Nanostructured Biopolymeric Materials - Diva1286485/FULLTEXT01.pdf · I. Flow-Assisted Organization...

Nanostructured Biopolymeric Materials

Hydrodynamic Assembly, Mechanical Properties and Bio-

Functionalities

NITESH MITTAL

DOCTORAL THESIS IN ENGINEERING MECHANICS

STOCKHOLM, SWEDEN 2019

Nanostructured Biopolymeric Materials

Hydrodynamic Assembly, Mechanical Properties and Bio-

Functionalities

Nitesh Mittal

Doctoral Thesis

Department of Mechanics

School of Engineering Sciences

Wallenberg Wood Science Center

KTH Royal Institute of Technology

Stockholm, 2019

Principal advisor

Docent Daniel Söderberg

Co-advisor

Professor Fredrik Lundell

Copyright © Nitesh Mittal, Stockholm, 2019

All Rights Reserved. No part of this thesis may be reproduced by any means

without permission of the author.

Cover Image © 2018 American Chemical Society

Paper I © 2018 American Chemical Society

Paper II © 2018 American Chemical Society

Paper III © 2019 American Chemical Society

Paper IV © 2017 American Chemical Society

Paper V © 2017 National Academy of Sciences USA

ISBN: 978-91-7873-088-9

TRITA-SCI-FOU 2019:5

AKADEMISK AVHANDLING

Som med tillstånd av Kungliga Tekniska Högskolan i Stockholm

framlägges till offentlig granskning för avläggande av teknologie

doktorsexamen fredag den 15e mars 2019, kl. 10.00 i sal F3,

Lindstedtsvägen 26, KTH, Stockholm. Avhandlingen försvaras på

engelska.

Opponent: Professor Chinedum Osuji, University of Pennsylvania, USA.

To my parents

मरी माता और पिता को समपिित

“Two things are infinite: the universe and human stupidity;

and I’m not sure about the universe”

−Albert Einstein

ABSTRACT

The need for high-end multifunctional materials from renewable resources has

evolved given a rapidly increasing population and accompanying environmental

concerns. Scalable assembly methods are and will be imperative in designing high-

performance environmentally friendly materials, requiring new processes allowing

control on all hierarchical levels. In this thesis, engineering concepts for

manipulation of nanoscale components from biopolymeric resources have been

applied to achieve extraordinary macroscale performance. The route chosen has been

fluid-phase assembly as it is one of the most promising methods for producing large,

ordered structures from nanoscale objects.

The thesis has three main parts; assembly of cellulose nanofibrils (CNFs) and

fundamentals associated with the processing technique, the combination of CNFs

with silk fusion proteins and finally the assembly of amyloid-like protein nanofibrils

(PNFs). In the CNFs assembly part, we have pursued the challenge of transferring

the full potential of CNFs to macroscale materials. CNFs are the most abundant

structural elements in biological systems and have impressively high strength and

stiffness, yet natural or man-made cellulose composites are much weaker than the

CNFs. We fabricated nanocellulose fibers in pursuit of maximal mechanical

performance by hydrodynamically controlling the structural ordering of nanofibrils,

resulting in continuous fibers with mechanical properties higher than any natural or

man-made macroscale biopolymeric material (Young’s modulus 86 GPa and a

tensile strength 1.57 GPa). As the hydrodynamic assembly process is largely

dependent on fundamental phenomenon controlling rotational and translational

diffusion, we have applied a novel methodology based on birefringence allowing

time-resolved in-situ investigations of diffusion and network dynamics of nanofibrils

including effects of anisotropic orientation distributions.

Genetic engineering enables the synthesis of bioengineered silk fusion proteins that

can serve as a foundation of new biomaterials. However, silk proteins are difficult

to process and cannot be obtained in large quantities from spiders. By combining

CNFs with recombinant spider silk proteins (RSPs) we have fabricated strong, tough

and bioactive nanocomposites. We demonstrate how small amounts of silk fusion

proteins added to CNFs give advanced bio-functionalities unattainable to wood-

based CNFs alone. Finally, flow-assisted assembly is applied to fabricate a material

from 100% non-crystalline protein building blocks with whey protein, a mixture with

β-lactoglobulin as the main component, which self-assemble into amyloid-like PNFs

stabilized by hydrogen bonds. We show how conditions during the fibrillation

process affect properties and morphology of the PNFs. Furthermore, we compare the

assembly of whey PNFs of distinct morphologies and show that PNFs can be

assembled into strong microfibers without the addition of plasticizers or crosslinkers.

Keywords: cellulose nanofibrils, recombinant spider silk proteins, whey protein

nanofibrils, fiber, polarized optical microscopy, orientation, diffusion, extensional

flow, tensile mechanical properties, bio-functionalities

SAMMANFATTNING

Behovet av avancerade multifunktionella material från förnyelsebara råvaror ökar

med en snabbt växande befolkning i världen och behovet av att värna om vår miljö.

Skalbara framställningsprocesser är och kommer att vara avgörande för

utformningen av dessa hållbara och miljövänliga material med hög prestanda, vilket

kräver nya processer som möjliggör kontroll av materialens struktur på alla

hierarkiska nivåer. I denna avhandling har ny teknologi för kontrollerad

sammanfogning av nanokomponenter baserade på biopolymerer tillämpats för att

uppnå extraordinära egenskaper på makroskopisk nivå. Den valda teknologin är

baserad på strömningsmekanisk sammanfogning, vilket är en lovande metod för att

kontinuerligt framställa makroskopiska strukturerade material från

nanokomponenter.

Avhandlingen består av tre huvudspår; strömningsmekanisk sammanfogning av

nanocellulosa (här syftar vi främst till cellulosananofibriller, CNF) och därtill

associerade grundläggande frågeställningar kring processen kopplat till detta;

framställning av material baserat på kombinationen av nanocellulosa med

silkesproteiner och slutligen framställning av material bestående av amyloidliknande

proteinnanofibriller (PNF). När det gäller sammanfogning av nanocellulosa har

målet vara att förstå de förhållanden under vilka framställningen av makroskopiska

material av CNF resulterar i materialegenskaper som förmår utnyttja de mekaniska

egenskaperna hos nanokomponenterna. Nanocellulosa är det mest förekommande

strukturella bio-baserade elementet på jorden och har en imponerande hög hållfasthet

och styvhet, och naturliga och framställda cellulosamaterial har hittills varit mycket

svagare än nanocellulosa. Genom att strömningsmekaniskt påverka hur

nanocellulosan bildar nanostrukturen hos kontinuerliga fiberr, med målet om

maximal mekanisk prestanda, har vi framställt fiberr som är starkare och styvare än

något annat naturligt eller konstgjort biomaterial (elasticitetsmodul 86 GPa och

draghållfasthet 1.57 GPa). Eftersom den strömningsmekaniska

framställningsprocessen i stor utsträckning är beroende av grundläggande fenomen

som rotationsdiffusion och nätverksgenerering av nanopartiklar, har vi även

tillämpat en ny metod baserad på dubbelbrytning som möjliggör tidsbesparande in-

situ karakterisering av rotationsdiffusion och nätverksdynamik för nanofibrillära

komponenter, inklusive effekterna på anisotropa orienteringsfördelning.

Genteknik möjliggör syntes av silkesproteiner som kan tjäna som grund för nya

biomaterial. Silkesproteiner är emellertid svåra att hanteras och kan bara erhållas i

små mängder från spindlar. Genom att kombinera nanocellulosa med rekombinant

spindelsilkeprotein har vi tillverkat starka, tåliga och bioaktiva nanokompositer. Vi

visar hur små mängder silkesfusionsproteiner kan kombineras med nanocellulosa för

att ge avancerad biofunktionalitet som inte kan uppnås med nanocellulosamaterial.

Slutligen har den strömningsbaserade framställningsprocessen använts att tillverka

material från 100% icke-kristallina byggstenar av vassleprotein. Dessa byggstenar

består av en blandning med p-laktoglobulin som huvudkomponent, vilken

självgenererar amyloidliknande proteinnanofibriller stabiliserade av vätebindningar.

Vi visar hur förhållandena under fibrilleringsprocessen påverkar egenskaper och

morfologin hos proteinnanofibrillerna. Slutligen framställs material av

proteinnanofibriller med olika morfologier, där vi visar hur dessa kan sammanfogas

i starka kontinuerliga fiberr utan tillsats av mjukgörare eller tvärbindande

komponenter.

Nyckelord: cellulosananofibriller, spindelsilkeproteiner, proteinnanofibriller,

fiberr, polariserad optisk mikroskopi, fiberorientering, rotationsdiffusion, töjflöden,

mekaniska egenskaper, biofunktionalitet.

LIST OF PAPERS

This thesis is a summary of the following five appended papers:

I. Multiscale Control of Nanocellulose Assembly: Transferring Remarkable

Nanoscale Fibril Mechanics to Macroscale Fibers.

Nitesh Mittal, Farhan Ansari, Krishne Gowda V., Christophe Brouzet, Pan Chen, Per

Tomas Larsson, Stephan V. Roth, Fredrik Lundell, Lars Wågberg, Nicholas A.

Kotov, and L. Daniel Söderberg.

ACS Nano, 2018, 12 (7), 6378-6388.

II. Size-Dependent Orientational Dynamics of Brownian Nanorods.

Christophe Brouzet, Nitesh Mittal, L. Daniel Söderberg and Fredrik Lundell.

ACS Macro Letters, 2018, 7 (8), 1022-1027.

III. Characterizing the Orientational and Network Dynamics of Polydisperse

Nanofibers at the Nanoscale.

Christophe Brouzet, Nitesh Mittal, Fredrik Lundell and L. Daniel Söderberg.

Macromolecules, 2019. (Accepted)

IV. Ultrastrong and Bioactive Nanostructured Bio-Based Composites.

Nitesh Mittal, Ronnie Jansson, Mona Widhe, Tobias Benselfelt, Karl M. O.

Håkansson, Fredrik Lundell, My Hedhammar, and L. Daniel Söderberg.

ACS Nano, 2017, 11 (5), 5148-5159.

V. Flow-Assisted Assembly of Nanostructured Protein Microfibers.

Ayaka Kamada,ǁ Nitesh Mittal,ǁ L. Daniel Söderberg, Tobias Ingverud, Wiebke

Ohm, Stephan V. Roth, Fredrik Lundell, and Christofer Lendel.

Proceedings of the National Academy of Sciences of the United States of America

(PNAS), 2017, 114, 1232-1237. (ǁCo-first author)

Author's contribution to the appended papers:

- Paper I, IV, V. Principal author: participated in planning and designing, performed

substantial set of experiments and part of manuscript preparation.

- Paper II, III. Co-author: participated in planning and designing, part of

experimental work (sample preparation and characterization) and part of manuscript

preparation.

Other publications not included in the thesis

VI. Ion-Mediated Control of Mechanical Properties of Nanocellulose

Filaments.*

Nitesh Mittal, Farhan Ansari, and L. Daniel Söderberg.

Manuscript in Preparation.

VII. Understanding the Mechanistic Behavior of Highly Charged Cellulose

Nanofibers in Aqueous Systems.

Lihong Geng, Nitesh Mittal, Chengbo Zhan, Farhan Ansari, Priyanka R. Sharma,

Xiangfang Peng, Benjamin S. Hsiao, and L. Daniel Söderberg.

Macromolecules, 2018, 51 (4), 1498-1506.

VIII. Dynamic Characterization of Cellulose Nanofibrils in Sheared and

Extended Semi-Dilute Dispersions.*

Tomas Rosén, Nitesh Mittal, Stephan V. Roth, Peng Zhang, L. Daniel Söderberg and

Fredrik Lundell.

Manuscript in Preparation. (Available on arXiv:1801.07558)

IX. Wettability and Humidification Induced Buried Morphological

Rearrangement at Nanoscale in Thin Films.*

Calvin J. Brett, Nitesh Mittal, Wiebke Ohm, Marc Gensch, Lucas P. Kreuzer, Volker

Körstgens, Martin Månsson, Henrich Frielinghaus, Peter Müller-Buschbaum, L.

Daniel Söderberg, and Stephan V. Roth.

Submitted Manuscript.

X. Printing of Supramolecular 3D Networks for Biosensing and Smart Wound

Dressings.* ǁ

Moien Alizadehgiashi, Daniel Pinto Ramos, Nitesh Mittal, L. Daniel Söderberg, and

Eugenia Kumacheva.

Manuscript in Preparation.

ǁThe study has been carried out at University of Toronto, Canada. The author spent

21 weeks in the group of Prof. Eugenia Kumacheva (Polymers, Interfaces and

Materials Science lab) at the Department of Chemistry as an International Visiting

Graduate Student.

*Titles and authors list may differ in the finally accepted and published version of

manuscripts that are presently under preparation/submitted.

Conference Presentations

I. Flow-Assisted Organization of Nanostructured Bio-Based Materials.

Nitesh Mittal, Fredrik Lundell, Lars Wågberg, My Hedhammar, and Daniel

Söderberg.

255th ACS National Meeting & Exposition, New Orleans, LA, USA, 2018.

II. Hydrodynamically Assembled Nanostructured Bio-Based Materials.

Nitesh Mittal, Fredrik Lundell, and Daniel Söderberg.

MRS Fall Meeting & Exhibition, Boston, MA, USA, 2017.

III. Effect of Cellulose Nanofibril Morphology on the Strength and Stiffness of

Macroscopic Filaments.


253rd ACS National Meeting & Exposition, San Francisco, CA, USA, 2017.

IV. Flow Based Approach for the Fabrication of Super-Strong Nanocellulose

Filaments- An Attempt to Utilize its Potential.



V. Elongational Flow Assists with the Assembly of Protein Nanofibrils.

Nitesh Mittal, Ayaka Kamada, Christofer Lendel, Fredrik Lundell, and Daniel

Söderberg.

APS 69th DFD Meeting, Portland, Oregon, USA, 2016.

VI. Hydrodynamic Alignment and Assembly of Carboxymethylated Cellulose

Nanofibrils: Effect of Concentration.



VII. Hydrodynamic Alignment and Assembly of Nanofibrillated Cellulose in

the Laminar Extensional Flow: Effects of Solidifying Agents.


APS 68th Annual DFD Meeting, Boston, MA, USA, 2015.

Abbreviations and Symbols

CNF Cellulose nanofibril or cellulose nanofiber

CNC Cellulose nanocrystal

PNF Protein nanofibril

RSP Recombinant spider protein or recombinant spider silk protein

TEMPO 2,2,6,6-tetramethylpiperidinyl-1-oxyl

IgG Immunoglobulin G

HDF Human dermal fibroblast

HDMEC Human dermal microvascular endothelial cells

AFM Atomic force microscopy

FE-SEM Field-emission scanning electron microscopy

TEM Transmission electron microscopy

FT-IR Fourier-transform infrared spectroscopy

SAXS Small-angle X-ray scattering

WAXS Wide-angle X-ray scattering

POM Polarized optical microscopy

MD Molecular dynamics

RH Relative humidity

wt Weight

Dr Rotary diffusion coefficient

G’ Storage modulus

fc Orientation index

E Young’s modulus

σ , σu Ultimate tensile strength

ε Strain to failure

Nc Crowding number

CONTENTS

1. Introduction…………………………………………………….... 1

1.1 Biological Materials……………………………………………1

1.1.1 Nanocellulose from Wood……………………….………1

1.1.2 Spider Silk……………………………………………….4

1.1.3 Protein Nanofibrils………………………………………8

1.2 Materials Fabrication Strategies…………………………….... 9

1.2.1 Microchannels In-Situ Structural Dynamics……………13

2. Scope and Objectives……………………………………………14

3. Materials and Methods………………………………………….15

4. Nanostructured CNF Fibers…………………………………… 26

4.1 Assembly and In-Situ Orientation Dynamics of CNFs………26

4.2 Mechanical Properties……………………………………….31

4.3 Deformation Mechanism at the Nanoscale…………………. 34

4.4 Mechanical Properties in a Wider Perspective………………36

5. Fundamentals of the Hydrodynamic Assembly Process….…...39

5.1 Double Flow-Focusing Geometry………………………...…39

5.2 Shear vs Elongation…………………………………...……. 40

5.3 Rotary Diffusion and Size Dependency………………….… 41

5.4 Network Formation at the Nanoscale…………………….… 46

6. Nanostructured CNF/Silk Composites…………………………48

6.1 Fabrication of CNF/Silk Nanocomposites…………………...48

6.2 Mechanical Properties of Anisotropic Fibers…………….… 51

6.3 Bio-Functionalities…………………………………………. 54

7. Nanostructured Protein Microfibers…………………………...59

7.1 Assembly and In-Situ Orientation Dynamics of PNFs………60

7.2 Micro- and Molecular Structure in the Gel State…………….65

7.3 Hierarchical Protein Materials and Comparison with Other

Building Blocks……………………………………………...65

8. Conclusions……………………………………………………... 67

9. Future Work and Outlook………………………………….......70

Acknowledgements…………………………………………………. 73

References…………………………………………………………... 74

Appendix……………………………………………………………. 86

INTRODUCTION

1

1. Introduction

Synthesis of novel materials from biomass has substantial environmental and

economic impacts, particularly in present time of global warming, growing

population and depleting petroleum oil reserves. It is, therefore, the need to discover

material sources that can offer a sustainable path forward.1 In this direction,

industries are making attempts to develop “greener” materials; governments are

encouraging research on bioproducts; scientists in academia are searching for eco-

friendly materials. Till a certain extent, bio-based polymeric materials (also

represented here as biopolymeric materials or bio-based materials) from renewable

resources can contribute positively towards solving this complex problem.2

Natural (or biological) materials possess many inspiring properties such as

complexity, miniaturization, hierarchical organization, hybridization, and

malleability.3 These materials are typically formed through the assembly of

repetitive molecular to nano- and micron-sized building blocks into hierarchical

structures, which are stabilized by non-covalent interactions. Few examples of

biological materials with multiscale organization are silk, cellulose, chitin and

collagen.4-6 The highly ordered arrangement of the building blocks in these

biological materials such as CNFs in plants, chitin nanofibrils in animals, collagen

nanofibrils in ligaments and silk fibroins in spider silk and in silkworms7 provides

them with remarkable mechanical properties and flexibility, while also presenting

biological functions that interact with the surrounding environment.8 These features

make biopolymeric building blocks promising candidates for the development of

structures that are strong, lightweight and biodegradable for automobile, energy,

optics and biomedical applications. Lately, researchers have been seeking ideas of

designing novel materials by mimicking the architecture of natural materials which

is challenging from an engineering perspective. However, rational design

approaches, combining theory and the experimental tools, have aided in designing

bio-based materials through learnings from nature. In this direction, recent initiatives

by forest companies to invest in nanocellulose pilot plants, and by startups to upscale

the production of spider/silkworm silk to make threads for prototype garments are

encouraging for commercialization of bio-based materials.9

INTRODUCTION

2

1.1 Biological Materials

1.1.1 Nanocellulose from wood

Wood is a hierarchical architecture of nature that performs the functions of

providing mechanical support and water and nutrients transport through the tree.

Ibiot has a porous structure where the constituents are arranged at different length

scales to provide the mechanical integrity (Figure 1.1). Cellular architecture of wood

contains porous lumen and relatively dense cell wall.10 The cell wall of wood is made

up of a primary layer (P) and three secondary layers (S1, S2 and S3). These cell walls

are attached to each other by middle lamella (ML). Each layer in the cell wall

consists of semi-crystalline cellulose fibers (~40 wt.%) embedded in a matrix of

amorphous hemicelluloses and lignin (~30 wt.% each). The volume fraction and

orientation of cellulose fibers varies in each layer. For instance, S1 layer has random

orientation of cellulose fibers, whereas, the fibers are highly oriented in S2 layer.

Cellulose fibers can also be obtained from plant fibers, bacteria or tunicates.11 In

cellulose fibers, there is a base fundamental reinforcement unit that strengthen the

cellulose-based structures, known as “nanocellulose”. By extracting the

nanocellulose from the cellulose fibers, the defects associated with the hierarchical

structure can be removed.11

Figure 1.1. Hierarchical structure of wood with components at different length

scales. (Reproduced from Ref.12, Copyright 2013 Elsevier)

INTRODUCTION

3

Nanocelluloses are high aspect ratio nanoparticles (Figure 1.2) that can be

classified broadly in two categories (1) CNCs, also known as whiskers, which are

typically 150-250 nm long, rigid and rod-like with high crystallinity and (2) CNFs,

also known as NFC, MFC nanofibrillated cellulose and microfibrillated cellulose,

respectively, which are relatively longer (up to few microns) and more flexible.11-12

The exact dimension of nanocellulose particles depends on the source as well as on

the chemical pretreatment method. Nanocellulose from wood is typically obtained

from a pulp, which has long cellulose fibers with high cellulose content relative to

native wood. This is because the matrix components are partially separated during

the pulping process. The earliest reports on nanocellulose preparation date back to

early 1950s, where sulfuric acid hydrolysis of cellulose fibers was done to obtain

CNCs.13 Acid hydrolysis dissolves the disordered regions of cellulose fibers leaving

the highly ordered domains intact, as they are resistant to degradation from acids.

Sulfuric acid treatment induces surface charges to CNCs, which provides them the

colloidal stability. CNCs from wood can have width of 3-15 nm.11, 14 Extraction and

characteristic features of CNCs prepared by different methods have been reported

elsewhere.11, 15-16

Figure 1.2. TEM images of nanocelluloses.

CNFs are substantially longer than CNCs (Figure 1.2), and contain highly

ordered as well as less ordered regions (kinks).17 CNFs are typically obtained after

the mechanical disintegration of pulp through a high pressure homogenizer or by

sonication after a mild pre-treatment (enzymatic or chemical). Attempts to obtain

CNFs from pulp fibers were initiated in 1980s, where no pre-treatment on the pulp

fibers was done. The obtained CNFs had a broad size distribution high mechanical

energy was required to liberate the fibrils from the pulp fibers.18-19 Thereafter several

INTRODUCTION

4

pretreatments methods to assist the disintegration process were developed. Chemical

pre-treatment typically induces the charges on the surface of pulp fibers which

facilitates disintegration due to electrostatic repulsion. One widely used chemical

pre-treatment method is TEMPO (2,2,6,6-Tetramethylpiperidin-1-yl)oxyl mediated

oxidation that converts primary hydroxyl groups to carboxylates (negative

charges).20 Other methods include carboxymethylation,21 and cationic

functionalization.22-23 CNFs with surface charge density of up to 1500 µequiv g-1

have been reported with TEMPO mediated oxidation method.20 Due to the relatively

good colloidal stability and ease of processing, TEMPO-oxidized and

carboxymethylated CNFs are used in this thesis.

Molar mass of cellulose is reduced significantly during pretreatment and after

disintegration.20 CNFs of the order of 0.4 µm in length and 2-5 nm in width, may be

obtained when oxidized the pulp to a high surface charge.24 Crystallinity of wood

CNFs is >30%, depending on the type of wood and chemical pre-treatment method.

Availability of abundance of hydroxyl groups provides a functional surface that

plays a key role in determining the properties of CNF based nanocomposites. CNFs

can also form physically entangled network with strong interactions due to the

hydrogen bonds and Van der Waals forces.25

1.1.2 Spider Silk

Spider silk has impressive mechanical properties (strength and elasticity)

despite being spun at ambient temperatures and pressures with water as solvent

(Figure 1.3). The spider makes this fascinating material by processing the main

protein constituents along with auxiliary compounds and by controlling the

folding and crystallization simultaneously, to create a hierarchical structure with

ordered and repetitive arrangement of the building blocks.26 Spider does the

spinning of its fibers using minimal forces as the ‘dope’ (the constituent material

from which silk is spun) is liquid crystalline. Spiders can produce silks with

different characteristics. In materials perspective, spider dragline silk that comes

from the major ampullate gland of spiders, is one of the toughest bio-based

material with strength comparable to synthetic Kevlar and higher than that of

tendons.27 Apart from the remarkable mechanical properties, spider silk has the

ability to control bleeding and promote wound healing.28

INTRODUCTION

5

Figure 1.3. A schematic of the silk production system of a spider. The secondary

structure adopted by the spider proteins (spidroins) along the gland and duct with the

changes in pH, salt concentrations and water contents. (Reproduced from Ref.28,

Copyright 2011 Springer Publishing Group)

Dragline silk is mainly composed of two types of proteins (also refer as

spidroins), major ampullate spidroin (MaSp1 and MaSp2). Architecture of MaSp is

composed of alanine that provides the strength to fibers by the formation of β-sheets

and the Glycine-rich unit, provides elasticity with the presence of amorphous

regions. These repetitive sequences are covered by non-repetitive N-terminal and C-

terminal domains.29-30 The non-repetitive domains assist with the storage and

assembly of spidroins that are stored in the spiders’ glands as liquid crystalline dope.

The conformation of proteins inside the dope is primarily α-helical. Transition of α-

helix to β-sheets occurs inside the spinning duct, where the dope undergoes

conformational changes, in the presence of ions, change in pH, water removal and

the shear forces.31

It is difficult to farm spider because of their territorial and cannibalistic

behavior.32 From an industrial perspective, the amount of silk that can be harvested

directly from the spiders is not sufficient. Recombinant production of spider silk

proteins is considered as an alternate route to overcome this challenge.33 Tobacco,

potatoes, the yeast and E.coli bacteria have been utilized as hosts for synthetic

production of RSPs by the use of genetic engineering.32 However, E.coli is most

widely used host due to ease of genetic manipulation, short reproduction times, and

low-production cost.

INTRODUCTION

6

Figure 1.4. Schematic of IgG-fluorophore binding to Z-silk.34 (Image credit: Ronnie

Jansson)

Functionalization of silk is a valuable route forward to find the applications in

the fields of biotechnology and medicine.35 Several strategies have been applied to

functionalize RSPs to produce bioactive silk-based materials.36-37 The methods of

silk functionalization rely either on non-covalent or covalent interactions (Figure

1.4). Non-covalent functionalization techniques count typically on ionic interactions.

However, physical entrapment of functionalizing molecule is also possible. In

contrast, strong covalent chemical coupling between the material and the

functionalizing molecule is the basis of covalent functionalization. Bonds are stable

over time in the latter case.

INTRODUCTION

7

Figure 1.5. Spider silk proteins fibers. A, B, C are optical microscope images and D

is a SEM image. Scale bars are 2 mm in A and B, 100 µm in C and 10 µm in D.

(Reproduced from ref.28, Copyright 2011 Springer Publishing Group)

Recombinant spider silk scaffolds (Figure 1.5) functionalized with cell-

binding motifs have shown positive effects on cell characteristic features such as cell

viability, proliferation and viability.37-38 Cell binding motifs that are present in the

extracellular matrix (ECM) proteins (e.g., laminin, fibronectin, collagen,

vitronectin), mediate interactions between cells and ECM through recognition by

cell-surface integrins.39-40 To mimic the ECM environment, bioactive molecules like

growth factors and cell-binding peptides have been combined with silk building

blocks to create materials that have functional cell scaffolds with cell interacting

properties. Another example of silk for tissue engineering is for the cornea tissue

formation by the use of optically transparent silk films.41 Biodegradability and the

adjustment of drug release profile also make silk building blocks, a suitable carrier

for drug delivery.42 However, to facilitate wider use of the bio-functionalities

observed with the silk-based materials, there is a strong incentive to reduce the high

cost of materials based on RSPs. One potential route to achieve this is to combine

the silk proteins with a compatible building block that can be produced in larger

volumes and is available at lower cost.

INTRODUCTION

8

1.1.3 Protein Nanofibrils

Proteinaceous materials with the core structure of amyloids have recently

gained much attention and efforts are now turning more towards artificial self-

assembly of such structures for diverse set of applications. Utilizing the rich

chemistry and unique properties of proteins is however, not straightforward and thus

the generation of tailored protein materials is challenging due to understanding and

controlling protein self-assembly at multiple length-scales.43 A fruitful research

outcome to solve this challenge is the fibrillar self-assembly of proteins.44 Such near

one-dimensional protein nanostructures that are mostly homomolecular have been

used as load-bearing components in functional coatings, in catalytic scaffolds and

epigenetic information storage.45-47 Hierarchical artificial proteinaceous materials

for a variety of biological functions have also been developed.48-49

Self-assembly of globular proteins under denaturing conditions leads to the

formation of nanofibrils with several nanometers of thickness and micrometers of

length. Dense network of hydrogen bonds is the characteristic core of nanofibrils and

that provides a very high Young’s modulus to these nanofibrillar structures due to

strong intermolecular backbone-backbone interactions between polypeptide chains

within the nanofibrils.50 These PNFs share certain structural characteristics similar

with the pathogenic amyloid aggregations formed in vivo, hence they often called as

“amyloid-like”.51 The high aspect ratio and the surface charges help these nanofibrils

to form entangled networks or gels, which assists with the fabrication of hierarchical

structures.

Figure 1.6. AFM images of PNFs. (Image credit: Christofer Lendel)

INTRODUCTION

9

Proteins obtained from milk, soybean and eggs can self-assemble into

amyloid-like fibrils upon incubation at low pH far from protein’s isoelectric point

and high temperature conditions.52 An often used model protein for self-assembly

into PNFs is Bovine β-lactoglobulin because of its well-characterized structure, and

easiness of availability.51 Morphologies (Figure 1.6) and mechanical properties of

PNFs can be tuned in the presence of salts and also by varying the protein monomer

concentrations and ionic conditions.52 Thus, amyloid-like fibrils assembly is a

general form of nanoscale ordering of polypeptide molecules with impressive

mechanical properties that has a huge potential in the creation of future

proteinaceous functional materials.

1.2 Materials Fabrication Strategies

Multidimensional manufacturing strategies can be applied to mimic the

structure and properties of natural materials by assembling the building blocks. In

this direction, artificial spinning techniques are applied to regulate the orientation of

nanofibrils in 1D, where the in planes organization of fibrils can be achieved through

chemical modification,53 by applying magnetic54 or electric-fields55 and also through

spray-assisted layered assembly.56 In addition, biopolymer nanofibrils can also be

combined with inorganic building blocks such as nanoclays, graphene sheets etc. to

fabricate 2D materials (nanopapers and membranes) to achieve advanced optical and

electrical properties with superior biodegradability.57-58 These materials show

synergistic properties owing to the controlled arrangement of individual

components. The fabrication of 3D materials remains challenging, however,

advances in 3D printing technologies managed to produce biopolymer nanofibril

materials with ordered arrangements.59-61

Wet Spinning: Wet spinning is a widely used technique to control the

organization of nanofibrils in macro-sized fibers. In this process, a liquid jet

(typically a suspension) is extruded through a confined-space or a channel into a

coagulation bath, where the jet solidifies into a fiber after precipitation. The method

of wet spinning has been applied to suspensions or gels of biopolymer nanofibrils,

including nanocellulose,62 chitin63 and silk,64 and spun into oriented fibers with

diameters ranging from micrometers to millimeters and lengths up to several meters.

The alignment of the nanofibrils in the fibers is controlled by the flow conditions

and assembly is induced through the solvent exchange during or after extrusion

through the ducts, where the duct dimensions also play an important role on the

organization of the nanofibrils.62 The interactions between the nanofibrils inside the

duct can be strong as well as weak, depending on the network formed by the

nanofibrils based on the number of contact points, which in typically dependent on

nanofibrils length and concentration in the suspensions or gels.65-66

INTRODUCTION

10

Dry Spinning: Dry spinning technique has also been applied to silk fibroins to

mimic the properties of natural silk by spiders, where the silk dope is converted to a

fiber upon the evaporation of the volatile solvent.67 This technique has been used for

CNFs as well, to generate fibers with oriented nanofibrils.68 Post-processing steps

are typically not required in the dry spinning techniques. One major advantage of

dry spinning technique is that it can generate polymorphic arrangement of

nanofibrils within a single fiber, which is typically not the case with the wet spinning

techniques, where fibrils are oriented along the fiber axis.67 However, dry spinning

requires much higher dope concentrations to spun the fibers to provide a strong

enough network of nanofibrils to avoid the collapse.4

Wet and dry spinning are the efficient techniques to fabricate fibers with aligned

nanofibrils with improved mechanical properties. However, the mechanical

properties of these fibers are much inferior compared to the natural fibers with the

same building blocks. The highest obtained mechanical properties from the dry-spun

CNF fibers is 12.6 GPa of Young’s modulus and 222 MPa of tensile strength.62, 68 A

specific elastic modulus and strength of CNF-based fibers fabricated by wet spinning

were ~22 GPa and ~206 MPa,4 respectively, which are much less than the theoretical

and experimentally determined values characteristic of the individual CNF. Natural

wood pulp fiber with high orientation of nanoscale crystalline planes, free from

damage and natural imperfections, has a highest reported strength of ∼1600 MPa

(stiffness unknown).24, 69 Similarly, the fibers fabricated with the wet spinning of silk

fibroins or proteins are much inferior compared to the mechanical properties of

natural spider silk.70-71 The big gap in mechanical properties4 is due to the lack of

suitable processes and devices that can tailor the nanofibrils organization and thus

improve the mechanical performance.

Microfluidic Spinning: Microfluidic spinning is an advanced way to perform the

wet spinning. It is a promising technique to control the assembly of

nanoparticles/nanofibrils for the fabrication of load-bearing structures such as fibers

with improved mechanical properties (Figure 1.7). An intrinsic advantage of using

microfluidic tools is the ability to control the motion of nanoparticles in suspensions

by optimizing the processing parameters under the influence of flow-fields (shear or

elongation). These tools also provide opportunities to in-situ control the

physicochemical conditions such as pH, salt gradients etc. to induce controlled self-

assembly.4 Although microfluidic tools have demonstrated the ability to produce

fibers with enhanced mechanical properties,24, 34, 66 challenges associated to the

perfect nanostructuring of nanofibrils to realize the full potential of nanoscale

building blocks in macrostructures and also related of the scale-up of production still

need to be addressed.

INTRODUCTION

11

Figure 1.7. Different microfluidic platforms used to fabricate fibers.72 (A) A triple-

orifice spinneret with five inlets connecting to a single outlet . (B) Pulled glass

capillaries based double-coaxial microfluidic device. (C) Glass capillaries integrated

in Polydimethysiloxane (PDMS). (D) A PDMS device fabricated using soft

lithography having rectangular channels. (E) A PDMS device with cylindrical

channels. ECM- extracellular matrix, CAC- collagen-alginate composite, WSC-

water-soluble chitosan, NaA- sodium alginate, PGA- propylene glycol alginate.

(Reproduced from Ref.72 with permission from the Royal Society of Chemistry).

INTRODUCTION

12

Figure 1.8. (A) Schematic of SAXS setup and microfluidic channel. Microfluidic

channel is placed between the source of X-ray beam and the detector. Scattered

photons are recorded by the detector in the form of 2D scattering patterns.

(Reproduced from Ref.73 with permission from the Royal Society of Chemistry) (B)

Quantification of in-situ orientation of nanofibrils using SAXS. (Reproduced from

Ref.74, Copyright 2014 Nature Publishing Group)

INTRODUCTION

13

1.2.1 Microchannels In-Situ Structural Dynamics

Obtaining in-situ structural information of anisotropic nanoparticles under

dynamic flow conditions is challenging. The macroscopic behavior of hierarchical

materials fabricated from such nanoparticles via hydrodynamic approaches, largely

depends on the nanoscale network and interparticle contacts in the microchannels.

These features often associate with changes in the structure of individual phases due

to the flow-fields, temperature, confinement, concentration and channel geometry.

Quantification of in-situ structural dynamics is not an easy task as characterization

techniques which allow spatial and time-resolved measurements are limited. This

largely limits our fundamental understandings on the motion and the behavior of

anisotropic macromolecules.75 Therefore, in-depth characterization of nanoparticle

dynamics and a thorough understanding of the nanoscale physics will pave the way

for flow-based fabrication processes.

Existing characterization techniques such as rheology measurements, and

electron microscopy,76 are predominantly limited to observing static systems (

solid/liquid systems with fixed volume). Other technique such as dynamic light

scattering is hard to couple with microfluidic channels and have problems with the

light absorption by the dispersants and is unable to characterize the polydisperse

samples, which is typically the case with building blocks of biological origin.14

Continuous monitoring of structural changes at the nanoscale in out of equilibrium

microfluidic systems is feasible through SAXS (Figure 1.8). Numerous anisotropic

nanoparticles such as amyloids, CNFs in microfluidic channels have recently been

investigated using SAXS.73-74 While SAXS sources offer high brilliance (1012–14

photons s-1 at the sample) and small beam sizes (100 µm to sub-µm),77 one big

problem is the availability of high intensity X-rays that is primarily limited to

synchrotron sources. New lab-scale methods and tools to characterize nanoparticle

dynamics under flow will thus contribute critically to the development of knowledge

and technology for controlled assembly of colloidal systems.

OBJECTIVES

14

2. Scope and Objectives

Biopolymeric building blocks constitute the basic architectural element of many

biological (or natural) materials that fulfil structural functions. Tightly controlled

and well-ordered arrangements of these building blocks at multiple length scales -

from molecular to nano- and micron level, provides exceptional properties to these

biological materials. However, the complexity of artificially designing new

hierarchical materials by realizing the potential of these natural building blocks at

length scales most suitable to human technologies is a big challenge. The basic

knowledge required to meet this challenge is incomplete in many dimensions, in

particular at the nanoscale. This emphasizes the need for the development of novel

processing strategies and the fundamental understanding thereof, to address multiple

physical and mechanical factors that impose tremendous constraints on designing

nanostructured materials with improved performance.

The scientific objective of this thesis is primarily related to fabrication of

nanostructured biopolymeric materials (mainly fibers) with high mechanical

properties, fundamental understanding of process parameters that affect the materials

characteristics and establishing structure-property relationships in the engineered

materials. By utilizing microfluidic tools, the goal is to improve the understanding

of organization and assembly of nanoscale biopolymeric building blocks in

hierarchical materials. In addition, the possibilities to combine basic knowledge from

different scientific fields such as fluid physics, colloidal chemistry and genetic

engineering to design new high-performance structural and functional biopolymeric

materials are explored. The cases investigated in the present thesis provide

encouraging insights with respect to the general potential of biopolymers.

EXPERIMENTAL SECTION

15

3. Materials and Methods

Detailed information on the chemicals, materials and methods are available in

the appended papers (I to V) and associated supporting information. This chapter

briefly summarizes material preparation and characterization techniques.

Preparation of CNFs

TEMPO-oxidized CNFs were used for the studies reported in paper I to III and

carboxymethylated CNF (commercially available) for the study in paper IV (Figure

3.1). For the CNF in paper IV, 1 wt.% gel was kindly provided by Innventia AB

(now RISE Bioeconomy), Stockholm, Sweden. The CNF was derived after the

defibrillation of a dissolving grade pulp that had been carboxymethylated to a charge

density of 600 ± 50 μequiv g−1, according to a procedure reported before.21

For the studies reported in papers I to III, chemically bleached wood fibers (a mixture

of 60% Norwegian spruce and 40% Scots pine, provided by Domsjo Fabriker AB,

Sweden) were used to prepare CNF suspensions. The relative glucose content of the

pulp was >96%. Pulp fibers were chemically treated with a 2,2,6,6-

tetramethylpiperidinyl-1-oxyl (TEMPO)-mediated oxidation reactions as reported in

detail in paper I. Oxidized pulp fibers were thoroughly washed with DI water.

Suspensions of the chemically pretreated pulp fibers were then passed through a

high-pressure homogenizer. Gel-like CNF suspensions, with a concentration of >5

g L−1 were obtained after homogenization. suspensions were diluted by adding DI

water followed by mixing and sonication. Unfibrillated precipitates were removed

by centrifugation at 5000 rpm for 60 min. The supernatant fluids were then used for

the studies. Dry content was determined using gravimetric analysis.


16

Figure 3.1. TEM images of TEMPO-oxidized CNFs of different surface charge

densities in µequiv g-1 (CNF-380, 550, 820, 980 and 1360) and carboxymethylated

CNFs obtained from Innventia (now RISE Bioeconomy). Scale bars are 500 nm.

Preparation of PNFs

Whey protein isolate (WPI) was obtained from Arla Food Ingredients. A ∼100 g/L

concentration of WPI was dissolved in hydrochloric acid (HCl) at pH 1 and the

mixture was dialyzed with HCl at pH 2 for 24 h at 23°C temperature using a 6-8 kDa

molecular weight membrane. The solution was then diluted to different


17

concentrations to obtain nanofibrils of different morphologies (40 g/L for straight

fibrils and 75–80 g/L for the curvy ones). PNFs were obtained after the incubation

of protein monomers at 90 °C for a period of 3 days. PNF suspensions were then

dialyzed using a 100-kDa molecular weight membrane for 7 days to remove the

unfibrillated proteins. The final concentrations were measured by UV-Vis and

gravimetric analysis of lyophilized materials.

Characterization of cellulose- and protein nanofibrils

Cellulose nanofibrils (CNFs): Length and height of CNFs used in paper I to IV are

measured using TEM and AFM (Figures 3.2 and 3.3). The lengths of TEMPO-

oxidized CNF (paper I to III) and carboxymethylated CNF (paper IV) (∼250 for each

charge density) are measured using TEM (JEOL JEM-1400 TEM) at an accelerating

voltage of 120 kV. The images were acquired using a ruby camera following

“systematic, uniform, random” rule to avoid bias. Droplets of CNF suspensions were

deposited on carbon-coated copper grids that were stained with 2% uranyl acetate

solution. The height of CNFs (∼150 for each charge density) were measured using

AFM (MultiMode 8, Bruker, Santa Barbara, CA, USA). Silicon wafers were

oxidized to form silica layer. The wafers were then washed with ethanol and milli-

Q followed by drying and plasma treatment to make the surface hydrophilic. Wafers

were then dipped in the CNF suspensions and air dried at room temperature.

A protocol based on conductometric titration78 was used to determine the carboxylate

contents of TEMPO-oxidized CNFs (paper I to III). 100 ml suspensions of 0.1 wt.%

CNFs were used, and the pH of solution was adjusted to 2.5 with HCl. 0.01 M

standardized NaOH was used to titrate the suspensions and pH was adjusted to 11,

and the conductivity was monitored with a benchtop meter (FE20 FiveEasy, Mettler-

Toledo).


18

Figure 3.2. Length distribution of TEMPO-oxidized CNFs of different charge

densities.

Figure 3.3. AFM image of TEMPO-oxidized CNFs and width distribution of

different surface charged nanofibrils measured with AFM.


19

Protein nanofibrils (PNFs): Morphology of PNFs were investigated using AFM

(Dimension FastScan Bruker) in tapping mode. Samples were diluted in HCl at pH

2 and droplets of PNF suspensions were placed on freshly cleaved mica surfaces

followed by the air drying. Experiments were also performed in liquid state using

ScanAsyst Liquid+ cantilevers (Bruker) operating in Peak Force mode. Both the

measurements from the dry and the liquid states showed similar differences between

the morphologies of straight and curved PNFs. For further analysis on the fibril’s

morphology, end-to-end lengths and contour lengths were calculated from AFM

images by using Image J software. For the persistence length (see SI or paper V),

~300 fibrils for each morphology were measured. Persistence lengths and fibril

heights measured with AFM were used to get a rough estimation of the relative

elastic moduli of the PNFs using a method reported elsewhere.79-80

Preparation of composite CNF/silk films and fibers

CNF films (paper IV) were casted from a CNF suspension onto the Teflon sheets,

and were air-dried at room temperature. CNF/silk composite films (CNF/Z-silk and

CNF/FN-silk) were also prepared in the same way. Z- and FN-silks were produced

in Escherichia-coli and purified essentially according to a procedure described

before by Hedhammar et al.81-83 CNF and CNF/silk fibers were prepared using a

double flow-focusing geometry. The hydrogel fibers were collected in a water bath,

winded up on a steel frame followed by air drying at room temperature.

Preparation of composite CNF/Z films and coatings

CNF/Z films were casted onto Teflon sheets and air-dried at room temperature in the

same manner as mentioned previously for CNF and CNF/silk. Films of CNF and

CNF/Z-silk were also prepared as controls. The amount of Z molecules in the CNF/Z

films was equimolar to the amount of Z-silk in the CNF/Z-silk films. For the

preparation of CNF/Z coatings, the casted CNF films were immersed (dip-coated) in

a Z protein solution (0.063 mg mL-1) for 15 min at room temperature. Following

removal of the unbound Z solution by thoroughly washing the films using Tris

buffer, films were incubated with 20 mM Tris buffer of pH 8 for 15 min at room

temperature. Controls of CNF/Z-silk and CNF coatings were prepared the same way.

Binding of IgG to CNF/Z-silk

Dried films and fibers of CNF/Z-silk were placed in phosphate-buffered saline of

PBS at pH 7.4, followed by the incubation with rabbit anti-mouse IgG-Alexa Fluor

488 conjugate (0.05 mg mL-1, Thermo Fisher Scientific) containing 0.5% (v/v)


20

Tween 20 for 30 minutes at room temperature. Composite films and fibers were

washed three times with a mixture of PBS and 0.5% Tween 20 buffer. Fluorescence

microscopy was used to check the binding of IgG to the films and fibers. Composites

of CNF and CNF/FN-silk were treated in the same way that were used as controls.

Binding of IgG to CNF/Z

Binding of rabbit anti-mouse IgG-Alexa Fluor 488 conjugate to as prepared CNF/Z

films and coatings was performed the similar way, as previously mentioned for

CNF/Z-silk. For controls, films and coatings of CNF/Z-silk and CNF were used.

Culture of Fibroblasts and Endothelial Cells

Primary human dermal fibroblasts (HDF) were cultured in Dulbecco's Modified

Eagle Medium/Nutrient Mixture F-12 together with fetal bovine serum (5%) and

penicillin−streptomycin (1%). Cells were expanded for a period of 7 days and

viability was checked with trypan blue. Experiments were carried out in passages of

8 to 13. For endothelial cells (EC) culture, the human dermal microvascular

endothelial cells (HDMECs) isolated from dermis were used. Gelatin-coated flasks

were used to expand and culture EC.

Live/Dead Staining

Living and dead cells were visualized with live/dead viability assay based on two-

color fluorescence for real-time identification of living and dead cells with two

probes (intracellular esterase activity and plasma membrane integrity). Enzymatic

conversion is used to determine the intracellular esterase activity in living cells.

Ethidium homodimer-1 enters the damaged membranes of cells and bind to nucleic

acids to undergoes a significant enhancement of fluorescence, resulting in a bright

red fluorescence in the nuclei of dead cells. Phosphate-buffered saline (PBS) was

used to wash the cell matrices.

Filamentous Actin Staining

Staining of filamentous actin in the cells was done using Phalloidin. Before staining,

cell matrices were washed with PBS together with paraformaldehyde (4%), and

permeabilized with Triton X-100 (0.1%). 4′,6-diamidino-2-phenylindole (DAPI)

was used for nuclear staining.


21

Enumeration of Cells Bound to Films

HDMECs were seeded onto CNF films with or without FN-silk proteins and were

cultured for 2, 4, and 6 after washing with PBS buffer. Nuclei and F-actin were

stained, and images were taken using an inverted Nikon Eclipse Ti fluorescence

microscope. Stained nuclei numbers were counted manually in each image and

expressed in terms of number of cells per area (A = 1 mm2). Statistical evaluation

was done using Student’s t-test.

Small Angle X-ray Scattering (SAXS)

SAXS was used to quantify orientation of CNFs and PNFs at different downstream

positions in the channel. Anisotropy of the orientation distribution of nanofibrils was

measured with the anisotropy of 2D scattering patterns. P03 beamline84 at PETRA

III storage ring (DESY, Hamburg) was used to perform the SAXS measurements.

The wavelength (𝜆 ) of the X-ray beam was 0.96 Å with a distance of 6950 mm

between the sample and the detector. The size of the X-ray beam was 20 × 10 µm2

(horizontal × vertical) and to monitor the scattering diffractograms, a single-photon

counting detector (Pilatus 300k by Dectris, Switzerland) with a pixel size of 172 ×

172 µm2 was used.

We used a slightly modified channel geometry for the in-situ fibrils orientation

measurements compared to the one that was used for the fibers fabrication. The

second focusing step from the double flow-focusing geometry was mimicked by a

contraction step, and the width of channel after the contraction was set to 0.5 mm.

Due to the limited access to the synchrotron and high exposure times, experiments

were carried out under dynamic flow conditions without inducing the transition to

gel.

Order parameters (S) were quantified based on a similar procedure as described

before.84-86 Scattering patterns with the flow of only DI water were used as

background. An assumption on zero-level intensity is made by considering that no

fibrils is aligned in the perpendicular direction to the flow at the position in the

channel, where highest intensity (orientation) is observed. The intensity at azimuthal

angle (𝜃) = 0 or π is therefore subtracted to all the distributions. The final intensity

distributions were averaged between q range of 0.1 to 0.5 nm−1, for each azimuthal

angle.

S is defined as:


22

𝑆 =<3

2𝑐𝑜𝑠2𝜃 −

1

2> (1)

Expanding the average gives:

𝑆 = ∫ 𝐼(𝜃) (3

2𝑐𝑜𝑠2𝜃 −

1

2) 𝑠𝑖𝑛𝜃 𝑑𝜃

𝜋

0 (2)

which is normalized according to:

∫ 𝐼(𝜃)𝑠𝑖𝑛𝜃 𝑑𝜃 = 1𝜋

0 (3)

where 𝐼(𝜃) is the intensity distribution averaged along examined q-value for each

azimuthal angle.

Fiber Characterization

Wide Angle X-ray Scattering (WAXS): WAXS measurements were performed at PO3

beamline at PETRA III storage ring at DESY (Hambury, Germany). The wavelength

(𝜆 ) of the X-ray beam was 0.96 Å with a distance of 71 mm between the sample and

the detector. The size of the X-ray beam was 6 × 14 µm2 (horizontal × vertical) and

to monitor the scattering diffractograms, a single-photon counting detector (Pilatus

300k by Dectris, Switzerland) with a pixel size of 172 × 172 µm2 was used. Order

parameter (S) and orientation index (fc) were calculated from the intensity

distribution profiles, according to the below equation:

𝑓𝑐 = (180° − 𝐹𝑤ℎ𝑚)/180° (4)

where, Fwhm is the full width at half-maximum of the azimuthal profiles.

Tensile test measurements: Tensile tests were carried out in the Solid Mechanics

Department at KTH on Instron E100 instrument equipped with a load cell of 5N

(Figure 3.4). Temperature and humidity inside the room were controlled based on

the requirements. The dimensions (test-section length and diameter) were measured

by an optical microscope at 20x magnification and few samples were also cross-

checked with SEM for the fiber diameters to validate the results from the optical

microscope and avoid errors. Diameters of individual CNF and PNF fibers were

uniform throughout the cross section; however, the diameters were slightly different

for the different fibers. Typical diameters of the CNF and PNF fibers were around

6.8 ± 0.9 μm and 40 ± 5 μm, respectively. End of the fibers were glued on the papers.

The span length was 9−12 mm, and measurements were performed at a crosshead

speed of 0.5 mm min−1. Cross-section of fibers were assumed to be circular.


23

Figure 3.4. Instron E100 instrument equipped with a 5N load cell used for the tensile

tests of CNF, CNF/silk and PNF fibers. Inset is the zoomed view of test section

where sample is mounted.

Scanning and Transmission Electron Microscopy (SEM & TEM): For the SEM in

papers I, III and IV, a 5 nm platinum−palladium layer was sputtered onto the fibers

surface. The imaging of CNF, CNF-silk and PNF fibers was performed by using a

Hitachi S-4800 Field Emission Scanning Electron Microscope operated at an

acceleration voltage of 1 kV. For information on the bulk structure of CNF fibers in

paper I, samples were polymerized overnight at 60 °C and embedded in epoxy-

embedding medium. Ultrathin sections of 80 nm were cut and placed on 75 mesh

Formvar coated copper grids. The imaging was performed using a JEOL 1400 plus

TEM at 120 kV, equipped with a Ruby camera.

Rheology

We used a parallel plate geometry (DHR-2 rheometer, TA Instruments) to

investigate the viscoelastic properties of CNF, CNF/silk and PNF hydrogels (papers

IV and V). A 25 mm diameter parallel plate setup with a truncation gap of 300 µm

(paper IV) and 500 μm (paper V), respectively was used. For the CNF and CNF/silk

experiments, suspension to gel transition was induced with HCl of pH 2 and for the


24

PNF experiments, we used pH 5.2 acetate buffer (Figure 3.5). Subsequently, 0.25%

strain was applied, and storage and loss moduli were measured for a period of 45

min or more (until a constant value is reached). Fracture stress was also measured at

a constant frequency of 1 Hz.

Figure 3.5. Schematic of the parallel-plate setup used for the rheology of CNF,

CNF/silk or PNF gels. Suspension was placed on the bottom plate of the rheometer

and gelling agent was added along the circumference to form gel.

Rotary Diffusion Measurements

The rotary diffusion constant (Dr), is highly dependent on the length distribution and

concentration of the fibrils, i.e. the physical crosslinking of the fibrils during their

rotation. Doi and Edwards87 developed a model to study the rotary diffusion of

monodisperse rods, which was extended to polydisperse systems by Marrucci and

Grizzuti88-89 (see papers II and III for details).

The diffusion coefficient of polydisperse suspensions in semi-dilute regime is given

by88-89:

𝐷𝑟(𝐿) =𝛽𝑘𝐵𝑇

ƞ𝐿4

1

∫ 𝑐(𝐿′)𝐿′𝑄(𝐿,𝐿′)𝑑𝐿′(∫ 𝑐𝐿

0(𝐿′)𝐿′4𝑄(𝐿,𝐿′)𝑑𝐿′+ 𝐿3 ∫ 𝑐(𝐿′)𝐿′𝑄(𝐿,𝐿′)𝑑𝐿′)′

∞𝐿

∞0

(5)

where β is a numerical factor, kB the Boltzmann constant, T the temperature, η the

viscosity of the solvent and �� the concentration distribution of the rods depending on

L. Q represents the anisotropic corrections depending on the orientation of rods

having a length L and 𝐿′ and is defined as:

𝑄(𝐿, 𝐿′) =4

𝜋 ∫ 𝑑𝑝 𝑑𝑝′𝛹(𝑝, 𝐿)𝛹(𝑝′, 𝐿′)|𝑝ᴧ𝑝′| (6)

where Ψ(𝑝, 𝐿) and Ψ(𝑝′, 𝐿′) being the orientational distributions of the rods of length

L and 𝐿′. Therefore, the diffusion coefficient of a rod of length L depends on its


25

interactions with other rods of different lengths, which may have different

orientation distributions. Hence the relaxation dynamics between the rods and

relative and coupled.

We assume that the anisotropic corrections are negligible, because the system is

always far from full alignment. This means that Q ≈ 1 and this leads to

𝐷𝑟(𝐿) =𝛽𝑘𝐵𝑇𝐿∗

4

ƞ𝐿7 Γ(L) (7)

with 𝐿∗ = (∫ ��(𝐿′)𝐿′𝑑𝐿′∞

0)-1/2

and

𝛤(𝐿) =∫ 𝑐(𝐿′)𝐿′𝑑𝐿′

∞0

∫ 𝑐𝐿

0(𝐿′)𝐿′(𝐿′/𝐿)3𝑑𝐿′+∫ 𝑐(𝐿′)𝐿′𝑑𝐿′)

∞𝐿

(8)

The factor Γ(L) depends on the length L and on the concentration distribution �� of

the nanorods of different lengths, related to the length distribution in case of

polydisperse nanoparticles suspensions (CNFs). It is the correction factor due to

interactions between nanorods of different lengths.

RESULTS & DISCUSSIONS

26

4. Nanostructured CNF Fibers

Excellent strength of nanoscale building blocks in materials is due to their defect-

free molecular structure. It is exceedingly challenging to devise nanostructuring

strategies for translating such high intrinsic properties to macroscopic materials.

Cellulose has theoretically been perceived as strongest and stiffest among the bio-

based materials in its nanoscale crystalline form. However, these superior properties

are yet to be realized for nanocellulose-based macroscopic structures. In this chapter,

our findings on the CNF fibers in pursuit of maximal mechanical performance, by

hydrodynamically controlling the structural ordering of nanofibrils, are reported. We

employ high intensity X-rays and POM to examine the behavior of CNFs in the

microchannel. The use of double flow-focusing channel results in fibers with highly-

ordered and strongly interacting arrangement of CNFs. The E of the fibers reaches

up to 86 GPa and a σu of 1.57 GPa, which are substantially higher than any natural

or synthetically prepared bio-based material. We also compare the mechanical

properties of our fibers with fossil-based synthetic fibers that are typically used.

4.1 Assembly and In-Situ Orientation Dynamics of CNFs

Figure 4.1. Schematic of double flow-focusing channel used for CNFs assembly.

The CNFs suspension is injected in the core flow (light brown color), DI water (blue

color) in the first sheath, and acid at low pH (light green color) in the second sheath

flows. Arrows demonstrate the flow direction. Hydrodynamic and electrostatic

interactions at different positions along the channel are illustrated schematically on


27

the right. Position 1, poor fibril alignment due to Brownian diffusion and

electrostatic repulsion (illustrated with the dashed arrows) caused by dissociated

carboxyl (−COOH) groups on the fibrils surface. Position 2 hydrodynamically

induced alignment (illustrated by solid, green arrow) occurs during

acceleration/extension. Position 3 further increase in alignment during

acceleration/extension, and in position 4, following the acid addition, Brownian

diffusion is minimized due to the transition of CNF suspension to an immobilized

volume filling arrested state due to protonation of the COO− groups. For illustration,

the relative size of the fibrils has been magnified around 300 times. The use of acid

for transforming the free-flowing CNF suspension to a fibrous colloidal glass

ensures that the electrostatic repulsions are replaced by van der Waals forces in the

protonated carboxyl groups. This complete removal of electrostatic repulsion is not

possible with simple electrolytes.90

Figure 4.2. (A) FE-SEM image of the fiber surface, where the dense fibrillar

network with well-preserved anisotropic arrangement can be seen. (B) FE-SEM

image of the cross-section of the fiber.


28

Figure 4.3. TEM images of a nanostructured CNF fiber highlighting a uniform

distribution of fibrils inside the bulk structure without any obvious voids or defects.

Scale bars are 100 nm. (A) and (B) correspond to different locations.

Fibers from TEMPO-oxidized CNFs are fabricated by the process of

hydrodynamic alignment and assembly. Fibrils in the suspension are aligned in the

microchannel before “locking” the nanostructure into metastable colloidal glass.91

This goal was accomplished using extensional flow-field-based double flow-

focusing channel (Figure 4.1). In the core flow, suspension of charged CNF fibrils

is used, where fibrils are free to rotate due to electrostatic repulsions and Brownian

motion, only restrained by fibril−fibril physical crosslinking (partial entanglements).

Note that the electrostatic repulsion caused by the dissociated COOH groups on the

surface of CNFs is much higher than the attractive van der Waals forces at neutral

to slightly alkaline pH. The first sheath flow of DI water supports electrostatic

repulsion and prevents transition into the colloidal glassy state in contact with

channel walls; it also aligns the fibrils toward the flow direction. Before the

alignment is diminished by the Brownian diffusion, the second flow of low pH acid

enhances the fibril alignment as observed with the in-situ SAXS experiments

discussed later in this chapter, while reducing the electrostatic repulsion between the

fibrils due to protonation of carboxyl (COO−) groups that allows the supramolecular

interactions between CNFs to self-organize the fibrils into a well-packed state with

maximized CNF−CNF contacts. The continuous threads obtained from the flow-

induced assembly were subsequently held at their ends and air-dried.

Characterization of the assembled structure with FE-SEM (Figure 4.2) sampled in

the longitudinal direction show uniform sized CNF fibers with dense and highly

oriented nanofibrils without obvious packing defects or voids (Figure 4.3).

Micrographs of fiber cross sections sampled in the transversal direction confirm the

dense fibrillar packing and reveal a well-defined layered structure (Figure 4.2B). In


29

this study (paper I), we used TEMPO-oxidized CNFs with three different surface

charge densities (550, 820 and 1360 µequiv. g-1) named as CNF-550, CNF-820 and

CNF-1360 with mean fibril lengths of 590, 683 and 391 nm, respectively (Figure

3.2).

Figure 4.4. In-situ study of the alignment, dealignment and rotary diffusion of

fibrils. (A) Microscopy image of the channel used for SAXS measurements, placed

between two cross-polarized filters rotated 45° from the vertical axis (white arrows).


30

White color corresponds to the birefringence signal obtained for the CNF-550

suspension. Numbers represent the positions where in-situ measurements were

carried out. Scale bar is 1 mm. (B) SAXS scattering diffractograms at different

positions along the channel for CNF-550 suspension (top row) and CNF-1360

suspension (bottom row). Curved lines represent the beam stopper. (C) Local order

parameters calculated from the SAXS scattering diffractograms as a function of

downstream position normalized with the channel width (h). (D) Birefringence

signal obtained in a single flow-focusing channel for the CNF-550 suspension. Scale

bar is 1 mm. The black squares along the center line represent the positions where

the Dr values are estimated. (E) Birefringence signal obtained at z/h = 3 as a function

of time when the flow is viciously stopped. The inset shows the initial decay of the

birefringence signal for the CNF-1360 suspension. The dashed line indicates the fit

of the initial decay to measure the Dr. Negative and positive time represents the

condition before and after stopping the flow, respectively. (F) Dr as a function of

downstream position along the center line. The vertical dashed line shows the

position z/h = 3 from where the birefringence decays are plotted in (E). This position

also corresponds to the filled black square at z/h = 3 in (D). The colors in (E) and (F)

correspond to different suspensions: CNF-550 (red) and CNF-1360 (blue).

(Reproduced from Ref.24. Copyright 2018 American Chemical Society)

In-situ monitoring of CNFs behavior in microchannel was essential for

successful adaptation of multiple environments. POM shows the orientation of fibrils

(Figure 4.4A), with gradually increasing birefringence from randomized suspensions

to an oriented state.74 In-situ alignment is quantified using synchrotron-based SAXS

(Figure 4.4B). Local order parameters of CNF suspensions for two different lengths

of the constituting fibrils (CNF-550 with 590 nm and CNF-1360 with 391 nm mean

lengths) are calculated (Figure 4.4C). An order parameter of 1 represents a fully

oriented state of the fibrils (in the direction of fiber preparation), and 0 corresponds

to an isotropic fibril distribution. Initially, the shear from channel walls gave rise to

some ordering (z/h = 1) in the CNF suspension. At the beginning of the focusing step

(z/h = 2), the order parameter decreases due to deceleration of the core flow followed

by a sudden increase after the focusing due to acceleration (3 ≤ z/h < 4).

Subsequently, the order parameter decreases slightly (4 ≤ z/ h < 7) before increasing

again during acceleration in the contraction (7 ≤ z/h < 10). The decrease in order

parameter after the focusing and contraction steps (4 ≤ z/h < 7; 10 ≤ z/h < 15)

indicates nanofibril relaxation toward isotropy, primarily due to Brownian

diffusion.74 Substantial differences between the orientation and disorganization

behavior of CNF-550 and CNF-1360 suspensions are observed: the shorter the fibrils

(CNF-1360), the faster the process of orientation and disorganization (Figure 4.4C).

This effect is due to the diffusivity based on length distributions of the nanofibrils as


31

given by Dr of CNF-1360, which is twice that to CNF-550 (Figure 4.4D-F) (Refer

papers I and III for details).

4.2 Mechanical Properties

Figure 4.5. Tensile mechanical properties and nanostructure characterization of the

CNF fibers. (A) Stress−strain curves for nanostructured fibers made from CNFs of

different using double flow-focusing channel. Experiments are conducted at 50%

RH. Effect of (B) physical (RH) and (C) chemical (cross-linking with BTCA)

approaches for tuning the tensile mechanical properties of fibers (prepared from

CNF-550 suspension). Plots comparing (D) E and σu of different fibers prepared in

this work. LHC and CL stand for low humidity condition and cross-linking,

respectively. Error bars correspond to the standard deviation obtained from 10

samples for each case. (E) Azimuthal integration of the (2 0 0) scattering plane of

the diffractograms for CNF-550 and CNF-1360 samples. Diffractogram

corresponding to a CNF-550 fiber is shown in the inset. (Reproduced from Ref.24.

Copyright 2018 American Chemical Society)

We establish a relationship between fibril characteristics (length, surface charge)

and tensile mechanical properties, that is important to provide the foundation for

future rational design of materials with targeted performance maxima. For a

thorough understanding on the structure-property relationship, we fabricated fibers

with CNF-550, CNF-820, and CNF-1360 and compared the mechanical properties.

The stress−strain curves in Figure 4.5A show an initial pseudo-elastic linear region,

followed by the deviation from linearity (plastic region). The “knee” in the tensile


32

test curves represent the elastic−plastic transition and is attributed to yielding

mechanisms related to relative sliding between the fibrils.92 CNF-550 and CNF-820

show similar stress−strain behavior with a E and σu of ∼70 GPa and ∼1200 MPa,

respectively (Figure 4.5A-D). This is probably due to the relatively small difference

in the mean length of the constituent nanofibrils and indicates that strength and

stiffness of the CNF fibers are independent of fibril surface charge. This is further

verified by preparing fibers with CNF-1360 with mean fibril length much lower than

CNF-550 and CNF-820, where we observed a significant decrease in σu (630 MPa).

Moreover, the relatively lower E of CNF-1360 (45 GPa) is also due to the low

nanofibrils orientation, as verified by WAXS (Figure 4.5E). The (2 0 0) reflection of

cellulose I, which generally is used to study the orientation of cellulose crystals, is

here used to track the orientation of CNFs, as crystals are aligned in the fibril

direction.93 The fibril orientation in the CNF fibers was calculated using the

following equation based on azimuthal breadth analysis:78

𝑓𝑐 = (180° − 𝐹𝑊𝐻𝑀)/180° (9)

Full width at half maximum (FWHM) is determined from the azimuthal profile of (2

0 0) plane (Figure 4.5E). The fc for CNF-550 fibers is 0.92 (order parameter = 0.70),

whereas CNF-1360 fibers show fc of 0.83 (order parameter = 0.53). In general, the

stiffest materials have tendency to be the strongest if provided with defect-free

structure and strong interfaces to ensure adequate load transfer and cohesion within

the material.5 The E and σu of our fibers vividly demonstrate the importance of

connectivity and bonding at the fibril−fibril interfaces facilitating multiscale stress

transfer. To further outspread the property range of these materials, we evaluated

different approaches based on physical (varying the ambient condition) and chemical

(covalent cross-linking of the fibrils) strategies. With CNF-550 fibers conditioned at

14% RH for a period of 40 h before testing, CNF-LHC exhibited a E and σu values

as high as 82 ± 4 GPa and 1320 ± 85 MPa, respectively (Figure 4.5D). As one might

partially expect from previous studies on bioinspired nanocomposites, strengthening

of interfibrillar interactions and removal of water molecules from the fibril surface

resulted in marked improvement of fiber mechanics with respect to other CNF

composites but also led to embrittlement and low failure strain. At high RH, water

acts as a plasticizer and allows extended elastic deformations, thus reducing the E

and increasing the ε.94 To reduce the humidity effect and potentially improve the

connectivity between the nanofibrils, the CNF-550 fibers were cross-linked by

1,2,3,4-butane tetracarboxylic acid (BTCA) that was used to neutralize the

suspensions during the flow-based assembly of the CNF fibers. BTCA creates

covalent bridges between CNF fibrils,95 replicating to some degree the cross-links

between cellulose and lignin/hemicellulose. The average strength of cross-linked

fiber (CNF-CL) tested at 50% RH increased to 1430 MPa (highest measured value


33

of 1570 MPa) with negligible change in modulus (Figure 4.5D). Chemical cross-

linking introduces covalent bonds between the fibrils, which improves the

interfibrillar connectivity and stress transfer.

To the best of our knowledge, even the mechanical properties of fiber fabricated

with CNF-1360 in this study, have previously been unachievable for CNF fibers

prepared with other approaches that rely mostly on the shear forces (Appendix Table

1). Hence, there is a strong and profound benefit of exploiting extensional flow-fields

for the orientation and assembly of CNFs (or elongated particles, in general), to bring

fresh insights in designing future high-performance load-bearing architectures by

getting closer to the theoretical limit of the building blocks. Further, the fc values for

highly oriented nanofibrils in CNF-550 fibers, reported in this study are only slightly

higher than those reported for CNF-based fibers fabricated with other approaches.62,

96 However, up to 6 times higher σu values of our fibers (Appendix Table 1) indicates

that apart from the orientation of nanoscale building blocks, interfaces and

interactions at the nanoscale play a key role in determining the mechanical properties

of macrostructures.

Other important observation is that the highest ever recorded strength of plant

fiber with high orientation of nanoscale crystalline planes, free from damage and

natural imperfections, is ∼1550 MPa (stiffness unknown).69 The plant fiber had

structural integrity provided by nature through centuries of evolution, where

nanocellulose are embedded in soft amorphous matrices of natural binders

(hemicelluloses and lignins) that probably assist with the distribution of the stress

and thus reduce the nanoscale defects. Interestingly, our TEMPO-oxidized CNF

fibers fabricated using double flow-focusing channel reach the σu equal to the highest

ever reported value for a plant fiber found in the nature with E on the order of

∼70−90 GPa.


34

4.3 Deformation Mechanism at the Nanoscale

Figure 4.6. Stress-strain data from loading-unloading experiments in uniaxial

tension for nanostructured CNF-550 fiber tested at 50 % RH at the same pulling rate

used for the tensile test of the fibers. Beyond the yield point, two cycles were

measured. Yellow and red lines represent the same slopes in different regimes

(before and after yield point).

Figure 4.7. Stress-strain data from loading-unloading experiments in uniaxial

tension for cellulose nanopaper with random-in-plane CNFs network. Reproduced

from Ref.92. (Copyright 2008 American Chemical Society)


35

Figure 4.8. MD simulation model for the relative sliding between neighboring CNF.

Variation of number of hydrogen bonds per unit area for two surfaces in direct

contact. Scale bar is 5 nm. (Image credit: Pan Chen)

High ε values (∼6%) obtained for the uncrosslinked CNF fibers is rather

uncommon for highly oriented cellulose structures. Structural changes and

interfibrillar interactions at the nanoscale are further investigated by cyclic

loading−unloading tests under uniaxial tension in the post-yield regime (Figure 4.6).

The post-yield E may increase due to reorientation of the fibrils as for the case of

random-in-plane CNFs network (Figure 4.7), along the test direction.92 Interestingly,

the fiber integrity is completely recovered upon unloading at post-yield strain values

in the present study. Several micro- and nanocomposites also show a decrease in

modulus due to the damages in the form of cracks at reinforcement−matrix interface.

However, unchanged E in the present case suggests a lack of further change in

orientation (due to the very high orientation of CNFs in the fibers) and the

mechanisms of plastic deformation must involve reformable secondary bonds as in

the case of stick−slip.97 The existence of stick-slip mechanism between the fibrils

was further verified by MD simulation (Figure 4.8), where the number of hydrogen

bonds per unit area remains constant during the relative sliding between the

nanofibrils.98 Plastic deformation of CNF fibers is substantially reduced upon cross-

linking, and the stress− strain curve becomes relatively linear as the secondary


36

interactions are replaced by covalent bonds (Figure 4.5C). Although interfibrillar

interactions are dominated by the relatively weak hydrogen bonds and van der Waals

forces,25 the high orientation of fibrils amplifies their effects due to collective

synergy of molecular interlocking, leading to stiffening and effective energy

dissipating mechanisms (stick− slip and molecular zip-up).97, 99

4.4 Mechanical Properties in a Wider Perspective

Figure 4.9. Tensile mechanical properties of bio-based and selected synthetic fiber

materials. Overview of specific σu vs specific E for a range of materials. The region

of fibers fabricated in the present work is shown in dark gray. (Reproduced from

Ref.24. Copyright 2018 American Chemical Society)


37

Figure 4.10. Tensile mechanical properties of bio-based and synthetic fiber

materials including aramids, carbon and ultra-high-molecular-weight polyethylene.

Dragline spider silk is considered as strongest biopolymeric fiber from past

several years. The comparison of the nanostructured CNF fibers reported in this

study, with the properties of dragline silk is also revealing. Notably, the stiffness of

CNF fibers is 8 times higher than dragline spider silk fibers with strengths higher

than most categories of dragline spider silks (Figure 4.9). The specific strength of

our CNF fibers is also higher than the metals, alloys, and silica-based E-glass fibers

(Appendix Table 2). Interestingly, the strongest CNF fibers have tensile strengths

1.2−1.5 times higher than wet-spun carbon nanotubes and graphene fibers100-101,

despite the fact that theoretical strength of crystalline cellulose is orders of

magnitude lower than that of a carbon nanotube and a graphene sheet. This is

possibly one of the limiting factors because the specific strength and stiffness of the

highly oriented CNF fibers are still lower than aramids, ultrahigh molecular weight

polyethylene (UHMWPE), and carbon fibers (Figure 4.10).

Carbon nanotubes as building blocks have much better mechanical properties

than CNFs as they contain 2D network of covalently linked carbon-carbon bonds, so

breaking a carbon nanotube requires high amount of energy to break many such

covalent bonds present together in a dense network. Whereas, UHMWPE consists

mostly of dense crystals (up to 99% crystallinity) where carbon chains are packed in

the most ideal scenario. Such crystals have high modulus and low density that results

in the elevated specific mechanical properties.102 Same is with aramids, where chains


38

inside the crystals are highly oriented. Nevertheless, the current results on tensile

mechanical properties of CNF fibers highlight the central role of processing

strategies and associated fundamental parameters (e.g., building blocks size,

interactions, and network formation) on realizing the true potential of nanoscale

building blocks.


39

5. Fundamentals of the Hydrodynamic Assembly

Process

Successful assembly of elongated nanoparticles such as CNFs, PNFs etc. into

high-performance macroscopic materials by using the microfluidic tools, depends

on fundamental phenomena controlling rotational and translational diffusion. The

elevated tensile mechanical properties of CNF fibers reported in chapter 4 and PNF

fibers reported in chapter 7, are results of our improved fundamental understanding

on the behavior of nanoscale building blocks in the microchannels and also due to

the understanding of the parameters associated with the processing technique. It is

rather challenging to understand and quantify the in-situ dynamics of nanoparticles

in the microfluidic channels. In this chapter, our studies on a new experimental

methodology to characterize the orientation and network dynamics of elongated

nanoparticles suspensions in assembly flows using orientation relaxation, are

reported. The relaxation is measured by birefringence and the methodology is

illustrated using nanocelluloses (CNFs and CNCs) as model systems. The advantage

of extensional flow in the organization and assembly of elongated nanoparticles over

the shear flow, are also discussed.

5.1 Double Flow-Focusing Geometry

Microfluidic tools based on co-flow,103-104 step-emulsification,105 T-junction106

or flow-focusing107 geometries have facilitated the development of both fundamental

and applied research in the field of chemistry, materials, biology and physical

sciences.108 In particular, the flow-focusing configuration (Figure 5.2A), where a

core fluid is focused by two sheath flows has been successfully employed for a

variety of applications.109 Flow-focusing relies on an extensional flow and have

demonstrated a great potential for the controlled assembly of elongated nanoparticles

into strong fibers or filaments.24, 34 Extensional flow is crucial for the orientation and

forces elongated nanoparticles to align in the flow direction (discussed in next

section). In the flow-focusing channel, the rate of extension must be larger than the

rotary Brownian diffusion of the nanoparticles to align them along the fiber axis for

the enhancement of uniaxial mechanical properties.74 This rate of extension can be

tuned with the flow-rates of the sheath fluids.

Assembly of elongated nanoparticles in the suspensions to solid fibers inside the

flow-focusing geometry typically relies on a transition to a gelled or an arrested-

state.24, 34, 110 For the CNFs assembly reported in chapter 4, a double flow-focusing

channel has been used (Figure 4.1) as single step flow-focusing channel has


40

disadvantages of channel clogging and gel build-up at solid interfaces (corners)

where the core and active sheath flow comes in contact.74, 111 This in turn result into

the lower orientation of CNFs inside the fibers. Hence, to obtain an enhanced and

uniform orientation of the nanofibrils throughout the fiber cross-section, we stick to

the multi-step flow-focusing channel (double flow-focusing) for most of the studies

on materials fabrication reported in this thesis. Double flow-focusing channel assists

in enhancing the orientation of nanofibrils inside the fibers (Figure 4.4C) and the use

of DI water into the first-focusing channel, solved the problem of gel build-up at the

channel walls and also prevent shear forces on the nanofibrils from the channel walls.

5.2 Shear vs Elongation

Figure 5.1. Illustration on the arrangement of fibrils while processing through

extensional and shear-based flow systems. Fibrils arrange uniformly in the liquid jet

of core flow of CNF in the extensional flow, which is not the case in shear dominant

systems, where fibrils can flip and rotate while processing. This leads to

inhomogeneous arrangement of fibrils across the cross-section of macroscopic

fibers, which in turn has implications on the mechanical properties.

Shear and extensional flows may give rise to the significant differences in the

orientation of nanofibrils inside the fibers. An extensional flow-field will align fibrils

without intermittent flipping and rotation, as is the case for shear where alignment is

a time-averaged quantity (Figure 5.1).24 The non-uniform arrangement of nanofibrils

from skin to core in the fibers will bring down the number of contact points between


41

two adjacent fibrils and hence, the structural integrity. This will have an influence

on the mechanical properties of nanostructured materials.

A thorough investigation on the differences in the behavior of CNFs under the

influence of shear and extensional forces by Rosén et al. is further insightful in this

regard.112 Extensional flow is found to be effective in the alignment of CNFs and the

alignment vanishes quickly due to the Brownian effects as soon as acceleration is

stopped, though strong shear is still present inside the microchannels. Rotary

diffusion (discussed in the next section) that is responsible for the de-alignment is

faster under the influence of shear compared to the extensional forces.112 Both shear

and extensional forces tend to align the nanofibrils in the flow direction. However,

in the extensional flows-based systems, the perpendicular distance between fibrils is

closer compared to the shear flow-based systems, that helps to reduce the rotational

diffusion of the nanofibrils.

5.3 Rotary Diffusion and Size Dependency

In this section, a new experimental methodology that allows real-time

assessment of orientation dynamics of elongated nanoparticles under dynamic flow

conditions, is described. The approach is based on a flow-stop methodology

developed by Rosén et al.112 coupled with thorough analysis of the orientation

relaxation.14, 113 Once the flow inside the microchannel reached a steady state, it is

rapidly stopped, and the relaxation of the oriented particle toward isotropy,

depending on the diffusion of the system constituents, is then measured. The

experiments are conducted at different locations in the channel. Combining these

measurements with a model of rotational diffusion,87, 89 this methodology provides

information about the typical size of the aligned nanoparticles before the flow is

stopped. Therefore, the evolution of the length distribution of the aligned

nanoparticles in the flow can readily be quantified. For the demonstration, we used

two types of nanocelluloses: CNCs, being almost monodisperse in length, and CNFs,

with a broad length distribution (polydisperse) (Figure 3.2).


42

Figure 5.2. (A) Typical birefringence signal for the CNF suspension. The signal is

normalized by its maximum value in the channel. The squares represent the different

locations where the birefringence is averaged in space. (B) Length distributions of

the CNC (dashed dotted line) and CNF (solid line) suspensions. The insets show

TEM images with the scale bars representing 200 nm. (Reproduced from Ref.14.


Experiments for rotary diffusion studies are conducted on a single flow-

focusing channel (refer paper II and III for details). The orientation of nanocelluloses

is visualized through birefringence using POM.112 A typical birefringence signal

from CNFs is shown in Figure 5.2A. Once the flow is stopped after reaching a steady

state, the birefringence decay B(t) is observed at multiple locations along the

centerline of the channel marked by the black squares in Figure 5.2A. The length

distributions of the corresponding nanocellulose samples measured using TEM are

plotted in Figure 5.2B.

It is important to model the Brownian effects inside the microfluidic channels

and correlate the return-to-isotropy of the aligned nanoparticles. Doi and Edwards87

developed a theory to study the Brownian motion of monodisperse rods which

depends on the length L, diameter d, and concentration c of the particles. For cL3 ≪


43

1 (dilute regime), the rods are free to rotate without any interparticle interactions,

while for 1 ≪ cL3 ≪ L/d (semi-dilute regime), the effects of particle interactions on

the particle dynamics become significant due to volume exclusion and network

formation.114-115 With a dry weight concentration of 41 g/L (cL3 ≈ 4.2), the CNCs

suspension belongs to the semi-dilute regime. The rotational diffusion coefficient of

a rod in an isotropic semi-dilute system is

𝐷𝑟 =𝛽

(𝑐𝐿3)2

𝑘𝐵𝑇 ln(𝐿

𝑑)

3𝜋ƞ𝐿3 (10)

where T is the temperature, kB is the Boltzmann constant, η is the solvent viscosity,

and β is a numerical factor. β is expected be of order 1−10, but several experiments

measured its magnitude in the range of 103−104.14, 116 This discrepancy has led the

need to redefine theories, model more precisely the diffusion of the nanorods through

the network, and possibly take into account nanoparticles flexibility. The

birefringence relaxation toward isotropy is predicted to be exponential with a typical

time scale of τ = 1/(6Dr).87 Figure 5.3 shows the birefringence relaxation of the CNC

suspension at different locations in the channel, both in a lin−lin and log−lin plot

(inset). All the curves are normalized by their initial value, B0. The curves almost

collapse, and the inset indicates an exponential decay. This illustrates that

independent of the initial alignment, only one timescale can be associated with the

relaxation of the monodisperse CNCs particle system, in agreement with Doi and

Edwards theory.87

Figure 5.3. Normalized birefringence decays for the CNC suspension at different

positions in the channel: x/h =−1 (blue), 0.75 (light blue), 1.25 (green), 1.75

(magenta), 4.75 (red), and 8 (black). The inset shows the same plot with a

logarithmic scale for the vertical coordinates.


44

The Brownian diffusion model of Doi and Edwards was primarily applicable

for monodisperse systems and hence, there was a need to extend it to polydisperse

systems, to understand the in-situ dynamics of CNFs or other similar biological

systems, where polydispersity is ubiquitous. Marrucci and Grizzuti proposed a

model for the polydisperse systems.89 In such systems, concentration of

nanoparticles in the suspensions depends on the entanglement length (L*) (see

methods for details).

Figure 5.4. Normalized birefringence decays for CNCs (empty diamonds) and CNFs

(filled circles), at x/h = 1.75. The fit of each curve using the inverse Laplace

transform method is shown as dashed dotted line for CNCs and solid line for CNFs.

The inset shows the normalized birefringence decay for CNF in a log−lin plot.

Figure 5.5. Normalized birefringence decays for CNFs. The symbols are

experimental data and the solid lines are the inverse Laplace transform fit. The colors

are for different positions in the channel: x/h = −1 (blue), 0.75 (light blue), 1.25

(green), 1.75 (magenta), 4.75 (red), and 8 (black). The two insets are schematics of


45

the orientation of the short (blue) and long (red) fibrils before the flow is stopped, in

the focusing region and far downstream in the channel.

𝐵(𝑡) = ∫ 𝐵0(𝐿)exp (−6𝐷𝑟∞

0𝐿(𝑡)𝑑𝐿 (11)

Here B0(L) is the contribution of the elongated nanoparticles of length L to the total

birefringence signal B0, before the flow is stopped. This leads to a relaxation toward

isotropy with multiple time scales, as can be seen for the CNFs in Figure 5.4 (filled

circles and inset). The difference with the exponential relaxation of the CNCs, also

plotted in Figure 5.4 (empty diamonds). These results highlight a strong coupling

between the orientation relaxation of a nanoparticle system and the length

distribution. This coupling allows us to extract information on the length distribution

of nanoparticles present in the suspension. Indeed, eq. 11 is a Laplace transform of

B0(L) and this quantity can be estimated by inverting B(t), using a method described

by Rogers et al.113 The numerical factor β given in eq. 10 has to be adapted. For the

CNFs, it has been set to β = 1.5 × 103, which is of the same order of magnitude as

reported by other experiments in the earlier studies.14 For the CNCs, β = 0.5 matches

well with the length distribution obtained using TEM in dry state (Figure 5.2B). The

discrepancy between the two values reflects that CNCs and CNFs suspensions form

3D networks differently based on their aspect ratio, and morphology. This analysis

further allows to quantify the evolution of length distribution of the aligned CNFs

along the channel and get information on the orientation dynamics of the nanofibrils

suspension in the flow (Figure 5.5) (refer paper II and III for details).

Above mentioned results allow us to establish a comprehensive understanding

of the orientation dynamics of CNFs in microfluidic channels, specifically under

extensional flow-fields. Shorter CNF nanoparticles are quickly oriented by the flow,

while more time is needed to align long ones. Furthermore, short particles de-align

faster and their motion is therefore difficult to control during the assembly. These

results are further supported by the in-situ SAXS measurements on CNF-550 and

CNF-1360 (Figure 4.4C). This scenario can be also deduced from Figure 5.5, used

as a diagnostic plot: aligned short fibrils in the focusing region exhibit short time

scales in the decay (light blue triangles), while aligned long fibrils further

downstream exhibit longer time scales (red squares and black diamonds). These

findings have strong implications on optimizing the mechanical performance of the

CNF fibers fabricated via flow-based assembly. The typical time for the transition

to gel or an arrested state, achieved with the double flow-focusing geometry in

chapter 4 is estimated to be of the order of a second (paper I),24 that is, longer than

the time found here for the dealignment of the shortest fibrils. Therefore, there is a

possibility that the CNF fibers fabricated using flow-focusing approach in chapter 4,


46

may composed of long aligned CNFs embedded into an isotropic matrix of shorter

CNFs.

5.4 Network Formation at the Nanoscale

Results in the previous section highlight that the orientational dynamics of CNFs

or other polydisperse nanofiber systems, depends strongly on the length distribution.

Polydispersity has significant influence on the dynamics of nanoscale networks and

in turn the mechanical properties of the hierarchical macroscale materials. A broad

length distribution further challenges our fundamental understanding on the physical

phenomenon at the nanoscale in an entangled network of nanofibers, experiencing

multiple regimes depending on their length and as polydispersity increases the mean

entanglement increases significantly.116 The nanofibers assembly under the

influence of Brownian effects can be controlled by tuning the concentration and

nanofiber dimensions in a way such that entanglement restrains Brownian motion

without impeding flow-induced alignment.24

We used Nc to parametrize the entanglement within the CNFs network in the

suspension. This dimensionless number is defined as the average number of

nanofibers in a sphere with a diameter equal to the length of the fibers.24, 65

𝑁𝑐 = 2

3∅ (

𝐿

𝑑)2 (12)

Here, L is the nanofiber length, d the diameter, and ∅ the volume fraction in the

suspension. The crowding number allows us to define three different concentration

regimes.117 For Nc < 16, the nanofibers are in the dilute regime, with no physical

crosslinking between them. In the semi-concentrated intermediate regime, where 16

< Nc < 60, the nanofibers have significant physical crosslinking but are not

immobilized. For suspensions with Nc > 60, the nanofibers are in a connected

network where free motion is hindered. The two thresholds, Nc = 16 and Nc = 60, are

respectively called connectivity and rigidity thresholds.116 In the literature on

polymer-based rigid rods, the crowding number is defined as the number of

nanofibers in a cube of side L.118 This creates a small discrepancy between the two

definitions, with a factor of one half. However, similar concentration regimes can be

defined in both the cases.

In the case of a network with a nanofiber length distribution, polydispersity

gives rise to a distribution of crowding numbers. Different approaches have been

suggested to estimate this distribution and the mean crowding number <Nc>, for a


47

log-normal or a general length distribution.116 In all the cases, polydispersity

increases the <Nc>, and therefore the mean entanglement. All the nanofiber

suspensions with concentrations of 0.3 wt.% used for the CNFs fiber fabrication in

chapter 4 exhibit a mean crowding number higher than the percolation threshold

(<Nc> = 16) showing that the interactions between the nanofibers are important and

mechanical interlocking between the nanofibers is present during all stages of flow-

focusing. At the concentration of 0.3 wt.%, most of the nanofiber suspensions are

above the rigidity threshold (<Nc> = 60). This clearly emphasizes that the nanofibers

are highly entangled and constrained in a 3D network. Note that the <Nc> does not

represent the proper entanglement for all nanofibers. Indeed, the entanglement of a

nanofiber originates from the number of contact points with its neighbors. Therefore,

the shorter the nanofibers, the smaller the number of contacts and less the

entanglement. This is further supported by a recent study from Nechyporchuk et

al.119 , where almost five times higher concentration of CNCs compared to CNFs

was required to fabricate nanostructured CNC fibers with the similar approach

demonstrated in chapter 4 for the CNFs and even that high concentration of CNCs

resulted in much weaker mechanical properties of CNC fibers compared to the CNF

fibers. Nevertheless, the entanglement length L∗ defined in the previous section

represents the threshold between small nanofibers in the dilute regime (not

entangled) and the larger ones in the semi-concentrated regime (entangled) (refer

paper III for details).

To maximize nanofiber orientation in the fabrication of materials with enhanced

mechanical properties, it is therefore necessary, to work with long and partially

entangled nanofiber systems, i.e. with broad length distributions and concentrated

suspensions slightly above the rigidity threshold. Indeed, long nanofibers own a de-

alignment timescale larger than the gel transition time scale and large concentrations

lead to high entanglement and thus maximize the flow-induced alignment. Too high

concentrations are however limited by the rigid volume-spanning arrested states,91

where the suspension becomes like a gel due to entanglement. In such gel-like

suspension, flow alignment is completely restricted. Hence, it is highly important to

optimize the nanoparticle systems based on the dimensions and concentrations to

obtain best possible mechanical properties at the macroscale.


48

6. Nanostructured CNF/Silk Composites

This chapter focuses on the supramolecular assembly of CNF and silk fusion

proteins for the fabrication of strong and bioactive nanocomposites. Cellulose and

silk are two excellent biological materials- silk from spider being tough due to high

extensibility whereas cellulose possess very high strength and stiffness. Along with

the elevated mechanical properties, building blocks of spider silk (spidroins) have

demonstrated potential as biomaterials, which is typically not the case with wood-

based CNFs. Unfortunately, silk proteins cannot be obtained in large quantities from

spiders and process of recombinant production is rather expensive, Therefore, in this

study, we combined small amounts (10%) of functionalized silk proteins with

nature’s most abundant structural component (CNF) to fabricate composites that can

have properties and functionalities of both silk and cellulose. We combined CNFs

with two types of silk fusion proteins: Z-silk and FN-silk.34 In Z-silk, the silk protein

is fused to an affinity domain Z,120 which adds the ability to selectively bind

immunoglobulin (IgG).83 In FN-silk, the material is instead functionalized with a

tripeptide Arg-Gly-Asp (RGD) containing cell-binding motif from the extracellular

matrix (ECM) protein fibronectin.121 The results presented in this chapter are

selected from paper IV. CNF and silk contents are referred in wt.% unless otherwise

stated.

6.1 Fabrication of CNF/Silk Nanocomposites

We mixed CNFs and RSPs in a hydrocolloidal suspension, where the stability of

the suspension is controlled by supramolecular interactions.122 A schematic of the

atomic model used for comparison of the molecular sizes of different components,

is shown in Figure 6.1A. We fabricated isotropic films and anisotropic fibers from

the CNF and CNF/silk proteins. For the films, a straightforward preparation process

based on solvent casting is used, where free-standing films with planar fibril

networks are obtained after evaporation of water (Figure 6.1B; Figure 1(b) of paper

IV). Except the simplicity, a key advantage of this technique is the capability to

prepare large-area films with well-defined dimensions, which is not possible using

other film forming methods like dip or spin coating.123 Morphology of the films

surface show a dense web-like fibrillar network, where the CNFs from the wood

pulp fibers are clearly visible. Individual CNFs are wriggled and physically

entangled to each other, indicating strong fibril-fibril interactions (e.g. hydrogen

bonds, Van der Waals interactions92), and a result of fibril flocculation due to the

crowding factor/number density. AFM images reveal the presence of proteins in the


49

form of small clusters (3-10 nm) homogeneously distributed throughout the dense

CNFs network.

Figure 6.1. (A) Schematic of functionalized RSPs, CNFs and CNF decorated with

RSP. (B) AFM images of CNF and CNF/silk film surfaces, where homogeneous

distribution of proteins throughout the surface can be seen. (C) FE-SEM image of a

CNF fiber prepared with double flow-focusing geometry.


50

Figure 6.2. (A) Average G’ of CNF and CNF/silk gels obtained from rheology

measurements. (B) FT-IR of CNF and CNF/silk fibers.

The orientation of fibrils in the fabricated films is isotropic in the place. To

fabricate anisotropic fibers, nanofibrils were oriented along the longitudinal

direction. We utilized the same approach of elongational flow-field based double

flow-focusing geometry that was applied for the fabrication of CNF fibers from

TEMPO-oxidized CNFs. For the experiments, we used a core flow of

carboxymethylated CNF or CNF/silk suspensions. The first sheath consists of DI

water (pH ~ 5.6) and the second sheath consists of HCl (pH ~ 2). Hydrogen ions

from the acid neutralizes the carboxyl groups on the fibrils surface to induce self-

assembly. Parallel and bundled arrangement of nanofibrils can be seen in the FE-

SEM image (Figure 6.1C) of CNF fiber. No difference in the surface morphology of

CNF and CNF/silk fibers is observed with FE-SEM.

Rheological measurements (Figure 6.2A) showed that the proteins increase the

gel stiffness of CNF significantly. Inclusion of Z- and FN-silk to CNF, resulted in

30-35% higher average G’ values than the corresponding CNF gels. These results

suggest better crosslinking of CNFs in the presence of proteins is probably due to

the formation of hydrogen bonds between the -NH groups of proteins and hydroxyl

groups of cellulose.124 However, more experimental investigations are required to

get information on the nanoscale interactions between the RSPs and CNFs. The

integration of proteins with CNF was also confirmed in the air dried fibers using FT-

IR. Proteins were not washed away in the water bath, where hydrogel fibers were

collected. Peaks appeared at wavelengths ~1650 cm-1 and ~1530 cm-1 in the

CNF/silk fibers correspond to the amide I and amide II bands of proteins,

respectively, which are related to the secondary structures of proteins and the

formation of β-sheets.125


51

Another interesting observation is that RSPs can form stable suspension with

charged CNFs that can be processed into hierarchical structures subjected to

hydrodynamic forces in the microchannels, despite the fact such forces are known

to promote uncontrolled aggregation of the silk proteins.126 The capacity of CNF as

a stabilizing agent has also been observed for carbon nanotubes127 and thus

demonstrates a viable route of using it as a common building block and /or as a

stabilizer for the fabrication of high-performance nanocomposites.

6.2 Mechanical Properties of Anisotropic Fibers

Figure 6.3. Mechanical and structural characterization of the CNF and CNF/silk

fibers. (A) Tensile mechanical properties measured at 50% RH. (B) E, (C) σu and

(D) ε of CNF and CNF/silk fibers. (E) WAXS diffractogram of a CNF fiber, where

IDG denotes the intermodular detector gap. (F) Azimuthal integration of the (2 0 0)

scattering plane of the diffractogram shown in Figure 6.3E.


52

The tensile mechanical properties of CNF and CNF/silk fibers are reported in

terms of E, σu and ε. CNF fibers prepared with carboxymethylated CNF revealed an

average E of 52.8 ± 0.7 GPa, σu of 830 ± 24 MPa, and ε of around 6% with a

toughness of 37 MJ m-3 (Figure 6.3). The highest obtained σu for a CNF fiber is 854.6

MPa. E and σu values obtained with the carboxymethylated CNF are slightly lower

than the values obtained with TEMPO-oxidized CNF fibers prepared with the same

strategy. This could be due to the difference in pulp properties that was used to

liberate the nanofibrils.

The orientation of nanofibrils in the CNF fiber is characterized using WAXS

(Figures 6.3E and F). The (2 0 0) reflection of cellulose I, that in general is used to

study the orientation of crystal planes, is here used to track the orientation of CNFs,

as crystals are aligned in the fibril direction.34 The fibril orientation in the CNF fiber

was calculated using the eq. 9. The value of fc in this study (0.86) is slightly higher

than previously reported values for cellulose fibers formed by the assembly of wood-

based CNF (fc varies from 0.55 to 0.83) and lower than the value of TEMPO-

oxidized CNF-550 fiber reported in chapter 4.24 This could also be a reason for

slightly lower values of mechanical properties obtained in this study. The scatter in

data between E and fc values obtained in this study compared to the other studies

reported earlier, suggest that apart from the degree of orientation, other structural

characteristics of the macroscopic fibers as well as chemical interactions between

the constituent fibrils play an important role in determining the macroscale

mechanical properties. The higher values of strength obtained for the fibers in this

study compared to the other techniques used to orient and assemble CNFs, could

also be related to the smaller dimensions of the current fibers, which in turn decreases

the probability of defects inside the macrostructures.98

We also tested the mechanical properties of CNF/silk fibers to observe any

potential difference in the mechanical properties in the presence of small amounts of

silk fusion proteins. Fibers of CNF/FN-silk showed comparable properties, with E

of 50.2 ± 1.3 GPa, σu of 812.5 ± 25.4 MPa, and ε of 6−8% with toughness of 43 MJ

m−3 (Figure 6.3). Surprisingly, fibers of CNF/Z-silk show higher ε (10%) with

toughness of 55 MJ m−3 and σu on the order of ∼1 GPa (978.8 ± 37.5 MPa) compared

to the pure CNF fibers, whereas the E (53.6 ± 1.8 GPa) is similar. One possible

explanation is that the addition of Z-silk proteins, which harbor a stably folded

helical Z-domain fused to the silk, promotes crosslinking between CNF fibrils and

enhances the overall strength. However, this effect is less evident for the CNF/FN-

silk fibers, in which the silk modules are instead fused to only a short peptide.

Nevertheless, in the seminal work of Xia et al.71 and Lazaris et al.,70 a E of 21 and

12 GPa and σu of 508 and 259 MPa have been reported for artificially spun fibers of


53

pure silk proteins alone. It is noteworthy that the E and σu of the CNF/silk fibers

designed in this study are 2−4.5 times higher, compared to the pure silk protein fibers

due to the presence of crystalline CNFs, whereas at the same time, the toughness of

artificial silk protein fibers is maintained.70-71 Tensile tests results of fibers are in

good agreement with the rheology measurements on corresponding gels of CNF and

CNF/silk proteins (Figure 6.2A), where increased values of G′ indicate the enhanced

interactions or crosslinking between the fibrils in the presence of silk proteins. The

tensile test results of the air-dried fibers along with FTIR measurements (Figure

6.2B) also indicate that RSPs are not merely present along with the CNFs in the

suspensions, rather they interact and stick well to the surface of CNFs.

Figure 6.4. Schematic representation of a CNF/silk fiber. Green-CNFs; purple-

RSPs.

Friction between the nanofibrils decorated with RSPs and a decrease in the pH

during the self-assembly could lead to the transition of the silk proteins from α-helix

to β-sheets and thereby promote crosslinking due to hydrogen bonds formation.124

Apart from the interfibrillar interactions and the network formation that contribute

to the stiffness of a pure CNF fiber, the presence of silk proteins enhanced the

crosslinking between the CNF (Figure 6.4), which could be a reason for enhanced

mechanical properties. To reveal the molecular origin of enhanced mechanical

properties, a thorough understanding of the conformational changes, intermolecular

interactions and heterogeneous microstructure at the nanoscale in the dried fibers, is

required.


54

6.3 Bio-Functionalities

Figure 6.5. Functional studies of composite films show presence of bioactivity. (A)

Optical microscopy images of CNF and composite CNF/silk films (top panel). Scale

bar is 4 mm. Images from fluorescence microscopy visualizing the binding of

fluorophore-labeled IgG (green, bottom panel). Scale bar is 1 mm. (B) Signal from

IgG-fluorophore bound to a film casted from a mixture of CNF and the Z domain

alone (CNF/Z), CNF premixed with Z-silk (CNF/Z-silk) or pure CNF, respectively,

visualized by fluorescence microscopy (top panel). Visualization of IgG-fluorophore

bound to a CNF film coated with the Z domain alone (CNF/Z), Z-silk (CNF/Z-silk),

or noncoated (CNF) (bottom panel). Scale bars are 1 mm. (C) HDF cultured on films

of CNF or CNF/FN-silk for 7 days. Top panel: live/dead staining (green, living cells;

bright red, dead cells; dim red, autofluorescence from CNF). Bottom panel: F-actin

staining (green, F-actin; blue, nuclei/ autofluorescence from CNF). Scale bars 500

μm for 2× and 100 μm for 10×. (D) Number of HDMECs (mean + SD) per mm2 on

freestanding films of CNF (open bar) and CNF/FN-silk (gray bar), at different time

points during culture, as evaluated by DAPI staining and manual counting of nuclei.

Statistically significant differences according to Student’s t test: **p < 0.01, ***p <

0.001, and *****p < 0.00001. (E) Representative Alamar blue viability graph of

three independent experiments with HDMEC growing on CNF and CNF/FN-silk

films (n = 3) for 1 week. Statistically significant differences according to Student’s

t test: *p < 0.05. Error bars show standard deviation. (Reproduced from Ref.34.


Specific affinity is the ability to selectively sort out one type of molecule from a

complex. It is essential in applications such as purification of antibodies

biomolecules. IgG-binding ability of CNF films in combination with Z-silk (Figure


55

6.5A and B) was evaluated. Macroscopic appearance of CNF and CNF/silk films are

nearly identical (Figure 6.5A, upper panel), however, CNF/silk films spread more

on the Teflon sheets where the films were casted, which might be due to the changes

in the wetting properties of CNFs in the presence of silk proteins. Fluorophore-

labeled IgG was used to evaluate the IgG-binding ability of the films (Figure 1.4),

which exhibits a clear green signal upon binding to the CNF/Z-silk film (Figure

6.5A, lower-middle panel). The signal is uniform over the entire surface area of the

CNF/Z-silk film, indicating a homogenous distribution of silk proteins in the

composite films. No green signal is observed for either the CNF or from CNF/FN-

silk films used for the comparison (Figure 6.5A, lower-left and right panel,

respectively). This indicates that the addition of Z-silk into the CNF adds a

prominent IgG-binding ability. This also highlights that the Z domains of Z-silk are

exposed outwards in the CNF/Z-silk film and retain the native structure (3D helical

fold) that is required for IgG-binding.

To investigate whether the combination of silk module (4RepCT) in Z-silk is

necessary to achieve IgG-binding ability, a control experiment was performed,

where CNFs were mixed with Z domain alone and casted into the films with silk

module being excluded. While performing the analysis on such films (CNF/Z), a

weak fluorescence signal is detected, indicating significantly low binding of the

labeled IgG (Figure 6.5B, upper-right panel). These results strongly emphasize that

covalent attachment of the Z domain to the silk module is necessary to achieve the

IgG-binding ability. This is probably due to poor adhesion of Z domain alone to the

CNF in the absence of silk module, that acts as an efficient mild glue. It can be

speculated that remaining Z domains will not be able to keep the 3D folded structure

and will be denatured upon direct contact with CNFs. Further investigation on the

role of silk module, is carried out on the CNF films that were coated with a solution

of the Z domain alone and the Z-silk fusion protein, respectively. Again, CNF coated

with the Z domain alone and CNF coated with Z-silk protein, was found to bind a

little IgG like the control film made of pure CNF (Figure 6.5B, lower-right and left

panel, respectively). This concluded that a covalent crosslinking between Z domain

and the silk module is needed, to strongly bind folded Z domains to the composite

materials.

In the recent years, attempts have been made to use wood-based nanocellulose

scaffolds for cell-culture.128 Mostly hydrogels of CNFs have been used for cell

encapsulation in a 3D micronenvironment.129-131 Typically, a good biocompatibility

of CNF with a maintained viability of the encapsulated cells is observed. However,

direct adhesion of cells with cellulose surfaces in a cell-matrix like fashion, is not

shown. This is very likely due to the lack of bioactive motifs in wood-based


56

cellulose. By glutaraldehyde cross-linking of CNF extracted from trees, Mertaniemi

and co-workers132 could show an improved spreading of cells on the threads after 1

week of culture.

We therefore evaluated the ability of CNF-based materials to support the

adhesion and proliferation of primary cells with the inclusion of a silk protein fused

to a cell binding motif from fibronectin (FN-silk). First, HDFs were seeded onto the

different films and adhesion of cells was traced by viability staining after 1 week

(Figure 6.5C, upper panel). A comprehensive coverage of living fibroblasts was

observed on the CNF/FN-silk films, with limited number of cells found on the films

made of pure CNF, although the present cells were also viable. Fibroblasts cultured

on CNF/FN-silk films exhibited well developed stress fibers and elongated

morphology (Figure 6.5C, lower panel). This verifies that the cell binding motif FN

fused to silk is presented uniformly in the composite films that can be recognized by

cell surface receptors mediating adhesion, e.g. integrins. To challenge the system

further, the films were seeded with HDMECs, a cell type known to be dependent of

ECM proteins for adhesion. The number of cells adhered to the CNF and CNF/silk

films were recorded 2, 3 and 6 days after seeding (Figure 6.5D). Results confirm that

inclusion of the FN-silk protein is mandatory to obtain a material which promotes

cell adhesion. HDMECs were well spread out after one-week culture on the

CNF/FN-silk films (Figure 6.6). In contrast, cells did not show these characteristics

on CNF films. Surface roughness that could be an additional factor towards the

adhesion of cells on any surface is also found to be very similar for the films

fabricated with CNF and CNF/silk proteins. To verify the proliferative nature of the

adhered cells, an Alamar blue viability assay was also carried out during 1 week of

cell culture (Figure 6.5E).

Figure 6.6. Freestanding CNF films with and without FN-silk after 7 days culture

with HDMECs and staining for f-actin (green) and nuclei (blue). Scale bars are 100

µm. (Image credit: Mona Widhe)


57

Figure 6.7. Affinity properties of the nanostructured fibers. (A) Optical microscopy

images of CNF and CNF/silk fibers. Scale bar is 4 mm. (B) CNF/Z-silk fibers ability

to bind IgG-fluorophore (green) visualized by fluorescence microscopy. Scale bar is

0.5 mm. (Reprinted from Ref.34. Copyright 2017 American Chemical Society)

Figure 6.8. Cell adhesion to fibers. (A) Live/dead staining of HDFs growing on CNF

(left panel) or CNF/FN-silk fibers (right panel) after 2 days (top panel) or 7 days

culture (middle and bottom panels). Green, living cells; bright red, dead cells; dim

red, autofluorescence from CNF. (B) F-actin staining (green) of HDMECs after 2

days (top panel) and 4 days (bottom panel) of culture adhering to CNF (left panel)

or CNF/FN-silk fibers (right panel). Dim blue: autofluorescence from CNF. Scale

bars 250 μm for 4× and 100 μm for 10×. (Reprinted from Ref.34. Copyright 2017

American Chemical Society)


58

We also evaluated the bio-functionalities in anisotropic fibers. Macroscopic

appearance of air-dried CNF/Z-silk and CNF/FN-silk fibers is like the corresponding

CNF fiber (Figure 6.7A). Regarding the IgG binding ability, a distinct green

fluorescence signal highlighting the binding of IgG, is observed in the CNF/Z-silk

fiber (Figure 6.7B, middle panel), unlike pure CNF or CNF/FN-silk fibers (Figure

6.7B, left and right panel, respectively). This indicates the visibility of appropriately

folded Z domains in the CNF/Z-silk fibers.

When cells were cultured on the fibers, a slightly different behavior was

observed compared to the isotropic films. Cells did not adhere to the CNF films,

however, fibers of CNFs supported some attachment and spreading, though to a

lesser extent than the CNF/FN-silk fibers that contain cell-binding motifs. HDFs

spread well along the CNF/FN-silk fiber, after 2 days in culture media, whereas on

CNF fibers, a few scattered rounded colonies of cells with no signs of uniform

spreading were found at this early point (Figure 6.8A and B, upper panels). However,

after 7 days, the HDFs also found to spread out along the CNF fibers, still in the

form of small clusters indicating rather poor attachment (Figure 6.8A, middle and

lower panels).

Better attachment of HDFs on CNF fibers compared to the CNF films might be

explained partly by the differences in surface morphology; the films have random

orientation of nanofibrils with nano-sized roughness of similar dimensions

throughout, while the CNF fibers have bundled uniaxial arrangement of nanofibrils

(0.2−0.3 µm) that give rise to distinct roughness values within the same material

(Figure 6.1C). It could be due to the enhanced adhesion of ECM proteins secreted to

the culture media, alignment and extension of cells on the oriented structures and

subsequent adhesion and proliferation of cells.133 In this direction, a thorough

investigation of surface micro- and nanoscale topography and correlation with the

cell functions is required. We also observe a high confluence and even distribution

of fibroblasts on the CNF/FN-silk fibers, already from start.

Further, the fibers were seeded with HDMECs, and the adhesion and spreading

of cells were analyzed by staining for F-actin. After 2 days, very few small and

rounded colonies of HDMECs were observed on CNF fibers (Figure 6.8B), while

the cells were observed to be spread out along the CNF/FN-silk fibers. After 4 days,

some HDMECs were observed to spread out also on the pure CNF fibers, though

much less in number compared to the CNF/FN-silk fibers. Thus, in both isotropic

and anisotropic hierarchical CNF structures, the presence of the cell binding motif

FN linked to the silk, promote significantly the cell adhesion to the cellulose-based

materials.


59

7. Nanostructured Protein Microfibers

Many natural materials are made from proteins. The properties of these

proteinaceous materials are closely linked to the hierarchical assembly of the protein

building blocks. Proteins from egg, soybean etc. can be assembled into PNFs that

have potential to be a building block of the macroscale materials. Recent advances

have revealed the molecular structure of PNFs, the mechanisms associated with

interfibrillar interactions while assembling into macrostructures, remain unexplored.

In addition, previous attempts to fabricate PNF-based materials largely rely on the

use of plasticizers due to weak interfibrillar interactions. In this chapter, our results

on the assembly of PNFs into strong microfibers without adding any plasticizer or

cross-linker by using double flow-focusing channel are summarized. The assembly

of PNFs of two distinct morphologies (straight and curved) is compared. Insights on

the differences in the assembly of PNFs compared to more crystalline biopolymer

building block i.e. CNF, are also provided.

Figure 7.1. Schematic of the double flow-focusing device used for PNFs assembly.

The PNF suspension is injected in the core flow. DI water injected in the first sheath

and acetate buffer (pH 5.2) in the second sheath (top left). The channel ends in a bath

containing acetate buffer (pH 5.2) (top right). Photograph of a hydrogel fiber in the

bath solution after spinning in the double flow-focusing setup (top right). Scale bar

is 1 cm. FE-SEM image of a dried fiber and a zoomed view in the inset (bottom

right). Scale bars are 30 µm and 1 µm, respectively. Fluorescence microscope image

of a fiber stained with thioflavin T (bottom left). Scale bar is 30 µm.


60

7.1 Assembly and In-Situ Orientation Dynamics of PNFs

The whey protein β-lactoglobulin self-assemble into amyloid-like nanofibrils at

high concentrations of urea,134 in the presence of alcohols,135 and at low pH and

elevated temperature.136 In the latter case, fibrillation occurs after the hydrolysis of

the protein into smaller peptide fragments which then spontaneously assemble into

PNFs. In the present study, WPI is used as a starting raw material due to an advantage

of abundant availability. Details on the assembly of WPI into PNFs of two distinct

morphology are mentioned in paper V. The measured E of straight fibrils is ∼5 GPa,

and for the curved ones is ~0.2−0.4 GPa. The solutions of obtained PNFs were then

extensively dialyzed to remove unfibrillated peptides and proteins. The dialyzed

samples of PNFs undergo a phase transition to a gel (arrested state) at a pH of 5.2

i.e. close to the isoelectric point of β-lactoglobulin.137 However, non-dialyzed

samples resulted in phase separation with opaque aggregates at this pH, highlighting

the importance of proper dialysis as unfibrillated proteins can impede the protein

fibrillar network in the gel state.

Similar to CNFs, the strategy used to fabricate pure protein fibers also rely on

suspension to gel transition during the flow-induced orientation process. We used

the same double flow-focusing device (Figure 7.1) with a core flow of PNF

suspension and two focusing steps of the sheath flows. In the first focusing channel,

DI water was used and in the second channel, we use acetate buffer at pH 5.2. Acetate

buffer neutralizes the charges on the surface of PNFs and induce the self-assembly.

We first examined the assembly of straight fibrils, where the concentration of PNF

in the core flow varied from 0.45% to 1.8% (the concentrations are lower than the

above-mentioned initial concentrations as they refer to the pure PNFs content of the

dialyzed suspensions). Surprisingly, no conditions were found where fibrils could be

assembled into fibers that were strong enough to be picked out from the bath

solution. Weak hydrogel fibers were obtained that were unable to withstand the

surface tension at the liquid–air interface while being pulled out. The curved fibrils,

on the other hand, were able to produce hydrogel fibers that were strong enough to

overcome the liquid-air surface tension (Figure 7.1). Notably, the fiber formation

with curved fibrils was possible at a concentration of 1.8%, where straight fibrils

were not able to produce strong enough hydrogel fibers. The extracted hydrogel

fibers were air-dried (Figure 7.1) and surface morphology was investigated using

FE-SEM. The images reveal a smooth surface with constant cross-section along the

fiber-axis (Figure 7.1). In the zoomed view of the fiber surface, graupel-like units of

the dimensions of a few hundred nanometers were observed. Interestingly,

morphology of the assembled PNFs observed in this study is similar to previous


61

studies on the assembly of RSPs,71 where colloidal aggregates clustered together to

form weak structures in the solution. However, surface morphology of non-

crystalline PNFs is very different from crystalline CNFs, where we observed

oriented bundles of fibrils on the fiber surface.24 The presence of β-sheet secondary

structures in the fibers was confirmed by fluorescence microscopy (Figure 7.1), by

using the amyloidophilic dye thioflavin T. We further characterized the tensile

mechanical properties of the fabricated fibers, where stress–strain curves reveal a

rather brittle material with E up to 288 MPa and ε of ∼1.5% (paper V). Interestingly,

the measured E values of fibers are in good agreement with the estimated values for

the curved nanofibrils.


62

Figure 7.2. In-situ study of PNFs in the microchannel. (A) Schematic of the channel

geometry used for the SAXS experiments. White circles show the locations where

in-situ measurements were carried out. (B) SAXS diffractograms at selected

locations in the channel for straight (Upper Row) and curved (Lower Row) fibrils.

(C) Local order parameters calculated from SAXS diffractograms as a function of

downstream position z in the channel. The locations marked with circles in A

correspond to the data points shown in C. (D) Images of the channel placed between

two crossed polarization filters rotated 45° from the vertical axis (white arrows). For

straight fibrils, birefringence due to the alignment is observed after the focusing step.

In contrast, no birefringence is detected for the curved fibrils. (Reproduced from

Ref.66. Copyright 2017 National Academy of Sciences)


63

To understand the behavior of straight and curved PNFs in the microchannel,

we carried out similar SAXS experiments as reported earlier in chapter 4 for CNFs.

Same channel geometry was used for the SAXS experiments, where the second

focusing step was mimicked by a contraction (Figure 7.2A), as suspension to gel

transition was not required for these experiments. For the PNFs, similar to the CNFs,

an increased alignment of fibrils immediately after the focusing and the contraction

steps was observed (Figure 7.2B and C). Subsequently, relaxation of the aligned

fibrils toward isotropy due to Brownian motion result in decrease in order parameter.

Considerable differences in the behaviors of straight and curved PNFs were

observed. Very low orientation of curved fibrils in the flow direction is observed, as

indicated by the lower values of order parameter (Figure 7.2C). The maximum order

parameter recorded at z = 8 for the straight fibrils is 0.41, which is almost two times

the order parameter of curved fibrils (0.21) at the same channel location. SAXS

results are further confirmed qualitatively by POM (Figure 7.2D), where straight

fibrils displayed birefringence highlighting the ordering of PNFs and no

birefringence was observed for the curved fibrils that further confirms a lower degree

of alignment.

Figure 7.3. Behavior of PNF gels subjected to oscillatory strain in a parallel plate

rheometer. (Reproduced from Ref.66. Copyright 2017 National Academy of

Sciences)


64

Figure 7.4. Schematic of the interactions of straight and curved PNFs in bulk and in

the flow-focusing channel. (Left) In bulk, the deformation of gels of straight fibrils

under strain is governed by the joint motion (rotation and sliding), whereas curved

fibrils resist the strain due to the higher degree of entanglement. (Right) In the

channel, straight fibrils have higher orientation compared to the curved fibrils in the

flow direction as noticed with SAXS. In contrast, curved fibrils with shorter

persistence length are physically entangled resulting in a strong network between the

fibrils to form a fiber. (Reproduced from Ref.66. Copyright 2017 National Academy

of Sciences)

Figure 7.5. FT-IR spectrum of non-gelled suspension, hydrogel and dried fiber of

curved fibrils.


65

7.2 Micro- and Molecular Structure in the Gel State

Interfibrillar interactions are investigated for the straight and curvy fibrils by

studying the fracture behavior of the corresponding gel networks under macroscopic

strains using parallel plate rheometer (Figure 3.5). At similar fibril concentrations, a

noticeable difference in the value of G′ was observed (Figure 7.3). Initial alignment

of straight fibrils in a nematic phase could be the reason for the higher G′ values.138

Rheology results also reveal differences in the gel fracture behavior in response to

the increasing oscillatory strains. Gel from straight fibrils exhibits a slow decrease

in G′ before it fractures, whereas for the curved fibrils, a sudden drop in the G′ values

is observed. The differences in the gel fracture behavior is associated with the gel

networks, where the slow deformation for straight fibrils is presumably associated

with rotation and sliding of the nanofibrils, which plastically deforms their network

at high oscillatory strains. On the other hand, network of curved fibrils fractured

without plastic deformation, indicating that there might be an interlocking due to the

entanglements and eventually the entangled network breaks at high strain values.

Rheology results throw insights on the potential differences in the behavior of

morphologically distinct fibrils in the microchannels. When straight fibrils are

aligned by the extensional flow, the possibility to further optimize interfibrillar

network decreases due to higher alignment, resulting in suboptimal physical cross-

linking of nanofibrils. In contrast, when the entangled network of curved fibrils is

extended and thus redistributed under the influence of extensional flow-fields, the

physical crosslinking might be enhanced, and a stronger entangled network is

achieved (Figure 7.4). This is also clear from the FE-SEM image in Figure 7.1, where

dense packing of PNFs can be seen. To get information on the molecular structure

of PNFs in fibers, we acquired FT-IR spectra in amide I and amide II band regions

that correspond to the protein secondary structures.125 The amide I and II bands are

centered at ∼1,625 cm−1 and ∼1,517 cm−1, respectively, which reflects the presence

of β-sheets (Figure 7.5).

7.3 Hierarchical Protein Materials and Comparison with Other

Building Blocks

Typically, the orientation of nanofibrils or nanotubes along the fiber axis results

in strong nanoparticles network that improves the unidirectional mechanical

properties. For instance, the orientation of CNFs and carbon nanotubes inside the

macro-sized fibers resulted in enhanced tensile properties.24, 100 The results presented

in this chapter show a different behavior of PNFs compared to other building blocks.

Although the morphology and dimensions of PNFs resemble largely with CNFs, the


66

nanoscale structural features and chemical compositions are indeed different. The

differences in the origin of PNFs and CNFs can further elucidate the differences into

the hierarchical structures. CNFs from wood are highly crystalline and are extracted

through chemical pretreatment followed by mechanical decomposition of cellulose

fibers in trees. PNFs, on the other hand, are self-assembled from protein monomers

that do not correspond to any hierarchical architecture already exist in nature. Hence,

PNFs (non-crystalline) and CNFs (crystalline) are two distinct classes of

nanomaterials and have different assembly patterns into macroscopic materials.

Interestingly, our results on the assembly of PNFs into microfibers have some

similarities with the assembly mechanism of RSPs.71 Similar to the PNF microfibers,

spider silk also has a high degree of β-sheet secondary structure and the silk fibers

are formed in a duct inside the spiders, by using the elongational flow followed by

pH alterations.28 However, the nanoscale structural characteristics of amyloid-like

nanofibrils and spider silks are different. In spider silks, β-strands and β-sheets are

parallel to the main fiber axis, in a string-of-bead fashion, in which anti-parallel β-

sheets and amorphous regions co-exist and alternate, providing high strength and

extensibility than amyloids, where β-strands are arranged perpendicular to the fiber

axis while the β-sheets are parallel.139 Other important difference in the assembly of

spider proteins into silk, is the simultaneous formation and control of the

supramolecular structures inside the spider glands. This allows the formation of

entangled structures with optimized intermolecular interactions at the same time

highly ordered β-sheet regions are created. Considering that not even artificially

assembled silk protein fibers reach the mechanical properties on the same level to

the natural spider silk, it is indeed fascinating that we can produce pure protein fibers

from a rather crude raw material and with a simple and potentially scalable

experimental setup. Our insights on the PNF assembly will provide a basis for

rational design of novel materials with tailored properties.

CONCLUSIONS

67

8. Conclusions

The main findings from the studies reported in this thesis are summarized in this

chapter.

Flow-assisted assembly is a promising route for the fabrication of structural

materials with highly ordered arrangement of nanoscale building blocks such as

polymer- and protein-based nanofibrils, nanotubes, and inorganic nanorods.24, 34, 66,

140-143 Mechanical properties of our CNF and PNF fibers further support this

conclusion. The elevated mechanical performance of CNF fibers is a result of

understanding and controlling the key fundamental process parameters (nanoparticle

size, interactions, alignment, diffusion, network formation and assembly), that are

central to obtain a near-perfect nanostructuring. Length of the nanofibrils and

enhancement in interfibrillar interactions by physical and/or chemical crosslinking,

largely controls the mechanical properties of CNF fibers in the dry state. Conditions

such as fibrils length and concentration must be optimized such that there is a

restraint on the Brownian motion of fibrils inside the microchannels by mechanical

interlocking without hindering the movement that in turn will further promote the

flow-induced alignment. Above a certain fibril concentration in the suspensions,

fibril-fibril physical crosslinking not only impedes mobility, but also provide

necessary structural integrity during the assembly. However, too high concentrations

of fibrils in the suspensions results in volume-spanning arrested states due to

entanglements, which restrict the flow-induced alignment.

It is important to understand the effect of flow-fields on the organization of

building blocks, specifically in microfluidics-based design of new materials.

Typically, an extensional flow-field aligns fibrils without intermittent flipping and

rotation, which is the case for shear flow where alignment is a time-averaged

quantity. This enhances the physical crosslinking between the elongated

nanoparticles and thus results in structural integrity and improved mechanical

properties. Diffusion due to Brownian motion inside the microchannels is also a

crucial parameter at the nanoscale. However, the ability to quantify the diffusion

dynamics in microchannels remains challenging as the tools that can provide time-

resolved measurements under dynamic conditions are limited. Using POM, we build

a methodology to trace the rotational Brownian motion of nanocelluloses with the

help of birefringence measurements that further expands the capabilities to probe in-

situ dynamics of flowing suspensions. The results reveal that characteristic length

scales are strongly dependent on the nanoparticles (CNFs) length distributions and

entanglements.24 The knowledge from the diffusion and network dynamics of

CONCLUSIONS

68

nanocelluloses in this thesis can also be used to optimize the mechanical properties

hierarchical materials from other building blocks, apart from nanocelluloses. Our

methodology of dynamic characterization of nanocelluloses should be extendable to

other nanoscale building blocks such as amyloids, chitin, metal nanoparticles etc.

This tool will further help to get a facile and easy access to nanoscale dynamics for

a variety of biophysical systems that are presently limited by the access of

synchrotron-based X-rays and neutron techniques.

The study on CNF/silk composites shows that coupling of genetically engineered

silk proteins and CNFs leads to functional effects in both aqueous dispersed state

and dried state. Charged CNFs enhanced the dispersibility of silk proteins and

thereby the processability, which is not typical otherwise as hydrodynamic forces

tend to agglomerate silk proteins.126 Inclusion of Z-silk proteins gives rise to the

affinity properties in cellulose materials, while the incorporation of FN-silk proteins

in CNFs resulted in ideal matrix for cell culturing that enabled cell attachment,

migration, proliferation and differentiation. The efficient crosslinking between RSPs

and CNFs in the dry state enhanced the interactions between the nanofibrils, thereby

resulted in increased toughness and an elevated strength at break. Nonetheless, based

on the results presented in chapter 6, it can be expected that mechanical properties

and bio-functionalities in bio-based materials can be tuned through the coupling of

proteins and crystalline nanofibrils by interlinking architectures. This study further

shows a functional basis for the modularity found in natural materials and suggests

a biological approach for molecular-level genetic engineering of the soft-matrix in

nanocomposites.

The high order parameter of CNFs in the microchannels resulted in the

enhancement of uniaxial mechanical properties of macroscale fibers.24, 74 However,

the results obtained with the PNFs show a different behavior. Although the

morphology (shape and size) of PNFs might appear like CNFs, the intrinsic

properties are indeed different due to the differences in the origin. CNFs from wood

are extracted by mechanical fibrillation of long cellulose fibers. Crystalline part of

cellulose fibers highlights the intrinsic ability of trees in nature to form crystals by

highly ordered arrangement of the polymer chains. PNFs, on the other hand, are

obtained from the self-assembly of protein monomers that do not correspond to any

highly ordered structures that can lead to the formation of strong nano-sized crystals.

Hence, despite the similarities in size and aspect ratio, both the biopolymeric

building blocks (PNFs and CNFs) are two distinct classes of nanomaterials where

CNFs are crystalline and PNFs are non-crystalline. This is the reason the mechanical

properties of CNF fibers are orders of magnitude higher than PNF fibers. However,

our results with PNFs show some similarities between the microfibers fabricated

with whey proteins reported in this thesis and the assembly of RSPs reported

CONCLUSIONS

69

earlier.144 Like the PNFs, spider silk has β-sheet structures and the natural spider silk

fibers are also formed by a process inside the ampullate glands of spiders which is

similar to the spinning process here, where spidroins (spider proteins) are subjected

to elongational flow and pH alterations.145 However, the orientation of the β-sheets

relative to the fiber axis is perpendicular in the silk fibers to that expected for aligned

PNFs139, which leads to differences in structural characteristics and the mechanical

properties of dragline silk and amyloid-like whey protein nanofibril-based materials.

Furthermore, the molecular- and supramolecular structures that are formed

simultaneously during the spinning process inside the spider glands allows the

formation of entangled structures with much optimized intermolecular interactions.

This argument is further supported by the differences in the mechanical properties

observed between the natural dragline silk and man-made RSPs-based fibers. Our

study with PNFs shows the importance of interfibrillar entanglements over the

orientation to fabricate strong enough free-standing coherent protein fibers.

OUTLOOK

70

9. Future Work and Outlook

Mechanical properties of CNF fibers

Mechanical properties of CNF fibers obtained with our process is significantly

higher compared to other previous attempts made in the field. However, the true

potential of CNF is still unknown and a question of debate. Further developments in

the processing techniques are required to realize the full potential of CNFs in all-

cellulose-based macrostructures. It may be possible that the values reported in this

thesis are already close to the limit, however this still needs to be verified. One

potential way to achieve this is by tuning the length and crystallinity of CNFs. CNFs

from different wood species need to be tested. New ways to in-situ crosslink CNFs

in the microchannels may also contribute positively in this direction.

Role of ions on the mechanical properties of CNF fibers

In this thesis, we mostly used H+ ions from HCl to neutralize the charges of CNFs to

induce the self-assembly of nanofibrils. Apart from H+ ions from the acids, other

monovalent, divalent and trivalent ions such as Na+, Ca+2, Fe+2, Al+3 etc. can also be

used to lock the CNFs in the aligned state in the microchannels. It would be

interesting to see the influence of these ions on the mechanical properties and to

compare the nanoscale network and orientation of nanofibrils in the dried fibers with

WAXS.

Wet strength of CNF fibers

In our studies, we tested only the dry strength of CNF fibers (≤50% RH) as the focus

of our studies was to check the maximum potential of CNFs in all-cellulose

macrostructures. For industrial applications, one major challenge that needs to be

addressed is the sensitivity of CNF towards water. Water acts as a plasticizer for

polysaccharides such as cellulose146 and this drastically decreases the mechanical

properties when exposed to water or moist conditions. Hence, it would be

worthwhile to test the mechanical properties in wet conditions, and potentially

combine CNFs with other building blocks to bring down its sensitivity towards

water.

PNF fibers

Further exploration of processing conditions and/or protein building blocks is

required to fill the parameter space (mechanical entanglements, fibril orientation,

spatial distribution of fibrils, tuning the surface chemistry of PNFs and covalent

OUTLOOK

71

crosslinking) of PNFs assembly in detail. This will provide the basis for design of

novel protein-based materials with improved mechanical properties.

CNF/silk composites

Mechanical properties of CNF/silk composites can be tuned further by changing the

concentration of silk proteins with respect to CNFs. New silk proteins with stronger

ionic interactions and /or new bio-functionalities can also be synthesized and tested

with CNFs. Novel CNF/silk materials with improved mechanical properties or

advanced bio-functionalities can be designed. More on the fundamental side, it

would be interesting to find out the conformation of proteins at the nanoscale in the

dried structures (films and fibers) with respect to CNFs. Techniques like AFM and

neutron scattering may be able to provide more information.

Nanocomposites fabrication

Another direction is to fabricate high-performance nanocomposites with the process

of hydrodynamic orientation and assembly. Other building blocks such as chitosan,

chitin, metal nanoparticles, nanosheets etc. can be tried individually or in

combination and their mechanical properties can be optimized with more

fundamental understanding on their behavior in the suspension state.

Industrial production of fibers

Though we greatly manage to realize the potential of CNFs and PNFs in macro-sized

fibers, it is necessary to maximize the production speed in order to increase

scalability. Parallelization of flow-focusing channels can contribute positively in

scaling up the production. Potential ways to simultaneously dry the hydrogel threads

to be able to roll them up continuously, also need to be determined.

Rotary diffusion measurements using POM

We tried to quantify the in-situ orientational and network dynamics of

nanocelluloses using POM and birefringence measurements. Same technique can

also be used to get fundamental information on the behavior of other building blocks

such as amyloids, chitin etc. in the microchannels to optimize the performance of

nanostructured materials fabricated via hydrodynamic approaches. This tool can also

be used to check batch to batch variations in CNFs and other building blocks that

will further help to adapt these nanomaterials for industrial scale production. With

respect to the nanocellulose suspensions that are used as model systems in this thesis,

we used water as solvent. However, viscosity can be tuned by introducing low or

high molecular weight polymers and their effect on the relaxation dynamics

OUTLOOK

72

(characteristic length scales of diffusion) can also be studied. This information is

relevant for many polymer and particle-based dynamic systems.

Validation with simulations

Experiments performed in this thesis can be the test cases for simulations. It is

important to find out whether the models available today can predict and reproduce

the experimental results presented here. Most of the numerical models consider

elongated nanoparticles (CNFs, PNFs) as nanorods. It is challenging to consider the

periodic kinks or curviness of the fibrils into the numerical models. New models,

therefore, need to be developed to study the behavior of elongated nanoparticles in

the channels based on the differences in the morphology and the information on the

nanoscale physics can in turn be used in the industrial processes dealing with fibers

and suspensions.

ACKNOWLEDGEMENTS

73

Acknowledgements

The Knut and Alice Wallenberg foundation is greatly acknowledged for providing

funding through Wallenberg Wood Science Center.

I would like to express the deepest gratitude to my advisor Daniel for providing me

an opportunity to start this journey 4 years ago and for his continuous guidance,

much patience and support. I could have hardly imagined of meeting such a humble

person and getting a better advisor for my PhD studies. I would like to thank Farhan

for his support and guidance from the very initial days I started at KTH. I would

further like to thank my co-advisor Prof. Fredrik Lundell for the learnings that I have

had from the discussions with him. I thank Prof. Eugenia Kumacheva, University of

Toronto (UofT) for accepting me to her group as an International Visiting Graduate

Student and Prof. Lars Berglund (Director of WWSC) for kindly providing me the

financial assistance through WWSC for an international exchange.

A special thanks to my colleagues at WWSC and all members of WWSC Academy

for the wonderful and unforgettable moments spent together during the winter and

summer schools. It is important highlighting the names of Lilian, Nicola, Tobias B.

and Tahani from WWSC and Daniel and Mahshid from UofT for all the nice

discussions, providing a pleasant working environment and always being very

supportive. I would like to thank my colleagues Christophe, Krishne, Ayaka, Tomas,

and Calvin at the Department of Mechanics for sharing their knowledge and

providing the generous support. Thanks to Prof. Lars Wågberg, Prof. My

Hedhammar, Prof. Christofer Lendel, Prof. Benjamin Hsiao and Prof. Nicholas

Kotov for all the great insights and ideas. Thanks to Prof. Stephan Roth for kindly

providing the beamtime at DESY. Thanks to Prof. Henrik Alfredsson at KTH

Mechanics for always being very kind and helping me with the courses. Thanks to

Irene Linares Arregui from the Solid Mechanics Department at KTH for assistance

with tensile test measurements and helpful discussions.

I am much obliged to Prof. Shervin Bagheri, the reviewer of this work as well as

Lilian and Farhan for their valuable suggestions while proof reading this thesis. Most

importantly, I am forever grateful to my family, especially my parents and brother

for their unconditional love, encouragement, constant motivation and support not

only by numerous sacrifices but also by providing me with the best instead of theirs.

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74

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APPENDIX

86

Appendix

Table 1. Comparison of mechanical properties of our CNF fibers (paper I) with the

previously reported studies on the fabrication of nanocellulose based fibers using

different approaches.

Fiber Materials Strength*

(MPa)

Stiffness*

(GPa)

Year

(Ref.)

Cellulose

nanofibers

Cellulose

nanofibers

306

321

15.3

23.6

2016132

2011147

Cellulose

nanofibers

222 12.6 201568

Cellulose

nanocrystal-

cellulose acetate

220 16 2014148

Nanofibrillated

cellulose

323 37.7 2014149

Cellulose

nanofibrils

275 22.5 2011150

Cellulose

nanofibrils

576 18.8 201485

Cellulose

nanofibrils

Cellulose

nanofibrils

297

850

21

53.5

2016151

201734

Cellulose

nanofibrils

1570 86 Present

work

(paper I)

*Strength and stiffness reported here corresponds to the highest values obtained in

the reference studies. Terminologies used for CNFs are the ones reported in the

source articles.

APPENDIX

87

Table 2. Source references for data points used in Figure 4.9 and 4.10.

Fiber Reference

Coir Bledzki et al.152

Cotton Li et al.153

RSP Xia et al154; Elices et al.155; Plaza et al.156

Nylon Guan et al.157

Sisal Inacio et al.158

Jute Paul et al159; Bledzki et al.152

Hemp Li et al.153

Ramie Kalia et al.160

Dragline Liu et al.161; Cunniff et al.162

E-Glass Li et al.153; Wambua et al.163

S-Glass Li et al.153

Metals &

Alloys

http://www.grantadesign.com/download/pdf/teaching_resource_b

ooks/2-Materials-Charts-2010.pdf

*Additional data points are adapted from Ashby and Wegst.164

http://www.grantadesign.com/download/pdf/teaching_resource_books/2-Materials-Charts-2010.pdf

http://www.grantadesign.com/download/pdf/teaching_resource_books/2-Materials-Charts-2010.pdf

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