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Towards mechanical robust yet self-healing polyurethane elastomers via combination of dynamic main chain and dangling quadruple hydrogen bonds
Polymer 183 (2019) 121912
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Towards mechanical robust yet self-healing polyurethane elastomers via combination of dynamic main chain and dangling quadruple hydrogen bonds Jin Hu a, Ruibin Mo a, Xiang Jiang a, Xinxin Sheng b, **, Xinya Zhang a, * a
School of Chemistry and Chemical Engineering, South China University of Technology, Guangzhou, 510640, China Guangdong Provincial Key Laboratory of Functional Soft Condensed Matter, Department of Polymeric Materials and Engineering, School of Materials and Energy, Guangdong University of Technology, Guangzhou, 510006, China
b
A R T I C L E I N F O
A B S T R A C T
Keywords: Self-healing Dynamic covalent bonds Quadruple hydrogen bonds
One of the greatest challenges of robust self-healing materials is the confliction between high chain mobility for self-healing and a stable structure for mechanical strength. Herein, dangling 2-ureido-4[1H]-pyrimidione (UPy)functionalized side groups were introduced into the hard segments of thermoplastic polyurethane (TPU) elas tomers, where embedded the dynamic disulfide bonds in the main chain. The strong quadruple H-bonding interaction between UPy side groups acts as supramolecular crosslinkers enabling the TPU elastomer to have improved mechanical properties (tensile strength up to 25 MPa and toughness ~100 MJ m 3), and simulta neously the plasticizer effect of dangling side chain endows it with efficient healing ability at elevated tem peratures (80–100 � C) comparable to its linear analogues. This strategy shows great potential in designing robust self-healing TPU elastomer employing weak dynamic covalent bonds, with wide promising applications as wearable electronics, coatings and adhesives.
1. Introduction Inspired by nature, self-healing materials capable of repairing dam ages and restoring their original properties show great potential in numerous applications, especially in the fields of protective coatings [1–3], wearable electronics [4–6] and biomedical materials [7–10]. Among various approaches to develop self-healing polymers, the intro duction of dynamic covalent bonds and non-covalent interactions to chain molecules, termed as intrinsic self-healing, is one of the most powerful approaches. Intrinsic self-healing material was first introduced by Chen et al. [11], made by reversible Diels-Alder (DA) reaction of tetra-maleimide and tetra-furan. After that, people have successfully fabricated self-healing polymers based on various dynamic covalent bonds, such as DA reaction [11], disulfide [12–14], diselenide [15–17] and ditelluride [18] metathesis, imine metathesis [19–21], olefin metathesis [22–24], hindered urea bonds [25–27], transesterification [28–30], trans carbonation [31] of carbonates, transcarbamoylation [32,33] of ure thanes, transamination [34,35] of vinylogous urethanes, dioxaborolane
metathesis [36–38] of boronic esters, siloxane equilibration [39–41], alkoxyamine metathesis [42–44] and so on. Supramolecular interactions are widely employed to construct self-healing materials as well, including hydrogen bonds [45–48], metal-ligand coordination [49–51], ionic interactions [52], π-π stacking [53–55], host-guest interactions [56–58] and even only van der Waals forces [59], etc. In addition to the dynamic bonding required for structure reorgani zation, chain diffusion plays an important role in the healing process as well. According to Wool’s model [60], a healing process goes through five phases: (i) surface rearrangement, (ii) surface approach, (iii) wet ting, (iv) diffusion and (v) randomization, during which diffusion is a critical process for crack healing. Thus, it requires the polymer chain to be soft enough or has relatively low molecular weight, which is the main reason of compromised mechanical properties for self-healing materials. In general, self-healing ability and mechanical robustness of the material are contradictory and difficult to be optimized simultaneously. Thermoplastic polyurethane (TPU) is a thermal processable and solvent dissolvable elastomer that is widely used in dozens of fields due to its high resilience and resistance to impacts, abrasions and tears. As is
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (X. Sheng), [email protected] (X. Zhang). https://doi.org/10.1016/j.polymer.2019.121912 Received 22 July 2019; Received in revised form 18 September 2019; Accepted 15 October 2019 Available online 17 October 2019 0032-3861/© 2019 Elsevier Ltd. All rights reserved.
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commonly known, TPU is a segmented block copolymer that contains a low Tg soft block and a physically crosslinked hard domain dispersed in the soft matrix. And its performance and morphology can be easily tuned by selecting suitable diisocyanates, polyols and chain extenders. Taking these characteristics of TPU, researchers have designed various selfhealing TPU elastomers by introducing dynamic covalent bonds and non-covalent interactions into TPU chains. Normally, to impart self-healing ability to TPU elastomers, dynamic covalent linkages are introduced into loosely packed TPU backbones in the form of chain extenders. As a result, intrinsic self-healing materials that can heal under mild conditions always have compromised me chanical properties [61–64], due to the weak nature of dynamic bonds and increased mobility of chain structure. Addition of covalent cross linking points can enhance the mechanical properties of PU elastomers, but in the sacrifice of efficient self-healing ability. Yang et al. [65] used hexamethylene diisocyanate (HDI) trimer as a crosslinker, and aromatic disulfide as a dynamic bond to fabricate a self-healing PU elastomer, finding that increasing HDI trimer content would suppress the self-healing ability. The ordering structure of hard domains with labile dynamic covalent bonds embedding in may affect the self-healing properties of TPU elastomers as well. Kim et al. [61] found that the self-healing efficiencies of the TPUs decreased in the following order: IPDI-SS > IPDI-EG � HMDI-SS > HDI-SS � MDI-SS. HDI and MDI based TPU elastomer cannot heal up to 80 � C due to their tightly packed and crystalline structures of hard domains, hindering the effective disulfide bond exchanging between different hard domains. The volume content of hard phases is another key factor, revealed by Lai [66] et al., who discovered that increased content of hard segments will enhance the mechanical properties. 2-ureido-4[1H]-pyrimidinone (or UPy) moiety shows a strong ten dency for self-complementary dimerization involving four hydrogen bonds in a donor-donor-acceptor-acceptor (DDAA) array, offering a relatively high dimerization energy (44–50 kJ mol 1) and a rapid kinetic reversibility (koff ¼ ca. 8 s 1) [67,68]. The quadruple hydrogen bonding is thermally reversible above 80 � C, which is an apparent advantage for building self-healing materials. Yan et al. [69] introduced UPy moeity into the backbone of IPDI-based TPU elastomer, resulting a highly stretchable (up to 17000% strain) self-healing elastomer. However, its strength is very weak (0.91 MPa) due to the sterically hindering iso phorone backbone. Song et al. [70] introduced dangling UPy side group into the HMDI-based TPU elastomer, obtaining a strong (44 MPa) yet tough (345 MJ m 3) mechanical property and an improved healablity compared with 1,4-butanediol chain extended analogue. However, its complete healing still needs relatively high temperature (above 100 � C) and long time (over 24 h). Here we reported a strategy of combining dynamic main chain and dangling UPy side group to fabricate a robust TPU elastomer with effi cient healing ability. The widely researched disulfide bond was selected as a typical example to build a dynamic main chain. The dimers of UPy groups act as strong physical crosslinking points and sacrificial bonds to enhance the mechanical property, while their location at the dangling side chains would accelerate the diffusion of the UPy groups at elevated temperatures, in favor of self-healing processes. Hence, this work ex pands upon this notion by evaluating the tensile mechanical properties and self-healing ability of TPU elastomers with different contents of dangling UPy side groups.
(DMF) were purchased from Aladdin (Shanghai). Guanidinium carbon ate, ethyl acetoacetate, 2-mercaptoethanol and 2-amino-1,3-propane diol (Serinol) was purchased from Macklin (Shanghai). DMF was dried by the adsorption of 4 A molecular sieves for 24h before use. 2.2. Synthesis of PU-UPy-SS elastomers The typical synthetic procedure of PU-UPy-SS elastomers was described as follows (take PU-U2S8 as the example): PTMEG (50 g, 50 mmol) was stirred under vacuum and heated at 120 � C for 2 h to remove moisture residue and then cooled to 80 � C. Then HMDI (26.2 g, 100 mmol) and 1 drop of DBTDL was then added into the PTMEG. The mixture was stirred for 2 h under N2 atmosphere. Then, anhydrous DMF (80 g), UPy-diol (3.5 g, 9.1 mmol) and DBTDL (0.25 g) were fed into the flask and stirred until it was dissolved totally. After that, DMF (120 g), SS-diol (5.6 g, 36.3 mmol) was added into the mixture and further stir red for another 2 h. A clear resin was then poured into Teflon mold and dried under 80 � C in the oven for 24 h. The obtained film was further dried under vacuum at 70 � C for 48 h. 2.3. Instrumentation and characterization Nuclear magnetic resonance (NMR) were carried out using a Brucker AVANCE III HD 600 MHz spectrometer. The deuterium solvent used was dependent on the product and has been detailed in each case. Chemical shifts are given in parts per million (ppm) relative to TMS with s for singlet, d for doublet, t for triplet, q for quadruplet and m for multiplet. Gel permeation chromatography (GPC) coupled with refractive index (RI) and UV detectors was conducted in N,N-dimethylformamide (DMF) with LiBr (0.05 M) at 50 � C and a flow rate of 1.0 mL min 1 using three successively connected Styragel columns (HR2, HR4, HR6). A se ries of narrowly dispersed polystyrene standards were used for calibra tion to obtain accurate number-average molar weight. Fourier transform infrared spectroscopy (FTIR) The FTIR spectra were recorded on a Vertex 70 (Bruker, German) equipped with a ZnSe ATR attachment. All TPU films were scanned 64 times within a range of 400–4000 cm 1 at a resolution of 4 cm 1. Raman Spectra were measured by LabRaM Aramis (H.J.Y, France) using a laser excitation wavelength of 532 nm. The spectra were generally collected over a spectral range from 4000 to 50 cm 1 with an accumulation number of 4 cm 1. Transparency tests were conducted on a Hitachi U-3010 UV–Vi sible spectrophotometer with an integrating sphere attachment. Spectra were generally collected over a spectral range from 300 nm to 800 nm, and the scanning speed was medium mode. Thermogravimetric analysis (TGA) was performed on a Netzsch STA449 instrument using 5–10 mg of sample. Samples were heated at a linear heating rate of 10 � C min 1 from 30 to 700 � C under N2 atmosphere. Differential scanning calorimetry (DSC) was performed on Melt tler Toledo DSC 3. Samples were heated from 150 � C to 150 � C at a 20 � C min 1 heating rate under nitrogen flow. The glass transition temperature (Tg) was selected as the medium point of the glass transition. Dynamic mechanical thermal analysis (DMA) was performed on a TA Q800 instruments using rectangular films (ca. 20 � 5 � 0.5 mm). Samples were loaded in tension mode and a temperature ramp was performed from 70 � C to 150 � C at a rate of 3 � C min 1, with an oscillating strain of 0.15% and an angular frequency of 1 Hz. A static preload of 0.01 N and a static force track of 125% was applied. The glass transition temperature (Tg) was calculated from the maximum value of the loss factor (tanδ). Atomic force microscopy (AFM) was performed using a Bruker Multimode 8 instruments, operated at room temperature on the free air and performed in the tapping modes.
2. Experimental section 2.1. Materials Polytetramethylene ether glycol (PTMEG) (Mn ¼ 1000 g mol 1) purchased from BASF. Hydrogenated 4,40 -methylenediphenyl diisocya nate (HMDI) was supplied by Coverstro. Hexamethylenediisocyanate (HDI) was supplied by Wanhua Chemical Co., Ltd. (The catalyst ditu byltin dilaurate (DBTDL, tech. 97.5%), and N,N0 -dimethylformamide 2
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Small-angle X-ray scattering (SAXS) was performed on a Xenocs Xeuss 2.0 (France) SAXS instrument at 50 kV operating voltage and 0.6 mA current, with an Excillum MetalJet-D2 X-ray source and a de tector (Pilatus 3R 1M, Dectris, Swiss). The wavelength generated is 0.134144 nm and the resolution is 1043 � 981 pixels (pixel size: 0.172 � 0.172 mm2). The scattering intensity were corrected for the background prior to evaluations and were normalized according to the incident primary beam intensity. The exposure time for each sample was 300 s. The 2D SAXS patterns were transformed into the 1D intensity profiles by Fit2D software in the azimuthal range of 0–360� . The scat tering intensity I(q) is profiled as a function of scattering vector q ¼ 4 π λ 1 sin θ, where λ is the wavelength of the incident X-rays and θ is half of the scattering angle. The sample-to-detector distance is 2500 mm, covering a q range of 0.07–2.5 nm 1. Tensile test was carried out by a universal tensile machine (Zwick010). The specimens were cut into a small dumbbell-shaped specimen (GB/T 528–2009, Type 3) with a thickness of approximately 0.5–0.7 mm. If not specified, the tensile tests were performed under the following conditions: a strain rate of 500 mm/min, a gauge length of 16 mm, room temperature (23 � 2 � C) and a humidity of 60 � 10%. The average of at least three individual tensile tests was recorded for each sample. The toughness is calculated by integration of the area below stress-strain curves. Rheology test was performed by an Anton Parr MCR302 rheometer. Samples were measured using a 25 mm parallel plate and interplate gap of 0.5–0.7 mm within the linear viscoelastic regime. Temperature scans were carried out from 30 to 160 � C with a healing rate of 5 � C min 1, at an angular frequency of 10 rad/s and strain amplitude of 1%. Frequency sweeps were conducted at 100 � C in a range from 628 rad/s to 0.05 rad/s with an applied strain of 1%. Scratch self-healing test was carried out on an Axioskop 40 POL (CEISS, German) equipped with a Linkam (THMS600E) hot stage.
Self-healing efficiency measurements: To evaluate the selfhealing ability of the samples, specimens were completely cut in half with a clean scalpel. Then, the two pieces of the samples were put together and gently pressed for 10 s at room temperature. Then they were placed in an oven at specified temperatures for a preset time. During the self-healing process, no external stress was applied to the interface. The healed samples were again subjected to tensile tests after equilibrated at room temperature for 24 h healing efficiency was defined as the percentage of property (tensile strength, elongation at break, toughness, etc) of healed sample over the corresponding properties of original sample. 3. Results and discussion 3.1. Synthesis of TPU elastomers SS-diol is a diol chain extender with a dynamic disulfide bond and was prepared from oxidation of 2-mercaptoethanol. UPy-diol is a diol chain extender with a pendent UPy side group. UPy-diol was success fully prepared from the reaction of Serinol and UPy-NCO, which was synthesized from the reaction of 6-methylisocytosine (MIC) and excess HDI. MIC was produced by the condensation reaction of ethyl acetoa cetate and guanidine carbonate. Detailed preparation procedures and reaction formula were included in Supporting Information and their characterization are detailed in Fig. S1 to Fig. S4. Scheme 1 illustrates the synthetic routes of TPU samples using HMDI as diisocyanate, PTMEG as polyol and UPy-diol/SS-diol as chain extender. Several samples with varying molar ratios of UPy-diol/SS-diol (Table S1) were successfully prepared. A reference sample without dy namic bond was prepared using 1,4-butanediol (BDO) as chain extender as well. All TPU samples were fully dissolvable in DMF, meaning they were solution processable even with the strong association of UPy
Scheme 1. The synthetic routes of PU-UPy-SS samples with dynamic main chain and self-complementary quadruple hydrogen bonds between dangling UPy side chains, acting as thermo reversible cross linkers. An illustrative depiction of the structure of PU-UPy-SS samples is presented as well. 3
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groups. The molecular weight of sample was measured by GPC and the data are summarized in Fig. S5 and Table 1. The GPC data present a decreased molecular weight with increasing UPy-diol contents. The reason is that molecular chains of TPU were partially looping at low concentrations [71–73] due to the quadruple hydrogen bonding be tween UPy groups within the same chain, which decreased the measured molecular weight calibrated by linear polystyrene standards. Films for testing were prepared by simple solution casting method. The as prepared PU-UPy-SS solution in DMF was casting on Teflon mold, and dried in an oven for 24 h at 80 � C. The solvent was completely removed at 70 � C for 48 h under vacuum. Finally, a series of colorless transparent self-healing TPU samples were obtained. UV–vis trans mission spectra of samples with a thickness about 500 μm were shown in Fig. S6. All the samples were highly transparent, with a transmittance up to 93% in visible light spectrum. As shown in Raman spectra in Fig. S7, the peaks appearing at 510 cm 1 (S-S) and 640 cm 1 (C-S) confirm the existence of disulfide bond in PU-SS and PU-UPy-SS samples. The NMR spectra of PU-SS and PU-UPy-SS and their peak assignments are shown in Fig. 1. Three characteristic peaks appearing at 13.20, 11.82 and 10.00 ppm are attributed to quadruple hydrogen bonded UPy moeity [67], confirming the successful incorporation of UPy motif into the PU chains and formation of quadruple hydrogen bonds in chloroform.
the derivative TGA curves displayed that the degradation of TPU sam ples consist of three weight loss regions, which are 250–300 � C, 310–380 � C and 390–460 � C. The peak temperature of the first degra dation region shifts to lower temperature with increasing UPy contents. These results demonstrate that the existence of dangling UPy side groups slightly decrease the thermostability of TPU samples. This behavior may be attributed to the enhancement of the chain mobility by dangling UPy side chain at elevated temperatures, which accelerates the heat diffusion in the polymer matrix, resulting in the reduction of onset thermal degradation. And this effect becomes insignificant as the temperature rises too high, so the latter two weight loss regions show little change as UPy content increases. The glass transition temperature (Tg) was investigated by DSC and DSC thermograms for different UPy contents appear in Fig. S9. The Tg of all the TPU samples are nearly the same value of 45 � C, which is the Tg of the PTMEG soft segments, confirmed by DMA curves (see Fig. 4) as the onset of the α transition. The Tg of soft segments is higher than that of pure PTMEG at 75 � C, indicating that parts of hard segments were mixed into the soft phases. The small endothermic peak at about 100 � C in DSC curve of PU-U5S5 in Fig. S9 is related to the partially stacked UPy dimers. This stacking structure only appears at higher contents of UPy side groups, due to the restriction of the steric repulsion caused by the four chain strands of side-group structure.
3.2. FTIR analysis
3.4. Dynamic mechanical thermal behaviors
The structures of all the TPU films were characterized using ATRFTIR technique and the spectra are shown in Fig. 2 and Fig. S8. Two bands were obtained in the region of stretching vibrations of N-H groups: H-bonded N-H groups at 3325 cm 1 and free N-H groups at 3450 cm 1. The predominance of the peak at 3325 cm 1 indicates nearly all the N-H groups are H-bonded. The stretching vibration region – O group is appearing between 1600 cm 1 and 1750 cm 1, and can of C– be divided into five bands (see Fig. 2b and Table S2). This absorption region was well fitted using Gauss Lorenz curves. It is consistent with literatures [74–76], except for the bands appearing at 1612 cm 1 – O groups located on the six-membered assigned to the H-boned C– pyrimidone ring of UPy group. The stretching vibration of – O groups appears at 1720 cm 1 while the non-H-boned urethane C– – O groups (and free urea C– – O groups) absorbs at H-bonded urethane C– 1695 cm 1. The band at 1664 cm 1 is assigned to the disordered – O groups, and the absorption band of ordered H-bonded urea C– – O groups is shift to 1640 cm 1 due to the stacking of H-bonded urea C– bifurcated urea hydrogen bonds. As shown in Fig. 2a, the adsorption – O groups (disordered or ordered H-bonded) and strength of urea C– characteristic band of UPy groups at 1612 cm 1 increase with the feeding ratio of UPy-diol/SS-diol, confirming the successful incorpora tion of UPy moiety in the polymer chain.
Dynamic thermo-mechanical analysis (DMA) can provide insights into the relaxation modes and viscoelastic behaviors of the TPU samples, and the results are depicted in Fig. 4. The storage modulus of all samples starts dropping at the same temperature of about 45 � C, which is consistent with the Tg measured by DSC. The peak temperature of tan δ is chosen as Tg, as summarized in Table 2. Without dangling UPy motifs, the storage modulus of PU-SS sample shows a clear rubbery plateau between 10 and 60 � C, indicating that a clear phase separation exists, and the hard phases acting as plastic fillers to form physical crosslinked networks. And the hard domains start their transition at about 80 � C and at the same temperature, dynamic exchanging of disulfide bonds begin to happen, entering the flow tran sition or terminal relaxation of TPU samples. For a very low content of UPy groups, like below 10 mol%, the flow transition temperature cuts down. The reason is that low concentration of UPy groups resulting in a rare possibility for UPy groups to encounter each other and bind together. Therefore, the dangling UPy side chains function only as an internal plasticizer decreasing the stability of hard domains as the feeding ratio of UPy-diol/SSdiol is less than 1:9. As a result, the onset of flow transition is advanced to 50 � C. When the UPy-diol content is higher than 20 mol%, UPy groups start to form quadruple hydrogen bonds, behaving as thermal reversible crosslinkers reinforcing the hard domains. As the UPy contents continue increasing, there are two opposite effects on the polymer matrix, one is the internal plasticizer effect at elevated temperatures, increasing the chain mobility of hard segments. This is beneficial for self-healing pro cess. Another is the cross-linker effect at lower temperatures, increasing the crosslinking density, thereby strengthening the mechanical properties. The increased cross-linking density is evidenced by the enhancement of plateau modulus between 0 and 60 � C and the raising of Tg. The widening of tan δ peak and the depressed tan δ peak is due to the extra friction imposed by the dangling UPy dimers sticked into the soft matrix. Nonetheless, due to the facts that quadruple hydrogen bonds between UPy groups start to dissociate at above 80 � C, and that the exchange reaction of disulfide bonds occurs at above 80 � C as well, the storage modulus of all the samples drops quickly above 80 � C, which is benefi cial for self-healing process. Thus, by varying the feeding ratio of UPydiol/SS-diol, we can obtain tunable mechanical properties at room temperature without affecting the self-healing properties at higher
3.3. Thermal properties The thermal stability of as-prepared TPU samples was investigated through thermogravimetric analysis (TGA), as shown in Fig. 3. The onset degradation temperature (T5, the temperature at 5 wt% of weight loss) is slightly decreased as the feeding ratio of UPy-diol/SS-diol increases. And Table 1 The GPC results of TPU samples. Sample Name
Molar ratio of (UPy-diol:SS-diol)
Mn (Da)
Mw (Da)
Polymer Dispersity Index
PU-0 PU-SS PU-U1S9 PU-U2S8 PU-U3S7 PU-U4S6 PU-U5S5
0:0 0:1 0.1 : 0.9 0.2 : 0.8 0.3 : 0.7 0.4 : 0.6 0.5 : 0.5
9.63 � 104 5.48 � 104 3.31 � 104 2.78 � 104 2.50 � 104 2.23 � 104 2.14 � 104
1.59 � 105 7.20 � 104 5.04 � 104 4.63 � 104 4.06 � 104 3.75 � 104 3.43 � 104
1.65 1.31 1.52 1.66 1.62 1.64 1.60
4
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Fig. 1. The 1H NMR spectra of TPU-elastomers: (a) PU-U2S8, (b) PU-SS.
Fig. 2. (a) FTIR spectra of different TPU samples, and the absorptance is normalized according to the peak height at 1448 cm the band region of stretching vibration.
temperatures.
1
(δa(CH2)). (b) Peak fitting curves in
U2S8, PU-U5S5 clearly show this transition as well. Among PU-UPy-SS samples, the long period interdomain distances of phase separation for PU-U2S8, PU-U3S7, PU-U4S6, PU-U5S5 are 14.9, 13.0, 12.1, 11.2 nm, respectively. As the content of UPy groups is raised, the magnitude of the scattering maximum increases and its location shifts to a higher scat tering vector. A trend of decreasing scale of phase separation is man ifested as the content of UPy groups increases. But on the contrary, the electron density contrast between two phases is increasing, indicating the difference between two phases become more distinguishing with UPy side chains. These results suggest that the incorporation of dangling UPy groups leads to an increase of the ordering of hard segments, while the sizes of hard domains become smaller.
3.5. Morphological structures The microphase separated structure of TPU samples was investigated by atomic force microscopy (AFM) (Fig. 5) and small angle X-ray scat tering (SAXS) (Fig. 6). The AFM image of PU-SS clearly shows the phase segregated hard segments (bright areas) and soft segments (dark areas). The non-structured phase separation scale is about 100–300 nm, out of the measuring limit of SAXS, which is the main reason of absence of scattering peak for PU-SS and PU-U1S9 films. Upon introduction of UPy side groups, the scale of phase separation was reduced to 10–15 nm (according to d ¼ 2π/q, where q ¼ 4π λ 1 sin θ). AFM images of PU5
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Fig. 3. TGA (a) and DTG (b) curves of TPU samples with varying contents of UPy motifs.
Fig. 4. Storage modulus (E0 ), loss modulus (E’’) and tan δ versus temperature of TPU samples.
break (εb) and toughness (integral area of a stress-strain curve) of each film are summarized in Table 2. The PU-SS without the UPy groups exhibits a typical elastomer profile, strengthened by strain induced hardening. Due to the relatively weak nature of the hard phase based on HMDI and SS-diol, PU-SS sample is very soft till the elongation reached a critical point (1000% in Fig. 7a), after which the strain-induce hardening mechanism enhances its me chanical property. Upon introduction of the dangling UPy motifs, the mechanical property of PU-UPy-SS sample is obviously influenced. For example, compared with PU-SS film, the tensile strength of PU-U1S9 is much weaker due to the weakened stability of hard domains by dangling UPy side chains. It is worthy to point out that the low concentration of UPy groups reduces their probability of complimentary associating by quadruple hydrogen bonds. Thus, the adverse effect of dangling side chains plays a major role resulting an excessively weakened tensile strength. A further increase of UPy contents above 20% leads to an increased tendency of forming quadruple hydrogen bonds between UPy motifs in the hard segments, raising the tensile strength of PU-U2S8 and PU-U3S7. Again, strain induced hardening of soft segments still plays an important role after the sample underwent a large deformation. It is noted that this effect is gradually suppressed according to the observa tion of a lowered slope of stress strain curves when the strain is larger
Table 2 The mechanical properties of TPU samples with different contents of UPy groups. Sample name
Tg (� C)
Tensile strength (MPa)
Elongation at break (%)
Toughness (MJ m 3)
PU-SS PU-U1S9 PU-U2S8 PU-U3S7 PU-U4S6 PU-U5S5
7.0 0.1 0.6 1.8 12.8 18.6
13.0 � 0.6 2.1 � 0.3 25.2 � 0.7 22.8 � 0.7 14.6 � 0.1 11.6 � 0.2
1964 � 163 2943 � 142 1463 � 32 1406 � 33 1406 � 27 1215 � 10
64.7 � 0.9 39.9 � 1.5 101.0 � 0.7 108.8 � 0.8 90.3 � 0.2 69.6 � 0.4
3.6. Mechanical properties The mechanical properties of the PU-UPy-SS samples were evaluated by tensile testing. For this aim, 0.4–0.7 mm thick films were cast in a PTFE mold, which were then cut into dumbbell-shape for further tensile measurements. The high mechanical properties of polyurethanes are originated from phase separation, so thermal history will influence their mechanical properties. Thus, all the samples were equilibrated at room temperature for at least 7 days before tensile testing, to reduce the ef fects of thermal history. The typical tensile stress-strain curves are exhibited in Fig. 7a and the obtained tensile strength (σb), elongation at 6
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Fig. 5. SAXS intensity curves of different TPU samples.
To further elucidate the strengthening effect of self-complimentary UPy groups, cyclic loading-unloading test was conducted for TPU films. The typical stress-strain curves were shown in Fig. S10 and the evolving of hysteresis areas with the increasing cycle number was pre sented in Fig. 7b. The films were stretching to a large deformation of 1000% in order to destroy the strengthening sacrificial interactions in the first few cycles. As is shown in Fig. 7b, the first hysteresis area in creases with the increasing UPy content, while the hysteresis area of the last cycle is almost the same. It suggests that the quadruple H-bonding between UPy side chains is effective in reinforcing the tensile strength at small deformation as sacrificial crosslinking point, when putting aside the strain induced hardening effects at large deformation. And due to the dynamic nature of quadruple H-bonding, the opened quadruple hydrogen bonds are able to re-form again as time goes by, restoring the original mechanical properties, which can be accelerated by increasing temperature. 3.7. Self-healing abilities
Fig. 6. AFM phase images of TPU samples: (a) PU-SS, (b) PU-U2S8, (c) PU-U5S5.
The self-healing abilities of PU-UPy-SS samples were first evaluated intuitively by simple healing and stretching test. A dumbbell-shaped sample (PU-U2S8) was fully cut in half and healed at 100 � C for 2 h. As shown in Fig. 8, the specimen was fully healed and was able to sustain a high stretching ratio. And the fracture position was not located at the healing surface once it was snapped again, suggesting its mechanical property was full recovered. Then the self-healing process was observed by a scratch healing test, monitored by an optical microscope equipped with a Linkam hot stage. Scratches of 10–20 μm wide and 200–400 μm deep were made by scalpel on the samples. These samples were then
than 1000%. There are two possible reasons, which are the increased structure irregularity and the retarding effect of strong quadruple bonding. Therefore, the final strength shows a decreasing trend when the UPy content is above 30 mol%, even though the stress at lower strain is increasing. These observations indicate that only incorporation of a suitable amount of UPy side chains can effectively in enhance the me chanical properties of intrinsically weak self-healing TPU samples.
Fig. 7. (a) Typical stress-strain curves of TPU samples; (b) Hysteresis area of TPU samples under five successive loading-unloading cycles. 7
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Fig. 8. Optical microscope images after different healing times at 100 � C of scratched samples: (a) PU-SS, (b) PU-U2S8, (c) PU-U5S5.
placed on a hot stage to observe the healing of the cracks at constant temperatures. The optical images are shown in Fig. 9 and Fig. S10. The cracks of the surface closed in 10 min and then the traces become shallower and gradually disappear after 30 min. This is because the exchanging of disulfide bonds, the reversible associating of quadruple hydrogen bonds and the molecular chain diffusion of hard segments were all fully activated at 100 � C. For comparison, no sign of healing of the cracks on the film of PU-0 sample was observed after 30 min at 100 � C. As for PU-U1S9, the self-healing processes proceeds quickly even at 50 � C as shown in Fig. S10a. The main reason is the hard domains are
destroyed by the non-associating dangling UPy side chains at low con centration, accelerating the molecular chain diffusion. To quantitively evaluate the self-healing efficiencies, tensile testing was further applied to investigate the self-healing efficiencies of the samples at different temperatures. For this aim, a 0.4–0.7 mm thick dumbbell-shaped film was fully cut in half using a clean scalpel, and then gently pressed together at room temperature. They were allowed to heal at 40, 60, 80 and 100 � C for 2 h without any external pressure. The stress strain curves obtained for different samples at various healing temperatures were displayed in Fig. 10 and Fig. S11. The healing
Fig. 9. Two cut pieces coming from two different dog-boned films, one stained with methyl blue for clarifying, were allowed to heal at 100 � C for 2 h, and then stretched to observe the effectiveness of self-healing, after which the healed specimen was snapped again. The healed specimen did not break at the original healing surfaces, suggesting it was fully healed. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.) 8
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efficiency of PU-U1S9 was not included, because valid data could not be obtained due to the soft, sticky nature and unacceptable repeatability of mechanical properties after healing. First, after 2 h of self-healing at 100 � C, all samples were completely healed (healing efficiencies are greater than 90%). And they showed a significant extent of healing ef ficiency at 80 � C for 2 h, meaning the healing processes begins at 80 � C. But the full healing requires a prolonged time. It was noted that even at 40 � C, the samples were partially healed. The main driving force behind it was the weak hydrogen bonds between urethane or urea bonds. Therefore, besides quadruple hydrogen bonds and disulfide bonds, these weak hydrogen bonds contribute to the self-healing process as well. As shown in Table 3, compared with PU-SS, the healing efficiency of PU-U2S8 and PU-U3S7 were lowered. It can be attributed to the extra diffusion friction brought by quadruple hydrogen bonds. At lower temperatures, such as 40, 60 and 80 � C the healing efficiency of tensile strength gradually increases with the UPy contents while the healing efficiency of tensile strain drops down. These results indicate that UPy groups accelerate the healing of tensile strength but hampers the healing of stretching ability. This is due to the re-forming of strong quadruple hydrogen bonds contributing to the effective binding of damage sur faces. However, the high dimerization constant of UPy groups slows down the chain diffusion and randomization processes of molecular chains from both sides of damage surfaces, resulting in the decreased healing efficiency of tensile strain. To reduce this effect, UPy groups were introduced as side chains to accelerate chain diffusion when the dimers of UPy dissociate at elevated temperatures. For further understanding of the self-healing mechanism of this system, the rheology behaviors of PU-SS, PU-U2S8 and PU-U5S5 were measured as depicted in Fig. S12. The flow transition temperatures (the crossover temperature of G’and G’’ in the temperature ramping data) of PU-SS, PU-U2S8 and PU-U5S5 were 100 � C, 110 � C and 114 � C, respec tively. The flow transition temperature was only slightly increased. To better understand the effect of UPy side chain on the flow behaviors and relaxation time, the rheology frequency sweeping tests were then con ducted at 100 � C. As indicated in Fig. S12, the relaxation times (τ ¼ 2π/
ωc at the crossover point of G0 and G’’ in the frequency sweeping data) of
PU-SS, PU-U2S5 and PU-U5S5 were 0.28 s, 0.95 s and 1.45 s, respec tively. The relaxation time was increasing with introduction of UPy groups, but to a negligible extent. The small relaxation time of all samples at high temperatures is due to the low bond life time of disulfide bond and quadruple hydrogen bond, providing the PU-UPy-SS samples with efficient healing ability at elevated temperatures. Therefore, the self-healing ability of TPU sample was not significantly compromised while the mechanical property was strengthened by the introduction of UPy side groups. 4. Conclusions In this work, we have successfully fabricated mechanically enhanced self-healing TPU elastomers by integration of dangling UPy side group into the dynamic disulfide main chain. The variation of the feeding ratio of UPy-diol/SS-diol could be used to tailor the properties of this mate rial. The introduction of UPy side groups possess two opposite effects. One is the plasticizer effect of side chain, resulting a decreased onset degradation temperature, and increased chain mobility at elevated temperatures. The other is the thermo-triggered reversible crosslinking effect originating from the self-complimentary quadruple H-bonding of UPy motifs, increasing the crosslinking density of hard domains and thereby enhancing its mechanical robustness at room temperature. The optimization results revealed that using 20–30 mol % of UPy-diol chain extender led to a better result in balancing the mechanical robustness and self-healing properties, obtaining a tensile strength up to 25 MPa and a healing efficiency over 92% in 2 h at 100 � C. For future investi gation, the enhancing strategy of introducing UPy side groups can be applied to various self-healing systems incorporating different dynamic covalent bonds in the main chain. Moreover, the wide commercial availability of raw materials and the simple synthetic processes make this system very attractive for real industrial applications.
Fig. 10. The typical stress-strain curves of different TPU samples after 2 h healing at various temperature: (a) PU-SS; (b) PU-U2S8; (c) PU-U3S7; (d) PU-U4S6; (e) PU-U5S5. 9
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Table 3 The healing efficiencies of TPU samples healed at various temperatures for 2 h. Sample name PU-SS PU-U2S8 PU-U3S7 PU-U4S6 PU-U5S5
Healing efficiency of tensile strength (%)
Healing efficiency of elongation at break (%)
Healing efficiency of toughness (%)
40
60
80
100
40
60
80
100
40
60
80
100
25.0 15.0 17.8 28.3 49.0
52.7 39.3 41.2 46.5 69.3
93.8 76.0 79.1 90.1 93.3
98.5 92.4 94.2 99.7 96.6
71.1 46.9 44.3 34.2 33.2
87.0 73.6 72.7 58.5 53.2
96.8 92.3 90.0 89.3 86.6
102.9 99.9 98.3 95.9 99.9
33.8 14.6 13.1 12.9 18.9
55.8 35.1 41.0 32.0 49.3
97.4 74.6 74.4 78.2 87.0
105.2 99.5 94.5 94.3 99.4
Declaration of competing interestCOI
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Polymer journal homepage: http://www.elsevier.com/locate/polymer
Towards mechanical robust yet self-healing polyurethane elastomers via combination of dynamic main chain and dangling quadruple hydrogen bonds Jin Hu a, Ruibin Mo a, Xiang Jiang a, Xinxin Sheng b, **, Xinya Zhang a, * a
School of Chemistry and Chemical Engineering, South China University of Technology, Guangzhou, 510640, China Guangdong Provincial Key Laboratory of Functional Soft Condensed Matter, Department of Polymeric Materials and Engineering, School of Materials and Energy, Guangdong University of Technology, Guangzhou, 510006, China
b
A R T I C L E I N F O
A B S T R A C T
Keywords: Self-healing Dynamic covalent bonds Quadruple hydrogen bonds
One of the greatest challenges of robust self-healing materials is the confliction between high chain mobility for self-healing and a stable structure for mechanical strength. Herein, dangling 2-ureido-4[1H]-pyrimidione (UPy)functionalized side groups were introduced into the hard segments of thermoplastic polyurethane (TPU) elas tomers, where embedded the dynamic disulfide bonds in the main chain. The strong quadruple H-bonding interaction between UPy side groups acts as supramolecular crosslinkers enabling the TPU elastomer to have improved mechanical properties (tensile strength up to 25 MPa and toughness ~100 MJ m 3), and simulta neously the plasticizer effect of dangling side chain endows it with efficient healing ability at elevated tem peratures (80–100 � C) comparable to its linear analogues. This strategy shows great potential in designing robust self-healing TPU elastomer employing weak dynamic covalent bonds, with wide promising applications as wearable electronics, coatings and adhesives.
1. Introduction Inspired by nature, self-healing materials capable of repairing dam ages and restoring their original properties show great potential in numerous applications, especially in the fields of protective coatings [1–3], wearable electronics [4–6] and biomedical materials [7–10]. Among various approaches to develop self-healing polymers, the intro duction of dynamic covalent bonds and non-covalent interactions to chain molecules, termed as intrinsic self-healing, is one of the most powerful approaches. Intrinsic self-healing material was first introduced by Chen et al. [11], made by reversible Diels-Alder (DA) reaction of tetra-maleimide and tetra-furan. After that, people have successfully fabricated self-healing polymers based on various dynamic covalent bonds, such as DA reaction [11], disulfide [12–14], diselenide [15–17] and ditelluride [18] metathesis, imine metathesis [19–21], olefin metathesis [22–24], hindered urea bonds [25–27], transesterification [28–30], trans carbonation [31] of carbonates, transcarbamoylation [32,33] of ure thanes, transamination [34,35] of vinylogous urethanes, dioxaborolane
metathesis [36–38] of boronic esters, siloxane equilibration [39–41], alkoxyamine metathesis [42–44] and so on. Supramolecular interactions are widely employed to construct self-healing materials as well, including hydrogen bonds [45–48], metal-ligand coordination [49–51], ionic interactions [52], π-π stacking [53–55], host-guest interactions [56–58] and even only van der Waals forces [59], etc. In addition to the dynamic bonding required for structure reorgani zation, chain diffusion plays an important role in the healing process as well. According to Wool’s model [60], a healing process goes through five phases: (i) surface rearrangement, (ii) surface approach, (iii) wet ting, (iv) diffusion and (v) randomization, during which diffusion is a critical process for crack healing. Thus, it requires the polymer chain to be soft enough or has relatively low molecular weight, which is the main reason of compromised mechanical properties for self-healing materials. In general, self-healing ability and mechanical robustness of the material are contradictory and difficult to be optimized simultaneously. Thermoplastic polyurethane (TPU) is a thermal processable and solvent dissolvable elastomer that is widely used in dozens of fields due to its high resilience and resistance to impacts, abrasions and tears. As is
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (X. Sheng), [email protected] (X. Zhang). https://doi.org/10.1016/j.polymer.2019.121912 Received 22 July 2019; Received in revised form 18 September 2019; Accepted 15 October 2019 Available online 17 October 2019 0032-3861/© 2019 Elsevier Ltd. All rights reserved.
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commonly known, TPU is a segmented block copolymer that contains a low Tg soft block and a physically crosslinked hard domain dispersed in the soft matrix. And its performance and morphology can be easily tuned by selecting suitable diisocyanates, polyols and chain extenders. Taking these characteristics of TPU, researchers have designed various selfhealing TPU elastomers by introducing dynamic covalent bonds and non-covalent interactions into TPU chains. Normally, to impart self-healing ability to TPU elastomers, dynamic covalent linkages are introduced into loosely packed TPU backbones in the form of chain extenders. As a result, intrinsic self-healing materials that can heal under mild conditions always have compromised me chanical properties [61–64], due to the weak nature of dynamic bonds and increased mobility of chain structure. Addition of covalent cross linking points can enhance the mechanical properties of PU elastomers, but in the sacrifice of efficient self-healing ability. Yang et al. [65] used hexamethylene diisocyanate (HDI) trimer as a crosslinker, and aromatic disulfide as a dynamic bond to fabricate a self-healing PU elastomer, finding that increasing HDI trimer content would suppress the self-healing ability. The ordering structure of hard domains with labile dynamic covalent bonds embedding in may affect the self-healing properties of TPU elastomers as well. Kim et al. [61] found that the self-healing efficiencies of the TPUs decreased in the following order: IPDI-SS > IPDI-EG � HMDI-SS > HDI-SS � MDI-SS. HDI and MDI based TPU elastomer cannot heal up to 80 � C due to their tightly packed and crystalline structures of hard domains, hindering the effective disulfide bond exchanging between different hard domains. The volume content of hard phases is another key factor, revealed by Lai [66] et al., who discovered that increased content of hard segments will enhance the mechanical properties. 2-ureido-4[1H]-pyrimidinone (or UPy) moiety shows a strong ten dency for self-complementary dimerization involving four hydrogen bonds in a donor-donor-acceptor-acceptor (DDAA) array, offering a relatively high dimerization energy (44–50 kJ mol 1) and a rapid kinetic reversibility (koff ¼ ca. 8 s 1) [67,68]. The quadruple hydrogen bonding is thermally reversible above 80 � C, which is an apparent advantage for building self-healing materials. Yan et al. [69] introduced UPy moeity into the backbone of IPDI-based TPU elastomer, resulting a highly stretchable (up to 17000% strain) self-healing elastomer. However, its strength is very weak (0.91 MPa) due to the sterically hindering iso phorone backbone. Song et al. [70] introduced dangling UPy side group into the HMDI-based TPU elastomer, obtaining a strong (44 MPa) yet tough (345 MJ m 3) mechanical property and an improved healablity compared with 1,4-butanediol chain extended analogue. However, its complete healing still needs relatively high temperature (above 100 � C) and long time (over 24 h). Here we reported a strategy of combining dynamic main chain and dangling UPy side group to fabricate a robust TPU elastomer with effi cient healing ability. The widely researched disulfide bond was selected as a typical example to build a dynamic main chain. The dimers of UPy groups act as strong physical crosslinking points and sacrificial bonds to enhance the mechanical property, while their location at the dangling side chains would accelerate the diffusion of the UPy groups at elevated temperatures, in favor of self-healing processes. Hence, this work ex pands upon this notion by evaluating the tensile mechanical properties and self-healing ability of TPU elastomers with different contents of dangling UPy side groups.
(DMF) were purchased from Aladdin (Shanghai). Guanidinium carbon ate, ethyl acetoacetate, 2-mercaptoethanol and 2-amino-1,3-propane diol (Serinol) was purchased from Macklin (Shanghai). DMF was dried by the adsorption of 4 A molecular sieves for 24h before use. 2.2. Synthesis of PU-UPy-SS elastomers The typical synthetic procedure of PU-UPy-SS elastomers was described as follows (take PU-U2S8 as the example): PTMEG (50 g, 50 mmol) was stirred under vacuum and heated at 120 � C for 2 h to remove moisture residue and then cooled to 80 � C. Then HMDI (26.2 g, 100 mmol) and 1 drop of DBTDL was then added into the PTMEG. The mixture was stirred for 2 h under N2 atmosphere. Then, anhydrous DMF (80 g), UPy-diol (3.5 g, 9.1 mmol) and DBTDL (0.25 g) were fed into the flask and stirred until it was dissolved totally. After that, DMF (120 g), SS-diol (5.6 g, 36.3 mmol) was added into the mixture and further stir red for another 2 h. A clear resin was then poured into Teflon mold and dried under 80 � C in the oven for 24 h. The obtained film was further dried under vacuum at 70 � C for 48 h. 2.3. Instrumentation and characterization Nuclear magnetic resonance (NMR) were carried out using a Brucker AVANCE III HD 600 MHz spectrometer. The deuterium solvent used was dependent on the product and has been detailed in each case. Chemical shifts are given in parts per million (ppm) relative to TMS with s for singlet, d for doublet, t for triplet, q for quadruplet and m for multiplet. Gel permeation chromatography (GPC) coupled with refractive index (RI) and UV detectors was conducted in N,N-dimethylformamide (DMF) with LiBr (0.05 M) at 50 � C and a flow rate of 1.0 mL min 1 using three successively connected Styragel columns (HR2, HR4, HR6). A se ries of narrowly dispersed polystyrene standards were used for calibra tion to obtain accurate number-average molar weight. Fourier transform infrared spectroscopy (FTIR) The FTIR spectra were recorded on a Vertex 70 (Bruker, German) equipped with a ZnSe ATR attachment. All TPU films were scanned 64 times within a range of 400–4000 cm 1 at a resolution of 4 cm 1. Raman Spectra were measured by LabRaM Aramis (H.J.Y, France) using a laser excitation wavelength of 532 nm. The spectra were generally collected over a spectral range from 4000 to 50 cm 1 with an accumulation number of 4 cm 1. Transparency tests were conducted on a Hitachi U-3010 UV–Vi sible spectrophotometer with an integrating sphere attachment. Spectra were generally collected over a spectral range from 300 nm to 800 nm, and the scanning speed was medium mode. Thermogravimetric analysis (TGA) was performed on a Netzsch STA449 instrument using 5–10 mg of sample. Samples were heated at a linear heating rate of 10 � C min 1 from 30 to 700 � C under N2 atmosphere. Differential scanning calorimetry (DSC) was performed on Melt tler Toledo DSC 3. Samples were heated from 150 � C to 150 � C at a 20 � C min 1 heating rate under nitrogen flow. The glass transition temperature (Tg) was selected as the medium point of the glass transition. Dynamic mechanical thermal analysis (DMA) was performed on a TA Q800 instruments using rectangular films (ca. 20 � 5 � 0.5 mm). Samples were loaded in tension mode and a temperature ramp was performed from 70 � C to 150 � C at a rate of 3 � C min 1, with an oscillating strain of 0.15% and an angular frequency of 1 Hz. A static preload of 0.01 N and a static force track of 125% was applied. The glass transition temperature (Tg) was calculated from the maximum value of the loss factor (tanδ). Atomic force microscopy (AFM) was performed using a Bruker Multimode 8 instruments, operated at room temperature on the free air and performed in the tapping modes.
2. Experimental section 2.1. Materials Polytetramethylene ether glycol (PTMEG) (Mn ¼ 1000 g mol 1) purchased from BASF. Hydrogenated 4,40 -methylenediphenyl diisocya nate (HMDI) was supplied by Coverstro. Hexamethylenediisocyanate (HDI) was supplied by Wanhua Chemical Co., Ltd. (The catalyst ditu byltin dilaurate (DBTDL, tech. 97.5%), and N,N0 -dimethylformamide 2
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Small-angle X-ray scattering (SAXS) was performed on a Xenocs Xeuss 2.0 (France) SAXS instrument at 50 kV operating voltage and 0.6 mA current, with an Excillum MetalJet-D2 X-ray source and a de tector (Pilatus 3R 1M, Dectris, Swiss). The wavelength generated is 0.134144 nm and the resolution is 1043 � 981 pixels (pixel size: 0.172 � 0.172 mm2). The scattering intensity were corrected for the background prior to evaluations and were normalized according to the incident primary beam intensity. The exposure time for each sample was 300 s. The 2D SAXS patterns were transformed into the 1D intensity profiles by Fit2D software in the azimuthal range of 0–360� . The scat tering intensity I(q) is profiled as a function of scattering vector q ¼ 4 π λ 1 sin θ, where λ is the wavelength of the incident X-rays and θ is half of the scattering angle. The sample-to-detector distance is 2500 mm, covering a q range of 0.07–2.5 nm 1. Tensile test was carried out by a universal tensile machine (Zwick010). The specimens were cut into a small dumbbell-shaped specimen (GB/T 528–2009, Type 3) with a thickness of approximately 0.5–0.7 mm. If not specified, the tensile tests were performed under the following conditions: a strain rate of 500 mm/min, a gauge length of 16 mm, room temperature (23 � 2 � C) and a humidity of 60 � 10%. The average of at least three individual tensile tests was recorded for each sample. The toughness is calculated by integration of the area below stress-strain curves. Rheology test was performed by an Anton Parr MCR302 rheometer. Samples were measured using a 25 mm parallel plate and interplate gap of 0.5–0.7 mm within the linear viscoelastic regime. Temperature scans were carried out from 30 to 160 � C with a healing rate of 5 � C min 1, at an angular frequency of 10 rad/s and strain amplitude of 1%. Frequency sweeps were conducted at 100 � C in a range from 628 rad/s to 0.05 rad/s with an applied strain of 1%. Scratch self-healing test was carried out on an Axioskop 40 POL (CEISS, German) equipped with a Linkam (THMS600E) hot stage.
Self-healing efficiency measurements: To evaluate the selfhealing ability of the samples, specimens were completely cut in half with a clean scalpel. Then, the two pieces of the samples were put together and gently pressed for 10 s at room temperature. Then they were placed in an oven at specified temperatures for a preset time. During the self-healing process, no external stress was applied to the interface. The healed samples were again subjected to tensile tests after equilibrated at room temperature for 24 h healing efficiency was defined as the percentage of property (tensile strength, elongation at break, toughness, etc) of healed sample over the corresponding properties of original sample. 3. Results and discussion 3.1. Synthesis of TPU elastomers SS-diol is a diol chain extender with a dynamic disulfide bond and was prepared from oxidation of 2-mercaptoethanol. UPy-diol is a diol chain extender with a pendent UPy side group. UPy-diol was success fully prepared from the reaction of Serinol and UPy-NCO, which was synthesized from the reaction of 6-methylisocytosine (MIC) and excess HDI. MIC was produced by the condensation reaction of ethyl acetoa cetate and guanidine carbonate. Detailed preparation procedures and reaction formula were included in Supporting Information and their characterization are detailed in Fig. S1 to Fig. S4. Scheme 1 illustrates the synthetic routes of TPU samples using HMDI as diisocyanate, PTMEG as polyol and UPy-diol/SS-diol as chain extender. Several samples with varying molar ratios of UPy-diol/SS-diol (Table S1) were successfully prepared. A reference sample without dy namic bond was prepared using 1,4-butanediol (BDO) as chain extender as well. All TPU samples were fully dissolvable in DMF, meaning they were solution processable even with the strong association of UPy
Scheme 1. The synthetic routes of PU-UPy-SS samples with dynamic main chain and self-complementary quadruple hydrogen bonds between dangling UPy side chains, acting as thermo reversible cross linkers. An illustrative depiction of the structure of PU-UPy-SS samples is presented as well. 3
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groups. The molecular weight of sample was measured by GPC and the data are summarized in Fig. S5 and Table 1. The GPC data present a decreased molecular weight with increasing UPy-diol contents. The reason is that molecular chains of TPU were partially looping at low concentrations [71–73] due to the quadruple hydrogen bonding be tween UPy groups within the same chain, which decreased the measured molecular weight calibrated by linear polystyrene standards. Films for testing were prepared by simple solution casting method. The as prepared PU-UPy-SS solution in DMF was casting on Teflon mold, and dried in an oven for 24 h at 80 � C. The solvent was completely removed at 70 � C for 48 h under vacuum. Finally, a series of colorless transparent self-healing TPU samples were obtained. UV–vis trans mission spectra of samples with a thickness about 500 μm were shown in Fig. S6. All the samples were highly transparent, with a transmittance up to 93% in visible light spectrum. As shown in Raman spectra in Fig. S7, the peaks appearing at 510 cm 1 (S-S) and 640 cm 1 (C-S) confirm the existence of disulfide bond in PU-SS and PU-UPy-SS samples. The NMR spectra of PU-SS and PU-UPy-SS and their peak assignments are shown in Fig. 1. Three characteristic peaks appearing at 13.20, 11.82 and 10.00 ppm are attributed to quadruple hydrogen bonded UPy moeity [67], confirming the successful incorporation of UPy motif into the PU chains and formation of quadruple hydrogen bonds in chloroform.
the derivative TGA curves displayed that the degradation of TPU sam ples consist of three weight loss regions, which are 250–300 � C, 310–380 � C and 390–460 � C. The peak temperature of the first degra dation region shifts to lower temperature with increasing UPy contents. These results demonstrate that the existence of dangling UPy side groups slightly decrease the thermostability of TPU samples. This behavior may be attributed to the enhancement of the chain mobility by dangling UPy side chain at elevated temperatures, which accelerates the heat diffusion in the polymer matrix, resulting in the reduction of onset thermal degradation. And this effect becomes insignificant as the temperature rises too high, so the latter two weight loss regions show little change as UPy content increases. The glass transition temperature (Tg) was investigated by DSC and DSC thermograms for different UPy contents appear in Fig. S9. The Tg of all the TPU samples are nearly the same value of 45 � C, which is the Tg of the PTMEG soft segments, confirmed by DMA curves (see Fig. 4) as the onset of the α transition. The Tg of soft segments is higher than that of pure PTMEG at 75 � C, indicating that parts of hard segments were mixed into the soft phases. The small endothermic peak at about 100 � C in DSC curve of PU-U5S5 in Fig. S9 is related to the partially stacked UPy dimers. This stacking structure only appears at higher contents of UPy side groups, due to the restriction of the steric repulsion caused by the four chain strands of side-group structure.
3.2. FTIR analysis
3.4. Dynamic mechanical thermal behaviors
The structures of all the TPU films were characterized using ATRFTIR technique and the spectra are shown in Fig. 2 and Fig. S8. Two bands were obtained in the region of stretching vibrations of N-H groups: H-bonded N-H groups at 3325 cm 1 and free N-H groups at 3450 cm 1. The predominance of the peak at 3325 cm 1 indicates nearly all the N-H groups are H-bonded. The stretching vibration region – O group is appearing between 1600 cm 1 and 1750 cm 1, and can of C– be divided into five bands (see Fig. 2b and Table S2). This absorption region was well fitted using Gauss Lorenz curves. It is consistent with literatures [74–76], except for the bands appearing at 1612 cm 1 – O groups located on the six-membered assigned to the H-boned C– pyrimidone ring of UPy group. The stretching vibration of – O groups appears at 1720 cm 1 while the non-H-boned urethane C– – O groups (and free urea C– – O groups) absorbs at H-bonded urethane C– 1695 cm 1. The band at 1664 cm 1 is assigned to the disordered – O groups, and the absorption band of ordered H-bonded urea C– – O groups is shift to 1640 cm 1 due to the stacking of H-bonded urea C– bifurcated urea hydrogen bonds. As shown in Fig. 2a, the adsorption – O groups (disordered or ordered H-bonded) and strength of urea C– characteristic band of UPy groups at 1612 cm 1 increase with the feeding ratio of UPy-diol/SS-diol, confirming the successful incorpora tion of UPy moiety in the polymer chain.
Dynamic thermo-mechanical analysis (DMA) can provide insights into the relaxation modes and viscoelastic behaviors of the TPU samples, and the results are depicted in Fig. 4. The storage modulus of all samples starts dropping at the same temperature of about 45 � C, which is consistent with the Tg measured by DSC. The peak temperature of tan δ is chosen as Tg, as summarized in Table 2. Without dangling UPy motifs, the storage modulus of PU-SS sample shows a clear rubbery plateau between 10 and 60 � C, indicating that a clear phase separation exists, and the hard phases acting as plastic fillers to form physical crosslinked networks. And the hard domains start their transition at about 80 � C and at the same temperature, dynamic exchanging of disulfide bonds begin to happen, entering the flow tran sition or terminal relaxation of TPU samples. For a very low content of UPy groups, like below 10 mol%, the flow transition temperature cuts down. The reason is that low concentration of UPy groups resulting in a rare possibility for UPy groups to encounter each other and bind together. Therefore, the dangling UPy side chains function only as an internal plasticizer decreasing the stability of hard domains as the feeding ratio of UPy-diol/SSdiol is less than 1:9. As a result, the onset of flow transition is advanced to 50 � C. When the UPy-diol content is higher than 20 mol%, UPy groups start to form quadruple hydrogen bonds, behaving as thermal reversible crosslinkers reinforcing the hard domains. As the UPy contents continue increasing, there are two opposite effects on the polymer matrix, one is the internal plasticizer effect at elevated temperatures, increasing the chain mobility of hard segments. This is beneficial for self-healing pro cess. Another is the cross-linker effect at lower temperatures, increasing the crosslinking density, thereby strengthening the mechanical properties. The increased cross-linking density is evidenced by the enhancement of plateau modulus between 0 and 60 � C and the raising of Tg. The widening of tan δ peak and the depressed tan δ peak is due to the extra friction imposed by the dangling UPy dimers sticked into the soft matrix. Nonetheless, due to the facts that quadruple hydrogen bonds between UPy groups start to dissociate at above 80 � C, and that the exchange reaction of disulfide bonds occurs at above 80 � C as well, the storage modulus of all the samples drops quickly above 80 � C, which is benefi cial for self-healing process. Thus, by varying the feeding ratio of UPydiol/SS-diol, we can obtain tunable mechanical properties at room temperature without affecting the self-healing properties at higher
3.3. Thermal properties The thermal stability of as-prepared TPU samples was investigated through thermogravimetric analysis (TGA), as shown in Fig. 3. The onset degradation temperature (T5, the temperature at 5 wt% of weight loss) is slightly decreased as the feeding ratio of UPy-diol/SS-diol increases. And Table 1 The GPC results of TPU samples. Sample Name
Molar ratio of (UPy-diol:SS-diol)
Mn (Da)
Mw (Da)
Polymer Dispersity Index
PU-0 PU-SS PU-U1S9 PU-U2S8 PU-U3S7 PU-U4S6 PU-U5S5
0:0 0:1 0.1 : 0.9 0.2 : 0.8 0.3 : 0.7 0.4 : 0.6 0.5 : 0.5
9.63 � 104 5.48 � 104 3.31 � 104 2.78 � 104 2.50 � 104 2.23 � 104 2.14 � 104
1.59 � 105 7.20 � 104 5.04 � 104 4.63 � 104 4.06 � 104 3.75 � 104 3.43 � 104
1.65 1.31 1.52 1.66 1.62 1.64 1.60
4
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Fig. 1. The 1H NMR spectra of TPU-elastomers: (a) PU-U2S8, (b) PU-SS.
Fig. 2. (a) FTIR spectra of different TPU samples, and the absorptance is normalized according to the peak height at 1448 cm the band region of stretching vibration.
temperatures.
1
(δa(CH2)). (b) Peak fitting curves in
U2S8, PU-U5S5 clearly show this transition as well. Among PU-UPy-SS samples, the long period interdomain distances of phase separation for PU-U2S8, PU-U3S7, PU-U4S6, PU-U5S5 are 14.9, 13.0, 12.1, 11.2 nm, respectively. As the content of UPy groups is raised, the magnitude of the scattering maximum increases and its location shifts to a higher scat tering vector. A trend of decreasing scale of phase separation is man ifested as the content of UPy groups increases. But on the contrary, the electron density contrast between two phases is increasing, indicating the difference between two phases become more distinguishing with UPy side chains. These results suggest that the incorporation of dangling UPy groups leads to an increase of the ordering of hard segments, while the sizes of hard domains become smaller.
3.5. Morphological structures The microphase separated structure of TPU samples was investigated by atomic force microscopy (AFM) (Fig. 5) and small angle X-ray scat tering (SAXS) (Fig. 6). The AFM image of PU-SS clearly shows the phase segregated hard segments (bright areas) and soft segments (dark areas). The non-structured phase separation scale is about 100–300 nm, out of the measuring limit of SAXS, which is the main reason of absence of scattering peak for PU-SS and PU-U1S9 films. Upon introduction of UPy side groups, the scale of phase separation was reduced to 10–15 nm (according to d ¼ 2π/q, where q ¼ 4π λ 1 sin θ). AFM images of PU5
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Fig. 3. TGA (a) and DTG (b) curves of TPU samples with varying contents of UPy motifs.
Fig. 4. Storage modulus (E0 ), loss modulus (E’’) and tan δ versus temperature of TPU samples.
break (εb) and toughness (integral area of a stress-strain curve) of each film are summarized in Table 2. The PU-SS without the UPy groups exhibits a typical elastomer profile, strengthened by strain induced hardening. Due to the relatively weak nature of the hard phase based on HMDI and SS-diol, PU-SS sample is very soft till the elongation reached a critical point (1000% in Fig. 7a), after which the strain-induce hardening mechanism enhances its me chanical property. Upon introduction of the dangling UPy motifs, the mechanical property of PU-UPy-SS sample is obviously influenced. For example, compared with PU-SS film, the tensile strength of PU-U1S9 is much weaker due to the weakened stability of hard domains by dangling UPy side chains. It is worthy to point out that the low concentration of UPy groups reduces their probability of complimentary associating by quadruple hydrogen bonds. Thus, the adverse effect of dangling side chains plays a major role resulting an excessively weakened tensile strength. A further increase of UPy contents above 20% leads to an increased tendency of forming quadruple hydrogen bonds between UPy motifs in the hard segments, raising the tensile strength of PU-U2S8 and PU-U3S7. Again, strain induced hardening of soft segments still plays an important role after the sample underwent a large deformation. It is noted that this effect is gradually suppressed according to the observa tion of a lowered slope of stress strain curves when the strain is larger
Table 2 The mechanical properties of TPU samples with different contents of UPy groups. Sample name
Tg (� C)
Tensile strength (MPa)
Elongation at break (%)
Toughness (MJ m 3)
PU-SS PU-U1S9 PU-U2S8 PU-U3S7 PU-U4S6 PU-U5S5
7.0 0.1 0.6 1.8 12.8 18.6
13.0 � 0.6 2.1 � 0.3 25.2 � 0.7 22.8 � 0.7 14.6 � 0.1 11.6 � 0.2
1964 � 163 2943 � 142 1463 � 32 1406 � 33 1406 � 27 1215 � 10
64.7 � 0.9 39.9 � 1.5 101.0 � 0.7 108.8 � 0.8 90.3 � 0.2 69.6 � 0.4
3.6. Mechanical properties The mechanical properties of the PU-UPy-SS samples were evaluated by tensile testing. For this aim, 0.4–0.7 mm thick films were cast in a PTFE mold, which were then cut into dumbbell-shape for further tensile measurements. The high mechanical properties of polyurethanes are originated from phase separation, so thermal history will influence their mechanical properties. Thus, all the samples were equilibrated at room temperature for at least 7 days before tensile testing, to reduce the ef fects of thermal history. The typical tensile stress-strain curves are exhibited in Fig. 7a and the obtained tensile strength (σb), elongation at 6
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Fig. 5. SAXS intensity curves of different TPU samples.
To further elucidate the strengthening effect of self-complimentary UPy groups, cyclic loading-unloading test was conducted for TPU films. The typical stress-strain curves were shown in Fig. S10 and the evolving of hysteresis areas with the increasing cycle number was pre sented in Fig. 7b. The films were stretching to a large deformation of 1000% in order to destroy the strengthening sacrificial interactions in the first few cycles. As is shown in Fig. 7b, the first hysteresis area in creases with the increasing UPy content, while the hysteresis area of the last cycle is almost the same. It suggests that the quadruple H-bonding between UPy side chains is effective in reinforcing the tensile strength at small deformation as sacrificial crosslinking point, when putting aside the strain induced hardening effects at large deformation. And due to the dynamic nature of quadruple H-bonding, the opened quadruple hydrogen bonds are able to re-form again as time goes by, restoring the original mechanical properties, which can be accelerated by increasing temperature. 3.7. Self-healing abilities
Fig. 6. AFM phase images of TPU samples: (a) PU-SS, (b) PU-U2S8, (c) PU-U5S5.
The self-healing abilities of PU-UPy-SS samples were first evaluated intuitively by simple healing and stretching test. A dumbbell-shaped sample (PU-U2S8) was fully cut in half and healed at 100 � C for 2 h. As shown in Fig. 8, the specimen was fully healed and was able to sustain a high stretching ratio. And the fracture position was not located at the healing surface once it was snapped again, suggesting its mechanical property was full recovered. Then the self-healing process was observed by a scratch healing test, monitored by an optical microscope equipped with a Linkam hot stage. Scratches of 10–20 μm wide and 200–400 μm deep were made by scalpel on the samples. These samples were then
than 1000%. There are two possible reasons, which are the increased structure irregularity and the retarding effect of strong quadruple bonding. Therefore, the final strength shows a decreasing trend when the UPy content is above 30 mol%, even though the stress at lower strain is increasing. These observations indicate that only incorporation of a suitable amount of UPy side chains can effectively in enhance the me chanical properties of intrinsically weak self-healing TPU samples.
Fig. 7. (a) Typical stress-strain curves of TPU samples; (b) Hysteresis area of TPU samples under five successive loading-unloading cycles. 7
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Fig. 8. Optical microscope images after different healing times at 100 � C of scratched samples: (a) PU-SS, (b) PU-U2S8, (c) PU-U5S5.
placed on a hot stage to observe the healing of the cracks at constant temperatures. The optical images are shown in Fig. 9 and Fig. S10. The cracks of the surface closed in 10 min and then the traces become shallower and gradually disappear after 30 min. This is because the exchanging of disulfide bonds, the reversible associating of quadruple hydrogen bonds and the molecular chain diffusion of hard segments were all fully activated at 100 � C. For comparison, no sign of healing of the cracks on the film of PU-0 sample was observed after 30 min at 100 � C. As for PU-U1S9, the self-healing processes proceeds quickly even at 50 � C as shown in Fig. S10a. The main reason is the hard domains are
destroyed by the non-associating dangling UPy side chains at low con centration, accelerating the molecular chain diffusion. To quantitively evaluate the self-healing efficiencies, tensile testing was further applied to investigate the self-healing efficiencies of the samples at different temperatures. For this aim, a 0.4–0.7 mm thick dumbbell-shaped film was fully cut in half using a clean scalpel, and then gently pressed together at room temperature. They were allowed to heal at 40, 60, 80 and 100 � C for 2 h without any external pressure. The stress strain curves obtained for different samples at various healing temperatures were displayed in Fig. 10 and Fig. S11. The healing
Fig. 9. Two cut pieces coming from two different dog-boned films, one stained with methyl blue for clarifying, were allowed to heal at 100 � C for 2 h, and then stretched to observe the effectiveness of self-healing, after which the healed specimen was snapped again. The healed specimen did not break at the original healing surfaces, suggesting it was fully healed. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.) 8
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efficiency of PU-U1S9 was not included, because valid data could not be obtained due to the soft, sticky nature and unacceptable repeatability of mechanical properties after healing. First, after 2 h of self-healing at 100 � C, all samples were completely healed (healing efficiencies are greater than 90%). And they showed a significant extent of healing ef ficiency at 80 � C for 2 h, meaning the healing processes begins at 80 � C. But the full healing requires a prolonged time. It was noted that even at 40 � C, the samples were partially healed. The main driving force behind it was the weak hydrogen bonds between urethane or urea bonds. Therefore, besides quadruple hydrogen bonds and disulfide bonds, these weak hydrogen bonds contribute to the self-healing process as well. As shown in Table 3, compared with PU-SS, the healing efficiency of PU-U2S8 and PU-U3S7 were lowered. It can be attributed to the extra diffusion friction brought by quadruple hydrogen bonds. At lower temperatures, such as 40, 60 and 80 � C the healing efficiency of tensile strength gradually increases with the UPy contents while the healing efficiency of tensile strain drops down. These results indicate that UPy groups accelerate the healing of tensile strength but hampers the healing of stretching ability. This is due to the re-forming of strong quadruple hydrogen bonds contributing to the effective binding of damage sur faces. However, the high dimerization constant of UPy groups slows down the chain diffusion and randomization processes of molecular chains from both sides of damage surfaces, resulting in the decreased healing efficiency of tensile strain. To reduce this effect, UPy groups were introduced as side chains to accelerate chain diffusion when the dimers of UPy dissociate at elevated temperatures. For further understanding of the self-healing mechanism of this system, the rheology behaviors of PU-SS, PU-U2S8 and PU-U5S5 were measured as depicted in Fig. S12. The flow transition temperatures (the crossover temperature of G’and G’’ in the temperature ramping data) of PU-SS, PU-U2S8 and PU-U5S5 were 100 � C, 110 � C and 114 � C, respec tively. The flow transition temperature was only slightly increased. To better understand the effect of UPy side chain on the flow behaviors and relaxation time, the rheology frequency sweeping tests were then con ducted at 100 � C. As indicated in Fig. S12, the relaxation times (τ ¼ 2π/
ωc at the crossover point of G0 and G’’ in the frequency sweeping data) of
PU-SS, PU-U2S5 and PU-U5S5 were 0.28 s, 0.95 s and 1.45 s, respec tively. The relaxation time was increasing with introduction of UPy groups, but to a negligible extent. The small relaxation time of all samples at high temperatures is due to the low bond life time of disulfide bond and quadruple hydrogen bond, providing the PU-UPy-SS samples with efficient healing ability at elevated temperatures. Therefore, the self-healing ability of TPU sample was not significantly compromised while the mechanical property was strengthened by the introduction of UPy side groups. 4. Conclusions In this work, we have successfully fabricated mechanically enhanced self-healing TPU elastomers by integration of dangling UPy side group into the dynamic disulfide main chain. The variation of the feeding ratio of UPy-diol/SS-diol could be used to tailor the properties of this mate rial. The introduction of UPy side groups possess two opposite effects. One is the plasticizer effect of side chain, resulting a decreased onset degradation temperature, and increased chain mobility at elevated temperatures. The other is the thermo-triggered reversible crosslinking effect originating from the self-complimentary quadruple H-bonding of UPy motifs, increasing the crosslinking density of hard domains and thereby enhancing its mechanical robustness at room temperature. The optimization results revealed that using 20–30 mol % of UPy-diol chain extender led to a better result in balancing the mechanical robustness and self-healing properties, obtaining a tensile strength up to 25 MPa and a healing efficiency over 92% in 2 h at 100 � C. For future investi gation, the enhancing strategy of introducing UPy side groups can be applied to various self-healing systems incorporating different dynamic covalent bonds in the main chain. Moreover, the wide commercial availability of raw materials and the simple synthetic processes make this system very attractive for real industrial applications.
Fig. 10. The typical stress-strain curves of different TPU samples after 2 h healing at various temperature: (a) PU-SS; (b) PU-U2S8; (c) PU-U3S7; (d) PU-U4S6; (e) PU-U5S5. 9
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Polymer 183 (2019) 121912
Table 3 The healing efficiencies of TPU samples healed at various temperatures for 2 h. Sample name PU-SS PU-U2S8 PU-U3S7 PU-U4S6 PU-U5S5
Healing efficiency of tensile strength (%)
Healing efficiency of elongation at break (%)
Healing efficiency of toughness (%)
40
60
80
100
40
60
80
100
40
60
80
100
25.0 15.0 17.8 28.3 49.0
52.7 39.3 41.2 46.5 69.3
93.8 76.0 79.1 90.1 93.3
98.5 92.4 94.2 99.7 96.6
71.1 46.9 44.3 34.2 33.2
87.0 73.6 72.7 58.5 53.2
96.8 92.3 90.0 89.3 86.6
102.9 99.9 98.3 95.9 99.9
33.8 14.6 13.1 12.9 18.9
55.8 35.1 41.0 32.0 49.3
97.4 74.6 74.4 78.2 87.0
105.2 99.5 94.5 94.3 99.4
Declaration of competing interestCOI
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