Structural and Mechanical Properties of ... - ACS Publications

Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

Structural and Mechanical Properties of Supramolecular Polyethylenes Jérémie Lacombe,† Samuel Pearson,‡ Franz Pirolt,§ Sébastien Norsic,‡ Franck D’Agosto,‡ Christophe Boisson,‡ and Corinne Soulié-Ziakovic*,† †

Laboratoire Matière Molle et Chimie, UMR 7167 CNRS-ESPCI Paris, Ecole Supérieure de Physique et de Chimie Industrielles de la Ville de Paris, PSL Research University, 10 rue Vauquelin 75005 Paris, France ‡ Laboratoire Chimie, Catalyse, Polymères et Procédés (C2P2), Equipe LCPP Bat 308F, Université de Lyon, Université Lyon 1, CPE Lyon, CNRS UMR 5265, 43 Boulevard du 11 Novembre 1918, F-69616 Villeurbanne, France § Institute for Chemistry and Technology of Materials (ICTM), Graz University of Technology, Stremayrgasse 9, 8010 Graz, Austria S Supporting Information *

ABSTRACT: Thymine (Thy) or 2,6-diamino-1,3,5-triazine (DAT) end-groups were efficiently installed on well-defined polyethylenes (PEs) synthesized by catalyzed chain growth (CCG) polymerization. Mono- and bifunctional low-molar mass PEs (1200−1500 g·mol−1) formed lamellar morphologies with long-range order upon cooling from the melt due to microphase segregation of polar supramolecular units and apolar PE chains. Crystallization of Thy functions into rigid planes at 180 °C induced very long-range lamellar order in Thy-functionalized PEs and dramatically suppressed PE crystallization (from 67% to 19%). DAT-functionalized PEs, whose end-groups do not crystallize, showed slightly reduced PE crystallinity (62%) and less long-range order, since assembly was instead driven by PE crystallization. Mechanical analysis of the bifunctional PEs demonstrated high moduli roughly proportional to PE crystallinity but low strains at break due to the absence of chain entanglements and/or tie chains between crystalline lamellae. This work offers important insights for designing supramolecular systems with tunable thermal and mechanical properties.

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INTRODUCTION Supramolecular chemistry allows the conception of exciting new materials, such as self-healing polymers1−5 and stimuliresponsive drug delivery systems.6−8 In supramolecular polymers, the combination of supramolecular interactions and phase separation can lead to mesoscopic organizations in the bulk.9−19 In particular, with crystalline polymers, long-range organizations have also been evidenced but with a strong perturbation of chain crystallinity.11,20−23 In a previous study, we hypothesized that low molar mass polyethylene (PE) bearing end groups capable of supramolecular interactions may behave like high density PE (HDPE) with high crystallinity at low temperatures, where end-group associations are strong, and as low molar mass PE at high temperatures, where supramolecular interactions are broken. Such reversible behavior would provide a significant processing advantage over traditional HDPE, which is highly viscous in the melt.24 We demonstrated that low molar mass well-defined polyethylenes obtained by catalyzed chain growth polymerization (CCG) and equipped at one extremity with thymine (Thy) or 2,6-diamino1,3,5-triazine (DAT) end groups were microphase-separated into lamellar structures.25 However, it was not possible to study these organizations in detail because of their low regularity induced by incomplete chain-end functionalization. In partic© XXXX American Chemical Society

ular, the driving force for the organizations with noncrystalline extremities and the parameters which control the interlamellar distances still have to be identified. In this paper, we study the structural and mechanical properties of highly functionalized Thy and DAT supramolecular PEs. Thanks to improved chain-end functionalization, very regular lamellar microphase separations were unambiguously evidenced with both crystalline (Thy) and non crystalline extremities (DAT or Thy/DAT). More precisely, we have determined the effect of chain-end crystallization, molar mass and functionality on the supramolecular organization and PE crystallization. In addition, the driving force for the very regular organizations obtained with noncrystalline chain-ends has been demonstrated. Finally, mechanical tests have been performed in order to determine the impact of the long-range mesoscopic organizations on the macroscopic properties of supramolecular PEs such as rigidity or flexibility. Received: February 5, 2018 Revised: March 20, 2018

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DOI: 10.1021/acs.macromol.8b00270 Macromolecules XXXX, XXX, XXX−XXX

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0.14° to 7°, applying the conversion q = 4π(sin θ)/λ with 2θ being the scattering angle with respect to the incident beam and λ the wavelength of the X-rays. Capillaries of 1.0 mm thickness were used and exposure was set to 60 s and repeated five times. A 15 min equilibration time was set at every temperature before starting the measurements. 3-Point Bending Tests. PE powders were compression molded under 3 T at 20 °C above their melting temperature in a parallelepipedic shape. Tests were performed on a DMA TA Q800 in 3-point bending. A strain ramp was applied at 0.05%·min−1 at 20 °C, and the stress was measured. Synthesis. The molar masses of the polymers were determined by the number-average polymerization degree (n) using NMR (see Supporting Information). A polyethylene (PE) with a polymerization degree of 43 or 54 is noted PE1200 or PE1500, respectively. Synthesis of H2N-PE1200-NH2. Telechelic amino-functionalized PE was obtained according to a previously published procedure.28 A functionality of 87% was determined by 1H NMR (details in Supporting Information). 1H NMR (2/1 v/v TCE/C6D6, 400 MHz, 363 K), δ: 2.67 (t, 4H), 2.33−2.40 (m, 8H), 1.47 (m, 4H), 1.24 (br m, 171H), 0.95 (br s, 4H). 13C NMR (101 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 37.34, 32.28, 30.22, 30.00, 29.56, 29.18. Molecular weight found by 1H NMR is consistent with SEC (Mn = 1129 g/mol, D̵ = 1.058). Synthesis of PE1500-NH2. Amino-functionalized PE was obtained with a functionality of 100% following a recently published procedure using a transfer agent bearing a disilazane protected amino function (details in Supporting Information).29 1H NMR (2/1 v/v TCE/C6D6, 400 MHz, 363 K), δ: 2.53 (t, 2H), 1.24 (br m, 217H), 0.83 (t, 3H). 13 C NMR (101 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 42.61, 34.40, 32.24, 30.00, 29.58, 27.24, 22.93, 14.02. Synthesis of Thymine-Activated Pentafluorophenol (Thy-PFP). the protocol of Egholm et al. was used.30 The product was obtained with a 76% yield. 1H NMR (DMSO-d6, 400 MHz, 363 K), δ: 11.55 (s, 1H), 7.64 (s, 1H), 4.99 (s, 2H), 1.77 (s, 3H). Synthesis of PE1500-Thy. In a 250 mL bicol flask fitted with a reflux condenser was dissolved 2.74 g (1.8 mmol) of PE1500-NH2 under argon in 150 mL of toluene at 100 °C. A solution of 1.27 g (3.6 mmol) of Thy-PFP in 10 mL of dried DMF was added to the flask. The reaction mixture was stirred at 100 °C for 3 h and then cooled and diluted with methanol (300 mL).The resulting suspension was filtered and the solid collected was then washed with methanol (3 × 100 mL) and dried to a colorless powder. Functionality was estimated from 1H NMR to 100% (details in Supporting Information). 1H NMR (400 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 6.75 (s, 1H), 6.47 (br s, 1H), 3.95 (s, 2H), 3.08 (q, 2H), 1.69 (s, 3H), 1.25 (br m, 199H), 0.83 (t, 3H). 13C NMR (101 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 166.27, 164.37, 152.03, 140.62, 110.78, 50.77, 40.00, 32.21, 30.00, 27.78, 22.90, 14.04, 12.02. Synthesis of PE1500-DAT. In a 1 L bicol flask fitted with a reflux condenser were dissolved 5.66 g (3.8 mmol) of PE1500-NH2 and 1.09 g (7.5 mmol) of DAT-Cl under argon in a mixture of 300 mL of toluene and 150 mL of DMF at 110 °C. 5.8 mL (33.3 mmol) of N,Ndiisopropylethylamine (DIEA) were then added and the solution was stirred at 110 °C for 64 h. The reaction mixture was then cooled and diluted with methanol (300 mL), and the resulting suspension was filtered. The colorless solid obtained was added to toluene (250 mL), and the suspension was heated to 90 °C and then hot filtered. The filtrate was cooled, diluted with methanol (300 mL), and filtered. The solid collected was then washed with methanol (3 × 100 mL) and dried to a colorless powder. Functionality was estimated from 1H NMR to 100% (details in the Supporting Information). 1H NMR (400 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 4.59 (br s, 1H), 4.36 (br s, 3H), 3.23 (q, 2H,), 1.25 (br m, 203H), 0.84 (t, 3H). 13C NMR (101 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 168.09, 167.64, 40.88, 32.21, 30.00, 27.26, 22.89, 14.03. Synthesis of Thy-PE1200-Thy. In a 250 mL bicol fitted with a reflux condenser was dissolved 1.00 g of H2N-PE1200-NH2 (0.81 mmol) in 55 mL of toluene under argon at 100 °C. A solution of 0.57 g (1.62 mmol) of Thy-PFP in 10 mL of dried DMF was added to the flask.

EXPERIMENTAL SECTION

Materials. All solvents and reactants were purchased from SigmaAldrich, TCI or Alfa Aesar. All chemicals were used as received unless otherwise stated. Tetrachloroethylene (Sigma-Aldrich) was purified by distillation and stored over 3 Å molecular sieves. Deuterated benzene (Eurisotop) was dried over 3 Å molecular sieves. Analytical Techniques. High-resolution liquid nuclear magnetic resonance (NMR) spectroscopy was carried out with a Bruker DRX 400 spectrometer operating at 400 MHz for the 1H nucleus and 101 MHz for 13C. Spectra were recorded at 363 K using a 5 mm QNP probe for 1 H NMR and a PSEX 10 mm probe for 13C NMR. Polymer samples were examined as 5−15% (w/v) solutions. A mixture of tetrachloroethylene (TCE) and deuterated benzene (C6D6) (2/1 v/v) was used as solvent. Chemical shift values (δ) are given in ppm with reference internally to tetramethylsilane for 1H NMR and to the methylenes of the PE chain at 30.00 ppm for the 13C NMR. Size exclusion chromatography (SEC) was performed using a Viscotek system (from Malvern Instruments) equipped with three columns (Waters Styragel HT6E 300 mm × 7 mm i.d. from Agilent Technologies). Portions (200 μL) of sample solutions with concentrations of 5 mg mL−1 were eluted in 1,2,4-trichlorobenzene using a flow rate of 1 mL min−1 at 150 °C. The mobile phase was stabilized with 2,6-di-tert-butyl-4-methylphenol (200 mg L−1). The OmniSEC software was used for data acquisition and data analysis. The molar mass distributions were calculated with a calibration curve on the basis of narrow polyethylene standards (338; 507; 1,010; 2,030; 17,000; 27,300; 43,400; 53,100; 65,700; 78,400 g mol−1) from Polymer Standards Service (Mainz, Germany). Fourier transform inf rared (FTIR) spectroscopy was carried out at different temperatures with a Bruker-Tensor 37 spectrometer at a resolution of 4 cm−1 equipped with a thermally controlled SPECAC Goldengate ATR accessory (32 scans, range 600−4000 cm−1). Dif ferential scanning calorimetry (DSC) experiments were performed under helium on a TA Q250 instrument. All samples were melted at 200 °C for 2 min to erase their thermal history, then cooled to 0 °C, and heated to 200 °C at a rate of 10 °C·min−1. The degree of crystallinity χc was calculated by the equation: χc =

ΔHm φ . ΔHm100%

Here, ΔHm and ΔHm100% are the measured melting enthalpy and the calculated melting enthalpy for 100% crystalline PE. φ is the weight fraction of PE in the sample, taking into account the DAT and Thy weight fraction. The enthalpy of 100% crystalline PE depends on the melting temperature of the PE chain according to the following formula:26

ΔHm100% = 93.22 + 4.249 × 10−3 × Tm 2 − 7.413.10−6 × Tm 3 The melting enthalpy can then be linked with the PE molar mass with27 Tm(x) = 414.3

x − 1.5 x + 5.0

x being the number of methylene units in the PE chain. Polarized optical microscopy (POM) was performed at different temperatures in transmission using a Leica Leitz DM RD light microscope and a heating stage Linkam Scientific LTS350. Small angle X-ray scattering (SAXS) utilized a high-flux SAXSess camera (Anton Paar, Graz, Austria) connected to a Debyeflex 3003 Xray generator (GE-Electric, Germany), operating at 40 kV and 50 mA with a sealed-tube Cu anode. The Goebel-mirror focused and Kratkyslit collimated X-ray beam was line shaped (17 mm horizontal dimension at the sample). The scattered radiation was measured in the transmission mode and recorded by a one-dimensional MYTHEN-1k microstrip solid-state detector (Dectris Switzerland), within a q-range (with q being the magnitude of the scattering vector) of 0.01 to 0.5 Å−1. Using Cu Kα radiation of wavelength 1.54 Å and a sample-todetector distance of 307 mm, this corresponds to a total 2θ region of B


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Macromolecules Scheme 1. Thy and DAT Functionalization of Amino-Terminated Mono- and Bifunctional PEs

The reaction mixture was stirred at 100 °C for 3 h and then cooled and diluted with methanol (300 mL). The resulting suspension was filtered, and the solid collected was then washed with methanol (3 × 100 mL) and dried to a colorless powder. Functionality was estimated from 1H NMR to 87% (details in the Supporting Information).1H NMR (400 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 9.78 (s, 2H), 7.34 (s, 2H), 6.85 (s, 2H), 4.09 (s, 4H), 3.27 (q, 4H), 2.50 (t, 4H), 2.39 (t, 4H), 1.71 (s, 6H), 1.50 (m, 4H), 1.24 (br m, 171H). 13C NMR (101 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm), δ: 166.69, 140.58, 110.22, 50.18, 39.83, 32.39, 32.09, 30.24, 30.00, 29.61, 29.27, 12.13. Synthesis of DAT-PE1200-DAT. In a 500 mL bicol fitted with a reflux condenser were dissolved 1.80 g of H2N-PE1200-NH2 (1.45 mmol) and 0.85 g of DAT-Cl (5.82 mmol) in a mixture of 150 mL of toluene and 75 mL of DMF under argon at 110 °C. Then 4.5 mL of DIEA (29.1 mmol) was added, and the solution was stirred at 110 °C for 64 h. The reaction mixture was then cooled and diluted with methanol (300 mL), and the resulting suspension was filtered. The colorless solid obtained was added to toluene (250 mL), and the suspension was heated to 90 °C and then hot filtered. The filtrate was cooled, diluted with methanol (300 mL), and filtered. The solid collected was then washed with methanol (3 × 100 mL) and dried to a colorless powder. Functionality was estimated from 1H NMR to 76% (details in Supporting Information). 1H RMN (2/1 v/v TCE/C6D6, 400 MHz, 363 K), δ: 5.02 (s, 2H), 4.46 (s, 8H), 3.36 (m, 4H), 2.51 (t, 4H), 2.37 (t, 4H), 1.45 (m, 4H), 1.24 (br s, 171H). 13C NMR (101 MHz, 363 K, TCE/C6D6 2/1 v/v, ppm) ,δ: 167.76, 167.06, 40.88, 32.67, 30.23, 30.00, 29.60, 29.22. Preparation of the Mixture PE1500-DAT/Thy-PE1500. An equimolar mixture was prepared by mixing 72 mg of PE1500-DAT and 75 mg of PE1500-Thy in 10 mL of boiling toluene. Toluene was slowly evaporated at 70 °C under a saturated atmosphere. The mixture was then dried overnight at 80 °C under vacuum.

agent25) which resulted in 80% functionalities, Thy grafting onto amino-terminated PEs was almost quantitative. Hence, PE1500-Thy and Thy-PE1200-Thy were obtained with functionalities of 100% and 87% respectively. 2,6-Diamino-1,3,5-triazine (DAT) was grafted according to a previously described protocol but with a longer reaction time to increase the conversion (Scheme 1, 64 h instead of 24 h).25 PE1500-DAT and DAT-PE1200-DAT were obtained with functionalities of 100% and 76% respectively. Long-Range Lamellar Organizations Driven by the Thy Crystallization. It has already been demonstrated that the crystallization of Thy units induces a long-range lamellar organization for telechelic poly(propylene oxide) (PPO) oligomers9 and monofunctional PE oligomers.25 In the latter system, Thy-bearing PE chains were confined between the thymine crystalline planes while a minority of unfunctional chains was phase separated. In this part, we analyze how the functionality, i.e. mono vs bifunctionality, impact Thy crystallization, PE crystallinity, and the long-range organization. Thermal Properties of PE1500-Thy and Thy-PE1200-Thy. Thermal analyses of PE1500-Thy and Thy-PE1200-Thy are presented in Figure 1. For both, an endotherm and an exotherm are observed at high temperatures (183 °C/180 °C for PE1500-Thy and 191 °C/184 °C for Thy-PE1200-Thy) with corresponding melting enthalpies of 24 J.g−1 and 32 J.g−1. These peaks are attributed to the melting and the crystallization of Thy units.25 At lower temperatures, the monofunctional PE shows an endotherm at 121 °C on first heating and at 111 °C on second heating, which correspond to the melting of PE chains. On cooling, the exotherm at 102 °C corresponds to the chain crystallization (Figure 1A). After the grafting of Thy units on PE chains, the functionalized polymer crystallized out of solution upon cooling, and it is this material which was analyzed in the first heating. In contrast, the second heating was preceded by crystallization of the sample from the melt. These different thermal histories are responsible for the PE melting temperature differing between the first and second heating. The first melting temperature is in agreement with the prediction of Broadhurst for PE chains of 1500 g·mol−1 (118 °C)27 while the second melting temperature is lower. The lower melting temperature can be explained by the confinement of PE chains through nanoscale phase segregation. By cooling from the melt, Thy units first crystallize into rigid planes, which confine and restrict the motion of molten PE chains.25 The confinement of polymer chains in nanometer-scale domains is known to decrease their crystallization (and corresponding melting) temperature.31,32 The crystallization of Thy prior to PE was verified by polarized optical microscopy (Supporting Information). Upon cooling from the melt, crystalline spherulites,

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RESULTS AND DISCUSSION Synthesis of Thy and DAT End-Functionalized Mono and Telechelic PEs. Quantitatively functionalized monofunctional amino-terminated PE (PE1500-NH2) was synthesized by catalyzed chain growth (CCG) polymerization of ethylene using a transfer agent bearing a disilazane-protected amino function.29 The telechelic amino-terminated PE, NH2-PE1200NH2, was synthesized following a published procedure to give a functionality of 87%.28 This last synthesis involved three reaction steps starting from a heterobifunctional vinyl-PE-I: first, elimination of HI to give vinyl-PE-vinyl, followed by a thiol−ene reaction using boc-cysteamine, and finally removal of the boc groups to yield H2N-PE-NH2. Thy was attached on PE chains using a pentafluorophenol ester of thymine-1-acetic acid. This activated ester allowed high yield and rapid grafting of Thy (3 h) onto amino-terminated polyethylenes without the need for catalyst or other reactants (Scheme 1). Compared with the reaction previously used to graft Thy (direct amidification using TBTU as coupling C


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Figure 1. DSC of (A) PE1500-Thy and (B) Thy-PE1200-Thy (exo down).

Figure 2. SAXS of (A) PE1500-Thy and (B) Thy-PE1200-Thy at different temperatures on cooling.

the melt, PE chains melt between 20 and 80 °C with an enthalpy of 43 J.g−1 which corresponds to a crystallinity of 19% (second heating), dramatically lower than for the initial H2NPE1200-NH2 (67%, Supporting Information). Chain-end mobility was restricted at both extremities by crystalline Thy planes, which drastically impeded the ability of the PE chains to align and crystallize. Structural Analysis of PE1500-Thy and Thy-PE1200-Thy. In order to study the impact of Thy crystallization on the longrange organization of Thy-functionalized PEs, SAXS experiments were performed on PE1500-Thy and Thy-PE1200-Thy at various temperatures while cooling from the melt (Figure 2). In the melt (≥200 °C), neither of the Thy- functionalized PEs were totally disordered. The SAXS plots present a broad peak at 1.27 nm−1 (49 Å) which correspond to the correlation hole peak commonly observed for block copolymers in the disordered state,33 and also for telechelic supramolecular PPO.34 This peak indicates that there is a typical distance between the high electron density Thy units and the PE molten chains, meaning that Thy units are somewhat clustered in the melt. By fitting this peak for Thy-PE1200-Thy with the Leibler’s theory on block copolymers,33 the radius of gyration of PE chains can be estimated as 21.6 Å and the Flory interaction

characterized by Maltese crosses under crossed polarizers, appeared at 170 °C with diameters up to 300μm. Below 100 °C, an additional birefringence resulting from the crystallization of PE chains appeared. Further support for Thy-induced confinement of PE chains being responsible for suppressing PE crystallization (and melting) temperature was found by comparing with the thermal behavior of the precursor polymer PE1500-NH2. PE1500-Thy exhibited a PE crystallinity of 46% based on the PE melting enthalpy, which is significantly lower than for PE1500-NH2 (67%, Supporting Information), which melts and crystallizes at 117 and 108 °C respectively (Supporting Information). The absence of crystallizable end groups in PE1500-NH2 gave a melting temperature very close to that predicted by Broadhurst even after cooling from the melt. Thus, the lowering of PE crystallinity and corresponding crystallization temperature when cooling PE1500-Thy from the melt can be unequivocally attributed to prior crystallization of the Thy units. The effect of Thy crystallization was even more marked for the telechelic Thy-PE1200-Thy (Figure 1B). For a sample crystallized from solution (first heating), a broad melting peak centered at 113 °C with an enthalpy of 114 J.g−1 corresponding to a crystallinity of 50% is observed. After crystallization from D


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Macromolecules parameter between PE and thymine χPE/Thy as 7.2 (Supporting Information). This last value is close to the one calculated using Fedor parameters25 (8.5) and reflects the high incompatibility, even at high temperatures, between Thy units and PE chains. For both Thy-functionalized PEs, at temperatures below the Thy crystallization (180 °C), SAXS shows clear evidence of lamellar structures characterized by diffraction orders positioned with integer qi/q0 ratios. Hence, the organization is driven by Thy crystallization. It is very long-range ordered, especially for PE1500-Thy which shows 9 orders of diffraction (Figure 2A). The characteristic distance of the lamellar organizations at 120 °C is 102 Å for PE1500-Thy and 57 Å for Thy-PE1200-Thy. At 40 °C, after PE chains have crystallized, the interlamellar distances have slightly decreased to 93 and 51 Å respectively. This probably results from the higher density of the crystalline phase of PE compared to its amorphous phase (1.00 versus 0.85 g·cm−3).35 We also note that the full width at half-maximum is smaller for molten chains (180 °C) than for crystalline chains (40 °C). Hence, the higher regularity of lamellae is favored by the increased mobility of PE molten chains. It is noteworthy that the first diffraction order q0 observed at both 180 and 120 °C is missing from the SAXS plot of PE1500-Thy at 40 °C (Figure 2A). Indeed, in the 40 °C plot, the first peak position (1.35 nm−1) and the second peak position (2.02 nm−1) differ by a factor of exactly 1.50, which means that they correspond respectively to the second and third-order diffraction peaks of a lamellar structure with an interlamellar distance of 93 Å. The absence of the q0 peak, expected in this case at 1.35/2 = 0.67 nm−1, has already been observed in the literature for certain triblock copolymers in a certain temperature range.36,37 Gobius du Sart et al. explained it by a peculiar configuration in which the electron densities and the sizes of the different blocks of the lamellae (Thy and PE in our case) lead to a null diffraction intensity for q0.36 SAXS data of PE1500-Thy at 180 °C were fitted with a lamellar structure having an interlamellar thickness L with a standard deviation σD, and crystalline planes of thickness lc with a standard deviation σc (using notations of Sun38). A good fit was obtained with L = 100 Å, σD = 3.5 Å, lc = 7.0 Å and σc = 1.0 Å (Supporting Information). These results highlight the very low standard deviations of the interlamellar and crystalline thicknesses, meaning that the organization is very regular. The crystalline plane thickness suggests that the Thy units are not aligned perpendicular to the lamellar planes but are more likely tilted, since two associated Thys exhibit a length of 11 Å.25 Such a tilted conformation has already been demonstrated for alkyl-Thy derivatives with an alkyl,39 ester,40 or urethane link.41 Figure 3 summarizes the nanostructure of PE-Thy and Thy-PEThy in the melt and at room temperature. Interlamellar Distance Controlled by the Thy Volume Fraction. After characterizing the lamellar dimensions in PE1500-Thy, we were interested in whether the interlamellar distance could be predicted by some characteristic(s) of the polymer, and whether such a model could be generalized to other polymers. When the lamellar structure is formed (at 180 °C on cooling), PE chains are molten so it is not relevant to compare the size of extended chains with the experimental interlamellar distances. Instead, since the lamellar order is driven by Thy crystallization, we propose that the volume fraction of Thy is the main parameter controlling the interlamellar distance. Thy volume fraction f Thy corresponds to the ratio of Thy volume VThy to the total volume (Thy + PE) VTot:

Figure 3. Scheme of (A) PE1500-Thy and (B) Thy-PE1200-Thy in the melt and at 40 °C.

fThy =

VThy VTot

Lamellae are considered as infinite in the y and z directions (i.e., perpendicular to the PE/Thy interfaces), giving: L=

d Thy fThy

with L the interlamellar distance and dThy the Thy plane thickness. To test the validity of this hypothesis, the interlamellar spacings of Thy-functionalized PEs are plotted as a function of Thy volume fraction in Figure 4. To supplement the data set, experimental lamellar distances of alkyl-Thy C12H25-Thy and C18H37-Thy have been added (syntheses and characterizations in Supporting Information). The model fits well with the experimental results of mono and bifunctional Thy-PEs and monofunctional alkyl-Thys when a Thy plane thickness of 7 Å is imposed (Figure 4). In addition, this model also fits well with the interlamellar distances previously reported for Thyfunctionalized PPO.9 Hence, based on the volume fraction of the crystallizing end-group, this model is suitable to predict the lamellar spacing of any amorphous or molten chains, whatever their chemical nature. Moreover, it does not depend on chain functionality since it fits both mono and bifunctional chains. Indeed, because they have the same volume fraction of Thy, a bifunctional chain of molar mass 2 M will have the same interlamellar distance as a monofunctional chain of molar mass M. For Thy volume fractions below 0.05, the lamellar distances diverge, suggesting a loss of the long-range ordered organization. As shown with Thy-functionalized PPO chains, if Thy volume fraction is too small, the Thy crystalline planes E


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increase. Finally, the experimental interlamellar distances are ruled by the Thy volume fraction rather than the size of PE extended chains. Long-Range Lamellar Organizations Driven by PE Crystallization. It has been previously shown that the longrange lamellar organization of telechelic PPO oligomers was totally suppressed by replacing crystallizable Thy units by noncrystallizable DAT units.9 In contrast, a lamellar structure has been observed for monofunctional PE-DAT oligomers, albeit no long-range order.25 Here, we analyze in more detail the organization of PE chains with non-crystallizable extremities (DAT or Thy/DAT) displaying very high functionalities. Thermal Properties of PE1500-DAT and DAT-PE1200-DAT. Thermal analyses of PE1500-DAT and DAT-PE1200-DAT are presented in Figure 5. The DSC thermogram of PE1500-DAT exhibits an endotherm and an exotherm corresponding to the melting and crystallization of PE chains. The melting temperature is 119 °C, in good agreement with the values predicted by Broadhurst for a PE chain of 1500 g·mol−1.27 Hence, even though two chains may associate with each other by their DAT units, PE chains do not melt as a chain of double (or more) molar mass, since for 3000 g·mol−1, Tm is expected to be 129 °C.27 After cooling from the melt, the second heating of DAT-PE1200-DAT reveals a bimodal melting at 113 and 121 °C. These are most likely associated with DAT-grafted PEs (113 °C) and non grafted PEs (121 °C), given that the DAT functionality was established on 85% of the chains (experimental section). PE crystallinities are 59% for PE1500-DAT and 62% for DAT-PE1200-DAT, slightly lower than those of the parent amino PEs (both 67%, Supporting Information), indicating that grafting DAT disturbs PE chain crystallization less than grafting Thy (46% and 19% for mono and bifunctional derivatives, respectively). This can be explained by the fact that DAT units are not crystallizable.25,34 This is confirmed by polarized optical microscopy: above the melting temperature of PE, samples are fully isotropic (Supporting Information). Finally, the high PE crystallinities with DAT evidence that DAT units are not included into the crystalline lamellae contrary to other systems in the literature.19,43 Structural Analysis of PE1500-DAT and DAT-PE1200-DAT. In order to study the structures of DAT-functionalized PEs, SAXS experiments on cooling from the melt were performed on PE1500-DAT and DAT-PE1200-DAT (Figure 6). In the melt

Figure 4. Interlamellar distance with the volume fraction of Thy for mono and bifunctional Thy-PEs and Thy-PPO.9 Data corresponding to C12-Thy and C18-Thy have also been added. The dotted line corresponds to the proposed model with Thy crystalline planes of 7 Å thickness.

required to obtain the mesoscopic order can no longer be formed.42 In contrast, PE1500-NH2 shows no organization at room temperature even though a very weak peak at 117 Å in the SAXS analysis emerges from the background at 110 °C just after PE crystallization (Supporting Information). H2N-PE1200NH2 presents a very weakly ordered lamellar structure (only two diffraction orders) with a characteristic distance of 84 Å at 40 °C (Supporting Information). The terminal amino functions probably interact with each other through H-bonds to induce this weakly regular organization. To sum up, the crystallization of Thy drives the long-range lamellar organization of Thy functionalized PEs. The PE chains are then confined between Thy crystalline planes, hindering their crystallization as shown by lower PE crystallization/ melting temperatures and crystallinities. This perturbation is even more significant when the motions of PE chains are restricted on both extremities as in the telechelic Thy-PE1200Thy. Yet, with longer PE chains, i.e., a lower volume fraction of Thy, the confinement should decrease and the PE crystallinity

Figure 5. DSC of (A) PE1500-DAT and (B) DAT-PE1200-DAT (exo down). F


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Figure 6. SAXS of (A) PE1500‑DAT and (B) DAT-PE1200-DAT at different temperatures on cooling from the melt. Modeling of the electron density profile across the lamellae for (C) PE1500-DAT and (D) DAT-PE1200-DAT at room temperature.

(140 °C), as already observed for Thy homologues (Figure 2), both DAT functionalized PEs present a strong correlation hole peak around 1 nm−1. By fitting the peak of DAT-PE1200-DAT (q = 1.27 nm−1, 49 Å) with Leibler’s model, the radius of gyration of PE molten chains is estimated as 20.5 Å and the Flory interaction parameter between PE and DAT χPE/DAT as 7.1, which is similar to χPE/Thy estimated for Thy-PE1200-Thy. Hence, the incompatibilities of Thy with PE and DAT with PE seem comparable. At 100 °C, when PE chains have crystallized, SAXS patterns reveal lamellar organizations with a characteristic distance of 215 Å for PE1500-DAT and 85 Å for DAT-PE1200-DAT. It is possible to calculate the molecular length d of an extended PE chain functionalized by DAT tilted by an angle φ from the normal layer:25 d = dPE,c + dPE,a + dDAT = nc

DAT is approximately half the experimentally observed value, meaning the PE chains are organized as a bilayer within the interlamellar space (Figure 6C). The calculation of d assuming a bilayer gives a value of 206 Å, in good agreement with the experimental result. It should be noted that diffraction peaks of PE1500-DAT (0.29, 0.88, and 1.47 nm−1, Figure 6A) are indexed as q0, 3q0, and 5q0. The lack of even-order diffraction peaks results from a particular symmetry of the electron density across the lamellae, courtesy of the PE chains in PE1500-DAT forming a bilayer. Defining the axis perpendicular to the lamellae as the x-axis, extinction of even-order diffraction peaks has been shown to occur if the electron density is an odd function relative to the positions x = L/4 and x = 3L/4.46 For this condition to be met in PE1500-DAT, the reduced electron density within the CH3/CH3 thickness at x = L/2 (i.e., the PE bilayer junction) would have to be exactly compensated by the higher electron density within the DAT/DAT thicknesses at x = 0 and x = L. Indeed, the electron densities of DAT-NH (molar volume 91.3 cm 3 ·mol −1 and 65 electrons),35 CH3 (molar volume of 33.5 cm3·mol−1 and nine electrons)47 and PE chain with 59% crystallinity were estimated as 0.43, 0.16, and 0.32 e·Å−3 respectively. These calculated values were used to graphically describe the electron density profile across the lamellae as shown in Figure 6C, illustrating that the special symmetry condition required for even-order extinction is indeed met. The absence of even-order diffraction peaks therefore also confirms the bilayer organization for PEDAT. On the other hand, the electron density profile across the lamellae of bifunctional DAT-PE1200-DAT displays only regions

c PE cos φ + nalPE,a + dDAT 2

where nc and na are the number of crystalline and amorphous carbons in the chain, cPE is the c-parameter of the orthorhombic unit cell of PE, lPE,a is the thickness per carbon in the amorphous PE layer, and dDAT is the length of a DAT unit. Taking into account the PE crystallinity, and applying a tilted angle of 35° (which is classical for PE44,45), a DAT length of 5.5 Å,25 cPE = 2.55 Å, and lPE,a = 0.712 Å, the model gives a theoretical lamellae thickness of 103 Å for PE1500-DAT and 90 Å for DAT-PE1200-DAT. This estimate for bifunctional DAT is very close to the 85 Å obtained by SAXS, meaning that the PE chains in DAT-PE1200-DAT are organized as a monolayer as depicted in Figure 6D. In contrast, the theoretical d for PE1500G


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Macromolecules

Thermal and Structural Analysis of PE1500-DAT/Thy-PE1500. An equimolar mixture of PE1500-Thy and PE1500-DAT was prepared in hot toluene (experimental section). The thermal analysis reveals a similar behavior to PE1500-DAT, displaying PE melting and crystallization temperatures at 119 and 112 °C, respectively (Supporting Information). Since Thy/DAT pairs are not crystalline,34 the absence of Thy crystallization suggests the association of Thy units with DAT, which is in accordance with the much higher association constant of Thy/DAT compared to Thy/Thy and DAT/DAT.49 The PE crystallinity of 56% in the mixture is also very similar to PE1500-DAT, evidencing that PE crystallization is not much altered. Just after PE crystallization, SAXS analysis shows the appearance of a lamellar morphology with a characteristic distance of 224 Å (Supporting Information). As with PE1500DAT, we observe extinction of even diffraction peaks proving that a bilayer morphology is formed. Hence, even though the association constant is much higher for Thy/DAT than for DAT/DAT (835 M−1 vs 4 M−1),49 the same nanostructure is formed. We can conclude that supramolecular PEs with noncrystallizable end groups also form highly ordered lamellar organizations driven by PE crystallization which confines the phase-segregated supramolecular units into well-defined planes. Compared with Thy-functionalized PE, DAT-PEs and PE-Thy/ DAT-PE mixtures have a much higher crystallinities but a less ordered organization. This results from the fact that the PE, although well-defined, is not monodisperse, and that its interface with DAT (or Thy/DAT) is far less rigid than that with Thy crystalline planes. Mechanical Properties of Long-Range Ordered Supramolecular Polyethylenes. One of the main interests of telechelic supramolecular polymers is to attain polymers with large apparent molar masses and consequent mechanical properties thanks to supramolecular interactions, yet presenting low molar mass polymer properties, i.e., low viscosity at high temperatures and/or solubility when supramolecular bonds are broken. The Thy and DAT telechelic PEs are indeed soluble at rather low temperature in toluene (80 °C). We study in this section the impact of the mesoscopic organization of supramolecular PEs on their mechanical properties at room temperature. A nonfunctional PE of 1200 g·mol−1, PE1200, is used as a reference. Samples were prepared by compression molding and analyzed by three-point bending at 20 °C (Figure 8). All PEs are brittle solids since they break in the linear regime. Reference PE1200 has a flexural modulus of 2.5 GPa resulting from the high crystallinity of the material (71%). However, its strain at break is only 0.30% (σr = 7.4 MPa). This result can be explained by its molar mass being lower than the entanglement molar mass for polyethylene (Me = 1850 g· mol−1).50 In addition, there are no, or not enough, tie chains to link crystalline lamellae as in traditional high-density polyethylene. Indeed, Brown and Ward have shown that the higher the molar mass, the higher the number of tie molecules.51 DAT-PE1200-DAT properties are close to those of PE1200. The flexural modulus is slightly increased to 3.3 GPa and the strain at break is slightly decreased to 0.24% (σr = 7.9 MPa), indicating that its lamellar organization scarcely impacts the mechanical properties of PE. In contrast, the flexural modulus of Thy-PE1200-Thy falls to 0.40 GPa while its strain at break increases to 1.9% (σr = 7.3 MPa), meaning this material is slightly more flexible. The low crystallinity of PE chains in ThyPE1200-Thy (19%) explains both the lower modulus (less rigid)

of higher electron density (relative to the PE chains) at each extremity of the lamellae (Figure 6D), meaning that both odd and even diffraction peaks are observed in the SAXS analysis. The driving force for the long-range organization of DATfunctionalized PE is not the crystallization of the end-groups as observed in the Thy-functionalized PEs. To further probe the organization process, infrared spectra of PE1500-DAT at different temperatures on cooling were performed (Figure 7),

Figure 7. FT-IR of PE1500-DAT at different temperatures on cooling from 200 °C.

and compared to the complementary SAXS data. At 200 °C, a band at 1569 cm−1 corresponding to the quadrant stretching vibrations of the triazine ring (predominantly CN stretching) is observed.48 On cooling to 160 °C, a band at 1552 cm−1 emerges, corresponding to shorter CN bonds resulting from H-bond interactions with amino groups of another DAT unit. The intensity of this band steadily increases when moving to 140 and 126 °C (just above the PE crystallization temperature of 124 °C), suggesting the existence of DAT clusters in the melt. Further intensification of this band at 124 and 120 °C, i.e., during and after PE crystallization, indicates that the DAT units become further confined and exhibit stronger interactions among themselves. The fact that the PE organizes into lamellae upon crystallization (as shown in the SAXS analyses) suggests that DAT units more likely form DAT/DAT pairs rather than large clusters. PE crystallization brings them closer to each other and drives the long-range ordered bilayer. A similar result was obtained by Ungar et al. with a PE chain of 2680 g·mol−1 functionalized with an acid group at one end.46 This behavior is in contrast to the Thy-functionalized PEs, where it was crystallization of the Thy end groups, rather than crystallization of the PE chains, which drove the long-range order. The fact that long-range ordering is attainable in the DATfunctionalized PEs is also probably due to the narrow molar mass distributions of the precursor PEs obtained by CCG polymerization. Finally, just after PE crystallization, the SAXS analysis of PE1500-DAT shows that weak even-order diffraction peaks do appear (120 °C, Figure 6A). This might be explained by the formation of lamellae comprising of PE monolayers with a characteristic distance of 114 Å. Their almost complete absence at 80 °C (and at 30 °C) suggests a reorganization of these chains to a bilayer structure during the cooling. H


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Macromolecules

As a perspective, we could use the concept of supramolecular block copolymers to improve the mechanical properties of lowmolar mass PEs. By mixing a supramolecular crystalline PE with a complementary telechelic supramolecular polymer whose chains are incompatible and amorphous, we might obtain a microphase-separated material. We might expect formation of nanodomains separated through supramolecular complementary interactions (the Thy/DAT interaction is stronger and more selective than the Thy/Thy and DAT/DAT interactions).49 Since the Thy/DAT pair is not crystalline, the crystallinity of PE domains should be preserved to confer toughness to the materials, while amorphous domains should confer flexibility. This work will be reported separately in a future paper.

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ASSOCIATED CONTENT

S Supporting Information *

Figure 8. Stress−strain curve by 3-point bending for PE1200, DATPE1200-DAT, and Thy-PE1200-Thy.

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b00270. Materials and characterization data for all compounds, fitting SAXS data of PE1500-Thy with a model of lamellar structure, fitting correlation hole peaks for Thy-PE1200Thy and DAT-PE1200-DAT, experimental procedures, thermal and structural characterization of C12−Thy and C18−Thy, and thermal and structural characterization of PE1500-Thy/DAT-PE1500 (PDF)

and the increase in strain at break (larger amorphous domains). This is not observed with DAT-PE1200-DAT for which DATDAT layers do not perturb the PE crystallinity (62%). To sum up, mechanical properties of long-range organized PEs are not very different from nonfunctional PE. Because of the phase separation of polar end-groups, PE chains do not behave like high molar mass analogues in bulk. In addition, with a crystallizable end-group, the flexibility increases due to the lower PE crystallinity, but the modulus strongly decreases at the same time.

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AUTHOR INFORMATION

Corresponding Author

*(C.S.-Z.) E-mail: [email protected]. Telephone: +33 1 40 79 46 30.

CONCLUSION The excellent control over the molar mass distribution in the CCG process and the high chain end fidelity facilitated by improved postpolymerization functionalizations allowed the self-organization of narrowly distributed, highly end-functional mono- and bifunctional PEs to be studied. We observed the formation of lamellar structures with Thy, DAT, and Thy/DAT end groups. With Thy/Thy systems, the driving force for the organization is the end-group crystallization while with noncrystallizable pairs, DAT/DAT and Thy/DAT, it is the chain crystallization. Simple models to describe the interlamellar distances with Thy and DAT units supported these mechanisms. We also revealed that low-molar mass end-functional supramolecular PEs do not behave as high molar mass chains in the bulk. In fact, due to their mesoscopic organizations, oligomeric PE chains remain confined within their own lamellae, with the end groups acting more as boundaries rather than bonds between PE chains. To improve mechanical properties, supramolecular PEs should have longer chains, at least 2Me (3700 g·mol−1), and noncrystallizable supramolecular units such as DAT or Thy/ DAT. As we observed for PE and PPO42 supramolecular polymers, the longer the polymer chain, i.e. the lower the supramolecular unit fraction, the less the material is phase separated (if at all). The loss of organization could be beneficial to the mechanical properties but long chains would also decrease the probability of end-groups associations. Adding small amounts of multifunctional low molar mass supramolecular PEs could also provide cross-links improving the toughness of the material while maintaining mesoscopic organizations as already observed with self-healing materials.52

ORCID

Franck D’Agosto: 0000-0003-2730-869X Corinne Soulié-Ziakovic: 0000-0001-6949-3630 Author Contributions

The manuscript was written through contributions of all authors. Funding

The authors received funding from the French National Agency for Research (ANR SUPRA PE 13-BS08−0006). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Prof. Ludwik Leibler and Prof. Otto Glatter for very helpful discussions. We are also grateful to Dr. Winnie Nzahou for the synthesis of PE-1500-NH2. Finally, we thank the NMR Polymer Center of Institut de Chimie de Lyon (FR5223) for assistance and access to the NMR facilities.

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REFERENCES

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