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Multilayered Scaffolds for Osteochondral Tissue Engineering
Based on Bioactive Glass and Biodegradable Polymers
Mehrlagige Gerüststrukturen für das Knochen-Knorpel Tissue
Engineering basierend auf bioaktivem Glas und bioabbaubaren
Polymeren
Der Technischen Fakultät
der Friedrich-Alexander-Universität
Erlangen-Nürnberg
zur
Erlangung des Doktorgrades DOKTOR – INGENIEUR
vorgelegt von
Patcharakamon Nooeaid
aus Trang
Als Dissertation genehmigt
von der Technischen Fakultät
der Friedrich-Alexander-Universität Erlangen-Nürnberg
Tag der mündlichen Prüfung: 15.05.2014
Vorsitzende des Promotionsorgans: Prof. Dr.-Ing. habil. Marion Merklein
Gutachter: Prof. Dr.-Ing. habil. Aldo R. Boccaccini
Prof. Dr. rer. nat. Peter Greil
CONTENTS
ZUSAMMENFASSUNG ..........................................................................................................vii
CHAPTER 1 Introduction ........................................................................................................ 1
CHAPTER 2 State of the Art and Literature Review ............................................................ 3
2.1 Characteristics of the osteochondral interface .............................................................. 3
2.2 Scaffolds for osteochondral tissue engineering ............................................................. 6
2.2.1 Scaffold materials ................................................................................................. 6
2.2.2 Scaffold fabrication techniques ........................................................................... 16
2.2.3 Strategies of multilayered scaffold ...................................................................... 26
2.3 Cells and bioactive molecules/growth factors for osteochondral tissue engineering .. 35
CHAPTER 3 Objectives and Outline ....................................................................................39
CHAPTER 4 Preparation and Characterization of Biodegradable Polymer Coated 45S5
Bioglass-Based Scaffolds for Subchondral Bone Tissue Engineering Applications ..43
4.1 Introduction .................................................................................................................. 43
4.2 Materials and methods................................................................................................. 44
4.2.1 Fabrication of 45S5 Bioglass®-based scaffolds ................................................. 44
4.2.2 Preparation of biodegradable polymer coated 45S5 Bioglass®-based scaffolds 45
4.2.3 Characterization and mechanical testing ............................................................ 46
4.2.4 Statistical analysis............................................................................................... 48
4.3 Results and discussion ................................................................................................ 48
4.3.1 Morphology ......................................................................................................... 48
4.3.2 Mechanical properties ......................................................................................... 57
4.3.3 Degradation behavior ......................................................................................... 59
4.3.4 In vitro bioactivity ................................................................................................ 62
4.4 Conclusions ................................................................................................................. 70
CHAPTER 5 Development of 45S5 Bioglass®-based Scaffolds for Controlled Antibiotic
Released in Bone Tissue Engineering via Biodegradable Polymer Layered Coating ...71
5.1 Introduction .................................................................................................................. 71
iv
5.2 Materials and methods................................................................................................. 72
5.2.1 Fabrication of TCH-loaded layered biodegradable polymer coated Bioglass®-
based scaffolds ............................................................................................................ 72
5.2.2 Characterization and testing ............................................................................... 73
5.2.3 Statistical analysis............................................................................................... 75
5.3 Results and discussion ................................................................................................ 75
5.3.1 Surface property of polymeric coatings .............................................................. 75
5.3.2 Morphology ......................................................................................................... 78
5.3.3 Mechanical properties ......................................................................................... 79
5.3.4 Chemical structure .............................................................................................. 81
5.3.5 In vitro drug release ............................................................................................ 82
5.4 Conclusions ................................................................................................................. 87
CHAPTER 6 Porous Biodegradable Polymer-based Scaffolds for Cartilage Tissue
Engineering Applications .....................................................................................................89
6.1 Introduction .................................................................................................................. 89
6.2 Materials and methods................................................................................................. 91
6.2.1 Fabrication of Alg-foams ..................................................................................... 91
6.2.2 Fabrication of PLLA fibers and Alg-based fibers ................................................ 92
6.2.3 Characterization and testing ............................................................................... 92
6.2.4 Statistical analysis............................................................................................... 95
6.3 Results and discussion ................................................................................................ 96
6.3.1. Effect of processing conditions on the physical and mechanical properties of the
foams ........................................................................................................................... 96
6.3.2 Effect of electrospinning conditions on the properties of fibers ........................ 108
6.4 Conclusions ............................................................................................................... 121
CHAPTER 7 Multilayered Scaffolds Suitable for Osteochondral Tissue Engineering 123
7.1 Introduction ................................................................................................................ 123
7.2 Materials and methods............................................................................................... 125
7.2.1 Fabrication of multilayered scaffolds ................................................................ 125
v
7.2.2 Characterization and testing ............................................................................. 127
7.3 Results and discussion .............................................................................................. 128
7.3.1 Microstructure ................................................................................................... 128
7.3.2 Interfacial strength of multilayered scaffolds .................................................... 132
7.3.3 Mechanical properties of integrated bilayered scaffolds .................................. 133
7.3.4 In vitro bioactivity .............................................................................................. 134
7.4 Conclusions ............................................................................................................... 139
CHAPTER 8 Biological Response of Osteoblasts Culturing on Bioglass-based
Scaffolds for Bone Regeneration ......................................................................................141
8.1 Introduction ................................................................................................................ 141
8.2 Material and methods ................................................................................................ 142
8.2.1 Fabrication of Bioglass®-based scaffolds ......................................................... 142
8.2.2 In vitro cell culture ............................................................................................. 143
8.2.3 Characterization techniques ............................................................................. 143
8.2.4 Statistical analysis............................................................................................. 146
8.3 Results and discussion .............................................................................................. 146
8.3.1 LDH activity ....................................................................................................... 146
8.3.3 Metabolic activity............................................................................................... 149
8.3.4 Osteoblastic activity .......................................................................................... 150
8.3.5 Cell morphology ................................................................................................ 152
8.4 Conclusions ............................................................................................................... 155
CHAPTER 9 Biological Response of Chondrocytes and Mesenchymal Stem Cells on
Alginate/Chondroitin Sulfate Scaffolds for Cartilage Regeneration ..............................157
9.1 Introduction ................................................................................................................ 157
9.2 Materials and methods............................................................................................... 159
9.2.1 Fabrication of Alg/ChS-foams ........................................................................... 159
9.2.2 Characterization and testing ............................................................................. 159
9.2.3 Release of ChS ................................................................................................. 161
9.2.4 In vitro culturing of primary porcine chondrocytes and human MSCs .............. 162
vi
9.2.5 Statistical analysis............................................................................................. 165
9.3 Results and discussion .............................................................................................. 165
9.3.1 Characterization of Alg/ChS-foams .................................................................. 165
9.3.2 Release profile of ChS molecules..................................................................... 174
9.3.3 The influence of culturing conditions on the primary porcine chondrocytes
activity ........................................................................................................................ 175
9.3.4 The influence of ChS molecules on chondrogenic differentiation of chondrocytes
and MSCs .................................................................................................................. 178
9.3.5 The influence of chondrogenic induction (TGF-1) on the activity of MSCs .... 183
9.4 Conclusions ............................................................................................................... 185
CHAPTER 10 Summary and Future perspectives ...........................................................187
REFERENCES ......................................................................................................................193
LIST OF FIGURES .................................................................................................................... I
LIST OF TABLES ................................................................................................................... XI
ABBREVIATIONS AND SYMBOLS ..................................................................................... XIII
ACKNOWLEDGEMENTS ..................................................................................................... XXI
LIST OF PUBLICATIONS .................................................................................................. XXV
ZUSAMMENFASSUNG
In der vorliegenden Dissertation wurden mehrlagige Scaffolds, die für die
Gewebeentwicklung an Grenzflächen geeignet sind, z. B. Knochen-Knorpel-Regeneration,
hergestellt und im Detail diskutiert. Ihre Bauweise, poröse Struktur, physiochemische und
mechanische, sowie biologische Eingeschaften wurden umfassend betrachtet.
Bioglas®-basierte Schäume wurden als Grundlagematerial der Scaffold für den
unterknorpeligen Knochenteil ausgewählt. 3D Bioglas®-basierte poröse Scaffold, die eine
der Spongiosa ähnliche Bauweise und poröse Struktur aufweisen, wurden mit der
Schaumnachbildungsmethode hergestellt. Außerdem konnte die mechanische Festigkeit
und strukturelle Stabilität der Scaffold durch die Beschichtung mit bioabbaubaren Polymeren
verbessert werden. Es wurden verschiedene Polymerbeschichtungen untersucht, dazu
gehören Alginat (Alg), Gelatine (Gel), PDLLA und PHBHHx Beschichtungen, die im
Vergleich zu unbeschichteten Bioglas®-basierten Scaffolds zu einer Erhöhung des E-Moduls
und der Druckfestigkeit führen. Ergänzend dazu haben solche Scaffolds die Bioaktivität
unterstützt, welche durch die Bildung von HA in der SBF Lösung nachgewiesen wurde.
Demzufolge stellen alle in dieser Arbeit entwickelten Scaffolds geeignete Kandidaten für die
Knochenregeneration dar. Darüber hinaus können die polymerbeschichteten Bioglas®-
basierten Scaffolds als Träger von Medikamenten / Biomolekülen, z.B. Überbringer von
Antibiotika, als Anwendung in der Knochengewebeentwicklung dienen. Multifunktionelle
Scaffolds, die auf auf TCH-beladenen Polymerschichten basieren und mit dem Bioglas-
basiertem Scaffolds beschichtet sind, haben im Vergleich zu unbeschichteten Scaffolds eine
verbesserte mechanische Festigkeit aufgewiesen und eine kontrollierte
Medikamentenausscheidung über 14 Tage nach Eintauchen in PBS. Die biologischen
Eigenschaften von Alg-beschichteten Bioglas®-basierten Scaffolds wurden durch die
Züchtung mit Osteoblasten (MG-63) ausgewertet, mit der Zielsetzung ihre Biokompatibilität
und Fähigkeit zur Knochenmineralisierung zu bestätigen. Im Vergleich zu unbeschichteten
und RGD-modifizierten Alg-beschichteten Bioglas®-basierten Scaffolds, weisen Alg-
viii
beschichtete Scaffolds eine gute Biokompatibilität auf und fördern das Zellwachstum und die
Osteoblastenaktivität.
In Bezug auf die Knorpelphase im Knochen-Knorpel-Gewebe war Alg
höchstinteressant, weil seine chemische Struktur jener von der Hyaluronsäure (HyA) ähnelt,
wenn in Erwägung gezogen wird, dass HyA die Hauptkomponente in knorpligen ECM ist.
Poröse Säulenstrukturen aus 3D-Schäumen wurden erfolgreich hergestellt, um die Migration
und Anordnung von Zellen, und den anschließenden Aufbau von neuem Gewebe durch die
Optimierung der Polymerkonzentration und der Gefriertrocknungsbedingungen zu
unterstützen. Optimierte 3 Gew./Vol.% Alg-Schäume wiesen eine Porengröße im Bereich
von 125 - 325 µm auf, was für die Unterstützung der Besämung und Migration von
Chondrozyten geeignet ist. Ergänzend dazu waren die Schäume in der Lage Wasser in der
gleichen Größenordnung wie der native Knorpel (~ 80 %) zu absorbieren. Die mechanische
Festigkeit und strukturelle Stabilität der Schäume wurde durch den Einsatz ionischer
Vernetzung verbessert. Der E-Modul und die Druckfestigkeit der Schäume lagen
entsprechend bei 0.220 ± 0.009 and 0.14 ± 0.02 MPa, was im Größenbereich des nativen
Knorpels liegt (E-Modul von 0.24 – 0.85 MPa and Druckfestigkeit von 0.01 – 3 MPa).
Wesentlich dabei ist, dass die Schäume nicht mineralisiert waren, was bedeutet, dass sie in
Kontakt mit Körperflüssigkeiten keine Knochenbildung hervorrufen können, was aber für die
Knorpelregeneration notwendig ist. Gemäß des Diagramms in Abbildung 10.1, welches die
entsprechenden Kernpunkte bei der Herangehensweise in der Knorpelgewebeentwicklung
zusammenfasst, haben die in der vorliegenden Arbeit angefertigten porösen Schäume die
meisten gerüstbezogenen Kriterien erfüllt. Allerdings mangelt es den Schäumen an
Zelladhäsion, was sich auf die Vermehrung und Differenzierung von Zellen negativ auswirkt.
Daher wurden die Alg-Schäume erstmals durch die Einbindung von biologischen
Signalstoffen, z. B. Chrondroitinsulfat (ChS), modifiziert, mit dem Ziel die Zelladhäsion und
das Verhalten der Zellen zu verbessern. ChS ist eines von natürlichen
Glycosaminoglycanen im Knorpel, welches die Funktion hat den Metabolismums von
Chondrozyten durch die Induzierung der Synthese vom Typ II Kollagen (Col II) und
Proteglykanen (PGs) zu stimulieren. Die Alg/ChS-Schäume haben entweder die
ix
Chondrozyten oder die MSCs unterstützt und die Zellvermehrung und Zelldifferenzierung.
Die Ausscheidung von Col II und PGs von Chondrozyten- und MSCs, die auf Alg/ChS-
Schäumen besamt wurden, wurden als Marker der Knorpelregeneration charakterisiert.
Diese Ergebnisse haben die wichtigsten zellenbezogenen Anforderungen erfüllt (Abbildung
I). Außerdem haben die Alg/ChS-Schäume, in welchen das ChS als biologischer Signalstoff
dient, assoziiert durch den Einsatz von TGF-β1 eine nennenswerte Steigerung der
Chondrogenese von MSCs begünstigt. Dieses Ergebnis deutete darauf hin, dass die
Einbingung von Biomolekülen (ChS) in Kombination mit Wachstumsförderern (TGF-β1) eine
wichtige Rolle in Hinsicht auf die knorpelige Differenzierung and anschließende
Matrixproduktion spielt. Trotzdem hat das ChS die Zelladhäsion weniger als erwartet
verbessert. Das liegt möglicherweise an der geringen Menge von ChS, welches in die Alg-
Schäume eingebunden wurde. Diese Menge vermag nicht ausreichend genug sein, um von
den Zellen erkannt zu werden. Folglich tendierten die Chondrozyten und MSCs innerhalb
der Poren Klumpen zu bilden, aber haben sich kaum an die Porenwände der Schäume
angehaftet. Deswegen verbleiben einige herausfordernde Fragestellungen hinsichtlich der
drei Ecksteine bei der Vorangehensweise der Gewebeentwicklung offen. Als erstes ist es
notwendig die Schäume zu modifizieren (z. B. durch eine Oberflächenfunktionalisierung),
um die Zelladhäsion zu steigern. Zweitens ist es empfehlenswert die Auswirkung der ChS-
Freisetzungsrate auf die Zellabspaltung und Matrixproduktion weiterhin intensiv zu
untersuchen, in Verbindung mit der Funktion des zugegebenen Wachstumsförderers.
x
Abbildung I Zusammenfassung der wichtigsten herausfordernden Aspekte im Bereich der
Knorpelgewebeentwicklung, die in der vorliegenden Dissertation untersucht wurden, sowie
Andeutung der Kriterien, welche mit den entwickelten Scaffold erfüllt wurden.
xi
Mehrlagige Scaffolds wurden entwickelt, die auf optimierten Scaffold für den
Unterknorpelknochen und Knorpel basieren. Da das ideale Scaffold für die Knochen-
Knorpel-Reparatur noch nicht existiert, gewinnt die Entwicklung von Strategien, welche
längerfristig ein ausgezeichnetes Ergebnis bieten, immer mehr an Aufmerksamkeit und
erhalten einen beträchtlichen Forschungsaufwand. Der Schwerpunkt der vorliegenden Arbeit
lag daher, in Hinsicht auf das Material, auf modernen Strategien der zwei- oder mehrlagige
Scaffolds, einschließlich der integrierten zweilagigen und monolitischen zweiphasigen
Scaffolds, die auf dem Alg-Schaum und Alg-beschichteten Bioglas® Scaffolds basieren.
Obwohl es naheliegend ist, dass integrierte zweilagige Scaffolds aufgrund einer möglichen
Delamination an der Grenzfläche zwischen den Schichten einen Schwachpunkt bieten
können, hat unsere Studie nachgewiesen, dass die Delamination durch das Einfügen einer
adhäsiven Zwischenphase, welche als Grenzfläche zwischen dem ausgeprägten Knorpel
und der Knochenschichten dient, überwunden werden kann. Hinsichtlich der
Herstellungsparameter und der später insbesondere an der Grenzfläche auftretenden
porösen Struktur, konnten im Gegensatz dazu monolithische zweiphasige Scaffolds kaum
kontrolliert werden. Zusätzlich wurde im Rahmen von Grenzflächenuntersuchungen mit Hilfe
von Mikrozugversuchen nachgewiesen, dass die Grenzflächenbruchfestigkeit der
integrierten zweilagigen scaffold höher ist im Vergleich zu jener von monolitischen
zweiphasigen Scaffolds. Daher kann im Bereich der Materialien vorläufig zusammengefasst
werden, dass das in der vorliegenden Arbeit entwickelte integrierte zweilagige
Scaffoldsystem ein geeigneter Ansatz ist, um weiter als Knochen-Knorpel-Konstruktion
entwickelt zu werden.
Zusätzlich wurden mit Hilfe des Elektrospinnens Fasergewebenetze (PLLA und
Alg/Gel) hergestellt und als Knorpelphase in zweilagigen Scaffold untersucht. Die
Fasernetze hatten eine Porengröße von 50 µm (mit einer Dicke von bis zu 500 µm), wobei
mit zunehmender Netzdicke eine Abnahme der Porengröße festgestellt wurde. Die kleine
Porengröße der elektroversponnenen Fasern schränkt bekanntlich die Anwendung dieser
Fasern in GewebeentwicklungsScaffolds ein. Die kleine Porengröße der Fasernetze vermag
nach der Implantation die Zellmigration und den Nährstofftransfer hemmen. Die bedingte
xii
dreidimensionale Struktur könnte für die Regeneration des Neo-Knorpels nicht passend
sein. Die Fasernetze wären dennoch ein interessanter Kandidat für die Anwendung als
kalzifizierte Knorpelschicht an der Knorpel-Knochen-Grenzfläche. Dichte Fasernetze mit
kleiner Porengröße können als Barriere gegen die Infiltration von Knochenzellen vom
Unterknorpel aggieren und in vivo die Gefäßneubildung von der Knorpelphase verhindern.
Daher ist die Strategie der mehrlagige Scaffolds für die Knochen-Knorpel-
Gewebeentwicklung vielversprechend und sollte weiter verfolgt werden, um eine mehr den
nativen Geweben ähnliche biomimetische Struktur zu erreichen.
Wie in Abbildung 10.2 gezeigt wird in der vorliegenden Arbeit ein neuartiges
mehrlagiges Scaffold mit einer anspruchsvolleren Struktur als eine Perspektive aufgezeigt,
mit dem Schwerpunkt auf funktionalen Knochen-Knorpel-Scaffolds. In der gegenwärtigen
Arbeit erfolgte die Gestaltung des Scaffold als eine Kombination von Alg-Schaum, PLLA
Fasern und PDLLA-c-BG Scaffold für Knorpel, Grenzflächen- und Unterknorpel-
Knochenphasen. Kurz dargestellt, wurde die Alg-Lösung auf die PLLA Fasernetze/PDLLa-c-
BG zweilagiges Scaffold angewandt und es kam zur Gelbildung durch den Zusatz von
CaCl2▪H2O. Nach der Gefriertrocknung bildete sich eine poröse Alg-Phase. Diese wurde auf
das Oberteil des Fasernetzes aufgebracht (Abbildung II). In diesem Fall hat das Fasernetz
als intermediäre Schicht zwischen dem Bioglas®-basierten Scaffolds und dem Alg-Schaum
aggiert, welches als Grenzfläche zwischen Knorpel und Unterknorpelknochen gedacht war
und ein dichtes ECM aufwies. Dieses Konzept wurde durch eine frühere Studie von Yunos
et al. [94] inspiriert und unsere in Abschnitt 7 präsentierten Befunde haben aufgezeigt, dass
die Bildung von HA an der Grenzfläche zwischen dem Bioglas®-basierten Scaffold und den
PLLA Fasern die mechanische Stabilität der Grenzfläche weiter verbessern. Es wird
zusätzlich empfohlen osteokonduktive Hybridfasernetze (z. B. PLLA/ Bioglas® Hybridfasern)
als Grenzflächenphase zu verwenden, anstatt von einzelnen Polymerfasern. Dieses Design
imitiert kalzifizierten Knorpel, welcher eine hohe Fähigkeit zur lokalen Mineralisierung an der
Grenzfläche bietet. Anschließend sollte eine starke Grenzfläche gebildet werden, die in der
Lage ist den Knorpel und die knochenähnlichen Schichten während der in vitro Zellzucht
and in vivo Züchtungsbedingungen zu integrieren.
xiii
Es wird erwartet, dass die Entwicklungen der Scaffold und das Wissen, welches
während der vorliegenden Arbeit gewonnen wurde, hilfreich sein werden und in naher
Zukunft zu Fortschritten im Bereich der Knochen-Knorpel-Geweberegeneration führen.
Abbildung II SEM Abbildung, die ein empfohlenes mehrlagiges Gerüstmodel für
Anwendungen in der Knochen-Knorpel-Gewebeentwicklung aufzeigt, inklusive Alg-Schaum
für die Knorpelphase, PLLA Fasernnetz für die kalzifizierte Grenzflächenphase und PDLLA-
c-BG Gerüst für die Knochenphase.
xiv
CHAPTER 1
Introduction
The number of research studies in the field of interface tissue engineering,
especially the cartilage-bone (osteochondral) interface [1,2], is continuously increasing given
the need for treatment large sections of the population worldwide suffering from
osteoarthritis (OA). OA, the degeneration of osteochondral tissue at the joint (Fig. 1.1),
affects around 630 million people worldwide and continues to expand as populations age [4].
Osteochondral repair remains a challenge for surgeons and researchers due to the poor
capacity of self-repair of cartilage and the limitations of present surgical techniques. Indeed
current surgical procedures, including debridement, microfracture, mosaicplasty,
periosteal/perichondrial transplantation and autologous chondrocyte implantation (ACI), are
limited in their long-term benefit and usually require second surgery [1,2]. Therefore,
scaffold-based tissue engineering is a promising approach aimed at supporting the
generation of new tissues in order to fulfill unmet clinical demands. Scaffolds based on
tailored combination of biomaterials were investigated in this study targeting the structure
and properties required for their use in osteochondral regeneration. In general,
osteochondral defects affect both the articular cartilage and the underlying subchondral
bone, which are distinct in compositional, biological, physio-chemical and mechanical
properties [3]. Moreover, cartilage tissue exhibits intrinsic complexity along its distinct four
zones (superficial, middle, deep, calcified zones), with each zone being defined by a
particular composition and organization of cells and extracellular matrix (ECM) [3–7].
Therefore, scaffold design in recent reported research are becoming more sophisticated in
terms of the combination of various biomaterials and fabrication techniques in order to
mimic the specific characteristic features of both tissue types (instead of using a single
biomaterial) [5–7]. It has become apparent that in order to design the appropriate
osteochondral scaffold, it is essential to understand well the anatomical structures and
properties of the native tissues to be regenerated. The anatomical structure and
2
characteristic properties of the tissues to be regenerated will guide the design of the scaffold
and will dictate the specific requirements of ideal scaffolds. The design of the scaffolds will
lead to the selection of suitable biomaterials and fabrication techniques, which will
prominently influence the structural and mechanical properties of the designed scaffolds.
The design of scaffolds and the fabrication techniques were critically considered as a core of
this research. According to the characteriatics of osteochondral scaffold, the research tasks
were devided into two main parts, namely (i) scaffolds for subchondral bone layer and (ii)
scaffolds for articular cartilage layer. Bioactive glass (type 45S5) and alginate are mainly
investigated for used as scaffolds for bone and cartilage, respectively. In addition, an extra
functionality was incorporated by developing an antibiotic drug releasing capability into the
bone scaffold was included. At the same time, biological molecule, i.e. chondroitin sulfate,
was incorporated into the cartilage scaffold in order to enhance cell adhesion, proliferation
and dfiiferentiation. The bi- or multilayered scaffolds was manufactured to replicate the
characteristics of the cartilage-bone tissue interface and the different design strategies were
evaluated and compared. Finally, in vitro cell culture on the scaffolds were reported, in order
to confirm their biological and cellular responses and to assess the relative advantages and
disadvantages of the different concepts proposed, highlighting promising avenues for further
research and the clinical demand.
Figure 1. 1 The scheme of knee osteoarthritis (joint degeneration disease), which is the
result of cartilage wearing out in the load bearing joint (Image courtesy of
Commonsensehealth.com [3]).
CHAPTER 2
State of the Art and Literature Review
2.1 Characteristics of the osteochondral interface
Since the causes leading to osteochondral defect remain elusive, it is necessary to
understand the anatomical structure of the tissues involved in order to gain knowledge about
the mechanisms involved in the disease [8]. Furthermore, the characteristic properties of the
natural tissues (bone and cartilage) are vital aspects to be understood in order to seek a
suitable scaffold for the repair of the defect, according to tissue engineering approach [4,8].
The osteochondral interface involves cartilage and subchondral bone with hyaline cartilage
lying on top of cancellous bone. Cartilage functions as a protector of bone from high
stresses and acts as a reducer of friction at the edge of bone [5], as shown as a common
case of knee articulating joint in Fig. 2.1. Moreover, highly flexible characteristics of articular
cartilage are relevant considering the ability of cartilage to withstand dynamic compressive
loads several times the body weight [6]. However, cartilage has a poor capacity of
regeneration due to its highly organized structure, low chondrocyte numbers, low metabolic
rate, and restricted rate of chondrocytes to divide and migrate due to the cartilage dense
matrix [7]. On the other hand, there are numerous successful reparative approaches
available for bone [9]. Unlike cartilage, bone can be self-repaired involving induction of
vascularity and bone remodeling by osteoblasts and osteoclasts [9]. In bone defect sites,
bone marrow stem cells (BMSCs) can differentiate into bone cells which require the support
of extensive vascularity to provide nutrients and proteins to stimulate bone tissue repair [9].
Bone is a complex tissue consisting of water, type I collagen (Col I) and
hydroxyapatite (HA: Ca10(PO4)6(OH)2) crystals, in which Col I and HA provide the tissue’s
stiffness and compressive strength [10,11]. Subchondral bone (SB) is mainly composed of
bulk of Col I and it has a loosely organized porous structure, these features demonstrate the
characteristics of cancellous bone, as shown in Fig. 2.1 [12]. SB is composed of osteoblasts,
which actively secrete ECM components in order to build up the bone tissue, and
4
osteoclasts which are indicative of bone resorption activities [12]. In order to engineer a
bone scaffold with an attempt to mimic the natural bone, bone tissue generation and
mineralization process in bone must be understood. Bone matrix maturation involves the
expression of alkaline phosphatase (ALP) and non-collagenous proteins (i.e. osteocalcin,
osteopontin and bone sialoprotein) [12–14]. Calcium and phosphate-binding proteins
regulate deposition of minerals by the regulation of the amount and size of HA crystals
formed [13,14]. Collagen (Col), a major component in the ECM, functions as a
microenvironment for apatite nucleation [12–14]. In general, bone can form by two different
pathways, including endochondral ossification for long bone and intramembrane ossification
for flat bone [12]. Both pathways origin from precursor cells, which follow the condensation
of the mesenchyme (cartilaginous template only occurs in the case of endochondral
ossification) and finally bone formation occurs [12]. Bone formation is an ongoing process
that alters the size and shape of bone by partial resorption of preformed bone tissue
(modeling) and simultaneous deposition of new bone (remodeling) [12]. When bone is
broken, inflammation occurs; blood is supplied to the channels of the broken area causing
swelling and bruising, which is known as hematoma [12]. Dead cells then release cytokines,
which initiate the healing process [12,15]. In concert, osteoclasts remove the dead cells and
fibroblasts form fibrocartilage as spongy material, which is called soft callus formation stage
[12,15]. Afterwards, the hard callus formation stage starts and the soft callus (cartilage)
transforms into woven bone. This stage is guided by the release of minerals such as calcium
and phosphate into the cartilage tissue [15].
Cartilage is composed of four zones, which each zone has different organizations of
chondrocytes and different orientation of Col fibrils (Fig. 2.1) [12,16–19]. First, superficial
zone, in contacting with superficial fluid, is composed of flattened chondrocytes and Col
fibrils, which Col fibrils are parallelly aligned to the articular surface. Second, middle zone
contains rounded chondrocytes and randomly aligned Col fibrils. Third, deep zone is
composed of vertical columns of chondrocytes, while Col fibrils align perpendicularly to the
articular surface. Finally, calcified zone has specific hypertrophic chondrocytes with low cell
density. The hypertrophic chondrocytes have unique ability to synthesize type X collagen
5
(Col X) and calcified ECM [12,16–19]. The thin layer between non-calcified cartilage and
calcified cartilage is called tidemark, which is believed to act as nutrient diffusion through the
cartilage structure and serve as an attachment of the Col fibrils [20]. Moreover, the series of
interdigitation, which connect the calcified cartilage with SB, support the transformation of
shear stress from the articulation into tensile and compressive stresses [20]. Type II collagen
(Col II) fibrils (up to 60 % dry weight of cartilage) provide high tensile strength and withstand
shear stresses [6]. Moreover, proteoglycans (PGs) embedded within Col II fibrils ( 35 % dry
weight of cartilage) provide the ability to withstand high compressive stress.
Figure 2. 1 Anatomy of the knee joint, which is the most common case found in joint
degeneration disease (according to [21]), demonstrating arrangement of ECM and
organization of chondrocytes along different zones in cartilage (Reproduced from Nooeaid et
al. [19] with the permission of John Wiley and Sons) and showing the structure of cancellous
bone as subchondral bone (Reproduced from Meyer et al. [12] with the permission of
SPRINGER-VERLAG BERLIN/HEIDELBERG).
6
2.2 Scaffolds for osteochondral tissue engineering
The design of scaffolds for osteochondral regeneration is based on the
consideration of physical and biochemical properties of the cartilage-bone interface [16].
Biochemical properties, including chemical composition, and 3D structure of scaffold
material, mostly affect the cellular behavior [16]. Physical properties, including structural
architecture and geometry, biodegradation behavior and mechanical properties, influence
both cellular activity and mechanical stability [16,22]. In regard to the zonal distinct layers of
osteochondral tissue with complex compositions and property variation, the recent trends of
osteochondral scaffolds are based on bi- or multi-layered structures [5,19,23,24]. Each layer
is customized to closely replicate the features of the specific tissues to be regenerated [24].
Singular scaffold materials, on the other hand, for example single phase either bioceramic or
polymer, have not been reported to successfully support the regeneration of osteochondral
tissue. The reason is that the single phase scaffolds lack the inherent physical structure
required and individual materials cannot achieve all requirements for osteochondral
regeneration [25,26]. The development of osteochondral scaffolds must therefore focus on
the use of composite-based structure based on multilayered and gradient structures
[6,20,27].
2.2.1 Scaffold materials
The current selection of the scaffold material is one of the main factors to be
considered in scaffold-based tissue engineering. The material needs to match the criteria of
the tissues to be regenerated or repaired. In general, suitable materials should primarily
exhibit biocompatibility, controlled biodegradability and sufficient mechanical strength
[28,29]. It should also provide a desirable environment for cell attachment, proliferation and
differentiation [28,29], which are cell functions mainly altered by the intrinsic properties of the
material. In addition, the selected biomaterial should be able to be processed economically
into desired shapes and dimensions [29]. Promising osteochondral scaffold materials will be
based on two different materials, including bioceramics and bioactive glasses, and
biodegradable polymers.
7
Bioceramics and bioactive glasses are widely used as the potential candidates for
bone tissue engineering applications [30,31], since bioceramics and bioactive glasses (i.e.
synthetic HA, calcium phosphate (CaP) and Bioglass®) are bioactive, leading to bone-like
apatite formation after immersion in body fluids [30,31]. Moreover, these artificial substrates
can bond to natural bone after implantation [28,32,33]. For instance, CaP with different Ca/P
ratio such as HA, tricalcium phosphate (TCP) and biphasic calcium phosphate (BCP)
provide excellent physical properties in terms of stability, degradation rate and processibility.
Moreover, CaP shows biocompatibility and bioactivity, which is the ability to bind to bone by
the release of Ca and P ions, and they also enhance bone tissue formation. However, CaP
has low mechanical strength and brittleness, which limits their application in load bearing
devices. Bioactive glasses are experiencing increasing research efforts due to their high
bioactivity and for having significantly higher mechanical strength in comparison to most
CaP ceramics. According to the literature, bioactive glasses containing Ca- or Si-based are
the most promising for bone scaffolds due to their ions-releasing ability, high
osteoconductive and osteoinductive properties, and hydrophilic behavior [34,35]. On the
other hand, polymer-based scaffolds are not suitable for bone repair because polymers do
not exhibit osteoconductive properties as bioceramics and bioactive glasses do. In addition,
the stiffness and fracture strength of polymers are insufficient for applications in bone tissue
regeneration, compared to those of bioceramics [36].
2.2.1.1 Bioglass® and its composites
with biopolymers as bone scaffolds
45S5 Bioglass® (45 % SiO2, 24.5 % CaO, 24.5 % Na2O and 6 % P2O5 by wt.)
discovered by Hench in 1971 [31], provides excellent osteoconductivity, osteoinductivity,
bioactivity, controlled degradability and ability to deliver cells for bone tissue regeneration
[31,33,34]. 45S5 Bioglass is also able to form interfacial bonding to soft and hard tissues
[31,37]. Especially in bone bonding, 45S5 Bioglass® can bond to bone in vitro, which has
been described by the formation of a dual layer, including a silica rich layer and
hydroxycarbonate apatite (HCA) layer, on its surface in contact with body fluids [31,38]. The
formation of HCA in vitro is suggested to occur also in vivo leading to bone bonding
[31,33,39]. Hench et al. [40] have described the bonding mechanisms at the interface of
8
Bioglass® substrate and body fluids via HCA formation. The reactions between Bioglass
®
surface and body fluids can be summarized into 5 stages, as described in Table 2.1 (Stage
1-5) [29]. Afterwards, the HCA layer on the Bioglass® surface supports the cellular reactions
in order to form bone (Stage 6-12). At the same time, dissolution products from Bioglass®
surface can up-regulate gene expression of osteoblasts, which this phenomenon controls
osteogenesis, leading to faster bone formation in comparison with synthetic HA [34,35].
Moreover, it has been reported that the ionic dissolution products of 45S5 Bioglass® may
induce an expression of osteoblast genes (i.e. insulin-like growth factor-2 (IGF-2)), leading to
increased cell proliferation and stimulated new bone formation [31,41]. By this fact, 45S5
Bioglass does not only exhibit osteoconductive property, but it is also an active stimulation
of osteoblasts. For instance, alginate/45S5 Bioglass composite scaffolds exhibited better
osteoblastic differentiation of osteosarcoma (MG-63) cells compared to pure alginate
scaffolds because the dissolution products of Bioglass can stimulate osteoblast proliferation
and differentiation, as evidenced by increased osteoblast markers (i.e. ALP, osteocalcin and
osteopontin) [42]. Similarly, El-Gendy et al. [43] have confirmed the osteoblastic
differentiation of human dental pulp stromal cells cultured on 3D porous 45S5 Bioglass®-
based scaffolds.The Bioglass composition and the dissolution products are the key factor
that enables the osteostimulation of osteoblasts and promotes proliferation and
differentiation of the bone cells. Price et al. [44] presented that the surface of Bioglass®
efficiently supports osteoblasts (MG-63 cells) proliferation and their functions in comparison
with the surface of titanium and cobalt chrome. This phenomenon was suggested by the
reason of different surface chemistry (related to chemical composition) of the materials. This
property affects protein adsorption and cell adhesion. In addition, since Bioglass exhibits
fast biological response in culture medium, the ion exchange takes place and induces
alkalinization of the medium, which this phenomenon has an influence on cell metabolism
[45].
9
Table 2. 1 Mechanisms of bioactivity and bone bonding of Bioglass®, according to
[29,31,34,39,46].
Stages Surface reactions
1 Exchange of Na+ and K
+ with H
+ and H3O
+ from body fluids, leading to hydrolysis of silica
groups and formation of silanol (Si-OH) groups: Si-O-Na+ + H
+ Si-OH + Na
+
2 Network dissolution of silica (SiO2) in the form of silicic acid (Si(OH)4) and continued
formation of Si-OH groups: Si-O-Si + H2O 2Si-OH
3 Condensation and polymerization of silica-gel on glass surface: Si-OH + Si-OH Si-O-Si
4 Further dissolution of glass: chemisorption of amorphous Ca2+
, PO43-
ions from the glass
and the solution through silica-gel, leading to formation of amorphous CaP on the surface
of silica-gel
5 Crystallization of HCA layer: continued dissolution of glass, amorphous CaP incorporates
OH- and CO3
2- from the solution and crystallizes as HCA
6 Adsorption of growth factors on HCA layer
7 Action of macrophages
8 Attachment of osteoprogenitor cells
9 Proliferation and differentiation of osteoprogenitor cells
10 Generation of ECM by osteoblasts
11 Crystallization of ECM, forming nanocrystalline mineral and collagen on the surface of
glass
12 Proliferation of bone
Even though Bioglass has excellent bioactivity and osteoconductivity as required
for bone regeneration, in form of porous structure, Bioglass has limitations such as low
mechanical strength and intrinsic brittleness [29,33,47]. Therefore, Bioglass®-based
scaffolds for load bearing applications are mostly fabricated in composite-form by a
combination with biodegradable polymers [28,48,49]. As reviewed by Chen et al. [28],
Rezwan et al. [48] and Chen et al. [49], biodegradable polymers have been incorporated into
bioactive glass-based scaffolds in order to enhance the mechanical integrity and flexibility in
dynamic environments of injured bone. Aliphatic polyesters such as poly(L-lactide) (PLLA),
poly(D-lactide) (PDLA), poly(D, L-lactide) (PDLLA), poly(lactic-c-glycolic acid) (PLGA),
polycaprolactone (PCL) and poly(hydroxyalcanoate) (PHAs) family, and natural
10
biodegradable polymers such as chitosan (CS), gelatin (Gel), collagen (Col) and alginate
(Alg) coated Bioglass®-based composite scaffolds have been reported by Boccaccini and
co-workers [50–57]. It has been established that polymer coated Bioglass-based scaffolds
exhibit significant improvement of the mechanical properties [50–57]. In addition, acidic
degradation products from synthetic biodegradable polyesters can be buffered by the
dissolution products from Bioglass®-based scaffolds [28,48,49]. On the other hand, natural-
derived biodegradable polymers provide some additional benefits such as excellent
biocompatibility, non-toxicity and ability to favor cell interaction [58].
Alginate (Alg) is a polysaccharide-based polymer obtained from marine brown
algae [59,60]. Alg is an unbranched binary copolymer consisting of (1→4) linked β-D-
mannuronic acid (M) and -L-guluronic acid (G) residues of varying composition and
sequence [61]. Alg provides useful properties for tissue engineering applications, it is
biocompatible, biodegradable, non-immunogenic, low-toxic, abundant in sources and it can
be obtained at low prices [60]. Alg can be mostly processed in the form of cation-crosslinked
Alg-gel beads/capsules to encapsulate living cells serving as a cell delivery vehicle in vivo
[62–65]. In this study, Alg was chosen as one of the suitable polymer coatings for Bioglass®-
based scaffolds. As indicated above, polymer coating was proposed to improve the
mechanical strength and fracture toughness of Bioglass®-based scaffolds. As presented by
Erol et al. [54], homogeneous Alg coating on Bioglass-based scaffolds can improve the
mechanical properties, while maintaining scaffold bioactivity. Moreover, Alg, which is
compatible for a variety of cells such as chondrocytes and osteoblasts, has been shown to
maintain the phenotype either of seeded cells or encapsulated cells [62–66]. This
characteristic is crucial for specific tissue regeneration, in particular cartilage regeneration, in
order to avoid de-differentiation of cells and subsequently to avoid the formation of
unspecified tissue like fibrocartilage [67,68]. However, Alg has no adhesive sites to cells and
does not adsorb serum proteins due to its high hydrophilicity [61,69]. Therefore, peptides
with a cell adhesive sequence-modified Alg (i.e. Arg-Gly-Asp (RGD) containing peptide)
(RGD-Alg) have been used to enhance cell adhesion on Alg [61,70]. Since amino acid
sequence RGD in fibronectin acts as a primary cell attachment cue, it has been
11
demonstrated that RGD linear peptide coupling to Alg can enhance osteoblasts adhesion for
4 times when compared to unmodified Alg [71]. In addition, in the form of hydrogels, it has
been previously reported that MC3T3-E1 cells seeded into RGD-modified Alg hydrogels
promoted higher osteoblastic differentiation and mineralization compared to MC3T3-E1 cells
seeded into unmodified hydrogels, as evidenced by higher ALP activity and osteocalcin level
[72].
Gelatin (Gel) is a biomacromolecule derived from Col which is the most abundant
protein in the ECM of connective tissue, such as skin, bone and cartilage [73–75]. Recently,
Gel is used in both soft and hard tissue engineering applications, i.e. in forms of
microcapsule, microsphere, wound dressing and scaffold [73]. The use of Gel-based
scaffolds is an appealing approach in bone tissue engineering due to Gel’s biodegradability,
biocompatibility, non-immunogenic properties and relatively low cost [59,76]. When
compared with Col, Gel does not exhibit antigenicity under physiological conditions [74].
However, the poor mechanical properties and water sensitivity of Gel limit its use to non-
load-bearing applications only [59]. In order to overcome these limitations, Gel has been
used either in combination with synthetic polymers or with the application of chemical
crosslinking, leading to the improvement of the thermal and mechanical properties, and to
increased water resistance [77,78]. In the present work, Gel is one of the degradable
polymers used as a polymer coating on Bioglass-based scaffolds, which was aimed to
improve the mechanical strength of the scaffolds. Metze et al. [53], Erol et al. [74] and
Desimone et al. [79] have presented significantly improved compressive strength and
fracture toughness of bioactive glass-based scaffolds by Gel coating. In addition, Gel coating
does not induce negative effects on bioactivity of bioactive glasses. This result was
confirmed by the formation of HCA after immersion in simulated body fluid (SBF) [53] . In
addition to providing a benefit in the mechanical properties, Gel coated scaffolds (i.e. Gel
coated TCP scaffolds) were reported to support MC3T3-E1 cell adhesion and subsequently
to promote cell proliferation and differentiation compared to uncoated scaffolds, as
presented by Kim et al. [80].
12
Poly(lactic acid) (PLA) is an aliphatic polyester derived from renewable resources
[81]. PLA is a semi-crystalline polymer exhibiting high tensile strength and elongation
compared to natural polymers [60,81]. This character of PLA makes it suitable for low load
bearing applications [60]. The linear structure of PLA has methyl (-CH3) side groups, leading
to a hydrophobic feature and an ability to be soluble in organic solvents (such as chloroform,
dimethylene chloride (DMC), dichloromethane (DCM), methanol (MeOH), ethanol (EtOH),
benzene, acetone, etc.) [81,82]. PLA can be degraded by the mechanism of homogeneous
hydrolysis erosion [83], leading to lactic acid obtained as a degradation product. The
degradation product helps to reduce the pH of the environment and induce further
degradation [83,84]. The physical properties and degradability of PLA depend on the
racemization of D- and L-isomers. Semi-crystalline PLLA is synthesized from L-lactide, while
amorphous PDLLA is obtained from DL-lactide [81]. As a result, PLLA and PDLLA exhibit
different mechanical properties and degradation rate. PLA is one of the most widely used
polymers in tissue engineering applications due to its biocompatibility, biodegradable control
and suitable mechanical properties [52,56,60,83,84]. Moreover, PLA can be processed
readily and reproducibly [60]. In terms of composite scaffolds, polyesters have been used as
polymer coatings on porous bioceramic and bioactive glass-based scaffolds. For example, it
was shown that PDLLA coating reduced the brittleness of ceramic scaffolds, as presented
by Yunos et al. [52,85,86], Chen et al. [56], Bretcanu et al. [55] and Novak et al. [87].
Moreover, biocompatibility of either PLLA or PDLLA coated Bioglass®-based scaffolds has
been confirmed by in vitro culturing with human osteosarcoma cell line (HOS-TE85) [55], for
example. It has been found that the polymer coating, scaffold microstructure and surface
roughness influenced the cell behavior. In addition, cell differentiation has been confirmed by
culturing mesenchymal stem cells (MSCs) on PDLLA/Bioglass composite scaffolds [88].
Poly(3-hydroxybutyrate-co-3-hydroxyhexanoate (PHBHHx) is a member of PHA
biopolyester family [89]. PHBHHx has higher elastomeric mechanical properties compared
to poly(3-hydroxybutyrate) (PHB) and poly(3-hydroxybutyrate-co-valerate) PHBV [90,91].
PHBHHx, which is synthesized by microorganisms, is a copolymer of hydroxyl butyrate (HB)
and hydroxyl hexanoate (HH) with the adjustable content of HH represented by ‘x’ [92].
13
Increased content of HHx leads to reduced crystallinity and subsequently the tensile strength
and the elongation at break increased compared to PHB [89,90]. Thus molecular weight
(Mw) and chemical composition of PHBHHx can be tailored to meet the physical properties
required for various tissue-engineered scaffolds [89–91]. Recently, PHBHHx has become a
promising candidate for tissue engineering scaffolds due to good mechanical properties, low
toxicity, biodegradability and biocompatibility with various cell types, i.e. fibroblasts,
osteoblasts, chondrocytes and MSCs [91]. The variety of PHBHHx or PHAs used in
biomedical applications has been reviewed by Chen et al. [92].
2.2.1.2 Composite-based scaffolds as controlled drug-delivery systems
In general, scaffolds are used as a template able to support the growth and repair of
tissues. In recent research, the scaffolds are being enhanced to form multifunctional
systems, which are able to combine tissue regeneration and local drug delivery [93,94].
According to a convenient type of composite scaffolds (biodegradable polymer coated
bioactive glass scaffolds) developed for bone tissue engineering, the polymer coating layer
can act as a carrier of bioactive molecules such as drugs and growth factors [95,96]. At the
same time, such polymer coating can improve the mechanical properties of porous bioactive
glass scaffolds [52,57]. As reported by Yaylaoglu et al. [97], CaP/Gel scaffolds have been
loaded with gentamicin for in-situ drug delivery combined with tissue engineering.
Continuous release of the drug upon 4 weeks in vivo was observed with the release rate
depending on the degradation rate of the Gel component. Kim et al. [98] developed HA-
based scaffolds with controlled tetracycline release function by using PCL/HA hybrid coating.
The scaffolds presented improved mechanical properties due to the presence of PCL hybrid
coating, while the drug entrapped in the polymeric coating exhibited a sustained release
profile. Moreover, improved mechanical properties and sustained drug release function have
been confirmed by developing vancomycin-loaded PHBV coated 45S5 Bioglass-based
scaffolds, as reported by Li et al. [57]. The coated scaffolds provided a lower initial burst
release when compared to the drug release of uncoated scaffolds. In addition, a controlled
drug release over 6 days in phosphate buffer saline (PBS) was measured. Francis et al. [95]
have reported gentamicin-loaded PHB microsphere coated 45S5 Bioglass-based scaffolds,
14
which not only presented controlled drug release, but also maintained the bioactivity of
Bioglass scaffolds. Similarly, multifunctional scaffolds based on vancomycin-loaded poly(n-
isopropylacryliamide-c-acrylic acid) microgels dispersed in PLGA coated 45S5 Bioglass-
based scaffolds (Olalde et al. [96]) exhibited improved mechanical properties and
maintained bioactivity. In addition, they exhibited controlled release rate from the drug-
loaded microgels. The polymer coatings protect drug molecules from the aqueous
environment and inhibit the fast dissolution of drugs, subsequently the slow release is
achieved [99]. In this case, the dissolution of the drug is caused by the degradation of the
polymer carrier associated with the diffusion of the drug through voids in the carrier [99].
2.2.1.3 Biodegradable polymers as cartilage scaffold
To engineer cartilage-like tissue that mimics the complex and unique structure of
natural cartilage, the focus here is the design of scaffolds with chondroinductive and
chondroconductive properties [100–102]. Recently, it has been shown that cartilage
scaffolds with rather sophisticated 3D architecture can be developed for example starting
from fibrin and agarose-based materials [24]. The ideal scaffold for cartilage tissue
engineering should be biocompatible, biodegradable and show sufficient mechanical
properties in order to resist mechanical forces. Moreover, it should exhibit appropriate
structural and geometrical properties for supporting cell proliferation and differentiation. The
dedifferentiation of cells must be avoided. In addition to these requirements, cartilage
scaffolds should achieve the tissue-like elastic properties, which can tolerate shock
absorption and deformation [24,103].
Biodegradable polymers are widely used in cartilage regeneration, since they can be
fabricated in the forms of hydrogels, porous foams and fibers, which are suitable structures
for scaffolds [104]. Rationale of using polymers as a cartilage scaffold is their intrinsic
elasticity, controlled degradability and sufficient mechanical strength close to the physical
characteristics of native cartilage [105]. In addition, polymers can be customized in terms of
their physical properties by the regulation of Mw and crystallinity [81]. In terms of chemical
design, current research trends focus mostly on natural polymers, i.e. Col and hyaluronan
(HyA), considering that both are components of the cartilaginous ECM [105,106]. Even
15
though Col and HyA-based scaffolds have been investigated for new cartilage generation,
they exhibit drawbacks concerning the mechanical properties and cost. Thus the
development of alternative, cost-effective materials are of great current interest. For
instance, Gel and Alg exhibit chemical structures similar to Col and HyA, respectively, and
both polymers are inexpensive [107]. Nevertheless, the physical and chemical crosslinking is
crucial for natural polymer scaffolds because they are not mechanically stable in aqueous
environments [65,108,109]. Consequently, biodegradable synthetic polymers represent
another group of materials of choice for cartilage scaffolds. Synthetic biodegradable
polymers are beneficial in terms of mechanical properties, controlled biodegradability and
processibility [104]. Regarding biomimetic approaches to tailor chemical composition and
mechanical stability, researches in last decade have focused on the combination of distinct
polymers, including blending of natural polymers [110–116], synthetic polymers [117,118],
and natural and synthetic polymers [58,119–123].
Alg is a highly interesting polysaccharide, which is widely used as a scaffold in
cartilage regeneration. Alg has a chemical structure similar to HyA and it is cost-effective
compared to HyA [124]. In vitro and in vivo studies [59,69,73,109,125–129] have shown that
Alg is able to support the viability, maintaining the round phenotype of chondrocytes and
promoting the formation of Col II and glycosaminoglycans (GAGs). The production of Col II
and GAGs is an indication of cartilage regeneration [130,131]. In terms of manufacturing, Alg
scaffolds can be easily fabricated via mild gelation via interaction with cations (Ca2+
, Cu2+
,
Zn2+
, Sr2+
and Fe2+
), according to the so-called egg-box mechanism (Fig. 2.2) [109,127,128].
Such crosslinking process via ionic interaction involving anionic chains of Alg and cations
leads to the formation of water-insoluble Ca-Alg gels [109,127,128]. The Alg-gels can be
used as an encapsulation device for living cells [127,132,133]. Since the released cations
exchange with Na+ in the culture medium and in body fluids, ionic crosslinking has been
proved to be non-toxic in vitro and after implantation [125]. Moreover, the Ca-Alg gels can be
transformed to a foam-like structure by the application of a lyophilization process
[63,128,134], which will be detailed in a later section. According to the characteristic
properties stated above, Alg was focused in the present wotk.
16
Figure 2. 2 Schematic diagram showing the gelation-mechanism of alginate and calcium
cations by the formation of egg-box structure (Image courtesy K. Kashima and M. Imai
[135]).
2.2.2 Scaffold fabrication techniques
The scaffold fabrication technique is another important factor, which affects the
structural architecture and geometry of scaffolds. According to physical considerations for
suitable osteochondral scaffolds, pore size, porosity and interconnectivity are crucial for
supporting cell migration and tissue regeneration. In order to fabricate porous structures
suitable for bone- and cartilage-repair, targeted porosity and pore sizes of scaffolds, specific
for bone and cartilage regeneration, must be taken into account. The appropriate pore size
of bone scaffold is considered to be in the range of 100 - 600 µm, which has been confirmed
to be sufficient for cell migration, nutrient and waste transportation, vascularization, and
tissue ingrowth [136]. In contrast, vascularization does not occur in cartilage, which is
composed of dense connective ECM, thus pore sizes around 50 - 300 µm are sufficient for
chondrocytes proliferation and ECM secretion [6,137]. This requirement is attributed to the
fact that chondrocytes show high tendency of differentiation when the pore size is around 30
times the cell diameter (diameter of chondrocytes 10 - 15 µm) [6]. Highly porous scaffolds
17
allow for more cell attachment and consequently result in more tissue formation compared to
less porous scaffolds, which is linked to a greater transportation of nutrient and metabolic
waste products [6,73]. Microstructures exhibiting larger surface area are additionally
required for supporting cell attachment and ECM regeneration [103]. A high pore
interconnectivities are required for homogeneous cell seeding, which influences the quality
of the formed tissue [6].
The mechanical properties of scaffolds are influenced by their microstructure as well
as by the intrinsic properties of the material used [6]. As porosity compromises the
mechanical strength of scaffolds, increasing porosity leads to the reduction of strength [6].
Thus, the porous structure of scaffolds should be tailored to achieve sufficient porosity and
pore interconnectivity with the maintenance of suitable mechanical properties. The
mechanical properties of scaffolds should match those of native osteochondral tissues in
order to withstand local loads in the joint by in vivo studies [6]. The mechanical properties of
natural human cartilage-bone tissues are summarized in Table 2.2. In general, cartilage has
the function to transform compressive forces into tension mode and to further transfer loads
to the underlying SB [138]. At the same time, cartilage is able to withstand shear forces by
supplying a intrinsic low friction-surface [8,9,118,139]. On the other hand, the underlying SB
mainly supports compressive and tension loads [9,140].
Table 2. 2 Mechanical properties of natural healthy human osteochondral tissues
[6,12,73,76,141–143].
Mechanical properties (MPa) Articular cartilage Subchondral bone
Compressive modulus
Compressive strength
Young’s modulus (Tension)
Ultimate Tensile strength
Shear modulus
0.24 - 0.85
0.01 - 3
5 - 25
3.7 - 10.5
0.2 - 2
0.05 - 0.6
2 - 12
445
3 - 20
No report
18
Importantly, the biodegradability of scaffold materials influences the formation and
functionality of new tissues [126]. An appropriate scaffold should exhibit degradation rate
matching the formation of new tissues and must maintain the structural stability until the new
tissue fully assumes the load-bearing function [9,144]. The degradation rate of the scaffold
can be altered by the variation of the material used, namely composition, chemistry and
porous structure [22,105]. In particular, the architecture and topography of scaffolds, which
greatly affected cell attachment, proliferation and differentiation, are partly influenced by the
fabrication techniques. The appropriate fabrication technique needs to be able to generate a
porous scaffold with reproducible architecture and provide mechanical functions for load-
bearing environment [104]. In order to fabricate scaffolds for osteochondral repair, the
combination of different fabrication techniques is crucial for the achievement of sophisticate
structures such as multilayered scaffolds [50,145]. Currently, the variety of fabrication
techniques available, including solvent casting and particle leaching, melt molding, freeze-
drying, thermal induced phase-separation, electrospinning and rapid prototyping techniques
[6,146], are all considered in the fabrication of polymer-based scaffolds. On the other hand,
bioceramic/bioactive glass-based scaffolds are frequently fabricated by foam-replication,
rapid prototyping, fused deposition remodeling, robocasting, stereolithography and 3D-
printing [146]. Rapid prototyping can be used to fabricate both polymer- and ceramics-based
scaffolds [146]. The advantages and disadvantages of current fabrication techniques applied
for manufacturing both bone and cartilage scaffolds are summarized in Table 2.3.
Particularly, freeze-drying, electrospinning and foam-replica techniques, focusing on recent
work, will be discussed. Foam-replication, freeze-drying and electrospinning techniques
were chosen in this study according to the required physical properties, such as porosity,
pore size, architecture and geometry of the scaffolds, and mechanical properties of the
specific tissues. In addition, all techniques are simple cost-effective.
19
Table 2. 3 Current 3D scaffold fabrication techniques for polymers and ceramics.
Fabrication techniques Pros(+)/Cons(-) References
Polymeric scaffolds
Solvent casting/particle
leaching
+ Pore size 30 - 300 µm
+ Controlled pore sizes by particle size of salt/porogen
- Limit for thin membrane with thin wall section
- Low porosity 20 - 50 % and insufficient pore
interconnectivity
- Required toxic solvents
- Remained salt particles in matrix
- Time consumer
[19,58,119,123
,147,148]
Melt molding + Solvent-free method
+ Pore size 50 - 500 µm
+ Controlled macropore geometry
- Porosity 80 %
- Suitable for thermoplastics
[19,119,149]
Freeze-drying + Can be incorporated in conjunction with thermal
induced phase separation
+ Controlled pore size and pore orientation
+ Porosity 90% and pore size 50 - 400 µm
+ High pore interconnectivity
[110,112,114,1
16,150–156]
Thermal induced phase-
separation (TIP)
+ Extensively applied in the fabrication of
microspheres for drug-delivery system
+ Suitable to fabricate porous polymer/ceramic
composite-based scaffolds
+ High porosity 97 %
+ Pore size 200 µm
+ Obtained high volume of interconnected micro-
pore structure
[51,93,157]
Electrospinning + Can be used to fabricated hybrid fibers (organic-
inorganic mixture)
+ High porosity
- Small pore size
- Limit designed architecture – needs post-
[86,117,158–
174]
20
Fabrication techniques Pros(+)/Cons(-) References
fabrication techniques
- Insufficient mechanical properties
Rapid prototyping/solid free
form (SFF)
+ Manufactured by computer - generated design
+ Optimized microstructure and mechanical
functions
- Low porosity 60 %
[138,175,176]
Fused deposition
remodeling (FDM)
+ Reproducibility
+ Controlled structure by computer-controlled
method
[102,177]
3D printing + Precise deposition of cells and matrix in layer-by-
layer fashion
+ Highly reproducible architecture
+ Easily tailored porosity
- Limits for used in load-bearing applications,
especially in the cases of natural-derived
polymeric matrices.
[27,178]
Ceramic-based scaffolds
Replication + Conventional technique
+ High porosity 90 %, pore size 100 - 700 µm
+ Achieved architecture similar to that of cancellous
bone
[32,85,150,179
,180]
Rapid prototyping/solid free
form
+ Controlled architecture
+ Designed scaffolds can be fit on the defect site
[138,175,181]
Stereolithography
+ Versatile with respect to the freedom of design
and scale (submicrons-decimeters)
+ Manufactured in layer-by-layer fashion by
computer- controlled method
+ Fabricated gradient scaffolds in porosity and pore
size
[6,182]
3D printing/ robocasting + Achieved thick struts
+ Pore size 500 µm
+ Sufficient compressive strength
- Porosity 60 %
[27,33,93,183]
21
2.2.2.1 Foam replication technique
The foam fabrication technique was first developed in 1963 for ceramic foam
manufacturing [28]. This technique (Fig. 2.3) involves the production of ceramic foams by
coating a polymer template (i.e. polyurethane (PU) foam) with a ceramic slurry (ceramic
powder/binder/water mixture). Then the sacrificed template is burnt out and the ceramic particles
are sintered by using proper heat treatment. As-sintered foams exhibit high porosity ( 80 %)
and pore size in hundreds microns, depending on the pore size of the used polymer template.
However, highly porous scaffolds are obtained with relatively low mechanical properties,
exacerbated by the intrinsic brittleness of ceramics, which are difficult to handle [28,184].
The foam replication method to fabricate Bioglass-based scaffolds was patented by
Boccaccini group at Imperial College London in 2006 [185]. The technique is currently widely
used to fabricate bioactive glass-based scaffolds in the field of bone tissue engineering. It has
been proved that 45S5 Bioglass-based scaffolds, for example, supported osteoblasts activities
[28]. Cells migrated efficiently and proliferated into entire porous structure [28]. In addition, the
low mechanical properties of Bioglass-based scaffold can be overcome by the incorporation of
polymer phases, forming composite-based scaffolds [48], as mentioned previously. Compared to
other fabrication techniques, such as rapid prototyping, stereolithography, etc., the foam
replication technique is more cost-effective and less time-consuming necessitating simple
equipment [28].
22
Figure 2. 3 Schematic diagram of the foam replication technique employed to produce 3D
porous bioceramics- and bioactive glass-based scaffolds (according to [28,184]).
2.2.2.2 Freeze-drying technique
Freeze-drying technique is an attractive dehydration method, known as lyophilization
[186]. It is well known in the food industry [187] and it has become widely used in biomedical
applications, in particular scaffolds for tissue regeneration [188]. It basically works by
freezing the solution at a temperature below the freezing point of the solvent used following
by the reduction of the surrounding pressure below atmosphere pressure to allow the frozen
solvent in the material bulk to sublimate directly from solid phase to gas phase [105]. The
basic principle of freeze-drying process can be explained with reference to a simple water
phase diagram, as demonstrated in Fig. 2.4. The process basically consists of three stages.
The first freezing stage involves a fast decrease of material temperature at temperature
underneath the freezing point (TC). The next stage is drying the material by heating below
the triple point (TA) and under vacuum conditions (below PA) to force sublimation, leading to
the formation of an interconnected pore structure. After this stage, the water ( 7 - 8 %) still
bound to the porous material can be desorbed by increasing temperature [189]. Pore size
and orientation of pores are mainly influenced by the freezing temperature [189–191]. If the
freezing temperature is lowered (rapid freezing rate), for example in liquid nitrogen, the
formed nuclei of ice crystallization are small, leading to small pore size of samples after
drying [156,191]. In another case, at - 20 C freezing temperature (i.e. in a freezer), the pore
23
size of the dried sample is larger compared to the case of rapid freezing rate [156,191]. For
instance, in the case of gelatin foams, which are frozen at - 20 °C, their pore size has been
reported in the range of 250 - 300 µm, while at freezing temperatures - 80 C, smaller pore
sizes are obtained in the range of 45 - 50 µm [156]. In addition, the smaller pore size exhibits
thicker pore walls and subsequently higher mechanical properties [156,192]. Thus the
freezing temperature is the most important factor on determining the microstructure of
freeze-dried samples. The balance between pore size and mechanical properties must
therefore be optimized in each case for specific tissue regeneration. Scaffold architectures
fabricated by using freeze-drying technique are highly interconnected, which is necessary for
tissue ingrowth and regeneration [105]. Moreover, the scaffolds show achievable pore size
up to 300 µm and the porosity up to 97 % [105,193,194]. Freeze-drying also causes less
damage to the material and does not cause shrinkage or toughening of the material being
dried [105].
Figure 2. 4 (A) The schematic diagram of the freeze-drying (lyophilazation) process showing
also the phase diagram of water representing the mechanism of freeze-drying ([186,189]).
24
3D porous alginate foams as scaffolds for cartilage regeneration are being
extensively researched currently [59,62,63,134,195–201]. Most previous studies have
shown that freeze-dried Alg scaffolds with suitable porous structure can be used for culturing
with chondrocytes and MSCs. The proliferation and differentiation of cells, and the formation
of Col II and GAGs have been confirmed. As reported in the study of Lee et al. [195], Alg
foams promoted the adhesion, proliferation and differentiation of human chondrocytes and
the formation of specific cartilaginous matrices was detected. More recently, Wan et al. [197]
have prepared Alg foams in combination with CS by using freeze-drying technique. In vitro
culture of chondrocytes-seeded scaffolds was developed by on-site gelation (chondrocytes
embedded Alg gelation) in order to promote functional restoration and maintenance of the
round phenotype of chondrocytes. Petrenko et al. [134] have used freeze-drying method
with Ca-Alg hydrogel to develop porous scaffolds with wide pores for culturing with MSCs.
By this approach, cell adhesion and proliferation were not observed because Alg has
basically limited cellular interaction. Another approach has been developed by the
incorporation of Gel as a surface grafting onto the inner pore walls. As a result, the
adhesion, proliferation and differentiation of MSCs were improved [134]. In contrast, Miralles
et al. [202] have confirmed that Alg freeze-dried sponges promoted a favorable environment
for the growth of chondrocytes compared to Alg beads. Consequently, PGs rich matrix was
significantly detected by qualitatively histological evaluation in the case of Alg sponge. This
result has been suggested by the macroporous structure of the sponges indicating that
macro-channels allow better cell seeding and migration, compared to micro-porous gels
[202]. In addition, Wan et al. [196] confirmed that chondrocytes-embedded Alg hydrogels
did not exhibit the organization of synthesized Col in layer-anisotropic manner as in native
cartilage. Chondrocyte-clusters, chondrocyte proliferation and chondrogenic gene
expression (i.e. Col II, transcription factor Sox-9 and aggrecan) were observed in a porous
Alg sponge cultured with chondrocytes after 4 weeks, which indicates cartilage regeneration,
as evaluated by Yen et al. [203].
25
2.2.2.3 Electrospinning technique
Electrospinning involves the induction of static electric charges on the molecules of
solution at the level that causes the self-repulsion of charges [204]. Once the repulsion force
overcomes the force of surface tension of the solution, a jet of the solution forms, known as
Taylor cone, and stretches fibers toward the grounded collector [170,205]. Therefore, a
typical electrospinning set up consists of syringe containing solution, syringe pump, voltage
generator and collector, as demonstrated in Fig. 2.5. The fiber diameter can be controlled by
the variation of polymer solution concentration, flow rate, applied voltage and distance
between needle tip and collector [170,206]. Electrospinning has gained high popularity in the
field of tissue engineering due to its capability of fabricating fibers in submicron scales (30
nm – 10 µm in diameter). The nanotopographical features of electrospun fibers can for
example mimic Col fibrils in connective tissues [207,208]. Numerous materials, especially
polymers (synthetic and natural-derived polymers), have been successfully electrospun into
porous fibrous scaffolds for tissue engineering applications. Synthetic polymer fibers exhibit
sufficient mechanical properties and controlled biodegradability [117,208,209]. Particularly,
the biodegradation of scaffolds directly relates to the ability of the scaffold to maintain its
structure and subsequently to support cellular activity [208]. As investigated by Li et al.
[208], electrospun PLLA and PCL fibers exhibit higher proliferation of either chondrocytes or
MSCs compared to PLGA and PDLLA fibers. PLLA and PCL fibers maintained their fibrous
architecture over the culture time (21 days), while the degradation of PLGA and PDLLA
fibers was detected after 3 days in culture [208]. However, synthetic polymer fibrous
scaffolds usually lack the appropriate surface properties, i.e. hydrophilicity, to support cell
attachment and proliferation [117]. Natural polymer-based electrospun fibers are another
promising group of materials for scaffold development due to their excellent biocompatibility
[210–212]. CS/PEO fibers [213], for instance, were seen to support the adhesion of
chondrocytes and to maintain the round phenotype throughout the period of study.
Moreover, the study of Skotak et al. [214] proved that crosslinked Gel fibers cultured with
chondrocytes showed cell viability and supported the round phenotype of chondrocytes over
7 days in culture. In addition, Col II was observed in a high ratio to Col I [214].
26
Figure 2. 5 Schematic diagram of the electrospinning process in horizontal direction, which
is composed of voltage supply, syringe and needle, syringe pump and collector; the SEM
image shows electrospun PLLA fibers.
2.2.3 Strategies of multilayered scaffold
Current scaffold materials and designed strategies being applied in the field of
osteochondral tissue engineering are summarized in Table 2.4. The scaffold strategies can
be categorized into three available systems, namely single phase, bi- or multi-layered and
gradient structures, involving a variety of biocompatible materials and designs. Findings of
previous studies are expected to be a guideline for an improvement of scaffold designs to
get closer to an ideal osteochondral scaffold of relevance for clinical practice. Multilayered
and gradient composite scaffolds are being widely researched for osteochondral repair
instead of using single phase materials due to the fact that they can be designed and
fabricated to mimic the complex zonal structure of the native tissue [152,215–220]. As
reported in several studies, single material-based scaffolds could not provide the formation
of hyaline cartilage, instead fibrocartilage was generated. This is one reason of the
insufficient mechanical stability of the new tissues grown, leading to unsuccessful long term
27
repair. Moreover, the incorporation of cells and/or growth factors is essential for improved
scaffold performance.
Indeed, each scaffold strategy has its own advantages and disadvantages, which
are the result of either the used scaffold material or the design. The ideal osteochondral
scaffold has not been developed as yet. Additionally, the biological responses of different
engineered scaffold (in vitro and in vivo) are also different, also different animal models
usually show different outcomes. Also, in vitro studies of scaffolds mimicking the complex
structure of natural osteochondral tissue require specific culture systems. For example,
specifically designed bioreactors, i.e. double-chamber bioreactors [221], are required due to
different cell types and different culture media. According with this requirement, it becomes
apparent that osteochondral scaffolds need to avoid the migration of osteoblasts and
vascular tissue from the subchondral phase to the upper chondral phase, which is risky for
osteogenesis in the layer of cartilage. However, it has not been confirmed whether the
intermediate layer between cartilage and SB needs to be dense or porous. Taken into
consideration the anatomical structure of natural tissue, the intermediate layer called
calcified cartilage is very dense due to a mineralized matrix containing hypertropic
chondrocytes and packed extracellular fibrils. Therefore, the available free space is
assumed to be much smaller in comparison with that in cartilage. In addition, calcified
cartilage is connected to the subchondral bone by interdigitation. This might be the natural
barrier to avoid the migration of bone cells and vascular diffusion from SB. Importantly, new
tissues regenerated from implantation of scaffolds must be healthy and durable for long
term, which is initially confirmed by the organization of the new tissue in zonal arrangement
mimicking the structure of natural cartilage tissue.
28
Table 2. 4 Summary of current strategies in osteochondral tissue engineering.
Scaffold Materials Notable findings References
Cartilage Interface SB
Single phase
PLGA + Cell-free PLGA sponges showed a
possibility to repair full-thickness
defects in rabbits by absorbed local
cells from the erupted underlying bone
marrow.
- Small pore size ( 100 µm), porosity (
83 %) and hydrophobic property of the
scaffolds
Nagura et al.
(2007) [222]
PCL
+ PCL scaffolds in combination with
periosteal grafts provided excellent
integration with SB.
- Deficient development of hyaline
cartilage.
Mrosek et al.
(2008) [101]
PCL/F127 + PCL/F127 scaffolds culturing with
ADSCs and TGF-β1/BMP-7 improved
gross appearance of osteochondral
defect in rabbits.
- By histological results, the growth
factors did not significantly improve the
ability of the scaffold to repair the
defects due to high degree of foreign
body reaction.
Im et al.
(2009) [223]
29
Scaffold Materials Notable findings References Notable Findings References
Cartilage Interface SB Interface SB
Bi-/multi-layered
Fibrin
Fibrin glue PCL
- Fibrin scaffolds for cartilage phase
exhibited rapid degradation.
- Lack of mechanical support for
cellular development and excretion
of ECM.
Swieszkows-ki
et al. (2007)
[138]
PCL
Fibrin glue PCL/
TCP
+ PCL scaffolds for cartilage phase
prolonged degradation and
enhanced the mechanical
properties in comparison with fibrin
scaffolds.
+ PCL/TCP scaffolds for SB phase
promoted better BMSCs
proliferation compared to pure PCL
scaffolds after implantation in
rabbits.
PCL - PCL/
TCP
+ MSCs-seeded PCL constructs
promoted cartilaginous production
after implantation in pigs.
+ Resurfacing cartilaginous MSCs-
seeded PCL constructs with
electrospun Col I mesh could
protect cell leakage, reduce
generation of fibrocartilage and
enhance content of GAGs.
+ PCL/TCP-SB phase promoted high
mineralization.
Ho et al.
(2010) [224]
30
Scaffold Materials Notable findings References
Cartilage Interface SB
Col I Activated
plasma
Col I/β-
TCP
+ MSCs-seeded triphasic constructs
promoted cartilaginous production
and osseointegration.
- The constructs did not provide
better outcomes compared to the
conventional treatments such as
osteochondral autograft transfer
system, as confirmed by
histological results.
Marquass et al.
(2010) [225]
Agarose
Agarose/
PLGA/ 45S5
BG
PLGA/
45S5
BG
+ Three-layered scaffolds supported
co-culture of chondrocytes and
osteoblasts, and promoted the
generation of continuously distinct
hyaline cartilage-calcified cartilage-
SB.
+ BG presenting in intermediate and
SB phases enhanced the formation
of CaP, indicating ability of
mineralization.
Jiang et al.
(2010) [226]
Col I - Col I/
HA
+ Layered scaffolds promoted
chondrogenic and osteogenic
differentiation of MSCs.
- Culturing different cell types in the
layered scaffolds at the same time
was limited. Designed double-
chamber bioreactor is thus
required.
Zhou et al.
(2011) [221]
31
Scaffold Materials Notable findings References
Cartilage Interface SB
PA6 Non-porous
PVA
PVA/H
A
+ In vivo studies of BMSCs-seeded
PA6 and PVA/HA separately in
rabbit muscle pouch showed that
PA6 constructs provided a
cartilaginous marker (i.e. Col II);
and PVA/HA constructs provided
an osteogenic marker (i.e. Col I).
- Non-porous PVA intermediate layer
was expect to function as a native
calcified cartilage but it is not a
clear requirement for ideal
cartilage-bone interface that
whether the interface layer is dense
or porous.
Qu et al.
(2011) [152]
Col I
- Col I + MSCs-seeded Col I scaffolds
promoted both cartilage- and bone-
like tissues during in vitro culture,
respectively.
+ After formed bilayered constructs,
undifferentiated MSCs in the
intermediated layer were
differentiated into hypertrophic
chondrocytes and produced Col X.
+ Zonal organization of osteochondral
tissue was formed by the optimal
designed scaffold and controlled
cell density.
- Mechanical environments must be
taken into account for better zonal
organization of tissues.
Cheng et al.
(2011) [100]
32
Scaffold Materials Notable findings References
Cartilage Interface SB
CS/Gel Fibrin glue CS/
Gel/HA
+ TGF-β1 and BMP-2 enhanced
chondrogenesis and osteogenesis
of MSCs on CS/Gel and CS/Gel/HA
constructs in vitro, respectively, and
also supported the regeneration of
osteochondral tissue in vivo.
Chen et al.
(2011) [114]
PDLLA - PDLLA/
45S5
BG
+ In vitro culture with ACDC5 showed
cell attachment, proliferation and
migration through PDLLA fibers,
which are suitable for cartilage
regeneration.
Yunos et al.
(2012) [86]
Poly HEMA/
HyA
- Poly
HEMA/
HyA/
nHA
+ HyA incorporated into polyHEMA
layer encouraged chondrogenesis
of chondrocytes and played an
important role in maintenance of
chondrocytes phenotype.
+ nHA in SB phase induced osteo-
conductivity and -inductivity of
MSCs- seeded constructs.
- Small pore size of polyHEMA/HA-
SB phase ( 38 µm) was conflicted
to the ideal pore size suitable for
vascularization and bone ingrowth
( 300 - 500 µm).
Galperin et al.
(2012) [227]
33
Scaffold materials Notable findings References
Cartilage Interface SB
Gradient/graded
HyA/
Col I
Col I/HA
(40/60)
Col I/HA
(30/70)
+ Gradient in compositions provided
specific proper environment for
specific differentiated cells and
supported the generation of zonal
organized tissues.
+ By using freeze-drying technique, the
obtained scaffold also exhibited
gradient in porous structures.
- MSCs are seemed to be only cell
source suitable for the gradient
strategy because it is extremely
complicated to perform in vitro co-
culture of different cell types in the
single gradient structure.
Tampieri et al.
(2008) [110]
PLGA/
TGF-β1
- PLGA/
HA/
BMP-2
+ Gradient in material compositions
and growth factors promoted the
regeneration of cartilage and SB
after implantation in rabbit knees.
+ Growth factors triggered the
differentiation of progenitor cells into
chondrogenic and osteogenic cells
after implantation.
+ The gradient scaffolds based on the
combination of BMP-2 and nHA
promoted fast bone formation in vivo
and restored new cartilage.
- The manufacture of the gradient in
growth factors is only possible in the
form of microsphere-based scaffolds,
while it is complicated in the form of
sponge-like scaffolds.
Mohan et al.
(2011) [228]
34
Scaffold materials Notable findings References
Cartilage Interface SB
Agarose/
MSCs in
osteoge-
nic
medium
Agarose/
MSCs in
basic
culture
medium
Agarose/
MSCs in
osteoge-
nic
medium
+ Gradient in MSCs within different
culture media stimulated
differentiation of MSCs into
chondrocytes and osteoblasts on
cartilage and SB phases,
respectively.
+ Gradient in cell types was observed
at intermediate phase.
- Gradient-generating culture device is
required.
Shi et al.
(2012) [229]
Col I Col I/nHA
(40/60)
Col I/nHA
(30/70)
+ After implantation in sheep for 6
months, either cell-free scaffolds or
chondrocytes-seeded scaffolds
provided the formation of
osteochondral tissue and filled the
full-thickness defect.
+ No difference of tissue regeneration
was found in both, cell-free and
chondrocytes-seeded scaffolds, due
to the recruitment of local BMSCs.
+ Cell-free gradient scaffolds were
tested in human patients: at 2 years
follow-up, 70 % of the defects were
completely filled with new tissues.
- Low number of patients and scarcity
of evaluation cases were limited in
the comparison of the results.
Kon et al.
(2010, 2012)
[230,231]
35
2.3 Cells and bioactive molecules/growth factors for osteochondral
tissue engineering
A current approach in osteochondral tissue engineering is to reconstruct the
functional engineered cartilage–bone interface by co-culturing chondrocytes and osteoblasts
into multilayered scaffolds [232]. Chondrocytes and osteoblasts are an obvious choice for
cell sources because they are found in native cartilage and bone tissues, respectively.
Autologous chondrocytes and osteoblasts have been successfully used in the regeneration
of cartilage [131,209,233–235] and bone [44,171,236–238], respectively. In addition, co-
culture of chondrocytes and osteoblasts has been extensively studied for osteochondral
interface [102,111,232,239]. Cao et al. [102] investigated porous PCL scaffolds co-culturing
with chondrocytes and osteoblasts. Both cell types produced specific ECM in each own
compartment and they also migrated and integrated at the interface of the scaffold. In this
approach, the important aspect is the interaction between osteoblasts and chondrocytes
associating with scaffold material during co-culture, which leads to the proper formation of
osteochondral interface [239]. However, it is critical to control specific osteogenic and
chondrogenic phenotypes during co-culture in order to mimic each tissue zone to native
tissues.
Even though the utilization of autologous cells is conducive to the growth of
functional tissues, they are difficult to isolate and easily change their phenotypes during the
culture process due to their dedifferentiation capacity [25,240]. Therefore, another available
cell source, i.e. progenitor cells (stem cells), is a valid alternative. Stem cells are promising
cells for tissue engineering due to their multipotent nature and self-renewal capacity [241–
243]. Focusing on osteochondral repair, stem cells as single cell source, which can be
differentiated into chondrogenic and osteoblastic cell lines, also overcome the limited supply
of primary cells. Since embryonic stem cells (ESCs) have a limitation in differentiation
capacity [103], bone marrow mesenchymal stem cells (BMSCs) and adipose-derived stem
cells (ADSCs) are the most widely used in current research because they are abundant in
the human body and can be isolated from bone marrow stroma and adipose tissue,
respectively. In particular, BMSCs are promising cells for regeneration of tissues due to
36
their rapid proliferation and easy differentiation into osteogenic lineage by the use of
supplemented dexamethasone, ascorbic acid and β-glycerophosphate in vitro [25]. It has
been shown in in vitro study that BMSCs can be differentiated into the osteogenic lineage by
the presence of growth factor genes [33]. BMSCs can also undergo chondrogenic
differentiation when cultured in the presence of transforming growth factors (TGF) [4]. The
specific signaling molecules, such as TGF-β family, insulin-like growth factors (IGF), bone
morphogenetic proteins (BMP) and fibroblast growth factors (FGF), are extensively
employed to facilitate tissue growth, by promoting tissue specific proteins and simulating the
differentiation of stem cells [244]. They bind to cell surface receptors and activate
intracellular signaling pathways, which affect cell proliferation, differentiation and ECM
synthesis during tissue regeneration [5,23,244]. The study of Re’em et al. [245] showed that
BMSCs-seeded RGD-modified Alg sponge culturing in the presence of TGF-β1 provided
appropriate progression of BMSCs differentiation. More recently, the combination of TGF-β1
and BMP-2 has been confirmed to activate BMSCs differentiation and to promote the
formation of hyaline cartilage, as shown in the study of Toh et al. [246]. Focusing on
cartilage tissue engineering, growth factors in combination (i.e. TGF-β1/BMP-7 and TGF-
β1/IGF-1) are synergistically high effective for hyaline cartilage regeneration compared to a
single growth factor [244]. For instance, MSCs cultured with the introduction of TGF-β1 were
able to promote an increased cartilaginous ECM synthesis and a decreased Col I synthesis
when TGF-β1 was combined with BMP-7 [244].
Even though BMSCs are widely used in osteochondral repair, the yield of cells from
bone marrow harvest is small because only 10 - 25 ml of bone marrow can be obtained
from a human [5,247–249]. Alternatively, synovial membrane derived cells are gaining
consideration due to their great chondrogenic potential, which is suitable for cartilage tissue
engineering [250–253]. The synovial membrane is the supportive layer of a joint, which is
composed of cellular lining layer [5]. The biopsy of the synovial membrane provides the
accessibility of autologous MSCs [5]. Human synovial MSCs have shown evidence of
chondrogenesis when cultured in 3D Alg scaffolds without the presence of growth factors
[250]. In addition to BMSCs and synovial MSCs, ADSCs obtained from lipoaspirates provide
37
several advantages related to their abundance, suitable accessibility and relatively low donor
morbidity [223]. However, it was evidenced that ADSCs promoted lower chondrogenic and
osteogenic potential compared to BMSCs [23,254]. This limitation could be overcome by
using the combination of growth factors (i.e. TGF-βs/BMPs) [223,254]. Growth factors have
been used with BMSCs in order to stimulate cell proliferation and the synthesis of
extracellular molecules in vitro and in vivo [244,253]. Since it is commonly found that in vitro
culture of either cell-free scaffolds or cells-seeded scaffolds promotes fibrocartilage
formation instead of the formation of hyaline cartilage, certain growth factors have been
included to induce cell growth into a specific tissue formation [25]. For example, Huang et al.
[177] investigated cartilage formation induced by TGF-β1-loaded fibrin glue incorporated into
PCL scaffold after in vitro culturing with BMSCs and implantation in a lapine model. The
constructs were richly populated with chondrocytes after 4 weeks of implantation, while
immature bone was identified at week 6 of implantation [177]. Therefore, TGF-β1 is
regarded as a cartilage-inducing factor as well as a bone-inducing factor [19].
Generally, growth factors can be introduced to promote the repair of osteochondral
tissue by different methods, either by incorporating them into the scaffold [121,255,256] or
by adding them in the culture medium as a supplement [70,114,245,257]. Cui et al. [118]
investigated the chondrogenic capacity of ADSCs-seeded PGA/PLA fibrous scaffolds
cultured in the presence of TGF-β1 as a medium supplement. Col II and GAGs were
observed after 2 weeks of culture. Similarly, BMSCs-seeded PCL fibrous scaffolds were
cultured in chondrogenic medium (in the presence of TGF-β3) in order to analyze
chondrogenesis and in osteogenic medium in order to analyze mineralization [258]. After 3
weeks in culture, formation of cartilaginous tissue and mineralization were observed in the
newly formed ECM by day 45 [258]. In a recent approach, osteochondral scaffolds have
been designed creating a gradient of encapsulated growth factors, as reported in the study
of Mohan et al. [228]. Chondral PLGA microspheres loaded with TGF-β1 and osseous
PLGA/HA microspheres loaded with BMP-2 were implanted in rabbit knees. It was found
that the gradient scaffolds promoted the complete bone ingrowth and cartilage regeneration
with high GAGs content [228]. Re’em et al. [259] showed that BMSCs seeded onto both,
38
TGF-β1-loaded porous Alg scaffolds as a chondroinductive phase and BMP-4-loaded
porous Alg scaffold as an osteoinductive phase, promoted the differentiation into
chondrocytes and osteoblasts, respectively. Moreover, biomolecules, for example HyA and
chondroitin sulfate (ChS), have been incorporated into either scaffolds or culture medium to
stimulate chondrogenesis of MSCs [68,198,260].
CHAPTER 3
Objectives and Outline
The main goal of this research is to study and develop biomaterials-based scaffolds
suitable for osteochondral tissue engineering applications, focusing on combination of
inorganic and polymer materials, and including novel fabrication techniques. This work will
include the characterization of the physic-chemical and mechanical properties as well as the
biological response, e.g. in vitro cell culture, of the new scaffolds. The different tasks carried
out to achieve the final thesis objectives are summarized in Fig. 3.1.
The fabrication of an appropriate scaffold for bone regeneration based on
biodegradable polymer coated 45S5 Bioglass®-based scaffolds to satisfy the structural,
physico-chemical, mechanical and biological properties of native bone-tissue. The variety of
biodegradable polymer coated Bioglass®-based scaffolds investigated is reported in
Chapter 4, based on the basic requirements for suitable bone scaffolds such as mechanical
and in vitro biological properties.
The biodegradable polymer coated Bioglass®-based scaffolds were developed
further for imparting a local drug-delivery capability for bone tissue engineering. The
combination of synthetic and natural polymers is carried out as suitable layered coating on
the Bioglass®-based scaffold, which provides a dual functions of improving the mechanical
properties and also acting as a drug carrier. The details are discussed in Chapter 5.
The fabrication and characterization of engineered scaffolds for cartilage
regeneration concerning the selection of biomaterials and fabrication techniques to achieve
the required structural architecture, morphology, physico-chemical and mechanical
properties is detailed in Chapter 6. In this part of the study, alginate-based scaffolds are
fabricated by two different techniques, including freeze-drying and electrospinning
techniques, in order to study the effects of two different architectures and geometries of the
obtained scaffolds. In addition, their morphology, physico-chemical and mechanical
properties are reported and discussed.
40
Chapter 7 discusses the design and fabrication of bi- or multi-layered scaffolds
suitable for osteochondral tissue engineering based on the optimized scaffold materials and
fabrication methods. Recent strategies discussed previously, considering the biomimetic
approach leading to integrated- and monolithic-layered scaffolds are included in this work to
elucidate the advantages and disadvantages of each design of multilayered scaffolds.
The cellular response and activity on the optimized designed scaffold for bone and
cartilage regeneration are discussed in Chapters 8 and 9, respectively. In detail, MG-63
osteoblast-like cells are seeded on Bioglass®-based composite scaffolds to study cell
proliferation and metabolism. While chondrocytes and MSCs are seeded on 3D
alginate/chondroitin sulfate-based foams to evaluate cell adhesion, proliferation and
differentiation, and cartilaginous matrix formation, which provides basic information about
the suitability of the scaffolds for cartilage tissue regeneration.
In order to improve further this research work and to reach closer to the ideal
osteochondral scaffold highly feasible for in vitro and in vivo culture studies and available for
use in clinical practices, the summary and future perspectives are highlighted in Chapter 10.
41
Figure 3. 1 Schematic diagram of entire tasks carried out in the dissertation thesis.
42
CHAPTER 4
Preparation and Characterization of Biodegradable Polymer Coated
45S5 Bioglass-Based Scaffolds for Subchondral Bone Tissue
Engineering Applications
4.1 Introduction
Osteochondral tissue engineering involves the regeneration of both cartilage and SB
tissues, which are connective tissues providing the body with mechanical support and
protection [16,139,140,261]. Morphologically, bone is classified into two forms with different
structural and functional properties: cortical and cancellous bone [12,262]. SB demonstrates
the characteristics of cancellous bone, as described in Chpater 2. In order to engineer
scaffolds for SB regeneration, an attempt to mimic the natural tissue structure is
fundamental. In general terms, biomaterials-based scaffolds, cells and active molecules or
growth factors are three essential components for bone tissue engineering [263]. Especially,
scaffolds as the temporary framework of tissue regeneration, e.g. artificial ECM, are crucial
to achieve the required properties of new bone tissue, including structural and mechanical
properties [34]. By this approach, optimized scaffolds must be developed via the selection of
suitable materials and fabrication methods. Regarding ideal scaffolds for bone regeneration,
3D highly porous templates with tailored porosity, pore size and high interconnectivity are
generally required. In addition, bone tissue engineering requires biocompatible and
biodegradable scaffolds made from biomaterials exhibiting biodegradation matching the rate
of bone tissue formation. The biomaterials should also enable cell attachment, proliferation
and differentiation, ideally exhibiting osteoconductive and osteoinductive properties.
Moreover, the bone tissue scaffolds must have sufficient mechanical properties (mimicking
mechanical properties of natural bone) for load bearing applications and adequate structural
stability for use in clinical practice.
44
According to these required properties, expecially in relation to bone bonding, 45S5
Bioglass was chosen material for bone scaffold in the present study. Highly porous 45S5
Bioglass®-based scaffolds can be fabricated by foam replication technique [32]. The
architecture of such scaffolds is similar to that of cancellous bone, being also similar to the
structure of the underlying SB in the osteochondral interface tissue. However, porous 45S5
Bioglass®-based scaffolds are characteristically brittle and exhibit low mechanical stability,
also the relatively high porosity of scaffolds results in reduced mechanical properties [34].
Therefore, biodegradable polymer coatings are applied to improve the mechanical properties
and structural stability of scaffolds, while their porosity and bioactivity are maintained [52,56].
Several biodegradable polymers, including natural polymers and synthetic polymers, were
studied in this area as suitable coatings for highly porous bioceramic scaffolds [52–
57,87,96,98,142,264–270]. Different polymers provide different characteristics, degradation
rate, and mechanical properties. Both natural (Alg and Gel) and synthetic (PDLLA and
PHBHHx) biodegradable polymers are therefore considered in this study in order to compare
their performance and efficiency as a coating on Bioglass®-based scaffolds, mechanical
properties and in vitro bioactivity are mainly considered. Improvement of the mechanical
strength and structural stability of scaffolds was aimed at achieving as the polymer would fill
and bridge the microcracks on the struts of Bioglass®-based scaffolds. In addition, the
polymer coating should increase the thickness of the struts without closing the pores.
Another purpose of this work is to optimize the coating conditions of each investigated
polymer and to compare their performances for appropriate scaffolds in osteochondral tissue
engineering.
4.2 Materials and methods
4.2.1 Fabrication of 45S5 Bioglass®-based scaffolds
45S5 Bioglass®-based scaffolds were prepared by the foam replication method, as
described in Chapter 2. In brief, this technique involves coating PU foam with a Bioglass®
slurry. The slurry for the immersion of PU foam was prepared as follows. PVA, purchased
from Merck KGaA, Germany, was dissolved in DI H2O with concentration of 3.5 wt/v %.
45
Then 45S5 bioactive glass powder, of 45S5 Bioglass® composition (Schott electronic
packaging GmbH, Germany) was added to PVA solution with concentration of 40 wt/v %.
The whole procedure was carried out at 80 C under vigorous magnetic stirring for 2 h. The
“Eurofoam” PU foam with 45 ppi (pore per inch) served as a sacrificial template. It was cut to
size 10 mm × 10 mm × 10 mm. Then PU foams were immersed in the prepared slurry for 1
min. The foams were then removed and the extra slurry was completely squeezed out
manually. The samples were then dried in an oven at 60 C for 12 h. The coating thickness
of the samples (green bodies) was increased by repeating the slurry coating procedure for
three times. Heat treatment was carried out for burning out PU templates and for sintering
45S5 Bioglass® structure. The burning condition of PU templates and sintering condition
were designed to be 450 C for 1 h and 1100 C for 2 h, respectively. The heating and
cooling rates were 2 and 5 C/min, respectively.
4.2.2 Preparation of biodegradable polymer coated 45S5 Bioglass®-based scaffolds
Different types of polymer coated 45S5 Bioglass®-based scaffolds were fabricated.
Natural biodegradable polymers, including Alg and Gel, and biodegradable synthetic
polymers, including PDLLA and PHBHHx, were chosen to coat 3D 45S5 Bioglass®-based
scaffolds by using dipping technique. The characteristics of the various polymer solutions
used are shown in Table 4.1. A simple manual dip coating method was used for infiltrating
polymer into the scaffolds’ structure and to adhere the polymers to the surface of struts. The
procedure of polymer coating of scaffolds is explained for each polymer used in the following
paragraph.
For Alg coating, coating solution was prepared as follows: Alg was dissolved in DI
H2O with the concentration of 2 wt/v % at room temperature, under magnetic stirrer for 2 h.
The 2 wt/v % solution was diluted to 0.5 and 1.5 wt/v % with DI H2O. Then 45S5 Bioglass®-
based foam was immersed in Alg solution for 5 min and dried at room temperature for 24 h.
The coating process of Gel and PDLLA coated 45S5 Bioglass®-based scaffolds was carried
out using the same procedure, following the conditions shown in Table 4.1. In contrast, in
case of PHBHHx coated 45S5 Bioglass®-based scaffolds, PHBHHx was dissolved in
chloroform (Merck KGaA, Germany) with concentration of 1 and 5 wt/v % at 50 C, under
46
magnetic stirring for 2 h. 45S5 Bioglass®-based scaffolds were soaked in PHBHHx solution
for 10 sec and 5 min, respectively, and dried at room temperature for at least 30 min before
starting another coating layer. In our experiments, the coating process was carried out 30
times in the case of 1 wt/v % solution. Each coating sample was labeled as Alg-c-BG, Gel-c-
BG, PDLLA-c-BG and PHBHHx-c-BG for Alg, Gel, PDLLA and PHBHHx coated scaffolds,
respectively.
Table 4. 1 Polymer coating conditions for polymer coated 45S5 Bioglass®-based scaffolds.
An as-sintered rectangular shaped 45S5 Bioglass®-based scaffold, with the dimensions of 8
mm × 8 mm × 8 mm, was soaked in 5 ml of each polymer solution.
Biodegradable
polymers
Coating conditions
Concentrations
(wt/v %) Soaking times No. of dipping cycles
Alg 0.5, 1.5, 2 5 min 1
Gel 1.5, 3, 5 5 min 1, 3
PDLLA 2, 5, 8 5 min 1
PHBHHx1 1, 5 10 sec, 5 min 30, 1
4.2.3 Characterization and mechanical testing
(i) Porosity
The percent porosity of uncoated (P1) and polymer coated (P2) scaffolds was
calculated from Eq. 4.1 and 4.2 [56]:
P1 (%) = [1 – (Wscaffold/BG/Vscaffold)] × 100 (4.1)
P2 (%) = [1 – (Wscaffold/BG/Vscaffold) + (Wpolymer coating/polymer/Vpolymer coating)] × 100 (4.2)
; where Wscaffold, Wcoated scaffold and Wpolymer coating are the weight of sintered scaffold, polymer
coated scaffold and polymer coating (Wcoated scaffold – Wscaffold), respectively, BG and Polymer
are the density of solid 45S5 Bioglass® (BG = 2.7 g/cm
3) and solid polymer (Alg = 1.02
1 30 coating cycles were done in the case of 1 wt/v % PHBHHx solution and one coating cycle was done in the case
of 5 wt/v % PHBHHx solution.
47
g/cm3, Gel = 0.98 g/cm
3, PDLLA = 1.26 g/cm
3 and PHBHHx = 1.2 g/cm
3), respectively, and
Vscaffold, Vcoated scaffold and Vpolymer coating are the volume of sintered scaffold, polymer coated
scaffold and polymer coating (Vcoated scaffold – Vscaffold), respectively, determined from the
dimensions of scaffolds.
(ii) Microscopy
The microstructure of the scaffolds was characterized by scanning electron
microscopy (SEM; LEO 435 VP), before and after coating. Samples were sputter coated and
observed at an accelerating voltage of 10 kV. The pore size was measured from images
taken from SEM by using software Image J (Version 1.42S).
(iii) XRD analysis
The scaffolds were characterized by using X-ray diffraction (XRD) analysis in order
to assess the crystallinity after sintering. The scaffolds were ground into powder. Then 0.1 g
of the powder was collected for XRD (Siemens D500) analysis, employing Cu k radiation.
Data were collected over the range of 2 = 15 - 70 using a step size of 0.02 and a counting
time of 25 sec per step.
(iv) In vitro acellular bioactive study
The bioactivity of uncoated and polymer coated scaffolds was investigated by
immersion in SBF solution (pH 7.4 at 37 °C) for 1, 3, 7, 14 and 28 days. The SBF solution
was prepared, following Kokubo et al. [271]. Each sample (with the dimensions of 8 mm × 8
mm × 8 mm) was placed in polystyrene bottle containing 50 ml SBF, then it was incubated in
orbital shaker (IKA RS 4000i) at 37 C with the shaker speed of 90 rpm. The SBF solution
was replaced twice a week in order to avoid the changes in the chemistry of samples. At
each time point, the sample was removed, cleaned with DI water and dried at room
temperature for 24 h. The morphology of scaffolds after immersion in SBF was observed by
SEM. The formation of HCA crystals was characterized with the use of Fourier-transform
infrared (FTIR) spectroscopy and x-ray diffraction (XRD) analysis.
(v) FTIR analysis
The possible presence of HA on the scaffolds upon immersion in SBF was identified
with the use of Fourier-transform infrared spectroscopy (FTIR; Nicolet Nexus 6700, Thermo
48
Scientific, Waltham, MA). For these measurements, the scaffolds were grinded and mixed
with anhydrous potassium bromide (KBr) powder in the ratio of 1/300 by wt. The mixture was
pressed to a pellet by using an electro-hydraulic press (MAUTHE MASCHINENBAU PE-010;
Wesel, Germany) with a pressure of 10 × 104 N. Then the pellet was analyzed by using
transmission mode with a resolution of 4 cm-1
, in the wavenumber range of 4000 - 400 cm
-1
and applying 64 scans.
(vi) Mechanical testing
The porous scaffolds prepared with dimensions of 8 mm × 8 mm × 8 mm were
tested using a universal testing machine (Zwick Z050) by applying a compression load at a
cross-head speed of 5 mm/min, preload at 0.1 N and maximum load at 1 kN. Stress-strain
curves were recorded to determine the relevant mechanical properties, compressive
strength and to assess the work of fracture. Six specimens were tested for each condition
and data were presented as mean ± standard deviations (SD).
4.2.4 Statistical analysis
Statistical comparisons were carried out using one-way ANOVA method, which p =
0.05 was considered to be a significant difference.
4.3 Results and discussion
4.3.1 Morphology
(i) General characterization
45S5 Bioglass®-based scaffolds fabricated by foam replication technique exhibited
porosity of 92 % with the pore size in the range of 100 - 700 μm. A highly porous
trabecular structure and high pore interconnectivity are obtained and this can be appreciated
in Fig. 4.1 (A). This porosity is desired for facilitating high ability of cell seeding and to enable
efficient cell proliferation and bone ingrowth, as discussed in Chapter 2. Multiple slurry
coatings were carried out to increase the thickness of struts, relative density, and
consequently to improve the strength of scaffolds, compared with that of scaffolds prepared
from only one coating. The scaffolds made by a triple coating process were sufficiently
strong for handling safely in the laboratory by forceps or fingers, while scaffolds made by a
49
single coating were not. The strut of scaffold shown in Fig. 4.1 (B) exhibits quite a dense
surface indicating good densification achieved at high sintering temperature (1100 C),
which was caused by well bonding of Bioglass® particles. The extensive densification and
the presence of crystalline phase in as-sintered scaffolds are expected to lead to
improvement of the mechanical properties [32].
The crystalline phase of the as-sintered scaffolds was confirmed by the results of
XRD analysis (Fig. 4.2), which shows the sharp crystalline peaks of Na2Ca2Si3O9 compared
to the pattern of as-received 45S5 Bioglass®
powder, matching the same angular location
and the same form of peaks as in a previous study on sintered 45S5 Bioglass® [32].
Figure 4. 1 SEM images of 45S5 Bioglass®-based scaffolds fabricated by foam replication
technique: (A) 3D porous structure and (B) surface of scaffold struts.
50
Figure 4. 2 X-ray patterns of as-received Bioglass® and as-sintered 45S5 Bioglass
®-based
scaffolds. The major peaks of the phase Na2Ca2Si3O9 are marked by ■.
However, bioceramics and bioactive glasses in a porous form have usually low
strength and fracture toughness because of their intrinsic brittleness, leading to a hindrance
in their clinical applications and commercial uses. As discussed in Chapter 2, polymeric
coating on porous foams is one method being increasingly investigated to improve the
mechanical properties of scaffolds [52,55]. Different polymers, including natural polymers
(Alg and Gel) and synthetic polymers (PDLLA and PHBHHx), were infiltrated in this study
into the porous scaffolds to test the hypothesis that the polymer coating can improve the
mechanical strength and toughness of the porous scaffolds by filling the cracks and open
pores on the struts of scaffolds. For each polymeric coating, the polymer concentration,
immersion times and number of dipping cycles were varied (Table 4.1).
(ii) Alginate coated scaffolds
First, Alg coated Bioglass®-based scaffolds (Alg-c-BG) were studied. At 2 wt/v % Alg
concentration, it was found that some pores were clogged with a polymer membrane, as
shown in Fig. 4.3 (E), which is not satisfactory as it could hinder cell migration during in vitro
cell culture study. Additionally, the surface of coated scaffolds was not homogeneous due to
51
a too high concentration used and subsequent high viscosity of the coating solution (Fig. 4.3
(F)), leading to hindrance of polymer infiltration. In contrast, coating with a 0.5 wt/v % Alg
solution led to scaffolds, which were not stable due to the fact that the polymer did not fully
cover the surface of the struts (Fig. 4.3 (B)), even though pores were not clogged (Fig. 4.3
(A)). In case of 1.5 wt/v % Alg concentration, the coating was homogeneous and no pores
were seen to be closed by a polymer membrane (Fig. 4.3 (C)). Fig. 4.3 (D) shows the
smooth surface of a coated scaffold using a 1.5 wt/v % Alg solution, the polymer coating
layer covered efficiently the entire surface of the struts. Thus qualitatively optimized Alg-c-
BG scaffolds were obtained using a solution with 1.5 wt/v % Alg concentration. It can be
concluded that a too high Alg concentration ( 2 wt/v % concentration) is not suitable for
infiltrating into the scaffold’s porous structure due to the high viscosity. Therefore, the
viscosity of Alg solution plays a key role in the quality of coating, as also reported in the
previous study [54].
(iii) Gelatin coated scaffolds
Fig. 4.4 (A-F) shows structural and morphological features of Gel coated Bioglass®-
based scaffolds (Gel-c-BG), which were obtained from 1.5, 3 and 5 wt/v % Gel
concentrations, 5 min of immersion time and single dip coating. It was found that 3 and 5
wt/v % Gel concentrations exhibited many clogged pores due to the high solution viscosity
(Fig. 4.4 (C and E)). Gel coatings with lower concentration (1.5 wt/v %) were found to lead to
non-clogging coatings, while the scaffolds obtained were not sufficiently strong because the
coating layer was too thin. Taking into account the thickness of coating layer can be
increased by increasing the number of coating cycles. The morphology of the triple coated
scaffold is shown in Fig. 4.4 (G and H). In this study, three dipping cycles were qualitatively
considered to lead to suitable mechanical properties and to maintain the suitable porous
characteristics of the scaffolds.
Preliminarily, the natural-derived polymers (Alg and Gel) can be recommended to be
used as coating of Bioglass-based scaffolds using relatively low polymer concentrations in
order to obtain homogeneous coatings. In addition, the mechanical strength can be
52
enhanced by increasing the number of dipping cycles, as discussed in the case of Gel
coatings.
(iv) PDLLA coated scaffolds
For PDLLA coated scaffolds (PDLLA-c-BG), it is seen from Fig. 4.5 that their open
pores were still maintained at PDLLA concentration of 5 wt/v % (Fig. 4.5 (C)). In contrast, the
pores were blocked with the 8 wt/v % coating (Fig. 4.5 (E)), considering that this polymer
solution was highly viscous. This phenomenon is also related to the fact that a polymer
solution with higher concentration takes a longer time for infiltrating through the porous
structure in comparison with lower concentration polymer solutions. Moreover, the used
organic solvent (i.e. DMC) is fast evaporated at room temperature, leading to non-
homogeneous coating on the whole surface area of the struts, as shown in Fig. 4.5 (E and
F). Fig. 4.5 (D) shows that the PDLLA coating integrated well on the strut’s surface, also
filling cracks and pores on the struts’ surface and it leads to a smooth surface. In contrary to
the struts’ surface of 3 wt/v % PDLLA-c-BG scaffold (Fig. 4.5 (A)), some pores on the struts
were observed after coating (Fig. 4.5 (B)).
(v) PHBHHx coated scaffolds
In case of PHBHHx coated scaffolds (PHBHHx-c-BG), to avoid the phenomenon of
closed pores, the optimized conditions should be the use of low polymer concentration and
short immersion time, and the thickness of the struts was increased by multiple coatings.
This hypothesis was thus investigated by coating the scaffolds with 1 wt/v % concentration,
10 sec of immersion time and 30 dipping cycles in comparision with coating the scaffolds
with 5 wt/v % concentration, 5 min of immsion time and one dipping cycle. It was found that
low concentration and short immersion time supported better homogeneous coating surface
in comparison with higher concentration, longer immersion time with one coating cycle
(compare between Fig. 4.6 (A and B) and Fig. 4.6 (C and D)). This result can be described
by the fact that PHBHHx has elastomeric properties (rubber-like behavior), which its
solutions with higher concentration and longer soaking time tend to stick at the intersection
of struts impeding efficient infiltration, leading to clogged pores (Fig. 4.6 (A)). In contrary to
the use of low concentration and short soaking time, the thickness of coating was built up by
53
30 dipping cycles without closing pores, as shown in Fig. 4.6 (C). This effect thus causes
different coating morphologies.
(vi) Summary of results on polymer coatings
The optimized coating parameters of different biodegradable polymers are
summarized in Table 4.2. By Alg coating, the porosity of scaffolds was reduced from 92 to
84 %. Moreover, Alg formed thin film well adhering on the struts of the scaffold, showing a
smooth surface. The Gel-c-BG scaffolds (porosity 82 %) obtained by using a 1.5 wt/v %
solution, 5 min soaking time and 3 dipping cycles showed the morphology in a similar
manner to the case of Alg coating. These results can be described by the fact that since
both, natural polymer coatings (Alg and Gel) and the strut surface of Bioglass-based
scaffolds, exhibit hydrophilic feature, they are chemically compatible. According to this fact,
the viscosity of natural polymer solutions mainly influences the quality of coating. In contrast,
the optimized conditions of synthetic polymer coatings provided a thinner coating layer on
the scaffolds than that observed in the cases of natural polymer coating. The porosity of
PDLLA-c-BG and PHBHHx-c-BG scaffolds was 85 % and 86 %, respectively. By
increasing concentrations 5 wt/v %, the PDLLA and PHBHHx solutions tended to form as a
membrane covering the pores of the scaffold, instead of coating onto the struts. It seems
that not only the viscosity of polymer solutions plays a key role in the obtained morphology
of synthetic polymer coated scaffolds, but also the different surface chemistries of synthetic
polymer and Bioglass surface influence the quality of coating.
54
Figure 4. 3 SEM images of Alg coated 45S5 Bioglass®-based scaffolds with variable
concentrations, including (A, B) 1 wt/v %, (C, D) 1.5 wt/v % and (E, F) 2 wt/v %, the scaffolds
were coated once for 5 min soaking time.
55
Figure 4. 4 SEM images of Gel coated 45S5 Bioglass®-based scaffolds with variable
concentrations, including (A, B) 1.5 wt/v %, (C, D) 3 wt/v %, (E, F) 5 wt/v %, the scaffolds
were coated once on for 5 min soaking time and (G, H) 1.5 wt/v % coated for 5 min and 3
dipping cycles.
56
Figure 4. 5 SEM images of PDLLA coated 45S5 Bioglass®-based scaffolds with variable
concentrations, including (A, B) 3 wt/v %, (C, D) 5 wt/v % and (E, F) 8 wt/v %, the scaffolds
were coated once for 5 min soaking time.
Figure 4. 6 PHBHHx coated 45S5 Bioglass®-based scaffolds with variable concentrations,
soaking times and number of dipping times, including (A, B) 5 wt/v %, 5 min of soaking and
one coating cycle, and (C, D) 1 wt/v %, 15 sec of soaking time and 30 coating cycles.
57
Table 4. 2 Optimized polymer coating conditions for different biodegradable polymer coated
45S5 Bioglass®-based scaffolds.
Biodegradable
polymers
Coating conditions
Concentrations (wt/v %) Soaking times No. of dipping cycles
Alg 1.5 5 min 1
Gel 1.5 5 min 3
PDLLA 5 5 min 1
PHBHHx 1 15 sec 30
4.3.2 Mechanical properties
Although as-sintered scaffolds can be handled without crumbling, they are not
particularly strong. As shown in Fig. 4.7 (A), the typical compressive stress-strain curve of
uncoated BG scaffold exhibited a jagged behavior due to the presence of micro-cracks on
the as-sintered struts. The compressive strength and modulus were low, in the average of
0.021 ± 0.003 MPa and 0.05 ± 0.01 MPa, respectively. By coating Bioglass®-based scaffolds
with different polymers, the compressive strength and elastic modulus were significantly
improved (* p 0.05) due to the reduction of micro-cracks on the struts’ surfaces by polymer
filling. The compressive stress-strain curve of all types of polymer coated scaffolds showed
larger area under the curve compared to that of uncoated scaffolds. This behavior means a
higher mechanical stability of the scaffolds (in comparison to uncoated scaffolds) for
handling and further in vitro acellular and cellular evaluations. All biodegradable polymers
led to improvement of the mechanical properties of the base scaffolds in the similar range of
0.1 - 0.3 MPa for the compressive strength. The optimized synthetic polymer coatings
(PDLLA and PHBHHx) led to significantly higher scaffold elastic modulus (# p 0.05) than
the optimized natural polymer coatings (Alg and Gel) (Fig. 4.7 (B)), which is the result of the
intrinsic characteristics of the synthetic polymers concerning higher mechanical properties
compared to natural polymers. Even though the natural polymers, both Alg and Gel,
exhibited an homogeneous coating and also a thicker coating layer onto the scaffolds, the
58
thin PDLLA and PHBHHx coatings might be more effective in terms of crack-filling behavior
[272]. Generally, natural polymers have a limitation due to their relatively low mechanical
properties and they are not suitable for use in load bearing applications. The suggested
behavior was supported by the appearance of the specimens after the compression strength
test (Fig. 4.7 (C)). Uncoated Bioglass-based scaffolds, which are brittle and have jagged
stress-strain curve, were completely crushed and reduced to powder after compression test,
while the polymers coated Bioglass-based scaffolds did not crumble and showed though
fracture behavior. For example, PDLLA- and PHBHHx-c-BG exhibited higher elasticity
(higher elastic modulus). After compressive load, the scaffolds could recover partially their
initial dimensions (i.e. up to 70 % by height of the specimen). Especially, PHBHHx coated
scaffolds showed elastomeric mechanical behavior, similarly to the study of Deng et al.
[233]. The present PHBHHx-c-BG scaffolds could recover suddenly their initial dimensions
after removing the force. In contrast, natural polymer coated Bioglass®-based scaffolds
showed different fracture behavior, e.g. after loading, the struts were broken and compacted.
This behavior means that these specimens absorbed force but they could not transfer force.
The specimens were separated apart in small pieces in the cases of Alg- and Gel-c-BG
scaffolds. These qualitative results thus clearly confirm that the PDLLA and PHBHHx
coatings in combination with the optimized coating conditions impart better improvement in
terms of mechanical properties of scaffolds due to the characteristics and mechanical
properties of polymers coating layer themselves and the ability to fill the cracks.
59
Figure 4. 7 (A) Representative compressive stress-strain curves of biodegradable polymer
coated Bioglass®-based composite scaffolds in comparison with the curve of uncoated
Bioglass®-based scaffolds, (B) normalized mechanical properties (compressive modulus and
strength) of coated scaffolds compared to those of uncoated scaffolds (as reference) and (C)
appearance of coated scaffolds after compression load. * indicates significantly different
mechanical properties of coated scaffolds in comparison with those of uncoated scaffolds,
and # indicates significantly different mechanical properties of synthetic polymer coated
scaffolds in comparison with the properties of natural polymer coated scaffolds (p 0.05).
4.3.3 Degradation behavior
The biodegradation behavior of Bioglass®-based scaffolds was investigated. This
parameter is very important because the biodegradation rate of scaffolds has an impact on
the success of tissue regeneration. In this study, the degradation behavior of uncoated and
coated Bioglass®-based scaffolds was monitored by weight loss and pH variation upon
immersion in SBF solution. The weight loss of uncoated Bioglass-based scaffolds (Fig. 4.8
(A)) increased with culturing time. Upon immersion for 7 days, uncoated Bioglass®-based
60
scaffolds lost gradually their weight up to 25 %. From 7 to 28 days, the rate of weight loss
was reduced to 10 %. This result can be explained that at the initial immersion time, the
dissolution of Bioglass surface took place, leading to the significant weight loss. On the
other hand, HA formed on the surface of Bioglass after longer immersion in SBF, which is
initiated by ionic exchange between Bioglass surface and SBF (as detailed in Chapter 2).
Therefore, the weight loss of Bioglass-based scaffolds was compensated with the formation
of HA after 7 days of immersion. When compared to the pH study in SBF (Fig. 4.8 (B)), the
pH value of the SBF solution after 1 day of immersion was seen to be significant increase
from 7.4 to 8.5, which was attributed to the ionic exchange between the ions from Bioglass
(such as Si4+
and Na+) and H
+ and H3O
+ from SBF. Subsequently, the pH value was reduced
to 7.8 after 7 days of immersion and it remained constant in the range of 7.6 – 7.8 upon 28
days immersion. This pH variation corresponded to the trend of weight loss.
The weight loss of biodegradable polymer coated scaffolds increased with culturing
time and it was seen to vary with the types of polymer coating. All scaffold types could be
classified into two main trends, including water soluble polymers (Alg and Gel) and water
insoluble polymers (PDLLA and PHBHHx). Both Alg- and Gel-c-BG scaffolds showed high
resorption rate in the immersion time investigated. After 14 days immersion in SBF, Gel-c-
BG scaffolds maintained fast resorption rate and lost around 70 % of their weight at day 28,
Fig. 4.8 (A), whereas Alg-c-BG scaffolds maintained constant resorption rate in the period
between 14 days until 28 days immersion, which showed a weight loss of 60 %. This result
demonstrated that Gel coating showed the fastest bioresorption rate among all polymers
investigated. This behavior is likely caused by the hydrophilic nature of Gel, leading to high
ability of water uptake, which is one of the important factors influencing biodegradation rate.
In addition, the incubation conditions at 37 °C can influence the dissolution of gelatin. On the
other hand, Na-Alg can be partly crosslinked with Ca2+
in SBF forming Ca-Alg; therefore, this
phenomenon enhances the Alg network stability and retards the dissolution of Alg coating
during immersion in SBF [273].
In contrast, PDLLA and PHBHHx, which exhibit hydrophobic characteristic, tend to
degrade initially by a hydrolysis process. Both PDLLA and PHBHHx showed lower
61
degradation rate compared to Alg and Gel. The biodegradation behavior of PDLLA- and
PHBHHx-c-BG scaffolds exhibited the same trend (Fig. 4.8 (A)), indicating that both aliphatic
polyesters provide the same degradation profile. Also, the degradation profile indicated by
weight loss as function of immersion times corresponded to pH variation in all polymers
investigated, as shown in Fig. 4.8 (B). At day 1 in immersion in SBF, pH values of all
investigated coated scaffolds were significant lower compared to the value of uncoated
scaffolds. This phenomenon can be explained that the polymer coating may inhibit the ion
release of Bioglass, and therefore lead to reduced pH value. Furthermore, it is likely that
after 3 days of immersion, acidic groups resulting from degradation of each polymer, in
particular PDLLA and PHBHHx coatings, decreased the pH value of SBF soultion, while Si
and Ca ions released from the Bioglass® surface compensated the pH decrease. The pH
values in the case of both PDLLA- and PHBHHx-c-BG scaffolds maintained in the range of
7.1 – 7.2 after 7 days of immersion. In contrast, Gel and Alg coatings led to higher pH values
than PDLLA and PHBHHx coatings, which is likely related to the different acidicity of each
degradation product resulting from each polymer. Therefore, the pH of the SBF solution was
influenced by the degradation rate of the polymer coating and the dissolution behavior of the
Bioglass®. However, different polymers provided different degradation rates and the range of
pH values of different polymer coated Bioglass® scaffolds was different during the immersion
time.
62
Figure 4. 8 (A) % weight loss of biodegradable polymer coated 45S5 Bioglass®-based
scaffolds after 1, 3, 7, 14 and 28 days of immersion in SBF and (B) variation of the pH of the
SBF solution.
4.3.4 In vitro bioactivity
(i) Uncoated Bioglass scaffolds
The in vitro bioactivity of as-sintered Bioglass®-based scaffolds was assessed by
investigating the formation of HCA on scaffolds’ surfaces upon immersion in SBF for 1, 3, 14
and 28 days. Fig. 4.9 shows the formation of HCA on struts’ surface after immersion in SBF
for 28 days. It was observed that apatite started to form after immersion in SBF for 1 day
(Fig. 4.9 (A)). HCA formation was more clearly observed on the surface of the scaffolds after
immersion in SBF for 3 days (Fig. 4.10 (B)). After 14 days of immersion in SBF, aggregation
of HCA was observed and the surface of the scaffolds was completely covered after 28 days
in SBF (Fig. 4.9 (C and D)).
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Figure 4. 9 SEM images of as-sintered 45S5 Bioglass®-based scaffolds after (A) 1, (B) 3,
(C) 14, and (D) 28 days of immersion in SBF, showing possible formation of HA indicated
qualitatively by the morphology of the deposited structures.
Apatite formation was confirmed by FTIR spectroscopy (Fig. 4.10 (A)). Possible
apatite formation was evidenced by generation of amorphous and crystalline HA phase,
which was expressed as a double peak at 605 cm-1
and 565 cm-1
(P-O bending)
[37,274,275], as depicted by ▲ in Fig. 4.10. Moreover, a double broad peak at 1488 cm-1
and 1420 cm-1
corresponding to amorphous CaP phase was detected after 1 days of
immersion in SBF [37], as depicted by ■ in Fig. 4.10 (A).
In addition, XRD analysis confirmed the formation of HA (Fig. 4.10 (B)). The peak of
HA phase was detected at 2 32 [32,37,275]. The height of the HCA peak increased with
increasing times in SBF, e.g. between 14 and 28 days, indicating that higher amount of HCA
occurred with increasing immersion time in SBF. In contrast, sharp diffraction peaks of
Na2Ca2Si3O9, which is the crystalline phase in as-sintered scaffolds reduced with immersion
time and they nearly disappeared after 28 days of immersion in SBF. This result indicates
the reduction of the crystallinity of as-sintered 45S5 Bioglass®–based scaffolds with
increasing culturing time in SBF, which has been previously reported in literature [32,275].
64
Figure 4. 10 (A) FTIR spectra of as-sintered 45S5 Bioglass®-based scaffolds after different
immersion times in SBF, in comparison with the scaffolds before immersion. The
characteristic peaks of HA are marked by ▲ and ■, and (B) XRD patterns of as-sintered
45S5 Bioglass®-based scaffolds after different immersion times in SBF, in comparison with
the scaffolds before immersion. The major peak of HA is marked by , while the crystalline
peaks of Na2Ca2Si3O9 were marked by ■.
(ii) Alg coated scaffolds
After coating with biodegradable polymers, each type of polymer coatings showed
slightly different behavior when immersed in SBF. For example, after 1 day immersion in
SBF, the Alg coating started to peel off the struts, as depicted by the arrow in Fig. 4.11 (A).
There was no apatite formation observed after 1 day on the Alg coated scaffolds, which was
confirmed by FTIR spectroscopy, as shown in Fig. 4.12 (A). After 1 day immersion in SBF,
there was no appearance of P-O bending mode (at 560 - 550 cm-1
and 610 - 600 cm-1
),
which are the characteristic peaks of HA. This behavior was not different from that of
uncoated scaffolds, e.g. the possible apatite formation was clearly observed after 3 days
immersion in SBF (Fig. 4.9 (B)). It can be concluded that Alg coating does not affect the
bioactivity of Bioglass®. Moreover, after 3 days immersion in SBF, the HCA formation was
detected by FTIR spectroscopy, as depicted by ▲ in Fig. 4.12 (A), even though it was not
clearly observed by SEM (Fig. 4.11 (B)). Finally, after 28 days, it was found that the Alg
65
coating still maintained the adherence to the scaffold, which is probably caused by
crosslinking between Alg coating and Ca ions in SBF, and thus retarding the dissolution of
Alg, as previously described in Section 4.3.3. Also, HA covered the struts’ surfaces (Fig.
4.11 (C)). The FTIR spectra of coated scaffolds after 28 days exhibited high intensity of the
HA characteristic peaks at 604 and 563 cm-1
attributed to P-O bending, as marked by ▲ in
Fig. 4.12 (A). In addition, the presence of a double broad peak at 1453 and 1420 cm-1
attributed to amorphous CaP phase [37], confirmed the HA formation after 28 days
immersion in SBF. The characteristic peak of Alg coating, in particular at 1638 cm-1
(-C=O
stretching) as marked by ■, maintained after 28 days in SBF, indicating the lasting of Alg
coating. The results of the in vitro bioactivity study of Alg-c-BG scaffolds agreed with the
data reported by Erol et al. [54] in that Alg coating has no negative effect on the bioactivity of
Bioglass®-based scaffolds.
(iii) Gel coated scaffolds
The Gel coating was dissolved rapidly in SBF due to the high bioresorption rate of
gelatin [276]. Therefore, the Gel coating should be completely degraded after only a few
days in SBF as stated by Bellucci et al. [276]. However, FTIR spectra shows that after 3
days immersion in SBF, Gel was detected by the presence of the characteristic peak of –
C=O and –C-N- stretching at 1637 cm-1
(as marked by ■ in Fig. 4.12 (B)) [78,276]. With
increasing time in SBF, the intensity of the characteristic peak of Gel was reduced with time
and the peaks almost disappeared after 28 days of immersion. At the same time, HA was
formed following the fracture of the coating, which allowed ions release from the Bioglass®
substrate. Apatite crystals were not clearly observed by SEM after immersion for 3 days,
while after 28 days immersion such crystals were visible (Fig. 4.11 (D, E and F)). Moreover,
apatite formation was confirmed by the presence of peaks at 604 and 565 cm-1
attributed to
the P-O bending mode of crystalline apatite and amorphous CaP, respectively. The peaks of
P-O bending were clearly detected on samples after 14 days immersion in SBF,
incorporation with an existence of double broad at 1464 and 1420 cm-1
attributed to
amorphous CaP phase (Fig. 4.12 (B)).
66
(iv) PDLLA coated scaffolds
Synthetic polymer coated Bioglass®-based scaffolds exhibited different behavior
related to HA formation in comparison to natural polymer coated scaffolds. Due to the fact
that synthetic polymers have slower degradation rate in comparison with natural polymers,
the PDLLA and PHBHHx coatings maintained adherence with the scaffold during the
complete period of 28 days immersion in SBF. This behavior should extensively retard the
rate of ion release from the bioactive phase and consequently HA formation should be
delayed. It is likely that apatite crystals formed on both, PDLLA-c-BG and PHBHHx-c-BG
scaffolds, after 3 days immersion in SBF (Fig. 4.11 (G and J)). The reason for this behavior
could be the hydrolytic (erosion) mechanism [268,277]; active in these polymers, as shown
in Fig. 4.11 (G) in the case of PDLLA-c-BG scaffolds. The polymer coating was perforated in
contact with SBF, hence providing channels for leakage of ions from the Bioglass® substrate
and exposing nucleation sites for HA formation. After 28 days immersion, HA covered the
entire surface of the scaffold, as reported also by Bretcanu et al. [55] for PDLLA coated
scaffolds. The appearance of PDLLA and HCA formation after immersion in SBF for different
time points was confirmed by FTIR spectroscopy (Fig. 4.12 (C)). Upon 28 days in SBF, the
characteristic peaks of PDLLA were maintained, including the peaks at 1750 cm-1
attributed
to -C=O and at 1659 cm-1
attributed to –COO- [278], thus confirming the existence of PDLLA
coating. HCA phase was evidenced by the presence of the characteristic peaks at 603 and
565 cm-1
attributed to P-O bending, as marked by ▲ in Fig. 4.12 (C). In addition, a broad
double peak at 1454 and 1422 cm-1
attributed to amorphous CaP phase was gradually
increased, while the –C=O peak at 1751 cm-1
, as marked by ■, indicating PDLLA coating
decreased in the intensity with increasing time in SBF.
(v) PHBHHx coated scaffolds
The PHBHHx-c-BG scaffolds exhibited possible HA formation similarly to the
PDLLA-c-BG scaffolds. PHBHHx is a polyester-based polymer, which initially degrades by
hydrolytic mechanism [277,279]. In Fig. 4.11 (J), the channels caused by the fracture of the
polymer coating are visible, which were enabled the nucleation of HCA crystals. After 28
days immersion in SBF, the PHBHHx coating maintained adherence to the scaffold, as
67
showed by an inset in Fig. 4.11 (L). HA fully covered on the PHBHHx coated Bioglass®
surface. The bioactivity of PHBHHx-c-BG was confirmed by FTIR spectroscopy (Fig. 4.12
(D)); the characteristic peaks of HA was detected at 603 and 564 cm-1
, according to P-O
bending mode and a broad peak at around 1480 - 1420 cm-1
, according to amorphous CaP
phase. It was obviously detected after 3 days immersion in SBF and the intensity of the
peaks increased with increasing immersion time. At the same time, the peaks of PHBHHx
coating at 1750 and 1645 cm-1
, attributed to –C=O and –COO-
[280,281], respectively,
broadened with increasing immersion time, indicating partly degradation of PHBHHx coating.
(vi) Summary of results on the bioactivity
All investigated biodegradable polymer coated Bioglass®-based scaffolds were
confirmed to maintain the bioactive properties of Bioglass® by the formation of HA in contact
with SBF, even though they exhibit different behavior among them in terms of degradation
rate and formation of apatite. The PDLLA and PHBHHx coatings showed fast rate of HA
formation due to the fact that these coating were very thin and due to the occurrence of
hydrolytic biodegradation in contact with SBF. On the other hand, the resorption of natural
polymer coatings was initiated by water absorption, subsequent swelling and polymer
detaching from Bioglass® substrate occurred. HA formed by the channels/holes formed after
the polymer coating detached, which resulted in an open surface for ion release. This result
demonstrated that the selection of the polymer coating must be taken into consideration to
tailor degradation kinetics of composite scaffolds. Another factor is the quality of coating, i.e.
inhomogeneity of polymer coating, which initiates the pathway for ion release of Bioglass
substrate. Moreover, a suitable composite scaffold can be selected for controlled drug
release by tailoring the polymer coating layer as a drug carrier-system, thus leading to a
multifunctional scaffold for tissue engineering therapeutics [96,282]. This investigation is
reported in Chapter 5.
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Figure 4. 11 SEM images of biodegradable polymer coated 45S5 Bioglass®-based
composite scaffolds after immersion in SBF, showing formation of HA: (A, B, C) Alg-c-BG,
(D, E, F) Gel-c-BG, (G, H, I), PDLLA-c-BG and (J, K, L) PHBHHx-c-BG scaffolds for 1, 3 and
28 days, respectively.
69
Figure 4. 12 FTIR spectra of biodegradable polymer coated 45S5 Bioglass®-based scaffolds
after 28 days of immersion in SBF: (A) Alg-c-BG, (B) Gel-c-BG, (C) PDLLA-c-BG, and (D)
PHBHHx-c-BG scaffolds. The characteristic peaks of HA are marked by ▲ and the
characteristic peaks of polymer coatings are marked by ■.
70
4.4 Conclusions
Biodegradable polymer coated 45S5 Bioglass®-based scaffolds were successfully
fabricated in order to develop scaffolds with improved mechanical properties and stability. All
polymer coatings studies led to significant improvement of compressive modulus, strength
and toughness of the bare scaffolds. By comparison between biodegradable natural (Alg
and Gel) and synthetic (PDLLA and PHBHHx) polymers, each coating type showed
differences in the morphology of the scaffolds’ surface, in biodegradation behavior and
subsequently in vitro bioactivity. However, it was proven that all investigated polymer
coatings did not affect the interconnectivity, pore structure and even the in vitro bioactivity of
scaffolds. Coating of scaffolds by using PDLLA and PHBHHx led to structures of higher
mechanical properties in comparison to those of scaffolds coated by Alg and Gel. Therefore,
it can be concluded that the composite scaffolds developed in this study show the potential
for use in bone regeneration based on their physical and mechanical properties, and
bioactivity. In addition, biodegradable polymer coated Bioglass-based composite scaffolds
are attractive for controlled drug release applications, which is another aspect of this project.
Finally, in vitro cell culture study should be performed as reported in Chapter 8, in order to
evaluate the biological properties of the new scaffolds.
CHAPTER 5
Development of 45S5 Bioglass®-based Scaffolds for Controlled
Antibiotic Released in Bone Tissue Engineering via Biodegradable
Polymer Layered Coating
5.1 Introduction
Recent developments in bone tissue engineering provide alternative approaches for
the repair of bone defects caused by trauma and infection [93,283]. Bone repair scaffolds
associated with the application of drugs (i.e. antibiotics and antitumoral medicaments) and
growth factors attract increasing attention to avoid infections, regulate cell growth and
develop bone regeneration [268,284,285]. Basically, bone scaffolds should be
biocompatible, bioresorbable, osteoconductive and possibly they should act as a drug carrier
[283,285–289]. Scaffolds are usually made from tailored combinations of inorganic and
organic phases, which are chosen aiming at replicating the structure and composition of
bone tissue [283,289,290]. Bioactive glasses [95,286–288] and bioceramics [97,216,291–
293], as the inorganic component in composites satisfy the desirable property of bioactivity
for the application in bone tissue engineering approaches. Several natural- and synthetic-
derived biodegradable polymers have been explored as the organic component for
development of composite scaffolds. Specific studies on composite scaffolds with drug
delivery ability have been discussed in Chapter 2.
According to the study presented in Chapter 4, the tailored biodegradable polymer
coated Bioglass-based scaffolds were aimed to develop in as a multifunctional scaffold for
bone tissue engineering applications and were presented in this chapter. By combining the
advantages of synthetic and natural-derived polymers, two polymer coating layers were
applied onto the Bioglass-based scaffold with different purposes, including (i) to improve
the mechanical properties of Bioglass-based scaffolds by first coating with synthetic
PDLLA/P123 copolymer and (ii) at the same time to add extra-functionality on the scaffolds
72
by second coating with tetracycline antibiotic-loaded natural polymers such as Alg and Gel.
PDLLA was chosen as a coating material on the scaffold due to its good mechanical
properties as shown in Chapter 4. However, PDLLA, which is a hydrophobic polyester, is not
chemically compatible with water soluble antibiotics, thus it exhibits a low drug entrapment
efficiency [294]. Therefore, natural-derived biodegradable polymers, including Alg and Gel,
were selected as the drug carriers. The polymer networks of both Alg and Gel are able to
load a wide range of bioactive substances, cells and drug molecules with minor interaction
between the drug and the polymer matrix [295]. In addition, natural polymers are superior
biocompatible, which facilitate the adhesion and proliferation of cells (i.e. osteoblasts) [289].
Based on the fact mentioned above, Alg and Gel were investigated as a drug carrier in the
present study. However, the layered coating cannot be formed as desired because of the
incompatibility of PDLLA, and Alg and Gel coating layers. Therefore, we aimed to overcome
this problem by modifying surface chemistry of PDLLA coating with an amphiphilic polymer
blending (i.e. P123 copolymer). Furthermore, the mechanical properties and drug release of
the multifunctional scaffold (Alg- vs. Gel-drug carriers) were investigated.
5.2 Materials and methods
5.2.1 Fabrication of TCH-loaded layered biodegradable polymer coated Bioglass®-
based scaffolds
45S5 Bioglass®-based scaffolds were prepared following the same procedure
detailed in Chapter 4 by using the foam replication method. For the preparation of polymer
coatings, PDLLA (Purac Biomaterials, Gorinchem, Netherland) was dissolved in DMC with a
concentration of 5 wt/v % at room temperature while stirred for 2 h. Then PEG-PPG-PEG
triblock copolymer (Pluronic P123, Mn 5800 Da; Sigma) was added into the PDLLA
solution with the PDLLA to P123 ratio of 90/10 wt %. The mixture was continuously stirred
until P123 was completely dissolved. Afterwards, a 45S5 Bioglass® scaffold was immersed
in the polymer solution during 5 min. Subsequently the scaffold was removed from the
solution and dried at room temperature for 24 h. An obtained coated scaffold was labeled as
PL/P123-c-BG.
73
Tetracycline hydrochloride (TCH)-loaded Alg and Gel solutions were prepared as
follows: Na-Alg (Mw 200000 Da; Sigma) was dissolved in DI H2O with a concentration of
1.5 wt/v % at room temperature while stirred for 2 h. A concentration of 1.5 wt/v % Gel (type
A from porcine skin with 300 g bloom; Sigma) was dissolved in DI H2O at 50 C while stirring
for 1 h. Then 375 µg/ml TCH (C22H24N2O8 · HCl; Appli Chem GmbH, Darmstadt, Germany)
was added into both Alg and Gel solutions. Finally, the PL/P123-c-BG scaffold was
immersed in the TCH-loaded Alg and Gel solutions contained TCH (5 ml of solution/scaffold)
for 5 min and dried at room temperature for 24 h. The drug-loaded scaffolds were labeled as
T-Alg-c-(PL/P123-c-BG) and T-Gel-c-(PL/P123-c-BG) for Alg and Gel as the drug carriers,
respectively. TCH loaded uncoated scaffolds were prepared as control, by dipping uncoated
scaffolds in TCH/DI H2O with a concentration of 375 µg/ml for 5 min. Then the scaffold was
taken out and dried at room temperature for 24 h. These samples were labeled as T-BG.
5.2.2 Characterization and testing
(i) Capillarity test
In order to evaluate the porosity and surface property of the polymeric coatings, a
qualitative capillarity test was performed according to [296]. Briefly, the TCH-loaded
polymeric coating solution, which served as a testing fluid, was prepared as described in
section 5.2.1. In this case, TCH-loaded Alg solution, which exhibits a yellow color, was
added in a petridish. Then a coated scaffold was slowly placed on the surface of the
solution, while the testing time was recorded until the scaffold was completely wet (the fluid
went up through the entire porous network of the scaffold). PDLLA-c-BG and PL/P123-c-BG
scaffolds were tested in order to compare the surface property of the different polymeric
coatings.
(ii) Contact angle measurement
In order to evaluate the hydrophilicity of each polymeric coating, the wettability of
pellets prepared with the same conditions as the 3D scaffolds was measured using a water
contact angle instrument (DSA30, Kruess, Germany). The pellets were prepared as follows:
0.3 g of Bioglass® powder were added in a stainless steel die (diameter: 10 mm) and pellets
were obtained by cold uniaxial pressing using an electro-hydraulic press (MAUTHE
74
MASCHINENBAU PE-010; Wesel, Germany) working at a load 4 × 104 N. The obtained
pellets were sintered using the same conditions used for porous Bioglass® scaffolds. As-
sintered Bioglass® pellets were coated with TCH drug, PL/P123, and TCH loaded Alg and
Gel following the same procedure described above for the scaffolds.
(iii) Microscopy
The microstructure of the scaffolds was characterized by SEM (LEO 435VP from
Zeiss Leica). The scaffolds were cross-sectioned by using a razor blade. The samples were
then sputter-coated and observed at an accelerating voltage of 10 kV.
(iv) Chemical analysis
The chemical structure of the scaffolds was investigated by using FTIR (Nicolet
6700). Bioglass®-based scaffolds were grinded and the obtained powder was mixed with
potassium bromide (KBr) powder in a weight ratio of 1/300 (scaffold/KBr). The mixture was
pressed into a pellet by using an electro-hydraulic press (MAUTHE MASCHINENBAU PE-
010; Wesel, Germany) at a load of 105 N. Then, pellets were measured by using FTIR in
transmission mode with the resolution of 4 cm-1
in the wavenumber range of 4000 - 400 cm-1
.
(v) Mechanical testing
Polymer coated cubic shaped Bioglass® scaffolds with dimensions of 8 mm × 8 mm
× 8 mm were tested under compression deformation mode by using a universal testing
apparatus (Zwick Z050). The cross-head speed used was 2 mm/min, the preload was 0.1 N
and the maximum load was 50 N. Stress-strain curves were recorded to determine the
mechanical properties. The elastic modulus was calculated from the initial linear slope of the
compressive stress-strain curves, while the compressive strength was obtained from the
maximum stress before the specimens collapse. Eight specimens were tested and the
results are presented as average ± SD.
(vi) In vitro drug release profile
The in vitro drug release behavior of the scaffolds, including T-BG, T-Alg-c-
(PL/P123-c-BG) and T-Gel-c-(PL/P123-c-BG) scaffolds, with dimensions of 8 mm × 8 mm ×
8 mm were evaluated. Each scaffold was immersed for up to 14 days in a glass vial
containing 5 ml of PBS (0.1 M; Sigma) solution at 37 C and pH 7.4. At given interval times,
75
2 ml of PBS was taken and replaced with fresh PBS. The absorbance of the drug containing
PBS at the wavelength of 362 nm was measured by using a UV spectrophotometer
(Specord 40; Analytikjena, Germany). Then, the amount of drug released was determined by
using a linear relation between absorbance and known concentrations of TCH (2.5 - 100
µg/ml), as given in Eq. 5.1:
Absorbance = [0.0268 x concentration (µg/ml)] – 0.1206, R2 = 0.99. (5.1)
The amount of drug release is reported as a percentage of cumulative drug release ± SD
with respect to the immersion time.
At the same time, in order to confirm the bioactivity of scaffolds after coating with
two polymer layers, T-Alg- and T-Gel-c-(PL/P123-c-BG) scaffolds with dimensions of 8 × 8 ×
8 mm3 were investigated. Each scaffold was placed in a polystyrene bottle containing 50 ml
of SBF solution at 37 C and pH 7.4 [297]. After 14 days of immersion, the scaffold was
taken, washed twice with DI water and dried at room temperature. Afterwards, possible HA
formation and also the morphological changes were analyzed by using SEM.
5.2.3 Statistical analysis
The data were analyzed by using one-way ANOVA analysis and Turkey’s multiple-
comparison test to determine statistical differences. A confidence interval of 95 % (p = 0.05)
was used for all analysis.
5.3 Results and discussion
5.3.1 Surface property of polymeric coatings
A double coating layer based on PDLLA, and Alg and Gel was designed to be
applied on the 45S5 Bioglass®-based scaffolds to impart multifunctionality, e.g. drug delivery
capability. The challenge in developing such synthetic-natural polymer layered coatings is
the difference of surface chemistry between the hydrophobic PDLLA and hydrophilic Alg or
Gel limits their compatibility. As a result of polymer incompatibility, the TCH-loaded Alg and
Gel could not infiltrate the porous structure of unmodified PDLLA-c-BG scaffold. This
observation was qualitatively confirmed by the capillarity test [296] (Fig. 5.1), using the
TCH-loaded Alg solution (yellow color) as a test solution. The test involves the determination
76
of porosity and capillarity of surface [296], and it revealed the lack of capillarity (wettability) in
the case of PDLLA-c-BG scaffolds. The scaffolds were seen to remain on the surface of the
coating solution (see Fig. 5.1 (Aa and Ba)). Consequently, the TCH-loaded coating solution
could not infiltrate the pore structure and the coating of the struts of PDLLA-c-BG scaffold
was not successful. In order to overcome this problem, a modification of the surface
chemistry of PDLLA-c-BG scaffold was necessary. The approach developed in this study
involved the addition of P123 copolymer in order to increase the hydrophilicity of the PDLLA-
c-BG scaffold, leading to hydrophilicity matching to that of Alg and Gel. The P123 copolymer
was utilized because it contains both hydrophilic (PEG) and hydrophobic (PPG) chains,
which can be homogeneously blended with PDLLA in DMC solution. By using this approach,
the capillarity was obvious in the case of PL/P123-c-BG scaffolds, showing wettability
increase and that the TCH-loaded coating solution ascended through the whole pore
network of the scaffold in few seconds (see Fig. 5.1 (Ab and Bb)).
The increase in the hydrophilicity of PDLLA-c-BG scaffolds was also confirmed by
the water contact angle values, as shown in Fig. 5.2. After coating with PDLLA/P123 blend,
the contact angle value of the scaffolds was significantly decreased (from 74.3 ± 0.2 for
pure PDLLA-c-BG to 31 ± 2 for PDLLA/P123-c-BG) to nearly the values of T-Alg and T-Gel
coatings (38.0 ± 0.5 and 39 ± 0 , respectively). These results prove that the blend of
PDLLA and P123 copolymer can drastically modify the surface chemistry of pure PDLLA
and thus the T-Alg and T-Gel solutions can efficiently infiltrate the porous structure of
Bioglass-based scaffolds, forming layered polymer coatings on the Bioglass-based
scaffolds, as desired in this study.
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Figure 5. 1 (A) Scheme of the capillarity test of Bioglass®-based scaffolds, showing the
effect of surface chemistry on the permeability of the porous scaffolds and (B) photographs
representing the coated scaffolds during the capillarity test.
Figure 5. 2 Contact angles of Bioglass®-based scaffolds showing the surface wettability of
different coatings. * indicates the significant difference (p 0.05) of the modified coatings on
the Bioglass-based scaffolds in comparison with PL-c-BG scaffolds.
78
5.3.2 Morphology
Fig. 5.3 shows SEM micrographs at different magnifications, of different coated
scaffolds. The morphology of the scaffolds after coating with PDLLA/P123 blend is shown in
Fig. 5.3 (A and B). The surface of the coated scaffold was homogeneous and smooth
compared to the surface of T-BG scaffolds (Fig. 5.3 (B) and Fig. 5.3 (D)), which might be the
result of the used P123 copolymer, considering that P123 copolymer shows an ability to
enhance the rheological property of polymers [298]. It is thus obvious that the polymer
homogeneously covered the entire strut even though some uneven areas could be
observed, as shown in Fig. 5.3 (B). After coating with TCH-loaded Alg and Gel as second
coating layers, as shown in Fig. 5.3 (E and G), even though the color of the scaffolds
became yellow, no morphological changes of the struts were observed by SEM (both T-Alg-
and T-Gel-c-(PL/P123-c-BG) scaffolds) compared to the PL/P123-c-BG scaffolds (Fig. 5.3
(B)). In detail, a fairly homogeneous coating not inducing blocking of pores was observed
(see Fig. 5.3 (E and G)). However, at higher magnification (Fig. 5.3 (F and H)), the rougher
surfaces of the TCH-loaded polymer coatings could be observed in comparison with
PL/P123-c-BG scaffolds (Fig. 5.3 (B)). This observation was the same as reported in the
previous study of Mouriño et al. [299]. It was found that the surface of scaffolds became
rougher after coated with the second layer of Alg [299]. It is probably caused by the fact of
polymer agglomeration during drying. In detail, the alginate coating, for example, took longer
time to be dried compared to the PDLLA/P123 coating. By this fact, the Alg coating may tend
to gather together during drying. However, the obtained rough surface is believed to be
suitable for cell adhesion and proliferation.
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Figure 5. 3 SEM images of the scaffolds showing the morphological porous structure and
morphology of a coating surface of: (A, B) PL/P123-c-BG scaffolds, (C, D) T-BG scaffolds,
(E, F) T-Alg-c-(PL/P123-c-BG) scaffolds and (G, H) T-Gel-c-(PL/P123-c-BG) scaffolds.
5.3.3 Mechanical properties
The typical stress-strain curves in Fig. 5.4 (A) as well as the normalized compressive
strength and modulus in Fig. 5.4 (B) illustrate the improvement of the elastic modulus and
the compressive strength of Bioglass®-based scaffolds by layered polymeric coating. In
detail, the scaffolds coating with PDLLA/P123 (curve b in Fig. 5.4 (A)) exhibited an
improvement of the mechanical properties up to 10 times in comparison with the uncoated
scaffolds as depicted as curve (a) in Fig. 5.4 (A). This can be assigned to the formation of a
uniform PDLLA/P123 coating on the struts, as well as to the filling of any cracks on the
80
surface of the struts, ceasing the crack propagation occurring on Bioglass®-based scaffolds
under load. It is likely that the mechanical strength of polymer coated Bioglass-based
scaffolds in the present study was higher, compared to the values reported in the previous
studies [57,282]. It can be described by the fact that the Bioglass-based scaffolds in the
study of Li et al. [57], for example, were partially coated with polymer, while the polymer was
fully covered the struts of Bioglass-based scaffolds in the present work. Moreover, the
second layer coated with the TCH-loaded Alg and Gel did not enhance further the
mechanical properties of the scaffold. As represented in the compressive stress-strain
curves of both T-Alg-c- and T-Gel-c-(PL/P123-c-BG) scaffolds (curves (c) and (d) in Fig. 5.4
(A)), that their stresses responsed to the strain in the similar trend to the curve of PL/P123-c-
BG scaffolds. The reason for this can be that only a thin layer of Alg and Gel is formed due
to the low polymer concentration used. According to these results, the mechanical properties
of the layered polymer coated Bioglass-based scaffolds predominantly occupied by the
coating layer of PDLLA/P123 blend are confirmed.
Figure 5. 4 Mechanical properties of polymer coated Bioglass-based scaffolds: (A)
representative compressive stress-strain curves and (B) average elastic modulus and
average compressive strength of the scaffolds. * (p 0.05) indicates the statistical
significance of compressive mechanical properties of coated scaffolds, compared to those of
uncoated T-BG scaffolds.
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5.3.4 Chemical structure
FTIR analysis was performed on coated scaffolds to confirm the existence of the
polymeric coating and the drug entrapment. First, the spectra of the Bioglass-based
scaffolds before and after drug loading without polymer carrier were considered (Fig. 5.5). In
detail, the spectrum of the T-BG scaffolds presents the characteristic peaks of Bioglass®,
including the double peaks at the wavenumber 1100 - 1040 cm-1
attributed to Si-O-Si
stretching mode [46,180], and the peak at the wavenumber of 458 cm-1
attributed to the Si-
O-Si bending mode [180]. The characteristic peaks of Bioglass were not changed after
loading with TCH, indicating that loaded TCH molecules do not initiate a chemical reaction
with Bioglass. This result means that TCH molecules loaded on the Bioglass scaffolds do
not lose their activity. Moreover, the double peaks at 3482 and 3350 cm-1
in the spectrum of
T-BG scaffold (see the inset (I) in Fig. 5.5), can be understood as an overlapping effect
between the –OH stretching broad peak of Bioglass (3700 - 3000 cm-1
), and double peaks
of TCH (3363 and 3304 cm-1
) and –CH stretching of phenol framework in TCH [287]. In
contrast, the spectrum of T-Alg-c-(PL/P123-c-BG) scaffold presents the peaks at 1620 and
1420 cm-1
assigned to –COO- asymmetric and symmetric stretching of Alg, respectively,
suggesting the existence of Alg in the coated scaffold. Moreover, the peak at 1753 cm-1
observed in the spectrum of T-Alg-c-(PL/P123-c-BG) scaffold is attributed to -C=O stretching
of PDLLA coating. As observed also in the spectrum of T-Gel-c-(PL/P123-c-BG) scaffold, the
-C=O stretching of PDLLA coating exists at the wavenumber of 1759 cm-1
. In addition, the
peak at 3435 cm-1
is assigned to –NH stretching of gelatin. Other detectable peaks at 1642
and 1456 cm-1
are assigned to –C=O and –C-N- stretching, and the peak at 1383 cm-1
is
assigned to –N-H- bending, as shown in the inset (II) in Fig. 5.5, which confirms the
existence of Gel coating. However, the characteristic peaks of TCH in the finger print region
(1500 - 500 cm-1
) are not obvious in the spectra of the coated scaffolds. Also, the change in
the peak position, in either Alg or Gel, is not represented in any spectra of the coated
scaffolds. Therefore, possible molecular interaction between the drug and the polymer
coating cannot be confirmed based on FTIR results.
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Figure 5. 5 FTIR spectra of TCH, BG, TCH-loaded Bioglass scaffolds and TCH-loaded
polymer coated Bioglass®-based scaffolds.
5.3.5 In vitro drug release
(i) Release profile
Fig. 5.6 (A) shows the cumulative percentage of TCH release from the Bioglass-
based scaffolds for up to 14 days of immersion in PBS. T-BG scaffolds showed an initial
burst release of 53 % at 1 h, which increased to 99 % in 4 h. This result confirms the low
drug binding affinity of the uncoated Bioglass-based scaffolds. In contrast, in polymer
coated scaffolds, lower initial burst release values (1 h) at 27 % and 22 % for Alg and Gel
coatings, respectively, were measured. Afterwards, both TCH-loaded polymer coated
scaffolds provided almost complete drug release ( 99 %) over 14 days in a sustained
behavior. This drug release kinetic is favorable as it should not only facilitate an effective
initial antibacterial effect but also promote long term protection against infection. Both Alg
and Gel carriers provided a similar release profile, including (i) an initial burst release as a
result of the release of the free drug presenting at the polymer coating surface and (ii) a
further relatively slow release induced by the drug molecules “protected” by the polymer
coating. As described in the literature [216,268,300], drug molecules embedded in the
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coating can diffuse through available pathways, i.e. pores and channels, into the medium.
Diffusion pathways can be influenced by the presence of an inhomogeneous coating
accompanied with the intrinsic degradation of the coating. Considering the result of the
degradation study (Fig. 5.6 (B)), Gel coated scaffolds were seen to exhibit slightly faster
degradation rate compared to Alg coated scaffolds. This result can be explained by the fact
that the Alg coating might partly crosslink with Ca ions in PBS [212], which should lead to
slower the degradation of the Alg coating. Also, the initial burst release of T-Alg-c-(PL/P123-
c-BG) scaffolds was slightly lower than that of T-Gel-c-(PL/P123-c-BG) scaffolds. In addition,
the possible interaction of the negatively charged drug (TCH) and the positively charged
polymer (Gel) was not observed in this study, and a superior binding affinity of TCH and Gel
cannot be confirmed by the results of drug release. Therefore, it seems that the key factors
influencing the release of drug from the natural polymer coated Bioglass-based scaffolds
are mainly related to coating homogeneity and dissolution/degradation of the polymer
coating. Compared to previous recent studies [57,282], the initial burst release of the TCH-
loaded Alg coating (22 %) in the present study was significantly lower, e.g. it was 63 % in
vancomycin-loaded CS coating [282] and 33 % in vancomycin-loaded PHBV [57]. However,
the different drug used in the present study should be taken into consideration. On the other
hand, TCH-loaded PCL/HA coated HA scaffolds have released, at the initial stage (1 h), 44
% of the load [291], indicating that the natural coating developed in the present study
enables better release control, reducing the initial burst release. The present approach
using natural polymers such as Alg and Gel as drug carrier seems to lead to superior
performance of the scaffold as drug delivery device in terms of water soluble drug
entrapment and protection of the drug, in comparison with systems based on synthetic
polymers. Another important feature of the present approach is that it is possible to de-
couple the mechanical stability function (provided by the synthetic polymer) from the drug
release function (provided by the natural polymer).
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Figure 5. 6 (A) Drug release profile and (B) degradation behavior of TCH-loaded polymer
(Alg and Gel) coated Bioglass scaffolds.
(ii) Morphology after drug release
SEM analysis was used to observe the morphological change of the scaffolds after
14 days of drug release (Fig. 5.7). The polymeric coating was partly maintained on both
scaffolds. In Fig. 5.7 (A)), it can be observed that the surface of the drug-loaded Alg coated
scaffold became rougher by the creation of many holes and channels, which indicates the
degradation of polymeric coating. The appearance of underneath smoother surface was also
observed, which probably is the PDLLA/P123 coating, as depicted by the arrow in Fig. 5.7
(A). It is also possible that the PDLLA/P123 coating started to be eroded during the
degradation of the outer drug-loaded alginate coating. These holes serve as a connection
pathway between the PDLLA/P123 coating and the liquid medium (PBS). The SEM image of
the T-Gel-c-(PL/P123-c-BG) scaffold in Fig. 5.7 (B) shows the generation of holes on the
coating surface, as shown previously on the surface of the T-Alg-c-(PL/P123-c-BG) coating
(Fig. 5.7 (A)). In contrast, the residual polymeric coating (depicted by a solid arrow) was the
PDLLA/P123 layer, while the outer drug-loaded Gel coating could not be distinguished. It is
likely that the Gel coated has degraded after the 14 d immersion in PBS. In addition, the
PDLLA/P123 coating was partly degraded and the surface of the strut could be also
observed (dashed arrow in Fig. 5.7 (B)). The release of TCH-loaded Alg and Gel carriers
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predominantly influenced by the degradation of the polymer coating is thus confirmed.
Moreover, it seems that the bioactivity of Bioglass was not inhibited by the layered polymer
coating, since amorphous PDLLA coating used exhibited partly degradation during
experimental study in PBS, as confirmed by the holes observed in Fig. 5.7 (A and B). By this
fact, dissolution of Bioglass can take place and this phenomenon leads to formation of HA
in SBF. As shown in Fig. 5.7 (C and D), formation of possible apatite crystals was observed
on the strut of the both scaffolds after 14 days of immersion in SBF. In particular, it was
obvious in the case of T-Gel-c-(PL/P123-c-BG) scaffold (Fig. 5.7 (D)) that the crystals
covered the strut of the scaffold, while the degradation of polymer coating took place (as
depicted by a solid arrow).
Figure 5. 7 SEM images of TCH-loaded polymer coated Bioglass-based scaffolds after in
vitro release in PBS for 14 days: (A) T-Alg-c-(PL/P123-c-BG); the arrows predicting the
PL/P123 coating and (B) T-Gel-c-(PL/P123-c-BG) scaffolds; dashed arrow depicting the
Bioglass® struts and solid arrow line predicting PL/P123 coating, and SEM images of TCH-
loaded polymer coated Bioglass-based scaffolds after immersion in SBS for 14 days: (C)
T-Alg-c-(PL/P123-c-BG and (D) T-Gel-c-(PL/P123-c-BG) scaffolds; solid arrows depicting
polymer coating.
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(iii) Chemical structure after drug release
The FTIR spectra shown in Fig. 5.8 enable to detect the chemical changes of the
scaffolds after 14 days of immersion in PBS. First, a new peak at wavenumber 1475 cm-1
is
observed as a broad double peak close to the peak of -COO- of Alg at 1420 cm-1
, in the
spectra of T-Alg-c-(PL/P123-c-BG) scaffold after immersion (see Fig. 5.8 (A)). The broad
double peak is assigned the overlapping of –COO- stretching band of Alg with –CH2-
bending band of PDLLA [301]. Moreover, the peak at 1620 cm-1
, ascribed –COO- stretching
in Alg, is reduced in intensity, indicating that the Alg content is reduced after immersion.
Moreover, the peak is seen to be shifted to 1643 cm-1
, which is possibly due to protonation
of carboxylate groups [295]. Therefore, these results confirm that Alg coating remains on the
scaffold after 14 days of immersion in PBS. The FTIR spectra of the T-Gel-c-(PL/P123-c-
BG) scaffold are reported in Fig. 5.8 (B). The double peak at wavenumber 1479 and 1424
cm-1
appears after immersion. Similarly to the T-Alg-c-(PL/P123-c-BG) coated scaffold
discussed above. The absorption bands corresponding to PDLLA are stronger detectable in
the spectrum of the scaffold after immersion, which is the result of degradation of the Gel
coating. The degradation of the Gel coating is evidenced by the broader peak of –NH
stretching (3700 - 3000 cm-1
).
Figure 5. 8 FTIR spectra of T-Alg- and T-Gel-c-(PL/P123-c-BG) scaffolds after 14 days of
immersion in PBS.
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5.4 Conclusions
Multifunctional layered polymer coated Bioglass-based scaffolds with drug delivery
capability were fabricated by coating Bioglass foams with two different polymer coatings,
including PDLLA/P123 blend and Alg or Gel. The scaffolds exhibited improved mechanical
properties and superior drug delivery function characterized by a low initial burst release and
subsequent controlled drug release. In addition, both Alg and Gel served as a drug carrier
and they did not show significantly different performances in their degradation and release
behaviors. The multifunctional scaffolds fabricated, exhibiting improved mechanical
properties and controlled drug release coupled with high bioactivity characteristic of
Bioglass, belong to a growing family of advanced scaffolds for bone tissue engineering.
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CHAPTER 6
Porous Biodegradable Polymer-based Scaffolds for Cartilage Tissue
Engineering Applications
6.1 Introduction
OA affects the musculoskeletal system, especially the joints in the hip, hand, knee
and spin cord [302]. At the joints, the cartilage tissue covering the end of long bones inhibits
direct rubbing of bones against each other [23,302] and it allows the joint to work smoothly
and without causing pain [302]. When OA occurs, the cartilage, which works as a shock
absorber, becomes worn off in some areas, finally leading to loss of elasticity of the cartilage
tissue [125]. In this case, the bones may rub against each other causing very severe pain
[303]. Due to its complex and unique structure, and exposure to high pressure and motion,
the repair of cartilage is one of the most challenging areas in tissue engineering
[23,241,304]. Moreover, the avascular and aneuronal nature of cartilage limits the ability to
deliver signal molecules, growth factors or cellular components for the tissue repair process
[232,241,304]. In order to repair osteochondral defects, a cell-scaffold-based repair strategy
is currently being favored instead of using arthroscopic debridement, microfracture, autograft
and autologous implantation. The reason is that cell-scaffold-based implantation exhibits the
potential to regenerate hyaline cartilage without creating defects at other sites of the joint
[23,241,305]. Therefore, the goal of this approach is to engineer a durable cartilage tissue
that provides a smooth joint surface reconstruction and is resistant to high loads, shear
stresses, combined with good friction properties. Besides the cell type and source that
certainly affect the outcome of the procedure [305], both the scaffold composition and
architecture are relevant for the repair of osteochondral defects. Furthermore, growth factors
and/or cytokine have to be integrated in order to support cell differentiation [24,305]. For
osteochondral defect regeneration, besides the required properties of scaffolds for cartilage,
which were previously mentioned in Chapter 2, cartilage scaffolds must be well connected
90
with the underlying SB or bone engineered scaffold, in order to enhance in-situ integration of
the osteochondral system [125].
This study is focused on the preparation of distinct layered scaffolds or multilayered
scaffolds, based on bioactive glasses and biodegradable polymers. For the cartilage side of
the scaffold, which will be the focus of this chapter, the elastic properties are a crucial
requirement for mimicking the natural cartilage tissue. Therefore, soft polymeric materials
are materials of choice for cartilage scaffolds in comparison with rigid bioceramics and
metals. In addition, the physical properties (such as structural and architectural) of scaffolds
are crucial for cell behavior. It is well established that a suitable scaffold should exhibit a
similar structure to that of the natural ECM of the tissue to be regenerated, which provides a
suitable microenvironment for cell adhesion, proliferation and differentiation. Based on the
required properties above, Alg was chosen as a scaffold material for cartilage in this study
due to its interesting properties, as mentioned in Chapter 2. In addition, the porous structure
and pore geometry of the Alg-based scaffolds can be tailored by adjusting the parameters of
the fabrication techniques used. Two different fabrication techniques, including freeze-drying
and electrospinning, are focused in the present work. Both techniques provide the ability to
create scaffolds with tailored porous structure, as discussed in Chapter 2.
In the scope of this chapter, Alg was initially fabricated as 3D porous foam by a
combination of gelation and freeze-drying techniques. In order to get more insights on the
scaffold characteristics and the mechanism of the used technique, parameters such as the
polymer and crosslinking agent concentrations are investigated. The first aim is to optimize
processing conditions, so that the fabricated scaffold can achieve the basic requirements as
a suitable cartilage scaffold. The optimized Alg scaffold was further modified in order to
enhance the cell adhesion, proliferation and differentiation based on in vitro chondrocytes
and stem cells culture.
In another approach, fabrication of Alg fibers by using electrospinning technique was
attempted. Few successful studies of electrospun Alg [212,306–308] have been reported
and the electrospinning of Alg aqueous solution is still challenging. Blending solutions of
Alg/PEO [212,306], Alg/PVA [307] and Alg/Gel [308] have been fabricated into fibers, while
91
individual Alg aqueous solution cannot be used to fabricate fibers by electrospinning
technique due to several reasons. First, Alg is not able to be dissolved in organic solvents,
leading to the lack of conductivity of the solution. Second, the Alg/H2O solution is not able to
be stretched under electrostatic field during electrospinning, due to its weak intermolecular
interaction and subsequently insufficient mechanical properties [173,309,310]. Under these
circumstances, the influence of electrospinning processing conditions and the fundamental
mechanism of the electrospinning technique were investigated in detail. Therefore, a
synthetic biodegradable polymer (poly(L-lactide) (PLLA)) and natural-derived polymer (Alg)
were electrospun and comparatively investigated regarding the basic desired properties for
cartilage regeneration.
Finally, the comparison between the obtained scaffolds is discussed in terms of their
microstructure, physical and mechanical properties. The possible advantages and
disadvantages of the different scaffold types were investigated, in order to seek a suitable
scaffold to be developed further for cartilage regeneration.
6.2 Materials and methods
6.2.1 Fabrication of Alg-foams
Na-Alg (Sigma, = 1.02 g/cm3) with a concentration of 4 wt/v % was prepared in DI
H2O and stirred at room temperature for 2 h. The Na-Alg solution was diluted to 2 and 3 wt/v
%, respectively. 1 ml of each of these Alg solutions was added in a 48 well-plate, followed by
the addition of 100 µl of 0.1, 0.5 and 1 M CaCl22H2O per well, respectively. The mixture was
kept at room temperature for gelation as maximum as 30 min. As next, the Alg-gel was
frozen at - 20 C for 24 h. Then frozen samples were sublimated and a porous structure was
generated after 24 h lyophilizing at temperature of - 50 C under vacuum conditions. As a
result cylindrical 3D porous Alg sponges were obtained, having a diameter of 8 mm and a
height of 8 mm. Finally, for ionic crosslinking, the Alg sponges were immersed in 0.5 M
CaCl22H2O for 4 h and dried at room temperature for at least 24 h.
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6.2.2 Fabrication of PLLA fibers and Alg-based fibers
Poly (L-lactide) (PLLA) was dissolved in DCM/methanol (80/20 vol %) with
concentrations of 5 and 7.5 wt/v % at room temperature for 2 h. A PLLA solution was
fabricated as a fibrous mesh by using electrospinning technique. The processing parameters
consisted of 8.5 - 20 kV of applied voltage, 0.5 - 2 ml/h of flow rates, 15 - 18 cm of tip-
collector distances and 2 h of deposition time, as detailed in Table 6.1.
Na-Alg was blended with PVA, P123 and gelatin in DI H2O for different weight ratios
and solvent systems, as detailed in table 6.2. The mixture solution of Alg/Gel blend was kept
at 50 C in order to avoid gel-formation of Gel. During the electrospinning process, hot air
conditions (i.e. a heat gun) were applied to the set-up in order to maintain a constant
temperature of the process at 50 C. The variety of electrospinning processing conditions of
Alg-based solutions is detailed in Table 6.2.
Electrospun Alg/Gel fibers were chemically crosslinked by using glutaraldehyde
(GA) exposure. The GA crosslinking agent was prepared at concentration of 30 vol % in
H2O/EtOH (70/30 vol %) and was transferred to a 1 l glass beaker. The fiber mesh was
attached on the wall of beaker by using double-side adhesive tape. The beaker was
suddenly wrapped with aluminium foil tightly and was kept at room temperature for 48 h for
crosslinking with GA vapor.
6.2.3 Characterization and testing
(i) Porosity and density
The porosity (P) of Alg-foams was calculated by using Eq. 6.1, while the density ()
of the foams was determined by using Eq. 6.2:
P = 1 – [Wfoam/(Alg x Vfoam)] (6.1)
(g/cm3) = Wfoam/Vfoam (6.2)
; where Wfoam is the weight of the Alg freeze-dried foam, Alg is the density of solid Alg (1.02
g/cm3), and Vfoam is the volume of Alg freeze-dried foam, determined from the dimensions of
the foam.
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(ii) Microscopy
The cross-section appearance of scaffolds was investigated by using light
microscopy (light microscope Axioplan Zeiss), equipped with a digital camera system and
operated with the software LEICA DFC290.
The microstructure of the freeze-dried foams and electrospun fibers was
characterized by SEM. The SEM (LEO 435 VP) was operated with a tungsten filament, at an
acceleration voltage of 10 kV. Cross-section and plan-view SEM imaging was performed in
order to observe the pore structure. For SEM investigations the samples were sputter-
coated. The pore size of the foams and the diameter of fibers were evaluated from SEM
images using the free available software Image J (version 1.42S).
(iii) Chemical analysis
The chemical structure of the foams and fibers was investigated by using attenuated
total reflectance-FTIR (ATR-FTIR) in transmission mode with the resolution of 4 cm-1
, the
collected signals in 32 scans and the spectrum range of 4000 - 525 cm-1
.
(iv) XRD analysis
The scaffolds were characterized by using XRD (Siemens D500) analysis in order to
assess the crystallinity by employing Cu k radiation. Data were collected over the range of
2 = 10 - 70 using a step size of 0.02 and a counting time of 25 sec per step.
(v) Thermal analysis
Thermal transition features, including glass transition temperature (Tg) and melting
temperature (Tm) of the scaffolds were measured by using differential scanning calorimetry
(DSC; DSC Q 2000). Measurements were performed in the temperature range of (-50) - 200
C with heating rate of 10 C/min and double runs. Tg and Tm were determined by the
minimum value of the first and second endothermic transition peaks, respectively, while cold
crystallization temperature (Tc) was determined by the maximum value of the exothermic
transition peak. The percent crystallinity (Xc) was determined by using Eq. 6.3 [311,312]:
% Xc = [(∆Hm - ∆Hc)/ ∆Hm] × 100 (6.3)
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; where ∆Hm is the heat of melting (J/g), ∆Hc is the heat of cold crystallization (J/g)
and ∆Hm is a reference value, representing the heat of melting of 100 % crystalline
polymer (∆Hm of PLLA 93 J/g [312]).
(vi) In vitro biodegradation study
The water absorption of alginate freeze-dried scaffolds was investigated by
immersion in PBS (pH 7.4) for 45 min, 2 h, until 4 months, respectively. Each sample was
placed in a polystyrene bottle containing 50 ml of the PBS medium and incubated in an
orbital shaker at 37 C and 90 rpm. The medium was replaced two times per week. At
interval time, the sample was removed and blotted with filtered paper before weighed. The
obtained weight is wet weight (Wwet). The water absorption at each time point was
determined by using Eq. 6.4 [295]:
Water absorption (%) = [(Wwet – Wdry)/Wwet] x 100 (6.4)
; Wwet is the wet weight (g) of sample after soaking in medium and Wdry is the weight (g) of
the sample before soaking in the medium, in the following referred as the dry weight. Four
samples were investigated and the results are reported as average values ± SD.
The degradation profile and structural stability of the crosslinked Alg-foams were
evaluated by investigating the weight change of the scaffolds in different media, including DI
H2O, PBS and SBF with respect to the immersion times. These experiments were carried
simultaneously to the water absorption experiments. The weight change at each time point
and media was determined by using Eq. 6.5 [219]:
Weight change (%) = [(Wwet – Wdry)/Wdry] × 100 (6.5)
(vii) Mechanical testing
Alg-foams (dimensions of 8 mm in diameter and 8 mm in height) were compressed
by using a universal testing machine (Zwick Z050) for mechanical testing. Thereby, the
cross-head speed during compression testing was 2 mm/min. The samples were pre-loaded
at 0.1 N, while a maximum load of 50 N was used. A stress-strain curve was used to
determine the mechanical properties. The elastic modulus was calculated from the initial
slope of the elastic regime of the stress-strain curves, while the compressive strength was
95
obtained from the maximum stress before the samples collapse. The results are averaged
from eight specimens that were tested for each condition.
In addition, dynamic mechanical analysis (DMA; Mark IV) was performed in order to
study the mechanical response of the foams under the replication of real forces, which are
expected to dominate in vivo [313]. Alg-foams (dimensions of 8 mm in diameter and 3 mm in
height) were tested in both dry and wet states, under a sinusoidal load with a rate defined by
a frequency (in Hz). Before testing in the wet state, the foams were immersed in PBS
solution until reaching equilibrium conditions. The measurement was performed in
compression mode of deformation, with the cycle frequency varied in the range 0.1 - 10 Hz
(according to the range of typical skeletal movement in vivo [313]) and with the maximum
strain amplitude of 1 %. The viscoelastic properties of the foams are represented as storage
modulus (E’) and loss factor (tan ) as a function of frequency.
The fibrous meshes in the form of thin films (5 mm in width, 30 mm in length and
500 µm in thickness) were tested under tension mode by using a universal testing machine
(Frank) by applying a force at a cross-head speed of 1 mm/min, preload at 0.1 N and
maximum load at 50 N. A tensile stress-strain curve was recorded to determine mechanical
properties. The Young’s modulus was calculated from the initial linear slope of stress-strain
curves and the tensile strength was obtained from the maximum value of stress before
fracture. Cast films fabricated by using the same solution employed for the electrospun
solution were also investigated as control. Eight specimens were tested and the results were
presented as average ± SD.
6.2.4 Statistical analysis
The data were analyzed by using one-way ANOVA analysis and Turkey’s multiple-
comparison test to determine statistical differences. A confidence interval of 95 % (p = 0.05)
was used for all analysis.
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6.3 Results and discussion
6.3.1. Effect of processing conditions on the physical and mechanical properties of
the foams
(i) Microstructure
3D porous Alg-foams were fabricated by means of gelation and freeze-drying
technique. Since Na-Alg is naturally water soluble, the structural stability of Alg-foams
cannot be maintained in contact with biological environments. Therefore, it is necessary to
transform soluble Na-Alg into an insoluble Ca-Alg configuration, which can be achieved by
ionic crosslinking, as described in Chapter 2. This crosslinking results in the formation of
Alg-gel, which is water resistant. The slow diffusion of Ca ions containing solution from the
superficial surface toward the bottom of the bulk solution led to the generation of tubular-like
structures along the gel [314], as observed in Fig. 6.1 (B)). A plan-view image in Fig. 6.1 (A)
shows circular-shaped structures spreading around the surface area, whose formation is
attributed to the addition of the CaCl22H2O agent [314]. After freezing the gel, the
propagation of a planar ice front occurred dependent on freezing rate. During lyophilizing,
the ice was sublimated and this process induced a horizontal flat ladder feature, as observed
by plan-view SEM imaging (Fig. 6.1 (C)) and tubular-like structures, as shown in the cross-
section SEM image in Fig. 6.1 (D). The structure of the foams is suitable to match the
required pore structure of scaffolds in order to support cell organization, as observed in
natural cartilage tissue, in which chondrocytes align in a columnar manner [24,125,128,129].
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Figure 6. 1 Optical microscopy images showing 3 wt/v % Alg-gel after the introduction of 0.1
M CaCl22H2O into Alg solution: (A) superficial surface (plan-view image) and (B) cross-
section of Alg-gel. SEM images of 3 wt/v % Alg-foam after lyophilized: (C) plan-view and (D)
cross-section SEM image (Reproduced from Nooeaid et al. [145] with the permission of
John Wiley and Sons).
(ii) Pore geometry and structural stability
The pore geometry and the mechanical properties of Alg-foams depend strongly on
concentration of both Alg solution and crosslinking agent used. Alg solution at
concentrations of 2, 3 and 4 wt/v %, and CaCl22H2O at concentrations of 0.1, 0.5 and 1 M
were varied in order to find out optimum conditions, which are used to fabricate the foams
suitable for the application of cartilage regeneration. Therefore, porosity, pore size, water
absorption, physical and mechanical properties, and reproducible efficiency of obtained
forms were investigated.
First, the influence of crosslinking agent concentration on gel-formation and porous
structure of the foams was studied. It was found that at given concentrations of 0.5 and 1 M
CaCl22H2O fast gel-formation (in few min) occurred for each Alg concentration. This
phenomenon led to a hard and inhomogeneous gel. As shown in Fig. 6.2, uncontrolled
shape was obtained after lyophilization. In contrast to this, a lower concentration of
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crosslinking agent (0.1 M CaCl22H2O) slowed down the gelation rate and consequently
homogeneous foam were formed exhibiting flexibility and softness when they were handled
with fingers. This result is in an agreement with the previous study that adding 0.075 - 0.1 M
CaCl22H2O into 2 wt/v % Alg solution resulted in a mechanically more stable gels, in
comparison to the one fabricated by using higher concentration of CaCl22H2O [127].
Moreover, it has been shown that the gel, obtained by using lower CaCl22H2O
concentration, remained cohesive upon handling. This result could be due to the fact that
low amount of Ca ions diffuse slowly through the Alg solution and they induce gel formation
at the same time. This phenomenon is known as ionotropic gelation [314]. As a result, the
solution has enough time to form a gel homogeneously. Therefore, 0.1 M CaCl22H2O was
chosen as the crosslinking agent for the gelation of Alg in the present study.
However, obtained Alg-foams did not exhibit sufficient mechanical stability and did
not last for even 1 h long in aqueous solutions. This is supported by the result of weight loss
of Alg-foams with respect to immersion time in DI H2O (at 37 °C) (Fig. 6.3). The foams were
completely decomposed in DI H2O in only 45 min, indicating their insufficient time-dependent
mechanical stability. In order to overcome this limitation, physical crosslinking was applied
by soaking the foams in 0.5 M CaCl22H2O for 4 h. As a result, the foams after crosslinking
lasted longer in DI H2O (i.e. up to 50 % weight loss was observed after 4 weeks of
immersion), indicating their suitability for the intended application in osteochondral tissue
engineering.
Regarding the influence of Alg concentration on foam structure, it was found that
increased Alg concentration led to increased density and relatively reduced porosity of the
foams. The porosity decreased from 97 % down to 93 % with increased concentration from
2 to 4 wt/v % (Fig. 6.4).
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Figure 6. 2 Optical photograph showing the appearance of 3 wt/v% Alg-foams obtained with
variation of CaCl22H2O concentrations (0.1 – 1 M).
Figure 6. 3 Comparison of weight loss as a function of immersion time in DI H2O of Alg-
foams without (w/o) and with (w) crosslinking by immersion in 0.5 M CaCl22H2O for 4 h (the
inset represents the weight loss of Alg-foams without crosslinking) (Reproduced from
Nooeaid et al. [145] with the permission of John Wiley and Sons).
100
Figure 6. 4 The effect of Alg concentrations on the porosity and density of Alg-foams.
SEM images in Fig. 6.5 show the comparison between the microstructure of Alg-
foams fabricated from using different Alg concentrations. 2 and 3 wt/v % Alg-foams exhibited
uniform porous structure with open pores (Fig. 6.5 (A and B) and Fig. 6.5 (D and E)), while
non-uniform porous structure with some closed pores were observed in the case of 4 wt/v %
Alg-foams (Fig. 6.5 (G and H)). Closed pores are not desirable because they inhibit
migration of cells and further limit cellular activity. Moreover, the pore size of foams is an
important factor for cell migration. Appropriate pore size for supporting chondrocytes and
stem cells has been reported in the range of 50 - 300 μm [137,315,316]. In the present work,
Alg-foams, which were fabricated by using different concentrations, exhibited pore sizes in a
wide range of 30 - 500 μm, as shown in Fig. 6.5 (C, F and I). This pore size range overlaps
the required pore size for a suitable cartilage scaffold. In addition, the increase of Alg
concentration resulted in the reduction of pore size. For instance, 4 wt/v % Alg-foams
exhibited pore size in the range of 30 - 325 µm (average pore size 180 ± 47 µm), while 3
wt/v % Alg-foams showed pore size in the range of 125 - 325 µm (average pore size 237 ±
48 µm). In the case of 2 wt/v % Alg-foams, a wide pore size in the range of 150 - 500 µm
(average pore size 305 ± 55 µm) was observed. Thus concerning foam pore size, 3 wt/v %
101
Alg-foams delivered the closest value to the target pore size, which is desired for suitable
cartilage scaffolds.
Figure 6. 5 SEM images (in plan-view) of (A, B) 2 wt/v %, (D, E) 3 wt/v % and (G, H) 4 wt/v
% Alg-foams, included distribution of pore size: (C) 2 wt/v %, (F) 3 wt/v % and (I) 4 wt/v %
Alg-foams.
(iii) Mechanical properties
In addition to the pore configuration, the mechanical properties of the foams are one
of the important factors, which is necessary to be considered for use in load bearing
applications. The Alg concentration used provides the main impact on the mechanical
properties of the obtained foams. As shown in Fig. 6.4, the increase of the Alg concentration
leads to increased density of Alg-foams, which results in the enhancement of the mechanical
102
properties under compression load, as shown in Fig. 6.6 (A). In the representative
compressive stress-strain curves of 2, 3 and 4 wt/v % Alg-foams (Fig. 6.6 (A)), all samples
present the same trend, which is defined by sigmoidal shape (a characteristic of polymeric
cellular solids [317]). In addition, the stress-strain curves of 4 wt/v % foam shows drop
features, attributed to the occurrence of inhomogeneous deformation of the foam under
compression load. In contrast, 2 and 3 wt/v % foams exhibited smooth curves, indicating
more homogeneous deformation during compression. The shape of these curves manifests
three deformation mechanisms, including: (I) linear elastic region, in which the foam shows
elasticity and the instant rise of stress represents the deformation of the intact foam; (II)
brittle crushing region, which represents progressive rupture and collapse of the cell walls;
(III) densification region, which shows the densification of the compressed specimen
[317,318]. Elastic modulus was 0.220 ± 0.009 MPa in the case of 3 wt/v % Alg-foams,
while 4 wt/v % Alg-foams exhibited the modulus in the similar range ( 0.22 ± 0.05 MPa)
(Fig. 6.6 (B)). Importantly, 4 wt/v % Alg-foams did not provide an improvement of elastic
modulus. This is probably related to the formation of a non-uniform foam when the
concentration reached 4 wt/v %. The compressive strengths of 2, 3 and 4 wt/v % Alg-foams
were 0.047 ± 0.004, 0.14 ± 0.02 and 0.15 ± 0.02 MPa, respectively. 3 wt/v % Alg-foams
exhibited significant improvement of the elastic modulus and compressive strength
compared to 2 wt/v % Alg-foams, while 4 wt/v % Alg-foams did not lead to further
improvement of the mechanical properties (see Fig. 6.6 (B)). This result suggests that at
higher concentration (above 3 wt/v %) a critical viscosity is reached, which inhibits the
diffusion of Ca ions through Alg solution during the gelation process, leading to formation of
an inhomogeneous gel. Subsequently, an inhomogeneous foam is obtained after
lyophilization, which negatively influences the mechanical properties.
103
Figure 6. 6 (A) Representative compressive stress-strain curves of 2, 3 and 4 wt/v % Alg-
foams and (B) the mechanical properties, including elastic modulus2 and compressive
strength3 of the foams as a function of concentrations.
(iv) Viscoelastic properties
Since cartilage is exposed to dynamic compression forces, DMA was used to
investigate the viscoelastic properties of the foams in both, dry and wet state. In particular,
the tests of wet foams (by immersion in PBS solution) were performed with the aim to
replicate the realistic conditions in the joint [313]. Fig. 6.7 shows the storage modulus (E’) as
a function of frequency (Hz) in the dry and wet state. In detail, the values of E’ of dry foams
were higher than those of wet foams. This result suggests that water molecules absorbed in
the foam act as a plasticizer in the polymer network, leading to molecular mobility and
subsequently reduced stiffness. E’ values of wet foams ( 0.5 MPa; 0.01 - 10 Hz) were
exhibited in the range of reference values of native cartilage (0.01 - 1.5 MPa; 0.01 - 10 Hz
[313]). In addition, tan values of wet foams were higher than those of dry foams, which is
suggested by the contribution of friction between water molecules and scaffold [319]. Finally,
E’ and tan values of both, dry and wet foams, remained constant along the given
2 Modulus of elasticity was determined at the region of elastic deformation in the stress-strain curve, according to
Hooke’s law ( = E) [383]. 3 Compressive strength was obtained at the yield point, where the onset of plastic deformation takes place.
Therefore, the represented compressive strength in this study is equivalent to yield strength [383].
104
frequency range. This result indicates that the viscoelastic properties of the foams are in the
steady state within the measured time scale.
Figure 6. 7 Dynamic mechanical properties of 3 wt/v % Alg-foams in compression mode
presenting the storage modulus (E’) and the loss factor (tan ) as a function of frequency, in
both dry and wet state.
(v) Thermal properties
Effect of crosslinking on the thermal transition of 3 wt/v % Alg-foams was evidenced
by DSC results in Fig. 6.8. The thermal transition of both Alg-foams with and without
crosslinking shows only one significant endothermic peak detected from the first heating
ramp (Fig. 6.8 (A)). The peak minimum refers to the melting temperature (Tm) [320]. Tm of
Alg-foams increased from 78.1 °C to 96.6 °C with crosslinking. This result indicates that the
crosslinking enhances intermolecular interaction in Alg network. The increase of enthalpy
(∆Hm) required to break the interaction and to melt the foam was observed after crosslinking.
In detail, ∆Hm required for melting Alg-foam with crosslinking was 560.5 J/g, while ∆Hm
required for melting the foam without crosslinking was 495.5 J/g. In addition, another
important feature obtained from the DSC thermogram in Fig. 6.8 (B) is the glass transition
temperature (Tg), which primarily refers to the softness and flexibility of the polymer chains
105
[317]. Tg is determined from the second heating ramp in an endothermic transition. It was
found that the introduction of crosslinking did not play a main role in Tg, as shown in Fig. 6.8
(B). Moreover, the detectable Tg of both foams with and without crosslinking at around 23 -
24 °C indicates that the foams might be softer at 37 °C, either in the in vitro or in vivo culture
conditions. However, it can be confirmed that the foams are not decomposed at that
temperature.
Figure 6. 8 DSC thermogram: (A) the 1st heating and (B) the 2
nd heating cycle runs of Alg-
foams with and without crosslinking.
(vi) Swelling and degradation profile
In general, natural cartilage resides in contact with synovial fluid within the synovial
membrane, and the high water content in cartilage provides sufficient softness and lowers
the friction during joint motions [16,19,24]. It has been reported that natural articular cartilage
consists of 80 % water with respect to the wet weight [6]. In the present work, water
absorption of Alg-foams reached a value of 82 ± 2 % (determined after 4 h of immersion in
PBS).
Apart from the ability to absorb water, the structural stability of scaffolds must be
maintained in physiological conditions until the regeneration of new tissue takes place. With
this aim, weight change of the foams in aqueous media (i.e. DI H2O, PBS and SBF, all at 37
106
°C) with respect to immersion time is presented in Fig. 6.9 (A). At the initial stage ( 4 h), the
weight change of the foams in each medium was drastically increased with increasing
immersion time. This is justified by the fact that Alg-foams tend to absorb water (free and
bulk water) to fill the void region inside the polymer network until reaching the equilibrium
state [295,321–323]. This phenomenon follows the mechanism of hydration initiated by the
hydrophilic groups of Alg [295]. In this case, the foams reached equilibrium state at week 2,
in which water cannot be further gained by the foams. Afterwards, the weight change
remained constant until the end of the experiment (4 months of immersion time). This result
indicates that neither decomposition nor degradation of the foams immersed in DI H2O has
taken place yet. The structural stability of Ca-Alg network can be maintained until the
osmotic pressure overcomes the force of ionic-crosslinking interaction.
In contrast, the weight change of Alg-foams during immersion in PBS is mainly
driven by Ca2+
-Na+ exchange (Ca ions interacting with Alg chains and Na ions in the
medium) [295,324]. The initial phenomenon in alkaline media such as PBS and SBF (both at
pH 7.4) follows Eq. 6.6 [325]:
Ca-Alg2 + 2NaCl 2Na-Alg + CaCl2 (6.6)
When a Ca ion, linking with two Alg chains, exchanges with two Na ions from NaCl in the
medium, Ca-crosslinked Alg network (Ca-Alg2) transforms to Na-Alg. Since this reaction is
reversible, Na-Alg is able to converse to Ca-Alg in contact with excess CaCl2 in the medium.
Alg-foams immersed in PBS showed the highest weight change among all media used (see
Fig. 6.9 (A)). It is considered that more free hydrophilic groups (i.e. COO-) exist in the Alg
network due to loss of Ca-interaction. Thus the network is able to gain more water to fill the
voids. After 4 weeks of immersion, the trend of weight change was significantly reduced,
indicating possible occurrence of decomposition and the foams completely decomposed in
PBS at week 16.
In the case of immersion in SBF, the trend of weight change in the initial stage was
similar to the case of immersion in PBS. However, after 3 weeks in SBF, the weight change
did not vary until the end of the investigation, while the foams in PBS have decomposed.
107
This observation is suggested to be the result of the reversible reaction which is induced by
the existence of CaCl2 in SBF.
The loss of Ca-interaction inside the Alg network was confirmed by FTIR
spectroscopy (Fig. 6.9 (B)). The characteristic peaks of Alg in the region of 1150 - 900 cm-1
exhibited reduced intensity after 7 days of immersion in all media, indicating deformation of
the Alg network. In particular, the reduced intensity of Ca-O peak at 1009 cm-1
confirmed the
loss of Ca-interaction.
In summary, Alg-foams immersed in different media exhibit slightly different
decomposition rates and therefore different life-times. This result is likely due to the different
compositions of the media used, which directly influences the structural stability of the Alg
network. Therefore, the degradation study of Alg-foams performed in either PBS or SBF in
the present work is only used as a prediction for further in vitro cell culture. In order to
precisely replicate the real body system, degradation studies should also investigate
enzymatic degradation (i.e. with lysozyme and hyaluronidase solution [313]), which play a
fundamental role in the degradation of polysaccharides in vivo.
Figure 6. 9 (A) Weight change of 3 wt/v % Alg-foams as a function of immersion time in
different media, including DI H2O, PBS and SBF. (B) ATR-FTIR spectra of the foam after 7
days of immersion in the media.
108
6.3.2 Effect of electrospinning conditions on the properties of fibers
(i) Spin-ability of PLLA solution
The electrospinning of biodegradable polymers in the present study was initially
started with a synthetic-derived polymer (PLLA), in order to understand fundamentals of the
influence of polymer solutions and processing conditions using a well-investigated polymer.
First, PLLA fibers were electrospun according to the previous studies of Yunos et al. [85,86].
The optimum conditions have been reported as 5 wt/v % of PDLLA/DMC solution, 8.5 kV of
applied voltage, 15 cm of deposition distance and 0.9 ml/h of feed rate. In our experiment, it
was found that using the solution of 5 wt/v % PLLA/(DCM/MeOH) (80/20 vol %) (Trial 1;
Table 6.1) at the same conditions as Yunos et al. [85], it was not possible to obtain fibers.
With decreasing feed rate (0.5 ml/h; Trial 2), fibers were obtained with the formation of
beads. Even though the applied voltage and deposition distance were adjusted (Trials 3 - 8),
uniform fibers were not achieved by using 5 wt/v % PLLA solution. The presence of beads is
shown in Fig. 6.10 (A), which is a material obtained by using the processing condition in Trial
7. This result can be explained by the fact that the concentration 5 wt/v % PLLA in
DCM/MeOH has insufficient viscosity and low strength in order to be stretched under the
effect of the electrostatic field.
When the concentration was increased to 7.5 wt/v % with the same used feed rate
of 0.5 ml/h and deposition distance of 15 cm, electrospinning at voltages less than 15 kV
(Trials 9, 10) did not lead to fiber formation due to the insufficient strength of electrostatic
field. Under such low strength of electric field, the stable Taylor cone is not able to form and
to accelerate the polymer jet to be drawn from the needle tip to the collector. According to
this, the voltage was raised up (15 - 20 kV; Trials 11 - 15), and in this case the uniform fibers
were obtained with diameters in the range 1 - 1.3 µm, as shown as an example obtained
from using the conditions in Trial 13 (Fig. 6.10 (B)).
In summary, the polymer solution (concentration and solvent) plays the most
important role in the spin-ability. As observed, the PLLA solution started to exhibit suitable
electrospinning behavior at a concentration of 7.5 wt/v %, which is in agreement with the
study of Kenawy et al. [326], who showed that poly(lactic acid) solutions can be electrospun
109
well at concentrations of around 7 - 8 wt/v %. In addition, the external factors affecting
electrospinning, including applied voltages, deposition distance and feed rate, have an
influence on the formation of electrospinning jet. These factors have an effect on the fiber
morphology and diameter. The applied voltage affects the strength of the electric field, the
acceleration of the electrospinning jet and the stretching ability of the polymer solution. The
deposition distance relates to the flight time of the polymer jet, which plays a role in the
evaporation rate of the solvent. Finally, the feed rate determines the controlled volume of the
polymer solution in order to maintain a stable Taylor cone at the given voltage and it mainly
affects the diameter of obtained fibers.
Table 6. 1 Electrospinning conditions and primary observations of PLLA fibers ( indicates
no fiber formation, indicates beads incorporated into fibers and indicates uniform fibers).
Concentrations
(wt/v %) Trails
Processing parameters Results
Voltage (kV) Distance (cm)) Feed rate (ml/h)
5 1 8.5 15 0.9
2 8.5 15 0.5
3 15 15 0.5
4 18 15 0.5
5 20 15 0.5
6 20 18 0.5
7 15 15 1.15
8 15 15 2
7.5 9 8.5 15 0.5
10 12.5 15 0.5
11 15 15 0.5
12 15 18 0.5
13 18 15 0.5
14 20 15 0.5
15 20 18 0.5
110
Figure 6. 10 Optical microscopy images of electrospun PLLA fibers obtained by using the
conditions in (A) Trial 4 and (B) Trail 13, according to electrospinning conditions used in
Table 6.1.
(ii) Spin-ability of Alg solution
Up to date, the electrospinning of Alg solution has been studied by considering
variations of the solution systems [173,307–310,327–329]. One of the approaches is to
incorporate water soluble synthetic-derived biodegradable polymers (i.e. PVA and PEO) into
the Alg solution. In the first attempt, Alg was blended with PVA, glycerol, P123 and Gel in
aqueous solution. The series of solution, composition and processing conditions are
reported in Table 6.2. The results of the preliminary study showed that the Alg solutions
blended with PVA and glycerol are not able to form fibers even with variations of
composition and processing conditions (Trials 1 - 4). Regarding the study of Safi et al. [327],
the spin-ability of the Alg solution was seen to improve by increasing hydrogen bonding
between the hydroxyl groups in Alg and either the hydroxyl groups in PVA or the ether
groups in PEO. By this approach, Alg fibers were formed by using a concentration of 7 wt/v
% PVA blended with 2 wt/v % Alg in the volume ratio of 70/30 and 50/50 vol %. The
concentration used in Trail 2 (Table 6.2) was relatively the same as the optimum condition
according to the study of Safi et al. [327]. By using the same electrospinning conditions,
including 12 kV of voltage and 10 cm of needle-collector distance, no fibers were obtained
even with variation of conditions. A reason for this behavior might be the insufficient viscosity
of the used solution, which is caused by the different Mw of the polymers used.
111
The addition of glycerol in the Alg solution was also attempted (Trials 3 and 4; Table
6.2). Glycerol was aimed to function as a co-solvent, which is expected to enhance the
flexibility of chain conformation and to increase entanglement of Alg molecules [309].
However, fibers were not achieved by this approach, even by the adjustment of processing
conditions. The spin-ability of the Alg solution was further investigated by blending with P123
and the introduction of MeOH as a co-solvent (Trials 5 - 8). P123 copolymer was used in
order to reduce the viscosity of the solution, while MeOH was introduced to enhance the
conductivity of the solution. However, an improvement of spin-ability was not observed. In
detail, as observed from the experiment, no jet was formed during electrospinning. Only
beads were deposited on the collector. This result indicates that the investigated solutions
were not sufficient suitable in terms of interchain interaction to be drawn out and travelled
forward to the collector under the given electrostatic field.
By using the solution of Alg/Gel blending (Trials 9 - 14), a suitable system was found
(Trial 12), which was the only studied system able to form fibers under the investigated
electrospinning conditions (Fig. 6.11 (B)), including voltage of 12 kV, needle-collector
distance of 12 cm and feed rate of 0.1 ml/h. In addition, the set-up was subjected to a heat
gun in order to control the temperature of the process at 50 C. Thus the spin-ability of the
Alg solution can be conveniently improved by the addition of Gel. However, it was still not
possible to achieve high reproducibility, in comparison to the electrospinning of PLLA
solution. Importantly, the electrospinning of Alg/Gel/H2O (1/15/84 wt %) solution was well
performed under controlled temperature conditions. When the content of Gel was decreased
(Trials 9 - 11), the fibers were obtained with the appearance of beads, as shown in Fig. 6.11
(A; Trail 9). It can be concluded that the Gel content plays the key factor in the chain
entanglement of the Alg solution. It is probably due to positively charged Gel forming ionic
interaction with negatively charged Alg. Nevertheless, it is necessary to maintain the
temperature of the solution above the gelling temperature of Gel ( 50 C) in order to avoid
the gelation of the solution during electrospinning. In addition, the applied heat accelerates
the evaporation of the solvent (H2O), leading to skin formation out of the solution jet, which
helps to stabilize the jet being able to travel forward to the collector. However, it was
112
observed that the Taylor cone was not stably maintained and the jet was not continuously
formed. This is believed to be one of the reasons of the poor spin-ability of the Alg solution.
Therefore, the optimization of electrospinning of Alg solutions requires further research
efforts, in order to enhance both the spin-ability and the reproducibility. As mentioned above,
Alg fibers can be only electrospun with the addition of high amount of Gel. Thus the main
component of the obtained fibers became Gel instead of Alg.
Table 6. 2 Electrospinning conditions and primary observations of alginate fibers ( indicates
no fiber, indicates beads and indicates uniform fibers).
Solutions Compositions
(wt. %) Trials
Electrospinning parameters
Results Voltage
4 (kV)
Distance
(cm)
Feed rate
(ml/h)
Alg/PVA/H2O 1.4/2.1/96.5 1 12 - 20 8 - 15 0.01
3/2/95 2 12 - 20 8 - 15 0.01
Alg/Glycerol/
H2O
2/49/49 3 12 - 20 8 - 15 0.01
4/48/48 4 12 - 20 8 - 15 0.01
Alg/P123/ H2O 3/2/95 5 12 - 20 8 - 15 0.01
1.5/2/96.5 6 12 - 20 8 - 15 0.01
Alg/P123/
H2O/MeOH
3/2/94/1 7 12 - 20 8 - 15 0.01
1.5/2/95.5/1 8 12 - 20 8 - 15 0.01
Alg/Gel/H2O5 1/10/89 9 12 12 0.01
10 12 8 0.01
11 15 12 0.01
1/15/84 12 12 12 0.01
13 12 8 0.01
14 15 12 0.01
4 The applied voltage in the case of Alg solution was fixed at 20 kV because the Alg-based aqueous solutions are
sensitive to high electrostatic field, possibly leading to an occurrence of electric spark during electrospinning. 5 The heat gun was subjected to the electrospinning set-up in order to keep the solution at around 50 °C.
113
Figure 6. 11 Optical microscopic images of electrospun Alg/Gel fibers obtained by using the
conditions in (A) Trial 9 and (B) Trail 12, according to electrospinning conditions used in
Table 6.2.
(iii) Microstructure
SEM images in Fig. 6.12 show the microstructure of electrospun PLLA and Alg/Gel
fibers, which were obtained by using the optimum conditions (Trail 13; Table 6.1 in the case
of PLLA and Trial 12; Table 6.2 in the case of Alg/Gel). The uniform PLLA fibers, as shown
in Fig. 6.12 (A), were obtained with average diameter of 0.4 ± 0.1 µm, while Alg/Gel fibers
were obtained as a very thin mesh by using the same deposition time as in the case of PLLA
(2 h), as shown in Fig. 6.12 (B). Alg/Gel fibers were uniform without the formation of beads
and they have an average diameter of 0.5 ± 0.2 µm. This result confirms the low spin-ability
of the Alg solution when compared to PLLA solution.
Additionally, when the Alg/Gel fibers were deposited for longer time (9 h) in order to
obtain a thicker mesh, the morphology of Alg/Gel fibers was different from the case of PLLA
fibers. A ribbon-like structure formed in the case of Alg/Gel-electrospinning (see the inset in
Fig. 6.12 (D)), while a common round morphology was found in the case of PLLA fibers (the
inset in Fig. 6.12 (C)). This phenomenon can be described by the reason that the jet of
Alg/Gel solution tends to form a skin on the jet during electrospinning and consequently the
skin collapses into a ribbon as it dried [330,331].
114
Figure 6. 12 SEM images of PLLA fibers, which were deposited for (A) 2 h and (C) 9 h, and
Alg/Gel fibers, which were deposited for (B) 2 h and (D) 9 h. The distribution of fiber
diameters of both fiber types is included.
115
(iv) Chemical structure
The presence of both Alg and Gel in the obtained fibers was investigated by using
ATR-FTIR spectroscopy (Fig. 6.13 (A)). The characteristic peaks of Gel were detected at
1637 and 1540 cm-1
, attributed to amide I and II of pure Gel type A, respectively [332]. The
characteristic peaks of Alg were detected as a doublet peak with low intensity at 1082 and
1035 cm-1
, attributed to C-C and C-O stretching of Alg, respectively [332]. Thus the presence
of both Alg and Gel in the electrospun fibers is confirmed.
Moreover, Alg/Gel fibers were chemically crosslinked by exposing to GA vapor, in
order to improve their water resistance and structural stability. The exposure of GA vapor
was chosen in this study instead of soaking in GA solution in order to avoid the potential
toxicity from GA solution. The ATR-FTIR spectrum in Fig. 6.13 (A) of the Alg/Gel fibers after
GA crosslinking is presented in comparison to uncrosslinked fibers to confirm the occurrence
of crosslinking. The detectable peak at 2369 cm-1
in the spectrum of crosslinked fibers
confirmed the occurrence of crosslinking, according to the literature [77]. In addition, the
characteristic peak of amide I (at 1637 cm-1
) in the spectrum of uncrosslinked fibers was
shifted to 1647 cm-1
(see the inset in Fig. 6.13 (A)). This result indicates that a chemical
reaction has taken place by the crosslinking. The effective crosslinking was additionally
confirmed by the detection of ethylene groups (–CH2-) of GA at 2930 and 2880 cm-1
[214],
which exhibited higher intensity in comparison to the spectrum of fibers without crosslinking.
The mechanism of GA crosslinked gelatin is shown in Fig. 6.13 (B). The carbonyl groups in
GA tend to react with reactive amine groups (–NH2) in Gel, forming terminated -C=N- bonds
on a GA molecule.
116
Figure 6. 13 (A) FTIR spectra of Alg/Gel fibers before and after GA crosslinking (The inset
represents the change of amide I peak before and after crosslinking) and (B) mechanism of
the crosslinking reaction between Gel with GA, according to [78].
(v) Crystalline structure
The crystalline structure of electrospun fibers was analyzed by using XRD analysis
in comparison with the cast film used as control (Fig. 6.14). The -crystalline phase of PLLA
cast film appeared as a sharp and narrow peak at 17.3 and 19.7 2 (as marked by ■ in
Fig. 6.14) and the β-crystalline phase appeared at 30 2 (as marked by in Fig. 6.15),
according to reference [333]. In contrast, the crystalline peaks almost disappeared in the
case of PLLA fibers, indicating that the reduction of crystallinity in PLLA has been taken
place during electrospinning. Consequently, an amorphous phase was dominant in the
fibers. The reason for this behavior is likely the fact that the jet travelling toward the collector
has no enough time for crystal growth during electrospinning [172,238]. As a result, this
phenomenon directly influences the mechanical properties of the fibers. The reduction of
crystallinity leads to the decrease of the mechanical strength, which is discussed in Section
(vii).
117
Figure 6. 14 XRD patterns of PLLA fibers showing the reduction of crystallinity in
comparison with PLLA cast films (■ and indicate - and β-crystalline phase of PLLA,
respectively).
(vi) Thermal properties
The thermal properties of the electrospun PLLA fibers are reported as the DSC
curves in Fig. 6.15. Tg of PLLA fibers is detected at 63.7 C, which is comparable to the
value of as-received PLLA pellets (67 °C) [334] and of porous PLLA scaffolds fabricated by
phase separation technique [320,335]. This result indicates that the molecular mobility of
PLLA fibers does not occur at the constant temperature (37 C) in body fluids. In contrast,
Alg-foams exhibited Tg values lower than 37 °C. Thus under in vitro culture conditions, it can
be stated that the molecular structure of PLLA fibers is more stable than that of Alg-foams.
An exothermic transition peak is observed at 79.2 C as a cold crystallization. Tm as
endothermic transition peak is shown at 177.6 C, which is not remarkably different from Tm
of PLLA pellets at 170 °C [334]. These results confirm the molecular stability of PLLA fibers
under the physiological conditions (37 °C), without the occurrence of melting or re-
crystallization.
118
Regarding the XRD results in Fig. 6.14, the reduction of crystallinity of PLLA after
electrospinning is confirmed by quantitative measurement of the degree of crystallinity,
which is based on the DSC result (Fig. 6.15). According to Eq. 6.3, the percent crystallinity
(Xc) is determined by considering the heat energy of melting (∆Hm) and the heat energy of
cold crystalline (∆Hc), which are obtained from the area under the peaks. The crystallinity
(Xc) of PLLA fibers is 28.5 %, while Xc of the cast film has been reported in the range of 35
- 37 % [301]. Therefore, this result confirms the weak crystalline structure found by XRD,
which is the result of the electrospinning process.
Figure 6. 15 DSC thermogram of PLLA fibers, indicating Tg, Tc, and Tm.
119
(vii) Mechanical properties
Electrospun fibrous scaffolds are interesting for application in cartilage tissue
engineering because the fibers provide a high surface area to volume ratio and allow cells to
grow while providing also sufficient mechanical support [336]. In terms of the mechanical
properties, the high flexibility of the fibers is relevant for achieving low friction and adequate
load transfer, which are characteristic features of cartilage tissue, as discussed in Chapter 2.
In order to evaluate the mechanical properties of the fibers, tensile testing was performed to
compare also with PLLA cast films. As shown in Fig. 6.16 (A), a representative tensile
stress-strain curve of PLLA fibers shows the typical behavior of synthetic polymers, which is
composed of both elastic and plastic regions, as observed also for PLLA cast films (see the
inset in Fig. 6.16 (A)). Importantly, the fibers exhibited higher elongation at break ( 19.4 ±
0.3 %) when compared to the value of the cast films (11 ± 2 %), as reported in Fig. 6.17 (B).
This behavior indicates that PLLA fibers are more flexible than the cast films, which can be
ascribed to the fact that randomly oriented fibers generate highly intersectional interaction
along the mesh. In contrast, the elastic modulus and tensile strength of the fibers ( 18.7 ±
0.6 and 2.3 ± 0.1 MPa, respectively) are lower than those of the cast films ( 47 ± 10 and 21
± 3 MPa, respectively) (Fig. 6.16 (B)), which is likely due to the reduction of the crystallinity
in the fibers. When compared to the reference values reported in Table 2.2; Chapter 2, the
elastic modulus of PLLA fibers is confirmed to be in the range of the values obtained from
the native cartilage tissue (5 - 25 MPa). The tensile strength of PLLA fibers is closer to the
value range of the strength of native cartilage (3.7 - 10.5 MPa).
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Figure 6. 16 (A) Representative tensile stress-strain curve of PLLA fibrous meshes in
comparison with the typical curve of PLLA cast films (the inset) and (B) the values of elastic
modulus1, tensile strength
2 and elongation at break
3 of PLLA under tension deformation
mode.
The mechanical properties of crosslinked Alg/Gel fibers could not be tested due to
the very low yield of electrospun Alg/Gel fibers. Since the reproducibility is low, Alg/Gel
fibers were not able to be obtained as a stripe which is generally required as specimen for
tensile testing. Up to date, the mechanical properties of Alg/Gel fibers have not been
reported. Nevertheless, the range of mechanical properties available from the previous
studies with respect to the fiber diameter is summarized in Fig. 6.17, showing the trend of
the mechanical properties of fabricated synthetic and natural polymer-based fibers. In brief,
the elastic modulus of Alg/PEO fibers fabricated in the study of Bhattarai et al. [306] is
significantly lower than the values of synthetic-derived polymer fibers. This result is
suggested by the weak intrinsic strength of Alg itself and the lack of entanglement in the
electrospun Alg solution. In contrast, Gel fibers with and without crosslinking (Huang et al.
[337] and Panzavolta et al. [338]) exhibit the elastic modulus in the same range as the
1 Elastic modulus or Young’s modulus (E), according to tension deformation, was determined by using the same
procedure as compressive deformation. 2 Tensile strength in this case was represented as the values of strength at fracture or strength at break [383].
3 Elongation at break or elongation of fracture, which indicates the ductility of materials, was determined from the
strain at fracture in the stress-strain curve under tension force [383].
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values of PLLA fibers (Wang et al. [238] and Cao et al. [278]). In this case, the fibers of
Gel/PLLA (Yan et al. [301]) exhibited a suitable combination of good mechanical properties
and controlled degradation rate of PLLA, in addition to the convenient biological properties of
Gel.
Figure 6. 17 Summary of Young’s modulus values of electrospun fibers with respect to fiber
diameter, which were collected from recent literature reports [238,278,301,306,337–340],
mainly on polyesters, Gel and Alg. The star indicates the position of PLLA fibers obtained in
the present work.
6.4 Conclusions
Porous Alg-foams were successfully fabricated by the combination of gelation,
freeze-drying and physical crosslinking methods. The optimum conditions are composed of
3 wt/v % Alg solution, 0.1 M CaCl22H2O gelation agent and 0.5 M CaCl22H2O crosslinking
agent. The obtained foams exhibited high porosity 95 %, suitable pore size of 237 ± 48
µm and relative high water absorption ( 82 %). Elastic modulus and compressive strength
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were 0.22 ± 0.09 MPa and 0.14 ± 0.02 MPa, respectively. The foams showed high
interconnected pore structure and tubular-like porosity, which is suitable for cell arrangement
mimicking the organization of chondrocytes in native articular cartilage. Therefore, Alg-
foams fabricated in the present study provide the potential for use as a cartilage scaffold,
according to porous structure, pore geometry and physico-mechanical properties.
Alg fibers were fabricated by electrospinning of Alg/Gel/H2O solution (1/15/84 wt %).
The spin-ability of Alg-based solution was low in comparison with that of PLLA solution,
leading to low yield of obtained fibers. The reason is likely the low chain entanglement of the
alginate solutions used. The obtained fibers lack suitable structural stability and they
exhibited low water resistance, which was modified by chemical crosslinking (i.e. exposure
of GA vapor). The electrospinning of Alg solution requires further investigation in order to
enhance the spin-ability. For instance, proper solution systems, which could be based on
Gel addition are recommended to be further developed.
Here, Alg scaffolds based on two different structures, namely porous freeze-dried
foams and electrospun fibers, are comparatively studied. Alg-foams fabricated by freeze-
drying technique provide advantages in that they exhibit suitable porosity and pore size for
cartilage regeneration. The pore size in the range of 50 - 300 µm has been suggested to be
feasible for culturing with chondrocytes and for transporting nutrients to cells. In contrast,
fibrous scaffolds exhibit relatively small pore size ( 50 µm in the case of Alg/Gel fibers). In
addition, the columnar pores of Alg-foams can be tailored by the variation of the solution
used and conditions in the process of freeze-drying. On the other hand, Alg/Gel fibers with
fiber diameter of submicron dimension and relative large area per volume ratio exhibit
advantages to structurally mimic the native ECM and to support cell adhesion. In term of
manufacturing, Alg-foams show better reproducibility when compared to Alg/Gel-fibers.
From the materials point of view based on the studies presented in this chapter, it can be
preliminarily concluded that Alg-foams represent a more convenient scaffolds for further
investigations than electrospun structures in the context of their application in cartilage
regeneration strategies , e.g. as the cartilage-side scaffold in a multilayered system for
osteochondral regeneration.
CHAPTER 7
Multilayered Scaffolds Suitable for Osteochondral Tissue Engineering
7.1 Introduction
Osteochondral repair is one of the most challenging areas in the general field of
interface tissue engineering due to the complexity of tissues involved and limitation in self-
repair capability of cartilage [16,19,139,140,261]. Therefore, scaffolds based on biomimetic
approaches for osteochondral tissue engineering are becoming more sophisticated
[20,24,139]. Numerous research groups are concentrated in bi- or multilayered stratified
scaffolds in order to assemble the structure, architecture and functional properties of the
complex cartilage-bone interface tissues with the purpose to achieve the organization of
cells and new tissues by both in vitro and in vivo approaches. In addition, the formation of a
stable interface between cartilage and subchondral bone remains a significant challenge [8,
9], since the interface must exhibit sufficient structural integrity for both cartilage and SB to
become effectively connected. By employing a stratified scaffold approach, suitable scaffold
manufacturing techniques are necessary, representing an important factor to obtain optimal
porous structures and suitable mechanical properties of the rather complex resulting
composite scaffolds. In addition, manufacturing techniques are mainly dependent on the
selected biomaterials as well as on the applications. From a mechanical standpoint, cartilage
and bone are referred to as soft and hard tissues, respectively [261]. Natural polymers (such
as Col and polysaccharides) or water-soluble low molecular weight synthetic polymers are
suitable for fabricating cartilage scaffolds, whereas high molecular weight synthetic
polymers, ceramics, composites and metals are widely used for bone tissue engineering
scaffolds, which has been discussed extensively in Chapter 2. In order to fabricate a suitable
scaffold for the application of osteochondral tissue regeneration, the characteristics and
properties of involved tissues must be initially taken into account. The design of scaffolds,
selection of biomaterials and use of fabrication techniques are a crucial conceptual step. In
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the context of the present project, the specific topics related for selected scaffold materials,
the chosen fabrication techniques and suitable scaffold characterization approaches have
been discussed previously in Chapters 4 and 6 for the applications of the scaffolds in SB
and cartilage regeneration, respectively. The main goal of the investigation reported in the
present chapter was the development of multilayered scaffold approaches, which were
based on the optimum processing conditions reported in the previous chapters. The scope
of the work presented thus includes: (i) to develop novel multilayered composite scaffolds
with different approaches, according to biomimetic considerations and (ii) to overcome the
weak and unstable interface of the multilayered composite scaffolds in order to avoid
delamination during in vitro studies and for safe further implantation.
The multilayered scaffolds in the present work were categorized into four different
systems (Fig. 7.1), including (A) monolithic biphasic scaffolds (Alg-foam/Alg-c-BG scaffold),
(B) integrated bilayered scaffolds (Alg-foam/Alg–c-BG scaffold), and (C and D) integrated
electrospun fiber mesh/coated Bioglass® bilayered scaffold (PLLA fiber mesh/PDLLA-c-BG
and Alg/Gel fiber mesh/Alg-c-BG scaffold). It was noted that the polymer coating onto the
Bioglass-based scaffolds for bone phase was the same material as used to fabricate
cartilage scaffold due to the fact of chemical compatibility between dintinct phases. All
investigated systems were mainly different in the utilization of fabrication techniques and
consequently a variety of structural properties of the scaffolds was obtained. In detail, the
interfacial phase was generated by using two different methods, including (i) polymer
infiltration leading to a single biphasic scaffold with continuously formed interface (according
to system A) and (ii) polymeric adhesive serving as an interface, leading to integrated
bilayered scaffold (according to system B). On the other hand, the interface of scaffolds in
the systems C and D was created by directly electrospinning the polymer solution onto
Bioglass scaffolds. The fibers anchored onto the scaffold struts were aimed at enabling a
suitable integration between the two scaffold layers.
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Figure 7. 1 The schematic diagram of the four types of multilayered scaffolds for
osteochondral tissue engineering developed in this project.
7.2 Materials and methods
7.2.1 Fabrication of multilayered scaffolds
System A: Single biphasic scaffolds were fabricated as follows: a cylindrical Alg-g-
BG scaffold (fabricated as reported in Chapter 4; with the dimensions of 8 mm width, 8 mm
length and 5 mm height) was placed into a custom cylindrical mold (with the dimensions of 8
mm width, 8 mm length and 10 mm height) and then 800 µl of 3 wt/v % Alg solution was
added on top of the scaffold inside the mold. The Alg solution was allowed to infiltrate into
the porous scaffold for 10 sec, in order to form an intermediate phase. Then, 100 µl of 0.1 M
CaCl22H2O agent was added into the mold, in order to induce the gelation of the Alg
component. After 30 min gelation, the mold was frozen at – 20 C overnight and was
lyophilized at – 50 C under vacuum for 24 h. Finally, the obtained biphasic scaffold was
gently removed from the mold and crosslinked by immersion in 0.5 M CaCl22H2O agent for
4 h before being dried at room temperature for 24 h. The scaffolds were of the same
dimensions as the dimensions of the used mold.
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System B: 3 wt/v % Alg-foam (fabricated as reported in Chapter 6) was integrated
with Alg-c-BG scaffold (fabricated as reported in Chapter 4) by incorporation of 2 wt/v %
Alg/45S5 Bioglass® (Alg/BG 1:3 by wt.) solution, acting as an adhesive layer. In detail, a
small amount of Alg/BG adhesive was applied on one side of Alg-c-BG scaffold (with the
dimensions of 8 mm width, 8 mm length and 5 mm height) by using a painting brush
(Perikan No. 23). Then, a cubic shaped Alg-foam (without crosslinking; dimensions of 8 mm
width, 8 mm length and 5 mm thickness) was suddenly placed on the adhesive coated Alg-c-
BG scaffold and pressed manually. The obtained multilayered scaffold was immersed in 0.5
M CaCl22H2O agent for 4 h for crosslinking. In this step, the layers of Alg-foam and the
interface were exposed to the crosslinking agent. Finally, the crosslinked multilayered
scaffold (with the dimensions of 8 mm diameter and 10 mm height) was dried at room
temperature for 24 h.
System C: 7.5 wt/v % PLLA solution was electrospun directly onto PDLLA-c-BG
scaffolds (with the dimensions of 8 mm width, 8 mm length and 5 mm height), which were
adhered on the collector by using double-side adhesive tape. PDLLA-c-BG scaffolds were
fabricated by the same procedure described in Chapter 4. The electrospinning conditions
followed the optimum conditions described in Chapter 6 (18 kV supplied voltage, 15 cm of
needle tip-collector distance, 0.5 ml/h of feed rate and 2 h of deposition time). The PLLA
fiber mesh/PDLLA-c-BG bilayered scaffolds were gently removed from the collector and
were dried at room temperature for 24 h in order to evaporate remained solvent.
System D: The solution of Alg/Gel/DI H2O was prepared in the ratio of 1/15/84 wt %.
The Alg/Gel solution was directly electrospun onto Alg-c-BG scaffolds (with the dimensions
of 8 mm width, 8 mm length and 5 mm height), which were adhered on the collector. The
optimum processing conditions achieved in Chapter 6 were used (12 kV of voltage supplied,
12 cm of needle tip-collector distance and 0.01 ml/h of feed rate). Since the yield of Alg/Gel
electrospinning was very low, the Alg/Gel solution was allowed to be spun for 9 h in order to
obtain a fibrous mesh in the thickness of few hundred microns. Then, the Alg/Gel fiber
mesh/Alg-c-BG bilayered scaffolds were gently removed from the collector and the meshes
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were chemically crosslinked by exposing to GA vapor for 48 h. The crosslinking procedure
was detailed in Chapter 6.
7.2.2 Characterization and testing
(i) Microscopy
The appearance of the cross-sectioned multilayered scaffold was characterized
under light microscopy (LEICA M50) with camera operation of LEICA IC80 HD. The
microstructure of the multilayered scaffold was characterized by SEM (LEO 435 VP).
Samples were sputter coated and observed at an accelerating voltage of 10 kV.
(ii) Mechanical testing
Bilayered scaffolds (System B) with the dimensions of 8 mm width, 8 mm length and
10 mm height were tested in compression using a universal testing machine (Zwick Z050) by
applying the compression load at a cross-head speed of 2 mm/min, the preload at 0.1 N and
the maximum load at 50 N. The stress-strain curves were recorded in order to determine the
mechanical behavior of bilayered scaffolds and to compare it with the mechanical properties
of Alg-c-BG scaffolds and Alg-foams. The elastic modulus was calculated from the initial
linear slope in the stress-strain curve and the compressive strength was obtained from the
maximum stress before the sample collapsed. Eight specimens were tested and data were
presented as mean ± SD.
The mechanical strength at the interface of the bilayered scaffolds (System A versus
system B) was quantitatively investigated by using a micro-tensile testing machine
(Zwick/Roell 1120) by applying a maximum force of 100 N and a speed rate of 1 mm/min.
Specimens with dimensions of 3 mm width, 3 mm length, 10 mm height were placed in a
sample holder and fixed with the application of ethyl-cyanacrylate (Loctite® 454) glue and
activated with a glue activator (Loctite® 7455, Loctite Deutschland GmbH, Munich,
Germany). At least eight specimens were tested for each scaffold system and the data were
presented as mean ± SD.
(iii) In vitro acellular bioactivity
The bioactivity and delamination of multilayered scaffolds (Systems C-D) were
qualitatively investigated by immersion in SBF solution (according to Kokubo et al. [297]) (pH
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7.4) for 28 days. Each scaffold (System B: the dimensions of 8 mm diameter and 10 mm
height, and Systems C and D: the dimensions of 8 mm diameter and 5 mm height in height)
was put in a polystyrene bottle containing 50 ml SBF solution, which was then incubated in
an orbital shaker (IKA RS 4000i) at 37 C using a rotating speed of 90 rpm. The SBF
solution was replaced twice a week. At every time point, the sample was removed, cleaned
with DI H2O and dried at room temperature for 24 h. The morphology and microstructure of
scaffolds after immersion in SBF were investigated by SEM. The formation of HA was
characterized by using FTIR (Nicolet Nexus 6700, Thermo Scientific, Waltham, MA). The
samples for both, SEM and FTIR spectroscopy, were prepared following the same
procedure as described in Chapter 4 for Bioglass-based scaffolds and in Chapter 6 for Alg-
foams and fibers.
7.3 Results and discussion
7.3.1 Microstructure
Four different approaches for multilayered scaffolds were investigated for
osteochondral tissue regeneration. According to the literature [19,342], two different designs
of multilayered scaffolds suitable for osteochondral tissue engineering have been presented,
namely biphasic but monolithic scaffold and integrated bilayered scaffold. Up to date, it has
not been confirmed which of the designs, either monolithic biphasic scaffold or integrated
bilayered scaffold, is the optimal for this application. Therefore, both scaffold systems were
investigated and comparatively discussed in terms of structure, mechanical integrity at the
interface and in vitro bioactivity.
(i) System A: Monolithic Alg-foam/Alg-c-BG biphasic scaffolds
Fig. 7.2 (A) shows the cross-section of a monolithic Alg-foam/Alg-c-BG biphasic
scaffold, which forms a single continuous macroscopic structure including two distinct
phases. The biphasic scaffold was formed by the infiltration of Alg solution into the porous
Bioglass structure and applying fast gelation process, which leads to fusion of the Alg-gel
with the 3D Bioglass-based scaffold by forming an interconnected interface, as shown in
Fig. 7.2 (a). After lyophilizing, the interface (up to 500 µm in thickness; Fig. 7.2 (a))
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generated a porous structure with pore size 100 µm (Fig. 7.3 (A)). Some closed pores
were observed in some regions along the underlying Bioglass-based scaffold (Fig. 7.3 (A
and a)). This phenomenon is related to the relatively low controlled amount of Alg solution
during the infiltration process. It is likely that the excess amount of infiltrated solution can be
avoided by decreasing the infiltration time ( 10 sec), however the interface would not be
sufficiently strong by using shorter time (tested by trial-and-error).
(ii) System B: Integrated Alg-foam/Alg-c-BG bilayered scaffolds
In the case of system B (Fig. 7.2 (B)), Alg-foam served as the scaffold for the
cartilage side and Alg-c-BG scaffold served as the bone scaffold. These scaffolds were
prepared separately and integrated by using Alg/Bioglass adhesive. The intermediate
phase is observed as a dense thin layer in Fig. 7.2 (b). From SEM images (Fig. 7.3 (B and
b)), it was confirmed that Alg-foam and Alg-c-BG scaffold were well integrated by applying
an intermediate layer. It is likely that the intermediate layer provides a suitable connection
between the walls of the Alg-foam and the struts of the underlying Bioglass-based scaffold.
(iii) System C: Integrated PLLA fiber mesh/PDLLA-c-BG bilayered scaffolds
The bilayered scaffolds in system C (Fig. 7.2 (C)) were fabricated by electrospinning
PLLA on top of PDLLA-c-BG scaffolds. During electrospinning, the fibers tended to randomly
align onto the struts of the Bioglass scaffold, used as substrate forming the connecting
points between the two parts (as marked by a dashed circle in the inset; Fig. 7.2 (C)). In
addition, some polymer jets were able to infiltrate through the pores of the Bioglass-based
scaffold, leading to additional integration between both phases, as observed in Fig. 7.2 (c).
As shown in the SEM images, the connection points were obvious between both phases (as
marked by dashed circles in Fig. 7.3 (C)). Moreover, the electrospun PLLA fiber mesh was
obtained as a relatively dense mesh with thickness of around 50 µm (Fig. 7.3 (c)).
(iv) System D: Integrated Alg/Gel fiber mesh/Alg-c-BG bilayered scaffolds
In this approach, random Alg/Gel fibers did not closely pack and became a dense
mesh like in the case of PLLA fibers (Fig. 7.3 (D)). Subsequently, the Alg/Gel fiber mesh
provided larger pore sizes and interconnectivity, which were maintained even after applying
chemical crosslinking, when compared to the PLLA fiber mesh (Fig. 7.3 (d)). In this case in
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fact, a fluffy mesh was obtained. In addition, it was found that the fibers did not infiltrate into
the pores of the underlying Bioglass-based scaffold (Fig. 7.3 (D)). This behavior is likely
the result of the intrinsic low strength of the electrospun Alg/Gel solution, as discussed in
Chapter 6. This result might lead to low strength at the interface of the bilayered scaffold.
Figure 7. 2 Optical microscopic images showed the appearance of three different
approaches of multilayered scaffolds, including (A, a) system A: monolithic Alg/Alg-c-BG
biphasic scaffold, (B, b) system B: integrated Alg/Alg-c-BG bilayered scaffold (Reproduced
from Nooeaid et al. [145] with the permission of John Wiley and Sons) and (C, c) system C:
integrated electrospun PLLA fibers/PDLLA-c-BG bilayered scaffold (the inset shows a plan-
view of fibers integrated on the struts of the Bioglass-based scaffold).
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Figure 7. 3 SEM images showing cross-sections of four different types of multilayered
scaffolds: (A, a) system A, (B, b) system B (Reproduced from Nooeaid et al. [145] with the
permission of John Wiley and Sons) (C, c) system C and (D, d) system D (the dashed line
marks the interface between the two phases).
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7.3.2 Interfacial strength of multilayered scaffolds
Since the interface has a significant impact on the mechanical integrity of bilayered
scaffolds, the interfacial strength of each scaffold system was quantitatively investigated by
application of micro-tensile testing. However, the testing apparatus could not be used on the
scaffolds belong to systems C and D due to the limited 3D structure of the fiber mesh. Thus,
the relative tensile stress-strain curves of bilayered scaffolds in systems A and B were
comparatively investigated in the present work (Fig. 7.4 (A)). Under tension, the interface of
scaffolds exhibited elastic deformation at the initial stage (I) until reaching the yield point,
where the interface-layer starts to break. Afterwards, at stage (II) deformation progresses
continuously until the breaking point (III). Moreover, the stress-strain curves of the interface
followed mainly a typical profile of polymeric materials. The integrated bilayered scaffolds
(System B), which include the adhesive-layer, exhibited higher Young’s modulus (linear
slope at stage (I)) and exhibited higher strength at break, compared to the case of monolithic
layered scaffolds, which have continuous interface. The strength at break of the integrated
bilayered scaffolds was determined in the range of 0.08 - 0.2 MPa, while the strength at
break of the monolithic biphasic scaffolds was lower (up to 0.1 MPa) (Fig. 7.4 (B)). It can be
concluded that the interface formed by the application of an adhesive layer was stronger
than the continuous interface formed in-situ during scaffold fabrication. The reason is likely
the good interconnection between the two distinct phases as a result of the presence of
dense crosslinked Alg/Bioglass adhesive-layer. In contrast, the continuous interface formed
by polymer infiltration through the pores of Alg-c-BG scaffold was not as strong as expected.
This result is explained by the fact that the continuous interface was highly porous, leading
to relative lower strength in comparison with the dense interface-layer formed in the case of
integrated bilayered scaffolds.
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Figure 7. 4 (A) Representative stress-strain curves of the bilayered scaffolds (System A vs.
System B) and (B) distribution of the strength at break values of the scaffolds in systems A
and B (the red dashed line is included for the visual aid) (Reproduced from Nooeaid et al.
[145] with the permission of John Wiley and Sons).
7.3.3 Mechanical properties of integrated bilayered scaffolds
In addition to the quantification of the interface strength, the mechanical behavior of
bilayered scaffolds under compressive loading is essential to be investigated. Since joints in
the skeletal system (i.e. knee joint) have to withstand local compression forces during
motion, the compressive stress-strain curves of bilayered scaffolds was monitored in order
to study their mechanical response under compression. According to the better results
achieved from the mechanical interface testing (Fig. 7.4), the system B scaffolds were
chosen to be investigated in this study, in comparison with individual Alg-foam and Alg-c-BG
scaffold. The typical stress-strain curve of the bilayered scaffolds (curve a; Fig. 7.5)
exhibited a sigmoid shape, which represents the same mechanical behavior as the Alg-foam
(curve b; Fig. 7.5). Exceptionally, at the initial stage (in the region up to 20 % deformation)
the curve of bilayered scaffolds showed limited jagged behavior. This is indeed the typical
characteristic of brittle cellular solids (i.e. ceramic foams) [56,317], which was also presented
in the curve of Alg-c-BG scaffold (curve c; Fig. 7.5). These results indicate that the
mechanical response of the system B bilayered scaffolds is predominantly influenced by the
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Alg-foam. It can be concluded that Alg-foam mainly acts as a load absorber and load
transferor for the bilayered scaffold, which is supported by the high deformability of the foam
in comparison to brittle Bioglass-based scaffold. Therefore, Alg-foam is confirmed as being
suitable for use as a cartilage scaffold according to the required mechanical performance,
including elasticity and softness [24,343,344].
Figure 7. 5 Representative compressive stress-strain curve of integrated bilayered scaffold
(system B) in comparison with the curves of Alg-foam and Alg-c-BG scaffold (Reproduced
from Nooeaid et al. [145] with the permission of John Wiley and Sons).
7.3.4 In vitro bioactivity
(i) System B: Integrated Alg-foam/Alg-c-BG bilayered scaffolds
It was confirmed that the bilayered scaffold maintained the structural integrity over
28 days in SBF without delamination (Fig. 7.6). This result indicates a strong adhesion
between the two distinct phases. Moreover, the porous structure of the Alg-foam was also
maintained without deformation during immersion in SBF. Importantly, HA did not form on
Alg-foams after 28 days in SBF, which is a desirable behavior avoiding mineralization of the
cartilage side of the scaffold [24,103,125,345]. In the case of Alg-c-BG scaffold, the results
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are in an agreement with the bioactivity assessment in Chapter 4. Alg-c-BG scaffold
exhibited the expected ability of HA formation, with HA deposited on the struts of Bioglass-
based scaffold after 28 days of immersion in SBF. However, at day 1 HA formation had not
been observed yet due to the Alg coating, as discussed in Chapter 4. In addition, HA
formation was observed at the interface of the bilayered scaffold (see the inset; Fig. 7.6),
even on the pore walls of the Alg-foam that are in contact with the struts of the Bioglass-
based scaffold. This phenomenon is attributed to the effect of the thin Alg/Bioglass
adhesive layer. It is likely that Bioglass® incorporated into the adhesive improved the
bioactivity of the interface, mimicking the structure of highly mineralized calcified cartilage
[5,8,226].
The no formation of HA on the cartilage side of scaffold (Alg-foams) was confirmed
by ATR-FTIR spectroscopy, as shown in Fig. 7.7. The characteristic peaks of HA were not
detected. In addition, the intensity of the Ca-O peak at 1009 cm-1
was reduced after 28 days
in SBF (the inset in Fig. 7.7), indicating the loss of crosslinking and consequently the
reduction of the mechanical stability of the foam. In contrast, the peak at 1022 cm-1
attributed
to O-H bending increased in intensity after 28 days in SBF. This result indicates that free
negative groups (COO-) are present after losing ionic-interaction during immersion in SBF,
leading to interaction with water molecules and subsequently more O-H groups are
generated, compared to the foam before immersion.
In summary, the desired bioactivity of the bone side of the scaffold was achieved in
addition to the required non-mineralization of the cartilage-side of the scaffold. Moroever, the
bilayered structure of scaffolds in system B was maintained after 28 days of immersion in
SBF, indicating that the fabricated bilayered scaffolds in the present study exhibited the
strong interface and they can avoid delamination during experimental period.
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Figure 7. 6 SEM images of integrated bilayered scaffold (system B) after immersion in SBF
for 1 and 28 days (the dashed line indicates the interface between the two phases)
(Reproduced from Nooeaid et al. [145] with the permission of John Wiley and Sons).
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Figure 7. 7 ATR-FTIR spectra of Alg-foam after immersion in SBF for 28 days in order to
confirm non-mineralization of the foam (the inset shows the absorption bands in the
wavenumber 1200 - 850 cm-1
).
(ii) System C: Integrated PLLA fiber mesh/PDLLA-c-BG bilayered scaffolds
The bioactivity assessment of system C bilayered scaffolds is presented in Fig. 7.8
by SEM images. After 28 days of immersion in SBF, no delamination was observed in all
investigated scaffolds, indicating suitable adhesion strength at the interface between the
PLLA fiber mesh and the PDLLA-c-BG scaffold. As shown in the inset, some fibers were
adhered to the strut of the PDLLA-c-BG scaffold. In the PDLLA-c-BG layer, HA formation
was not obvious after 1 day in SBF, while HA almost completely covered the struts of
PDLLA-c-BG scaffold after 28 days in SBF. In contrast, HA did not form on the PLLA fiber
mesh after 28 days of immersion in SBF. Hence, PLLA fiber mesh intended for the cartilage
side of the scaffold exhibited no mineralization as required for its use as a cartilage scaffold.
Moreover, the fiber mesh still maintained the interconnectivity during the entire immersion
period in SBF due to the relatively slow degradation of PLLA.
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Figure 7. 8 SEM images of integrated bilayered scaffold (system C) after immersion in SBF
for 1 and 28 days.
(iii) System D: Integrated Alg/Gel fiber mesh/Alg-c-BG bilayered scaffolds
In these scaffolds, after 28 days of immersion in SBF, it was observed that the layer
of Alg/Gel fibers became a thin membrane (Fig. 7.9; interface), when compared to the SEM
image of the scaffold before immersion in SBF (Fig. 7.3 (D and d)). It is likely that
interconnection points formed among the fibers after immersion in SBF (as marked by the
dashed circles in Fig. 7.9 (b)). Since the melting temperature of gelatin ( 250 bloom) is
close to the SBF temperature (at 37 °C) [301,346], the Alg/Gel fibers (300 bloom of used
Gel) can be partly melted at the incubation conditions, forming a stronger interconnected
network. Therefore, this effect is suggested to be the reason for the obtained denser layer of
the mesh after 28 days of immersion in SBF. In addition, some cracks were found along the
fiber mesh (Fig. 7.9), which this behavior is likely caused by the breakage of some fibers, as
shown in Fig. 7.9 (b). This result indicates that the Alg/Gel fibers are not sufficiently strong in
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order to maintain their structure and to avoid deformation during incubation, even though
chemical crosslinking was applied. In addition, the Alg/Gel fibers did not induce HA
formation, leading to a non-mineralized phase, which is desired for the cartilage side of the
scaffold. In contrast, HA was formed and fully deposited on the strut surfaces of the Alg-c-
BG scaffold after 28 days in SBF, as shown in Fig. 7.9.
Figure 7. 9 SEM image of integrated bilayered scaffold (system D) after immersion in SBF
for 28 days (the dashed line indicates interface between the Alg/Gel fiber mesh and the Alg-
c-BG scaffold and the dashed circles indicate the interconnection between the fibers formed
after immersion in SBF).
7.4 Conclusions
Four approaches of bilayered scaffolds were successfully fabricated for use as a
scaffold in osteochondral tissue engineering applications, according to the porous structure,
mechanical properties and bioactivity. Interestingly, integrated bilayered scaffolds exhibited
higher interfacial strength compared to monolithic but biphasic scaffolds. In addition, the
formation of a controllable interface (by the process of polymer infiltration) of monolithic
biphasic scaffolds was not effectively achieved. In the case of application of electrospinning
of PLLA directly on the Bioglass scaffold, the fibrous layer formed was densely packed.
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This leads to small pore size with increasing the thickness of the mesh. In contrast, the
electrospinning of Alg/Gel led to fiber mesh, which was loosely packed. This mesh, however,
exhibited low interconnectivity and subsequently led to low mechanical properties. The GA
crosslinked Alg/Gel fibers were partly deformed after 28 days of immersion in SBF.
Moreover, the Alg/Gel mesh was challenging to produce due to its low spin-ability.
Therefore, the system B – (Alg-foam/Alg-c-BG) bilayered scaffolds are the promising
approach for further investigation and for its use as a suitable scaffold for osteochondral
tissue regeneration. This scaffold type was therefore considered for subsequent studies
discussed in Chapter 9.
CHAPTER 8
Biological Response of Osteoblasts Culturing on Bioglass-based
Scaffolds for Bone Regeneration
8.1 Introduction
An approach in the field of bone tissue engineering requires suitable biomaterial
scaffolds, cells and biomolecules and/or growth factors [55,247]. The 3D scaffold has to
support cell adhesion, proliferation and differentiation, and induce ECM deposition, and
therefore new bone can be grown in 3D [247,262]. The scaffold biomaterials for bone
regeneration are required to provide osteoconductive or even osteoinductive characteristics
in order to accelerate osteogenic differentiation of stem cells [28,347]. 45S5 Bioglass® has
been well established as a bone scaffold due to its osteoconductivity, because of its ability to
bond to bone [31,34,93,276]. In addition, Bioglass can up-regulate the expression of
osteoblast genes in response to dissolution products (i.e. Ca, P and Si), which plays a role in
controlled and enhanced osteogenesis [35,41]. However, highly porous Bioglass®-based
scaffolds are rather brittle and mechanically weak, thus they are not suitable for use in load
bearing applications, as mentioned in Chapter 4. Therefore, biodegradable
polymer/Bioglass® composite scaffolds become a target of interest in bone tissue
engineering approaches [52]. Since the integrated bilayered scaffolds (Alg-foam/Alg-c-BG)
developed in Chapter 7 showed a promising potential for use in osteochondral tissue
engineering, in this chapter Alg-c-BG as SB scaffolds were targeted to study the biological
responses by cultured with bone-like cells. By the fact that Alg has no adhesive sites to cells
and does not adsorb serum proteins due to its high hydrophilicity [61,69], peptides with a cell
adhesive sequence-modified Alg (i.e. Arg-Gly-Asp (RGD) containing amino acid) (RGD-Alg)
was therefore used to enhance cell adhesion on Alg coating in this study. Since amino acid
sequence RGD in fibronectin acts as a primary cell attachment cue, it has been
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demonstrated that RGD linear peptide coupling to Alg can enhance osteoblasts adhesion for
several times when compared to unmodified Alg [71].
Therefore, in vitro culture of MG-63 osteoblast-like cells on RGD-Alg and Alg coated
Bioglass®-based scaffolds was investigated in comparison with uncoated Bioglass
®-based
scaffolds, in order to confirm the potential of the developed composite scaffolds for use in
osteochondral tissue engineering applications. Qualitative and quantitative cell adhesion and
proliferation, and cell activity of MG-63 cultured on the coated scaffolds for 3, 7, 14 and 21
days were discussed in this chapter, in comparison with equivalent results on uncoated
Bioglass-based scaffolds.
8.2 Material and methods
8.2.1 Fabrication of Bioglass®-based scaffolds
The fabrication of uncoated Bioglass scaffolds has been described in Chapter 4.
As-fabricated Bioglass scaffolds were sterilized by heat treatment. In detail, the scaffolds
were placed in a furnace (Model B180; Nabertherm GmbH, Germany) and were heated at
160 °C for 7 h. In the case of Alg-c-BG scaffolds, 1.5 wt/v % Alg in DI H2O was prepared as
a coating solution and it was filtered by using a sterilized mesh (pore sizes ~ 0.45 µm). After
that, the sterilized scaffolds were coated with the Alg solution by dipping using a solution of 5
ml per scaffold. The coated scaffolds were then dried under sterilized hood for 24 h. After
drying, Alg-c-BG scaffolds were immersed in 0.5 M CaCl22H2O sterilized solution (1 ml per
scaffold) for 4 h in order to crosslink the Alg coating, as optimized previously in the case of
Alg-foams for cartilage scaffolds. Finally, crosslinked Alg-c-BG scaffolds were washed with
sterilized DI H2O and dried in the sterilized hood for 24 h. RGD-modified Alg coated
Bioglass-based scaffolds (RGD-Alg-c-BG) were prepared by dipping sterilized Bioglass-
based scaffolds in 1.5 wt/v % RGD-Alg solution. Briefly, 100 µM RGD-coupled Alg
(NovaMatrix, FMC Biopolymers, Norway) was mixed with 1 ml Alg solution (1.5 wt/v %).
Then, the RGD-Alg solution was filtered using the same procedure as in the case of the pure
Alg solution. Finally, the sterilized scaffolds were coated and crosslinked as previously
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mentioned in the case of Alg-c-BG scaffolds. It is noted that the coating, drying and
crosslinking processes were performed under sterilized hood.
8.2.2 In vitro cell culture
The sterilized scaffolds were pre-treated in culture medium for 24 h, which were
incubated at 37 C, 5 % CO2 and 95 % humidity, in order to stabilize the pH variation due to
the ionic exchange process between scaffolds’ surface and medium. After that, the scaffolds
were gently washed with PBS before cell seeding.
MG-63 osteoblast-like cells were grown in T-150 culture flasks containing culture
medium and kept at 37 C in an atmosphere of 5 % CO2. The cells were trypsinized and
collected as a cell pellet by 5 min of centrifugation at 1200 rpm. Then, the cells were
suspended in culture medium. The culture medium is composed of Dulbecco’s modified
eagle medium (DMEM) with low glucose with L-glutamine supplemented with 10 v/v % fetal
bovine serum (FCS), 1 v/v % penicillin-streptomycin and ascorbic acid. The scaffold was
placed in a 48 well-plate and subsequently seeded with 1 ml cell suspension containing 1
million cells. The cell-seeded scaffolds were cultured in an incubator (at 37 °C, 5 % CO2 and
95 % humidity) for 3, 7, 14 and 21 days. The culture medium was renewed twice a week.
8.2.3 Characterization techniques
(i) Lactate dehydrogenase (LDH) assay
The relative number of cells was evaluated by in vitro toxicology assay kit; lactate
dehydrogenase (TOX7) (Sigma-Aldrich). This assay involves the measurement of the
number of cells via total cytoplasmic LDH [348]. Four scaffolds per each type were placed in
a 48-well plate. Then, 1 ml of lysis buffer was added in each well and incubated at room
temperature for 30 min. After that, 1 ml solution was transferred into a 1.5 ml reaction
vessel. The cell lysis supernatant (140 ml) was obtained after centrifuging the reaction
vessel for 5 min at the speed of 1200 rpm, and it was transferred into a 1 cm cuvette. 60 µl
of LDH substrate mixture (20 µL of each LDH assay dye, LDH assay substrate and LDH
assay growth factor) was added into the cuvette, then the mixture was incubated at room
temperature in dark condition for 30 min. The enzymatic reaction was stopped by addition of
1 N HCl (300 µl per cuvette). Finally, 500 µl of DI H2O was added before the measurement
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by UV-Vis spectrophotometer at wavelength 490 and 690 nm was performed. The LDH
activity is presented as the difference between the absorbance at wavelength 490 nm and
the absorbance at wavelength 690 nm.
(ii) Cell imaging
Cell viability, distribution and formation of HA were qualitatively investigated by using
fluorescence microscopy (ZEN with fluorescent lamp HXP120C). DAPI (4, 6-diamino-2-
phenylindole) staining was used to visualize the cell nuclei. The fluorescent bright blue is
observed when the dye binds selectively to double stranded DNA. Encountered OsteoImage
(Lonza) dye was used to visualize HA of the bone-like nodules, which was deposited by
cells, in green fluorescence.
Before staining, the cells-seeded scaffolds were washed with PBS and cultured with
staining reagents starting from OsteoImage mineralization assay for 30 min, followed by
fluorescent fix agent for 15 min. The scaffolds were washed with PBS and were
counterstained with DAPI for 10 min. Finally, the stained scaffolds were kept in PBS before
analyzing.
(iii) Cell morphology
The cell morphology of MG-63 seeded on BG, Alg-c-BG and RGD-Alg-c-BG
scaffolds after 14 days in culture was analyzed by SEM (LEO 435 VP). After the cells-
seeded scaffolds were washed with PBS, they were fixed with 3 v/v % GA in 0.1 M sodium
cacodelate for 1 h and subsequently they were dehydrated by using the series of EtOH
solutions (30, 50, 70, 80, 90, 100 %, respectively) for 30 min. Finally, the fixed samples
were dried by using critical point drying (Leica EM CPD300, Germany). All the samples were
sputter coated and observed under SEM at 10 kV.
(iv) AlamarBlue assay
Cell metabolic activity was quantitatively evaluated by using AlamarBlue (AB) cell
reagent (Invitrogen). AB assay is based on oxidation-reduction reaction, which indicates the
cells’ metabolism activity [349]. Six scaffolds per each type were investigated at 3, 7, 14 and
21 days culture. The scaffolds were placed in a 48-well plate, then 500 µl of 10 v/v % AB dye
in culture medium was added in each well. The samples were placed in the incubator at 37
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C and 5 % CO2 atmosphere for 4 h. After that, 500 µl of medium from each well was taken
into a 1 cm cuvette. 500 µl of DI H2O was then added into the cuvette. The absorbance
value of solution was measured at wavelength 570 and 600 nm by using UV-Vis
spectrophotometer (Specord 40; Analytikjena, Germany). The results are presented in % AB
reduction with respect to culture time, which were determined by using Eq. 8.1 [349]:
% AB reduction = [ALW – (AHW × R0)] × 100 (8.1)
; where ALW and AHW are the difference between the absorbance of sample and the
absorbance of medium blank at 570 and 600 nm, respectively, and R0 is correction factor,
which was determined by using Eq. 8.2 [349]:
R0 = AOLW/AOHW (8.2)
; where AOLW and AOHW are the difference between the absorbance of AB mixture and the
absorbance of medium blank at wavelength 570 and 600 nm, respectively.
(v) Alkaline phosphatase (ALP) assay
The ALP assay was used to evaluate the osteoblastic activity, considering that
osteoblasts are very rich in ALP enzyme [350]. The ALP enzyme catalyzes the hydrolysis of
p-nitrophenyl phosphate (p-NPP; transparent) into p-nitrophenol (p-NP; yellow) [275,350].
Four scaffolds of each type were investigated. Briefly, 150 µl of cell supernatant was added
into a 1 cm cuvette. 100 µl of DI H2O was added, then 100 µl of ALP buffer at pH 9.8 (0.1 M
Tris; Mw 12114 g/mole, 2 nM MgCl2; 95.3 g/mole, 9 mM p-NPP) was added. The cuvette
was incubated at 37 C in dark condition. The reaction was stopped by addition of 1 N NaOH
(300 µl), while the reaction time was recorded. The cuvette was filled with DI H2O (350 µl)
and then the absorbance was measured by using UV-Vis spectrophotometer at wavelength
405 and 690 nm. The relative ALP activity was determined by using Eq. 8.3:
ALP activity (OD/min/g) = ∆A/(T × P) (8.3)
; where ∆A is the difference between the absorbance values at 405 nm and at 690 nm, T is
reaction time in min and P is total protein concentration (g/ml). The protein content was
determined by the method of Bradford protein assay with respect to the calibration of
standard protein solutions. The standard solutions were prepared at concentrations of 0,
100, 200, 400, 600, 800 and 1000 µg/ml. The absorbance of the standard solution was
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measured by using UV-Vis spectrophotometer at a wavelength 595 nm. The calibration
curve was obtained as Eq. 8.4:
Total protein concentration (µg/ml) = (A595 – 0.014)/0.001; R2 = 0.99 (8.4)
The total protein amount in each sample (25 µl of supernatant and 975 µl of Bradford
solution in a 1 cm cuvette) was obtained by measuring the absorbance at wavelength 595
nm by using UV-Vis spectrophotometer.
8.2.4 Statistical analysis
All data were analyzed by using one-way analysis of variance (ANOVA) and turkey’s
multiple-comparison test to determine statistical differences. A confidence interval of 95 % (p
= 0.05) was used for all analyzes. Mean values and SD are presented.
8.3 Results and discussion
8.3.1 LDH activity
In vitro LDH activity describes the relative number of cells over culture time, which
was determined by measuring relative LDH activity from cell lysates [348]. In Fig. 8.1, it is
seen that the number of cells on uncoated BG scaffolds significantly increased over the
culture time, in particular after cultivation for 14 days (* p 0.05). This result indicates the
proliferation of MG-63 cells cultured on uncoated BG scaffolds. At the initial culture period
(up to 7 days), cells seeded on both Alg-c-BG and RGD-Alg-c-BG scaffolds exhibited higher
LDH activity when compared to cells seeded on uncoated BG scaffolds, in particular cells
seeded on RGD-c-BG scaffolds showed the highest LDH activity. On the other hand, LDH
activity of cells seeded on uncoated BG scaffolds was significantly higher than that of cells
seeded on coated scaffolds after 14 days of culture. This result can be explained that at the
initial time of culture (up to 7 days), intense ionic exchange between the surface of uncoated
BG scaffolds and the culture medium induced increased pH of the medium, leading to
decreased cell proliferation. This response is in agreement with the studies of in vitro cell
culture of PDLLA coated and P(3HB) coated Bioglass-based scaffolds [55,351]. It can be
speculated that polymer coating plays a role in surface reactivity of Bioglass-based
scaffolds, which the coating inhibits the ion exchange at the Bioglass surface in the initial
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culture. Furthermore, at day 14 in culture, the cells seeded on uncoated BG scaffolds
exhibited the highest LDH activity compared to the cells seeded on coated scaffolds. The
reason is likely that over a longer period of time, the dissolution of the polymer coating took
place in contact with the culture medium, leading to pH variation and subsequently inhibition
of cell growth on the coated scaffolds. This explanation can be confirmed by low LDH activity
of cells seeded on coated scaffolds at day 21 of culture, compared to cells seeded uncoated
scaffolds. In addition, the result of LDH activity confirms more cells of RGD-Alg-c-BG
scaffolds compared to Alg-c-BG scaffolds. It is likely that RGD-Alg-c-BG scaffolds supported
better the growth of MG-63 cells over culture time compared to Alg-c-BG scaffolds.
Figure 8. 1 Relative LDH activity of MG-63 osteoblast-like cells cultured on uncoated BG,
Alg-c-BG and RGD-Alg-c-BG scaffolds. The results presenting the difference of optical
densities are presented as mean ± SD (n = 4). * (p 0.05) indicates significant difference of
different scaffolds at different culture times.
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8.3.2 Cell imaging
Cell adhesion and distribution were qualitatively evaluated by DAPI staining, giving a
blue color by using fluorescence microscopy. Fluorescent DAPI staining images in Fig. 8.2
confirms effective infiltration of seeded MG-63 cells through the porous structure of all
investigated scaffolds and confirms the adhesion of cells on the surface of all the scaffolds
after 3 and 14 days in culture. It is seen that MG-63 cells exhibited quite homogeneous cell
distribution on the struts of all uncoated BG, Alg-c-BG and RGD-Alg-c-BG scaffolds. By this
result, the number of cells was difficult to distinguish among all 3D scaffold types. In
addition, it can be concluded that even though Alg lacks cell adhesion moieties, the present
study proved that porous Alg-c-BG scaffolds can support adhesion of MG-63 cells. Similar to
the previous study of Srinivasan et al. [42] that porous Alg-based scaffolds are
biocompatible with MG-63 cells and can support the growth of cells.
Figure 8. 2 Fluorescent microscopic images of MG-63 osteoblast-like cells-seeded BG, Alg-
c-BG and RGD-Alg-c-BG scaffolds after 3 and 14 days in culture by using DAPI stain for cell
nuclei.
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8.3.3 Metabolic activity
The metabolic activity of MG-63 cells seeded on BG, Alg-c-BG and RGD-Alg-c-BG
scaffolds was presented in % AB reduction. Fig. 8.3 shows that in the case of uncoated BG
scaffolds, % AB reduction increased over culture time and in particular % AB reduction
significantly increased (* p 0.05) after 14 days of culture. MG-63 cells maintained their
metabolic activity on uncoated BG scaffolds over 21 days in culture, which is indicated by
the reached values of % AB reduction. This result indicates that uncoated BG scaffolds can
activate metabolic activity of osteoblast-like cells over culture time. This phenomenon is
supported by the effect of dissolution product of Bioglass. In contrast, MG-63 cells seeded
on both Alg-c-BG and RGD-Alg-c-BG scaffolds did not show increased % AB reduction over
the culture time. It confirms the LDH activity (Fig. 8.1) that progressive cell proliferation was
not observed on the coated scaffolds, which is indicated by no changes in cell metabolic
activity over the culture time. In addition, it is likely that both coated scaffolds, Alg-c-BG and
RGD-Alg-c-BG, did not lead to significant difference in cell metabolic activity.
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Figure 8. 3 Cell metabolic activity of MG-63 osteoblast-like cells cultured on uncoated BG,
Alg-c-BG and RGD-Alg-c-BG scaffolds. The results in % AB reduction are presented as
mean ± SD (n = 6). * (p 0.05) indicates significant difference of different scaffold types at
different culture times.
8.3.4 Osteoblastic activity
Fig. 8.4 shows ALP activity results of MG-63 cells on the surface of BG, Alg-c-BG
and RGD-Alg-c-BG scaffolds. ALP activity increased in all types of scaffolds with increasing
culture time. After 14 days in culture, all investigated scaffolds promoted significant increase
of ALP activity compared to the culture at day 3 (* p 0.05). At day 7 in culture, increased
osteoblastic activity was not significantly detected. In addition, the results of ALP activity are
in agreement with the results of LDH activity and cell metabolic activity. The increase in LDH
activity cell and metabolic activity leads to an increase in total protein adsorption and
subsequently ALP activity [275]. At day 3 of culture, cells seeded on Alg-c-BG and RGD-Alg-
c-BG scaffolds exhibited significantly lower ALP activity than that detected in cells seeded on
uncoated BG scaffolds (* p 0.05). This result can be explained by the fact that the
dissolution products from uncoated BG scaffolds can activate gene expression in
osteoblasts [42], while the release of ions from the Bioglass surface in the case of coated
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scaffolds may be inhibited by the polymer coating at the initial stage of culture, as previously
discussed in the result of LDH activity. However, during 7 and 21 days in culture, there is no
significant difference among all scaffold types, indicating that by longer period of culture
time, the polymer coatings did not negatively affect the osteoblastic activity. It is suggested
that the resorption of the coating layer occurs with increasing culture time and thus ions
released from the coated scaffolds has an effect, as supported by the in vitro bioactive study
of Alg-c-BG scaffolds in Chapter 4. In addition, it is likely that the uncontrollable release of
Ca ions from crosslinked Alg and RGD-modified Alg coatings did not negatively influence the
ALP production of cells cultured on both coated scaffolds. Also, the Ca ions seem to support
the ALP activity of the cells cultured on both coated scaffolds. After 14 days of culture, there
is no further increase of ALP activity in all scaffolds, indicating complete osteoblastic activity.
As also reported by the previous study on Alg/nano Bioglass composite scaffolds seeded
with MG-63 cells, the cells exhibited maximum ALP activity at day 7 and the ALP activity
further decreased after prolong culture, which this phenomenon correlates to the maturation
of the cells [42]. Moreover, at day 14 and day 21 in culture there is no significant difference
of ALP activity between Alg-c-BG and RGD-Alg-c-BG scaffolds. This result indicates that
RGD-Alg-c-BG scaffolds did not improve the activity of the cells in terms of ALP production.
In summary, all scaffold types, including uncoated BG, Alg-c-BG and RGD-Alg-c-BG
scaffolds, can support the activity of ALP production of MG-63 cells. Since ALP is an early
indicator of mineralization of osteoblasts, further tests such as Col production, osteocalcin
and osteopontin should be evaluated in order to confirm the mineralization of MG-63 seeded
on the scaffolds.
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Figure 8. 4 ALP activity up to 21 days of MG-63 osteoblast-like cells cultured on uncoated
BG, Alg-c-BG and RGD-Alg-c-BG scaffolds. The results are reported as mean ± SD (n = 4).
* p 0.05 indicates significant difference of results for different scaffold types at different
culture times.
8.3.5 Cell morphology
SEM images in Fig. 8.5 confirm the cell growth of MG-63 cells on all scaffold types
after 14 days in culture. The SEM images (Fig. 8.5 (A-C)) showing overview of porous
scaffolds confirm that the cells were able to infiltrate into the porous structure of all scaffold
types. This observation indicates that the fabricated Bioglass-based scaffolds exhibited
suitable pore size and porosity for supporting the cell seeding and infiltration, and further cell
growth. Polymer coatings (both Alg and RGD-Alg) did not inhibit infiltration of cells, since
open pores still maintained. In addition, the cells are seen to be elongated on the struts of
scaffolds and the cells grew reaching tens of microns in all investigated scaffold types, as
shown in Fig. 8.5 (D-F). The well flattened cells covering the scaffold struts tended to group
and formed a monolayer in all scaffold types. It is likely that there is no significant difference
of cell morphology and cell distribution among all scaffold types. In addition, the crystals
were formed on the struts of all investigated scaffolds, in particular the uncoated scaffolds.
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The crystals are likely to be HA crystals that were initiated by ionic exchange between
Bioglass surface and the culture medium. This behavior has been previously reported
[55,351,352]. In the case of coated scaffolds, it can be explained that the dissolution of
polymer coating enables direct contact of the Bioglass surface with the medium, leading to
formation of HA crystals. It can be also seen that the cells grew and covered the crystals, as
obviously shown in Fig. 8.5 (E and F).
The HA formation can be confirmed by fluorescent images of stained OsteoImage
(green color) on MG-63-seeded scaffolds after 3, 7, 14 and 21 days in culture (Fig. 8.6). The
formation of HA was observed in all investigated scaffolds after 3 days in culture and it
tended to increase with culture time.
Figure 8. 5 SEM cross-sectioned images showing MG-63 cells-seeded BG, Alg-c-BG and
RGD-Alg-c-BG scaffolds after cultured for 14 days.
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Figure 8. 6 Confocal microscopic images of MG-63 cells cultured on BG, Alg-c-BG and
RGD-Alg-c-BG scaffolds, stained with OsteoImage (green), after 3,14 and 21 days.
All the results indicate high biocompatibility and the ability to stimulate metabolic
activity of MG-63 osteoblast-like cells on BG, Alg-c-BG and RGD-Alg-c-BG scaffolds. The
present study proves that Alg and RGD-Alg coatings did not provide negative effect on cell
activity. In addition, Alg-c-BG and RGD-Alg-c-BG scaffolds could support the adhesion,
growth and activity of osteoblast-like cells similar to previous studies of polymer coated
Bioglass-based scaffolds, for example PDLLA-c-BG and PHB-c-BG scaffolds [55,351].
Even though the polymer coating generally reduces surface roughness of Bioglass-based
scaffolds and therefore inhibits cell adhesion, this phenomenon did not appear in the present
study by Alg coating, as evidenced by SEM images in Fig. 8.5. In addition, negatively
charged Alg coating did not negatively affect cell adhesion. Therefore, it can be speculated
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that dissolution products of Bioglass and pH variation of the culture medium in the initial
culture plays the important role in this study. By this fact, the study of pH variation of the
culture medium over culture time should be taken into account.
8.4 Conclusions
From the performed in vitro biological investigation, all scaffolds investigated,
namely uncoated BG, Alg-c-BG and RGD-Alg-c-BG scaffolds, supported MG-63 osteoblast-
like cells adhesion and osteoblastic metabolic activity. The results indicate that all types of
investigated scaffolds are compatible with MG-63 cells and can support their growth. Given
high porosity and macro-pores of both Alg-c-BG and RGD-Alg-c-BG scaffolds promoted
effectively cell infiltration through the porous structure of the scaffolds similar to the case of
uncoated scaffolds. RGD-modified Alg coated onto BG scaffolds led to higher number of
cells over culture time compared to pure Alg coated BG scaffolds, as evidenced by LDH
activity. In addition, cell metabolic activity, proliferation and ALP of all scaffold types were
quantitatively confirmed. The ALP activity results showed that there is no significant
difference of osteoblastic activity among all investigated scaffolds after 14 days in culture.
Moreover, HA increasingly deposited on the surface of all scaffolds with increasing culture
time, which was confirmed by OsteoImage stained fluorescent images. By the present
results, uncoated BG and RGD-Alg-c-BG scaffolds are attractive for bone tissue engineering
applications, while Alg-c-BG scaffolds are limited in their adhesion ability. In addition, further
evaluations should be determined such as Ca content, Col I production, osteocalcin and
osteopontin in order to confirm the mineralization ability of osteoblast-like cells seeded on
the scaffolds. Since the use of MG-63 osteoblast-like cells cultured on the scaffolds in the
present study preliminarily confirmed their osteoblastic activity, MSCs cultured on the
scaffolds should be further studied in order to confirm the osteoblastic differentiation.
Here, the present chapter has confirmed the ability of fabricated Bioglass-based
scaffolds for use in subchondral bone tissue engineering applications, while the porous Alg-
foams fabricated by using freeze-drying technique (as presented in Chapter 6) were aimed
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to investigate the potential for use in cartilage regeneration and were presented in the next
chapter.
CHAPTER 9
Biological Response of Chondrocytes and Mesenchymal Stem Cells on
Alginate/Chondroitin Sulfate Scaffolds for Cartilage Regeneration
9.1 Introduction
In tissue engineering approaches, scaffolds, cells and biomolecules are used as
main components for tissue regeneration [19,163,247]. Besides the requirement of a
scaffold material to provide a proper environment for cell growth, the cell source has to be
considered as well in order to develop a specific strategy for tissue regeneration. In the
cartilage engineering approach based on MSCs, MSCs have to be retained within the
scaffold and be capable to undergo chondrogenic differentiation in the specific scaffold
environment [163,242]. In this chapter, porous Alg-foams fabricated by freeze-drying
technique (as detailed in Chapter 6) were aimed to confirm the potential for use as a
cartilage scaffold. Alg scaffolds have been shown to promote proliferation and differentiation
of MSCs to chondrocytes and subsequently to provide their expression of Col II and PGs
[198,234,250,353]. However, Alg allows only a limited cell adhesion, because it does not
exhibit functional groups in order to be recognized by the cells [69]. Therefore, Alg scaffolds
should be modified to enhance the cell adhesion. This functionalization has been done by
grafting with RGD, which facilitates integrin recognition and binding [70,245], grafting with
Gel [134], and blending with fibrin [200], CS [201,354,355] and Col [356,357]. In addition,
chondrocytes and MSCs have been extensively considered to study the optimization of the
scaffold properties and culture conditions. Autologous chondrocytes have a capacity to be
expanded in vitro and match the host’s immunological system; however, chondrocytes tend
to lose their specific phenotype during monolayer expansion [131,234,358]. Another
approach for cartilage tissue engineering is based on MSCs, which have the ability to
differentiate into multiple tissues, including bone and cartilage [242,359]. MSCs exhibit less
donor-site morbidity, they are cost-effective, they can easily be expanded due to their high
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proliferation capacity and yield to equal or better long-term outcomes in comparison to
chondrocytes [305]. Moreover, biomolecules and growth factors, which are supplemented
into scaffolds, are crucial for the success of differentiation of MSCs and may influence the
regeneration of cartilage [198,358]. For instance, Chang et al. [360] fabricated
Col/HyA/Chondroitin sulfate (ChS) scaffolds by using the same composition as reported for
the natural cartilage. The scaffolds have shown chondrocytes adhesion and good cell
distribution, and supported the secretion of cartilaginous ECM. Coates et al. [198] have
shown that the incorporation of ChS into HyA and Alg hydrogels has a positive effect on
chondrogenesis, because it up-regulates Sox-9 mRNA and down-regulates Col I. This was
explained by the fact that ChS, which was used for cartilage repair, helps to regulate the
metabolic activity of chondrocytes [198] and stimulates the generation of PGs [361].
Therefore, ChS tends to act as a biological additive in order to regulate chondrocyte
phenotype [198] and increase cell proliferation [361,362]. This can be achieved either by
loading the molecules into the scaffolds [234,361] or by using them as a supplement in the
culture medium [358]. ChS introduced into the culture medium have been shown to increase
the production of sulfated mucopolysaccharides by cultured chondrocytes [363]. In addition,
Steinmetz et al. [364] reported that the presence of ChS in culture medium reduced the
production of Col I during the terminal differentiation of MSCs encapsulated in PEG
hydrogel, in particular when the dynamic culture was applied.
In the present work, in a first attempt ChS was incorporated into the Alg porous
scaffolds, with the aim to improve their cell response and cell activity. First, the impact of
ChS on the physical and mechanical properties of the scaffolds was investigated. Another
aim was to study the influence of ChS on the biological properties of scaffolds by in vitro
porcine chondrocytes and MSCs culturing. Two different culture conditions, including static
and dynamic cultures, were applied in the case of chondrocytes culturing Alg-foams. This
experiment was carried out with the aim to study the effects of different culture conditions on
chondrogenic differentiation. The cell viability, retention of cell phenotype and chondrogenic
differentiation were investigated.
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9.2 Materials and methods
9.2.1 Fabrication of Alg/ChS-foams
A mixture of Na-Alg (purchased from Sigma Aldrich) and chondroitin-4-sulfate A
sodium salt (ChS; purchased from Sigma Aldrich) was prepared in DI H2O with the
concentration of 3 wt/v % and the ratio of 85/15 % by wt. of Alg/ChS. The mixture was stirred
at room temperature for 2 h. 1 ml of the Alg/ChS solution was added into a 48 well-plate,
while 100 µl of 0.1 M CaCl22H2O solution was added per well. The gelation reaction was
finished after 30 min at room temperature. After this, the gel was placed into the freezer at -
20 C and it was frozen after 24 h. Then the frozen samples were lyophilized by using the
freeze-drying technique for 24 h at - 50 C under vacuum conditions (as explained in
Chapter 6). Cylindrical 3D porous Alg/ChS-foams were obtained, having dimensions of 8
mm in diameter and 8 mm in height. Then the foams were immersed in 0.5 M CaCl22H2O
solution (pH 2) for 4 h, in order to achieve ionic crosslinking. The crosslinked foams were
dried at room temperature for 24 h. Pure Alg-foams were fabricated by using the same
procedure as the one for the fabrication of Alg/ChS-foams. Thereby a solution of 3 wt/v %
Na-Alg/H2O was used.
9.2.2 Characterization and testing
(i) Porosity
The porosity of Alg/ChS-foams (P) was calculated from Eq. 9.1:
% P = [1 – (Wfoam/(Alg/ChS × Vfoam))] × 100 (9.1)
Here, Wfoam is the weight of the Alg/ChS-foam, Alg/ChS is the density of solid Alg/ChS blend
(Alg/ChS 1.02 g/cm3), and Vfoam is the volume of Alg/ChS-foam, which was determined from
the dimensions of the foam.
(ii) Microscopy
The morphology of Alg/ChS-foams was characterized by SEM (LEO 435 VP), by
using plan-view and cross-section imaging in order to observe the features of the pores.
Samples were sputter-coated and observed at an accelerating voltage of 10 kV. The pore
size of foams was evaluated from SEM images by using the free available software Image J.
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(iii) Chemical structure
The existence of ChS in the foams was identified by using ATR-FTIR (Nicolet 6700)
spectroscopy in the transmission mode. Thereby the wavenumber resolution was 4 cm-1
and
the range of 4000 - 525 cm-1
. Moreover, the existence of ChS was confirmed by using
energy-dispersive X-ray (EDX) spectroscopy, being available in the SEM instrument (SEM-
EDX; Inca analyzer, Oxford instruments).
(iv) Thermal properties
The thermal properties of Alg/ChS-foams were analyzed by using DSC (Q2000).
The measurements were performed in the temperature range of (- 50) - 200 C, by using a
heating rate of 10 C/min. The results of Alg/ChS-foams were compared to those of Alg-
foams.
(v) In vitro biodegradation
The biodegradation and water absorption of Alg/ChS- and Alg-foams were
investigated by immersion in PBS solution (pH 7.4, at 37 C) for 6 weeks. Each sample was
placed in a polystyrene bottle containing 50 ml of PBS and incubated in an orbital shaker
(IKA RS 4000i) at 37 C, using a speed of 90 rpm. The PBS solution was replaced twice a
week. At interval of immersion time, the sample was removed, blotted with filtered paper
before the weight was determined. The water absorption and weight change were
determined by using Eq. 9.2 and 9.3, respectively:
% Water absorption = [(Wwet – Wdry)/Wwet] x100 (9.2)
% Weight change = [(Wwet – Wdry)/Wdry] x100 (9.3)
Here, Wwet and Wdry is the weight after and before immersion in PBS, respectively. The
presented results are averaged over four samples.
(vi) Mechanical testing
Compression strength experiments were performed on Alg/ChS- and Alg-foams (8
mm in diameter and 8 mm in height), by using a universal testing machine (Zwick Z050).
The following parameters were used: cross-head speed of 2 mm/min, pre-load of 0.1 N and
maximum load of 50 N. The stress-strain curves were recorded in order to determine the
relevant mechanical properties, including the elastic modulus and compressive strength. The
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elastic modulus was calculated from the initial linear slope of the stress-strain curves and the
compressive strength was obtained from the maximum stress before the cell walls of the
foam were collapsed. The presented results are averaged over six samples.
In addition, a DMA (Mark IV) was performed in order to study the mechanical
response of the foams under the application of forces similar to those dominating in vivo
[313] . The foams (8 mm in diameter and 3 mm in thickness) were tested in both, the dry and
wet state, by applying a sinusoidal load with a rate defined by the frequency (in Hz). Before
testing in the wet state, the foams were immersed in PBS solution until reaching equilibrium
conditions. The foams were tested in the compression mode as a function of the frequency
cycle. The used frequency was in the range of 0.1 - 10 Hz, according to the range of a typical
skeletal movement in vivo [313]. The tests were performed under the maximum strain
amplitude of 1 %. The viscoelastic properties of the foams - the storage modulus (E’) and
loss factor (tan = E’’/E’) - are presented by using the log scale and plotting versus the log
scale of frequency.
9.2.3 Release of ChS
The relative ChS content in the foams was determined according to the modified
Dische’s carbazole reaction, which has been modified by Bitter and Muir [365]. In detail,
each of Alg/ChS- and Alg-foams (8 mm in diameter and 3 mm in thickness) was immersed in
5 ml of PBS solution (pH 7.4, 37 C) in a glass vial. At each time point of immersion, 2 ml
of PBS solution was taken out and replaced by the same amount of fresh PBS. Afterwards,
0.2 ml of the obtained solution was given for reaction with 1 ml of 0.025 M sodium
tetraborate decahydrate (Borax; ACS reagent 99.5 %, Sigma-Aldrich) in concentrated
sulfuric acid (H2SO4, Sigma-aldrich), with the aim to initiate degradation of ChS. The mixture
was then cooled down to room temperature and heated up in a boiling water bath for 10 min.
After that, the mixture was again cooled down to room temperature. 40 µl of 0.125 wt/v %
carbazole ( 95 % GC, Sigma Aldrich) in EtOH was added into the mixture and heated up in
the boiling water bath for further 15 min, until pink chromogen was formed. Finally, the
absorbance of the pink colored complex was measured at 530 nm by using the UV
spectrophotometer (Specord 40; Analytikjena, Germany). The amount of the released ChS
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was calculated by using a calibration curve of known ChS concentrations (0 - 120 µg/ml), as
shown in equation 9.4:
Absorbance = (0.0018 x ChS concentration) – 0.00497, R2 = 0.99. (9.4)
9.2.4 In vitro culturing of primary porcine chondrocytes and human MSCs
This part of the study was carried out in Laboratory for Experimental Trauma
Surgery, Department of Orthopedic, Trauma and Reconstructive Surgery Charité-University
Medicine Berlin, Campus Benjamin Franklin.
(i) Isolation of porcine chondrocytes
The chondrocytes were isolated according to the method reported by Lohan et al.
[131]. Briefly, porcine cartilage was harvested from the articular cartilage in the knee joint of
pigs (3 - 6 months old). The removed cartilage samples were minced into 1 mm slices. Then,
the samples were enzymatically digested with 0.4 % pronase (7 U/mg, Roche, Basel,
Switzerland) and diluted in Ham’s F-12/Dulbecco’s modified Eagle’s (DMEM) medium (1:1)
(Biochrom AG, Berlin, Germany) for 1 h at 37 C. Then the samples were subsequently
digested with 0.2 wt/v % collagenase ( 0.1 U/mg, SERVA Electrophoresis GmbH,
Heidelberg, Germany) and diluted in the growth medium for 16 h at 37 C. The isolated
chondrocytes were suspended in the growth medium [Ham’s F-12/DMEM (1:1) containing 1
ml/100 ml fetal calf serum (FCS) (Biochrom AG), 25 µg/ml ascorbic acid (Sigma-Aldrich), 50
IU/ml streptomycin, 50 IU/ml penicillin, 2.5 µg/ml amphotericin B and 1 ml/100 ml essential
amino acids (all from Biochrom AG)]. The cell suspension was seeded at a density of 2.8 ×
104 cells/cm
2 in culture flasks.
(ii) Isolation of human bone marrow derived MSCs
Human MSCs were isolated from human femoral head spongiosa (obtained from
patients undergoing joint replacement surgeries of the hip joints) using density gradient
centrifugation with a separating solution, biocoll (Biochrom AG, Berlin, Germany). The
spongiosa of a femoral head was minced and pressed through a cell sieve. Bone spongiosa
fragments were removed and the liquid rest was pressed through a 140 μm pore diameter
filter membrane. To remove the remnants of the particles the isolated cell suspension was
washed with PBS and centrifuged at 200 g in 4 °C. The purified pellet was mixed with the
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biocoll solution (Biochrom AG) and centrifuged at 200 g in 4°C. After 20 min, the interphase
containing MSCs was extracted, washed with PBS and centrifuged at 200 g in 4 °C.
Subsequently, MSCs were re-suspended in stem cell growth medium [DMEM 51 ml/100 ml
(Biochrom AG, Berlin, Germany) containing selenium (5 ng/ml; Aldrich), transferring (5
µg/ml; Sigma), linoleic acid (4.7 µg/ml; Sigma), insulin (5 µg/ml; Sigma), ascorbic acid
(1µg/ml; Sigma), dexamethasone (1 µg/ml; Sigma D4902), MCDB 201 with L-glutamine
solution (34 ml/100 ml; Sigma), Foetal calf serum (FCS; 15 ml/100 ml; Biochrom),
streptomycin (50 IU/ml) and penicillin (50 IU/ml)] and seeded in culture flasks (Cell plus
culture flask, Sarstedt, Nümbrecht, Germany). The cultivation proceeded at 37 °C, 90% air
humidity and 5% CO2. The growth medium was changed every 2 - 3 days.
(iii) Static and dynamic cultures
Two types of cells, namely primary porcine articular chondrocytes and human
MSCs, were used. The Alg/ChS- and Alg-foams (both: 8 mm in diameter and 3 mm in
thickness) were sterilized by using a plasma treatment, respectively. The sterilized Alg-
foams were then pre-conditioned for 72 h, while the sterilized Alg/ChS-foams were pre-
conditioned for 24 h, in the basal medium. After that, the foams were soaked in a
suspension of either chondrocytes or MSCs in the basal medium (10 million cells of passage
2 - 3 of chondrocytes and passage 4 - 5 of MSCs). The cell-seeded foams were then placed
in an incubator at 37 C and 5 % CO2. The medium was changed 3 times a week.
The induction of in vitro chondrogenesis of MSCs seeded on the foams was
evaluated by culturing with the presence of TGF-1. MSCs-seeded Alg- and Alg/ChS-foams
were cultured in chondrogenic medium [DMEM with 3.7 g/l NaHCO3 and 4.5 g/l glucose
(Biochrom AG, T041-01) containing 10 µg/ml L-glutamine (Biochrom AG, K0282), 25 µg/ml
HEPES (Biochrom AG, L1613), 10 µg/ml sodium pyruvate (Sigma, S8636), 0.1 µl/ml
dexamethasone (Sigma, D4902), 1.7 µl/ml ascorbic acid (Sigma, A8960), 1 µg/ml prolin
(Sigma, P8865), 1 µl/ml ITS+1 (Sigma, I2521), 50 IU/ml streptomycin, 50 IU/ml penicillin and
10 ng/ml TGF-1 (Petro Tech)], which were placed in an incubator at 37 C and 5 % CO2.
The medium was changed 3 times a week.
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During dynamic culture conditions the growth medium and the material degradation
products have to be exchanged after 2-3 days. However, dynamic culturing provides a more
homogeneous cell distribution, compared to static culturing [131,366]. Therefore, two
different culture conditions were comparatively studied in the case of chondrocytes culturing
on the Alg-foams. The dynamic culturing was performed by stirring the cell suspension in a
bioreactor filter tube (TPP, Switzerland) using an orbital shaker at 15 °C and the speed of 12
rpm (digital type roller Struart SRT9D, Bibby Scientific, USA). The medium was changed 3
times a week.
(iv) Cell viability
Cells-cultured for 7 and 14 days in Alg- and Alg/ChS-foams were washed with PBS
and incubated in fluorescein diacetate (FDA; Sigma-Aldrich) (3 µg/ml dissolved in acetone
and diluted 1:1000 in PBS) for 15 min at 37 C. Then, the samples were rinsed three times
with PBS and counterstained with a propidium iodine (PI) solution (1 mg/ml dissolved in PBS
and diluted 1:100 in PBS) (Sigma-Aldrich) for 1 min under dark conditions at room
temperature. The green and/or red fluorescence was visualized by using fluorescence
microscopy (Axioskop 40, Carl Zeiss, Jena, Germany) or confocal laser microscopy (TCS
SPE II, Leica Microsystems, Wetzlar, Germany). Fluorescence images were taken using an
XC30 camera (Olympus Soft Imaging Solution GmbH, Germany).
(v) Histology
For hematoxylin eosin (HE) staining, the samples were stained with a Harris
hematoxylin solution (Sigma-Aldrich) for 4 min, rinsed with water and counterstained in eosin
(Carl Roth GmbH, Karlsruhe, Germany) for 4 min.
For alcian blue (AB) staining, the samples were incubated in a 1 % acetic acid for 3
min and then stained with 1 % AB (Karl Roth, Karlsruhe, Germany) for 30 min. After that, the
stained samples were rinsed with 3 % acetic acid and counterstained with a nuclear fast red
aluminum sulfate solution (Carl Roth) for cell nuclei staining for 5 min. Finally, the samples
were rinsed with water and subsequently dehydrated in an ascending EtOH series. Then,
the samples were observed histologically by using light microscopy (Axioskop 40, Carl
Zeiss) and images were taken using a XC30 camera.
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(vi) Immunohistology
The samples were rinsed with Tris buffered saline (TBS: 0.05 M Tris, 0.015 M NaCl,
pH 7.6) and incubated with 5 mg/ml pronase (7 U/mg; Roche, Basel, Switzerland, diluted in
TBS) for 5 min at 37 C. The samples were subsequently rinsed with TBS and blocked with
a protease-free donkey serum (5 % diluted in TBS) for 30 min at room temperature. Then
the samples were rinsed again with TBS and incubated with the polyclonal rabbit anti-type II
or anti-type I collagen antibodies (27.5 µg/ml) (both: Acris antibodies, Herford, Germany), by
using a humidifier chamber overnight at 4 C. After that, the samples were rinsed with TBS
and incubated with a donkey-anti-rabbit-Alexa-Fluor® 488 (10 mg/ml, Invitrogen) secondary
antibody for 30 min at room temperature. Negative controls included omitting the primary
antibody during the staining procedure. Finally, cell nuclei were counterstained with DAPI
(0.1 µg/ml) (Roche). The labeled samples were rinsed several times with TBS, embedded
with Fluoromount G (Southern Biotech, Biozol Diagnostica, Birmingham, USA) and
examined using fluorescence microscopy (Axioskop 40).
9.2.5 Statistical analysis
The data were analyzed by using the one-way analysis of variance (ANOVA) and
the Turkey’s multiple-comparison test in order to determine statistical differences. The
confidence interval of p = 0.05 was used for all analysis. The results are reported as mean
values, with SD.
9.3 Results and discussion
9.3.1 Characterization of Alg/ChS-foams
(i) Morphology
The plan-view SEM image of an Alg/ChS-foam in Fig. 9.1 (A) shows a uniform
distribution of pores, while the cross-section SEM image in Fig. 9.1 (B) reveals a columnar
structure of pores with ladders, which was generated by the same mechanism as discussed
in the case of pure Alg-foams in Chapter 6. Alg/ChS-foams exhibited slightly decreased
porosity ( 93 %) compared to pure Alg-foams (porosity 95 %). The pore size of Alg/ChS-
foams (197 ± 61 µm) was slightly smaller than the one of Alg-foams (237 ± 48 µm) (Fig. 9.3
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(C)). In addition, micropores ( 50 µm) were observed on the pore walls of the Alg/ChS-foam
(Fig. 9.1 (B)). This result confirms that the Alg/ChS-foams exhibited a combination of micro-
and macro-pores. Even though only 15 wt % of ChS was incorporated into the foam, ChS
may influence the molecular arrangement of the Alg network. According to the literature
[367,368], the reptation model involves an arrangement of two different anionic
polysaccharides, in which two different types of polymer cannot intersect each other. Alg,
which is the main component of the foam, forms an egg-box network by ionic crosslinking. A
Ca ion interacts with four carboxylate groups of the Alg chains. The crosslinked Alg network
is ordered in a tube-like structure, as shown in Fig. 9.3 (D). The ChS molecules are confined
into the Alg network. Therefore, in a proper environment the ChS molecules move inside the
Alg-network like a snake [367,368]. Fajardo et al. [368] have reported that the ChS
molecules maintain inside the tubular structure of Alg under high acidic conditions (pH 2),
get mobilized and further diffuse through the network, causing an increase of the pH (pH
5).
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Figure 9. 1 SEM images of Alg/ChS-foams in (A) plan-view and (B) cross-section; (C)
shows a histogram of the pore size distribution of Alg/ChS-foams and (D) is a scheme
showing the reptation model in the case of Alg/ChS blend, according to [367,368].
(ii) Chemical structure
In order to confirm the incorporation of ChS into the Alg-foams, the chemical
structure of the Alg/ChS-foams was investigated by using ATR-FTIR spectroscopy. The
ATR-FTIR spectrum of an Alg/ChS foam in Fig. 9.2 (A) exhibited peaks at identical positions
as observed for the Alg-foam. The characteristic peaks of Alg (marked by ▲) at 1600 and
1427 cm-1
can be attributed to vibrations of -C=O and -COOH groups [368,369]. The
characteristic peaks of ChS (marked by ■) at 1605 and 1558 cm-1
can be attributed to -C=O
and -N-H- stretching, respectively [370]. The peak of ChS at 922 cm-1
was attributed to the
vibration of C-O-C groups, while the peak at 853 cm-1
assigned the vibration of C-O-S
groups [370], which were both not present in the spectrum of Alg/ChS-foams. This can be
explained by the low content of ChS incorporated into the foam. Only a broad peak at 1258
cm-1
was detected for the Alg/ChS-foam, which can be attributed to the -S=O stretching
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vibration of sulfate groups in ChS [368,370,371]. Therefore, it can be concluded that the
incorporation of ChS into the foam was successful.
The EDX results in Fig. 9.2 (B) additionally confirm the presence of ChS in the foam.
The peaks of sulfur (S) were detected in the case of the Alg/ChS-foam, while they were not
pronounced in the case of the pure Alg-foam. The FTIR and EDX results confirm the
presence of ChS within the Alg-foam. In addition, the ionic crosslinking in the foams was
confirmed by strong peaks of calcium (Ca), detected by EDX (see Fig. 9.2 (B)).
Figure 9. 2 (A) ATR-FTIR spectrum and (B) EDX spectra of the Alg/ChS-foam in
comparison to the spectra of pure Alg- and pure ChS-foams, which both results confirm the
existence of ChS in the foam.
(iii) Thermal properties
The thermal properties of the foams were characterized by DSC. Fig. 9.3 (A) shows
the heat transition in endothermic direction, indicating the melting of the material [372,373].
The melting temperature (Tm) of the Alg/ChS-foam is at 83.4 °C, which was lower than the
Tm of the Alg-foam, being at 96.59 °C. At this point, the polymer molecules of the Alg/ChS-
foam needed a higher temperature for melting, compared to the Alg-foam. Regarding
molecular point of view, the Alg/ChS-foams exhibited a lower chain mobility [374–376],
169
compared to Alg-foams. The reason for this behavior is the presence of ChS, which may
inhibit the interaction between Alg and Ca ions, leading to less intermolecular interactions,
compared to a network without ChS. Nevertheless, this hypothesis has to be proved by the
determination of the crosslinking degree.
Fig. 9.3 (B) shows the endothermic transition of the foams, obtained from the 2nd
heating cycle. This transition results from the energy absorption, required for the change of
the glassy state (hard, brittle) to the rubbery state (soft, flexible), also known as the glass
transition [373]. The Tg of Alg/ChS-foams (29.9 C) was slightly higher than the Tg of Alg-
foams (24.1 °C). The increase in Tg can be the result from the main chain rigidity and
crosslinking. Hence, this result suggests a possible bulk rigidity and brittleness, caused by
the incorporation of ChS.
In summary, even though the incorporation of ChS is believed to influence the
thermal properties of the foams, it can be ensured that the Alg/ChS-foams are not able to be
melted and decomposed under in vitro culture conditions at 37 C. In addition, the results of
the Tg confirm that at 37 C (culture conditions) the foams only become softer, while the
structural integrity is still maintained.
Figure 9. 3 DSC thermograms: (A) the 1st heating and (B) the 2
nd heating cycle of Alg/ChS-
and Alg-foams.
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(iv) Water absorption
The water absorption of scaffolds used for cartilage regeneration is an important
factor related to the local pressure resistance of the structure [313,377]. In general, the
negative charges of GAGs in the cartilaginous ECM mainly cause the absorption of water
molecules, which leads to a reduced friction at the osteochondral joint [361,362]. It has been
reported that articular cartilage is composed of 80 % water with respect to the total wet
weight [6]. In the present work, a water absorption of 82 ± 2 % was observed for the Alg-
foams (see Fig. 9.4), which is similar to the reference value of 80 % [43]. The water
absorption of Alg/ChS-foams was measured to be 88 ± 2 % and thus it is higher than the
water absorption value of the Alg-foams. The higher water absorption of the Alg/ChS-foams
is caused by the negative charges of ChS, which enhance the capacity of a foam to absorb
water [362].
In summary, both, the Alg/ChS- and Alg-foams did not exhibit a significant difference
of water absorption from the reference value reported in the literature [43]. Therefore, in
terms of high water absorption, both types of foams are suitable for cartilage regeneration.
Figure 9. 4 Water absorption in % of Alg- and Alg/ChS-foams, in comparison to the water
absorption of natural articular cartilage (*), as referenced in the literature [6].
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(v) Biodegradation behavior
In addition to the water absorption, the structural stability of scaffolds has to be
maintained until new tissue grows. Therefore, in order to evaluate the structural stability of
the scaffolds, a biodegradation study was performed. The weight change of the foams,
which is caused by water absorption, was determined during immersion in PBS. The weight
change of Alg/ChS-foams in PBS ( 750 %) was slightly higher than the weight change of
Alg-foams ( 650 %), while the trend of the weight change in both cases is the same (Fig.
9.5). In detail, the foams initially absorb a high amount of the PBS solution and start to
maintain an increased weight after 1 day of immersion. Importantly, after 3 weeks of
immersion, the weight change of the Alg/ChS-foams slightly decreases, while the weight
change of the Alg-foams is maintained until the end of the experiment (6 weeks). It can be
concluded that the Alg/ChS-foams start to lose their stability after 3 weeks of immersion in
PBS, which might be a sign of decomposition.
Figure 9. 5 Degradation profile of Alg/ChS- and Alg-foams in PBS, evaluated by % weight
change with respect to the immersion time.
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(vii) Mechanical properties
Representative compressive stress-strain curves of the Alg/ChS- and Alg-foam are
shown in Fig. 9.6 (A). The mechanical properties (elastic modulus and compressive
strength) of the Alg/ChS-foams are significantly higher than those of the Alg-foams (* p
0.05) (Fig. 9.6 (B)). It is believed that the incorporation of ChS into the Alg-foam reinforces
the Alg matrix. According to the reptation model shown in Fig. 9.1 (D), the molecules of ChS,
which are confined in the Alg network, decrease the void space in the network. The
incorporation of ChS significantly increases the elastic modulus of the Alg-foams from 0.22 ±
0.09 to 0.7 ± 0.1 MPa and increases the strength from 0.11 ± 0.01 to 0.8 ± 0.2 MPa.
Especially, the elastic modulus of the Alg/ChS-foams reaches the upper reference value for
healthy human cartilage, which has been reported to be 0.24 - 0.85 MPa [6,12,73,76,141–
143]. Alg/ChS-foams, which exhibit a high elastic modulus (high stiffness), are therefore
suitable in load bearing applications, i.e. cartilage and bone tissue engineering.
Fig. 9.7 presents the storage modulus (E’) and tan of the dry and wet foams as a
function of frequency. It is obvious that the presence of ChS increases the viscoelastic
properties of the dry foams. The E’ values of the dry foams are higher than the values of the
wet foams, for both, Alg- and Alg/ChS-foams. This behavior is caused by the entrapped
water in the wet foams, which acts as a plasticizer and decreases the stiffness of the foams
[313,319,332]. In addition, the E’ values of the wet Alg/ChS-foams are lower than the values
of the Alg-foams. This result is caused by the higher PBS absorption of the Alg/ChS-foams
(see Fig. 9.5) and the incorporation of negatively charged ChS. The results correlate with tan
results, where both, the wet Alg- and Alg/ChS-foams exhibit higher values compared to the
dry foams. Since in the wet foams the pores are filled with PBS, it is difficult to deform them
by applying a dynamic force. Moreover, the Alg/ChS-foams exhibit higher tan values than
the Alg-foams in both, the wet and dry states. This behavior can be explained by the theory
of dynamic mechanical properties of polymers that the loss factor is defined as the loss
modulus (E’’) divided by the storage modulus (E’): tan = E’’/E’. Since the absorption of
water decreases the stiffness (E’ decreases), tan increases at constant E’’, which means
that the internal friction of the material increases. In addition, the E’ values of both, the Alg-
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and Alg/ChS-foams, obtained in the wet state are in the range of values reported for the
cartilage tissue (0.1 - 1 MPa, 0.1 - 10 Hz) [313]. These results confirm that the viscoelastic
properties of Alg- and Alg/ChS-foams (either in dry or wet state), do not change with the
variation of frequency and time, which indicates that no deformation occurs during oscillatory
testing [343,378].
Figure 9. 6 (A) Representative compressive stress-strain curves of the Alg/ChS- and Alg-
foam; (B) mechanical properties of Alg/ChS and Alg-foams.
Figure 9. 7 Dynamic mechanical analysis of Alg/ChS- and Alg-foams: storage modulus (E‘)
and tan as a function of frequency; note the logarithmic scaling.
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9.3.2 Release profile of ChS molecules
The release of ChS, which is incorporated into the foams, was investigated by
immersion in PBS solution (pH 7.4 at 37 C). By using carbazole reaction described by
Dische [379], a pink colored complex of ChS can be detected at 530 nm wavelength by
using UV-VIS. Fig. 9.8 represents the cumulative release of ChS with respect to the
immersion time in PBS for 14 days. The ChS was gradually released with increasing time in
PBS. However, after day 7 of immersion, no further ChS was released. As confirmed by
immersion for 14 days, the content of the released ChS was not changed compared to day 7
of immersion. This result confirms that at a pH 5 the molecules of the ChS were released
from the tubular network of Alg, as in agreement with [368]. In addition, the ChS release rate
observed in the present work can be used as a guideline for further in vitro cell culture
studies. The ChS delivered to the culture medium is aimed to act as a biological cue, in
order to up-regulate the chondrogenic differentiation.
Figure 9. 8 Cumulative release of ChS from Alg/ChS-foams immersed in the PBS solution
(pH 7.4, 37 °C) measured by using carbazole reaction, as described in Section 9.2.3.
175
9.3.3 The influence of culturing conditions on the primary porcine chondrocytes
activity
(i) Cell viability
The influence of two types of culture conditions (static and dynamic) on the behavior
of primary porcine chondrocytes-seeded Alg-foams was investigated. The aim was to study
the influence of the culture conditions on the cell activity, in particular the cell viability and
production of cartilaginous ECM. The cell viability was observed in both, the static and
dynamic cultures (see Fig. 9.9 (A) and (B), respectively). In addition, aggregates of
chondrocytes were seen to form in both cases (static and dynamic cultures) after culturing
for 7 days, as exemplarily depicted by white arrows in Fig. 9.9. These observations confirm
the biocompatibility of the scaffolds. The two seeding conditions did not provide a significant
difference in terms of cell viability.
Figure 9. 9 The viability of primary porcine chondrocytes on Alg-foams after 7 days of
culture: (A) static culturing and (B) dynamic culturing (on a rotatory device, 12 rpm at 15 C).
Viable cells are green, dead cells appear red.
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(ii) Histological evaluation
The influence of the culturing systems on the production of cartilaginous ECM was
investigated. The results are shown in Fig. 9.10, where the histological evaluation of HE and
AB is depicted. Formation of cell clusters was observed histologically in both cases, static
and dynamic culture conditions (see Fig. 9.10 (A-D)). The results confirm that chondrocytes
still retained a round phenotype over 7 days of culture. In addition, the ECM was formed,
which surrounds the chondrocytes (as depicted by the arrows in Fig. 9.10 (B)), indicating the
deposition of cartilage PGs after 7 days of culture.
Figure 9. 10 Histological evaluation of HE and AB staining of porcine chondrocytes-seeded
Alg-foams after 7 days of static and dynamic culturing. Using AB staining, the foam stains
unspecifically blue. Cell nuclei are violet, sulfated GAGs within the cell clusters are
characterized by a faint blue staining.
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(iii) Collagen immunolabeling
The production and deposition of cartilage-specific Col II by chondrocyte-seeded on
Alg-foams was confirmed by the immunohistological evaluation (Fig. 9.11 (B and E)). In both
cases, static and dynamic culture conditions, the cell clusters produced Col II (green-
stained), which is a marker of cartilaginous ECM. Chondrocytes produced a minimum
amount of Col I under both culture conditions (Fig. 9.11 (C and F)).
In summary, the culture conditions (static and dynamic) did not provide significant
differences in terms of cell viability and the production of cartilaginous ECM. The
biocompatibility of the Alg-foams was confirmed by these results. In addition, the Alg-foams
induced the aggregation of chondrocytes and promoted the production of cartilaginous ECM
in both static and dynamic culture conditions. Hence, the results confirmed that given
porosity and pore size of the present Alg-foams exhibited a proper microenvironment in
order to support chondrocyte infiltration and the growth of chondrocytes, and provided an
attractive matrix for cartilage regeneration.
Figure 9. 11 Col II and Col I immunolabeling of porcine chondrocytes-seeded Alg-foams
after culturing for 7 days, observed by confocal microscopy. Cell nuclei are stained in blue,
Col is stained in green color.
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9.3.4 The influence of ChS molecules on chondrogenic differentiation of
chondrocytes and MSCs
(i) Cell viability
Since the different culture conditions (i.e. static and dynamic) did not lead to
differences in the growth of chondrocytes, MSCs-seeded scaffolds were cultured under
static conditions. In a first set of experiments, MSCs were simply cultured under non-
chondrogenic conditions for 7 days. By using the live-dead cell assay (see Fig. 9.12), it was
found that mostly viable cells were presented in both, Alg/ChS- and Alg-foams, after 7 days
of culture. MSCs were seen to form clusters (as depicted by the arrows in Fig. 9.12), as
observed in the case of porcine chondrocytes, which resemble pre-cartilage condensation of
MSCs [359]. The cell clusters showed a homogeneous distribution in the entire scaffold (Fig.
9.12). These results confirm the biocompatibility of Alg/ChS-foams. In addition, the highly
porous foams with pore size in the range of 50 - 300 µm support the formation of cell
clusters. However, the presence of ChS did not show any significant improvement of cell
adhesion onto Alg-foams, as originally expected.
Figure 9. 12 Cell viability (FDA/PI live-dead assay) of MSCs-seeded (A, C) Alg- and (B, D)
Alg/ChS-foams after 7 days of culture. Viable cells are green, dead cells appear red.
179
In the case of chondrocytes culturing, mostly viable cells were obvious after 14 days
in culture. In particular, it is likely that cell adhesion was improved by the incorporation of
ChS into Alg-foam after culturing with chondrocytes for 14 days. Fig. 9.13 shows the viability
of chondrocytes seeded on the foams (as indicated by green color). After 14 days in culture,
Alg/ChS-foams were superior in supporting chondrocyte adhesion compared to Alg-foams.
In addition, chondrocytes were seen to spread on the Alg/ChS-foam (Fig. 9.13 (B)), while
chondrocytes seeded on Alg-foam tended to form clusters inside the pores of the foam after
7 days (as mentioned previously in Fig. 9.9 (A)) and the clusters were maintained even after
14 days in culture (Fig. 9.13 (A)). Therefore, this qualitative observation indicates that the
incorporation of ChS into Alg-foam can enhance the adhesion of chondrocytes after culturing
for 14 days.
Figure 9. 13 Cell viability (FDA/PI live-dead assay) of porcine chondrocytes-seeded (A) Alg-
and (B) Alg/ChS-foams after 14 days of culture. Viable cells are green, dead cells appear
red.
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(ii) Histological evaluation
The histological evaluation (Fig. 9.14 (A and B)) shows the formation and deposition
of abundant ECM in both, the Alg/ChS- and Alg-foams, after 7 days of culture. The ECM
was presented in the cell clusters. In addition, the cartilage-specific sulfated PGs containing
ECM was detected by AB assay as a faint blue staining incorporated in the cell clusters (Fig.
9.14 (C and D)). These results confirm the high potential of chondrogenic differentiation of
MSCs in both scaffolds. Moreover, many MSCs acquired a mostly rounded phenotype after
7 days in culture, indicating that both scaffolds might allow further chondrogenic
differentiation of MSCs. Cells revealed no features of hypertrophy such as enlarged lacunes
and bullous cells.
Figure 9. 14 Histological evaluation (HE and AB staining) of MSCs-seeded Alg- and
Alg/ChS-foams after 7 days of culture (without chondrogenic induction). For alcian blue
staining, sulfated cartilage PGs are stained in blue, cell nuclei are stained in red and the
foams appear blue.
181
(iii) Collagen immunolabeling
MSCs-seeded Alg/ChS- and Alg-foams revealed clearly Col II immumolabeling after
7 days of culture (Fig. 9.15 (A and C)), in comparison to the negative control (Fig. 9.15 (A)).
It is obvious that Col II, which is indicated by green staining, was surrounding the cell nuclei
(stained in blue). This phenomenon is proposed as the onset of chondrogenesis of MSCs
[359]. In particular, the culturing of MSCs on Alg/ChS-foams produces an intensive Col II
(see Fig. 9.15 (C) and Fig. 9.16 (B)). However, Col I could also be detected in both,
Alg/ChS- and Alg- foams (see Fig. 9.15 (E and F)) and supporting result in Fig. 9.16 (C and
D)). A comparable result was observed in the case of chondrocytes, cultured on Alg-foams.
It is likely that the incorporation of ChS plays an important role in the chondrogenic
differentiation of MSCs, which proceeds by the aggregation of MSCs. According to the
literature, the incorporation of ChS into either porous scaffolds or hydrogels enhances both,
the chondrogenic gene expression and cartilage-specific matrix production [359,364,380]. In
addition, ChS has been found to down-regulate the expression of Col X [359,364,380]. In
order to prove this hypothesis, a longer experiment (i.e. 14 days of culture) is desired and
quantitatively biochemical assays are required to confirm the influence of ChS on the
production of Col II and PGs.
182
Figure 9. 15 Col II and Col I immunolabeling of MSCs-seeded Alg- and Alg/ChS-foams after
7 days of culture, observed by fluorescence microscopy. Col is stained in green and cell
nuclei are stained in blue.
Figure 9. 16 Col II and Col I immunolabeling of MSCs-seeded Alg- and Alg/ChS-foams after
7 days of culture, observed by confocal microscopy. Collagen is stained in green and cell
nuclei are stained in blue.
183
9.3.5 The influence of chondrogenic induction (TGF-1) on the activity of MSCs
MSCs-seeded Alg- and Alg/ChS-foams were cultured with the induction of TGF-1
growth factor in order to evaluate its effect on chondrogenic differentiation of MSCs.
According to the literature, the TGF- family plays an important role in cartilage regeneration
[246,381,382]. In vitro studies have been demonstrated that TGF-1, which serves as a
chondroinductive factor, induced early Col II expression and PGs accumulation during
chondrogenesis [246,381].
(i) Cell viability
The spreading cells were particularly detected on Alg/ChS-foams after 17 days of
non-induced culture (Fig. 9.17 (B)), while cell clusters were observed on Alg-foam (Fig. 9.17
(A)). Viable cells were mostly visible on both, Alg- and Alg/ChS-foams, appearing in green
color after 14 days of culture with the presence of TGF-1 (Fig. 9.17 (C and D)). In addition,
cell aggregation was detected in higher density with the presence of TGF-1, particularly on
Alg/ChS-foam (Fig. 9.17 (D)), compared with the absence of TGF-1.
Figure 9. 17 Cell viability (FDA/PI live-dead assay) of MSCs-seeded Alg- and Alg/ChS-
foams after culture with and without chondrogenic induction (+TGF-1). Viable cells are
green, dead cells appear red.
184
(ii) Collagen immunolabelling
With the presence of TGF-1, a superior Col II (green color) surrounding cell nuclei
(blue color) was detected in both MSCs-seeded Alg- and Alg/ChS-foams (Fig. 9.18 (C and
D), respectively), compared to the absence of TGF-1 (Fig. 9.18 (A and B)). It is likely that
the induction of TGF-1 to the culture medium stimulated the chondrogenic differentiation of
MSCs, leading to the production of specific Col II in both Alg- and Alg/ChS-foams.
Figure 9. 18 Col II and Col I immunolabelling of MSCs-seeded Alg- and Alg/ChS-foams after
14 days of culture with and without chondrogenic induction, observed by confocal
microscopy. Col II is stained in green and cell nuclei are stained in blue.
185
9.4 Conclusions
The Alg- and Alg/ChS-foams provided a good biocompatibility with porcine
chondrocytes and MSCs. The foams fabricated in the present work are confirmed to be a
proper environment for the support of both, porcine chondrocytes and MSCs. Chondrocytes
and MSCs culturing on foams produced abundant Col II and PGs after static culturing for 14
days. Some Col I was also produced. In chondrocytes cultures, Col I appeared to be lower
expressed than Col II. However, in MSCs cultures, substantial amount of Col I was detected
as unexpected in vitro. In particular, the incorporation of ChS into Alg-foams enhanced
chondrocytes and MSCs adhesion, and promoted better spreading of cells compared to Alg-
foams without incorporation of ChS. It can be concluded that the ChS, which was
incorporated into the foams, plays a significant role for the biological properties of the foams,
acting on the response of chondrocytes and MSCs. The results confirmed that an
enhancement of cell adhesion occurs with addition of ChS. In addition, the influence of ChS
on the production of cartilaginous ECM as well as on TGF-1-induced chondrogenesis of
MSCs-seeded foams must be further investigated by quantitative biochemical assays. Also,
the chondroinductive features of Alg/ChS-foams must be confirmed for applications of the
foams in cartilage regeneration and as the cartilage compartment of bilayered scaffolds.
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CHAPTER 10
Summary and Future perspectives
Multilayered scaffolds suitable for interface tissue engineering, i.e. osteochondral
regeneration, were fabricated and intensively discussed in the dissertation. Their
architecture, porous structure, physico-chemical and mechanical properties, and biological
properties were all comprehensively considered.
Bioglass®-based foams were chosen as the scaffold material for the subchondral
bone part. 3D Bioglass-based porous scaffolds exhibiting architecture and porous structure
similar to those of natural cancellous bone were obtained by using foam replication
technique. Moreover, the mechanical strength and the structural stability of the scaffolds
could be improved by coating with biodegradable polymers. Different polymer coatings,
including alginate (Alg), gelatin (Gel), PDLLA and PHBHHx coatings, were investigated
which lead to the enhancement of elastic modulus and compressive strength compared to
uncoated Bioglass-based scaffolds. In addition, such coated scaffolds maintained the
bioactivity of 45S5 Bioglass, which was evidenced by the formation of HA after immersion
in SBF solution. Therefore, all biodegradable polymer coated Bioglass®-based composite
scaffolds developed in the present study can be an appropriate candidate for use in bone
regeneration. In addition, the polymer coated Bioglass-based scaffolds can be used as a
drug/biomolecule carrier, i.e. antibiotic delivery, in bone tissue engineering applications.
Multifunctional scaffolds based on TCH-loaded polymer layers coated Bioglass-based
scaffolds exhibited improved mechanical strength compared to uncoated scaffolds as well as
controlled drug release over 14 days after immersion in PBS. The biological properties of Alg
coated Bioglass-based scaffolds only have been evaluated by culturing with osteoblasts
(MG-63) in order to confirm their biocompatibility, and ability to support cell metabolic activity
and ALP production. When compared to uncoated Bioglass-based scaffolds and RGD-
188
modified Alg coated Bioglass-based scaffolds, Alg coated scaffolds exhibited good
biocompatibility with MG-63 cells, and promoted the cell growth and osteoblastic activity.
Concerning the cartilage phase in osteochondral tissue, Alg was highly interested
because its chemical structure is similar to hyaluronic acid (HyA), considering that HyA is the
main component in cartilaginous ECM. Columnar porous structure of 3D Alg-foams was
successfully fabricated in order to support the migration and arrangement of cells, and
subsequently organization of new tissue by optimizing polymer concentrations and freeze-
drying conditions. Optimum 3 wt/v % Alg-foams exhibited pore size in the range of 125 - 325
µm, which is suitable for supporting chondrocytes seeding and migration. In addition, the
foams were able to absorb water in the same range as in native cartilage ( 80 %). The
mechanical strength and the structural stability of the foams were improved by the
application of ionic crosslinking. The elastic modulus and compressive strength of the foams
at 0.220 ± 0.009 MPa and 0.14 ± 0.02 MPa, respectively, were in the range of targeted
values of native cartilage (elastic modulus at 0.24 - 0.85 MPa and compressive strength at
0.01 - 3 MPa). Importantly, the foams were non-mineralized, meaning that they are not able
to induce bone formation in contact with body fluids, which is a requirement for cartilage
regeneration. According to the diagram summarizing the related issues in cartilage tissue
engineering approach in Fig. 10.1, the porous foams fabricated in this study have achieved
most of scaffold related criteria. However, the foams still lack cell adhesion, which negatively
affects cell proliferation and differentiation. Therefore, the Alg-foams were modified for the
first time by incorporating a biological cue, i.e. chondroitin sulfate (ChS), in order to improve
cell adhesion and cell behavior. ChS is one of natural glycosaminoglycans in cartilage,
which has a function to stimulate chondrocytes’ metabolism, inducing the synthesis of type II
collagen (Col II) and proteoglycans (PGs). Alg/ChS-foams supported either chondrocytes or
MSCs, and promoted cell proliferation and differentiation. The expression of Col II and PGs
of chondrocytes- and MSCs-seeded on Alg/ChS-foams was characterized as a marker of
cartilaginous regeneration. These results fulfilled the major cell related requirement (Fig.
10.1). In addition, Alg/ChS-foams, in which ChS serves as a biological cue, associated with
introduction of TGF-β1, significantly promoted the enhancement of chondrogenesis of
189
MSCs. This result indicates that the incorporation of biomolecule (ChS) combined with
growth factor (TGF-β1) plays an influential factor in cartilaginous differentiation and
subsequently matrix production. However, ChS slightly provided the expected enhancement
of cell adhesion. This result is probably due to the low amount of ChS incorporated into the
Alg-foam. This amount might not be sufficiently effective to be recognized by the cells.
Consequently, both chondrocytes and MSCs tended to form clusters inside the pores, but
rarely adhered on the pore wall of the foams. Therefore, some challenging issues
concerning three cornerstones of tissue engineering approach remain to be further studied.
First, it is necessary to modify the foams (i.e. surface functionalization) in order to enhance
cell adhesion. Second, the effect of ChS release rate on the cell differentiation and the
matrix production is suggested to be intensively investigated further in association with the
function of the added growth factor.
190
Figure 10. 1 Summary of major challenging issues in the area of cartilage tissue
engineering, as investigated in this dissertation, indicating the criteria that have been fulfilled
by the developed scaffolds.
The multilayered scaffolds were designed based on the optimized scaffolds for
subchondral bone and cartilage. Since the ideal scaffold for osteochondral repair does not
exist yet, the development of scaffold strategies, which provide superior long-term outcome,
are gaining more attention and receiving considerable research efforts. Therefore, recent
strategies of bi- or multi-layered scaffolds, including integrated bilayered and monolithic
biphasic scaffolds based on Alg-foam and Alg coated Bioglass® scaffold, were the focus of
this research from the materials point of view. Even though it seems likely that integrated
bilayered scaffolds may exhibit a weak point due to possible delamination at the interface
between the layers, our present study proved that the delamination can be overcome by the
incorporation of an adhesive-intermediate phase, which serves as an interface between
distinct cartilage and bone layers. In contrast, monolithic biphasic scaffolds were hardly
191
controlled in terms of fabrication parameters and subsequently porous structure in particular
at the interface. In addition, the interfacial testing by using a micro-tensile testing machine
proved that the interfacial strength at break of integrated bilayered scaffolds was higher
compared to that of monolithic biphasic scaffolds. Therefore, it can be preliminarily
summarized in the realm of materials that the integrated bilayered scaffold system
developed in this study is an appropriate approach to be further developed as an
osteochondral construct.
Additionally, fiber meshes (PLLA and Alg/Gel) were electrospun and investigated as
the cartilage phase in bilayered scaffold. The fiber meshes exhibited pore size up to 50 µm
(with up to 500 µm in thickness) and reduction of the pore size was found with increasing
thickness of the mesh. The small pore size is known to be a limitation of electrospun fibers
for use as tissue engineering scaffolds. The small pore size in fiber meshes may inhibit cell
migration and nutrient transfer after implantation. A limited 3-dimensional structure may not
be suitable for regeneration of new cartilage. Nevertheless, the fiber meshes would be an
interesting candidate for use as a calcified cartilage layer at the cartilage-bone interface.
Dense fiber meshes with small pore size can act as a barrier against the infiltration of bone
cells from subchondral bone and to avoid vascularization of the cartilage phase in vivo.
Therefore, this strategy of multilayered scaffolds for osteochondral tissue engineering
applications is promising and it should be developed further in order to achieve more closely
biomimetic structure to the structure of native tissues.
For instance, a novel multilayered scaffold with more sophisticated structure is
offered by the present research as a new perspective focusing on functional osteochondral
scaffolds, as shown in Fig. 10.2. The scaffold is created based on the present work by the
combination of Alg-foam, PLLA fibers and PDLLA-c-BG scaffold for cartilage, interface and
subchondral bone – phases, respectively. Briefly, Alg solution was applied onto PLLA fiber
mesh/PDLLA-c-BG bilayered scaffold and formed gelation by the addition of CaCl2H2O
agent. A porous Alg phase was formed after lyophilization. It was integrated on top of fiber
mesh (Fig. 10.2). In this case, the fiber mesh acted as an intermediate layer between
Bioglass®-based scaffold and Alg-foam, which was aimed as a cartilage-subchondral bone
192
interface exhibiting dense ECM. This concept was inspired by the previous study of Yunos et
al. [86] and our finding presented in Chapter 7 that the formation of HA at the interface
between Bioglass-based scaffold and PLLA fibers enhanced further the mechanical stability
of the interface. Osteoconductive hybrid fiber meshes (i.e. PLLA/Bioglass hybrid fibers) are
additionally recommended to be utilized as the interface phase instead of single polymeric
fibers. This design mimics calcified cartilage, which provides high ability for local
mineralization at the interface. Subsequently, a strong interface should be formed which is
able to integrate the cartilage and bone-like layers during in vitro cell culture and in in vivo
culture conditions.
It is expected that the scaffold developments and knowledge gained during the
present investigation will lead to advances in the field of osteochondral tissue regeneration
in the near future.
Figure 10. 2 SEM image showing a recommended multilayered scaffold – model for
osteochondral tissue engineering applications, including Alg-foam for cartilage phase, PLLA
fiber mesh for calcified interface phase and PDLLA-c-BG scaffold for bone phase.
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LIST OF FIGURES
Figure 1. 1 The scheme of knee osteoarthritis (joint degeneration disease), which is the
result of cartilage wearing out in the load bearing joint [3]. ............................. 2
Figure 2. 1 Anatomy of the knee joint, which is the most common case found in joint
degeneration disease (Source: lpch.org [21]), demonstrating arrangement of
ECM and organization of chondrocytes along different zones in cartilage
[19], and showing the structure of cancellous bone as subchondral bone
[12]. .................................................................................................................. 5
Figure 2. 2 Schematic diagram showing the gelation-mechanism of alginate and
calcium cations by the formation of egg-box structure (Courtesy K. Kashima
and M. Imai [135])..........................................................................................16
Figure 2. 3 Schematic diagram of the foam replication technique employed to produce
3D porous bioceramics- and bioactive glass-based scaffolds (according to
[28,184]). .......................................................................................................22
Figure 2. 4 (A) The schematic diagram of the freeze-drying (lyophilazation) process
showing also the phase diagram of water representing the mechanism of
freeze-drying (modified from [186,189]). .......................................................23
Figure 2. 5 Schematic diagram of the electrospinning process in horizontal direction,
which is composed of voltage supply, syringe and needle, syringe pump and
collector; the SEM image shows electrospun PLLA fibers. ...........................26
Figure 3. 1 Schematic diagram of entire tasks carried out in the dissertation thesis. .....41
Figure 4. 1 SEM images of 45S5 Bioglass®-based scaffolds fabricated by foam
replication technique: (A) 3D porous structure and (B) surface of scaffold
struts. .............................................................................................................49
Figure 4. 2 X-ray patterns of as-received Bioglass® and as-sintered 45S5 Bioglass
®-
based scaffolds. The major peaks of the phase Na2Ca2Si3O9 are marked by
■. ....................................................................................................................50
II
Figure 4. 3 SEM images of Alg coated 45S5 Bioglass®-based scaffolds with variable
concentrations, including (A, B) 1 wt/v %, (C, D) 1.5 wt/v % and (E, F) 2 wt/v
%, the scaffolds were coated once for 5 min soaking time. ..........................54
Figure 4. 4 SEM images of Gel coated 45S5 Bioglass®-based scaffolds with variable
concentrations, including (A, B) 1.5 wt/v %, (C, D) 3 wt/v %, (E, F) 5 wt/v %,
the scaffolds were coated once on for 5 min soaking time and (G, H) 1.5
wt/v % coated for 5 min and 3 dipping cycles. ..............................................55
Figure 4. 5 SEM images of PDLLA coated 45S5 Bioglass®-based scaffolds with variable
concentrations, including (A, B) 3 wt/v %, (C, D) 5 wt/v % and (E, F) 8 wt/v
%, the scaffolds were coated once for 5 min soaking time. ..........................56
Figure 4. 6 PHBHHx coated 45S5 Bioglass®-based scaffolds with variable
concentrations, soaking times and number of dipping times, including (A, B)
5 wt/v %, 5 min of soaking and one coating cycle, and (C, D) 1 wt/v %, 15
sec of soaking time and 30 coating cycles. ...................................................56
Figure 4. 7 (A) Representative compressive stress-strain curves of biodegradable
polymer coated Bioglass®-based composite scaffolds in comparison with the
curve of uncoated Bioglass®-based scaffolds, (B) normalized mechanical
properties (compressive modulus and strength) of coated scaffolds
compared to those of uncoated scaffolds (as reference) and (C) appearance
of coated scaffolds after compression load. * indicates significantly different
mechanical properties of coated scaffolds in comparison with those of
uncoated scaffolds, and # indicates significantly different mechanical
properties of synthetic polymer coated scaffolds in comparison with the
properties of natural polymer coated scaffolds (p 0.05). ............................59
Figure 4. 8 (A) % weight loss of biodegradable polymer coated 45S5 Bioglass®-based
scaffolds after 1, 3, 7, 14 and 28 days of immersion in SBF and (B) variation
of the pH of the SBF solution. .......................................................................62
III
Figure 4. 9 SEM images of as-sintered 45S5 Bioglass®-based scaffolds after (A) 1, (B)
3, (C) 14, and (D) 28 days of immersion in SBF, showing possible formation
of HCA indicated qualitatively by the morphology of the deposited structures.63
Figure 4. 10 (A) FTIR spectra of as-sintered 45S5 Bioglass®-based scaffolds after
different immersion times in SBF, in comparison with the scaffolds before
immersion. The characteristic peaks of HCA are marked by ▲ and ■ and (B)
XRD patterns of as-sintered 45S5 Bioglass®-based scaffolds after different
immersion times in SBF, in comparison with the scaffolds before immersion.
The major peak of HCA is marked by , while the crystalline peaks of
Na2Ca2Si3O9 were marked by ■. ...................................................................64
Figure 4. 11 SEM images of biodegradable polymer coated 45S5 Bioglass®-based
composite scaffolds after immersion in SBF, showing formation of HCA: (A,
B, C) Alg-c-BG, (D, E, F) Gel-c-BG, (G, H, I), PDLLA-c-BG and (J, K, L)
PHBHHx-c-BG scaffolds for 1, 3 and 28 days, respectively. ........................68
Figure 4. 12 FTIR spectra of biodegradable polymer coated 45S5 Bioglass®-based
scaffolds after 28 days of immersion in SBF: (A) Alg-c-BG, (B) Gel-c-BG, (C)
PDLLA-c-BG, and (D) PHBHHx-c-BG scaffolds. The characteristic peaks of
HCA are marked by ▲ and the characteristic peaks of polymer coatings are
marked by ■. ..................................................................................................69
Figure 5. 1 (A) Scheme of the capillarity test of Bioglass®-based scaffolds, showing the
effect of surface chemistry on the permeability of the porous scaffolds and
(B) photographs representing the coated scaffolds during the capillarity test.77
Figure 5. 2 Contact angles of Bioglass®-based scaffolds showing the surface wettability
of different coatings. * indicates the significant difference (p 0.05) of the
modified coatings on the Bioglass-based scaffolds in comparison with PL-
c-BG scaffolds. ..............................................................................................77
Figure 5. 3 SEM images of the scaffolds showing the morphological porous structure
and morphology of a coating surface of: (A, B) PL/P123-c-BG scaffolds, (C,
IV
D) T-BG scaffolds, (E, F) T-Alg-c-(PL/P123-c-BG) scaffolds and (G, H) T-
Gel-c-(PL/P123-c-BG) scaffolds. ...................................................................79
Figure 5. 4 Mechanical properties of polymer coated Bioglass scaffolds: (A)
representative compressive stress-strain curves and (B) average elastic
modulus and average compressive strength of the scaffolds. * (p 0.05)
indicates the statistical significance of compressive mechanical properties of
coated scaffolds, compared to those of uncoated T-BG scaffolds. ...............80
Figure 5. 5 FTIR spectra of TCH, BG, TCH-loaded Bioglass scaffolds and TCH-loaded
polymer coated Bioglass®-based scaffolds. ..................................................82
Figure 5. 6 (A) Drug release profile and (B) degradation behavior of TCH-loaded
polymer (Alg and Gel) coated Bioglass scaffolds. .......................................84
Figure 5. 7 SEM images of TCH-loaded polymer coated Bioglass-based scaffolds after
in vitro release in PBS for 14 days: (A) T-Alg-c-(PL/P123-c-BG); the arrows
predicting the PL/P123 coating and (B) T-Gel-c-(PL/P123-c-BG) scaffolds;
dashed arrow depicting the Bioglass® struts and solid arrow line predicting
PL/P123 coating, and SEM images of TCH-loaded polymer coated
Bioglass-based scaffolds after immersion in SBS for 14 days: (C) T-Alg-c-
(PL/P123-c-BG and (D) T-Gel-c-(PL/P123-c-BG) scaffolds; solid arrows
depicting polymer coating. .............................................................................85
Figure 5. 8 FTIR spectra of T-Alg- and T-Gel-c-(PL/P123-c-BG) scaffolds after 14 days
of immersion in PBS. .....................................................................................86
Figure 6. 1 Optical microscopy images showing 3 wt/v % Alg-gel after the introduction of
0.1 M CaCl22H2O into Alg solution: (A) superficial surface (plan-view image)
and (B) cross-section of Alg-gel. SEM images of 3 wt/v % Alg-foam after
lyophilized: (C) plan-view and (D) cross-section SEM image. ......................97
Figure 6. 2 Optical photograph showing the appearance of 3 wt/v% Alg-foams obtained
with variation of CaCl22H2O concentrations (0.1 – 1 M). .............................99
Figure 6. 3 Comparison of weight loss as a function of immersion time in DI H2O of Alg-
foams without (w/o) and with (w) crosslinking by immersion in 0.5 M
V
CaCl22H2O for 4 h (the inset represents the weight loss of Alg-foams
without crosslinking). .....................................................................................99
Figure 6. 4 The effect of Alg concentrations on the porosity and density of Alg-foams.100
Figure 6. 5 SEM images (in plan-view) of (A, B) 2 wt/v %, (D, E) 3 wt/v % and (G, H) 4
wt/v % Alg-foams, included distribution of pore size: (C) 2 wt/v %, (F) 3 wt/v
% and (I) 4 wt/v % Alg-foams. .....................................................................101
Figure 6. 6 (A) Representative compressive stress-strain curves of 2, 3 and 4 wt/v %
Alg-foams and (B) the mechanical properties, including elastic modulus and
compressive strength of the foams as a function of concentrations. ..........103
Figure 6. 7 Dynamic mechanical properties of 3 wt/v % Alg-foams in compression mode
presenting the storage modulus (E’) and the loss factor (tan ) as a function
of frequency, in both dry and wet state. ......................................................104
Figure 6. 8 DSC thermogram: (A) the 1st heating and (B) the 2
nd heating cycle runs of
Alg-foams with and without crosslinking. ....................................................105
Figure 6. 9 (A) Weight change of 3 wt/v % Alg-foams as a function of immersion time in
different media, including DI H2O, PBS and SBF. (B) ATR-FTIR spectra of
the foam after 7 days of immersion in the media. .......................................107
Figure 6. 10 Optical microscopy images of electrospun PLLA fibers obtained by using the
conditions in (A) Trial 4 and (B) Trail 13, according to electrospinning
conditions used in Table 6.1. .......................................................................110
Figure 6. 11 Optical microscopic images of electrospun Alg/Gel fibers obtained by using
the conditions in (A) Trial 9 and (B) Trail 12, according to electrospinning
conditions used in Table 6.2. .......................................................................113
Figure 6. 12 SEM images of PLLA fibers, which were deposited for (A) 2 h and (C) 9 h,
and Alg/Gel fibers, which were deposited for (B) 2 h and (D) 9 h. The
distribution of fiber diameters of both fiber types is included. .....................114
Figure 6. 13 (A) FTIR spectra of Alg/Gel fibers before and after GA crosslinking (The
inset represents the change of amide I peak before and after crosslinking)
VI
and (B) mechanism of the crosslinking reaction between Gel with GA,
according to [78]. .........................................................................................116
Figure 6. 14 XRD patterns of PLLA fibers showing the reduction of crystallinity in
comparison with PLLA cast films (■ and indicate - and β-crystalline phase
of PLLA, respectively). ................................................................................117
Figure 6. 15 DSC thermogram of PLLA fibers, indicating Tg, Tc, and Tm. .......................118
Figure 6. 16 (A) Representative tensile stress-strain curve of PLLA fibrous meshes in
comparison with the typical curve of PLLA cast films (the inset) and (B) the
values of elastic modulus, tensile strength and elongation at break of PLLA
under tension deformation mode. ................................................................120
Figure 6. 17 Summary of Young’s modulus values of electrospun fibers with respect to
fiber diameter, which were collected from recent literature reports
[238,278,301,306,337–340], mainly on polyesters, Gel and Alg. The star
indicates the position of PLLA fibers obtained in the present work. ............121
Figure 7. 1 The schematic diagram of the four types of multilayered scaffolds for
osteochondral tissue engineering developed in this project. ......................125
Figure 7. 2 Optical microscopic images showed the appearance of three different
approaches of multilayered scaffolds, including (A, a) system A: monolithic
Alg/Alg-c-BG biphasic scaffold, (B, b) system B: integrated Alg/Alg-c-BG
bilayered scaffold and (C, c) system C: integrated electrospun PLLA
fibers/PDLLA-c-BG bilayered scaffold (the inset shows a plan-view of fibers
integrated on the struts of the Bioglass-based scaffold). ..........................130
Figure 7. 3 SEM images showing cross-sections of four different types of multilayered
scaffolds: (A, a) system A, (B, b) system B, (C, c) system C and (D, d)
system D (the dashed line marks the interface between the two phases). .131
Figure 7. 4 (A) Representative stress-strain curves of the bilayered scaffolds (System A
vs. System B) and (B) distribution of the strength at break values of the
scaffolds in systems A and B (the red dashed line is included for the visual
aid). ..............................................................................................................133
VII
Figure 7. 5 Representative compressive stress-strain curve of integrated bilayered
scaffold (system B) in comparison with the curves of Alg-foam and Alg-c-BG
scaffold. .......................................................................................................134
Figure 7. 6 SEM images of integrated bilayered scaffold (system B) after immersion in
SBF for 1 and 28 days (the dashed line indicates the interface between the
two phases). ................................................................................................136
Figure 7. 7 ATR-FTIR spectra of Alg-foam after immersion in SBF for 28 days in order
to confirm non-mineralization of the foam (the inset shows the absorption
bands in the wavenumber 1200 - 850 cm-1
). ...............................................137
Figure 7. 8 SEM images of integrated bilayered scaffold (system C) after immersion in
SBF for 1 and 28 days. ................................................................................138
Figure 7. 9 SEM image of integrated bilayered scaffold (system D) after immersion in
SBF for 28 days (the dashed line indicates interface between the Alg/Gel
fiber mesh and the Alg-c-BG scaffold and the dashed circles indicate the
interconnection between the fibers formed after immersion in SBF). .........139
Figure 8. 1 Relative LDH activity of MG-63 osteoblast-like cells cultured on uncoated
BG, Alg-c-BG and RGD-Alg-c-BG scaffolds. The results presenting the
difference of optical densities are presented as mean ± SD (n = 4). * (p
0.05) indicates significant difference of different scaffolds at different culture
times. ...........................................................................................................147
Figure 8. 2 Fluorescent microscopic images of MG-63 osteoblast-like cells-seeded BG,
Alg-c-BG and RGD-Alg-c-BG scaffolds after 3 and 14 days in culture by
using DAPI stain for cell nuclei. ...................................................................148
Figure 8. 3 Cell metabolic activity of MG-63 osteoblast-like cells cultured on uncoated
BG, Alg-c-BG and RGD-Alg-c-BG scaffolds. The results in % AB reduction
are presented as mean ± SD (n = 6). * (p 0.05) indicates significant
difference of different scaffold types at different culture times. ...................150
Figure 8. 4 ALP activity up to 21 days of MG-63 osteoblast-like cells cultured on
uncoated BG, Alg-c-BG and RGD-Alg-c-BG scaffolds. The results are
VIII
reported as mean ± SD (n = 4). * p 0.05 indicates significant difference of
results for different scaffold types at different culture times. .......................152
Figure 8. 5 SEM cross-sectioned images showing MG-63 cells-seeded BG, Alg-c-BG
and RGD-Alg-c-BG scaffolds after cultured for 14 days. ............................153
Figure 8. 6 Confocal microscopic images of MG-63 cells cultured on BG, Alg-c-BG and
RGD-Alg-c-BG scaffolds, stained with OsteoImage (green), after 3,14 and
21 days. .......................................................................................................154
Figure 9. 1 SEM images of Alg/ChS-foams in (A) plan-view and (B) cross-section; (C)
shows a histogram of the pore size distribution of Alg/ChS-foams and (D) is
a scheme showing the reptation model in the case of Alg/ChS blend,
according to [367,368]. ................................................................................167
Figure 9. 2 (A) ATR-FTIR spectrum and (B) EDX spectra of the Alg/ChS-foam in
comparison to the spectra of pure Alg- and pure ChS-foams, which both
results confirm the existence of ChS in the foam. .......................................168
Figure 9. 3 DSC thermograms: (A) the 1st heating and (B) the 2
nd heating cycle of
Alg/ChS- and Alg-foams. .............................................................................169
Figure 9. 4 Water absorption in % of Alg- and Alg/ChS-foams, in comparison to the
water absorption of natural articular cartilage (*), as referenced in the
literature [6]. ................................................................................................170
Figure 9. 5 Degradation profile of Alg/ChS- and Alg-foams in PBS, evaluated by %
weight change with respect to the immersion time. ....................................171
Figure 9. 6 (A) Representative compressive stress-strain curves of the Alg/ChS- and
Alg-foam; (B) mechanical properties of Alg/ChS and Alg-foams. ...............173
Figure 9. 7 Dynamic mechanical analysis of Alg/ChS- and Alg-foams: storage modulus
(E‘) and tan as a function of frequency; note the logarithmic scaling. ......173
Figure 9. 8 Cumulative release of ChS from Alg/ChS-foams immersed in the PBS
solution (pH 7.4, 37 °C) measured by using carbazole reaction, as
described in Section 9.2.3. ..........................................................................174
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Figure 9. 9 The viability of primary porcine chondrocytes on Alg-foams after 7 days of
culture: (A) static culturing and (B) dynamic culturing (on a rotatory device,
12 rpm at 15 C). Viable cells are green, dead cells appear red. ...............175
Figure 9. 10 Histological evaluation of HE and AB staining of porcine chondrocytes-
seeded Alg-foams after 7 days of static and dynamic culturing. Using AB
staining, the foam stains unspecifically blue. Cell nuclei are violet, sulfated
GAGs within the cell clusters are characterized by a faint blue staining.....176
Figure 9. 11 Col II and Col I immunolabeling of porcine chondrocytes-seeded Alg-foams
after culturing for 7 days, observed by confocal microscopy. Cell nuclei are
stained in blue, Col is stained in green color. ..............................................177
Figure 9. 12 Cell viability (FDA/PI live-dead assay) of MSCs-seeded (A, C) Alg- and (B,
D) Alg/ChS-foams after 7 days of culture. Viable cells are green, dead cells
appear red. ..................................................................................................178
Figure 9. 13 Cell viability (FDA/PI live-dead assay) of porcine chondrocytes-seeded (A)
Alg- and (B) Alg/ChS-foams after 14 days of culture. Viable cells are green,
dead cells appear red. .................................................................................179
Figure 9. 14 Histological evaluation (HE and AB staining) of MSCs-seeded Alg- and
Alg/ChS-foams after 7 days of culture (without chondrogenic induction). For
alcian blue staining, sulfated cartilage PGs are stained in blue, cell nuclei
are stained in red and the foams appear blue. ............................................180
Figure 9. 15 Col II and Col I immunolabeling of MSCs-seeded Alg- and Alg/ChS-foams
after 7 days of culture, observed by fluorescence microscopy. Col is stained
in green and cell nuclei are stained in blue. ................................................182
Figure 9. 16 Col II and Col I immunolabeling of MSCs-seeded Alg- and Alg/ChS-foams
after 7 days of culture, observed by confocal microscopy. Collagen is
stained in green and cell nuclei are stained in blue. ...................................182
Figure 9. 17 Cell viability (FDA/PI live-dead assay) of MSCs-seeded Alg- and Alg/ChS-
foams after culture with and without chondrogenic induction (+TGF-1).
Viable cells are green, dead cells appear red. ............................................183
X
Figure 9. 18 Col II and Col I immunolabelling of MSCs-seeded Alg- and Alg/ChS-foams
after 14 days of culture with and without chondrogenic induction, observed
by confocal microscopy. Col II is stained in green and cell nuclei are stained
in blue. .........................................................................................................184
Figure 10. 1 Summary of major challenging issues in the area of cartilage tissue
engineering, as investigated in this dissertation, indicating the criteria that
have been fulfilled by the developed scaffolds. ...........................................190
Figure 10. 2 SEM image showing a recommended multilayered scaffold – model for
osteochondral tissue engineering applications, including Alg-foam for
cartilage phase, PLLA fiber mesh for calcified interface phase and PDLLA-c-
BG scaffold for bone phase. ........................................................................192
LIST OF TABLES
Table 2. 1 Mechanisms of bioactivity and bone bonding of Bioglass®, according to
[29,31,34,39,46]............................................................................................... 9
Table 2. 2 Mechanical properties of natural healthy human osteochondral tissues
[6,12,73,76,141–143]. ...................................................................................17
Table 2. 3 Current 3D scaffold fabrication techniques for polymers and ceramics. .......19
Table 2. 4 Summary of current strategies in osteochondral tissue engineering. ...........28
Table 4. 1 Polymer coating conditions for polymer coated 45S5 Bioglass®-based
scaffolds. An as-sintered rectangular shaped 45S5 Bioglass®-based
scaffold, with the dimensions of 8 mm × 8 mm × 8 mm, was soaked in 5 ml
of each polymer solution. ..............................................................................46
Table 4. 2 Optimized polymer coating conditions for different biodegradable polymer
coated 45S5 Bioglass®-based scaffolds. .......................................................57
Table 6. 1 Electrospinning conditions and primary observations of PLLA fibers (
indicates no fiber formation, indicates beads incorporated into fibers and
indicates uniform fibers). .............................................................................109
Table 6. 2 Electrospinning conditions and primary observations of alginate fibers (
indicates no fiber, indicates beads and indicates uniform fibers). ........112
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ABBREVIATIONS AND SYMBOLS
Abbreviations:
3D Three-dimentional
45S5 BG 45S5 Bioglass®
5S5 BG 5S5 Bioglass®
AB assay Alamar Blue assay
AB Alcian blue
ACI Autologous chondrocytes implantation
ACDC5 Chondrocyte-like cells
ADSCs Adipose-derived stem cells
Alg Alginate
ALP Alkaline phosphatate
ATR-FTIR Attenuated total reflectance-Fourier transform infrared spectroscopy
BCP Biphasic calcium phosphate
BMP-2, -7 Bone morphogenetic protein-2, -7
BMSCs Bone marrow mesenchymal stem cells
Ca-Alg Calcium alginate
CaP Calcium phosphate
Col Collagen
Col I, II, X Type I, II, X collagen
ChS Chondroitin sulfate
CS Chitosan
DAPI 4,6-diamidino-2-phenylindole
DCM Dichloromethane
DMA Dynamic mechanical analysis
DMC Dimethyl carbonate
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DMEM Dulbecco’s modified eagle medium
DSC Differential scanning calorimetry
E´ Storage modulus
E´´ Loss modulus
ECM Extracellular matrix
ESCs Embryonic stem cells
EtOH Ethanol
F127 Pluronic F127 (Blockcopolymer of poly(ethylene glycol))
FDA Fluorescein diacetate
FGF-1 Fibroblast growth factor-1
Fig. Figure
FTIR Fourier transform infrared spectroscopy
GA Glutaraldehyde
GAGs Glycosaminoglycans
Gel Gelatin
HA Hydroxyapatite
HB Hydroxy butyrate
HCA Hydroxycarbonate apatite
HE Hematoxylin eosin
hMSCs Human mesenchymal stem cells
HH Hydroxyl hexanoate
HOBs Human osteoblasts
HOS-TE85 Human osteosarcoma cell line
HyA Hyaluronic acid, Hyaluronan
IGF-1, -2 Insulin-like growth factor-1,-2
LDH Lactate dehydrogenase
MC3T3-E1 Mouse osteoblast cell line
MeOH Methanol
MG-63 Human osteoblast cell line
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MSCs Mesenchymal stem cells
Mw Molecular weight
nHA Nanohydroxyapatite
Na-Alg Sodium alginate
OA Osteoarthritis
OD Optical density
P Porosity
P123 Poly(ethylene glycol)-c-poly(propylene glycol)-c-poly(ethylele
glycol)-triblock copolymer
PA6 Polyamide 6
PBS Phosphate buffer saline
PCL Polycaprolactone
PEG Poly(ethylene glycol)
PEO Poly(ethylene oxide)
PGs Proteoglycans
PGA Poly(glycolic acid)
PHA Poly(hydroxyalcanoate)
PHB Poly(3-hydroxybutyrate)
PHBHHx Poly(3-hydroxybutyrate-c-3-hydroxyhexanoate)
PHBV Poly(3-hydroxybutyrate-c-valerate)
PLA Poly(lactic acid)
PLLA Poly(L-lactic acid)
PDLA Poly(D-lactide)
PDLLA Poly(D,L-lactide)
PI Propidium iodine
PLGA Poly(lactic-c-glycolic acid)
p-NP Para-nitrophenol
p-NPP Para-nitrophenylphosphate
PolyHEMA Poly(2-hydroxyethyl methacrylate)
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PVA Poly(vinyl alcohol)
PU Polyurethane
SB Subchondral bone
SEM Scanning electron microscopy
SBF Simulated body fluids
SD Standard deviations
SOX-9 Protein transcription factor, acting during chondrocyte differentiation
TC Cold crystalline temperature
Tg Glass transition temperature
Tm Melting temperature
Tan Loss factor
TBS Tris buffered saline
TCH Tetracycline hydrochloride
TCP Tricalcium phosphate
TGF-β1, 3 Transforming growth factor- beta 1, 3
Xc Degree of crosslinking
XRD X-ray Diffraction
Chemical symbols:
-CH2- Ethyl group
-CH3 Methyl group
-C=O Carbonyl group
-C-O-C- Ether bond
-COO- Carboxylate group
-COOH Carboxyl group
-OH Hydroxyl group
-C=N- Imine bond
-N-H- Amine bond
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Al2O3 Aluminium oxide
Ca2+
Calcium ion
CaCO3 Calcium carbonate
CaCl22H2O Calcium chloride dehydrate
CaO Calcium oxide
Ca10(PO4)6(OH)2 Hydroxyapatite
CO2 Carbon dioxide
CO32-
Carbonate ion
Cu2+
Copper ion
DI H2O Deionized water
Fe2+
Iron ion
H+ Hydrogen ion
HCl Hydrochloric acid
H2SO4 Sulfuric acid
H3O+ Hydronium ion
K+ Potassium ion
KBr Potassium bromide
MgCl2 Magnesium dichloride
Na+ Sodium ion
Na2CO3 Sodium carbonate
NaCl Sodium chloride
Na2O Disodium oxide
NaOH Sodium hydroxide
OH- Hydroxide ion
P2O5 Phosphorous pentoxide
PO43-
Phosphate ion
S Sulfur
Si Silicon
SiO2 Silicon dioxide
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Si-OH Silanol group
Si-O-Si Silica gel
Si(OH)4 Silicic acid
Sr2+
Strontium ion
Zn2+
Zinc ion
Symbols:
% Percent
wt % Weight in percentage
vol % Volume in percentage
wt/v % Weight per volume in percentage
v/v % Volume per volume in percentage
° Degree
°C Degree Celsius
°C/min Degree Celsius per minute
Density
g Gram
g/mole Grams per mole
g/cm3 Grams per cubic centimeter
ng/ml Nano-grams per milliliter
µg/ml Micro-grams per milliliter
M Molar
mM Millimolar
N Normal
sec Second
min Minute
h Hour
µl Microliter
XIX
ml Milliliter
l Liter
ml/h Milliliters per hour
mm/min Millimeters per minute
cm Centimeter
cm-1
Reciprocal centimeter
µm Micrometer
nm Nanometer
N Newton
kN Kilo Newton
MPa Mega Pascal
kV Kilo Volt
P Porosity
ppi Pores per inch
rpm Rounds per minute
J/g Joules per gram
Hz Hertz
U/mg Units per milligram
XX
ACKNOWLEDGEMENTS
I would like to thank my funding from the Royal Thai Government Scholarship
granted by the Office of the Civil Service Commission (OCSC), Bangkok, Thailand. Also
thanks to Faculty of Agricultural Product Innovation and Technology, Srinakharinwirot
University, Bangkok, Thailand for providing me a next future journey in the research area.
I also would like to thank all the people who help me along the journey to carry out
this work. My particular gratitude is dedicated to:
First of all, Prof. Dr.-Ing. habil. Aldo R. Boccaccini, who gave me the great
opportunity to work at the Institute of Biomaterials, Department of Materials Science and
Engineering, University of Erlangen-Nuremberg. Thanks for giving me the chance to work
with variety of researches in the field of tissue engineering and in collaboration with many
research groups. I learned a lot of knowledge and achieved plenty of experiences.
Thanks to all collaborations for the kind exchange of knowledge and scientific
discussion, including:
- Prof. Dr. rer. nat. habil. Dirk W. Schubert (Institute of Polymer Materials,
Department of Materials Science and Engineering, University of Erlangen-Nuremberg) for
the close collaboration and for the opportunity to use the facilities during my project.
- PD Dr. med.-vet. Gundula Schulze-Tanzil and her co-workers in Charité-
Universitätsmedizin Berlin, Campus Benjamin Franklin Klinik für Orthopädische for very nice
collaboration and scientific discussion in the field of cell biology for cartilage tissue
engineering. I also would like to thank for her dedication in revising a part of this dissertation.
- Prof. Dr. Ulrich Lohbauer and Dr. Andrea Wagner in the Dental Clinic – Operative
Dentistry and Periodontology, University of Erlangen-Nuremberg, for giving me the
opportunity to use microtensile testing.
- Prof. Dr. Dirk Höfer, Dr. Timo Hammer, and Marina Müller from Hehenstein
Institutes, Institute for Hygiene and Biotechnology, Boennigheim, Germany, who
XXII
collaborated with us on the topic of adipose-derived stem cells culturing on Bioglass-based
scaffolds for vascularization.
- Dr. Subha Narayan Rath (Nikolaus Fiebiger Zentrum for molecular medicine,
University of Erlangen-Nuremberg), and Prof. Dr. med. Ulrich Kneser and his co-workers
(Department of Plastic and Hand Surgery, University Hospital of Erlangen, University of
Erlangen-Nuremberg), who are the close collaboration in the field of osteogenic
differentiation of MSCs on Bioglass-based scaffolds. Special thanks to Dr. Subha Narayan
Rath for kindly discussion about cell biology.
- Both co-authors in my first review paper about osteochondral tissue engineering,
PD Dr. med. Justus Beier (Department of Plastic and Hand Surgery, University Hospital of
Erlangen, University of Erlangen-Nuremberg) and Prof. Dr. Vehid Salih (Eastman Dental
Institute, UCL, London, United Kingdom), who kindly advised and supported during the
process.
- Dr. Beatriz Olalde and co-workers (TECNALIA Health Division and Ciber-BNN,
San Sebastián, Spain) for the fruitful collaboration in the field of drug delivery applications.
Numerous thanks to Dr. Judith A. Roether for her friendly scientific advices and
spiritual support. I am appreciate for her sincerely helps since the first day of my staying in
Erlangen, Germany.
I would like to thank to Dr. Rainer Detsch and Ms. Alina Grünewald for their kindly
help to carry on the bone cell culture on my Bioglass-based scaffolds. I am thankful for their
big attempt. My work would not have been possible without their assistance.
Thanks to Dr. Ing. Joachim Kaschta (Institute of Polymeric Materials, Department of
Materials Science and Engineering, University of Erlangen-Nuremberg) and Dr. Raquel Silva
Lourenço (Institute of Biomaterials) for dynamic mechanical analysis - training. Also thanks
to Dr. Menti Goudouri for the analysis of FTIR spectroscopy and for kindly advising a part of
my thesis.
Special thanks to Dr. Ranjana Rai, who is the first friend and one of my best friends
during my staying in Germany. She has been like my sister, who is always ready to help me.
I would not have passed hard times of my life without her friendship.
XXIII
I also would like to thank to my international friends, Wei Li, Yaping Ding, Qiang
Chen, Qingqing Yao, Bapi Sarker, Preethi Balasubramanian, Fereshteh Zeinab and Kai
Zheng. They are lovely friends who never let me feel lonely during my study abroad. Also
thanks for useful scientific discussion and helping several experimental works.
In addition, I would like to thank to my office mates, Anahi Phillippart, Valentina
Miguez Pacheco, Marwa Tallawi, Dr. Sandra Cabañas Polo, Dr. Judith Bortuzzo and Samira
Tansaz, and also Rama Krishna Chinnam and Luis Cordero for their friendship and their
kindly help in many issues.
Special appreciation and respect are dedicated to Heinz Mahler and Bärbel Wust,
who have helped me a lot concerning my livelihood.
Also thanks to other Biomat-members, Sigrid Seuß, Jasmin Hum and Alexander
Hoppe for experimental assistance. All of my bachelor, master and mini-project students,
Alexander Ritter, Birgitta Carlé, Mani Diba, Anke-Lisa Metze, and Eva Weber, who have
usually brought up fresh ideas and have been willing to learn and to achieve new
experiences together with me.
Especially, I would like to thank to my best friends in Thailand, Wanruedee Temnil,
Patchamon Duangnum, Sudarat Jindabot, Worratai Panichnitinon, Pimpaporn Paebamrung,
Nattaporn Aimampaiwong and Piyachat Chuysrinual, who beside me in all sufferings and
happiness.
Very special thanks are dedicated to Dr. Ing. Mirza Mackovic for his best support.
Thanks for always expecting the best for me, and supporting me in every life situation.
My particular gratitude is dedicated to my parents, Mr. Prasop and Mrs. Theeranuch
Nooeaid. Thanks for their unconditional love, care and support, and thanks for growing me
with gentleness and modesty. They inspire me to keep reaching for my goals and they are
my power behind my every success. Also, I would like to thank to Nooeaid and Suwanarat
families for their care and support.
XXIV
LIST OF PUBLICATIONS
(1) Nooeaid P., Roether J.A., Weber E., Schubert D.W., Boccaccini A.R. Technologies for
multilayered scaffolds suitable for interface tissue engineering. Adv. Eng. Mater.
2013, DOI: 10.1002/adem.201300072.
(2) Nooeaid P., Salih V., Beier J. P., Boccaccini A. R. Osteochondral tissue engineering:
scaffolds, stem cells and applications. J. Cell. Mol. Med. 2012, 16, 2247-70.
(3) Nooeaid P., Roether J.A., Schubert D.W., Boccaccini A.R. Polymer coated bioactive
glass foams: toughened scaffolds for bone tissue engineering. Biofoams 2011, 21-4.
(4) Nooeaid P., Schulze-Tanzil G., Boccaccini A.R. Stratified scaffolds for osteochondral
tissue engineering. A book chapter in Methods in Molecular Biology. Submitted.
(5) Li W., Nooeaid P., Roether J.A., Schubert D.W., Boccaccini A.R. Preparation and
characterization of vancomycin releasing PHBV coated 45S5 Bioglass-based glass-
ceramic scaffolds for bone tissue engineering. J Eur Cer Society 34, 2014, 505-14.
(6) Subha R., Nooeaid P., Arkudas A., Beier J.P., Strobel L.A., Brandl A., Horch R.E.,
Boccaccini A.R., Kneser U. Adipose-Derived and Bone Marrow-Derived Mesenchymal
Stem Cells Display Different Osteogenic Differentiation Patterns in 3D Bio-active
Bioglass Scaffolds. J Tissue Eng Regen Med 2013, DOI: 10.1002/term.
(7) Yao Q., Nooeaid P., Detsch R., Roether J.A., Dong Y., Goudouri M.O., Schubert
D.,Boccaccini A.R. Bioglass/chitosan-polycaprolactone bilayered composite scaffolds
intended for osteochondral tissue engineering. J. Biomed. Mater. Res. Part A 2014,
In-press.
(8) Yao Q., Nooeaid P., Roether J.A., Dong Y., Zhang Q., Boccaccini A.R. Bioglass®-
based scaffolds incorporating polycaprolactone and chitosan coatings for controlled
vancomycin delivery. Ceramic Inter. 2013, In-press.
(9) Hum J., Luczynski K.W., Nooeaid P., Newby P., Lahayne O., Hellmich C., Boccaccini
A.R. Stiffness improvement of 45S5 Bioglass®-based scaffolds through natural and
synthetic biopolymer coatings: an ultrasonic study. Strain 2013, 49, 431-9.
XXVI
(10) Liverani L., Roether J.A., Nooeaid P., Trombetta M., Schubert D.W., Boccaccini A.R.
Simple fabrication technique for multilayered stratified composite scaffolds suitable for
interface tissue engineering. Mater. Sci. Eng. A 2012, 557, 54–8.
(11) Metze A., Grimm A., Nooeaid P., Roether J.A., Hum J., Newby P.J., Schubert D.W.,
Boccaccini A.R. Gelatin coated 45S5 Bioglass®-derived scaffolds for bone tissue
engineering. Key Eng. Mater. Vol. 2013, 541, 31-9.
(12) Handel M., Hammer T.R., Nooeaid P., Boccaccini A.R., Hoefer D. 45S5 Bioglass-
based 3D-scaffolds seeded with human adipose tissue-derived stem cells (hASC)
induce in vivo vascularization in the CAM angiogenesis assay. Tissue Eng Part A.
2013, 19, 2703-12.
(13) Olalde B., Garmendia N., Sáez-Martínez V., Argarate N., Nooeaid P., F. Morin,
Boccaccini A.R. Multifunctional bioactive glass scaffolds coated with layers of
poly(D,L-lactide-co-glycolide) and poly(n-isopropylacrylamide-c-acrylic acid) microgels
loaded with vancomycin. Mater. Sci. Eng. C Vol. 2013, 33, 3760-7.
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