Nonlinear Mechano-Optical Behavior and Strain-Induced Structural

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

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Nonlinear Mechano-Optical Behavior and Strain-Induced Structural Changes of L-Valine-Based Poly(ester urea)s Keke Chen,† Nathan Z. Dreger,† Fang Peng,† Bryan D. Vogt,† 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 Macromolecules Downloaded from pubs.acs.org by UNIV OF NEW ENGLAND on 10/04/18. For personal use only.

†

S Supporting Information *

ABSTRACT: The uniaxial mechano-optical behavior of a series of amorphous L-valine-based poly(ester urea) (VAL-PEU) with varying diol lengths was studied to elucidate the molecular mechanism associated with their thermal shape memory properties. A custom, real-time measurement system was used to capture the true stress, true strain, and birefringence during the temporary shape programming at stretching temperatures above the glass transition temperature (Tg). The mechanooptical response of VAL-PEUs exhibits an initial photoelastic behavior related to enhanced segmental correlation at low temperatures above the Tg. A characteristic temperature, defined as the liquid−liquid (Tll) transition (rubbery−viscous transition), was found at about 1.05 Tg (K) (at Tg + 15 °C) at strain rate of 0.017 s−1, above which the segmental contacts largely “melt” and the initial slope of the stress-optical curves becomes temperature independent. This temperature corresponds to the temperature where mean relaxation time for the polymer is maximized. Real-time infrared spectroscopy (IR) and in situ wideangle X-ray scattering (WAXS) revealed a strain-induced intersegmental structural change during stretching. The intermolecular hydrogen bonding between the urea−urea and/or urea−carbonyl groups was found to adopt different bonding modes at the onset of strain hardening with a concurrent increased separation distance between adjacent segments. The hydrogen bonding strengthened supramolecular packing. The compromised shape recovery is a result of the disruption of the rigid segmental correlation, induced either by large extension at T < Tll or by progressive thermal effects at T > Tll, both of which may be minimized at Tll.

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INTRODUCTION Shape memory polymers (SMPs) have increasingly found practical applications in medicine, including self-deployable medical devices,1,2 expandable implants,3,4 self-tightening sutures,5 and on-demand drug delivery systems.6 Design considerations of SMPs in biomedical applications extend beyond their shape memory properties but include tunable mechanical properties,7 controllable actuation temperature,8 suitable biodegradation profile, etc.2,9 α-Amino acid-based poly(ester urea)s (PEUs) have demonstrated tunable mechanical properties ranging from hard (bone)10 to soft (vascular)11 tissues, controllable degradation rate by changing the diol length,10,12 and nontoxic degradation byproducts that lead to limited inflammatory response in vitro and in vivo.13,14 In addition, these materials possess functionalizable handles to attach functionalities such as osteogenic growth peptide,15 adhesive catechol groups,16 and radiopaque iodines.17 More recently, PEUs have been demonstrated to exhibit thermal shape memory properties that arise from a broad glass © XXXX American Chemical Society

transition range and a strong physical cross-linking (hydrogen bonding).18,19 The previously demonstrated versatility makes PEUs promising candidates for SMP-based biomedical applications. A typical dual-shape thermal SMP requires a thermally reversible phase transition to program a temporary shape and cross-linked net points to maintain the permanent shape.5,20 For polymers that use glass transition (Tg) for shape memory, increasing the temperature above Tg provides polymer chain mobility in the rubbery state to enable deformation of the polymer chains to the temporary shape, which is kinetically trapped upon quenching below Tg. Shape recovery upon heating above Tg is driven by the release of this kinetically trapped chain conformation. Significant efforts have been focused on modifying the chemical composition of SMPs to Received: June 3, 2018 Revised: September 21, 2018

A

DOI: 10.1021/acs.macromol.8b01176 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

Scheme 1. A Two-Step General Synthetic Route of L-Valine-Based Poly(ester urea)s with Different Diol Chain Lengths (n = 4, Poly(1-VAL-8); n = 5, Poly(1-VAL-10); n = 6, Poly(1-VAL-12))

tailor their reversible thermal phase transition.21−24 However, the shape memory performance of SMPs depend on both the chemical composition and the physical programming history.25,26 To understand the effects of these programming history, generalized linear viscoelastic models have been employed,27−30 including a simplified Maxwell−Weichert model with three parallel units to model the stress relaxation behavior31 and a multibranch model to describe the energy storage and release of multi-SMP.32 These theoretical approaches to predict shape memory behavior provides potential pathways for designing SMPs, but these models depend on assumptions about structural evolution of polymer chains during the temporary shape programming. Knowledge from experimental studies about the underlying structural evolution is limited. One way to understand the structural changes during the deformation is by studying the real-time birefringence. Birefringence is a measure of the optical anisotropy of polymers, which provides molecular level insights into the polymer chain orientation and structural evolution processes.33−37 Many studies have utilized the real-time captured birefringence along with the state of stress and strain to track the structural evolution associated with the mechanistic changes during processing.38−47 While the linear stress-optical relationship (SOR) governs at small to moderate deformation in true fluids,48−50 SOR deviates from linearity beyond a critical state of deformation.44,51−55 SOR has gained industrial importance in optimizing process conditions for desired properties.56,57 The main body of the work presents a real-time mechanooptical study of a series of amorphous L-valine (VAL)-based PEUs to reveal the structural evolution during the uniaxial deformation. VAL-PEUs were chosen as they have demonstrated promising results as surgical meshes for soft tissue repairs,58,59 where their shape memory properties will further provide the potential pathways for laparoscopic deployment and facilitate minimally invasive approaches for the implantation. In such applications, the ability for a device to precisely recover from the deformed state to the original shape is of vital importance.60 A range of stretching temperatures above Tg was employed in the study to understand the temperaturedependent molecular interactions during the process and to study their effect on the recoverable strain. Real-time infrared spectroscopy (IR) and in situ wide-angle X-ray scattering (WAXS) were used to reveal the molecular rearrangement at different length scales during the deformation. The real-time studies are further coupled with off-line dynamic mechanical analysis to provide insights into the molecular mechanism associated with the incomplete strain recovery. The systematic understanding of the structural evolution provides molecular

insights for materials manipulation for improved shape memory properties. These findings also expand fundamental knowledge about the role of physical cross-links, in this case hydrogen bonding, on strengthening amorphous shape memory materials.

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EXPERIMENTAL SECTION

Materials. 1,8-Octanediol, 1,10-decanediol, 1,12-dodecanediol, triphosgene, sodium carbonate, 1,1,1-tris(hydroxymethyl)ethane, and p-toluenesulfonic acid monohydrate were purchased from Sigma-Aldrich (Milwaukee, WI). Toluene, chloroform, and N,Ndimethylformamide were purchased from Fisher Scientific (Pittsburgh, PA). L-Valine was purchased from Bachem (Torrance, CA). All solvents were reagent grade, and all chemicals were used as obtained without further purification. L-Valine-based poly(ester urea)s (PEUs) were synthesized using a two-step process, and the synthesis details were reported elsewhere.59 In brief, an esterification was first performed between a L-valine and a diol (1,8-octanediol, 1,10-decanediol, and 1,12-dodecanediol) under acidic conditions (in the presence of p-toluenesulfonic acid (pTsOH) with toluene). Next, an interfacial polymerization of the respective monomers with triphosgene yielded the poly(ester urea) homopolymers (Scheme 1). Table 1 summarizes the number-average

Table 1. Characteristics of the L-Valine-Based Poly(ester urea)s

a

samplesa

Mn

Mw

ĐM

Tg (°C)

poly(1-VAL-8) poly(1-VAL-10) poly(1-VAL-12)

42000 46000 51000

71000 71000 75000

1.7 1.6 1.4

48 35 27

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

molecular mass (Mn), the weight-average molecular 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 compressionmolded using a vacuum compression press (TMP Technical Machine Products Corp.). Prior to compression, PEU polymers were dried under vacuum for 24 h at 50 °C. The compression molding temperatures were set at 132 °C for poly(1-VAL-8), 122 °C for poly(1-VAL-10), and 116 °C for poly(1-VAL-12). (The set temperature was calculated based on the equation of T (K) = 1.25Tg + 10.) Pressures of 35, 70, and 35 MPa were applied for 5 min in sequence. After that, the system was cooled with a pressure of 6.9 MPa to prevent wrinkle formation on the film surface. 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 Q-200) was used to measure the glass transition temperatures (Tg) of the as-synthesized 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 B


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Real-Time Infrared-Optical Behavior during Uniaxial Stretching. Poly(1-VAL-10) films of uniform thickness were prepared from solution (16 wt % solids in dimethylformamide (DMF)) by blade casting. Film samples were dried for 48 h at 50 °C to a final thickness of 11.5 μm and cut into dumbbell shape (40 mm in width and 40 mm in length between the clamps). Poly(1-VAL-10) film sample was uniaxially stretched at a constant stretching speed of 10 mm/min at Tg + 10 °C. A real-time mechano-optical measurement integrated with an ultrarapid-scan polarized Fourier-transform infrared spectrometer was used for this study. Details of the instrument are described elsewhere.64 The real-time stress, strain, and spectral birefringence were recorded and calculated in the same manner as described in the previous section. Additionally, a disk mirror interferometer was used to record the infrared (IR) absorbance at a rate of 300 spectra/s. Offline Wide-Angle X-ray Scattering (WAXS). Equatorial wideangle X-ray scans of uniaxially stretched samples were obtained at room temperature using a Bruker D8 Quest single crystal diffractometer with Cu micro source, equipped with a Photon-II detector. Samples were prepared by stretching to 100% engineering stretch ratios followed by rapid air cooling without relaxation period. In Situ WAXS during Uniaxial Deformation. Two-dimensional (2D) WAXS patterns were obtained using 11- Beamline for Complex Materials Scattering (CMS) at the National Synchrotron Light Source II (NSLS-II) in Brookhaven National Laboratory (BNL) using an Xray wavelength, λ, of 1.307 Å and a Photonic Sciences in-vacuum CCD detector. The sample-to-detector distance was 0.23 m, and the exposure time for each image was 10 s. Air and other background scattering were subtracted from the experimental data for all data analysis. A Linkam TST-350 high-temperature tensile stage modified for in situ WAXS measurement was used to perform the uniaxial deformation. PEU samples (25 × 5 × 0.5 mm3) of a gauge length of 20 mm were uniaxially stretched at a constant rate of 10 mm/min at Tg + 5 °C and Tg + 20 °C. Shape Recovery Characterization. Rectangular VAL-based PEU samples (25 × 5 × 0.5 mm3) were prepared by compression molding following the same procedures as described above. Dynamic mechanical analysis (TA, DMA-Q800) was used in controlled force mode with heating and cooling rates of 10 °C min−1. Dual-shape memory testing was completed in two sequential runs. VAL-based PEU specimens were first deformed to 100% strain at four different deformation temperatures (Td) (Tg + 5 °C, Tg + 10 °C, Tg + 20 °C, and Tg + 30 °C). The strain recovery profiles were then recorded during heating samples to the recovery temperatures (Tr) that are the same as Td and isothermally held for 60 min.

atmosphere. As glass transition temperature can be defined in different ways,61 here the onset temperature was used to determine the Tg, which is defined from the point at which the first deviation from the baseline on the low-temperature side is observed. Stress−Strain Behavior with Online Birefringence. Dumbbell-shaped PEU samples were cut from the compression-molded films. The narrowest region was 21 mm wide, and the sample was 20 mm long between the clamps. A custom-designed real-time mechanooptical measurement instrument was used for this study.41,62,63 In brief, the system simultaneously records optical retardation Γt, using a 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 as

ji W zy Dt = jjj t zzzD0 j W0 z k {

(1)

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

(2)

ij W yz elongation L = t − 1 = jjj 0 zzz − 1 j Wt z initial length L0 k { 2

true strain =

true stress =

Ft = WD t t

(3)

Ft 2

( )D Wt W0

0

(4)

where Γt, Dt, 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 100% strain in their rubbery state at temperatures above Tg as listed in Table 2. Samples were then held at

Table 2. Stretching Temperatures for the L-Valine-Based Poly(ester urea)s samples

Tg (°C)

poly(1-VAL-8)

48

poly(1-VAL-10)

35

poly(1-VAL-12)

27

stretching temp (°C) Tg Tg Tg Tg Tg Tg Tg Tg Tg Tg Tg Tg Tg Tg Tg

+ + + + + + + + + + + + + + +

5 10 15 20 30 5 10 15 20 30 5 10 15 20 30

■

53 58 63 68 78 40 45 50 55 65 32 37 42 47 57

RESULTS AND DISCUSSION True Mechanical Behavior. In one shape memory cycle, the sample is first heated above Tg and deformed by applying a tensile load, and then temperature is decreased below Tg and the load is removed. A typical shape memory cycle of poly(1VAL-8) can be found in Figure S1. The deformation process above Tg can be visualized from the true stress−strain curves as shown in Figure 1. Figure 1A depicts the representative true stress−strain curves of poly(1-VAL-12). As seen in Figure 1, PEUs deform elastically in the initial regime (true strain < ∼0.05), beyond which a plastic deformation, depicted by a slow increase in true stress with true strain, is observed. The elastic deformation becomes smaller with increasing temperature and eventually becomes indistinguishable at Tg + 15 °C. As strain is increased at lower temperatures above Tg, strain hardening occurs, indicated by a rapid rise in true stress versus true strain as shown in Figure 1B at Tg + 5 °C. While this strain hardening behavior is observed at Tg + 5 °C and Tg + 10 °C for poly(1-VAL-12), it is less obvious for poly(1-VAL-10) and not observed for poly(1-VAL8) at Tg + 10 °C within the investigated strain. Table 3

the final strain for 30 min at the same stretching temperatures, during which the isothermal stress relaxation was monitored. 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 to room temperature before being removed from the clamps. C


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Figure 1. (A) True stress−true strain curves of poly(1-VAL-12) (Tg ∼ 27 °C) uniaxially stretched at different temperatures (dashed line determines the onset of strain hardening at the point intersects the true stress−train curve, and arrow indicates the onset of strain hardening). (B) True stress−true strain curves of poly(1-VAL-8) (Tg ∼ 48 °C), poly(1-VAL-10) (Tg ∼ 35 °C), and poly(1-VAL-12) (Tg ∼ 27 °C) stretched at Tg + 5 °C and Tg + 30 °C.

primary and secondary bond lengths.44,67,68 The molecular deformation mechanisms at this level has been related to an increase in internal energy and thus leads to limited change in birefringence.67 Once the materials were stretched beyond this regime, the stress-optical curve exhibits a positive deviation. This region is identified as regime I (marked in Figure 2A), where birefringence increases faster with stress. The faster increase in birefringence implies a molecular alignment and at the same time the configurational entropy of the material is altered. At the end of regime I, the increase of birefringence with increasing stress decelerates, where the polymer chains approach their finite extensibility. As temperature increases, the initial photoelastic regime reduces to a smaller stress level and is replaced with a linear stress-optical behavior at Tg + 15 °C. All three VAL-PEUs exhibit this phenomenon (Figure 2B). Figure 2C plots the photoelastic constant and regime I slope as a function of temperature. The photoelastic constant is determined from the slope of the initial photoelastic regime at T < Tg + 15 °C, whereas the stress-optical constant (SOC) is determined from the initial linear portion of the stress-optical curve at T > Tg + 15 °C. The stress-optical constant is defined as the slope of the linear stress-optical law.69 The linear stress-optical law is obeyed when birefringence arises from molecular orientation.70 A characteristic temperature can be identified for all VALPEUs, at which the photoelastic constant and the regime I slope coincide, and the stress-optical constant becomes invariant with further increasing temperature. This temperature is defined as the “liquid−liquid transition temperature” (Tll) and has been observed in other amorphous polymer systems such as polystyrene, polycarbonate, and amorphous poly(ethylene oxide)/poly(vinyl acetate) blends.44,71−75 This transition has been attributed to the thermal disruption of intersegmental contacts, which led to cooperatively rearranging regions and formed “a liquid of fixed structure” at temperatures below Tll.75,76 The Tll transition occurs at a wide range with respect to Tg for different polymer systems.74 With the strain rate of 0.017 s−1 (stretching speed of 20 mm/min), Tll is determined to be 1.047Tg (K) for poly(1-VAL-8), 1.049Tg (K) for poly(1-VAL-10), and 1.050\Tg (K) for poly(1-VAL-12).

Table 3. True Stress and True Strain at the Onset of Strain Hardening for VAL-PEUsa samples poly(1-VAL-8) poly(1-VAL-10) poly(1-VAL-12)

stretching temp (°C) Tg Tg Tg Tg Tg

+ + + + +

5 5 10 5 10

σ (MPa)

ε

4.58 3.43 3.29 5.20 4.16

1.05 1.03 1.25 0.80 1.13

σ = true stress, ε = true strain.

a

summarizes the onset of strain hardening for VAL-PEUs, which is determined by tracing a straight line from the origin to where the point intersects with the true stress−strain curve (as shown by the dashed lines in Figure 1A). Increasing temperature shifts the onset of strain hardening to a larger strain. While the onset of strain hardening occurs at a similar strain level for poly(1-VAL-8) and poly(1-VAL-10), the onset for the PEU with the longest diol appears at a lower strain, reflecting a more mobile structure due to longer diol chain length. At the onset of strain hardening, the polymer chains begin to approach their finite extensibilities, with entanglements and the interpenetration of polymer coils leading to a steep rise in stress upon straining.65 In the case of poly(1-VAL12) as shown in Figure 1A, the slope of strain hardening decreases with increasing temperature, implying a less extended-chain configuration during deformation.66 These PEU materials exhibit “taffy-pull” behavior at higher temperature, where chain relaxation and chain slippage are less restricted and the strain hardening phenomenon is not observed within the investigated strain and strain rate. Mechano-Optical Behavior. The molecular orientation developed under applied stress can be visualized by plotting the birefringence with true stress as shown in Figure 2. Figure 2A shows the representative mechano-optical behavior for poly(1-VAL-12) stretched above Tg. At Tg + 5 °C and Tg + 10 °C, an initial photoelastic behavior is observed, indicated by limited increase in birefringence with large increase in stress. Typical for amorphous polymers, this photoelastic behavior has been associated with the intermolecular resistance to segment rotation, valence bond bending, and/or changes in D


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Figure 2. (A) Mechano-optical behavior of poly(1-VAL-12) stretched at various temperatures (the scale magnified plots are shown inside). (B) Compared curves of poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12) stretched at Tg + 5 °C and Tg + 30 °C (the scale magnified plots are shown inside). (C) Photoelastic and stress-optical coefficient as a function of temperature. (D) Stress-optical constant (■), Tg (●), and Tll (▲) as a function of diol length.

Figure 3. (A) Strain-optical behavior of poly(1-VAL-12) stretched at various temperatures (the scale magnified plots are shown in the inset). (B) Compared curves of poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12) stretched at Tg + 5 °C and Tg + 30 °C (the scale magnified plots are shown in the inset). (C) Strain-optical constant as a function of temperature for poly(1-VAL-8) (▲), poly(1-VAL-10) (■), and poly(1-VAL-12) (●).

The influence of the diol chain length in backbone on SOC, Tg, and Tll is shown in Figure 2D. The SOC is found to be 0.0016 MPa−1 for poly(1-VAL-8), 0.0017 MPa−1 for poly(1VAL-10), and 0.0018 MPa−1 for poly(1-VAL-12). These results indicate a slight increase in the intrinsic birefringence of VAL-PEUs with increasing diol length. This may be attributed to the presence of additional alkyl groups with increasing diol

in the backbone, which results in higher difference in the polarizabilities for PEUs with longer diol length. Strain-Optical Behavior. Strain-optical behavior quantifies the relationship between the macroscopic deformation (strain) with the microscopic deformation (represented by birefringence), which is essential to understand the relationship between the chain conformation developed with an external deformation in a shape memory cycle. These E


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Figure 4. (A) X-ray scattering curves and their corresponding 2D patterns for poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12) prior to stretching. (B) Stress−strain curves (solid lines) and their corresponding Bragg spacings (LVDW spacing, solid symbols) as a function of strain for poly(1-VAL-8) (▲), poly(1-VAL-10) (●), and poly(1-VAL-12) (■) stretched at Tg + 20 °C. (C) Stress−strain curves (solid lines) and their corresponding Bragg spacings (LVDW spacing, solid symbols) as a function of strain for poly(1-VAL-10) stretched at Tg + 5 °C (●) and Tg + 20 °C (■).

In Situ WAXS. In addition to birefringence measurements that represent the anisotropy of the entire chain during deformation, WAXS was used to examine the molecular rearrangement at a different length scale during the uniaxial stretching of PEUs. As can be seen in Figure 4A, the X-ray scattering patterns of pristine poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12) samples show two broad amorphous peaks, indicating the materials are amorphous. The main amorphous peak displays a peak maximum at q ∼ 1.31 Å−1 (“equivalent Bragg spacing” d ∼ 4.79 Å) and is relatively invariant with changing diol length in the backbone. This peak has been correlated to the van der Waals contacts of nearest adjacent nonbonded atoms (VDW peak) and is relatively insensitive to the polymer chain packing.78 An extra amorphous peak is observed at a lower scattering vector and shifts from q ∼ 0.56 Å−1 to q ∼ 0.45 Å−1 (from d ∼ 11.25 Å to d ∼ 13.90 Å) with increasing diol length. This peak arises from the electron density fluctuations at distance larger than that of van der Waals distances (LVDW peak) and has been attributed to the intersegmental packing of supramolecular chain and/or the separation of closest neighboring chain segments (some degree of near-range, intersegmental order).78,79 This LVDW peak represents a distribution of supramolecular chain packing distance, with the peak maximum indicating the preferred packing distance. The position of the LVDW peak has been found to increase with the size of side group, where a larger pendant group results in an increased “diameter” for a polymer molecule viewed as a “cylinder”, occupying larger space.78,79 The intrinsic molecular volume, so-called van der Waals volume (VvdW), is defined as the space occupied by a molecule, which is impenetrable to other molecules with normal thermal energies.80 With increasing diol length in the main chain, the VvdW increases by 32.84 Å3/unit (20.46 per cm3/mol) per additional two alkyl groups. (Detailed calculation of VvdW using the Bondi group contribution method is given in the Supporting Information.) Given the presence of one urea group per one PEU unit in regardless of diol length, the increased VvdW with additional diol results in a lower fraction of the urea groups for the PEUs with longer diol. Because PEUs are strengthened by the hydrogen bonding associated with the urea groups,19,77 the lower fraction of the urea group leads to a decreased hydrogen bonding density for the PEUs with longer diol. It has been found that the hydrogen bond formed between the adjacent chain segments reduces the imperfect packing and increases

temperature-dependent relationships are depicted in Figure 3. A nonlinear strain-optical behavior is observed for all VALPEUs as shown in Figure 3. This implies that the macroscopic deformation does not necessarily lead to the molecular strain exactly. Figure 3A shows the representative strain-optical curves for poly(1-VAL-12). The strain-optical curve initially shows a linear region at small deformation, followed by a nonlinear region where the development of birefringence transitions to a slower rate with increasing true strain. The linear region decreases to lower strain level with increasing temperature. Figure 3B compares the strain-optical behavior of poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12). While the materials generally show similar behavior at temperature above Tg, the birefringence of poly(1-VAL-12) increases more efficiently with true strain. From the slope of the linear portion (true strain < ∼0.05) of the strain-optical curve, a strain-optical constant can be determined and is plotted in Figure 3C. As temperature increases, the strain-optical constant decreases. This implies that the chain relaxation starts to dominate, and the chain orientation is suppressed with increasing temperature. Compared with PEUs with shorter diol length, poly(1-VAL12) with the longest diol shows the highest level of birefringence at the same strain level. The theoretical calculation of the optical anisotropy of poly(ester urea)s can be obtained empirically by a simple summation of bond polarizabilities and has been previously explained in detail.77 Following the same method, the intrinsic optical anisotropy for VAL-PEUs has been calculated to be 29.70 × 10−25 cm3 for poly(1-VAL-8), 36.51 × 10−25 cm3 for poly(1-VAL-10), and 66.68 × 10−25 cm3 for poly(1-VAL-12). The observed birefringence under applied sample deformation corresponds to the calculated intrinsic optical anisotropy multiplied by the chain deformation. Because the chain relaxation seems to be not significant during the sample deformation (as suggested in Figure 2), the highest level of birefringence at the same level of sample deformation can be explained by the highest intrinsic optical anisotropy of poly(1-VAL-12). In addition, compared with poly(1-VAL-10), poly(1-VAL-8) has a larger strainoptical constant at Tg + 5 °C but a slightly decreased optical anisotropy. This suggests a higher efficiency in orienting the polymer chains of poly(1-VAL-8) at Tg + 5 °C, which may be a result of the more effectively suppressed segmental motion and chain relaxation. F


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Figure 5. (A) Real-time evolution of IR absorbance for the N−H and C−H stretching for poly(1-VAL-10) band during stretching and holding. (B) True stress−strain curves () and full width at half-maximum (fwhm) of the N−H (■, || direction; □, ⊥ direction) and C−H (●, || direction; ○, ⊥ direction) stretching band as a function of true strain during stretching (a, dashed line determines the onset of strain hardening at the point intersecting the true stress−strain curve, and arrow indicates the onset of strain hardening) and as a function of time during holding (b).

packing density.81−83 Therefore, a larger intersegment separation and looser packing, and thus a larger LVDW spacing, are observed for PEUs with longer diol length. Further examination of VAL-PEUs with in situ WAXD (Figure 4B) reveals an increase in the LVDW spacing during uniaxial deformation at Tg + 20 °C. As the deformation starts, where the stress−strain curve enters the elastic region, the LVDW spacing (intersegmental distance) does not seem to change. With further stretching the sample into the plastic region, a limited increase of the LVDW spacing (intersegmental distance) follows. At the onset of strain hardening, a significant increase of the LVDW spacing (intersegmental distance) is observed. Once the materials approach their finite extensibilities, the external deformation may lead to further disintegration/weakening of segment−segment interaction, resulting in large separation between segments and thus increased intersegmental distance. The increase is more prominent for poly(1-VAL-12). At a lower stretching temperature (at Tg + 5 °C in Figure 4C), the fast increase of the LVDW spacing (intersegment packing distance) is shifted to a lower strain, where the onset of strain hardening simultaneously occurs at a lower strain level. Compared with the relatively faster increase of birefringence (chain orientation) with increasing strain, the intersegment packing distance does not experience a significant increase until the onset of strain hardening. While the polymer chain orientation is developed with a limited increase of the intersegment distance (and thus a less disruption of the segment−segment interaction) during deformation at Tg + 20 °C, it occurs together with a larger intersegmental separation at a lower strain at Tg + 5 °C. Real-Time IR Analysis. As mentioned in the previous section, the hydrogen bonding associated with the urea groups plays a crucial role in strengthening poly(ester urea)s to resist mechanistic deformation.19,77 In this regard, the real-time infrared (IR) spectroscopy was used to probe the molecular nature that is responsible for the intersegmental changes during the deformation process. Understanding the strengthening mechanism and their dynamic changes is crucial to understanding the shape imprinting of PEUs, which is associated with the hydrogen bonding structures.18,19 Figure 5A shows the time-resolved IR spectra of poly(1-VAL-10) uniaxially stretched at Tg + 10 °C. A significant broadening of the N−H stretching band at 3200−3500 cm−1 is observed during the deformation. This peak is a collective reflection of

different N−H stretching modes and is related to the hydrogen bonding nature of the materials. The breadth of the N−H spectra represents a distribution of hydrogen bonded N−H groups at different distances and geometries.84 In Figure 5B, the full width at half-maximum (fwhm) of the N−H stretching band (urea groups) and the C−H stretching band (alkyl groups) is plotted as a function of true strain and time, along with the true stress−strain curve during the stretching and holding. The C−H stretching band maintains nearly the same fwhm during the entire process. The fwhm of the N−H band does not change as the deformation starts until the onset of the strain hardening (true strain ∼0.95), where an abrupt increase of the fwhm is observed. The significant increase of the fwhm at onset of the strain hardening indicates an increasing number of hydrogen bonding species. Beyond the onset point of the strain hardening, the dislocation of the adjacent polymer chains is likely to occur, together with the larger intersegmental separation, force the hydrogen bonds to adopt different bonding configurations. In addition, the fwhm of N−H band is observed to increase in both parallel (stretching) and perpendicular (shrinking) directions during stretching. This implies that the simultaneous rearrangement of hydrogen bonding is not one directional but occurs in both directions. However, the fwhm increases more significantly in parallel than in perpendicular direction, suggesting a higher level of hydrogen bonding distortion in the stretching direction. This phenomenon is consistent with the dichroic ratio of N−H band (Figure S7), which decreases initially during stretching (true strain < ∼0.95) and starts to plateau at the onset of strain hardening (true strain > ∼0.95). Associated with backbone orientation, N−H groups first orient in the perpendicular direction at true strain < ∼0.95, and the disruption of hydrogen bonding is only observed when further orientation cannot be developed after the onset point. Additional details of dichroic ratio calculation and discussion can be found in the Supporting Information. Moreover, it is shown in Figure 5B that the fwhm of the N− H spectra remains at the increased level whereas the true stress exhibits a decrease during holding. This can be attributed by the chain relaxation and recoil that are more contributed by the C−H alkyl chains (Figure S7), and thus the broader distribution of the hydrogen bonded species (associated with N−H urea groups) is less altered. Isothermal Relaxation Behavior. A quenching step is required to complete the shape memory cycle, in which the G


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Figure 6. Isothermal stress relaxation (A), birefringence relaxation (B), and normalized stress relaxation curves (the solid red line is the leastsquares fit of eq 5) (C) for poly(1-VAL-12) stretched at various temperatures; the fitted mean relaxation times as a function of temperatures for poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12) (D).

motions against rotational energy barrier (not related to the entropy−elastic stresses).67 At low stretching temperature above Tg, the enhanced intermolecular resistance to segment rotation results in an increased energy barrier and thus a larger distortional stress. The stress relaxation spectra of polymers correspond to a broad size distribution of polymer chain lengths and the topological constraints of segments. A stretched exponential (Kohlrausch−Williams−Watts)85,86 form has been employed to model the relaxation process in complex polymer systems87−93

deformation-induced entropic change is kinetically trapped to later trigger the recovery upon heating. In this step, relaxation of this entropic energy may simultaneously occur. The isothermal relaxation behavior of the polymers can provide insights into the kinetics of this quench process. Obviously the stretching temperature affects the relaxation behavior. Figures 6A and 6B show the stress relaxation and birefringence relaxation curves of poly(1-VAL-12) after a 100% strain with a stretching speed of 20 mm/min at different temperatures above Tg. The relaxation spectra for poly(1-VAL-8) and poly(1-VAL-10) can be found in Figures S6-A and S6-B (Supporting Information). The stress relaxation spectra at all temperatures show a timedependent behavior: a fast exponential decay for a short period of time, followed by a slow decay for a few of decades of time. The birefringence relaxation curves also follow a similar twostage decrease. According to de Vries,67 the fast decrease in birefringence is related to local distortion, particularly due to side group twisting and aligning in the direction of deformation. Only the slow decay in birefringence is more associated with molecular chain orientation. For VAL-PEUs, the valine side group may contribute to the former part and the polymer backbone with long chain diol may contribute to the latter component. At Tg + 5 °C and Tg + 10 °C, a larger portion of true stress is relaxed in the initial region of fast decay. At temperatures above Tg + 15 °C, the large decrease of true stress is suppressed and gives way to a smoother transition to the slow region, as can be seen in Figure 4C where the stress relaxation curve is normalized to the maximum true stress after the step strain. As mentioned above, the region of fast relaxation comes from the distortional internal stress, which is mainly due to intermolecular energy effects that is related to segmental

β σ (t ) = e−(t / λ) σ0

(5)

where λ is the mean relaxation time and β can be related to the number of relaxation moments, n, by ⟨τ n⟩ =

0

∫∞ t n−1e−(t /λ)

β

dt

(6)

In particular, Baeurle et al. have successfully employed this form to model the stress relaxation behavior of homogeneously cross-linked thermoplastic elastomers with transient cross-links when subjected to a nonlinear extensional deformation.92 In their system, the transient cross-links form domains of different number of relaxing elements under deformation, and these domains are caused by thermal fluctuations that induce domain size fluctuations.92 As revealed by WAXS, VALPEUs exhibit some preferred intersegmental packing (nearrange order) that may be attributed the hydrogen bonding which acts as transient physical cross-links to enhance segmental correlation and produce cooperatively rearranging regions.77 The randomness of the relaxation processes H


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representative curves of poly(1-VAL-8) in Figure 7A, the total recoverable strain changes with temperature. Figure 7B plots the recovery ratio of poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12) as a function of temperature. Interestingly, the recovery ratio is found to show an opposite trend at the temperature below and above Tg + 15 °C. This temperature has been determined as the Tll transition for VAL-PEUs as well as the characteristic temperature at which the materials experience the longest relaxation time. In addition, the interchain (LVDW) spacing is found to return to the same state as that of the materials prior to deformation (Figure S8 shows the LVDW spacing of PEUs), implying the restoration of the preferred intersegmental packing (near-range, intersegmental order). However, owing to the dynamic nature, the hydrogen bonding structure is less likely to return to the same local site after the reconfiguration. In other words, the hydrogen bonds do not always re-form or not re-form at the same original positions. These microscopic changes are reflected in the macroscopic nonrecoverable strains. The structural evolution of poly(ester urea)s under uniaxial deformation is schematically illustrated in Figure 8. The preferred intersegmental packing is strengthened by the hydrogen bonding at the urea−urea and/or the urea−carbonyl site, which also gives rise to the cooperatively rearranging regions at a larger length scale. At T < Tll, the strengthened segmental contacts lead to suppressed chain mobility and enhanced strain hardening. The increase in chain orientation is accompanied by the straininduced structural changes in the intersegmental distance and hydrogen bonding. At T > Tll, the hydrogen bonding interactions become weaker,77 and the cooperatively rearranging regions are less topologically constrained. This allows chain relaxation and leads to less microscopic chain orientation with the macroscopic straining. This indicates that the thermal shape memory properties of poly(ester urea)s are possibly defined by a competition between two mechanisms: the hydrogen bonding species that undergo bonding rearrangement/exchange induced by rigid segmental contacts at T < Tll, and the diffused segmental interaction at T > Tll. The Tll transition indicates an optimal temperature range for deformation (the temperature to program a temporary shape), at which temperature the rigid segmental structures (that lead to suppressed chain mobility) transform into a mobile structure while the segment−segment

contributed by the segmental correlation will result in a number of the relaxation moments. The relaxation data were fit to the stretched exponential stress relaxation model92,93 as shown by the red solid lines in Figure 6C with the fitting parameters summarized in Table 4. Table 4. Fitting Parameters Calculated from Eq 5 for the Data in Figure 6 sample

Tg (°C)

T − Tg (°C)

poly(1-VAL-8)

53 58 63 68 78 40 50 55 65 75 32 37 42 47 57

5 10 15 20 30 5 10 15 20 30 5 10 15 20 30

poly(1-VAL-10)

poly(1-VAL-12)

λ (s)

β

× × × × × × × × × × × × × × ×

0.18 0.22 0.29 0.32 0.47 0.23 0.26 0.31 0.44 0.42 0.20 0.27 0.31 0.37 0.31

2.31 4.27 1.04 7.50 1.61 5.60 6.32 9.64 2.03 1.06 6.66 1.02 1.87 1.20 5.53

102 102 103 102 102 102 102 102 102 102 102 103 103 103 102

Figure 6D indicates that there is a maximum in the mean relaxation time λ at a characteristic temperature of Tg + 15 °C, which was also determined to be the Tll for VAL-PEUs. Tll is the temperature above which the material transforms from a liquid of “fixed structure” to a “true liquid”. The shorter relaxation time and thus fast energy dissipation at T < Tg + 15 °C may be explained by the presence of the rigid segmental correlations in the structure. Increasing temperature above Tg + 15 °C disrupts the segment−segment contacts. This progressive thermal disruption of segment−segment interactions reduces the number of cooperatively rearranging regions, allowing for chain relaxation and leading to shorter relaxation time at T > Tg + 15 °C. Recoverable Strain and Structural Interpretation. To further study the effect of deformation history in a SMP cycle and the associated molecular basis, the samples were cooled below their respective Tgs after deformation and then heated to the corresponding deformation temperature to record the macroscopic recoverable strain. As can be seen from the

Figure 7. (A) Recovery profile of poly(1-VAL-8) after stretching at different temperatures. (B) Recovery ratio of poly(1-VAL-8), poly(1-VAL-10), and poly(1-VAL-12) after 100% strain as a function of temperature. I


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Figure 8. Schematics of the structural evolution of poly(ester urea)s during uniaxial stretching at T < Tll and T > Tll.

subject to deformation and provide insights into the kinetics of the Tg-based shape memory polymers.

interactions can still give rise to cooperatively rearranging regions. In addition, the onset of strain hardening indicates the extension limit that allows the macroscopic deformation with less disruption of the intersegmental structure. Beyond the onset of strain hardening, the larger interchain separation forces the bond exchange in the hydrogen bonding structure. The reconfigured bonding structure results in the diminished strain recovery.

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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.8b01176. Synthesis and characterization details of all compounds; the stress−strain, stress-optical, strain-optical, and isothermal relaxation curves of poly(1-VAL-8) and poly(1-VAL-10); X-ray scattering curves for poly(1VAL-8) uniaxial stretched at Tg + 5 °C, poly(1-VAL-10) at Tg + 5 °C and at Tg + 20 °C, and poly(1-VAL-12) at Tg + 5 °C; X-ray scattering curves and equivalent Bragg spacings for poly(1-VAL-8), poly(1-VAL-10), and poly(1- VAL-12) before, after stretching, and after recovery; dichroic ratio of N−H and C−H stretching bands for poly(1-VAL-10); Bondi group contribution method to calculate the van der Waals volume of L-valine-based poly(ester urea)s (DOCX)

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CONCLUSION The uniaxial mechano-optical behavior of a series of amorphous L-valine-based poly(ester urea) (VAL-PEU) with varying diol lengths was studied. At low temperatures above Tg, the stress-optical behavior was characterized by an initial photoelastic regime caused by the glassy behavior associated with the rigid chain correlations of the heterogeneous structure. The liquid−liquid (Tll) transition (rubbery−viscous transition) was determined at about 1.05 Tg (K) (at Tg + 15 °C), where the initial glassy behavior gives way to a linear stress-optical curve. The isothermal relaxation behavior of VAL-PEUs exhibits a maximum value of the mean relaxation time at Tg + 15 °C. Tg + 15 °C is also found to show the maximum strain recovery upon heating from the deformed state. The interchain packing distance displayed an increase at the onset of strain hardening, which was enhanced at T < Tg + 15 °C. Concurrently, the intermolecular hydrogen bonds, which enhances the segmental correlations in the VAL-PEU structure, was found to adopt different bonding modes at the onset of strain hardening. Once disrupted, the reconfigured bonding structure results in an incomplete strain recovery upon heating at T < Tg + 15 °C. The real-time mechanooptical results provide a molecular-level understanding of the stress-optical relationship and structure development of poly(ester urea)s during deformation. Together with the in situ WAXD and IR results, the studies further decouple the strain and structural changes at different length scales. The findings further expand fundamental knowledge about the effects of physical interaction (i.e., hydrogen bonding) on the mechanical and physical properties of amorphous polymers

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AUTHOR INFORMATION

Corresponding Authors

*(M.C.) E-mail [email protected]; Tel 765-494-1539. *(M.L.B.) E-mail [email protected]; Tel 330-972-2834. ORCID

Bryan D. Vogt: 0000-0003-1916-7145 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. The authors are grateful to Dr. S. Shams Es-haghi for critical review of the use of stress relaxation fitting model. The authors thank Dr. Masafumi Fukuto and Ruipeng Li for their help in performing J


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the in situ wide-angle X-ray scattering experiments at Brookhaven National Laboratory. This research used resources of the Complex Materials Scattering (CMS/11-BM) beamline, operated by the National Synchrotron Light Source II, which are U.S. DOE Office of Science Facilities, at Brookhaven National Laboratory under Contract DE-SC0012704. The authors also thank Dr. Matthias Zeller and Megan Forshey for their help with the wide-angle X-ray scattering experiments in the X-ray Diffraction Laboratory in Purdue University.

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