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Understanding toughness in bioinspired cellulose nanofibril/polymer nanocomposites Alejandro J. Benitez, Francisco Lossada, Baolei Zhu, Tobias Rudolph, and Andreas Walther Biomacromolecules, Just Accepted Manuscript • Publication Date (Web): 15 Jun 2016 Downloaded from http://pubs.acs.org on June 15, 2016

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Understanding toughness in bioinspired cellulose nanofibril/polymer nanocomposites Alejandro J. Benítez, Francisco Lossada, Baolei Zhu, Tobias Rudolph, Andreas Walther*. DWI ─ Leibniz-Institute for Interactive Materials, Forckenbeckstr. 50, 52056 Aachen, Germany E-mail: [email protected]

KEYWORDS: nanocellulose, nanopaper, glass transition temperature, copolymer, toughness.

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Abstract: Cellulose nanofibrils (CNFs) are considered next generation, renewable reinforcements for sustainable, high-performance bioinspired nanocomposites uniting high stiffness, strength and toughness. However, the challenges associated with making well-defined CNF/polymer nanopaper hybrid structures with well-controlled polymer properties have so far hampered to deduce a quantitative picture of the deformation mechanisms, and limits the ability to tune and control the mechanical properties by rational design criteria. Here, we discuss detailed insights on how the thermo-mechanical properties of tailor-made copolymers govern the tensile properties in bioinspired CNF/polymer settings, hence at high fractions of reinforcements and under nanoconfinement conditions for the polymers. To this end, we synthesize a series of fully water-soluble and non-ionic copolymers, whose glass transition temperatures (Tg) are varied from – 60 °C to 130 °C. We demonstrate that well-defined polymer-coated core/shell nanofibrils form at intermediate stages and that well-defined nanopaper structures with tunable nanostructure arise. The systematic correlation between the thermal transitions in the (co)polymers, as well as its fraction, on the deformation mechanisms of the nanocomposites is underscored by tensile tests, SEM imaging of fracture surfaces and dynamic mechanical analysis. An optimum toughness is obtained for copolymers with a Tg close to the testing temperature, where the soft phase possesses the highest molecular mobility and cohesive strength. New deformation modes are activated for the toughest compositions. Our study establishes quantitative structure/property relationships in CNF/(co)polymer nanopapers and opens the design space for future, rational molecular engineering using reversible supramolecular bonds or covalent crosslinking.

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INTRODUCTION Cellulose nanofibrils (CNFs) have been emerging in the last decade as a renewable, high performance reinforcement due to their versatile processing into films, nanocomposites or fibers, ease of chemical functionalization, hybridization with functional nanoparticles, and most importantly for realizing materials uniting high stiffness, toughness and strength.1-10 In plants, the mechanical performance is linked to the size, orientation, and interactions of the hard reinforcing cellulose microfibrils and the soft hydrated lignin-hemicellulose matrix.11,

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Bioinspired nanocomposites aim at mimicking the order found in natural system and focus on ordered arrangements of hard and soft phases at high fractions of reinforcements and with precisely engineered energy-dissipation mechanisms.13-15 The interplay of the different length scales during harsh mechanical loading serves to minimize failure, to reduce crack propagation and aims for the implementation of synergistic mechanical properties combining high stiffness and toughness.11-13,

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Sophisticated control over length scales and interactions is crucial to

achieve outstanding mechanical performance. CNFs are challenging to implement as a reinforcement into highly reinforced, bioinspired nanocomposites using classical hydrophobic polymers, because the unlike polarities prevent ordered nanocomposite formation and efficient stress transfer. Hence the combination with water-soluble polymers in aqueous suspensions and subsequent film casting has been the route of choice for bioinspired processing of such nanocomposites.17-20 In such a colloidal preparation route it is important to consider the long-range electrostatic interactions between the (typically negatively) charged CNFs and the polymer of choice during nanocomposite formation.21-23 For instance, strongly interacting polymers that lead to flocculation with CNF will not allow to make ordered, nanostructured composites, and thus, quantitative structure/property relationships based 3


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on nano-/mesoscale structures and interactions are overruled by microscale structures (flocculates).6, 9, 23 This imposes great constraints on the selection of polymers to derive detailed structure/property relationships in order to progress fundamentally and substantially in the understanding of the behavior of such bioinspired, highly reinforced CNF/polymer nanocomposites. Most success in terms of understanding and toughness enhancements has been reported by combination of CNFs with non-ionic polysaccharides, such as hydroxyethylcellulose (HEC) or plasticized starch polymers.17, 18, 24, 25 In addition, literature contains materials showing singular, and often serendipitously discovered mechanical properties that cannot be understood based on classical nanocomposite theory. This includes for instance non-linear toughness enhancements by adding small amounts of solubilizing polymers or special block copolymer micelles to NFC.19, 20, 26, 27 In addition to the importance of polymers, we could recently demonstrate that multiple yield points can occur in pure CNF nanopapers. Such additional toughness gains are related to enhanced interfibrillar sliding overcoming additional structural elements, and can only be observed from colloidally well-dispersed CNFs.6 Considering these studies it becomes obvious that a detailed understanding of the mechanical properties of highly reinforced, bioinspired CNF nanopapers would greatly profit from a systematic approach, ideally varying the thermo-mechanical polymer properties (energy dissipation), polymer fraction (interfibrillar spacing and interfibrillar slippage), and at some point also changing in depth the CNF/polymer interface. The latter is for most water-soluble polymers however very strong as CNFs can make multiple hydrogen bonds. Due to the restrictions of bioinspired CNF nanocomposites, that is well-controlled structure formation and absence of flocculation at high fractions of reinforcements, commercial water-soluble polymers are hardly

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applicable because those are mainly based on high Tg materials, e.g. polyelectrolytes, polysaccharides. To address this problem and understand in depth how the fraction and thermo-mechanical properties of a polymer govern the mechanical response in biomimetic CNF nanocomposites, we here synthesize a wide-ranging series of eight water-soluble and non-ionic copolymers with widely different glass transition temperatures, ranging from –60 °C to 130 °C. We demonstrate that they can be integrated via polymer-coated core/shell nanofibrils into highly reinforced CNF/polymer nanopapers leading to defined interfibrillar periodicities in bulk. This allows us to link the macroscopic mechanical response of the highly reinforced CNF/polymer network in detail to the fraction and thermo-mechanical properties of the tailor-made copolymers and the nano- and mesoscale structures. Optimum levels of toughness can be identified and closely correlated to the polymer properties. The overall study establishes first quantitative structure/property relationships in such well-defined CNF/polymer nanopapers with controlled thermo-mechanical properties of the soft phase, and opens a predictive design space for future implementation of rationally designed covalent and non-covalent interactions. EXPERIMENTAL SECTION Preparation of Cellulose Nanofibrils (CNFs) A 1.4 wt% suspension of TEMPO-oxidized Kraft pulp oxidized under neutral conditions was set to pH = 9 with NaOH and homogenized in a microfluidizer MRT CR5 applying four shear cycles (1400, 3 x 1000 bar).1 The content of carboxyl groups is 0.44 mmol/g.

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Polymer Synthesis Poly(ethylene glycol methyl ether methacrylate) (EG, Mn = 475 g/mol) and N,Ndimethylacrylamide (DMAm) were copolymerized by free radical copolymerization at room temperature in 140 mL of Milliq water with a total monomer or comonomer concentration of 0.20 mol/L and redox initiation employing a 2:1 molar ratio of potassium peroxodisulfate (KPS) and N,N,N’,N’-tetramethylethylenediamine (TEMED) (all Aldrich, highest purity available), respectively. The concentration between [comonomer] and [KPS] was set to 1000 : 1. The initial molar ratios of [EG] : [DMAm] were 90/10, 80/20, 60/40, 20/80 and 10/90, respectively. Homopolymers were synthesized following the same procedure. After 3 hours, the reaction mixture was dialyzed against water to remove residual monomer and initiator, and freeze-dried. The comonomer ratio was determined by 1H NMR and the molecular weights by gel permeation chromatography (GPC) in dimethylformamide using poly(methyl methacrylate) calibration. The copolymers are abbreviated with EGxDMAmy, where the subscripts denote the molar composition of the final copolymer determined by 1H NMR. Atomic Force Microscopy (AFM) Atomic force micrographs were acquired using a NanoScope V AFM (Digital Instruments Veeco Instruments, Santa Barbara, CA) operating in tapping mode. The samples were obtained by dipcoating from diluted suspensions in water (ca. 0.005 wt%) onto freshly cleaved mica. Differential Scanning Calorimetry (DSC) The samples were encapsulated in aluminum pans and ultra-high purity dry nitrogen was used as an inert atmosphere for all tests in a PerkinElmer DSC8500 calibrated with cyclohexane, indium and zinc. DSC thermograms were recorded at 20 °C/min. The pans containing the samples were

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punctured and held at 160 °C for 3 min to erase all previous thermal history and remove water residues. A second scan spanned a temperature range from -80 to 160 °C and the Tg was taken at the point where the change in specific heat (Cp) is half of the change in the completed transition. The sample weight was approximately 20 mg. All measurements were repeated twice. Nanocomposite Preparation Appropriate amounts of a diluted CNFs suspension (pH = 9; adjusted by NaOH; 0.25 wt%)6 were added to different volumes of the EGxDMAmy copolymers solution (0.25 wt%) under rapid stirring to reach the final targeted compositions. The polymer content in the nanocomposites was 1, 5, 10, 20, 35 and 50 wt%. Then the CNF/EGxDMAmy suspensions were filtered using a 50 µm pore size mesh sieve to remove dust and other large particles, followed by passing the solution through a 20 µm pore size mesh. Finally, 60 mL of the suspensions were added to petri dishes at room temperature to give films of ca. 22 µm thickness. Dynamic Mechanical Analysis (DMA) ISO 527-2 type 1BB specimens of the nanocomposites (12 mm gauge length and 2 mm width) were prepared using a Tox Pressotechnik cutter. A TA Q800 device was used in tensile loading with strain of 0.07%, preload of 0.001 N, force track of 125%, and frequency of 1 Hz. First, the samples were held at 160 °C during 5 min to remove residual water molecules. Then, equilibrated at -100 °C for 3 min and scanned up to 200 °C at a rate of 5 °C /min. Two samples were measured for each composite. Tensile Tests Tensile tests were performed on a Zwick Test Control II Z005 with a 100 N load cell at 21 °C and 55% relative humidity. Samples were conditioned at 55 % relative humidity for 2 days. ISO 7


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527-2 type 1BB were tested at a strain rate of 1.8 mm/min (strain = 0.15 1/s) following ASTM 638-14 standard. The data present an average of 7 specimens. Scanning electron Microscopy (SEM) Cross section fracture surfaces were observed with a Hitachi S4800 FE-SEM at 5 kV. The samples were sputter-coated in a Leica EM ACE600 with a thin Au/Pd layer of ca. 8 nm. X-Ray diffraction (XRD) Diffractograms were recorded from 5 to 35° (2Θ) in Bragg−Brentano geometry mode with X-φ-Z stage at 240 mm (reflection) configuration using an Empyrean diffractometer from PANalytical. Data was recorded using Cu Kα (0.1541 nm) radiation from an Empyrean Cu LFF HR X-ray tube at 40 kV voltage and 40 mA current and using a PIXcel3D detector. The Segal crystallinity was calculated using the 2 0 0 reflection.28 Fourier Transform Infrared Spectroscopy (FTIR). FTIR spectra were recorded using a Thermo Nicolet Nexus 470 spectrometer with a smart split ATR single reflection Si crystal with a resolution of 4 cm−1. RESULTS AND DISCUSSIONS The starting point for the rational design approach to understand toughening and inelastic deformation in CNF-based nanopapers is the design of an energy-dissipating soft phase providing strong interactions with the nanofibrils and promoting large deformation while maintaining high strength. Previously, we and others observed that nanopapers undergo large deformations once they are soaked in water.6,

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This is due to water molecules acting as

plasticizer between the CNFs, reducing their interfibrillar hydrogen bonding and promoting

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multiscale deformation mechanisms such as pull-out phenomena of mesoscale layers and of individual nanofibrils. We initially hypothesized that the introduction of a polymeric toughening agent of high molecular weight and with a tailored thermal transition could enhance frictional sliding of nanofibril junctions without excessively compromising strength by providing sufficient chain mobility and molecular cohesion. Unfortunately, the detailed tuning of the thermal transition, that is the glass transition temperature, Tg, of a polymeric system suitable for homogeneous integration into CNF nanopapers is in fact not straightforward. The most important prerequisites for the selection of the (co)polymer are high degrees of water-solubility and the absence of strong, long-range polyelectrolyte attraction to ensure smooth integration into the nanofibrillar network without direct coagulation of the CNF dispersion. Both objectives are of paramount importance to deduce quantitative structure/property relationships, because ill-defined network structures, in particular formed by coagulation, can overcompensate the influence of the detailed molecular engineering of the fracture energy-dissipation properties of the copolymer. Furthermore, the copolymer should still form sufficiently strong interfacial adhesion to the nanofibrils in the dried state to allow efficient stress transfer. To meet these diverse requirements and achieve a tunable Tg, we synthesized eight water-soluble poly[(ethylene glycol methyl ether methacrylate)-co-N,Ndimethylacrylamide] (EGxDMAmy) copolymers with different molar ratios through free radical copolymerization (Figure 1a). Those copolymers are non-ionic, amorphous in bulk, and have the potential to be widely tuned in the Tg (Tg (EG) = -62 °C and Tg (DMAm) = 130 °C). Additionally, they can form tight cohesion with cellulose in the bulk phase by hydrogen bonding.29, 30 We denote these copolymers as EGxDMAmy, whereby the subscripts denote the molar fractions in the final material as determined by 1H NMR. The resulting copolymers were characterized by gel

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permeation chromatography (GPC), 1H NMR, and differential scanning calorimetry (DSC; Table 1, Figures 1a-c and S1). As desired, the free radical polymerization produces high molecular weight copolymers with apparent weight average molar masses (Mw) in the range of 1630─2330 kg/mol (Table 1). The developed synthetic protocols deliver a set of copolymers with widely different compositions from pure poly(ethylene glycol methyl ether methacrylate) (EG100) to pure poly(N,N-dimethylacrylamide) (DMAm100). A monitoring of the incremental monomer consumption during the polymerization indicates a slightly faster integration of EG, which leads to a slight drift in the composition as the synthesis advances. However, a homogeneous bulk material with a single glass transition is obtained in all cases and there is no phase segregation (Figure 1; see also Figure S2 for DSC of a phase-segregated blend of EG100 and DMAm100 for comparison). Table 1. Molar compositions obtained from 1H NMR of our (co)polymers with their respective weight-average molecular weight (Mw) and molecular weight distribution (Ð) measured by GPC and glass transition temperature (Tg) obtained from DSC. Mw Tg (Co)polymer Ð [kg/mol] [°C] DMAm100 1648 1.35 129.8 ± 0.5 EG06DMAm94 1750 1.38 112.2 ± 0.8 EG11DMAm89 1684 1.42 104.5 ± 0.8 EG21DMAm79 1633 1.61 72.2 ± 0.4 EG54DMAm46 2327 1.30 25.6 ± 0.5 EG72DMAm28 1680 1.80 -2.9 ± 0.8 EG89DMAm11 2181 1.55 -61.5 ± 0.7 EG100 1846 2.60 -62.1 ± 0.6 The physical properties of the resulting (co)polymers are distinctly different as can be seen by macroscopic observation as presented in Figure 1b. The homopolymer DMAm100 (Tg = 130 °C). is brittle and can be cast into an amorphous and transparent film, while the pure EG100 is a very flexible and tacky polymer melt (Tg = -62 °C). As a comparison, the copolymer EG54DMAm46

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(Tg = 26 °C) is less flexible than EG100, but still not rigid enough to produce a solid film similar to DMAm100. Representative DSC heat flow traces are displayed in Figure 1c for various EGxDMAmy copolymers together with their respective homopolymers. The Tgs can be determined from the steps as second order phase transitions. A single thermal transition is observed for all (co)polymers, spanning the full temperature range from very low Tg (EG100) to very high Tg (DMAm100). The absence of other transitions further confirms a homogeneous melt and a smooth copolymerization, as the continuous shift is indicative of a mixing on a molecular level.

Figure 1. (Co)polymer synthesis and thermal characterization. (a) Synthesis of poly[(ethylene glycol methyl ether methacrylate)-co-N,N-dimethylacrylamide] copolymers (EGxDMAmy) using aqueous redoxinitiated free radical copolymerization. (b) Photographs of representative examples of the resulting (co)polymers. (c) DSC thermograms and (d) rationalization of the evolution of Tg using the models from Fox, Gibbs-di Marzio (GDM) and Brostow with different polynomial degrees. The molar ratios in (c) correspond to the notation of the EGxDMAmy in Table 1.

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Interestingly enough, a non-linear dependence can be observed upon plotting the Tg versus the weight fractions of the comonomers (Figure 1d). Hence the system deviates from the commonly used Gibbs-di Marzio model using the ideal volume of mixing, and the Fox equation considering the weight fraction by assuming a similar specific volume (Equation 1 and 2, respectively):31, 32 33-36

𝑇! = 𝑥! 𝑇! ! + 𝑥! 𝑇! ! ! !!

=

!! !!

!

+

!! !!

(1) (2)

!

where x1, x2, φ1, φ2, Tg1 and Tg2 are the weight fractions, specific volumetric fractions and the glass transition temperatures in Kelvin of a binary blend, respectively. The reason behind the deviation is the presence of non-negligible interchain interactions in EGxDMAmy. Hence, we applied a more elaborate model to allow a predictive interpretation of the copolymer systems. A recent model from Brostow et al.32 (Equation 3) explains deviations from classical Fox and Gibbs-di Marzio behavior by defining a quadratic polynomial centered in 2x1 ─ 1 = 0. The amount of necessary fitting coefficients ai (up to three: a0, a1 and a2), as well as the polynomial degree required to describe the experimental function Tg(x1), together with their magnitude, sign and associated deviation, are a measure of the complexity of the system. 𝑇! = 𝑥! 𝑇! ! + 1 − 𝑥! 𝑇! ! + 𝑥! 1 − 𝑥! [𝒂𝟎 + 𝒂𝟏 2𝑥! − 1 + 𝒂𝟐 2𝑥! − 1 ! ]

(3)

Figure 1d shows the regressions applied to our EGxDMAmy system and Table 2 summarizes the coefficients with their respective correlation parameter (R2) for the Brostow model with different polynomial degrees (constant, linear and quadratic). The Brostow coefficients are indicative of the type of interactions present in the system. A negative deviation (ai < 0) generally indicates

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the absence of strong intersegmental interactions and/or free-volume enhancements promoting chain segment mobility. In contrast, our EGxDMAmy system shows a positive deviation (ai > 0) originating from strong interactions (such as hydrogen bonding or dipolar interactions), which strengthens miscibility and translates into higher Tg. FTIR confirms such interactions by signal displacements in the copolymer EG54DMAm46 spectra for the -CONR2 groups and a slight shift for –COOR compared to their respective homopolymers (see Figure S3a-b). A reliable prediction of the Tg of the herein discussed (co)polymer system requires to at least consider the a0 and a1 parameter in the polynomial regression to reach high confidence with R2 ≈ 0.98 of the fitting parameters for future design. Further studies are however necessary to disclose a more precise origin and meaning of the fitting parameters and their relation to molecular details.32-34 Table 2. Fitting coefficients (in Kelvin) for different polynomial degrees of the Brostow model of the studied EGxDMAmy system (see Equation 3) . Polynomial Degree Constant Linear Quadratic

a0

a1

a2

R2

229 ± 40 238 ± 27 200 ± 32

221 ± 69 177 ± 65

234 ± 134

0.952 0.979 0.985

After having established a detailed characterization and prediction of the thermal properties, we next turn to the nanocomposite preparation by solvent casting of CNF/EGxDMAmy dispersions according to Scheme 1.

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Scheme 1. Formation of copolymer-coated core/shell nanofibers and their packing into CNF/polymer nanopapers through film casting.

The CNFs used in this study were prepared by TEMPO-mediated oxidation1 and subsequent microfluidization, yielding micrometer-long nanofibrils with diameters of 2.0-2.5 nm (atomic force microscopy, AFM in Figures 2a and S4a-b) and having a content of surface carboxylic acid groups of 0.44 mmol/g. Earlier we demonstrated that highest colloidal stability is important to reach CNF nanopapers of high toughness.6 Therefore, we adjusted the aqueous CNF dispersions to pH = 9 with NaOH (at 0.25 wt%), and added this dispersion with different volumes to the EGxDMAmy (co)polymer solutions (0.25 wt%) under rapid stirring to reach the final targeted compositions. The polymer content in the nanocomposite was fixed to 1, 5, 10, 20, 35 and 50 wt%, hence from very high to intermediate CNF concentrations and adhering to bioinspired nanocomposite design principles operating at high fractions of reinforcements. The addition of CNF to the copolymer follows considerations of classical layer-by-layer coating of colloids and minimizes potential bridging or even metastable flocculation if interactions between both

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components were too strong.37, 38 Note that this process has proven very useful to make welldefined, self-assembled nacre-inspired nanocomposites based on polymer-coated core/shell nanoclay.29,

39, 40

Macroscopically, all hydrocolloid dispersions are homogeneous and do not

show any flocculation. A representative atomic force microscopy (AFM) image of a CNF/EG54DMAm46 dispersion shows astonishingly well-defined core/shell nanofibrils, where the soft copolymer shell is clearly seen surrounding the CNF core (phase image in Figure 2c). The height image in Figure 2b allows determining the changes in dimensions. An increase in the average height (diameter) from pure CNFs (2.0-2.5 nm) to polymer-coated CNFs (3.0-3.5 nm) confirms the coating with the polymer. The polymer shell is even visible as lateral shoulders spreading onto the surface as pointed by the black arrows (Figure 2d and Figure S4 for height average determination). The presence of these core/shell building blocks is important as they predefine the nanostructure found in the final nanopaper to the highest possible extent. It also confirms interactions between the copolymers and the CNF even though weak and dynamic, as flocculation does not occur. Further evidence of interactions is found by FTIR analysis (Figure S3c-d).

After film casting, the mixtures develop transparent and homogeneous composite

nanopapers (Figure 2e).

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Figure 2. Characterization of CNFs, their copolymer-coated core/shell nanofibrils and subsequent nanocomposites. (a) AFM height image of CNFs, (b) AFM height image of CNFs after mixing with EG54DMAm46, and (c) corresponding phase image showing the nanofibrils coated by the soft copolymer. (d) Averages of five cross sectional profiles of uncoated and coated CNFs. Black arrows indicate areas where the polymer coating spreads on the surface of the mica (see Figure S4). All samples were prepared by deposition from dilute aqueous dispersions onto freshly cleaved mica (z-scale = 15 nm). (e) Transparent nanocomposite of CNF/EG54DMAm46 at 50 wt% of polymer content (thickness = 20 µm).

We evaluate the tensile mechanical properties in a two-step fashion by first comparing the effect of different fractions of the copolymer in individual sets of nanopapers, and later on comparing the effect of the copolymer properties at constant fractions of it. All measurements are done at 55% relative humidity (RH). Figure 3a-d depict the series of tensile curves for CNF/DMAm100, CNF/EG21DMAm79, CNF/EG54DMAm46 and CNF/EG100 at different polymer contents separately. All mechanical parameters of the nanocomposites are summarized in Table S1. Pure CNF nanopaper is

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presented as black curve for comparison, while increasing brightness in red indicates increasing content of the soft (co)polymer phase up to 50 wt%. In general, the inclusion of the (co)polymer within the interstitial space of the nanofibrillar network reduces systematically the Young’s modulus, yield point and tensile strength (E, σy and σb, respectively). This result is indicative of a well-controlled additive behavior in a nanocomposite, because the mechanical consistency of the (co)polymers is lower than pure CNF nanopaper.14 It overall originates from the high control of the CNF/copolymer nanocomposite structure by appropriate selection of the polymer system. Strikingly, the most interesting understanding can be gained by looking at the capability for inelastic deformation, and hence toughness, which is highly sensitive to the polymer content and moreover to the Tg of the copolymer. Looking at one extreme, DMAm100 with a Tg of 130 ºC, it can be observed that all nanopapers fracture at the same deformation εb = 5 - 6 % independent of the polymer content. Additionally, the primary yield point, where the transition from elastic to inelastic deformation occurs, stays on nearly the same level up to a polymer content of ca. 35 wt%, indicating a strong nanofiber network. The corresponding scanning electron microscopy (SEM) image of the cross section shows a semi-rough surface with large bundles and holes, reminiscent of an untoughened CNF nanopaper structure (Figure 3e; Figure S5 shows a SEM micrograph of pure CNF). The DMAm100 is a brittle polymer at ambient conditions with poor elasticity due to its high Tg, and therefore the rupture of the nanocomposite is majorly controlled by the CNF because the polymer phase lacks flexibility and does not aid in interfibrillar movement during deformation.

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Figure 3. Mechanical tensile properties for CNF/EGxDMAmy nanocomposites with different weight fractions of (co)polymer. Tensile curves for (a) CNF/DMAm100, (b) CNF/EG21DMAm79, (c) CNF/EG54DMAm46 and (d) CNF/EG100 with different polymer contents. Gradient arrows show the increase in polymer content and black arrows indicate a second yield point. (e-h) Cross section micrographs of fractured nanocomposites with 50 wt% (co)polymer, respectively. The micrographs show a clear transition of the deformation mechanism from (d) a relatively brittle fracture of CNFs and high Tg component to (f-h) pull-out phenomena (mesoscale layers and nanofibrils) stimulated by the increase in EG content and the reduced Tg in the copolymer. Scale bars are 5 µm.

The inclusion of the EG comonomer in the copolymer has a decisive effect on the nanocomposite behavior by reducing the Tg and allowing the polymer chains to undergo segmental relaxation at room temperature on the nanoscale and activating other deformation mechanisms on the mesoscale. While the reduction in strength is more abrupt in the depicted CNF/EGxDMAmy and CNF/EG100 nanocomposites compared to CNF/DMAm100, the strains at break show a continuous extension upon increase of the polymer content and reach to between 17 and 26% upon incorporation of 50 wt% copolymer. Interestingly, the primary yield points stay on a relatively high level for CNF/EG21DMAm79 with high Tg (72 ºC, Figure 3b) at low fractions of copolymer (5 – 10 wt%) as compared to the samples based on copolymers with lower Tg, and the CNF/EG21DMAm79 nanocomposites show some onsets of a non-linear, synergistic behavior (see also σy and σb values in Table S1). More importantly though, a new 18


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second yield point appears at ε = 6─8%. It is located in the same region as the fracture point of the nanocomposites with low polymer content, or which are based on high Tg nanocomposites CNF/DMAm100 or pure CNF nanopapers (black arrows in Figure 3b-d). The second yield point is characterized by a change in the slope of the strain hardening (Figure S6 and Table S2). This second yield point corresponds to motion of structures with different length scales, and can only be observed in presence of sufficient polymer that mediates the movement of the structures without excess stress concentration and thereby preventing breakage (around ca. 35 wt% polymer). Consistently, the corresponding micrographs in Figure 3f-h show a different cross section, in which pull-out of mesoscale layers and individual fibrils are now visible as strong deformation mechanisms. The transition is evident by comparison to the original CNF/DMAm100 (Tg = 130 °C). In a previous study,6 we observed a similar relationship between dried and fully hydrated nanopapers indicating that a second kind of motion is activated when the interfibrillar interactions are sufficiently weakened. The structures associated to this movement are possibly related to strong cooperative entanglements in mesoscale sheets. Of course, nanopapers soaked in water can achieve large deformation (εb ≈ 17 %) but drastically loose mechanical strength (σb ≈ 230 MPa at 60 %RH to 5 MPa when soaked in water)6, while our nanocomposites are able to retain at least 50% of the original strength at even higher levels of ductility. Hence, these results indicate a higher cohesive energy between the copolymer EGxDMAmy chains and the nanofibrils, which serve to toughen the materials through increased energy dissipation during large strain deformation. The second yield point in the CNF nanocomposites occurs at a certain polymer volume fraction and it is important to correlate it to the structural dimensions6, as approximated by an ordered lattice as a function of the composition and the diameter of the one-dimensional (1D) CNFs

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(Figure 2b). For simplicity, we assume an ideally ordered hexagonally packed phase of infinitely long 1D nanofibrils (here, CNFs elongated to infinity along their main axis).

Figure 4. Nanoconfinement in CNF/polymer nanocomposites. (a) Nanofibril separation as a function of the volume fraction of the nanofibrils, for different CNF radii as indicated within the figure. (b) Structural models rationalizing the geometrical constraints in oriented CNF/polymer nanocomposites assuming (gray) infinitely long cylindrical reinforcements embedded in a hexagonal fashion within a (blue) polymeric matrix. For calculations see supporting information (Figure S7).

The plot in Figure 4a demonstrates that a certain minimum fraction of polymer (>9 vol%) is necessary to prevent direct contact of the CNF, by first integrating polymer into the interstitial phase in a hexagonal lattice.41 While this is to some extent idealized for semiflexible CNFs with some amorphous domains, this fraction is independent of the diameter of the CNFs. On the contrary, the thickness of the interstitial matrix, d, between the CNFs depends very strongly on the diameter of the reinforcements. This is very important to consider when discussing mechanical properties of differently prepared and sourced CNFs, which may not have the same diameter. For instance TEMPO oxidation typically leads to the thinnest nanofibrils while enzymatic pretreatments or direct defibrillation may lead to thicker ones.42-44 The calculations exemplify however the level of nanoconfinement that a 1D CNF network imposes on the

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polymer phase in any case. It can be concluded that a minimum separation, d, is necessary for enhanced sliding of the CNFs and observation of a second yield point. Considering our nanocomposites (ρCNF ≈ 1.45 g/cm3, 6, 7 and ρcopolymer ≈ 1 g/cm3; CNF diameter = 4 nm by AFM), a second yield point is observed starting from a CNF volume fraction around 60 vol% (35 wt% polymer content), which corresponds to a minimum separation distance between two CNFs in an ideal hexagonal lattice of just around 2 nm. Hence the polymer exhibits strong nanoconfinement, which is fundamentally important to understand the final mechanical response. Next, we discuss in more detail the effect of the Tg on details of the toughening effect. Figure 5ab present a comparison of the tensile curves between CNF nanopapers and the nanocomposites at 35 and 50 wt% of polymer content, respectively. Furthermore, we calculate the work of fracture or toughness (Ut) from the area under the tensile σ-ε curves for all compositions (Figure 5c). While the relationship between Ut and polymer content for an individual set of nanopapers for one copolymer is fairly linear (orange-red vertical gradient), the overall trend inverts upon incorporation of EG units and when going from high to low Tg. A maximum in toughness occurs for intermediate Tgs. The inversion can be seen by comparing the series of CNF/DMAm100 and CNF/EG100. The high Tg polymer, DMAm100, leads to lower Ut at higher polymer fractions, because the materials is not able to undergo inelastic deformation due to lack of sufficient polymer dynamics. On the other side, EG100 with the lowest Tg shows the inverse behavior and higher toughness can be achieved upon incorporation of larger amounts of EG. Most importantly, however, an intermediate Tg near room temperature, as e.g. found in CNF/EG54DMAm46 (Tg = 26 °C) leads to the highest extensibility and to the highest toughness of all hybrid nanopapers. This is due to an optimum combination of molecular mobility and

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cohesive strength, and clearly confirms that the polymer dynamics need to be in an appropriate window to allow highest deformation in CNF hybrid nanopaper settings. To provide a balanced discussion on the toughness gains, we also report the changes in stiffness, Young’s modulus (E; Figure 5d), obtained from the slope of initial linear deformation, as a function of composition and Tg. Similarly, we find a consistent reduction in stiffness for increasing polymer content, as expected, exemplifying our achieved levels of predictive design. Additionally, we provide the dependence of E and Ut with the volume fraction in the SI (see Figure S8).

Figure 5. Glass transition temperature controls toughness. Tensile curves of nanocomposites with (a) 35 wt% and (b) 50 wt% polymer content, (c) work of fracture (Ut) calculated from the area under the stress-strain curves and (d) Young’s modulus, E, as a function of Tg of the copolymers and for the various compositions within sets of hybrid nanopapers.

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Then, to further evaluate the effect of (co)polymer properties on the deformation within the nanofibril networks, we use dynamic mechanical analysis (DMA) in tensile mode to link the thermal transitions in DSC to thermo-mechanical response accessible by DMA. Additionally, DMA benefits from higher sensitivity and allows to monitor relaxation processes even in highly reinforced bioinspired nanocomposites.45-48 Herein, we now focus on selected hybrid nanopapers with widely different Tg and at 50 wt% polymer content. Figure 6 shows the dependence of the storage modulus (E’) and loss modulus (E’’) against temperature for the pure CNF nanopaper and selected CNF/EGxDMAmy nanocomposites. Looking first at pure CNF, it is possible to observe a high and persistent storage modulus, E’ (black line), which can be attributed to the stiff percolating network and strong hydrogen bonds between adjacent nanofibers. The E’ modulus of pure CNF decreases slightly from 21 to 16 GPa spanning a temperature range between -100 and 120 °C in Figure 6a, followed by a more abrupt decrease to 13 GPa at 200 °C. The behavior is strictly related to the increase of free volume and mobility of the amorphous sites as the temperature rises.49 The correspondingly low values for the loss modulus of CNF nanopapers, E’’, are due to the high crystallinity within the nanofibrils of 77% (Figure S9) and the lack for energy dissipation in crystalline areas (Figure 6b).8, 50, 51 Finally, an increase in viscous dissipation occurs for high temperatures above 100 °C, originating from increased molecular motion in the amorphous parts inside the CNFs.

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Figure 6. Thermo-mechanical properties analyzed by DMA. (a) Storage modulus E’ and (b) loss modulus E’’ for CNF/EGxDMAmy nanocomposites with 50 wt% polymer content.

The DMA curves of the nanocomposites show softening by reduced E’ values with the inclusion of the (co)polymer, and as a function of the Tg (Figure 6a). This is in agreement with the trend of the Young´s modulus (E) in tensile testing curves (Figure 5d). Focusing first on the CNF/DMAm100 (50 wt%) hybrid nanopaper, it can be seen that the storage modulus, E’, shows a similarly constant, yet lower plateau up to ca. 120 °C. Thereafter, a rapid decrease occurs, which is reminiscent of the thermal transition of the polymer, that is its Tg at 130 °C, which contributes to viscous dissipation of fracture energy. This can also be seen by the concurrent increase in E’’ at high temperatures, and the occurrence of a broad peak. Given the sensitivity of DMA to details of thermal transitions, this transition peak is much broader compared to the second order phase transitions, Tg, obtained by DSC for the pure polymer.52 Moreover, a comparison of DMA analysis of CNF/DMAm100 and the pure homopolymer DMAm100 shows a substantial broadening of the thermal transition in this area within the nanocomposites as opposed to a sharp peak for the pure polymer (see Figure S10). Hence, the (co)polymers inside the interstitial sites of the nanofibril network show a wide distribution of segmental dynamics. This originates from the nanoconfinement and high level of interface. The broadness of the transition reflects different

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propensities for thermal relaxation due to the presence of adsorbed (immobilized) chain segments and rather unhindered, non-adsorbed segments, which overall can lead to a gradient distribution of relaxation times.53-55 Again the behavior of the nanocomposite with lowest Tg is fundamentally different. CNF/EG100 (50 wt%) shows a strong loss in E’ at very low temperatures, reminiscent of its Tg, and a continuous decrease of E’ across all temperatures. This steady decrease originates from the increased mobility of the polymer above its Tg. In accordance, the loss modulus, E’’, shows a wide and strong peak centered around Tg EG. Finally, at high temperature (ca. 150 °C), the loss modulus, E’’, displays a slight peak originating from the underlying CNF network. Based on the ranges obtained from pure CNF and its nanocomposites formed with the homopolymers, we separate the DMA spectra into different regions: Region I is dominated by EG100 transition in the range -75─-48 °C. Region III corresponds to the high temperature region characteristic of the pure CNF and contains features of the DMAm100. Region II presents the intermediate temperatures where the copolymers show major transitions in the DSC (Figure 1c). Table 3 lists the onset temperatures and peak temperatures for the different transitions of the nanocomposites evaluated in Figure 6. Table 3. Compilation of thermal transitions obtained from E’’ for the CNF/EGxDMAmy nanocomposites together with Tg obtained for pure (co)polymers by DSC. (Co)polymer DMAm100 EG54DMAm46 EG100

Tg,DSC

a

129.8 ± 0.5 25.6 ± 0.5 -62.1 ± 0.6 Pure CNFs

Region Ib Tα onsetb Tα peakb

Region IIb Tβ onsetb Tβ peakb

Region IIIb Tγ onsetb Tγ peakb

-75 -75

-56 -48

4 -

22 -

110 124 127

160 168 157

-

-

-25

19

110

-

(a)

Temperatures are reported in [°C]. (b) Error values were obtained from standard deviation and estimated to be ca. ±3 °C for all samples.

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Looking at the CNF/EG54DMAm46 sample, a high level of complexity is observed. It displays thermo-mechanical transitions at various positions, and overall an intermediate behavior in terms of the scaling of E’ and E’’ versus the temperature compared to the two nanopapers CNF/EG100 and CNF/DMAm100. The transition in Region I is related to the side group relaxation of the EG comonomer, which is based on a small oligomeric polyethylene glycol species itself, and which has some considerable impact on the thermo-mechanical properties Then, the important, yet broad transition of the copolymer main chain segment is obtained in the area of the Tg of the copolymer (26 ºC). This transition located near to the testing temperature allows the (co)polymers to have an excellent combination of molecular mobility and cohesive strength at this point, therefore providing a maximum ability for energy dissipation leading to high deformation and toughness in the nanocomposites. Finally, the transition in Region III relates to the relaxation observed for pure CNF, or potentially to some relaxation of tightly adsorbed chain segments, which also become dynamic at this point as the amorphous part of the cellulose experiences sufficient thermal energy for its transition. The overall response of the material is a consequence of the contribution of all relaxation modes, and the design can be tailored to specific application temperatures. CONCLUSIONS The knowledgeable chemical design of the thermo-mechanical properties of fully water-soluble and non-ionic copolymers and their smooth integration into CNF nanopapers allowed us for the first time to systematically study and comprehensively understand the occurrence and extent of toughening in CNF/(co)polymer nanocomposites at high fractions of reinforcements. There is a systematic correlation between the thermal transition temperature obtained in the pure soft phase of the synthesized (co)polymers and the deformation mechanisms, which the nanofibrillar

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network undergoes under mechanical stress. An optimum toughness can be obtained upon tuning the Tg of the copolymer to the service temperature of a given application. There are threshold values for the polymer content which need to be reached to observe this toughening, which is roughly located at 35 wt% of (co)polymer, and which corresponds to a small separation of the CNF nanofibers in an idealized network structure of only a few nm. The (co)polymers are strongly nanoconfined, but their well-balanced interactions allow a predictive design of mechanical performance. Based upon these tailor-made copolymers, an integration of supramolecular bonding motifs within the matrix or at the CNF/matrix interface may allow for the rational design of sacrificial bonds to enable synergetic mechanical properties by high levels of design in the future. Supporting Information Supporting 1H NMR for EGxDMAmy copolymers, DSC of EG100/DMAm100 blends, FTIR analysis, AFM height profile, tensile properties of CNF/EGxDMAmy nanocomposites, SEM micrograph of pure CNF, second yield point calculation, nanoconfinement calculations, toughness dependence with volume fraction, XRD pattern for pure CNF and DMA measurement for homopolymer DMAm100 are available free of charge via the Internet at http:// pubs.acs.org. Acknowledgments We thank Alexander Eckert and Khosrow Rahimi for help with the XRD measurements. We thank the BMBF for funding the AquaMat research group in the framework of the NanoMatFutur program. This work was performed in part at the Center for Chemical Polymer Technology, supported by the EU and North Rhine-Westphalia (EFRE 30 00 883 02). References 1. Isogai, A.; Saito, T.; Fukuzumi, H. Nanoscale 2011, 3, 71-85. 2. Klemm, D.; Kramer, F.; Moritz, S.; Lindström, T.; Ankerfors, M.; Gray, D.; Dorris, A. Angew. Chem., Int. Ed. 2011, 50, 5438-5466.

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Understanding toughness in bioinspired cellulose nanofibril/polymer nanocomposites Alejandro J. Benítez, Francisco Lossada, Baolei Zhu, Tobias Rudolph, Andreas Walther*. DWI ─ Leibniz-Institute for Interactive Materials, Forckenbeckstr. 50, 52056 Aachen, Germany E-mail: [email protected]

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Polymer

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