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BIOCOMPOSITES BASED ON
CORE-SHELL CELLULOSE
NANOFIBERS
Preparation, Structure and Properties
Kasinee Prakobna
Doctoral Thesis
Wallenberg Wood Science Center
Department of Fiber and Polymer Technology
School of Chemical Science and Engineering
KTH Royal Institute of Technology
Stockholm, 2015
Supervisor
Professor Lars A. Berglund
Copyright © 2015 Kasinee Prakobna
All Rights Reserved.
ISBN 978-91-7595-513-1
ISSN 1654-1081
TRITA-CHE Report 2015:18
Tryck: US-AB, Stockholm 2015
AKADEMISK AVHANDLING
Som med tillstånd av Kungliga Tekniska Högskolan i Stockholm framlägges till
offentlig granskning för avläggande av teknologie doktorsexamen den 26 maj
2015, kl. 10.00 i sal F3, Lindstedtsvägen 26, KTH, Stockholm. Fakultetsopponent:
Dr. David Plackett, The University of British Columbia, Canada. Avhandlingen
försvaras på engelska.
“Strive not to be a success, but rather to be of value.”
Albert Einstein
ABSTRACT
Cellulose nanofibers (CNFs) are of interest as load-bearing components for
polymer matrix nanocomposites. A wide range of nanostructured materials
including nanopaper/films, foams, aerogels, and hydrogels can be prepared from
CNFs. The material performance can be fine-tuned when CNFs are combined
with different polymer matrices. The main idea of the present study is to test the
hypothesis that the concept of Core-shell CNFs can provide processing and
performance advantages for nanocomposites through improved nanostructural
control.
The problems of matrix distribution and interface structure at nanoscale are
investigated. The first part of this thesis (Paper I-III) describes an alternative
preparation procedure for biocomposites based on the Core-shell concept.
Inspired by the structural framework and mechanical function of the primary cell
wall in plants, Core-shell CNFs are formed by coating wood CNFs with a
polysaccharide matrix, and are subsequently used for fabrication of biocomposite
films. The nanostructure of Core-shell CNFs and their nanocomposites is
characterized. Mechanical properties of the biocomposites at various hydration
conditions are investigated. A study on molecular water mobility and moisture
stability of the materials is conducted. In the present thesis, three different
biological polysaccharides including amylopectin, xyloglucan and
galactoglucomannan have been used for man-made nanocomposites based on
Core-shell nanofibers.
The later sections of the thesis (Paper IV-V) describe an alternative method to
disintegrate holocellulose nanofibers from wood chips. These CNFs based on
holocellulose possess Core-shell structure with native hemicelluloses as “shell”.
Peracetic acid pretreatment is used for the preparation of holocellulose. This
procedure is a promising single-step pulping as preparation for nanofibrillation.
The nanostructure of the holocellulose nanofibers are characterized, and
compared with CNFs prepared from enzymatic pretreatment. The holocellulose
nanofibers are used for preparation of films and porous materials. The influence
of nanostructural characteristics of two different nanofibers (CNFs from
enzymatic pretreatment vs holocellulose nanofibers from peracetic acid
pretreatment) on final properties of the nanocomposites is clarified.
Favorable characteristics of Core-shell fibrils are reported in terms of colloidal
stability, controlled matrix distribution, improved interface characteristics, and
improved hygromechanical properties of the biocomposites. In both a scientific
and industrial context, the Core-shell nanofiber concept thus offers great potential
for the materials design of new cellulose nanomaterials with unique
characteristics.
Keywords: biocomposites, cellulose nanofibers, polysaccharide, holocellulose,
nanostructure, dispersion, mechanical properties, hygromechanical properties,
moisture stability, film, porous material
SAMMANFATTNING
Nanofibrer från cellulosa (CNF) är av stort intresse som lastbärande
komponent i nanokompositer och ett flertal nanostrukturerade material såsom
nanopapper, nanoskum, aerogeler eller hydrogeler kan framställas från CNF.
Materialegenskaperna kan finjusteras då CNF tillsätts till en polymermatris.
Avhandlingen handlar om ett koncept för att kontrollera matrisfördelning och
gränsyta mellan CNF och polymermatrisen, i form av sk ”Core-shell”-CNF. Första
delen (artikel I-III) handlar om ”Core-shell”-CNF framställda genom att adsorbera
en polysackarid på ytan av CNF. Nanostrukturen hos ”Core-shell”-CNF och dess
kompositmaterial undersöks, liksom mekaniska egenskaper vid olika
hydreringstillstånd. Molekylrörlighet hos vatten studeras med NMR. Tre olika
polysackarider undersöks; amylopektin, xyloglukan och galaktoglukomannan.
Avhandlingens andra del (artikel IV och V) handlar om CNF från holocellulosa,
med vedflis som råvara. Holocellulosa innebär att huvudsakligen lignin har
extraherats från veden, medan cellulosa och hemicellulosa är välbevarade. Denna
CNF har också ”Core-shell”-struktur med hemicellulosa som ytskikt. Perättiksyra
används som en mild framställningsmetod. Det innebär en process i ett enda steg
för att framställa massa med bevarad cellulosastruktur. Nanostruktur och kemisk
sammansättning karakteriseras och jämförs med enzymatiskt framställd CNF
från sulfitmassa. CNF från båda typerna av nanocellulosa (holocellulosa och
enzymatisk CNF) används för att framställa kompositer och porösa material
(honeycombs och skum). Sambanden mellan struktur och materialegenskaper
klarläggs, liksom effekterna från olika typer av CNF. Avhandlingen visar att
”Core-shell”-CNF är ett fungerande koncept för nanostrukturerade kompositer
genom påvisande av ett flertal lovande och fördelaktiga egenskaper såsom
kolloidal stabilitet, kontrollerad matrisfördelning och förbättrade gränsyte- samt
hygromekaniska egenskaper. Såväl vetenskapligt som industriellt kan ”Core-
shell”-CNF bidra till utvecklingen av nya nanomaterial baserade på cellulosa.
LIST OF PAPERS
This thesis is based on the following five appended papers:
High-Performance and Moisture-Stable Cellulose-Starch Nanocomposites I.
Based on Bioinspired Core-Shell Nanofibers
Kasinee Prakobna, Sylvain Galland, and Lars A. Berglund (2015),
Biomacromolecules, 16, pp 904-912.
Core-Shell Cellulose Nanofibers for Biocomposites – Nanostructural II.
Effects in Hydrated State
Kasinee Prakobna, Camilla Terenzi, Qi Zhou, István Furó, and Lars A.
Berglund (2015), Carbohydrate Polymers, 125, pp 92-102.
Strong Effects from Galactoglucomannan Hemicelluloses on Mechanical III.
Behavior of Wet Cellulose Nanofiber Gels
Kasinee Prakobna, Victor Kisonen, Chunlin Xu and Lars A. Berglund,
Submitted manuscript.
Holocellulose Nanofibers of High Molar Mass and Small Diameter for IV.
High-Strength Nanopaper
Sylvain Galland, Fredrik Berthold, Kasinee Prakobna, and Lars A. Berglund,
Submitted manuscript.
Mechanical Performance and Architecture of Biocomposite Honeycombs V.
and Foams from Core-Shell Holocellulose Nanofibers
Kasinee Prakobna, Fredrik Berthold, and Lars A. Berglund, Manuscript.
The author’s contributions to the appended papers are as follows:
I. All of the experimental work and data analysis except TEM observation,
and all of the manuscript preparation.
II. All of the experimental work and data analysis except 2H NMR relaxometry,
and most of the manuscript preparation.
III. All of the experimental work and data analysis except extraction and molar
mass analysis of GGM, and all of the manuscript preparation.
IV. Part of the experimental work and data analysis include preparation of
holocellulose nanofibers (peracetic acid pretreatment), observation on
nanostructure and light transmittance of the materials, and most of the
manuscript preparation.
V. All of the experimental work and data analysis, and all of the manuscript
preparation.
Other relevant papers not included in this thesis:
Composite Films of Nanofibrillated Cellulose and O-Acetyl-VI.
Galactoglucomannan (GGM) Coated with Succinic Esters of GGM
Showing Potential as Barrier Material in Food Packaging
Victor Kisonen, Kasinee Prakobna, Chunlin Xu, Arto Salminen, Kirsi S.
Mikkonen, Dimitar Valtakari, Patrik Eklund, Jukka Seppälä, Maija
Tenkanen, and Stefan Willför (2015), Journal of Materials Science, 50, pp 3189-
3199.
Nanostructural Effects on Polymer and Water Dynamics in Cellulose VII.
Biocomposites – 2H and 13C NMR Relaxometry Camilla Terenzi, Kasinee Prakobna, Lars A. Berglund, and István Furó
(2015), Biomacromolucules, DOI: 10.1021/acs.biomac.5b00330.
LIST OF ABBREVIATIONS
AFM Atomic force microscopy
AP Amylopectin
CNF(s) Cellulose nanofiber(s)
DP Degree of polymerization
DLS Dynamic light scattering
DMAc Dimethyl acetamide
DSP Digital speckle photography
DVS Dynamic vapor sorption
E Young’s modulus
Es Young’s modulus of the solid cell wall material
EIC Ethyl isocyanate
FE-SEM Field-emission scanning electron microscopy
GGM O-acetyl-galactoglucomannan
HPAEC High performance anion exchange chromatography
MALLS Multi angle laser light scattering
Mn Number average molar mass
Mw Weight average molar mass
NMR Nuclear magnetic resonance
PAD Pulsed amperometric detector
PDI Polydispersity index
PTFE Polytetrafluoroethylene
SEC Size exclusion chromatography
SC-CO2 Supercritical carbondioxide
TEM Transmission electron microscopy
TEMPO Tetramethylpiperidine-1-oxyl
TOCNs TEMPO-oxidized cellulose nanofibers
XG Xyloglucan
CONTENTS
INTRODUCTION .................................................................................................. 1 1
1.1 Objective .......................................................................................................... 2
BACKGROUND .................................................................................................... 3 2
2.1 Wood cell walls .............................................................................................. 3
2.1.1 Structural biopolymers .................................................................... 3
2.1.2 Ultrastructure of the cell walls ........................................................ 5
2.2 Cellulose nanofibers from wood .................................................................. 7
2.2.1 Pretreatment approaches ................................................................. 8
2.2.2 Reinforcement materials .................................................................. 9
2.3 Biocomposite films based on cellulose nanofibers .................................. 10
2.4 Core-shell cellulose nanofibers and their biocomposite films ............... 13
2.5 Holocellulose nanofibers and their biocomposite films ......................... 15
2.6 Application of nanofibers in porous materials ........................................ 16
EXPERIMENTAL................................................................................................. 19 3
3.1 Material components ................................................................................... 19
3.1.1 Cellulose nanofibers (CNFs) prepared by enzymatic pretreatment
(Paper I-V) ......................................................................................19
3.1.2 Holocellulose nanofibers prepared by peracetic acid pretreatment
(Paper IV-V) ...................................................................................19
3.1.3 Polysaccharide matrices (Paper I-III) .............................................20
3.2 Preparation of biocomposite materials ..................................................... 21
3.2.1 Cellulose and holocellulose nanopaper (Paper I-IV) ....................... 21
3.2.2 Neat polysaccharides (AP and XG) films (Paper I-II) ................... 21
3.2.3 Adsorption experiment and preparation of Core-shell nanofibers
(Paper I-III) .................................................................................... 21
3.2.4 Nanocomposite films (Paper I-III).................................................. 22
3.2.5 Porous materials based on nanofibers (Paper V) ............................ 23
3.3 Characterization methods ........................................................................... 25
3.3.1 Determination of matrix contents by sugar analysis (Paper I-III). 25
3.3.2 Colloidal behavior by dynamic light scattering (DLS) (Paper I-III)25
3.3.3 Total charge determination (Paper IV) .......................................... 25
3.3.4 Nanostructure of nanofibers and their networks (Paper I-V) ........ 26
3.3.5 Density and porosity (Paper I-V) ................................................... 26
3.3.6 Tensile mechanical properties (Paper I-IV) .................................... 27
3.3.7 Mechanical properties in compression (Paper V) ........................... 27
3.3.8 Dynamic vapor sorption (DVS) (Paper I-III)................................. 27
3.3.9 Molecular moisture dynamic by 2H solid-state NMR (Paper II) ... 28
3.3.10 Molar mass analysis of nanofibers by SEC (Paper IV) .................. 28
3.3.11 Molar mass analysis of matrices by SEC-MALLS (Paper I-II) ...... 29
3.3.12 Light transmittance (Paper IV) ...................................................... 29
RESULTS AND DISCUSSION ........................................................................... 31 4
4.1 Core-shell nanofibers and their biocomposite films (Paper I-III) ........... 31
4.1.1 Preparation of Core-shell nanofibers with polysaccharide matrices31
4.1.2 Nanostructure of the nanofibers and their networks ..................... 33
4.1.3 Hygromechanical properties of Core-shell vs “Mixed”
nanocomposites ...............................................................................41
4.1.4 Moisture sensitivity of Core-shell vs “Mixed” nanocomposites ....49
4.2 Conclusions (from Paper I-III) .................................................................... 54
4.3 Holocellulose nanofibers and their biocomposite films (Paper IV-V) .. 56
4.3.1 Holocellulose nanofibers .................................................................56
4.3.2 Films from holocellulose nanofibers ................................................60
4.3.3 Porous materials from holocellulose nanofibers ..............................62
4.4 Conclusions (from Paper IV-V) .................................................................. 68
FUTURE WORK .................................................................................................. 71 5
ACKNOWLEDGEMENTS ................................................................................. 73 6
REFERENCES ...................................................................................................... 75 7
INTRODUCTION 1
INTRODUCTION 1
Wood fibers, which are a widely available bioresource, have been exploited
for sustainable resource products such as paper packaging, textiles and wood
composites. To extend the traditional use of the wood fibers towards high-
performance materials, it is essential to utilize the full potential of their nanoscale
building blocks. Interest in cellulose nanofibers (CNFs) has increased as is
apparent both in academic and industrial research.1, 2 The growing interest in
CNFs is primarily due to their small dimensions, high aspect ratio and ability to
form entangled network in addition to outstanding mechanical properties. The
potential of wood CNFs has therefore been considered for advanced load-bearing
applications. Various approaches to disintegrate CNFs from wood fibers have
been explored.3-7 However, special attention to CNF characteristics and their
reinforcement potential is required. Inherent properties and the nanostructure of
disintegrated CNFs will decisively influence the final properties of
nanocomposites.
Nanocomposites of CNF reinforced natural biopolymers such as starch and
hemicelluloses have been progressively developed,8-13 primarily to overcome
poor mechanical behavior of the neat biopolymers (brittleness), low moisture
stability and inferior film-forming properties. Potential application of these
polysaccharide-based nanocomposites includes packaging materials.8, 11, 14-17
Favorable interfacial interactions between CNFs and the polysaccharide matrix
could be expected due to chemical similarities of the polymers. Since CNFs can
potentially offer large specific interface area in the nanocomposites, the
molecular and nanoscale control of interfacial phenomena between components
is necessary. In biological materials, plant cell walls can be viewed as models for
polysaccharide nanocomposites. The combination of mechanical stiffness,
strength, toughness and interfacial stability despite a highly hydrated state is
achieved within their nanostructural network. As compared with the natural
plant cell walls, it is important to note that typical man-made nanocomposites
possess poor nanostructural control of matrix distribution. In the present thesis,
an attempt has been made to coat wood CNFs with the polysaccharide matrix,
forming “Core-shell” nanofibers. Therefore, better control of the matrix
2 INTRODUCTION
distribution at nanoscale, as in the plant cell walls, is expected in the resulting
nanocomposites. The reinforcement effect and moisture stability of the Core-shell
nanocomposites have been compared with “Mixed” nanocomposites, prepared by
conventional mixing, and with a polymer matrix less favorably distributed.
1.1 Objective
The scientific objective of the present thesis is to develop preparation methods
for Core-shell nanofibers and characterize their structure. In addition, the objective
is to investigate processing-structure-property relationships for biocomposite
materials based on the Core-shell CNFs, with emphasis on nanostructural aspects.
In the first section (Paper I-III), the main purpose is to address the problem of
how nanostructural features of polymer matrix distribution and
interface/interphase is influencing hygromechanical properties and moisture
stability for the Core-shell nanocomposites.
The major focus in the second section (Paper IV-V) is related to wood
holocellulose nanofibers, which possess native Core-shell structure. Peracetic acid
pretreatment as well as key characteristics of the obtained holocellulose
nanofibers is of particular interest. In addition, the potential of how these
holocellulose nanofibers can be utilized for high-performance film and porous
materials is investigated.
BACKGROUND 3
BACKGROUND 2
2.1 Wood cell walls
Wood is of interest for many branches of industrial applications. The
molecular interactions among wood components and their arrangement on the
nanolevel have important implications and contribute to the remarkable
mechanical properties of wood.18, 19 The wood cell walls can be considered as a
naturally occurring biocomposite materials, possessing exceptional mechanical
properties of high modulus and strength combined with toughness and
flexibility. It can be transformed into nanofibers and that serve as building blocks
for nanomaterials. The architectural details of their structural biopolymer
organization and the nanostructural design resulting from biosynthesis are truly
remarkable and can inspire the development of man-made nanocomposites.
2.1.1 Structural biopolymers
Polysaccharides are the major biopolymeric components in the wood cell
walls. Cellulose, which is the primary load-bearing biopolymer in plants, is a
major structural component of the cell wall. Cellulose is a linear chain of many
thousand glucose units. They are covalently linked by β-1,4-glycosidic bonds,
and their repeat unit is cellobiose consisting of two glucose rings (Figure 1(a)).
The degree of polymerization (DP) for cellulose from wood can be as high as
10,000.20 The cellulose chains are arranged in parallel with an extended-chain
conformation, and are co-crystallized, forming microfibrils.21 The diameter of an
individual microfibril is about 3-4 nm, and the fibrils form microfibril aggregates
with diameters of approximately 20 nm.22, 23 Extensive hydrogen bonding, which
is formed within and between the cellulose chains of microfibrils, contributes to
their unique properties of mechanical strength and chemical stability. Specific
tensile strength of the microfibrils surpasses that of many natural and man-made
polymeric fibers.24 Thus, the microfibrils, which can be extracted from the cell
wall in the form of CNFs, offer the possibility to improve the mechanical
properties of a wide range of polymeric materials.
4 BACKGROUND
Figure 1. Chemical structure of (a) cellulose, (b) xyloglucan (XG) and (c) O-acetyl-
galactoglucomannan (GGM). They are the principal structural polysaccharides in
spruce cell walls.
The second most abundant group of polysaccharides in wood cell walls is
hemicelluloses. The content of hemicelluloses in wood varies in the range of 25-
40 wt%.25 In contrast to cellulose, hemicelluloses are hetero-polysaccharides,
consisting of linear homo- or copolymer backbone with varying degrees of
branching.26 They can consist of different monosaccharides such as glucose,
mannose, xylose, galactose and arabinose. The degree of polymerization (DP) for
hemicelluloses in woods probably varies between 50-200,20 which is much lower
than for cellulose. This is consistent with their function to physically link
cellulose microfibrils in hydrated state.
Type and content of hemicelluloses vary depending on the plant. For
softwoods such as spruce, xyloglucan (XG) and galactoglucommannan (GGM)
are the major hemicelluloses in primary and secondary cell walls, respectively.
XG has a backbone resembling cellulose with the exception that it is regularly
substituted with α-D-xylopyranose branches at the O-6 position. A minority of
this branched moiety is further substituted with β-D-galactopyranose.27 An
BACKGROUND 5
example of chemical structure for XG is presented in Figure 1(b). The ratio of
these monosaccharide constituents and their pattern of substitution varies
depending on plant origin of XG.28 Apart from XG in the primary cell wall,
storage XG (for instance present in tamarind seeds) is available in larger quantity
and shows potential for industrial applications. XG from tamarind seed can be
employed as an alternative high performance biopolymer.29
In the secondary cell wall, galactoglucomannan (GGM) is a typical
hemicellulose for softwoods. In the case of spruce, the GGM content ranges
between 20-25 wt%.20, 25 The main chain of GGM is built from randomly
distributed β-1,4-D-mannopyranosyl and β-1,4-D-glucopyranosyl units. The
hydroxyl group at C-2 or C-3 position of the mannose unit is partly substituted
by acetyl groups. Some of the mannose units are further substituted by α-1,6-D-
galactopyranosyl units (Figure 1(c)). The ratio of theses monosaccharide
constituents (mannose:glucose:galactose) varies in the range of 3.5-4.5:1:0.5-1 for
spruce GGM.20, 30, 31 Large fractions of water-soluble GGM can be extracted from
spruce wood.32, 33 There is potential to use GGM in packaging films due to its low
oxygen34 and water permeability.35
2.1.2 Ultrastructure of the cell walls
Wood cell walls have a hierarchical structural network consisting of several
layers and different constituents. The exceptional mechanical properties of the
cell walls are governed by the structural arrangement at multiple scales.18 The
ultrastructure on a nanoscopic to microscopic scale is suggested in Figure 2. The
cell walls provide a supportive framework assembled from reinforcing semi-
crystalline cellulose microfibrils. They are embedded within an amorphous and
hydrated matrix of hemicelluloses and lignin. These polymers are intimately
mixed on a molecular and nanostructural scale. The principal framework is
analogous to man-made composites such as glass fibers in a plastic matrix,36
although the molecular and nanostructural design features are much more
sophisticated in the biological composites.
6 BACKGROUND
Figure 2. Schematic illustration of the hierarchical structure in wood cell walls and in
the primary plant cell walls (adapted from Carpita and McCann 2000, and Gibson
2012).18, 36
The chemical composition and nanostructure of the cell walls vary in different
wood species and cell wall types. Primary cell walls consist of a thin layer of
randomly arranged microfibrils in a network, which provides greater ductility to
the cell walls. The structural network of primary cell walls may contain
approximately equal contents of cellulose and xyloglucan (XG), embedded in an
independent network of pectin.36 XG coats and provides physical cross-links to
cellulose microfibrils. There are three distinct domains of XG in the cell walls.37 It
corresponds to the domain where XG forms the cross-links between cellulose
microfibrils. In addition, XG is intimately associated on the surface, and may
even be entrapped within the microfibrils. The arrangement of microfibrils and
XG in the primary cell walls is depicted in Figure 2. The supramolecular
organization of XG plays an important role in regulating mechanical properties
of the primary cell walls.38 XG shows strong molecular interaction and adhesion
to cellulose even at hydrated conditions.39 The cell walls can become extensible
when XG is incorporated into the network.38, 40
Secondary cell walls are rigid and composed of three different layers (S1, S2,
S3). The mechanical properties of wood cell walls are largely governed by the
cellulose-rich middle secondary layer (S2), where the microfibrils are highly
oriented. Considering polymeric matrix in the secondary cell walls, a network of
glucomannan-lignin-xylan is formed in softwoods.41, 42 Hemicelluloses are,
however, closely associated with cellulose through hydrogen bonding.43 They
influence flexibility and rigidity of the structural network.44 Hemicelluloses may
BACKGROUND 7
also serve as a nanostructural interphase between cellulose and lignin. A
variation in chemical structure and arrangement of hemicelluloses is important
for tailoring physical properties of the cell walls.
The favorable interactions between cellulose and hemicelluloses at the
ultrastructural scale can also be considered in the context of cell wall assembly
and design. Cellulose chains are synthesized within rosettes at the plasma
membrane surface.45 The rosettes act as nanospinnerets extruding the cellulose
chains to form a self-assembling microfibril. Hemicelluloses are synthesized in
the Golgi apparatus, and diffuse towards the plasma membrane. After secretion,
the cellulose and hemicelluloses are associated into an integrated network of cell
walls.46 The coalescence and further aggregation of microfibrils is prevented by
the presence of hemicellulose as a coating layer,46 which also serves as a hydrated
lubricating boundary layer at the nanoscale.
The assembly of cell wall components takes place in an aqueous environment.
It is considered that the primary cell walls contain up to 75-80 wt% of water,
which is primarily located in the matrix region.46 This is advantageous from the
viewpoint of structural features. The hemicelluloses have enough mobility to
arrange for proper conformation.36 It is important for regulating their role in the
cell walls and also for the physical properties such as cell wall extension.47
2.2 Cellulose nanofibers from wood
Cellulose nanofibers (CNFs) are present in the structural network of wood
cell walls in the form of microfibrils. CNFs, which are classified within a family
of nanocellulose, are of interest as a strong reinforcement in nanocomposites. The
key characteristics of high aspect ratio,3 outstanding mechanical properties,48, 49
good network formation,4 high surface area50 and versatile chemical
functionality51 encourage the utilization of CNFs for novel applications. CNFs are
commonly disintegrated from wood fibers. The production of CNFs primarily
consists of a pretreatment step followed by mechanical disintegration to liberate
the microfibrils. Typical morphology of wood fibers and CNFs are shown in
Figure 3.
8 BACKGROUND
Figure 3. Light microscopy image of spruce wood fibers and atomic force microscopy
(AFM) image of the obtained CNFs via enzymatic approach.
2.2.1 Pretreatment approaches
Various pretreatment approaches have been developed for CNF production.
The general aim of the pretreatment is to improve nanofibrillation level and
reduce energy consumption during the subsequent mechanical disintegration.
Turbak et al.52 originally adopted a direct mechanical disintegration approach
without any pretreatment procedure. CNFs were liberated from a dilute
suspension of wood fibers. Non-uniform nanofibers having width in the range of
25 to 100 nm were obtained. To achieve sufficient degree of nanofibrillation, this
approach required a high number of passes through the mechanical
homogenizer. Consequently, the resulting nanofibers were shortened and
damaged, and high energy was consumed during mechanical disintegration.
Enzymatic pretreatment in combination with mechanical disintegration was
found to facilitate nanofibrillation by Henriksson et al.3 and Pääkkö et al.4 The
obtained CNFs have high average molar mass (DP ~ 1,000) and high aspect ratio
(typical diameter between 5-20 nm, and lengths of several micrometers). The
nanofibers are long, flexible and entangled, leading to formation of strong
networks and gels.4
Chemical pretreatment in combination with mechanical disintegration is
another route towards uniform CNFs. TEMPO-mediated oxidation was carried
out on the cellulose fiber surface, where primary hydroxyl groups at C-6 position
were converted to aldehyde and carboxylic groups.5, 53 After the chemical
BACKGROUND 9
pretreatment, mild mechanical treatment was adequate for nanofibrillation.
TEMPO-oxidized CNFs (TOCNs), bearing negative charge due to the introduced
carboxylic group, are individualized and have a uniform diameter of
approximately 3-4 nm. Depending on the oxidization level, nanofibers with low
molar mass (DP of 250-300) might be obtained. The length of TOCNs is about 0.5
µm.54 For TOCNs with high carboxylate content (~1.5 mmol/g), their suspension
is highly transparent.15 Similar concepts of surface modification were used to
facilitate nanofibrillation of wood fibers. For instance, positively charged CNFs
were prepared by surface quaternization.7 Another method of surface
modification using carboxymethylation was explored by Wågberg et al.6
Pulp fibers obtained after delignification with a considerable amount of native
hemicelluloses preserved are termed “holocellulose fibers”. To obtain nanofibers,
the holocellulose fibers could be prepared from chlorite delignification and
subjected directly to mechanical grinding.55 The high hemicellulose content is
positive for the ease of nanofibrillation. The holocellulose nanofibers contain up
to 20 wt% of hemicelluloses. The typical diameters of holocellulose nanofibers are
approximately 10-20 nm, which is related to the original size of fibril aggregates.
However, the wood cell walls contain individualized nanofibers (microfibrils)
with a diameter of approximately 3 nm.56 Therefore, it may be possible to convert
the holocellulose fibers into thinner nanofibers, as is the intention of the present
study.
2.2.2 Reinforcement materials
CNFs have received significant interest as high-performance reinforcement
materials. Their remarkable mechanical properties provide potential for load-
bearing applications. CNFs are also of obvious interest in the context of
sustainable development. The mechanical properties of CNFs are substantially
improved compared with wood fibers, having 10 microns scale diameter. The
strong effect of reinforcement for composites based on CNFs was first revealed
by Nakagaito and Yano.57 The strength of the CNF-based composite was
comparable to some metal alloys.
From the viewpoint of engineering design, a low-density material with
favorable mechanical properties is attractive. CNFs are likely to have advantages
10 BACKGROUND
over man-made fibers due to their intrinsic properties of low density, high
specific modulus and high specific strength, as presented in Table 1. It should be
noted that CNFs are assembled from semi-crystalline extended chains, and thus
the reinforcing effect of CNFs is expected to be somewhat lower than that of
cellulose crystals. The reinforcing capability of CNFs is also controlled by several
parameters such as their aspect ratio, degree of dispersion, orientation
distribution, the extent of mechanical damage and surface chemistry attributes.
The resulting nanostructural characteristics and mechanical properties of the
CNFs are of particular importance and efficient preparation routes for CNF
production are obviously also desirable.
Table 1 Mechanical properties of engineering fibers and cellulose-based
reinforcement materials.
Density
(g/cm3)
Young’s
Modulus
(GPa)
Tensile
Strength
(GPa)
Specific
Modulus
(GPa/gcm-3)
Specific
Strength
(GPa/gcm-3)
Glass fiber1 2.5 70 2.4 28 0.95
Carbon fiber58 1.8 150-500 1.5-5.5 80-280 0.83-3.1
Wood fiber58 1.5 14-27 0.3-1.4 9-18 0.20-0.93
Cellulose
nanofiber59-61 1.5 > 55 1.6-3.0 > 35 1.0-2.0
Crystalline
cellulose58, 62 1.6 130-150 7.5-7.7 80-95 4.7-4.8
2.3 Biocomposite films based on cellulose nanofibers
Long nanofibers and their entangled network allow the formation of flexible
CNF-based films. A rapid and convenient preparation for large CNF-based films
with high mechanical properties was reported by Sehaqui et al.63 The procedure
mainly involves water-based filtration and constrained drying. Upon drying,
strong interfibril interaction is developed.64 The obtained CNF film, referred to as
CNF nanopaper, shows high modulus (13.4 GPa) and strength (232 MPa)63 due to
good random-in-plane orientation of the nanofibers.48 The orientation can be
further improved by partial alignment of the nanofibers in the wet state by cold-
BACKGROUND 11
drawing.49 CNF-based films with improved modulus (> 30 GPa) and tensile
strength (575 MPa) were reported61 when the surface of nanofibers was grafted
with soft polymer to facilitate the alignment. The nanofibers have ability to form
a dense network, which is promising for barrier properties in packaging
application.14, 65 Owing to the unique characteristics of individualized CNFs in a
densely packed network, CNF-based films with optical transparency could also
be prepared,66 which is useful for electronic applications such as display screens,
etc.
In practice, CNFs, especially those with low surface charge, have a tendency
for self-aggregation due to strongly interacting surfaces.67 The state of CNFs
dispersion plays an important role in controlling the final structure of the
nanofibrillar network (nanopaper). Aqueous suspensions of CNFs can be
stabilized by electrostatic repulsion and steric hindrance.68 Good dispersion of
CNFs during processing is essential to achieve desirable mechanical
reinforcement.69
Polysaccharide biopolymers such as hemicelluloses and starch have been
widely investigated as a candidate component for biomass-based materials. To
transform these polysaccharides into green materials for wide-spread industrial
applications, several limitations on their inherent properties such as
hygromechanical properties and poor film formation need to be solved. Their
mechanical performance is very sensitive to moisture content, especially when
plasticizers are introduced. Substantial reduction in tensile modulus and strength
under high humidity conditions is a common problem.29, 70 The attractive
characteristics of CNFs are particularly interesting and may serve to overcome
these challenges. Strong molecular interactions between CNFs and other
polysaccharides at the interface as well as the semi-crystalline nature of CNFs are
advantages. For example, the addition of CNFs leads to strong reinforcement and
reduced water sensitivity for starch films.8, 9, 12, 71 There have been a considerable
number of studies devoted to hemicelluloses such as galactoglucomannan and
xylan,10, 11, 72, 73 and similar positive effects from CNF addition have been observed.
The neat films obtained from this class of polysaccharides are generally brittle
due to their insufficient chain length and high glass transition temperature.74 The
12 BACKGROUND
film forming capability of hemicelluloses and starch can be improved by the
flexible and entangled network of CNFs.13, 75
CNFs and other polysaccharides are usually compatible. In the context of
processing, they can be mixed in aqueous suspension, and easily shaped into the
final materials. In most studies, the effect of CNFs on mechanical properties of
the biocomposite films is studied at low CNFs content. By using filtration, well-
dispersed biocomposites with high CNFs content can successfully be prepared76
and strong reinforcement effects can be achieved.77
It has also been reported that various sources of CNFs such as tunicates and
bacteria can be used, although the reinforcement effect varies due to structural
organization and mixing effects.78, 79 The network formation is governed by the
geometry and stiffness of the nanofibers. For wood CNFs and polysaccharide
matrices, strong reinforcement is likely due to favorable interface bonding and
nanofiber entanglements.78 However, the reinforcing and film-forming potential
from CNFs could also vary due to the difference in nanostructural morphology,9
which is related to the extent of nanofibrillation, preparation and degree of
dispersion.
Biocomposite films based on CNFs and CNF/Amylopectin (AP) are shown in
Figure 4. CNFs are obtained as a gel-like suspension (~2 wt%) after enzymatic
pretreatment and disintegration. A translucent film of CNFs and CNF/AP can be
prepared by filtration of a hydrocolloidal suspension followed by drying. The
nanofibrillar network forms a lamellar structure at microscale. The mechanical
properties of CNFs are strongly influenced by the surrounding humidity due to
adsorbed moisture. The mechanical properties (modulus, strength, and ductility)
of CNF/AP at elevated humidity are improved compared with neat AP, but it is
still far below neat CNF. In the present study, we will investigate if the
mechanical properties of CNF/AP films can be improved by using a Core-shell
nanofiber approach.
BACKGROUND 13
Figure 4. Typical characteristics of biocomposite film based on enzyme-pretreated
CNFs, and the hygromechanical properties of CNF/Amylopectin biocomposite film.
2.4 Core-shell cellulose nanofibers and their biocomposite films
Fiber/matrix interfacial interaction is important for the reinforcement of the
composite materials.80 The localized region where the molecular conformation
and mobility of the matrix is influenced by the rigid surface of the fiber is defined
as an interphase.80 The molecular nature of the nanoscale interphase can affect the
final properties of the nanocomposites.81 Thus, it becomes important to have
better control of materials at the molecular and nanostructural scale.
The design of the cell walls has been studied at molecular level.36, 46 The simple
model of their organization for cellulose microfibrils (reinforcing element) and
hemicelluloses (polysaccharide matrix) is proposed by Fengel22 (Figure 5(a)). In
the primary cell walls, the surface of cellulose microfibrils is coated by XG.46
More specifically, the basic structural unit is assembled from the microfibrils
where the surface is coated by a monolayer of XG.82 An interesting consequence
is related to the role of interface hemicelluloses, which serve as a glue between
the microfibrils.83 The strong binding of hemicelluloses to cellulose is beneficial to
the molecular interaction and bond recovery during mechanical deformation of
the cell walls.84
14 BACKGROUND
Assemblies in the form of biomimetic composites, where CNFs and ductile
polymers serve as load-bearing and energy-dissipating components, have been
investigated.76, 85-87 The design principle for these materials is based on a
composite with a high content of CNFs, and the CNF surface is modified through
polymer adsorption. A study of bacterial cellulose synthesized in the presence of
hydroxyethyl cellulose (HEC), which is a hydrated amorphous polysaccharide,
were conducted.85 Biocomposites of cellulose fibrils coated by a monolayer of
HEC were obtained, where fibril dispersion and mechanical properties were
improved remarkably. A biomimetic composite with high toughness was
obtained in an analogous study on wood-based CNFs and HEC.76 The coating
layer of HEC on CNFs showed advantages for nanofiber alignment during tensile
stretching, resulting in biocomposites with high modulus and strength.88 The
effect of nanoscale interactions between the components on the mechanical
properties of biomimetic composites has been investigated.86 Adsorption of
polymers such as cellulose derivatives86 and xyloglucan89 can reduce friction on
cellulose surfaces, resulting in favorable nanofiber slippage and toughness
enhancement.90 The biocomposite films with small amount of these adsorbed
polysaccharides (approximately 2 wt%) become more ductile, indicated by
increased strain-to-failure and work of fracture.86, 87 This positive effect could also
be observed on biomimetic gels under wet conditions.87
Figure 5. Model of (a) ultrastructure organization of cellulose microfibrils within the
wood cell walls (Fengel 1971) and (b) ideal nanostructure of Core-shell cellulose
nanofibers.
BACKGROUND 15
To mimic the design strategies used by nature, Core-shell nanocomposites
with high CNF content and adsorbed polysaccharide matrix have been
developed.77, 91 The ideal model of Core-shell cellulose nanofibers is proposed in
Figure 5(b). The polysaccharide matrix is coated on the surface of CNFs, and
presented in a form of “shell” layer on “core” nanofiber surface. The unbound
matrix is removed from the mixed CNF/polysaccharide suspension, and Core-
shell nanocomposites are formed by filtration. Biomimetic approaches of nacre-
like composite prepared by polymer adsorption and filtration have been
developed for nanoclay-based materials.92, 93 In Core-shell nanocomposites, one
goal is to enhance and control the distribution of polymer matrix at
nanostructural level. The presence of the adsorbed matrix may inhibit self-
aggregation of CNFs in colloidal suspension. Steric stabilization of the
suspension is expected through polymer chain extension.94 The network with less
aggregated nanofibers is essential for ductility and strength of the final
nanocomposites. Moreover, the adsorbed matrix, which is confined on the CNF
surface, behaves differently from the bulk matrix. Strong interfacial interactions
between the components maximize the reinforcing effect in Core-shell
nanocomposites,77 especially under hydrated conditions.91 The hygromechanical
properties of the Core-shell nanocomposites are improved as a result of better
control of matrix distribution at the nanoscale.
2.5 Holocellulose nanofibers and their biocomposite films
Holocellulose is typically defined as a total polysaccharide fraction of
cellulose and hemicelluloses from wood cell walls. In fact, the native structural
organization of holocellulose, where the cellulose microfibrils are coated by
hemicellulose matrix,46 is the ideal structure for Core-shell nanofibers.
Holocellulose nanofibers with diameters of 10-20 nm have been isolated from
spruce wood powder via chlorite treatment and mechanical grinding by Iwamoto
et al.55 The high hemicellulose content is beneficial for nanofibrillation55 and the
coalescence of microfibrils into aggregates is inhibited.95 The physical properties
of holocellulose films were compared between once-dried and never-dried
nanofibers, and no significant difference was observed. The holocellulose films
showed tensile modulus of approximately 16 GPa and strength of 280 MPa,55
16 BACKGROUND
which is higher than that of enzyme-pretreated CNFs.63 The light transmittance
was slightly lower than for transparent acrylic resins, indicating good
transparency of the holocellulose film.55 Thus, holocellulose nanofibers are of
great interest for the preparation of biocomposites with good properties. Another
advantage of the holocellulose is the high yield of polysaccharide components
from starting material.
2.6 Application of nanofibers in porous materials
In the context of engineering design, porous materials can be classified into
two main categories; foams and honeycombs, depending on the architecture of
their solid cells.96 Foam is a cellular solid where the interconnected network of
the solid cells is built-up in three dimensions, while a honeycomb is the two-
dimensional arrays. The porous materials can possess low density; i. e., low
relative density – (ρ*/ρs), where ρ* is the density of the porous material and ρs is
the density of the bulk material.96 They have attracted attention for a wide range
of applications such as packaging,97, 98 insulating materials99 and biomedical
scaffolds.100, 101
There have been many studies on the potential use of CNFs as a reinforcing
component in porous materials, particularly for foams.98, 102-104 Since CNFs can
form stable suspensions in water, freeze-drying, also termed lyophilization,105 can
be utilized. The freeze-drying is a process, where a CNF suspension is subjected
to freezing prior to sublimation of the solidified phase under reduced pressure.101
The architecture of the porous materials such as pore morphology, pore size and
organization can be controlled during the freezing step, and depends on
nucleation and growth of ice-crystal templates.101, 103 Moreover, the cell wall
thickness and density can be altered by varying solid content and composition of
the cell wall components.103
The unique porous structure results in favorable mechanical response for the
foams. The CNF-based foams with a wide range of densities (7-103 kg/m3) and
mechanical properties can be tailored by varying the solid content of the mixture
suspension before freezing (Figure 6(a)).104 Mixtures of CNFs and water-soluble
polymers such as XG and starch can be prepared to improve mechanical
BACKGROUND 17
properties of the CNF foams. Ultra-high porosity CNF-based foams demonstrate
high mechanical properties and energy absorption capability, which is promising
for substitution of synthetic polymer-based polystyrene foams.102, 104
Compression tests are particularly interesting to investigate the mechanical
behavior of porous materials. Figure 6(b) shows a typical compressive stress-
strain curve of the CNF foams. There are three regions, which are correlated to
different deformation mechanisms, including (I) linear elasticity, (II) collapse
plateau, and (III) densification.96 Linear elasticity is controlled by cell wall
bending and stretching, and Young’s modulus (E*) is calculated as the initial
slope of the stress-strain curve. Collapse plateau is related to cell collapse, where
the strain is no longer recoverable. Based on the study on structure and
deformation mechanism of CNF foams by Ali and Gibson,106 the plateau is
associated with plastic deformation of the cell wall as well as some cracking. The
modulus and plateau stress of the foams are dependent on relative density as
described by Gibson and Ashby.96 Finally, densification occurs once the opposing
cell walls come into contact and completely collapse.
Figure 6. (a) Foam materials prepared from enzyme-pretreated CNFs (Sehaqui et
al. 2010)107 and (b) Compressive stress-strain curve for CNF foam (density 50
kg/m3), showing the three regimes of mechanical deformation.96
A honeycomb structure is a porous two-dimensional material formed with
well-defined pore connectivity and open porosity along one direction. The
18 BACKGROUND
cellular structure of wood is an example of honeycombs in nature (Figure 7(a)).
An ideal model of honeycomb structure is presented in Figure 7(b). The cell
structure of honeycombs is highly aligned, and thus offers good mechanical
response when loading along the cell, which is promising for load-bearing
applications. Both Young’s modulus (E*) and compressive strength (σ*) of
honeycombs vary linearly with relative density (ρ*/ρs):96
𝐸∗
𝐸𝑠=
𝜌∗
𝜌𝑠 (1)
𝜎∗
𝜎𝑠=
𝜌∗
𝜌𝑠 (2)
where Es and σs are the modulus and the compressive strength of the bulk
material. Relative density (ρ*/ρs) is the volume fraction of solid material.
One requirement for controlled fabrication of honeycomb-like structures is
associated with the control of cell wall alignment. From a practical point of view,
freeze-drying can be applied for the preparation of honeycombs. One way to
achieve CNF-based honeycombs is to control the direction of ice-crystal growth
during freezing. The concept is similar to freeze-casting technique used for
porous ceramics,100 where the material suspension is subjected to unidirectional
freezing before sublimation. As a consequence, the porous structure with
unidirectional channels,108 referred as honeycombs, is obtained.
Figure 7. SEM images of (a) cross-section and (b) longitudinal section of wood cell
wall structure (Gibson 2012)18 together with (c) ideal honeycomb structure.
EXPERIMENTAL 19
EXPERIMENTAL 3
The materials and overview of methods involved in the present thesis are
summarized below. More details are available in the appended papers.
3.1 Material components
3.1.1 Cellulose nanofibers (CNFs) prepared by enzymatic pretreatment
(Paper I-V)
To prepare enzyme-pretreated CNFs, never-dried spruce sulfite pulp (kindly
provided by Nordic Paper Seffle AB, Sweden) was treated with an endoglucanase
enzyme (Novozymes, Denmark), as previously reported by Henriksson et al.3, 48
For nanofibrillation, the enzyme-treated pulp was subjected to mechanical
disintegration using a microfluidizer (Microfluidics Ind., USA) for 3 passes
through big chambers (diameter of 400 and 200 µm), and 5 passes through small
chambers (diameter of 200 and 100 µm). A translucent suspension of enzyme-
pretreated CNFs (solid content ~ 2 wt%) was successfully prepared.
3.1.2 Holocellulose nanofibers prepared by peracetic acid pretreatment
(Paper IV-V)
To prepare holocellulose nanofibers, spruce wood chips were cut into fine
“match sticks”, and subjected to delignification using peracetic acid solution (32
wt% in acetic acid, Sigma-Aldrich) followed by mechanical disintegration. To
avoid decomposition of peracetic acid, metal ions in the wood sticks were
removed by treating with a DTPA/sodium sulfite solution before the
delignification step. The delignified wood fibers obtained from peracetic acid
pretreatment were slushed using a pulp slusher (10000 revolutions). For
nanofibrillation, the sample was subjected to mechanical disintegration using the
microfluidizer. It was sufficient to liberate the holocellulose nanofibers after
passing the sample for 1 pass through the big chambers and 3 passes through the
small chambers. A transparent suspension of holocellulose nanofibers (solid
content ~ 1 wt%) was successfully prepared. The holocellulose nanofibers are also
20 EXPERIMENTAL
referred to as Core-shell nanofibers (Paper IV-V), having the “shell” layer of native
hemicelluloses coating on the “core” nanofiber surface.
3.1.3 Polysaccharide matrices (Paper I-III)
Maize amylopectin (AP) was supplied from Sigma-Aldrich, Germany. Besides
amylopectin component, the purchased AP contains 7.4 wt% of amylose. An
aqueous suspension of AP (solid content ~ 1 wt%) was prepared with deionized
(DI) water. Several steps of heating (at 65 and 135 °C) were applied to achieve
solubilized AP. The obtained AP suspension was optically transparent. The
degree of polymerization (DP) for this AP is 1.03x106.
Tamarind seed xyloglucan powders (Glyloid 3S), which are a commercially
purified product, were provided by Dainippon, Japan. An aqueous suspension of
XG (solid content ~ 0.5 wt%) was prepared by dispersing the XG powders in DI
water, and heated at 50 °C overnight. To remove water-insoluble impurities, the
XG suspension was subsequently subjected to centrifugation at 4500 rpm for 30
min.29 The obtained XG suspension was freeze-dried and collected in a form of
dried XG for long time storage. The dried XG can be re-dispersed in DI water for
further experiments. The weight average molar mass (Mw) of XG is 493,000 Da,
and the number average molar mass (Mn) is 394,800 Da.
Spruce galactoglucomannan (GGM) can be isolated from process water of
thermomechanical pulping, as previously reported by Willför et al.32 and Xu et
al.109 The GGM powders used in the present work were provided by Willför
group. The spruce GGM from process water was purified by ethanol
precipitation. The obtained GGM has the weight average molar mass (Mw) of 39
Da. An aqueous suspension of spruce GGM (solid content ~ 0.1 wt%) was
prepared by dispersing the GGM powders in DI water, and heated at 90 °C for 4
h.
EXPERIMENTAL 21
3.2 Preparation of biocomposite materials
3.2.1 Cellulose and holocellulose nanopaper (Paper I-IV)
Cellulose and holocellulose nanopapers were prepared by a papermaking-like
process.63 A diluted suspension of nanofibers (solid content ~ 0.2 wt%) was mixed
using an Ultra Turrax mixer (IKA, T25 Digital) at 8000 rpm for 3 min.
Subsequently, the nanofibers suspension was degassed for 10 min. The network
of nanofibers was formed as a wet cake on a filter membrane (0.65 µm, DVPP
Millipore) during vacuum filtration, as described by Henriksson et al.48 To obtain
the nanopaper, the wet cake was subjected to vacuum drying using a Rapid
Köthen sheet dryer at 93 °C for 10 min.63
3.2.2 Neat polysaccharides (AP and XG) films (Paper I-II)
The preparation of Neat AP and Neat XG films was based on a general method
of film casting. Prior to film casting, aqueous suspensions of AP (solid content ~ 1
wt%) and XG (solid content ~ 0.5 wt%) were degassed. The samples were
carefully transferred to petri dishes coated with Teflon FEP film (Bytac®, USA)
and dried in an oven at 40 °C for a few days. A typical thickness of the resulting
films after drying was approximately 30-40 µm.
3.2.3 Adsorption experiment and preparation of Core-shell nanofibers
(Paper I-III)
For adsorption study, a series of CNF/AP, CNF/XG, and CNF/GGM mixtures
with varying matrix content were prepared from aqueous suspensions of CNFs
and the matrices. To allow adsorption, the mixed suspensions of CNF/AP and
CNF/XG (solid content ~ 0.2 wt% based on nanofibers) were thoroughly mixed
by magnetic stirring for 24 h at room temperature, while the mixing was
performed at 90 °C for 4 h in the case of CNF/GGM. After the adsorption period,
the excess/unbound matrix was removed from the mixture by centrifugation at a
speed of 4500 rpm for 10 min. The sediment was collected and washed several
times. The washing was performed by re-dispersing the sediment in DI water,
and repeating the centrifugation step in order to remove the unbound matrix.
22 EXPERIMENTAL
The obtained material contains CNFs coated with adsorbed polysaccharide
matrix, referred to as Core-shell nanofibers in Paper I-III.
3.2.4 Nanocomposite films (Paper I-III)
In the present thesis, nanocomposite films of CNFs with polysaccharide
matrices including AP, XG, and GGM were prepared by two different
approaches (Figure 8).
For Core-shell nanocomposites, diluted suspensions (solid content ~ 0.2 wt%)
of Core-shell nanofibers (CNF/AP, CNF/XG, CNF/GGM) with varying matrix
content were thoroughly dispersed using the Ultra Turrax, and then exposed to
degassing. Thereafter, the films were formed by vacuum filtration and drying,
according to the method for CNF nanopaper.
For “Mixed” nanocomposites, a series of aqueous mixtures for CNF/AP,
CNF/XG, and CNF/GGM suspensions (solid content ~ 0.2 wt%) with varying
matrix content were prepared. The thorough mixing was performed similar to
the procedure for adsorption experiment. The main difference between Core-shell
and “Mixed” nanocomposites is related to the unbound matrix. For “Mixed”
nanocomposites, the mixture suspensions of CNF/AP, CNF/XG, and CNF/GGM
were directly subjected to filtration without the removal of unbound matrix.
Other experimental details related to film formation were performed similar to
that for Core-shell nanocomposites.
EXPERIMENTAL 23
Figure 8. Schematic illustration of the difference between Core-shell and “Mixed”
approaches, used for the preparation of nanocomposite films. The figure is taken from
Paper I.
3.2.5 Porous materials based on nanofibers (Paper V)
Porous materials based on cellulose and holocellulose nanofibers were
prepared by freeze-drying, as previously reported by Sehaqui et al.104 In the
present thesis, both honeycombs and foams were developed by directional and
non-directional freezing, respectively. To obtain the porous materials with
varying densities, aqueous suspensions of cellulose and holocellulose nanofibers
with different concentrations were prepared. For the porous materials with low
density, the nanofiber suspensions were diluted with DI water. For the porous
materials with high density, the nanofiber suspensions were concentrated using
high-speed centrifugation (Avanti® J-E centrifuge, Beckman Coulter with a JLA-
10.500 rotor). By varying speed and time for the centrifugation, the nanofiber
suspensions with various concentrations were obtained. After the centrifugation,
the bottom phase, containing concentrated nanofibers, was collected. The
obtained nanofiber suspensions with different concentrations were carefully
transferred to aluminum molds (diameter of 50 mm) for freeze-drying.
24 EXPERIMENTAL
Prior to the freeze-drying, the samples were kept at 4 °C overnight in order to
avoid microscopic cracking during freezing. During freezing, the aluminum
molds were directly in contact with the liquid nitrogen (-196 °C), and thus the
samples were frozen. The preparation of foams and honeycombs are different
(Figure 9). For foams, the bottom and the side of the molds were directly in
contact with the liquid nitrogen, which allow non-directional freezing. In the case
of honeycombs, the samples were prepared by directional freezing. The side of
aluminum molds was covered with an insulating material, preventing ice-crystal
growth due to thermal gradient during the freezing. As a consequence, the ice-
crystal growth was allowed in only one direction, particularly from bottom to top
of the aluminum mold. Subsequently, the frozen samples were freeze-dried (3-4
days) in a benchtop freeze-dryer system (6 liter FreeZone, Labconco Corporation,
USA), as previously described by Sehaqui et al.104 Porous materials with different
densities varying between 6-55 kg/m3 were successfully prepared.
Figure 9. Schematic illustration of freezing step for the preparation of foams and
honeycombs. The figure is modified from Paper V.
EXPERIMENTAL 25
3.3 Characterization methods
3.3.1 Determination of matrix contents by sugar analysis (Paper I-III)
The content of polysaccharide matrix in the nanocomposites was determined
by sugar analysis, according to a standard method of SCAN-CM 71:09. Five main
sugar references of glucose, mannose, xylose, galactose, and arabinose were
selected for the analysis. To obtain dissolved monosaccharides for the analysis,
acid hydrolysis was performed on the freeze-dried sample of neat CNF and the
nanocomposites. Thereafter, the sugar composition was determined using high-
performance anion exchange chromatography system (Dionex, USA) equipped
with pulsed amperometric detection (HPAEC-PAD). Data processing was
performed with Chromeleon software.
3.3.2 Colloidal behavior by dynamic light scattering (DLS) (Paper I-III)
To characterize the aggregation of nanofibers in aqueous suspension, DLS
measurement was performed using a Zetasizer Nano ZEN3600 (Malvern
Instruments Ltd., UK). The light source was operated at 633 nm. The nanofiber
suspensions were adjusted to a final concentration of 100 mg/L in DI water. To
avoid interference from dust particles and large aggregates, the samples were
filtered through 5 µm membrane (Acrodisc) prior to the measurement. All
measurements were performed at 25 °C.
3.3.3 Total charge determination (Paper IV)
The total charge content of pretreated pulp fibers (enzymatic and peracetic
acid pretreatments) was determined by conductometric titration, according to a
standard method SCAN-CM 65:02. A diluted suspension of pretreated pulp
fibers was prepared in DI water. Prior to the titration, some NaCl was added in
the pulp suspension, and the pH was adjusted to 3 with 0.01 M HCl. During the
titration, a solution of 0.1 M NaOH was added and the conductivity of the pulp
suspension was followed. The total charge content was calculated based on the
sum of strong (sulphonic) and weak (carboxylic) acids on the pulp fibers.
26 EXPERIMENTAL
3.3.4 Nanostructure of nanofibers and their networks (Paper I-V)
Atomic force microscopy (AFM)
AFM was performed to analyze the nanostructure of nanofibers for Neat CNF
and Core-shell nanofibers. For all samples, a diluted suspension was spin-coated
on a mica disc. The samples were scanned using Nanoscope IIIa AFM (Veeco
Ltd., USA) in a tapping mode. The employed cantilevers had a tip radius of 8 nm
and spring constant of 5 N/m. Height images were recorded and used for the
analysis of nanofibers diameter.
Field-emission scanning electron microscopy (FE-SEM)
The morphology of nanofibrillar network and fracture surface of the
nanocomposite films in addition to the cellular structure of CNF-based porous
materials were observed using FE-SEM (Hitachi, S-4800, Japan). The sample was
placed on the surface of carbon tape mounted on a metal stub. Prior to
observation, the sample was coated with graphite and gold-palladium using
Cressington 208HR sputter coater for 5 and 30 s, respectively.
Transmission electron microscopy (TEM)
Structural information on the arrangement of nanofibers compared between
Core-shell and “Mixed” CNF/AP nanocomposites was provided by TEM
observation (FEI Tecnai Spirit, USA). The nanocomposites were embedded in
epoxy resin, and then cross-sectioned at room temperature using an ultra-
microtome (Leica EM UC7, Germany) equipped with a diamond knife (Diatome,
Switzerland). The obtained samples were stained by exposure for 1 h to RuO4
vapors.
3.3.5 Density and porosity (Paper I-V)
Density of the materials was determined based on their air-dried mass and
volume. The volume of each specimen was calculated by measuring thickness
and surface area of 3-5 set of data. At least five specimens were used for the
calculation of density for each sample.
EXPERIMENTAL 27
Porosity of the material was calculated based on the density according to the
following equation:
𝑃𝑜𝑟𝑜𝑠𝑖𝑡𝑦 = 1 − 𝜌𝑚𝑎𝑡𝑒𝑟𝑖𝑎𝑙
𝜌𝑐𝑒𝑙𝑙𝑢𝑙𝑜𝑠𝑒 (3)
where the density of cellulose 𝜌𝑐𝑒𝑙𝑙𝑢𝑙𝑜𝑠𝑒 was considered to be 1500 kg/m3.48
3.3.6 Tensile mechanical properties (Paper I-IV)
Tensile measurements of Neat CNF and nanocomposite films were performed
using a universal testing machine Instron 5944 (Instron, USA) equipped with a
500 N load cell. The samples were conditioned at 23 °C under controlled relative
humidity of 50 or 85 RH% for at least one week prior to the tensile testing. The
samples were cut into thin rectangular strips (25x5 mm2). The gauge length was
set to 20 mm. The tensile behavior was recorded at a strain rate of 10 %/min. For
each sample, an average value of at least 5 specimens was reported for tensile
properties. Digital Speckle Photography (DSP) was performed to obtain an
accurate value of strain. A correlation of strains determined from the DSP and the
grip displacement during tensile testing was established. The reported modulus
was calculated based on the corrected value obtained from the DSP
measurement.
3.3.7 Mechanical properties in compression (Paper V)
The mechanical properties of porous materials (foams and honeycombs) were
measured by compression testing. Prior to the measurement, the sample was
conditioned at 23°C and 50 RH% for at least one week, and cut into 1 cm cubes.
The compression testing was performed using the Instron 5944 equipped with a
500 N load cell. The measurement was carried out on the samples both parallel
and perpendicular to the direction of ice-crystal growth. All specimens were
tested at a strain rate of 10%/min. For each sample, an average value from 3-5
specimens was reported for compressive properties.
3.3.8 Dynamic vapor sorption (DVS) (Paper I-III)
Moisture sorption isotherm of the materials was determined with a DVS
instrument (Surface measurement Systems Ltd, UK). Prior to DVS measurement,
28 EXPERIMENTAL
moisture was removed from the dried samples by additional drying in a vacuum
oven at 40 °C for 24 h. The DVS experiments were conducted in an adsorption
mode, starting from 0 to 90 RH%. Each RH condition was maintained for at least
400 min to ensure steady state. The moisture uptake at a particular RH% was
calculated according to the equation below:
𝑀𝑜𝑖𝑠𝑡𝑢𝑟𝑒 𝑢𝑝𝑡𝑎𝑘𝑒 = 100𝑥𝑊𝑚𝑜𝑖𝑠𝑡−𝑊𝑑𝑟𝑦
𝑊𝑑𝑟𝑦 (4)
where 𝑊𝑚𝑜𝑖𝑠𝑡 is the sample weight equilibrated at particular RH%, and 𝑊𝑑𝑟𝑦 is the
weight of the dried sample.
3.3.9 Molecular moisture dynamic by 2H solid-state NMR (Paper II)
2H solid-state NMR was used as a tool to investigate moisture-related
properties of Neat CNF, Neat XG and their nanocomposites at molecular level.
NMR relaxometry was performed to analyze molecular dynamics and molecular
distribution of moisture in the materials. Prior to the NMR experiment, all film
samples were conditioned at 92 RH% (saturated KNO3 solution in 2H2O) for 3
weeks. 2H2O was exploited in order to separate signals among the polymer and
water molecules. The NMR measurement was performed using a Bruker
Advance II spectrometer operating at a frequency of 46.1 MHz and a short (2.3
µs) 90° pulse. The distribution of 2H transverse relaxation time was obtained by
Inverse Laplace Transform (ILT).
3.3.10 Molar mass analysis of nanofibers by SEC (Paper IV)
Prior to molar mass analysis, the nanofibers were solvent-exchanged three
times with dry dimethyl acetamide (DMAc) at ambient temperature.
Subsequently, LiCl/DMAc and EIC reagents were added to dissolve the
nanofibers. The dissolved samples were treated in a Retsch vibratory mill type
MM-2 for 30 min followed by filtration through 0.45 µm PTFE-filter (Advantec,
MFS). The chromatographic system consists of a 2690 Separation Module (Waters
Corp.) equipped with a guard column (Mixed-A 20 µm 7.5x50 mm, Polymer
Laboratories) followed by four column (Mixed-A 20 µm 7.5x300 mm, Polymer
Laboratories) connected in series. The mobile phase of LiCl/DMAc was used. The
separations were performed at 80 °C. Molar mass calibration of SEC was
EXPERIMENTAL 29
performed using a linear pullulan standard set with Mw ranging from 738 to
1,660,000 Da (Polymer Laboratories). Detection was performed using a 2478 dual
wavelength absorbance detector (Water Corp.) at 295 nm coupled with a 410
differential refractometer (Waters Corp.).
3.3.11 Molar mass analysis of matrices by SEC-MALLS (Paper I-II)
Size exclusion chromatography coupled with a multi-angle laser light
scattering detector (SEC-MALLS) (PSS/Agilent SECurity GPC system, USA) was
employed to determine the molar mass of AP and XG matrices. The samples
were dissolved with a solvent system of DMSO/LiBr. Molar mass calibration of
SEC was performed using a linear pullulan standard set with Mw ranging from
342 to 805,000 Da (PSS, Germany). The solvent system of DMSO/LiBr was also
used for SEC analysis. Data processing was performed with PSS WinGPC
UniChrom software.
3.3.12 Light transmittance (Paper IV)
Light transmittance of nanopaper was measured using an UV spectrometer
(UV-2550, Shimadzu, Japan). The wavelengths ranging from 200 to 900 nm were
scanned. The transmittance spectra were recorded at room temperature, with air
as a reference.
RESULTS AND DISCUSSION 31
RESULTS AND DISCUSSION 4
The first section will describe Core-shell cellulose nanofibers with
polysaccharide matrices. The nanostructure, hygromechanical properties and
moisture sorption behavior of the obtained nanocomposite films will be
discussed. The second section will describe another type of Core-shell nanofibers,
referred to as native holocellulose nanofibers. Peracetic acid pretreatment will be
emphasized as an alternative method for preparing the holocellulose nanofibers.
The key characteristics of the holocellulose nanofibers as well as the materials
including film and honeycombs will be compared with those obtained from
enzyme-pretreated CNFs.
4.1 Core-shell nanofibers and their biocomposite films (Paper I-III)
4.1.1 Preparation of Core-shell nanofibers with polysaccharide matrices
To address the importance of matrix distribution at nanoscale, CNF-based
nanocomposites were prepared by two different methods, see Figure 8 and 10. In
the first approach, Core-shell cellulose nanofibers were prepared by coating the
polysaccharide matrix on the CNF surface. The bound matrix serves as a “shell”
layer for the “core” nanofiber. The excess matrix, termed “unbound” matrix, was
carefully removed from a colloidal suspension of Core-shell nanofibers prior to
film formation, as described in the experimental section. For “Mixed” method, it
is a traditional procedure where the colloidal suspension of CNFs and matrix is
subjected to simple mixing. The mixing step was allowed to continue until an
equilibrium state was reached. Thereafter, all nanocomposite films including
Core-shell and “Mixed” were formed by a papermaking-like process, where a wet
gel of nanofibers was obtained by filtration and then subjected to rapid water
evaporation.63 Strong secondary interactions including hydrogen bonding were
developed between the nanofibers after drying.
Different polysaccharide matrices including amylopectin (AP), xyloglucan
(XG), and galactoglucomannan (GGM) were exploited as the polymer matrix for
the CNF-based nanocomposites. The interactions between CNFs and these
32 RESULTS AND DISCUSSION
structurally different polysaccharides are responsible for their unique binding
capability.110 For Core-shell nanocomposites, the molecules of polysaccharide
matrix are tightly bound to the CNF surface. In contrast, “Mixed”
nanocomposites also have a fraction of unbound matrix. As shown in Figure 10,
the amount of remaining polymer matrix in Core-shell and “Mixed”
nanocomposites is plotted versus amount of polymer added to the suspension,
and the difference between two curves indicates the amount of unbound matrix.
A large fraction of unbound matrix is present in the case of AP matrix. The
difference in molecular structure among the matrix polysaccharides affects the
level of their interaction with cellulose. This also influences the fraction of bound
and unbound matrix in the final nanocomposites. Favorable association is
expected in both XG and GGM, which is related to the molecular rigidity of their
backbone.111-113 Also, adsorption to cellulose is part of their function in the plant
cell walls. XG can adsorb strongly to cellulose surfaces even in the hydrated
state.89 It should also be pointed out that molar mass can influence the observed
differences in the amount of remaining matrix. Higher molar mass may correlate
with a higher amount of adsorbed polymer.
RESULTS AND DISCUSSION 33
Figure 10. Composition of remaining matrix in the nanocomposites for (a) CNF/AP,
(b) CNF/XG, and (c) CNF/GGM. The data are for both “Mixed” and Core-shell
nanocomposites. The matrix contents in the obtained nanocomposite films are
presented next to the corresponding data points. The figure is a combination of data
from Paper I-III.
4.1.2 Nanostructure of the nanofibers and their networks
Nanofibers in aqueous suspension
It is essential to utilize the unique properties of individualized nanofibers in
nanocomposites. The preparation of CNF-based nanocomposites is typically
carried out in aqueous suspension. Therefore, the dispersion state of the colloidal
nanofibers is important to the nanostructure formation in the final
nanocomposites. This is eventually related to the material performance such as
mechanical properties.69 CNFs have a tendency to self-aggregate, which is an
unfavorable characteristic since agglomerates can act as stress concentrations and
cause premature nanocomposites fracture at small strain. To obtain a well-
34 RESULTS AND DISCUSSION
dispersed CNF suspension, electrostatic forces or steric hindrance can help to
separate the nanoparticles.114
Dynamic light scattering (DLS) data presents estimated hydrodynamic size of
nanofibers in aqueous suspension. As shown in Figure 11, distributions of the
size for Neat CNF and their Core-shell nanofibers with different polysaccharide
matrices are presented. Neat CNF exhibits a broad size distribution, with some
large size aggregates. In most cases, the size of large aggregates is reduced when
a polysaccharide matrix is adsorbed to the CNF surface. It is believed that the
presence of side chains or branches, for instance in AP, contributes to steric
stabilization68 of the CNFs in suspension.
It is important to note that DLS data is calculated based on diffusion
coefficients of spherical particle,115 which is not expected to correlate with the
nanofiber width. Hence, additional experiments were performed to observe
morphology and size distribution of the nanofibers. In the following section,
determination of the nanofiber width by atomic force microscopy (AFM) will be
presented.
Figure 11. Distribution of estimated hydrodynamic size of nanofibers in aqueous
suspension compared between Neat CNF and Core-shell nanofibers of (a) CNF/AP and
CNF/XG, and (b) CNF/GGM. The distributions are based on DLS data. The figure is a
combination of data from Paper I-III.
RESULTS AND DISCUSSION 35
Size distribution of nanofibers
The morphology and size distribution of nanofibers were observed from AFM
height images in order to avoid the tip broadening artifact. AFM section analysis
was pursued to determine the height of nanofibers, which in turn correlate to
their width. The histograms in Figure 12 show a comparison of height
distributions between Neat CNF and the Core-shell nanofibers with three different
polysaccharide matrices. The morphology of their nanofibers is also presented in
the inset. It is evident from the histograms that the diameter of nanofibers was in
the range of 2 to 20 nm. There are some differences between the Neat CNF used
for Paper I-II and III, perhaps due to the origin of the starting pulp fibers.
For all samples, the majority of nanofibers (80-95%) were in the range below
10 nm. The nanofiber aggregates with diameter above 10 nm were excluded in
part of the analysis. A summary of number average and weight average
diameters of all samples in addition to the percentage of nanofibers in the ranges
below 10 nm is listed in Table 2. Considering the values of number and weight
average diameters, Core-shell nanofibers tend to be wider in diameter, as
compared with their reference (Neat CNF). This is in line with the change in
percentage of nanofibers in the fine range below 10 nm. For Core-shell nanofibers,
the fraction of individual nanofibers with diameter of 3-5 nm decreased, while
the fraction of nanofibers with diameter of 5-8 nm and 8-10 nm increased
significantly as compared with Neat CNF. In this context, Core-shell CNF/GGM
shows the strongest effects.
The morphology of nanofibers for Core-shell CNF/GGM is depicted in Figure
12. FE-SEM images show the nanofibrillar network for Neat CNF and Core-shell
CNF/GGM hydrogels, which are dried from supercritical carbon dioxide (SC-
CO2). The nanofibers with larger diameter are clearly visible in Core-shell
CNF/GGM image. Since the GGM content is rather small (10.6 wt%), it is likely
that the GGM matrix is binding individual nanofibers together, which leads to
the increase in diameter of the observed nanofibers.
36 RESULTS AND DISCUSSION
Figure 12. Height distribution of nanofibers based on AFM measurement for (a) Neat
CNF, (b) Core-shell CNF/AP 23 wt%, (c) Core-shell CNF/XG 31 wt%, (d) Neat CNF for
GGM, and (d) Core-shell CNF/GGM 10.6 wt%. The light-colored region on the AFM
images represents the nanofibers with large diameter in the range of 10-14 nm. In the
case of CNF/GGM, nanofibers with increased diameters are also observable from the
FE-SEM images. The figure is a combination of data from Paper I-III.
RESULTS AND DISCUSSION 37
Table 2 Number average and weight average diameters (dn and dw), and the
percentage of nanofibers with focus on a small diameter fraction (< 10 nm). PDI is
polydispersity index dw/dn. The result is a combination of data from Paper I-III.
Nanofibrillar network in nanocomposites
In order to further study the difference in matrix distribution between Core-
shell and “Mixed” nanocomposites, the characteristics of their nanofibrillar
network were observed in dried hydrogels using FE-SEM. The CNF/XG
nanocomposite was selected to make a fair comparison between the two cases,
since their matrix content is comparable (XG ~ 31-32 wt%). As shown in Figure
13, there is a large difference in appearance among the nanofibrillar networks of
Neat CNF, Core-shell CNF/XG and “Mixed CNF/XG”. For Neat CNF, a highly
porous nanofibrillar network composed of uniformly sized nanofibers is
observed. For Core-shell CNF/XG, a network structure of nanofibers with larger
bundle size, as compared with Neat CNF, is noticeable. Based on molecular mass
analysis, the extended chain length of XG is as high as 800 nm in this study. It is
most likely that the assembly of nanofibers is formed via physical linkages of
adsorbed XG.116, 117 Instead, “Mixed” CNF/XG provides evidence for matrix-rich
regions in the presence of XG coated nanofibers. The XG matrix is also entrapped
within the pore space. As a consequence, an advantage of Core-shell over “Mixed”
nanocomposites is that the matrix distribution is more homogeneous and
controllable at nanoscale (Table 3).
Sample dn
(nm)
dw
(nm) PDI
% Fraction of all nanofibers
(unit: nm)
2.5-5.5 5.5-8.5 8.5-10.5 0-10
Neat CNF (Paper I-II) 5.7±2.2 6.3 1.1 39.2 35.1 10.8 86.5
Core-shell CNF/AP 23 wt% 6.1±2.0 6.6 1.1 30.0 44.7 11.3 87.3
Core-shell CNF/XG 31 wt% 6.3±2.2 6.7 1.1 27.1 38.2 14.7 80.0
Neat CNF (Paper III) 4.4±1.7 5.9 1.3 64.3 25.9 4.9 95.1
Core-shell CNF/GGM
10.6 wt% 5.4±1.7 7.5 1.2 35.6 45.3 8.7 89.6
38 RESULTS AND DISCUSSION
Figure 13. FE-SEM images of the nanofibrillar network for Neat CNF, and
nanocomposites of Core-shell CNF/XG 31 wt% and “Mixed” CNF/XG 32 wt%. Scale
bars are 300 nm. To preserve the nanofibrillar structure from aggregation during
typical drying, all samples were prepared from hydrogels (solid content ~ 15 wt%)
and subjected to SC-CO2 drying. The figure is taken from Paper II.
Table 3 Overview of preparation methods and characteristics of the nanofibrillar
network obtained from Core-shell vs “Mixed” nanocomposites
Core-shell nanocomposite “Mixed” nanocomposite
Preparation Polymer matrix is
adsorbed to CNF in
aqueous suspension.
Excess polymer is
removed prior to
filtration and drying.
Polymer and CNF are
mixed and then subjected
to filtration and drying.
Characteristic Polymer matrix is coated
on individual CNF.
Polymer matrix
distribution is more
homogeneous
CNF and polymer matrix
distributions are more
inhomogeneous.
RESULTS AND DISCUSSION 39
Fracture surface of nanocomposites
In addition to the observation of the nanofibrillar network, the differences in
fracture surface of the final films between Core-shell and “Mixed” nanocomposites
were also investigated. As shown in Figure 14, the film of CNF and all CNF/AP
nanocomposites (AP content ~ 23-25 wt%) showed lamellar structure.8, 48
However, there was a significant difference in the nanofibrillar nature among
them. For Core-shell CNF/AP (Figure 14(b)), the smooth surface of the nanofibers
indicates a homogeneous coating layer of AP. It is apparent that the network of
Core-shell CNF/AP is assembled from large bundles of the AP coated nanofibers,
instead of small nanofibril aggregates as can be observed in Neat CNF (Figure
14(a)). For the Mixed CNF/AP (Figure 14(c)), it can be seen that the lamellar
structure is inhomogeneous, apart from having smaller thickness for individual
layers as compared with Core-shell samples. This is most likely due to the
presence of unbound matrix in-between the CNF-rich laminae. Further
observation at fine scale was performed using TEM (Figure 14 (d,e)), and it
confirms the significant difference in nanostructure between Core-shell and
“Mixed” samples. Within a large thickness of individual lamina observed in
Figure 14(b), the layered structure of Core-shell nanocomposite is highly
organized at fine scale (Figure 14(d)). In contrast, larger agglomerated structures
are apparent in “Mixed” nanocomposites, see Figure 14(e).
The comparison of nanostructures between Core-shell and “Mixed” CNF/XG
nanocomposites with higher matrix content (31-32 wt%) is shown in Figure 15.
The fracture surfaces were observed after tensile testing. Features in line with
Core-shell nanocomposites, showing XG coated nanofibers, are noticeable; see
Figure 15(b). In Figure 15(c), “Mixed” CNF/XG also exhibits features of matrix-
rich regions within the fibrous network, which is in accordance with “Mixed”
CNF/AP nanocomposites.
40 RESULTS AND DISCUSSION
Figure 14. FE-SEM images of fracture surface for (a) Neat CNF, and nanocomposites of
(b) Core-shell CNF/AP 23 wt% and (c) “Mixed” CNF/AP 25 wt% films. These samples
were fractured under liquid N2. TEM images of ultra-thin sections for (d) the Core-
shell CNF/AP 23 wt% and (c) the “Mixed” CNF/AP 25 wt% also show the nanostructure
of the films at fine scale. The figure is modified from Paper I.
Figure 15. FE-SEM images of tensile fractured surface for (a) Neat CNF, (b) Core-shell
CNF/XG 31 wt% and (c) “Mixed” CNF/XG 32 wt% nanocomposite films. The figure is
taken from Paper II.
RESULTS AND DISCUSSION 41
4.1.3 Hygromechanical properties of Core-shell vs “Mixed” nanocomposites
The main focus in this section is related to the importance of matrix
distribution for hygromechanical properties of the biocomposites. The
mechanical properties of polysaccharide-based materials, including CNF
networks,69 are known to be very sensitive to moisture. The effect of matrix
distribution as well as interface structure in Core-shell and “Mixed”
nanocomposites is investigated with respect to mechanical properties, especially
at highly hydrated conditions.
Wet conditions
Let us consider the severe case when water is saturated in the system. The
comparison of Core-shell and “Mixed” nanocomposites was performed at two
different hydration states. First, wet gels of Neat CNF in addition to Core-shell and
“Mixed” CNF/XG were prepared and subjected to tensile testing (Figure 16(a) and
Table 4). The CNF/XG nanocomposites have common designs with the primary
cell walls. Data shows that XG provides physical linkages within the CNF
network,46 leading to an improvement in extensibility of the two CNF/XG
nanocomposites in wet condition.38 This is analogous to its biological function in
the wood cell walls.118 There is; however, a significant difference in tensile
properties in terms of Young’s modulus and strength for the two CNF/XG
nanocomposites. Depending on the preparation procedure, which is related to
the matrix distribution and interface structure, the mechanical properties of
CNF/XG wet gels are regulated in different manners. Core-shell CNF/XG
nanocomposites show unique improvements in mechanical properties.
Interestingly, the wet gel of Core-shell CNF/XG offers better performance in terms
of high modulus, strength, and extensibility compared with Neat CNF (at the
same dry content), despite XG matrix content up to 30 wt%. In addition to the
wet gels, soaked films of Core-shell and “Mixed” CNF/XG nanocomposites were
prepared and subjected to tensile testing (Figure 16(b) and Table 4). The effect of
different matrix distributions on the mechanical properties is very pronounced.
The soaked film of Core-shell CNF/XG shows better tensile properties than
“Mixed” nanocomposites. This is in good agreement with the study on wet gels.
Extensive cross-linking as well as effective matrix distribution at nanostructural
scale is an advantage for Core-shell CNF/XG nanocomposites.
42 RESULTS AND DISCUSSION
The study on mechanical properties in wet state of Core-shell nanocomposites
was further developed for the GGM matrix. This is mimicking the main
hemicellulose component of the secondary cell walls of spruce. As can be seen in
Figure 16(c), Core-shell CNF/GGM also showed significant improvement in tensile
properties as compared to Neat CNF. The corresponding tensile data of Core-shell
CNF/GGM is listed in Table 5. The development of the mechanical performance is
consistent with GGM content, which is an indirect indication of good dispersion
and homogeneous nanocomposite structure. The final GGM content was varied
in the narrow range between 5.6-13.4 wt% due to its low binding capability. The
adsorption results indicate that the majority of GGM exists in the form of bound
matrix (Figure 10(c)), presumably as a GGM coating of CNF. According to results
from AFM measurement, nanostructure of the nanofibers (Figure 12(d,e)) in
addition to considerations of the GGM content, it suggests that the presence of a
GGM coating facilitates physical linkages between adjacent nanofibers.113 The
GGM coating is closely integrated within the CNF network, resulting in Core-shell
CNF/GGM nanocomposites with superior mechanical properties under wet
conditions.
RESULTS AND DISCUSSION 43
Figure 16. Tensile stress-strain curves for CNF/XG and CNF/GGM nanocomposites
under wet conditions. The data are obtained from different testing conditions for (a)
wet gels of CNF/XG (after filtration without drying, see Table 4), (b) soaked films of
CNF/GGM (solid films are first dried and then soaked in water for 30 min, see Table
4) and (c) wet gels of CNF/GGM (after filtration without drying, see Table 5). The
figure is a combination of data from Paper II-III.
44 RESULTS AND DISCUSSION
Table 4 Tensile properties of Neat CNF, Core-shell CNF/XG 31 wt% and “Mixed”
CNF/XG 32 wt% nanocomposites as wet gels and soaked films. For the wet gel, the
suspension of nanofibers was filtered, and then the resulting material was subjected
to mechanical testing. When the wet gel was first dried and soaked in water for 30
minutes, this material is referred to as “Soaked films”. The data is extracted from
Paper II.
Wet gels (no drying) Soaked films (30 min)
Neat CNF Core-shell
CNF/XG
“Mixed”
CNF/XG
Core-shell
CNF/XG
“Mixed”
CNF/XG
Dry content (%) 22.7 24.0 21.7 23.7 22.7
Young’s modulus
(MPa)
9.5 ± 0.6 27.1 ± 4.8 3.3 ± 0.6 32.7 ± 8.1 7.2 ± 1.3
Tensile strength
(MPa)
0.4 ± 0.1 1.3 ± 0.4 0.2 ± 0.1 3.0 ± 0.8 1.1 ± 0.4
Elongation (%) 15.6 ± 3.8 41.9 ± 7.4 36.4 ± 7.0 32.2 ± 4.7 27.8 ± 4.8
Work of fracture
(kJ/m3)
34.8 ± 9.7 370.1 ± 93.9 47.5 ± 18.6 559.6 ± 112.3 148.1 ± 25.7
Table 5 Tensile properties for wet gels of Neat CNF and Core-shell CNF/GGM
nanocomposites with various GGM contents. The data is taken from Paper III.
Neat CNF
Core-shell CNF/GGM
(GGM content – wt%)
5.6 10.6 13.4
Solid content (%) 14.9 15.1 18.0 16.9
Young’s modulus (MPa) 4.1 ± 0.5 5.9 ± 2.2 21.3 ± 2.0 28.8 ± 6.6
Tensile strength (MPa) 0.11 ± 0.04 0.15 ± 0.03 0.33 ± 0.04 0.65 ± 0.16
Elongation (%) 9.2 ± 2.1 8.3 ± 1.0 15.8 ± 3.0 10.9 ± 1.3
Work of fracture (kJ/M3) 8.2 ± 3.0 7.8 ± 1.7 46.0 ± 10.1 50.8 ± 12.8
RESULTS AND DISCUSSION 45
Hydrated conditions (at relative humidity of 50 and 85 RH%)
There have been many studies on the improvement in mechanical properties
of CNF-based polymer films.48, 57, 76, 119, 120 The mechanical properties of
CNF/biopolymer films at high humidity are often critical for final engineering
applications. The aim of this experiment is to develop polysaccharide-based
biocomposite films with improved hygromechanical properties. The idea is
related to the outstanding mechanical behavior of Core-shell nanocomposites in
wet conditions, as shown in the previous section.
Starch is a polysaccharide, which has been widely investigated for potential
applications such as film and foam materials. There are, however, some
drawbacks with respect to the material performance and processing.
Amylopectin (AP) is the major constituent of starch. Previous studies in our
research group have demonstrated the potential of CNFs for plasticized AP. The
obtained CNF/plasticized AP films showed superior mechanical properties8 with
reduced moisture sensitivity,71 compared to the pure plasticized AP film. The
improvement of CNF/plasticized AP was primarily due to the inherent
characteristics of CNFs and their network. As the content of plasticized AP was
increased, the resulting materials showed strong reduction in mechanical
properties and moisture resistance.
In this experiment, Core-shell and “Mixed” CNF/AP nanocomposites were
prepared in the absence of plasticizer. Tensile behavior and corresponding
properties of the CNF/AP nanocomposites conditioned at 50 and 85 RH% are
presented in Figure 17 and Table 6. The difference in matrix distribution between
Core-shell and “Mixed” CNF/AP nanocomposites shows remarkable effects on
their hygromechanical properties. Outstanding mechanical properties can be
achieved in the case of Core-shell CNF/AP nanocomposites, which are comparable
to Neat CNF, despite variation in AP content. The tensile behavior of Neat CNF
and Core-shell CNF/AP in elastic and plastic regions is always consistent at
particular relative humidity (50 and 85 RH%). It is evident that the CNF network
completely dominates the properties of Core-shell CNF/AP nanocomposites, while
no such effect is observed for the “Mixed” CNF/AP. The reason is primarily due to
the control of bound matrix, which is coated on the nanofiber surface, in the case
of Core-shell nanocomposites.
46 RESULTS AND DISCUSSION
Figure 17. Tensile stress-strain curves for solid films of Neat CNF, Neat AP, Core-shell
CNF/AP 23 wt%, and “Mixed” CNF/AP 25 wt% nanocomposites at (a) 50 RH% and (b)
85 RH%. The figure is modified from Paper I.
RESULTS AND DISCUSSION 47
Table 6 Tensile properties for solid films of Neat CNF, Core-shell CNF/AP 23 wt% and
“Mixed” CNF/AP 25 wt% nanocomposites at different relative humidities (50 and 85
RH%). The values in parentheses are standard deviations. The data is extracted from
Paper I.
50 RH% 85 RH%
Neat
CNF
Neat
AP
Core-shell
CNF/AP
“Mixed”
CNF/AP
Neat
CNF
Neat
AP
Core-shell
CNF/AP
“Mixed”
CNF/AP
Density
(kg/m3)
1375 1445 1449 1415 - - - -
Porosity
(%)
8.5 NA 3.4 5.7 - - - -
Modulus
(GPa)
14.6
(0.7)
2.3
(0.2)
13.6
(0.2)
11.0
(0.7)
12.3
(0.8)
1.7
(0.3)
10.8
(0.5)
8.6
(0.4)
Tensile
strength
(MPa)
227
(26)
43
(4)
221
(17)
172
(20)
155
(26)
36
(3)
151
(12)
100
(12)
Yield
strength
(MPa)
121
(2)
46
(2)
117
(2)
90
(8)
60
(2)
33
(2)
69
(2)
50
(3)
Elongation
(%)
7.3
(1.5)
6.1
(1.8)
7.9
(1.5)
8.2
(1.3)
7.8
(1.5)
6.3
(0.6)
7.8
(0.6)
7.0
(1.7)
Work of
fracture
(MJ/m3)
11.4
(3.8)
2.4
(0.7)
12.1
(3.1)
9.7
(2.3)
8.0
(3.8)
1.9
(0.2)
8.0
(1.3)
4.9
(1.8)
In addition to AP, two structurally different hemicelluloses, XG and GGM,
were employed as a polysaccharide matrix for Core-shell nanocomposites. The
effects of matrix distribution and nanostructural interface on the
hygromechanical properties of Core-shell and “Mixed” CNF/XG conditioned at 50
and 85 RH% are demonstrated in Figure 18(a,b). Stress-strain curves of the two
CNF/XG nanocomposites in comparison with Neat CNF and Neat XG films are
provided. Compared with Neat CNF, the two CNF/XG nanocomposites showed
48 RESULTS AND DISCUSSION
lower stiffness, strength and elongation at 50 RH%. The result is in accordance
with the mechanical properties of biocomposites based on wet gel of bacterial
cellulose and XG.38 There is no significant difference between Core-shell and
“Mixed” CNF/XG at this condition (50 RH%). Perhaps strong interfacial
interaction121 between CNFs and XG in addition to the domains of entangled
and/or cross-linked fibrils38 are responsible for reduced ductility, reflecting
premature failure in the two CNF/XG films. At elevated humidity (85 RH%), the
two CNF/XG nanocomposites exhibited similar characteristics of lower stiffness
and strength than that of Neat CNF, whereas greater extensibility was observed in
the case of Core-shell CNF/XG film.
Tensile stress-strain curves for Core-shell CNF/GGM films, with varying GGM
contents in the range of 5.6-13.4 wt%, conditioned at 50 and 85 RH% are
presented in Figure 18(c,d). It is interesting to note that Neat CNF and Core-shell
CNF/GGM undergo similar elastic deformation. At 50 RH%, the modulus of Neat
CNF and Core-shell CNF/GGM films was varied within the same range of 13.5-13.8
GPa. When the humidity was elevated from 50 to 85 RH%, the moisture uptake of
Neat CNF and Core-shell CNF/GGM films was increased from approximately 7 to
13 wt%. Thus, the modulus was slightly decreased to the range of 12.0-12.8 GPa.
Strong strain-hardening48 was observed for Neat CNF and all Core-shell CNF/GGM
films, in spite of premature failure in the case of Core-shell CNF/GGM. The slope
observed during plastic deformation was comparable in all samples, and seems
insensitive to the GGM content. Strong interfacial adhesion between CNFs and
GGM is a likely explanation.
In brief, a significant difference in hygromechanical properties compared
between Core-shell and “Mixed” nanocomposite films is observed in the case of
the CNF/AP system. The evidence of nanostructual and interface effects is less
pronounced in CNF/hemicellulose films (CNF/XG and CNF/GGM). The reason is
primarily due to strong interfacial adhesion between CNFs and hemicelluloses.122,
123 The majority of hemicellulose matrix is intimately associated with the CNF
surface and confined within their nanofibrillar network. Only trace amount of
unbound hemicellulose is present in the CNF/hemicellulose nanocomposites.
Thus, the effect from “bulk” hemicellulose in “Mixed” nanocomposite films is
weak, despite strong effects in water saturated conditions.
RESULTS AND DISCUSSION 49
Figure 18. Tensile stress-strain curves for solid films of (a,b) CNF/XG and (c,d)
CNF/GGM nanocomposites at 50 and 85 RH%, respectively. The figure is a
combination of data from Paper II-III.
4.1.4 Moisture sensitivity of Core-shell vs “Mixed” nanocomposites
The inherent hygroscopic properties of polysaccharides are attributed to polar
hydroxyl groups and many polysaccharides have high solubility in water. CNFs,
however, possess attractive characteristics of semi-crystalline fibrillar structure.124
Moreover, they can form a dense nanofibrillar network,48 and have the ability to
provide strong molecular interaction via stabilized hydrogen bonding. Therefore,
there has been considerable interest to reduce moisture sensitivity of
polysaccharide-based materials by utilizing CNFs.10, 11, 70, 71 It has been suggested
that the existence of strong adhesion of the CNF/polysaccharide matrix interface
contributes to a reduction in moisture sensitivity.71, 125, 126 The aim of this study is
to investigate differences in interfacial effects at molecular and nanoscale.
50 RESULTS AND DISCUSSION
Moisture sorption isotherms compared between the nanocomposite films of
Core-shell CNF/XG 31 wt% and “Mixed” CNF/XG 32 wt% are plotted in Figure 19.
Prediction for moisture sorption of the CNF/XG nanocomposite based on rule of
mixtures; i.e. calculated on the basis of the moisture uptake contribution from the
individual components, is also displayed in the figure. It is apparent that the two
CNF/XG nanocomposites adsorb less moisture than expected, suggesting strong
molecular interactions between CNFs and XG in nature. Core-shell CNF/XG
nanocomposite exhibits the most promising result, where the moisture sorption
is almost identical to Neat CNF, despite an XG content of 31 wt%. This indicates
that the available sites for moisture sorption in the amorphous polymer are more
restricted in the case of Core-shell nanocomposites.
The moisture sorption data for Core-shell and “Mixed” CNF/AP
nanocomposite films are presented by showing the difference between
experimental data and predicted values based on the rule of mixtures, see Figure
20. Negative value of the differences means that the material adsorbs less
moisture than expected. The result consistently reveals that Core-shell CNF/AP
nanocomposite is less susceptible to moisture sorption than the “Mixed”
nanocomposite. This is in support of the difference in moisture sorption behavior
for the two CNF/XG nanocomposites. In Figure 20, data for Core-shell CNF/GGM
with varying GGM contents is also displayed. It is verified that Core-shell
nanocomposites generally adsorb less moisture than expected. This effect can be
correlated to the unique polymer matrix distribution for the Core-shell
nanocomposites. The more controlled nanostructural interphase is probably one
of the key reasons.
RESULTS AND DISCUSSION 51
Figure 19. Dynamic vapor sorption isotherms for solid films of Neat CNF, Neat XG,
“Mixed” CNF/XG 32 wt%, and Core-shell CNF/XG 31 wt%. Predictions for CNF/XG 31
wt% (MCNF/XG) is based on sorption in neat components and rule of mixtures; MCNF/XG =
MCNFWCNF + MXGWXG where M is measured moisture content, and W is weight
fraction of the component. The inset shows isotherms for CNF-based films in the
range 40-90% RH. The figure is taken from Paper II.
Figure 20. Comparison of moisture sorption between experimental data and
calculated values (based on rule of mixtures) for CNF/AP and CNF/GGM
nanocomposites. Negative values mean that the material adsorbs less moisture than
predicted from rule of mixtures. The figure is extracted from Paper I and II.
52 RESULTS AND DISCUSSION
To investigate the difference in material-moisture interaction between Core-
shell and “Mixed” CNF/XG nanocomposites at the molecular scale, measurement
of moisture mobility derived from 2H solid state NMR relaxation was performed.
The 2H NMR analysis of magnetization (M0r) and relaxation time (T2r) in addition
to moisture content (MCr) for Neat XG, Neat CNF, Core-shell CNF/XG 31 wt%, and
“Mixed” CNF/XG 32 wt% at 92 RH% are presented in Figure 21. In Figure 21(a),
the reported values are normalized to the values for Neat XG. M0r correlates with
the number of 2H atoms, which is related to the moisture content (MCr) in the
material. The moisture mobility in the Core-shell nanocomposites is significantly
reduced from that of “Mixed” nanocomposites, indicated by lower T2r in the Core-
shell CNF/XG film. In spite of differences in the water and matrix contents (Figure
21(b)) between Core-shell nanocomposite and Neat CNF, it is interesting to observe
that the water mobility in these materials is comparable. This indicates that water
clustering for Core-shell nanocomposites is localized at fibril surfaces, similar to
the Neat CNF.126 In “Mixed” nanocomposites, the water clustering results in faster
NMR-dynamics due to the presence of “unbound/bulk” matrix. This study
provides improved understanding of the mechanisms associated with the
nanostructural effects in the nanocomposites.
RESULTS AND DISCUSSION 53
Figure 21. (a) Ratios of gravimetric deuterated water content (MCr), NMR
magnetization (M0r) and relaxation time (T2r) for Neat XG, Neat CNF, “Mixed” CNF/XG
32 wt%, and Core-shell CNF/XG 31 wt% at 92 RH%. The reported values are
normalized to the values for Neat XG. The values in parentheses next to sample label
are the average moisture content from the NMR conditioning protocol. (b) A plot of 2H NMR T2 relaxation time as a function of gravimetric heavy water content for all
samples. Note that larger values for T2 correlate with increased water mobility. The
figure is taken from Paper II.
54 CONCLUSIONS
4.2 Conclusions (from Paper I-III)
We have investigated an alternative route for high-performance and moisture-
stable nanocomposites based on Core-shell concepts. Inspired by the cell wall
biosynthesis, Core-shell nanofibers have been prepared by coating wood CNFs
with a thin layer of polysaccharide matrix. This allows us to create nanofibers
with Core-shell structure, where the CNFs serve as the “core” reinforcing element
with the “shell” layer of polymer matrix. The preparation procedure in
combination with the data obtained from adsorption, colloidal particle size,
nanofiber diameter and high-resolution microscopy images are in support of the
Core-shell hypothesis. The nanocomposite network of Core-shell nanofibers has
been formed through a water-based papermaking-like process without polymer
loss. The Core-shell nanocomposites have advantages with respect to
nanostructural control of the polymer matrix distribution and interface structure
over a reference “Mixed” nanocomposites obtained from conventional mixing.
The nanostructural effects on mechanical and moisture-related properties for
the two nanocomposites have been investigated at different hydration states. The
Core-shell nanocomposite concept results in unique improvements in mechanical
properties and moisture stability compared with “Mixed” nanocomposites.
However, the extent of improvement in material performance for Core-shell
nanocomposites depends on the nature of the polymer matrix. For CNF/AP films,
the Core-shell nanocomposite conditioned at 50 and 85 RH% reveals strong
improvements in mechanical properties compared with the reference material.
This is due to nanostructural effects in terms of matrix distribution and interface
structure. For the CNF/hemicellulose samples (CNF/XG and CNF/GGM), the
Core-shell nanocomposites show a strong increase in mechanical properties
compared with “Mixed” nanocomposite and Neat CNF in wet conditions. The
enhancement arises from combined effects of favorable matrix distribution,
physical crosslinking, and strong interfacial adhesion between CNFs and
hemicelluloses. The results from 2H solid state NMR relaxation support the
difference in matrix distribution between Core-shell and “Mixed” nanocomposites
at the molecular level. In the Core-shell nanocomposites, strongly reduced water
mobility is obtained, possibly due to the confined XG matrix associated with the
CONCLUSIONS 55
CNF surface. Significant increase in water mobility is noticed for the “Mixed”
nanocomposites due to the presence of “bulk” or “unbound” matrix. The
increased moisture sorption in “Mixed” nanocomposites is in good agreement
with the NMR results. This can be explained in terms of the different states of
polymer matrix distribution in Core-shell and “Mixed” nanocomposites at
nanoscale. The Core-shell approach allows preparation of nanocomposite
materials with controlled nanostructure and improved hygromechanical
properties. Thus, the influence of polymer matrix distribution and interphase
structure has been clarified.
56 RESULTS AND DISCUSSION
4.3 Holocellulose nanofibers and their biocomposite films (Paper
IV-V)
4.3.1 Holocellulose nanofibers
Fundamental properties of nanofibers such as morphology, molecular weight,
surface chemistry and ease of nanofibrillation are importance for the preparation,
structure and properties of the nanocomposite materials. In addition to CNFs
obtained by enzymatic3 or chemical pretreatments,5 holocellulose nanofibers
containing high fraction of native hemicelluloses (up to 20 wt%) can be prepared.
Delignification processes such as chlorite treatment55 can be applied to obtain
holocellulose fibers prior to mechanical disintegration. Nevertheless, extensive
degradation of cellulose can be a drawback during chlorite delignification.127 It
suggests that the decrease in molecular weight of cellulose has a negative
influence on strength of the resulting material.48
The motivation of Paper IV is to introduce an alternative route for preparation
of holocellulose nanofibers. Peracetic acid pretreatment was performed directly
on wood chips prior to mechanical disintegration. In Figure 22, it is apparent that
the starting wood fibers are still intact after the peracetic acid pretreatment. The
fibers were progressively disintegrated into nanofibers by increasing the number
of passes through the microfluidizer. To evaluate the ease of disintegration for
nanofibers, the obtained materials are observed using FE-SEM. Sufficient
nanofibrillation of the holocellulose fibers was achieved after 1+5 passes (1 and 5
passes through the big and small chambers, see experimental section). Thus, the
disintegration of holocellulose nanofibers is more efficient in terms of nanofiber
structure and energy input than that required for the enzyme-pretreated CNFs
(3+5 passes). As shown in Figure 22, the obtained suspension of holocellulose
nanofibers shows good optical transparency, suggesting more uniform and
individualized nanofibers than the enzymatically pretreated CNFs commonly
used in our laboratory. The nanostructure observed by TEM and AFM in
addition to the distribution of holocellulose nanofiber diameters obtained by
AFM height analysis is presented in Figure 23. It is confirmed that the
holocellulose nanofibers are well-individualized. The diameter of holocellulose
RESULTS AND DISCUSSION 57
nanofibers is in the range of 2-11 nm. The majority of the holocellulose nanofibers
have diameters around 3-6 nm, which is closely related to the diameter of
individualized nanofibers (3-4 nm).5, 22 A significant decrease in large fibril
aggregates is observed, as compared to the reported data for CNFs in Paper I-II.
Figure 22. FE-SEM images of peracetic acid treated aspen wood after passing through
the microfluidizer with varying number of passes. The nomenclature of nanofibers is
referred to the number of pass through big + small chambers (1+5 means 1 and 5
passes through the big and small chambers, respectively). The photograph shows the
obtained suspension of holocellulose nanofibers compared with the reference
(enzyme-pretreated CNFs). The solid content of the suspensions is approximately 1-2
wt%. The figure is modified from Paper IV.
58 RESULTS AND DISCUSSION
Figure 23. Micrographs from (a) TEM and (b) AFM images of holocellulose nanofibers
obtained from aspen wood prepared by peracetic acid pretreatment and mechanical
distintegration. The sample was obtained after 6 passes (1+5) through the
microfluidizer. (c) Distribution of diameter for the holocellulose nanofibers based on
AFM height images. The figure is taken from Paper IV.
Two different sources of wood fibers, including aspen and spruce, were
subjected to peracetic acid pretreatment and nanofibrillation. Molar mass
analysis of these holocellulose nanofibers is compared with spruce enzyme-
pretreated CNFs in Table 7. The result reveals that holocellulose nanofibers of
high molar mass could be obtained from peracetic acid pretreatment. The
obtained nanofibers also show considerably lower values of polydispersity index
(PDI). It indicates a more narrow distribution of molar mass, which correlates
with nanofiber length.54 Compared with CNFs partially disintegrated with 1+0
passes, it is apparent that peracetic acid pretreatment shows promising result
with respect to high and preserved cellulose molar mass. Thus, the molar mass of
cellulose in the nanofibers is strongly influenced by the pretreatment approach.
The ease of nanofibrillation relies on chemical characteristics such as
hemicelluloses content55, 128 and surface charge.5 Table 8 summarizes these
characteristics of the nanofibers obtained from peracetic acid and enzymatic
pretreatments. It is evident that a considerable amount of native hemicelluloses
(up to 24 wt%) are preserved during the peracetic acid pretreatment. The charge
content of holocellulose nanofibers was approximately 200 µeq/g. Overall, the
hemicellulose and charge content of the holocellulose nanofibers is roughly twice
RESULTS AND DISCUSSION 59
as high as enzyme-pretreated CNFs. This plays an important role in the enhanced
nanofibrillation of the holocellulose nanofibers.
Table 7 Molar mass analyses of nanofibers obtained from peracetic acid and
enzymatic pretreatments determined using SEC. A different number of passes during
mechanical disintegration was applied to the nanofibers. The nomenclature of
nanofibers is referred to the number of passes through big + small chambers,
respectively. The data is taken from Paper IV.
Wood species Pretreatment Number of
passes
Molar mass
x 106 (g/mol)
Polydispersity
index (PDI)
Aspen Peracetic acid 1+0
1+1
1+3
1+5
1.3
1.3
0.83
0.83
19
15
11
7.0
Spruce Peracetic acid 1+0
1+1
1+3
1+5
1.1
0.83
0.62
0.54
12
9.4
6.5
6.3
Spruce Enzymatic 1+0
3+0
3+1
3+5
0.80
0.73
0.67
0.47
26
19
19
13
Table 8 Hemicellulose and charge contents of nanofibers obtained from peracetic acid
and enzymatic pretreatments. Different number of passes during mechanical
disintegration was applied to the nanofibers. The data is taken from Paper IV.
Wood species Pretreatment Number of
passes
Hemicellulose
content (wt%)
Charges
(µeq/g)
Aspen Peracetic acid 1+5 23.9 193
Spruce Peracetic acid 1+5 23.1 207
Spruce Enzymatic 3+5 12.3 101
60 RESULTS AND DISCUSSION
4.3.2 Films from holocellulose nanofibers
Optical transparency of holocellulose nanopaper film
Individualized nanofibers with uniform diameter is a key characteristic to
achieve high optical transparency.65 The light transmittance spectrum of the dried
films obtained from holocellulose nanofibers and enzyme-pretreated CNFs is
shown in Figure 24(a). The corresponding photograph of the obtained films with
a thickness of approximately 45 µm is presented in Figure 24(b). It is apparent
that the level of light transmittance for the holocellulose film was significantly
higher than that from the enzyme-pretreated CNFs, which can be explained in
terms of their differences in nanostructural characteristics. Interestingly, the
results are approaching the reported data for TEMPO-oxidized nanocellulose
films.65
Figure 24. (a) Light transmittance spectrum of spruce nanopaper prepared by
peracetic acid and enzymatic pretreatments and (b) photograph shows their light
transmittance behavior. The figure is taken from Paper IV.
Mechanical properties of holocellulose film
Tensile stress-strain curves of the holocellulose films prepared by peracetic
acid pretreatment are plotted in Figure 25, along with the corresponding tensile
properties in Table 9. The results are also compared with a typical film obtained
from enzyme-pretreated CNFs. An increase in mechanical properties in terms of
high Young’s modulus, tensile strength, and elongation was observed for the film
obtained from holocellulose nanofibers. The modulus was as high as 16.8 GPa
and 15.7 GPa for the holocellulose film from spruce and aspen, respectively. The
RESULTS AND DISCUSSION 61
fracture surface of holocellulose films exhibited a very dense structure. In
contrast, the film from enzyme-pretreated CNFs showed a typical layered
structure with porosity.48 Most likely, the unique architecture of holocellulose
films is influenced by the high fraction of native hemicelluloses. The tensile
strength of holocellulose films ranged from 290 to 327 MPa, depending on the
wood origin. Compared to the strength (243 MPa) and elongation (6.0%) of the
reference film (spruce enzyme-pretreat CNFs), the spruce holocellulose film
showed a strength of 290 MPa and elongation of 6.6%. These values reflect the
advantage of holocellulose nanofibers with high molar mass and uniform
diameter, obtained from peracetic acid pretreatment.
Figure 25. Comparison of tensile stress-strain curves and FE-SEM images of tensile
fracture surfaces for nanocellulose films prepared by peracetic acid and enzymatic
pretreatments. The figure is taken from Paper IV.
62 RESULTS AND DISCUSSION
Table 9 Tensile properties of nanocellulose films prepared by peracetic acid and
enzymatic pretreatments. The data is taken from Paper IV.
Wood
species
Pretreatment Number
of passes
Young’s
modulus
(GPa)
Tensile
strength
(MPa)
Elongation
(%)
Aspen Peracetic acid 1+5 15.7 ± 0.5 327 ± 12 7.2 ± 0.3
Spruce Peracetic acid 1+5 16.8 ± 1.4 290 ± 18 6.6 ± 0.4
Spruce Enzymatic 3+5 14.9 ± 1.5 243 ± 9 6.0 ± 0.4
4.3.3 Porous materials from holocellulose nanofibers
Honeycombs vs Foams
CNFs can serve as a reinforcement network for novel cellulosic materials such
as porous foams.102, 104 The nanoscale fibrils allow for tailoring cellular structures
of the porous materials at a fine scale. The purpose of this study is to prepare
CNF-based honeycombs, where the cellular structure of the porous materials is
better controlled. Honeycombs with different densities were prepared from
colloidal suspensions of holocellulose nanofibers with varying concentrations.
The relationship between the nanofiber concentration in suspension and density
of the resulting porous material is shown in Figure 26(a). Holocellulose
honeycombs with a limited range of densities (8-21 kg/m3) were obtained. It is
challenging to achieve high-density holocellulose honeycombs, due to the
difficulty to concentrate the suspension of holocellulose nanofibers. Compressive
stress-strain curves for the obtained holocellulose honeycombs are displayed in
Figure 26(b). The mechanical behavior of holocellulose honeycombs is strongly
dependent on the density. High density honeycombs resulted in improved
mechanical properties with respect to modulus, plateau stress and energy
absorption. Most likely, the cell wall thickness is increased resulting in increasing
density of the structural honeycombs.96
RESULTS AND DISCUSSION 63
Figure 26. (a) Concentration of nanofiber suspension before freezing and the resulting
density of honeycombs. (b) Longitudinal compressive stress-strain curves for
holocellulose honeycombs with varying densities. The figure is taken from Paper V.
Holocellulose honeycombs with highly aligned cell walls are expected due to
controlled directional freezing during preparation. Figure 27 illustrates the
hierarchical structure of the obtained holocellulose honeycombs. At
nanostructural level, the structural cell walls were built from a nanofibrillar
network of the holocellulose nanofibrils. At microstructure level, an elongated
and continuous porous network with well-defined pore connectivity was created.
It is evident that the cell geometry is highly anisotropic, and the alignment of the
cells is parallel to the freezing direction.
To investigate the relationship between cell wall structure and mechanical
behavior of the honeycombs, a comparison of Young’s modulus along
longitudinal and transverse directions is presented in Figure 28(a). There is a
substantial difference in modulus between the two directions. This unique
response confirms the anisotropic structure of the holocellulose honeycombs. In
contrast, for 3-dimensional cellular holocellulose foams (Figure 28(b)), only a
slight difference in modulus between the two directions was observed. This
reflects the low degree of anisotropy in cell shape of the foams. For the
honeycombs (Figure 28(a)), it is also interesting to observe a remarkable
improvement of the modulus in the longitudinal direction when the density is
increased. The modulus of the honeycombs was raised by two orders of
magnitude when the density was increased from 6 to 21 kg/m.3 This is
considerably improved compared to the reported data for CNF-based foams104
due to the anisotropy of cell geometry in the honeycombs.
64 RESULTS AND DISCUSSION
Figure 27. FE-SEM images of hierarchical structure for holocellulose honeycombs
with a density of 12 kg/m3. The yellow arrow indicates the freezing direction. The
figure is taken from Paper V.
Figure 28. Young’s modulus of holocellulose (a) 2-dimensional honeycombs and (b) 3-
dimensional foams compared in longitudinal and transverse loading directions. The
figure is taken from Paper V.
The micromechanics scaling models for honeycombs presented by Gibson
and Ashby96 allows us to predict the modulus of solid cell wall material (Es)
obtained from the holocellulose nanofibers. In Figure 29(a), the data for Young’s
modulus of the holocellulose honeycombs are plotted against relative density
(ρ*/ρs). By fitting the modulus data of holocellulose honeycombs to the equation
in Figure 29(a), a modulus of 133 MPa is estimated for the solid cell wall. This
value is two orders of magnitude lower than that for nanopaper structures of
holocellulose nanofibers (16.8 GPa), which should be representative for the solid
cell wall. There are several possible explanations. Most likely, the degree of
orientation of the cells is substantially off-axis. The relationship between cell axis
orientation and axial modulus is highly sensitive to orientation angle at small
RESULTS AND DISCUSSION 65
angles. In addition, defects in the form of large cell wall pores and smaller scale
porosity may also contribute to the reduced modulus.
To demonstrate the influence of microstructural control on mechanical
behavior of the porous materials, the reported data of modulus and energy
absorption for holocellulose foams in addition to nanocellulose foams from
previous study104 are also included in Figure 29(a,b). The modulus of the present
honeycombs is improved considerably compared to the foams, as expected from
theoretical predictions.96
Figure 29. (a) Young’s modulus and (b) energy absorption along longitudinal
direction of the honeycombs prepared from holocellulose nanofibers, plotted against
relative density. The result is presented with the data of holocellulose foams and
CNF-based foams from previous study by Sehaqui et al. (2010).104 The figure is taken
from Paper V.
Holocellulose vs Cellulose nanofibers
Inherent characteristics of the nanofibers, such as their mechanical properties
and ability to fuse and form a network are important for tailoring the porous
materials with high-performance and well-controlled structure. The focus of this
study is to compare the mechanical properties and cell wall structure of
honeycombs obtained from holocellulose versus cellulose nanofibers. The
holocellulose nanofibers were prepared from peracetic acid pretreatment,
whereas the cellulose nanofibers were prepared from enzymatic approach.
In Figure 30(a,b), the modulus and energy absorption of honeycombs
obtained from the two different nanofibers are plotted against relative density.
66 RESULTS AND DISCUSSION
The modulus of holocellulose honeycombs was significantly increased
(approximately 1 order of magnitude) from that of nanocellulose honeycombs,
although there was no difference in the energy absorption.
To learn about the details of the cellular structure, the microstructure of
honeycombs obtained from holocellulose and cellulose nanofibers was observed
using FE-SEM (Figure 31). The unique structure of honeycombs is apparent in
both cases. However, the cellular structure of holocellulose honeycombs is built
from a well-connected and more regular cell network. The microstructure of
holocellulose honeycombs partly explains the higher modulus. In contrast, the
honeycombs obtained from enzyme-pretreated CNFs (right image) contain large
pores, which explain lower honeycomb stiffness.
Figure 30. (a) Young’s modulus and (b) energy absorption along longitudinal
direction of the honeycombs prepared from holocellulose and cellulose nanofibers,
plotted against relative density. The figure is taken from Paper V.
RESULTS AND DISCUSSION 67
Figure 31. FE-SEM images of cell wall structure for honeycombs prepared from
(a) holocellulose nanofibers and (b) enzyme-pretreated CNFs. The materials have
similar density (approximately 11-12 kg/m3). The figure is taken from Paper V.
68 CONCLUSIONS
4.4 Conclusions (from Paper IV-V)
Holocellulose nanofibers are a promising nanocellulose reinforcement for
biocomposites, where the matrix is included as a CNF coating. The
corresponding nanocomposite films have improved mechanical properties
compared with previous CNF films. The holocellulose nanofibers can be
prepared directly from wood chips based on a combined process of peracetic acid
pretreatment and mechanical disintegration. The peracetic acid pretreatment
results in holocellulose nanofibers with high cellulose molar mass and uniform
CNF diameters. At the nanofibril level, the cellulose nanofiber “core” is coated by
the “shell” layer consisting of mixed native hemicelluloses. Nanofibrillation of the
pulp is facilitated, possibly due to the high content of hemicelluloses and the
presence of charged groups. The holocellulose nanofibers appear to mainly
consist of nanofibrils with a diameter as small as 3-6 nm.
By exploiting the unique characteristics of the obtained holocellulose
nanofibers, nanocellulose films with remarkable mechanical properties and
optical transparency were obtained. With respect to the mechanical properties,
the holocellulose films show Young’s modulus as high as 16.8 GPa, and tensile
strengths in the range of 290-327 MPa. The holocellulose films offer superior
mechanical properties over that obtained from enzyme-pretreated CNFs.
Moreover, the optical transparency of the holocellulose film is approaching that
obtained from TEMPO nanofibers (TOCNs). From the composite materials
perspective, structural control of nanofibers and polymer matrix distribution is
important in order to achieve desirable properties.
The potential to prepare CNF-based porous materials with improved
performance was investigated. Porous materials, including honeycombs and
foams, were prepared from nanofiber suspensions using freeze-drying
techniques. To prepare the honeycomb structure, a simple and versatile strategy
to control unidirectional freezing was applied for cell wall alignment. The
compressive modulus of the honeycombs is significantly improved compared
with foams, as expected from theoretical predictions. The effect is primarily
attributed to the cell alignment. A large difference of the modulus between
CONCLUSIONS 69
longitudinal and transverse directions is in support of a highly anisotropic
structure of the honeycombs, which is consistent with the microstructure
observed by FE-SEM. The predicted cell wall modulus was very low compared to
the modulus of holocellulose nanopaper. This indicates that improved cell
alignment and less cell wall defects would improve the properties even further.
The presence of native hemicelluloses facilitated the dispersion of holocellulose
nanofibers. Thus, the holocellulose honeycombs showed a more uniform cell
structure with less defects, as compared with honeycombs obtained from
enzyme-pretreated CNFs. This is also reflected in the increased out-of-plane
modulus for holocellulose honeycombs. The presence of native hemicelluloses in
the holocellulose nanofibers, acting as the “shell” layer, is of considerable
importance during preparation of the nanocomposite materials. In addition, the
interfiber bonding function of the hemicellulose “shell” has positive effects on the
final properties.
FUTURE WORK 71
FUTURE WORK 5
In order to enhance the mechanical properties of CNF-based films, there has
been an effort to align polymer coated88 or grafted61 nanofiber networks by
stretching. Better nanofibril alignment is achieved, resulting in CNF-based films
combining high modulus and strength. For Core-shell CNFs, the “shell” polymer is
promising and may provide nanoscale lubrication, facilitating nanofibril sliding
and reorientation in hydrated state. Wet gels of Core-shell nanocomposites can be
prepared and subjected to nanofibril alignment by stretching. As shown in the
present thesis, the wet gels of Core-shell nanofibers with hemicellulose matrices
(CNF/XG and CNF/GGM) exhibit improved stability compared with Neat CNF.
There may therefore be potential to align the wet gel of Core-shell
nanocomposites to a larger extent, and this would improve modulus and
strength.
This thesis has demonstrated that more inhomogeneous nanostructure (resin-
rich regions, CNF aggregates) has negative effects on the nanocomposite
properties. Improved hygromechanical properties and moisture stability of Core-
shell nanocomposite make these materials interesting as packaging materials.
Barrier properties of the Core-shell nanocomposite films need to be investigated. It
is possible that polymers strongly adsorbed to the cellulose nanofibers may show
improved gas barrier performance.
Core-shell nanofibers possess a large number of available hydroxyl groups
obtained from both CNFs and polysaccharide matrices. To extend the range of
properties for Core-shell nanocomposites, the surface hydroxyl groups can be
used as a molecular anchor for further functionalization or modification with
reactive chemical groups, for instance epoxides. The “shell” of neighboring CNFs
can be chemically cross-linked for improved hygromechanical stability. It may be
useful to achieve the materials with high-density of functional groups in
combination with improved moisture stability and reinforcing capability.
A successful development of CNF-based porous materials with highly
anisotropic structure (honeycombs) was demonstrated. From the viewpoint of
mechanical properties, honeycombs provide better out-of-plane performance
72 FUTURE WORK
than foams. However, the predicted modulus of the solid cell wall (Es) for
honeycombs is still far below that measured in tension for the nanopaper. It
suggests that the alignment of the structural cell wall is off-axis and that the cell
wall has defects. Therefore, better control of the cell wall alignment and cell
structure is necessary for improved mechanical properties.
ACKNOWLEDGEMENTS 73
ACKNOWLEDGEMENTS 6
First of all, I would like to express my sincere gratitude and appreciation to
my supervisor Prof. Lars Berglund for his encouragement and scientific guidance
throughout my study period. It has been a real pleasure working with him.
Special thanks are given to SCG Paper PCL (Thailand) in The Siam Cement
Group for the company’s scholarship and giving me the opportunity to
strengthen my research skills.
The Wallenberg Wood Science Center (WWSC) is gratefully acknowledged for
providing lab facilities.
I would like to acknowledge all my co-authors for their contributions to the
work. Former and present colleagues and friends in the Biocomposites group,
WWSC and the Department of Fiber and Polymer Technology are gratefully
acknowledged for making a pleasant working atmosphere. I would also like to
thank Assoc. Prof. Qi Zhou for his helpful advice. Aihua, Andong, Houssine,
Marielle and Sylvain for their kind support at the beginning of my study. Núria,
Michaela, Farhan, Gisela, Ivón, Surinthra (Janne), Anna, Mikaela, and Linn for
being really good friends and standing beside me. Ramiro is thanked for his
invaluable support in the lab. Malin and Jakob for their useful suggestions in the
group meetings. Seira, Ngesa, Joby, Qiliang, Nicholas, Andreas, Emma, and
Martin for their kinds and making the workplace more enjoyable. Special thanks
go to Shoaib for his useful suggestion for thesis preparation. Prof. Monica Ek is
acknowledged for giving support for quality check of the thesis.
I am also grateful to Khun Wichan Charoenkitsupat for giving me the
opportunity to become part of SCG Paper and always encouraging me to pursue
my ambitious goals. It meant a lot to me.
Finally, I am especially grateful to my dear family and boyfriend for offering
endless moral support and boundless love. All of their love, support and patience
give me the inner strength to go through all challenges. Thank you for always
being there for me.
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