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Role of Hydrogen Bonding on Nonlinear Mechano-Optical Behavior of L‑Phenylalanine-Based Poly(ester urea)s Keke Chen,† Jiayi Yu,‡ Gustavo Guzman,∥ S. Shams Es-haghi,∥ Matthew L. Becker,‡,§ and Miko Cakmak*,†,∥,⊥ †

Department of Polymer Engineering, ‡Department of Polymer Science, and §Department of Biomedical Engineering, The University of Akron, Akron, Ohio 44325, United States ∥ School of Materials Engineering and ⊥School of Mechanical Engineering, Purdue University, West Lafayette, Indiana 47907, United States S Supporting Information *

ABSTRACT: The uniaxial mechano-optical behavior of a series of amorphous L-phenylalanine-based poly(ester urea) (PEU) films was studied in the rubbery state. A custom, real-time measurement system was used to capture the true stress, true strain, and birefringence during deformation. When the materials were subjected to deformation at temperatures near the glass transition temperature (Tg), the photoelastic behavior was manifested by a small increase in birefringence with a significant increase in true stress. At temperatures above Tg, PEUs with a shorter diol chain length exhibited a liquid−liquid (Tll) transition (rubbery−viscous transition) at about 1.06Tg (K) under the tested strain rate of 0.017 s−1 (stretching speed of 20 mm/min), above which the material transforms from a heterogeneous “liquid of fixed structure” to a “true liquid” state. The initial photoelastic behavior disappears with increasing temperature, as the initial slope of the stress optical curves becomes temperature independent. Fourier transform infrared spectroscopy (FTIR) was used to study the effect of hydrogen bonding on the physical properties of PEUs as a function of temperature. The average strength of hydrogen bonding diminishes with increasing temperature. For PEUs with the longest diol chain length, the area associated with N−H stretching region exhibits a linear temperature dependence. However, a three-stage temperature dependence was observed for PEUs with shorter diol chain length. The presence of hydrogen bonding enhances the “stiff” segmental correlations between adjacent chains in the PEU structure. As a result, the photoelastic constant decreases with increasing hydrogen bonding strength. repair of bone defects.12 The synthesis and characterization of these materials have been studied extensively.9−13 The elastic moduli of poly(ester urea)s are tunable and similar to those of commercially available poly(lactic acid).14 The properties of this family of materials are affected by chain architecture and the use of different functional groups.9,10,12,15 For example, PEU has been functionalized with different iodine content to increase radiopacity for enhanced contrast in fluoroscopic imaging for medical devices.13 The introduction of branching units also changes the mechanical properties.10 The mechanical and physical properties of PEUs are significantly influenced by processing conditions, specifically by the chain orientation development during processing. The barrier properties and mechanical performance of the films can be improved by inducing the preferential chain orientation by uniaxial (machine direction orientation MDO) and biaxial stretching (tenter frame, double bubble, etc.)16 Harnessing this performance

1. INTRODUCTION Significant efforts have focused on the development of degradable polymers for biomedical applications. Degradable polymers such as poly(ε-caprolactone) (PCL), poly(lactide) (PLA), poly(glycolide) (PGA), and their copolymers have been used widely for applications in regenerative medicine and drug delivery.1−4 Some other clinically used degradable polymers include poly(carbonates) and poly(urethanes).5,6 While their properties have been studied extensively in vitro and in vivo, these materials are limited by their relatively narrow range of physical and chemical properties.7,8 Amino acid-based poly(ester urea)s (PEU)s have been utilized in a number of biomedical applications due to their tunable mechanical properties, reactive functionalization handles, and nontoxic degradation byproducts.9−12 The incorporation of L-phenylalanine amino acids in backbone increases the mechanical strength and retains the noncytotoxic properties required for regenerative medicine applications.9 The phenylalanine-based PEU system has been functionalized with osteogenic growth peptide (OGP) to impart bioactive signaling and showed promising results in osteogenic differentiation in the surgical © XXXX American Chemical Society

Received: November 7, 2016 Revised: January 13, 2017

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Scheme 1. A Two-Step General Synthetic Route of L-Phenylalanine-Based Poly(ester urea)s with Different Diol Chain Lengths

requires an in situ method for assessing chain orientation during film formation or in a postformation processing step. Birefringence17,18 provides useful information regarding polymer chain orientation and the associated molecular mechanism.19−24 When captured in real time along with the evolution of true stress and strain, birefringence provides molecular insights into the mechanistic changes during the structural evolution during processing.25−31 Numerous studies have shown a direct relationship between the measured birefringence and the state of stress of materials.17,18,32 The birefringence is found to follow the linear stress optical rule (SOR) at small to moderate deformations in true fluids:24,32−38 Δn = Cσ

Table 1. Summary of the L-Phenylalanine-Based Poly(ester urea)s with Different Diol Chain Lengths

a

samplesa

Mn

Mw

ĐM

Poly(1-PHE-8) Poly(1-PHE-10) Poly(1-PHE-12)

48000 35000 22000

91000 71000 38000

1.9 2.0 1.7

Sample notation represents the number of carbon atoms in the diol.

Poly(1-PHE-10), and 103 °C for Poly(1-PHE-12). Pressures of 35, 70, and 35 MPa were applied for 5 min in sequence. After that, the system was cooled down with a pressure of 6.9 MPa to prevent wrinkle formation on the film surface. The vacuum was maintained during the entire process to minimize hydrolytic degradation. Compression molded PEU films were visually inspected to ensure that there were no defects or bubbles present. Thermal Analysis. Differential scanning calorimetry (TA, DSC Q200) was used to measure the glass transition temperatures (Tg) of assynthesized PEU polymers and compression molded PEU films. Samples (∼5 mg) were sealed in hermetic aluminum pans and were scanned at a heating rate of 10 °C/min in a dry nitrogen atmosphere. The glass transition temperature can be defined in different ways.43 In the present study, the Tg of PEU films refers to the onset temperature instead of the inflection point temperature, defined from the intersection of the first deviation from the baseline on the lowtemperature side. Stress−Strain Behavior with Online Birefringence. Dumbbellshaped PEU samples were cut from the compression molded films. The narrowest region of the film samples was 21 mm wide and 20 mm long between the clamps. A real-time mechano-optical measurement instrument designed by our group was used for this study.27,44,45 In brief, the system simultaneously records optical retardation Γt, using spectral birefringence method, and continuously monitors sample width Wt and force Ft during stretching. This allows continuous determination of instantaneous cross-sectional area. With the assumption of (a) simple uniaxial extension and (b) incompressibility, the real-time sample thickness Dt is calculated using

(1)

where Δn is the birefringence, C (Pa−1) is the stress optical constant, and σ (Pa) is the true stress. However, it has been shown that the SOR deviates from the linearity beyond low to moderate stress levels or during phase transition stages such as strain-induced crystallization.32,33,39,40 The SOR has been widely applied in many polymer systems such as polymer melts and amorphous polymers in order to optimize process conditions for desired properties.41,42 In this paper, we present a real-time study of the mechanooptical birefringence of a series of PEU polymers during uniaxial deformation over a range of temperatures. Special focus is placed on the effect of diol chain length in the PEU structure on material properties. The real-time study of the mechanooptical behavior is also coupled with the infrared spectra obtained as a function of temperature to study the effect of hydrogen bonding on the mechano-optical properties of PEUs.

2. EXPERIMENTAL SECTION Materials. The potassium bromide (KBr) window was purchased from Pike Technologies (Pike Technologies, Fitchburg, WI). 1,1,1,3,3,3 Hexafluoro-2-propanol (HFIP) was purchased from CovaChem (CovaChem, LLC., Loves Park, IL). The synthesis of L-phenylalanine-based poly(ester urea)s (PEUs) containing different diol chain lengths was reported previously.9 Briefly, L-phenylalanine-based poly(ester urea)s were synthesized in a two-step process. First, an esterification was carried out between Lphenylalanine and 1,8-octanediol, 1,10-decanediol, and 1,12-dodecanediol under acidic conditions. Next, an interfacial polymerization of the respective monomers with triphosgene yielded the poly(ester urea) homopolymers (Scheme 1). Table 1 shows the number-average molecular mass (Mn), the weight-average mass (Mw), and postprecipitation molecular mass distribution (ĐM) for the respective PEUs. Additional synthesis and characterization details can be found in the Supporting Information. Sample Preparation. PEU films (∼500 μm) were compression molded using a vacuum compression press (TMP Technical Machine Products Corp.). Prior to compression, PEU polymers were dried under vacuum for 24 h at 60 °C. The compression molding temperatures were set at 135 °C for Poly(1-PHE-8), 122 °C for

⎛W ⎞ Dt = ⎜ t ⎟D0 ⎝ W0 ⎠

(2)

where W0 is the initial sample width, Wt is the real-time sample width, and D0 is the initial sample thickness. Thus, the real-time spectral birefringence, true stress and true strain are determined using Δnt =

Γt Dt

true strain =

true stress =

B

(3)

⎛ W ⎞2 elongation L = t − 1 = ⎜ 0⎟ − 1 initial length L0 ⎝ Wt ⎠

Ft = WD t t

(4)

Ft Wt 2 W0

( )D

0

(5) DOI: 10.1021/acs.macromol.6b02415 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules where Γt, Dt, and Lt are the real-time optical retardation, sample thickness, and sample length, respectively; L0 is the initial sample length, and Ft is the real-time force. PEU samples were uniaxially stretched at a constant stretching speed of 20 mm/min to 300% strain in their rubbery state at temperatures above the glass transition temperature. Table 2 lists the

Table 2. Stretching Temperatures for the L-PhenylalanineBased Poly(ester urea)s with Different Diol Chain Lengths samples

Tg (°C)

stretching temp (°C)

Poly(1-PHE-8)

45

Poly(1-PHE-10)

35

Poly(1-PHE-12)

20

55 60 65 75 45 50 55 65 25 35 45 50

stretching temperatures for the PEU polymers during the experiments. Prior to each experiment, PEU samples were thermally equilibrated in the environmental chamber at the desired temperature for at least 10 min. After each experiment, the samples were cooled down to room temperature by air before being removed from the clamps. Fourier Transform Infrared (FTIR) Spectroscopy. Samples for FTIR analysis were prepared by casting thin films (∼3 μm) onto KBr windows (32 × 3 mm, Pike Technologies, Fitchburg, WI) from PEU solution (2% w/v) in 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP) (CovaChem, LLC., Loves Park, IL) at room temperature. Samples were kept in a vacuum oven for about 48 h at 60 °C to remove the residual solvent and moisture. Infrared spectra were recorded using a Thermo Electron Nicolet 4700 Fourier transform infrared spectrometer. The resolution was set at 2 cm−1, and a minimum of 64 scans were used to acquire the signal. An Omega temperature controller was used to obtain the spectra as a function of temperature, ranging from 30 to 180 °C, at 10 °C increments. After the sample was mounted inside the cell, the samples were allowed a thermal equilibrium time of 15−20 min before each scan.

Figure 1. True stress−true strain curve of Poly(1-PHE-8) (Tg ∼ 45 °C) (A), Poly(1-PHE-10) (Tg ∼ 35 °C) (B), and Poly(1-PHE-12) (Tg ∼ 20 °C) (C) uniaxially stretched at different temperatures.

3. RESULTS AND DISCUSSION True Mechanical Behavior. The true stress−true strain curves of PEUs with different diol chain lengths are shown in Figure 1 during uniaxial stretching. The mechanical behavior of Poly(1-PHE-8), Poly(1-PHE10), and Poly(1-PHE-12) exhibits a similar temperature dependence when stretched at temperatures above Tg. At temperatures near Tg, the deformation of PEU initially shows an elastic region (true strain < ∼0.05), followed by a plastic deformation region where the stress−strain curve enters a regime with a slow increase in true stress with true strain. The elastic region becomes less prominent with increasing temperature. At temperatures near Tg, strain hardening is observed at the end of the plastic region, indicated by a steep rise in true stress with true strain. The onset of strain hardening is determined by tracing a straight line from the origin to where the point intersects with the stress−strain curve. The illustrative description of the method is marked in Figure 1. Table 3 summarizes the onset of strain hardening for all three PEU

Table 3. True Stress and True Strain at the Onset of Strain Hardening for PEU Polymersa

a

samples

σ (MPa)

ε

Poly(1-PHE-8) Poly(1-PHE-10) Poly(1-PHE-12)

4.36 4.50 6.38

1.45 1.23 0.71

σ = true stress, ε = true strain.

polymers at temperature near Tg. The increase of diol chain length shifts the onset point of strain hardening to a smaller strain. During strain hardening, the polymer chains become aligned and together with entanglements and approach their finite extensibilities leading to rapid rise in stress.46 Higher temperature significantly suppresses the onset of strain hardening but promotes the plastic deformation region to a longer range. Eventually the material exhibits “taffy pull” behavior as the chains simply slide passed one another without showing any sign of strain hardening. While Poly(1-PHE-8) and Poly(1-PHE-10) show “taffy pull” behavior at a similar C

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Figure 2. Mechano-optical behavior of Poly(1-PHE-8) (A), Poly(1-PHE-10) (B), and Poly(1-PHE-12) (C, the scale magnified plot is shown inside) stretched at various temperatures; photoelastic and stress optical coefficient as a function of temperature for Poly(1-PHE-8), Poly(1-PHE-10), and Poly(1-PHE-12) (D).

temperature of Tg + 30 °C, Poly(1-PHE-12) exhibits this behavior at a lower temperature of Tg + 15 °C, reflecting a more flexible and mobile structure due to their longer diol chain length. Mechano-Optical Behavior. The mechano-optical behavior of Poly(1-PHE-8), Poly(1-PHE-10), and Poly(1-PHE-12) stretched at different temperatures is shown in Figure 2. At temperatures near Tg (below Tg + 20 °C for Poly(1-PHE8) and Poly(1-PHE-10)), PEUs initially exhibit glassy behavior as deformation starts, depicted by little increase in birefringence with a significant increase in true stress. This initial region, denoted as photoelastic regime in Figure 2A−C, corresponds to the glassy behavior and associated primarily with bond bending upon the application of stress. This is a common behavior for amorphous polymers. For example, Muller and Pesce have observed similar glassy behavior in the initial region of the stress−optical curve of amorphous polycarbonate when stretched at temperature near Tg.32 When PEU films were further stretched beyond the photoelastic regime, a positive deviation could be seen for the slope of the stress−optical curve. This region is identified as regime I, where birefringence increases mostly linearly at a faster pace with stress. Following the regime I, the increase in birefringence transitions to a slower rate with increasing stress. The photoelastic regime, regime I, and regime II are identified from the differences in the slopes of mechano-optical curves. The slopes are marked by black dotted lines in Figures S2−S4 (Supporting Information), and the partition of different regimes is determined from the intersection of the lines. In the case of Poly(1-PHE-12) at the temperature near Tg, the regime II levels off and reaches a

plateau, where the polymer chains reach the maximum extensibility. Consistent with previous observations,30,47 the initial glassy regime is suppressed with increasing temperature. For Poly(1PHE-8) and Poly(1-PHE-10), the initial photoelastic regime is not observed above Tg + 20 °C but is replaced by a linear stress−optical behavior. As the temperature is further increased, relaxation dominates. Therefore, the stress−optical behavior only shows a linear relationship in the initial region and then is followed by positive deviation from the stress−optical curve, which occurs at 65 °C for Poly(1-PHE-8) and 55 °C for Poly(1-PHE-10). However, for Poly(1-PHE-12), the initial glassy behavior is observed even at Tg + 30 °C. The photoelastic constant is determined as the slope of the initial photoelastic regime, and the stress−optical constant (SOC) is obtained from the initial linear portion of the stress− optical curve above Tg + 20 °C for Poly(1-PHE-8) and Poly(1PHE-10). Figure 2D plots photoelastic constant and regime I slope as a function of temperature. The SOC is found to be 0.0032 MPa−1 for Poly(1-PHE-8) and 0.0023 MPa−1 for Poly(1-PHE-10). These results indicate that the intrinsic birefringence of Poly(1-PHE-10) is lower than that of Poly(1-PHE-8). It should be noted in Figure 2 that Poly(1-PHE-8) and Poly(1-PHE-10) both have a critical temperature (at Tg + 20 °C), beyond which the slope of the initial glassy regime and the slope of regime I becomes constant and linear. This temperature is the so-called “liquid−liquid transition temperature” (Tll), defined by Boyer as the temperature above which the polymers transforms from “a liquid of fixed structure” to “a true liquid”.48−50 There has been significant controversy about D

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Figure 3. Strain−optical behavior of Poly(1-PHE-8) (A), Poly(1-PHE-10) (B), and Poly(1-PHE-12) (C, the scale magnified plot is shown inside) stretched at various temperatures. Strain optical coefficient as a function of temperature for Poly(1-PHE-8), Poly(1-PHE-10), and Poly(1-PHE-12) (D).

Strain−Optical Behavior. The temperature dependence of the strain−optical behavior for Poly(1-PHE-8), Poly(1-PHE10), and Poly(1-PHE-12) is shown in Figure 3. If all the macroscopic deformation on the true strain were reflected on the microscopic deformation as measured by birefringence, a linear strain−optical curve should be expected. However, a nonlinear strain−optical behavior is observed for PEU polymers. At small deformations, a linear dependency of birefringence on true strain is observed. As the temperature is increased, the initial linear portion decreases to a lower strain level, and the nonlinear strain−optical behavior dominates. A strain−optical constant can be determined as the slope of the linear region in the birefringence−true strain curve. Figure 3 shows the strain optical constant of Poly(1-PHE-8), Poly(1-PHE-10), and Poly(1-PHE-12) as a function of temperature. The decreasing trend of strain optical constant with increasing temperature suggests that the polymer chains are oriented less efficiently, as chain relaxation increasingly dominates as temperature increases. At temperatures near Tg, the strain optical constant of PEU with longer diol chain length is found to be a larger value (0.0412 for Poly(1-PHE-12), 0.0281 for Poly(1-PHE-10), and 0.0262 for Poly(1-PHE-8)). The intrinsic optical anisotropy varies in both sign and magnitude from polymer to polymer, depending on the orientation of the dominating polarizable group relative to the chain backbone.55,56 Intrinsic optical anisotropy values for PEUs were theoretically calculated to be 77.10 × 10−25 cm3 for Poly(1-PHE-8), 62.25 × 10−25 cm3 for Poly(1-PHE-10), and 55.32 × 10−25 cm3 for Poly(1-PHE12) (Supporting Information). While the optical anisotropy shows a decreasing trend the with increasing diol chain length,

the existence and nature of Tll. It is also sometimes called the “rubbery−viscous transition” or “rubbery−viscous transformation.51,52 Despite the controversy, Tll is now considered a molecular level relaxation (not a thermodynamic transition) associated with thermal disruption of segment−segment contacts. These intermolecular segment−segment interactions produce cooperatively rearranging regions that “melt” at Tll52,53 and act to suppress the relaxation during deformation. Increasing temperature and decreasing deformation rates tend to eliminate this effect. Similar observations were also found by Mulligan et al. in poly(lactic acid)47 and by Kanuga et al. in poly(ethylene naphthalate)/poly(ether imide) blends.30 For different polymer systems, the Tll transition occurs at a different range with respect to Tg.54 In the case of PEU, Tll is determined at 1.063Tg (K) for Poly(1-PHE-8) and 1.065Tg (K) for Poly(1PHE-10) under the tested strain rate of 0.017 s−1 (stretching speed of 20 mm/min). The liquid−liquid transition (rubbery−viscous transition) was not observed for Poly(1-PHE-12) up to the temperature at Tg + 30 °C. Dynamic mechanical analysis also shows the absence of Tll transition in the loss modulus vs temperature curve of Poly(1-PHE-12) (Supporting Information). However, the exact reason is unclear. Upon many studies of elastomers, Boyer summarized that Tll should be weak and may not have a great influence on their physical properties since elastomers tend to be flexible hydrocarbons.48 One possible explanation is that the structure of Poly(1-PHE-12) backbone is elastomeric in nature due to the increased fraction of flexible hydrocarbons in the main chain and thus exhibits some similarities in the findings. E

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Macromolecules the strain constant at temperatures near Tg shows an increasing trend. This suggests the longer diol chain increases the flexible segments and thus results in a higher efficiency in orienting the polymer chains. With further deformation, birefringence increases continuously with increasing strain, but at a slower rate compared to the initial linear region. At high enough temperatures, chain relaxation dominates and the polymer chains exhibit extensive mobility. Therefore, the birefringence forms a plateau as the material is strained. Because of the flexible hydrocarbons in the main chain, Poly(1-PHE-12) exhibits this behavior at a lower temperature of 45 °C with respect to Tg (at Tg + 25 °C) compared to Poly(1-PHE-8) and Poly(1-PHE-10) which show similar behavior at Tg + 30 °C. FTIR Analysis. Like many polymer systems in which hydrogen bonds form between adjacent main chains,57−61 poly(ester urea)s are also strengthened by hydrogen bonds associated with the urea groups, which play a major role in reinforcing the material by linking adjacent chains. In order to probe the molecular basis associated with the deformation process in PEU, infrared spectroscopy was used as an effective way to elucidate the intermolecular interactions occurring in polymer chains. In an infrared spectrum, the bands at 3200− 3500 cm−1 which corresponds to the N−H (urea) stretching are the central focus to study the temperature effect on the hydrogen bonding in all three PEU polymers. The C−H stretching band of alkyl groups in the PEU structure is found at 2800−3000 cm−1. These bands do not change with temperature and thus is used to normalize the IR spectra in order to eliminate the thickness effects during the temperature studies. Poly(ester urea)s have the possibility of multiple hydrogen bonded structures in the urea groups. Hence, the N−H stretching region is a collective reflection of different N−H stretching modes. Figure 4 shows the scale expanded spectra in the N−H region at selected temperatures for Poly(1-PHE-8), Poly(1-PHE-10), and Poly(1-PHE-12). It is noted that with increasing temperatures the peak maximum of the N−H band shifts to higher wavenumber, the N−H band broadens, and a continuous decrease in the area under the spectrum is observed. Similar spectral changes were reported by Skrovanek and Coleman et al. in the infrared temperature studies of the polyamides and polyurethane.62−64 Figure 5 shows the peak maxima of Poly(1-PHE-8), Poly(1PHE-10), and Poly(1-PHE-12) as a function of temperature. As the temperature increases, the spectrum of Poly(1-PHE-8) first is centered at 3325 cm−1 at 30 °C but gradually shifts to a higher wavenumber of 3348 cm−1 at 180 °C. A similar trend is found in the spectrum of Poly(1-PHE-10), the peak maximum of which is centered at 3318 cm−1 at 30 °C and shifts to 3345 cm−1 with increasing temperature. A significant shifting occurs as the temperature increases from 120 to 160 °C. However, the trend for the infrared spectrum of Poly(1-PHE-12) is quite different. The peak maximum of Poly(1-PHE-12) initially occurs at a higher wavenumber of 3376 cm−1 and shifts only by a few wavenumbers to 3383 cm−1 as the temperature increases. The shifting in the wavenumber of the hydrogen bonded N−H stretching is interpreted as a change in the average strength of the hydrogen bonded N−H groups with temperature.62−64 Besides the temperature dependence, it is clearly indicated in Figure 5 that the average strength of the hydrogen bonded N− H groups of Poly(1-PHE-8) and Poly(1-PHE-10) is higher than that of Poly(1-PHE-12).

Figure 4. Selected scale expanded spectra of the N−H stretching region for Poly(1-PHE-8) (A), Poly(1-PHE-10) (B), and Poly(1PHE-12) (C) as a function of temperature from 30 to 170 °C.

Figure 5. Peak maximum of N−H stretching band for Poly(1-PHE-8) (C8), Poly(1-PHE-10) (C10), and Poly(1-PHE-12) (C12) as a function of temperature from 30 to 180 °C.

The broadening of the spectra is shown in Figure 6 by plotting the full width at half-maximum (fwhm) of each band as a function of temperature. Poly(1-PHE-8) and Poly(1-PHE-10) exhibit a broadening at the temperature range from 80 to 140 °C, while Poly(1-PHE-12) maintains the same width over the whole temperature range. The breadth of the N−H spectra mainly reflects the distribution of hydrogen bonded N−H groups at different distances and geometries.63 Skrovanek found that the completely amorphous nylon has the N−H band with F

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of the N−H stretching band, as detailed in Table 4, provides information regarding the distribution of the average strength Table 4. Normalized Area of the N−H Stretching Band normalized areaa

a

Figure 6. Full width at half-maximum (fwhm) of N−H stretching band for Poly(1-PHE-8) (C8), Poly(1-PHE-10) (C10), and Poly(1PHE-12) (C12) as a function of temperature from 30 to 180 °C.

a fwhm of 112 cm−1 at 30 °C while the fwhm of the semicrystalline nylon is of approximately 52 cm−1 at 30 °C.65 As illustrated in Figure 6, the fwhm of PEUs appears to be a larger value with increasing diol chain length at 30 °C, with Poly(1-PHE-12) showing the largest fwhm at 130 cm−1. This suggests a broad distribution of the hydrogen bonded species in PEU, which may be partly due to the presence of both urea groups and ester groups in the PEU structure. Figure 7 depicts how the area under N−H stretching bands of Poly(1-PHE-8), Poly(1-PHE-10), and Poly(1-PHE-12) changes from 30 to 180 °C. The plot of the normalized area

temp (°C)

Poly(1-PHE-8)

Poly(1-PHE-10)

Poly(1-PHE-12)

30 50 70 90 110 130 150 170

1 0.964 0.929 0.840 0.820 0.776 0.748 0.667

1 0.959 0.974 0.963 0.868 0.772 0.694 0.678

1 0.929 0.907 0.865 0.767 0.701 0.662 0.582

Area at 30 °C = 1.00.

of hydrogen bonded structures of PEU. For Poly(1-PHE-8) and Poly(1-PHE-10), the normalized area of N−H stretching region changes with temperature in three stages, while a quasilinear temperature dependence is observed in the case of Poly(1-PHE-12). This clearly indicates that the average strength of hydrogen bonding diminishes with increasing temperature. The decrease in the total area may be attributed not only to the transformation of hydrogen bonded to nonhydrogen bonded N−H groups as the temperature increases but mainly to a decrease in the absorption coefficient as the average strength of the hydrogen bonds decreases.64

Figure 7. Normalized area of N−H stretching band for Poly(1-PHE-8) (A), Poly(1-PHE-10) (B), and Poly(1-PHE-12) (C) as a function of temperature from 30 to 180 °C. G

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connection between the strength of hydrogen bonding and photoelastic constant is shown. Figure 9 provides a schematic description for the predictive structural evolution of poly(ester urea)s in the photoelastic regime during uniaxial stretching. An exponential decrease in the photoelastic constant is observed with increased average hydrogen bonding strength for all PEUs. The photoelastic constant reaches a quasi-constant value at a critical hydrogen bonding strength, beyond which a further increase in the hydrogen bonding strength has little effect on photoelasticity. It is also noticed that this critical hydrogen bonding strength increases with increasing diol chain length in the PEUs. The distance between a donor and an acceptor is one of the measurements of the strength of the hydrogen bonds: the shorter the distance, the stronger is the hydrogen bond.67 At lower average hydrogen bonding strengths (higher temperatures), chains are further apart, and progressive thermal disruption of segment−segment interactions reduces the number of cooperatively rearranging regions, allowing for chain relaxation and higher photoelastic constants. At the liquid−liquid transition (rubbery−viscous transition) limit, cooperatively rearranging regions completely “melt”, as hydrogen boding interactions become weak, and photoelasticity is not observed. As photoelasticity is a relaxation process and not a thermodynamic transition, for PEUs with longer diol chain lengths a stronger hydrogen bonding is required to achieve the same photoelastic constant as compared to the ones with shorter alkyl chains. The role of hydrogen bonding and the cooperatively rearranging regions at the temperatures near Tg further explain the thermal shape memory behavior of this polymer system, which requires the activation of the chain mobility above Tg and the presence of intermolecular hydrogen bonds for shape recovery.68 A more detailed study exploring the shape memory behavior of the amino acid-based PEU system with diols of different lengths is under preparation.

Many studies solely attribute the change in the total area of the N−H stretching region to the transformation of a fraction of hydrogen bonded N−H groups to “free” groups.59,66 Unlike their previous colleagues, Skrovanek and Coleman et al. found that the change in the total N−H area is mainly a reflection of the difference in the average absorption coefficients for the hydrogen bonded N−H stretching vibration.62 They found a strong dependence of the absorptivity coefficient with hydrogen bond strength and attributed the decrease in the total area of N−H stretching with increasing temperature to the change in the absorption coefficient as the average strength of the hydrogen bonds diminished.62 H-Bonding Effect on Materials’ Photoelasticity. Figure 8 shows the correlation between the photoelastic constant and

Figure 8. Photoelastic constant as a function of normalized area in N− H stretching band for Poly(1-PHE-8) (C8), Poly(1-PHE-10) (C10), and Poly(1-PHE-12) (C12).

the normalized area of the N−H stretching band, giving the molecular insights into the influence of hydrogen bonding on the stress−optical behavior of poly(ester urea)s. As the average strength of hydrogen bonded N−H groups increases, a continuous decrease is observed in the photoelastic constant. To the best of our knowledge, this is the first time such a

Figure 9. Schematic of the structural evolution of poly(ester urea)s in the photoelatic regime during uniaxial stretching. H

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4. CONCLUSIONS The mechano-optical behavior of PEU with different diol chain length was studied. At low stretching temperatures, the stress− optical behavior was characterized by an initial photoelastic regime contributed by the glassy behavior caused by rigid chain correlations of the heterogeneous structure. The liquid−liquid (Tll) transition (rubbery−viscous transition) for Poly(1-PHE8) and Poly(1-PHE-10) was determined at the temperature where the photoelastic behavior disappeared and PEU transformed from “a liquid of fixed structure” to “a true liquid”. The presence of intermolecular hydrogen bonds enhances the segmental correlations between adjacent PEU chains, which is consistent with the initial photoelastic behavior for PEU under deformation. The average strength and the distribution of the hydrogen bonded structures in PEU exhibit strong temperature dependence. PEUs with the longest diol chain length exhibits a linear temperature dependence for the area associated with N−H stretching region, whereas PEUs with shorter diol chain length show a three-stage development with temperature. It was also found that the photoelastic constant decreases as the average strength of hydrogen bonding increases. The real-time mechano-optical results provide a deep and systematic understanding of the stress−optical relationship and structure development of poly(ester urea)s at the molecular level. Coupled with IR results, these findings expand fundamental knowledge about the effects of hydrogen bonding interaction on the mechanical and physical properties of polymers.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b02415. Synthesis and characterization details of all compounds, the descriptive figures for the slopes of different regimes in the mechano-optical curves, dynamic mechanical analysis, and theoretical calculation of the optical anisotropy of L-phenylalanine-based poly(ester urea)s (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; Ph 765-494-1539 (M.C.). ORCID

Matthew L. Becker: 0000-0003-4089-6916 Miko Cakmak: 0000-0001-9128-7833 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Ohio Department of Development’s Innovation Platform Program and The National Science Foundation (DMR-1507420). The authors are grateful to Dr. Greg Peterson for critical review of this manuscript.

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REFERENCES

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DOI: 10.1021/acs.macromol.6b02415 Macromolecules XXXX, XXX, XXX−XXX