High-Density Molded Cellulose Fibers and Transparent

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High-Density Molded Cellulose Fibers and Transparent Biocomposites Based on Oriented Holocellulose Xuan Yang, Fredrik Berthold, and Lars A. Berglund ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b22134 • Publication Date (Web): 14 Feb 2019 Downloaded from http://pubs.acs.org on February 15, 2019

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ACS Applied Materials & Interfaces

High-Density Molded Cellulose Fibers and Transparent Biocomposites Based on Oriented Holocellulose Xuan Yang†, Fredrik Berthold‡, and Lars A. Berglund†* †

Wallenberg Wood Science Center, Department of Fiber and Polymer Technology, KTH Royal Institute of Technology, SE-10044 Stockholm, Sweden

‡

RISE - Research Institutes of Sweden, Mäster Samuelsgatan 60, SE-11121 Stockholm, Sweden

KEYWORDS. wood, nanocellulose, high strength, modulus, PMMA, interface

ABSTRACT.

Eco-friendly materials based on well-preserved and nanostructured wood cellulose fibers are investigated for the purpose of load-bearing applications, where optical transmittance may be advantageous. Wood fibers are subjected to mild delignification, flow orientation, and hot-pressing to form an oriented material of low porosity. Biopolymer composition of the fibers is determined. Morphology is studied by SEM, cellulose orientation is quantified by x-ray diffraction, and effect of beating is investigated. Hot-pressed networks are impregnated by MMA monomer and

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polymerized to form thermoplastic wood fiber/PMMA biocomposites. Tensile tests are performed, as well as optical transmittance measurements. Structure-property relationships are discussed. High-density molded fibers from holocellulose have mechanical properties comparable with nanocellulose materials, and are recyclable. The thermoplastic matrix biocomposites showed superior mechanical properties (Young’s modulus of 20 GPa and ultimate strength of 310 MPa) at a fiber volume fraction of 52%, with high optical transmittance of 90%. The study presents a scalable approach for strong, stiff and transparent molded fibers/biocomposites.

INTRODUCTION One of the more interesting materials design challenges with cellulosic materials, is to combine high mechanical performance with high optical transparency. Yano et al showed the concept for thin nanocellulose paper films,1,2 and there is an excellent review on the subject, with emphasis on electronics and optoelectronics applications of flexible cellulose films.3 If cellulosic materials of high mechanical performance could be processed into thicker structures rather than thin films, load-bearing applications could be considered in construction and transportation (i.e. automotive industry). Molded cellulosic materials for semi-structural load-bearing applications would be attractive since they are from renewable resources, could be recyclable and also provide new functionalities related to optical transmittance. The term “semi-structural” is often used for injection or compression molded glass fiber composites, which have a modulus of 5-13 GPa and an ultimate strength of 50-150 MPa.4 Although transparent wood from chemically treated veneer5 could be used, the use of chemically processed wood cellulose fibers has advantages. Such fibers are used in paper and board applications, and are widely available at low cost. Shaping methods


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commonly used for polymer composites, such as compression molding or resin transfer molding, could be readily adapted for scalable processing. Nanocellulose has many advantages in terms of intrinsic properties, and the mechanical performance reported for nanocellulose materials is very high.6,7 Since the viscosity of nanocellulose suspensions is very high,8,9 and the cost may also be high,10 it is interesting to explore the potential of lower cost wood cellulose fibers. Holocellulose fibers from wood have therefore been prepared by mild delignification of wood.11,12 Because of the mild chemical and mechanical treatment, such fibers have well-preserved cellulose molar mass, hemicellulose content and very high tensile strength.12 It is of great interest to further explore the characteristics of these holocellulose fibers, since the nanostructure of the cell wall in the fibers is expected to be much better preserved compared with industrially processed wood fibers. Yet the concept of mild delignification could well be of industrial interest, provided the materials performance is promising. Molded fibers is a term used for industrial materials where a three-dimensional structure is prepared from pulp fibers, and used in some load-bearing packaging material function. Typical applications include egg cartons and trays. In the research literature, high-density molded fiber materials (also termed binderless fiberboard) have been prepared by compression molding of fibers at elevated temperature under high pressure.12–16 These materials have low porosity, do not have additives, are from renewable resources and can have a modulus approaching 20 GPa, with strength values around 120 MPa. They are biodegradable and should potentially be recyclable. Recently, it was found that such high-density molded fibers can have improved mechanical properties if they are prepared from nanostructured holocellulose fibers.12 In addition, the materials showed high optical transparency. Annual plant fibers, like flax, ramie, hemp, and jute, have better


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mechanical properties than wood cellulose fibers because of higher cellulose content and lower microfibril angles.17–19 Still, in the context of sustainable development, wood cellulose fibers have the advantage of low cost and availability. Possibly, the milder chemical treatment of wood fibers and the molded fiber concepts explored here, can inspire the development of new materials and improved bio-based products. The optical transmittance of regular paper is normally much lower than for transparent plastic materials. This is because regular paper is composed of wood cellulose fibers with diameters at the scale of tens of microns, the surfaces are often rough and the material structure is highly porous. This results in light scattering at cellulose/air interfaces, which leads to low transmittance.3 For the high-density molded holocellulose fibers,12 the porosity was reduced to 20%, the agglomeration of cellulose fibrils typical for industrial wood cellulose fibers was prevented,20 and thus the transmittance was as high as 55% at a thickness of 170 µm.12 Although high-density molded plant fibers have strong eco-friendly characteristics, transparent polymer matrix biocomposites also have advantages. Yano et al. showed the attractive combination of optical transmittance and mechanical properties for bacterial cellulose (BC)/acrylic nanocomposite films.1,2 The films were prepared by soaking the dried, nanoporous BC paper in acrylic resin precursors, followed by curing the acrylic into a chemically crosslinked polymer. For the related transparent wood materials, delignified veneer can be impregnated with methyl methacrylate (MMA), followed by thermal polymerization into a PMMA thermoplastic matrix.21 Although the refractive index match with cellulose is not perfect (cellulose ~1.53 and PMMA ~1.49), the material is highly transparent, as well as thermoformable. High molar mass PMMA can also have a good combination of modulus, strength and toughness, which is interesting for biocomposite matrices.


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In the present study, novel concepts for high-performance molded holocellulose fibers and thermoplastic polymer matrix biocomposites are developed for oriented and hot-pressed holocellulose networks. The term “holocellulose” refers to cellulosic wood fibers with wellpreserved cellulose and hemicellulose biopolymers, but with most of the lignin removed.22 The objective is to compare with nanocellulose films as well as industrial wood cellulose fibers, and to analyze structure-property relationships for two classes of materials (high-density molded cellulose fibers, and thermoplastic matrix holocellulose biocomposites). Molded fibers based on oriented holocellulose fibers are shown to have excellent mechanical properties, competing with nanocellulose materials. The high performance means that the materials are of interest for semistructural applications. Such binderless materials also have great potential to be recyclable. The thermoplastic biocomposites have much improved optical transmittance up to 90 %, with exceptional Young’s modulus reaching 27 GPa, and ultimate strength exceeding 300 MPa. The nanostructural preservation of fiber structure, high amount of native hemicellulose, and the high density makes it possible to translate the high intrinsic properties of nanostructured fibers into favorable material properties.

EXPERIMENTAL SECTION Materials. Spruce wood veneer was purchased from Calexico Wood AB, Sweden. Peracetic acid (PAA, 38-40%), sodium hydroxide (>98%), Methyl methacrylate (MMA, 99%), and 2,2′Azobis(2-methylpropionitrile) (AIBN, ≥ 98%) were purchased from Sigma-Aldrich. “Södra green”, a commercially available fully bleached Kraft pulp from soft wood (Södra, Sweden) was used as a reference fiber sample.


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Holocellulose Fiber Preparation. Holocellulose fibers were prepared using PAA delignification according to a previously published work.12 In brief, softwood spruce was delignified under mild conditions (4 wt. % PAA, pH = 4.8 before reaction) for 45 min at 85 °C, without mechanical stirring. Multiple rounds of PAA treatment were performed at the same reaction conditions, using fresh batches of PAA solutions, until all individual fibers were liberated. Beating of Fibers. In order to evaluate the effect of beating on the properties of the fibers and molded fibers, holocellulose fibers were subjected to different extents of beating in a PFI mill, from 125, 250, to 500 revolutions. The holocellulose fibers discussed in this work were labeled according to the revolution number (R): “Fiber-0R”, “Fiber-125R”, “Fiber-250R” and “Fiber500R”. Beating were set relatively low because mat formation was insufficiently dense in the dynamic sheet former after higher extent of beating. Fully bleached Kraft pulp fibers were labeled as “K-Fiber-0R”. Chemical Composition. The lignin content (Klason lignin) was determined using the TAPPI T 222 om-2 method. Carbohydrate analysis was carried out after the total hydrolysis of each fiber, using a Dionex ICS-3000 ion chromatography system (Thermo Fisher Scientific Inc., USA), and the results are shown in Table S1. Based on these results, glucomannan content was calculated based on galactose, glucose, and mannose contents using a 1:3 ratio of glucose:mannose,23,24 while the xylan content was calculated based on xylose and arabinose. Fiber Length and Width. A fiber tester (Lorenzen & Wettre AB, Sweden) was used to determine the weighted fiber length and width, which is calculated as the squared sum of individual fiber lengths (or widths) divided by the sum of individual fiber lengths (or widths). X-Ray Diffraction (XRD). XRD measurements were carried out using a PANalytical X’Pert alpha1 powder diffractometer (Netherlands) with CuKα radiation generated at 45 kV and 40 mA.


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Crystallinity was calculated by peak deconvolution, fitting all individual crystalline peaks and the broad amorphous peak located at approximately 21.5º, until an iteration with an R2 value of 0.997 was reached.25 The crystallite sizes (D) from the 110, 1-10, 200 lattice planes were calculated in Scherrer Equation (Equation 1), where λ is the X-ray wavelength, β is the angular full width at half maximum intensity (FWHM), and Ɵ is the scattering Bragg angle.26,27 0.9𝜆

𝐷 = 𝛽 cos 𝜃

- Equation 1

Wide-Angle X-ray Diffraction (WAXD). WAXD was performed using a single crystal X-ray diffractometer (Bruker D8 VENTURE, USA). Two-dimensional diffraction patterns were recorded by mounting the film perpendicularly to the incident beam. Based on the azimuthal intensity distribution from the equatorial (200) reflection in the diffractograms, the degree of orientation (П) and Hermans orientation parameter (𝑓) were calculated according equations 2-4.28 In equation 2, the FWHM is the full width at half maximum based on a Gaussian peak fitting. In equations 3 and 4, 𝜑 is the azimuthal angle (angle between the axis of drawing and the axis of the cellulose crystallites), and 𝐼(𝜑) is the intensity along the Debye-Scherrer ring. 𝜑 and 𝐼(𝜑) were then used to calculate the f value, where an 𝑓 = 1 corresponds to maximum orientation parallel to the direction of drawing and an 𝑓 = 0 corresponds to random orientation of the fibrils. An average of five different points is reported for each type of sample. П= 𝑓=

180−𝐹𝑊𝐻𝑀 180 3〈𝑐𝑜𝑠2 𝜑〉−1 2

〈𝑐𝑜𝑠 2 𝜑〉 =

⁄2 2 ∑𝜋 0 𝐼(𝜑) sin 𝜑𝑐𝑜𝑠 𝜑𝑑𝜑 ⁄2 ∑𝜋 0 𝐼(𝜑) sin 𝜑 𝑑𝜑

- Equation 2 - Equation 3 - Equation 4

Preparation of molded fibers. Figure 1 shows the procedure used to manufacture the final holocellulose molded fibers from spruce. Handsheets with oriented holocellulose fibers were


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prepared using a dynamic sheet former (FiberTech AB, Sweden) with a targeted weight per surface area (grammage) of 110 g/m2, using a starting pulp concentration of 2 g/L water and a drum rotation speed of 1066 m/min. Molded holocellulose fibers were obtained using a previously reported two-stage hot-pressing approach: first, 1 MPa at room temperature for 5 min, followed by compression molding at 105 °C at 15 MPa for 15 min.12 A 5 kg weight was placed on top of the final holocellulose material to limit warping during cooling. Final molded fibers were termed “MF-0R”, “MF-125R”, “MF-250R” and “MF-500R”, according to the different beating extents described above. Molded fibers based on fully bleached Kraft pulp fibers after different extents of beating were termed “K-MF-1000R” and “K-MF-3000R”. Note that a minimum beating of 1000 revolutions was used, since the mechanical properties of Kraft papers produced from lightly beaten fibers are too poor to be tested. The density and grammage of the final molded fibers were determined by measuring their dimensions and weight. For comparison, low density fiber networks were prepared by freeze-drying of wet handsheets directly after dynamic sheet formation. Porosity was estimated based on Porosity = 1 − 𝜌∗ /𝜌𝑠

- Equation 5

Where 𝜌∗ is the density of the final molded fiber, and 𝜌𝑠 is the density of the solid holocellulose estimated to 1.5 g/cm3. Preparation of PMMA/Fiber Composites. Following a reported method,21 composites were obtained by infiltrating the holocellulose MF-0R with a methyl methacrylate (MMA) monomer and oligomer solution (AIBN initiator) followed by oven curing at 50, 60, and then 70 °C for 12 h respectively. As a reference, neat PMMA material was prepared using the same procedure. Mechanical Properties. Tensile tests were performed using an Instron 5944, which is equipped with a 500 N load cell and a video extensometer. Molded fibers and composites were cut into strips


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along the desired direction (longitudinal or transverse). Molded fibers have a width of 5 mm, the gauge length was set to 25 mm and the strain rate was set to 10% per minute. Composites have a width of 7mm, the gauge length was set to 50 mm and the strain rate was set to 10% per minute. Prior to testing, the specimens were conditioned at a relative humidity of 50% and 23 °C for a minimum of 3 days. Zero-span tensile tests were performed on the freeze-dried wet handsheets using the same testing conditions described above. A minimum of 5 specimens were tested for each type of sample. Molder fiber have high in-plane orientation. Composite Mechanics. By treating the MF as laminate structure, with each ply as a crossgrained wood, and assuming no deformation mechanism effect from porosity, the effective modulus of holocellulose fibers can be estimated based on the following equations:29,30 1 𝐸𝑀𝐹

=

𝑐𝑜𝑠4 𝜃 𝐸𝐿

+

𝑠𝑖𝑛4 𝜃 𝐸𝑇

1

+ (𝐺 −

2𝑣 𝐸𝐿

) 𝑠𝑖𝑛2 𝜃𝑐𝑜𝑠 2 𝜃

2𝐺 (1 + 𝑣) = 𝐸𝑟𝑎𝑛𝑑𝑜𝑚

- Equation 6

- Equation 7

where 𝐸𝑀𝐹 is the MF longitudinal effective modulus, 𝐸𝐿 is the longitudinal fiber effective modulus, 𝐸𝑇 is the transverse fiber modulus (12 GPa),31,32 𝑣 is the Poisson ratio (0.47).31,32 𝐺is the shear modulus, which is calculated to be 6.1 GPa based on molded fibers with random-in-theplane fiber orientation (𝐸𝑟𝑎𝑛𝑑𝑜𝑚 = 18 GPa).12 Although the estimates for 𝐸𝑇 , 𝑣 and 𝐺 can be debated, the approach is still helpful, since 𝐸𝑀𝐹 depends strongly on 𝐸𝐿 . Equation 6 is based on laminate plate theory and Equation 7 is valid for isotropic materials. Scanning Electron Microscope (SEM). The nanostructure of single fibers, freeze-dried wet handsheets, molded fibers and composites were observed by field emission scanning electron microscopy (FE-SEM, Hitachi S-4300, Japan) after palladium sputtering (30 seconds to give a ca. 5 nm palladium conductive layer; sputter coater from Cressington 208HR, UK).


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Optical Properties. According to ASTM Standard D1003–13, transmittance was measured using an integrating sphere, equipped with a high brightness light source, whose spectrum spans the UV to near-IR wavelengths (170−2100 nm; EQ-99 from Energetiq Technology Inc.).

RESULTS AND DISCUSSION Figure 1 summarizes the molded fibers preparation stages. Spruce veneer was cut to smaller pieces, followed by mild PAA delignification and facile disintegration to individualized, discrete holocellulose fibers.12 Fibers are then oriented in a dynamic sheet former, where a dilute water suspension of fibers is sprayed on the wall of a rotating drum. Water is removed in two steps of compression molding. A molded fiber network of high density is formed. Finally, the network is impregnated by mildly pre-polymerized MMA, which is then cured so that a thermoplastic polymer matrix biocomposite is formed. Note that fibers were beaten to different extents (125, 250, 500 revolutions). Beating is an industrial procedure to increase inter-fiber bond strength for improved mechanical properties of paper.33,34

Figure 1. Schematic representation of the processing route from spruce to transparent molded fibers and biocomposites with highly oriented fibers. Processes include PAA delignification, PFI beating, dynamic sheet former, two compression molding steps, and polymer impregnation. To


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illustrate transparency, molded fibers and biocomposite samples were placed on top of text printed to a sheet of regular paper.

Chemical composition, fiber aspect ratio and degree of crystallinity. After beating, negligible changes were observed in chemical composition and aspect ratio (length/diameter), (Table 1). A slight decrease in the contents of lignin and hemicellulose was observed, leading to a small increase in the cellulose percentage. Detailed length and width distributions are plotted in Figure S1; lengths (~2.41 mm) and widths (~34.3 µm) were barely affected by beating. From the XRD diffractogram (Figure S2), the crystallinity of the holocellulose fibers decreased from 64% to 57% after mechanical beating (Table 1). Other studies report an initial decrease followed by an increased crystallinity with the PFI beating.35,36 Initially increased crystallinity is attributed to the removal of amorphous lignin and hemicellulose. Here, lignin was almost completely removed during the PAA delignification process. Thus, the observed decrease in crystallinity is due to mechanically induced destruction in the cellulose crystallites. The cellulose crystallite sizes from the 200, 110 and 1-10 planes were calculated by the peak deconvolution method using the Scherrer formula (Table 1). The crystallite size in the 200 plane increased slightly, whereas the crystallite sizes in the 110 and 1-10 planes increased significantly. Zhao et al. reported a similar result for PFI milled tobacco pulp fibers.36 Mechanical beating not only damages fiber surfaces, but also the cell wall structure is delaminated. Hemicellulose redistribution increases cellulose-cellulose interactions, so that nanofibrils coalescence during drying, and cellulose crystallite size is increased.37,38 These changes may decrease the mechanical properties of individual fibers, as will be discussed in further detail below.


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Table 1. Chemical composition, aspect ratio and crystallinity of holocellulose fibers and Kraft fibers after different extents of beating Crystallite size (nm)

Lignin

Cellulose

Hemicellulose

Lengtha (mm)

Widtha (µm)

Crystalli nity

200

110

1-10

Fiber-0R

2.3%

70.8%

26.9%

2.41

34.4

63%

4.0

4.5

13

Fiber-125R

1.9%

72.2%

25.9%

2.41

34.3

62%

4.2

5.1

15

Fiber-250R

1.9%

72.3%

25.8%

2.40

34.2

60%

4.3

5.4

18

Fiber-500R

1.9%

72.5%

25.7%

2.40

34.2

57%

4.6

6.4

21

K-Fiber-0R

2.1%

80.5%

17.4%

2.56

31.4

68%

4.8

7.5

20

a

Fiber lengths and widths are calculated as weight average.

Fiber Morphology. Although there is no fiber shortening after beating, fiber surface morphology changed dramatically. Figure 2A shows the original holocellulose fiber (Fiber-0R) with a smooth surface and preserved “pit” structure, whereas substantial fiber damage is observed with increased beating, due to partial unwinding of the outer-layer (mainly S1) of the fiber cell wall. Molded fiber density and porosity. The holocellulose fibers were used to prepare wet “paper” materials using a dynamic sheet former, as shown in Figure 1. All materials reached 110 g/m2 weight per unit area (“grammage”), thus indicate negligible material loss during manufacturing (Table 2). Also, all materials had a similar thickness (95 µm), density (~1.18 g/cm3) and porosity (21%), thus these parameters do not influence the comparison of mechanical and optical properties.


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Figure 2. SEM images of A) fibers after different extents of beating, B) cross-section of freezedried wet handsheets, C) surface and D) cross-section of molded fibers.

Table 2. Grammage, thickness, density, porosity and orientation of molded fibers prepared from fibers beaten to different extents. Standard deviations are reported in parentheses. Grammage Thickness (g/m2) (µm) MF-0R MF-125R MF-250R MF-500R K-MF-1000R K-MF-3000R

110 (5) 112 (5) 114 (6) 113 (5) 107 (5) 109 (5)

96 (5) 95 (5) 95 (5) 94 (5) 90 (5) 90 (5)

Density (g/m3)

Porosity

1.15 (0.05) 1.18 (0.06) 1.20 (0.05) 1.20 (0.06) 1.19 (0.05) 1.21 (0.05)

23% 21% 21% 20% 21% 19%

Degree of orientation ,П 0.78 (0.02) 0.74 (0.01) 0.73 (0.02) 0.70 (0.03) 0.65 (0.03) 0.64 (0.02)

Hermans orientation parameter, f 0.56 (0.01) 0.54 (0.01) 0.52 (0.01) 0.51 (0.01) 0.42 (0.02) 0.40 (0.02)


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Fiber Orientation in Molded fibers. The representative diffraction patterns from WAXD measurements are presented in Figure 3. The calculated results (Table 2) show that fiber orientation decreased slightly after beating, with the degree of orientation (Π) and the Hermans orientation parameter (f) reduced from 0.78 to 0.70 and from 0.56 to 0.52, respectively. The reason might be that fibers become more fibrillated i.e. hairier after beating (Figure 2A), making alignment more difficult as they may entangle more easily with other fibers. The f values presented here can be considered as a laminated structure based on holocellulose fibers with rotating angle of 20°, after eliminating the contribution from the mircrofibril angle (MFA ~ 15°) of wood single fiber itself (detailed calculations in Supporting Information).12 Figure 2B show that most fibers are highly oriented in the cross-sectional view. Although the f data for holocellulose molded fibers is lower than for oriented cellulose nanopaper (f = 0.72, cold drawing)28 and oriented cellulose nanocrystal (CNC) films (f = 0.69, wet shear casting),39 the results here are still highly notable, since the alignment of micron-scale fibers over large areas (20 cm × 100 cm) is non-trivial and has direct industrial relevance. Note that the use of dynamic sheet former to make oriented molded fibers has been reported,40,41 but a quantitative measure of fiber orientation is rarely provided.

Figure 3. WAXD diffraction patterns of molded holocellulose fibers prepared from fibers beaten to different extents.


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Structure of Molded fibers. The surface morphology and inner structure of holocellulose molded fibers were studied using SEM (Figure 2). In the freeze-dried wet handsheets, individual fibers are less distinguishable with more fibrillated entities from fiber surfaces in the inter-fiber region after beating (Figure 2B). These results are consistent with the individual fiber morphology images (Figure 2A), where fibers have a roughened, fibrillated outer surface or “shell” after beating. This also explains why beaten holocellulose molded fiber surfaces becomes smoother (Figure 2C). Rough “hairy” fibrillar surfaces can more easily fill gaps during hot-pressing from the wet-state. Also, the cross-section of molded fibers (Figure 2D) show fewer gaps in the interfiber area after beating. This microstructural feature influences mechanical and optical properties to a large extent, which will be discussed further.

Figure 4. Stress-strain curves of holocellulose molded fibers along the A) longitudinal direction and B) transverse direction; C) Zero-span tensile strength of molded fibers; D) SEM images of fracture surfaces after tensile tests; E) total optical transmittance of molded fibers.


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Table 3. Mechanical and optical properties of molded fibers based on holocellulose fibers and Kraft fibers. Standard deviations are reported in parentheses. Transmittance data are for molded fibers of 95µm thickness. Longitudinal

Transverse

Transmittance (%)a

MF-0R

Yield Strength (MPa) 87 (7)

Young’s Modulus (GPa) 27 (1)

Ultimate Strength (MPa) 260 (15)

Strain at Break (%) 1.7 (0.2)

Yield Strength (MPa) 20 (2)

Young’s Modulus (GPa) 5.5 (0.4)

Ultimate Strength (MPa) 43 (2)

Strain at Break (%) 3.1 (0.4)

MF-125R

70 (10)

22 (2)

240 (20)

2.1 (0.2)

22 (3)

4.8 (0.4)

52 (4)

3.6 (0.6)

68

MF-250R

65 (7)

20 (1)

220 (20)

2.2 (0.2)

23 (3)

4.2 (0.3)

54 (3)

4.1 (0.3)

71

MF-500R

48 (7)

17 (1)

220 (10)

2.7 (0.2)

20 (3)

4.2 (0.3)

51 (3)

5.7 (0.3)

74

K-MF-1000R

22 (4)

5.2 (0.3)

60 (5)

3.6 (0.4)

8 (1)

2.1 (0.1)

21 (1)

5.1 (0.6)

47

K-MF-3000R

40 (5)

11 (0.5)

130 (10)

3.4 (0.3)

8 (1)

1.6 (0.1)

24 (1)

5.4 (0.5)

60

a

67

Transmittance is reported as average values in the wavelength range of 550−900 nm

Mechanical Properties of Molded Fibers. Figures 4A and 4B and Table 3 show longitudinal (parallel to preferred fiber direction) and transverse stress-strain behavior of molded fibers loaded in uniaxial tension. In the longitudinal direction, MF-0R has a yield strength of 87 MPa, a Young’s modulus of 27 GPa and an ultimate strength of 260 MPa. This exceptionally high mechanical performance surpasses earlier data including “molded” fibers based on high purity sulfite fibers13 or Kraft pulp fibers,42 binderless fibreboard based on microfibrillated cellulose,14 and Zelfo type materials based on hemp/flax.16 Moreover, it is comparable with oriented cellulose nanopaper.28,40 Holocellulose molded fibers with random-in-the-plane fiber orientation has a Young’s modulus of 18 GPa and an ultimate strength of 195 MPa.12 The present study thus shows 50% increase in Young’s modulus, and 33% increase in tensile strength compared with identical materials in ref 11, but with random-in-the-plane orientation. One reason for high mechanical properties is in the intrinsic properties of the fibers. Based on the WAXD data for cellulose orientation, the present material can be viewed as a mechanical analogue to a laminate of unidirectional lamellae, with an orientation of ± 20° to the direction of


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tensile testing. The 27 GPa modulus of this laminate, corresponds to an “effective” longitudinal fiber modulus of 36 GPa, see the composite mechanics analysis in the Experimental section. Based on this model, if we could improve the alignment of the present fibers, the maximum modulus we can expect is 36 GPa. This estimate seems reasonable, since it is similar to previous data reported for softwood cell walls with a microfibril angle MFA of 15°.43 As another reference, densified delignified spruce veneer was reported to have a Young’s modulus of 35 GPa (~20% porosity).44 The well-preserved native cell wall morphology and well-preserved cellulose/hemicellulose components are also important. The present holocellulose fibers have high crystallinity (Table 1), high cellulose molar mass,12 and no visible cell wall delamination. The preservation of native hemicellulose helps to maintain cell wall integrity, and to prevent cellulose nanofibril aggregation after drying, as supported by preserved cellulose crystallite size (Table 1). Here, cellulose fibrils are well-dispersed and surrounded by hemicellulose in the cell wall. In comparison, fibril agglomeration is often a problem for a high-purity cellulose nanofibril (CNF) film. The role of hemicelluloses is interesting. They have been used as an added bonding agent to paper.45,46 The fiber pull-out lengths in holocellulose MF-0R fracture surfaces are very short (Figure 4D). This correlates with strong fiber-fiber bonding. Furthermore, it is likely that the fiber network integrity during progressive failure events is better preserved due to the role of hemicelluloses. Beating is commonly used to improve paper properties,33–35 but for the present materials the effect was the reverse. After beating, the yield strength decreased from 87 to 48 MPa, Young’s modulus decreased from 27 to 17 GPa, and the ultimate strength decreased from 260 to 220 MPa. The decrease in Young’s modulus after beating is due to cell wall delamination, and possibly decreased fiber crystallinity (Table 1). Decreased ultimate MF strength after beating is a result of


17


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decreased fiber intrinsic strength, as confirmed by Zero-span tensile tests (Figure 4C). Although the present extent of beating is much lower (≤ 500 revolutions) than typical in industry, it was sufficient to significantly damage the fibers. In the transverse direction, MF-0R has a Young’s modulus of 5.5 GPa and an ultimate strength of 43 MPa (Figure 4B and Table 3), which is lower than expected. The reason is the presence of fiber-fiber interface debond cracks, see Figure 2D. With increased beating, the molded fibers exhibit a decreased “transverse” Young’s modulus, which is attributed to cell wall damage so that the intrinsic transverse cell wall modulus is decreased. The ultimate transverse strength of molded fibers increases slightly after beating, which may be due to increased fiber orientation in the transverse direction (Table 2). Optical properties of molded fibers. Molded holocellulose fibers have an optical transmittance of ~70% at 95 µm thickness (Figure 4E and Table 3), and their appearance is presented in Figure 1. Beating improves the transmittance of the molded fibers, since the density of the material is increased as fibers become more deformable from beating and scattering voids are removed. Beaten fibers have roughened surfaces and this also reduces the presence of interface gaps/voids (Figure 2C and 2D). Comparison with molded Kraft fibers. Hot-pressed industrial pulp fibers (bleached Kraft) were selected as a molded fiber reference material. The Kraft fibers have similar aspect ratio as holocellulose fibers, but the hemicellulose content is lower (Table 1 and 2). This molded Kraft fiber reference (K-MF-3000R) has a Young’s modulus of only 11 GPa and an ultimate strength of only 134 MPa even after beating for 3000 revolutions (Table 3), yet porosities are comparable. Although the orientation is not as strong as for molded holocellulose fibers (MF-0R, Table 2), other important reasons for inferior performance are lower intrinsic fiber strength and lower fiber-


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fiber interface bonding due to chemical degradation during Kraft pulping. Moreover, K-MF-3000R has lower optical transmittance (60%) than molded holocellulose fibers at the same thickness. Most likely, reasons include more aggregated cell wall fibrils, and more interface gaps at microscale. Thermoplastic MF-0R/PMMA biocomposites. Highly transparent MF/PMMA biocomposites were prepared based on molded holocellulose fibers (MF-0R) as a reinforcing fiber network. The MF was simply impregnated by MMA monomers and oligomers, followed by in-situ polymerization (Figure 1). The volume fraction of holocellulose fibers in the MF/PMMA biocomposite was 52%, corresponding to a cellulose content of 58 wt% with an estimated porosity of only 4%.

Figure 5. A) Total optical transmittance of PMMA, molded holocellulose fibers (MF-0R) and MF0R/PMMA biocomposite; B) Fractography images of MF-0R/PMMA biocomposite fracture surfaces, C) SEM images of freeze-fractured cross-sections of MF-0R/PMMA; D) Stress-strain curves of PMMA, MF-0R molded fibers and MF-0R/PMMA biocomposite in the longitudinal (L) and transverse (T) directions.


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Table 4. Literature data for thickness, cellulose content, transmittance, Young’s modulus and ultimate strength of neat PMMA and various cellulose/PMMA biocomposites.

Neat PMMA

140

Cellulose content (wt%) -

MF-0R/PMMA

140

58

90%

20 (2)

310 (30)

This study

K-MF-3000R/PMMA

140

58

87%

9 (1)

180 (10)

This study

Wood/PMMA

1200

23

80%

4

90

21

Pulp Fibers/PMMA

50

28

88%

-

-

47

Cotton powder/PMMA

-

56

88%

-

-

48

CNF/PMMA

200

5

90%

2.1

57

49

CNF/PMMA

100

38

90%

5.1

88

50

CNF/PMMA

53

62

30%

-

-

2

CNC/PMMA

-

0.5

80%

-

-

51

CNC/PMMA

500

20

80%

2.5

50

52

BC/PMMA

65

65

80%

21

325

1

Thickness (µm)

96%

Young’s modulus (GPa) 3.2 (0.4)

Ultimate strength (MPa) 61 (4)

This study

Transmitt ance

REF

Biocomposite Optical Properties. The thermoplastic MF biocomposites show significantly improved optical transmittance (90% at 140 µm thickness) compared with the neat molded fiber material MF-0R (67% at 95 m thickness) (Figure 5A). Neat PMMA has a transmittance of 96% at the same thickness, which means that the fiber network has only a small effect on transmittance. Three main reasons are contributing: First, PMMA forms smooth surface layers on MF-0R to reduce surface reflection (Figure 5B). Secondly, PMMA fills gaps/voids throughout the fiber network including the middle section (Figure S3), especially within inter-fiber regions (Figure 5C), resulting in biocomposites with porosities below 4%. The freeze-fractured surfaces in Figure 5C indicate minimal debonding between the polymer phase and fibers. PMMA penetrates inside the cell wall lumen (probably through pits), see oval contours of fibers in Figure 5C. Less light scattering occurs as voids are removed, resulting in high transmittance. It is also possible that holocellulose fibers themselves show lower cell wall scattering than Kraft fibers due to less


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aggregation of cellulose fibrils.20 Thirdly, the refractive index of PMMA (~1.49) is a fairly good match of cellulose (~1.53), which helps to reduce scattering at the cellulose/polymer interface. Table 4 shows that the present biocomposites have the highest transparency and one of the highest cellulose contents (58 wt%) when compared with literature data for other cellulose/PMMA biocomposites such as delignified wood veneer,21 acetylated cotton powder,48 acetylated pulp fibers,47 BC,1 CNFs,2,49,50 and CNCs.51,52 Biocomposite Tensile Tests. Figure 5D presents stress-strain curves for MF-0R/PMMA biocomposites, molded holocellulose fibers (MF-0R) and neat PMMA. The polymer matrix biocomposite has a Young’s modulus of 20 GPa and an ultimate strength of 310 MPa at a fiber content of 52 vol%. In transverse direction, Young’s modulus is 10 GPa and strength 85 MPa. The biocomposite shows higher ultimate strength compared with MF-0R (260 MPa in the longitudinal direction and 43 MPa in the transverse direction). Although the yield strength is lower for the polymer matrix biocomposite compared with MF-0R, the strain to failure becomes higher with the polymer. PMMA delays final fracture, possibly by increasing the crack growth toughness for microscale debond cracks so that damage development is shifted to higher strains. The interface between the fibers and polymer matrix appears to be intact in Figure 5C. Strong fiber/polymer interface is also indicated by short fiber pull-out lengths (below 100µm) at the fracture surfaces from room temperature tensile tests (Figure 5B). Additionally, a recent simulation study indicates that greater cellulose/PMMA adhesion energy occurs at amorphous cellulose surface other than crystalline cellulose surface. In present study, holocellulose fibers with high hemicellulose content may promote better cellulose/PMMA interface.53 Note that the Kraft-3000R/PMMA biocomposite based on molded Kraft fibers has much lower mechanical properties, although it has similar cellulose volume fraction and porosity (Table 4 and Figure S4). Moreover, the present


21


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biocomposites have much better mechanical properties than most of other cellulose/PMMA biocomposite, 21,49,50,52 and are comparable to BC/PMMA biocomposites.1 This is related to the high mechanical properties of the molded, oriented holocellulose fiber network, and the strong fiber/matrix interface.

CONCLUSIONS Highly oriented, molded holocellulose wood fiber materials, with a porosity of around 20%, showed exceptionally high Young’s modulus (27 GPa) and ultimate strength (260 MPa), combined with an optical transmittance of ~70% at 95 µm thickness. These materials have potential in semistructural load-bearing applications, are fully native, recyclable and biodegradable, with delignification as the only chemical treatment. Based on the hot-pressed molded fiber networks, thermoplastic and thermoformable PMMA biocomposites with 52 vol% of nanostructured holocellulose fibers are prepared with only a few percent porosity. Ultimate strength is increased (310 MPa) and Young’s modulus is still 20 GPa. With this polymer matrix, optical transmittance becomes as high as 90% at 140µm thickness, since refractive indices are fairly similar for polymer and cellulose, and the light-scattering from voids is strongly reduced. Although the measured average fiber orientation can partly explain the mechanical property data, a key is in the preservation of native cellulose fibril structure (molar mass, long-range order), and hemicellulose distribution in the wood fiber cell wall (no fibril-fibril aggregation). Furthermore, fiber-fiber joints are strong and hemicellulose-rich in the compressed wood fiber network, so that the high intrinsic fiber properties are favorably transferred into the mechanical properties of the molded fibers. When a continuous thermoplastic polymer matrix phase is added, stress transfer is further improved and micro-scale damage development is delayed in the thermoplastic


22

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biocomposite. As a consequence, strain to failure is increased from 1.7% for the molded fibers to 3.1% for the biocomposite. Since stress keeps increasing with increasing strain, the biocomposite reaches an ultimate strength of 310 MPa, which is exceptionally high for a thermoplastic matrix wood fiber biocomposite. For these high-density molded fibers, established structure-property relationships from the paper and board literature no longer apply. The reason is that the porosity of the present materials is so much lower, and stress transfer mechanisms are different. As an example, beating of the fibers decreases rather than increases mechanical properties. The reason is mechanical fiber damage, and this shows that intrinsic mechanical properties of the fibers are critical for the present materials. The holocellulose materials are not only vastly superior to comparable materials from regular bleached wood fibers, but they are also comparable with data for materials from cellulose nanofibrils. One reason is the well-preserved nanostructural organization of the fiber cell wall, in contrast to industrial fibers with aggregated cellulose fibrils. Cellulose crystallite size is unchanged after chemical processing, and nanoscale fibrils are well-dispersed in a continuous matrix phase of well-preserved hemicellulose. From processing and economical points of view, wood fibers have strong advantages compared with cellulose nanofibrils, due to much lower cost and much lower suspension viscosity. The concepts presented in the study do not only show the potential of wood fibers in semi-structural and ecofriendly materials, they also suggest milder processing routes towards novel grades of wood cellulose fibers.


23


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ASSOCIATED CONTENT The Supporting Information is available free of charge on the ACS Publications website. Relative carbohydrate composition; length and width distributions of fibers after beating; XRD diffractograms of fibers after beating; calculation of fiber orientation from WAXD results; SEM image of MF-0R/PMMA composite; stress-strain curves of neat PMMA and biocomposites.

AUTHOR INFORMATION Corresponding Author *Email: [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT We acknowledge funding from the Knut and Alice Wallenberg foundation through the Wallenberg Wood Science Center at the KTH Royal Institute of Technology. We thank Kerstin Slettengren (RISE) and Hui Chen for sample preparation and analysis. Useful discussions with Tiffany Abitbol (RISE), Michael Reid and Erik Jungstedt are gratefully acknowledged.


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