Self-Healable Covalently Adaptable Networks Based on Disulfide Exchange

Introducing dynamic covalent bonding into thermoset polymers has received considerable attention because they can repair or recover when damaged, thereby minimizing waste and extending the service life of thermoset polymers. However, most of the yielded dynamic covalent bonds require an extra catalyst, high temperature and high-pressure conditions to trigger their self-healing properties. Herein, we report on a catalyst-free bis-dynamic covalent polymer network containing vinylogous urethane and disulfide bonds. It is revealed that the introduction of disulfide bonds significantly reduces the activation energy (reduced from 94 kJ/mol to 51 kJ/mol) of the polymer system for exchanging and promotes the self-healing efficiency (with a high efficiency of 86.92% after being heated at 100 °C for 20 h) of the material. More importantly, the mechanical properties of the healed materials are comparable to those of the initial ones due to the special bis-dynamic covalent polymer network. These results suggest that the bis-dynamic covalent polymer network made of disulfide and inter-vinyl ester bonds opens a new strategy for developing high-performance vitrimer polymers.


Introduction
Traditional thermosets typically exhibit excellent mechanical properties, robust chemical resistance and credible structural stability because of the three-dimensional (3D) permanently cross-linked network structure, and they are widely applied in many fields, such as insulation and packaging materials, textiles, diaphragms and sealants [1]. However, the 3D cross-linked network formed in thermoset polymers renders them infusible and insoluble, hindering their repair and recovery when damaged; thus, the polymers exhibit low reprocessability and recyclability, resulting in significant waste [2]. To overcome this disadvantage, minimize waste and extend the service life of thermoset polymers, a new material that can be repaired while maintaining the properties imparted by the cross-linked network structure is urgently required.
Introducing dynamic covalent bonding into thermoset polymers can be a promising strategy to address these challenges. The dynamic covalent bonding, where the covalent bonds can be reversibly broken-reconfigured in response to external stimuli (e.g., light [3], heat [4,5], pH [6] and chemical stimuli [7]), rearranges the formed cross-linked network, making the network reproducible. In 2011, Leibler et al. [8] first introduced dynamic epoxy-ester bonds into the cross-linked polymer network and demonstrated the recyclability of the thermoset polymers through multiple heat treatments, terming it as a vitrimer. Thereafter, a variety of dynamic covalent bonds were introduced into cross-linked networks [9], including allyl sulfide-thiol exchange [10], boronic ester and oxime ester [11,12], transesterification of (thio) ester [13], transesterification of silyl ether [14], transalkylation The synthesis of monomer TMPTAA and PETMPA was performed as follows: TMPTA (or PETMP), UV1173 (5 wt.%), and acetone (30 mL) were added to a flask wrapped in tin foil and stirred under nitrogen protection for 10 min. Afterwards, allyl acetoacetate was added into the mixture under a nitrogen atmosphere and stirred for 5 min. Thereafter, the flask was irradiated under an ultraviolet (UV) lamp (8 mW, 365 nm) for 8 h. At the end of the reaction, the solvent was removed and a yellow oil liquid was obtained.

Synthesis of Cystamine (Cys)
Cystamine dihydrochloride (8.00 g, 35.5 mmol) and water (80 mL) were added to a 250-mL beaker. After dissolving, potassium hydroxide (KOH) (6.00 g, 107.0 mmol) was added, and the mixture was stirred for 1 h. The resulting mixture was extracted with 4 × 150 mL of dichloromethane; the collected organic layer was passed through anhydrous MgSO 4 and dried via spin evaporation at 45 • C to afford 4.3 g of light-yellow oil (yield 80%).

Synthesis of Polymer Films
The synthesis route and composition of the polymer films are shown in Scheme 1 and Table 1, and the experimental procedure is as follows: TMPTAA (or PETMPA), amine curing agent and acetone (6 mL) were added to a 10-mL weighing bottle and stirred for 5 min. Afterwards, the mixture was poured into a polytetrafluoroethylene (PTFE) mold (8 × 8 × 2 cm) and cured at 30 • C for 12 h. Finally, the polymer films with a thickness of 500 µm were obtained.

Characterization
Nuclear Magnetic Resonance (NMR) Spectroscopy. 1 H NMR (400 MHz) and 13 C NMR (100 MHz) spectra were recorded by Bruker AV-400 spectrometer. Chemical shift (δ) was expressed in parts per million (ppm), and each sample was tested with CDCl 3 as a solvent.
FT-IR Spectroscopy. The structure of prepolymers and polymers were characterized on a Nicolet iS20 FT-IR spectrometer from 4000 cm −1 to 500 cm −1 with setting resolution to 0.25 cm −1 , each sample was scanned 32 times using a diamond ATR probe.
Thermal Gravimetric Analysis (TGA). Thermogravimetric analysis (TGA) experiments were carried out on a TA-Q50 thermogravimetric analyzer, ramped up from 25 • C to 750 • C at a rate of 10 • C/min in a nitrogen atmosphere.
Differential Scanning Calorimetry (DSC). DSC experiments were performed using DSC-Q2000 under a nitrogen atmosphere from −40 • C to 120 • C at a rate of 20 • C/min. Dynamic Mechanical Analysis (DMA). Dynamic Mechanical Analysis (DMA) was conducted with a TA Q800 (New Castle, DE, USA) instrument. Each sample had a size of a 7 mm × 3 mm × 0.5 mm rectangular shape. For the DMA experiments, a preloaded axial force of 0.01 N, a displacement amplitude of 10 µm, and a regular frequency of 1 Hz were preloaded. Each sample was heated to 180 • C at −3 • C/min at −20 • C. The tensile test was performed using DMAQ800 and the applied force was increased with a rate of 3 N/min to 18 N. Stress Relaxation Analysis (SRA) experiments were conducted at the setting strain of 1% and a temperature range of 60 • C-130 • C. All of the DMA experiments results were collected through TA's Thermal Advantage for Q Series software New Castle, DE, USA.
Self-Healing Experiments. The film (long × wide × high= 7 mm × 3 mm × 0.5 mm) was cut in half using a knife, put into contact without further pressure and placed in a 100 • C environment for self-healing experiments. The optical micrographs of the self-healed samples were obtained during the healing process using a digital camera attached to a Meiji optical microscope. The tensile test was performed using the DMAQ800, the self-healing efficiency was measured and were as detailed in the previous section. All measurements were averaged by at least four measurements. The self-healing efficiency (η) of the polymer was determined by the ratio of the tensile strength of the self-healed sample (σ healed ) to that of the initial sample (σ initial ), and the self-healing efficiency (η) was calculated as follows [31]: Swelling experiments. The swelling ratio (SR) of the polymer films was detected as follows. A dried polymer film (2 × 1 cm 2 ) was first weighed and recorded as W 1 and immersed in an Ace solvent for several hours. The weight of the wet polymer film was recorded as W 2 after quickly wiping its surface with a filter paper. The SR of the films in the Ace solvent was calculated as follows: M (%) = [(W 2 − W 1 )/W 1 ] × 100%. The gel content (GC) was determined by cutting a polymer film of mass W 1 , soaking it in the Ace solvent for 48 h, and drying it at 80 • C for 48 h; the sample weight was denoted as W 3 . The GC was obtained with the following equation: M (%) = W 3 /W 1 × 100%. The cross-link density (v e ) of the polymer film was calculated using the Flory-Rhener relationship [40][41][42][43]: where v s is the molar volume of the solvent; v 2 is the volume fraction of the swollen polymer; and χ is the polymer-solvent interaction parameter. The volume fraction (v 2 ) of the dissolved polymer was calculated according to the following formula: where ρ acetone is 0.788 g/cm; and ρ polymer is the density of the polymer. The polymer-solvent interaction parameter χ was calculated as follows: where R is the gas constant (8.314 J/K/mol); and T is the absolute temperature (273 K).
Additionally, for acetone, v s is 73 cm 3 /mol, δ polymer is the solubility parameter of the polymer, δ acetone is the solubility parameter of acetone and δ acetone is 20.4 (J/cm 3 ) 0.5 , while δ polymer can be determined by the solubility method.

Synthesis of Acetoacetic Acid Esterified Functional Monomers (TMPTAA and PETMPA)
TMPTAA was synthesized from TMPTA and allyl acetoacetate by the thiol-click reaction under 365 nm UV light. This preparation method is a highly efficient, direct and fast method to obtain the functional monomers of acetoacetate without by-products. Figure 1 shows the 1 H NMR spectra of the monomers. For TMPTAA (Figure 1a), the -CH 2 -proton peak h' next to the thiol group of TMPTA shifted from 1.59 to 2.62 ppm (proton peak h, -CH 2 -group adjacent to the thioether), and the -CH 2 -proton peak c' of AER shifted from 4.63 to 4.23 ppm (proton peak c, -CH 2 -group adjacent to the ether) after the thiol-click reaction. Combined with the disappearance of the proton peaks e' and d' assigned to the double bond, the TMPTAA was successfully prepared. The same chemical shift change is observed in Figure 1b, implying that the monomer PETMPA was obtained. thiol-click reaction. Combined with the disappearance of the proton peaks e' and d' assigned to the double bond, the TMPTAA was successfully prepared. The same chemical shift change is observed in Figure 1b, implying that the monomer PETMPA was obtained. Further, the 13 C NMR spectra of TMPTAA and PETMPA is shown in Figure 2. After the thiol-click reaction between AER and TMPTA, the peak position changed. Specifically, the C=C bond signal peaks g' and f' of AER at 131 ppm and 181 ppm disappeared, while the methylene signal peak e' adjacent to the AER ester group at 65 ppm shifted to 63 ppm. Further, the methylene signal peak h' next to the TMPTA thiol group at 19 ppm shifted to 28 ppm, and the methylene signal peak i' near the TMPTA carbonyl group at 38 ppm Further, the 13 C NMR spectra of TMPTAA and PETMPA is shown in Figure 2. After the thiol-click reaction between AER and TMPTA, the peak position changed. Specifically, the C=C bond signal peaks g' and f' of AER at 131 ppm and 181 ppm disappeared, while the methylene signal peak e' adjacent to the AER ester group at 65 ppm shifted to 63 ppm. Further, the methylene signal peak h' next to the TMPTA thiol group at 19 ppm shifted to 28 ppm, and the methylene signal peak i' near the TMPTA carbonyl group at 38 ppm shifted to 35 ppm. Similar changes in the NMR signal peaks occurred in PETMPA. These results showed that the thiol-click reaction was performed, and the acetoacetate-functionalized TMPTAA and PETMPA were prepared.

Synthesis and Characterization of the Polymer Films
The chemical structure of the TMPTAA, PETMPA and four polymer networks we characterized by FT-IR and are shown in Figure 3. Although all films showed similar cha acteristic IR absorption peaks, there were changes in the density and position of sever absorption peaks, when compared with monomers. Typically, the peak at 1736 cm −1 (C= stretching vibration in TMPTAA and PETMPA) was dramatically weakened, and new a sorption peaks appeared at 1650 cm −1 (C=C stretching vibration) and 1598 cm −1 (C= stretching vibration) [32,[44][45][46]. Further, the detection of the cross-link structure was ve ified by acetone swelling experiments, where ca. 0.1 g of the polymer network was su merged in acetone, and the swelling rate of the film was measured over time (Figure S The films were not dissolved, only swollen, after being immersed for 48 h (see Table  demonstrating that the crosslinking films were successfully prepared.

Synthesis and Characterization of the Polymer Films
The chemical structure of the TMPTAA, PETMPA and four polymer networks were characterized by FT-IR and are shown in Figure 3. Although all films showed similar characteristic IR absorption peaks, there were changes in the density and position of several absorption peaks, when compared with monomers. Typically, the peak at 1736 cm −1 (C=O stretching vibration in TMPTAA and PETMPA) was dramatically weakened, and new absorption peaks appeared at 1650 cm −1 (C=C stretching vibration) and 1598 cm −1 (C=O stretching vibration) [32,[44][45][46]. Further, the detection of the cross-link structure was verified by acetone swelling experiments, where ca. 0.1 g of the polymer network was submerged in acetone, and the swelling rate of the film was measured over time ( Figure S1). The films were not dissolved, only swollen, after being immersed for 48 h (see Table 2), demonstrating that the crosslinking films were successfully prepared.    Figure 4a shows the mechanical properties of the polymer films. The mechanical properties gradually increased with the increase in the cross-link density ( Figure S1; Tables 2 and 3). Crosslinking density showed a trend of P3 > P4 > P1 > P2, and the Young's Modulus from 6.0 MPa (P2) to 31.3 MPa (P3). All of the tensile data are recorded in Table  3 and it was observed that the Young's Modulus increased as the crosslinking density increased with more acetoacetate-functional groups in the crosslinking substrate (P3 > P1 and P4 > P2). Moreover, the films prepared from 1,6-hexanediamine posed better Young's Modulus than the cystamine (P1 > P2 and P3 > P4). Figure 4b shows the DSC curves of the films; the Tg of the films were in the order of P3 (25 °C) > P4 (22 °C) > P1 (19 °C) > P2 (11 °C). Note that the films prepared with cystamine had a lower Tg than those made with 1,6hexanediamine, while those synthesized with trifunctional TMPTA had a lower Tg than those with PETMPA. The increased cross-link density of the polymer films resulted in a corresponding increase in Tg. The TGA curves showed the thermal stability of the films, as shown in Figure 4c, where the cystamine-prepared films, P2 and P4, exhibited a 10% weight loss at 210 °C, while the 1,6-hexanediamine-prepared films, P1 and P3, exhibited 10% weight loss at 250 °C. The relatively low thermal-induced weight loss temperature of the P2 and P4 films was attributed to the unstable breakage of the disulfide bond at 210 °C.   Table 3 and it was observed that the Young's Modulus increased as the crosslinking density increased with more acetoacetate-functional groups in the crosslinking substrate (P3 > P1 and P4 > P2). Moreover, the films prepared from 1,6-hexanediamine posed better Young's Modulus than the cystamine (P1 > P2 and P3 > P4). Figure 4b shows the DSC curves of the films; the T g of the films were in the order of P3 (25 • C) > P4 (22 • C) > P1 (19 • C) > P2 (11 • C). Note that the films prepared with cystamine had a lower T g than those made with 1,6-hexanediamine, while those synthesized with trifunctional TMPTA had a lower T g than those with PETMPA. The increased cross-link density of the polymer films resulted in a corresponding increase in T g . The TGA curves showed the thermal stability of the films, as shown in Figure 4c, where the cystamine-prepared films, P2 and P4, exhibited a 10% weight loss at 210 • C, while the 1,6-hexanediamine-prepared films, P1 and P3, exhibited 10% weight loss at 250 • C. The relatively low thermal-induced weight loss temperature of the P2 and P4 films was attributed to the unstable breakage of the disulfide bond at 210 • C.
The dynamic mechanical properties of the films were investigated by DMA experiments. As shown in Figure 4d, all of the samples showed an obvious rubber platform after the glass transition. We calculated the crosslinking density of each film by using the cross-linking density calculation formula: v e = E'/3RT [47], where E' represents the energy storage modulus under T g + 50 • C, R refers to the gas constant, and T is T g + 50 • C. T g is closely correlated with the cross-linking density and transferred from 30 • C to 46 • C, with T g increasing with the number of prepolymer acetoacetate groups, and lower cystamineprepared films compared to 1,6-diaminohexane-prepared films, which is consistent with the results obtained by DSC (with the same trend as T g ). This suggests that the mechanical properties of the film can be effectively adjusted according to the amount of regulated prepolymer acetoacetate groups of flexible chain-segment disulfide monomers. In addition, the addition of cystamine reduces the mechanical properties of the film.   The dynamic mechanical properties of the films were investigated by DMA experiments. As shown in Figure 4d, all of the samples showed an obvious rubber platform after the glass transition. We calculated the crosslinking density of each film by using the crosslinking density calculation formula: ve = E'/ 3RT [47], where E' represents the energy storage modulus under Tg + 50 °C, R refers to the gas constant, and T is Tg + 50 °C. Tg is closely correlated with the cross-linking density and transferred from 30 °C to 46 °C, with Tg increasing with the number of prepolymer acetoacetate groups, and lower cystamine-prepared films compared to 1,6-diaminohexane-prepared films, which is consistent with the results obtained by DSC (with the same trend as Tg). This suggests that the mechanical properties of the film can be effectively adjusted according to the amount of regulated prepolymer acetoacetate groups of flexible chain-segment disulfide monomers. In addition, the addition of cystamine reduces the mechanical properties of the film.

Stress-Relaxation Experiments
Stress-relaxation measurement was taken to investigate the viscoelasticity of the polymer films and the rate of exchange reactions within the material. The relaxation time τ is defined as the time taken for the material to relax from the initial modulus to 1/e of the

Stress-Relaxation Experiments
Stress-relaxation measurement was taken to investigate the viscoelasticity of the polymer films and the rate of exchange reactions within the material. The relaxation time τ is defined as the time taken for the material to relax from the initial modulus to 1/e of the initial modulus. Figure 5a-d shows the stress-relaxation curves for the polymer films. The relaxation time for P1 decreased from 1997 s to 343 s as the test temperature increased from 60 • C to 90 • C (Figure 5a), while the relaxation time for P2 decreased from 1211 s to 250 s (Figure 5b). Additionally, as the test temperature increased from 100 • C to 130 • C, the relaxation time decreased from 3487 s to 373 s for P3 (Figure 5c), and from 1110 s to 174 s for P4 (Figure 5d). To deeply study the energy required for the macroscopic flow of the material, the activation energy (E a ) of polymer films was calculated using the Arrhenius equation: τ = τ 0 × exp (E a /RT) [48][49][50][51] and is shown in Figure 5c,d. The Ea of the polymer films was in the following order: P3 (94 kJ/mol) > P4 (72 kJ/mol) > P1 (61 kJ/mol) > P2 (51 kJ/mol), demonstrating that the disulfide bond decreased the activation energy for the same type of film system.  (Figure 5b). Additionally, as the test temperature increased from 100 °C to 130 ° the relaxation time decreased from 3487 s to 373 s for P3 (Figure 5c), and from 1110 s 174 s for P4 (Figure 5d). To deeply study the energy required for the macroscopic flow the material, the activation energy (Ea) of polymer films was calculated using the Arrh nius equation: τ = τ0 × exp (Ea /RT) [48][49][50][51] and is shown in Figure 5c,d. The Ea of t polymer films was in the following order: P3 (94 kJ/mol) > P4 (72 kJ/mol) > P1 (61 kJ/m > P2 (51 kJ/mol), demonstrating that the disulfide bond decreased the activation ener for the same type of film system.

Self-Healing Property of the Polymer Films
To determine the effect of the disulfide in the film system on the self-healing properties of the films, the strip samples were cut and heated in a 100 • C oven for 2 h, 8 h and 20 h. Thereafter, their morphological changes during the repair process were visualized using an optical microscope to ascertain the variation in the self-healing abilities over time. Figure 6a,b, depict the photographs of healed P2 and P1, respectively. The scratch on the surface of P2 was visible when it was first cut off, and became smaller as the healing time increased. When the healing time reached 20 h, most scratches on the P2 surface disappeared, and healed P2 could suspend a 500 g weight (Figure 6c). However, no significant scratch healing was observed after P1 was treated at 100 • C for 20 h. These results demonstrated that the introduction of the disulfide bond results in the excellent self-healing ability of the films. Figure 6a,b, depict the photographs of healed P2 and P1, respectively. The scratch on the surface of P2 was visible when it was first cut off, and became smaller as the healing time increased. When the healing time reached 20 h, most scratches on the P2 surface disappeared, and healed P2 could suspend a 500 g weight (Figure 6c). However, no significant scratch healing was observed after P1 was treated at 100 °C for 20 h. These results demonstrated that the introduction of the disulfide bond results in the excellent self-healing ability of the films. To further investigate the self-healing properties and healing mechanisms of the polymer films, their self-healing efficiency was quantified by stress-strain experiments. Four film samples were cut and healed by heating in an oven at 100 °C. Figure 7 and Table 4 show the self-healing efficiency of P1, P2, P3 and P4 healed at 100 °C at different times. A significant difference was observed between the self-healing efficiency of the films containing the disulfide bonds and those without the disulfide bonds at different healing times. With the extension of the healing time from 2 h to 20 h, the molecular chains diffused into the damaged area more deeply, facilitating the reorganization of dissociated disulfide and hydrogen bonds [21]. In particular, the P4 film healed at 100 °C for 20 h with a high self-healing efficiency of 86.92%, which was 50.57% higher than that for 2 h. Moreover, P4 showed the highest stress recovery in the same healing time, although the difference in the self-healing efficiency between P2 and P4 was only 0.67% (20 h healing). As shown in Figure 7a-d, the self-healing efficiency of P4 and P2 showed an increasing trend, while P1 and P3 were not significantly healed for 20 h. This indicates that the introduction of the disulfide bonds enhanced the movement of the elastomer molecular chain segments and increased the self-healing efficiency [43,52]. To further investigate the self-healing properties and healing mechanisms of the polymer films, their self-healing efficiency was quantified by stress-strain experiments. Four film samples were cut and healed by heating in an oven at 100 • C. Figure 7 and Table 4 show the self-healing efficiency of P1, P2, P3 and P4 healed at 100 • C at different times. A significant difference was observed between the self-healing efficiency of the films containing the disulfide bonds and those without the disulfide bonds at different healing times. With the extension of the healing time from 2 h to 20 h, the molecular chains diffused into the damaged area more deeply, facilitating the reorganization of dissociated disulfide and hydrogen bonds [21]. In particular, the P4 film healed at 100 • C for 20 h with a high self-healing efficiency of 86.92%, which was 50.57% higher than that for 2 h. Moreover, P4 showed the highest stress recovery in the same healing time, although the difference in the self-healing efficiency between P2 and P4 was only 0.67% (20 h healing). As shown in Figure 7a-d, the self-healing efficiency of P4 and P2 showed an increasing trend, while P1 and P3 were not significantly healed for 20 h. This indicates that the introduction of the disulfide bonds enhanced the movement of the elastomer molecular chain segments and increased the self-healing efficiency [43,52].

Conclusions
In conclusion, a series of novel polymer films containing vinylogous urethane and disulfide bonds were developed from multifunctional acetoacetate (TEMPTAA and PETMPA) and modified Cys. The experimental results indicated that the high functional group equipped with acetoacetate is good for improving the crosslinking density of the

Conclusions
In conclusion, a series of novel polymer films containing vinylogous urethane and disulfide bonds were developed from multifunctional acetoacetate (TEMPTAA and PETMPA) and modified Cys. The experimental results indicated that the high functional group equipped with acetoacetate is good for improving the crosslinking density of the polymer network. The stress-relaxation results showed that the Cys films prepared from the same functional monomer yielded relatively low activation energy and high chain segment movement because of the low crosslinking density resulting in low T g and Young's modulus. The equipped low-activation energy (e.g., P2: E a = 51 kJ/mol) ensured a fast exchange of the dynamic disulfide bonds. Thus, the disulfide-containing films presented excellent self-healing ability, with a repair efficiency of 86.92% for P4 heated at 100 • C for 20 h. Additionally, all of the polymer films containing vinylogous urethane double bonds exhibited excellent mechanical properties. It can be concluded that the strategy of introducing dynamic disulfide and vinylogous urethane bonds can provide excellent mechanical properties and self-healing abilities to the materials, which will provide a potential possibility for future research on vitrimer polymers.

Institutional Review Board Statement: Not applicable.
Data Availability Statement: All data generated or analyzed during this study are included in the published articles included in this article.

Conflicts of Interest:
The authors declare no conflict of interest.