A Multiscale Investigation on the Mechanism of Shape Recovery for

Article pubs.acs.org/Macromolecules

A Multiscale Investigation on the Mechanism of Shape Recovery for IPDI to PPDI Hard Segment Substitution in Polyurethane Wei Xu,† Ruoyu Zhang,*,† Wei Liu,† Jin Zhu,† Xia Dong,‡ Hongxia Guo,‡ and Guo-Hua Hu§ †

Ningbo Key Laboratory of Polymer Materials, Ningbo Institute of Material Technology and Engineering, Chinese Academy of Sciences, Zhongguan West Road 1219, Ningbo 315201, People’s Republic of China ‡ Beijing National Laboratory for Molecular Sciences, Institute of Chemistry, Chinese Academy of Sciences, Beijing 100190, People’s Republic of China § Laboratory of Reactions and Process Engineering, CNRS-University of Lorraine, Nancy 54001, France S Supporting Information *

ABSTRACT: Shape memory thermoplastic polyurethane (SMTPU) containing isophorone diisocyanate (IPDI) in hard segment has excellent shape recoverability even after large strain deformation. However, the underlying mechanism of shape recovery remains unclear. In this study, 1,4-phenylene diisocyanate (PPDI) in the polyurethane is gradually substituted by IPDI, and multiscale effects are examined by normal and dichroic Fourier transform infrared spectroscopy (FTIR), smallangle X-ray scattering (SAXS), single-molecule force spectroscopy (SMFS), and mechanical test. Contradictory to the traditional conclusion, the degree of microphase separation decreases as the content of IPDI increases, while the macroscopic shape recoverability is largely improved. With dichroic FTIR and SAXS, we find that the morphology of hard phases changed from lamellar-like to fibrillar-like, which is more stable under stretching. SMFS experiments discover that IPDI could increase the elasticity of polymer chain and could endow the hard phases with “elastic” under stress. With these two factors, we are able to explain the high recoverability of the SMTPU containing IPDI.

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INTRODUCTION Thermoplastic polyurethane (TPU) elastomer is a type of thermoplastic elastomer that has many advantageous properties including good processability, excellent elasticity, and appreciable biocompatibility, thus affording it a wide range of applications, such as Spandex fiber, automotive bumpers, cable jackets, biomedical products, etc.1−7 TPU is composed of soft segments, usually polyether or polyester, and hard segments, which typically consist of diisocyanate and chain extender. This segmented structure is the key feature of TPU molecular chain. Because of the thermodynamic incompatibility, the hard segments phase-separate into hard domains, while soft segments aggregate into deformable regions. As a result, the existence of multiple phases creates a microphase separation morphology which permits TPU’s elastomeric performance. The hard domains could act as the physical linkages and are the frame of the material’s structure, while the soft domains provide elasticity and changeability. This sort of structure is required in recently developed shape memory materials, thus making TPU a suitable candidate for such use.8−15 One of the most crucial parameters for shape memory materials is recoverability, which is usually not very high for thermoplastic materials, especially after large strain stretching. The physical linkages could be easily destroyed by stress; thus, polymer chains would have directional reptation during © XXXX American Chemical Society

elongation. Thermoset materials, on the other hand, could avoid such problem by forming cross-linked chemical bonds. Researchers have investigated several methods to address this deficiency.16−21 For example, Yang et al.20 studied the effect of the regularity of hard segment structure on shape memory properties by comparing two series of thermoplastic polyurethanes synthesized from 1,4-phenylene diisocyanate (PPDI, planar shape) and diphenylmethane diisocyanate (MDI, bent shape). The results indicated that PPDI-based polyurethane exhibited better shape recoverability than its MDI-based counterpart by ∼10% in the entire range of hard segments content. They believed that the planar structure of PPDI-based hard segments allowed for greater stability of physical crosslinking points under external force in comparison to a bent hard segment structure. Additionally, Saralegi et al.21 investigated the effect of the addition of cellulose nanocrystals on thermoplastic polyurethane shape memory properties. Results showed that the crystallinity of the hard phase was improved by adding cellulose nanocrystals, thus leading to the enhancement of recoverability. Yet despite this, the improvement on the strain recoverability is still limited, usually at ∼85% recover ratio under strain of ∼300%. Received: June 2, 2016 Revised: July 17, 2016

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DOI: 10.1021/acs.macromol.6b01172 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Scheme 1. Multiscale Behavior of TPU Containing Nonplanar Structure during Stretching

In this paper, as shown in Scheme 1, we will combine various techniques, from nanoscale to macroscale, to investigate the effect of the nonplanar structure on multiscale properties and study the structural mechanism of the largely improved recovery ratio. Here, SAXS is used to study the phase structure, dichroic FTIR is adopted to investigate the orientation of hydrogen bonding, and the technique of SMFS is applied to study the single chain mechanical property.

In contrary, and surprisingly, we have found that some nonplanar rings could tremendously improve the strain recoverability to ∼98% or even higher under strain of 500− 800%.22,23 The common feature of these structure is that they have at least two conformations and are capable of transforming from one to another under stress. Similar behavior has been observed in polysaccharides by Marszalek and Li.24,25 We also noticed that the addition of nonplanar structures decreased the degree of hydrogen-bonding association, which coincides with the findings of other researchers.26,27 Therefore, the conventional conclusion that the stereocomplex structure could interfere with the formation of the hydrogen bonding and consequently reduce strain recoverability cannot be applied here. Nevertheless, the underlying mechanism of shape recovery remains unclear to this point. In order to understand such an extraordinary phenomenon, we need to design specific experiments while combining multiscale techniques to find out the underlying mechanism. The structure of polyurethane and its dynamics within are complex, and many research groups have devoted significant effort to studying potential driving factors. These studies often employ two techniques: small-angle X-ray scattering (SAXS) and Fourier transform infrared spectroscopy (FTIR). For instance, Hsiao et al.28,29 used the former to study the phase behavior of microphases and carefully examine structural models. At the same time, Hsu et al.30−34 used time-resolved FTIR to study the phase separation kinetics and variation of hydrogen bonding during microphase separation. These two techniques cover the length scale of 1−10 nm, a range with particular importance for the examination of TPU materials. On the other hand, we introduced the usage of single-molecule force spectroscopy (SMFS) to investigate the difference between TPU chains with and without nonplanar rings.35−39 It showed that with the replacement of aromatic ring by nonplanar rings the elasticity of single polymer chain varies significantly. However, little research has been done using these techniques in conjunction to understand the multiscale structural features during the stretching and recovery processes of TPU.

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EXPERIMENTAL SECTION

Materials. Poly(tetramethylene glycol) (PTMG, Mn = 2000 g/ mol) and isophorone diisocyanate (IPDI) were purchased from Aladdin (Shanghai, China). Hydroquinone bis(2-hydroxyethyl) ether (HQEE) and 1,4-phenylene diisocyanate (PPDI) with 99.5% purity were kindly supplied by Yantai Yusheng Chemical Company (Yantai, China) and Hangzhou Yilian Chemical Company (Hangzhou, China) separately. All the other reagents and solvents of analytical grade were purchased from Aladdin (Shanghai, China) and used as received except dimethylformamide (DMF). DMF was dried with calcium hydride and vacuum distilled before use. Polymer Synthesis. The prepolymer method was adopted to synthesize TPU samples. In the prepolymer preparation step, poly(tetramethylene glycol) (PTMG) with Mn of 2000 g/mol was dried in a vacuum oven to remove residual water. Then, in a 100 mL three-necked round-bottomed flask with nitrogen atmosphere, 4 mmol of PTMG2000 was adequately mixed with 16 mmol of diisocyanates either isophorone diisocyanate (IPDI), 1,4-phenylene diisocyanate (PPDI), or a mixture of the two. The NCO/OH ratio was 4. Next, the mixture was heated for about 3 h at 45 °C to prepare the prepolymer with isocyanate group end-capped. After adding 20 g of dimethylformamide (DMF) as solvent, the NCO content in prepolymer was determined by titration with di(n-butyl)amine. The NCO content in the solution was generally 0.6−0.7 mmol/g. In the chain extension step, equivalent hydroquinone bis(2-hydroxyethyl) ether (HQEE) was added to react with residual NCO groups. The chain extension reaction was performed at 65 °C with catalyst, dibutyltin dilaurate (DBTDL), for several hours. The disappearance of NCO peak, which located at 2270 cm−1 on the Fourier transform infrared spectrum (FTIR), marked the termination of chain extension. Film Preparation. To prepare uniform polyurethane films, bubbles in polymer solution were pumped out in a vacuum oven before being poured into a smooth Teflon mold. The Teflon mold was flat rectangle shape with dimensions of 160 × 120 × 2 mm (length × width × B


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Macromolecules thickness). The mold was placed in an oven maintained at 60 °C to evaporate off the solvent and then placed in a vacuum oven at 60 °C for complete drying. For mechanical test and small-angle X-ray scattering experiment, the thickness of the polyurethane films was about 0.5 mm. For in-situ infrared dichroic measurements, the film thickness was 10−20 μm. GPC. The weight-average molecular weight and polydispersity index of TPU samples were measured by a HLC-8320 gel permeation chromatograph. The standard sample was monodisperse polystyrene, and the mobile phase was DMF with a flow rate of 1 mL/min. SAXS. Small-angle X-ray scattering (SAXS) experimental data were collected at the beamline BL16B1 in Shanghai Synchrotron Radiation Facility (SSRF). The measurements were carried out at room temperature (about 25 °C). We use an X-ray wavelength of 1.24 Å and a sample-to-detector distance of 2030 mm. A Linkam TST350 tensile hot stage was used to stretch sample films from the initial length, L0 = 4 mm, with a speed of 4 mm/min. The detection spot on the sample remained fixed, as the sample film was stretched on both sides simultaneously. The recovery of specimens was also performed at a constant rate of 4 mm/min. Images were recorded by a MAR165CCD detector with a resolution of 2048 × 2048 pixels (pixel size = 79 × 79 μm2). The acquisition time for imaging was approximately 10 s. Raw scattering patterns were corrected for background noise through background subtraction. The scattering curves were presented as intensity (I) versus scattering vector (q = 4π sin θ/λ). The interdomain spacing (D) was calculated with Bragg’s equation:40 D=

2π qmax

f = − 2.3

R0 + 2 R − 1 R0 − 1 R + 2

(3)

The value of f has been found to be dependent on the orientation of the molecular chain axis with respect to the deformation direction. Consequently, the f values of 1, −0.5, and 0 correspond to a perfectly parallel, perfectly perpendicular, and random orientation, respectively.42 As the existence of hydrogen bonding, the peaks of the hydrogen-bonded carbonyl group (1700 cm−1) and free carbonyl group (1730 cm−1) are overlapped in the carbonyl group absorption band (1670−1760 cm−1).43 The FTIR spectrum between 1670 and 1760 cm−1 was deliberately resolved into two independent spectral peaks, which correspond to a hydrogen-bonded carbonyl group and a free carbonyl group separately. In this way, the orientation function could be calculated from the separated curve of the hydrogen-bonded carbonyl group. Mechanical Measurements. Mechanical measurements were performed on an Instron 5567 apparatus at room temperature. The shape recoverability after tensile break was measured at a constant tensile rate of 100 mm/min by recording the recovery ratio at 3 min after breakage. Dumbbell-shaped splines with dimensions of 35 mm in length, 2 mm in width, and ∼0.5 mm in thickness were used. SMFS. Single molecule force spectroscopy (SMFS) was performed on an Agilent 5500 controlled atmosphere scanning probe microscopy. Samples were first dissolved in DMF and then diluted to 10−4 mg/mL. Neat silicon wafers, which had been treated with hot piranha solution (70:30 v/v 98% H2SO4/30% H2O2), were immersed in the diluted solution overnight. Molecular chains of the sample were adsorpted onto the surface of silicon wafers through physical absorption. The silicon wafers were then flushed smoothly by the solvent to wash away unstable molecules before use, after which the SMFS experiment was conducted at room temperature. In the test, the silicon wafers were placed on the sample table and then scanned by the force curve mode of the Agilent 5500 controlled atmosphere scanning probe microscopy. The tensile speed of the cantilever was 1.0 μm/s.

(1)

Here, qmax corresponds to the peak position in the scattering curves. The equatorial direction intensities and the meridional direction intensities were obtained by cake integration in the azimuthal angular range of −3° ≤ Φ ≤ 3° and 87° ≤ Φ ≤ 93°, respectively. The equatorial direction was parallel to the deformation direction, while the meridional direction was perpendicular to it. XRD. X-ray diffraction (XRD) data were recorded by a Bruker AXS D8 Advance with λ = 1.541 Å, 2.2 kW. All peak areas were in the 2θ range from 5° to 60°. FTIR. Fourier transform infrared (FTIR) characterization was performed on a Thermo Nicolet 6700 Fourier transform infrared spectrometer from Thermo-Fisher Scientific, scanning from 500 to 4000 cm−1. The sample was prepared by casting the polyurethane solution on a KBr disk. Infrared Dichroic Measurements. In-situ infrared dichroic measurements of the PU films were performed using a Thermo Nicolet 6700 Fourier transform infrared spectrometer (Thermo-Fisher Scientific) and a Linkam TST350 tensile hot stage (Linkam Scientific). The stretching and recovery process were both carried out at 25 °C. Specimens with a thickness of 10−20 μm were uniaxially stretched from an initial length of 4 mm to 500% strain, at a rate of 4 mm/min. Specimen recovery was also performed at a rate of 4 mm/min until the applied stress reached zero. Infrared spectra were collected incrementally at every 20% increase in strain with polarized IR radiation parallel and perpendicular to the deformation direction, respectively. The orientation function ( f) of an absorption band was calculated as follows:41

f=

R−1 R+2

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RESULTS AND DISCUSSION Sample Preparation. Polyurethane samples are synthesized by the prepolymer method, as displayed in Scheme 2. The

Scheme 2. Synthesis of Polyurethane Samples

(2)

Here, the dichroic ratio is R = A∥/A⊥, where A∥ and A⊥ represent the absorbance area of the target band measured with IR radiation polarized parallel and perpendicular, respectively, to the deformation direction. In addition, R0 = 2 cot2 α, where α is the angle between the molecular chain axis and the transition moment associated with the absorption band. For carbonyl group on the molecular chain, α is generally cited as 78°.42 Thus, eq 2 could be simplified as C


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Macromolecules Table 1. Composition of Polyurethane Samplesa composition (relative mole ratio) IPDI-0 IPDI-20 IPDI-40 IPDI-60 IPDI-80 IPDI-100

IPDI

PPDI

PTMG

HQEE

0 0.8 1.6 2.4 3.2 4

4 3.2 2.4 1.6 0.8 0

1 1 1 1 1 1

3 3 3 3 3 3

HS (wt %) 36.6 38.8 37.7 40.1 40.2 42.7

(38.2) (39.1) (40.0) (40.9) (41.8) (42.6)

nonplanar structure content in HS(mol %) 0.0 13.9 25.0 38.2 45.4 60.0

(0.0) (11.4) (22.9) (34.3) (45.7) (57.1)

Mw (×104)

PDI

8.7 9.4 22 26 11.3 21

3.8 3.5 3.8 3.6 3.3 4

a

The HS (wt %) means the weight percent of hard segment, which is defined as the percent by weight of diisocyanates and chain extender in the polyurethane. Nonplanar structure content in HS (mol %) is calculated as the mole percent of IPDI in the diisocyanates and chain extender. The value in the parentheses is the theoretical value; the outside one is the practical value, calculated by 1H NMR spectroscopy (Supporting Information Figure S1).

Figure 1. 2D-SAXS patterns of all samples at 0−500% strain and recovery.

soft segment is polyether polyol PTMG, with Mn = 2000 g/ mol, while the hard segment is composed of IPDI, PPDI, and HQEE, as the chain extender. Data concerning composition of raw materials, weight-average molecular weight (Mw), polydispersity index (PDI), hard segments content, and nonplanar structure content in hard segments are given in Table 1. Polyurethane samples are named as IPDI-0, IPDI-20, IPDI-40, IPDI-60, IPDI-80, and IPDI-100 according to the varying molar ratio of IPDI to total diisocyanates, 0%, 20%, 40%, 60%, 80%, and 100%, respectively. SAXS Analysis of Phase Domains during Stretching and Recovery. The stability of hard domains under 500%

strain was studied by small-angle X-ray scattering (SAXS) experiment. The evolution of 2D-SAXS patterns, including the stretching and recovery processes, are displayed in Figures 1 and 2, respectively. The stretching direction is consistent with the equatorial direction in the 2D-SAXS patterns. In the unstretched state (0% strain), depending on the IPDI content, a broad or narrow circular scattering halo can be seen in the first column of Figure 1. This circular scattering ring is the typical feature of microphase separation between soft and hard domains and indicates that the microphase is randomly oriented.44 D


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Figure 2. 2D-SAXS patterns of all samples from 500% strain to zero stress.

strain, it can be surmised that the correlation between hard domains in the direction perpendicular to stretching becomes weaker, though it still exists. As deformation continues, the two arcs are gradually squished. No streak is observed on the meridian even at larger strain, as was the case with IPDI-0, indicating no stress-induced needle-like crystals. The hard domains of IPDI-100 sample are thus more likely to undergo a deformation process similar to the affine deformation.47 Inferred from the rhombus-shaped pattern and very weak scattering intensity, the hard domains of IPDI-0 are largely destroyed after the recovery from 500%. However, for the IPDI-100 sample, the recovery state pattern is almost identical to the unstretched state patterns, which indicates that the hard phases mostly recover to their original position. On the basis of the above comparisons between the IPDI-0 and IPDI-100 samples, we can deduce that these two polyurethanes, containing either PPDI-HQEE or IPDI-HQEE as hard segments, undergo completely different hard domain deformation processes. The evolution process of IPDI-20 is similar to IPDI-0 based on the 2D-SAXS patterns. Nevertheless, some differences are still perceivable. At 200%−300% strain, the scattering pattern shows an appearance of joints on the meridian, signifying the correlation between hard domains perpendicular to the stretching direction. Moreover, the strong streak appears at larger strain in comparison to IPDI-0. Finally, after recovery,

IPDI-0 sample shows an H-shaped pattern with rather weak scattering intensity on the meridian at 200% strain. It indicates that hard domains tilt along the stretching direction under small strain and gradually become uncorrelated in the direction perpendicular to stretching direction. When strain reaches 300%, the H-shaped pattern evolves into a pattern of two lobes on the equator and a strong streak on the meridian. The appearance of the latter is probably caused by the formation of needle-like crystals oriented along the stretching direction, supported by the studies conducted by other researchers using AFM and WAXS.45,46 The lobe patterns indicate that stacks of hard domains are oriented along the stretching direction. With strain increasing from 300% to 500%, the 2D-SAXS pattern of IPDI-0 shows a similar shape only with the reduction of both lobes’ size and the scattering intensity, which indicates the permanent breakup happened in hard domains under large strain stretching. For the IPDI-100 sample, the evolution process of 2D-SAXS patterns is significantly different from that of IPDI-0. At the beginning of deformation, the circular scattering halo becomes elliptical, indicating a slight change of interdomain spacing that increases in the stretching direction and decreases in the direction perpendicular to stretching. At 200% strain, the pattern evolves into a shape with two arcs on the equator and edge joints on the meridian. In comparison to IPDI-0 which shows little scattering intensity on the meridian under the same E


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Figure 3. 1D-SAXS intensity profiles in the parallel direction (||) and the perpendicular direction (⊥), including two different states (before stretching and after release): (a) IPDI-0, (b) IPDI-20, (c) IPDI-40, (d) IPDI-60, (e) IPDI-80, and (f) IPDI-100. The parallel and perpendicular directions are the directions parallel or perpendicular to stretching, respectively.

direction after release. These phenomena indicate that the hard domains are seriously damaged, especially in the parallel direction. For the IPDI-20 sample, the intensity difference between each state is also large, although the q peak can be determined in both directions. The q peak in the parallel direction clearly shifts toward lower values of q, indicating a huge increase of interdomain spacing in this direction. On the other hand, the q peak in the perpendicular direction shifts moderately toward higher values of q, which is to suggest a reduction in interdomain spacing in the perpendicular direction. For the IPDI-40 sample, the intensity difference and q-peak shifting are minute in comparison to those of IPDI0 and IPDI-20an observation that suggests better elasticity and restoration capability. Notably, for IPDI-60, -80, and -100, the two curves overlap well in both directions, with only a tiny q peak shifting. Thus, to summarize, with increasing IPDI content, the stability of hard domains gradually improves.

the state of IPDI-20 is also completely different from the unstretched pattern, meaning that the hard domain structure is permanently altered substantially. For IPDI-40, IPDI-60, and IPDI-80 samples with increasing IPDI content, the evolution processes of 2D-SAXS patterns are similar to the process of IPDI-100 and can recover the initial state to some extent after release. To further characterize the stability of the hard domain, we plot 1D-SAXS intensity profiles to compare unstretched state and recovery status of different samples, separately. As shown in Figure 3, the parallel direction is consistent with the stretching direction, while the perpendicular direction is transverse to it. The first-order scattering peaks in the 1D-SAXS intensity profiles are evident, suggesting the presence of the microphaseseparated structure in these PU samples.48 For the IPDI-0 sample, the intensity difference between the two states is evident in both directions as shown in Figure 3a. Moreover, the scattering vector (q) peak, which corresponds to interdomain spacing, is no longer visible in the parallel F


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Macromolecules On the other hand, the interdomain spacing ratio (Dr) can be used to describe the strain-induced instability as well. Dr is defined by the following equation: D Dr = after release D before stretch (4) Here, Dafter release and Dbefore stretch represent the interdomain spacing of polyurethane samples after release and before stretch, respectively. The greater deviation of Dr from 1, the more destruction caused by stretching. Figure 4 shows the

Figure 4. Tendency of interdomain spacing ratio (Dr) with growing IPDI content. || and ⊥ represent the parallel and perpendicular directions, respectively. Figure 5. XRD profiles of PU samples (a) before and (b) after stretching.

tendency of the interdomain spacing ratio (Dr) in the parallel and perpendicular directions with increasing IPDI content. It indicates that Dr in the parallel direction is always greater than 1, while Dr in the perpendicular direction remains below 1, meaning that hard domains are more or less oriented in the stretching direction. Remarkably, the Dr∥ ratio of IPDI-0 is missing because the hard domains are largely destroyed in the stretching direction, hardly showing a peak or shoulder in the intensity profile. For the IPDI-40, -60, -80, and -100 samples, Dr is close to 1 in both directions, which suggests that the introduction of IPDI weakens the orientation of hard domains after release and improves their stability. XRD Study of Crystallization. Since the crystallization of the hard domains would seriously affect the shape recovery results, we carried out the XRD experiments for specimens before and after 500% elongation. The XRD patterns of the samples before stretching are shown in Figure 5a. All onedimensional curves feature one broad peak around 21°, suggesting the presence of amorphous state without obvious crystal structure.49 By comparison, results of the released samples, having been stretched to 500% strain and maintained for 5 min, are displayed in Figure 5b. These curves show a similar shape as that of unstretched samples, with broad peak around 21° and only a slight difference in peak height. Combined with the SAXS scattering pattern, it is implied that strain-induced crystallization in hard domains of IPDI-0 and IPDI-20 may happen, but very weakly. The crystallization of PTMG is impossible because its melting point is below room temperature, as shown in Figure S2. FTIR Analysis of Hydrogen-Bonding Association and Orientation. The variation of hydrogen-bonding association with IPDI concentration is monitored by the IR spectra (Figure 6a), and the result is shown in Figure 6b. As in polyether polyurethane, the CO···H−N hydrogen bonds only exist in

hard domains, while the free carbonyl group only dispersed in soft domains.31,32,50 Accordingly, hydrogen-bonding association of carbonyl group can be utilized to characterize the microphase separation between hard and soft domains. The specific calculation method and spectra analyses are explained in the Supporting Information (Figure S3). Figure 6b shows that from IPDI-0 to IPDI-40 the content of associated carbonyl group decreases rapidly as IPDI content increases, whereas from IPDI-40 to IPDI-100, the association concentration stays relatively constant. The dependence of hydrogen-bonding association on IPDI component is thus nonlinear. Such an observation is qualitatively consistent with the results obtained by other researchers; that is, segmented polyurethanes, as well as polyureas, with stereoregularity hard segments show sharper microphase separation than those constructed from less symmetric hard segments.26,27 A schematic diagram of microphase separation is shown in Figure 6c. With benzene rings gradually substituted by cyclohexane rings, the amount of hydrogen bonding among hard segments decreases; as a consequence, more hard segments are mixed into soft domains.51 The stability of hydrogen bonding is characterized by in-situ infrared dichroic measurements. Figure 7 displays the orientation function (f) of the hydrogen-bonded carbonyl group in all samples under a 500% strain stretching and recovery process. As polyether glycols were used as soft segments, the hydrogen-bonded carbonyl group (1700 cm−1) is located only in hard domains.32 For the IPDI-0 and IPDI-20 samples, f of the hydrogen-bonded carbonyl group decreases until strain reaches about 140%, suggesting that hard segments become perpendicularly oriented with respect to the deformation direction at small strain. With larger deformation, G


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Figure 6. (a) FTIR spectra of polyurethane samples. (b) Hydrogen-bonding association of polyurethane samples. (c) Schematic diagram of microphase separation in polyurethane elastomers containing different contents of nonplanar rings.

f of hydrogen-bonded carbonyl group begins to increasea change that is attributable to the parallel orientation of hard segments with respect to the stretching direction. Such a phenomenon implies the breaking and reorganization of hard phases. During the recovery process, f of the hydrogen-bonded carbonyl group continues to grow, eventually preventing a complete return to zero. High residual f after recovery indicates the destruction of original hard segments and poor hydrogenbonding stability. However, for the IPDI-60, IPDI-80, and IPDI-100 samples, a linear-like increase of f is observed during stretching, indicating a gradual increase of hard segment orientation as strain is induced. At the end of the releasing process, f almost decreases to zero, suggesting that hydrogen bonding is fully recovered to the initial randomly oriented state. For the IPDI-40 sample, the stability of hydrogen bonding is intermediate. The degree of orientation decreases during recovery, though a small fraction with permanent orientation still exists. Macroscopic Recovery Property. To test recoverability, samples are stretched until break occurs to observe their recovery ratio. After breaking, the shape recovery ratio is measured within 3 min. As shown in Figure 8, the recovery ratio rises rapidly from 75% of IPDI-0 to 97% of IPDI-100 and reaches over 95% with IPDI content greater than 40%. Thus, it is evident that greater stability of hard domains results in better recovery of macroscopic mechanical property. Such properties, including multiple stretching and recovery curves, can be found in Figure S4, and similar trends are observed. Structural Mechanism during the Gradually Replacement of PPDI with IPDI. The above results can be interpreted using Bonart’s model, in which there are two different hard

domain types in thermoplastic polyurethanes at low content of hard segments.52 As shown in Figure 9, the first type is the lamellar hard domain model wherein the long axis is perpendicular to the polymer chain axis, denoted as model 1. Importantly, the concept of lamellar here is completely different from the lamellar in polymer crystals. Here, the lamellar domain is amorphous and linked by hydrogen bonding. The second type is the fibrillar hard domain model, in which the long axis is parallel to the polymer chain axis, shown in model 2. Furthermore, the orientation of hydrogen bonding is approximately perpendicular to the hard segment axis in each model. When external tensile stress is applied, these two hard domain models undergo different orientation processes, forming hydrogen bondings with opposite orientations at the same time. The evolution process of 2D-SAXS patterns for IPDI-0 and IPDI-20 in our work is similar to the results found for segmented poly(urethane−urea) elastomer researched by Hsiao et al.,29 whose model could explain the SAXS and insitu infrared dichroic results of IPDI-0 and IPDI-20. The evolution process of lamellar hard domains is presented by model 1 in Figure 9. In the initial state, the lamellar hard domains, together with hard segments, are randomly oriented. Consequently, a circular scattering halo appears in the 2DSAXS patterns, and f of in-situ infrared is clear to zero. At a small strain, lamellar hard domains begin to rotate along the long axis parallel to the stretching direction, with hard segment axis perpendicular to the stretching direction. As a result, f of insitu infrared is in the negative region at the beginning of deformation. Under larger strain, a part of lamellar hard domains fails to afford external stress, causing it to break, while H


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Figure 7. Orientation function of the hydrogen-bonded carbonyl group upon 500% strain (stretching and recovery processes): (a) IPDI-0, (b) IPDI20, (c) IPDI-40, (d) IPDI-60, (e) IPDI-80, and (f) IPDI-100.

scattering pattern and f of in-situ infrared continue to increase. On the contrary, the lobe in scattering pattern shrinks and its intensity decreases, indicating that the lamellar structure becomes less and less. In the recovery process, the broken lamellar domains of IPDI-0 and IPDI-20 are incapable of recovering their initial state and position. Stress tears the majority of the lamellar domains into pieces, and very small parts of them crystallize into needle-like crystals, leading to the rhombus-shaped pattern and weak scattering intensity in the 2D-SAXS patterns. The survived lamellar domains are tiny and may recover back to the initial state. Therefore, f of in-situ infrared continues to grow slightly, which can be affected by two aspects: unbroken lamellar hard domains return to the initial state and newly formed fibrillar hard domains with orientation. For the IPDI-60, IPDI-80, and IPDI-100 samples, we infer that there is only slight breakage of hard domains during the stretching process. It is likely that nonplanar rings hinder the stacking of hard segments and impede the formation of lamellar hard domains. As a result, the hard segments of these samples can only form a fibrillar structure. The evolution process of fibrillar hard domains, presented by model 2 in Figure 9, is relatively simple. At the beginning of deformation, fibrillar hard

Figure 8. Recovery ratio after breaking in 3 min of all samples.

the orientation of the rest lamellar domains keeps increasing. In the meantime, needle-like crystals containing hard segments form with the long axis parallel to the stretching direction, showing a streak on the meridian of the 2D-SAXS patterns. When the strain is large enough to destroy most of the lamellar hard domains, more fibrillar hard domains and needle-like crystals form. Correspondingly, both the intensity of streak in I


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Figure 9. Different deformation processes of lamellar and fibrillar hard domain models during stretching. Model 1: lamellar hard domain; model 2: fibrillar hard domain. For simplicity, hard segments with different length are presented with black bars of the same size.

domains rotate along both the long axis and hard segment axis parallel to the stretching direction. Consequently, the 2D-SAXS patterns are elliptical, and f of in-situ infrared continues to rise without a negative orientation region. On the other hand, the interdomain spacing in the stretching direction increases, while it decreases in the perpendicular direction. As a result, the 2DSAXS patterns are compressed due to deformation. Finally, in the recovery process, both 2D-SAXS patterns and f of in-situ infrared return to the initial state, demonstrating excellent hard domain recovery and hydrogen-bonding stability. For the IPDI-40 sample, according to the 2D-SAXS patterns and f of in-situ infrared results, the evolution of hard domain involves a process that combines model 1 with model 2, as shown in Figure 9. Because PPDI is gradually substituted with IPDI, both lamellar and fibrillar hard domain structure exist in the IPDI-40 sample. During deformation, most parts of the lamellar hard domains are destroyed under external force; however, the majority of hard domains, including the remaining part of lamellar hard domains and most fibrillar hard domains, maintain their structural integrity. As a result, due to the damage of most lamellar domains, there is perceivable intensity difference in the 2D-SAXS intensity profiles of IPDI-40, and f of in-situ infrared is still positive after recovery, suggesting the existence of permanently orientated hydrogen bondings. Although the morphological explanation seems to provide a physical picture during the shape transformation and recovery, we must keep in mind that hard phases work as net points and maintain the dimensional frame of the elastomer. Since the external stress would apply and concentrate on these hard domains, the glass state of the hard phases is required to keep shape memory behavior. The glass transition temperatures of the hard phases of each sample are listed in Table 2, and the

Table 2. Glass Transition Temperatures of the Hard Phases Tg (°C) IPDI-0 IPDI-20 IPDI-40 IPDI-60 IPDI-80 IPDI-100

69 60 43 28

details of DMA experiment are described in the Supporting Information (Figure S5). As the content of nonplanar structures increases, the Tg of hard domains decreases rapidly. Notably, the Tg of IPDI-100 is about 28 °C, very close to room temperature. Since IPDI-0, IPDI-20, and IPDI-40 with very high Tg could not withstand the external stress, the excellent recovery ratio of IPDI-60, IPDI-80, and IPDI-100 cannot be attributed to the rigid glass state. On the other hand, the binding energy of hydrogen bonding is not very high, normally around ∼40 kJ/mol. With the nonplanar structure of IPDI, the hydrogen-bonding energy would probably be even lower, which could lead to a reduction in the stability of hard phases. Therefore, the existence of stable hard domains under large strain may have other underlying mechanisms. In this work, we adopt SMFS to investigate the possible mechanism. SMFS Investigation of Single Chain Mechanical Properties. Figure 10 displays the single molecule force spectroscopy results and their dependence on the synthesized polyurethanes. The modified freely jointed chain (M-FJC) model is used to fit the force−extension curves of all samples.53−55 The M-FJC model treats a polymer chain as an aggregate of Kuhn segments with a length of lk (Kuhn length) and segment elasticity, Ksegment. The model is shown as follows: J


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Figure 10. (a−f) Fitting results of the single molecular force spectrum by the modified freely joined chain (M-FJC) model. The black dotted lines in (a−f) are the fitting results, and the red lines are the experimental results: (a) IPDI-0 with lk = 0.40 nm, Ksegment = 12 000 ± 1300 pN nm−1; (b) IPDI-20 with lk = 0.36 nm, Ksegment = 10 000 ± 1000 pN nm−1; (c) IPDI-40 with lk = 0.36 nm, Ksegment = 9000 ± 900 pN nm−1; (d) IPDI-60 with lk = 0.35 nm, Ksegment = 8000 ± 900 pN nm−1; (e) IPDI-80 with lk = 0.35 nm, Ksegment = 6400 ± 600 pN nm−1; (f) IPDI-100 with lk = 0.37 nm, Ksegment = 5200 ± 500 pN nm−1 (lk, Kuhn length, represents chain flexibility and Ksegment is chain elasticity). (g) lk of all samples. (h) Ksegment of all samples and linear fit of Ksegment.

⎡ ⎛ Fl ⎞ k T ⎤⎛ nF ⎞⎟ x(F ) = ⎢coth⎜ k ⎟ − B ⎥⎜⎜Lcontour + ⎢⎣ Flk ⎥⎦⎝ K segment ⎟⎠ ⎝ kBT ⎠

obvious impact on the flexibility of the polyurethane chains in our work. Nevertheless, different IPDI content in polyurethane chains greatly influence main chain elasticity of the samples. Ksegment of IPDI series are shown in Figure 10a−f. The value of Ksegment continuously decreases with increasing IPDI content (Figure 10h) and, more specifically, with increasing content of nonplanar rings. Employing a least-squares linear fit, the value of Ksegment decreases by approximately 64 pN nm−1 for every 1 mol % increase in IPDI content. The low value of Ksegment designates better elasticity and easier elongation under external force. As the structure and content of soft segments are identical for all samples, the results indicate that differing content of nonplanar ring in the hard segments leads to the variation of chain elasticity. Marszalek et al.25 have demonstrated that the chair-to-boat transition of nonplanar glucopyranose ring causes the elastic deformation of the polysaccharide

(5)

Here, F is the external load, x is the length of the polymer chain under external load, lk is the Kuhn length, kB is the Boltzmann constant, T is the absolute temperature, Lcontour is the contour length of the polymer chain, n is the number of Kuhn segments in the polymer chain (equals to Lcontour/lk), and Ksegment is the segment elasticity (describing the deformability of Kuhn segments). The normalized force−extension curves superimpose well, as demonstrated in Figure S6, which signifies that the force curves reflect the single-chain elongation of the TPUs.56 As presented in Figure 10g, the values of lk are all in the range of 0.35−0.40 nm, indicating that the different chain structures have no K


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could relax the stress concentration and preserve the hard domains. Nevertheless, we currently lack direct evidence for “elastic” hard domains, and thus further work will be required in the future to confirm this hypothesis.

chain. When loading the same external force, polymer chains with more nonplanar rings are easier to deform.57 Similarly, IPDI could contribute its enthalpic elasticity here. The SMFS results and analysis above imply a possible mechanism in which hard domains may possess some degree of “elasticity”, even when they are in the glass state. The nonplanar ring structure may act as the six-membered ring in a polysaccharide chain, which could exhibit a conformational transition from chair to boat, and thus extend under stress.24,25,58 Consequently, it is reasonable to suspect that the fibrillar hard domain with IPDI could elongate along the stretching direction and partially relieve the stress concentration. The hard domains are therefore preserved by the nonplanar ring structures. Still, such a hypothesis lacks direct experimental evidence, and further work is needed to substantiate it.

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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.6b01172. 1 H NMR, DSC, and DMA spectra for TPU samples; a detailed description of FTIR peak-fit process in the C O stretching region; recovery ratio of the cyclic tensile testing of 500% strain; normalized force−extension curves of single molecule force spectroscopy (PDF)

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CONCLUSION In order to understand the super recoverability induced by nonplanar ring structure after large strain stretching, we designed a series of TPU samples, with aromatic rings in hard segment gradually replaced by nonplanar structures, to study their multiscale properties. In IPDI-0 and IPDI-20 samples, the hard phases were seriously ruptured after 500% strain elongation as shown in SAXS results. Therefore, the frame of the TPU was destroyed, and recoverability is quite limited. In contrast, the hard domains remained intact after 500% stretching for IPDI-60, IPDI-80, and IPDI-100 samples, with more than 95% shape recovery. IPDI-40, an intermediate sample, showed moderate recoverability. Dichroic FTIR experiments showed a similar trend in hydrogen-bonding association. Hydrogen bonding in IPDI-60, IPDI-80, and IPDI100 was well preserved, which allows both deformation along the stretching direction under increasing strain and recovery to the initial state after release. However, in IPDI-0, IPDI-20, and IPDI-40 samples, the orientation of hydrogen bonding was initially parallel to the strain direction before turning into vertical, while accompanying the break of hard domains. In the restoring process, the hydrogen bonding remained vertical to the stress direction and was incapable of recovery. With these two facts, we proposed that in IPDI-0, IPDI-20, and IPDI-40 samples the hard phases are lamellar-like structures which break into pieces under large strain. This is followed by the orientation of smaller hard domains along the elongation direction. However, they maintain this orientation or even increase the tilt angle during the macroscopic shape recovery process. Although DSC and XRD did not show strong evidence of crystallization, the SAXS results led us to believe that partial crystallization occurred during the elongation process. On the other hand, the hard domains in IPDI-60, IPDI-80, and IPDI100 samples are fibrillar-like and may easily orient along the stress direction. Such a structure did not show any evidence of crystallization but demonstrated an outstanding recovery to the initial state, at almost 100% recovery ratio. However, the integrity of hard phases with high nonplanar structure content cannot be explained by the fibrillar model. Because of the stress concentration effect, the fibrillar structure can break into even smaller pieces under stress, especially when Tg is close to room temperature. The SMFS test revealed that the elasticity of the polymer chain increases as the nonplanar structure content increases. On the basis of this, we proposed a possible mechanism in which hard domains containing high IPDI content possess some degree of deformability, which

AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (R.Z.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is supported by Zhejiang Provincial Natural Science Foundation of China (Grant LY15B040006), the Open-end Funds of Beijing National Laboratory for Molecular Sciences (Grants 20140147 and 20150115), and the Open Fund of Zhejiang Provincial Top Key Discipline of Aquaculture (Grant XKZSC05). The authors thank Nissim Ray in University of Michigan, Ann Arbor, Dan Zhou in California Institute of Technology, and Prof. Charles C. Han in Shenzhen University for their assistance in revising the paper. We also appreciate the support of small-angle X-ray scattering from the Shanghai Synchrotron Radiation Facility (SSRF).

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REFERENCES

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