Fire-retardant and ductile clay nanopaper biocomposites based on montmorrilonite in matrix of cellulose nanofibers and carboxymethyl cellulose

European Polymer Journal 49 (2013) 940–949

Contents lists available at SciVerse ScienceDirect

European Polymer Journal journal homepage: www.elsevier.com/locate/europolj

Fire-retardant and ductile clay nanopaper biocomposites based on montmorrilonite in matrix of cellulose nanofibers and carboxymethyl cellulose Andong Liu, Lars A. Berglund ⇑ Department of Fiber and Polymer Technology, Royal Institute of Technology, SE-10044 Stockholm, Sweden Wallenberg Wood Science Center, Royal Institute of Technology (KTH), SE-10044 Stockholm, Sweden

a r t i c l e

i n f o

Article history: Received 24 September 2012 Received in revised form 20 December 2012 Accepted 24 December 2012 Available online 8 February 2013 Keywords: Nanocellulose Nanofibrillated Hybrid Barrier Mechanical Nanocomposite

a b s t r a c t Nacre-mimetic clay bionanocomposites of high clay content show interesting properties although low strain to failure is a limitation. For this reason, three-component nanocomposite films were prepared based on sodium montmorrilonite clay (MTM), a water-soluble cellulose derivative (CMC) of fairly high molar mass, in combination with nanofibrillated cellulose (NFC) from wood pulp. The nanocomposite is cast from an aqueous colloidal dispersion. First, the effects of CMC content on CMC/MTM compositions with high volume fraction of MTM (36–83 vol.%) were studied by FE-SEM, XRD, UV, DMTA and TGA. In addition, fire retardance and oxygen permeability characteristics were measured. The effect of NFC nanofiber addition to the matrix phase was then evaluated. This two-phase CMC/NFC matrix phase results in significantly improved modulus, strength but also strain to failure. NFC has a favorable effect by shifting catastrophic failure mechanisms to higher strains. Ó 2013 Elsevier Ltd. All rights reserved.

1. Introduction Polymer/clay nanocomposites (PCNs) have been actively studied since 1967 [1], and were then developed for automotive applications [2,3]. Polyamide 6/clay nanocomposites were successfully developed at Toyota, and applications were demonstrated [2,3]. More than 10,000 papers related to PCN have since been published since. The incorporation of less than 5 wt.% clay dramatically improves gas barrier properties [4], flame retardancy [5], and mechanical properties [3]. Numerous efforts have been carried out in order to increase the clay content in the polymer matrix, but most of them failed because of the very high aspect ratio of clay platelets. If the processing approach leads to random in space platelet orientation, then ⇑ Corresponding author. Address: Wallenberg Wood Science Center and Department of Fiber and Polymer Technology, Royal Institute of Technology, SE-10044 Stockholm, Sweden. Tel.: +46 8 7908118; fax: +46 8 7908101. E-mail address: [email protected] (L.A. Berglund). 0014-3057/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.eurpolymj.2012.12.017

the maximum theoretical clay content is very low and strong agglomeration effects are obtained. In addition, the hydrophilic surface and nanoscale size of clay platelets could make it difficult to achieve good dispersion of clay, in particular in hydrophobic polymer matrices. Some papers have reported on PCNs with very high clay contents, from about 50 wt.% to 80 wt.%. The strength and modulus can be quite high. For example, Kotov et al. reported on MTM/polyelectrolyte nanocomposites with about 50 wt.% MTM (montmorrilonite) [6] prepared by elegant but fairly time-consuming layer-by-layer (LbL) deposition techniques [7]. Using a similar technique, greatly improved modulus and strength was reported for MTM– PVA (polyvinylalcohol) composites where the MTM–polymer interaction is strong [8]. In other studies [9,10], a much simpler preparation route was used which was inspired by papermaking filtration procedures. A hydrocolloidal dispersion of exfoliated silicate platelets was mixed with a water-soluble polymer solution. The polymer adsorbed to the clay platelets, which were coated by the polymer. In order to achieve control of the polymer matrix

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

distribution, the excess polymer present in solution was removed. The stiffest and strongest composition with about 50 wt.% MTM showed a modulus of 45 GPa, tensile strength of 250 MPa and strain to failure of 0.9%. MTM platelets were also combined with nanofibrillated cellulose (NFC) from wood to produce 200 mm diameter flat clay nanopaper sheets by a semi-automatic paper-making procedure [11]. These NFC nanofibers from wood pulp have a diameter of around 15 nm and a length of several micrometers, and are flexible in bending although stiff and strong in tension. Clay nanopaper has considerable ductility due to the NFC network [12]. Addition of chitosan reduces filtration time and modifies mechanical properties [13]. The polymer adsorption method developed for true nanostructural control [9,10], was also used to prepare MTM/chitosan bionanocomposite films [14]. The hybrid films prepared by first adsorbing chitosan show improved optical transparency and mechanical properties compared with hybrid films prepared by directly mixing MTM and chitosan (without separate adsorption step). Although direct mixing and filtration is disadvantageous in terms of nanostructural control, this processing route is simpler. In early work, PCNs with high content of synthetic smectite saponite clay were combined with CMC [15]. CMC (carboxymethyl cellulose) in sodium salt form is a water-soluble cellulose derivative used as a negatively charged polymeric additive in industrial papermaking. Flexible and transparent clay/polymer films showed high thermal resistance and good gas barrier properties with only 10 or 20 wt.% water-soluble polymer as a binder [15]. Clay platelets were highly ordered parallel to the film surface, and the oxygen transmission rate was extremely low at dry conditions. However, the tensile strength was only around 25 MPa with a strain to failure of 1.8%. Although most PCNs with high content of ordered clay are stiff and strong, a remaining challenge is to design materials with increased strain to failure. PCNs with high clay content rely on matrix failure processes [16] and platelet pull-out [17] as important mechanisms, which can increase strain to failure. Nanofibrillated cellulose (NFC) from plant cell walls is an interesting building block for new nanocomposites. NFC networks have been used to reinforce a variety of polymer matrices [18–21]. In contrast to pure bacterial cellulose [22], NFC based on enzymatically pretreated wood pulp [23] forms stable hydrocollodial dispersions [24]. The reason is not completely clear although it may be due to weakly charged hemicelluloses at the surface of NFC. Due to the stability of the NFC hydrocolloid, watersoluble polymers can be readily combined with NFC. In a previous study [12], NFC nanofibers were directly used to provide a nanofibrous matrix to MTM platelets and formed clay nanopaper structures. The clay content was as high as 89% by weight with maintained mechanical robustness. This clay composition had a strain to failure as high as 2%, due to the load-carrying function of the continuous, web-like cellulose nanofiber network. A later study combined electronegative NFC with clay in exceptionally thin films of very high mechanical performance [25]. The very low film thickness may contribute to the favorable properties obtained.

941

Encouraged by the interesting characteristics of threecomponent clay nanopaper containing clay, NFC and chitosan [13], (strength 103 MPa and strain to failure 2.4%) we proceed to study a nanocomposite system based on MTM, NFC and CMC. The main technical objective is to aim for improved ductility, since it is very difficult to combine high MTM content (>25 vol%) with ductile deformation behavior. This has not been achieved in previous studies. Since both CMC and the MTM surface are electronegative, weak interaction is expected which may increase strain to failure. First, structure–property relationships in MTM–CMC nanocomposite films are studied with MTM content in the range 36–83% by volume. The properties will be compared with a previous chitosan study [14] (strength 99 MPa, strain to failure 2.3%). Secondly, the effect of adding a small amount of NFC is studied. Mechanical properties, thermal stability, gas barrier properties, optical transparency and fire-shielding characteristics are evaluated. 2. Experimental 2.1. Materials Carboxymethyl cellulose sodium salt (CMC) with an average molecular weight of 250,000 and a DS of 0.7 was obtained from Sigma Aldrich Co., 1.0 wt.% CMC was dissolved in de-ionized water by vigorous stirring before using. The MTM (Cloisite Na+, Southern MTM Products) was a sodium montmorillonite (MTM) with a cation-exchange capacity (CEC) of 92 meq/100 g. The average width of the platelets is 110 nm as described by the manufacturer, with a thickness of about 1 nm. 1.0 wt.% hydrocolloidal MTM dispersion was prepared by dispersing 10 g of MTM in 1 L of de-ionized water under vigorous stirring. A nanofibrillated cellulose (NFC) hydrocolloid dispersion was prepared as described in previous work [26]. Degree of polymerization (DP) was estimated to be 480 after homogenization, as determined from the average intrinsic viscosity of the dissolved polymer [26]. An NFC aqueous suspension with a concentration of 1.6 wt.% was obtained and stored at 4 °C. The NFC suspension was diluted to 0.2 wt.% in de-ionized water and stirred before use. 2.2. Preparation of CMC/MTM hybrid films CMC/MTM hybrid films were prepared as follows: 1.0 wt.% CMC solution was added slowly to the MTM dispersion and the mixture was stirred for 24 h. Then it was poured into a Teflon mold and dried in the oven at 55 °C. CMC/MTM hybrid films with a thickness in the range 60– 70 lm were obtained with initial CMC–MTM weight ratio of 10:90, 20:80, 30:70, 40:60 and 50:50, coded as CMC10–MTM90, CMC20–MTM80, CMC30–MTM70, CMC40–MTM60, CMC50–MTM50, respectively. 2.3. Preparation of NFC/CMC nanocomposite films CMC nanocomposite films were prepared as follows: 0.2 wt.% NFC solution was added slowly into the 1.0 wt.%

942

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

CMC solution and was stirred for 24 h. Then it was poured into a Teflon mold and dried in the oven at 55 °C. Finally, NFC/CMC films with a thickness in range 60–70 lm were obtained with initial CMC–NFC weight ratio of 100:0, 100:4, 100:8, 100:12, coded as CMC, CMC100–NFC4, CMC100–NFC8, CMC100–NFC12, respectively.

2.4. Preparation of three-component CMC/MTM/NFC nanocomposite films CMC/MTM/NFC nanocomposite films were prepared as follows: According to the method of preparation of CMC/ MTM hybrid films, CMC/MTM mixture with a weight ratio of 50:50 and 20:80 were prepared first. Then a desired quantity of 0.2 wt.% NFC solution was added slowly into the CMC/MTM mixture and was stirred for 24 h. Then it was poured into a Teflon mold and dried in the oven at 55 °C. Finally, CMC/MTM/NFC nanocomposite films with thickness in the range of 60–70 lm were obtained with initial CMC–MTM–NFC weight ratio of 50:50:4, 50:50:8, 50:50:12, 20:80:4, 20:80:8, 20:80:12, coded as CMC50– MTM50–NFC4, CMC50–MTM50–NFC8, CMC50–MTM50– NFC12, CMC20–MTM80–NFC4, CMC20–MTM80–NFC8, CMC20–MTM80–NFC12, respectively. Table 1 shows all information of the initial components of samples in aqueous solutions/dispersion.

capture secondary electron images of the surface of fractured cross-sections. 3.2. X-ray diffraction (XRD) Wide angle XRD patterns were recorded by an X’pert Pro diffractometer at room temperature. The Cu Ka radiation source was operated at 40 kV and 40 mA. Patterns were recorded by monitoring diffractions from 1° to 10°. The scan speed was 2°/min. 3.3. Tensile testing Uniaxial tensile tests of the films were performed with a Universal Materials Testing Machine from Instron, USA, equipped with a 100 N load cell. Specimens of 40 mm in length and 60–70 lm in thickness and 5 mm in width were tested at a crosshead speed of 4 mm/min. The relative humidity was kept at 50% and the temperature at 23 °C. The specimens were conditioned for at least 48 h in this environment prior to testing. The displacement was measured by Digital Speckle Photography (DSP). A pattern was prepared for the DSP by applying printer toner to the sample surface. During tensile testing, images were obtained of the whole specimen. The frame rate was set to 5 fps. The results for each material are based on at least six specimens, unless reported differently. 3.4. Thermogravimetric analysis (TGA)

3. Characterization of materials 3.1. Scanning electron microscopy (SEM) The film morphology was examined with a JEOL JSM820 Field Emission Scanning Microscope. The samples were immersed in liquid N2 and brittle fracture was induced by bending. The specimens were fixed on a metal stub using carbon tape and coated with a double-layer coating (5 nm) consisting of graphite and gold–palladium using Agar HR sputter coaters prior to imaging. A Hitachi S-4800 scanning electron microscope operated at 1 kV to

Table 1 The notation used for different materials, and the corresponding composition expressed as parts by weight (pbw). Sample codes

CMC (pbw)

MTM (pbw)

NFC (pbw)

CMC CMC100–NFC4 CMC100–NFC8 CMC100–NFC12 CMC50–MTM50 CMC50–MTM50–NFC4 CMC50–MTM50–NFC8 CMC50–MTM50–NFC12 CMC40–MTM60 CMC30–MTM70 CMC20–MTM80 CMC20–MTM80–NFC4 CMC20–MTM80–NFC8 CMC20–MTM80–NFC12 CMC10–MTM90

100 100 100 100 50 50 50 50 40 30 20 20 20 20 10

0 0 0 0 50 50 50 50 60 70 80 80 80 80 90

0 4 8 12 0 4 8 12 0 0 0 4 8 12 0

Thermogravimetric analysis (TGA) was conducted on a Perkin–Elmer TGA 7-thermal analyzer from 25 to 800 °C with a heating rate of 10 °C/min under oxygen with flow rate 50 ml/min. 3.5. Oxygen transmission rate The oxygen permeability of the material at 23 °C was determined using a Mocon OX–TRAN TWIN equipped with a coulometric oxygen sensor. Degassed film samples with thickness of 60–70 lm were mounted in an isolated diffusion cell and were subsequently surrounded by flowing nitrogen gas to remove sorbed oxygen from the samples. The sample had a circular exposure area of 100  104 m2 obtained by covering part of the film with a tight aluminum foil with an adhesive on its surface. One side of the sample was initially exposed to flowing oxygen containing 1% hydrogen at atmospheric pressure. The oxygen pressure was zero on the other side. The flow rate (Q) through the sample was measured and, from the steady-state flow rate (Q1), the oxygen permeability coefficient (P) was calculated. 3.6. Volume fraction calculations The porosity of the nanocomposite films was assumed to be zero, since microscopy revealed no indications of voids. The density of CMC is assumed to be 1590 kg m3. The density of cellulose is assumed to be 1460 kg m3, while the density of MTM is assumed to be 2860 kg m3. The volume fractions (Vi, i = 1, 2, . . ., n) of MTM, NFC and

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

CMC are calculated from the weight fraction of NFC and from the density of the films according to the following equations [27]:

V i ¼ W i xðqct =qi Þ where qct is the density of the solid constituting the composite, which was calculated from the densities, qi, of the constituents (i = 1, 2, . . ., n), Wi:

1 ðW i =qi Þ i¼1

qct ¼ Pn

943

1.3–1.4 nm. It means the gallery distance increased very little compared with the pure MTM. This may indicate that CMC molecules only adsorb at the edges of MTM platelets by electrostatic interaction. MTM platelets have negative charge on the surface and positive charge at the edge. CMC is a negatively charged polymer, should be readily adsorbed to the edge of MTM. As CMC content was increased to 40 and 50 wt.%, the d-space values increased to 1.7 and 1.8 nm, respectively. This is due to intercalation of CMC molecules in the galleries between individual MTM platelets. 4.2. CMC/MTM nanocomposites films – mechanical properties

4. Results and discussion 4.1. CMC/MTM nanocomposites – lamellar structure by FESEM and XRD Fig. 1 presents a photograph of the free-standing CMC20/MTM80 nanocomposite film with 69% volume fraction of MTM (70 lm film thickness), demonstrating optical transparency. The yellowish tone is probably related to the color of the native MTM influenced from the presence of metal ions. Despite the high MTM content, the film is non-brittle and quite flexible in bending. FE-SEM micrographs of cross-sections of cryo-fractured surfaces are presented in Fig. 2. For the pure CMC film in Fig. 2a, there is no sign of lamellar order, as expected. However, for the CMC/MTM hybrid films, a layered structure is apparent. The layered structure is fairly regular and the smallest apparent layer thickness is at a scale of about 100 nm and below. The sheet-like lamellae are parallel to the film surface. The layers have a slightly wavy appearance and are interpenetrating into neighboring layers. This phenomenon was also observed in biomimetic organic– inorganic composites made from PVA/MTM by LbL or MTM/polymer self-assembly [9,14,17,28]. XRD data (Fig. 3) from CMC/MTM hybrid nanocomposites with increasing CMC content from 0 wt.% to 50 wt.% show that MTM basal-space reflection peaks move towards lower angle. According to Bragg’s formula, -space values (d001) for CMC/MTM with 10–30 wt.% CMC are around

In Fig. 4A, stress–strain curves are presented for CMC/MTM with increasing CMC content. Both modulus (E) and tensile strength (r) increase markedly with increasing CMC content. E and r for CMC50–MTM50 (VMTM = 35.7%) are 5.2 GPa and 95 MPa, as compared with the much lower values for pure CMC (E = 2.0 GPa and r = 56 MPa). The high clay content composition CMC10–MTM90 (VMTM = 83.3%) has E = 3.1 GPa and r = 42 MPa, see also Table 2 for a summary of the data. The low modulus and strength of this composition require some discussion, It is encouraging that nanocomposite films with such high MTM content can be prepared and physically handled, since one can expect good fire retardancy and barrier property characteristics. However, the stress transfer between the CMC matrix and the MTM platelets is non-ideal. The low modulus is a strong indication. The MTM platelets form agglomerates with low load transfer efficiency. MTM-rich entities also lead to stress concentration effects and premature fracture at low strain and stress. At very high MTM content, CMC is not uniformly distributed with the present preparation method. The stiff MTM platelets and platelet aggregates are not fully loaded since they interact weakly with each other and with the matrix. As a consequence, E and r show positive linear correlation with CMC volume fraction for the range of compositions studied (Fig. 4). Higher CMC content leads to better reinforcement efficiency for the present compositions. Addition of increasing amounts of the anionic CMC polyelectrolyte matrix improves stress transfer so that modulus and strength are increased. An advantage with the present system is the comparably high strain to failure despite high inorganic content. At a VMTM of 69%, the present strain to failure is 2.9%. 4.3. CMC/MTM nanocomposites – thermogravimetric analysis, flame-shielding and oxygen barrier properties

Fig. 1. Free-standing CMC20–MTM80 (20 wt.% CMC polymer, 80 wt.% MTM clay) nanocomposite film of high optical transparency.

Results from thermogravimetric analysis (TGA) of CMC/ MTM nanocomposites with different content of CMC are presented in Fig. 5. For pure CMC, line (a), there is only one degradation stage and the decomposition rate is very fast. For CMC/MTM nanocomposites (line b–f), the decomposition rate was dramatically reduced with decreasing CMC content. The main reason is that the MTM nanoplatelets formed continuous protective solid layers, so that

944

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

Fig. 2. FE-SEM micrographs of the cross-section of cryo-fractured surfaces of pure CMC (a), CMC50–MTM50 (b), CMC20–MTM80 (c) and CMC10–MTM90 (d).

Fig. 3. XRD patterns of pure MTM (a), CMC10–MTM90 (b), CMC20– MTM80 (c), CMC30–MTM70 (d), CMC40–MTM60 (e) and CMC50–MTM50 (f). d-Space data are provided next to each curve.

oxygen diffusion is hindered and oxidation kinetics becomes slower. One may speculate that silicate layer fusion

and/or char formation of the CMC are contributing mechanisms. These observations are very interesting since the fire retardance characteristics of this system are likely to be favorable. Fig. 6A–C, illustrates the fire retardant characteristics of CMC50–MTM50. Initially, the CMC at the surface caused quick burning (see the video in Supporting information). After a few seconds, the burning flame was extinguished. The shape of the film was preserved even after exposure to a flame for one minute. The integrity of the CMC/MTM film after burning can be attributed to char formation of the CMC. From FE-SEM micrographs (Fig. 6D–F), it is apparent that CMC50–MTM50 shows intumescent behavior (swells as it is exposed to heat). The film expands more than five times in thickness direction due to exposure to the flame (Fig. 6C). The outer MTM layers formed dense fireproofing layers during the process, which hindered flame penetration and gas diffusion. The heat was not sufficient to support the burning process itself and the film self-extinguished. It is worth noting that the charred CMC–MTM film kept its layered structure. The oxygen transmission rate (OTR) of CMC50–MTM50 at dry conditions is about 0.008 cc/m2 day atm. This very low value is due to the highly ordered organization of the impenetrable layered silicate platelets. The path for gas diffusion becomes tortuous and the path length increases dramatically. However, the CMC50–MTM50 suffers from moisture sensitivity due to the presence of the hydrophilic CMC matrix. The OTR at 50% relative humidity is about 2 cc/m2 day atm.

945

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

Fig. 5. Data from TGA experiments in O2-environment for pure CMC film (a) CMC10–MTM90 (b), CMC20–MTM80 (c), CMC30–MTM70 (d), CMC40–MTM60 (e) and CMC50–MTM50 (f).

Fig. 4. Mechanical properties. (A) Stress–strain curves in tension for pure CMC (a) CMC10–MTM90 (b), CMC20–MTM80 (c), CMC30–MTM70 (d), CMC40–MTM60 (e) and CMC50–MTM50 (f). (B) Tensile strength (j) and modulus (s) as a function of CMC volume fraction for CMC/MTM films.

4.4. NFC reinforcement effects on CMC/MTM nanocomposites The purpose of adding NFC nanofibers to the CMC matrix is to increase the strain to failure of the clay nanocomposites. The CMC50–MTM50 composition (36% MTM by

volume) has 95 MPa strength and a strain to failure of 2.5%. If the strain to failure could be increased to around 6%, then the material would be quite tough and easy to handle in film form. CMC/MTM with high volume fraction of MTM (36–83 vol.%) show many interesting features including optical transparency, very low oxygen transmission rate, good thermal stability and attractive mechanical properties considering the high inorganic content. However, the strain to failure of CMC/MTM is low as was presented in Fig. 4A. Previous work on MTM nanopaper based on MTM and a nanofibrillated cellulose (NFC) matrix [12] showed non-linear stress–strain behavior and favorable strain to failure even at MTM contents as high as 83 vol% (2.5% strain to failure). The reason is that the NFC nanofibers formed a continuous web-like nanofiber network ‘‘matrix’’ in which MTM platelets and aggregates were embedded. NFC was therefore added to the present nanocomposite. First, the effect of NFC on the pure CMC films was investigated and tensile curves are shown in Fig. 7. The strength and strain to failure of CMC films

Table 2 Summary information of weight fraction, volume fraction and mechanical properties of hybrid films. Sample codes

Wf of CMC/MTM/NFC (%)

Vf of CMC/MTM/NFC (%)

Tensile strength (MPa)

Tensile modulus (GPa)

Strain to failure (%)

CMC CMC100–NFC4 CMC100–NFC8 CMC100–NFC12 CMC50–MTM50 CMC50–MTM50–NFC4 CMC50–MTM50–NFC8 CMC50–MTM50– NFC12 CMC40–MTM60 CMC30–MTM70 CMC20–MTM80 CMC20–MTM80–NFC4 CMC20–MTM80–NFC8 CMC20–MTM80– NFC12 CMC10–MTM90

100/0/0 96.2/0/3.8 92.6/0/7.4 89.3/0/10.7 50/50/0 48.1/48.1/3.8 46.3/46.3/7.4 44.7/44.7/10.6

100/0/0 95.9/0/4.1 92/0/8 88.5/0/11.5 64.3/35.7/0 60.9/33.8/5.3 57.8/32.1/10.1 55/30.6/14.4

54.0 ± 2.7 67.4 ± 5.1 71.2 ± 0.8 84.6 ± 2.9 94.9 ± 4.1 105.5 ± 1.3 112.3 ± 2.9 130.4 ± 5.3

2.00 ± 0.10 2.04 ± 0.09 2.26 ± 0.15 2.45 ± 0.11 5.51 ± 0.09 5.78 ± 0.12 5.77 ± 0.21 6.24 ± 0.31

8.8 ± 0.4 12.0 ± 0.8 14.9 ± 1.7 16.3 ± 1.0 2.5 ± 0.3 3.2 ± 0.3 4.1 ± 1.1 5.2 ± 0.6

40/60/0 30/70/0 20/80/0 19.2/77.0/3.8 18.5/74.1/7.4 17.9/71.4/10.7

54.6/45.4/0 43.5/56.5/0 31/69/0 29.1/64.6/6.3 27.3/60.8/11.9 25.7/57.4/16.9

85 ± 2.1 75.8 ± 3.6 60.1 ± 0.7 64.8 ± 5.1 76.5 ± 1.3 82.6 ± 2.0

4.93 ± 0.21 4.81 ± 0.13 3.80 ± 0.04 5.04 ± 0.20 5.36 ± 0.09 5.53 ± 0.14

3.6 ± 0.2 3.1 ± 0.3 2.9 ± 0.2 2.7 ± 0.3 3.4 ± 0.8 4.2 ± 0.3

10/90/0

16.7/83.3/0

42.0 ± 1.6

3.13 ± 0.19

1.8 ± 0.2

946

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

Fig. 6. Flame-shielding characteristics of CMC50–MTM50. Images before burning (A), during exposure to flame (B) and after burning (C); FE-SEM images of cross-section of CMC50–MTM50 after burning (D–F). Arrow direction indicates the direction of burning flame propagation.

Fig. 7. Stress–strain curves of nanocomposites based on CMC polymer and NFC nanofibers, with CMC and NFC content expressed as parts by weight.

increased strongly with NFC content, although a maximum of only 11.5 vol.% was added. The strong effect of NFC on these properties is due to good dispersion of NFC and the formation of a nanofiber network [19,29]. The strain to failure improves by extension of the strain-hardening region of the stress–strain curve. The NFC network not only increases the load-bearing capacity of the material in this region, but also delays critical failure phenomena to higher strains. The present data show better properties than recent work based on hydroxypropyl cellulose [21]. The main

reason may be that the present CMC has intrinsically better mechanical properties. A high content of NFC in a matrix of hydroxyethyl cellulose (HEC) results in even better mechanical properties than present data [30]. The effect of a small amount of NFC nanofibers on the mechanical properties of CMC/MTM was studied further. Fig. 8A–C shows tensile strength, modulus, and strain to failure of neat CMC films, CMC50–MTM50 and CMC20– MTM80 three-phase nanocomposites with increasing NFC content. MTM platelets are embedded in a matrix of NFC and CMC. Compared with the previous chitosan–NFC– MTM study [13], CMC–NFC–MTM nanocomposites show lower modulus but much larger strain to failure. The strength is about the same for comparable compositions. Most likely, the reason for lower modulus with electropositive CMC combined with electropositive MTM is that chitosan interacts more strongly with MTM. Chitosan (electropositive) and MTM surfaces (electronegative) have opposite charges. It is likely that weak interaction between the present CMC–NFC matrix and the MTM is actually positive for strain to failure, since MTM pull-out from the matrix may be facilitated during the failure process. Table 2 summarizes weight fractions, volume fractions and mechanical properties of nanocomposites. Tensile strength, modulus and strain to failure increase with NFC volume fraction. The CMC50–MTM50–NFC12 composition shows 130 MPa strength, a modulus of 6 GPa and a strain to failure of 5.2%. This attractive combination of properties

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

947

higher resolution (Fig. 9C) showed some fiber-like bridges between layers indicating that NFC nanofibers are linking the layers. This may contribute to improved strain to failure of nanocomposites with NFC in the matrix. 5. Conclusions

Fig. 8. Tensile strength (A), modulus (B) and strain to failure (C) of CMC films, CMC50–MTM50 hybrid films and CMC20–MTM80 hybrid films with different weight % of NFC. The volume fractions of each component are provided in Table 2.

is due to the synergy of NFC and MTM. Damage processes become less localized, are delayed and the corresponding increase in strain to failure extends the strain-hardening plastic deformation region so that also strength is increased. From XRD results (Fig. 9A), the d-space of MTM in CMC50–MTM50 hybrid films with varying NFC content was almost constant. This indicates that NFC nanofibers did not affect the dispersion state of MTM during processing. However, the structure in the FE-SEM image of CMC50/MTM50 hybrid films in Fig. 2B is different from the structure in Fig. 9B where 12 wt.% NFC has been added. Fig. 9B shows less sign of a layered structure. Images with

Optically transparent three-phase nanocomposites were prepared based on high content of MTM clay platelets (up to 65% by volume) in a matrix of CMC polymer and an NFC nanofiber network. Clay platelets were highly ordered in a direction parallel to the film surface. Due to the network forming characteristics, NFC nanofibers strongly improved the toughness of CMC–MTM nanocomposite films, where toughness is interpreted as the area under the stress–strain curve from tensile tests. In particular, strain to failure and tensile strength were increased with addition of small amounts of NFC nanofibers. The best previously reported biopolymer matrix clay nanocomposites with high clay content (>25 vol%) are based on chitosan and show strengths around 100 MPa and strain to failure of 2.3–2.4%. Here the best composition had a strength of 130 MPa, and a strain to failure as high as 5.2%. This composition had 31 vol.% MTM, and the matrix had 55 vol% CMC and 14 vol% NFC. The NFC network increases the strain to failure, possibly by prevention of premature cracking at small scale. The weak CMC–MTM interaction is another factor, caused by electronegative charge at both MTM surfaces and in CMC. This may cause a larger extent of MTM pull-out during failure and contribute to ductility. The modulus of the composition is only 6.2 GPa at an MTM volume fraction as high as 31%, which is in support of weak CMC–MTM interaction. Compared with data for chitosan matrix MTM bionanocomposites (50/50 by weight) reported in the literature, the optical transparency was much better. This indicates better dispersion of the clay in the present study, possibly due to repulsion between CMC and MTM (both electronegative) in the colloidal state. The orientation and high clay content result in high thermal stability and an oxygen transmission rate (OTR) of CMC50–MTM50 in the dry state as low as 0.008 cc/m2 day atm. The flame retardant characteristics were highly favorable. The gas barrier properties contribute to this and possibly also silicate layer fusion and charring of the polysaccharides. Intumescent behavior was observed so that the clay nanopaper structure expanded significantly in thickness direction during exposure to the flame. The combination of water-based processing, a high content of clay platelets ordered parallel to the surface, and a two-phase polymer–NFC network matrix is an attractive concept for new nanocomposites. It results in a combination of functionalities such as fire retardancy, optical transparency and, as demonstrated in the present study, high strength and ductility in combination with low density. Acknowledgements The BiMaC Innovation Center is gratefully acknowledged for financial support of Dr. Andong Liu during the very early part of the project. WWSC then funded the

948

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949

Fig. 9. (A) XRD patterns of pure MTM (a) and CMC50–MTM50 (b), CMC50–MTM50–NFC4 (c), CMC50–MTM50–NFC8 (d) and CMC50–MTM50–NFC12 (e). (B) and (C) Images of the cross-section of cryo-fractured surfaces of CMC50–MTM50–NFC12.

major experimental work. The SSF consortium Fire Foam is also acknowledged since data analysis and writing of the manuscript was funded by this program.

[11]

[12]

Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/ j.eurpolymj.2012.12.017.

[13]

[14]

References [15] [1] Dekking HGG. Propagation of vinyl polymers on clay surfaces. II. Polymerization of monomers initiated by free radicals attached to clay. J Appl Polym Sci 1967;11:23–6. [2] Usuki A, Kojima Y, Kawasumi M, Okada A, Fukushima Y, Kurauchi T, et al. Synthesis of nylon 6/clay hybrid. J Mater Res 1993;8:1179–84. [3] Kojima Y, Usuki A, Kawasumi M, Okada A, Fukushima Y, Kurauchi T, Kamigaito O. Mechanical-properties of nylon 6–clay hybrid. J Mater Res 1993;8:1185–9. [4] Triantafyllidis KS, LeBaron PC, Park I, Pinnavaia T. Epoxy–clay fabric film composites with unprecedented oxygen-barrier properties. J Chem Mater 2006;18:4393. [5] Gilman JW, Jackson CL, Morgan AB, Harris Jr R. A study of the flammability reduction mechanism of polystyrene-layered silicate nanocomposite: layered silicate reinforced carbonaceous char. Chem Mater 2000;12:1866. [6] Tang Z, Kotov NA, Magonov S, Ozturk B. Nanostructured artificial nacre. Nat Mater 2003;2:413–8. [7] Decher G. Fuzzy nanoassemblies: toward layered polymeric multicomposites. Science 1997;277:1232–7. [8] Podsiadlo P, Kaushik AK, Arruda EM, Waas AM, Shim BS, Xu J, et al. Ultrastrong and stiff layered polymer nanocomposites. Science 2007;318:80–3. [9] Walther A, Bjurhager I, Malho JM, Pere J, Berglund LA, Ikkala O. Large-area, lightweight and thick biomimetic composites with superior material properties via fast, economic, and green pathways. Nano Lett 2010;10:2742–8. [10] Walther A, Bjurhager I, Malho JM, Berglund LA, Ikkala O. Supramolekulare Kontrolle der mechanischen eigenschaften

[16] [17]

[18]

[19]

[20]

[21]

[22]

[23]

feuerabschirmender biomimetischer Perlmuttanaloga. Angew Chem 2010;122:6593–9. Sehaqui H, Liu AD, Zhou Q, Berglund LA. Fast preparation procedure for large, flat cellulose and cellulose/inorganic nanopaper structures. Biomacromolecules 2010;11:2195–8. Liu AD, Walther A, Ikkala O, Belova L, Berglund LA. Clay nanopaper with tough cellulose nanofiber matrix for fire-retardancy and gas barrier functions. Biomacromolecules 2011;12:633–41. Liu AD, Berglund LA. Clay nanopaper composites of nacre-like structure based on montmorrilonite and cellulose nanofibers— Improvements due to chitosan addition. Carbohydr Polym 2012;87:53–60. Yao HB, Tan ZH, Fang HY, Yu SH. Artificial nacre-like bionanocomposite films from the self-assembly of chitosan– montmorillonite hybrid building blocks. Angew Chem Int Ed 2010;49:1–6. Ebina T, Mizukami F. Flexible transparent clay films with heatresistant and high gas-barrier properties. Adv Mater 2007;19:2450–3. Munch E, Launey ME, Alsem DH, Saiz E, Tomsia AP, Ritchie RO. Tough, bio-inspired hybrid materials. Science 2008;322:1516–20. Bonderer LJ, Studart AR, Gauckler LJ. Bioinspired design and assembly of platelet reinforced polymer films. Science 2008;319:1069–73. Fernandes SCM, Freire CSR, Silvestre AJD, Neto CP, Gandini A, Berglund LA, et al. Transparent chitosan films reinforced with a high content of nanofibrillated cellulose. Carbohydr Polym 2010;81:394–401. Svagan AJ, Azizi Samir My AS, Berglund LA. Biomimetic polysaccharide nanocomposites of high cellulose content and high toughness. Biomacromolecules 2007;8:2556–63. Yano HB, Nakahara S. Bio-composites produced from plant microfiber bundles with a nanometer unit web-like network. J Mater Sci 2004;39:1635–8. Zimmerman T, Bordeanu N, Strub E. Properties of nanofibrillated cellulose from different raw materials and its reinforcement potential. Carbohydr Polym 2010;79:1086–93. Zhou Q, Malm E, Nilsson H, Larsson P, Iversen T, Berglund T, et al. Nanostructured biocomposites based on bacterial cellulosic nanofibers compartmentalized by a soft hydroxyethylcellulose matrix coating. Soft Matter 2009;5:4124–30. Henriksson M, Henriksson G, Berglund LA, Lindström T. An environmentally friendly method for enzyme-assisted preparation of microfibrillated cellulose (MFC) nanofibers. Eur Polymer J 2007;43:3434–41.

A. Liu, L.A. Berglund / European Polymer Journal 49 (2013) 940–949 [24] Pääkkö M, Ankerfors M, Kosonen H, Nykänen A, Ahola S, Österberg M, et al. Enzymatic hydrolysis combined with mechanical shearing and high-pressure homogenization for nanoscale cellulose fibrils and strong gels. Biomacromolecules 2007;8:1934–41. [25] Wu CN, Saito T, Fujisawa S, Fukuzumi H, Isogai A. Ultrastrong and high gas-barrier nanocellulose/clay-layered composites. Biomacromolecules 2012;13:1927–32. [26] Henriksson M, Berglund LA, Isaksson P, Lindström T, Nishino T. Cellulose nanopaper structures of high toughness. Biomacromolecules 2008;9:1579–85.

949

[27] Agarwal BD, Broutman LJ. Analysis and performance of fiber composites. 2nd ed. John Wiley and Sons, Inc.; 1990. [28] Bhatnagar A, Sain M. Processing of cellulose nanofiber-reinfroced composites. J Reinf Plast Compos 2005;24:1259–68. [29] Favier V, Chanzy H, Cavaille JY. Polymer nanocomposites reinforced by cellulose whiskers. Macromolecules 1995;28:6365–7. [30] Sehaqui H, Zhou Q, Berglund LA. Nanostructured biocomposites of high toughness-a wood cellulose nanofiber network in ductile hydroxyethylcellulose matrix. Soft Matter 2011;7:7342–50.