Cycloreversion-enhanced toughness and degradability in mechanophore-embedded end-linked polymer networks

Domain: chemistry
Published: 2026-01-03
Editorial outcome: major_revision

Abstract
Enhancing the toughness while achieving triggerable degradation in single-network polymer systems without modifying their inherent chemical composition or network architecture remains a significant challenge. Here we demonstrate a smart end-linked polymer network that “self-strengthen” during use yet “self-destruct” upon certain stimuli. Embedding nonscissile cyclobutane-fused tetrahydrofuran mechanophores within the middle of end-linked polymer networks significantly enhances both toughness and degradability. Under mechanical stress, the force-coupled cycloreversion of these mechanophores releases concealed chain segments, enabling single-network materials to exhibit threefold toughness and tenfold tear energies compared to conventional counterparts. Additionally, ball-milling griding of the bulk material unveils acid-sensitive enol ether units, leading to a markedly improved degradation profile under acidic conditions. This dual effect—originating from the force-coupled cycloreversion of cyclobutane-fused tetrahydrofuran mechanophores—provides an ideal combination of superior mechanical performance and on-demand degradability.

Introduction
As crosslinked polymer networks become increasingly prevalent in biomedical devices and flexible electronics1,2, there is a growing demand for materials that offer both superior mechanical strength and triggerable, on-demand degradation. Polymer mechanochemistry uses mechanical force to trigger chemical reactions in polymers, offering great promise for improving material performance3. Recently, increasing attention has been focused on using mechanochemistry to modulate the structure and properties of polymer networks4,5. By integrating stress-responsive mechanophores into polymer chains, researchers can precisely control force transmission and target crack initiation. One approach is to harness the intermediates generated by force-induced bond scission to initiate chemical reactions within the material, thereby endowing the polymer networks with self-strengthening properties6–12. However, this approach generally alters the polymer network architecture and introduces complexities in large-scale manufacturing.

In addition to force-induced chemical transformations, toughness can be enhanced by dissipating energy through premature mechanophore activation within a polymer network. In this context, incorporating scissile mechanophores as cross-linkers in randomly cross-linked networks produces markedly tougher materials13,14. During the stretching process, mechanophore cross-links act as sacrificial bonds, enabling the primary polymer strands chains to stretch rather than fracture. Moreover, embedding nonscissile cyclobutane mechanophores into the primary network of double-network hydrogels has also been shown to double their tearing energies by unveiling hidden hydrocarbon chain segments15. However, owing to the inherent brittleness of the primary network, this reinforcing effect only becomes apparent upon introduction of a secondary network.

Embedding scissile mechanophores such as cyclobutane derivatives into the middle of end-linked polymer networks generally yields similar swelling ratios and moduli but significantly lower fracture strains and tearing energies compared to gels without mechanophores (Fig. 1a)16–18. In contrast to randomly cross-linked systems, premature mechanophore activation usually creates defects to weaken the materials due to the highly cross-linked architecture and brittle property of end-linked networks. Until now, conventional mechanophores have failed to enhance the toughness of end-linked polymer networks, underscoring the need for innovative mechanophore designs19. Inspired by the releasing hidden length effect in double-network hydrogels15, we therefore envision whether a bicyclic cyclobutane mechanophore might endow a highly cross-linked, single-network polymer with enhanced toughness20.

Here we demonstrate that embedding nonscissile cyclobutane-fused tetrahydrofuran mechanophores into the middle of end-linked polymer networks significantly enhances both the toughness and degradability (Fig. 1c). The dual enhancing effect is attributed to the force-coupled cycloreversion of cyclobutane-fused tetrahydrofuran mechanophores, which released both hidden lengths and acid-sensitive enol ether units to the polymer chains21. Compared to the conventional analogues, embedding nonscissile cyclobutane-fused tetrahydrofuran mechanophores lead to networks that stretch 100% further and exhibit tear energies that are nine times larger. We have observed that releasing hidden lengths of nonsciccile mechanophores in highly cross-linked single-network systems enhances the strength of the materials. Notably, integration of bicyclic cyclobutane mechanophores have been reported to maintain material stability until an external force triggers degradation on demand21–28. In addition to the toughening effect, the incorporation of nonscissile cyclobutane-fused tetrahydrofuran mechanophores also improved the degradability when the materials were treated with ball-milling grinding under bulk conditions21. The unveiled enol ether units enabled the polymer network materials with significantly accelerated degradation rate under acidic conditions. Therefore, the described end-linked polymer network possesses an ideal combination of desirable mechanical properties and on-demand degradability.

Results
Photopolymerization is a fast and versatile method used in many applications—from fabricating high-resolution 3D printed objects to producing microfluidic devices29. Initially, we employed the well-established thiol-ene photopolymerization30 to access end-linked polymer networks with nonscissile cyclobutane-fused tetrahydrofuran mechanophores. A diallyl ether monomer bearing a cyclobutane-fused tetrahydrofuran mechanophore (C3) was designed and easily synthesized from dimethyl acetylenedicarboxylate, 2,3-dihydrotetrahydrofuran and allyl bromide in large scale. Meanwhile, C1 bearing a scissile mechanophore and C2 without mechanophore were prepared to understand the influence of cyclobutane-fused tetrahydrofuran mechanophore on the following formed end-linked polymer networks. Since the preparation of polymer networks through step-growth polymerization requires the participation of monomers with a functional degree greater than two, a commercially available tetrafunctional monomer, pentaerythritol tetra(3-mercaptopropionate) (PETMP), was selected as the crosslinker. Using diphenyl(2,4,6-trimethylbenzoyl)phosphine oxide (TPO) as an initiator, photoinduced thiol-ene step-growth polymerization was carried out with the synthesized diene monomers (Fig. 2a). Initially, solutions of [C1]0:[PETMP] = 2:1 for PN1, [C2]0:[PETMP] = 2:1 for PN2, [C2]0:[C3]0:[PETMP] = 1:1:1 for PN3 and [C3]0:[PETMP] = 2:1 for PN4 were prepared, followed by the addition of 0.5 wt% TPO. After thorough dissolution, the mixtures were added to Teflon molds and subjected to photocuring under the irradiation of a 405 nm purple LED. Cross-linking was successfully achieved within minutes. To ensure high monomer conversion, the irradiation time was extended to one hour, resulting in complete and highly crosslinked films PN1, PN2, PN3, and PN4 (Supplementary Fig. 11). FTIR spectroscopy confirmed the completion of the photopolymerization process, as the characteristic absorption peaks at 1646 cm−1 and 2568 cm−1 corresponding to C=C and S–H bonds completely disappeared after the photocuring process (Supplementary Figs. 12–15). The structural similarity of C1, C2, and C3 led to comparable polymerization reactivity. Therefore, when resin solutions of the same ratio were prepared, PN1-4 with identical polymer network structures were expected to be obtained. The four polymer networks differed only in the mechanochemically responsive structural units at the bridging junctions.

To determine the influence of cyclobutane-fused tetrahydrofuran mechanophores on the strength of polymer networks, uniaxial tensile testing experiments were performed on PN1-4. As expected, the representative stress-strain curves of dumbbell-shaped samples of PN1-4 exhibited marked differences (Fig. 2b). The maximum stress value at fracture went in the order PN1 < PN2 < PN3 < PN4, indicating that PN3 and PN4 had greater strength. The introduction of the cyclobutane-fused tetrahydrofuran mechanophores enhanced the tensile strength of the material, meaning that PN4 could withstand greater external stress before reaching the breaking point. Additionally, PN4 exhibited a higher elongation at break than PN1-3 (33% vs. 24% vs. 19% vs. 15%) and demonstrated better ductility. Notably, the stress-strain curves revealed that PN1-4 had similar Young’s moduli of 4.6 ± 0.1, 5.3 ± 0.2, 5.1 ± 0.2, and 5.0 ± 0.2 MPa, respectively (Fig. 2c). This indicated that, apart from the differences in the embedded mechanical mechanophores structures, PN1-4 had similar crosslinking densities and network structures, showing no essential difference under stress conditions insufficient to activate the mechanophores. In comparison to PN2 (88.2 ± 6.4 kJ/m3), PN1 exhibited reduced toughness (46.1 ± 0.9 kJ/m3), whereas PN4 demonstrated a significant increase in toughness (258.5 ± 24.1 kJ/m3) (Fig. 2d). Polymer networks with different mechanophore contents exhibit varying levels of toughness (PN4 > PN3 > PN2), suggesting a positive correlation between toughness and the number of cyclobutane-fused tetrahydrofuran mechanophores. The difference in toughness between PN1-4 could be attributed to distinct mechanophores at the bridging junctions. Given that the cyclobutane unit constitutes a scissile mechanophore, premature mechanophore activation usually creates defects to weaken the polymer network. However, the bicyclic structure of the cyclobutane-fused tetrahydrofuran mechanophore could dissipate part of the energy through ring cleavage and reduction under certain stress, followed by the release of the length stored in the cyclobutane structure15. This process promoted chain extension within the single-network materials, thereby achieving higher fracture stress and elongation at break.

To further quantitatively determine the strength difference between PN1-4, the tear energies of PN1-4 were characterized using the Rivlin-Thomas method for notched films in pure shear geometry31. Unfortunately, the inherent brittleness of PN1 precludes reliable tear resistance characterization. The representative stress-strain curves of unnotched and notched samples for PN2-4 under uniaxial tensile conditions are shown in Fig. 2e, f. The stress-strain behavior of the unnotched samples for PN2-4 was identical before reaching the ultimate fracture point, confirming that PN2-4 had similar polymer network structures. However, the critical stress and strain required for fracture were both lower for PN2 than for PN3 and PN4, indicating that PN3 and PN4 exhibit superior toughness compared to PN2. Although these polymer networks had similar network structures, the introduction of the cyclobutane-fused tetrahydrofuran mechanophores in PN4 led to significant changes in the tear resistance of the material. The representative stress-strain curves under uniaxial tension of the notched samples for PN2-4 are shown in Fig. 2f, and the stress-strain behavior of notched PN2-4 was highly reproducible (Supplementary Fig. 22). The critical tensile strain for crack growth in PN4 was 7.9 ± 0.3%, which was greater than the observed critical tensile strain of 2.4 ± 0.6% for PN2 and 5.7 ± 0.1% for PN3 (Supplementary Table 2). The tear energies of PN2-4 were determined from the critical strain required for crack growth in the notched samples, yielding values of 3.8 ± 1.8, 18.3 ± 0.6, and 38.2 ± 2.8 J/m2, respectively (Fig. 2g). The tear energy of PN4 was tenfold compared to PN2, demonstrating that the introduction of the cyclobutane-fused tetrahydrofuran mechanophores into the end-linked polymer network significantly improved the tear energy, resulting in a tougher material.

Subsequently, the shear modulus and viscoelastic properties of polymer networks were characterized using small-amplitude oscillatory rheology tests. Frequency sweep rheology from 0.1 to 100 rad/s at 25 °C and 1% strain showed that PN1-4 exhibited similar storage moduli (G′) of 380 ± 9, 353 ± 8, 380 ± 8, and 370 ± 23 kPa, respectively, and G′ was markedly higher than the loss modulus (G″), as evidenced by a loss factor (tan δ) < 1 (Fig. 3a, b), indicating that both polymer networks exhibited good and stable elastic behavior. Additionally, temperature sweep rheology from 25 °C to 140 °C at 1% strain revealed that the G′ and G″ values of PN1-4 were relatively stable, while G′ values were consistently higher than G″ values, indicating that the polymer networks retain excellent elasticity and thermomechanical stability (Fig. 3c). As expected, the following consistent features were observed across PN1-4: (i) similar shear moduli, with nearly identical trends in G′ as a function of frequency or temperature; (ii) similar swelling ratios (Fig. 3d); (iii) similar gel contents (Fig. 3e). These results further confirmed that apart from the differences in the embedded mechanophores structures, PN1-4 possessed similar network structures. Moreover, the glass transition temperature (Tg) of PN1-4 exhibited a range of values (Fig. 3f). The scissile cyclobutane and nonscissile cyclobutane-fused tetrahydrofuran, which have a strained ring structure, likely made the polymer chains stiffer and limited their movement, which in turn increased Tg.

Polymer networks are known to have poor solubility in organic solvents, making their degradation particularly challenging32. Until now, it remains unclear how incorporating cyclobutane-fused mechanophores will influence the degradability of highly cross-linked polymer networks33. The introduction of the cyclobutane-fused tetrahydrofuran mechanophores not only enhanced the toughness of the polymer network through chain extension but also introduced acid-sensitive enol ether units into the main chain after the force-induced cycloreversion, effectively promoting its acid degradation (Fig. 4a). To verify the feasibility of bulk mechanical activation and acid degradation of the end-linked polymer networks, ball-milling activation experiments were conducted on the films PN2 (Fig. 4b, top) and PN4 (Fig. 4b, bottom). After 8 h of ball-milling at 50 Hz under room temperature, the morphology of the films changed significantly, resulting in light gray solids (Fig. 4c). Since both PN2′ and PN4′ obtained after ball-milling grinding were insoluble in THF (Supplementary Fig. 29), the size exclusion chromatography (SEC) elution curves showed low signal intensity under the premise of ensuring the same concentration of tested samples (Supplementary Fig. 30, red and blue lines). When the ball-milled materials were added to a TFA/H2O (10/1 v/v) solution at room temperature, PN4′ produced a clear and transparent solution within 10 min (Fig. 4d, bottom), while PN2′ remained a turbid solution (Fig. 4d, top). Notably, the SEC elution curve of hydrolyzed PN4′ displays a pronounced oligomer peak in the low molecular weight region, indicative of nearly complete hydrolysis (Supplementary Fig. 30, green line). In contrast, the SEC elution curve of hydrolyzed PN2′ shows a much smaller oligomer peak, likely due to partial hydrolysis of the ester groups in the PETMP crosslinkers under acidic conditions (Supplementary Fig. 30, orange line).

To more clearly illustrate the difference in degradability, we quantified the insoluble solid residues and soluble small-molecule products obtained after ball-milling activation and subsequent acid hydrolysis of PN2′ and PN4′. PN2′ showed only a slight decrease in insoluble content, whereas PN4′ exhibited a dramatic reduction (Fig. 4e). After 8 h of ball-milling, hydrolysis of PN4′ produced completely soluble small molecules or oligomers, while approximately 80% of PN2′ remained insoluble. Meanwhile, SEC analysis of soluble products obtained from acid hydrolysis of PN2′ and PN4′ (Fig. 4f) revealed that the SEC profile of hydrolyzed PN2′ changed minimally with activation time, whereas that of hydrolyzed PN4′ developed a strong oligomer peak in the low molecular weight region. Different from that of hydrolyzed PN2′, the ¹H NMR spectrum of the acid degradation products of PN4′ showed a characteristic peak at 9.8 ppm, indicating the formation of aldehydes after hydrolysis (Fig. 4g). Furthermore, high resolution electrospray ionization mass spectrometry (HR-ESI-MS) analysis of the small-molecule products from PN4′ identified the species originating from acid-hydrolyzed enol ethers (Fig. 4h). These findings confirmed that the cycloreversion of the cyclobutane-fused tetrahydrofuran mechanophores through ball-milling activation, thereby providing direct evidence for mechanophore activation. Control experiments, where PN2 and PN4 were not subjected to ball-milling activation, showed no oligomer formation via SEC analysis (Supplementary Fig. 34). These results demonstrate that the cyclobutane-fused tetrahydrofuran mechanophores undergo cycloreversion to form enol ethers after ball-milling, significantly enhancing the degradability of the end-linked polymer network.

In summary, we have demonstrated the dual role of nonscissile cyclobutane-fused tetrahydrofuran mechanophores in enhancing toughness while enabling triggerable degradation in single-network materials. This dual effect stems from the force-coupled cycloreversion of the mechanophores, which releases hidden lengths and acid-sensitive enol ether units within the polymer chains. Compared to conventional analogues, end-lined polymer networks incorporating these mechanophores exhibit threefold toughness and tenfold tear energies. Furthermore, the cyclobutane-fused tetrahydrofuran mechanophores also retain the elasticity and thermomechanical properties of the materials. Additionally, the incorporation of nonscissile cyclobutane-fused tetrahydrofuran mechanophores improves the degradability of the materials when subjected to ball-milling grinding under bulk conditions. Upon ball-milling, the cycloreversion of the cyclobutane-fused tetrahydrofuran mechanophores generates acid-sensitive enol ether units, which significantly accelerate the hydrolytic degradation of the polymer network under acidic conditions. Therefore, the described end-linked polymer network with nonscissile cyclobutane-fused tetrahydrofuran mechanophores offers an ideal combination of desirable mechanical properties and on-demand degradability.

Methods
Materials
All of the commercially available reagents were used as received, unless otherwise noted. 2,3-dihydrofuran, N,N-diisopropylethylamine, 3-oxabicyclo[3.2.0]heptane-2,4-dione, allyl bromide and trifluoroacetic acid, sodium hydroxide, and sodium sulfate (Na2SO4) were purchased from Shanghai Titan Scientific Co., Ltd. Dimethyl acetylenedicarboxylate, palladium on activated carbon (10%, wetted with ca. 55% Water), tetrabutylammonium iodide, and 1,4-butanediol were purchased from Bidepharm Co., Ltd. Lithium aluminium hydride (LiAlH4, 2.5 mol/L in THF) was purchased from Energy Chemical Co., Ltd. PETMP was purchased from Shanghai Aladdin Biochemical Technology Co., Ltd. TPO was purchased from TCI Shanghai Chemical Industrial Development Co., Ltd. Hydrogen (H2) was purcased from Shanghai Pujiang Special Gas Co., Ltd. Ethyl acetate (EtOAc) and petroleum ether (PE) were purchased from Sinopharm Chemical Reagent Co., Ltd. Toluene (Sinopham Chemical Reagent Co., Ltd.) and tetrahydrofuran (THF, J&K Scientific Co., Ltd.) were dried over Vigor VG-P7 solvent purification system under an argon atmosphere.

Nuclear magnetic resonance
Nuclear magnetic resonance spectra were recorded on Bruker Ascend 500 (1H: 500 MHz, 13C: 126 MHz) NMR spectrometers at ambient temperature unless otherwise noted. Chemical shift values for protons are referenced to the residual proton resonance of chloroform-d (CDCl3-d: 7.26 ppm). Chemical shift values for carbons are referenced to the carbon resonances of CDCl3-d (δ: 77.0 ppm). NMR assignments were assisted by 1H-13C HSQC and 1H-13C HMBC data.

Size exclusion chromatography
SEC was performed at 35 °C with THF as mobile phase at a flow rate of 1.0 mL/min, using an Agilent 1260 Infinity LC system. The system was equipped with a PL1110-1520 guard column (PLgel 5 μm Guard 50 × 7.5 mm) and two PL1110-6500 columns (PLgel 5 μm MIXED-C 300 × 7.5 mm within molecular weight of 200–2,000,000 × g/mol) connected in series and fitted with a 1260 infinity II Refractive Index detector. Determination of molecular weights and polydispersities were calibrated against polystyrene standards.

Differential scanning calorimetry
Differential scanning calorimetry measurements of polymers were performed on a DSC 250 analyzer at a heating and cooling rate of 10 °C/min. Glass transition temperature (Tg) values were determined from the second heating scan.

Photocuring procedures
Mixtures were prepared in brown vials according to the following molar ratios: [C1]0:[PETMP] = 2:1 for PN1, [C2]0:[PETMP] = 2:1 for PN2, [C2]0:[C3]0:[PETMP] = 1:1:1 for PN3 and [C3]0:[PETMP] = 2:1 for PN4, followed by the addition of 0.5 wt% TPO to each mixture. The mixture was heated on a hot plate to accelerate dissolution and then vortexed mixing for a few seconds. The film samples (a length of 60 mm, a width of 60 mm, and a thickness of 1 mm) were prepared by adding the uniformly mixed solution into a Teflon mold and irradiated with a purple LED light (λmax = 405 nm, 35 W) for one hour at 25 °C. Upon completion of photocuring procedures, the film samples were stored under light-shielded conditions in preparation for subsequent FTIR-ATR, swelling ratio, gel content, thermal, and mechanical characterizations.

Fourier-transform infrared attenuated total reflectance
Fourier-transform infrared attenuated total reflectance measurements were performed on a Thermo Fisher Nicolet iS50.

Rheology
Rheological frequency sweep and temperature sweep experiments were performed on a HAAKE MARS 60 rheometer. A parallel-plate geometry with a radius of 10 mm was used and coupled with a bottom plate, maintaining a typical gap between 1.00 and 1.15 mm. Frequency sweep experiments were conducted from 0.1 to 100 rad/s with a constant oscillatory strain of 1% at 25 °C. Temperature sweep experiments were conducted at a heating rate of 5 °C/min with a constant oscillatory strain of 1%.

Tensile measurements
Tensile measurements were conducted on an Instron 3343 machine testing instrument at ambient conditions (25 °C). Upon completion of photocuring procedures, the film samples were carefully demolded and subsequently sectioned into dumbbell-shaped samples (a length of 35 mm, a width of 6 mm, and a thickness of 1 mm) using a custom-designed die-cutting apparatus. Tensile measurements were conducted on an Instron 3343 machine testing instrument at ambient conditions, stretched at 2.4 mm/min until fracture. Young’s modulus was calculated as the slope of the stress-strain curve within the strain range of 0%–1%. Toughness was calculated by integrating the stress-strain curves up to the point of fracture.

Tearing energy
Tearing energy measurements were conducted using the Rivlin-Thomas method. Upon completion of photocuring procedures, the film samples were carefully demolded and subsequently sectioned into several rectangular samples (a length of 20 mm, a width of 16 mm, and a thickness of 1 mm) using a razor blade. For notched rectangular samples, a 5 mm notch was introduced in the center of one side of the rectangle using a razor blade, perpendicular to the edge. Tearing energy measurements were conducted on an Instron 3343 machine testing instrument at ambient conditions, using the pure shear geometry with both notched and unnotched rectangular samples. The notched and unnotched rectangular samples were put into the tensile grips with a gauge length of 3 mm and then stretched at 0.66 mm/min until fracture. Tearing energy (Г) was calculated by integrating unnotched stress-strain curves up to the critical strain (εp) at which the crack of the notched samples began to propagate to determine strain energy density W(εp) and then multiplying by the sample gauge length (h0).
Γ = W(εp) × h0

Swelling ratios and gel contents
The disc samples (×3) of PN1-4 (10 mm diameter, 1 mm thick) were weighted and submerged in THF for 24 h, and then the mass of the swollen samples was measured. Subsequently, the samples were dried in a vacuum oven until a constant weight was achieved and the mass of the dried samples was recorded. The swelling ratios were calculated by (mass of the swollen sample)/(mass of the dry sample). The gel contents were calculated by (mass of the dry sample)/(initial mass) × 100%.

Ball-milling activation and degradation experiments
Experimental procedure for the kinetic study of the ball-milling activation process: Ball-milling activation experiments were conducted using a 25 mL stainless steel grinding jar with stainless steel grinding balls on Tissuelyser-II (Shanghai Jingxin®) with a frequency of 50 Hz. 500 mg polymer networks (PN2 or PN4) were initially cut into small pieces and then placed into a 25 mL stainless steel grinding jar containing a single stainless steel grinding ball with a diameter of 20 mm. The stainless steel grinding jar with polymers was subjected to grinding for 2 h at a frequency of 50 Hz. Subsequently, the 20 mm grinding ball was replaced with one 14 mm stainless steel grinding ball and 50 stainless steel grinding balls, each with a diameter of 4 mm. The grinding process was then continued at the same frequency of 50 Hz for an additional 6 h. Among them, each 30-min ball-milling period is immediately followed by a 30-min pause, repeating in this cyclic manner. Experimental procedure for the kinetic study of the degradation process: During the ball-milling activation process, polymer network films (20 mg) were taken every two hours and dissolved in 5 mL of a 10/1 (v/v) TFA/H2O mixture. The mixture was stirred at room temperature for 10 min. Subsequently, all volatile components were removed under vacuum. An appropriate amount of THF was added to dissolve the residue, followed by centrifugation to separate the insoluble fraction. The supernatant was concentrated under reduced pressure to obtain the dissolved portion (small-molecule products). After drying in a vacuum oven, both fractions were weighed separately to calculate insoluble solids contents and soluble small-molecule products contents after hydrolysis. The obtained oligomers were then redissolved in 4.5 mL of THF for SEC analysis.

High resolution electrospray ionization mass spectrometry
High resolution electrospray ionization mass spectrometry spectra were recorded on a Waters Synapt HDMS quadrupole/time-of-flight (Q-ToF) mass spectrometer (Waters, Beverly, MA) in positive ion mode.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "Li and Tang take inspiration from previous work demonstrating that a scissile cyclobutane mechanophore toughens polymer\nnetworks. The authors present a clever design incorporating Xia’s non-scissile cyclobutene-fused tetrahydrofuran into end-\nlinked polymer networks, thus combining Craig’s toughening effect with acid degradability post-mechanophore activation.\nThey compare the thermomechanical and tensile properties of these networks to a control system, confirming that while\nthese materials have different Tg’s (both well below room temperature), the rubbery plateau moduli are similar. However, the\nmaterial with the mechanophore is significantly tougher and has much higher tearing energies than the control. Moreover, the\nactivated mechanophore is acid-labile, and the authors demonstrate degradation by first ball-milling followed by acidic\nhydrolysis.\nThis work is notable in demonstrating a significant toughening effect from reactive strand extension in a single network. The\nacid degradability of the activated mechanophores is an added bonus that increases the broad appeal of this manuscript.\nThis manuscript will be suitable for publication in Nature Communications following major revisions.\n1. It may not be necessary to use 100% mechanophore to obtain significant mechanical effects; Beech et all previously\nshowed that the effect of the weak linker dominated up to the percolation threshold (40-50%). Both to compare to the theory\npresented in that work and to make these networks more practically appealing, the authors should mix C1 and C2 to see if\nthe toughening effect persists even when the mechanophore is effectively diluted.\n2. Line 135: “PN2 135 demonstrated superior elasticity and more stable dynamic mechanical properties during shear strain.”\nWhat evidence is provided to support this statement? The plateau moduli of PN1 and PN2 are not significantly different,\nsuggesting similar network topologies.\n3. Figure 3d-e: Why do G’ and G” increase above 140 ºC for both networks? This behavior is atypical of permanent (and\ndynamic) polymer networks and suggests some sort of decomposition that creates more crosslinks. However, no new\nfeatures are seen in the DSC between 140–200 ºC. After the networks are heated above 140 ºC and cooled to room\ntemperature, do they regain their original properties?\n4. A key claim of this work is a mechanically induced degradation of the polymer network. To support this claim, authors take\nGPC of polymers PN1 and PN2 and NMRs. However, the authors don’t identify the likely products of degradation or quantify\nthe extent of degradation. Comparing the intensity of the SEC peaks is not a reliable approach to compare the degradation\nunless all aspects of sample preparation are highly controlled; if normalized, the molecular weights of the hydrolyzed peaks\nfrom PN1 and PN2 may be similar. The authors should (a) quantify the remaining solids after acidic hydrolysis, (b)\ncharacterize the small-molecule products and estimate their yield (based on mass balance).\n5. While the authors conduct ball milling grinding for mechanical activated degradation, they should conduct control thermal\nexperiments. The seminal paper from Xia (ref. 21) states that this mechanophore is thermally unstable beyond 65 C. As\nsuch, they determine whether ball milling under these conditions heats the sample above 65 C.\n6. Line 162: Why are 8 h of ball-milling required? This seems significantly longer than comparable examples in the literature\n(e.g. ACS Macro Lett. 2020, 9, 6, 836–842, ACS Appl. Bio Mater. 2021, 4, 5, 4176–4183). If more quantitative analysis is\nperformed, the extent of degradation over time can be monitored.\n7. Following ball-milling, the authors should perform FTIR analysis to confirm formation of an alkene consistent with\nmechanophore activation.\n8. Figure captions are duplicated for figures 3 and 4.", "reviewer_2": "The work reported by Li and Tang utilized a non-scissile cyclobutene mechanophore to enhance both mechanical\ntoughness and degradability through mechanically triggered cycloreversion, resulting in the formation of an enol ether\nmoiety. The specific cyclobutene mechanophore with mechanically gated degradability has been reported, and the strategy\nof dictating network mechanical properties through incorporation of scissile and non-scissile mechanophores has also been\ndocumented. The use of this non-scissile mechanophore in enhancing the single network toughness is novel. The improved\ntoughness was demonstrated by tensile and fracture tests, and the enhanced degradability was investigated though ball\nmilling experiments following by acid-triggered enol ether degradation. The methods and experiments details are sound.\nThese results are interesting but preliminary, and additional, important experiments are necessary to draw/validate these\nconclusions.\n1. The swelling and sol-gel fraction of PN1 and PN2 need to be characterized to ensure the similarity in these two networks.\n2. The structural difference in C1 and C2 crosslinkers could lead to the distinct mechanical properties in PN1 and PN2\nelastomers, as evidenced by the difference Tg values in Fig 3f. Therefore, comparison for gel systems (as in previous report:\nJ. Am. Chem. Soc. 2021, 143, 10, 3714–3718) are more convincing. The rigid C2 crosslinker lead to the increase in Tg (or\nlower segmental mobility, Fig 4a and 4b) and potentially mechanical toughness.\n3. A control design with scissile cyclobutane analog is necessary to validate the effect of non-scissile mechanophore on\nenhanced mechanical toughness.\n4. Can the author quantify the mechanophore activation post the ball milling? As ball milling is strong enough to break the C-\nC bonds in PN1, the better degradability in PN2 after ball milling suggested good extent of mechanophore activation.\n5. The balling milling time of 8h is long, how does the heat generation during the ball milling affect the mechanophore ring\nopening?\nMinor thing:\n1. The mechanophore abundance in previous double-network hydrogels (Science 2021, 374, 193–196) is actually not low,\ncompared to the single mechanophore per strand length in this report."}, "author_response": "Response to the reviewers’ comments and description of changes made to the manuscript for our\nsubmission on 21-May-2025 (Manuscript numbmer: NCOMMS-25-38649-T). The modified parts of the\nmanuscript with respect to the original versions are highlighted on a yellow background.\nReviewer #1 (Remarks to the Author):\nLi and Tang take inspiration from previous work demonstrating that a scissile cyclobutane mechanophore\ntoughens polymer networks. The authors present a clever design incorporating Xia’s non-scissile\ncyclobutene-fused tetrahydrofuran into end-linked polymer networks, thus combining Craig’s\ntoughening effect with acid degradability post-mechanophore activation. They compare the\nthermomechanical and tensile properties of these networks to a control system, confirming that while\nthese materials have different Tg’s (both well below room temperature), the rubbery plateau moduli are\nsimilar. However, the material with the mechanophore is significantly tougher and has much higher\ntearing energies than the control. Moreover, the activated mechanophore is acid-labile, and the authors\ndemonstrate degradation by first ball-milling followed by acidic hydrolysis.\nThis work is notable in demonstrating a significant toughening effect from reactive strand extension in a\nsingle network. The acid degradability of the activated mechanophores is an added bonus that increases\nthe broad appeal of this manuscript. This manuscript will be suitable for publication in Nature\nCommunications following major revisions.\nOur response: We sincerely thank you for recognizing the significance of our work and for\nproviding constructive comments and suggestions to further improve the manuscript. In response to\nyour concerns, we have conducted additional experiments and integrated the new results into the\nrevised version. Detailed responses to each comment are provided below.\n1. It may not be necessary to use 100% mechanophore to obtain significant mechanical effects; Beech et\nall previously showed that the effect of the weak linker dominated up to the percolation threshold (40-\n50%). Both to compare to the theory presented in that work and to make these networks more practically\nappealing, the authors should mix C1 and C2 to see if the toughening effect persists even when the\nmechanophore is effectively diluted.\nOur response: In response to your suggestion, we prepared a polymer network (PN3) containing\n50% cyclobutane-fused tetrahydrofuran mechanophores by blending C2 (previously referred to as\nC1) and C3 (previously C2). Stress-strain and tearing tests were conducted on PN3 (Fig. 2). The\nresults indicate that both the toughness and tear energy of PN3 are enhanced to some extent, with\nvalues lying between those of PN2 (previously PN1) and PN4 (previously PN2). Specifically, the\ntoughness of PN3 was 30% higher that of PN2, and its tear energy was four times greater. These\nfindings demonstrate that even at 50% incorporation, the cyclobutane-fused tetrahydrofuran\nmechanophores still contribute effectively to toughening, although the enhancement is less\npronounced than in the network with 100% mechanophore content (PN4).\nAs discussed by Beech et al. in their work on how reactivity-guided depercolation processes\ndetermine fracture behavior in end-linked polymer networks (ACS Macro Lett. 2023, 12, 1685-\n1691), chain scission during tearing tends to avoid strong bonds and preferentially proceeds through\nconsecutive cleavage of weak linker, these results in relatively low toughness—a behavior dictated\nby the difference in reactivity between weak linker and strong linker, with the effect of the weak\nlinker dominated up to the percolation threshold (40-50%). In contrast, our findings reveal certain\ndeviations from their conclusions. This discrepancy may stem from variations in fracture\nmechanisms. In our work, the cleavage of the nonscissile cyclobutane-fused tetrahydrofuran\nmechanophore via cycloreversion is followed by chain extension, ultimately resulting in the rupture\nof conventional covalent bonds. This mechanism contributes to the toughening of the polymer\nnetwork. Polymer networks with different mechanophore contents exhibit varying levels of\ntoughness (PN4 > PN3 > PN2), suggesting a positive correlation between toughness and the number\nof cyclobutane-fused tetrahydrofuran mechanophores. Therefore, the difference in fracture behavior\nis likely a key factor accounting for the observed discrepancies. We have added these data (Fig. 2)\nand corresponding discussion (page 3, paragraph 2) in the revised manuscript.\n2. Line 135: “PN2 135 demonstrated superior elasticity and more stable dynamic mechanical properties\nduring shear strain.” What evidence is provided to support this statement? The plateau moduli of PN1\nand PN2 are not significantly different, suggesting similar network topologies.\nOur response: Indeed, there is no significant difference in the modulus between PN2 (previously\nPN1) and PN4 (previously PN2). Therefore, we have removed the inappropriate description.\n3. Figure 3d-e: Why do G’ and G” increase above 140 ºC for both networks? This behavior is atypical of\npermanent (and dynamic) polymer networks and suggests some sort of decomposition that creates more\ncrosslinks. However, no new features are seen in the DSC between 140–200 ºC. After the networks are\nheated above 140 ºC and cooled to room temperature, do they regain their original properties?\nOur response: Thank you for raising this important point. In accordance with your suggestion, we\nperformed two consecutive rheological temperature sweeps on PN2 (previously labeled PN1) and\nPN4 (previously PN2) under identical conditions. The results indicated that both G′ and G″\ngradually increase above 140 °C, followed by an abrupt decrease. We attributed this behavior to\nthermal expansion of the polymer networks during heating, which led to a sharp increase in the\nnormal force and ultimately exceeded the measurable range of the instrument.\nFurthermore, during the second rheological temperature sweeps, the initial shear moduli of PN2\nand PN4 remained nearly unchanged, and the evolution of the modulus with temperature closely\nfollowed the trend observed in the first cycle, as shown in the figures below (a and b). This\nconfirmed that the polymer networks fully recovered their original properties after being heated\nabove 140 °C and cooled back to room temperature. To prevent potential misinterpretation by\nreaders, we have retained only the data below 140 °C in the revised Figure 3c.\n4. A key claim of this work is a mechanically induced degradation of the polymer network. To support\nthis claim, authors take GPC of polymers PN1 and PN2 and NMRs. However, the authors don’t identify\nthe likely products of degradation or quantify the extent of degradation. Comparing the intensity of the\nSEC peaks is not a reliable approach to compare the degradation unless all aspects of sample preparation\nare highly controlled; if normalized, the molecular weights of the hydrolyzed peaks from PN1 and PN2\nmay be similar. The authors should (a) quantify the remaining solids after acidic hydrolysis, (b)\ncharacterize the small-molecule products and estimate their yield (based on mass balance).\nOur response: Thank you for the suggestion. We quantified the contents of insoluble solids and\nsoluble small-molecule products after ball-milling activation and acid hydrolysis for PN2\n(previously labeled PN1) and PN4 (previously PN2) (Fig. 4e). When the ball-milling activation\ntime reached 8 hours, hydrolysis of PN4′ yielding complete soluble small-molecule products,\nwhereas that of PN2′ still remained in approximately 80% of insoluble solids. Meanwhile, SEC\nanalysis of soluble small-molecule products obtained from acid hydrolysis of PN2′ and PN4′ (Fig.\n4f) clearly showed that with increasing ball-milling activation time, the SEC elution curve for\nhydrolyzed PN2′ changed only slightly, while that for hydrolyzed PN4′ exhibited a pronounced\noligomer peak in the low molecular weight region, demonstrating a substantial contrast.\nFurthermore, HR-ESI-MS analysis of the small-molecule products from hydrolyzed PN4′ identified\nthe species originating from acid-hydrolyzed enol ethers (Fig. 4h). This finding confirmed that the\ncycloreversion of the cyclobutane-fused tetrahydrofuran mechanophores through ball-milling\nactivation, thereby providing direct evidence for mechanophore activation. The quantitative data for\ninsoluble solids and soluble products have been included in the revised supplementary information\n(Supplementary Tables 3 and 4). Corresponding discussion of these results has been added to the\nrevised manuscript (Page 7, paragraph 2).\n5. While the authors conduct ball milling grinding for mechanical activated degradation, they should\nconduct control thermal experiments. The seminal paper from Xia (ref. 21) states that this mechanophore\nis thermally unstable beyond 65 C. As such, they determine whether ball milling under these conditions\nheats the sample above 65 C.\nOur response: Xia et al. reported the synthesis of a polymer containing cyclobutane-fused\ntetrahydrofuran units via ring-opening metathesis polymerization, the mechanophore in the main\nchain was thermally unstable due to the presence of cyclobutane adjacent to unsaturated double\nbonds, which necessitated a subsequent hydrogenation step to stabilize the material. In contrast, our\npolymer networks feature cyclobutane mechanophores flanked by methylene groups, which\nsignificantly enhance their thermal stability.\nTo further address your concern, we carefully monitored the thermal conditions during the ball-\nmilling activation process. During the ball-milling activation process, we employed a cyclic\nprocedure in which ball-milling was conducted for 30 minutes followed by an immediate 30-minute\npause. This intermittent approach was intentionally adopted to prevent excessive heat buildup and\navoid non-specific chain scission or other side reactions that could interfere with the mechanophore\nactivation outcome. Additionally, we measured the internal temperatures of the ball-milling jars\nimmediately after ball-milling using an infrared thermometer. The temperatures were consistently\nmaintained below 40 °C (See Supplementary Fig. 25). These results confirm that the experimental\nconditions during ball milling did not exceed 65 °C, and therefore, the observed degradation\noriginates from mechanochemical activation rather than thermal effects.\n6. Line 162: Why are 8 h of ball-milling required? This seems significantly longer than comparable\nexamples in the literature (e.g. ACS Macro Lett. 2020, 9, 6, 836–842, ACS Appl. Bio Mater. 2021, 4, 5,\n4176–4183). If more quantitative analysis is performed, the extent of degradation over time can be\nmonitored.\nOur response: Thank you for your suggestion. To clarify this point, we conducted a systematic\nkinetic study of the ball-milling activation process for the polymer networks (PN2 and PN4)\n(Supplementary Fig. 26). We quantified the amounts of insoluble solids and soluble small-molecule\nproducts after ball-milling activation and acid hydrolysis for PN2′ and PN4′ (Fig. 4e). The results\nindicated that with prolonged ball-milling activation time, the amounts of insoluble solids showed\nslight decrease for PN2′ while that of PN4′ decreased dramatically. When the ball-milling time\nreached 8 hours, hydrolysis of PN4′ yielding complete soluble small-molecule products, whereas\nthat of PN2′ still remained approximately 80% of insoluble solids.\nMeanwhile, SEC analysis of soluble small-molecule products obtained from acid hydrolysis of\nPN2′ and PN4′ (Fig. 4f) clearly showed that the SEC elution curve for hydrolyzed PN2′ changed\nonly slightly, while that for hydrolyzed PN4′ exhibited a pronounced oligomer peak in the low\nmolecular weight region, demonstrating a substantial contrast. This suggested that after 8 hours of\nball-milling activation, PN4′ underwent nearly complete hydrolysis, resulting in a much higher\nsmall-molecule products content. The extended milling time was therefore necessary to ensure\nbetter hydrolytic degradability of PN4. Corresponding discussion of these kinetic results has been\nadded to the revised manuscript (Page 7, paragraph 2).\n7. Following ball-milling, the authors should perform FTIR analysis to confirm formation of an alkene\nconsistent with mechanophore activation.\nOur response: Following ball-milling, we conducted FTIR analysis on PN4 before and after the\nprocess. Regrettably, no distinct differences were observed around 1650 cm⁻¹, as shown in the figure\nbelow. However, HR-ESI-MS analysis of the small-molecule products from PN4 definitively\nidentified the structures as acid-hydrolyzed enol ethers (Fig. 4h). This finding confirmed that the\ncycloreversion of the cyclobutane-fused tetrahydrofuran mechanophores through ball-milling\nactivation generated these acid-sensitive intermediates, thereby providing direct evidence for\nmechanophore activation.\n8. Figure captions are duplicated for figures 3 and 4.\nOur response: We appreciate your feedback and the issue of duplicated captions for Figures 3 and\n4 has been corrected.\nReviewer #2 (Remarks to the Author):\nThe work reported by Li and Tang utilized a non-scissile cyclobutene mechanophore to enhance both\nmechanical toughness and degradability through mechanically triggered cycloreversion, resulting in the\nformation of an enol ether moiety. The specific cyclobutene mechanophore with mechanically gated\ndegradability has been reported, and the strategy of dictating network mechanical properties through\nincorporation of scissile and non-scissile mechanophores has also been documented. The use of this non-\nscissile mechanophore in enhancing the single network toughness is novel. The improved toughness was\ndemonstrated by tensile and fracture tests, and the enhanced degradability was investigated though ball\nmilling experiments following by acid-triggered enol ether degradation. The methods and experiments\ndetails are sound.\nThese results are interesting but preliminary, and additional, important experiments are necessary to\ndraw/validate these conclusions.\nOur response: Thank you very much for your recognition of this research work, as well as your\nconstructive suggestions and recommendations. Regarding the accuracy of the test results, we have\nconducted additional measurements and provided more detailed characterization and analyses to\nstrengthen the evidence supporting our findings. These updates have been incorporated into the\nrevised manuscript and Supporting Information. Please refer to our detailed point-by-point\nresponses below.\n1. The swelling and sol-gel fraction of PN1 and PN2 need to be characterized to ensure the similarity in\nthese two networks.\nOur response: We appreciate your suggestion. Accordingly, we measured the swelling ratios and\ngel fractions of PN1, PN2, PN3, and PN4 under identical conditions. The results demonstrated\nsimilar swelling ratios and gel fractions across all samples, confirming that these polymer networks\npossess similar crosslinking densities and network structures. The corresponding data have been\nadded to the revised manuscript (Figs. 3d and 3e).\n2. The structural difference in C1 and C2 crosslinkers could lead to the distinct mechanical properties in\nPN1 and PN2 elastomers, as evidenced by the difference Tg values in Fig 3f. Therefore, comparison for\ngel systems (as in previous report: J. Am. Chem. Soc. 2021, 143, 10, 3714–3718) are more convincing.\nThe rigid C2 crosslinker lead to the increase in Tg (or lower segmental mobility, Fig 4a and 4b) and\npotentially mechanical toughness.\nOur response: Thank you for this thoughtful comment. We performed comprehensive\ncharacterization of the polymer networks (PN1, PN2, PN3, and PN4), including uniaxial tensile\ntests, tear tests, rheological measurements, swelling ratios, and gel fractions. The results\ndemonstrated that, apart from the differences in the embedded mechanophores structures, these\npolymer networks possessed similar network structures, as evidenced by the following: (i) similar\nYoung’s moduli (Fig. 2c); (ii) similar shear moduli, with nearly identical trends in the storage\nmodulus (G′) as a function of frequency or temperature (Figs. 3a, 3b and 3c); (iii) similar swelling\nratios (Fig. 3d); and (iv) similar gel contents (Fig. 3e). As you pointed out, differences in glass\ntransition temperature (T ) were observed, which we attributed to the incorporation of the embedded\ng\nmechanical mechanophores. To address your concern regarding the potential increase in mechanical\ntoughness with higher T , we synthesized a control polymer network containing scissile cyclobutane\ng\nanalog (PN1). Although PN1 exhibited a higher T than the conventional covalently polymer\ng\nnetwork without mechanophores (PN2), its toughness was actually lower than that of PN2.\nFurthermore, to eliminate the influence of T variations on the mechanical properties—such as the\ng\npossibility that a higher T could lead to a higher modulus at the same test temperature—all\ng\nmechanical tests were conducted at temperatures well above the T of each polymer network. Under\ng\nthese conditions, all four networks showed similar Young’s moduli, confirming that T variations\ng\ndid not influence the measured mechanical performance.\n3. A control design with scissile cyclobutane analog is necessary to validate the effect of non-scissile\nmechanophore on enhanced mechanical toughness.\nOur response: In accordance with your suggestion, we synthesized a polymer network containing\nscissile cyclobutene mechanophores (PN1) as a control system. This material was thoroughly\ncharacterized through a series of experiments, including uniaxial tensile testing (Fig. 2), rheology\nand thermal properties, as well as measurements of swelling ratios and gel fractions (Fig. 3). The\ncorresponding results and discussions have been included in the revised manuscript (Page 3,\nparagraph 2).\n4. Can the author quantify the mechanophore activation post the ball milling? As ball milling is strong\nenough to break the C-C bonds in PN1, the better degradability in PN2 after ball milling suggested good\nextent of mechanophore activation.\nOur response: Due to severe signal overlap in 1H NMR spectrum (Supplementary Fig. 31), it was\nnot possible to directly quantify the exact percentage of mechanophore activation after ball milling\nand subsequent hydrolysis. To better evaluate the activation process, we conducted a systematic\nkinetic analysis of the ball-milling activation for PN2 and PN4 (Supplementary Fig. 26). We\nquantified the insoluble solids and soluble small-molecule products after ball-milling activation\nfollowed by acid hydrolysis (Fig. 4e). With increasing ball-milling activation time, hydrolyzed PN2′\nshowed only a slight decrease in insoluble content, while hydrolyzed PN4′ exhibited a pronounced\nand continuous decrease. When the ball-milling time reached 8 hours, hydrolysis of PN4′ yielding\ncomplete soluble small-molecule products, whereas that of PN2′ still remained approximately 80%\nof insoluble solids. Moreover, 1H NMR and HR-ESI-MS analyses of the small-molecule products\nfrom hydrolyzed PN4′ confirmed the formation of acid-hydrolyzed enol ethers, providing direct\nevidence for the cycloreversion of the cyclobutane-fused tetrahydrofuran mechanophore (Figs. 4g\nand 4h). These results clearly demonstrate that the enhanced degradability of PN4 originates from\nefficient mechanophore activation under ball-milling conditions. The corresponding discussion has\nbeen added to the revised manuscript (Page 7, paragraph 1).\n5. The balling milling time of 8h is long, how does the heat generation during the ball milling affect the\nmechanophore ring opening?\nOur response: During the ball-milling activation process, we employed a cyclic procedure in which\nball-milling was conducted for 30 minutes followed by an immediate 30-minute pause. This\nintermittent approach was intentionally adopted to prevent excessive heat buildup and avoid non-\nspecific chain scission or other side reactions that could interfere with the mechanophore activation\noutcome. Additionally, we measured the internal temperatures of the ball-milling jars immediately\nafter ball-milling using an infrared thermometer. The temperatures were consistently maintained\nbelow 40 °C (See Supplementary Fig. 25).\nMinor thing:\n1. The mechanophore abundance in previous double-network hydrogels (Science 2021, 374, 193–196)\nis actually not low, compared to the single mechanophore per strand length in this report.\nOur response: We appreciate the comment and have modified the inappropriate description\naccordingly (See Page 1, paragraph 2)."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have thoroughly addressed the reviewers' concerns, adding new data and insights that strengthen the\nmanuscript and its conclusions. Publication is recommended, though the following relevant manuscripts have come out\nsince the original submission and should be cited since they also feature mechanophore-enhanced toughness in single\nnetworks: 10.1021/acscentsci.5c00932 and 10.1038/s41557-025-01984-9", "reviewer_2": "The authors performed additional control experiments, including a scissile cyclobutane mechanophore (C1) and blends of\nthe C2 and C3 mechanophores, to clarify how activation of a non‑scissile mechanophore contributes to the material\nproperties. They also carried out ball‑milling experiments with temperature monitoring and kinetic analysis,\ndemonstrating that mechanophore activation leads to polymer degradation after acid treatment.\nThe authors have satisfactorily addressed all previous concerns, and the manuscript is acceptable for publication."}, "author_response": "Response to the reviewers’ comments and description of changes made to the manuscript for our\nsubmission on 08-Nov-2025 (Manuscript numbmer: NCOMMS-25-38649A).\nReviewer #1 (Remarks to the Author):\nThe authors have thoroughly addressed the reviewers' concerns, adding new data and insights that\nstrengthen the manuscript and its conclusions. Publication is recommended, though the following\nrelevant manuscripts have come out since the original submission and should be cited since they also\nfeature mechanophore-enhanced toughness in single networks: 10.1021/acscentsci.5c00932 and\n10.1038/s41557-025-01984-9\nResponse: We wish to express our sincere gratitude to you for your thorough and insightful\nfeedback, which has greatly improved the quality of this manuscript. The recently published relevant\nmanuscripts have been cited as References 19 and 20 in the revised manuscript.\nReviewer #2 (Remarks to the Author):\nThe authors performed additional control experiments, including a scissile cyclobutane mechanophore\n(C1) and blends of the C2 and C3 mechanophores, to clarify how activation of a non-scissile\nmechanophore contributes to the material properties. They also carried out ball-milling experiments with\ntemperature monitoring and kinetic analysis, demonstrating that mechanophore activation leads to\npolymer degradation after acid treatment.\nThe authors have satisfactorily addressed all previous concerns, and the manuscript is acceptable for\npublication.\nResponse: Thank you for your thorough evaluation of our manuscript and your kind recognition of\nour revision efforts. We appreciate you for recommending our paper for publication."}]

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