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Table of Contents
Abstract
1 | INTRODUCTION
2 | EXPERIMENTAL
2.1 | Materials
2.2 | Nanocomposites&Apos; Preparation
2.3 | Characterization
2.3.1 | NMR Spectroscopy
2.3.2 | Small-Angle X-Ray Scattering
2.3.3 | Light Microscopy
2.3.4 | Transmission Electron Microscopy—Nanofiller Dispersion
2.3.5 | Thermogravimetric Analysis
2.3.6 | Infrared Spectroscopy
2.3.7 | Thermo-Mechanical Properties Of Elastomers
2.3.8 | Tensile Tests
3 | RESULTS AND DISCUSSION
3.1 | Synthesis Of The Nanocomposites
3.1.1 | Completion Of Monomer Conversion And The Degree Of Nanofiller Incorporation
3.1.2 | Length (Molecular Mass) Of The Elastic Chains
3.1.3 | Nanofiller Distribution
3.2 | Tuning Of The Chain Length Of The PolyMEA Matrix
3.3 | Thermo-Mechanical Properties Of Elastomers
3.4 | Tensile Properties
3.4.1 | High Reversibility Of The Stretching Deformation
3.4.2 | Effect Of Clay Content On The Tensile Properties
3.4.3 | Effect Of PolyMEA Chain Length On Tensile Properties
3.5 | Thermal Stability Of The Nanocomposite Elastomers: TGA
4 | CONCLUSIONS
ACKNOWLEDGMENTS
Abstract
Funding information Grantová Agentura České Republiky, Grant/Award Number: 19-04925S Physically cross-linked solvent-free supramolecularly assembled nanocomposite elastomers were prepared, which displayed exceptionally high extensibility (up to 6000%), besides generally high mechanical properties (G' in rubber region between 1.5 and 40 MPa). The nanocomposites are based on linear poly(2methoxyethylacrylate) (polyMEA) and montmorillonite clay (physical crosslinker), and were obtained via free-radical polymerization of the monomer in the presence of the exfoliated nanofiller dispersed in water. The mechanical properties of the studied products were varied in a very wide range by changing the concentrations of the radical redox co-initiator pair, at given constant nanofiller loadings. The latter in turn also strongly altered the product properties. This applied synthesis approach, aimed at obtaining longer elastic chains, made possible to increase the elongation at break up to six times, and also to tremendously and simultaneously increase the toughness (effect of entanglements), as well as to shift the tensile curves between “plastic-like elastomer” and “simple elastomer.” In all cases, however, the nanocomposite samples displayed a highly efficient recovery, even after very high deformations. The structure–property relationships were deeper elucidated by thermo-mechanical analysis (DMTA), TGA (thermal stability, elastic chains' immobilization), TEM and X-ray diffraction.
Institute of Macromolecular Chemistry, Czech Academy of Sciences, Praha, Czech Republic Correspondence Beata Strachota, Institute of Macromolecular Chemistry, Czech Academy of Sciences, Heyrovskeho nam. 2, CZ-162 06 Praha, Czech Republic. Email: beata@imc.cas.cz Funding information Grantová Agentura České Republiky, Grant/Award Number: 19-04925S Physically cross-linked solvent-free supramolecularly assembled nanocomposite elastomers were prepared, which displayed exceptionally high extensibility (up to 6000%), besides generally high mechanical properties (G' in rubber region between 1.5 and 40 MPa). The nanocomposites are based on linear poly(2- methoxyethylacrylate) (polyMEA) and montmorillonite clay (physical cross- linker), and were obtained via free-radical polymerization of the monomer in the presence of the exfoliated nanofiller dispersed in water. The mechanical properties of the studied products were varied in a very wide range by changing the concentrations of the radical redox co-initiator pair, at given constant nanofiller loadings. The latter in turn also strongly altered the product proper- ties. This applied synthesis approach, aimed at obtaining longer elastic chains, made possible to increase the elongation at break up to six times, and also to tre- mendously and simultaneously increase the toughness (effect of entanglements), as well as to shift the tensile curves between “plastic-like elastomer” and “simple elastomer.” In all cases, however, the nanocomposite samples displayed a highly efficient recovery, even after very high deformations. The structure–property relationships were deeper elucidated by thermo-mechanical analysis (DMTA), TGA (thermal stability, elastic chains' immobilization), TEM and X-ray diffraction. KEYWORD S elastomers, nanoparticles, nanowires, and nanocrystals, nanostructured polymers, supramolecular structures, mechanical properties
1 | INTRODUCTION
This work is dedicated to the synthesis and to property tuning of solvent-free nanocomposite elastomers (xerogels), by which an exceptionally high elongation at break was achieved in combination with a markedly increased toughness. Traditional commercial elastomers like natural or industrial rubber display elongations at break in the range of typically 100 to 600%, exceptionally extensible ones up to 1000%.1 In the recent years, however, several types of advanced ultra-stretchable elastomers were developed, which are of interest for soft robotics, or for biomedical engineering. Their functionalized derivatives are of interest as actuators, J Appl Polym Sci. 2020;e49836. wileyonlinelibrary.com/journal/app © 2020 Wiley Periodicals LLC 1 of 24 https://doi.org/10.1002/app.49836 sensors or sensor skins, as well as wearable electronics parts. These ultra-stretchable materials include both solvent-free elastomers, as well as gels—usually hydrogels. Highly extensible (hydro)gels are the most studied sub-group of the ultra-stretchable elastomers. In hydrogels the elastic chains are uncoiled and hence partly stretched due to the swelling, which should reduce the elongation at break, if compared to fully analogous nonswollen structures. Nevertheless, the ultra-extensible hydrogels were described earlier and more often in the literature, and achieved higher elongations at break than the solvent-free ultra-extensible elastomers. The ultraextensibility of these special gels is based on their relatively sophisticated structures, especially in case of materials which possess moduli in the MPa range (see2), in contrast to the structurally simple “classical” rubbery polymer networks. The advanced ultra-stretchable hydrogels always contain nano-objects of different shape and composition, form which numerous very long polymer chains extend, thus forming brush-like units. These latter chains are connected via entanglements and by some permanent cross-links (covalent bonds, strong entanglements, strong adsorption of chain segments), thus forming the ultra-extensible material. Similar structural patterns are also found in the most advanced solvent-free ultra-extensible elastomers discussed further below. As elastomers, however, the ultra-extensible hydrogels have the disadvantage of containing solvent, the evaporation of which changes their properties, up to the final stage of an often glassy material. The ultrastretchable hydrogel structures include linear poly(Nisopropyl acrylamide) (“PNIPAm”) cross-linked by PNIPAm nano-gel particles3 (maximum elongations εbr around 1400%), or polyacrylamide hydrogels cross-linked by a small amount of an 8-armed star-like macro-crosslinker (εbr = 15,000% – a record value; fairly reversible deformation up to 4000%).4 Most of the highly extensible hydrogels are cross-linked physically, by inorganic nanofillers, while the above-discussed examples can be regarded as organic–organic nanocomposites. Inorganic nanofillers can greatly improve the mechanical properties of a given polymer matrix, also in the swollen state, via specific interface interactions, for which their high specific surface is very advantageous.5 If all nanofiller dimensions are sufficiently small, optical transparency additionally can be preserved,6,7 while specific chemical,8–17 optical,18,19 electrical,20,21 magnetic,22,23 or gas barrier24–26 properties can be lent to the matrix. As a nanofiller phase introduces heterogeneity, albeit only on the nanometre scale, it seldom improves the extensibility of simple materials. However, the discussed ultrastretchable hydrogels are notable exceptions, due to their complex structures, as well as due to small and finely dispersed nanofiller. Such systems include polyethylene glycol/hydroxyapatite nanocomposite gels (εbr = up to 2000%),27 or polyacrylamide/Ca(OH)2 nano-spherulite gels (εbr = 6000–12,000%).7 Related to the latter are environmentally relatively stable gels, in which the swelling fluid is ethylene glycol instead of water (εbr = 4500%). 28 2D-nanofillers also were employed to prepare ultraextensible hydrogels, based e.g. on polyacrylamide/ graphene oxide (εbr = 3500% 29), on polyacrylamide/layered double hydroxide (εbr = 6000%), 30 or on PNIPAm/ clay (this was historically the first discovered ultraextensible hydrogel with high mechanical properties31: εbr = up to 1400%). Related poly(dimethylacrylamide)/ clay gels achieved εbr = 2000–2500% (and high toughness).2 Via modified synthesis of the latter nanocomposite gel, record εbr values of up to 12,000% were achieved, in combination with attractive recovery behaviour.32 Self-healing is a very attractive property discovered in the polyacrylamide/clay hydrogels,33,32 and is possible due to the physical nature of the crosslinking33–36 in them. Highly extensible solvent-free and environmentally stable elastomers were more difficult to obtain, and are hence much less studied in literature, than ultra-extensible hydrogels. Two different groups of these materials can be differentiated: super-soft, as well as rubbery products, the latter with moduli in the MPa range. Super-soft elastomers (modulus: 5000 down to 500 Pa, similar like hydrogels) are based on long linear bottlebrush structures crosslinked either by entanglements or by covalent bonds (εbr = 1000%), and can be of interest for soft implants, ultra-compressible sensors, or in soft robotics.37 In such elastomers, the highly flexible pendant chains take the role which solvent molecules play in hydrogels. The first stiffer, “rubber-like” super-extensible elastomers38 were based on a solvent-free system structurally similar to the above-discussed polyacrylamide/clay hydrogels, but with hydrophobic and flexible (glass transition near −20 C) poly(2-methoxyethyl acrylate) (polyMEA) elastic chains. In the polyMEA/clay nanocomposites, the solvent present in its hydrogel analogues was replaced by a higher amount of polyMEA chains. Achieved were εbr values between 1000 and 3000%, in combination with strengths from 6 down to 1.2 MPa, respectively, decreasing with εbr. Their potential applications include substrates for cell harvesting, regenerative medicine, or tissue engineering.39–41 In the present work, the polyMEA/clay nanocomposite elastomers were considerably further developed. A different rare example of “rubber-like” super-extensible elastomers are the highly extensible polystyrene-poly(butyl acrylate)-polystyrene triblock copolymers42 (εbr = 2000%, strength up to 3.6 MPa). Selfassembly (nano-phase-separation) leads to a morphology somewhat similar to all the above-discussed nanocomposites: The polystyrene blocks undergo aggregation, thus forming small hard nano-domains acting as strong physical cross-links. The soft poly(butyl acrylate) block is a highly flexible chain, acting as a spring which connects the hard domains of polystyrene. Importantly, the flexible poly(butyl acrylate) chains were “doped” by some amount of statistically dispersed acrylamide repeat units, the hydrogen bridging of which formed dynamic (soft) cross-links between the chains, thus greatly increasing the toughness of the whole elastomer. In their previous work, the authors studied stimuliresponsive nanocomposite PNIPAm hydrogels reinforced with in-situ-nano-silica.43–46,21 This filler considerably improved the mechanical properties and the rate of temperature-response. A similar system with nano-TiO2 also was developed.47 More recently, the authors turned their attention to poly(N-isopropylacrylamide)/clay hydrogels (Haraguchi gels),48–51 including a comprehensive study of the formation process. They found out, that variation of the concentrations and ratios of the radical redox co-initiators, used during the gels' synthesis, made possible to tune the material properties of the hydrogels in a very wide range. The principle underlying the changes in properties was the achievement of very different chain lengths in the polymer matrix of these nanocomposites (molecular masses between 0.5 and 5 MDa). In this work, the aim was to achieve record values of ultra-extensibility in solvent-free elastomers, simultaneously with high moduli and high toughness, namely in the polyMEA/clay nanocomposite, while applying the authors' previous experience with the successful synthesis of very long elastic chains in polyacrylamide/clay gels. Application possibilities include advanced structural materials with self-healing potential, soft robotics, implant technology or regenerative medicine.
2 | EXPERIMENTAL
2.1 | Materials
2-Methoxyethyl acrylate (MEA) was obtained from Tokyo Chemical Industry Co., Ltd., ammonium persulfate (APS), and N,N,N0,N0-tetramethylethylenediamine (TEMED) were purchased from Sigma-Aldrich and used as received without further purification. The synthetic hectorite clay, “Laponite RDS” (chemical composition: Na0.7[(Si8Mg5.5Li0.3)O20(OH)4]) was friendly donated by BYK Additives & Instruments.
2.2 | Nanocomposites' preparation
PolyMEA/clay nanocomposite elastomers were prepared by in situ free-radical polymerization of MEA in water, in the presence of exfoliated clay (Laponite RDS) platelets. The homogenous aqueous dispersion of the RDS clay was obtained by stirring it in water for 24 h. Afterwards, the monomer MEA was added to the dispersion and the mixture was purged with argon. Next, the redox co-initiator TEMED was added and further stirring for ca. 1 min followed. Subsequently, the second co-initiator APS (as a 1% aqueous solution) was admixed, which started the polymerization reaction. After a brief final stirring, the reaction mixture was transferred into an argon-filled mould (internal dimensions: 100 x 50 x 5 mm3), which consisted of two glass sheets enclosing a rubber spacer. The reaction was left to run at 25 C for 24 h. The resulting hydrated product (opaque white hydrogel) was washed with pure water and subsequently dried at room temperature for 24 h, and finally at 50 C under vacuum for 24 h, in order to obtain the final solvent-free (and transparent) elastomer. The amounts of components used to prepare all the studied samples are given in Table 1. The neat matrix (polyMEA) was also prepared as a reference material: Its synthesis was analogous like in case of the nanocomposites, only the RDS clay dispersion step was skipped. The concentration of the MEA monomer (and hence of C C bonds) in the reaction mixture always was 0.75 mol/L. Several clay concentrations were tested: Cclay = 2, 4, 5, 6 and 10 wt.% in dry nanocomposite. The concentrations of the initiator APS and of the activator TEMED were characterized by the following molar ratios: r(APS) = [APS]/[MEA monomer] and r(TEMED/APS) = [TEMED]/[APS]. At “standard conditions”, the ratios were the following: r(APS) = 0.0087 and r(TEMED/APS) = 3.2. For the study of the effect of the co-initiators' concentrations on nanocomposite properties, either both concentrations were lowered at r(TEMED/APS) = const. = 3.2, or [APS] was kept constant and [TEMED] was varied (so that r(TEMED/APS) was changing), as is listed in Table 1.
2.3 | Characterization
2.3.1 | NMR spectroscopy
1H NMR spectra were recorded on a Bruker Avance III 600 spectrometer operating at 600.2 MHz. All NMR spectra were measured on samples placed in 5 mm NMR tubes, which were degassed and sealed under nitrogen; sodium 2,2-dimethyl-2-silapentane-5 sulfonate (DSS) was used as an internal NMR standard.
2.3.2 | Small-angle X-ray scattering
The experiments were performed using a pinhole camera (Molecular Metrology SAXS System) attached to a microfocused X-ray beam generator (Osmic MicroMax 002) operating at 45 kV and 0.66 mA (30 W). The camera was equipped with a multiwire, gas-filled area detector with an active area diameter of 20 cm (Gabriel design). Two experimental setups were used to cover the range of the scattering vector q from 0.004 to 1.1 Å−1, where q = (4π/λ)sin θ (λ is the wavelength of the X-rays, and 2θ is the scattering angle). The scattering intensities were put on an absolute scale using a glassy carbon standard.
2.3.3 | Light microscopy
The phase separation in the reaction mixture was investigated by standard wide-field light microscopy using the transmitted bright field mode, on a Nikon80i (from Nikon). The samples were put between microscopy glasses and observed directly without special preparation.
2.3.4 | Transmission electron microscopy—nanofiller dispersion
In order to characterize the nanofiller dispersion, transmission electron microscopy (TEM) was employed. Ultrathin slices (approximately 60 nm thick) of the solvent-free elastomers were cut using the Ultracut UTC ultramicrotome from Leica. The slices were put on supporting Cu grids and observed with the Tecnai G2 Spirit Twin 12 microscope (from FEI, Czech Republic) in the bright field mode at the acceleration voltage of 120 kV.
2.3.5 | Thermogravimetric analysis
TGA was performed using a Pyris 1 TGA thermogravimetric analyzer (Perkin Elmer, USA). in a temperature range from 35 to 750 C, at the heating rate of 10 C/min under a constant gas flow of 20.0 mL/min. All analyses were carried out in nitrogen, as well as in air.
2.3.6 | Infrared spectroscopy
The infrared (FTIR) spectra of the elastomeric films were measured using a Spectrum 100 spectrometer (from PerkinElmer) equipped with a mercury–cadmium– telluride (MCT) detector and an universal ATR (attenuated total reflectance) accessory with a diamond prism.
2.3.7 | Thermo-mechanical properties of elastomers
Dynamic-mechanical thermal analysis (DMTA) was carried out using an ARES-G2 apparatus from TA Instruments. The analysed temperature range was from −80 to +100 C, the heating rate + 3 C min−1. The applied oscillatory deformation had the constant frequency of 1 Hz, while the deformation amplitude was varied between 0.01 and 5% (regulated by the “auto-strain” function). The geometry of standard specimens was 20 mm × 6 mm × 1 mm. The temperature dependences of the storage shear modulus (G'), of the loss modulus (G") and of the loss factor (tan δ) were recorded.
2.3.8 | Tensile tests
The tensile properties of small and thin samples of the prepared elastomers were measured using an ARES-G2 from TA Instruments (maximum allowed force: 20 N), at room temperature, with a cross-head speed of 0.25 mm/s. Rectangular specimens (total specimen length: typically 15 mm, length between jaws: 3 mm, width of the specimens: 2 mm, thickness: 1 mm) were used, in order to prevent unwanted fracture close to the clamping region and to avoid grip slippage. At least three measurements were carried out for each sample. Presented are tensile curves closest to the average one.
3 | RESULTS AND DISCUSSION
3.1 | Synthesis of the nanocomposites
The studied nanocomposite elastomers were prepared (see Scheme 1) by free-radical polymerization of an aqueous solution of methoxyethyl acrylate (MEA), in the presence of dispersed commercial clay nano-platelets (“RDS”: nearly circular shape, diameter: ca. 25 nm, thickness: ca. 1 nm; see Scheme 1 bottom left). The nanofiller, which simultaneously served as physical cross-linker (see Scheme 2 and34,35), contained pyrophosphate as dispersion-enhancing agent (adsorbed on platelet edges, see Scheme 1 bottom left) and was fully exfoliated by stirring for 24 h in the water amount planned as reaction solvent, prior to the addition of further synthesis components. The polymerization was initiated by the redox pair ammonium peroxodisulfate (APS)/tetramethyl ethylene SCHEME 1 Synthesis of the studied nanocomposites [Color figure can be viewed at wileyonlinelibrary.com] diamine (TEMED) (see mechanism in Scheme 1 bottom right). After the reaction was finished (24 h process time at room temperature), the product was obtained as a very soft hydrogel. Subsequently, it was dried, in order to obtain the final solvent-free elastomer. The material properties of the nanocomposite elastomers were varied in a wide extent by changing the content of the RDS clay nanofiller (2–10 wt.% in dry gel), as well as by changing the concentrations of the co-initiators (and thus the average elastic chain length; see Experimental Part, and the discussion further below). The lower and upper limits of the tested clay concentrations arose from practical considerations: Clay amounts below 2 wt. % were not tested, because already at 2 wt.% of RDS the nanocomposite is very sticky and is difficult to handle, similarly like the neat polyMEA matrix. On the other hand, at 10 wt.% of the nanofiller, the reaction mixture becomes relatively viscous, so that increasing challenges arise with the removal of bubbles (important for realistic tensile tests), or with the homogeneous admixing of the last added synthesis component, the co-initiator APS (crucial at reduced co-initiator concentrations). Figure 1 illustrates the appearance of the reaction mixture and of the product during different stages of the preparation procedure: At first, a transparent mixture can be seen, which becomes turbid after several minutes of reaction. This phase-separation is caused by the hydrophobicity of polyMEA, in contrast to MEA monomer, which is water-soluble. The polyMEA/clay product in its hydrogel form hence is heterogeneous and composed of water-rich and polymer-rich domains, sized up to several micrometres, as can be observed by optical microscopy (see inlay in Figure 1). The mechanical properties of the hydrogels obtained as intermediates are poor. They are not well-cross-linked, due to the early precipitation of polymer/clay brush-like particles as the (hydrophobic) polymer chains start to grow from clay surface (see discussion further below). Efficient physical cross-linking of the whole material via polymer chains' entanglement (as shown in Scheme 2) is achieved only after product drying (24 h at room temperature in air +24 h at 50 C in vacuum): The previous turbidity disappears (see Figure 1), and the final solvent-free elastomers display excellent mechanical properties, including a very high extensibility. In this context it should be noted, that the drying process corresponds to the self-healing of a highly fragmented material (the SCHEME 2 Physical cross-linking in Haraguchi-type gels [Color figure can be viewed at wileyonlinelibrary.com] individual hydrophobic domains). The final dried nanocomposites hence can be expected to display selfhealing as well. Indeed, the neat matrix, as well as the nanocomposite with 2 wt.% of clay, were observed to display a high tendency to grow-together even at room conditions, without any “activation” treatment (glue-like behaviour). The self-healing was not further studied in this already extensive work.
3.1.1 | Completion of monomer conversion and the degree of nanofiller incorporation
An important aspect of the nanocomposite syntheses was the completion of the monomer conversion, as well as an efficient incorporation of the nanofiller. This was especially of interest in the cases, where the material properties were very widely varied via changing the coinitiators' concentration. The absence of un-reacted monomer was confirmed by means of FTIR spectroscopy: There is practically no trace of the intense C C stretching peaks of monomeric MEA, which would occur at 1638 and 1620 cm−1 (see Reference 52) in the FTIR spectra of the prepared solvent-free nanocomposites (see examples in the Supplementary Information file, SI-Figure 5). However, eventual unreacted monomer could have been lost during work-up of the precursor hydrogel. Thermogravimetric analyses in air (discussed further below), during which the organic part of the nanocomposites was burned-off at higher temperatures, nevertheless indicate, that the experimental weight fractions of the organic (lost weight) and inorganic (ash residue) fractions of the studied nanocomposites are very close to the expected ones (compare further below Figures 13C and 14C vs. Table 1). Hence, nearly quantitative MEA conversion, as well as RDS incorporation both can be assumed.
3.1.2 | Length (molecular mass) of the elastic chains
The chain length (molecular mass) of polyMEA, which was widely varied by the co-initiators concentrations, played an important role in this study. Its direct analysis was hence of considerable interest, and was attempted using a modified procedure, similar to the one used by the authors for the analysis of Haraguchi-type hydrogels48: Nanocomposite specimens were put into a mixture of THF and 40% hydrofluoric acid (1:1) and dissolved to a homogeneous solution. The clay platelets were dissolved by HF, while the predominantly hydrophobic polyMEA also dissolved in the THF-rich mixture. Without the addition of HF, the nanocomposites swell in THF, but do not dissolve. The organic polymer was then isolated by repeated precipitation/dissolution, using H2O/HF (later pure H2O) and THF. In spite of the successful isolation, repeated several times, the polyMEA solutions yielded no signals if subjected to GPC analysis (with any detector). The solutions also showed a tendency to gelation upon relatively short storage. The achieved change in polyMEA chain length hence had to be assessed only indirectly, in view of the tensile properties. In a previous work,48 a variation of molecular mass between 0.5 and 5 MDa was achieved by the authors for linear polyacrylamide chains grown on clay, leading to a nearly two-fold increase of elongation at break,49 via a similar synthesis approach like is used in the present work.
3.1.3 | Nanofiller distribution
Clay platelets observed by transmission electron microscopy The morphology of the solvent-free (and outwardly homogeneous) elastomers containing different amounts of clay is illustrated in Figure 2 (TEM). It can be observed, that the clay platelets (RDS) are generally wellexfoliated in the dry nanocomposites, especially at low clay contents. Further it can be seen in Figure 2, that the nano-platelets are not evenly distributed, but arranged to patterns which are separated by regions with little clay. Especially at the highest clay contents, e.g. at 6 and 10 wt.%, the cell-like (diameter: ca. 200 nm) nanofiller pattern becomes more distinct and the nanofiller more enriched in it, including occasional loose aggregates of intercalated filler (see Figure 2). The observed morphology obviously is a result of the above-mentioned phase separation during the nanocomposites' synthesis. An idealized structure is shown in Scheme 3, with precipitated polymer-covered clay platelets and with clay-free domains filled by polymer grown in the later reaction stages. X-ray diffraction (SAXS/WAXS) on the polyMEA/clay nanocomposites The morphology of the prepared nanocomposites was also investigated by means of X-ray diffraction. Figure 3 compares the diffractograms of the neat nanofiller, of its dispersion in water (as used to prepare the synthesis mixture for the sample “6 RDS” with 6 wt.% of clay in the dry nanocomposite), of the hydrogel form of the “6 RDS” product, as well as of the final (dried) “6 RDS” sample. The XRD profile of the neat RDS clay displays a shoulder at the value of the scattering vector q = 0.558 Å−1 (≡ 15.9 , corresponding to d = 0.56 nm) and several relatively sharp reflections at 1.45 Å−1 (20.5 ≡ 0.43 nm), 1.88 Å−1 (26.7 ≡ 0.33 nm) and 2.46 Å−1 (35.1 ≡ 0.26 nm), which all can be assigned to internal distances in the RDS (montmorillonite) platelets. In addition to this, RDS also displays a broad flat (amorphous) peak centred on 0.5 Å−1 (7 ≡ 1.3 nm), which is generated by the somewhat varying distance between the irregularly stacked nanoplatelets. The XRD profile of neat RDS further contains a scattering feature which step-wise rises already below the mentioned flat peak at 0.5 Å−1 and continues to grow approximately linearly with decreasing q (down to q = 0.005 Å−1), with an approximate slope of −2 (log/log scale). This SAXS pattern can be assigned to smaller and larger continuous domains of co-planarly arranged RDS platelets (2D objects). SCHEME 3 Idealized representation of the arrangement of clay platelets and of polyMEA chains in the studied solvent-free nanocomposite elastomers, with “freely” mobile elastic chains, and with strongly immobilized ones (close to the platelets) [Color figure can be viewed at wileyonlinelibrary.com] The aqueous dispersion of RDS displays a similar but simpler pattern than the neat clay: In the wide-angle region, a broad “amorphous” peak is observed with local maxima near 2.05 Å−1 (29 ≡ 0.31 nm) and 2.88 Å−1 (41.4 ≡ 0.22 nm), which are both characteristic of pure water, and which overlap eventual “amorphous” reflections of the exfoliated and randomly oriented clay platelets. The RDS dispersion also displays a distinct SAXS scattering feature (like dry RDS), but without the stepwise scattering increase near 0.5 Å−1 and without the flat peak at this position generated by the inter-platelet stacking distance: this latter disappeared due to exfoliation. The scattering intensity grows with decreasing q over less than two decades, after which the curve “flattens.” The shape and the slope (very close to −2.0 in log/log scale) of the SAXS pattern are consistent with well-exfoliated and randomly oriented 2D platelets: no long-distance coplanarity, hence shorter linear range in the SAXS region. Guinier analysis of the pattern yields a radius of gyration of 12.1 nm (platelet width: 24.2 nm). The reaction mixture for the synthesis of “6 RDS” (without the last additive, the co-initiator APS) generates a highly similar XRD pattern like the aqueous RDS dispersion. The slight differences between the two mentioned diffractograms (more extended SAXS feature of the “6RDS synthesis mixture”) can be assigned to some clay platelet association caused by the added co-initiator TEMED, which has a high affinity for the clay platelets, as was first observed by the authors in Reference 48. This is also illustrated by the NMR spectra in the Supplementary Information file, SI-Figures 1 and 2, which prove the strong adsorption/immobilization of TEMED on the clay. The small difference between the scattering curves in the region from 0.4 to 2 Å−1 possibly is caused by monomer adsorption on the platelets. The hydrogel form of the nanocomposite “6 RDS” has a relatively similar XRD pattern (Figure 3) like its synthesis mixture (without APS), but it additionally shows three “amorphous” peaks characteristic of the neat polyMEA matrix (compare Figure 3 with Figure 4), namely at 0.67 Å−1 (9.39 ≡ 0.94 nm), at 1.49 Å−1 (21.0 ≡ 0.42 nm) and at 3.01 Å−1 (43.3 ≡ 0.21 nm). The SAXS region in the hydrogel is somewhat more irregular than in the reaction mixture, suggesting the simultaneous presence of isolated randomly oriented, as well as of associated more or less co-planar platelets. This irregularity seems to be the result of the further-above mentioned precipitation of polymer-covered clay platelets. The X-ray diffraction pattern of the solvent-free nanocomposite elastomer “6 RDS” is very similar to the one of the “6 RDS hydrogel”: One small difference is, that the first peak (at 0.67 Å−1) of the polyMEA matrix is somewhat more distinct in the dried nanocomposite. Another difference is the very linear shape of the SAXS scattering pattern in the dried sample and its extent over two decades of q, with a slope relatively close to −2 (in log/log scale; exact value: −2.5). This suggests a more extended arrangement of the 2D clay platelets in the dried nanocomposite, than in the hydrogel or in the reaction mixture, e.g. like the patterns observed by TEM in Figure 2 in clay-rich domains. The high negative value of the slope could be a consequence of non-sharp borders of the polymer-covered (hairy) nanofiller platelets. The Figure 4 compares the diffractograms of the neat polyMEA matrix and of the solvent-free nanocomposite elastomers containing 2–10 wt.% of the RDS clay. As already mentioned above, the polyMEA matrix displays three characteristic “amorphous” peaks at 0.67, 1.49 and 3.01 Å−1. The matrix additionally displays a SAXS scattering feature (region: 0.005–0.2 Å−1), which is much less intense than in the nanocomposites. The SAXS scattering feature in all the nanocomposites is nearly perfectly linear (slope: −2.5), like in case of the above-discussed sample “6 RDS”, and it decreases by four orders over two decades of q. The linear scattering features in the SAXS region are nearly perfectly parallel for all the nanocomposites, and their intensity expectedly increases with nanofiller content. In the region between q = 2 and 3 Å−1, small reflection peaks of the RDS platelets barely can be discerned (overlaid over the region of the amorphous peaks of polyMEA), which are more intense at higher filler loadings (better visible in detail images of the WAXS region in SI-Figure 4). Another interesting feature appears at 10 wt.% of RDS: In the region q = 0.2–0.45 Å−1, the course of the scattering curve can be interpreted as an overlaid very flat and small peak centred around 0.3 Å−1, which would correspond to a distance d = 2.1 nm of polymer-intercalated RDS platelets (the distance was 1.3 nm in the neat clay). To a much smaller extent, this feature might be present also in the samples with 6 and 5 wt.% of RDS. To sum up, it can be concluded, that the XRD analyses are in good agreement with the morphologies observed by TEM and with the schematic nanocomposite structure suggested in Scheme 3. Also the average RDS platelet width of ca. 25 nm, as stated by the producer, can be confirmed.
3.2 | Tuning of the chain length of the polyMEA matrix
In this work, the main aim was to vary the mechanical properties and the extensibility of the polyMEA/clay nanocomposites by changing the chain length of the polyMEA component. In the authors' previous work concerned with poly(N-isopropylacrylamide)/clay hydrogels, which were prepared using the same redox initiating system (APS/TEMED, see Scheme 1 bottom right), the authors found that, the chain length can be widely varied (by one order: between 0.5 and 5 MDa) by changing the concentration and/or the ratio of the co-initiators,48 thus achieving up to a two-fold increase in extensibility.49 The latter work highlighted the key role of the adsorption of the co-initiator TEMED on the clay platelets. This latter effect was expected to be important also in the present work (see Scheme 4). During the synthesis (see Section 2), first the clay platelets (RDS) were thoroughly dispersed in water, after which the monomer MEA was added. Subsequently, the co-initiator TEMED was added, which was found to adsorb very strongly on RDS, albeit the efficiency of the adsorption was smaller with RDS, than with the “XLS” clay studied in Reference 48, because RDS was added in markedly smaller amounts to the reaction mixtures studied in the present work. The adsorption on RDS was demonstrated by 1H-NMR (see SI-Figure 1 and 2) as an apparent decrease of signal intensity (as well as signal broadening) of TEMED (at unchanged TEMED concentration) if RDS was added, due to TEMED fixation on clay surface, and hence its slow relaxation (this leads to decrease in signal intensity under the standard pulse sequence regime in the NMR experiment). Similarly, the MEA monomer was also found to absorb on RDS (see SIFigure 3), although most of it stays in solution, due to relatively high MEA concentration vs. few places on RDS surface – as consequence of the synthesis amounts of clay and MEA. At the begin of the polymerization process, the situation on (and around) the clay platelets hence is as depicted in Scheme 4: In the moment, when the APS peroxo initiator is added as the final component to the mixture, TEMED is nearly quantitatively adsorbed on the platelets, along with much more numerous monomer molecules, which additionally richly occur also in the surrounding solution. The admixed APS molecules diffuse to the platelets and generate radicals upon collision with TEMED, thus starting the polymerization, at first at the platelet surface. In the early reaction stage, clay platelets covered by some polyMEA still are homogeneously dispersed in the reaction mixture (Scheme 4, right), but later they precipitate due to polyMEA hydrophobicity. The chain ends continue to grow also after phase separation, on the interface between precipitate and the aqueous solution of MEA monomer. In case that the termination reactions are sufficiently suppressed, the chain lengths achieved by free-radical polymerization are controlled by the ratio of the monomer molecules to the initiating centres. Hence, a smaller number of initiating radicals ideally would lead to longer elastic chains, if the amounts of monomer and clay nanofiller remain unchanged (see Scheme 5). As already SCHEME 4 Start of the MEA polymerization on the RDS clay platelets and the early reaction stage [Color figure can be viewed at wileyonlinelibrary.com] mentioned above, in a previous work48 the authors achieved a ten-fold change in the elastic chain length in a polyacrylamide/clay nanocomposite by varying the coinitiators concentrations, hence this approach was expected to be efficient also for polyMEA/clay. Longer elastic chains in simple polymer networks should lead to a higher material extensibility (elongation at break), but also to lower moduli (smaller concentration of elastic chains which are larger). In case of the physically cross-linked elastomers of Haraguchi type, which are composed of linear polymers and clay platelets, the situation is more complex, as polymer chain entanglements (besides polymer adsorption on clay, and chain end recombinations) are likely to play an important role (see Schemes 2 and 5). In the latter case, long polymer chains could support more numerous entanglements, so that not only elongation at break, but also the modulus might grow with increasing chain length. In view of their previous experience from References 48 and 49, the authors in this work chose to vary the concentration of the initiating redox pair TEMED/APS in two ways: 1. by simultaneously decreasing the concentrations of both co-initiators, which leads to a smaller number of polymerization-starting sites (adsorbed TEMED) on the clay platelets, and additionally to a markedly slower rate of radical production (lower value of the kinetic term c(TEMED) * c(APS)); 2. by decreasing only the concentration of TEMED, which also leads to a smaller number of radical sites on RDS platelets, but to a less dramatic decrease in the rate of radical production (as c(APS) stays unchanged). Both methods for obtaining longer polyMEA chains were applied on samples with different clay content, in order to vary the mechanical properties and extensibility in the widest possible range, and to evaluate the effects of chain length at different nanofiller loadings. As mentioned in the synthesis discussion further above, it was not possible to determine the molecular mass of the polyMEA component in the prepared polyMEA/clay nanocomposites, due to branching reactions, but it could be expected, that comparable effects were probably achieved like in the authors' previous work48 concerned with the chemically related polyacrylamide/clay (10-fold chain length increase).
3.3 | Thermo-mechanical properties of elastomers
The thermo-mechanical characteristics of the prepared polyMEA/clay nanocomposites are presented in Figures 5–7. Figure 5 analyses the effect of the RDS clay content in samples prepared at standard concentrations of the coinitiators (TEMED and APS). It can be observed, that the nanocomposites with 2 and 4 wt.% of RDS clay, especially “2RDS”, display similar DMTA profiles (G' = f(T), see Figure 5(a)) and relatively similar tan δ delta curves (Figure 5(b)) like the neat matrix. In 2RDS, the increase in modulus in the rubbery region is negligible, while in 4RDS, the reinforcing effect of the clay nanofiller already is apparent. In case of 2RDS, and even more so of 4RDS, the glass transition temperature (Tg) moderately shifts to lower values (see maxima of tan δ peaks in Figure 5(b)), probably due to a somewhat more disordered packing of polyMEA chains in both nanocomposites, in comparison to the neat matrix. The neat matrix and 2RDS (but not 4RDS) are sticky, glue-like materials. A further increase (from 4 to 6 wt.%) of the content of RDS clay, which acts both as nanofiller and as physical cross-linker, leads to a SCHEME 5 Tuning polyMEA chain length via variation of the number of active initiating centres per one RDS platelet [Color figure can be viewed at wileyonlinelibrary.com] marked change in the DMTA profile: The rubbery modulus increases by nearly one order of magnitude in the low-temperature region of the rubbery plateau, and the glass transition becomes divided into two distinguishable steps (see Figure 5(a)). Accordingly, two peaks are observed in the tan δ = f(T) curve for both transitions (Figure 5(b)). The appearance of the second glass transition step can be assigned to a prominent immobilized fraction of the chains in the polymer matrix, most likely such ones between closely neighbouring platelets, as suggested in Scheme 6 as well as in the discussion of the X-ray diffraction (see further above: Figure 4: flat peak near 0.3 Å−1 assigned to polymerintercalated platelets). Further increase of the clay FIGURE 6 Thermomechanical properties (DMTA) analysis) of 4RDS nanocomposites (a) prepared at reduced concentrations of both co-initiators; (b) prepared at reduced concentrations of TEMED only; in both graphs: left y axes: G' = f(T); right y axes: tan δ = f(T) [Color figure can be viewed at wileyonlinelibrary.com] FIGURE 7 Thermo-mechanical properties (DMTA) analysis) of 10RDS nanocomposites (a) prepared at reduced concentrations of both co-initiators; (b) prepared at reduced concentrations of TEMED only; in both graphs: left y axes: G' = f(T); right y axes: tan δ = f(T) [Color figure can be viewed at wileyonlinelibrary.com] content from 6 to 10 wt.% leads only to small changes in the G' = f(T) and tan δ = f(T) curves: Significantly, the tan δ peak of the immobilized polymer fraction becomes the most prominent, and the rubbery modulus moderately rises. To sum up, the polyMEA/clay nanocomposites prepared with standard initiator concentrations are typical rubbery materials (G' = ca. 1 MPa) at clay contents of 2 and 4 wt.%, and stiff elastomers (G' = ca. 20 MPa) at clay contents of 6 wt.% and higher. The former nanocomposites (2RDS, 4RDS) display a single sharp glass transition and tend to simple behaviour during tensile tests (as will be discussed further below), while the latter group (6RDS, 10RDS) displays a glass transition in two steps and tends to a distinct neck formation upon tensile testing (these latter materials also display elastic recovery, however). The effects of the increase in chain length of the polyMEA matrix on the thermo-mechanical properties of the polyMEA/clay nanocomposites were evaluated on the samples 4RDS (soft elastomer) and 10RDS (stiff elastomer). Longer polyMEA chains were generated either by simultaneously reducing the concentration of both coinitiators by the factor of 2, 4 and 10, or by reducing only the concentration of the TEMED co-initiator by the factor of 2, 4 and 8. In case of the sample 4RDS, the effects of the chain length variation on the DMTA profiles are relatively moderate: The strongest effect is the increase of the rubbery modulus by the factor of 4 (if concentrations of both co-initiators are reduced, see Figure 6(a)) or 3 (if concentration of TEMED alone is reduced, see Figure 6(b)). Both methods for increasing the chain length lead to slightly up-shifted glass transition temperatures. If the concentration of both co-initiators is lowered two times, the mentioned fourfold increase in rubbery modulus is achieved (Figure 6(a)), simultaneously with the slight upshift of Tg. Subsequent further reduction of the concentration of both co-initiators leads to a slight decrease in rubbery modulus (seemingly a trend reversal, albeit weak). In case of the chain length variation by reducing the TEMED concentration only, the general trend is similar: the maximum modulus increase and most of the Tg up-shift is achieved when c(TEMED) is reduced by the factor of 2 (see Figure 6(b)). Further decrease of c (TEMED) leads to a decrease of rubbery modulus, especially at higher temperature where the modulus nearly returns to its initial value (achieved with standard synthesis). The Tg values further slightly increase, if c (TEMED) is lowered (Figure 6(b)). A comparison with the results of tensile tests (discussed further below) shows, that the structure–property relationships in the studied nanocomposites are relatively complex: In case of the 4RDS samples prepared at reduced concentrations of both co-initiators, the elongation at break initially increases (1/2 [TEMED, APS]) but then decreases upon further reduction of the co-initiators concentrations (while toughness steadily increases, see further below: Figure 11(a), (b)), even more markedly than the rubbery moduli in Figure 6(a). On the other hand, if c(TEMED) only is reduced, the elongation at break steadily increases (simultaneously with toughness, see Figure 11(c),(d)), while the rubbery modulus reaches a maximum at 2x less TEMED (Figure 6(b)) and subsequently decreases upon further reduction of c(TEMED). In this latter case, the combination of both mechanical tests suggests, that longer polyMEA chains are formed, which are not enough entangled to cause an increasing trend in rubbery moduli, but which still are sufficiently entangled to generate a higher modulus than with standard chain length (in a simple network, the longer chains would reduce the modulus). In case of the sample 10RDS, the effects of the changes in co-initiator concentrations on the DMTA profiles are more dramatic than in case of 4RDS (see Figure 7). If the concentration of both co-initiators is reduced simultaneously, the rubbery modulus strongly and gradually increases (no trend reversal). At the lowest coinitiators concentration (“1/10 (TEMED, APS),” see Figure 7(a)), the two-step glass transition is replaced by a SCHEME 6 Adsorbed, immobilized and unhindered (free) polymer chain segments in polyMEA/clay nanocomposites: (a) with low clay content, (b) with high clay content (leading to a larger fraction of chain segments immobilized in different degree) [Color figure can be viewed at wileyonlinelibrary.com] simple one (see G' = f(T) curves) and the rubbery modulus reaches its highest value, six times higher than in case of standard synthesis. If the tan δ = f(T) curves are compared (Figure 7(a)), the peaks of the moderately and of the strongly immobilized matrix fractions gradually decrease in intensity and nearly disappear at the co-initiators concentration “1/10 (TEMED, APS),” leaving only one intense glass transition peak of the unhindered matrix chains. The effect of the chain length on the glass transition behaviour, namely the decreasing fraction of the immobilized segments, is illustrated in Scheme 7. In case that the length of polyMEA chains in “10RDS” is varied by reducing c(TEMED), the effects are somewhat smaller (see Figure 7(b)): Most of the changes occurs if reducing c(TEMED) by the factor of 2: The rubbery modulus increases four times and the tan δ peak of the immobilized matrix fraction markedly flattens and shifts to higher temperatures. Further decrease of c (TEMED) causes little effect in the G' = f(T) and tan δ = f (T) curves.
3.4 | Tensile properties
The high tensile properties are the most attractive property of the elastomeric nanocomposite polyMEA/clay gels, which were previously known38 to display an extensibility comparable to the Haraguchi's hydrogels (also physically cross-linked by clay, as discussed in the Section 1, see Reference 31). However, the synthesis modification done in this work, which was aimed at obtaining longer elastic chains, made it possible to improve the extensibility of the solvent-free elastomers several times (6× in the extreme case). Maximum elongations of up to 6000% (60-fold stretching) were achieved, which is a record value for solvent-free elastomers and which also markedly exceeds the mentioned Haraguchi hydrogels (as discussed in the Introduction, ultraextensible hydrogels were generally more easily obtained than solvent-free ultra-extensible elastomers). Interestingly, the highest extensibility values were obtained for the nanocomposites with the highest clay content. This finding seems to correlate with the authors' previous observation in Reference 48 that high clay contents favoured the growth of longer polyacrylamide chains. In general, products with widely different deformation behaviour, average moduli, as well as maximum extensibilities were obtained via changing the concentration of the co-initiators.
3.4.1 | High reversibility of the stretching deformation
The studied polyMEA/clay nanocomposites nearly completely recover (after sufficiently long time) even from high non-destructive deformations, as noticed already by Haraguchi and co-workers in Reference 38, and as illustrated in Figure 8. This property was SCHEME 7 Immobilized and unhindered (free) polymer chain fractions in polyMEA/clay nanocomposites: (a) with short polymer chains, (b) with long polymer chains (for simplicity, only one clay platelet and one polymer molecule is shown in each case) [Color figure can be viewed at wileyonlinelibrary.com] characteristic for all the studied samples, independently of the shape of their tensile curve. The degree of recovery of the deformation can be seen in detail in Figure 8 (d)–(f). If only the deformed part of the sample is considered (3 mm long; the flattened and more opaque sample ends did not take part in the deformation; the inclusion of the non-deformed parts into analysed sample length would numerically but not correctly improve the recovery degree), the exemplary samples recover from maximum deformation of 2000% back to 200% residual strain (90% of the elongation is recovered: Figure 8(e)), or from 3000% back to 633% (80% of the elongation is recovered: Figure 8(f)). The contraction of the previously stretched samples is somewhat delayed, but most of it occurs in several minutes and most of it in 30 min, although some slower contraction still occurs after this period of time (see Figure 9(a)). The Figure 9(b) additionally shows that the tensile properties of the stretched samples are approximately regenerated after a certain “healing time.” This healing process seems to be slower than the deformation retraction (see Figure 9(b)). The delayed but nearly complete elastic recovery of the polyMEA/clay nanocomposites can be explained by the structure suggested in Scheme 8: In addition to physical cross-links formed by reversible entanglements of polymer chains, there are also permanent (covalent) cross-links which lead to branched structures extending over the whole sample (or over large parts of it). These permanent cross-links (branched polymer molecules) are then responsible for the highly efficient shape memory of the samples, and also for the problems with analysing the molecular mass of the polyMEA component, as discussed further above (Synthesis). The reversible entanglements (physical cross-links) raise the “dynamic” modulus value, as well as the samples' toughness, and are also responsible for the delayed course of elastic recovery. The high reversibility even of very large deformations and the relatively small residual strain (in comparison to the maximum deformation) indicates that the studied polyMEA/clay nanocomposites are rather to be classified as true elastomers rather than so called “plastomers” (materials displaying elastic behaviour at small deformations and predominantly plastic one at large strains). The dependence on synthesis parameters of the deformation recovery, as well as of the time-dependent re-generation of the tensile properties will be described in a separate follow-on work.
3.4.2 | Effect of clay content on the tensile properties
The results of the characterisation of the tensile properties of the studied nanocomposite elastomers are summarized in Figures 10–12. The stress–strain curves were analysed in two formats: first in the “engineering format,” which is most frequently used in the literature (it has a deep sense, if characterizing construction parts) and which assumes the nominal constancy of the sample cross-section. As an alternative depiction, graphs plotting an idealized “true stress” are also shown in the Discussion and in the Supporting Information File. In that second case, volume constancy (Poisson ratio of 0.5) is assumed for the deformed part of the specimen, and hence a decrease of cross-section which is equal to the elongation ratio (= actual_length/initial_length). Effects like neck formation in elastomer samples (see e.g. Figure 10 and inlay with green frame in it) are not well-covered by this model, which calculates an average cross-section for the whole deformed specimen. This second format of the stress–strain curves focuses the interest on the intrinsic material properties in the deformed state, instead of describing the properties of a mechanical part. The “true stress” curves hence are of academic interest if evaluating structure–property relationships. The Figure 10 shows the effect of clay content on nanocomposite samples prepared at standard initiating conditions: The glue-like neat polyMEA matrix displays an elongation at break of ca. 1000%, in combination with low moduli (initial region, as well as average), low yield stress, very low stress at break, and with low toughness. The rubbery behaviour of polyMEA (see the discussion of DMTA further above), as well as the tensile properties seem to be the result of entanglements of the long polymer chains of the neat matrix. Addition of 2 wt.% of RDS much increases the stiffness (especially the initial modulus—curve slope), the toughness, as well as the yield stress and the stress at break. Similarly like in case of the neat matrix and even more distinctly, the sample 2RDS displays a deformation curve similar to thermoplastics, but in contrast to thermoplastics, it recovers nearly completely even from high non-destructive deformations (after some delay, as mentioned above). The sample 2RDS displays a similar gluing tendency like the neat matrix. The small amount of physical cross- linking by clay seems to highly improve the extensibility, which for the sample 2RDS reproducibly is ca. two times higher (1800%) than the extensibility of the neat matrix, or of the higher-reinforced nanocomposites. In case of higher RDS loadings, the maximum extensibility recedes back to 1000%, but the stiffness and toughness of the nanocomposites strongly increase, as well as the yield stress and stress at break. Also the “thermoplastic-like” shape of the tensile curves in the engineering format always is preserved (although the samples fairly efficiently recover from deformations, like the exemplary specimens in Figures 8 and 9). It can be noted, that the plastic-like tensile curves display a maximum (yield point) followed by stress decrease in the “engineering format,” while in the “true stress format” (see SI-Figure 6), there is a relatively abrupt decrease of slope (to a less positive value, but not negative) at the yield point. The trends can be explained as follows: At low clay content, there are more numerous polymer chains per one clay platelet (physical cross-linker). This favours a structure in which a shorter segment of the polyMEA polymer chain is adsorbed, while a longer segment is free to participate in entanglements (see principle in Scheme 7 further above) and in occasional branching (see Scheme 8 further above) via chain transfer reactions. Additionally, at very low clay concentrations like in case of 2RDS (2 wt.% in dry elastomer = ca. 0.2 wt.% in the initial synthesis solution), some polyMEA molecules likely at first form in solution, due to some non-adsorbed initiators' molecules (see NMR spectra of TEMED adsorption in SI- Figure 1–2), and are later attached to the clay-anchored elastically active chains, via the mentioned chain transfer reactions. In this way, the elastic effect of polyMEA chains is greatly enhanced at very low clay contents, like in 2RDS. At higher clay concentrations, on the other hand, the adsorbed segment of the polyMEA chains is longer, as there are fewer polyMEA chains per clay platelet, and hence the “free” flexible segment of these chains which undergoes entanglements and occasional cross-linking is shorter. Secondly, due to less initiator molecules per clay platelet, also the formation of additional chains apart of the platelets' surface is much reduced, which possibly is the main reason for the abrupt decrease in elongation at break between 2RDS and 4RDS, and for the subsequent saturation of the clay effect. This saturation, namely the absence of significant further decrease of the elongation at break with increasing filler loading at the highest clay contents, may be further favoured by the fact, that the presence of clay platelets under certain conditions helps the formation of longer polyacrylate chains (see Reference 48), which would balance the “adsorption effect” at high clay concentrations.
3.4.3 | Effect of polyMEA chain length on tensile properties
The effect of varying the chain length via changing the concentrations of the co-initiators was investigated in detail for the soft elastomeric sample 4RDS (4 wt.% of clay) and the stiff elastomer 10RDS (10 wt.% of RDS). The results obtained for the sample 4RDS are summarized in Figure 11. As mentioned further above, the polyMEA chain length was varied by either by simultaneously reducing the concentration of both coinitiators (TEMED and APS) by the factor of 2, 4 and 10 (effects: see Figure 11(a),(b)), or by reducing only the concentration of TEMED by the factor of 2, 4 and 8 (see Figure 11(c),(d)). It can be seen that the changes in the concentrations of the initiating system make possible a very broad variation of the tensile properties, which were tremendously improved, especially the elongation at break and the toughness (see Figure 11), and even a change of the type of the tensile curve, from quasi-plastic to elastomer-like was achieved. Simultaneous decrease of the concentration of both initiators, “c(TEMED, APS)”, while keeping constant the TEMED:APS molar ratio, leads to somewhat more dramatic effects, than the decrease of concentration of TEMED alone. Lowering c(TEMED, APS) leads to a great increase in elongation at break and to a tremendous increase in toughness. In case of 10-fold decrease of c (TEMED, APS), where the highest toughness is achieved, also the shape of the tensile curve (in the “engineering format”) changes from “thermoplastic-like” to “elastomer-like.” In case of the “true-stress-format,” all the curves are elastomer-like, except the “standardly” prepared sample: see Figure 11(a). The elongation at break upon 2-fold reduction of c(TEMED, APS) increases 3.6 times (to 3600%), while further reduction of c(TEMED, APS) leads to a moderate decrease of the elongation at break, down to 2500%, but still 2.5x better than in “standard” sample. On the other hand, the toughness (and also the yield stress and the stress at break) further dramatically increases with decreasing c(TEMED, APS) (see Figure 11 (a),(b)). The 10-fold reduction of c(TEMED, APS) means a 10-fold reduction of the number of TEMED molecules per RDS clay platelet, as well as a 100-fold reduction of the rate of the radical initiation reaction (TEMED + APS). The increase in toughness and (generally) in the elongation at break appears to be a consequence of the formation of longer polyMEA chains. The decreasing elongations at break (with falling c(TEMED, APS)) are possibly a consequence of inhibition by traces of oxygen (some syntheses of “4RDS - 1/10 c(TEMED, APS)” were observed to fail), of the resulting irregularity in chain lengths, and maybe also of the less efficient entanglement of very long matrix chains, due to precipitation in course of the synthesis (as discussed further above) in the nanocomposite 4RDS, which is not very rich in nanofiller. Reducing the concentration of the TEMED coinitiator alone, c(TEMED), leads to comparably strong effects like the decrease of c(TEMED, APS), but the changes in the shapes of the tensile curves are less dramatic (Figure 11(c),(d)): The curves in the “engineering format” always keep the thermoplastic-like shape, although a long-range and increasingly marked positive slope is observed in the post-yield-point region for the samples synthesized with 1/4 and with 1/8 of c(TEMED)standard. The maximum achieved elongation at break (at the lowest c(TEMED) value) is even slightly higher than the one obtained by reducing c(TEMED, APS): 3750 vs. 3600%. The values of the yield stress are somewhat smaller, and the values of stress at break are markedly smaller (e.g. with “1/8 c(TEMED)”) if compared to samples with reduced c(TEMED, APS). The elongation at break grows steadily, if c(TEMED) alone is reduced: The inhibition effect of oxygen seems to be less pronounced, as the concentration of the peroxo-component APS remains on the standard level. The initiation rate also is decreased less dramatically in this series: Eight times in the extreme case, if c(TEMED) was reduced 8-fold. If optimized 4RDS specimens are compared, which were synthesized at theoretically the same (slowed-down) rate of initiation, “4RDS 1/2 (TEMED, APS)” vs. “4RDS - 4x less TEMED”, the sample with “4x less TEMED” is much tougher, has a dramatically higher stress at break but a somewhat smaller elongation at break. Due to two times fewer chain-starting nitrogen centres (from TEMED) in the sample with “4x less TEMED”, this latter should contain fewer polyMEA elastic chains, which should hence be longer and possibly more entangled, which would explain the higher stiffness. If the “4RDS” nanocomposites are compared, which were prepared using the same (reduced) concentration of the TEMED co-initiator, e.g. “4RDS 1/2 (TEMED, APS)” vs. “4RDS - 2x less TEMED”, it can be observed that their tensile curves mostly overlap, the yield stress and the stress at break are similar, but the elongation at break of the sample with “2x less TEMED” is smaller by ca. 1/3, if compared with “4RDS 1/2 (TEMED, APS)”. The difference might be tentatively explained by the faster radical production rate in the sample “2x less TEMED”, which might favour shorter chains (due to side-reactions). In case of the sample 10RDS, the concentration of the initiating system has even more dramatic effects than in 4RDS (see Figure 12). The changes in the shapes of the tensile curves of 10RDS (Figure 12) are much more sensitive to the initiators' concentrations, and elastomer-like curves are more easily achieved, in spite of the higher content of the nano-reinforcement/cross-linker (RDS clay). Record-high elongations at break also were achieved at 10 wt.% RDS (ca. two times higher than with 4RDS), namely up to ca. 6000%. As mentioned further above in view of previous work,48 the higher-concentrated clay platelets likely favour the formation of longer polyMEA chains, thus boosting the extensibility. Simultaneous decrease of the concentration of both initiators, “c(TEMED, APS)”, during the synthesis of 10RDS, leads to a tremendous increase in the elongation at break (see Figure 12(a),(b)), which reaches its maximum (ca. 6000%) at “1/2 c(TEMED, APS)”. A further decrease of the co-initiators' concentration also reduces the elongation at break (like in 4RDS), but only slightly, while the toughness continues to strongly grow. This indicates that the inhibition effect is smaller in 10RDS, possibly due to oxygen adsorption on RDS. While the curve of “10RDS 1/2 (TEMED, APS)” still is thermoplastic-like (albeit without the distinct peak at the yield point), already the next curve, the one of “10RDS 1/4 (TEMED, APS)” is elastomer-like. Generally, all the 10RDS samples optimized by reducing c(TEMED, APS) displayed neck formation at the begin of the tensile test. The ones formed at 1/4 and 1/10 c(TEMED, APS) additionally displayed the formation of several necks (see inlay in Figure 12), which manifested itself as occasional distinct irregularities in the tensile curves in Figure 12. As mentioned further above and illustrated in Figures 8 and 9, also these samples display a highly efficient deformation recovery, independently of the tensile curve shape. The necking also does not occur in case of repeated tensile test, only during the first one. Reducing the concentration of the TEMED co-initiator, c(TEMED), in 10RDS leads to analogous effects like the reduction of c(TEMED, APS), but the improvement in tensile properties (see Figure 12(c),(d))—though excellent—is visibly smaller, than in case that c(TEMED, APS) was lowered. Multiple neck formation is even more pronounced than in the c(TEMED, APS) series, especially in case of the samples “10RDS - 4x less TEMED” and “10RDS - 8x less TEMED”, which nevertheless display elastomer-like curves. This latter behaviour correlates with the more prominent presence of immobilized polyMEA fractions in optimized 10RDS samples obtained by lowering c(TEMED) alone as demonstrated by DMTA (see Figure 7(b) further above: tan δ = f(T) curves). In contrast to the analogous series based on 4RDS, the elongation at break (of 10RDS) somewhat decreases if c (TEMED) is lowered more than two times (but is still much better than in the “standardly prepared” sample). This trend seems to correlate with the rapidly increasing tendency to multiple neck formation during the tensile test, which might favour earlier sample destruction. To sum up it can be concluded, that the wide variation of mechanical properties of the polyMEA/clay nanocomposite elastomers via changing the co-initiators' concentrations was highly successful. Both synthesis approaches, the lowering of the concentration of both coinitiators, as well as the lowering of the concentration of TEMED only, yielded promising products if soft nanocomposite elastomers with low clay contents are concerned. In case of stiff elastomers rich in clay, the first method yielded similar but always better products than the second.
3.5 | Thermal stability of the nanocomposite elastomers: TGA
The effect of the nanofiller content and of the synthesis conditions on the thermal stability of the prepared nanocomposites was evaluated by means of thermogravimetric analysis (TGA), both in nitrogen and in air. As will be discussed below, the inorganic RDS nanofiller was found to markedly improve the stability of the nanocomposites, if heated in air. On the other hand, even a marked increase in polyMEA chain length caused only a minimal (but detectable) change in the TGA traces. The effect of the content of clay (RDS) is shown in Figure 13: In case of the TGA tests conducted in nitrogen atmosphere (Figure 13(a),(b)), there is no visible water release from the nanocomposites, maybe a slight one from the neat MEA matrix (MEA: at 100 C: loss of 1.5%, at 200 C: 2.6%; nanocomposites at 200 C: 1–1.8%). Increasing clay content only slightly and gradually up-shifts the temperature of maximum decomposition rate in nitrogen Tmax (main peak in dTGA graph, Figure 13(b)), from 395 C to 400 C. In contrast to Tmax, the RDS nanofiller strongly upshifts the temperature of the onset of decomposition, Tonset in nitrogen (defined as temperature at which 5 wt. % loss is achieved), if going from neat matrix to the nanocomposites, from 286 C to 344 C–355 C (nanocomposites). This effect is mainly due to the absence in the nanocomposites of a small decomposition step between 250 C and 295 C (local Tmax of small dTGA peak at 279.5 C), which is characteristic for the neat matrix. The shift of Tonset is the sole marked effect caused by the RDS nanofiller if the samples are heated under nitrogen. This effect might be caused by adsorption or by the hindrance of diffusion by the RDS platelets of some reactive decomposition intermediate. From Figure 13(b), it can be further noted, that the MEA matrix to a small extent undergoes carbonization: The mass residues at 700 C are 2.4% (ideal: 0%) for neat polyMEA, 4.06% for 2RDS (id.: 2.0%), 6.11% for 4RDS (id.: 4.0%) and 13.3% for 10RDS (id. 10.0%). The results of the TGA tests conducted in air are illustrated in Figure 13(c), (d). It can be observed, that the addition even of the lowest tested amount of clay (2% in 2RDS) causes a strong up-shift of the main decomposition temperature Tmax, from 343 C (neat polyMEA) to 398 C (2RDS). Further increasing clay content does not significantly shift Tmax: it rises only slightly from 398 C (2RDS) to 403 C (10RDS). This might indicate, that the principle of the stabilizing effect of the RDS nanofiller in air is the “barrier effect,” namely the hindrance of the oxygen diffusion into the sample (which would already be marked at 2 wt.% RDS). The lower-temperature decomposition process of the neat matrix is practically absent in air (in contrast to TGA in nitrogen), only a very small step (and dTGA peak), which is invisible without zooming the Figure 13(c),(d), can be found near 287 C. The temperature of decomposition onset Tonset (at 5% loss) is higher for the neat polyMEA matrix in air (305 C) than in nitrogen (286 C), while in case of the nanocomposites, weight loss in air starts at lower temperatures than in nitrogen (305 C–319 C vs 344 C–355 C). An interesting feature are the char residues temporary occurring in all samples (polyMEA and nanocomposites) in the range 400 C–550 C (see Figure 13(c)). Above 500–550, the char burns-off, thus causing small steps in the traces in Figure 13(c) and small flat dTGA peaks with Tmax values at 499 C–511 C (Matrix), 503 C (2RDS), 515 C and 556 C (4RDS) and 502 C–516 C (10RDS). Expectedly, no carbonized material survives in air at 700 C: the final residues are 0.44% (polyMEA), 2.69% (2RDS), 4.18% (4RDS), 9.57% (10RDS). The effects like the temporary formation of the char fraction in all samples, or the suppression of the small low-temperature decomposition step in polyMEA, as well as the higher Tonset of polyMEA, can all be attributed to radical reactions with oxygen which induce radical cross-linking: the temporary char fractions are even higher than the carbonized fraction obtained under nitrogen (a certain tendency to cross-linking was already observed during the synthesys – polymer chain length analysis discussed further above). On the other hand, oxidation-induced degradation clearly prevails in all samples at higher temperatures, and it also likely causes the somewhat lower Tonset value of the nanocomposites. If the TGA traces in nitrogen and air are compared (see especially Figure 13(b),(d)), it can be noted that the RDS clay stabilizes the nanocomposites much more distinctly in air (shift in Tmax in comparison to neat matrix). The neat matrix is much less stable in air than in nitrogen (Tmax = 343 C vs. 395 C). If the Tmax values are compared, which were obtained for the nanocomposites only (2RDS, 4RDS and 10RDS), in nitrogen and in air, it can be concluded, that these temperatures are practically the same in both atmospheres (396 C–400 C in N2 vs. 398 C–403 C in air). Tmax of the nanocomposites hence apparently is determined by the thermally-induced fragmentation process of polyMEA, and does not depend on the surrounding atmosphere. In Figure 14, the effect of the length of the elastic chains on the TGA traces is evaluated. In a simple case, the TGA traces in N2 or in air should not depend on the chain length, as far as the volume fraction of the filler and of polyMEA remains unchanged. However, as was discussed further above, the different chain length influences the fraction of the polymer topologically immobilized by the proximity of the rigid nanofiller. Hence, the effect of chain length on the TGA traces was evaluated for the sample, where it was the strongest in mechanical analyses, namely for 10RDS prepared at a 10 times lower concentration of both initiators. In Figure 14 it can be noted, that the effect of the increased chain length is small but visible: In nitrogen, and even more so in air, Tonset is shifted to lower temperature in case of the product with longer elastic chains, which also causes a visible change in shape of the dTGA peaks related to Tmax. Also the char fraction formed in nitrogen is smaller for the long-chain product, and the “temporary” char fraction burns-off more efficiently in air. All the mentioned small differences likely are caused by a much diminished fraction of immobilized polymer in the product with long chains (as was discussed further above: DMTA), which means a higher mobility and reactivity in degradation processes of the polyMEA phase, and possibly a less dense char phase formed from it.
4 | CONCLUSIONS
• physically cross-linked solvent-free nanocomposite elastomers were prepared, which possess an exceptionally high extensibility (up to 6000%), besides generally high mechanical properties (G' in rubber region between 1.5 and 40 MPa); • synthesis consisted of free-radical polymerization of 2-methoxyethylacrylate (MEA) in an aqueous dispersion of exfoliated montmorillonite clay; • the mechanical properties of the nanocomposites were varied in a very wide range by changing the concentrations of the radical redox co-initiator pair, at given constant nanofiller loadings (the filler strongly alters the product stiffness); this synthesis approach was generally aimed at obtaining longer elastic chains; • two synthesis approaches for obtaining longer polyMEA chains were tested: the lowering of the concentration of both co-initiators (ammonium peroxodisulfate (APS) and N,N0-tetramethyl ethylene diamine (TEMED)), as well as the lowering of the concentration of TEMED only (with constant concentration of APS at “standard” level); in the applied synthesis procedure, TEMED is practically quantitatively adsorbed on clay prior to polymerization start, so that polyMEA starts to grow from the clay platelets; • both methods yielded promising products, if the softer nanocomposite elastomers with low clay contents were optimized; in case of the stiffer elastomers with high clay content, the reduction of the concentration of both co-initiators, yielded similar but always better products than the other method; • the tested chain length extension made possible to increase the elongation at break of the studied nanocomposites by up to six times, to tremendously (and often simultaneously) increase their toughness (20–30 times), as well as to shift their tensile curves between “plastic-like elastomer” and “simple elastomer” type; • in order to elucidate structure–property relationships, the nanocomposites were also characterized by thermo-mechanical analysis (DMTA) and TGA (thermal stability, elastic chains' immobilization), and their morphology was analysed by TEM and X-ray diffraction; • the nature of the physical cross-linking, the behaviour of the hydrated nanocomposites during their drying to the final product, as well as the properties of the dried products with the smallest clay contents indicate, that the polyMEA/calay nanocomposite elastomers are promising candidates for self-healing materials.
ACKNOWLEDGMENTS
The authors thank Ms. Jiřina Hromádková (electron microscopy), Ms. Eva Miškovská (SAXS) and Ms. Markéta Karbusická (TGA). The authors also thank the Czech Science Foundation, project Nr. 19-04925S for the financial support of this work. ORCID Beata Strachota https://orcid.org/0000-0001-5163-9304
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