Nanocomposites from Clay, Cellulose Nanofibrils, and Epoxy with

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Nanocomposites from Clay, Cellulose Nanofibrils, and Epoxy with Improved Moisture Stability for Coatings and Semi-Structural Applications Lilian Medina, Farhan Ansari, Federico Carosio, Michaela Salajkova, and Lars A. Berglund ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.9b00459 • Publication Date (Web): 17 Apr 2019 Downloaded from http://pubs.acs.org on April 17, 2019

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ACS Applied Nano Materials

Nanocomposites from Clay, Cellulose Nanofibrils, and Epoxy with Improved Moisture Stability for Coatings and Semi-Structural Applications Lilian Medina†,ǁ, Farhan Ansari†,┴,ǁ, Federico Carosioǂ, Michaela Salajkova‡, Lars A. Berglund†,* †Department

of Fiber and Polymer Technology, Wallenberg Wood Science Center, KTH Royal

Institute of Technology, 10044 Stockholm, Sweden ǂDipartimento

di Scienza Applicata e Tecnologia, Politecnico di Torino, Sede di Alessandria,

Alessandria, Italy ‡Department

of Biosciences, University of Oslo, Oslo, Norway

*Corresponding author: [email protected] KEYWORDS: biocomposite, nanocellulose, mechanical, nanofibril, montmorillonite

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ABSTRACT

A new type of high reinforcement content clay-cellulose-thermoset nanocomposite was proposed, where epoxy precursors diffused into a wet porous clay-nanocellulose mat, followed by curing. The processing concept was scaled to > 200 µm thickness composites, the mechanical properties were high for nanocomposites and the materials showed better tensile properties at 90% RH compared with typical nanocellulose materials. The nanostructure and phase distributions were studied using transmission electron microscopy; Young’s modulus, yield strength, ultimate strength and ductility were determined as well as moisture sorption, fire retardancy and oxygen barrier properties. Clay and cellulose contents were varied, as well as the epoxy content. Epoxy had favorable effects on moisture stability, and also improved reinforcement effects at low reinforcement content. More homogeneous nano- and mesoscale epoxy distribution is still required for further property improvements. The materials constitute a new type of three-phase nanocomposites, of interest as coatings, films and as laminated composites for semi-structural applications.

INTRODUCTION Cellulose biocomposites based on wood cellulose nanofibrils (CNF) and a polymer matrix show potential for improved materials performance and new functions, compared with more conventional plant fiber composites.1 Such CNFs are mechanically disintegrated from cellulosic wood fibers, often after chemical or enzymatic pretreatment.2 Typically, CNFs are semi-flexible and available as hydrocolloidal suspensions where the CNF dimensions are 3-10 nm in diameter and 0.5-2.0 µm in length. Nanoporous CNF films can be prepared by filtration and drying3-4 or by casting from suspension so that an intermingled, random-in-the-plane CNF network is formed. 2 ACS Paragon Plus Environment

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Neat CNF films can show high modulus and strength,5-6 optical transmittance,6 low thermal expansion,6 and low oxygen transmission rate.7 If a water-soluble polymer is mixed into the colloid, it is possible to prepare nanocomposite films, as was done with parenchyma cell CNF,8 and high content wood-based CNF.9 One may also impregnate nanoporous films with thermoset precursors, and cure the films into CNF/thermoset composites.10-12 In earlier work, CNF was combined with montmorillonite clay nanoplatelets (MTM) so that a ductile “clay nanopaper” inorganic/organic hybrid material with highly oriented clay platelets was formed.4,

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MTM/CNF materials can have high ultimate strength,14 Young’s modulus15 and

improved fire retardancy.13, 16 They are also recyclable, since the interaction between clay and CNF is fairly weak. Like for neat CNF, the addition of a water-soluble polymer offers even greater flexibility and control in manipulating the parameters (e.g. interfacial interactions) to obtain specific performance. For instance, thin MTM/CNF composites with poly(vinyl alcohol) (PVA) were found to have excellent combination of strength and ductility. The resulting MTM/CNF/PVA composites, prepared by casting from water, were reported to have a modulus of 23 GPa, ultimate strength of 300 MPa and a toughness twice higher than that of natural nacre at 63 vol% fillers (CNF/MTM ratio of 2:1).17 However, CNF-based materials are very sensitive to moisture. The CNF surface is rich in hydrophilic hydroxyl groups that adsorb significant amount of moisture. The surface hydration of the fibrils has a negative effect on the composite performance, especially in terms of mechanical18 and barrier properties.19 Various approaches have been used in order to improve properties of CNFbased materials in highly moist or fully wet environment. Most studies focus on water-based processing of CNF, emphasizing CNF network formation. One strategy is to mix the CNF and an “active” component in the dilute state, followed by cross-linking mechanisms. Some examples of such “active” components include vitrimer nanoparticles activated by thermal curing,20 algal 3 ACS Paragon Plus Environment


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polysaccharides entangled within the CNF network and locked in the wet state with cation treatment,21 or physical cross-linking by chitosan during base treatment.22 One may also include intrinsically hydrophobic, but colloidal stable particles such as water-based latexes,23-25 or positively charged species such as quaternary alkylammonium.26 Water repellence can also be achieved by chemical modification after materials preparation.27-28 The general concept of nacre-mimetic materials for clay nanoplatelets-neat polymer systems was introduced by Kotov and co-workers, who used sequential deposition of clay and polyelectrolytes to build micron-thick films.29 In a later study, exceptionally high mechanical properties were reached with glutaraldehyde cross-linking.30 Walther et al. devised a novel core-shell approach, where a polymer, such as PVA, is adsorbed on the clay nanoplatelets, and non-sorbing polymer is removed by centrifugation.31-32 This method allows faster, and scalable filtration preparation of nacre-mimetic materials, and the location of the polymer matrix is tailored. The effect of clay platelet aspect ratio and relative humidity was studied and increased platelet aspect ratio improves modulus and strength.33-34 However, the moisture sensitivity of CNF-based materials is also valid for nacre-inspired clay-polymer nanocomposites in general, since they tend to be based on watersoluble polymers. Although chemical crosslinking is improving properties,30-31,

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the intrinsic

hygrophilicity of water-soluble polymers still leads to high moisture sorption. An interesting insitu polymerization approach36 was attempted for a clay-polymer system, but the absolute properties of the resulting composites were not so high. These materials are probably not suitable for demanding engineering applications (e.g. automotive), where high mechanical performance is required at elevated relative humidity. Another strategy to address moisture sensitivity is to combine CNF or clay with a thermoset resin matrix. This relies on the proven performance of e.g., epoxies, in terms of much better stability than water-soluble polymers under moist conditions. Amine-cured epoxies also form covalent 4 ACS Paragon Plus Environment

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bonds with CNF.37-39 At 15 vol% CNF in epoxy, the effects from moist environment was almost negligible, since the epoxy matrix strongly reduced CNF fibril hydration.37 For clay composites, the use of a thermoset resin is also beneficial although MTM hygroscopicity is a challenge. For instance, the equilibrium moisture content increased linearly with MTM content in MTM-EP, although the rate of diffusion was reduced by almost 50% at 6 wt% of clay.40 Another study showed that the modulus and strength of MTM-EP were reduced by almost 50% when relative humidity was increased from 24% to 98%. Interestingly, the same decrease was observed for this particular neat epoxy due to plasticization effects.41 In the present study, porous clay/CNF films are impregnated by EP precursors to form films, followed by nanocomposites curing. The EP functions as a polymer binder. Three constituents are used: fibrillar CNF, MTM clay platelets and an epoxy polymer matrix. CNF forms a continuous nanofibril network with MTM platelets preferentially aligned in-the-plane. Epoxy has a binder function, and reduces moisture sensitivity. Relationships between composition and nanostructure on physical properties are investigated. Mechanical properties, moisture sorption, gas barrier and fire retardancy properties are characterized. Compared with earlier clay-epoxy nanocomposites, the mechanical performance is much superior due to oriented MTM platelets and the strong reinforcement from the cellulose nanofibrils. In the automotive industry, typical “semi-structural” applications for sheets and molded structures made from glass fiber composites have a Young’s modulus in tension of 5-10 GPa and a tensile strength of around 100 MPa,42 which is significantly lower than for the present composites. These three-phase nanocomposites can be used as fire retardant semi-structural laminated composites, even at elevated relative humidity, and are prepared by scalable processing approach. The critical role of mesoscale (≈100 nm scale) organization is apparent, which is an opportunity for improved processing concepts, structural homogeneity and enhanced properties. 5 ACS Paragon Plus Environment


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EXPERIMENTAL SECTION Materials. The epoxy monomer used was a DGEBA (diglycidyl ether of bisphenol A) purchased from TCI Chemicals (Japan), with molecular weight 340.42 g mol-1 and an epoxy equivalent weight of 170 g eq-1. The curing agent was an aliphatic difunctional polyetheramine, Jeffamine D400, provided by Huntsman Holland B.V (The Netherlands), with a molecular weight of 400 g mol-1 and an amine equivalent weight of 115 g eq-1. Acetone (VWR, Sweden) was ≥ 99.5% pure and used as received. The CNF suspension was prepared by enzymatic pretreatment and high-pressure homogenization of a sulphite pulp (Nordic AB, Sweden), as reported elsewhere.43 After homogenization, the resulting gel had a concentration of 1.5-2.0 wt%. The CNF properties have been extensively characterized in earlier works, with typical length of 1-3 μm, mean diameter of 6.6 ± 3.3 nm and slight negative charge of 110 μeq g-1.44 The MTM suspension was prepared by shear-mixing a dilute (0.5 wt%) suspension of MTM powder (Na+ Cloisite, BYK additives, Germany, 2.86 g cm-3) in water, followed by ultrasonication for 5 min, and 30 min centrifugation at 4500 rpm to remove large aggregates. The stable, supernatant fraction of the MTM suspension was kept and the sonication/centrifugation steps were repeated until no aggregates could be observed after centrifugation, typically 3 times. The solid content of the final suspension was 0.3 – 0.4 wt%. The clay platelets dimensions measured in our earlier work are 1.1 nm thickness and about 120 ± 75 nm length/width.35 Preparation of nanocomposites. The CNF was first diluted to about 0.1 wt% in Milli-Q water and sheared with an Ultra-Turrax probe mixer (IKA T25, Germany) for 10 min at 12500 rpm. A predetermined amount of MTM suspension was added to the CNF suspension and sheared for an additional 5 min. The CNF/MTM colloidal mixture was then vacuum-filtered using an ultrafiltration membrane (Merck-Millipore) with pore size 0.65 μm. When filtration was finished, 6 ACS Paragon Plus Environment

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the wet cake (about 85 wt% water) was solvent-exchanged in an acetone bath for 2 days, changing the solvent 2-3 times per day. CNF/MTM mats with two different ratios of CNF/MTM – 3/1 and 1/1 – were prepared. At higher MTM content, the wet mat was too fragile and difficult to handle. The obtained CNF/MTM/acetone mat was then impregnated with epoxy. The impregnating solutions were prepared in concentration of 10, 30 and 50 wt% in acetone and mass ratio of about 7:3 between DGEBA and the aliphatic diamine (stoichiometric ratio as indicated by DSC). The acetone cakes were placed in the impregnating solution overnight with mild stirring. They were then removed from the solution, left to dry in air for about 30 min and cured in an oven for 2 hours at 90 °C and 2 hours at 120 °C. After curing, they were hot-pressed for 30 min at 90 °C and 200 kN (Fontijne Presses, The Netherlands) in order to have flat specimens for testing. Note that since the original filtered mats are the same, the thickness of the final composites was higher for high-epoxy content specimens and varied between ≈60 to 250 μm. Field Emission-Scanning Electron Microscopy (FE-SEM). The cross-section of the samples was imaged using a Hitachi S4800 FE-SEM (Japan). The typical operating conditions were 1 kV accelerating voltage and short working distance. We imaged (a) flat cross-sections prepared from ultramicrotoming and (b) fractured surfaces after tensile testing. The samples were coated with a thin layer (2-4 nm) of Pt:Pd using a Cressington sputter coater prior to observation. Transmission Electron Microscopy (TEM). TEM analysis was performed using a transmission electron microscope JEOL 1400Plus equipped with a Ruby camera, both from Japan and an energydispersive X-ray Spectroscopy (EDX) detector (Oxford Instruments, United Kingdom) and operated at 100 kV. The samples were first embedded in an epoxy-embedding medium (Sigma Aldrich, Sweden) and cured overnight at 60 °C. The samples were then ultramicrotomed with a diamond knife (Diatome, Switzerland) to thin sections of about 80 nm. To avoid a background noise for the EDX analysis, analytical (Beryllium) sample holder was used. 7 ACS Paragon Plus Environment


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Tensile testing. The uniaxial tensile properties of the nanocomposites were determined using a Universal Testing Machine, model 5944 from Instron, USA. The strain was measured using a video extensometer and the load was measured with a 500 N load cell. The gauge length was 22 mm, and the strain rate 0.1 min-1. Prior to drying and curing, the tensile specimens were cut to strips of 3-4 mm width and about 60 mm length and uniform thickness (60-250 μm depending on the sample composition). The tensile specimens were conditioned in 50% relative humidity (RH), 22±1 °C room for at least 40 hours before testing. For the high-humidity testing, the sample were conditioned at 90% RH. The reported modulus, strength, and strain values are averaged over at least 4 specimens. For comparative analysis, the tensile modulus of the composites was also estimated by a simple rule of mixtures (ROM) model:

𝐸𝐶𝑂𝑀𝑃 = 𝐸𝑁𝑃𝑉𝑁𝑃 + 𝐸𝑚(1 ― 𝑉𝑁𝑃)

Equation 1

Here, VNP is the total volume fraction of reinforcement (MTM + CNF), ENP is the experimental modulus of the neat clay nanopaper (only MTM/CNF, no epoxy) and Em the experimental modulus of the epoxy matrix. The purpose is to estimate the reinforcement efficiency for various MTM/CNF nanopaper structures as they are provided with an EP matrix binder. The estimate of the composite modulus (ECOMP) in Equation 1 predicts a linear relationship between ECOMP and the content of MTM/CNF nanopaper, VNP. All volume fractions are calculated assuming densities of 1.1 for the EP, 2.86 for MTM and 1.5 g cm-3 for CNF. Moisture sorption. The moisture content was determined as the increase in weight (MettlerToledo precision balance ± 0.1 mg, Switzerland) from a dry sample (2 days in oven at 105 °C) and after equilibration in 50, respectively 90% RH atmosphere.


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Oxygen Transmission Rate (OTR). Oxygen permeability measurements in humid condition (23°C at 50 or 90 % RH) were performed using an Extraperm apparatus (Extra solutions, Italy). Samples have been tested using an aluminum mask to reduce the exposed area (i.e. 5 cm2). The permeability was calculated from the OTR and thickness of the samples. Vertical flammability. The flammability was tested in vertical configuration on 60x15 mm2 samples. A 20 mm blue methane flame was employed to ignite the specimen from its short side (flame application time: 5s). The test was repeated 3 times for each formulation evaluating parameters such as burning time and final residue. Prior to flammability tests, all specimens were conditioned in a climatic chamber (48 h at 23 ± 1 °C and 50% RH). Thermogravimetric analysis (TGA). Thermogravimetric analyses were performed on a TGADiscovery TA Instruments (USA) under nitrogen or synthetic air atmosphere. The sample (approx. 10±1 mg) was placed in open platinum pans and heated from 50 to 800 °C using a heating rate of 10 K min-1.



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RESULTS & DISCUSSION Processing The nanocomposites were prepared by a “prepreg” approach similar to sheet molding compounds42 or processes in the aerospace industry for carbon fiber/epoxy composites. The first stage is different; a stable hydrocolloidal suspension of MTM platelets and CNF is filtered to obtain a wet MTM/CNF/water cake. The residual water is exchanged to acetone and the MTM/CNF/acetone cake is placed in an impregnation bath containing epoxy precursors dissolved in acetone. Epoxy precursors diffuse into the MTM/CNF/acetone cake. The impregnated cakes are finally air-dried to remove the solvent, then cured and hot-pressed, see Figure 1a. This preparation approach allows for high volume fraction of comparably well-dispersed nanoscale reinforcement (CNF and MTM) oriented in-the-plane. CNF can also be considered a “processing aid” providing mechanical integrity to the fragile wet MTM cake, and facilitating epoxy impregnation.

Figure 1. (a) Schematic processing method for nanocomposites preparation, and (b) SEM micrographs of ternary Cellulose-Montmorillonite-Epoxy composites fracture surfaces.


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The composition was controlled in two different ways. First, the ratio of CNF and MTM reinforcement was varied to obtain a CNF:MTM ratio of 3:1 or 1:1 in the “prepreg”, termed C3M1 and C1M1. Secondly, the amount of epoxy was controlled by resin precursor concentration in the impregnation bath (3 different compositions each for C3M1 and C1M1). The compositions and nomenclature of the six materials are summarized in Table 1. Table 1. Sample nomenclature, wt% and vol% of the different phases. Sample

“prepreg” CNF, wt%

MTM, wt%

EP, wt%

CNF, vol%

MTM, EP, vol% vol%

CME 35/35/31

C1M1

35

35

31

36

19

45

CME 24/24/52

C1M1

24

24

52

22

12

66

CME 13/13/74

C1M1

13

13

74

11

6

84

CME 48/16/36

C3M1

48

16

36

50

9

41

CME 30/10/60

C3M1

30

10

60

26

4

70

CME 15/5/80

C3M1

15

5

80

12

2

86

Morphology The morphology of the prepared materials was characterized in an attempt to describe the nanocomposite structure at several length scales. This is important since larger scale aggregates are highly detrimental to mechanical properties of nanocomposites. SEM images of fracture surfaces show a layered nanocomposite structure (Figure 1b), with a layer thickness below 100 nm. The distribution of the components within the in-plane layers appears fairly homogeneous at scales of around 200 nm. There was no sign of larger micro-scale aggregates on the fracture surfaces, but the layered structure with MTM-rich regions becomes apparent at higher reinforcement content (Figure S1). In Figure S1, even though no large aggregates are visible, the samples show thick layered structures (few micrometers) with through-thickness concentration fluctuations. Similarly, 11 ACS Paragon Plus Environment


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high-resolution TEM micrographs of the composite cross-section do not show any microscale MTM aggregates (Figure 2). However, the higher magnification cross-section images (Figure 2 middle and bottom) show regions of strongly different density (dark: MTM-rich and bright: organic-rich regions). It means inhomogeneity in the distribution of inorganic and organic phases up to the scale of ≈100 nm for MTM tactoid thickness, and as much as a few micrometers for polymer-rich regions containing CNF and EP. From images at the highest magnification (Figure 2c, f, i, l), it is concluded that individual MTM platelets (≈1.2 nm thickness) form dense tactoids (stacks of MTM platelets) of about 30 up to 100 platelets in thickness. One may also note that the samples appear more homogeneous at lower MTM content. In EDX analysis images, dark regions consistently showed higher silicon content, which is associated with MTM (Figure S2). These local differences in silicon content are more pronounced for higher MTM content nanocomposites (larger for C1M1 compared with C3M1), and increase when the total volume fraction of reinforcement is increased. It is difficult to distinguish CNF from epoxy since they have similar contrast in TEM. Note that the composites show a “wavy” oriented cross-section (Figure 1b and Figure 2). The origin of this waviness is not clearly identified but may be related to shrinkage effects. A similar waviness was observed in Kotov’s original nacre article.29 The most likely explanation for the layered structure with “loose” clay tactoids at the 30-100 nm thickness scale is that the epoxy diffusion through the wet MTM/CNF mat is non-ideal. Although homogeneous epoxy impregnation has been achieved for pure CNF mats,37-38 clay platelets are impenetrable barriers. In earlier studies on impregnation of pure MTM clay layers by epoxy,45 it was found that platelet-shaped, hydrophilic clays of high aspect ratio can hinder epoxy transport. One may speculate that this would lead to epoxy-rich regions between MTM platelets. For instance, the light, high contrast areas in Figure 2c may correspond to epoxy-rich regions. The presence of clay-rich and clay-depleted layers in the filtered mat is unlikely, since the same MTM/CNF mat 12 ACS Paragon Plus Environment

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impregnated with higher epoxy content appears much more homogeneous (e.g. Figure 2 d-f and jl). Recent work on MTM/CNF filtered mats, using synchrotron X-ray scattering characterization method, was confirmed to show highly uniform nanostructure down to < 10 nm.15 Hydrophilic clay platelets appear to limit homogeneous epoxy precursor transport, so that micro- and nanostructures become inhomogeneous. Polymer-rich regions (CNF/EP mixtures) may extend for a few micrometers in both thickness and width directions, see Figure 2 and Figure S1. Modification of the preparation procedure could possibly improve the mesoscale homogeneity of the materials.

Figure 2. High-resolution TEM micrographs for (a-c) CME 35/35/31, (d-f) CME 48/16/36, (g-i) CME 24/24/52 and (j-l) CME 30/10/60 composite. Scale bars are 1 μm, 200 nm and 100 nm from top to bottom. 13 ACS Paragon Plus Environment


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Mechanical properties The stress-strain curves of neat epoxy and the composites are presented in Figure 3 a-b. The composite curves show phenomenologically distinct elastic and plastic regions, and the yielding transition is apparent as a “knee” in the curves. This is typical for CNF network films; for ductile polymer matrices, this CNF behavior is often retained in corresponding CNF-polymer matrix nanocomposites, especially when the CNF content is high.46 It is interesting to note that this typical behavior is present even in composites prepared with high MTM content (C1M1 series, Figure 3b). Composites with higher CNF content in the initial MTM/CNF network (C3M1 series) show higher strain to failure (up to ≈ 6 %) compared with the C1M1 samples (≈ 2 % strain to failure). CNF network deformation therefore seems to be the primary mechanism for nanocomposites ductility, and limits the brittleness otherwise expected for high clay content composites. Composites with highest reinforcement content (CNF and MTM) show the lowest yield strength. If EP is viewed as a polymer binder, this low yield strength suggests that yielding in low EP content composites is related to interfacial debonding between reinforcing constituents. CNF/MTM interfacial adhesion may be weak in the present composites.15


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Figure 3. Mechanical properties of composites at 50% relative humidity. Stress-strain curves of (a) “C3M1” + epoxy composite and (b) “C1M1” + epoxy composite. (c) Modulus as a function of the reinforcement content (sum of CNF and MTM) and (d) Ultimate strength as a function of the reinforcement content (sum of CNF and MTM). The solid lines show fit to experimental results, whereas dashed lines in (c) are obtained from a simple rule of mixtures (Equation 1). The Young’s modulus (Figure 3c) and ultimate strength (Figure 3d) increased significantly with the addition of just small amounts of MTM/CNF reinforcement to the neat epoxy. The modulus and strength of high reinforcement content (CNF and MTM) composites can reach 18 GPa and 136 MPa respectively (CME 35/35/31). The modulus scales with the combined (MTM+CNF) volume fraction of reinforcement VNP, although not perfectly linearly as would be expected from simple 15 ACS Paragon Plus Environment


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micromechanics (Equation 1).47 The separate modulus effects from MTM and CNF are difficult to distinguish in the complex compositions and nanostructures. Nevertheless, higher MTM content tends to lead to higher modulus. At similar total reinforcement content VNP, the C1M1 composites (1:1 weight fraction CNF:MTM) exhibit a systematically higher modulus compared with C3M1 (3:1 weight fraction CNF:MTM). This is also illustrated in a three-dimensional plot of modulus as a function of MTM and CNF content (Figure S3), where MTM addition provides increased modulus. One may note that an identical MTM grade was added to the same CNF in a separate study of MTM/CNF nanopaper, and similar effect was observed. Young’s modulus was increased from 17 GPa for neat CNF to 28 GPa at 50 wt% of MTM.15 The ultimate strength does not show linear dependence on total reinforcement content. A small reinforcement content has a strong effect, so that the ultimate strength is increased from 32 MPa for neat EP to approximately 104 MPa for CME 15/5/80 and CME 13/13/74. At higher reinforcement content, CME 35/35/31 and CME 48/16/36, the strength increases marginally to ≈136 MPa. The ultimate nanocomposites strength does not vary strongly between C1M1 and C3M1 samples. This is related to the reduced ductility of the C1M1 composites, since ultimate strength depends on the post-yield strain-hardening behavior. Mechanical behavior of the composites can be interpreted based on the nanostructural observations discussed earlier. With lower MTM content (C3M1 series), the epoxy is more homogeneously distributed, and ductility is also better preserved. The ductility of C1M1 composites is not strongly influenced by VNP, even though the microstructure appeared more homogeneous at higher epoxy content. One may note that even though the composite nanostructure showed heterogeneity at the scale of around 200 nm, the mechanical properties of the composites at ambient conditions (50% relative humidity and 23 °C) are so high that the composites could be considered for semi-structural 16 ACS Paragon Plus Environment

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applications where injection molded or compression molded glass fiber composites are often used.42,

48

An earlier report analyzed the potential of CNF composites for semi-structural

applications and summarized the exceptional performance in terms of ultimate strength and ductility. Modulus is largely controlled by the modulus of pure CNF nanopaper.48-49 The present addition of MTM clay platelets to a CNF/EP nanocomposite allows a modulus as high as 18 GPa, while maintaining high ultimate strength of 136 MPa. To the best of our knowledge, the present modulus is the highest reported for clay-epoxy nanocomposites. This is primarily due to high content of MTM clay oriented in the plane in the composite structure and the lack of significant microscale aggregates. Figure 2 still shows that the in-plane orientation of the MTM platelets can be further improved. It is also likely that improved MTM/EP and MTM/CNF interfacial strength, as well as increased clay aspect ratio would further improve strength properties. Note the importance of selecting a ductile epoxy of low yield strength, in order to utilize the ductility of the CNF network. As a rough estimate of composite modulus (ECOMP), we are assuming a simple rule of mixtures, see Equation 1 in Methods section. It is assumed that ECOMP depends on ENP, the modulus of the neat MTM/CNF nanopaper, and the modulus of the EP, scaled with respect to their volume fraction. Based on Equation 1 it is possible to estimate the effective ENP in the composite. Dashed lines in Figure 3c present predicted ECOMP using the experimental modulus for neat EP (Em ≈ 2.5 GPa) and pure MTM/CNF nanopaper (ENP ≈ 28 GPa for C1M1 and 24 GPa for C3M1, see Table S1). The experimental modulus (ECOMP) matches reasonably well with the prediction (dashed lines) for the C3M1 composites. The experimental composites data are close to the theoretical ones, suggesting that the nanopaper potential was well utilized. For C1M1 composites, the experimental values are even higher than the predicted modulus, particularly at low (≈ 15 vol%) and intermediate (≈ 32 vol%) reinforcement content. One may speculate that impregnation of epoxy resin to the relatively 17 ACS Paragon Plus Environment


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coarse structure of filtered C1M1 prepreg with high MTM clay content improves stress transfer between phases (MTM, CNF and EP) so that the reinforcement effect is improved. For all composites, the EP impregnation is also likely to reduce porosity.

Moisture sorption and hygromechanical performance The mass increase due to moisture sorbed is presented in Figure 4b. The moisture content correlates directly with relative humidity and the volume fraction of reinforcement (MTM/CNF). All the composites adsorbed less than 3% moisture at 50% RH, but this increases significantly from 2.5 to 8.3 wt% at 90% RH. At similar volume fraction, composites with higher CNF content (C3M1 series) show slightly higher moisture sorption than C1M1. This is due to the higher moisture sorption of CNF fibrils, due to the high hydrophilic surface area. Moisture sorption appears higher at higher reinforcement content (i.e. it does not increase linearly with reinforcement VNP). This indicates strong favorable effects from EP addition.


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Figure 4. Mechanical properties of CME composites at 50 and 90% relative humidity, and 23°C. Stress-strain curves of (a) C1M1 with 16 and 55 vol% reinforcement, C3M1 with 14 and 59 vol% reinforcement, respectively. (b) Moisture sorption at 50 and 90% RH for all samples. The neat clay nanopapers have 7 and 12 wt% moisture content at 50% RH and as high as 15 and 21 wt% at 90% RH for C3M1 and C1M1 respectively. (c, d) Modulus as a function of reinforcement vol% for (c) C1M1 and (d) C3M1. The solid lines in (c) and (d) are fits to experimental results, whereas dashed lines are obtained from the rule of mixtures (Equation 1). Mechanical performance at high relative humidity (90% RH) can be used to analyze effects from nanostructure and interfaces in these ternary (three-phase) composites. A comparison of stressstrain behavior at 50% and 90% relative humidity (Figure 4a) shows that the modulus, yield strength as well as the ultimate strength are generally reduced when exposed to high humidity, since the moisture content in the material is increased. For instance, the modulus and ultimate strength for the composites with highest total reinforcement decreased from 18 GPa to 12 GPa and 139 MPa to 91 MPa (for CME 35/35/31). In contrast, the moduli of composites with low reinforcement content (CME 13/13/74 and CME 15/5/80) appear to be practically unaffected by high relative humidity (Figure 4). This is quite remarkable, since CNF-based composites are expected to perform poorly at high humidity, due to surface hydration caused by moisture sorption. The moisture stability of the low VNP composites may arise from the covalent reactions between CNF and the epoxides.38 At low reinforcement content, epoxy is well-distributed with lower moisture content, and most MTM platelets and fibrils are probably interacting with epoxy, so that properties at 50% RH are retained at 90% RH. Composites at high reinforcement VNP are more moisture sensitive due to inhomogeneous distribution of the epoxy resin and higher moisture content. A similar effect was reported in less complex CNF/epoxy composites,37 where moisture 19 ACS Paragon Plus Environment


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stability at low CNF content (15 vol %) was attributed to strong CNF/EP interfacial interactions due to CNF-epoxide reactions. At higher CNF content, incomplete fibril coverage and higher moisture content caused the mechanical properties to decrease strongly with humidity. The C1M1 composites show higher modulus and comparable strength as C3M1 at similar total reinforcement content. Surprisingly, at high relative humidity (90% RH), the strength of C3M1 composites are higher than C1M1 composites, see Table S1. This may be related to unique CNF/CNF interactions in the C3M1 composites that allow higher ductility due to interfibrillar sliding. Although yielding occurs at a lower stress in C3M1, the post yield strain hardening causes the ultimate strength to be higher. Such strong plastic strain hardening is not observed in the higher clay content C1M1 series. The case of C3M1 composites is particularly interesting with respect to experimental and predicted values. At 50% RH, the experimental moduli match closely with predictions from the rule of mixtures (Figure 4d). In contrast, the experimental moduli at 90% RH are higher than the predicted values. This suggests that EP limits the deteriorating effect caused by moisture in a pure C3M1 nanopaper, due to lower moisture content and strong effects from favorable CNF/epoxy stress transfer. The EP is itself moisture stable, but also reduces moisture adsorption by CNF and MTM, by interactions with accessible surface hydroxyls.38 A similar deviation (higher experimental moduli than predicted) is observed for C1M1 composites, albeit at both relative humidities (50% and 90% RH). For composites used in other than marine applications, the performance at elevated relative humidity is of key relevance. As an even more severe test, the composites with the highest fraction of reinforcement were also immersed in Milli-Q water for 15 h and subjected to mechanical tensile testing. The water uptake of these composites was about 18-20 wt%, and 15 h was sufficient to provide a steady-state water content in the composite, as shown in Figure S4. Although the 20 ACS Paragon Plus Environment

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properties were significantly reduced, as expected, it is interesting to note that the composites still showed an elastic modulus of 3-3.5 GPa and ultimate strength of 50-70 MPa.

Barrier properties The oxygen permeability of the composites are presented in Figure 5. The addition of MTM/CNF reinforcement resulted in considerable improvement in the oxygen barrier performance (lower permeability) compared with the neat epoxy. The permeability was reduced by more than 50% with just 15 vol% reinforcement to the epoxy, and continued to decrease with increasing reinforcement content. This trend is expected since highly oriented MTM platelets are known to form strong gas barriers by forming a long tortuous path for the diffusing medium.50 Moreover, CNF is also a good oxygen barrier due to its semi-crystalline structure.19 It is interesting to note from Figure 5 that the absolute reinforcement content is the most important parameter controlling oxygen permeability. Surprisingly, C3M1 composites form a roughly similar barrier than C1M1 composites at 50% RH, even though the MTM content is lower for similar reinforcement VNP. This might be related to the structure of the initial prepreg, where the dispersion of CNF/MTM reinforcement is more favorable in C3M1 composites. At 90% RH, the permeability is generally slightly increased, since the nanoscale components adsorb water. The absolute values of permeability of the present composites at 50% RH (in the range 0.5-2 cm3 mm m-2 day-1 atm-1) are much higher than reported in literature for 60 vol% MTM/epoxy composites (< 0.003 cm3 mm m-2 day-1 atm-1),45 possibly due to quite low MTM volume content (see Table 1). Similarly, the oxygen barrier reported for a 50/50 CNF/MTM “clay nanopaper” was higher at 50% RH (permeability: 0.045 cm3 mm m-2 day-1 atm-1). However, when the clay



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nanopaper was subjected to 95% RH, its permeability increased by 2 orders of magnitude.13 This clearly illustrates the benefit of epoxy impregnation for moist environment applications.

Figure 5. Oxygen barrier properties of the composites at 23 °C, 50 and 90% RH. Fire retardancy and thermal degradation properties The compositions of highest mechanical performance (high CNF + MTM reinforcement content/low epoxy content) have been tested for thermal stability in inert and oxidative atmosphere, and their potential for fire retardancy was further evaluated by vertical flammability test, see Figure 6. In both air and nitrogen, TGA curves show one main degradation step, with maximum weight loss at ≈350 °C (Figure 6a) that result in the formation of a thermally stable carbonaceous residue containing MTM. The presence of a single step indicates the presence of interaction between the thermally degradable constituents (i.e. CNF and epoxy), possibly through covalent bonds.38 Indeed, as reported in supporting information, a more thorough TGA analysis in N2 consistently shows one degradation step for all compositions (Figure S5) with Tonset values that gradually increase towards neat EP by increasing EP content. This suggests that the impregnation and curing process can positively delay the degradation of the composites. 22 ACS Paragon Plus Environment

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In air, an additional step, starting at 450°C, is observed and is related to the oxidation of the organic residue produced during the first step. For both air and nitrogen, the total residue increases with MTM content, as expected, since MTM marginally loses weight in N2 and in air (≈10-15 wt% loss at 800 °C) mostly due to water removal and dehydroxylation.51 In air, the total amount of residue coincides almost perfectly with total MTM content for both samples. There is thus very little organic contribution to total residue (≈4-5 wt%, Table S2). This is slightly surprising, since synergy effects on thermal degradation of MTM/CNF nanopaper have been previously observed, particularly in oxidative atmosphere.16 Three mechanisms were articulated: (1) the organization of MTM platelets in the plane, which limits oxygen diffusion and slows down oxidative pyrolysis or combustion (2) low thermal conductivity in the out-of-plane direction, due to high amount of MTM/CNF interface and (3) favorable degradation pathway of cellulose in the presence of MTM, where a thermally stable char is preferentially produced. In the present nanocomposites, the oriented structure is retained, but presence of epoxy results in the formation of polymer-rich regions (see Figure 2) that are more easily oxidized and furthermore, also reduces the total area of MTM/CNF interface so that the favorable cellulose charring effect is reduced. The resulting structure is not as well-organized as MTM/CNF materials and can only partially slow down the oxidation step, as clearly observable by the prolonged weight loss step that spans from 450 to 800°C. In order to investigate the potential for fire retardancy, vertical flammability tests have been performed on the nanocomposites with the highest fraction of reinforcement, and on neat epoxy as a reference. The neat epoxy sample quickly ignites when the flame is applied, and undergoes complete combustion with heavy smoke production, melt-dripping behavior and no residue. The nanocomposites, on the other hand, do not show melt-dripping and leave some residues after combustion, despite the flame spreading quickly and reaching the clamp, see Table S3. Such 23 ACS Paragon Plus Environment


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residues are non-re-ignitable, and are also dramatically expanded (> 500%) in thickness direction, see Figure 6b-c. This type of “intumescent” behavior is an important flame retardant mechanism and could even be improved by a better control of the nanostructure.52 The performance of EP containing nanocomposites in terms of flammability, is lower than expected from neat MTM/CNF nanopaper, where 30 wt% MTM was sufficient to observe selfextinguishing behavior in the same testing conditions.16 Here, the sample with 33 wt% MTM burns completely and leaves a comparably low amount of residue. This is likely related to the structure of prepared composites, which hinders MTM/CNF charring synergy, as already observed by TGA. Indeed, the dispersion of clay platelets is important to reduce flammability of clay/polymer composites, primarily by formation of clay barriers to volatiles.53 Based on earlier TEM images (Figure 2), the dispersion of MTM is better in MTM/CNF nanopaper than in the current MTM/CNF/EP nanocomposites.

Figure 6. (a) TG curve in nitrogen and technical air atmosphere (b) residue of CME 33/33/33 after vertical flammability and (c) residue of CME 50/17/33 after vertical flammability. Scale bars in image (b) and (c) are 1 cm. 24 ACS Paragon Plus Environment

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CONCLUSIONS MTM-CNF “nanopaper” shows high mechanical properties, but cannot be used long-term at high relative humidity (e.g. 90% RH) due to moisture sorption and disintegration of MTM-CNF adhesion. A scalable concept was therefore introduced, where EP precursors diffused into wet, porous MTM-CNF “cakes”, followed by curing. EP acted as a polymer binder and strongly improved phase interaction in moist state as well as stress transfer between phases. At 35 wt% MTM and 35 wt% CNF with 30 wt% EP, the Young’s modulus and ultimate strength were 18 GPa and 139 MPa at 50% RH. To the best of our knowledge, these are the highest values reported for MTM-EP composites. At conditions of 90% RH and ambient temperature, mechanical properties still reached E=12 GPa and an ultimate strength of 91 MPa, which is sufficient for many semistructural composite applications. The MTM-CNF-EP nanocomposites showed considerable ductility, due to the ductility of the nanoscale CNF network and the EP. For ultimate strength, CNF content was important because of the strain-hardening behavior of the CNF network. As EP was added, the excellent fire retardancy of MTM/CNF was compromised since favorable thermochemical degradation mechanisms (MTM-CNF synergies) were reduced. The materials also showed good gas barrier properties, due to high MTM content and favorable in-plane orientation distribution. Although layered MTM/CNF/EP nanocomposites have comparably well-dispersed components, with high macro-scale ductility, it is concluded that there is room for improved nanoparticle dispersion. TEM studies at multiple scales reveal clay-rich layers, in particular in high MTM-content compositions. Within each lamina, polymer-rich cross-sectional regions (CNF+EP) could reach typically 0.5-1 µm in thickness and up to 3 µm in length. Small-scale MTM-rich regions could have a thickness of 100 nm and lengths up to 1 µm.



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ASSOCIATED CONTENT Supporting Information. Additional TEM and SEM images, EDX elemental analysis, detailed tensile mechanical data, three-dimensional plot of modulus against MTM and CNF content, tensile and water uptake data after immersion in water, TGA curves and associated peaks, and flammability tests burning parameters. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author *Email: [email protected] (Lars A. Berglund), Tel.: +46-8-7908118, Fax: +46-8-7906166. Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ǁ These authors contributed equally to the work. Present Address ┴ Department of Materials Science and Engineering, Stanford University, Materials Science and Engineering, Lomita Mall, Stanford, CA, USA 94305 Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS The authors acknowledge funding from the Wallenberg Wood Science Center and the Swedish Foundation for Strategic Research (SSF) through the FiReFoam project (RMA11-0065). 26 ACS Paragon Plus Environment

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