), a spherulitic morphology is observed with a noticeable increase in spherulite size compared with the control and Z5 series. The apparent trend in spherulite size could be the result of peptide secondary structure, intramolecular hydrogen bonding between PZLY and PEG, or molecular weight variations between the 5 and 20 PZLY block.11,12,44 For instance, the presence of a maltese cross pattern has also been seen in peptide−PEG−peptide triblocks containing either PBLG28,37 or PBLA12 blocks. In each example, the ratio of PEG to peptide influences the secondary structure and phase behavior. Floudas et al. demonstrated that a volume fraction of PBLG less than 0.26 resulted in a more phase-separated structure, allowing for PEG to crystallize and form long-range ordering.11 Additionally, Tanaka et al. altered the secondary structure of PBLA from α-helices to β-sheets, resulting in less ordered PEG domains.12 Thus, multiple factors contribute to the crystallinity of PEG and overall phase separation in solid-state hybrids. For this current work, further microscopic analysis via scanning electron microscopy (SEM) F


Article

Biomacromolecules

Figure 5. X-ray spectra of network PU hybrids. On the left, SAXS data indicate phase separation with associated d spacings. On the right, XRD data highlight PEG crystalline peaks.

disappearance of the 72 helix PEG structure and decrease in PEG crystallinity are attributed to the PZLY/PEG interactions. Furthermore, the crystalline structure of PEG and poly(ethylene oxide) (PEO) are well known,38,56 suggesting that the 33.0 and 47.4° peaks are due to intermolecular packing in PEG.37,42 The appearance of a modified PEG crystal structure is only apparent in the Z5-Y series, indicating that PEG crystal growth is influenced by the β-sheet structures. When PZLY20 was incorporated into the network PU, the 72 helix PEG structure was maintained, which is hypothesized to be the result of similar packing sizes of an α-helix and a 72 helix.11,12 This variation in crystal structure with peptide block length in PEG hybrids has been previously seen, suggesting that the relative ratio of peptide to PEG as well as the secondary structure influences the ability of PEG to crystallize.11,12 For example, Akashi et al. discovered that PBLA−PEG−PBLA triblocks adopted an α-helix/72 helix/α-helix conformation, resulting in cocrystallization due to the similar pitch within each helical structure.12 Whereas PZLY secondary structure had a strong influence on PEG crystal structure, the overall PZLY content resulted in a decrease in the height of the PEG crystalline peaks (22.2 and 27.0°). This reduction in peak intensity indicates an enhancement in PEG/PZLY interactions with increasing PZLY content. Thus altering the soft segment packing and phase behavior in PZLY−PU hybrids was achieved by tuning the physical associations within the peptide block. Additionally, it is noted that no well-defined secondary structures are seen in the XRD spectra. The absence of a peptide peak could be a result of structural irregularity in the PZLY segment or overlapping peaks for PEG crystals and secondary structures.12,57 To further probe the crystallinity within network PUs, differential scanning calorimetry was utilized. Second heating curves of PU films are shown in Figure 6 with corresponding percent crystallinity of the PEG block calculated using the melting transition. PEG−HDI had a melting PEG transition at 46 °C with a crystallinity of 50%, which was expected based on previously reported network PEG−PUs.39,58,59 Upon the addition of PZLY, the melting transition of PEG was still apparent, but the degree of crystallinity was decreased. For Z5−10 and Z5−20, the crystallinity decreased to 32 and 24%, respectively. When the PZLY block length was increased to 20 (Z20−10 and Z20−20), the crystallinity was higher than that of the Z5−Y hybrids with crystallinity values of 44 and 35%, respectively. This enhanced crystallinity for Z20−Y indicates that the secondary structure has an influence on the ability of

(PZLY organization), nano-(PEG lamellae), and micro(PZLY/PEG spherulite) assembly. These results motivated further investigation of phase behavior via thermal and X-ray analysis. 3.4. Crystallinity and Film Mechanics PZLY−PUs. To explore the phase behavior between hard and soft segments in network PUs, small-angle X-ray scattering (SAXS) and X-ray diffraction (XRD) studies were conducted. SAXS experiments allow for the observation of nanometer-scale organization, such as domain sizes of the hard phase in PUs. XRD examines the crystalline structure and packing of a material on the angstrom scale. To begin, SAXS analysis revealed a single, broad peak (Figure 5) observed for all PU hybrids and control, indicating some degree of phase separation present. PEG−HDI displays a well-defined peak at ∼0.42 nm−1 (∼15 nm) attributed to the incompatibility of the PEG segments and the multifunctional hard segment. Upon PZLY (10 wt %) incorporation, Z5−10 and Z20−10 display a similar domain spacing of ∼16 nm, indicating that the PZLY block did not influence phase separation. However, peak broadening was evident at higher PZLY loadings (Z5−20 and Z20−20). This peak broadening is indicative of a wider distribution of hard segment sizes or a more phase-mixed morphology in these samples, as seen in POM and AFM.48−50 Additionally, a slight shift was seen in Z20−20 to ∼0.35 nm−1 (∼19 nm). The increase in hard domain size, specific to the Z20−20 network hybrid, suggests that PZLY contributes to the hard block associations either by promoting hard segment self-association or by forming a pseudo-hard segment with the HDI trimer.16,18 Thus, the SAXS spectra confirm the phase behavior observed in microscopic studies and further support that the higher PZLY content films enhance the interactions between the PU blocks, facilitating phase mixing. Whereas SAXS probes the overall long-range ordering, XRD experiments were performed to explore the crystal structure of the PEG soft block and potential PZLY secondary structures. Understanding PEG crystallinity is critical to the examination of spherulite formation. As seen in Figure 5, prominent PEG crystallization peaks occur at 22.2 and 27.0° (0.46 and 0.38 nm).51,52 Additionally, a minor peak at 17.5° (0.59 nm) for the control PEG−HDI indicates that PEG forms a 72 helix structure.53−55 With the addition of PZLY (Z5−10), the PEG crystalline peak intensities were reduced, and the peak at 17.5° disappeared; however, two additional small peaks at 33.0 (3.15 Å) and 47.4° (2.22 Å) emerged. Similarly, in Z5−20, the prominent PEG crystallization peaks almost disappeared, and new peaks at 33 and 47° are now the most prominent. The G


Article

Biomacromolecules

permanent network may slow down the kinetics of spherulite growth.3 The morphology of polymeric materials plays a major role in film mechanics and failure mechanisms. For instance, semicrystalline polymers with a spherulitic structure can be brittle and easily crack under load due to break up at the spherulite boundary.63 Alternatively, the percolated fiber morphology of PDMS−peptide PUs facilitates stress transfer via fiber−fiber interactions.16 Thus it is anticipated that variations in morphological development between the network PUs will result in tunable mechanics. Figure 8 shows the stress−strain curves of network PUs and summarizes the average Young’s modulus and elongation-at-break values. The Young’s moduli correlated well with the degree of crystallinity of PEG. PEG− HDI films (50% crystallinity) exhibited a modulus of ∼200 MPa. A reduction in moduli to ∼100 MPa was observed in Z5−10 that exhibited a PEG crystallinity of only 32%, with a further reduction to ∼50 MPa for Z5−20 with 24% crystallinity. The same trend was observed in Z20−10 and Z20−20 hybrids. This observation is consistent with previously explored PBLA−PEG−PBLA film mechanics, suggesting that the peptide block had little effect on the moduli but instead PEG crystallinity enhanced film strength.12 The elongation-at-break of network PU films showed little correlation to PEG crystallinity but instead was related to the presence and size of spherulites. Higher extensibility was observed for network films with smaller PEG spherulites. Elongations of 350−400% were exhibited by both PEG−HDI and Z5−10 with a spherulite diameter of ∼115 μm. With an increase in diameter to ∼200 μm for Z20−10, the extensibility decreases to ∼150%. This trend has also been observed in PEO blends and polypropylene films, resulting from the interspherulitic strain being proportional to spherulite size.63,64 The strain-at-break for the network hybrids with 20 wt % PZLY content (Z5−20 and Z20−20) did not follow this trend but instead exhibited brittle failure. The brittleness observed at high peptide content may be due to several factors, such as an increase in rigidity with added PZLY content and weaker interactions between the ill-formed spherulites causing failure at lower strains. These weak interspherulite associations were probed via POM of network PU films post-uniaxial tension. In Figure S7, well-defined spherulites deformed and aligned into a fiber-like morphology in PEG−HDI, Z5−10, and Z20−10. For Z5−20 and Z20−20, uniaxial alignment did not result from fiber drawing; instead, voids developed early in the deformation process at the spherulite boundary. Thus, it can be concluded that the incorporation of PZLY in network PUs allows for the control of spherulite size and interspherulite associations, which, in turn, can be utilized as a mechanism to tailor mechanical properties.

Figure 6. DSC spectra of the second heating cycle for network PUs. The dashed line indicates the PEG Tm peak. This peak was used to calculate the percent crystallinity. Crystallinity (%) = (ΔHm/ ΔH0)*100. Exotherm is up.

PEG to crystallize, whereby higher β-sheet content in the Z5 series hinders PEG crystallinity. This observation was most noticeable in Z5−20, which displayed a 50% reduction in PEG crystallinity. Additionally, the ratio of PEG and peptide has a significant impact on phase separation, where increasing peptide volume fraction enhances phase mixing.11 Thus, a reduction in PEG crystallinity with increasing peptide content was anticipated due to the enhanced interaction between PEG and PZLY. Multiple PEG melting transitions are observed in the first heating cycle (Figure S5), indicative of broad crystal size distributions imparted by the thermal history of the samples.39,60 With this difference in first and second heating cycles, in situ monitoring of spherulite morphology at the PEG melting transition was obtained using POM to analyze crystal growth after a heating and cooling cycle.61 Films were examined at intervals of 5 °C within a temperature range of 25 to 65 °C. Figure S6 provides a visual representation of the disruption of the PEG spherulites upon melting, resulting in complete disappearance of birefringence. The temperatures at which this disappearance occurs match the PEG melting transition temperatures observed in the first heating cycle of DSC. Upon heating the films to 150 °C and rapid cooling to room temperature, spherulite structures are either minimized in size or completely destroyed (Figure 7). For instance, Z5− 10 still exhibits faint spherulites; however, Z5−20 only displays birefringence. This result suggests that the thermal history strongly influences the morphology of solid-state network PU films.44,46−48 Future work on the kinetics of spherulite formation will aid in the understanding of PZLY’s role in crystal growth; however, this is not within the scope of the current project.62 It is noted that the kinetics of spherulite growth are influenced by molecular mobility, suggesting that a

Figure 7. POM images of films after cooling to room temperature from 150 °C. Samples were held at this elevated temperature for 5 min before rapid cooling. Images are at 10× magnification. H


Article

Biomacromolecules

Figure 8. Stress−strain curves and associated mechanical data of network PU films. Additional insert is a magnified view, allowing for better visibility of the yield point. The table includes the percent crystallinity of PEG for the second heating cycle of DSC. Additionally, spherulite diameter is reported for comparison with mechanical data, calculated from SEM analysis.

4. CONCLUSIONS Here, we explored the fabrication and material properties of covalently cross-linked PEG−polyurea films with the addition of peptide blocks in the polymer backbone. Peptide content and repeat length of PZLY were varied within PU hybrids to explore the effects of peptide self-assembly and increased hydrogen-bonding sites on the resulting film morphology and mechanics. The resulting peptidic films displayed a range of molecular interactions, giving rise to the hierarchical organization of macroscopic films, composed of microscopic spherulite morphology, nanoscale fibers, and peptide secondary structures. On the molecular level, PZLY organized primarily into an αhelical formation, resulting in intramolecular interactions within the PZLY block and intermolecular interactions between PEG and PZLY. XRD analysis further supported the interactions of PZLY and PEG, highlighting a decrease in PEG crystallinity with increasing PZLY content and suggesting a modified PEG crystal structure with the increase in β-sheet PZLY structures. Microscopy studies of the network films using POM and SEM revealed the presence of spherulitic morphologies, ranging in size from 115 to 215 μm based on both the PZLY fraction and block length. Further exploration of the nanoscale structure revealed a fiber-like morphology resulting from PEG organization and hard segment aggregates of ∼16 nm in all hybrids except Z5−20. Z5−20 exhibited the highest level of phase mixing, as characterized by the lack of fiber formation and a reduction in PEG crystallinity due to the enhanced PZLY intermolecular hydrogen bonding. The morphology and spherulite size dramatically influenced the mechanical properties of uniaxially stretched network PUs. The elongation-at-break diminished with larger spherulite formation, and the modulus decreased with the reduction in PEG crystallinity upon incorporation of PZLY. Variations in peptide content and block length were utilized to achieve a

range in moduli (50 to 200 MPa) and elongations of 20− 400%. Ultimately, the addition of PZLY tuned the interfacial spherulite interactions, resulting in ductile to brittle mechanics upon increasing peptide fraction. Future exploration of spherulite kinetics is anticipated to enhance understanding of morphological development that can be utilized to achieve mechanics for a range of applications. This work has successfully manufactured network PZLY−PU hybrids with hierarchical organization, advancing our understanding of controlled architectures. Furthermore, these results will expand the use of solid-state polymer−peptide hybrids from medical applications to responsive shape-memory materials.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.biomac.8b00762.

■

Synthetic reaction quantities as well as 1H NMR, ATRFTIR, and morphological POM studies. Specifically, in situ POM of films as temperature was increased shows the disappearance of spherulites at defined temperatures. Additionally, post-strain POM analysis of PU hybrids. (PDF)

AUTHOR INFORMATION

Corresponding Author

* E-mail: [email protected]. ORCID

LaShanda T. J. Korley: 0000-0002-8266-5000 Present Address #

B.L.: Chemical Sciences Division, Oak Ridge National Laboratory, 1 Bethel Valley Road, Oak Ridge, TN 37830, USA.

I


Article

Biomacromolecules Author Contributions

(10) Hamley, I. W. PEG−Peptide Conjugates. Biomacromolecules 2014, 15, 1543−1559. (11) Floudas, G.; Papadopoulos, P.; Klok, H.-A.; Vandermeulen, G. W. M.; Rodriguez-Hernandez, J. Hierarchical Self-Assembly of Poly(γBenzyl-L-glutamate)−Poly(ethylene glycol)−Poly(γ-Benzyl-L-Glutamate) Rod−Coil−Rod Triblock Copolymers. Macromolecules 2003, 36, 3673−3683. (12) Tanaka, S.; Ogura, A.; Kaneko, T.; Murata, Y.; Akashi, M. Precise Synthesis of ABA Triblock Copolymers Comprised of Poly(ethylene Oxide) and Poly (β-Benzyl-L-Aspartate): A Hierarchical Structure Inducing Excellent Elasticity. Macromolecules 2004, 37, 1370−1377. (13) Hamley, I. W.; Krysmann, M. J. Effect of PEG Crystallization on the Self-Assembly of PEG/Peptide Copolymers Containing Amyloid Peptide Fragments. Langmuir 2008, 24, 8210−8214. (14) Matolyak, L. E.; Keum, J. K.; Van de Voorde, K. M.; Korley, L. T. J. Synthetic Approach to Tailored Physical Associations in PeptidePolyurea/polyurethane Hybrids. Org. Biomol. Chem. 2017, 15, 7607− 7617. (15) Cheng, J.; Deming, T. J. Synthesis of Polypeptides by RingOpening Polymerization of α-Amino Acid N-Carboxyanhydrides. In Peptide-Based Materials; Springer: Berlin, 2011; pp 1−26. (16) Johnson, J. C.; Wanasekara, N. D.; Korley, L. T. J. Influence of Secondary Structure and Hydrogen-Bonding Arrangement on the Mechanical Properties of Peptidic-Polyurea Hybrids. J. Mater. Chem. B 2014, 2, 2554. (17) Cho, C.-S.; Jo, B.-W.; Kwon, J.-K.; Komoto, T. Synthesis and Structural Study of an ABA Block Copolymer Consisting of Poly(εBenzyloxycarbonyl-L-Lysine) as the A Block and Poly(ethylene Oxide) as the B Block. Macromol. Chem. Phys. 1994, 195, 2195−2206. (18) Matolyak, L.; Keum, J.; Korley, L. T. J. Molecular Design: Network Architecture and Its Impact on the Organization and Mechanics of Peptide-Polyurea Hybrids. Biomacromolecules 2016, 17, 3931−3939. (19) Lecommandoux, S.; Achard, M.-F.; Langenwalter, J. F.; Klok, H.-A. Self-Assembly of Rod−Coil Diblock Oligomers Based on αHelical Peptides. Macromolecules 2001, 34, 9100−9111. (20) Schlaad, H.; Kukula, H.; Smarsly, B.; Antonietti, M.; Pakula, T. Solid-State Morphologies of Linear and Bottlebrush-Shaped polystyrene−poly(Z-L-Lysine) Block Copolymers. Polymer 2002, 43 (19), 5321−5328. (21) Schlaad, H.; Smarsly, B.; Losik, M. The Role of Chain-Length Distribution in the Formation of Solid-State Structures of Polypeptide-Based Rod−Coil Block Copolymers. Macromolecules 2004, 37, 2210−2214. (22) Naik, S. S.; Savin, D. A. Poly(Z-Lysine)-Based Organogels: Effect of Interfacial Frustration on Gel Strength. Macromolecules 2009, 42, 7114−7121. (23) Besteiro, M. C.; Guiomar, A. J.; Gonçalves, C. A.; Bairos, V. A.; de Pinho, M. N.; Gil, M. H. Characterization and in Vitro Hemocompatibility of Bi-Soft Segment, Polycaprolactone-Based Poly(ester Urethane Urea) Membranes. J. Biomed. Mater. Res., Part A 2009, 93A (3), 954−964. (24) Luo, S. Q.; Huang, C. Y. F.; McClelland, J. F.; Graves, D. J. A Study of Protein Secondary Structure by Fourier Transform Infrared/ Photoacoustic Spectroscopy and Its Application for Recombinant Proteins. Anal. Biochem. 1994, 216, 67−76. (25) Rathore, O.; Sogah, D. Y. Self-Assembly of β-Sheets into Nanostructures by Poly (Alanine) Segments Incorporated in Multiblock Copolymers Inspired by Spider Silk. J. Am. Chem. Soc. 2001, 123, 5231−5239. (26) Raimo, M. Estimation of Polymer Nucleation and Growth Rates by Overall DSC Crystallization Rates. Polym. J. 2011, 43, 78− 83. (27) Crist, B.; Schultz, J. M. Polymer Spherulites: A Critical Review. Prog. Polym. Sci. 2016, 56, 1−63. (28) Prud’homme, R. E. Crystallization and Morphology of Ultrathin Films of Homopolymers and Polymer Blends. Prog. Polym. Sci. 2016, 54−55, 214−231.

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. L.E.M. and L.T.J.K. conducted the majority of the experiments and wrote the paper with the corresponding author. Additional support of experiments, such as DSC, and editing was performed by C.B.T. B.L., J.K.K., J.E.C., and R.S.T. contributed by providing supplemental support for synthesis, SAXS, XRD, or AFM studies. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS Dr. J. Casey Johnson is acknowledged for his assistance with schematic design. We acknowledge funding support from the National Science Foundation (CAREER DMR-0953236, and DMR-1608441).

■

ABBREVIATIONS PZLY or Z, poly(ε-carbobenzyloxy-L-lysine); PEG, poly(ethylene glycol); PBLA, poly(β-benzyl-L-aspartate); PBLG, poly(γ-benzyl-L-glutamate); DMF, N-dimethylformamide; THF, tetrahydrofuran; DMAc, dimethylacetamide; N2, nitrogen; RT, room temperature; h, hour; HDI, hexamethylene diisocyanate; NCA, N-carboxyanhydride; NMR, nuclear magnetic resonance; PTFE, polytetrafluoroethylene; ATRFTIR, attenuated total reflectance Fourier transform infrared; POM, polarized optical microscopy; SEM, scanning electron microscopy; AFM, atomic force microscopy; SAXS, smallangle X-ray scattering; XRD, X-ray diffraction; DSC, differential scanning calorimetry; Et, elastic modulus; εb, elongationat-break; ASTM, American standards for testing and materials.

■

REFERENCES

(1) Johnson, J. C.; Korley, L. T. J. Enhanced Mechanical Pathways through Nature’s Building Blocks: Amino Acids. Soft Matter 2012, 8, 11431−11442. (2) Guzmán, D. L.; Randall, A.; Baldi, P.; Guan, Z. Computational and Single-Molecule Force Studies of a Macro Domain Protein Reveal a Key Molecular Determinant for Mechanical Stability. Proc. Natl. Acad. Sci. U. S. A. 2010, 107 (5), 1989−1994. (3) Grieshaber, S. E.; Farran, A. J. E.; Bai, S.; Kiick, K. L.; Jia, X. Tuning the Properties of Elastin Mimetic Hybrid Copolymers via a Modular Polymerization Method. Biomacromolecules 2012, 13, 1774− 1786. (4) Shu, J. Y.; Panganiban, B.; Xu, T. Peptide-Polymer Conjugates: From Fundamental Science to Application. Annu. Rev. Phys. Chem. 2013, 64, 631−657. (5) Top, A.; Roberts, C. J.; Kiick, K. L. Conformational and Aggregation Properties of a PEGylated Alanine-Rich Polypeptide. Biomacromolecules 2011, 12, 2184−2192. (6) Rathore, O.; Winningham, M. J.; Sogah, D. Y. A Novel SilkBased Segmented Block Copolymer Containing GlyAlaGlyAla βSheets Templated by Phenoxathiin. J. Polym. Sci., Part A: Polym. Chem. 2000, 38, 352−366. (7) Rathore, O.; Sogah, D. Y. Nanostructure Formation through βSheet Self-Assembly in Silk-Based Materials. Macromolecules 2001, 34, 1477−1486. (8) Desai, M. S.; Wang, E.; Joyner, K.; Chung, T. W.; Jin, H.; Lee, S. Elastin-Based Rubber-Like Hydrogels. Biomacromolecules 2016, 17, 2409−2416. (9) Lee, H. J.; Lee, J. S.; Chansakul, T.; Yu, C.; Elisseeff, J. H.; Yu, S. M. Collagen Mimetic Peptide-Conjugated Photopolymerizable PEG Hydrogel. Biomaterials 2006, 27 (30), 5268−5276. J


Article

Biomacromolecules (29) Brandrup, J.; Immergut, E. H.; Grulke, E. A. Polymer Handbook; Wiley-Interscience, 1999. (30) Matolyak, L. E.; Keum, J. K.; Van de Voorde, K. M.; Korley, L. T. J. Synthetic Approach to Tailored Physical Associations in PeptidePolyurea/polyurethane Hybrids. Org. Biomol. Chem. 2017, 15, 7607− 7617. (31) Dou, Q.; Wang, C.; Cheng, C.; Han, W.; Thüne, P. C.; Ming, W. PDMS-Modified Polyurethane Films with Low Water Contact Angle Hysteresis. Macromol. Chem. Phys. 2006, 207, 2170−2179. (32) Javni, I.; Zhang, W.; Petrović, Z. S. Effect of Different Isocyanates on the Properties of Soy-Based Polyurethanes. J. Appl. Polym. Sci. 2003, 88, 2912−2916. (33) Andrady, A. L. Science and Technology of Polymer Nanofibers; Wiley, 2008. (34) Fetters, L. J.; Lohse, D. J.; Richter, D.; Witten, T. A.; Zirkel, A. Connection between Polymer Molecular Weight, Density, Chain Dimensions, and Melt Viscoelastic Properties. Macromolecules 1994, 27, 4639−4647. (35) Chirgadze, Y. N.; Nevskaya, N. A. Infrared Spectra and Resonance Interaction of Amide-I Vibration of the Parallel-Chain Pleated Sheet. Biopolymers 1976, 15, 627−636. (36) Ishimura, Y.; Hamada, F.; Nakajima, A. Conformation Study of Poly(Nε-Carbobenzoxy-L-Lysine) in Helicogenic Solvents by SmallAngle X-Ray Scattering. Macromolecules 1978, 11, 382−387. (37) Min, X.; Fang, M.; Huang, Z.; Liu, Y.; Huang, Y.; Wen, R.; Qian, T.; Wu, X. Enhanced Thermal Properties of Novel ShapeStabilized PEG Composite Phase Change Materials with Radial Mesoporous Silica Sphere for Thermal Energy Storage. Sci. Rep. 2015, 5, 12964. (38) Takahashi, Y.; Tadokoro, H. Structural Studies of Polyethers, (-(CH 2)M-O-) N. X. Crystal Structure of Poly(ethylene Oxide). Macromolecules 1973, 6, 672−675. (39) Meng, Q.; Hu, J. A Poly(ethylene Glycol)-Based Smart Phase Change Material. Sol. Energy Mater. Sol. Cells 2008, 92, 1260−1268. (40) Keith, H. D.; Padden, F. J.; Russell, T. P. Morphological Changes in Polyesters and Polyamides Induced by Blending with Small Concentrations of Polymer Diluents. Macromolecules 1989, 22, 666−675. (41) Padden, F. J.; Keith, H. D. Spherulitic Crystallization in Polypropylene. J. Appl. Phys. 1959, 30, 1479−1484. (42) Hu, J.-L.; Mondal, S. Structural Characterization and Mass Transfer Properties of Segmented Polyurethane: Influence of Block Length of Hydrophilic Segments. Polym. Int. 2005, 54, 764−771. (43) Carlton, R. A. Thermal Microscopy. In Pharmaceutical Microscopy; Springer New York: New York, 2011; pp 65−84. (44) Gbabode, G.; Delvaux, M.; Schweicher, G.; Andreasen, J. W.; Nielsen, M. M.; Geerts, Y. H. Unique Crystal Orientation of Poly(ethylene Oxide) Thin Films by Crystallization Using a Thermal Gradient. Macromolecules 2017, 50, 5877−5891. (45) Li, Y.; Wang, Z.; He, T. Morphological Control of Polymer Spherulites via Manipulating Radial Lamellar Organization upon Evaporative Crystallization: A Mini Review. Crystals 2017, 7, 115. (46) Kaneko, T.; Shimokuri, T.; Tanaka, S.; Akashi, M. Micromorpholoy Memory in Amphiphilic Polypeptides. Polym. J. 2007, 39, 737−744. (47) Woo, E. M.; Lugito, G. Origins of Periodic Bands in Polymer Spherulites. Eur. Polym. J. 2015, 71, 27−60. (48) Sheth, J. P.; Wilkes, G. L.; Fornof, A. R.; Long, T. E.; Yilgor, I. Probing the Hard Segment Phase Connectivity and Percolation in Model Segmented Poly(urethane Urea) Copolymers. Macromolecules 2005, 38, 5681−5685. (49) Wang, L.; Hickner, M. A. Low-Temperature Crosslinking of Anion Exchange Membranes. Polym. Chem. 2014, 5, 2928. (50) Chang, S. L.; Yu, T. L.; Huang, C. C.; Chen, W. C.; Linliu, K.; Lin, T. L. Effect of Polyester Side-Chains on the Phase Segregation of Polyurethanes Using Small-Angle X-Ray Scattering. Polymer 1998, 39, 3479−3489. (51) Lee, S.; Tong, X.; Yang, F. The Effects of Varying Poly(ethylene Glycol) Hydrogel Crosslinking Density and the Crosslinking

Mechanism on Protein Accumulation in Three-Dimensional Hydrogels. Acta Biomater. 2014, 10, 4167−4174. (52) Vidotto, G.; Lévy, D.; Kovacs, A. J. Cristallisation et Fusion Des Polymères Autoensemencés. Colloid Polym. Sci. 1969, 230, 289−305. (53) Lomax, G. R.; Manorama, S. V.; Singh, R. K.; Parkash, O.; Russell, T. P.; Halley, J. W.; Nielsen, B.; Weitz, D. A. Breathable Polyurethane Membranes for Textile and Related Industries. J. Mater. Chem. 2007, 17, 2775. (54) Tadokoro, H.; Chatani, Y.; Yoshihara, T.; Tahara, S.; Murahashi, S. Structural Studies on Polyethers, [-(CH2)m-O-]n. II. Molecular Structure of Polyethylene Oxide. Makromol. Chem. 1964, 73, 109−127. (55) Braun, H.-G.; Meyer, E. Structure Formation of Ultrathin PEO Films at Solid InterfacesComplex Pattern Formation by Dewetting and Crystallization. Int. J. Mol. Sci. 2013, 14, 3254−3264. (56) Brena, B.; Zhuang, G. V.; Augustsson, A.; Liu, G.; Nordgren, J.; Guo, J.-H.; Ross, P. N.; Luo, Y. Conformation Dependence of Electronic Structures of Poly(ethylene Oxide). J. Phys. Chem. B 2005, 109, 7907−7914. (57) Drummy, L. F.; Phillips, D. M.; Stone, M. O.; Farmer, B. L.; Naik, R. R. Thermally Induced α-Helix to β-Sheet Transition in Regenerated Silk Fibers and Films. Biomacromolecules 2005, 6, 3328− 3333. (58) Ju, H.; McCloskey, B. D.; Sagle, A. C.; Kusuma, V. A.; Freeman, B. D. Preparation and Characterization of Crosslinked Poly(ethylene Glycol) Diacrylate Hydrogels as Fouling-Resistant Membrane Coating Materials. J. Membr. Sci. 2009, 330, 180−188. (59) Graham, N. B.; McNeill, M. E. Hydrogels for Controlled Drug Delivery. Biomaterials 1984, 5, 27−36. (60) Lee, D.; Lee, S.-H.; Kim, S.; Char, K.; Park, J. H.; Bae, Y. H. Micro-Phase-Separation Behavior of Amphiphilic Polyurethanes Involving Poly(ethylene Oxide) and Poly(tetramethylene Oxide). J. Polym. Sci., Part B: Polym. Phys. 2003, 41, 2365−2374. (61) Fu, G.; Kyu, T. Role of Molecular Complexation in SolidLiquid Phase Diagram of Poly-Ethylene Oxide/urea Mixture. Polymer 2017, 116, 350−356. (62) Singfield, K. L.; Chisholm, R. A.; King, T. L. A Physical Chemistry Experiment in Polymer Crystallization Kinetics. J. Chem. Educ. 2012, 89, 159−162. (63) Pavlov, V. I. Investigation of the Effect of Spherulite Size on the Strength and Deformation Characteristics of Isotactic Polypropylene Films. Sov. Mater. Sci. 1971, 4, 438−440. (64) Li, J.-Z.; Schultz, J. M.; Chan, C.-M. The Relationship between Morphology and Impact Toughness of Poly(l-Lactic Acid)/poly(ethylene Oxide) Blends. Polymer 2015, 63, 179−188.

K


Secondary-Structure-Mediated Hierarchy and ... - ACS Publications

Recommend Documents