External Stress-Free Reversible Multiple Shape Memory Polymers

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Applications of Polymer, Composite, and Coating Materials

External Stress-Free Reversible Multiple Shape Memory Polymers Ya Nan Huang, Long Fei Fan, Min Zhi Rong, Ming Qiu Zhang, and Yu Ming Gao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b10052 • Publication Date (Web): 05 Aug 2019 Downloaded from pubs.acs.org on August 11, 2019

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ACS Applied Materials & Interfaces

External Stress-Free Reversible Multiple Shape Memory Polymers Ya Nan Huang†, Long Fei Fan*‡, Min Zhi Rong*†, Ming Qiu Zhang*†, and Yu Ming Gao¶ †Key Laboratory for Polymeric Composite and Functional Materials of Ministry of Education, GD HPPC Lab, School of Chemistry, Sun Yat-Sen University, Guangzhou 510275, China ‡School of Textile Materials and Engineering, Wuyi University, Jiangmen, Guangdong 529020, China Guangdong JISU New Materials Co., Ltd, Dongguan 523527, China

¶

KEYWORDS: two-way shape memory effect, reversible shape memory effect, multiple shape memory polymer, hydrogen bonds, polyurethane

ABSTRACT: The present work is focused on developing external stress-free two-way triple shape memory polymers (SMPs). Accordingly, a series of innovative approaches are proposed for the material design and preparation. Polyurethane prepolymers carrying crystalline polytetrahydrofuran (PTMEG) and poly(ε-caprolactone) (PCL) as the switching phases are

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respectively synthesized in advance and then crosslinked to produce the target material. The stepwise method is believed to be conducive to manipulation of the relative contribution of PCL and PTMEG. Moreover, the chain extender, 2-amino-5-(2-hydroxyethyl)-6-methylpyrimidin-4ol (UPy), is incorporated to establish hydrogen bonds among the macromolecules. By straightforward stretching treatment at different temperatures, the hydrogen bond networks are successfully converted into internal stress provider, which overcomes the challenge of stress relaxation of the melted low melting temperature polymer (i.e. PTMEG) and increases the efficiency of stress transfer. Meanwhile, the contraction force of the switching phases is tuned to match the internal tensile stress. As a result, the internal stress provider can closely collaborate with melting/recrystallization of the crystalline domains, leading to the repeated multiple shape memory effects. The crosslinked polyurethane is thus able to reversibly morph among three shapes and displays its potentials as soft robot and actuator. The strategy reported here has the advantages

of

easy

accessible

raw

materials,

simple

reaction

and

facile

programing/deprograming/reprograming, so that it possesses wide applicability.

INTRODUCTION Shape memory polymers (SMPs) are one of the major classes of smart materials capable of fixing one or more deformations and recovering to permanent shape from temporary shapes in response to external stimuli. Based on whether the deformation is reversible, SMPs can be classified as one-way and two-way types. The latter exhibits more versatile merits and allows for advanced application prospects in soft robots1-7, actuators8-10, artificial muscles11-16, shape changing substrates17-21, intelligent fibers22, 4D printing23,24, etc. Unlike one-way SMPs, no re-


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programming step has to be applied to two-way SMPs for re-gaining shape memory effect (SME) after recovery. On the other hand, traditional one-way SMPs can only memorize one temporary shape and one permanent shape. In this case, physical or chemical crosslinking used to maintain the permanent shape, and a trigger mechanism (like glass transition and melting) fixes/releases the temporary shape by switching on/off the molecular mobility. In 2006, Lendlein and the co-workers25 firstly reported triple SMPs that can memorize two temporary shapes by introducing two switching units. The number of discrete switching phase transitions decided the number of temporary shapes. Later on, Xie et al.26 took advantage of a broad switching phase transition and the resultant SMPs exhibited dual-, triple- and quadruple-shape memory effects. Since then many studies of multiple shape memory polymers have been carried out in this aspect (Table S1). So far, however, most of the materials still belong to one-way SMPs, and there are no reports about external stress-free two-way multiple shape memory polymers, to our best knowledge. Only very few materials, like poly(ω-pentadecalactone) (PPD)-poly(ε-caprolactone) (PCL) copolymer27, crosslinked polyethylene/PCL blends28 and crosslinked ethylene vinyl-acetate copolymer (EVA)/PCL blends29, were able to realize quasi two-way multiple shape memory effect with the aid of external stress. In this context, the objective of the present work lies in the development of a two-way triple SMP that is able to take effect in the absence of applied force. The previous works on two-way dual SMPs30-37 indicate that the bi-directional variation of driving domain and internal stress is the key factor determining the functionality (Figure S1). Similar design should also be valid for making free-standing two-way multiple SMPs. According to this consideration, we plan to utilize two crystalline phases with different and separated melting temperatures as the switching units, so that two temporary shapes would be remembered.


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Under the circumstances, the corresponding internal stress should be tension as required by the orientational recrystallization of the crystalline phases during cooling from the molten state. Moreover, the following two difficult issues have to be overcome. (i) Since the multiple SME successively operates with changing temperature, the efficient transfer or maintaining of internal stress throughout the polymer is critical for oriented recrystallization of the melted phases from high to low temperatures. If inappropriate internal stress provider (e.g., the rubbery spheres that were successfully used in a two-way dual SMP29) were included in the envisaged two-way multiple SMPs, the stress would fail to be transmitted through the melt of the lower Tm crystalline phase when the higher Tm crystalline phase started to be crystallized in the course of cooling. The orientation degree of the recrystallized higher Tm crystalline phase had to be much lower than the expected as a result of the reduced internal stress. After a few cycles, the effect of the higher Tm crystalline phase would greatly decay and the two-way triple SMP would become two-way dual SMP. Similar event would not happen in the case of two-way dual SMPs because only one type of crystalline phase is included. (ii) The contraction force offered by the switching phases during melting should match the tensile stress given by the internal stress provider. Supposing that the internal tensile stress is much higher than the melting-induced contraction force, the crystalline phases would be hard to transform into entropically favorable random coils. Meanwhile, when the internal stress is rather low, the melted crystalline phases could not recreate the desired oriented crystalline structures along with cooling. All these would weaken or deteriorate reversibility of the SME. In this work, we propose a set of novel approaches to tackle the challenges discussed in the last paragraph for preparing external stress-free two-way triple SMPs. Basically, crosslinked polyurethane (PU) with crystalline PCL (Tm = 52.5 ℃38) and polytetrahydrofuran (PTMEG, Tm


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= 26.2 ℃39) is synthesized as the target polymer (Figures S2 and S3). Polyurethane is characterized by the hybrid structure of soft-/hard-segments and hence versatile properties. The crystalline regions of PCL and PTMEG serve as the switching phases in this case, while the crosslinking helps to prevent the structural collapse and secures the permanent shape even when both of the soft segments are melted. In particular, the PU prepolymers respectively containing PCL and PTMEG are produced in advance (Figure S2), and then crosslinked by the trifunctional crosslinking agent trimethylolpropane tris(3-mercaptopropionate) (TMPMP) (Figure S3). Such a step-by-step combination of PCL and PTMEG benefits the adjustment of the relative contents of the two switching phases, which in turn changes the balance of the force between the melted switching phases and internal stress provider. Besides, as the species of diisocyanate exerts an important influence on crystallization power of the PU containing crystalline soft segments, PU with structurally symmetric diisocyanate (e.g., hexamethylene diisocyanate (HDI), toluene diisocyanate (TDI) and 4,4'-diphenylmethanediisocyanate (MDI)) has higher crystallinity due to the higher regularity of the PU molecules than the PU with structurally asymmetric diisocyanate (e.g., isophorone diisocyanate (IPDI))40, we incline toward symmetric diisocyanate for constructing the PU. Compared with MDI and TDI, HDI excludes benzene ring and consists of aliphatic chain, which imparts attractive flexibility to HDI-based polyurethane so that the latter is more easily crystallized41,42. In consequence HDI is selected as the diisocyanate. Moreover, the chain extender 2-amino-5-(2-hydroxyethyl)-6-methylpyrimidin-4-ol (UPy)43 is employed, which enables the formation of inter- and intra-macromolecular hydrogen bonds of PU. The built-in hydrogen bond networks in PU prove to be deformable and feasible for playing the role of internal stress provider after simple programming30. The uniformly distributed and


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networked hydrogen bonds in the present polymer suggest that internal stress can be directly applied to the switching phases at the molecular level. The aforesaid possible relaxation of the internal stress transferred from the discrete internal stress providers due to the melted crystalline regions would thus be prevented. Because UPy dimers possess an excellent combination of high thermodynamic stability and rapid kinetic reversibility, which have been widely incorporated into polymeric materials for improving mechanical properties44-46, similar benefits would also be received by our PU.

2. EXPERIMENTAL SECTION 2.1 Materials 1,4-Butanediol (BDO, 99%), dibutyltin dilaurate (DBTDL, 95%), hexamethylene diisocyanate (HDI, 99%), trimethylolpropane tris(3-mercaptopropionate) (TMPMP), guanidine carbonate, PTMEG (Mn = 2900) and 2-acetyl butyrolactone were supplied by Aldrich. PCL (Mn = 3000) was purchased from Daicel Chemical Industries, Japan. 2.2 Synthesis Synthesis of UPy (Figure S4). 2-Acetyl butyrolactone (2.00 mL) and guanidine carbonate (3.30 g) were put to reflux in absolute ethanol (20.00 mL) in the presence of triethylamine (5.20 mL). After overnight refluxing at 80 oC, the solution became yellow and turbid. The solid content was filtered, washed with ethanol, and suspended in water. Then, the pH was adjusted to 6-7 with HCl-solution, and the mixture was stirred for a while. Finally, the pure product was obtained by filtration, rinsing of the residuum with water and ethanol, and drying. The proton nuclear magnetic resonance spectrum (1H NMR, 400 MHz, dimethyl sulfoxide (DMSO), 25 C, Figure S5) of the product shows the following characteristic peaks (δ/ppm):


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10.702, 6.284, 4.509, 3.361, 2.446, 2.062. 10.702 (1H, -NH-C(NH2)=N-), 6.284 (2H, -NHC(NH2)=N-), 4.509 (1H, -CH2-CH2-OH), 3.361 (2H, -CH2-CH2-OH), 2.446 (2H, -CH2-CH2-OH), and 2.062 (3H, -CO-C(CH2-CH2-OH)-C(CH3)-). The carbon-13 nuclear magnetic resonance spectrum (13C NMR, 400 MHz, DMSO, 25 C, Figure S6) with the following characteristic peaks also verifies the structure of the product (δ/ppm): 164.930 (-CO-C(CH2-CH2-OH)-C(CH3)), 164.17 (-CO-C(CH2-CH2-OH)-C(CH3)-), 155.029 (-NH-C(NH2)=N-), 109.93 (-CO-C(CH2CH2-OH)-C(CH3)-), 61.821 (-CO-C(CH2-CH2-OH)-C(CH3)-), 30.652 (-CO-C(CH2-CH2-OH)C(CH3)-), and 23.215 (-CO-C(CH2-CH2-OH)-C(CH3)-). Fourier transform infrared (FTIR) spectrum of the resultant is shown in Figure S7 (KBr/cm-1): 3376, 3118, 2954 and 2842. Mass spectrum of the exact mass calculated for [M+1]+ C7H11O2N3 169; found 168.0 (Figure S8). Synthesis of PU with UPy as the chain extender (PUUPy) (Figures S2 and S3). Five PUs with different ratios of the two switching phases (Table S2) were synthesized. Typically, HDI and two drops of DBTDL were dissolved in 100.00 mL N,N-dimethylformamide (DMF). The system was heated to 80 C under stirring in Ar. PCL was added with stirring for 1 h. Then, UPy was incorporated as chain extender under stirring for 1 h to obtain the PU prepolymer 1. To prepare PU prepolymer 2, the same procedures were followed except that PTMEG rather than PCL was incorporated. Eventually, PU prepolymer 1 and PU prepolymer 2 were mixed under stirring for 30 min and then the crosslinking agent TMPMP was added with stirring for 1 h (Figure S3). After that, the mixture was poured into a polytetrafluoroethylene (PTFE) mold to remove the solvent at room temperature for 96 h and PUUPy film was produced. Unless otherwise specified, the PUUPy specimens used in this study were made following formula #2 (Table S2).


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Synthesis of PU with BDO as the chain extender (PUBDO). To highlight the critical role of UPy, control sample of polyurethane excluding UPy was prepared following the above steps, except that UPy was replaced by equimolar BDO.

Figure 1. Programming for acquiring two-way triple SME. 2.3 Programming To prepare two-way triple SMPs, the as-synthesized PUUPy and the control PUBDO specimens were trained in accordance with the steps illustrated in Figure 1. Firstly, the dumbbell-shaped specimen was heated to 90 ºC for 10 min and stretched to 4 times the length of the original one. Next, the specimen was cooled down to room temperature and the applied force was removed. Finally, the specimen was heated up to 55 ºC for 10 min and cooled down to room temperature once more. 2.4 Characterization Proton nuclear magnetic resonance (1H NMR) and carbon-13 nuclear magnetic resonance (13C HMR) spectra were recorded by an AVANCE III (400 MHz, Bruker, Germany) by using dimethyl sulfoxide-d6 as the solvent.


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Mass spectra were collected by a LCMS-2010A spectroscopy (Shimadzu, Japan) by using dimethyl sulfoxide as the solvent. Dynamic mechanical analyzer (DMA, Q800, TA Instruments, USA) was employed to study dynamic mechanical properties at 1 Hz. The variation in the specimen’s length at zero force was measured under controlled forced mode. For each cycle of measurement, the sample was heated to 55 ºC at 1 ºC/min, kept at 55 ºC for 2 min, cooled down to -20 ºC at -1 ºC/min, and finally kept at -20 ºC for 2 min. Fourier transform infrared (FTIR) spectra were recorded by a Nexus 670 spectrometer (Thermo Nicolet, USA). The orientation of the polymers was characterized in terms of dichroism of δ(-CH2-) absorption of PCL and PTMEG at about 1470 cm-1 47. The extent of anisotropy in the oriented PU was assessed by the corresponding dichroic ratio, R = A∥/A⊥, where A∥ and A⊥ are the absorption intensities measured with radiation polarized parallel and perpendicular to the stretching direction, respectively. Accordingly, (R-1)/(R+2) gives a measure of the orientation degree47,48. Prior to the measurements, two background single-beam spectra were recorded with the polarizers parallel and perpendicular to the elongation direction with the same experimental parameters, respectively. Micro-infrared spectra were recorded by a Bruker EQUINOX55 spectrometer coupled with an infrared microscope (HYPERION A670-B). Micro-IR analyses of the thin samples were performed by placing the samples on the microscope motorized stage and selecting the area through a 15× IR objective at room temperature. Reflection sampling mode is selected without additional sample preparation prior to spectral analysis. All infrared spectra are acquired in the range of 4000-600 cm-1 with a spectral resolution of 4 cm-1. The measurement of the entire selected area was 10×10 μm2 and the size of each spot was 1 μm. The tested area was selected


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under a microscope, and then divided into small spots according to the aperture dimension. Every small spot was measured by moving the sample. Such movement was controlled by the computer. The melting and crystallization behaviors of the polymers were determined using differential scanning calorimeter (DSC, Q10, TA Instruments, USA) in nitrogen. Temperature was cyclically controlled. For each cycle of measurement, the sample was heated to 55 ºC at 1 ºC/min, kept at 55 ºC for 2 min, cooled down to -20 ºC at -1 ºC/min, and finally kept at -20 ºC for 2 min. Tensile strength was measured by a universal tester (CMT 6000, SANS, China) at a crosshead speed of 50 mm/min at 25 C according to ISO527-2.

3. RESULTS AND DISCUSSION Following the well-designed synthesis route (Figures S2 and S3), the target crosslinked polymer PUUPy was prepared. The FTIR spectra in Figure S9 verify the structure of the produced PU prepolymers. The obviously weakened absorptions of -NCO at 2272 cm-1 on the spectra of PU prepolymer 1 and prepolymer 2 in comparison with that on the spectrum of HDI indicate that PCL and PTMEG have completely reacted with the excessive HDI. Moreover, the absorptions of -NCO at 2272 cm-1 on the spectra of PU prepolymers 1 and 2 disappear when the mixture of the PU prepolymers is crosslinked by TMPMP, which means that the reaction between -NCO of the PU prepolymers and -SH of TMPMP must have taken place. The crystalline feature of the resultant PUUPy was characterized by DSC. Two endothermic and exothermic peaks are observed (Figure S10), representing that the crystalline PCL and PTMEG have been embedded in the polymer. The result coincides with the analysis of Figure S9. Accordingly, the melting temperatures, Tm, and the crystallization temperatures, Tc, are


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determined to be 49.3 °C, 22.6 °C, 17.0 °C and -20.6 °C, respectively. In comparison with the values of Tm and Tc of the oligomers PCL and PTMEG, those of the corresponding crystalline phases included in the PUUPy are lower, which is especially true for Tc. It is believed that the crosslinking reduces chain mobility of PUUPy and hence hinders the crystallization process. On the other hand, the DMA result of PUUPy in Figure S11a reveals that only a single Tg represented by the peak temperature of tan δ appears for PUUPy at about -47.9 C. Considering that HDI consists of aliphatic chains and the soft segments (PCL and PTMEG) also contain a large amount of methylene groups, the common micro-phase separation of polyurethane may be somehow suppressed so that the Tg of the hard segments is not perceived. In fact, the Tg of PUUPy (-47.9 C) is much higher than the glass transition temperatures of PCL (-72 C49) and PEMEG (-76 ~ -80 C50). It means that the strong macromolecular interaction owing to crosslinking must have introduced intensive constraints on the chain segments. Besides, PCL and PEMEG have similar Tg, which accounts for the relatively broad tan  peak of PUUPy. Figure S11b exhibits that the DMA curves of PUBDO resemble those of PUUPy, and Tg of the former is -50.4 C. Clearly, the hydrogen bonds originating from UPy do not greatly affect the movement of the chain segments of the polyurethanes. As mentioned in the Introduction, UPy is incorporated to establish hydrogen bond networks among the polyurethane for purposes of serving as internal stress provider after programming. To understand whether the planned hydrogen bonds have been formed, temperature dependent FTIR spectra of PUUPy are collected (Figure 2a). It is seen that with a rise in temperature the intensity of the peak at 3334 cm-1, which is contributed by the hydrogen bonded -NH-, decreases and that at 3451 cm-1, which is assigned to the non-hydrogen bonded -NH-, increases. Meanwhile, the peak of the hydrogen bonded -NH- shifts from 3334 cm-1 towards higher


11


wavenumber regime along with increasing temperature. The decrease of the peak height of the hydrogen bonded -NH- at 3334 cm-1 is insignificant from 30 to 80 ºC, but becomes significant from 80 to 150 ºC. The results indicate that hydrogen bonds are available in the polyurethane, and the dissociation of plenty of hydrogen bonds mainly takes place above 80 oC. That is, the hydrogen bond networks are basically stable below 80 oC. Furthermore, the micro-IR analysis of the surface of PUUPy (Figures 3a and 3b) reveals that the hydrogen bonds are homogeneously dispersed in the material. This ensures uniform distribution of internal stress at molecular level and hence uniform deformation driven by SME.

(a)

(b) 30 OC 40 OC 50 OC 60 OC 70 OC 80 OC 90 OC 110 OC 130 OC 150 OC

30 C 40 OC 50 OC 60 OC 70 OC 80 OC 90 OC 100 OC 110 OC 130 OC 150 OC

3500

3450

3400

3350

3300

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3200

Absorbance [a. u.]

O

Absorbance [a. u.]

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3500

-1

Wavenumber [cm ]

3450

3400

3350

3300

3250

3200

-1

Wavenumber [cm ]

Figure 2. Temperature dependences of the FTIR absorbance of -NH- group of (a) PUUPy and (b) PUBDO (ramp: 5 oC/min).

The plots in Figure 2b are similar to those in Figure 2a, except that the remarkable variation in the peak height of the hydrogen bonded -NH- starts at 70 ºC. Therefore, the relevant outcomes of the analysis of PUUPy are basically applicable to PUBDO. Figures 3c and 3d also exhibit the


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relatively even arrangement of the hydrogen bonds in PUBDO, which confirms the suitability of PUBDO for serving as the control of PUUPy. These information, together with the data of Figures S10 and S11 that describe the thermomechanical behaviors of the polyurethanes, is useful for formulating the programming procedures of the as-synthesized polyurethane towards SMP (Figure 1), and determination of the working temperature range of the programmed SMP.

Figure 3. Micro-infrared spectroscopy analysis of the surfaces of (a, b) PUUPy and (c, d) PUBDO. (a) Mapping of -CO- group in PUUPy. (b) Mapping of -NH- group in PUUPy. (c) Mapping of -COgroup in PUBDO. (d) Mapping of -NH- group in PUBDO. The color bars represent the intensity ratio of -CO- or -NH- to the reference -CH2-.


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According to the above characterization results, the structural variation of the PUUPy, especially the cultivation of the internal stress provider, along with the programming (Figure 1) is interpreted as follows. At the beginning, the specimen is heated to 90 ℃ (step ① in Figure 1), which is higher than the melting temperatures of the crystalline phases of PUUPy (Figure S10) and the dissociation temperature of hydrogen bonds as well (Figure 2). Since the crystalline phases are melted and the physical restriction offered by hydrogen bond network is largely eliminated, the specimen is allowed to be easily drawn to the length four times longer than the original (step ② in Figure 1). Then, the cooling under the constant straining (step ③ in Figure 1) results in appearance of the oriented crystalline phase from the melted PCL (Figure S12 and Table S3), and re-establishment of hydrogen bond networks. At this moment, the reconstructed hydrogen bond networks are not stressed. Next, the external stress is released at room temperature (25℃, step ④ in Figure 1), and the temporary shape of the elongated specimen has to keep unchanged because the macromolecules lack sufficient mobility. When the specimen is heated to 55 ℃ (step ⑤ in Figure 1), the oriented crystals of PCL melt again and transform into entropically favorable random coils. Accordingly, the specimen inclines to be shortened by the entropy increase induced contraction, while the hydrogen bond networks are gradually compressed as the specimen shrinks and tend to recover to the expanded state when they are reformed (i.e. step ③ in Figure 1). During shrinkage of the specimen, the compressed hydrogen bond networks start to serve as the internal stress provider, producing counterforce in the opposite direction (i.e. tensile stress), which increases with increasing the amount of melt PCL. When the tensile force provided by the hydrogen bond networks is approximately identical to the contraction force caused by entropy increase, the specimen basically reaches stress equilibrium and stops shrinking. Accordingly, the competition between the opposite forces leads


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to the fact that the ultimate specimen length is remarkably decreased, but slightly longer than the original one, as the hydrogen bond networks are still in a state of compression after step ⑤. The obvious shortening of the specimen at step ⑤ implies that the internal stress generated by the hydrogen bond networks is much smaller than the applied stress at step ②. Finally, the specimen is cooled down to room temperature (25 ℃) and the programmed version is available (step ⑥ in Figure 1). In the meantime, the contraction force caused by entropy increase no longer exists and the oriented crystalline of PCL is re-created under the tension exerted by the hydrogen bond networks, which results in slight increase of the specimen length. Since the orientation crystallization (step ⑥) proceeds under the internal tensile stress that is smaller than the external stress (step ②), the orientation degree of the resultant crystallites has to be lower in comparison with the case after step ③ or ④ (Figure S12 and Table S3). It is worth noting that the hydrogen bond networks in the programmed specimen are always under compression and the extent of compression becomes higher at elevated temperature when PCL melts. In addition, the crystalline phase of PTMEG is always molten throughout all stages of the programming as its Tm = 22.6 °C (Figure S10). Having been trained, an increase of temperature of the PUUPy specimen from Tlow (Tlow < Tm1 (melting temperature of PTMEG, see Figure 4), i.e. -20 oC) to Tmid (Tm1 < Tmid Thigh > Tm2, i.e. 55 oC), a continuous shortening of the specimen length (from shape II to shape III) should occur owing to melting of the crystalline domains of PCL. In the inverse process, when the specimen is


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cooled down from Thigh to Tmid, random coils of PCL return to oriented crystallization under the tension imposed by the hydrogen bond networks so that the specimen should grow longer (from shape III to shape II). A further cooling from Tmid to Tlow would result in orientational crystallization of PTMEG with the aid of the internal stress offered by the hydrogen bond networks and the length of the specimen increases from shape II back to shape I.

Figure 4. Mechanism of the temperature dependent two-way triple shape memory effect of the programmed PUUPy.

It means that so long as the cyclic melting/recrystallization of the crystalline components cooperate with the variation in the tension (tighter/looser tension) of the compressed internal force provider (i.e. the hydrogen bond networks), reversible triple shape transformation of the programed PUUPy would be enabled with no need for external force. The photos in Figure 5a demonstrate that the hypothesis is true. The specimen can reversibly deform among three lengths with temperature between -20 and 55 oC. Figures 5b and 5c record the corresponding


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thermomechanical behavior and thermogram of the programmed PUUPy accompanying the twoway multiple shape memory effect. Accordingly, the total average reversible strain is found to be 12.1 % on the basis of the data recorded during the first three heating/cooling cycles (Figure 5b), and the average reversible strains induced by the melting of PCL and PTMEG are 5.5 % and 6.7 %, respectively. Figure 5b further illustrates that there is slope change in the temperature dependence of strain. It may result from the fact that the contraction force aroused by entropy increment is slightly greater than the internal tensile stress provided by the hydrogen bond networks. During the cyclic heating and cooling for implementing the shape memory effect, the specimen length gradually decreases. Meantime, the entropy increase induced contraction decays, while the compression degree of the hydrogen bond networks is raised leading to increased reaction force (i.e. the internal stress). The trade-off between the opposite tendencies eventually results in the balance between the forces, so that the slope change turns to be nearly invisible along with the cycle number. In fact, the results shown in Figure 5 are obtained from the optimized recipe. As both crystalline PCL and PTMEG play the roles of switching phases, their relative fractions are critical for the memory of two temporary shapes. In extreme cases, the content of PCL or PTMEG is much higher than the other (refer to Formulae #1, #4 and #5 of Table S2 and Figure S13), the crystallinity of one of the polymers has also to be much greater than the other. As a result of the insufficient switching phase, the deformation of the soft segment with obviously lower content has to be nearly undetectable (Table S2 and Figure S14). The programed polyurethanes of Formulae #1, #4 and #5 can only memorize one temporary shape. On the other hand, under the circumstances that the contents of PCL and PTMEG are not greatly different


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from each other (refer to Formulae #2 and #3 of Table S2 and Figure S13), the reversible deformations of the two switching phases could be seen. Compared to the polyurethane made from Formula #3, the one from Formula #2 exhibits almost the same reversible strains driven by PCL and PTMEG (Table S2 and Figure S14), so that Formula #2 is assigned to be the optimum. Interestingly, the PUUPy specimen can reversibly morph between three different lengths by making use of its specific structure, which gives an added value to the material’s application. Depending on the actual needs, we can selectively excite one or two of the switching phases (PCL and PTMEG) to achieve the desired deformation mode. As shown in Figure 6, there are three possible routes to achieve reversible single or dual SME.


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Figure 5. (a) Length change of the programmed PUUPy during cyclic heating/cooling between 20 C and 55 C. (b) Thermomechanical behavior and (c) thermogram of the programmed PUUPy during cyclic heating/cooling between -20 and 55 C (ramp: 1 C/min).

Figure 6. (a) Effect of working temperature range on two-way triple shape memory effect of the programmed PUUPy. (b) Thermomechanical behaviors (measured by DMA) of the programmed PUUPy within different temperature ranges.


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Since the desired two-way triple SMP has been able to take effect in the absence of external force, the critical contribution of the internal stress provider (i.e. hydrogen bond networks) needs to be further investigated. As shown in Figure S15, when the hydrogen bonds in the programmed PUUPy are fully dissociated at 110 oC (Figure 2a) and the role of internal stress provider is no longer brought into play, no two-way shape memory effect can be seen. By using this habit, the internal stress exerted by the hydrogen bond networks can be estimated as follows. Firstly, in the case of cooling from 55 oC, the lowest average strain of the programmed PUUPy (1.1 %) is measured at 49.1 oC (refer to the first half of Figure S15). At this time, the expansion driven by re-crystallization has not yet been initiated, and the entropy increment aroused contraction force is equal to the tensile stress of hydrogen bond networks, while the latter reaches the maximum. As the average strain at 49.1 oC of the specimen after losing the reversible SME (refer to the latter half of Figure S15) is -57.3 %, the extensional strain caused by the tension exerted by the hydrogen bond networks at this temperature is determined to be 1.1 % - (-57.3 %), i.e. 58.4 %. By examining the tensile stress-strain curve of the polyurethane measured at 49.1 oC (Figure S16), which had been treated by lithium bromide solution to remove the inter- and intramolecular hydrogen bonds51-53, a stress of 0.63 MPa is found to be required for producing the strain of 58.4 %. In this context, it is known that the highest tensile stress offered by the hydrogen bond networks as internal stress provider is 0.63 MPa. In addition to the vital merit of the hydrogen bond networks, which is necessary for realization of the two-way SME, Figures S15 and S17 also illustrate another important function of the present material design. The programed polyurethane can be deprogramed simply in terms of heat treatment (Figure S15), and the deprogramed version can be reprogramed to gain the two-


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way triple SME once again (Figure S17). The habit must be conducive to renewal and reuse of decayed and defunctionalized materials even if they are crosslinked. The chain extender UPy plays two important roles in the system. On the one hand, it brings in much more hydrogen bonds to the polyurethane than BDO, which means higher internal stress and higher reversible strain. According to Figures S18 and S19, the highest tension offered by the internal stress provider in PUBDO is 0.23 MPa, which is significantly lower than that of PUUPy (0.63 MPa). As a result, the average total reversible strain of the programmed PUBDO is only 7.1 % (Figure S20) due to the relatively lower internal stress provided by the hydrogen bond networks, which forms striking contrast to that of PUUPy (12.1 %, refer to Table S2). On the other hand, the enhanced inter-molecular interaction resulting from the extensive hydrogen bonding would raise cohesion of the macromolecules and hence strength of the polymer. Compared to UPy-free PUBDO whose tensile strength is 19.9 MPa, tensile strength of PUUPy is remarkably higher (i.e. 30.7 MPa, Figure S21) due to the appearance of the UPy moieties in the backbones. The increased robustness is important for the usages as structural material. At last, the prototype applications of the programmed PUUPy are illustrated in Figure 7. Prior to the experiments, the programmed PUUPy strip is glued together with a piece of cardboard. The negligible variation in cardboard size in combination with the shape memory effect induced length change of the programmed PUUPy would lead to bending behavior of the composite. As shown in Figure 5, when the temperature decreases from 30 oC to -5 oC, the programmed PUUPy would expand. In contrast, when the temperature increases from -5 oC to 30 or 50 oC the programmed PUUPy would shrink. According to these habits, the bilayer composite reversibly bends to different directions within the temperature range of interests (Figure 7a). It is worth noting that the conventional multilayer composite materials[54-57] can only realize reversible


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bending deformation in the same direction because the thermal expansion degrees of the two layers are fixed within the same temperature range, and only one original shape and one temporary shape can be memorized. Meanwhile, the programmed PUUPy/cardboard bilayer composite is pre-deformed at 30 oC and sandwiched by crepe paper to mimic the curvature of lotus petals (Figure 7b). A plurality of these “petals” are then combined to produce a complete lotus flower (note: the bottommost petals do not contain the programmed PUUPy strips). Clearly, the lotus can be reversibly transformed with temperature showing the following three states: “bud” (55 oC), “fresh-blown” (30 oC) and “blooming” (-5 oC). These examples manifest that the two-way triple SMP can serve as soft robot (Figure 7a) and actuator (Figure 7b).


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Figure 7. The prototype applications of the programmed PUUPy as (a) soft robot and (b) actuator by taking advantage of the two-way triple shape memory effect driven multiple bending of bilayer structures.

4. CONCLISIONS External stress-free two-way triple shape memory polyurethanes have been prepared, which could reversibly morph among three shapes within different temperature ranges and proved to be able to be used as soft robot and actuator. The embedded crystalline phases of PCL and PTMEG successfully completed their task of serving as reversible domains, while the hydrogen bonds introduced by UPy demonstrated their comfort with the role of internal stress provider (the highest tensile stress created by which was determined to be 0.63 MPa). The core advances of this work lie in the subtle structural design of the target material and the programming procedures. Accordingly, the block against transferring of internal stress due to the melt of the lower Tm crystalline phase was removed, and the melting-induced contraction force managed to be compatible with the internal stress. The compressed hydrogen bond networks of the programed polyurethane closely cooperated with melting/recrystallization of the crystalline


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phases, producing reversible push-and-pull effects, so that the specimen length was allowed to reversibly and multiply change with temperature. Owing to the reversible nature of hydrogen bonds, moreover, the programed polyurethane can be deprogramed and re-programed whenever necessary. The synthesis route and training adopted above were straightforward and can be easily reproduced. There are no strict requirements for the basic materials, and routine semi-crystalline polymers containing hydrogen bonds would be qualified in principle. The proposed idea remarkably broadens the multiformity of authentic two-way shape memory polymers, which can operate without the aid of any applied force.

ASSOCIATED CONTENT Supporting Information. 1H NMR spectra,

13

C NMR spectrum, FTIR spectrum and Mass

spectrum of UPy; FTIR spectra of HDI, PCL, PTMEG, PU prepolymer 1, PU prepolymer 2 and PUUPy; DSC thermogram, thermomechanical behavior, polarized FTIR spectra and typical tensile stress-strain curve of PUUPy; thermomechanical behavior and typical tensile stress-strain curve of the crosslinked PUBDO. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected] (L. F. Fan) *E-mail: [email protected] (M. Z. Rong) *E-mail: [email protected] (M. Q. Zhang)


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Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT The authors thank the support of the National Natural Science Foundation of China (Grants: 51773229, 51333008 and 51673219).

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Table of Contents (TOC) Graphic


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External Stress-Free Reversible Multiple Shape Memory Polymers

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