Tuning the Nanoscale Properties of Phosphorylated Cellulose

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Functional Nanostructured Materials (including low-D carbon)

Tuning the Nano-Scale Properties of Phosphorylated Cellulose Nanofibril-Based Thin Films to Achieve Highly Fire-Protecting Coatings for Flammable Solid Materials Maryam Ghanadpour, Federico Carosio, Marcus Conny Ruda, and Lars Wågberg ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b10309 • Publication Date (Web): 27 Aug 2018 Downloaded from http://pubs.acs.org on August 28, 2018

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ACS Applied Materials & Interfaces

Tuning the Nano-Scale Properties of Phosphorylated Cellulose Nanofibril-Based Thin Films to Achieve Highly Fire-Protecting Coatings for Flammable Solid Materials Maryam Ghanadpour, †

†,*

‡

ξ

Federico Carosio, Marcus C. Ruda, and Lars Wågberg

†, §,*

Department of Fiber and Polymer Technology, KTH Royal Institute of Technology, 100 44

Stockholm, Sweden ‡

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

Viale Teresa Michel 5, 15121 Alessandria, Italy ξ

Cellutech AB, Greenhouse Laboratories, Teknikringen 38A, 114 28 Stockholm, Sweden

§

Wallenberg Wood Science Center at the Department of Fiber and Polymer Technology, KTH

Royal Institute of Technology, 100 44 Stockholm, Sweden *

Corresponding authors

Keywords: phosphorylated cellulose nanofibrils, nanocomposite, thin film, thermal stability, flame-retardant

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Abstract Ultrathin nanocomposite films were prepared by combining cellulose nanofibrils prepared from phosphorylated pulp fibers (P-CNF) with montmorillonite (MMT), sepiolite (Sep) clay or sodium hexametaphosphate (SHMP). The flame-retardant and heat-protective capability of the prepared films as casings for a polyethylene (PE) film was investigated. Heating the coated PE in air revealed that the polymer film was thoroughly preserved up to at least 300 °C. The P-CNF/MMT coatings were also able to completely prevent the ignition of PE film during cone calorimetry, but neither the P-CNF/Sep nor the P-CNF/SHMP coating could entirely prevent PE ignition. This was explained by the results from combined thermogravimetry infrared spectroscopy (TG-FTIR) which showed that the P-CNF/MMT film was able to delay the release of PE decomposition volatiles and shift its thermal degradation to a higher temperature. The superior flame-retardant performance of the P-CNF/MMT films is mainly attributed to the unique compositional and structural features of the film, where the P-CNF is responsible for increasing the char formation while the MMT platelets create excellent barrier and thermal shielding properties by forming inorganic lamellae within the P-CNF matrix. These films showed a tensile strength of 304 MPa and a Young’s modulus of 15 GPa with 10 wt.% clay, so that this composite film was mechanically stronger than the previously prepared CNF/clay nanopapers containing the same amount of clay. 1. Introduction During recent years, the use of hybrid inorganic-organic nanocomposite thin films/coatings as protection for flammable substrates in order to replace halogenated flame-retardant additives which have potential toxicity and inefficiency issues, has gained particular attention.1–5 Different substrates such as wood,4,6,7 cotton fabric,8 polyurethane foams9 and polymeric films5 have been 2 ACS Paragon Plus Environment

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coated either by direct attachment of thin films5–7 or by Layer-by-Layer (LbL) deposition of the flame-retardant components.10–12 Clay nanoparticles have been used extensively within the coatings as a green nano-filler mainly in the form of montmorillonite (MMT)13,14 or vermiculite (VMT).5 Synthetic or natural polymers have been employed as the continuous matrix for the preparation of nanostructured thin films characterized by a highly ordered distribution of clay within the structure. The inclusion of clay has been shown to be particularly unique in improving the gas barrier properties,4,15 flammability performance2,3,16 and mechanical strength17,18 of the nanocomposite coating, usually at an inorganic content above 30 wt.%.3,4,6 These nanoparticles contribute to the flame-retardant properties of the coated substrate mainly by reducing the rate of mass and heat transfer to and from the surface.4,19 Flame-retardant performance has also been achieved by incorporating a phosphorus-containing compound within the nanocomposite coating in order to improve the char formation of the polymer employed as the continuous matrix. 20–22 Due to the dehydration effect, in the presence of a carbon source, the phosphate groups can enhance char formation at high temperatures. The multicellular carbonaceous char produced acts as a physical barrier and thermally insulates the underlying substrate by limiting the heat transfer and preventing the release of combustible volatiles.23,24 Examples of these compounds are poly(phosphoric acid),23 poly(sodium phosphate),19,20 ammonium polyphosphate,25 phytic acid21 and sodium hexametaphosphate26 used in combination with different cationic polyelectrolytes to self-assemble ultrathin flame-retardant coatings by the LbL technique. The use of bio-based polymer matrices for the preparation of flame-retardant composite films is of great interest, since the use of renewable resources is necessary for a future sustainable society and since many of these compounds can act as carbon sources in fire-retardant compositions. In this respect cellulose, which is the most abundant biopolymer available on earth, is very



interesting, especially in the form of well-characterized cellulose nanofibrils (CNF)27,28 that can be used for the bottom-up tailoring of nanocomposites. CNF is conveniently disintegrated from the plant fiber cell wall and the nanofibrils have a square cross-section typically in the range of 25-100 nm2 and can be up to 1 µm in length.27,29 With their inherent excellent mechanical properties, high length-to-diameter ratio and high specific surface area, cellulose nanofibrils are industrially and scientifically very interesting.30 Previously, CNF has been used in combination with MMT to prepare nanocomposite films which have a nacre-like structure with high toughness and ductility7,18,31 and with significant oxygen barrier properties and flame-retardant performance.17,32,33 The films are prepared through an uncomplicated vacuum filtration of a water-based CNF/MMT colloidal suspension, leading to a specific structure, termed as brick and mortar, in which the clay nanoplatelets highly oriented parallel to the material surface are responsible for the unique set of properties. For example, the achieved structure creates a tortuous path for oxygen diffusion, thus maximizing the barrier effect.4,10 During combustion, the clay barrier hinders combustible volatiles from escaping from cellulose and it also catalyzes the char production of CNF, due to the presence of metal ions in its structure. Both effects contribute to the excellent flame-retardant performance of these composites.4,34 The literature on polymer flame-retardancy includes a vast number of studies on depositing flame-retardant thin films onto polymeric substrates and the subsequent characterization of the modified substrates.3,6,7,9,20 However, astonishingly few investigations have been conducted to relate the flame-retardation mechanisms to the properties of the thin films/coatings. Filling this gap of knowledge has been the initial focus of the present work, where we have used cellulose nanofibrils from wood fibers chemically modified by phosphorylation (P-CNF) in combination with a plate-like clay (MMT), a fibrous clay (sepiolite-Sep) or a phosphate-rich compound


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(sodium hexametaphosphate-SHMP) to prepare ultrathin (ca. 10 µm) flame-retardant nanocomposite films. Figure 1 shows a schematic presentation of the procedure used to prepare the nanocomposite films through a filtration/solvent casting technique.

Figure 1. Schematic illustration of the preparation of P-CNF based nanocomposite films, the constituents and SEM (scanning electron microscopy) images of the final film structures. The P-CNF has previously shown intrinsic flame-retardant properties as a result of the presence of phosphate groups in its structure, which enhances the char forming ability of the polysaccharide and it is expected that by combining the P-CNF with MMT, Sep or SHMP the flame-retardant ability can be significantly boosted. By a careful control of both the structural and compositional features of these films, it should be possible to identify the critical factors behind their flame-retardation mechanisms with respect to char-forming ability, gas barrier and



mechanical properties. To study the flame-retardant mechanisms, polyethylene (PE) substrates, which are highly flammable with no charring, were used and sealed between two P-CNF based films as schematically illustrated in Figure S1. The use of cone calorimetry in combination with thermogravimetry infrared spectroscopy (TG-FTIR) demonstrated that this sandwich structure could indeed be used to characterize the heat- and fire-protective capability of the films. The results showed that the P-CNF/MMT films had excellent flame-retardant properties compared to the rest of the used recipes since this composite prevented ignition of the PE film under the cone heat flux with a clay content of only 10 wt.%. Such behavior is attributed mainly to the ability of MMT platelets to form continuous inorganic nano-sheets within the P-CNF matrix with extremely high barrier properties and this in combination with the blocking of the release of volatiles, is responsible for the improved flame-retardant performance of the P-CNF-based thin films. The response of the sealed, sandwich structures to heating was further observed by using films recorded with a high speed camera. The mechanical properties of the nanocomposite films were also studied by tensile tests which indicated that the mechanical strength of the PCNF/MMT films was superior to that of the previously studied CNF/clay nanopapers. 2. Experimental section 2.1. Materials Di-ammonium hydrogen phosphate (ACS reagent grade) and sodium hydroxide (analytical grade) were supplied by Merck KGaA. Urea (ReagentPlus, ≥ 99.5 %), sodium hexametaphosphate (crystalline, 96 %) and sepiolite were supplied by Sigma Aldrich, and montmorillonite (Cloisite Na+) was purchased from BYK Additives & Instruments, Germany. Milli-Q water (18 MΩ) was used in all the experiments. The toluene (ACS grade) used for the phosphorylation reaction was supplied by Merck KGaA. The CNF used was prepared from 6 ACS Paragon Plus Environment

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phosphorylated fibers. A commercial never-dried softwood dissolving pulp (Domsjö Dissolving Pulp; Domsjö Fabriker AB, Aditya Birla, Sweden) with a solids content of 15 wt.% was used for this purpose. The phosphorylation was carried-out with (NH4)2HPO4 in the presence of urea using a Dean-Stark trap apparatus.35 The fibers were first suspended in toluene at a concentration of 2 wt.%, in a two-neck round-bottomed flask with the use of a magnetic stirrer. Thereafter (NH4)2HPO4 and urea were added to the suspension with the following molar ratio of anhydroglucose units (AGU) of cellulose chain to the reagents: AGU:(NH4)2HPO4:urea - 1:2:8 The flask was then connected to the trap. The suspension was heated to 150 °C, and the condensed water and toluene were collected from the trap in a beaker. A 10 min centrifugation step at 4,500 rpm was used to separate the collected water and toluene and the toluene was then poured back into the reaction flask. The heating at 150 °C was continued and the water/toluene collection-centrifugation procedure was repeated until the water has been completely removed from the reaction medium, which consisted of the initial water from the fibers and the water formed during the phosphorylation reaction. The fibers were then filtered from the suspension and solvent exchanged firstly to acetone and secondly to deionized water to completely remove the toluene. The phosphorylated fibers were washed with deionized water until the conductivity of the filtrate was less than 5 µS/cm. The bulk DSphosphate (degree of substitution of phosphate groups) of the modified fibers was 0.02 as determined by XPS analysis.29 In order to obtain a CNF gel, a 1 wt.% dispersion of the phosphorylated fibers was homogenized in deionized water using a high pressure homogenizer (Microfluidizer M-110EH, Microfluidics Corp.). A transparent gel (P-CNF) was achieved after a single pass through the large chamber (200 µm) and four passes through the smaller chamber (100 µm) of the homogenizer. The P-CNF dispersion



used for film preparation was obtained by diluting the P-CNF gel with Milli-Q water and sonicating the dispersion for 10 min at 40 % amplitude using a Sonics & Materials, Inc. Vibracell with a titanium probe (tip diameter, 13 mm). The dispersion was then centrifuged for 1 h at 4,500 rpm, in order to remove contaminants from the tip and aggregates and non-fibrillated fiber fragments, and the supernatant was collected as a colloidally stable P-CNF dispersion. The surface charge of the phosphorylated fibrils due to both phosphate and carboxylate groups was determined to be 1.34 meq/g using polyelectrolyte titration.29 The clay dispersions were prepared by dispersing either MMT or Sep in Milli-Q water using a high-speed disperser (Ultra-Turrax, IKA, Germany) at 12,000 rpm for 10-15 min. A procedure similar to that used for the P-CNF dispersions was used including ultra-sonication and centrifugation and the supernatant was collected as the MMT/Sep colloidally stable dispersion. SHMP was dissolved in Milli-Q water at a concentration of 0.1 wt.% by adjusting the pH of the solution to 9 using 1 M sodium hydroxide. 2.2. Methods 2.2.1. Nanocomposite film preparation. 50 ml aqueous composite suspensions of pure P-CNF, PCNF/MMT, P-CNF/Sep and P-CNF/SHMP were prepared at a concentration of 0.15 wt.%. The clay-containing suspensions were prepared by the addition of MMT or Sep to the P-CNF dispersion to a final composition of 10/90 wt.% clay/P-CNF. The suspension was then mixed with Ultra-Turrax at 15,000 rpm for 15 min, followed by 30 min degassing under slow stirring. The resulting suspensions were then vacuum filtered using a glass filter funnel fitted with a hydrophobic poly(vinylidene fluoride) membrane (HVLP, 0.45 µm, 90 mm diameter) from Merck Millipore Ltd. (Ireland), based on a previously described procedure.18 The films were dried under vacuum at 93 °C for 15 min. The same method was used to prepare the pure P-CNF 8 ACS Paragon Plus Environment

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film but with no addition of clay. To prepare P-CNF/SHMP films, the P-CNF dispersion and SHMP solution were mixed at pH 9.5 to a final composition of 10/90 wt.% SHMP/P-CNF. The suspension was then poured into a petri dish (85 mm diameter) and the film was formed by solvent casting. In this way, nanocomposite films with a final thickness of 10-15 µm were obtained. These are simply referred to as P-CNF, P-CNF/MMT, P-CNF/Sep and P-CNF/SHMP due to the similar contents of clay or oligomer in the composite film. 2.2.2. Sealed sandwich structure preparation. Polyethylene (PE) was used as the substrate for evaluating the protective ability and flame-retardant properties of the nanocomposite films. For this purpose, a three-layered sandwich structure was prepared by placing a piece of 50 µm thick PE film between two P-CNF based films. A small amount of P-CNF gel was used as the adhesive to stick the three layers together. The entire structure was dried under vacuum condition at 93 °C for 15 min (Figure S1). The sealed structures are referred to as PE-X where X represents the coatings used, i.e. P-CNF, P-CNF/MMT, P-CNF/Sep or P-CNF/SHMP films. 2.3. Characterization techniques 2.3.1. Scanning electron microscopy (SEM). The SEM imaging was performed with a Hitachi S4800 field emission scanning electron microscope (FE-SEM) and secondary electron images were obtained from the cross-section of the P-CNF-based films and residues from the heating element test. For imaging the films were sandwiched vertically between two complementary metal plates using conductive adhesive tape. The films were coated with a 5 nm thick platinum/palladium layer prior to imaging, using a Cressington 280HR high resolution Sputter Coater. 2.3.2. UV-vis measurement. The light transmittance of the P-CNF-based nanocomposite films was measured between wavelengths of 300 and 900 nm with a UV-2550 Spectrophotometer, 9 ACS Paragon Plus Environment


Shimadzu Corp. (Japan). The films were placed vertically in the sample compartment and the distance from the light source was 6 cm. 2.3.3. Thermogravimetric analysis (TGA). The thermal stability of the P-CNF-based films, MMT, Sep and SHMP were evaluated under nitrogen with a TGA/DSC 1 (Mettler Toledo, Leicester, UK). The specimens (approximately 10 mg) were placed in alumina crucibles and were heated from 50 to 800 °C in an inert atmosphere, at a heating rate of 10 °C/min. Parameters such as Tonset10% (temperature at 10 % weight loss), TMax (temperature at maximum rate of weight loss), the total and organic residue at 800 °C were obtained from the measurements. 2.3.4. Oxygen permeability. The oxygen-barrier properties of the films were determined at 23 °C and 0 % or 50 % RH using a MOCON OX-TRAN 2/20 equipped with a coulometric oxygen sensor. The 10-15 µm thick films were mounted in an isolated diffusion cell and were subsequently surrounded by a flow of nitrogen in order to remove the adsorbed oxygen from the samples. Except for the circular exposure area (5 cm2), the films were tightly covered with aluminium foil with an adhesive on their surfaces. The measurements were performed symmetrically, i.e. the same relative humidity was applied on both sides of the test pieces. 2.3.5. Mechanical testing. Tensile tests were performed for the films using an Instron 5944 (Instron, High Wycombe, UK) instrument equipped with a 500 N load cell. For each formulation, at least five specimens with a length of 20 mm, a thickness of 10-15 µm and a width of 5 mm were tested, applying a strain rate of 10 %/min. The films were conditioned at 23 ± 1 °C and 50 % RH for at least 48 h before the tensile test. 2.3.6. Combined thermogravimetry infrared spectroscopy (TG-FTIR). The uncoated PE and the sealed, sandwich structures (4x4 mm2 PE films sealed between two 7x7 mm2 P-CNF based films; 3 ± 0.2 mg) were placed in open alumina pans and heated under nitrogen from 50 to 800 °C at a 10 ACS Paragon Plus Environment

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rate of 10 °C/min in a Perkin Elmer thermobalance. The Tonset10% , TMax and residue at 800 °C were obtained from the TG and dTG/dT curves. The degradation products were analyzed by a Spectrum Two FT-IR spectrometer (Perkin Elmer, Waltham, MA, USA) coupled to the thermobalance; one spectrum being recorded every 6 sec. The transfer line was kept at 280 °C with a gas flux of 65 ml/min. 2.3.7. Heating element test. The reaction of the uncoated PE film and the sealed, sandwich structures to heating was further characterized by heating the samples up to 600 °C, using a flat ceramic heating element (d 10.8 mm/ 24 V/ 50 W/ 750 °C/ Button heater, Rauschert Steinbach GmbH, Germany). The change in structure of the samples during heating was recorded using a high speed camera (IDT N4M-S3) with a fixed focal length objective (Cosmicar 50 mm/F1.4). The samples were illuminated by a light emitting diode (LED) light source (IDT 7 LED) and the recording was performed from 300 °C to 350 °C, at a frame rate of 30 frames per second and a shutter speed of 8 s. 2.3.8. Cone calorimetry. The combustion behavior of the uncoated PE film and of the sealed, sandwich structures (a 50x50 mm2 PE film sealed between two round P-CNF based films with a diameter of 75 (P-CNF and P-CNF/clay) or 85 mm (P-CNF/SHMP)) was investigated by cone calorimetry from Fire Testing Technology, FTT, West Sussex, UK. The samples were tested in a horizontal configuration under an irradiative heat flux of 35 kW/m.28 Parameters such as Time to Ignition (TTI), peak of Heat Release Rate (pkHRR), Total Heat Release (THR) and Total Smoke Release (TSR) were evaluated from the measurements. The test was repeated three times for each formulation to ensure reproducibility. The samples were conditioned at 23 °C and 50 % RH for at least 48 h prior to the combustion test.



3. Results 3.1. Characterization of the P-CNF-based nanocomposite films 3.1.1. Morphology The morphology and optical properties of the films were evaluated and the results of these measurements are shown in Figure 2.

Figure 2. Optical and structural properties of the P-CNF based films: (a) digital photographs; (b) SEM cross-sectional micrographs at different magnifications; (c) UV-vis transmittance spectra of the films.


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Figure 2a shows the visual appearance of the P-CNF-based films. All the samples were very transparent and the text underneath the films is clearly visible. Figure 2b shows SEM crosssectional micrographs of P-CNF-based films at different magnifications. The films show a laminated cross-section and MMT platelets and Sep rods can be distinguished at higher magnifications as indicated by the white arrows. Despite the low content of clay or oligomer, the morphological differences among the composite films are quite apparent. The P-CNF/MMT showed a well-arranged lamellar structure (compared to P-CNF) in which the clay nanoparticles were mainly oriented parallel to the film surface. This structure resembles the nacre-like brick and mortar structure shown in previous studies of CNF/clay nanopapers.4,18,31 The arrangement of Sep nanorods shows no specific pattern, the clay particles being scattered within the P-CNF matrix. In contrast to the P-CNF sample, voids exist within the cross-section of the multilayers of the P-CNF/SHMP film. The voids can probably be ascribed to an aggregation of the phosphorylated fibrils during the long drying process during solvent casting.36 The light transmittance of the films over the visible range was further investigated by a UV-vis spectrometer (Figure 2c). The total light transmittance of the P-CNF/SHMP films was slightly higher than that of P-CNF film at 600 nm (85 % vs. 83 %), probably due to the lower thickness of the hybrid film. Since the P-CNF/SHMP films have excellent transmission, the voids must have dimensions smaller than the wave length of light, i.e. of nanodimensions. The addition of clay nanoparticles led to no significant decrease in film transparency and both P-CNF/MMT and PCNF/Sep had a light transmittance above 75 %, at 600 nm. This may be due to the low clay content and the use of colloidally stable dispersions from which aggregates were removed through ultra-sonication followed by centrifugation.



3.1.2. TGA The thermal stability of the pure MMT, Sep, SHMP and the P-CNF-based films were investigated by thermogravimetric analysis in nitrogen and the results are illustrated by the TG and dTG/dT curves in Figure 3 and the thermogravimetric data shown in Table 1.

Figure 3. (a) TG; (b) dTG/dT curves of P-CNF-based films, MMT, Sep and SHMP in nitrogen.


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Table 1. The thermogravimetric data for the P-CNF-based films, MMT, Sep and SHMP in nitrogen. Sample

T(onset10%)

T(Max)

Residue at

Organic residue

°C

°C

800 °C (%)

at 800 °C (%)

P-CNF

236

266

38.8

38.8

P-CNF/MMT

236

246

43.7

38.8

P-CNF/Sep

236

259

42.3

37.6

P-CNF/SHMP

270

323

42.4

36.1

MMT

648

99

88.5

-

Sep

271

99

84.8

-

-

99

98.6

-

SHMP

Apart from the minor weight loss up to 100 °C (Figure 3a) due to moisture removal from the materials, the films show a single step weight loss which is characteristic of cellulose pyrolysis. The thermal degradation of cellulose in nitrogen takes place according to two competitive pathways. The first is the depolymerization of glycosyl units to volatile species, mainly ascribed to the evolution of levoglucosan, and the second that occurs in competition with the depolymerization is the dehydration and decomposition of the same units to thermally stable aromatic char which forms the final residue at the end of the test.8,37 Compared to unmodified cellulose, the presence of phosphorus in the P-CNF film structure led to a lower onset degradation temperature towards char formation and a considerable reduction in the cellulose decomposition temperature.29 As shown in Table 1, the Tonset10% value is the same for the P-CNF and P-CNF/clay films but according to the dTG/dT curves, the maximum rate of weight loss was 10 and 32 % lower for the P-CNF/MMT and P-CNF/Sep films respectively, probably caused by



the reduced degradation kinetics in the presence of clay.4 The greater decrease in the case of PCNF/Sep may be due to water release from sepiolite that has retarded the cellulose degradation.38,39 The amount of organic residue at 800 °C (calculated based on the TG data of pure MMT, Sep and SHMP) is the same for the P-CNF and for the P-CNF/MMT or Sep samples, which illustrates the limited effect of the clay on char production of P-CNF at higher temperatures. As a comparison unmodified cellulose has an organic residue of around 16 %.29 The Tonset10% and TMax values of the P-CNF/SHMP were significantly higher than those of the pure P-CNF film, while the amount of organic residue left from the hybrid film at 800 °C was slightly lower. The shift of the thermal degradation of P-CNF/SHMP films to a higher temperature may be due to the high thermal insulating properties as a result of the microexpansion of the hybrid film upon heating. This is shown and described in more detail in section 3.2.3. According to the DSphosphate of P-CNF, the number of moles of phosphate groups is 9 times higher in the P-CNF/SHMP film than in the pure P-CNF film. However, interestingly the effect of SHMP on the dehydration and char-producing degradation of P-CNF was insignificant, as indicated by the values of the organic residue. The decrease in the maximum rate of weight loss was also minor (10 %) when P-CNF was combined with SHMP, also showing that the production of volatiles in the hybrid film is as significant as in the P-CNF film. 3.1.3. Oxygen permeability The gas barrier properties of the P-CNF-based films were measured at 0 and 50 % RH and the results are summarized in Table 2. All the films showed rather low oxygen permeability values under dry and humid conditions. The moisture sensitivity of the pure P-CNF seems however to be higher than that of the composite films as its O2 permeability at 50 % RH was about four times higher than that at 0 % RH. The change in the gas barrier properties on addition of clay is 16 ACS Paragon Plus Environment

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remarkably different at 50 % RH, with P-CNF/MMT showing a substantial decrease (60 %) in oxygen permeability compared to the pure P-CNF while the P-CNF/Sep shows a 23 % increase compared to the pure P-CNF film. The higher oxygen permeability in the presence of Sep can be due to the hygroscopic nature of this clay38,39 that causes swelling at high RHs and therefore creates additional sites for the permeation of the oxygen molecules.15 On the other hand, based on the morphological observations, it appears that the highly ordered organization of MMT platelets formed impenetrable layers which created a tortuous path for oxygen diffusion,40 resulting in much better gas barrier properties than those of the P-CNF/Sep films. The O2 permeability of PCNF/SHMP films was 33 % lower at 50 % RH than that of the P-CNF samples, which may be due to the higher density of this hybrid film. Table 2. Oxygen permeability of the P-CNF-based films at 0 and 50 % Relative Humidity (RH). O2 permeation (cm3 mm/m2 day atm) Sample

0 % RH

50 % RH

P-CNF

0.008

0.03

P-CNF/MMT

0.015

0.012

P-CNF/Sep

0.022

0.037

P-CNF/SHMP

0.016

0.02

3.1.4. Mechanical properties The tensile properties of the P-CNF-based nanocomposite films were compared with those of the pure P-CNF film at 23 °C and 50 % RH. Figure 4 shows typical stress-strain curves of the films and the average values of the mechanical properties are summarized in Table 3. To make a comparison between the mechanical properties of the phosphorylated CNF-based films and the 17 ACS Paragon Plus Environment


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films prepared from untreated CNF, the average values of nanopapers from enzymatic CNF as reported by Henriksson et al. are also summarized in Table 3.41

Figure 4. Typical stress-strain curves of the P-CNF-based films. Table 3. Mechanical properties of the enzymatic CNF and P-CNF-based films. Sample

Density

Tensile Strength

Tensile Strain

Young’s Modulus

(kg/m3)

(MPa)

(%)

(GPa)

1140

158.5 ± 18.2

6.4 ± 1.7

10.7 ± 1.3

P-CNF

1253.0

246.0 ± 16.4

5.5 ± 1.5

10.1 ± 1.7

P-CNF/MMT

1484.5

304.8 ± 27.5

4.4 ± 1.3

15.2 ± 1.9

P-CNF/Sep

1512.2

235.4 ± 30.8

4.5 ± 1.1

11.5 ± 1.4

P-CNF/SHMP

1589.4

186.2 ± 12.2

6.8 ± 1.3

9.5 ± 0.4

Enzymatic CNF41

The addition of MMT has significantly increased the tensile strength and the Young’s modulus of the P-CNF based nanocomposite film although the increase in film density is minor, due to the low clay content. It is known that, when low contents of exfoliated clay are used, a good dispersion of clay within the polymer matrix is achieved13,15 which is probably the case in the


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present study due to the ultra-sonication and centrifugation procedures used prior to the preparation of the P-CNF/clay suspension. This has probably improved the MMT-P-CNF matrix interaction so that the fibrils and clay nanoparticles are well connected, creating strong interfacial adhesion between the components of the film. Due to the remarkably high stiffness of the MMT nanoplatelets (elastic modulus 270 GPa),13 the composite film has a significantly higher modulus than the P-CNF film.13,42 The P-CNF may also interact with Sep, mainly through van der Waals interactions and hydrogen bonding between the clay surface silanol groups and the phosphate groups along the CNF chain.43 The increase in the Young’s modulus upon addition of Sep is however less than in the case of MMT and the tensile strength of the P-CNF/Sep has remained unchanged in comparison with that of the P-CNF film. Compared with the enzymatic CNF nanopaper, pure P-CNF films show a 55 % increase in tensile strength. Moreover, the PCNF/MMT films show a 27 % higher tensile strength than that of the previously studied enzymatic CNF/MMT nanopapers containing the same amount of clay.15 Higher strength of the P-CNF based films may be attributed to the high surface charge of the phosphorylated fibrils which lead to a lower ordering of the molecules on the surface of the fibrils and therefore more efficient contact between the fibrils themselves and with the clay particles in the dry film.29 The minor decrease in the tensile strain of the P-CNF/clay films compared to that of the pure P-CNF concurs with the increased brittleness of the composite films upon addition of only a few weight percentages of clay, as was also shown by earlier studies.44,45 As far as the P-CNF/SHMP films are concerned, the presence of the oligomer is responsible for the decrease in tensile strength and Young’s modulus and the slight increase in tensile strain. The SHMP has probably acted as a plastisizer, allowing the fibrils to slide against each other before fracture, and this results in an increase in tensile strain.46 This is also in line with the creation of nanopores in the P-CNF/SHMP films. 19 ACS Paragon Plus Environment


3.2. Flame-retardant and heat-protective properties of the sealed, sandwich structures 3.2.1. Cone calorimetry The flame-retardant properties of the P-CNF-based films was studied by cone calorimetry during which the resistance of the sealed sandwich structures to an irradiative heat flux of 35 kW/m2 representing the early stage of a developing fire was investigated in comparison with that of the uncoated PE film. The HRR and TSR curves and the images of the residues at the end of the tests are shown in Figures 5a-c, respectively and the cone calorimetry data are reported in Table 4. Upon exposure to the cone heat flux, which leads to an increase in the surface temperature of the PE film, the sample underwent thermal degradation with a concomitant release of volatile species which caused ignition and the subsequent substantial flaming of the polymer (pkHRR 62 ± 14 kW/m2). The PE film burned and left no residue at the end of the test, as shown in Figure 5c. The pure P-CNF film coating resulted in some flame protection for PE as illustrated by the small increase in the TTI and the 56, 61 and 39 % decreases respectively in the pkHRR, THR and TSR values (Table 4). The best fire protection properties were obtained with the P-CNF/MMT coatings for which the sealed structure did not ignite at all during the 180 s exposure to the cone heat flux. Unlike the PE-P-CNF sample, which left a damaged residue with macro-scale breakages that acted as preferential escape ways for volatiles, the PE-P-CNF/MMT sealed structure showed a very coherent residue with almost no cracks in the composite structure (Figure 5c). The release of PE decomposition gases outside the sealed structure and the diffusion of oxygen inside were in fact effectively blocked due to the excellent gas barrier properties of the PCNF/MMT films. As a result, the concentration of ignitable volatiles did not reach the lower flammability limit and the sample failed to ignite and showed an extreme reduction in the HRR.4,6,10 Among the three individual PE-P-CNF/Sep samples tested by cone calorimetry, only


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one showed ignition. Such a performance can be ascribed to the small defects and breaks in some P-CNF/Sep films probably caused by the insufficient integrity of the film. The TSR value of the PE-P-CNF/Sep sealed structures was higher than that of the PE-P-CNF samples as shown in Figure 5b. The presence of P-CNF/SHMP coatings has also resulted in a significant increase in the TTI of PE film, as expected according to the delayed thermal degradation of the PCNF/SHMP films as shown by the TG results. However, the sample eventually ignites and shows a higher TSR than the PE-P-CNF sealed structure. This can be explained by the severe damage to the P-CNF/SHMP films during combustion as shown by the residue; after the process of charring and expansion, the formation of cracks caused by the low mechanical strength of the PCNF/SHMP films allow the escape of PE degradation products and sample ignition.

Figure 5. The behavior of the sealed, sandwich structures during cone calorimetry testing: (a) HRR curves; (b) TSR curves; (c) residues of the uncoated PE film and sealed structures at the end of the test.



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Sealed, sandwich structures prepared from P-CNF-based films containing a high clay content (30 wt.% MMT or Sep) were also examined by cone calorimetry in order to obtain information about the optimal clay content. The HRR and TSR curves of the PE-P-CNF/MMT(30) and PE-PCNF/Sep(30) together with those of the uncoated PE are shown in Figure S2 and the cone calorimetry data are summarized in Table S1. Although the PE-P-CNF/MMT(30) did not ignite, most of the parameters including pkHRR, THR and TSR were inferior to those of the PE-PCNF/MMT(10). Previous studies have shown that large-scale MMT agglomerates are present in the structure of the nanocomposite films when the clay content exceeds 10 wt.%. This may lead to a deterioration in the gas barrier properties, which in turn affects the flame-retardant performance as observed in this case.15,33 This phenomenon was not observed for the PE-PCNF/Sep(30) sealed structures which showed a small improvement in the flame-retardant properties compared to those of the PE-P-CNF/Sep(10), as illustrated by the increase in TTI and the decrease in pkHRR (Table S1). This may be simply due to the higher inorganic content of the coatings rather than to any change in the structural properties of the films. Table 4. The cone calorimeter data for the uncoated PE film and the sealed, sandwich structures. TTI

pkHRR

THR

TSR

(s)

(kW/m²)

(MJ/m²)

(m²/m²)

PE film

35 ± 7

61.7 ± 13.6

1.8 ± 0.4

27.9 ± 3.0

PE-P-CNF

46 ± 29

26.8 ± 5.8

0.7 ± 0.3

16.7 ± 6.3

No ignition

4.0 ± 3.8

0.1 ± 0.1

10.9 ± 4.0

No ignition / 18 s for the

19.0 ± 26.8

0.8 ± 0.9

15.7 ± 2.0

7.5 ± 10.6

0.3 ± 0.4

19.6 ± 7.8

Sample

PE-P-CNF/MMT PE-P-CNF/Sep

one which ignited PE-P-CNF/SHMP

126 ± 28


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3.2.2. TG-FTIR spectroscopy In order to explain the flame-retardant behavior of the P-CNF-based films, the sealed, sandwich structures were further studied by TG-FTIR spectroscopy to determine the thermal degradation and composition of the volatile gases produced by the samples. The TG and dTG/dT curves are shown in Figure 6a and b, respectively, and the related thermogravimetric data are summarized in Table 5. The thermal degradation of the uncoated PE film proceeded through a steep single step weight loss starting at 440 °C, during which extensive depolymerization of the PE occurred, leaving almost no residue at 800 °C. In contrast, the sealed structures show two main steps of weight loss; the first one due to the depolymerization/charring of the P-CNF films, as the TMax1 values are quite close to the TMax values of the single films determined from the TGA measurement (Table 1), and the second step probably due to depolymerization of PE. According to the dTG/dT curves, the rate of PE film volatilization was extensively reduced when it was coated with the P-CNF-based films, as indicated by a decrease in the maximum in the 470-490 °C range. In addition, the degradation of PE was shifted towards higher temperatures in the presence of the P-CNF/MMT coatings, as shown by the higher TMax2 value compared to the TMax1 of PE. The PE-P-CNF/SHMP sealed structure also showed a minor shift of PE degradation to higher temperatures, probably due to the higher thermal stability of P-CNF/SHMP than that of the rest of the films.



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Figure 6. (a) TG; (b) dTG/dT curves of uncoated PE film and of the sealed sandwich structures in nitrogen. Table 5. The thermogravimetric data for the uncoated PE film and the sealed sandwich structures in nitrogen. Sample

T(onset10%)

T(Max1)

T(Max2)

Residue at

°C

°C

°C

800 °C (%)

PE film

440

477

-

1.0

PE-P-CNF

261

271

475

25.6

PE-P-CNF/MMT

245

266

498

30.3

PE-P-CNF/Sep

250

260

472

28.3

PE-P-CNF/ SHMP

297

325

484

16.4

The FTIR spectra of the gases evolved from the PE sample and from the sealed, sandwich structures between 50 and 800 °C are presented in Figure 7. The results indicate that the evolution of gases from the decomposition of PE starts at ca. 400 °C, as shown by the appearance of the bands at 2926 and 2855 cm-1 which can be assigned to asymmetrical and symmetrical C-H


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stretching of hydrocarbons formed by pyrolysis of polyethylene 47 (the accurate wavelength values of the two main bands can be observed from the rotated 3D spectra of pure PE film). The signals show a maximum intensity at ca. 500 °C in accordance with the PE TG curve which revealed that maximum decomposition occurs at 477 °C. The same signals can be observed in the spectra of the evolved gas from the sealed structures, but they start appearing at a higher temperature than the PE spectra (450 vs. 400 °C) and they show a reduction in signal intensity up to 800 °C. The intensity ratio of the band at 2926 cm-1 at 800 and 500 °C shows that in the case of the uncoated PE, the polymer is already mainly degraded and converted to gaseous species at 500 °C (ratio 0.1) but then when coated with a P-CNF-based film, some of the volatilization still occurs at temperatures above 500 °C (ratio 0.2). These changes, together with the intensity profile of the 2926 cm-1 band shown in Figure 8, indicate that there is a delay in the release of PE decomposition products as a result of the barrier effect of P-CNF-based coatings. Such a delay is particularly apparent in the case of PE-PCNF/MMT samples for which the maximum intensity of the C-H band was observed about 3 minutes later than that of the uncoated PE, as indicated by the intensity profiles.



Figure 7. FTIR spectra of the gases evolved from the pure PE film and from the sealed, sandwich structures as a function of temperature.

Figure 8. The intensity profile of the C-H band (2926 cm-1) as a function of time for PE and for the sealed, sandwich structures.


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3.2.3. Heating element test To further understand the heat-protective mechanism of the P-CNF-based films, the uncoated PE and the sealed, sandwich structures were heated by a ceramic heating element in air and the reaction of the samples during heating from 300 °C to 350 °C (150 °C to 200 °C for the pure PE film) was recorded using a high-speed camera. Videos of the behavior of the PE film, PE-P-CNF, PE-P-CNF/MMT, PE-P-CNF/Sep and PE-P-CNF/SHMP sealed structures can be found in the Supporting Information, Video S1-5. Snapshots taken from the recordings are included in Figure 9. Upon heating, the PE film starts to melt at around 150 °C and shrinks drastically as shown by Video S1 and the residue image in Figure 9a. The reaction to heating is however completely different when PE is coated with either the P-CNF or the composite films; where surprisingly PE films have been completely preserved up to at least 350 °C (Figure 9b-e). The slow-motion recordings in Video S2-5 show what happens during the heating of the sealed structures. There is basically an extensive expansion of the upper coating, probably due to the initial volatile release from P-CNF as already demonstrated by TG measurements. Despite a macroscopic expansion, the P-CNF-based coatings remain un-cracked up to 350 °C, and are able to encapsulate volatiles generated by PE degradation and to create an effective thermal insulation so that the melting of the PE film is delayed.33 The macro-expansion of the P-CNF and P-CNF/Sep coatings were more extensive than that of the P-CNF/MMT and P-CNF/SHMP coatings, as can be seen in the recordings and residue images in Figure 9. This indicates a better thermal insulation of the latter films which explains the larger temperature drop along the lower coating leading to a lower degree of decomposition and formation of volatiles. The phenomenon is particularly apparent in the PE-P-CNF/MMT sealed structure which, instead of visible expansion, showed the formation of macroscopic bubbles during heating (Video S3). It is further shown that the PE-P-CNF/SHMP



sealed structure charred only partially and that some of the coating was well preserved up to 350 °C, which is in agreement with the TGA studies showing a higher onset degradation temperature for the P-CNF/SHMP films than for the rest of the samples.

Figure 9. Digital photographs of the PE film and sealed, sandwich structures before and after heating with the heating element, together with snapshots from the recordings: (a) PE film; (b) PE-P-CNF; (c) PE-P-CNF/MMT; (d) PE-P-CNF/Sep; (e) PE-P-CNF/SHMP. The final residues, which consisted of sealed structures heated up to 600 °C, were further investigated by SEM imaging. The cross-sectional micrographs of the P-CNF, P-CNF/MMT, PCNF/Sep and P-CNF/SHMP film residues are shown in Figure 10. Compared to the fresh films 28 ACS Paragon Plus Environment

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(Figure 2b), the heated P-CNF and P-CNF/Sep film have shrunk considerably in thickness; from 8.1 ± 0.1 to 4.6 ± 0.03 µm and from 8.8 ± 0.1 to 4.4 ± 0.3 µm, respectively, which may be due to substantial charring in the case of the pure P-CNF film29 and to the degradation of the P-CNF matrix in the P-CNF/Sep film which has led to an entangled network of Sep nanorods as shown in the higher magnification image in Figure 10c. On the other hand, the P-CNF/MMT and PCNF/SHMP films have expanded more than 2 times (from 9.4 ± 0.2 to 21.2 ± 1.0 µm) and 5 times (from 7.1 ± 0.1 to 37.2 ± 3.8 µm), respectively. The change from a parallel to a more wavy orientation, as a cause of degradation of the intercalated P-CNF has probably caused the substantial expansion of the P-CNF/MMT film.33 In contrast, the expansion of the P-CNF/SHMP film may be related to the intumescent behavior induced by the presence of the phosphate-rich oligomer, as reported previously.26,48 The presence of bubbles on the cross-section microlayers which can be seen in Figure 10d is another indication of an intumescent system. The microexpansion and appearance of large voids along the cross-section of the P-CNF/MMT and PCNF/SHMP films in contrast to the dense residues of the P-CNF and P-CNF/Sep samples may be the reason for the improved thermal insulation of the former films which was revealed by the heating element test.2,4,33 In the case of P-CNF/SHMP films, this expansion is indeed responsible for the reduced rate of heat transfer and degradation, and it explains the higher onset degradation temperature of the films and the large increase in TTI of the coated PE during cone calorimetry.



Figure 10. SEM cross-sectional micrographs of P-CNF-based films heated with the heating element, at different magnifications: (a) P-CNF; (b) P-CNF/MMT; (c) P-CNF/Sep; (d) PCNF/SHMP 4. Discussion 4.1. Flame-retardant mechanisms of the P-CNF-based films in the sealed, sandwich structures The flame-retardant performance and the heat-protection ability of the P-CNF-based films shown by cone calorimetry, TG-FTIR spectroscopy, and the model heating element test on the sealed, sandwich structures, in relation to the film properties can explain the flame-retardant mechanism of the P-CNF-based films. According to the results, it is suggested that the two major flameretardant features of the films are: a) The charring ability of the films The intrinsic flame-retardant properties of the phosphorylated CNF are attributed to the high concentration of phosphate groups on the surface of the fibril aggregates, mainly attached to the cellulose C6 hydroxyl groups.29 As a result of this modification the formation of levoglucosan, a 30 ACS Paragon Plus Environment

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highly volatile monomer which is the 1,6-anhydro glucopyranose ring formed by glucose, is prevented during the pyrolysis and the main depolymerization mechanism of P-CNF is thus interrupted.49 The presence of phosphate groups can further improve the flame-retardant properties by catalyzing the dehydration of the phosphorylated cellulose towards char formation.8,37 For this reason SHMP, a phosphate-rich oligomer was mixed with P-CNF to investigate the flame-retardant performance in the presence of a higher concentration of phosphate groups. The TGA results for the P-CNF/SHMP samples showed however that the oligomer has a limited influence on the char-forming ability of P-CNF. Such behavior may be elucidated by reviewing the mechanism of the pyrolysis of phosphorylated cellulose. According to the pyrolysis mechanism proposed for flame-retarded cellulose by Kandola et al., the phosphoryl groups can be easily eliminated from the phosphorylated cellulose at temperatures below 400 °C.50,51 The dephosphorylation at C2 or C3 leads to the formation of double bonds inside the glucopyranose ring which can equilibrate with a carbonyl structure through keto-enol tautomerism as shown by the schematic illustration of the reaction routes in Figure 11a.52 These carbonyl groups can easily break into volatile acids, carbon monoxide and carbon dioxide at temperatures above 400 °C, causing substantial decomposition of the cellulose. On the other hand, the formation of conjugated double bonds between C5 and C6, as a result of dephosphrylation at C6, leads to a much lower concentration of carbonyl bonds and instead promotes the formation of char as demonstrated in Figure 11b.50,52



Figure 11. Schematic illustration of the reaction routes during pyrolysis of phosphorylated cellulose, rearranged from the pyrolysis routes suggested by Kandola et al. and by Aoki et al.51,52: (a) dephosphorylation at C2 or C3; (b) dephosphorylation at C6. During the pyrolysis of P-CNF/SHMP films, the phosphoric acid released from SHMP is probably able to further phosphorylate the P-CNF.48 However, it is expected that the phosphorylation occurs mainly on the C2/C3 hydroxyl groups since most of the C6 positions are already phosphorylated due to the rather high surface DSphosphate (0.2) of the P-CNF used here. As a result, in spite of increasing the number of moles of phosphates by 9 times in the composite compared to the number in the pure P-CNF film, the contribution to char formation is insignificant and volatilization of the P-CNF/SHMP is consequently as strong as for the P-CNF film. The proposed mechanisms thus strongly support the importance of the position of phosphoryl groups in the cellulose chain rather than their quantity for the flame-retardant performance of the phosphorylated CNF. It is further shown that the presence of metal ions in the structure of clay may enhance charring of cellulose.4,34 However the char formation of the P-CNF was not significantly affected in


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presence of clay, as revealed by the TG data of the P-CNF/clay films. This is another indication that the pure P-CNF has reached its maximum charring capacity due to phosphorylation of most of the available C6 hydroxyl groups. The thermally stable char produced by the P-CNF has glued together MMT or Sep in a hybrid protective structure where the effect of the employed nanoparticles has determined the barrier properties as discussed in the following section. b) The barrier properties of the films Among the recipes used in this work, the P-CNF/MMT films showed a delay in the release of PE decomposition gases and a shift of PE thermal degradation to a higher temperature as a result of their excellent gas barrier properties at high temperatures and the thermal insulation properties of the composite film which inhibited ignition of the coated PE during cone calorimetry. The low MMT content and its well dispersion in the P-CNF suspension may have assisted the regular orientation of the inorganic phase which contributed to reducing the gas diffusion across the composite film.31,53,54 This polymer/clay configurations also resembles the nano-structural brick and mortar organization,9,13 for which the z-directional thermal conductivity is reported to be extremely low.55 This characteristic, in combination with the excellent barrier properties, can result in strong thermal shielding effects, as were indeed observed for the P-CNF/MMT films.4,5 Sep was the other inorganic component used for the preparation of composite films, and one of the heat-protective mechanisms for this mineral may be the release of water from the clay at high temperatures to form monohydrated sepiolite and protoenstatite.38 The P-CNF/Sep films were nevertheless unable to completely prevent the ignition of the PE films under the cone heat flux, while the volatile release from the PE-P-CNF/Sep sealed structures was as great as that from the PE-P-CNF structures. Considering the morphological properties as being the main difference between the two types of clay, the inability of Sep nanoparticles, presumably due to their rod-like



shape, to form a continuous inorganic network within the P-CNF matrix and to act as a barrier to volatile/oxygen diffusion may be the main reason for the insufficient flame-retardancy of the PCNF/Sep films. Such structural difference has also led to the superior mechanical strength and stiffness of the P-CNF/MMT films in comparison with the P-CNF/Sep samples, where the continuous MMT lamellae have probably acted as a load-transferring phase. This is absent in the P-CNF/Sep system due to the random distribution of Sep nanorods within the P-CNF matrix.15 The mechanical performance can also affect the barrier properties and thus the flame-retardation of the thin films. High mechanical strength, mainly caused by efficient contact and a large number of interfaces between the film components,17,18 usually leads to good barrier properties due to the formation of a dense and impenetrable network.31,36 This is indeed the case with the PCNF/MMT films in which the optimal flame-retardant and gas barrier performance are associated with superior mechanical performance. The importance of the mechanical properties of the films was also clearly shown by the model heating studies in combination with the TG-FTIR measurements which clearly demonstrated the importance of the mechanical integrity of the films when heated both to prevent cracks and avoid the release of volatiles and penetration of oxygen and to withstand the increased pressure inside the sealed sandwich structures when volatiles are starting to evolve due to PE decomposition. It is important to stress that the sandwich structures were representative of the general situation when a thin protective film is used to provide fire retardancy, and these results are therefore of great general significance. It is also worth emphasizing that the excellent fire protection was achieved by the addition of only 10 wt.% MMT, to a great extent due to the intrinsic flame-retardant properties of the phosphorylated fibrils. Previous studies have shown, that in order to prevent ignition of the film during cone calorimetry, at least 30 wt.% of MMT is required in carboxymethylated CNF-based nanopapers.


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Thus, by employing P-CNF a reduction in MMT can be achieved with the advantages of lower density, higher optical transparency and greater ductility of the composite films.44,45 5. Conclusions Nanocomposite thin films were made by combining cellulose nanofibrils prepared from phosphorylated pulp fibers with 10 wt.% of MMT, Sep clay or SHMP in order to study the effects of structural and compositional features on the flame-retardation of P-CNF-based films. The heat-protective and flame-retardant properties of the films were studied in sealed sandwich structures, where PE was used as a substrate, laminated and sealed between two P-CNF-based films. Cone calorimetry and TG-FTIR showed the excellent barrier properties and flameretardant performance of the P-CNF/MMT films, with a delay in the release of PE decomposition gases and lack of ignition of PE under the cone heat flux. On the other hand, both PE-P-CNF/Sep and PE-P-CNF/SHMP structures ignited during cone calorimetry. Such behavioral differences are attributed to the structural features of the P-CNF/MMT composite film in which, in contrast to the random distribution of Sep nanorods, the regular orientation and ability of MMT nanoplatelets to form a continuous inorganic network within the P-CNF matrix led to optimal thermal shielding and gas barrier properties. To investigate the char-forming ability of P-CNF in the presence of a higher concentration of phosphate groups, SHMP was combined with the fibrils. The TG analysis showed that SHMP had a limited effect on char production of P-CNF during pyrolysis, probably because the position of the phosphoryl groups attached to cellulose after phosphorylation with SHMP, led to the formation of volatile species rather than thermally stable char at high temperatures. The mechanical properties of the P-CNF-based films were also studied and the best performance was shown by the P-CNF/MMT films with a tensile strength of 304 MPa and a Young’s modulus 35 ACS Paragon Plus Environment


of 15 GPa, resulting from the strong interfacial adhesion of the film components. In combination with excellent gas barrier and flame-retardant properties, this makes the P-CNF/MMT hybrids suitable for many applications such as interior construction and packaging materials. Associated content Supporting information Schematic illustration of the preparation of the sealed, sandwich structures, cone calorimetry results of the sealed structures prepared using P-CNF/MMT (30 wt.%) and P-CNF/Sep (30 wt.%) and videos of the uncoated PE and sealed structures tested with the heating element. Author information Corresponding authors *

Email: [email protected]

*

Email: [email protected]

Author contribution The manuscript was written through contribution of all authors. All authors have given approval to the final version of the manuscript. Notes The authors declare no competing financial interest.


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Acknowledgements The authors would like to acknowledge the Swedish Foundation for Strategic Research (SSF, grant number: RMA11-0065) for the financial support and Lars Wågberg also acknowledges the Wallenberg Wood Science Center for financial support.

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For Table of Contents Only


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Figure 1. Schematic illustration of the preparation of P-CNF based nanocomposite films, the constituents and SEM (scanning electron microscopy) images of the final film structures. 177x108mm (300 x 300 DPI)

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Figure 2. Optical and structural properties of the P-CNF based films: (a) digital photographs; (b) SEM crosssectional micrographs at different magnifications; (c) UV-vis transmittance spectra of the films. 177x143mm (300 x 300 DPI)


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Figure 3. (a) TG; (b) dTG/dT curves of P-CNF-based films, MMT, Sep and SHMP in nitrogen. 177x71mm (300 x 300 DPI)



Figure 4. Typical stress-strain curves of the P-CNF-based films. 84x68mm (300 x 300 DPI)


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Figure 5. The behavior of the sealed, sandwich structures during cone calorimetry testing: (a) HRR curves; (b) TSR curves; (c) residues of the uncoated PE film and sealed structures at the end of the test. 177x90mm (300 x 300 DPI)



Figure 6. (a) TG; (b) dTG/dT curves of uncoated PE film and of the sealed sandwich structures in nitrogen. 177x71mm (300 x 300 DPI)


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Figure 7. FTIR spectra of the gases evolved from the pure PE film and from the sealed, sandwich structures as a function of temperature. 177x93mm (300 x 300 DPI)



Figure 8. The intensity profile of the C-H band (2926 cm-1) as a function of time for PE and for the sealed, sandwich structures. 84x66mm (300 x 300 DPI)


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Figure 9. Digital photographs of the PE film and sealed, sandwich structures before and after heating with the heating element, together with snapshots from the recordings: (a) PE film; (b) PE-P-CNF; (c) PE-PCNF/MMT; (d) PE-P-CNF/Sep; (e) PE-P-CNF/SHMP. 140x133mm (300 x 300 DPI)



Figure 10. SEM cross-sectional micrographs of P-CNF-based films heated with the heating element, at different magnifications: (a) P-CNF; (b) P-CNF/MMT; (c) P-CNF/Sep; (d) P-CNF/SHMP 177x66mm (300 x 300 DPI)


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Figure 11. Schematic illustration of the reaction routes during pyrolysis of phosphorylated cellulose, rearranged from the pyrolysis routes suggested by Kandola et al. and by Aoki et al.50,51: (a) dephosphorylation at C2 or C3; (b) dephosphorylation at C6. 177x71mm (300 x 300 DPI)



Table of Contents graphic 84x36mm (300 x 300 DPI)


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Tuning the Nanoscale Properties of Phosphorylated Cellulose

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