A dynamic hydrogen-bonded crosslinked polymer, and a preparation method and application thereof

Abstract

The application discloses a dynamic hydrogen bond cross-linked polymer and a preparation method and application thereof. The structural formula of the dynamic hydrogen bond cross-linked polymer is as follows: wherein x is an integer of 30-200, and y is an integer of 100-800. The dynamic hydrogen bond cross-linked polymer has both low-speed high-efficiency energy dissipation and high-speed transient hardening mechanism, can maintain excellent impact resistance in a wide strain rate range, is suitable for being used as an impact-resistant protective material, and has simple preparation method, low production cost, and is suitable for large-scale industrial production and application.

Description

Technical Field

[0001] This invention relates to the field of functional polymer materials technology, specifically to a dynamic hydrogen-bonded crosslinked polymer, its preparation method, and its applications. Background Technology

[0002] In practical applications such as sports protection, automotive industry, and flexible bulletproofing, impact-resistant protective materials face impact velocities spanning multiple orders of magnitude (e.g., the impact velocity of a low-speed drop is approximately 0.1 s). -1 ~10s -1 The impact rate of a high-speed collision can reach as high as 100 seconds. -1 ~10000s -1 This raises the core performance requirement for impact protection materials: wide-rate adaptability protection.

[0003] Existing impact-resistant materials mainly include the following three systems, but these types of impact-resistant materials all have obvious limitations and cannot fully meet the needs of practical applications, as detailed below: 1) High-strength and high-toughness materials (e.g., chain network topology type, intermolecular force-dominated type, ordered microstructure type and orientation-induced type): These materials exhibit excellent impact resistance in a specific rate range, but their mechanical response mechanism is relatively fixed and it is difficult to achieve adaptive adjustment with changes in external impact rate. At low impact rates, there is insufficient activation of dissipation mechanism, while at high impact rates, there is a problem that the structure is prone to brittle fracture, and the rate adaptation boundary is narrow. 2) Shear-thickening fluid (STF): Although this type of material has rate response characteristics, it suffers from problems such as solvent evaporation and particle sedimentation. Its long-term stability is poor, and its liquid form limits its structural integration applications. 3) Shear-thickening gels (SSG) and dynamic hydrogen-bonded materials: These materials overcome the problem of poor stability in shear-thickening fluids. However, existing hydrogen-bonding units (such as urea, thiourea, 2-ureido-4[1H]-pyrimidinone, etc.) not only have an inherent contradiction between hydrogen bond binding energy and dynamic performance (for example, increasing hydrogen bond binding energy can enhance high-speed rigid response, but it will also simultaneously raise the dissociation energy barrier and reduce dynamic performance, thus suppressing low-speed dissipation ability; and vice versa), but also have the problem of low molecular chain orientation transfer efficiency, making it difficult to achieve uniform stress distribution and efficient energy dissipation simultaneously during impact, which makes the construction of wide-rate protection systems face a fundamental bottleneck.

[0004] Therefore, there is an urgent need to develop an impact-resistant material that can balance low-speed, high-efficiency energy dissipation with high-speed transient hardening and stable impact resistance over a wide strain rate range. Summary of the Invention

[0005] The purpose of this invention is to provide a dynamically hydrogen-bonded crosslinked polymer, its preparation method, and its application.

[0006] The technical solution adopted in this invention is: A dynamically hydrogen-bonded crosslinked polymer has the following structural formula: In the formula, x takes an integer from 30 to 200, and y takes an integer from 100 to 800.

[0007] Preferably, the number-average molecular weight of the dynamically hydrogen-bonded crosslinked polymer is 60,000 g / mol to 200,000 g / mol.

[0008] A method for preparing the dynamically hydrogen-bonded crosslinked polymer as described above includes the following steps: A reversible addition-fragmentation chain transfer radical polymerization was carried out in an organic solvent using a bis(trifluoromethyl)-containing squamamide, methoxydiethylene glycol methacrylate (MeO2MA), a chain transfer agent, and an initiator. The bis(trifluoromethyl)-containing squamamide was... (Sq-2CF3) was then separated and purified to obtain a dynamically hydrogen-bonded cross-linked polymer.

[0009] Preferably, the molar ratio of the bis(trifluoromethyl) squaramide, methoxydiethylene glycol methacrylate, chain transfer agent, and initiator is 1–2:1–20:0.02–0.7:0.002–0.35.

[0010] Preferably, the chain transfer agent is at least one of 4-cyano-4-[(dodecylthioalkylthiocarbonyl)thioalkyl]valerate (CAS No.: 870196-80-8), 4-cyano-4-(dodecyltrithiocarbonyl)valerate-2-(norbornene-1,2-dicarboximide)ethyl ester, 2-cyanopropyl-2-ylphenyl dithiocarbamate (CAS No.: 201611-85-0), and 4-cyano-4-(phenylthiocarbamate)valerate (CAS No.: 201611-92-9).

[0011] Preferably, the 4-cyano-4-(dodecyltrithiocarbonate)pentanoic acid-2-(norbornene-1,2-dicarboximide) ethyl ester is prepared by a method comprising the following steps: a) Reaction of cis-5-norbornene-endo-2,3-dicarboxylic anhydride and 3-amino-1-propanol yields N-(3-hydroxypropyl)-cis-5-norbornene-endo-2,3-dicarboximide; b) The reaction of N-(3-hydroxypropyl)-cis-5-norbornene-endo-2,3-dicarboximide and 4-cyano-4-[(dodecylthiocarbonyl)thioalkyl]valerate yields 4-cyano-4-(dodecyltrithiocarbonate)valerate-2-(norbornene-1,2-dicarboximide)ethyl ester.

[0012] Preferably, the initiator is at least one of azobisisobutyronitrile (AIBN), benzoyl peroxide (BPO), and azobisisoheptanenitrile (ABVN).

[0013] Preferably, the organic solvent is at least one selected from dimethylacetamide (DMAc), N,N-dimethylformamide (DMF), N-methylpyrrolidone (NMP), and dimethyl sulfoxide (DMSO).

[0014] Preferably, the reversible addition-fragmentation chain transfer radical polymerization is carried out at a temperature of 70℃ to 80℃ for a reaction time of 5h to 10h.

[0015] An impact-resistant protective material comprising the aforementioned dynamically hydrogen-bonded crosslinked polymer.

[0016] A composite impact-resistant protective material comprises an aramid fiber cloth and an impact-resistant film adhered to both sides of the aramid fiber cloth, wherein the impact-resistant film is made of the aforementioned dynamic hydrogen-bonded crosslinked polymer.

[0017] Preferably, the thickness of the impact-resistant film is 1 mm to 4 mm.

[0018] A method for preparing a composite impact-resistant protective material as described above includes the following steps: 1) The dynamically hydrogen-bonded cross-linked polymer is injected into a mold and then hot-pressed to obtain an impact-resistant film; 2) The impact-resistant film, aramid fiber cloth and impact-resistant film are stacked in sequence and then hot-pressed to obtain a composite impact-resistant protective material.

[0019] Preferably, the hot pressing in step 1) is carried out under conditions of pressure of 20MPa to 200MPa and temperature of 90℃ to 110℃, and the holding time is 10min to 30min.

[0020] Preferably, the hot pressing in step 2) is carried out under conditions of pressure of 50MPa to 200MPa and temperature of 90℃ to 110℃, and the holding time is 15min to 60min.

[0021] Applications of a dynamically hydrogen-bonded crosslinked polymer as described above in the fields of sports protection, automotive industry, or flexible bulletproofing.

[0022] The beneficial effects of the present invention are: the dynamic hydrogen bond crosslinking polymer of the present invention has both low-speed and high-efficiency energy dissipation and high-speed transient hardening mechanism, which can maintain excellent impact resistance over a wide strain rate range, making it suitable for use as an impact-resistant protective material. Moreover, its preparation method is simple and the production cost is low, making it suitable for large-scale industrial production and application.

[0023] Specifically: 1) The dynamic hydrogen bond crosslinking polymer of the present invention uses a square amide with a rigid planar four-membered ring structure as a dynamic hydrogen bond donor. After the formation of its bidentate hydrogen bonds, the aromaticity within the ring is enhanced, which can endow the hydrogen bond network with high binding energy and high dissociation energy barrier, thus overcoming the contradiction of "strong bonding - fast dynamic" in traditional hydrogen bonds: under low and medium speed impact, hydrogen bonds can undergo reversible dissociation and recombination to achieve efficient energy dissipation, while under high speed impact, hydrogen bonds cannot dissociate in time and are "frozen" to provide instantaneous rigidity. At the same time, the high dissociation energy barrier can lock the strain-induced molecular chain orientation structure, thereby achieving continuous adaptive impact protection from dissipation to hardening over a wide strain rate range. 2) The dynamic hydrogen-bonded crosslinking polymer of this invention uses a square amide with a rigid planar topology and strong directional hydrogen bonds, resulting in excellent orientation transfer efficiency and orientation stability: the orientation degree of the square amide groups (A0 / A0) 90 =40%) is significantly higher than that of the urea (80%) and thiourea (100%) systems, and the stretch-induced orientation structure can be effectively locked by the hydrogen bond network and the deorientation is suppressed, so that the polymer can achieve continuous and efficient stress transmission and energy dissipation under complex loads. 3) The dynamic hydrogen-bonded crosslinked polymer of this invention can leverage the synergistic effect of the dynamic hydrogen bond network of squamamide and the flexible polyether segments, resulting in the integrated application of multiple built-in functions: the polymer can heal micron-level scratches within 1 hour at room temperature without external stimulation (self-healing function), and it exhibits broad-spectrum strong adhesion to a variety of substrates from non-polar to polar (adhesion strength increases exponentially with increasing squamamide content). Furthermore, when combined with aramid fibers, it enables tear energy exceeding 600 kJ / m through interfacial bridging of hydrogen bonds. 2 It is significantly superior to pure polymer matrix or pure aramid fiber materials; 4) This invention employs reversible addition-fragmentation chain transfer radical polymerization to precisely control the content of squaramide and the molecular weight of the polymer, resulting in the optimal matching of chain entanglement and dynamic hydrogen bond network: Under the raw material ratio of this invention, the copolymer can exhibit a suitable glass transition temperature near room temperature, with moderate static mechanical properties, while the dynamic mechanical properties reach the optimal balance over a wide strain rate range, providing a clear structural parameter window for practical applications. Attached Figure Description

[0024] Figure 1The image shows the hydrogen nuclear magnetic resonance spectrum of the dynamically hydrogen-bonded crosslinked polymer of Example 1.

[0025] Figure 2 The image shows the hydrogen nuclear magnetic resonance spectrum of the dynamically hydrogen-bonded crosslinked polymer from Example 2.

[0026] Figure 3 The image shows the hydrogen nuclear magnetic resonance spectrum of the dynamically hydrogen-bonded crosslinked polymer of Example 3.

[0027] Figure 4 The graph shows the quasi-static mechanical property test results of the dynamically hydrogen-bonded crosslinked polymer in Example 1.

[0028] Figure 5 The results of polarization color testing of the dynamically hydrogen-bonded crosslinked polymers of Examples 1 and 3 after stretching are shown in the figure.

[0029] Figure 6 The figures show the test results of the falling ball impact protection of the dynamically hydrogen-bonded crosslinked polymers in Examples 1 and 3.

[0030] Figure 7 The figure shows the test results of the dynamic compressive mechanical properties of the dynamically hydrogen-bonded crosslinked polymer in Example 1 under high strain rate.

[0031] Figure 8 The figure shows the room temperature self-healing test results of the dynamically hydrogen-bonded crosslinked polymer in Example 1.

[0032] Figure 9 The figure shows the tear resistance test results of the composite impact-resistant protective material in Example 4. Detailed Implementation

[0033] The present invention will be further explained and described below with reference to specific embodiments.

[0034] Example 1: A dynamically hydrogen-bonded crosslinked polymer is prepared as follows: 1 mmol of Sq-2CF3, 15 mmol of MeO2MA, 0.032 mmol of 4-cyano-4-[(dodecylthioalkylthiocarbonyl)thioalkyl]valeric acid, and 0.016 mmol of AIBN were dissolved in 3.4 mL of DMAc and reacted at 75 °C for 6 h under an Ar atmosphere. After cooling to room temperature, the reaction solution was slowly added dropwise to diethyl ether to precipitate the product. After centrifugation, the solid was dissolved in DMAc, and the dissolution and precipitation operations were repeated twice each. After centrifugation, the solid was vacuum dried to obtain a dynamically hydrogen-bonded crosslinked polymer (denoted as PMOS-1; a light yellow solid with a number average molecular weight of 90,000 g / mol to 100,000 g / mol).

[0035] The 1H NMR spectrum of the dynamically hydrogen-bonded crosslinked polymer in this embodiment is shown below. Figure 1 As shown.

[0036] Depend on Figure 1 It can be seen that the molar ratio of Sq-2CF3 units to MEO2MA units in the polymer is 1:19, based on the ratio of protons on the Sq-2CF3 benzene ring (chemical shift value at 8.05, number 2, integral value divided by 2 represents the number of square amides) to the protons on the MEO2MA terminal methyl group (chemical shift value at 3.26, number 3, integral value divided by 3 represents the number of square amides). The small peak (number 2) near chemical shift value at 6.30 is a characteristic peak of the polymer's terminal group, with only one characteristic peak for a single terminal group per molecular chain. In summary, the molecular weight of the polymer can be further confirmed by the characteristic peaks and the molar ratio of Sq-2CF3 units to MEO2MA units.

[0037] Example 2: A dynamically hydrogen-bonded crosslinked polymer is prepared as follows: 1 mmol of Sq-2CF3, 15 mmol of MeO2MA, 0.05 mmol of 4-cyano-4-[(dodecylthioalkylthiocarbonyl)thioalkyl]valeric acid and 0.025 mmol of AIBN were dissolved in 3.4 mL of DMAc and reacted at 75 °C for 6 h under an Ar atmosphere. After cooling to room temperature, the reaction solution was slowly added dropwise to diethyl ether to precipitate the product. After centrifugation, the solid was dissolved in DMAc, and the dissolution and precipitation operations were repeated twice each. After centrifugation, the solid was vacuum dried to obtain a dynamically hydrogen-bonded crosslinked polymer (denoted as PMOS-2; a light yellow solid with a number average molecular weight of 60,000 g / mol to 70,000 g / mol).

[0038] The 1H NMR spectrum of the dynamically hydrogen-bonded crosslinked polymer in this embodiment is shown below. Figure 2 As shown.

[0039] Depend on Figure 2 It can be seen that the molar ratio of Sq-2CF3 units to MEO2MA units in the polymer is 1:19, based on the ratio of protons on the Sq-2CF3 benzene ring (chemical shift value at 8.05, number 2, integral value divided by 2 represents the number of square amides) to the protons on the MEO2MA terminal methyl group (chemical shift value at 3.26, number 3, integral value divided by 3 represents the number of square amides). The small peak (number 2) near chemical shift value at 6.30 is a characteristic peak of the polymer's terminal group, with only one characteristic peak for a single terminal group per molecular chain. In summary, the molecular weight of the polymer can be further confirmed by the characteristic peaks and the molar ratio of Sq-2CF3 units to MEO2MA units.

[0040] Example 3: A dynamically hydrogen-bonded crosslinked polymer is prepared as follows: 1 mmol of Sq-2CF3, 10 mmol of MeO2MA, 0.04 mmol of 4-cyano-4-[(dodecylthioalkylthiocarbonyl)thioalkyl]valeric acid and 0.02 mmol of AIBN were dissolved in 4 mL of DMAc and reacted at 75 °C for 6 h under an Ar atmosphere. After cooling to room temperature, the reaction solution was slowly added dropwise to diethyl ether to precipitate the product. After centrifugation, the solid was dissolved in DMAc, and the dissolution and precipitation operations were repeated twice each. After centrifugation, the solid was vacuum dried to obtain a dynamically hydrogen-bonded crosslinked polymer (denoted as PMOS-3; a light yellow solid with a number-average molecular weight of 90,000 g / mol to 100,000 g / mol).

[0041] The 1H NMR spectrum of the dynamically hydrogen-bonded crosslinked polymer in this embodiment is shown below. Figure 3 As shown.

[0042] Depend on Figure 3 It can be seen that the molar ratio of Sq-2CF3 units to MEO2MA units in the polymer is 1:12, based on the ratio of protons on the Sq-2CF3 benzene ring (chemical shift value at 8.05, number 2, integral value divided by 2 represents the number of square amides) to the protons on the MEO2MA terminal methyl group (chemical shift value at 3.26, number 3, integral value divided by 3 represents the number of square amides). The small peak (number 2) near chemical shift value at 6.30 is a characteristic peak of the polymer's terminal group, with only one characteristic peak for a single terminal group on each molecular chain. In summary, the molecular weight of the polymer can be further confirmed by the characteristic peaks and the molar ratio of Sq-2CF3 units to MEO2MA units.

[0043] Example 4: A composite impact-resistant protective material, the preparation method of which is as follows: 1) The dynamic hydrogen-bonded crosslinked polymer (PMOS-1) of Example 1 was injected into a mold, and then hot-pressed at a pressure of 20 MPa and a temperature of 100 °C for 10 min to obtain an impact-resistant film (thickness of 2 mm). 2) The impact-resistant film, aramid fiber cloth (1 mm thick) and the impact-resistant film are stacked in sequence, and then hot-pressed at a pressure of 50 MPa and a temperature of 100℃ for 15 min to obtain a composite impact-resistant protective material (denoted as PMOS / Kevlar).

[0044] Performance testing: 1) Following the ASTM D412 standard, the dynamic hydrogen-bonded crosslinked polymer (PMOS-1) from Example 1 was cast into dumbbell-shaped specimens (size of the tensile portion: width 10 mm, thickness 1 mm) and cylindrical specimens (size of the tensile portion: diameter 10 mm, height 5 mm). Quasi-static mechanical property tests were then conducted using an Instron universal testing machine (equipped with a force sensor of appropriate range). The tests included both tensile and compressive modes, and the specific testing procedures are as follows: a) Tensile test: Clamp the dumbbell-shaped specimen in the upper and lower wedge-shaped fixtures (ensuring the loading axis is strictly aligned with the specimen axis to avoid initial stress concentration caused by eccentricity). Set the strain rate as follows: In the Bluehill software, calculate and set the corresponding nominal strain rate to 0.05s based on the gauge length. -1 0.08s -1 0.67s -1 and 2s -1 Four sets of crosshead displacement rates were used, and at least three parallel specimens were tested for each rate. Data acquisition: The testing machine recorded the load and displacement in real time, converted them into engineering stress and engineering strain, until the specimen broke, and the fracture location and fracture morphology were recorded. b) Compression test: Place the cylindrical specimen in the center of the compression platform, ensuring the upper and lower pressure plates are in full contact with the specimen end faces (apply a thin layer of lubricating grease to the contact surfaces to reduce end-face friction). Set the strain rate as follows: set the corresponding nominal strain rate to 0.002 s. -1 0.008s -1 0.02s -1 and 0.1s -1 Four sets of compression rates (Note: the compression rate is generally lower than the tensile rate to accommodate the response characteristics of dynamically hydrogen-bonded crosslinked polymers in compression mode); Data acquisition: Force-displacement curves were continuously acquired, converted into engineering stress-strain, and unloaded after being compressed to a preset strain (approximately 0.8), and residual deformation was recorded; Environmental control: The entire process was conducted at room temperature (23℃±2℃) and standard humidity conditions, and sufficient time was given for the samples to recover between each set of tests to eliminate the influence of residual strain; Quasi-static mechanical property test results are as follows Figure 4 (a is the tensile test, b is the compression test) as shown.

[0045] Depend on Figure 4 From 'a', we can know that: Initial elastic segment: All four sets of curves showed a rapid linear increase in the small strain stage (strain < 2), and the slope increased significantly with the increase of strain rate, indicating that the initial stiffness of PMOS-1 has obvious rate correlation, reflecting the rate-dependent transformation of Sq-2CF3 hydrogen bonds from dynamic reversible dissociation to instantaneous frozen crosslinking. At higher loading rates, the molecular chains or internal microstructures do not have time to relax and rearrange, resulting in an increase in macro modulus. This is consistent with the precursor response of the impact hardening mechanism of PMOS-1 at the static level. Tough flow stage (0.05s) -1 ~0.08s -1 The stress reaches its peak around 0.2MPa to 0.4MPa and then slowly decreases, entering a low-stress steady-state platform. It does not fracture even when the strain exceeds 40, exhibiting extremely high ductility. The reason is that the hydrogen bond network and physical cross-linking points inside PMOS-1 have sufficient time to dissociate and reconstruct under low-speed loading. The energy dissipation mechanism is dominant, and the material absorbs energy through flow deformation. Softening and stretching stage (0.67s) -1 The peak stress rises to about 1.7 MPa, followed by significant stress softening and sustained large deformation extension, indicating that the network segments are fully extended, the crosslinks are gradually dissociated, and the network enters a steady-state rheological response, possessing both strength and ductility. The dynamic reconstruction and locking effect of the hydrogen bond network reach a balance. Brittle fracture stage (2s) -1 The stress rises sharply to about 3.0 MPa, and the fracture occurs directly after the peak (marked with "×" in the figure). There is no stress softening segment. The reason is that the hydrogen bond network does not have time to rearrange dynamically during rapid stretching. Local stress concentration leads to brittle fracture and a significant decrease in fracture toughness. This is consistent with the mechanism of necking concentration caused by local hardening in impact hardened materials under high-speed loading. The essence of the above three-stage transformation lies in the time-scale competition of the dynamic cross-linked network of hydrogen bonds: at low speeds, reversible network reconstruction dominates dissipation, and both modulus and stress are low; as the rate increases, the network reconstruction time window is compressed, hydrogen bonds gradually "freeze" and lock the chain segment orientation, and the energy storage mechanism replaces the dissipation mechanism, which is manifested as a significant increase in modulus and peak stress.

[0046] It is worth noting that the peak stress is from 0.05s -1 The pressure rose from approximately 0.3 MPa to 2 seconds. -1 The increase of approximately 3.0 MPa, representing an increase of about one order of magnitude, demonstrates significant positive strain rate sensitivity, which is a direct manifestation of the impact hardening nature of PMOS-1 in the quasi-static tensile mode.

[0047] Depend on Figure 4 From b, we can know that: Initial nonlinear rising segment (ε<0.6): Unlike the tensile curve, the compression curve does not have a clear yield plateau, but instead shows a continuous nonlinear hardening rising trend. The stress increases monotonically with strain, and the slope gradually increases. This reflects that in the compression mode, the internal hydrogen bond cluster network of PMOS-1 becomes increasingly dense with the accumulation of deformation. During the lateral convergence of chain segments, the hydrogen bonds dynamically dissociate and recombine, forming new lateral crosslinking points to resist slippage. The load-bearing network becomes more and more perfect, and the material continues to harden rather than soften, which is highly consistent with its intrinsic characteristics of impact hardening. Rapid hardening phase (ε>0.6): When the strain exceeds approximately 0.6–0.7, the stress in all curves increases sharply, reaching its highest rate (0.1 s⁻¹) at approximately 0.8 strain. -1 When the stress reaches about 60 MPa, the hydrogen bond network inside the material reaches a highly dense state under high compression ratio. The hydrogen bonds lock the chain segments in the compression-induced local conformation, preventing unloading and springback. The degree of freedom of the chain segments is extremely restricted, which is macroscopically manifested as a surge in stiffness. This mechanism of hydrogen bonds working together to lock the chain segment conformation and drive the dense network to resist volume compression under large deformation is the core mechanical response of PMOS-1 impact hardening. Strain rate effect: The compressive stress also increases monotonically with the strain rate, which is particularly significant in the densification stage. The stress difference between the highest and lowest strain rates is about 4 to 5 times, which further verifies the viscoelastic / impact hardening constitutive response of PMOS-1. Compared with the tensile mode, the separation of the various rate curves in the large strain range is more severe in the compression mode, indicating that the impact hardening effect of PMOS-1 is more prominent under compressive load. The root cause is that the hydrogen bond recombination time window is drastically compressed during high-speed compression, the network is more "frozen", the densified structure is more difficult to relax, and the energy storage efficiency is significantly improved.

[0048] Tensile and compressive tests together reveal the core mechanical characteristics of PMOS-1: Tensile-compressive asymmetry: The compressive strength (approximately 60 MPa) far exceeds the tensile strength (approximately 3 MPa). Compression is characterized by continuous hardening driven by the densification of the hydrogen bond network, while tension is characterized by yielding and softening mediated by hydrogen bond dissociation and recombination before entering a steady state. The dynamic behavior of hydrogen bonds is completely different in the two modes—during tension, dissociation and recombination are carried out along the loading direction to stabilize the chain segment orientation, while during compression, the chain segment conformation is locked laterally to resist volume densification—reflecting the anisotropic response behavior of the microstructure of PMOS-1 under tensile and compressive loads. Static precursor characteristics of rate-dependent impact hardening: Although this experiment is a quasi-static test, PMOS-1 has shown significant positive strain rate sensitivity, and the higher the rate, the more severe the compression hardening. This indicates that its impact hardening ability originates from the intrinsic viscoelastic mechanism of dynamic cross-linking of hydrogen bonds, rather than a simple inertial effect: during tension, hydrogen bond locking orientation enables chain segments to cooperate in bearing load at high speeds, and during compression, the hydrogen bond densification network causes the material stiffness to surge under volume compression. This conclusion provides an important static benchmark reference for the results of subsequent dynamic impact testing (SHPB). Energy absorption potential: The ultra-large elongation at break mediated by hydrogen bonds under tension and the wide hardening region driven by hydrogen bond densification under compression together endow PMOS-1 with excellent energy absorption capabilities. As a reversible cross-linking tool, hydrogen bonds dissipate or store energy in both deformation modes through orientation locking and network densification, respectively. Combined with its impact hardening characteristics, PMOS-1 has great application potential in the field of impact-resistant protective materials.

[0049] 2) Following ASTM D412 standards, the dynamically hydrogen-bonded crosslinked polymers (PMOS-1 and PMOS-3) from Examples 1 and 3 were cast into dumbbell-shaped specimens (the tensile portion was 10 mm wide and 1 mm thick). A birefringence observation device was then used in conjunction with an Instron universal testing machine to achieve real-time optical imaging of the orientation field during tensile testing. The optical path was as follows: white LED parallel light source → polarizer → dumbbell-shaped specimen → detector (90° orthogonal to the polarizer) → camera image acquisition, polarization... The polarizer and detector are fixed on the optical support, with the polarization axis direction fixed (45° to the tension axis). The dumbbell-shaped sample (before testing, confirm that the initial state of the sample is completely dark (extinct) between the orthogonal polarizers to prove that the material is optically isotropic when not under stress and there is no pre-orientation) is placed in the middle of the two polarizers (ensuring that parallel light uniformly illuminates the entire gauge length). Testing procedure: The dumbbell-shaped sample is clamped in the Instron testing machine fixture, the optical device is aligned with the gauge length, and the test is carried out under the following three conditions: PMOS-1 / 0.04s -1 PMOS-1 / 0.67s -1 PMOS-3 / 0.67s -1 Simultaneously with the initiation of tensile loading, the camera synchronously records and acquires birefringence images at specific tensile strain nodes to obtain real-time visualization results of the molecular chain orientation field. Semi-quantitative analysis is then performed using color scales (0–1500, corresponding to relative intensity of optical path difference). The polarization color test results are as follows: Figure 5 As shown.

[0050] Depend on Figure 5 It can be known that: a) PMOS-1 is uniformly low-grade blue-white in the gauge length section, and the orientation distribution is symmetrical along the axis with no obvious gradient. This indicates that under low strain rate, the molecular chain of PMOS-1 has enough time to slowly orient with external force while relaxing through dynamic dissociation and recombination of hydrogen bonds. The macroscopic orientation degree is maintained at a low level and is uniformly distributed. b) At higher speeds, the brightness and order of the PMOS-1 stripes are significantly increased, and the color is brighter bluish-white. This indicates that the molecular chain orientation is significantly higher at the same strain. At higher speeds, the hydrogen bonds mediated by the amide do not have time to dynamically recombine and relax, and the orientation is effectively "frozen". The birefringence signal is enhanced, which directly reflects the regulation of the degree of chain segment orientation retention by the strain rate. c) PMOS-3 exhibits rich high-order sub-color stripes at the same high speed, with color mark readings close to the upper limit of 500-650, far exceeding PMOS-1. Moreover, the stripe order is highest in the middle of the gauge segment and gradually decreases towards the clamping end, with a significant orientation gradient. In summary, the molecular chain orientation behavior of the dynamically hydrogen-bonded crosslinked polymer of this invention is synergistically regulated by the strain rate and the content of squamamide: the higher the strain rate, the more suppressed the dynamic relaxation of hydrogen bonds, and the more fully the orientation is preserved; the higher the content of squamamide, the stronger the hydrogen bond network constraint, the higher the degree of orientation, and the more significant the distribution gradient (mechanism: the squamamide group is the core donor / acceptor unit of the hydrogen bond network in the dynamically hydrogen-bonded crosslinked polymer system, and the higher the content, the greater the hydrogen bond density per unit volume and the more dynamic crosslinking points; during tensile deformation, the denser hydrogen bond network exerts a stronger constraint on the molecular chain, forcing the chain segments to align in the orientation direction). The more fully extended arrangement, and the higher density of hydrogen bonds after orientation, the stronger synergistic locking effect, which inhibits orientation relaxation, greatly enhances the birefringence signal. This is also the microstructure root of the dynamic hydrogen bond crosslinked polymer with high squamamide content having stronger impact hardening ability—the higher the hydrogen bond network density, the more stable the orientation-ordered state formed under rapid loading, and the more intense the macroscopic stiffening response. The birefringence test provides direct optical experimental evidence for the impact hardening mechanism of the dynamic hydrogen bond crosslinked polymer of this invention from the perspective of micro-orientation, revealing that the orientation locking driven by hydrogen bond network density is the structural basis of its macroscopic mechanical hardening response.

[0051] 3) The test materials POS (25%), OBC 9530 (Dow), PMOS-1, PMOU (8%), PMOS (2%), PMOS-3, silicone, NBR (nitrile rubber), and rubber were laid flat on top of the coverslip and placed on a uniform rigid base. Then, a steel ball with a mass of 111g was dropped from different heights to impact the material surface. The minimum critical height h at which the coverslip just breaks brittlely was determined by the lifting method. c (Unit: m), record the gravitational potential energy E of the steel ball at this height. c=m×g×h×c (m=0.111kg, g=9.81m / s 2 Then divide this energy by the material thickness t (in mm) measured by a thickness gauge with an accuracy of 0.01 mm to obtain the energy absorbed per unit thickness E. thick =E c / t is used to quantitatively evaluate the material's performance in low-velocity impact (strain rate 0.1s). -1 ~10s -1 The test results demonstrate the dynamic buffering and energy dissipation capabilities under the following conditions: At least 3 to 5 parallel samples were tested at each height, and the average value and standard deviation of the critical height, critical energy, and energy per unit thickness, as well as the breakage pattern of the cover glass, were recorded. The results of the falling ball impact protection test are as follows: Figure 6 As shown.

[0052] Note: The preparation process of POS(25%) is as follows: referring to the preparation process of PMOS-1 in Example 1, Sq-2CF3 and 2-[2-(2-methoxyethoxy)ethoxy]ethyl acrylate are copolymerized at a molar ratio of 1:3.5 to obtain POS(25%).

[0053] The preparation process of PMOU(8%) is as follows: referring to the preparation process of PMOS-1 in Example 1, the urea-containing monomer (replacing Sq-2CF3) and MEO2MA are copolymerized at a molar ratio of 1:10 to obtain PMOU(8%).

[0054] The preparation process of PMOS(2%) is as follows: referring to the preparation process of PMOS-1 in Example 1, Sq-2CF3 and MEO2MA are copolymerized at a molar ratio of 1:20 to obtain PMOS(2%).

[0055] Depend on Figure 6 It can be seen that PMOS-1 exhibits comparable protective performance to the currently leading protective material OBC 9530 (olefin block copolymer, density 0.887 g / cm³) in the falling ball protection performance test. 3 The material exhibits impact protection capabilities comparable to, but significantly superior to, other reference materials (hard-segment crystalline-soft-segment elastomer dual-phase structure). This result is of great significance. a) The unique mechanism by which hydrogen bonds in squamamide temporarily stabilize the molecular chain orientation is the core source of PMOS-1's high dissipation capability: During impact loading, the molecular chain rapidly aligns along the stress direction. The hydrogen bond network of squamamide temporarily locks this orientation state through dynamic dissociation and recombination, allowing the deformation work to be uniformly distributed and stored in the form of orientation elastic energy, rather than being concentrated in a local area and causing damage. When the external force is removed, the frozen hydrogen bonds gradually unlock, the oriented chain segments relax and rebound, completing the reversible release and dissipation of energy. This orientation-hydrogen bond coupling dissipation mechanism endows the material with adaptive energy management capabilities: at low speeds, the dynamic recombination of hydrogen bonds allows chain segment slippage, which is efficiently dissipated through internal friction; at high speeds, the hydrogen bonds are instantaneously frozen, fixing the orientation state into a transient high-modulus network to resist deformation. In contrast, the crystalline hard segments of OBC 9530 are static physical cross-linking points, and the chain segments cannot undergo this kind of orientation locking. Energy can only be passively dissipated through the high-elastic deformation of the soft segments, lacking the ability to actively regulate orientation-energy coupling. b) Mechanism comparison reveals the essential differences in chemical design: The protective performance of OBC 9530 stems from its microphase separation structure: the hard polyethylene blocks form crystalline microregions (melting point approximately 119°C) as physical crosslinking points, while the amorphous soft segments provide elastic recovery. Impact energy is primarily dissipated through the high elastic deformation of the soft segments and the stress transfer of the hard segment microcrystals. This is a static topology-dependent toughening strategy, whose energy dissipation capability is limited by the thermodynamic balance of phase region size and crystallinity. Furthermore, the rigidity of the crystalline hard segments prevents them from dynamically strengthening in response to high-speed impacts. PMOS-1, as a homogeneous linear copolymer, lacks this type of microphase separation structure. Its superior protective performance originates entirely from the aforementioned rate-dependent hydrogen bond lock-and-unlock mechanism, which can adaptively adjust the energy dissipation mode according to the impact intensity. In contrast, the crystalline hard segments of OBC 9530 do not possess this dynamic reversible response capability. c) The specificity of hydrogen bonding units determines the upper limit of performance: The superiority of PMOS-1 clearly demonstrates the irreplaceable nature of the squamamide hydrogen bonding unit. Compared with conventional hydrogen bonding units such as urea groups, squamamides have stronger dimer binding energy, a more regular linear quadruple hydrogen bond array, and faster dissociation-recombination kinetics, enabling them to rapidly form a transient physical cross-linked network under high-speed impact conditions, effectively dispersing stress and inhibiting crack propagation. More importantly, only squamamide hydrogen bonds can achieve temporary stability of molecular chain orientation—urea group hydrogen bonds are not strong enough to lock the orientation state within the high-speed deformation timescale; the rigidity of the OBC crystal hard segment restricts the segment orientation. This orientation locking ability is the root cause of PMOS-1's energy dissipation efficiency far exceeding that of other hydrogen bonding systems; d) Simplicity and Efficiency of Molecular Design: As a linear copolymer system, PMOS-1 achieves its superior protective performance without the complex segmental block copolymerization, precise phase separation control, or multiphase structure design required by OBC 9530. It achieves this through the self- and heterogeneous hydrogen bonds provided by the side-chain squamamide. This provides a simple and efficient molecular design strategy for developing next-generation lightweight, processable, and high-performance protective materials—directly embedding dynamic hydrogen bonds as programmable mechanical response units into the single-phase polymer backbone, rather than relying on thermodynamically driven microphase separation. This design not only simplifies synthesis and processing but also endows the material with intrinsically adaptive mechanical response characteristics, highlighting the application potential and irreplaceable nature of squamamide hydrogen bonds in the field of impact-resistant materials.

[0056] 4) The dynamic compressive mechanical properties of the dynamically hydrogen-bonded crosslinked polymer (PMOS-1) of Example 1 under high strain rates were tested using a split Hopkinson pressure bar (SHPB). (Test principle: High-pressure gas drives an impact bar to strike an incident bar, generating a very short-duration compressive stress pulse. This pulse propagates along the incident bar to the sample clamped between the incident bar and the transmission bar. Due to the difference in wave impedance between the sample and the bar materials, part of the pulse is reflected back to the incident bar, and the other part is transmitted into the transmission bar. The waveforms of the incident wave, reflected wave, and transmitted wave are recorded by strain gauges attached to the bar and a high-speed data acquisition system. Based on the one-dimensional elastic stress wave theory, the stress, strain, and strain rate of the sample during dynamic loading are directly calculated using the recorded waveform data, thereby obtaining the dynamic stress-strain curve, and then extracting parameters such as dynamic compressive strength, dynamic elastic modulus, peak strain, and strain rate sensitivity.) The sample is in the shape of a disc (size: diameter 10 mm, thickness 5 mm). Five parallel samples were tested at each strain rate. The test results are as follows: Figure 7 (a is the stress-strain relationship curve of PMOS-1 at different strain rates, b is a comparison of the energy absorption capacity of PMOS-1 and OBC 9530, and c is a snapshot of the high-speed deformation process of PMOS-1's SHPB at different strain rates.)

[0057] Depend on Figure 7 It can be known that: a) PMOS-1 at 6500s -1 Crack initiation has been observed, but at 7800s -1The peak stress further increases because hydrogen bonds freeze more fully at higher strain rates, molecular chain orientation and strain hardening occur earlier and at a faster rate, and a higher load-bearing capacity is established before the crack fully propagates, causing the stress to continue to rise with the increase of strain rate. This synergistic mechanism of "freezing-orientation-hardening" enables PMOS-1 to still exhibit the comprehensive advantages of high strength and high energy absorption at ultra-high strain rates. (PMOS-1's excellent performance under high-speed impact is due to the efficient coupling of group orientation and molecular chain orientation achieved by planar hydrogen bonds of square amides, which is highly synergistic with the hydrogen bond freezing effect at high strain rates. Its hydrogen bonds exhibit the dual characteristics of overall freezing and local buffering at high strain rates, which both enhances instantaneous strength and avoids brittle fracture.) (b) The energy absorption capacity of OBC 9530 under high-speed impact is significantly lower than that of PMOS-1, revealing the fundamental difference between the two in their high-strain rate response mechanisms. The hard segment of OBC 9530 is a crystalline polyethylene block, and its physical cross-linking depends on the unlocking of the crystalline region. Under high strain rates, the crystalline region cannot respond sufficiently and cannot freeze rapidly like the hydrogen bonds of squaramide to build an instantaneous rigid network. The chain segment movement of the amorphous region of the soft segment also cannot effectively dissipate energy due to the time scale mismatch, resulting in a significant decrease in the overall load-bearing capacity and energy absorption efficiency at high speeds. In addition, OBC 9530 lacks a dynamic non-covalent bond-mediated orientation-locking cooperative mechanism, and cannot form an ordered orientation-hardened structure under high strain rates, resulting in low stress wave transmission efficiency and difficulty in effectively dissipating local stress concentrations. In summary, the freezing effect of the dynamic hydrogen bond network and the synergistic effect of the molecular chain orientation induced by the squaramide are the decisive mechanisms by which the dynamic hydrogen bond crosslinked polymer of this invention achieves high energy absorption under high-speed impact. This mechanism is not possessed by OBC-type purely physically crosslinked elastomers, and it is also the core competitiveness of the dynamic hydrogen bond crosslinked polymer system in high-speed protection applications.

[0058] 5) The dynamic hydrogen-bonded crosslinked polymer (PMOS-1) from Example 1 was fixed on the stage of an inverted integrated fluorescence microscope. A 500 μm deep and 100 μm long scratch was then made on the sample surface using a utility knife at room temperature. The scratched area was located under the microscope, and an initial image was captured, recorded as time t=0. The scratch morphology was then continuously observed at set time intervals (1 min, 5 min, 10 min, 20 min, and 40 min) under the same microscope's bright-field or fluorescence mode. When the scratch completely disappeared from the field of view and could not be identified at different depths of focus, the cumulative time was recorded as the complete healing time of the scratch. The room temperature self-healing test results are shown in the figure below. Figure 8 As shown.

[0059] Depend on Figure 8 It can be seen that the room temperature self-healing of PMOS-1 can be divided into the following stages: Phase 1 (0 min ~ 5 min): Rapid physical closure of the interface At 0s, the scratch exhibits a clear, wide crack morphology, with debris and stress-damaged areas present on both sides. Within 1-5 minutes, the scratch width rapidly decreases, with the edges converging significantly towards the center, and the debris area noticeably diminishes. From a materials design perspective, the high driving force at this stage likely stems from two aspects: firstly, the newly formed crack surface possesses high excess surface energy at the moment of scratch formation, thermodynamically driving spontaneous contraction of the materials on both sides to reduce the system's energy; secondly, the residual dynamic hydrogen-bonded cross-linked network at the scratch edge undergoes entropic elastic rebound after local stress release, converting stored conformational entropy into macroscopic restoring force. This stage is primarily characterized by physical contact, with cross-interfacial molecular connections not yet effectively established. Second stage (5 min–20 min): Cross-interfacial molecular penetration and bonding After 5 minutes, the scratch outline continued to blur, and the fine branches on both sides of the main crack gradually disappeared. By 20 minutes, the scratch outline was no longer discernible. Considering the molecular structure design of this material, this stage likely corresponds to the core self-healing process: the low glass transition temperature of the MEO2MA segments provides sufficient segment mobility at room temperature, allowing molecular chains on both sides of the interface to diffuse across the repair interface; the rigid bidentate coplanar configuration of the squamamide groups facilitates the directional recombination of hydrogen bonds. Once the segments diffused into the interface region reach the hydrogen bond interaction distance with the opposite squamamide groups, they can form effective cross-interface molecular bridges. The formation of hydrogen bonds may further constrain the segment conformation, creating a synergistic traction effect and driving the repair interface from physical contact to molecular-level healing. Phase 3 (20 min–40 min): Network reconstruction and structural homogenization From 20 to 40 minutes later, the morphology of the scratched area and the surrounding matrix became consistent, and the repair interface basically disappeared. It is inferred that this stage was mainly characterized by the overall reconstruction of the hydrogen bond network and the rearrangement of chain entanglements: the cross-interface hydrogen bond connections gradually evolved from a localized dispersed state to a continuous network, and the orderly arrangement of molecular chains in the repair area was gradually restored; at the same time, the cross-interface diffused chain segments further locked the repair structure through topological entanglement, so that the mechanical response at the repair interface approached the level of the original matrix, and finally the macroscopic morphology and microstructure were restored simultaneously.

[0060] 6) The composite impact-resistant protective material (PMOS / Kevlar) from Example 4 was longitudinally cut to form a "trouser-shaped" specimen. The two legs of the specimen were then clamped in the upper and lower fixtures of a universal testing machine, and a tensile test was performed at a constant speed of 5 mm / min. The complete load was recorded. The displacement curve is used to calculate the tear energy. The formula for calculating the tear energy is as follows: T=U / (t·Δc), where U is the load. The area under the displacement curve (total absorbed energy), t is the sample thickness, Δc is the crack propagation length, and the resulting tear resistance test results are as follows: Figure 9 (a is the macroscopic morphology of the sample after tearing, b is the tearing process diagram, c is the force-displacement relationship curve: the peak tearing force and fracture displacement of the PMOS / Kevlar composite material are significantly improved compared with pure PMOS and pure Kevlar, d is the tearing energy calculated based on the integral of the force-displacement curve: the PMOS / Kevlar composite material achieves an order-of-magnitude improvement due to the hydrogen bond-fiber synergistic dissipation mechanism) as shown.

[0061] Depend on Figure 9 It is known that the tearing energy of PMOS / Kevlar can exceed 600 kJ / m. 2 (In the PMOS / Kevlar system, based on the different locations and modes of action of hydrogen bonds, they can be divided into three categories: interfacial adhesive hydrogen bonds, matrix-internal non-adhesive hydrogen bonds, and interfacial bridging hydrogen bonds. Interfacial bridging hydrogen bonds achieve crack path deflection and efficient energy dissipation through anisotropic constraints, which is the core of the tear resistance mechanism. Interfacial adhesive hydrogen bonds ensure the continuity and stability of interfacial stress transmission through dynamic dissociation-recombination, delaying interfacial debonding. Matrix-internal hydrogen bonds provide a continuous plastic dissipation channel, improving the overall toughness of the material. The three are highly synergistic and jointly determine the excellent tear resistance of the composite material.) It is far superior to that of pure PMOS matrix and pure aramid fiber, which shows the core position of interfacial bridging hydrogen bonds in interfacial reinforcement and tear resistance mechanism.

[0062] The above embodiments are preferred embodiments of the present invention, but the embodiments of the present invention are not limited to the above embodiments. Any changes, modifications, substitutions, combinations, or simplifications made without departing from the spirit and principle of the present invention shall be considered equivalent substitutions and shall be included within the protection scope of the present invention.

Claims

1. A dynamically hydrogen-bonded crosslinked polymer, characterized in that, The structure is as follows: In the formula, x takes an integer from 30 to 200, and y takes an integer from 100 to 800.

2. The dynamically hydrogen-bonded crosslinked polymer according to claim 1, characterized in that: The number-average molecular weight of the dynamically hydrogen-bonded crosslinked polymer is 60,000 g / mol to 200,000 g / mol.

3. A method for preparing the dynamically hydrogen-bonded crosslinked polymer as described in claim 1 or 2, characterized in that, Includes the following steps: A reversible addition-fragmentation chain transfer radical polymerization was carried out in an organic solvent using a bis(trifluoromethyl)-containing squamamide, methoxydiethylene glycol methacrylate, a chain transfer agent, and an initiator. The bis(trifluoromethyl)-containing squamamide was... The product is then separated and purified to obtain a dynamically hydrogen-bonded cross-linked polymer.

4. The preparation method according to claim 3, characterized in that: The molar ratio of the bis(trifluoromethyl) squaramide, methoxydiethylene glycol methacrylate, chain transfer agent, and initiator is 1–2:1–20:0.02–0.7:0.002–0.

35.

5. The preparation method according to claim 3 or 4, characterized in that: The chain transfer agent is at least one of 4-cyano-4-[(dodecylthioalkylthiocarbonyl)thioalkyl]valerate, 4-cyano-4-(dodecyltrithiocarbonyl)valerate-2-(norbornene-1,2-dicarboximide)ethyl ester, 2-cyanopropyl-2-ylphenyl dithiocarbamate, and 4-cyano-4-(phenylthiocarbamoylthio)valerate; and / or, the initiator is at least one of azobisisobutyronitrile, benzoyl peroxide, and azobisisoheptanenitrile.

6. The preparation method according to claim 3 or 4, characterized in that: The reversible addition-fragmentation chain transfer radical polymerization was carried out at a temperature of 70℃ to 80℃ for a reaction time of 5h to 10h.

7. An impact-resistant protective material, characterized in that, It includes the dynamic hydrogen-bonded crosslinked polymer as described in claim 1 or 2.

8. A composite impact-resistant protective material, characterized in that, The composition includes an aramid fiber cloth and an impact-resistant film adhered to both sides of the aramid fiber cloth; the impact-resistant film is made of the dynamic hydrogen-bonded crosslinked polymer as described in claim 1 or 2.

9. The composite impact-resistant protective material according to claim 8, characterized in that: The thickness of the impact-resistant film is 1mm to 4mm.

10. The application of a dynamically hydrogen-bonded crosslinked polymer as described in claim 1 or 2 in the fields of sports protection, automotive industry, or flexible bulletproofing.

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