Fluorinated vitrimer elastomers with a dual temperature response

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Fluorinated vitrimer elastomers with a dual temperature response Marc Guerre, Christian Taplan, Renaud Nicolay, Johan M. Winne, and Filip E Du Prez J. Am. Chem. Soc., Just Accepted Manuscript • DOI: 10.1021/jacs.8b07094 • Publication Date (Web): 19 Sep 2018 Downloaded from http://pubs.acs.org on September 19, 2018

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Journal of the American Chemical Society

Fluorinated vitrimer elastomers with a dual temperature response Marc Guerre1*, Christian Taplan1, Renaud Nicolaÿ,2 Johan M. Winne1 and Filip E. Du Prez1* 1

Polymer Chemistry Research Group and Laboratory for Organic Synthesis, Department of Organic and Macromolecular Chemistry, Ghent University Krijgslaan 281 S4-bis, B-9000 Ghent, Belgium. E-mail: fi[email protected], [email protected]. 2 Matière Molle et Chimie, UMR 7167 CNRS-ESPCI Paris, Paris Sciences et Lettres (PSL) Research University, 10 rue Vauquelin, 75005 Paris, France. Abstract Vitrimers are an emerging new class of permanently cross-linked polymeric materials that show a liquid behavior upon heating wherein the macroscopic deformation is controlled by the rate of internal chemical bond exchange reactions. Thus, quite uniquely among polymeric materials, flow rates and material viscosities can be enhanced or controlled by the addition of catalysts and additives. We now report a catalyst-free vitrimer system, prepared from mixing two simple components, wherein two competing bond exchange mechanisms co-exist, each showing a strikingly different temperature dependence, related to the large difference in activation energy for the different exchange pathways (60 vs 130-170 kJ/mol). The low barrier process is predominant at lower temperatures, but is outcompeted by the high barrier process that becomes dominant at higher temperatures because of its much more pronounced temperature dependence. The result is an interesting and highly unusual dual viscosity profile for this new class of vitrimer materials: a very gradual decrease in viscosity at lower temperatures, intercepted by a much sharper drop in viscosity at higher temperatures. The highly counterintuitive effect where a higher barrier pathway is dominant over a much lower barrier process, can be rationalized by the exchange mechanisms that involve different reactive species, but lead to the overall same exchange. We observed this unusual but highly promising behavior first for fluorinated vitrimer elastomers, aimed at high performance materials, but the effect was also shown to hold in related non-fluorinated elastomers. A new way to control and design the rheological behavior of vitrimers towards finely tuned and precisely controlled processing applications has thus been provided. Introduction Synthetic materials that are built up from high molecular weight polymer chains derive their material properties from intermolecular forces, chain mobility and physical chain entanglements. The resultant thermoplastic polymers are easily processed, but show a limited durability, and poor resistance to deformation (creep), weathering or dissolution. For applications where aspects such as high dimensional stability and integrity are important, covalent cross-linking of polymer chains is the desired strategy. Classically, this results in thermosetting polymers that need to be synthesized or at least fully hardened (or cured) in their final shape, prohibiting most of the standard ‘plastic’ processing possibilities. However, a clear paradigm shift has emerged around the turn of the century: reversible cross-linking of polymer chains, via strong supramolecular interactions1-3 or reversible covalent bonds,4-7 can combine the easy processing and high durability of thermoplastic and thermosetting polymers. The viscoelastic properties of such materials can be changed or introduced by exposure to a specific stimulus, such as light8 or heat,9 giving a temporarily de-cross-linked 1 ACS Paragon Plus Environment

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material that can then be reshaped, reprocessed, repaired or recycled. Given the growing global demands and societal incentives towards more sustainable materials, interest in this type of polymers has grown exponentially. Covalent cross-links that can be reversed to an open form, can self-evidently transform a thermosetting topology into a thermoplastic (processable) one, via a net-depolymerisation. In 2011,10 Leibler and coworkers introduced another paradigm shift within the field of reversibly cross-linked polymer networks. Somewhat counterintuitively, for macroscopic flow to occur in a polymer network, depolymerisation is not a condition. Polymer materials can indeed be designed wherein the network integrity is preserved under all conditions, but where macroscopic flow is enabled by covalent exchange reactions that rearrange the network connectivity. For this, network defects or reactive side chain groups must be able to associate themselves with existing network points, and dislodge another polymer chain from this bonding site. Interestingly, in such materials, the viscoelastic properties or the rate of macroscopic flow are not primarily governed by intrinsic chain mobility, or by the properties of the non-cross-linked material, but is controlled by the rate of the chemical exchange reaction, as the rate-limiting step in the physical process of macroscopic flow.11 The result is a polymer material that is completely insoluble at all temperatures, but that can be processed and shows stress relaxation with a temperature dependence of the viscosity, which aligns almost perfectly with the temperature dependence of the chemical reaction. In fact, it has been found consistently that measuring the stress relaxation times of these materials at different temperatures, readily allows the retrieval of the activation energy of the exchange reaction.12, 13 In contrast to all previously studied organic polymer liquids, these non-depolymerising but deformable networks show a very gradual viscosity decrease upon heating, a thermal behavior found previously only in certain inorganic materials (strong liquids) such as vitreous silica. Hence, upon the discovery of the first such material by Leibler and coworkers, i.e. a catalyst-doped polyester epoxy network, they introduced the term ‘vitrimers’ for this new class of polymer materials. 14-25 Fluorinated polymers show remarkable differences in properties compared to classical polymer materials, and have quickly found commercially relevant applications because of their high inertness and large impact on surface properties, most notably in inert rubbers,26 coatings27 and membranes.28 Among fluoropolymers, industrially produced perfluoropolyethers (PFPEs)29 are a class of growing interest: they are liquids at room temperature, exhibit low surface energy, good oxidative and thermal stability, are characterized by excellent lubricating properties, biological inertness, radiation resistance, pressure stability and are non-flammable. Today, PFPE building blocks are widely used as reactive intermediates in the synthesis of a large range of materials for the creation of robust fluorinated networks (e.g. polyurethanes, epoxy or polyester-based networks)30-33 for aerospace materials, microfluidic devices, protective coatings, optical applications,34-40 as well as electrolytes for lithium batteries.41, 42 While the introduction of reversible covalent bonds inside a fluoropolymer matrix could increase the mechanical properties and temperature resistance, and would in principle give easily (re)processable and repairable high performance fluorinated resins or elastomers, the combination of Covalent Adaptable Networks (CANs) with fluorinated materials has, to the best of our knowledge, never been explored so far. We found one report in which Li et al.43 2 ACS Paragon Plus Environment

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reported a PFPE-based elastomer cross-linked via supramolecular interactions, showing selfhealing behavior in combination with moderate mechanical properties, as expected for a weakly cross-linked system. Because of our ongoing research in the area of vitrimer-type chemistries, leading to several interesting chemistry platforms for material design, we were highly interested in implementing these chemistries to PFPE with the aim to create fluorinated vitrimers. We report here the full details of this study, in particular aimed at introducing the catalyst-free vinylogous urethane vitrimer chemistry44-47 into cross-linked PFPE elastomers. This study not only reveals novel promising fluorinated materials with excellent properties but also the full characterization of their visco-elastic properties, which yields a unique insight into the factors that govern their macroscopic flow. Our study is the first and very clear example of a dual vitrimer relaxation behavior wherein the covalent exchange reaction between network points follows two different chemical pathways with a high and low activation energy respectively. Counterintuitively, but actually within the expectations of the underlying chemistry, the high activation energy exchange can become the dominant process at higher temperatures (130-160°C), leading to a highly unusual and so far unprecedented thermal response of the viscosity of these materials: a gradual and moderate decrease of the viscosity in the lower temperature ranges, followed by a much sharper and pronounced temperature response at higher temperatures. Our indepth study of the PFPE elastomers focused in first instance on stress-relaxation experiments, but also included creep experiments, which fully complement our findings. Moreover, we could also demonstrate that the conceptually novel dualistic temperature response effect is not related to the fluoropolymer backbone or the presence of additives or impurities, but is rather related to the network architecture, combined with the intrinsic properties of the remarkable exchange reaction. Thus, this new and peculiar behavior should hold relevance for many vitrimer-type materials. Experimental part Material Benzylamine (>99%), hexylamine (>99 %), Polypropylene glycol (PPG, Mn = 400 and 2000 g.mol-1), tert-butylacetoacetate (TBAA, >98 %) and tris(2-aminoethyl)amine (TREN, 96%) were purchased from Sigma Aldrich and used without further purification. Fluorolink® E10H (PFPE-OH, Mn ≈ 1800 g·mol-1) was purchased from Acota (UK). 1,1,1,3,3-Pentafluorobutane (Solkane® 365-mfc) was purchased from Alpha Aesar. Tetrahydrofuran (THF), dimethylacetamide (DMAc), dimethylformamide (DMF) and methanol (MeOH) were purchased from Acros Organics. Sodium sulfate (Na2SO4) and sodium hydroxide (NaOH) were purchased from Carl Roth. Hydrochloric acid (HCl, 36 %) was purchased from Chem-lab. Instrumentation Nuclear magnetic resonance (NMR) spectra were recorded on a Bruker Avance I Ultrashield 300 MHz spectrometer. Deuterated acetone was used as the solvent in each sample. Coupling constants and chemical shifts are given in hertz (Hz) and parts per million (ppm), respectively. The experimental conditions for recording NMR spectra were as follows: 1D 1H: flip angle, 30°; acquisition time, 4.1 s; pulse delay, 7.75 µs; number of scans, 16; spectral window, 8 kHz. 1D 13C: jmod pulse program was used with acquisition time of 1.36 s and 3072 number of scans. 2D COSY 1H-1H: The cosygpppqf pulse sequence from Bruker 3 ACS Paragon Plus Environment


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catalogue was used without modification; acquisition time, F1 0.01 s and F2 0.17 s; spectra window, 6 kHz for both F1 and F2 dimensions; number of scans, 8. 2D HSQC 1H-13C: The hsqcedetgp pulse sequence from Bruker catalogue was used without modification; acquisition time, F1 7.68 ms and F2 0.16 s; spectra window, 6.4 kHz for F2 and 16.6 kHz for F1; number of scans, 8. Diffusion ordered NMR spectroscopy (DOSY) experiments were performed at 25 °C with 16 number of scans and 64 planes in F1 dimension. Infrared (ATR-FTIR) spectra were collected using a Perkin–Elmer Spectrum1000 FTIR infrared spectrometer with a diamond ATR probe. Size-Exclusion Chromatography (SEC) was recorded using a triple-detection SEC from Agilent Technologies with its corresponding Agilent software, dedicated to multidetector SEC calculation. The system used two PL1113-6300 ResiPore 300 × 7.5 mm columns with DMF (containing 0.1 wt. % of LiCl) as the eluent with a flow rate of 0.8 mL·min-1 and toluene as the flow rate marker. The detector suite was composed of a PL0390-06034 capillary viscometer detector. The entire SEC-HPLC system was thermostated at 70 °C. Thermogravimetric analyses (TGA) were performed with a Mettler Toledo TGA/ SDTA851e instrument under air or nitrogen atmosphere respectively at a heating rate of 10 K·min-1 from 25 °C to 600 °C for dynamic mode or at 100 °C and 150 °C during 120 min for isothermal mode. Differential scanning calorimetry (DSC) analyses were performed with a Mettler Toledo instrument 1/700 under nitrogen atmosphere at a heating rate of 10 K ·min-1. Dynamic mechanical analyses (DMA) was performed on a SDTA861e DMA from Mettler Toledo utilizing rectangular sample (ca. 2.1 mm (T), 2.0 mm (W) and 8.0 mm (L)) and using tension clamp. The axial force was adjusted to 0.1 N and a temperature ramp was then performed from -140 °C to 100 °C at a rate of 3 K·min-1 with an oscillating force of 0.1 N and a regular frequency of 1 Hz. Rheology experiments were performed on an Anton Paar MCR 302. The experiments were performed in parallel plate geometry using 8 mm sample disks. Unless specified, the experiments were performed using a normal force of 1 N, an oscillating frequency of 1 Hz and a strain of 1 %. For all rheology experiments, the applied stress was comprised in the linear viscoelastic region at the measured temperatures. For amplitude sweep experiments, the strain was varied from 0.01 to 100 %. For time sweep experiments, a frequency of 1 Hz and a strain of 1 % were applied and the storage modulus (G’) was followed over time at a constant temperature. For stress relaxation experiments, a strain of 1 % was applied to the material and the relaxation modulus (G(t)) was followed over time at a constant temperature. The series of stress relaxation experiments at different temperatures were performed successively on the same sample. Creep experiments were performed using a stress of 2800 Pa and the resulting strain was followed over time at constant temperature. Creep recovery experiments were performed using a stress alternating 0 Pa and 2800 Pa. In a typical test, a stress of 2800 Pa was applied for 300 s, before a 100 s-recovery period. This cycle was repeated multiple times and the corresponding strain profile was recorded over time. 4 ACS Paragon Plus Environment

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Uniaxial Tensile testing was performed on a Tinius-Olsen H10KT tensile tester equipped with a 100 N load cell using ASTM standard type IV dog bones (ISO 527-2-2B). The dog bone shaped samples had an effective gauge length of 12 mm, a width of 2 mm and a thickness of ± 2 mm and they were cut using a Ray-Ran hand operated cutting press. The tensile measurements were performed using a preload of 0.05 N and a pulling speed of 10 mm·min1 until sample failure. The stress was recorded as a function of strain . Reported values (elongation (%), stress at break (Pa) and young modulus (Pa) are the average and standard deviations of at least seven samples. Exhibited tensile curves were selected to be representative of the average values. Reprocessability. To reprocess the network, the polymer was broken into small pieces (< 2mm) and placed into a rectangular mold (70 mm x 40 mm x 2 mm) for compression molding. This assembly was placed in a 150 °C preheated compression press for 2 min under 1 metric ton of pressure. Then the pressure was increased to 3 tons and kept constant for an additional 3 min. After 5 min of pressing in total, the mold was cooled down for 30 min and the sample was removed from the metal sheet with a razor blade. Solubility and hydrolytic stability experiments. Solubility tests were carried out with samples of 4 mm diameter and 2 mm of thickness with a weight of around 60 mg and 5 mL of diverse solvents such as: DMF, MeOH, THF, and Solkane. Those tests were performed for 24 h at 25 °C in DMF, MeOH, THF, and Solkane, and also at 50 °C in THF and 80 °C in DMF. Then, the solvent was removed and the samples were dried under vacuum, overnight at 100 °C. The soluble fraction was calculated using equation (1), while the swelling ratio was calculated using equation (2). − 1 % =

% =

−

2

with mi, ms and md standing for initial, swollen and dry mass, respectively. As a control experiment, Soxhlet extraction was also carried out in refluxing THF for 24 h in THF. Typically, 1 g of sample was placed in an extraction unit of known weight and subsequently immersed in dried THF as solvent. Next, the sample was dried at 100 °C under vacuum overnight. The remaining dried sample was weighted and the soluble fraction was calculated using equation (1). Hydrolytic stability experiments were carried out using a similar sample size. Three different conditions were tested: neutral (deionized water), basic (NaOH(aq) 1M) and acidic (HCl(aq) 1M). Tests were performed during 24 h at both 25 °C and 80 °C. Soluble fractions and swelling ratios were calculated following equation (1) and (2), respectively. Synthetic Procedures Bis-perfluoropolyether acetoacetate (PFPE-AA): Fluorolink® E10H (PFPE, Mn ≈ 1800 g.mol-1, 100 g, 5.56x10-2 mol) and tert-butyl acetoacetate (19.3 g, 1.22x10-1 mol) were added in a 250 mL bottom flask. The viscous mixture was heated for 60 min at 130 °C, and tert-butanol product was continuously removed by distillation throughout the reaction. When the 5 ACS Paragon Plus Environment


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temperature of the still head dropped to 40 °C and tert-butanol condensation stopped, the unreacted tert-butyl acetoacetate was removed by distillation at 130 °C under 1 mbar vacuum. Yield = 99 %, functionality = 98 %, 1H NMR (300 MHz, (CD3)2CO, δ ppm): 1.92-1.97 (CH=C(OH)-CH3, enol form (8.5 %)), 2.19-2.26 (-CH2-(C=O)-CH3), 3.49-3.75 (OCF2CH2-OCH2CH2OCH2CH2- and -(C=O)-CH2-(C=O)-CH3), 3.76-3.96 (-OCH2CH2-O-(C=O)-CH2-), 3.98-4.17 (OCF2CH2-), 4.18-4.35 (-OCH2CH2-O-(C=O)-CH2-), 5.03-5.10 (-CH=C(OH)-CH3, enol form (8.5 %)). Bis-perfluoropolyether hexyl vinylogous urethane (PFPE-HVU): PFPE-AA (Mn ≈ 2000 g.mol-1, 3.0 g, 1.50x10-3 mol) and hexylamine (1.52 g, 1.50x10-2 mol) were mixed in a 25 mL bottom flask and stirred overnight at 60 °C. Then, water was added and the mixture was extracted with Solkane twice. The combined fluorinated phases were dried over Na2SO4 and the fluorinated solvent was evaporated in vacuo yielding the desired compound. Yield = 76 %, functionality = 98 %, 1H NMR (300 MHz, (CD3)2CO, δ ppm): 0.82-0.98 (-N(H)CH2(CH2)5CH3), 1.23-1.65 (-N(H)CH2(CH2)4CH3), 1.90-1.96 (-CH=C(NH-)-CH3), 3.20-3.34 (-N(H)CH2(CH2)5CH3), 3.44-3.92 (OCF2CH2-OCH2CH2-OCH2CH2-O-(C=O)-), 3.95-4.22 (-OCF2CH2- and -OCH2CH2-O(C=O)-), 4.34-4.44 (-CH=C(NH-)-CH3), 8.38-8.83 (-N(H)CH2(CH2)5CH3). PPG-AA building blocks and PPG-VU networks were prepared following the same conditions reported above for the PFPE-AA and PFPE-VU network. Model studies: Benzylamine (0.194 g, 1.82x10-3 mol) was added to a solution of PFPE-HVU model compound (0.400 g, 1.83x10-4) in DMAc (1.0 mL). Five equivalents of benzylamine were used to obtain a pseudo-first order reaction at low conversions. The mixture was heated (100 °C, 120 °C, and 140 °C) in a Schlenk tube equipped with a septum. Samples were taken at different time intervals, quenched in liquid nitrogen and analyzed by 1H NMR in (CD3)2CO. The reaction was followed by integration of the two distinct signals at 8.61 ppm and 8.96 ppm for the NH of hexyl and NH of benzyl model compounds, respectively. Network synthesis: PFPE-AA (fct. = 98 %, 20 g, 1.0x10-2 mol) and tris(2-aminoethyl)amine (1.00 g, 6.86x10-3 mol) were mixed in a 60 mL poly(propylene) vial using a DAC 150.1 FVZ speed mixer (typical conditions of mixing: 5 min with a speed of 2500 rpm). Then, after gelation, the polypropylene vial was broken, and the network was cured overnight at 80 °C under vacuum. Finally, the sample was cut into two pieces of 10 g, and pressed separately in steel molds, following the aforementioned reprocessability procedures. Results and Discussion PFPE functionalization and model reaction Bis-perfluoropolyether acetoacetate (PFPE-AA) was synthesized from Fluorolink® E10H (PFPE-OH), following a straightforward, high yielding acetoacetylation reaction as depicted in Scheme 1 (top).

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Scheme 1. a) Functionalization of Fluorolink E10H (PFPE-OH) into PFPE-AA, b) network synthesis via polycondensation reaction of acetoacetate and amine leading to vinylogous urethane bonds and c) general scheme for transamination of vinylogous urethanes.

The acetoacetylation reaction between the bulk chemical TBAA and polyol precursors allows for the simple preparation of multifunctional acetoacetate building blocks in large scale. This strategy has the advantage of releasing a low boiling point compound, tert-butanol, which can easily be removed from the reaction by distillation. Furthermore, the mixture does not require supplementary purification steps since the remaining excess of TBAA can be conveniently distilled under reduced pressure. In addition to this facile purification protocol, the reaction can be carried out in bulk, thus not requiring any organic or fluorinated solvent. Indeed, the viscous liquid Fluorolink E10H is miscible with tert-butylacetoacetate at 130 °C. This strategy differs from prior PFPE functionalization procedures, which report the use of fluorinated solvents, commonly classified as highly toxic and bio-accumulative. 31, 32, 48 The obtained PFPE-AA structure was confirmed by 1D and 2D NMR experiments (1H and 13C). Figure 1 displays the 1H NMR spectra of TBAA (1), Fluorolink E10H (2), and the prepared PFPE-AA (3). F FF F O

O

O

m

b

a O

q

c

F F F F

O n-1

O

e O

d

1H

O

g

f

g O

f d b+c

a

h

i

X

H2O

a

h

OH

(CD3)2CO

e 3)

O

i

X

b c b+c a

2)

X c

c a

b

a b

1)

X 5.0

4.5

4.0

3.5

3.0 (ppm)

2.5

2.0

1.5

1.0

1

Figure 1. H NMR spectra of: 1) tert-butylacetoacetate, 2) Fluorolink E10H and 3) PFPE-AA in (CD3)2CO at 25 °C.



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Upon functionalization, a significant downfield shift was observed for the ethylene glycol signals linked to the acetoacetate group. Indeed, signals assigned to the -CH2CH2-OH group observed in the area 3.50-3.90 ppm for Fluorolink E10H, were shifted to 3.76-3.96 ppm and 3.98-4.17 ppm for -CH2CH2-O-AA and -CH2CH2-O-AA, respectively. The characteristic acetoacetate signals were also shifted after the functionalization. The corresponding methyl signal can clearly be identified as a sharp singlet at 2.22 ppm, which has slightly shifted in comparison to the one of the starting TBAA. However, the signal assigned to the CH2 of the acetoacetate function is not clearly visible and possibly overlaps with signals of the ethylene glycol unit. The exact position of the CH2 signal (at 3.53 ppm) in the 3.49-3.75 ppm region was confirmed by a 2D correlation via COSY 1H-1H experiment (Figure S1). The chemical structure was also confirmed by 13C, 1H DOSY NMR and ATR-FTIR (Figure S2-S5). The DOSY 1 H NMR spectrum (Figure S4) shows the presence of a single diffusion coefficient for the PFPE and acetoacetate signals, which suggests a successful and quantitative functionalization of the alcohol moieties. The FTIR spectrum (Figure S5) also confirms the qualitative character of the acetoacetalization from the disappearance of the broad OH vibrational band of Fluorolink E10H above 3400 cm-1 and the appearance of new ketone vibrational bands at 1719 and 1740 cm-1. Considering the importance of chain-end functionality in stoichiometry calculations, the acetoacetate chain-end functionality was carefully quantified by 1H NMR. Special attention was directed to the reference peak (-OCF2CH2- signal, 4.18-4.35 ppm), since a small degree of coupling was reported for the industrial synthesis of Fluorolink E10H.42 Indeed, the SEC trace of Fluorolink E10H (Figure S6) clearly shows the presence of mono-, di-, tri- and tetraadducts at variant elution times. It is worth noting that the result can be severely affected if these coupling reactions are not adequately considered in functionality calculations. Thus, by simply considering the -OCF2CH2- signal as reference (four protons), and by assuming that this -OCF2CH2- group is always alcohol end-terminated, the amount of acetoacetate functions would be underestimated because the calculation does not take into consideration the proportion of -OCF2CH2- per polymer chain arising from coupling reactions. Nonetheless, this fraction can be estimated by deconvolution of the PFPE SEC trace. Thereby, the percentage of PFPE-AA acetoacetate functionality was finally evaluated to be more than 98 % (See Figure S7 and supporting information for calculation details). To further confirm the functionality of the PFPE precursor, PFPE-AA was reacted with a precise amount of mono-functional benzyl amine. After reaction of 0.90 and 1.00 eq of benzyl amine with respect to AA functionalities, a functionalization of 89 and 98 % was achieved, respectively (Figure S8). These results are in a good agreement with the calculation achieved by 1H NMR. The profound inductive effect of fluorine on the chemical reactivity of organic compounds is well-described in organic and polymer chemistry.49 Therefore, to dismiss a potential remote withdrawing effect of the fluorinated chain on the transamination exchange kinetics of vinylogous urethane, a model study with low MW compounds was conducted. In order to compare the results and the respective activation energy to a previously reported study by Denissen et al.,46 the kinetic study strictly followed the therein described conditions. Firstly, a PFPE-hexyl vinylogous urethane (PFPE-HVU) was synthesized by condensation reaction of PFPE-AA with hexylamine. Then, PFPE-HVU was reacted with five equivalents of benzylamine at different temperatures (100 °C, 120 °C and 140 °C), and samples were taken at different time intervals and subsequently analyzed by 1H NMR in deuterated acetone (Figure 2). The 8 ACS Paragon Plus Environment

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exchange kinetics can easily be followed by integration of the two distinct signals at 8.61 ppm (NH of hexyl-VU) and 8.96 ppm (NH of benzyl-VU), respectively (Figure S9).

Figure 2. Progression of the exchange reaction of PFPE-HVU and benzylamine at 100 °C, 120 °C and 140 °C.

Kinetics displayed in Figure 2 evolve as a first-order linear decay at low conversion and tend to a chemical equilibrium. Furthermore, the kinetic curves clearly show a temperature dependence of the exchange reaction. From these kinetic studies, resulting rate constants and Arrhenius law, an activation energy of 69 ± 2 kJ.mol-1 was calculated (Figure S10), which is in good agreement with those earlier reported (59 ± 6 kJ.mol-1 and 73 ± 12 kJ.mol-1).45, 46 In other words, the withdrawing effect of perfluorinated chains does not significantly influence the transamination exchange of vinylogous urethane groups and hence, can be neglected for the following discussion. Vitrimer synthesis and molecular characterization PFPE vinylogous urethane (PFPE-VU) vitrimers were prepared through the condensation reaction of PFPE-AA and cross-linker tris(2-aminoethyl)amine (TREN) (Scheme 1, b) in bulk at room temperature. Due to the high viscosity of PFPE-AA, which further increased upon the reaction, a speed mixer was used to enhance the homogeneity of the mixture. Since the polycondensation reaction between acetoacetate and primary amine readily occurs at room temperature, the gel point is reached after just a few minutes. Then, the vitrimer was isolated as an orange rubber (circular cube shape, Figure S11) before an additional postcuring step of 16 h at 80 °C under vacuum was performed. The curing step not only allows quantitative transformation of the acetoacetate groups into vinylogous urethane, but also removes water arising from the polycondensation. The obtained fluorinated vitrimer was subsequently pressed at 150 °C in its final shape (Figure S11). In the transamination reaction of vinylogous urethane moieties, free primary amines play a determinant role in the exchange mechanism (Scheme 1, c). Indeed, they not only promote exchange reactions without loss of integrity, but also allow a control over the exchange kinetics. In other words, an excess of free primary amines in vinylogous urethane vitrimers could significantly accelerate the exchange reactions, while an amine shortfall could 9 ACS Paragon Plus Environment


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drastically decrease or even stop them. Thus, in order to evaluate the effect of free amines with respect to exchange kinetics and to the resulting material properties, three PFPE-VU vitrimers with different compositions were made: PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10 with the subscript referring to 0, 5 and 10 mol% excess primary amine groups, respectively. FTIR-spectra of the obtained PFPE-VU vitrimers are displayed in Figure S12. The progress of the reaction was qualitatively confirmed by the conversion of the acetoacetate group into vinylogous urethanes as shown by the complete disappearance of ketone vibration bands at 1719 and 1740 cm-1 and by the appearance of new vibration bands at 1601 and 1650 cm-1, assigned to C=C and C=O of vinylogous urethane, respectively. No significant differences were noticed between the three PFPE-VU vitrimers and no vibration band assigned to residual water was observed around 3000 cm-1. The quantitative elimination of water was further confirmed by isothermal TGA thermograms, in which no weight loss was observed after two hours at 100 °C (Figure S13). Material characterization of vitrimers Notably, the displayed DSC thermograms (Figure S14) reveal two distinct glass transition temperatures: one transition at -41 °C assigned to the amine core and vinylogous urethane structures, and a second Tg around -100 °C assigned to the fluorinated PFPE domains (similar to the Tg of Fluorolink E10H, Figure S14 black curve). This observation was expected since fluorinated polymers are prone to phase separation in the presence of hydro-carbonated molecules and polymers.34, 35 DMA-experiments performed for PFPE-VU vitrimers also confirm the presence of two Tg values (Figure S15), i.e. a pronounced and broad glass transition around -40 °C and a second transition at -90 °C. Those values are in good agreement with the Tg of PFPE-methacrylate thermosets reported by Hu et al.35 via DMA (Tg1 = -55 °C and Tg2 = -80 °C). The appearance of a constant plateau modulus in the range of 1.6 MPa further confirms the cross-linked structure. Similar moduli were reported for permanently and non-permanently cross-linked elastomers. For instance, Ishibashi et al.50 found a modulus of 0.4 MPa for PDMS-based vitrimers, while Vitale et al.51 reported a modulus of 6.3 MPa for highly cross-linked PFPE networks. As reported by Desimone et al. for PFPE/PEG blends, the bulk morphologies of segregated networks are known to be heterogeneous and opaque.37 Indeed, although the PFPE-VU vitrimers appear transparent after their preparation, those with 5 mol.% and 10 mol.% of free amines turned turbid within three days at room temperature. This suggests a dynamic phase separation within the fluorinated matrix at room temperature from nano-sized domains, which do no scatter light, to micro-sized domains. Indeed, due to the low Tgs, the overall mobility of both hydrocarbonated and fluorinated domains is high at room temperature, which also implies that exchange reactions are taking place at this temperature on this time scale. Furthermore, after exposing the PFPE-VU vitrimers for two minutes at 150 °C, the turbidity vanishes and the materials become transparent again (Figure S16). This reversible transition is likely the result of hydrocarbonated domains rearrangement from microscopic into nanodomains because of dynamic exchanges. Finally, the contrasting evolution of turbidity for the vitrimers with different compositions can be rationalized by taking into account the different excess of free amine groups. Indeed, as will be reported in the creep analysis section (vide infra), an excess of primary amine groups promotes a faster deformation upon stress while the PFPE-VU1.00 vitrimer exhibits a much 10 ACS Paragon Plus Environment

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less pronounced creep. Thus, we assume that the amine groups enable a better migration and reorganization of the hydrogenated domains within the fluorinated matrix. The thermostability of all vitrimers was analyzed by TGA under both nitrogen and air atmosphere (Figure S17). With a heating rate of 10 K.min-1, the onset of degradation (5% weight loss) is observed around 280-290°C in nitrogen and 265 °C in air for the three samples. These results are consistent with the values obtained for robust48 permanently cross-link PFPE-based networks reported in the literature. Thus, the reversible vinylogous urethane chemistry does not have a significant impact on the thermal stability of PFPE-based materials. To estimate the effective thermal stability of PFPE vitrimers, isothermal experiments were also conducted. A weight loss of 1 % was observed after two hours in air at 150 °C (Figure S18) while the thermal degradation temperature of PFPE vitrimers in air was about 265 °C. In other words, such PFPE vitrimers cannot be continuously exposed over a longer period to temperatures exceeding 150 °C without taking into account a potential damage of the material. However, since the time required for reprocessing is in the order of minutes at 150 °C (vide infra), no significant influence is expected from network degradation. In the next step, viscoelastic properties of PFPE-VU vitrimers were in first instance studied by stress relaxation experiments (Figure 3, S19 and S20). For the rheological measurements, a constant strain of 1 % was applied, which is comprised in the linear viscoelastic region at the measured temperatures, and the relaxation modulus was followed as a function of time.

Figure 3. Representative stress relaxation experiments and Arrhenius plots of PFPE-VU vitrimers (a) Stress relaxations of PFPE-VU1.00 (black), PFPE-VU1.05 (red), and PFPE-VU1.10 (blue) vitrimers at 150 °C (b) Stress relaxation curves of PFPE-VU1.05 for temperature interval between 90 °C and 160 °C. The dotted line represents E/E0 = 1/e, which defines the characteristic relaxation time ԏ* (c) Arrhenius plots of PFPE-vitrimers with corresponding exchange mechanisms involved in the dual relaxation behavior.



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The stress relaxation curves shown in Figure 3a and b, show a full stress relaxation as a function of time. This relaxation behavior, which follows the Maxwell model, is the result of a vitrimer network that is entirely composed of reversible bonds without irreversible crosslinks. As expected, relaxation time values (ԏ*) at 150 °C decreased with an increasing content of free primary amines in the network (Figure 3a), from approximately 390 s (PFPEVU1.00), over 160 s (PFPE-VU1.05) to 57 s (PFPE-VU1.10). Also, when the temperature is increased from 90 °C to 160 °C for the PFPE-VU1.05 sample, ԏ* similarly decreased much from 4390 s to 33 s, which is consistent with regard to other reported vitrimer systems. From those relaxation times, the relationship between Ln(ԏ*) and 1000/T was fitted to the Arrhenius equation (Figure 3c) and the corresponding activation energy was calculated from the slope. In the initial study of vinylogous urethane vitrimers, Denissen et al.46 reported a good agreement between the Ea determined via chemical kinetic studies with low molar mass model compounds on the one hand, and the value determined via vitrimer stress relaxation measurements on the other hand. In contrast, within these systems containing much longer chain fragments, we observed a significant shift in Ea, which not only depends considerably on the network stoichiometry, but remarkably also on temperature: Ea of 58 ± 4 kJ.mol-1 for PFPE-VU1.00, 60 ± 2 and 169 ± 17 kJ.mol-1 for PFPE-VU1.05 as well as 58 ± 3 and 129 ± 3 kJ.mol-1 for PFPE-VU1.10 (from two different slopes). Previously, such shifts in activation energy had only been observed in catalyst- or additive-doped vinylogous urethane vitrimers,45 and they have never been observed within one and the same material. As no significant difference of shear modulus was noticed between all vitrimers (see time sweep experiments in Figure S21), this variation in activation energies can neither be related to a difference of cross-linking density nor to a change in chain mobility. Given what is known about the molecular basis of the rheology of vitrimers, we hypothesize that the observed peculiar shift in activation energy, with a stronger temperature dependence at higher temperatures, is related to the dynamic exchange mechanism. Indeed, the Arrhenius plots of Figure 3c can be explained by the presence of two distinct mechanisms (or pathways) for the same dynamic bond exchange process: one mechanism or pathway corresponding to an activation energy barrier of ca. 60 kJ.mol-1, and a second one corresponding to an enthalpic barrier of ca. 130-170 kJ.mol-1. Interestingly, PFPE-VU1.05 and PFPE-VU1.10 show this dual exchange mechanism on the same Arrhenius plot, while the transition point appears to depend on the network composition. Based on the Arrhenius plot, this transition was estimated to be ca. 145 °C and ca. 130 °C for PFPE-VU1.05 and PFPEVU1.10, respectively. To verify the reproducibility and the significance of this dual temperature response, stress relaxation experiments were repeated 3 times on the same PFPE-VU1.05 vitrimer (see error bars in Figure 3). Arrhenius plots of those experiments (Figure S22) clearly show that the dual-viscosity behavior remains distinctively visible. The herein described presence of a dual exchange mechanism within one material has never been reported so far for any vitrimer material, and thus deserves closer scrutiny. In previous studies, we have found45 that under neutral or slightly acidic conditions, the amine exchange invariably proceeds via a ‘protic pathway’ involving an activated iminium intermediate as electrophilic species, corresponding to an enthalpic barrier (Ea) of about 73-81 kJ.mol-1 (Figure 3c, pathway (i)). We have also found that adding a strong base blocks this pathway and rather promotes a much slower direct addition pathway between a neutral free amine 12 ACS Paragon Plus Environment

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and a neutral VU-moiety as a Michael-type acceptor (Figure 3c, pathway (ii)). This ‘background’ process was previously only revealed in the presence of a strong base, and was found to correspond to a much higher activation barrier (Ea) ranging from 102 to 122 kJ.mol1 . By proposing these known distinct transamination mechanisms as competing relaxation pathways within PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10, the observed rheological behavior can thus be satisfactorily understood at a molecular level, i.e. at sufficiently low temperatures, exchange will proceed via an iminium exchange, while at higher temperatures, the higher barrier process will become the dominant exchange pathway, as this reaction has a much more pronounced temperature dependence. Another clear indication of the two distinct pathways can be found in the observation that the high barrier (Michael-type) addition pathway is more strongly dependent on the free amine content of the material, as can be seen in Figure 3c: the steeper Arrhenius fitted lines, corresponding to the higher barrier exchange, drop more significantly with the increase of the free amine content. In the low barrier iminium pathway (i), attack of an amine species to the very electrophilic (and high energy) protonated intermediate will be rapid and is likely non-rate determining. In contrast, the bimolecular step between an unactivated, neutral VU and an amine will be rate determining for pathway (ii), as the only other reaction step involved is an internal proton transfer between the exchanging amine groups, which should be very fast. Remarkably, the high barrier mechanism is absent or kinetically insignificant in all other investigated vinylogous urethane based materials, while it is quite pronounced in the fluorinated vitrimers studied here. Because this unprecedented and quite unexpected double relaxation behavior can have serious implications for vitrimer processing applications, and could allow for the rational design of new materials with highly ‘tailored’ rheology profiles, we examined if this novel dual temperature response effect is related to the fluoropolymer backbone - or the presence of additives or impurities therein - or if it is rather related to the specific network architecture that includes long chain fragments between the exchangeable VU cross-links. All previous VU vitrimers were built up from small molecule monomers only, and did not contain larger non-dynamic fragments or chains. In fact, we found a very similar dual relaxation profile based on a VU vitrimer derived from a non-fluorinated poly(propylene glycol) backbone of comparable MW (PPG-VU1.05, Figure S23), implicating the molecular architecture as the main facilitator of the high energy pathway, and also pointing towards the general applicability of this dual rheology as a novel design concept in vitrimer chemistry. With respect to the established operative exchange mechanisms in these VU materials, it is worth noting that PFPE-VU1.00, which theoretically should not have free amine groups, also exhibits a full stress relaxation. It is hypothesized that this behavior can be related to minor stoichiometry discrepancies and network defects as ‘exchange hotspots’ in the vitrimer matrix. This further confirms the view that VU-protonation to the reactive iminium form is the most relevant step in the dynamic exchange mechanism, as a depletion of free amines does not shut down the exchange pathway. As expected, based on the bimolecular rate determining step, the second pathway is absent in stoichiometrical balanced VU networks. To verify if relaxation could also occur within a network without any free amines, the synthesis of a control sample was attempted with 5 % excess of acetoacetate moieties. 13 ACS Paragon Plus Environment


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However, even after 24 hours at 80 °C, the gel point for the reaction between PFPE-AA and TREN was not reached and the mixture remained liquid. In 2011,10 with the first vitrimer report, the correlation between the chemical relaxation and the viscoelastic properties resulted in the introduction of a specific temperature associated to vitrimer systems, i.e. the topology freezing transition temperature, labeled Tv. By definition, Tv is defined as the temperature above which a reversible network topology can be achieved through dynamic exchange. More precisely, the Tv has been defined as the temperature at which solid to liquid transition occurs with a defined viscosity of η = 1012 Pa·s.11 To estimate the related Tv values, the Maxwell equation is typically used to correlate the relaxation time to Tv (see SI for calculation details). A hypothetical Tv of around 10 °C was calculated for the fluorinated vitrimers (vide infra). Since for those elastomers, the Tv is much higher than the detected low Tg, the material viscosities will be mainly governed by the Tv,12 meaning that exchange reactions are expected to already occur at room temperature, which is perfectly consistent with the observed turbidity appearance after 3 days at ambient conditions. To evaluate the effect of such low Tv-value on the elastomeric material properties, cyclic creep recovery experiments, i.e. with the stress varying between 0 and 2800 Pa, have been conducted for PFPE-VU1.00, PFPE1.05 and PFPE-VU1.10 at 25 °C (Figure S24). The PFPE-VU vitrimers behave like a standard rubber at room temperature52 on these time scales, with a nearly full recovery of their initial shape after successive shearing. For the evaluation of creep at higher temperature, a similar experiment was conducted with temperature variations between 25 and 80 °C (Figure 4).

Figure 4. Cyclic creep recovery profiles of PFPE-VU1.05 vitrimer at 25 °C and 80 °C. The lower dashed black line corresponds to the applied force (2800 Pa) while the upper red line shows the strain response.

As expected from the acceleration of the dynamic exchange at high temperature, the materials’ response at 80 °C is completely different. Upon a first administration of 2800 Pa at 80 °C, the PFPE-VU1.05 elastomer instantaneously deforms about 1.2 % and does not return to its initial shape after removing the applied stress. After two cycles at 80 °C, a deformation of 1.5 % is observed, while a permanent deformation of 0.9 % was obtained after stress release. Nonetheless, upon cooling to 25 °C, the PFPE-VU1.05 elastomer regains its initial properties and behaves as a typical rubber with a full recovery of its shape after three cycles. 14 ACS Paragon Plus Environment

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To further explore the differences of creep in relation to the amine content for PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10 vitrimers, creep recovery experiments at different temperatures were achieved. (Figure 5, Figure S25).

Figure 5. (a,b) Creep recovery profiles of PFPE-VU1.00 (a) and PFPE-VU1.05 (b) at 20, 40, 60, 80 and 100 °C with an applied stress of 2800 Pa for 20 min and a recovery of 20 min (c) Creep rates of PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10 from 20 °C to 100 °C (d) Arrhenius plots of PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10 extracted from creep recovery experiments.

Figure 5a,b shows that PFPE-VU1.05 and PFPE-VU1.00 can withstand temperatures above 40 °C without significant deformation, while its response at high temperature is much more significant. A similar temperature dependence was observed for PFPE-VU1.10 vitrimers (Figure S25). From those creep recovery experiments, the rate of deformation, also called creep rate (ɛ),53, 54 was plotted as a function of the temperature (Figure 5c). A similar creep onset for the different PFPE-VU vitrimers can be observed, i.e. no significant creep rate is observed until 40 °C but it increases suddenly with the temperature and at different rates according to the vitrimer composition. This difference in creep behavior, especially at high temperature, can again be related to the amount of excess amine groups: an excess of amine promotes fast deformation upon stress as a result of better ability of the network to rearrange its structure through the underlying associative mechanism. However, although an increase of free primary amines allows a fast relaxation of the stress and thus fast reshaping/recycling of the network, it can also significantly affect the mechanical properties and oxidative sensitivity of the resulting materials. For that reason, the amount of free amines must be judiciously selected as a function of the desired applications and specifications.



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It is important to mention that the determination of Tv from the average slope of the Arrhenius data points, without taking the dual exchange mechanism into account, results in Tv values of 9 °C, 21 °C and 50 °C for PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10, respectively. However, such Tv values are in disagreement with the results of creep experiments, where the creep onset temperatures are the same for all PFPE-VU vitrimers. Moreover, from the creep recovery experiments, the relationship between ln (ɛ) and temperature (1000/T) was fitted to the Arrhenius equation55 (Figure 5d) and the corresponding activation energy was assessed from the slope. Activation energies of ca. 60 kJ.mol-1 were obtained in relatively good agreement with activation energies determined from stress relaxation experiments. This match in temperature dependence confirms the fact that at low temperature, exchange reactions continue to be governed by the low barrier process. Thus, even though the estimated Tv, obtained by extrapolation from a high temperature Arrhenius-fitted line, should give an approximate idea of the temperature where the dynamic exchange becomes relevant, this can be an ambiguous and flawed way to characterize the viscoelastic behavior of vitrimer materials. Indeed, it should always be taken into account that a lower barrier background exchange reaction can take over at lower temperatures, leading to a lower Tv than anticipated based on the measurement of the high temperature rheology. In this context, for low Tg materials, creep experiments should be regarded as a valuable and complementary analytical technique for a more rigorous characterization of vitrimers in addition to the high temperature stress relaxation experiments. Reprocessability, soluble fraction/swelling ratio and hydrolytic stability To evaluate the mechanical properties of the PFPE-VU vitrimers after successive recycling steps, the vitrimers were reprocessed five times using each time the same conditions, i.e. thermal compression molding at 150 °C for a duration of 5 minutes after a grinding process. Tensile strength experiments were achieved at 25 °C before each recycling step (Figure S26S29). The stress at break, Young modulus and elongation at break of these samples are given in Figure 6. Initially, PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10 exhibited similar properties. PFPE-VU1.05 and PFPE-VU1.10 recovered 95 % of their original properties and maintained this level, even after multiple recycling steps. Conversely, a noticeable decrease in stress at break from 0.93 MPa to 0.59 MPa upon the first reprocessing step of PFPE-VU1.00 was observed. After additional recycling steps, the stress at break progressively increased to a plateau of 0.73 MPa, but yet remained 22 % lower than the initial values. This drop in mechanical properties is probably the result of an insufficient healing step due to the lack of free amines, i.e. after grinding the network into small pieces, the low amount of free amines is not sufficient to efficiently fuse all pieces together and restore the original properties. The recycled PFPE-VU vitrimers were further analyzed by ATR-FTIR, DSC, TGA and rheology. On FTIR spectra, the recycled samples exhibit identical band vibrations without additional signals that could indicate the alteration of the network structure through the occurrence of side reactions (Figure S29). Accordingly, the same Tgs were observed on DSC thermograms (Figure S30). The detected thermal stability, however, showed a slight decreased from 286 °C to 276 °C compared to the non-reprocessed vitrimers (Figure S31). Stress relaxations experiments and the resulting Arrhenius plots were also reproduced for a recycled sample, again exhibiting the remarkable dual relaxation profile (Figure S32). Finally, time sweep experiments (Figure S33) showed identical shear modulus for the original and recycled 16 ACS Paragon Plus Environment

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samples, suggesting that the cross-linking density remains constant over successive recycling steps.

Figure 6. Elongation, Young modulus and stress at break values of PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10 after different recycling steps (up to five). Note: Full lines are a guide for the eyes.

The soluble fraction and swelling ratio of PFPE vitrimers are displayed in Table 1. The soluble fractions from swelling experiments performed in different solvents (THF, DMF, MeOH and Solkane) remained below 8 %, even at high temperature, which confirms the high and permanent cross-link density of the materials. At 25 °C, the soluble fractions of the original vitrimers were even lower than 1 %. The increase in free amine content in off-stoichiometric networks do not significantly affect the swelling of the network. As expected, the highest swelling ratio is obtained in the fluorinated solvent (i.e. Solkane, 100-120 %). However, the soluble fraction in such a solvent is similar to DMF and THF at high temperature, meaning that PFPE has also a good solubility in those solvents. Importantly, the soluble fraction and swelling ratio of PFPE vitrimers after five recycling steps are comparable to the original vitrimers, i.e. an increase of maximum 2 % soluble fraction and 60 % swelling ratio is observed. Additionally, the more significant increases are observed with to PFPE-VU1.00, which is consistent with a decrease of tensile strength and modulus in time sweep experiments. This again indicates the slight decrease of the cross-linking density, due to the incomplete healing of the material.



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Table 1. Soluble fraction and swelling ratio of PFPE-VU1.00, PFPE-VU1.05 and PFPE-VU1.10 vitrimers before and after five recycling steps. Soluble fraction 1.00 1.05 1.10 1.00 1.05 1.10 Swelling ratio (RC x5) (RC x5) (RC x5) (%) THF 25 °C 0.5 0.8 0.9 1.5 1.3 1.2 31 35 36 33 36 35 DMF 25 °C 0.8 1.0 1.0 0.9 1.0 1.1 18 19 19 18 20 19 MeOH 25 °C 0.6 1.0 1.0 1.9 1.1 1.1 6 8 8 6 7 7 Solkane 25 °C 3.8 5.0 5.2 7.0 7.0 6.7 101 120 119 160 127 124 DMF 80 °C 5.0 6.2 6.0 7.3 7.5 6.2 19 20 20 19 22 21 THF 50 °C 5.3 6.6 5.4 7.1 7.9 6.5 30 35 34 31 35 30 Note: A control experiment using Soxhlet extraction for PFPE-VU1.05 vitrimer in refluxing THF for 24 h resulted in a soluble fraction of 6.0 %. As expected for VU vitrimer materials, a treatment with an excess of primary amine in a good solvent at high temperature resulted in a complete dissolution of the vitrimers,46, 56 caused by dynamic amine exchanges leading to a net depolymerisation (Figure S34). Finally, in view of potential applications for fluorinated vitrimers, the first examples of which we report herein, the hydrolytic stability of our materials was tested by subjecting PFPEVU1.05 samples to three different hydrolytic conditions: 1M NaOH, neutral water and 1M HCl solutions (Table S2). After 24 h at 25 °C in those conditions, the general appearance of the samples remained the same (Figure S35). Generally, the PFPE-VU vitrimers were remarkably stable under both neutral and basic conditions ranging from 25 °C until 80 °C (See Table S1, soluble fraction and swelling ratio). However, the stability in acidic conditions is less pronounced. At 25 °C in 1M HCl solution, the network keeps its integrity but it does markedly deteriorate when heated to 80 °C in this strong acidic environment, as the sample becomes brittle and degrades into small fragments. Thus, the material’s stability is preserved under basic and neutral conditions, even at high temperature, but suffers from hydrolysis under harsh acidic conditions. Conclusion In summary, fluorinated vitrimers were prepared by implementing the catalyst-free transamination of vinylogous urethanes into a perfluoropolyether (PFPE) network matrix. These fluoroelastomers were synthesized in two simple steps from a commercially available ethoxylated prepolymer, using first a bulk end group functionalization with an acetoacetylating agent. This straightforward setup allows the preparation of PFPE precursors in very high yield through a cost-effective and solvent-free protocol. The final elastomers were prepared by simply combining the end-functionalized PFPEs with a commercial trifunctional amine, giving fluoroelastomers that, over the conducted experiments timescale, behave like a normal rubber with a full elastic response at room temperature. At 18 ACS Paragon Plus Environment

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higher temperatures, however, the herein presented fluoroelastomers exhibit a fast stress relaxation without the use of additives or catalysts and can be recycled and reshaped multiple times thanks to the dynamic exchange reactions. In this context, the reprocessability of PFPE vitrimers was also verified. Thermal, mechanical and rheological properties of the materials were investigated and compared before and after five recycling steps. It could be shown that the created vitrimers recovered 95 % of their initial mechanical properties. Besides, similar analyses such as thermal stability, stress relaxation and infrared spectroscopy demonstrated that the PFPE vitrimers can be recycled multiple times without noticeable degradation. In addition to that, very low soluble fractions and swelling ratios were obtained in common organic solvents. A first general conclusion is that this paper describes a promising substitute for currently available permanently cross-linked fluorinated elastomers. With respect to the observed relaxation times, highly remarkable and unprecedented effects were noted and verified in several ways. The rheological behavior of the materials can be readily controlled by adjusting the amount of free primary amine moieties, as observed in previous systems. However, these catalyst-free elastomer networks obtained from two simple components, showed an unexpected and quite counterintuitive double relaxation behavior, implicating different chemical exchange reactions that are operating within the same simple polymer matrix. At low temperature, the exchange reactions are governed by the normal, previously established mechanism for amine exchange with an activation energy of 60 kJ.mol-1. On the other hand, at higher temperature, a sharp and clear change in temperature dependence was noted, with a fitted activation energy that is more than double that of the low temperature exchange (130 kJ.mol-1). In light of our in-depth investigation of the unusual dual rheological behavior of these materials, we found that the activation energies at low temperatures could also be readily determined from creep recovery experiments. This suggests that creep recovery experiments for vitrimer elastomers can be considered as a rapid alternative to get a better view of the (complete) rheological profile, which we now speculate may often be more complicated than anticipated based solely on stress-relaxation experiments, which are often limited to high, relatively small temperature intervals. Our work also points out an obvious danger in calculating the topology freezing transition temperature, Tv, since extrapolation from a high temperature zone to a lower temperature may miss the fact that the higher temperature process will no longer be the dominant process at the lower temperature. In short, we can conclude that, as vitrimer rheology is a physical process essentially governed by chemical reactivity, polymer chemists should be prepared and ready to find and take into account all the peculiarities and extreme context-dependence that typify the study of chemical reactivity, even within a very clear-cut and chemically uncomplicated situation. Acknowledgements The authors thank Prof. Dr. Ludwik Leibler for his valuable comments. M. G. and C. T. acknowledge the Research Foundation-Flanders (FWO) for the Postdoctoral and Ph. D. fellowships. The authors would like to thank Bernhard de Meyer and Bastiaan Dhanis for the technical support. F.D.P thanks BOF-UGent (GOA-funding).



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Fluorinated vitrimer elastomers with a dual temperature response

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