Network Architecture and Its Impact on the ... - ACS Publications

Article pubs.acs.org/Biomac

Molecular Design: Network Architecture and Its Impact on the Organization and Mechanics of Peptide-Polyurea Hybrids Lindsay Matolyak,† Jong Keum,‡ and LaShanda T. J. Korley*,† †

Department of Macromolecular Science and Engineering, Case Western Reserve University, 2100 Adelbert Road, Cleveland, Ohio 44106-7202, United States ‡ Center for Nanophase Materials Sciences and Chemical and Engineering Materials Division, Oak Ridge National Laboratory, 1 Bethel Valley Rd, Oak Ridge, Tennessee 37831, United States S Supporting Information *

ABSTRACT: Nature has achieved controlled and tunable mechanics via hierarchical organization driven by physical and covalent interactions. Polymer−peptide hybrids have been designed to mimic natural materials utilizing these architectural strategies, obtaining diverse mechanical properties, stimuli responsiveness, and bioactivity. Here, utilizing a molecular design pathway, peptide−polyurea hybrid networks were synthesized to investigate the role of architecture and structural interplay on peptide hydrogen bonding, assembly, and mechanics. Networks formed from poly(β-benzyl-L-aspartate)−poly(dimethylsiloxane) copolymers covalently cross-linked with a triisocyanate yielded polyurea films with a globular-like morphology and parallel β-sheet secondary structures. The geometrical constraints imposed by the network led to an increase in peptide loading and ∼7x increase in Young’s modulus while maintaining extensibility (∼160%). Thus, the interplay of physical and chemical bonds allowed for the modulation of resulting mechanical properties. This investigation provides a framework for the utilization of structural interplay and mechanical tuning in polymer−peptide hybrids, which offers a pathway for the design of future hybrid biomaterial systems. hard segment.5 Likewise, incorporating poly(hexamethylene oxide) (PHMO) into the soft segment (SS) of a PDMS-based polyurethane increased the amount of phase mixing in the system, resulting in a decrease in storage modulus at room temperature.6 Further incorporation of covalent cross-links, most commonly through the integration of either polyfunctional polyol chain-extenders or polyfunctional isocyanates, has been shown to restrict mobility, resulting in either an increase in strength of amorphous polymers or disruption of the crystallinity in more ordered systems.4 For instance, Sheth et al. observed polyurethane/ureas with and without hard segment branching to investigate the importance of hydrogen bonding and chain architecture.7 Initially, the linear chains formed via hydrogen bonding yielded a fibrillar morphology. Upon increasing the amount of covalent cross-links via a mixture of di- and trifunctional isocyanates, the morphology shifted to a microphase-separated globular texture, which was attributed to the disruption of a well-structured hydrogen-bonded network. Thus, a bottom-up approach was achieved, whereby the incorporation and modulation of intermolecular interactions influences crystallinity, phase separation, and material function. With these Nature-inspired design principles in mind, a bottom-up strategy was employed to investigate the impact of network architecture on the organization and mechanics of polymer−peptide hybrids. Here, we present an investigation of

1. INTRODUCTION Natural materials, such as spider silk and elastin, achieve a range of mechanics that is specific to their material class by altering molecular design and allowing for hierarchical long-range ordering of peptides.1 For example, the Araneus spider produces silks with Young’s moduli ranging from 10 GPa for MA silk, to 0.003 GPa for viscid silk, by varying the size of the crystalline domains formed by physical cross-links in the form of β-sheet structures.2 By contrast, elastin achieves high elongations through covalently cross-linking of short peptide chains.3 Thus, Nature is skilled at harnessing the interplay of physical and covalent interactions to achieve tunable properties. As with natural materials, scientists and engineers employ physical and covalent bonds to design polymers with application-specific properties. Polyurethanes and polyureas are some of the most versatile examples, since incorporating physical associations, e.g., hydrogen bonding, and/or covalent cross-links is easy to accomplish synthetically. In terms of intermolecular forces, weak hydrogen bonding is present through urethane and urea linkages that self-organize and promote phase separation within the material by connecting a more rigid hard block to a more flexible soft block.4 By modifying the chemical structure of either the hard or soft segment, the domain assembly may be altered, allowing for changes in crystallinity, phase separation, and resulting mechanics. For instance, poly(dimethylsiloxane) (PDMS)based polyurethane/ureas with a crystalline hard segment (HS) exhibited larger domain sizes and enhanced storage modulus compared to polyurethane/ureas with an amorphous © XXXX American Chemical Society

Received: August 31, 2016 Revised: October 30, 2016

A

DOI: 10.1021/acs.biomac.6b01309 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules

Figure 1. Reaction scheme for covalently cross-linked polyurea-peptide network hybrids, where n = ∼5 and x = ∼24.

⎛ ⎞ xM pep ⎟⎟ wt%(peptide) = 100⎜⎜ ⎝ xM pep + yM PDMS + zMHDI ⎠

network polyurea−aspartate hybrids that incorporate physical interactions via self-organizing peptide units into a versatile polyurea platform. The peptide−polymer motif utilized in this work complements our previously reported linear polyureas containing poly(β-benzyl-L-aspartate)-b-PDMS-b-poly(β-benzyl-L-aspartate) (PBLA-b-PDMS-b-PBLA) soft segments.8 This work provides insight for additional molecular design strategies of synthetic peptide hybrid polymers with tunable material properties that will aid in further development of passive and stimuli responsive soft materials, such as implants, therapeutics, and tissue engineering.9,10

(1)

where x, y, and z are the molar quantities of the peptide, PDMS, and HDI, respectively, and Mpep, MPDMS, and MHDI are the molecular weights of the peptide, PDMS, and HDI, respectively. Synthetic details of the non-chain extended linear polyurea systems have been previously reported.8 Briefly, PBLA5-40 was synthesized in a N2 atmosphere glovebox by adding 1 g (0.21 mol) of PBLA5-b-PDMSb-PBLA5 and 50 mL of THF to a 150 mL round-bottom flask equipped with a Vigreux condenser. Next, PDMS (0.08 g, 0.03 mol) and HDI (0.04 g. 0.25 mol) were dissolved separately in 25 mL of THF each and added to the flask along with 5 drops of DBTDL catalyst. The solution was heated to 60 °C for 16 h. The final product was precipitated into distilled water, filtered, and washed with methanol. The resulting white solid was dried under vacuum until constant weight was achieved (∼24 h). To fabricate the cross-linked, polyurea networks, a commercially available, polyfunctional HDI (Desmodur N 3300 A) was utilized. Due to the polyfunctionality (f ∼ 3) 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 highly crosslinked networks (∼85% gel−sol). A flame-dried, 100 mL roundbottom flask with magnetic stirrer was placed inside a glovebox under N2 atmosphere. As an example, 0.25 g (0.05 mmol) of PBLA5-bPDMS-b-PBLA5 triblock and 0.71 g (0.28 mmol) of PDMS were dissolved in a flask containing 25 mL of THF. Upon complete dissolution, 0.13 g (0.25 mol) of predissolved HDI trimer in 15 mL of THF was added to the flask and allowed to stir at room temperature for 1 min prior to pouring the mixture into a round polytetrafluoroethylene (PTFE) mold (∼9 cm diameter). The mold was placed inside a desiccator under N2 for 1 h before heating in an oven at 60 °C under N2 purge. The reaction was held under these conditions for 18 h. After 18 h, the samples were equilibrated to room temperature under vacuum and maintained at RT for 24 h to remove any residual THF. The resulting films were clear, colorless, and flexible. Full synthetic quantities with corresponding amounts of each component (PDMS, HDI, and triblock) are given in Table S1 in the Supporting Information. Figure 1 details the reaction scheme of network polymer−peptide hybrids. 2.3. Characterization. 2.3.1. Molecular Weight. The molecular weight of the linear peptidic polyureas were characterized relative to narrow polystyrene standards utilizing gel permeation chromatography (GPC, Agilent 1200 series liquid chromatography system) equipped

2. EXPERIMENTAL SECTION 2.1. Materials. All materials were purchased from Sigma-Aldrich unless otherwise noted. Tetrahydrofuran (THF) was purified utilizing a Vacuum Atmosphere’s solvent purification system. α,ω-Bis(3aminopropyl)poly(dimethylsiloxane) (PDMS, 2500 g/mol) was dried at 100 °C under vacuum for a minimum of 16 h. 1,6hexamethylene diisocyanate (HDI) was vacuum distilled before use. Triphosgene, dibutyltin dilaurate (DBTDL), deuterated chlorofom, methanol, and β-benzyl-L-aspartate (BLA) were used as received. Desmodur N 3300 A (HDI trimer) with 21.8% isocyanate content was used as received from Bayer MaterialScience. β-benzyl-L-aspartate Ncarboxyanhydride (NCA) was prepared following literature procedure.8,11−13 2.2. Synthesis. 2.2.1. Poly(β-benzyl-L-aspartate)-block-Poly(dimethylsiloxane)-block-poly(β-benzyl-L -aspartate) (PBLA 5 -bPDMS-b-PBLA5) Triblocks. Triblocks were synthesized via amineinitiated NCA polymerization.8,11 As an example, the synthesis of PBLA5-b-PDMS-b-PBLA5 was performed in a flame-dried 100 mL round-bottom flask with a magnetic stirrer. In a nitrogen (N2) atmosphere glovebox, 2.5 g (0.01 mol) of BLA-NCA was dissolved with 25 mL of THF. To the solution, 2.5 g (0.001 mol) of PDMS predissolved in 25 mL of THF was added. The reaction was stirred at room temperature (RT) for 20 h and precipitated into distilled water; the precipitate was filtered, and washed with methanol. The final white solid was dried under vacuum until constant weight was achieved (∼24 h). 2.2.2. Peptidic Polyureas. The weight fraction of peptide was controlled in both the network and linear systems with the addition of excess PDMS. Calculated weight fractions were obtained using eq 1.8 B


Article

Biomacromolecules

2.3.7. Tensile Testing. A Zwick/Roell mechanical testing instrument equipped with a 500 N load cell was used to determine mechanical response. Films for tensile testing were cut using a dog bone 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 and 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. Hysteresis measurements were taken, cycling 20 intervals to a strain of 40%. This strain was the highest % that allowed for multiple cycles to be tested for both linear and network films. Hysteresis (%) is defined as the ratio of the area within the loading−unloading curve to the total area under the loading curve. The reported mechanical properties were an average of a minimum of three samples.

with refractive index and variable wavelength detectors and American Polymer Standard linear bed GPC columns with THF as the eluent at 25 °C. Absolute molecular weights of the triblocks were characterized via 1H nuclear magnetic resonance (NMR) end group analysis (Varian Inova 600 MHz, CDCl3) (Figure S1). The average molecular weight between cross-links (Mc) of the covalently cross-linked system was defined using eq 2.14

Mc =

Σ(m1 + m2 + m3 + ... + mn)

(1 )

fn − 2 m1 M1

+

(fn2 − 2)m2 M2

+

(f

n3

)

− 2 m3 M3

+ ... +

(f

nn

)

− 2 mn Mn

(2) where Σ(m1 + m2 + m3 + ... + mn) is the total mass of all the components present in the polymer, f n1, f n2, f n3... f nn is the functionality of each component, and M1, M2, M3 ... Mn is the molar mass of each component. For our system, both the PDMS and the peptide-based triblocks have a functionality of 2, simplifying to eq 3. Mc =

3. RESULTS AND DISCUSSION Linear polyurea peptide hybrid films were previously investigated for mechanical tunability through the interplay of peptide organization/hydrogen bonding arrangement.8 The use of PBLA as the peptide segment allowed for the control of secondary structure (α-helix, β-sheet, and random coil) as a function of repeat length as well as minimized any peptide sidechain hydrogen bonding interactions. The choice of a shorter peptide repeat length (5) yielded β-sheet secondary structures. The degree of hydrogen bonding increased with increasing peptide content, which allowed for more long-range ordering within the samples.8 We hypothesize that, by utilizing a covalent network architecture as the polyurea platform, additional hierarchy can be imposed via the overlay of covalent (e.g., chemical) cross-linking within the hard block and physical associations within the peptidic soft block. During network film fabrication, the peptide block length was limited to five repeats due to inadequate solubility with increasing molecular weight of the peptide block, where fully dissolved materials were necessary to yield homogeneous covalently cross-linked films. A hexamethylene polyisocyanate (HDI trimer; f ∼ 3), which has a triazine trione molecular structure, was chosen as the hard segment for comparison with the 1,6-hexamethylene diisocyanate (HDI; f = 2) utilized for the linear analogues.8,18,19 By introducing a covalent network via polyfunctionality during synthesis, as opposed to linking linear polymers together, the linear and network samples may be directly compared. Thereby, we can contribute any change in thermal and mechanical properties to the formation of a covalent network. 3.1. Synthetic Strategy. Polyurea-peptide hybrids were synthesized using PBLA5-b-PDMS-b-PBLA5 as the soft phase with either linear HDI or HDI trimer as the hard segment. Peptide content (10, 20, and 40 wt %) was varied within the network hybrids to observe the impact of peptide hydrogen bonding on nanostructure and mechanics. At 40 wt % PBLA, all the isocyanate groups were connected to PBLA units; however, assuming equal reactivity, only half of the isocyanate units in the 20 wt % hybrid were expected to be chemically reacted to the PBLA block. In contrast, the majority of the isocyanate segments were coupled to the PDMS blocks at 10 wt %. We chose a non-chain extended polyurea pathway to mitigate the influence of hard segment chemistry. PBLA5-b-PDMS-b-PBLA5 prepolymer was synthesized via ring opening polymerization of β-benzyl-L-aspartate N-carboxyanhydride (BLA-NCA) with a peptide block length of 5, promoting β-sheet conformation due to the short peptide repeat length. The PBLA-based prepolymer was incorporated into the polyurea hybrids (linear

∑ (mPDMS + mpep + mHDI) (fHDI − 2)mHDI MHDI

(3)

2.3.2. Degree of Covalent Cross-Linking. The gel−sol fraction of each covalently cross-linked film was determined via THF Soxhlet extraction for 24 h. A PTFE-coated fiberglass fabric with 18 1/2 gauge holes was used as a thimble. Each film was weighed before and after extraction, where the extracted film was collected and dried under vacuum until constant weight to remove any residual THF. The gel− sol fraction was calculated using eq 4.15

⎡ ⎛ mi − mf ⎞⎤ gel−sol fraction (%) = ⎢1 − ⎜ ⎟⎥100 ⎢⎣ ⎝ mi ⎠⎥⎦

(4)

where mi is the initial mass and mf is the final mass (mass after extraction) of the covalently cross-linked film. 2.3.3. Attenuated Total Reflectance Fourier Transform Infrared Spectroscopy (ATR-FTIR). The secondary structures of linear and cross-linked peptidic-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 assignment of secondary structures via calculation of the area under the peaks.16,17 2.3.4. 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 peptide−polyurea films of both linear and network systems. Phase images were collected in tapping mode with silicon tips from Bruker (333−380 kHz, 110−140 μm). 2.3.5. Small-Angle X-ray Scattering (SAXS). SAXS data was acquired at the Center for Nanophase Materials Sciences (CNMS) in Oak Ridge National Laboratory on an Anton Paar SAXSess mc2. 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 Kratky camera giving rise to a beam size of 18 mm (length) x 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 was processed using open access software (SASview version 3.1.1). 2.3.6. Dynamic Mechanical Analysis (DMA). Films were cut into approximately (9 × 3.5 × 0.2) mm specimens for thermomechanical analysis using a TA Instruments Q800 DMA. An operating frequency of 1 Hz and a heating rate of 3 °C/min were utilized in a temperature range of −145 to 150 °C. A minimum of three samples were tested for each specimen. Storage modulus and tan δ were measured for the loss coefficient and glass transition temperature (Tg) determination. C


Article

Biomacromolecules

Figure 2. Molecular structure comparison of the linear and network polyurea−PBLA hybrids.

Figure 3. ATR-FTIR region of secondary structure analysis for the linear and network polyurea−PBLA hybrids along with the relative amount of βsheet secondary structures.

system, and connectivity of either linear or network. For example, the sample PBLA5-10 (network) is a covalently crosslinked polyurea-PBLA hybrid containing five repeat units of PBLA and 10 wt % of PBLA. Control films without PBLA were also fabricated using the nomenclature PDMS−HDI (connectivity). Table S2 details the molecular weight (Mn) or molecular weight between cross-links (Mc) for the linear and network systems, respectively. Due to the synthetic design and resulting architecture of the hard domain obtained during covalent cross-linking, the molecular weight between cross-links increased with increasing peptide weight fraction, which is expected to result in enhanced elasticity.4,25 3.2. Secondary Structure Characterization. It was anticipated that a network architecture would influence the hydrogen bonding arrangement of the PBLA segments within the soft domain in comparison to a linear, physically crosslinked system. To examine the peptide secondary structure within these polyurea−PBLA hybrids, attenuated total reflectance Fourier transform infrared spectroscopy (ATRFTIR) was utilized. ATR-FTIR was chosen due to its reliability and sensitivity as it relates to peptide secondary structure determination and the ease of use for solid state film

and network) via step-growth polymerization with an isocyanate, and the peptide weight fraction was controlled by the addition of excess amine terminated-PDMS. A depiction of the resulting linear and network structures is shown in Figure 2. The network architecture was developed within the hard block through urea bond formation of the HDI trimer with either amine-terminated PDMS or PBLA5-b-PDMS-b-PBLA5. Several factors must be considered when network formation occurs simultaneously with film fabrication, including isocyanate to amine ratio ([NCO]/[NH2]),20,21 solubility, and solvent volatility.22 It is well-understood that the [NCO]/ [NH2] may influence material properties due to potential side reactions (e.g., biuret group formation)23 that lead to network defects, such as dangling chains.14 Both [NCO]/[NH2] and solvent type also influence the gel−sol fraction in the network24 and the degree of phase separation of the final film.22 Via optimization (Figure S2) of solvent type and [NCO]/[NH2], we obtained clear, uniform films with gel−sol fractions of ∼85% using THF as the reaction/casting solvent and a [NCO]/ [NH2] = 1.125. The following nomenclature was utilized: PBLA5-Z (connectivity), where Z is the peptide weight percentage within the D


Article

Biomacromolecules

Figure 4. Proposed secondary structure arrangement in polyurea−PBLA hybrids for both the linear and network samples. (A) PBLA5-10 linear, (B) PBLA5-40 linear, (C) PBLA5-10 network, (D) PBLA5-40 network. (E) Definition of the peptide secondary structures as well as the molecular structure of the linkages present.

Figure 5. Morphological characterization via AFM and SAXS. AFM phase images: a) PBLA5−10 (linear), b) PBLA5−20 (linear), c) PBLA5−40 (linear), d) PBLA5−10 (network), e) PBLA5−20 (network), and f) PBLA5−40 (network) of 1 μm × 1 μm and scale bar of 300 nm. SAXS plots: (g) linear and (h) network.

testing.8,26−29 Amide I bands (vCO) of 1620 to 1635 cm−1 were assigned to anti-parallel β-sheets, and amide I bands of 1635 to 1645 cm−1 were designated as parallel β-sheets.8,30 Figure 3 details the calculated content of parallel and antiparallel β-sheets for the polyurea-PBLA hybrids (network and linear). The full ATR-FTIR spectrum of the polyurea−peptide hybrids is shown in Figure S3. As expected, only β-sheet secondary structures were present in the linear and network polyurea-PBLA hybrids due to the short peptide repeat length.8,27 Within the linear polyureaPBLA systems, a mixture of parallel and anti-parallel β-sheets was observed, which was relatively unaffected by peptide content. The anti-parallel β-sheet motif was more favorable due to the higher stability of anti-parallel over parallel sheet structures. This structural observation is commonly exhibited in peptides and was also observed in the PBLA5-b-PDMS-bPBLA5 triblock.26,31 Upon introduction of covalent crosslinking via the trifunctional isocyanate, parallel PBLA β-sheets were exclusively formed, which is likely due to a combination of synthetic approach and chain mobility (Figure 4). During synthesis of both linear and network samples, the PBLA Nterminus was covalently attached to the isocyanate units,

resulting in the same peptide directionality (N-terminus reacted with isocyanate to form a urea linkage). This imposed directionality of the PBLA block requires mobility of the chains to induce an anti-parallel β-sheet conformation, which was easily achieved in the linear architecture. In the covalently cross-linked polyurea-PBLA hybrids, the network architecture restricted movement of the PBLA segments, limiting antiparallel β-sheet peptide interactions. This change in secondary structure, which was attributed to peptide segment ordering, is anticipated to alter the fibrillar morphology.27,32,33 Therefore, the morphology as a function of peptide content in network polyurea hybrids was investigated. 3.3. Influence of Network Architecture on Morphological Development. AFM phase images (Figure 5) were obtained to compare the surface morphology of the linear and network polyurea−PBLA systems. Nanofibrous or ribbon-like structures were observed in the linear peptidic films (Figure 5a−c). The observed fibrillar morphology was consistent with prior experimental investigations of polyurea systems and peptide−polymer conjugates as well as with theoretical studies utilizing coarse-grained modeling where the hydrogen bonding motifs played a dominant role in microstructure developE


Article

Biomacromolecules

aggregates analogous to the aggregated fibrils observed in AFM.24,48 From the AFM and SAXS analysis, the difference in assembly between the linear (i.e., physically cross-linked) and network (i.e., covalently cross-linked) polyurea−PBLA hybrids was attributed to spatial constraints, which disrupted the hydrogen-bonding propensity within the PBLA-containing soft domains. The interconnected fibrillar morphology in the linear hybrids is expected to aid in stress transfer during deformation,8,50 while the globular domain formation is expected to exhibit systematic softening of the sample and reduce the brittle character of the films.4,7,47 3.4. Influence of Network Architecture on Film Mechanics. 3.4.1. Dynamic Mechanical Characterization. Storage modulus (Figure 6) and tan δ (Figure S4) curves were

ment.7,8,28,34−38 In contrast, fiber-like nanostructures were not observed in the polyurea−PBLA networks (Figure 5d−f). At low peptide content (10 wt %), a smooth homogeneous surface was revealed. Above 10 wt %, a non-fibrous, globular-like microphase segregated morphology was visualized, where the soft segment appears dark and the hard domain is the lighter phase. It is proposed that the absence of ribbon-like structures in the covalently cross-linked polyurea-PBLA hybrids was due to the spatial constraints on the PBLA segments, reducing their ability to self-associate and inhibiting long-range order as a result of disrupted hydrogen bonding.28,39−42 This proposed confinement was supported by the observed differences in PBLA β-sheet conformation with covalent cross-linking (Figure 3). Small-angle X-ray scattering (SAXS) measurements were conducted to complement AFM analysis.5,7,8,43 As observed in Figure 5g,h, the presence of a covalent network hindered the packing of hard segments, which has been shown to alter phase separation and morphology.39,42,44 In the linear hybrids, the control PDMS-HDI film exhibited no phase separation (Figure 5g), evident by the lack of a SAXS correlation peak, which is expected for a non-chain extended systems with low hard segment content (18 wt %).4 In contrast, phase separation was observed in the network PDMS control, which exhibited a peak maximum at q = 0.9 nm−1 and was attributed to the incompatibility of the siloxane segments and the multifunctional hard segment. Upon peptide incorporation into the polyureas, both the linear and the network systems exhibited phase separation with a slight increase in domain spacing (Lp) as peptide fraction increases: Lp ∼ 8 nm (10 wt %) to Lp ∼ 10 nm (40 wt %). This shift in peak position to lower q values is most likely due to the incorporation of rigid PBLA segments, increasing the size of the “pseudo” hard block (peptide + HS) and interdomain spacing.8 We also examined the breadth of the scattering peak; peak broadening is indicative of an increase in phase mixing and/or a wider distribution of domain sizes and spacings.7,42,45 Differences in phase mixing between the linear and network hybrids were apparent. In the linear hybrids, the addition of PBLA sharpens the scattering peak (q = 0.9 nm−1), yielding a distinct peak at the highest peptide loading (PBLA5-40 linear). In contrast, peak broadening was observed for the peptide− polyurea network hybrids with increasing peptide content. Several factors may lead to peak broadening in the PBLA polyurea network, including an increase in Mc with increasing peptide content, which would alter the distribution of domain sizes and spacings,4 and/or partial miscibility of the PBLA segments in both the hard (HDI trimer) and soft (PDMS) phases. Shifts in microstructure upon chemical cross-linking have been noted previously, where either limited short-range ordering or complete disruption of microstructure development resulted.24,42,46−49 For example, Bos et al. reported a lack of a microphase segregated morphology in cross-linked poly(propylene glycol) based polyurethanes, which was attributed to the absence of hydrogen bonding between the urethane units.46 In a similar manner, the network induced-changes in PBLA secondary structure in this investigation inhibited the long-range fibrillar arrangement of the peptide motifs. Additionally, variations in the initial slope were seen at lower q values on the log−log plots; these variations were dependent on the peptide content in the linear systems, suggesting the existence of a larger, assembled super structure, such as fibrillar

Figure 6. DMA temperature sweep of linear and covalently crosslinked polyurea−PBLA hybrids.

obtained from a temperature sweep of the linear and network polyurea−PBLA hybrids utilizing dynamic mechanical analysis (DMA). In both the linear and network hybrids, an increase in plateau modulus with increasing peptide weight fraction was seen, most likely due to the “pseudo” hard segment character of the PBLA blocks.8 A reduction in the PDMS (Tg ∼ −125 °C) tan δ peak height was observed for both systems with increasing PBLA wt % due to a reduction in the amount of PDMS upon increasing peptide content (Figure S4, Table S1). With additional PBLA content, a higher modulus was maintained after the Tg (∼44 °C)8 of the peptide unit, which was attributed to more interchain hydrogen bonding of the βsheet PBLA domains. For the network polyurea−PBLA hybrids, the interchain hydrogen bonding of the parallel βsheet motif was more easily influenced by the ramp in temperature compared to the linear analogue, which contained more stable anti-parallel β-sheets. This disruption in PBLA hydrogen bonding organization led to a reduction in modulus for the network hybrids, most prominently seen above the peptide Tg in the DMA curve. For instance, the 10 wt % PBLA hybrids were able to maintain modulus in the linear sample up to 100 °C, while a drop in modulus at only 37 °C was observed for PBLA5-10 (network). Previous reports of network formation altering physical interactions within polyureathanes and polyureas have been shown to result in changes in material properties.7,41 Sheth et al. studied the importance of hydrogen bonding and chain architecture in mediating the long-range connectivity of the HS phase in an poly(ethylene oxide)/ poly(propylene oxide)-based poly(urethane/urea) system, F


Article

Biomacromolecules

Figure 7. (a) Representative stress−strain curves of linear and network polyurea−PBLA hybrids. (b) Calculated values of modulus, elongation-atbreak, and toughness.

Figure 8. Representative hysteresis curves and summary of quantified % hysteresis data for PBLA5-20 samples (both linear and network films) of select cycles.

formation had little effect on the elastic modulus as highlighted by DMA analysis. As expected, the elastic modulus increased with PBLA content in both the linear and covalently cross-linked materials. The network film PBLA5-40 exhibited the highest modulus at ∼70 MPa. This enhancement in modulus was attributed to peptide rigidity and increased β-sheet hydrogen bonding.51 The linear polyurea−PBLA hybrids displayed a higher elastic modulus compared to their network analogues at the same peptide fraction. It is proposed that the microphase segregated structure that induced fiber-like morphology observed in the linear hybrids facilitates stress-transfer via more fiber-to-fiber interactions with increasing peptide content.8,50 In comparison, the lack of fibrous domains and selectivity toward weaker parallel β-sheets in the polyurea−peptide hybrid networks resulted in lower elastic moduli. The reduction in extensibility within the linear hybrids as PBLA content increased was attributed to the increased peptide hydrogen bonding and rigidity. For the network systems, Mc increased with increasing PBLA content, enhancing extensibility.25,52 However, the overall extensibility was higher in the linear hybrids compared to the network polyurea-PBLA systems due the architectural constraints imposed by the multifunctional HDI cross-linking sites. For both the linear and network polyurea-PBLA systems, a defined yield point was not

concluding that these were critical factors in limiting ordering and percolation of the HS and achieving dimensional stability. Upon increasing network formation, the long-range connectivity of the HS was increasingly disrupted, which was accompanied by systematic mechanical softening.7 Likewise, as reported here, a reduction in storage modulus for the network hybrids was observed. Furthermore, the PBLA5-40 (linear) system was too brittle for evaluation via thermomechanical or tensile analysis, which is likely a result of its extensive hydrogen bonding. However, the PBLA5-40 (network) system exhibited systematic mechanical softening, which mitigated the brittle mechanics and yielded stable films suitable for testing. Focusing on the DMA findings, it is anticipated that variations in hydrogen bonding will affect the mechanical yielding and tensile hysteresis of the materials, depending on the dynamic character and prevalence of physical interactions. 3.4.2. Tensile Analysis. Stress−strain curves were obtained to compare the mechanical behavior, including elastic modulus (Et), toughness, and elongation-at-break (εb), of the linear and network polyurea−PBLA systems (Figure 7). The control polyurea PDMS-HDI network exhibited a significant reduction in extensibility compared to the linear polyurea PDMS-HDI, shifting from ∼1200% to ∼44% (Figure S5). The elastic modulus was relatively consistent (∼9 MPa) in both control architectures (linear, network), suggesting that the network G


Article

Biomacromolecules

phase separation was obtained in the network films, a globularlike microstructure was observed. As peptide content increased, the phase domain size increased slightly from 8 to 10 nm due to the addition of “pseudo” PBLA hard domains. A modulation in mechanics resulted from these variations in architecture via the incorporation of peptide hydrogen bonding and a permanent network. The absence of anti-parallel β-sheets upon covalent cross-linking weakened hydrogen bond formation, resulting in a reduction of storage modulus. However, the polyurea−PBLA network enhanced film stability, allowing for higher peptide loadings with improved Young’s modulus. An increase in Mc accompanied higher peptide amounts, which allowed the material to maintain extensibility as the modulus increased. Tensile hysteresis experiments confirmed that both physical and covalent bonds influence the deformation response; the network architecture limits energy absorption, while the peptide hydrogen bonding enhances dynamics. These bioinspired peptide hybrids offer a unique platform that overlays architecture and dynamics for responsive applications in therapeutics, scaffolds, and coatings.

observed, suggesting that the physical influence of adjacent hydrogen bonds was not well-dispersed with the exception of sample PBLA5-40 (network). The PBLA5-40 (network) hybrid exhibited a well-defined yield point, suggesting continuous hydrogen bonding throughout the film. The resulting stress− strain curves have shown that a combination of percolating physical interactions and covalent cross-links leads to both enhanced modulus and extensibility. Hysteresis experiments were conducted on the PBLA5-20 linear and network hybrids to explore the influence of structural interplay on unloading and loading mechanical response. A peptide content of 20 wt % was chosen because it contained the highest amount of peptide that maintained film stability in both the linear and network architectures. In the control films (Figure S6), a large hysteresis was noted in the first cycle (∼70%) that reduces to ∼35% hysteresis after 20 cycles. Upon covalent cross-linking, the hysteresis was decreased, where the first cycle exhibited a hysteresis of ∼30% that was reduced to ∼20% after 20 cycles. This reduction in hysteresis upon network formation suggests that more energy is stored compared to the linear control film.53 These control studies are supported by prior literature, in which the formation of chemical cross-links has been shown to reduce hysteresis by disrupting the hard domain structure.54 The addition of physical interactions via peptide incorporation (Figure 8) yielded higher mechanical energy storage for the linear systems compared to the network architecture due to the ability of the physical associations to be more easily reformed without geometric constraints. It is proposed that the dynamic physical interactions in the network film, due to the parallel β-sheets, are weaker, and reassociation after deformation is hindered by the mobility of network junction points. Furthermore, the domain structure of segmented polyureas may play a critical role in hysteresis.55,56 For example, Staudinger et al. studied the tensile hysteresis behavior of multigraft copolymers in which various morphologies (spheres, cylinders, lamellae) were achieved.55 Copolymers with lamellae structures exhibited minimal residual strain in hysteresis, while spherical structures displayed the highest hysteresis behavior. Similarly, the globular-like morphology of the PBLA5-20 (network) film exhibited a larger hysteresis loop than the fiber-like PBLA5-20 (linear) film. Overall, the tensile properties were dependent upon the interplay of both physical and covalent interactions. We have achieved a high peptide loading within these peptide-polyurea hybrids that resulted in elongations comparable to natural materials, such as elastin (150%),3 with enhanced Young’s modulus, and variable tensile hysteresis dependent on the presence of peptide and covalent cross-linking.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.biomac.6b01309. Synthetic reaction quantities as well as NMR, ATRFTIR, and molecular weight data; optimization of film casting for the network polyurea hybrids is also presented with corresponding images of film quality; additional tan δ, and mechanical tensile data of control films can be found in the Supporting Information (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. J.K. contributed by obtaining and aiding in the analysis of SAXS data. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors acknowledge funding support from the National Science Foundation (CAREER DMR- 0953236). The authors thank Richard Tomazin from the Swagelok Center for Surface Analysis of Materials at Case Western Reserve University for AFM image assistance on all film samples.

4. CONCLUSIONS In this work, a molecular design strategy for PBLA−polyurea hybrids was utilized to evaluate the role of architectural arrangements on peptide ordering and mechanics. The addition of a HDI trimer to polyurea step growth reactions served as the covalent cross-linker. By varying the PBLA content, variations in hydrogen bonding propensity were achieved within covalent networks, allowing for both physical and covalent cross-linking. Unlike the mixture of anti-parallel and parallel β-sheets present in the linear PBLA−polyurea hybrids, only parallel β-sheet formation was observed in the network hybrids, influencing the overall morphology. The fibrous morphology of the linear analogues was disrupted upon the formation of a network topology; although

■

REFERENCES

(1) Johnson, J. C.; Korley, L. T. J. Soft Matter 2012, 8, 11431−11442. (2) Gosline, J. M.; Guerette, P. A.; Ortlepp, C. S.; Savage, K. N. J. Exp. Biol. 1999, 202, 3295−3303. (3) Mithieux, S. M.; Rasko, J. E. J.; Weiss, A. S. Biomaterials 2004, 25, 4921−4927. (4) Szycher, M. Szycher’s Handbook of Polyurethanes, 2nd ed.; CRC Press: Boca Raton, FL, 2012; p 37−86. (5) Johnson, J. C.; Wanasekara, N. D.; Korley, L. T. J. Biomacromolecules 2012, 13, 1279−1286.

H


Article

Biomacromolecules (6) Choi, T.; Weksler, J.; Padsalgikar, A.; Runt, J. Polymer 2009, 50, 2320−2327. (7) Sheth, J. P.; Wilkes, G. L.; Fornof, A. R.; Long, T. E.; Yilgor, I. Macromolecules 2005, 38, 5681−5685. (8) Johnson, J. C.; Wanasekara, N. D.; Korley, L. T. J. J. Mater. Chem. B 2014, 2, 2554−2561. (9) Vandermeulen, G. W. M.; Klok, H. A. Macromol. Biosci. 2004, 4, 383−398. (10) Shu, J. Y.; Panganiban, B.; Xu, T. Annu. Rev. Phys. Chem. 2013, 64, 631−657. (11) Cheng, J.; Deming, T. J. Top. Curr. Chem. 2011, 310, 1−26. (12) Kricheldorf, H. R. Anionic Ring-Opening Polymerization: NCarboxyanhydrides. In Comprehensive Polymer Science; Allen, G., Bevington, J. C., Eds.; Pergamon Press: Oxford, U.K., 1989; p 531− 551. (13) Montalbetti, C.; Falque, V. Tetrahedron 2005, 61, 10827− 10852. (14) Fridrihsone, A.; Stirna, U.; Lazdiņa, B.; Misane, ̅ M.; Vilsone, D. Eur. Polym. J. 2013, 49, 1204−1214. (15) Besteiro, M. C.; Guiomar, a. J.; Gonçalves, C. a.; Bairos, V. a.; De Pinho, M. N.; Gil, M. H. J. Biomed. Mater. Res., Part A 2010, 93, 954−964. (16) Luo, S.; Huang, C.-Y. F.; McClelland, J. F.; Graves, D. J. Anal. Biochem. 1994, 216, 67−76. (17) Rathore, O.; Sogah, D. Y. J. Am. Chem. Soc. 2001, 123, 5231− 5239. (18) Dou, Q.; Wang, C.; Cheng, C.; Han, W.; Thüne, P. C.; Ming, W. Macromol. Chem. Phys. 2006, 207, 2170−2179. (19) Javni, I.; Zhang, W.; Petrović, Z. S. J. Appl. Polym. Sci. 2003, 88, 2912−2916. (20) Bagdi, K.; Molnár, K.; Wacha, A.; Bóta, A.; Pukánszky, B. Polym. Int. 2011, 60, 529−536. (21) Kasikci, H.; Pekel, F.; Ozkar, S. J. Appl. Polym. Sci. 2001, 80, 65− 70. (22) Dušková-Smrčková, M.; Dušek, K.; Vlasák, P. Macromol. Symp. 2003, 198, 259−270. (23) Dusek, K.; Spirkova, M.; Havlicek, I. Macromolecules 1990, 23, 1774−1781. (24) Krakovský, I.; Pleštil, J.; Baldrian, J.; Wübbenhorst, M. Polymer 2002, 43, 4989−4996. (25) Oprea, S. Adv. Polym. Technol. 2009, 28, 165−172. (26) Tanaka, S.; Ogura, A.; Kaneko, T.; et al. Macromolecules 2004, 37, 1370−1377. (27) Ibarboure, E.; Papon, E.; Rodríguez-Hernández, J. Polymer 2007, 48, 3717−3725. (28) Ibarboure, E.; Rodríguez-Hernández, J. Eur. Polym. J. 2010, 46, 891−899. (29) Kong, J.; Yu, S. Acta Biochim. Biophys. Sin. 2007, 39, 549−559. (30) Ma, M.; Zhong, J.; Li, W.; Zhou, J.; Yan, Z.; Ding, J.; He, D. Soft Matter 2013, 9, 11325. (31) Chou, K.-C.; Pottle, M.; Némethy, G.; Ueda, Y.; Scheraga, H. A. J. Mol. Biol. 1982, 162, 89−112. (32) Klinedinst, D. B.; Yilgör, I.; Yilgör, E.; Zhang, M.; Wilkes, G. L. Polymer 2012, 53, 5358−5366. (33) Termonia, Y. Biomacromolecules 2004, 5, 2404−2407. (34) Johnson, J. C.; Korley, L. T. J.; Tsige, M. J. Phys. Chem. B 2014, 118, 13718−13728. (35) McLean, R. S.; Sauer, B. B. Macromolecules 1997, 30, 8314− 8317. (36) Tiné, M. R.; Alderighi, M.; Duce, C.; Ghezzi, L.; Solaro, R. J. Therm. Anal. Calorim. 2011, 103, 75−80. (37) Versteegen, R. M.; Sijbesma, R. P.; Meijer, E. W. Macromolecules 2005, 38, 3176−3184. (38) Shao, H.; Parquette, J. R. Angew. Chem., Int. Ed. 2009, 48, 2525− 2528. (39) Janik, H.; Vancso, J. Polimery 2005, 50, 139−142. (40) Klinedinst, D. B.; Yilgör, I.; Yilgör, E.; Zhang, M.; Wilkes, G. L. Polymer 2012, 53, 5358−5366.

(41) Sheth, J. P.; Klinedinst, D. B.; Wilkes, G. L.; Yilgor, I.; Yilgor, E. Polymer 2005, 46, 7317−7322. (42) Wang, L.; Hickner, M. a. Polym. Chem. 2014, 5, 2928. (43) Das, S.; Yilgor, I.; Yilgor, E.; Inci, B.; Tezgel, O.; Beyer, F. L.; Wilkes, G. L. Polymer 2007, 48, 290−301. (44) Bouhelal, S.; Cagiao, M. E.; Benachour, D.; Djellouli, B.; Rong, L.; Hsiao, B. S.; Balta-Calleja, F. J. J. Appl. Polym. Sci. 2010, 117, 3262− 3270. (45) Chang, S. L.; Yu, T. L.; Huang, C. C.; Chen, W. C.; Linliu, K.; Lin, T. L. Polymer 1998, 39, 3479−3489. (46) Bos, H. L.; Nusselder, J. J. H. Polymer 1994, 35, 2793−2799. (47) Pichon, P. G.; David, L.; Méchin, F.; Sautereau, H. Macromolecules 2010, 43, 1888−1900. (48) Schmidt, K.; Tassone, C. J.; Niskala, J. R.; Yiu, A. T.; Lee, O. P.; Weiss, T. M.; Wang, C.; Fréchet, J. M. J.; Beaujuge, P. M.; Toney, M. F. Adv. Mater. 2014, 26, 300−305. (49) Pergal, M. V.; DŽ unuzović, J. V.; Ostojić, S.; Pergal, M. M.; Radulović, A.; Jovanović, S. J. Serb. Chem. Soc. 2012, 77, 919−935. (50) Fu, S. Y.; Lauke, B. Compos. Sci. Technol. 1996, 56, 1179−1190. (51) Keten, S.; Buehler, M. J. Nano Lett. 2008, 8, 743−748. (52) Genesky, G. D.; Cohen, C. Polymer 2010, 51, 4152−4159. (53) Nallicheri, R. A.; Rubner, M. F. Macromolecules 1991, 24, 526− 529. (54) Qi, H. J.; Boyce, M. C. Mech. Mater. 2005, 37, 817−839. (55) Staudinger, U.; Weidisch, R.; Zhu, Y.; Gido, S. P.; Uhrig, D.; Mays, J. W.; Iatrou, H.; Hadjichristidis, N. Macromol. Symp. 2006, 233, 42−50. (56) Zhu, Y.; Burgaz, E.; Gido, S. P.; Staudinger, U.; Weidisch, R.; Uhrig, D.; Mays, J. W. Macromolecules 2006, 39, 4428−4436.

I


Network Architecture and Its Impact on the ... - ACS Publications

Recommend Documents