Integrating transient and sacrificial bonds into biobased elastomers toward mechanical property enhancement and macroscopically responsive property

Journal Pre-proof Integrating transient and sacrificial bonds into biobased elastomers toward mechanical property enhancement and macroscopically responsive property Bin Liu, Zhenghai Tang, Zhao Wang, Liqun Zhang, Baochun Guo PII:

S0032-3861(19)30920-6

DOI:

https://doi.org/10.1016/j.polymer.2019.121914

Reference:

JPOL 121914

To appear in:

Polymer

Received Date: 12 September 2019 Revised Date:

14 October 2019

Accepted Date: 16 October 2019

Please cite this article as: Liu B, Tang Z, Wang Z, Zhang L, Guo B, Integrating transient and sacrificial bonds into biobased elastomers toward mechanical property enhancement and macroscopically responsive property, Polymer (2019), doi: https://doi.org/10.1016/j.polymer.2019.121914. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Ltd.

The incorporation of transient and sacrificial quadruple H-bonding motifs into bio-based elastomer that is synthesized from bio-based di-acids and diols can enhance the mechanical properties and bestow it with adaptive recovery.

Integrating transient and sacrificial bonds into biobased elastomers toward mechanical property enhancement and macroscopically responsive property Bin Liu,a Zhenghai Tang, a,* Zhao Wang,b Liqun Zhangb and Baochun Guoa,* a

Department of Polymer Materials Science and Engineering, South China University

of Technology, Guangzhou 510640, China. E-mail:

[email protected] ;

[email protected]. b

State Key Laboratory of Organic/Inorganic Composites, Beijing University of

Chemical Technology, Beijing 100029, China.

Abstract The development of biobased polymers from renewable resources offers a solution to the growing environmental concerns and scarcity of fossil feedstock. The synthesized biobased polymers, especially for biobased elastomers, are mechanically weak, which greatly restricts their applications. In the present work, we demonstrate that the integration of transient and sacrificial quadruple H-bonding motifs into as-synthesized biobased elastomer can enhance the mechanical properties and bestow it with adaptive performance. Specially, the biobased elastomer was synthesized through melting poly-condensation using biobased di-acids and diols as the starting materials. The biobased elastomer is covalently crosslinked and grafted with ureido-pyrimidinone (UPy) containing acrylate through thiol-ene click reaction by using pentaerythritol tetra(3-mercaptopropionate) as the linkers. Under external load, UPy aggregations based on H-bonding can function as sacrificial units through reversible rupture and re-formation events, leading to significant improvements on the modulus and strength of the biobased elastomer while maintaining the extensibility. In addition, the dissociation and re-formation of H-bonding under thermal stimuli impart the elastomer with thermo-activated shape memory behavior. Key Words: Biobased elastomer; Sacrificial bond; Dynamic bond

Research Highlight Quadruple H-bonding motifs are introduced into as-synthesized biobased elastomer through thiol-ene click reaction. UPy aggregations based on H-bonding can act in sacrificial manner to improve modulus and tensile strength of the elastomer while maintaining the extensibility. Dynamic nature of H-bonding allows the elastomer to access thermo-activated shape memory behavior. The structure-property relationships of the resulted elastomer materials are studied.

Introduction Recently, exploring biobased polymers from renewable resources to substitute petroleum-based polymers has increasingly gained attentions, because it offers a solution to the growing environmental concerns and scarcity of fossil feedstock.1, 2 Conventionally, the methods for producing biobased polymers can be divided into three categories: modification of naturally occurring polymers,3 microbial fermentation4 and polymerization of biobased monomers.5 In particular, the last approach is extensively studied as numerous biomass chemicals are potentially available in large amounts and can be ready to utilize as building blocks toward diverse polymeric architectures upon property demands.6, 7 Thus far, there are already a few successful commercial examples of biomass-derived plastics such as polylactic acid, poly(ɛ-caprolactone) and poly(butylene succinate). In contrast to biobased plastics, biobased elastomers are still in infancy although they find indispensable applications in seals, tires and biomedical materials.8 In our previous studies, a series of fully biobased polyester elastomers are synthesized through the polycondensation of itaconic acid with biobased diols and other biobased diacids, in which the C=C double bonds of itaconic acid provide active sites for crosslinking and the macromolecular chain flexibility can be readily adjusted by varying the composition.9-12 Unfortunately, these biobased elastomers having no crystallization or phase separation are soft and weak, which severely restricted their practical applications. Although the incorporation of fillers can enhance their mechanical properties, it also gives rise to limitations such as filler aggregation, interface control, and processing difficulty.11, 13

Many natural materials are featured with outstanding strength and toughness. A molecular mechanistic origin for the remarkable mechanical properties is benefited from the sacrificial bond in the natural materials, e.g., titin domains in muscle myofibrils14 and ionic bonds in bone.15 The sacrificial bond is able to withstand external force under small deformations and preferentially break over the covalent bonds at large deformations. Consequently, a large amount of energy can be dissipated and the material integrity is reserved.16 Inspired by nature, the sacrificial bond principle has been transferred into synthetic polymers, aiming to enhancing their mechanical properties.17-19 Compared with covalent sacrificial bonds, the noncovalent sacrificial bonds such as H-bonding,20-22 metal−ligand complex,23-25 and ionic interaction17 are particularly promising as it provides recoverable energy-dissipating mechanism through reversible bond breaking/reformation events. In addition, the ability of non-covalent bonds responds to external stimuli (heat, light, etc.) can be used to develop macroscopically responsive materials.26,27 For instance, mimicking modular architecture of titin, Guan et al. incorporated a reversibly unfolding cross-linker of 2-ureido-4[1H]-pyrimidone (UPy) dimer constrained within a macrocycle into synthetic polymer. The reversible breaking/reformation of UPy dimers under external force or at elevated temperature conferred the resulting material with excellent toughness, self-healability and shape memory behavior.28 Herein, we engineer quadruple H-bonding motifs into as-synthesized biobased elastomer to improve the mechanical properties and bestow it with adaptive performance.

The

biobased

elastomer

was

synthesized

through

melting

poly-condensation using industrially available biobased diols and diacids as the starting materials. Then the biobased elastomer is covalently crosslinked and grafted with UPy containing acrylate through thiol-ene click reaction by utilizing pentaerythritol tetra(3-mercaptopropionate) as the linkers. UPy dimers can function as sacrificial units through reversible rupture and reformation, leading to remarkable improvements on the mechanical properties. Additionally, the dynamic equilibrium nature of H-bonding allows the elastomer to access thermo-activated shape memory behavior.

Experimental

Materials 1,4-Butanediol

(BDO),

1,3-propanediol

(PDO),

tetrabutyltitanate

(TBT),

hexamethylene diisocyanate (HMDI), 2-amino-4-hydroxy-6-methylpyrimidine (MIS), hydroxyethyl

methacrylate,

dibutyltin

dilaurate,

pentaerythritol

tetra(3-mercaptopropionate) (PETMP) and 2,2-bimethoxy-2-phenylacetophenone (DMPA) were obtained from Alfa Aesar. Itaconic acid (IA), succinic acid (SA), sebacic acid (SeA), and tris(1-hydroxy-2,2,6,6-tetramethylpiperidin-4-yl) phosphate (inhibitor 705) were provided by Guangfu Fine Chemical Institute of Tianjin. Synthesis of Biobased Polyester Elastomer (BPE) BPE was synthesized according to our previous work, as illustrated in Figure 1a.29 Briefly, the monomers PDO (4.2 g, 55 mmol), BDO (4.9 g, 55 mmol), IA(1.3 g, 10 mmol), SA (4.3 g, 60 mmol), SeA (6.1 g, 30 mmol), and free radical inhibitor 705 (0.01 wt% of the reactants) were added into a three-necked flask. The pre-condensation was performed at 180 oC for 2 h under nitrogen atmosphere. After that, the catalyst TBT (0.05 wt% of the reactants) was added and the mixture was maintained at 220 oC for 4 to 6 h under stirring and reduced pressure until “Weisenberg effect” was observed. The resultant was dissolved in chloroform and precipitated with methanol, followed by vacuum dried at 40 oC overnight. The number average molecular weight and polydispersity index (Mw/Mn) are 48000 and 2.9. The composition of the BPE is characterized by 1H NMR (Figure S1). Synthesis of UPy-based methacrylate (UPyMA) UPyMA was synthesized by reacting UPy-hexamethylene diisocyanate (UPy-HMDI) with hydroxyethyl methacrylate according to the reported literature (Figure 1b).30 UPy-HMDI was prepared by reacting MIS (4.8 g) with HMDI (40 g) at 100 oC for 24 h under nitrogen atmosphere. The resulted UPy-HMDI (20 g, 68.2 mmol) was dispersed in anhydrous CHCl3 (200 mL), and hydroxyethyl methacrylate (12.3 mL, 102 mmol) was then added under nitrogen atmosphere. After stirring for 5 min, 4 drops of dibutyltin dilaurate were added and the mixture was heated to reflux overnight. The obtained cloudy suspension was poured into cold hexane, and the white product was collected and washed with hexane. The product was dried at 40 °C

under vacuum for 24 h, yielding UPy-based methacrylate (UPyMA) (26.10 g; 91%). FTIR and 1H NMR of UPyMA are shown in Figure S2-3. Synthesis of BPE/UPyMA-x A desired amount of UPyMA and BPE were dissolved in DMF, into which PETMP cross-linker (the molar ratio of thiol groups from PETMP and C=C double bonds from BPE and UPyMA is equal to 1) and DMPA photoinitiator (0.2 wt%) were added and stirred for 20 min. The mixture solution was poured into a PTFE mold and subjected to UV irradiation for 35 min (Light source: IntelliRay 400, Intensity: 34 mw/cm2). Finally, the resulted films were vacuum dried at 90 oC overnight. In this context, the sample was named as BPE/UPyMA-x, in which x represents the mass content of UPyMA relative to BPE.

Fig. 1 Synthesis of (a) BPE and (b) UPyMA. (c) Schematic for the preparation of BPE/UPyMA-x.

Characterizations Fourier transform infrared (FTIR) test was conducted using Bruker Vertex 70 FTIR spectrometer. 1H-NMR measurement was performed on Varian NMR spectrometer. Dynamic mechanical analysis (DMA) was carried out on TA DMA Q800 machine under a dynamic strain of 0.5%. The frequency was 1 Hz and the temperature was scanned from −50 to 150 °C at 3 °C/min. Small angle X-ray scattering (SAXS) was conducted on Bruker-AXS NanoSTAR instrument with X-ray wavelength of 0.154 nm. Differential scanning calorimetry (DSC) was conducted on a TA DSC Q20

machine by heating the samples from 30 to 120 oC with 1 oC/min. Uniaxial tensile test, cyclic tensile experiment and stress relaxation experiment were conducted on a U-CAN UT-2060 instrument at room temperature. Cyclic tensile test was conducted by stretching the sample to 100% strain and allowing them relaxation for a certain time before successive cycle. Stress relaxation test was carried out by measuring the stress under 50% strain at room temperature.

Results and Discussion Network structure of BPE/UPyMA-x The chemical crosslinking of BPE and grafting of UPy groups into the networks are clearly evidenced by FTIR spectra. As illustrated in Figure 2a, the as-synthesized BPE exhibits two characteristic absorptions at 1728 and 1640 cm-1, which are attributed to the stretching vibrations of C=O and C=C, respectively. In the FTIR spectra of BPE/UPyMA-0 and BPE/UPyMA-10, the absorption peak related to C=C of BPE at 1640 cm-1 is disappeared and the absorption peak related to thiols of PETMP at 2569 cm-1 is also absent. These results indicate the chemical reaction occurs between thiol groups and C=C, which leads to the covalent cross-linking of BPE. The covalently cross-linked structure of BPE/UPyMA-x is also confirmed by the fact that they are insoluble in toluene (Figure S4). In addition, compared with BPE/UPyMA-0, BPE/UPyMA-10 shows additional absorption peaks around 1669 (-HCCON-) and 1701 cm-1 (-NHCONH-), which are originated from the C=O stretching vibrations of the grafted UPyMA. The self-complementary UPy units can assemble into dimers through quadruple H-bonding. To verify the H-bonding interaction between UPy units in BPE/UPyMA-x, variable temperature FTIR test was performed by taking BPE/UPyMA-30 as an example (Figure 2b). As the temperature increases from 30 to 150 °C, the absorption peak for C=O stretching vibration shifts from 1669 to 1660 cm−1, and the absorption peak related to N-H bending vibration shifts from 1527 to 1518 cm−1. This is because the H-bonding are weakened as a result of dissociation of UPy dimers at elevated temperatures.31

Fig. 2 (a) FTIR spectra of PETMP, BPE, BPE/UPyMA-0 and BPE/UPyMA-10. (b) Variable temperature FTIR spectra for BPE/UPyMA-30 with a temperature increment of 10 oC.

As demonstrated above, BPE/UPyMA-x are bonded by dual cross-links of covalent cross-links and H-bonding based physical cross-links, as schematically illustrated in Figure 1c. To gain structure information at mesoscale, SAXS measurements of bulk BPE/UPyMA-x are conducted. As shown in Figure 3, SAXS profile of BPE/UPyMA-0 does not show any scattering peak, revealing no phase separation. In the case of BPE/UPyMA-x containing UPyMA, a broad scattering peak is observed, which indicates the presence of microphase-separated domains with inter-domain spacing ranging between 6.8 to 25.1 nm. The microphase separation likely stems from the aggregations of UPy groups driven by the combination of

-

interaction of the

UPy ring and the lateral hydrogen bonding between the ureido, which was also reported in other UPy group-containing polymers.32, 33 DSC curve of BPE/UPyMA-30 displays an obvious melting peak at 102 °C, which can be attributed to the melting of UPy stack hard-phase (Figure S5).32, 34 It provides complementary evidence for the presence of microphase separation in BPE/UPyMA-x samples.

Fig. 3 SAXS profiles of bulk BPE/UPyMA-x.

Origins for the improvements on the mechanical properties of BPE/UPyMA-x Figure 4a is the temperature dependence of storage modulus (E’) of BPE/UPyMA-x. As expected, the storage modulus of BPE/UPyMA-x is improved with the increase of UPyMA content due to the increased number of physical cross-links. After a sudden decrease in E’ around glass transition, E’ of BPE/UPyMA-x containing UPyMA gradually decreases and approaches that of BPE-UPyMA-0 as the temperature is increased beyond 150 oC. This is because the UPy aggregations based on H-bonding can serve as physical cross-links at relatively low temperatures to improve E’, while the aggregations are dissociated and no longer contribute to E′ at elevated temperatures. According to the loss factor versus temperature curves in Figure 4b, all BPE/UPyMA-x samples exhibit a distinct loss peak at around -20 °C, which is due to bulk segment relaxation. The glass transition temperature shifts to a higher temperature and tan δ peak value monotonically decreases with the increase of UPyMA content, revealing that the chain mobility is restricted by introducing UPy groups. In addition, a wide loss peak between 50 and 150 °C is observed in the BPE/UPyMA-x containing UPyMA, which is on account of the dissociation of UPy aggregations. 35

Fig. 4 (a) Storage modulus and (b) loss factor δ of BPE/UPyMA-x versus temperature.

The representative stress-strain profiles for BPE/UPyMA-x are illustrated in Figure 5a. BPE/UPyMA-0 is very weak with a tensile modulus (stress at 100% strain) of 0.6 MPa and an ultimate strength of only 0.9 MPa. The introduction of UPy groups into the networks gives rise to significant enhancements on the modulus and ultimate tensile strength without compromising the extensibility. For example, when compare to BPE/UPyMA-0, the modulus and ultimate strength of BPE/UPyMA-30 is increased from 0.6 to 3.5 and from 0.9 to 5.2 MPa, respectively, while the breaking strain is

retained. It is a huge challenge to improve the modulus and strength while maintaining the extensibility of polymers. Usually, there is a trade-off between modulus and extensibility as a ductile network may endure high deformation but shows shallow stress response, while a rigid network may show improved modulus but fails in short extension. Herein, BPE/UPyMA-x architectures are bonded by dual cross-links of covalent and physical cross-links, as shown in Figure 6a. Specifically, the covalent cross-links confer the networks with elasticity and maintains permeate shape, while H-bonding behave as physical cross-links to enhance the modulus. Under external force, the physical cross-links can undergo reversible breaking and re-formation events to dissipate energy. Meanwhile, accompanied with the breaking of physical cross-links, the polymer chains that are coiled up in loops and hidden by physical cross-links can be released to maintain large deformation (Figure 6b). As a result, more energy is necessary to decrease the entropy and increase the enthalpy upon stretching macromolecular chains. Such energy dissipating origins have been well documented in the natural materials having sacrificial bonds.15, 36

Fig. 5 (a) Stress-strain profiles for BPE/UPyMA-x. (b) Loading-unloading curves for BPE/UPyMA-x. (c) Loading-unloading cycles for BPE/UPyMA-20. (d) Stress-relaxation curves for BPE/UPyMA-x.

Fig. 6 (a) BPE/UPyMA-x is bonded by the covalent bonds and quadruple H-bonding. (b) Quadruple H-bonding can reversibly break and reform under external force. (c) Quadruple H-bonding can stabilize and release temporary shape through their reformation/dissociation under heat stimuli.

To disclose the sacrificial and reversible nature of physical cross-links based on UPy aggregations in BPE, loading-unloading tests were performed. Figure 5b is the loading-unloading curves of BPE/UPyMA-x with a strain of 100%. The hysteresis (area surrounded by loading-unloading curves) indicates the energy dissipation during stretching, which is associated with the fracture of the UPy aggregations. Compared with the small hysteresis loop area of BPE/UPyMA-0, BPE/UPyMA-x containing UPyMA exhibits more pronounced hysteresis, and hysteresis loop area consistently increases with the increase of UPyMA content (Figure S6). This finding indicates that the rupture of UPy aggregations causes large energy dissipation during stretching. Furthermore, cyclic tensile curves of BPE/UPyMA-20 are displayed in Figure 5c. After the first cycle, a notable residual strain is left and the hysteresis loop in the successive cycle becomes much smaller, which is because the temporarily re-formed H-bonding retards the resilience of covalent network. When the sample is settled at room temperature for a period of time before subsequent cycle, the elastic contraction imparted by covalent cross-links leads to the relaxation of temporarily re-formed H-bonding. As a result, the residual strain is decreased and the loading curve gradually approaches the first loading one with interval time. After keeping the sample at 80 °C for 180 s, the loading-unloading curve fully overlaps with the first cycle (Figure 5c). This is because the temporarily formed H-bonding that impedes the recovery of covalent network to the equilibrium state is dissociated at high

temperatures. The fully reversible feature suggests that the covalent network is reserved through the reversible rupture and reformation of H-bonding during deformation. Figure 5d is the stress relaxation curves for BPE/UPyMA-x. Compared with BPE/UPyMA-0, BPE/UPyMA-x containing UPyMA can relax the applied force much faster. The relaxation rate and relaxation degree are increased as UPyMA content increases, which provide implications that the H-bonding ruptures under loading. After a distinct reduction, the force is nearly invariable as the permanently crosslinked network can indefinitely keep the load. Due to the dynamic characteristics of H-bonding, the modulus of the BPE/UPyMA-x is thought to be strain rate-dependence. To explore this, the tensile curves for BPE/UPyMA-20 at various strain rates are illustrated in Figure 7a. Interestingly, in the intermediate strain rates between 0.003 and 0.35 s−1, the modulus gradually improves with strain rates. However, when the strain rates are lower than 0.003 s−1 or higher than 0.35 s−1, the modulus is strain rate independent (Figure 7b). These phenomena can be explained by the relaxation of reversible H-bonding. In the intermediate strain rate region, the fraction of un-relaxed H-bonding is increased with strain rates, namely, the fraction of the un-relaxed H-bonding at shorter time scale (higher strain rate) is higher than that at longer time scale (lower strain rate). Therefore, the contribution of H-bonding to the modulus increases with strain rate. However, at extremely high strain rate (>0.35 s−1), nearly all the H-bonding are un-relaxed. Consequently, the increase in the modulus becomes saturated and the modulus is nearly constant. Inversely, at very low strain rate (<0.003 s−1), almost all the H-bonding are relaxed at the timescale of tensile experiment and thus the modulus is unchanged. In addition, considering that the H-bonding provides additional energy-dissipating mechanism, the complete relaxation of H-bonding leads to the sample measured at very low strain rate having worse strength and extension ratio. To further disclose the dynamic nature of H-bonding, the reduced stress (f *) versus the reciprocal of extension ratio (λ) are plotted according to the well-known Mooney equation37 f * = σ / (λ - λ-2)=2C1 λ-1+2C2

where σ represents stress, C1 and C2 are constants. As shown in Figure 7c, f * of BPE/UPyMA-0 is a constant irrespective of λ, suggesting that in the absence of H-bonding, the network is featured with Gaussian chains and its uniaxial tensile can be described using rubber elasticity model. In the case of BPE/UPyMA-20, f * decreases with strain at small λ (1/λ > 0.5), which is caused by the breakage of UPy aggregations. However, at large λ (1/λ < 0.5), f * slightly increases with strain. This strain stiffening effect is due to the finite extensibility of polymer networks. By plotting f * versus stretching time, it can be seen that the curves are overlapped except for the strain stiffening under longer time regions (Figure S7). Considering that f (t)* is a sum contribution of chemical cross-link (fc*) and physical cross-link (fp*), a master curve of fp (t)* against time can be obtained by deducting the contribution of chemical cross-link (here 0.75 MPa determined from the value of f * with a strain rate of 0.0001 s-1 and λ-1 of 0.5). As shown in Figure 7d, the fp* in the overlapped zone is independent of λ but relies on time, suggesting that strain does not affect the breakage rate of H-bonding. In addition, the slope of the overlapped zone is -0.5, which is in accordance with the associative Rouse mode concerning relaxation caused by H-bonding dynamics. 37, 38

Fig. 7 (a) Stress-strain profiles and (b) modulus for BPE/UPyMA-20 at various strain rates. (c) f * versus λ−1 for BPE/UPyMA-20 at different strain rates and BPE/UPyMA-0 at 0.2 s−1 strain rate. (d) Time dependence of fp* for BPE/UPyMA-20.

Adaptive recovery of BPE/UPyMA-x

The dynamic features of physical interactions can be used to design macroscopically responsive polymers such as shape memory polymer that can memory the temporary shape and restore the permanent shape under external stimulus.39 In this context, a typical shape memory cycle of BPE/UPyMA-30 is depicted in Figure 8a. Under a tensile load of 0.8 N, the sample is stretched to 32% strain at 90 °C followed by cooling to 20 °C. During stretching at 90 °C, H-bonding are dissociated. As the temperature decreases to 20 °C, the free UPy groups re-form dimers and the dynamic of H-bonding is sluggish, which allows the sample keeping strained state to access temporary shape. Upon releasing external force, the strain is immediately decreased from 32% to 25% due to the partial dissociation of H-bonding caused by the entropic driving force of the permanent network. Accordingly, the shape fixing and recovery ratio are calculated to be about 80% and 95%, which can be comparable to previous literatures that also employ physical interactions to stabilize temporary shape.40 Finally, when the sample is heated to 90 °C, the H-bonding are dissociated to release the temporary shape, which enables rapid and complete recovery of the strain. Figure 8b displays the shape recovery process of BPE/UPyMA-30. It can be seen that the programmed spiral shape can restore the permanent linear shape within 30 s at 90 °C.

Fig 8. (a) Shape memory cycle curves for BPE/UPyMA-30. (b) Photos showing the shape recovery process of BPE/UPyMA-30. The sample is deformed at 90 °C and cooled to 20 °C to maintain the temporary spiral shape, which can recover the permanent linear shape by re-heating to 90 °C.

Conclusion In conclusion, we report a simple way to improve the mechanical properties of as-synthesized biobased elastomer and bestow it with adaptive performance by incorporating quadruple H-bonding motifs into the elastomer network. SAXS result reveals that the aggregation of UPy groups results in the formation of microphase

separation. According to DMA results, the introduction of UPy units restricts bulk segmental relaxation and a new relaxation process related to the dissociation of UPy aggregates is observed. UPy aggregations can serve as physical cross-links and act in sacrificial manner through reversible rupture and reformation event. As a result, the modulus and tensile strength of the biobased elastomer are significantly improved without compromising the extensibility. In addition, the dynamic nature of UPy aggregations allows the elastomer to access thermo-activated shape memory behavior. We envision that this work presents a successful paradigm to prepare mechanically robust and adaptive biobased elastomers by incorporating transient yet sacrificial bonds into the networks. Supporting Information. Molecular structure and 1H NMR of BPE, FTIR and 1H NMR of UPyMA, DSC curves of BPE/UPyMA-0 and BPE/UPyMA-30, swelling experiments

of

BPE/UPyMA-10,

hysteresis

energy

BPE/UPyMA-20,

time

dependence of f * for BPE/UPyMA-20 are listed in the Supporting Information.

Corresponding Author * Z. Tang. E-mail: [email protected]

* B. Guo. E-mail: [email protected] Notes The authors declare no competing financial interest.

Acknowledgment

This work was funded by the National Key Research and Development Program of China

(2017YFB0306900

(2017YFB0306905)),

China

National

Funds

for

Distinguished Young Scientists (51825303) and National Natural Science Foundation of China (51790503, 51703064, and 51673065). References [1]. Wilbon, P. A.; Chu, F.; Tang, C., Progress in renewable polymers from natural terpenes, terpenoids, and rosin. Macromol. Rapid Commun. 34 (2013) 8-37. [2]. Gandini, A., Polymers from renewable resources: a challenge for the future of macromolecular materials. Macromolecules 41 (2008) 9491-9504. [3]. Wang, S.; Lu, A.; Zhang, L., Recent advances in regenerated cellulose materials. Prog. Polym. Sci. 53 (2016) 169-206. [4]. Kawaguchi, H.; Ogino, C.; Kondo, A., Microbial conversion of biomass into bio-based polymers. Bioresour. Technol. 245 (2017) 1664-1673. [5]. Meier, M. A.; Metzger, J. O.; Schubert, U. S., Plant oil renewable resources as green alternatives in polymer science. Chem. Soc. Rev. 36 (2007) 1788-1802. [6]. Gandini, A.; Lacerda, T. M.; Carvalho, A. J. F.; Trovatti, E., Progress of Polymers from Renewable Resources: Furans, Vegetable Oils, and Polysaccharides. Chem. Rev. 116 (2016) 1637-1669. [7]. Sheldon, R. A., Green and sustainable manufacture of chemicals from biomass: state of the art. Green Chem. 16 (2014) 950-963. [8]. Wang, Z. K.; Yuan, L.; Tang, C. B., Sustainable Elastomers from Renewable Biomass. Acc. Chem. Res. 50 (2017) 1762-1773. [9]. Lei, W.; Zhou, X.; Russell, T. P.; Hua, K.-c.; Yang, X.; Qiao, H.; Wang, W.; Li, F.; Wang, R.; Zhang, L., High performance bio-based elastomers: energy efficient and sustainable materials for tires. J. Mater. Chem. A. 4 (2016) 13058-13062. [10]. Guo, B.; Chen, Y.; Lei, Y.; Zhang, L.; Zhou, W. Y.; Rabie, A. B. M.; Zhao, J., Biobased poly (propylene sebacate) as shape memory polymer with tunable switching temperature for potential biomedical applications. Biomacromolecules 12 (2011) 1312-1321. [11]. Wang, R.; Zhang, J.; Kang, H.; Zhang, L., Design, preparation and properties of bio-based elastomer composites aiming at engineering applications. Compos. Sci. Technol. 133 (2016) 136-156.

[12]. Tang, Z.; Kang, H.; Shen, Z.; Guo, B.; Zhang, L.; Jia, D., Grafting of Polyester onto

Graphene

for

Electrically

and

Thermally

Conductive

Composites.

Macromolecules 45 (2012) 3444-3451. [13]. Domun, N.; Hadavinia, H.; Zhang, T.; Sainsbury, T.; Liaghat, G.; Vahid, S., Improving the fracture toughness and the strength of epoxy using nanomaterials–a review of the current status. Nanoscale 7 (2015) 10294-10329. [14]. Rief, M.; Gautel, M.; Oesterhelt, F.; Fernandez, J. M.; Gaub, H. E., Reversible unfolding of individual titin immunoglobulin domains by AFM. Science 276 (1997) 1109-1112. [15]. Thompson, J. B.; Kindt, J. H.; Drake, B.; Hansma, H. G.; Morse, D. E.; Hansma, P. K., Bone indentation recovery time correlates with bond reforming time. Nature 414 (2001) 773-776. [16]. Zhou, X.; Guo, B.; Zhang, L.; Hu, G. H., Progress in bio-inspired sacrificial bonds in artificial polymeric materials. Chem. Soc. Rev. 46 (2017) 6301-6329. [17]. Sun, T. L.; Kurokawa, T.; Kuroda, S.; Bin Ihsan, A.; Akasaki, T.; Sato, K.; Haque, M. A.; Nakajima, T.; Gong, J. P., Physical hydrogels composed of polyampholytes demonstrate high toughness and viscoelasticity. Nat. Mater. 12 (2013) 932-937. [18]. Li, C. H.; Wang, C.; Keplinger, C.; Zuo, J. L.; Jin, L.; Sun, Y.; Zheng, P.; Cao, Y.; Lissel, F.; Linder, C.; You, X. Z.; Nan, B. Z., A highly stretchable autonomous self-healing elastomer. Nat. Chem. 8 (2016) 618-624. [19]. Ducrot, E.; Chen, Y. L.; Bulters, M.; Sijbesma, R. P.; Creton, C., Toughening Elastomers with Sacrificial Bonds and Watching Them Break. Science 344 (2014) 186-189. [20]. Huang, J.; Zhang, L. J.; Tang, Z. H.; Wu, S. W.; Ning, N. Y.; Sun, H. B.; Guo, B. C., Bioinspired Design of a Robust Elastomer with Adaptive Recovery via Triazolinedione Click Chemistry. Macromol. Rapid Commun. 38 (2017) 1600678. [21]. McKee, J. R.; Huokuna, J.; Martikainen, L.; Karesoja, M.; Nykanen, A.; Kontturi, E.; Tenhu, H.; Ruokolainen, J.; Ikkala, O., Molecular Engineering of Fracture Energy Dissipating Sacrificial

Bonds

Into Cellulose Nanocrystal

Nanocomposites. Angew. Chem. Int. Edit. 53 (2014) 5049-5053. [22]. Wu, J. R.; Cai, L. H.; Weitz, D., Tough Self-Healing Elastomers by Molecular Enforced Integration of Covalent and Reversible Networks. Adv. Mater. 29 (2017) 1702616.

[23]. Degtyar, E.; Harrington, M. J.; Politi, Y.; Fratzl, P., The Mechanical Role of Metal Ions in Biogenic Protein-Based Materials. Angew. Chem. Int. Edit. 53 (2014) 12026-12044. [24]. Neal, J.; Mozhdehi, D.; Guan, Z., Enhancing Mechanical Performance of a Covalent Self-Healing Material by Sacrificial Non-Covalent Bonds. J. Am. Chem. Soc. 137 (2015) 4846-4850. [25]. Tang, Z.; Huang, J.; Guo, B.; Zhang, L.; Liu, F., Bioinspired Engineering of Sacrificial Metal-Ligand Bonds into Elastomers with Supramechanical Performance and Adaptive Recovery. Macromolecules 49 (2016) 1781-1789. [26]. Kumpfer, J. R.; Rowan, S. J., Thermo-, Photo-, and Chemo-Responsive Shape-Memory

Properties

from

Photo-Cross-Linked

Metallo-Supramolecular

Polymers. J. Am. Chem. Soc. 133 (2011) 12866-12874. [27]. Roy, N.; Bruchmann, B.; Lehn, J. M., DYNAMERS: dynamic polymers as self-healing materials. Chem. Soc. Rev. 44 (2015) 3786-3807. [28]. Kushner, A. M.; Vossler, J. D.; Williams, G. A.; Guan, Z. B., A Biomimetic Modular Polymer with Tough and Adaptive Properties. J. Am. Chem. Soc. 131 (2009) 8766-8768. [29]. Wei, T.; Lei, L.; Kang, H.; Qiao, B.; Wang, Z.; Zhang, L.; Coates, P.; Hua, K. C.; Kulig, J., Tough Bio‐Based Elastomer Nanocomposites with High Performance for Engineering Applications. Adv. Eng. Mater. 14 (2012) 112-118. [30]. Berda, E. B.; Foster, E. J.; Meijer, E., Toward controlling folding in synthetic polymers: fabricating and characterizing supramolecular single-chain nanoparticles. Macromolecules 43 (2010) 1430-1437. [31]. Liu, Y.; Tang, Z.; Wu, S.; Guo, B., Integrating Sacrificial Bonds into Dynamic Covalent Networks toward Mechanically Robust and Malleable Elastomers. ACS. Macro. Lett. 8 (2019) 193-199. [32]. Sontjens, S. H. M.; Renken, R. A. E.; van Gemert, G. M. L.; Engels, T. A. P.; Bosman, A. W.; Janssen, H. M.; Govaert, L. E.; Baaijens, F. P. T., Thermoplastic elastomers based on strong and well-defined hydrogen-bonding interactions. Macromolecules 41 (2008) 5703-5708. [33]. Appel, W. P. J.; Portale, G.; Wisse, E.; Dankers, P. Y. W.; Meijer, E. W., Aggregation of Ureido-Pyrimidinone Supramolecular Thermoplastic Elastomers into Nanofibers: A Kinetic Analysis. Macromolecules 44 (2011) 6776-6784.

[34]. Kan, L.; Zhang, P.; Jiang, H.; Zhang, S.; Liu, Z.; Zhang, X.; Ma, N.; Qiu D.; Wei H., Microphase separation of a quadruple hydrogen bonding supramolecular polymer: effect of the steric hindrance of the ureido-pyrimidone on their viscoelasticity. RSC Adv. 16 (2019), 8905-8911. [35]. Ware, T.; Hearon, K.; Lonnecker, A.; Wooley, K. L.; Maitland, D. J.; Voit, W., Triple-shape memory polymers based on self-complementary hydrogen bonding. Macromolecules 45 (2012) 1062-1069. [36]. Becker, N.; Oroudjev, E.; Mutz, S.; Cleveland, J. P.; Hansma, P. K.; Hayashi, C. Y.; Makarov, D. E.; Hansma, H. G., Molecular nanosprings in spider capture-silk threads. Nat. Mater. 2 (2003) 278-283. [37]. Mayumi, K.; Marcellan, A.; Ducouret, G.; Creton, C.; Narita, T., Stress–strain relationship of highly stretchable dual cross-link gels: separability of strain and time effect. ACS. Macro. Lett. 2 (2013) 1065-1068. [38]. Narita, T.; Mayumi, K.; Ducouret, G.; Hébraud, P., Viscoelastic Properties of Poly(vinyl alcohol) Hydrogels Having Permanent and Transient Cross-Links Studied by

Microrheology,

Classical

Rheometry,

and

Dynamic

Light

Scattering.

Macromolecules 46 (2013) 4174-4183. [39]. Wojtecki, R. J.; Meador, M. A.; Rowan, S. J., Using the dynamic bond to access macroscopically responsive structurally dynamic polymers. Nat. Mater. 10 (2011) 14-27. [40]. Li, J. H.; Viveros, J. A.; Wrue, M. H.; Anthamatten, M., Shape-memory effects in polymer networks containing reversibly associating side-groups. Adv. Mater. 19 (2007) 2851-2855.

Declaration of interests ☐ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:

There is no conflict of interest.