Biodegradable polymers are increasingly used in biomedical applications. When implanted into a body, biodegradable material does not require removal surgery but will disappear through polymer degradation. In tissue engineering, porous biodegradable scaffolds are used to grow tissue structures that restore or improve natural tissue functions. Additive manufacturing techniques provide a feasible method to 3D fabricate patient-specific scaffolds that closely satisfy the requirements of their target tissue. In this dissertation, new biodegradable photocrosslinkable polymers were synthesized and used for stereolithography-based 3D fabrication of tissue engineering constructs. The development of new biodegradable polymers with great printing properties is a crucial step toward the successful use of additive manufacturing in more effective, personalized medicine.

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ISBN 978-952-60-6491-8 (printed) ISBN 978-952-60-6492-5 (pdf) ISSN-L 1799-4934 ISSN 1799-4934 (printed) ISSN 1799-4942 (pdf) Aalto University School of Chemical Technology Department of Biotechnology and Chemical Technology www.aalto.fi

BUSINESS + ECONOMY ART + DESIGN + ARCHITECTURE SCIENCE + TECHNOLOGY CROSSOVER DOCTORAL DISSERTATIONS

Laura E

lomaa

Synthesis of biodegradable photocrosslinkable polymers for stereolithography-based 3D

fabrication of tissue engineering scaffolds and hydrogels

Aalto

Unive

rsity

2015

Department of Biotechnology and Chemical Technology

Synthesis of biodegradable photocrosslinkable poly- mers for stereolithography-based 3D fabrication of tissue engineering scaffolds and hydrogels

Laura Elomaa

DOCTORAL DISSERTATIONS

Aalto University publication series DOCTORAL DISSERTATIONS 165/2015

Synthesis of biodegradable photocrosslinkable polymers for stereolithography-based 3D fabrication of tissue engineering scaffolds and hydrogels

Laura Elomaa

Doctoral dissertation for the degree of Doctor of Science in Technology to be presented with due permission of the School of Chemical Technology for public examination and debate in Auditorium KE2 (Komppa Auditorium) at the Aalto University School of Chemical Technology (Espoo, Finland) on the 2nd of December, 2015, at 12 noon.

Aalto University School of Chemical Technology Department of Biotechnology and Chemical Technology Polymer Technology

Supervising professor Prof. Jukka Seppälä, Aalto University, Finland Preliminary examiners Dr. Matthias Schnabelrauch, INNOVENT e.V., Germany Prof. Ann-Christine Albertsson, KTH Royal Institute of Technology, Sweden Opponent Prof. Peter Dubruel, Ghent University, Belgium

Aalto University publication series DOCTORAL DISSERTATIONS 165/2015 © Laura Elomaa ISBN 978-952-60-6491-8 (printed) ISBN 978-952-60-6492-5 (pdf) ISSN-L 1799-4934 ISSN 1799-4934 (printed) ISSN 1799-4942 (pdf) http://urn.fi/URN:ISBN:978-952-60-6492-5 Unigrafia Oy Helsinki 2015 Finland

Abstract Aalto University, P.O. Box 11000, FI-00076 Aalto www.aalto.fi

Author Laura Elomaa Name of the doctoral dissertation Synthesis of biodegradable photocrosslinkable polymers for stereolithography-based 3D fabrication of tissue engineering scaffolds and hydrogels Publisher School of Chemical Technology Unit Department of Biotechnology and Chemical Technology

Series Aalto University publication series DOCTORAL DISSERTATIONS 165/2015

Field of research Polymer Technology

Manuscript submitted 7 May 2015 Date of the defence 2 December 2015

Permission to publish granted (date) 23 June 2015 Language English

Monograph Article dissertation (summary + original articles)

Abstract Stereolithography (SLA) has become popular in 3D fabrication of tissue engineering (TE)

scaffolds due to its high resolution, mild building conditions, and fast production times. However, the availability of biodegradable polymers for SLA is very limited. The aim of this thesis was to synthesize new biodegradable, photocrosslinkable polymers for SLA-based scaffold fabrication, and to increase the understanding of how polymer properties and fabrication parameters affect the properties of the resulting TE scaffolds and hydrogels.

Star-shaped polycaprolactone (PCL) oligomers were synthesized via ring-opening polymerization and functionalized with methacrylic anhydride to yield photocrosslinkable macromers. The macromers were crosslinked by free radical polymerization using radical-forming photoinitiators. First, a photocrosslinkable, solid PCL macromer was used in SLA by simultaneously heating the macromer to decrease its viscosity. The PCL scaffolds prepared by SLA closely resembled their mathematically defined 3D models.

To improve the bioactivity of the scaffolds, a liquid, low-molecular PCL macromer was next combined with bioactive glass (BG) in SLA at room temperature. BG particles were homogeneously distributed within the resulting SLA fabricated scaffolds, increasing their mechanical strength. The formation of calcium phosphate deposits on the scaffold surface in simulated body fluid indicated the in vitro bioactivity of the composite material and increased the cell proliferation on the scaffold surface.

To tune the properties of the PCL macromer, a new photocrosslinkable poly(ester amide) was synthesized based on PCL and L-alanine-derived depsipeptide. Copolymerization of PCL with the depsipeptide increased the glass transition temperature and hydrophilicity of the polymer and accelerated its hydrolytic degradation. Also the compressive strength of the SLA fabricated scaffolds increased with depsipeptide content. The new copolymer significantly extended the repertoire of biodegradable polymers suitable for SLA-based scaffold fabrication.

Last, a photocrosslinkable poly(ethylene glycol-co-depsipeptide) was synthesized for SLA fabrication of cell-laden hydrogels. The stiffness of the hydrogels increased with increasing crosslinking time while the swelling degree and mass loss decreased. The encapsulated cells showed proliferation in the hydrogels, and tubular hydrogels were successfully fabricated as vascular graft models. Due to its good cell encapsulation capacity and printability, the new polymer was a highly desired addition to the very limited group of biodegradable polymers available for SLA fabrication of cell-laden TE hydrogels.

Keywords 3D fabrication; Photocrosslinking; Polycaprolactone; Polydepsipeptide; Stereolithography; Tissue engineering

ISBN (printed) 978-952-60-6491-8 ISBN (pdf) 978-952-60-6492-5

ISSN-L 1799-4934 ISSN (printed) 1799-4934 ISSN (pdf) 1799-4942

Location of publisher Helsinki Location of printing Helsinki Year 2015

Pages 136 urn http://urn.fi/URN:ISBN:978-952-60-6492-5

Tiivistelmä Aalto-yliopisto, PL 11000, 00076 Aalto www.aalto.fi

Tekijä Laura Elomaa Väitöskirjan nimi Biohajoavien valosilloitettavien polymeerien synteesi kudosteknologian tukirakenteiden ja hydrogeelien 3D-valmistukseen stereolitografialla Julkaisija Kemian tekniikan korkeakoulu Yksikkö Biotekniikan ja kemian tekniikan laitos

Sarja Aalto University publication series DOCTORAL DISSERTATIONS 165/2015

Tutkimusala Polymeeriteknologia

Käsikirjoituksen pvm 07.05.2015 Väitöspäivä 02.12.2015

Julkaisuluvan myöntämispäivä 23.06.2015 Kieli Englanti

Monografia Yhdistelmäväitöskirja (yhteenveto-osa + erillisartikkelit)

Tiivistelmä Stereolitografia (SLA) on valosilloitukseen perustuva 3D-valmistusmenetelmä, joka soveltuu

kudostekniikan tukirakenteiden valmistukseen sen korkean resoluution, mietojen olosuhteiden ja lyhyen valmistusajan vuoksi. Väitöstyön tavoite oli uusien biohajoavien valosilloitettavien polymeerien synteesi niin, että ne soveltuvat kudostukirakenteiden SLA-valmistukseen. Lisäksi tavoitteena oli tutkia polymeerien ominaisuuksien ja SLA-valmistusparametrien vaikutusta tukirakennekappaleiden ominaisuuksiin.

Työssä syntetisoitiin renkaanavaavalla polymeroinnilla biohajoavia oligomeerejä, jotka saatiin valosilloitettaviksi metakryloimalla oligomeerien pääteryhmät. Näin saadut makromeerit silloitettiin vapaaradikaalipolymeroinnilla valoherkkiä initiaattoreita hyödyntäen. Matemaattisesti mallinnetut kudostukirakenteet valmistettiin SLA:lla matalaviskositeettisistä esipolymeereistä. Kiinteä polykaprolaktoni (PKL) saatiin juoksevaksi lämmittämällä esipolymeeriä SLA:ssa, kun taas matalamman moolimassan PKL oli juokseva huoneenlämmössä. Valmiit 3D-rakenteet mukailivat tarkasti tietokonemallejaan.

Kudostukirakenteiden bioaktiivisuutta lisättiin sekoittamalla bioaktiivista lasia PKL-esipolymeeriin. Bioaktiivinen lasi jakautui tasaisesti 3D-valmistetuissa kappaleissa vahvistaen niiden mekaanista lujuutta, ja pinnalla olevat lasipartikkelit säilyivät avoimina ilman peittävää polymeerikerrosta. Kalsiumfosfaatin saostuminen tukirakenteiden pinnalle simuloidussa kudosnesteessä osoitti komposiittimateriaalin bioaktiivisuuden ja lisäsi solujen aktiivisuutta huokoisten tukirakenteiden pinnalla.

Valosilloitettavien biohajoavien makromeerien kirjoa laajennettiin kopolymeroimalla valosilloitettava polyesteriamidi käyttäen monomeereinä kaprolaktonia ja L-alaniinista johdettua depsipeptidiä. Depsipeptidi kasvatti PKL-pohjaisen polymeerin lasiutumispistettä ja hydrofiilisyyttä ja nopeutti polymeerin hajoamista. Se myös lisäsi SLA:lla valmistetun huokoisen tukirakenteen lujuutta. Uusi kopolymeeri laajensi näin mahdollisuuksia säätää makromeerin ominaisuuksia kohdesovelluksen vaatimusten mukaisiksi.

Lisäksi väitöstyössä syntetisoitiin valosilloitettava hydrofiilinen poly(etyleeniglykoli-ko-depsipeptidi). Sekoittamalla esipolymeeriliuos solususpension kanssa valmistettiin soluja sisältäviä hydrogeelejä SLA:lla. Silloitusajan lisääminen teki hydrogeelistä jäykemmän johtuen polymeeriverkoston tiivistymisestä, jolloin myös geelin laajeneminen ja hajoaminen vedessä väheni. Lisäksi polymeeristä 3D-valmistetttin onnistuneesti putkimainen verisuonimalli. Uusi polymeeri on merkittävä lisä materiaalijoukkoon, joka soveltuu soluja sisältävien biohajoavien hydrogeelien SLA-valmistukseen.

Avainsanat 3D-valmistus; Kudostekniikka; Polykaprolaktoni; Stereolitografia; Valosilloitus

ISBN (painettu) 978-952-60-6491-8 ISBN (pdf) 978-952-60-6492-5

ISSN-L 1799-4934 ISSN (painettu) 1799-4934 ISSN (pdf) 1799-4942

Julkaisupaikka Helsinki Painopaikka Helsinki Vuosi 2015

Sivumäärä 136 urn http://urn.fi/URN:ISBN:978-952-60-6492-5

Preface

I started the experimental work for this thesis as a visiting student in the De-

partment of Biomaterials Science and Technology at the University of Twente

and then continued the work in the Department of Biotechnology and Chemi-

cal Technology at Aalto University School of Chemical Technology and finally

completed the research in the Department of Orthopaedic Surgery at Stanford

University School of Medicine. I would like to gratefully acknowledge the

Graduate School in Chemical Engineering, the Finnish Foundation for Tech-

nology Promotion, the American-Scandinavian Foundation, the Research

Foundation of Helsinki University of Technology, the Emil Aaltonen Founda-

tion, and the Alfred Kordelin Foundation for funding this work.

My thesis is the result of a fruitful collaboration between interdisciplinary re-

search groups worldwide. I would like to thank Prof. Jukka Seppälä for offer-

ing me the opportunity to conduct my doctoral studies under his supervision

at Aalto University. I wish to express my warmest gratitude to Prof. Dirk

Grijpma for welcoming me to his research group at University of Twente and

giving me invaluable support at the beginning of my doctoral research. I am

deeply grateful to Prof. Yunzhi Peter Yang for warmly welcoming me into his

group at Stanford University and for being extremely supportive of my re-

search during the final two years of my thesis.

I would like to warmly thank all my co-authors, including Dr. Sandra Teixei-

ra, Dr. Risto Hakala, Dr. Harri Korhonen, Anne Kokkari, Chi-Chun Pan, Dr.

Yaser Shanjani, Dr. Andrey Malkovskiy, Prof. Timo Närhi, and Prof. Yunqing

Kang for their support and invaluable contributions to our publications. Also, I

warmly thank Dr. Jaana Rich, Lic. Tech. Minna Malin, and Sanja Asikainen for

their support during my doctoral studies. I am extremely grateful to all my

great colleagues and laboratory personnel whom I have been so lucky to work

with during these past few years.

Finally, I wish to express my warmest thanks to my family and my friends.

Thank you for the support and encouragement you have shown me. And Jiri,

my deepest gratitude belongs to you. You were there to belay me as I climbed

this route, and I have no words to thank you enough.

Stanford, March 2015

Laura Elomaa

“A scientist must also be absolutely like a child. If he sees a thing, he must

say that he sees it, whether it was what he thought he was going to see or not.

See first, think later, then test. But always see first. Otherwise you will only

see what you were expecting.”

Douglas Adams: So Long, and Thanks for All the Fish

Contents

List of Publications and Author’s Contribution List of Abbreviations and Symbols

1. Introduction ................................................................................ 1

1.1 Need for tissue engineering and additive manufacturing .................. 1

1.2 The scope of the thesis ..................................................................... 3

2. Background ................................................................................ 4 2.1 Requirements for porous TE scaffolds .............................................. 4

2.2 Stereolithography ............................................................................ 6

2.3 Biodegradable photocrosslinkable polymers for SLA ...................... 8

2.3.1 Synthetic biodegradable polyesters and poly(ester amide)s .......................... 8

2.3.2 Free-radical chain-growth photopolymerization ......................................... 10

2.3.3 Photocrosslinkable macromers ..................................................................... 12

2.3.4 Radical-forming photoinitiators ................................................................... 12

2.4 Material properties of TE scaffolds ................................................ 14

2.4.1 Overview ........................................................................................................ 14

2.4.2 Natural extracellular matrix as inspiration .................................................. 14

2.4.3 Polymer hydrophobicity ................................................................................ 15

2.4.4 Mechanical properties ................................................................................... 16

2.4.5 Polymer degradation and matrix remodeling ................................................ 17

2.4.6 Bioactive glass ............................................................................................... 18

3. Materials and Methods ............................................................ 20

3.1 Synthesis of photocrosslinkable macromers .................................. 20

3.2 Synthesis of water-soluble photoinitiator ....................................... 21

3.3 Photocrosslinking .......................................................................... 22

3.4 Stereolithography of TE scaffolds and hydrogels ............................ 23

3.5 Imaging and mechanical testing of the 3D structures ..................... 25

3.6 Cell cultures ................................................................................... 26

4. Results and Discussion ............................................................. 27 4.1 Synthesis of the monomer, oligomers, and macromers .................. 27

4.2 Photocrosslinking characteristics .................................................. 32

4.3 In vitro biodegradation of photocrosslinked films ......................... 34

4.4 Cytocompatibility of the polymers .................................................. 35

4.5 SLA fabricated scaffolds and hydrogels .......................................... 37

4.5.1 Surface morphology ....................................................................................... 37

4.5.2 Pore size distribution and pore interconnectivity ........................................ 38

4.5.3 Mechanical properties of porous scaffolds ................................................... 40

4.5.4 Bioactivity of BG/PCL composite scaffolds .................................................. 42

4.5.5 Characteristics of 3D fabricated hydrogels ................................................... 43

4.5.6 Cell proliferation on scaffolds and within hydrogels ................................... 46

5. Conclusions and Future Prospects ............................................ 51 References .................................................................................... 53

List of Publications

This doctoral dissertation consists of a summary and of the following publica-

tions that are referred to in the text by their Roman numerals.

I Elomaa L, Teixeira S, Hakala R, Korhonen H, Grijpma DW, Seppälä, JV. Preparation of poly(ε-caprolactone)-based tissue engineering scaf-folds by stereolithography. Acta Biomater 7 (2011) 3850-3856. DOI: 10.1016/j.actbio.2011.06.039.

II Elomaa L, Kokkari A, Närhi T, Seppälä JV. Porous 3D modeled scaf-folds of bioactive glass and photocrosslinkable poly(ε-caprolactone) by stereolithography. Compos Sci Technol 74 (2013) 99-106. DOI: 10.1016/j.compscitech.2012.10.014.

III Elomaa L, Kang Y, Seppälä JV, Yang Y. Biodegradable photocrosslink-able poly(depsipeptide-co-ε-caprolactone) for tissue engineering: Syn-thesis, characterization, and in vitro evaluation. J Polym Sci Part A Polym Chem 52 (2014) 3307-3315. DOI: 10.1002/pola.27400.

IV Elomaa L, Pan C-C, Shanjani Y, Malkovskiy A, Seppälä JV, Yang Y. Three-dimensional fabrication of cell-laden biodegradable poly(ethylene glycol-co-depsipeptide) hydrogels by visible light stereo-lithography. J Mater Chem B 3 (2015) 8348-8358. DOI: 10.1039/c5tb01468a.

Author’s Contribution

I Laura Elomaa was responsible for the experimental design, performed all the experimental work excluding cell experiments, and wrote the manuscript.

II Laura Elomaa was responsible for the experimental design, performed

all the experimental work excluding μ-CT scanning, cell experiments, and ion release measurements, and wrote the manuscript.

III Laura Elomaa was responsible for the experimental design, performed all the experimental work excluding GPC runs and mechanical testing, and wrote the manuscript.

IV Laura Elomaa was responsible for the experimental design, performed all the experimental work excluding AFM tests, and wrote the manu-script.

List of Abbreviations and Symbols

1H NMR proton nuclear magnetic resonance

3D three-dimensional

AFM atomic force microscopy

AMT additive manufacturing technology

BG bioactive glass

CAD computer-aided design

CL ε-caprolactone

CQ camphorquinone

DSC differential scanning calorimetry

ECM extracellular matrix

FTIR Fourier transform infrared spectroscopy

HA hydroxyapatite

HUVEC human umbilical vein endothelial cell

MMD 3-methyl-morpholine-2,5-dione

MMP matrix metalloproteinase

MSC mesenchymal stem cell

PCL poly(ε-caprolactone)

PDP polydepsipeptide

PEA poly(ester amide)

PEG poly(ethylene glycol)

PLA polylactide

PTMC poly(trimethylene carbonate)

RGDS arginine-glycine-aspartic acid-serine

SEM scanning electron microscopy

TEA triethylamine

TMP trimethylolpropane

ROP ring-opening polymerization

SLA stereolithography

Sn(Oct)2 stannous octoate

TE tissue engineering

UV ultraviolet

μ-CT micro-computed tomography

Terminology

Depending on authors and journals in the TE field, 3D fabrication methods are

interchangeably termed additive manufacturing technologies (AMT), rapid

prototyping (RP), or solid free-form fabrication (SFF). Currently, the term ‘3D

printing (3DP)’ is also commonly used to refer to the whole group of 3D fabri-

cation methods instead of only the powder/binder-based techniques as origi-

nally intended.

Symbols

Mn number average molecular weight (g/mol)

Tg glass transition temperature (°C)

Tm melting temperature (°C)

1. Introduction

1.1 Need for tissue engineering and additive manufacturing

Every year, millions of people require treatment for end-stage tissue or organ

failure resulting from disease, trauma, or aging. Every day in the USA, an av-

erage of 79 people receive an organ transplant, while an average of 21 people

dies waiting for a donor. 1 In the USA alone, more than 120,000 people were

waiting for an organ donor in 2013, while only 29,000 received a transplant;

this gap has grown each year for the past two decades due to the shortage of

suitable organ donors. 1 To address this challenge, the interdisciplinary field of

tissue engineering (TE) aims to provide patients with temporary biological

substitutes that mimic native tissue and can be used for restoring or improving

natural tissue function. 2, 3 The strategy behind TE aims to support cells that

are natively found in the body to induce natural healing of tissues and organs.

Unlike biomedical devices that cannot prevent the progressive deterioration of

injured tissue, TE materials are meant to substitute biological functions and

integrate into the healing tissue environment. 4

TE uses highly porous polymeric scaffolds to provide cells with a local envi-

ronment that enhances and regulates cell proliferation, guiding tissue for-

mation into a desired shape. These scaffolds mimic the functions of the native

extracellular matrix (ECM) where cell proliferation naturally occurs. 5 Upon

implantation into the body, the porous TE scaffolds can be acellular or contain

immature or mature cells, as shown in Figure 1. The biodegradability of the

polymer ensures that the scaffold degrades in the body while the damaged tis-

sue is repaired and no longer needs additional support. 6 Various additive

manufacturing techniques (AMT) have become increasingly popular for the

fabrication of TE scaffolds due to their increased automation and capacity for

highly controlled scaffold geometries. Combined with computer-aided design

(CAD) and various medical imaging techniques, AMTs can be used to fabricate

TE scaffolds that meet the individual needs of each patient, thus improving the

efficiency and comfort of treatment. 7, 8 In addition, use of the patient’s own

cells to form new tissue can help avoid problems related to immunological

rejection of allogeneic cells. 4

Figure 1. Additive manufacturing of TE scaffolds and their therapeutic targets in the body according to ref. 9

Even though additive manufacturing of complete, functioning human organs

has yet to be achieved, the development of new 3D fabrication techniques and

sophisticated materials allows for the fabrication of increasingly complex TE

constructs, such as skin grafts, blood vessel replicas, or vascularized bone

grafts. 10 Currently, there is particular interest in developing tissue grafts that

can restore vascular function within the graft. Adequate vascularization could

improve the integration between the TE scaffold and its host tissue and reduce

the risk of cell necrosis inside the scaffold. 11 The reduced likelihood of graft

failure would hasten the clinical translation of these TE constructs. Commer-

cial sales in the TE industry are rapidly increasing with the number of compa-

nies selling various low-cost TE products and services. 12 Besides the TE field,

pharmaceutical research can also benefit from the 3D fabrication of complex

artificial tissue systems, as they can help to screen drugs more efficiently and

decrease reliance on animal testing.

Since additive manufacturing has become increasingly popular in scaffold

fabrication and holds a great promise both in TE and pharmaceutical industry,

there is a significant need for biodegradable polymers that are suitable for the-

se fabrication techniques. However, the variety of such polymers is currently

very limited. The synthesis of new biodegradable and biocompatible polymers

is essential for the progress of the whole TE field. New custom-made polymers

with tunable properties will allow for 3D fabrication of highly defined TE scaf-

folds that can more closely satisfy the specific requirements of their target ap-

plications.

1.2 The scope of the thesis

This thesis aimed to expand the repertoire of biodegradable, photocrosslinka-

ble polymers that are suitable for 3D fabrication of TE scaffolds by stereo-

lithography (SLA), a photocrosslinking-based AMT that allows for high-

resolution scaffold fabrication without the need for harsh solvents or high pro-

cessing temperatures. Subsequently, this thesis sought to increase the under-

standing of how polymer properties and scaffold fabrication parameters affect

properties of the resulting TE scaffolds.

In the first publication, a group of photocrosslinkable poly(ε-caprolactone)

(PCL) macromers of different molecular weights were synthesized, and the

macromer with the lowest molecular weight was used for the fabrication of

mathematically defined TE scaffolds by visible light SLA. The star-shaped,

methacrylated macromer was hypothesized to be suitable for high-resolution

solvent-free scaffold fabrication by SLA, assuming that suitable viscosity can

be achieved by heating the solid macromer above its melting point. Then, the

use of potentially harmful solvents in scaffold fabrication could be avoided.

The second publication aimed to improve the bioactivity of the 3D fabricated

scaffolds by combining the photocrosslinkable PCL resin with bioactive glass

(BG). A liquid, low molecular weight PCL macromer was synthesized and sub-

sequently used with BG for scaffold fabrication by SLA. By using the liquid

PCL resin, the need for heating was avoided and SLA conditions were kept as

mild as possible. The scaffolds were characterized by mechanical tests and in

vitro bioactivity tests, and cell proliferation on the scaffolds was evaluated.

The ion release from BG was hypothesized to increase cell attachment and

proliferation on the composite scaffolds compared to neat polymer scaffolds.

In the third publication, a photocrosslinkable poly(ester amide) macromer

based on ε-caprolactone (CL) and the L-alanine-derived depsipeptide was syn-

thesized to expand the properties of available biodegradable scaffold polymers.

The newly synthesized liquid macromers were used for scaffold fabrication by

SLA. allow for tuning its degradation

and mechanical properties through the changes in its monomer ratio. The

chemical, thermal, and degradation properties of the new materials were char-

acterized, and their in vitro cytocompatibility was evaluated.

The fourth publication extended the use of biodegradable, photocrosslinka-

ble polymers to the 3D fabrication of cell-laden hydrogels by SLA. The aim was

to synthesize a new photocrosslinkable, water-soluble macromer based on

poly(ethylene glycol) (PEG) and L-alanine-derived depsipeptide to obtain a

macromer that is suitable for cell encapsulation during SLA. The chemical

structure and gelling properties and the hydrolytic mass loss and mechanical

properties as well as encapsulation capacity of the new polymer were charac-

terized. It was hypothesized that the new macromer that combines both natu-

rally derived and synthetic building blocks will take advantage of the biological

properties of peptides and the great controllability of the physical and me-

chanical properties of PEG. Subsequently, it was hypothesized that by control-

ling the light exposure time in SLA, the hydrogel properties can be tuned with-

out need for changing the parameters of the hydrogel solution in the process.

2. Background

2.1 Requirements for porous TE scaffolds

A porous TE scaffold should ideally replicate the native ECM structure and

allow cells to remodel the scaffold structure according to their own require-

ments. 13 The ECM is a highly porous 3D structure composed of interlocked

fibrous proteins and polysaccharides that are secreted by resident cells. 14 This

porous mesh provides the tissue with structural integrity and directs cell be-

havior through biomechanical interactions and mechanical cues. 15 As a substi-

tute for ECM, an essential function of TE scaffolds is to guide the proliferation

and growth of cells to form healthy new tissue. Since most of the primary or-

gan cells are anchorage-dependent, they cannot form new tissue by them-

selves, requiring the presence of a supporting structure to serve as a template

for cell growth. 16, 17 TE scaffolds mechanically support the regenerating tissue

and, if needed, can deliver growth factors or therapeutic agents to enhance

tissue growth or to treat diseases. 18, 19

Since each tissue type requires a specific matrix structure with certain mate-

rial properties, there is no single ideal scaffold design that can be defined. Ap-

propriate scaffold design requires the consideration of the matrix’s 3D archi-

tecture, including pore size and morphology, surface characteristics, mechani-

cal integrity, and the degradation behavior of the material. 20 Even though the

optimal properties of the scaffold depend on the tissue type to be mimicked,

some general criteria can be stated for an ideal scaffold. First, the material is

not allowed to be cytotoxic or systemically toxic, but needs to be biocompatible

with living cells. 17 This requirement also applies to the material’s degradation

products. Since biocompatibility depends not only on the material itself, but

also its biological environment, there is a need for both in vitro and in vivo

testing of the material. 6 In addition to biocompatibility, the scaffold needs to

be sterilizable and commercially reproducible. 17

Since the TE scaffold supports tissue growth within its 3D structure, the ide-

al scaffold is highly porous and its pores are well interconnected to ensure both

homogenous distribution of cells and a sufficient flow of oxygen and nutrients

inside the scaffold. Within the porous polymer scaffolds, two different porosity

levels exist. The total porosity consists of micropores on a scale of nanometers

to micrometers, indicating the volumes free of polymer chains, and of

macropores on a scale of micrometers to millimeters resulting from the scaf-

fold fabrication. 21 In TE applications, the microporosity is typically relevant

only in hydrogel matrices. The optimal size of macropores for efficient tissue

growth and vascularization depends on the tissue type: for example, the opti-

mal size for bone tissue scaffolds has been estimated in the range of 150-800 μm. 22-24 Even though high porosity is required, if the pore size and overall po-

rosity of the scaffold are high, they may dramatically decrease the mechanical

properties of the scaffold. 25 If mechanical support is inadequate, tissue for-

mation may fail due to excessive deformation of the scaffold. The mechanical

properties of the scaffold must closely match the properties of the target tissue

to prevent further degeneration of the defect area. 26 Therefore, finding a bal-

ance between the desired mechanical properties and good porosity is essential

for optimal performance of the TE scaffold.

While highly porous, mechanically strong TE scaffolds are desired in hard

tissue applications, flexible materials with high water intake are preferred in

soft tissue engineering. Hydrogels are a special group of TE scaffolds that con-

sist of water-absorbing, crosslinked polymer networks formed by either ionic,

covalent, or physical bonding of the polymer chains. 27 Natural hydrophilic

polymers like collagen, gelatin, and chitosan are widely used as hydrogel mate-

rials due to their inherent bioactivity; however, synthetic hydrophilic polymers

like PEG or polyacrylamides have the advantage of easy tuning of their chemi-

cal, mechanical, and viscoelastic properties. 28 Their high water content, per-

meability, 29 and soft tissue-mimicking viscoelastic characteristics 30 make hy-

drogels especially attractive for use in cell-laden constructs where the cross-

linked hydrogel matrix provides an ECM-like microenvironment for encapsu-

lated cells.

Like other TE scaffolds, hydrogels can contain both micropores due to the

void between polymer chains as well as macropores due to the hydrogel fabri-

cation process. The mesh size determining the microporosity of the photo-

crosslinked hydrogels depends on the crosslinking density of the gel, and that

in turn depends on the polymer molecular weight and concentration, 31 as well

as crosslinking time, 32 photoinitiator concentration, and light intensity. 33 The

nanoscale micropores allow for diffusion of small molecules such as oxygen

and glucose into the hydrogel, but may restrict the migration of cells and the

diffusion of larger ECM components in the 3D matrix. 34 These diffusion and

migration problems can be addressed by using either biodegradable polymers 35 or by introducing additional macroporosity into the hydrogel 36 to allow for

sufficient nutrient flow and cell proliferation. When a biodegradable polymer

is used, increase in the pore size and interconnectivity due to material degra-

dation supports cell ingrowth, angiogenesis, and tissue formation within the

hydrogel. 37

Porous scaffolds have been traditionally prepared using conventional tech-

niques, such as phase separation, 38 gas foaming, 39 or salt leaching. 40 Fast

eroding poly(ester anhydride) fibers have also been used to leach a pore net-

work within a more slowly degrading polymer matrix. 41 Even though these

conventional techniques can be relatively inexpensive, they are typically time-

consuming and suffer from poor control over pore size and distribution, there-

by causing poor pore interconnectivity and lack of mechanical strength. 42, 43 By

using additive manufacturing combined with defined pore architecture de-

signs, the mechanical strength of the scaffold can be increased and the pores

can be made highly interconnected, improving cell seeding and tissue growth

into the scaffold matrix. 44-46 AMTs, such as SLA, fused deposition modeling, or

selective laser sintering, are used for the rapid fabrication of complex free-

form parts defined by CAD models. 46-50 These techniques join liquid, powder,

or sheet materials in a layer-by-layer manner to form desired 3D structures.

Since no molds are used, the shape of the scaffold is not limited, and can be

designed to satisfy the specific requirements of the target tissue and custom-

ized to each individual patient. 46, 51 AMTs also allow for the fabrication of ani-

sotropic and mechanically graded scaffold structures, fulfilling the require-

ments for growing various cell and tissue types within the same scaffold. 52, 53

2.2 Stereolithography

SLA is a photocrosslinking-based additive manufacturing technology that al-

lows for the preparation of highly porous TE scaffolds by the solidification of a

liquid resin into well-defined geometries following a CAD file. 54 The advantage

of SLA is its capacity to build 3D structures in a spatially and temporally highly

controllable manner at room temperature, which also allows for the incorpora-

tion of living cells or heat sensitive peptides into the scaffold during fabrica-

tion. 36 The CAD file can be obtained by using mathematical equations and 3D

modeling software 52 or by utilizing clinical imaging techniques such as CT, 55

MRI, 56 or ultrasound imaging systems. 57

In the SLA process, the CAD model is sent to the SLA machine to fabricate

the 3D structure using UV or visible light in a layer-by-layer manner as shown

in Figure 2. 54 SLA requires information on the geometry and size of the 3D

structure in the form of a STL file, where the boundary surfaces of the object

are represented as numerous tiny triangles. 54 The STL file is sectioned into a

series of layers and transferred to the SLA machine; then a computer-

controlled laser beam or a digital light projector (DLP) with a digital mirror

device (DMD) is used to photocrosslink the desired pattern of the given layer

into the liquid resin. 50 In the case of the DMD, the picture is formed using an

array of rotatable micro-mirrors that reflect the light either to the focusing lens

(photocrosslinked voxels) or outside it (non-crosslinked voxels) according to

the designed pattern. 58 After photocrosslinking one layer, a build platform

where the 3D part is located moves away from the resin surface giving space

for the next layer. The photocrosslinking is repeated layer-by-layer to form the

final 3D structure. Depending on the SLA setup, the crosslinking light can

come from above the build platform to photocrosslink a thin layer of resin on

top of the 3D structure (bottom-up), or it can be projected from underneath

through a transparent plate containing the resin to crosslink a layer at the bot-

tom of the structure (top-down). 59

Figure 2. A schematic representation of top-down projection SLA with a digital mirror de-

vice.

To achieve high resolution, SLA requires a high level of control over the layer

thickness being crosslinked. This thickness, called cure depth (Cd), depends on

the energy of light exposure on the resin surface (E) and can be controlled by

adjusting the intensity of the light source and the exposure time. The depend-

ence can be derived from the Beer-Lambert law and described using Equation

(1) called the working curve equation: 60

(1)

In Equation (1), Ec is the critical exposure (mJ/cm2), meaning the minimum

energy level required to transform the photopolymer from liquid to solid, and

Dp is the penetration depth of the light into the resin. Both Ec and Dp of the

given resin can be determined by crosslinking various polymer layers at differ-

ent energy levels and plotting the layer thickness Cd as a function of the natural

logarithm of the crosslinking energy E. 60

Poor control over the cure depth will decrease the resolution of the scaffold

because the overlong crosslinking time results in over-cure. Consequently,

some voxels designed to remain uncrosslinked will polymerize, making the

pore size and porosity lower than designed. Exposure times that are too short

do not allow for proper crosslinking of the polymer layers, and the layer thick-

ness may remain inadequate for successful scaffold fabrication. To ensure

good attachment between subsequent layers, each layer should be crosslinked

for a slightly longer time than required for the given layer thickness. 61 If need-

ed, over-cure into the preceding layer can be reduced by increasing the pho-

toinitiator concentration or by adding dye into the resin. The dye competes

with the photoinitiator in absorbing light, decreasing the penetration depth of

the light due to the increased extinction coefficient of the resin. 54, 61 Thus, the

dye improves control over the cure depth, being especially useful in visible

light-sensitive resins where the extinction coefficient is typically lower when

compared to UV sensitive resins. 54

Compared with other AMTs, SLA shows its superiority in its high resolution.

The resolution of commercially available SLAs can be as high as tens of mi-

crometers, while other techniques can typically build structures with a resolu-

tion of only hundreds of micrometers. 49, 59, 62 Even higher resolutions can be

achieved using advanced micro- and two-photon SLA techniques. 63-67 The res-

olution of SLA depends on the technique used for picture generation. Typical-

ly, scanning-based laser SLAs can produce large 3D structures with lower reso-

lution, while projection SLAs are suitable for smaller parts requiring higher

resolution. 68 Besides the higher resolution, DLP techniques used in projection

SLAs produce the entire image at once, thus reducing production times com-

pared with scanning-based laser techniques. 54 Both laser SLAs 24, 69 and pro-

jection SLAs 70, 71 have been used for the fabrication of TE scaffolds. Recently, a

new type of SLA has been introduced that combines the benefits of both scan-

ning and projection SLA techniques. 68, 72 In this new technique, called scan-

ning-projection SLA, the whole DMD moves over the resin while projecting a

continuously updated pattern to photocrosslink large areas with high resolu-

tion. Another recent technique takes advantage of the usually undesired phe-

nomenon of oxygen inhibition to continuously build high-resolution 3D ob-

jects by UV crosslinking. 73 These new emerging techniques facilitate the fast

production of large 3D structures with small feature size, thus extending the

use of SLA to fabrication of large TE constructs.

2.3 Biodegradable photocrosslinkable polymers for SLA

2.3.1 Synthetic biodegradable polyesters and poly(ester amide)s

Biodegradability is one of the most significant properties of TE scaffolds; it

ensures that the material gradually disappears from the body and no addition-

al removal surgery is needed. Consequently, biodegradable polymers that con-

tain hydrolytically labile chemical bonds are highly desired as scaffold materi-

als, including polyesters, polyanhydrides, polycarbonates, and polyamides. 74-76

According to Vert et al., 77 biodegradable polymers are polymers that break

down enzymatically or non-enzymatically through polymer chain cleavage

when in contact with body fluids. In addition, bioresorbable polymers resorb

in the body through natural pathways. 77 Hydrolysis of a polymer occurs when

nucleophilic water or hydroxyl ion attacks the electron-withdrawing carbonyl

group and breaks the bond between the carbonyl and a heteroatom through

nucleophilic substitution with or without catalytic action of enzymes. 78, 79 Deg-

radation of the polymer chain decreases the polymer’s molecular weight, re-

sulting in mass loss of the polymer matrix and causing polymer erosion. 75 If

water diffusion into the polymer matrix is slower than polymer chain degrada-

tion, the matrix undergoes surface erosion resulting in mass loss on the matrix

surface. If water penetration is faster than polymer degradation, the matrix

undergoes bulk erosion, resulting in mass loss throughout the matrix. 80

Synthetic polyesters and polycarbonates, such as poly(lactic acid) (PLA),

PCL, and poly(trimethylene carbonate) (PTMC), are widely used biodegrada-

ble polymers in TE. PLA is a particularly popular biodegradable polymer that

can be synthesized in three isomeric forms, poly(L-lactide) (PLLA), poly(D-

lactide) (PDLA), and poly(D,L-lactide) (PDLLA). 81 High molecular weight

(HMW) PLLA is a comparatively strong, semi-crystalline polymer with a melt-

ing temperature (Tm) of 173-178 °C and a glass transition temperature (Tg) of

60-65 °C; it degrades in approximately two years. HMW PDLLA is an amor-

phous polymer with a Tg of 55-60 °C and degrades in approximately 12 to 16

months. 74 Due to their different characteristics, PLLA is typically used in load-

bearing applications and PDLLA in low strength scaffolds and drug release

applications. HMW PCL is a semi-crystalline polymer with a Tm around 60 °C

and Tg around -60 °C; due to its hydrophobicity and high crystallinity, it has a

comparatively long biodegradation time, taking two or three years. 74-76, 82 The

modulus of PCL is significantly lower than the modulus of PLA, and it is typi-

cally used in long-term implantable drug delivery systems and low-load TE

applications. PLA and PCL are both bioresorbable as their hydrolysis products,

lactic acid and 6-hydroxyicaproic acid, respectively, are metabolized via the

citric acid cycle and excreted naturally from the body. 82, 83

Even though the aliphatic polyesters have tunable mechanical properties,

and are sensitive to hydrolytic degradation, their lack of biological functionali-

ty may limit their use in biomedical applications. 84 To overcome these limita-

tions, the copolymerization of esters with α-amino acids resulting in poly(ester

amide)s (PEA)s may provide a feasible approach for tailoring polymer proper-

ties. Amino acids are attractive building blocks due to their abundant availabil-

ity and the large diversity of functional groups in their side chains. 85 As degra-

dation products, they are expected to be non-toxic and readily metabolized. 86

PEAs combine the biological functionality of amino acids with the hydrolyti-

cally labile ester groups of polyesters that ensure polymer biodegradability. 87

An important group of PEAs, polydepsipeptides (PDP)s, are polymers with

alternating ester and peptide bonds that can be synthetized via ring-opening

polymerization of morpholine-2,5-dione (MD) derivatives. 88, 89 The cyclic MD

monomer allows for easy copolymerization with various polyesters, such as

PLA and PCL, thus providing a facile method for modifying the properties of

these biodegradable homopolymers. 90-93

Even though most of the polymers in the biomedical field are thermoplastics

that can be melted and shaped repeatedly, the free-radical photopolymeriza-

tion in SLA requires the use of photocrosslinkable macromers that form net-

works under exposure to light where the covalent bonds crosslink the polymer

chains. 94 The resulting crosslinked thermoset with a molecular weight ap-

proaching infinity cannot be melted or dissolved, although it may swell in a

compatible medium. 95 While the thermoplastic polymers used in TE are usual-

ly HMW linear polymers, effective photocrosslinking typically requires the use

of low molecular weight, star-shaped macromers due to their high double

bond concentration. 96, 97 In the following chapters, the mechanism of free-

radical photopolymerization and characteristics of the photocrosslinkable

macromers and radical-forming photoinitiators are discussed more closely in

light of their significance in SLA-based scaffold fabrication.

2.3.2 Free-radical chain-growth photopolymerization

The SLA technique is based on free-radical photopolymerization as an efficient

method for converting a liquid prepolymer resin into a solid polymer network

under light exposure. 98 Photopolymerization requires three essential ele-

ments: electromagnetic radiation, photoinitiator molecules, and monomeric or

macromeric precursors with unsaturated vinyl moieties. 99 In free-radical

chain-growth photopolymerization, a photoinitiator absorbs light either in the

UV or visible light range, and the photons from the light source excite the pho-

toinitiator molecules into a high-energy radical state. 100, 101 Upon exposure to

light, the photoinitiator is promoted from its ground singlet state to its excited

singlet state, and then converted into its triplet state which yields reactive rad-

ical molecules. 102 Subsequently, these initiator radicals interact with precursor

molecules, forming the primary radicals that initiate the polymerization reac-

tion. 100, 102 Chain-growth polymerization then propagates through additional

vinyl groups. If a precursor molecule with more than two vinyl groups is used,

as shown in Figure 3, a complex crosslinked network is formed in a process

called photocrosslinking. 103

Figure 3. A schematic representation of free-radical chain-growth photopolymerization.

Photocrosslinking allows for the fast solidification of the liquid resin into a

polymer network at room temperature with high spatial and temporal control

over the polymerization.104 The rate of polymerization rapidly drops to zero

when the initiating light source is shuttered, so the reaction can be readily

stopped at any time. 103 During the polymerization, the material undergoes

deformation from a liquid to a loosely crosslinked rubbery gel, and then to a

highly crosslinked glassy network. 103 Thus, by controlling the light exposure

time, the desired material properties can be achieved.

In photopolymerization, two different phases can be distinguished: 103, 105

first, the polymerization is started with an autoacceleration phase during

which the polymerization rate increases strongly as a function of conversion.

In this stage, the mobility of the radicals in the increasingly crosslinked poly-

mer network decreases, resulting in an increase in radical concentration. Next,

an autodeceleration phase starts when the polymerization reaction reaches a

maximum rate and begins to decrease. During that stage, the propagation re-

action is first restricted and eventually stopped by vitrification. The rate of

free-radical chain-growth photopolymerization can be described by Equation

(2): 106

(2)

In this equation, [M] is the monomer concentration, kp and kt are the rate

constants for propagation and termination, respectively, Φ is the quantum

yield describing the number of propagating chains initiated per one photon

absorbed by the polymer, and Ia is the absorbed light intensity, which varies

with the depth of penetration. The surface area form of volumetric Ia (mol cm-3

s-1), termed I’a (mol cm-2 s-1), is obtained from the Beer-Lambert law that de-

scribes the exponential decay of the intensity of light passing through the me-

dium according to Equation (3): 106

(3)

In the equation, I0 is the light intensity on the material’s surface, D is the

travelled distance, [A] is the photoinitiator concentration, and α is the absorp-

tion coefficient of the photoinitiator (α= ε ln 10, where ε is the extinction coef-

ficient of the initiator). 106

Since light intensity decreases exponentially inside the material according to

the Beer-Lambert law (2), the rate of polymerization decreases with increasing

depth. Especially in thicker polymer samples, light attenuates dramatically

when going through the material. No polymerization occurs beyond the critical

depth, where the effective light intensity approaches zero. 103

Beside light attenuation, the presence of oxygen in the reaction can also limit

photocrosslinking. Oxygen reduces the overall efficiency of free radical photo-

polymerization by reacting with the radical molecules to form peroxy radicals.

107, 108 Peroxy radicals cannot readily reinitiate polymerization, and therefore

they effectively inhibit or retard the reaction. This oxygen inhibition may lead

to increased reaction times and tacky surfaces of the resulting polymers due to

the incomplete double bond conversion of the precursor molecules in contact

with the air. In SLA, the reduced crosslinking efficiency due to light attenua-

tion and oxygen inhibition decreases the resolution of the scaffold, and thus

requires special attention when designing the scaffold fabrication setup. In

addition, material shrinkage due to material phase changes during scaffold

fabrication may change the scaffold dimensions and thus decrease the resolu-

tion. The effect of material shrinkage on the resolution can be reduced by mak-

ing compensations to the CAD model before the scaffold fabrication. 109

2.3.3 Photocrosslinkable macromers

Free-radical photopolymerization has been widely applied in industrial appli-

cations, including coatings, adhesives, and inks, as well as in dental restora-

tives due to the good availability of double bond-containing precursors, most

typically various acrylates and methacrylates. 98 In tissue engineering, the pol-

ymers need to be biodegradable and biocompatible, which significantly limits

the selection of available materials. To date, a limited variety of biodegradable

photocrosslinkable polymers has been synthesized, including polymers based

on polyanhydrides 103 and various polyesters and polycarbonates, such as PLA,

PCL, and PTMC. 96, 110-113 To obtain a photocrosslinkable macromer with double

bonds, the biodegradable hydroxyl-ended polymer, called oligomer, must be

functionalized with double bond-containing end-groups. Most often, the high-

ly reactive acrylate or methacrylate groups used are derived from acryloyl or

methacryloyl chloride 97, 114 or methacrylic anhydride. 104, 115 The use of less re-

active maleate groups has also been reported. 116, 117 Polyfumarates that contain

double bonds along their backbone do not need further functionalization be-

fore photocrosslinking. 118 However, polyfumarate-based macromers typically

require the use of toxic solvents as a crosslinkers to obtain a sufficient rate of

photocrosslinking. 119, 120

For SLA-based scaffold fabrication, several biodegradable photocrosslinka-

ble macromers have been used. One of the first prototypes of biodegradable

SLA fabricated TE constructs was a disc-shaped sample based on

poly(propylene fumarate) (PPF) dissolved in diethyl fumarate (DEF). 24 A simi-

lar PEF/EF macromer system was used in several later studies. 119, 121-123 More

recently, resins based on methacrylated PDLLA dissolved in non-reactive ethyl

lactate diluent 61 and fumaric acid-functionalized PDLLA dissolved in reactive

N-vinyl-2-pyrrolidone 120 have been used for SLA fabrication of porous TE

scaffolds. Methacrylated PCL has been used by the author for SLA-based fabri-

cation of porous trachea scaffolds without need for diluents, 124 while methac-

rylated PTMC required the use of propylene carbonate as a non-reactive dilu-

ent. 125 In addition, (meth)acrylated PEG-based macromers have been used for

SLA fabrication of hydrogels in several studies. 34, 126-128 PEG is a hydrophilic

polymer that is widely used in the production of soft non-degradable hydro-

gels. Water-soluble polyether PEG can be defined as a bioabsorbable polymer

since it does not degrade through chain cleavage but can be excreted by the

kidneys up to high molecular weights. 129 However, due to the lack of degrada-

ble bonds, the crosslinked PEG hydrogel cannot be used as a self-degrading

temporary support in the body. 60 To address the lack of degradation, a photo-

crosslinkable resin based on biodegradable PEG-co-PLA macromer has been

previously formulated and used for SLA fabrication of degradable hydrogels

constructs. 71

2.3.4 Radical-forming photoinitiators

Conversion of a multifunctional macromer to a crosslinked polymer is initiated

by radical-forming photoinitiator molecules. Photoinitiators typically contain a

carbonyl group with non-bonding electrons that can be promoted to the π*

antibonding orbital by the absorption of light at the appropriate wavelength. 130

Free-radical photopolymerization can be initiated by producing radicals via

two different pathways. In a photocleavage mechanism, the photoinitiator un-

dergoes excitation by energy absorption and subsequently decomposes into

two initiating radicals; in a hydrogen-absorption mechanism, the molecule

undergoes excitation and subsequently interacts with a second compound, an

electron donor, by electron transfer to form a reactive radical. 106, 107, 131

Aromatic ketones are widely used photoinitiators due to their high quantum

yields of radical production. 106 They undergo a photocleavage reaction where

scission occurs at the bond between the carbonyl carbon and the α-carbon.

Similar photolytic reactions occur with benzoin, benzyl ketals, aroylphosphine

oxides, and α-aminoalkylphenones. 106 In the hydrogen-absorption mecha-

nism, where photoinitiators require a reducing agent to efficiently produce

free radicals to initiate polymerization, 132 tertiary amines with an abstractable

proton on the α-carbon are considered one of the most effective electron do-

nors. 106, 133

Most of the commercially available photoinitiators have traditionally been

sensitive to UV light. 101 However, UV light exposure is known to be harmful to

cells, causing damage to DNA and the cell membrane. 134 Consequently, signifi-

cantly lower intensities are acceptable for UV light than for visible light in a

clinical setting, thus impairing control of polymerization time by light intensi-

ty. 135 Another significant limitation of UV light is its poor depth of curing,

since many UV initiators strongly attenuate the polymerizing light. 104 In addi-

tion, UV light sources are more expensive than visible light sources. 106

The capacity of a photoinitiator to absorb light depends on its extinction co-

efficient (ε). A high ε value indicates a high probability of absorption at the

given wavelength, resulting in high quantum yields for the initiator. Typically,

many photoinitiators show the highest ε values at wavelengths near the UV-

visible light limit. 136 Visible light photoinitiators typically have a lower extinc-

tion coefficient, meaning that they do not absorb light as effectively as UV pho-

toinitiators do. 54 This results in deeper light penetration for visible light sensi-

tive resins than UV resins. However, the increased light penetration can cause

over-cure, making control over photocrosslinking more difficult; use of a dye

may be needed, as discussed before. 54

The biocompatibility of the photoinitiator is important because the unreact-

ed initiator molecules may leach from the photocrosslinked material into the

surrounding tissue. In addition, when photocrosslinking occurs in the pres-

ence of living cells, the photoinitiator cannot produce heat or high concentra-

tions of free radicals and the crosslinking needs to be fast. 137 The high-energy

radicals may cause cell death by attacking the double bonds of unsaturated

fatty acids and phospholipids in cell membranes. 138 They can also decrease the

bioactivity of essential proteins. 139 The cytotoxicity of photoinitiators is known

to depend on their hydrophobicity, presumably because the phospholipid bi-

layer of cell membranes is more permeable to hydrophobic molecules than

hydrophilic ones. 140 The initiator concentration can also have a critical effect

on cell survival since cytotoxicity of the initiator may dramatically increase

with its concentration. 141-143 In addition, biocompatibility studies have shown

that the cytotoxic effect of the same photoinitiator varies between different cell

lines. 140 In order to avoid potential cytotoxic effects of residual initiator, some

highly reactive macromers can be photocrosslinked even in the absence of

photoinitiators. However, the light intensity or exposure time have to be in-

creased significantly, which in turn may result in side reactions, such as photo-

lytic cleavage of ester bonds, and subsequently lead to changes in the polymer

physicochemical properties. 144

2.4 Material properties of TE scaffolds

2.4.1 Overview

When foreign material comes into contact with body fluids in vivo or cell cul-

ture medium in vitro, proteins adsorb to the material’s surface and cells inter-

act with the material through this adsorbed protein layer. 145 Since the main

function of TE scaffolds is to act as templates for tissue growth, good control

over protein adsorption and cell attachment is required. Surface properties,

such as chemical composition, roughness, and polarity of the scaffold material

are important in regulating the protein adsorption and cell attachment. 146

Since these material properties are typically interconnected, it is not always

possible to clearly identify the effect of one individual property on protein ad-

sorption and cell attachment; instead, the scaffold material needs to be con-

sidered as a complex combination of various interacting properties. Therefore,

the development of multifunctional biomaterials benefits from cooperation

between interdisciplinary research teams working together towards the specif-

ic biomedical applications. 147

2.4.2 Natural extracellular matrix as inspiration

The ECM surrounding cells in natural tissue consists of functional proteins,

glycosaminoglycans (GAG), and signaling molecules. 14 Cell adhesion to cell-

binding epitopes in the ECM coordinates cell fate processes including prolifer-

ation, differentiation, and migration. 148 The concentration of various proteins

and GAGs determines the geometry, porosity, and mechanical properties of

the ECM, all of which vary between different tissue types. 149

The major ECM proteins are elastin, collagen, fibronectin, and laminin, each

of which has its own important functions in the ECM. 14 Elastin is responsible

for the elastic properties of many tissues like arteries, while collagen gives ten-

sile strength to tissues like bone and skin. Fibronectin contains important col-

lagen-, heparin-, and cell-binding domains, and crosslinked laminin is known

to be the main component of basement membranes. 148 The main GAG poly-

saccharides include several sulfates, hyaluronic acid, and heparin. 150 In close

interaction with ECM proteins, sulfated and negatively charged GAGs provide

cells with a structural network and determine tissue resilience by maintaining

high water uptake. 148 In TE, ECM components have been widely used as both

biodegradable building blocks in hydrogel matrices and biomimetic implant

coatings. 26, 151-153

The exchange of information between cells and the ECM is mediated by

transmembrane receptors called integrins, which pass information about the

chemical composition and mechanical status of the ECM to cells and monitor

the status of cells to induce changes in the environment. 154 The attachment of

cells to the ECM occurs through interactions between integrins in the cell

membrane and integrin-specific ligands in ECM proteins. In TE scaffolds, in-

tegrin-binding proteins such as collagen and fibronectin have been widely

used to increase cell attachment to the material’s surface. 152, 155 However, cell

adhesion to ECM-like materials does not necessarily require the presence of

entire integrin-binding macromolecules: integrins can recognize and bind to

short peptide sequences alone. One of the most widely studied peptides, RGD

(arginine-glycine-aspartic acid), has been used with great success to improve

cell attachment on TE scaffolds. 156 However, the use of integrin-binding pep-tides in TE scaffolds does not guarantee improved cell attachment; the concen-tration and additional amino acid sequence of the peptide, as well as the length of the linker between the peptide and the scaffold, all significantly affect the peptide’s bioactivity. 157-160

2.4.3 Polymer hydrophobicity

The wetting properties of the polymer surface are an important factor in de-

termining protein absorption and cell adhesion to the polymer. Polymer hy-

drophobicity is defined by the surface free energies of solid-gas, liquid-gas,

and solid-liquid interfaces in polymer-water systems. 161 Liquid wets the poly-

mer surface if its molecules have a stronger attraction to the polymer mole-

cules than to each other. As a general rule, polymers exhibiting water contact

angles less than 65° are considered hydrophilic, while polymers with contact

angles greater than 65° are considered hydrophobic. 162

Hydrophobicity of the polymer surface strongly affects the extent of protein

adsorption. In an aqueous environment, proteins readily adsorb to hydropho-

bic surfaces since the replacement of water molecules at the hydrophobic pol-

ymer surface with adsorbed proteins results in a reduction of interfacial ener-

getics. 162 On hydrophilic surfaces, protein adsorption is energetically unfavor-

able. However, highly hydrophobic surfaces may induce strongly irreversible

protein adsorption and denature the native conformation of proteins. 102 As a

result, the proteins present their binding domains in a sterically unfavorable

manner and lose their bioactivity. 163 In addition, a highly hydrophobic poly-

mer surface leads to poor sample wetting and can result in poor cell attach-

ment. 164, 165 Even though the ideal hydrophobicity of a material is controversial

and depends on the target application, neither a highly hydrophilic nor a high-

ly hydrophobic polymer surface is optimal for cell attachment. 163

Unlike most of the natural polymers, synthetic biodegradable polyesters like

PLA and PCL are hydrophobic due to their low number of polar groups. 166

Since hydrophobicity is caused by non-polarity of the polymer, the polymer

can be made more hydrophilic by introducing polar groups, such as hydroxyl,

carboxyl, or amino groups, to the polymer surface by various techniques, in-

cluding plasma surface treatment and alkaline hydrolysis. 167-170 Another widely

used approach to increase polymer hydrophilicity is to copolymerize the hy-

drophobic polymer with a hydrophilic PEG polymer. 171-173 Even though PEG

does not contain polar side groups, it is a highly hydrophilic polymer due to its

high number of oxygen-containing ether groups that interact with water mole-

cules via hydrogen bonds, making it readily water-soluble. 174, 175

2.4.4 Mechanical properties

Mechanical properties of TE scaffolds are an essential factor in regulating cell

fate and tissue formation on the scaffold because cells on the scaffold surface

can feel and respond to substrate stiffness. 13 When cells anchor to the scaffold

surface, they probe its mechanical properties and transmit contractile forces to

the substrate through transcellular structures; cells respond to resistance from

the substrate by adjusting their adhesion and cytoskeleton organization. 176

Mechanical properties particularly affect stem cell differentiation on the scaf-

fold surface: 177 human mesenchymal stem cells (MSC) have shown the greatest

expression of neurogenic markers on soft matrices, myogenic markers on stiff-

er ones, and osteogenic markers on rigid matrices. 178 In another study, scaf-

folds of lower stiffness enhanced chondrogenesis of mesenchymal progenitor

cells, while stiffer scaffolds supported more osteogenesis. 179

For photocrosslinked polymers, modification of the molecular weight and ar-

chitecture of the precursor macromer provides a facile method for tuning the

mechanical properties of the resulting network. Typically, photocrosslinking of

a low molecular weight macromer leads to a highly crosslinked glassy material,

while photocrosslinking of higher molecular weight macromers results in a

more loosely crosslinked rubbery network. 97 Linear macromers typically form

networks with lower gel content than star-shaped macromers, thus affecting

the mechanical properties of the crosslinked polymer. 112 The low gel content

often results in decreased mechanical properties due to the plasticizing effect

of unreacted macromers in the resulting polymer. 104

Mechanical properties are especially important in cell-laden hydrogels where

the cells are surrounded by the polymer network. Cells and the surrounding

matrix interact in two ways: while the mechanical, structural, and chemical

composition of the matrix regulate intracellular processes, the cells are capable

of remodeling the matrix according to their own needs. 13 Changes in matrix

properties, such as stiffness, mesh size, and viscoelastic properties, affect cell

activity and phenotype13 as well as their morphology, 180 proliferation, and mi-

gration. 181 The spreading of encapsulated fibroblasts, 181 smooth muscle cells

(SMC), 180 and neuronal stem cells 182 is known to decrease with increasing

mechanical stiffness of the 3D matrix. These results are opposite to reports for

cells seeded on a hydrogel surface in 2D, where the spreading of SMCs, 183

MSCs, 184 and fibroblasts 185 increased with matrix stiffness. This difference in

cell behavior can be explained by the fact that forces transmitted to the cells

inside the 3D matrix differ significantly from forces on the 2D matrix surface.

146, 186 Unlike the 2D surface, the 3D hydrogel network can act as a physical

barrier to the cells, preventing them from proliferating and migrating inside

the stiff hydrogel matrix with a high crosslinking density. 181

The crosslinking density and mechanical stiffness of the photocrosslinked

hydrogels depend on various parameters, including polymer concentration

and molecular weight, 187 light intensity and photoinitiator concentration, 33

and crosslinking time. 185 Generally, when the crosslinking density is increased,

for instance, by decreasing molecular weight or increasing polymer concentra-

tion, a stronger gel with a decreased mesh size and lower swelling and diffu-

sion capacity is formed. 188 The decreased mesh size means a smaller pore size

and reduced permeability of the hydrogel matrix, which leads to impaired

transport of nutrients and metabolites in the gel. Consequently, the lack of

nutrients and free space prevents encapsulated cells from proliferating and

migrating inside the dense gel matrix. Therefore, hydrogels with a lower cross-

linking density combined with controlled biodegradation of the polymer are

typically desired instead of dense, non-degradable gels.

2.4.5 Polymer degradation and matrix remodeling

In TE scaffolds, the rate of polymer degradation is affected by several factors,

including the type and concentration of the hydrolytically labile bonds in the

polymer, the hydrophobicity of the polymer, and the composition and physical

properties of the polymer, as well as the porosity and pore size of the scaffold.

189-191 For crosslinked polymers, the crosslinking density significantly affects

the mechanism of polymer erosion: due to their higher water permeability,

loosely crosslinked networks tend to degrade via bulk erosion, losing their me-

chanical strength with a sudden drop, while highly crosslinked networks with

low water permeability typically exhibit surface erosion, losing their mechani-

cal strength linearly with degradation time. 192 The low crosslinking density

and increased hydrophilicity of a polymer can significantly accelerate network

degradation due to enhanced water intake and swelling of the polymer net-

work. 112, 193

The degradation behavior of photocrosslinked networks also depends on the

photopolymerization method used. When the polymer network is formed

through chain-growth free radical polymerization, the double bonds go

through addition type polymerization and form non-degradable carbon-

carbon chains called kinetic chains. 194 Kinetic chain length is the average

number of double bonds reacting with one radical prior to any chain transfer

or termination reaction, and it is affected by the initiator concentration and

initiation rates. 195, 196 The length of this polyaddition chain is important be-

cause it determines if the chain will be naturally removed from the body. To

prevent accumulation of high molecular weight kinetic chains in the body, the

kinetic chain length can be decreased by adding small amounts of transfer

agents, such as thiols, into the resin. 194, 195 However, the kinetic chain length of

photocrosslinked networks based on low molecular weight polyesters has

shown values significantly lower than the threshold value for renal clearance.

197

The biodegradation of the TE scaffold is especially critical in cell-laden hy-

drogels. To maintain the function of healthy tissue and to repair damaged tis-

sue, cells have the capacity to restructure their surrounding ECM by secreting

matrix metalloproteinases (MMPs) to digest the matrix. 13 In cell-laden hydro-

gels, this remodeling capacity is needed to allow the encapsulated cells to re-

shape the surrounding matrix into a natural tissue-like structure. 13 Since high

mechanical stiffness and a small mesh size can be barriers to cell proliferation

and migration in a 3D hydrogel matrix, the encapsulated cells are required to

remodel the matrix to make more space for their growth. Like other biode-

gradable TE scaffolds, hydrogels can degrade via hydrolysis or enzymatic pro-

teolysis of the polymer chains. When water molecules, alone or supported by

the cellular enzymes, cleave the hydrogel polymer, the crosslinking density

decreases, making the hydrogel softer and more porous. The accelerated deg-

radation leads to a matrix that is easier for cells to remodel, which in turn im-

proves cell spreading and migration in the 3D matrix. 198 A wide variety of

MMP-sensitive PEG hydrogels have been synthesized by adding MMP-

cleavable peptide sequences to the hydrogel polymer, helping cells to degrade

the 3D matrix and consequently increasing their proliferation and spreading

inside the gel. 36, 198, 199

2.4.6 Bioactive glass

Bioactive glass (BG) is an inorganic surface-active biomaterial that stimulates

bone regeneration by releasing soluble ionic species. It subsequently forms a

carbonated hydroxyapatite (HA) layer on its surface in biological fluids. This

HA layer is responsible for the strong interactions between BG and the colla-

gen fibrils of host bone. 200, 201 When exposed to tissue fluid, BG is soon cov-

ered by a silica-rich gel, after which a calcium phosphate-rich HA layer is

formed on top of the gel to promote the attachment of osteoblasts. 202 The HA

layer forms through five rapid stages, starting from fast ion exchange of Na+

and Ca2+ ions with H+ ions between the BG and biological fluid. 203, 204 High

local pH leads to attack of the silica glass by OH- ions, breaking its Si-O-Si

bonds. This creates silanol (Si-OH) groups that condensate near the glass sur-

face, where a silica-rich gel is polymerized. Next, Ca2+ and PO43- ions in the

medium form an amorphous CaO-P2O5 film on the silica-gel, and the film crys-

tallizes to form the carbonated HA layer. Proteins adsorb to this HA layer,

leading to the attachment of cells on the BG surface where they begin for-

mation of ECM. 203

The dissolution products of BG have been observed to stimulate osteopro-

genitor cells at the genetic level by up- and down-regulating genes related to

various cell functions including signal transduction, cell adhesion, and DNA

synthesis. 205 The changes in gene expression regulate the cell cycle towards

osteogenic differentiation and increased osteoblast proliferation. 206 In addi-

tion to increased bone formation, the dissolution products of BG are known to

play an important role in angiogenesis by stimulating fibroblasts to secrete

more vascular endothelial growth factors (VEGF), which is known to guide

neovascularization of new tissue constructs. 207, 208

The most widely studied bioactive silicate glasses are based on the formula 45S5, which is also referred to by its commercial name Bioglass®. 204 Its bio-activity is attributed to its high Na2O and CaO content, high CaO/P2O5 ra-tio, and low SiO2 content compared to chemically durable silicate glasses.

209 BG composition can be modified by doping it with small quantities of ele-

ments such as Cu, Zn, or Mg that are known to support bone growth. 206 For

instance, the incorporation of copper ions into BG has improved both its oste-

ostimulation and angiogenesis capacity as well as antibacterial properties, 210

whereas fluoride has improved its capacity to regulate osteoblast differentia-

tion. 211

In several studies, BG has stimulated more bone regeneration than synthetic

HA. 202, 212, 213 However, synthetic HA and tricalcium phosphate (TCP) are more

widely used in clinical applications due to the challenges in BG scaffold sinter-

ing. 203 To overcome the fragility of BG and the difficulties in its processing, an

attractive approach is to combine BG with biodegradable polymers to obtain

organic-inorganic composites. The polymer acts as glue and makes the BG

material suitable for a wide selection of scaffold fabrication methods. Several

biodegradable polymers have been used to fabricate BG/polymer composites,

including polyesters like polyglycolide, 214 PLA, 215, 216 and PCL. 217-219 These

composite materials mimic the composition of natural bone tissue that con-

sists of organic collagen and inorganic HA phases. 220 In TE applications,

BG/PCL-co-PLA scaffolds have been previously prepared by using thermal

crosslinking or photocrosslinking techniques followed by salt leaching, result-

ing in porous composite structures for cell seeding. 40, 221, 222 More recently,

FDM-based additive manufacturing has been applied to the fabrication of CAD

modeled skull bone reconstructions using BG/poly(methyl methacrylate)

composites 223 and to the fabrication of highly porous TE scaffolds of thermo-

plastic BG/PCL composites. 224, 225

3. Materials and Methods

3.1 Synthesis of photocrosslinkable macromers

Cyclic 3-methyl-morpholine-2,5-dione (MMD) monomer was synthesized by a

two-step reaction, starting from the Schotten-Baumann reaction of L-alanine

with chloroacetyl chloride followed by the cyclization of the resulting chloroa-

cetyl L-alanine (Scheme 1A). Briefly, chloroacetyl chloride dissolved in diethyl

ether was added dropwise to the L-alanine solution in a diethyl ether/water

mixture at -5 °C. Aqueous NaOH solution was added to the reaction solution to

prevent the amino acid from being protonated, which would prevent nucleo-

philic substitution. After extractions and purifications, the resulting cloroace-

tyl L-alanine dissolved in dimethylformamide was reacted in the presence of

triethylamine at 90 °C for 8 h. After solvent removal and purification, the cy-

clic MMD was ready for use in ring-opening polymerization (ROP).

Photocrosslinkable macromers were synthesized through ROP of cyclic

monomers to yield hydroxyl-ended oligomers followed by methacrylation to

yield macromers with double bond end groups as shown in Scheme 1B-D.

Star-shaped PCL oligomers with a targeted Mn of 750 g/mol (PCL750), 1,500

g/mol (PCL1500), 2,250 g/mol (PCL2250), 3,000 g/mol (PCL3000), or 6,000

g/mol (PCL6000) in publications I-III and poly(ε-caprolactone-co-

depsipeptide) (PCL-co-PDP) oligomers with a targeted Mn of 750 g/mol and

the MMD amount of 5 mol% (PDP5) or 10 mol% (PDP10) in publication III

were polymerized using three-armed trimethylolpropane as an initiator and

Sn(Oct)2 as a catalyst. Hydrophilic poly(ethylene glycol-co-depsipeptide)

(PEG-co-PDP) macromer with a targeted Mn of 11,000 g/mol in publication IV

was polymerized using a four-armed PEG as a macroinitiator. Sn(Oct)2-

catalysed ROP follows the coordination-insertion mechanism where the cata-

lyst first reacts with the hydroxyl groups of the initiator to form tin(II)alkoxide

that then acts as an actual initiator in polymerization. 226, 227 Star-shaped oli-

gomers were synthesized to obtain readily crosslinkable macromers. The mo-

lecular weight of the oligomers was determined by the monomer/initiator ratio

in the polymerization. 228 ROP of the star-shaped oligomers was run under an

argon or nitrogen atmosphere at 130 °C for 48 h (PCL in publication I), 140 °C

for 24 h (PCL in publication II), 120 °C for 64 h (PCL-co-PDP in publication

III), or 130 °C for 6 d (PEG-co-PDP in publication IV). The oligomers were

functionalized with an excess of methacrylic anhydride at RT (publications I-

II) or 60 °C (publications III-IV).

Scheme 1. Synthesis of A) cyclic 3-methyl-morpholine-2,5-dione (MMD) monomer, B) meth-

acrylated PCL macromer, C) methacrylated PCL-co-PDP macromer, and D) methacrylated PEG-

co-PDP macromer.

3.2 Synthesis of water-soluble photoinitiator

For photocrosslinking of PEG-co-PDP hydrogels in publication IV, a water-

soluble, visible light sensitive photoinitiator was required. Lithium phenyl-

2,4,6-trimethylbenzoylphosphinate (LAP) photoinitiator was synthesized as

previously described by Fairbanks et al. 229 Briefly, 2,4,6-trimethylbenzoyl

chloride was added dropwise to an equimolar amount of dimethyl phe-

nylphosphonite at RT under nitrogen atmosphere. After mixing for 18 h, a

fourfold excess of lithium bromide dissolved in 2-butanone was added to the

mixture, and the reaction was continued for 10 min at 50 °C. After heating, the

mixture was kept at RT for 4 h; the resulting precipitate was filtered and

washed once with 2-butanone, then twice with diethyl ether, and finally dried

under vacuum.

B) Methacrylated PCL homopolymer I-III

R' O

R'

OO

H

n

OH

HO

HO

O

+

R'' O

R''

OO

n

+ Sn(Oct)2

OOO+

+ Sn(Oct)2 O

OO

HN H

n

OOO+

R

R

R

OO

m

H

NH

O O

O+

R'

R'

R'

OO

m

O

OO

HN

nR''

R''

R''

OO

m

O

O

R' O

R' O

OO

HN O

O

H

m n

R'' O

R'' O

OO

HN O

O

nm

NH

O O

OO

O+OH

HO

HO+

OOO+

+ Sn(Oct)2

C) Methacrylated PCL-co-PDP copolymer III

D) Methacrylated PEG-co-PDP copolymer IV

OO

ClHN

OHO

O

ClCl

OH2N OH

O+

+ NaOH + TEA

NH

O O

O

A) Cyclic MMD monomer III, IV

3.3 Photocrosslinking

To photocrosslink polymer films for physicochemical characterization, PCL-

based macromers were mixed with 2 mol% of UV sensitive Irgacure 2959

(publication I) or 0.5 wt% of visible light sensitive camphorquinone (CQ) and

0.5 wt% of ethyl-4-N,N-dimethylaminobenzoate (EDMAB) (publication III).

Visible light sensitive resins used in SLA contained 5 wt% of Irgacure 369

(publication I) or 3 wt% of Lucirin® TPO-L (publications II and III). Hydrogels

in publication IV were prepared by mixing PEG-co-PDP solution (10% w/v)

with visible light sensitive LAP photoinitiator (0.25% w/v). PCL films in publi-

cation I were crosslinked for 30 min with UV light and PCL-co-PDP films in

publication III were crosslinked for 2 min using visible light. Scheme 2 shows

the mechanism of the photoinitiators used in this thesis. Photoinitiators other

than CQ/EDMAB underwent a photocleavage reaction, resulting in the scis-

sion of the bond between the carbonyl carbon and α-carbon. 142, 229 The

CQ/4EDMAB system, in turn, generated radicals by hydrogen abstraction,

where the amine interacted with the activated triplet state of the CQ carbonyl

to form the reactive radical. 104, 130 It has been stated that the CQH• radical

does not effectively initiate polymerization due to steric hindrance, but the

photocrosslinking is mainly initiated by the other, aminoalkyl radical. 230

Scheme 2. Radical formation of different photoinitiators under light exposure.

A) Irgacure 2959 I

1-[4-(2-Hydroxyethoxy)-

phenyl]-2-hydroxy-2-

methyl-1-propanone

D) LAP IV

Lithium phenyl-2,4,6-

trimethylbenzoylphosphinate

B) Irgacure 369 I

2-Benzyl-2-(dimethylamino)-

1-[4-(4-morpholinyl) phenyl]-

1-butanone

C) Lucirin TPO-L II, III

Diphenyl (2,4,6-

trimethylbenzoyl)-

phosphine oxide

P

O-

O

OLi+

P

O-

O

O

HOO

OHO

HOO OH

UV light

Visible lightLi+

NO N

OVisible light

NO

O

N

PO

O

O

Visible light O

PO

O

O

E) CQ + 4EDMAB III

1,7,7-Trimethylbicyclo[2.2.1]-

heptane-2,3-dione + ethyl-4-

N,N-dimethylaminobenzoate

ON O

O

O +

R

R

O

R

R

O* RN+

R

R

O-

RN

RNR

R

OH +Visible light

H2C

To characterize the degree of crosslinking of photocrosslinked samples, both

the gel content and double bond conversion of the macromers were analyzed.

For gel content studies, three photocrosslinked dry samples were weighed (m0)

and extracted with acetone (PCL-based polymers) or water (PEG-based poly-

mers), and then dried in vacuum until constant weight (md). Gel content G was

calculated using Equation (4). 231 Double bond conversion was studied with

FTIR analysis.

G = md / m0 × 100% (4)

The swelling degree of crosslinked samples was measured by weighting three

extracted and dried polymer samples (md), then immersing them in chloro-

form (PCL-based samples) or water (PEG-based samples) overnight. Swollen

samples were weighed after removal from the medium (msw). The density of

the photocrosslinked polymers (ρp) was measured using an analytical balance

at room temperature, while the density of the solvent (ρs) was 1.48 g cm-3

(chloroform) or 1.00 g cm-3 (water). The swelling degree Q was calculated us-

ing Equation (5): 231

Q = 1 + ρp × (msw/(md × ρs) – 1/ ρs) (5)

The in vitro degradation of photocrosslinked polymers in publications III

and IV was measured by following the mass loss of the photocrosslinked sam-

ples incubated in PBS (pH 7.4) at 37°C with shaking at 150 rpm. After a prede-

termined time, samples were removed from the media and subsequently

washed with deionized water and dried in a vacuum.

3.4 Stereolithography of TE scaffolds and hydrogels

The mathematically defined porous scaffolds in publications I-III were pre-

pared using an EnvisionTEC Perfactory® Mini Multilens SLA machine based

on visible light DLP technology. A diamond and gyroid pore architecture

shown in Figure 4 were modeled using the free K3DSurf surface generator

(k3dsurf.sourceforge.net), where the diamond (publication I) and gyroid (pub-

lications II and III) designs were obtained by generating the 3D surfaces de-

scribed by Equation (6) and Equation (7), respectively. The constant C was

varied to change the porosity of the scaffolds. These mathematically defined

pore architectures were selected to provide the scaffolds with a large surface

area and high pore interconnectivity, allowing for high cell seeding efficiency

and free nutrient flow inside the scaffolds.

f(x,y,z) = sin(x)sin(y)sin(z)+sin(x)cos(y)cos(z)

+cos(x)sin(y)cos(z)+cos(x)cos(y)sin(z)-C (6)

f(x,y,z) = cos(x)sin(y)+cos(y)sin(z)+cos(z)sin(x)-C (7)

Figure 4. Scaffold designs with a diamond (D) and gyroid (G) pore architecture.

The STL files for the porous scaffolds were prepared using the Rhinoceros®

software (McNeel). In publications II and III, the gyroid design was rotated to

make the pore channels open from corner to corner instead of top to down. 3D

designs were cut and scaled to meet the desired dimensions, and EnvisionTEC

Perfactory® software was used for finishing and slicing the STL files.

In publications I-III, porous 3D scaffolds were built using a photocrosslinka-

ble resin consisting of methacrylated PCL or PCL-co-PDP macromer mixed

with Irgacure 369 (5 wt%) photoinitiator and vitamin E (o.o6 wt%) inhibitor

(publication I) or Lucirin® TPO-L photoinitiator (3 wt%) (publications II and

III). Orasol Orange G dye (0.10 wt%) was used to prevent over-crosslinking

and thus improve the resolution of the scaffolds. Bioactive glass in publication

II was mixed with the PCL750 resin using 0, 5, 10, or 20 wt% of BG (BG0,

BG5, BG10, and BG20, respectively). In publication I, the solid resin

(PCL1500) was electrically heated above the melting point of the macromer

(43–46 °C) to decrease its viscosity. In publications II and III, the resins

(PCL750 and PCL-co-PDPs) were liquid at RT due to the low molecular weight

of the macromers and no heating was needed. In publications I-III, the 3D

modeled hydrophobic scaffolds were fabricated with the DLP-based Envision-

TEC Perfactory® SLA using blue light with a light intensity of 1600 mW/dm2.

In publication I, the layer thickness of 25 μm was crosslinked for 10 s/layer

and the xy-resolution was 32 μm. In publication II, the thickness was 50 μm,

the crosslinking time 12 s, and the resolution 42 μm. In publication III, respec-

tively, the thickness was 50 μm, the crosslinking time 12 s, and the resolution

32 μm. After SLA, the uncured resin was removed from the scaffolds by sol-

vent extraction, after which the scaffolds were dried in a vacuum.

To fabricate both cell-free and cell-laden hydrogels by SLA in publication IV,

10% w/v of the photocrosslinkable PEG-co-PDP macromer and 0.25% w/v of

the LAP photoinitiator were dissolved in EBM-2 medium. To prevent cells

from sedimenting in the polymer solution, the viscosity of the solution was

increased by adding 7% w/v of Percoll® medium to the solution. To link an

adhesion peptide (RGDS) to the hydrogel, the RGDS peptide was first coupled

with acryl-PEG-SVA (3,400 g/mol) in 50 mM NaHCO3 solution (pH 8.3) for 4

h at RT. The resulting acryl-PEG-RGDS was then dissolved in the hydrogel

solution to yield a 2 mM concentration. In cell-laden hydrogels, the hydrogel

solution was filtered and mixed with a HUVEC suspension to obtain a hydro-

gel solution with 5×106 cells/mL. The hydrogels were photocrosslinked in a

layer-by-layer manner using an in-house built visible light DLP-based SLA

(190 W, 3200 lumens). The layer thickness of 200 μm was crosslinked with

light exposure times of 100 s, 120 s, or 160 s/layer.

3.5 Imaging and mechanical testing of the 3D structures

The PCL scaffolds in publication I were visualized by scanning electron mi-

croscopy (SEM) employing a Philips XL30 device with 5.0 kV electron beam.

Characteristics of the pore network were determined by micro-computed to-

mography (μ-CT) using a GE eXplore Locus SP scanner equipped with an X-

ray tube operating at 90 kV and 80 μA. The resolution was 14 μm.

In publication II, the compression moduli of the PCL/BG scaffolds were

measured with a stress controlled rotational rheometer AR G2 (TA instru-

ments) using 20 mm plate-to-plate geometry. The samples were discoid scaf-

folds 20 mm in diameter and 3 mm in height. For comparison, similar nonpo-

rous PCL samples were also prepared and tested. Compression was performed

at a temperature of 23 °C and a speed of 10 μm/s until the stress normal was

45 N. To test wet samples, the scaffolds were immersed in SBF at 37 °C for 24

h before measurements. The water uptake of the scaffolds was defined as the

relative increase in the weight of the polymer fraction. A micro-CT system op-

erated by Scanco Medical AG was used to analyze the 3D morphology of the

scaffolds at a 6.6 μm voxel resolution.

To visualize the PCL-co-PDP scaffolds in publication III, samples were sput-

tered with a gold/palladium alloy and observed with a scanning electron mi-

croscope (SEM, FEI XL30 Sirion) at 5 kV. Mechanical properties of the photo-

crosslinked scaffolds were measured using an Instron 5944 mechanical tester

with a 2 kN load cell. A scaffold was placed between two parallel plates and

compressed at a rate of 1% strain/second using a 1 N preload.

In publication IV, the in vitro mass loss of the hydrogels was monitored by

measuring the dry weight of the samples after incubation in PBS (pH 7.4, with

0.02% sodium azide) at 37 °C. Triplicate samples were removed from the PBS,

weighed, and dried. After drying, the remaining mass was divided by the initial

mass to get the remaining mass percentage. Viscoelastic behavior of the 3D

fabricated hydrogels was analyzed with an Ares G2 rheometer (TA Instru-

ments) using 8-mm parallel plate geometry. The angular frequency sweep was

run from 0.1 to 10 rad/s at 0.5% strain and 25 °C. AFM force-distance meas-

urements were done with an Agilent 5500 AFM in a liquid cell with PBS. Tip

sensitivity calibration was performed on a glass surface at the end of each set

of experiments. The tip used for the experiments (NanoAndMore, k = 0.08

N/m, R = 950 nm) had a SiO2 sphere on its apex. FDS curve data analysis was

done using commercially available SPIP software from Image Metrology A/S.

In this thesis, the air-dried, Pd/Au-coated PEG-co-PDP hydrogels were imaged

by SEM (Zeiss Sigma series) at 4kV.

3.6 Cell cultures

For cytotoxicity and proliferation tests, several different cell lines were cul-

tured in appropriate cell culture medium. In publication I, NIH3T3 fibroblasts

were cultured in α-MEM medium supplemented with 10% fetal bovine serum

(FBS), 1% penicillin-streptomycin antibiotics, and 1% basic fibroblast growth

factor. In publication II, human gingival fibroblasts were cultured in DMEM

medium supplemented with 10% FBS and 100 U/μg penicillin-streptomycin

antibiotics. In publication III, C3H10T1/2 pluripotent mouse mesenchymal

cells were cultured in DMEM with 10% FBS and 1% antibiotic solution. In pub-

lications III and IV, human-umbilical-vein endothelial cells (HUVECs) were

cultured in EBM-2 medium supplemented with Single Quots (EGM-2) endo-

thelial growth supplement, 10% FBS, and 1% antibiotic solution. All cells were

cultured under standard conditions (5% CO2, 37 °C, 95% humidity).

Before cell seeding, photocrosslinked PCL-based samples were sterilized by

immersing them in ethanol and washing with PBS. Cytocompatibility of the

sterilized samples was evaluated by measuring the metabolic activity of the

seeded cells using either a colorimetric MTS assay (publication I), a colorimet-

ric AlamarBlue® assay (publication II), or a colorimetric MTT assay (publica-

tion III). Proliferation of the seeded pluripotent C3H10T1/2 cells in publication

III was monitored by measuring dsDNA content using a Pico Green assay, and

their osteogenic activity was evaluated by measuring their alkaline phospha-

tase specific activity using a colorimetric p-nitrophenyl phosphate (p-NPP)

method. In publication IV, the proliferation of HUVECs seeded or encapsulat-

ed in hydrogels was measured using a colorimetric AlamarBlue® assay.

4. Results and Discussion

4.1 Synthesis of the monomer, oligomers, and macromers

The cyclic MMD monomer used for ROP of PCL-co-PDP and PEG-co-PDP

copolymers was synthesized via a chloroacetyl alanine intermediate. In 1H

NMR, the formation of chloroacetyl alanine from chloroacetyl chloride and L-

alanine was seen as a shift of the quartet peak 2 from 3.66 ppm (Figure 5A) to

4.31 ppm (Figure 5B), and as the appearance of the singlet peak 1 at 4.05 ppm

(Figure 5B). Further formation of the cyclic MMD monomer was observed as a

shift of the singlet peak 1 at 4.05 ppm (Figure 5B) to two doublet peaks at

4.70–4.86 ppm (Figure 5C). The melting points of the resulting chloroacetyl

alanine and MMD were 97 °C and 151°C, respectively, which are comparable to

previously published values of 93 °C–96 °C for chloroacetyl alanine and 153.5

°C–154.5 °C for MMD. 232

Figure 5. 1H NMR spectra of A) L-alanine, B) chloroacetyl alanine, and C) cyclic MMD mon-

omer in D2O.III

Star-shaped hydroxyl-terminated PCL, PCL-co-PDP, and PEG-co-PDP oli-

gomers were synthesized through ROP of cyclic monomers. ROP of PCL and

PCL-co-PDP oligomers resulted in peaks a–e and 1–3 in the 1H NMR spectra,

as assigned in Figure 6A and 6B. After ROP, no remaining monomer peaks

were seen, while a small peak f’ attributed to the unreacted arms of TMP initia-

tor was detected. The hydrophilic PEG-co-PDP copolymer was synthesized by

ROP of the MMD monomer using a 4-arm PEG as an initiator. The ROP re-

sulted in a transfer of the monomer peaks 1-3 in Figure 5C to the correspond-

ing oligomer peaks in Figure 6C, and the appearance of peak b’ attributed to

the last CH2 groups of the PEG arms preceding the PDP units. Small MMD

peaks remained in the 1H NMR spectra despite the extended reaction time,

which presumably resulted from side reactions of ROP, including intramolecu-

lar back-biting and intermolecular transesterification. 233, 234 The residual

monomer was removed by dialysis and freeze-drying to ensure high purity of

the oligomer.

Figure 6. 1H NMR spectra of the oligomers in CDCl3: A) PCL,I B) PCL-co-PDP,III and C)

PEG-co-PDP.IV

To obtain the macromers with photocrosslinkable double bonds, the hydrox-

yl-ended oligomers were functionalized with methacrylic anhydride. The suc-

cessful methacrylation of PCL and PCL-co-PDP oligomers was seen as a shift

of the peak a’ attributed to the last PCL unit preceding the hydroxyl end-group

and as the appearance of new peaks i, j, and k attributed to the methacrylate

end-groups (Figure 7A-7C). Methacrylation was continued until no peak a’ of

the oligomer at 3.65 ppm was seen in 1H NMR spectra, indicating near 100%

substitution. PCL macromers with a molecular weight over 1500 g/mol were

successfully purified by precipitation in cold isopropanol (publication I), while

the ones with a lower molecular weight were soluble in polar isopropanol.

Therefore, in publication II the low molecular weight PCL macromer was

washed with saturated sodium bicarbonate solution, and in publication III, the

PCL and PCL-co-PDP macromers were precipitated in non-polar hexane. The

photocrosslinkable macromers were purified until no methacrylic anhydride

peaks were detected around 2.04, 5.85, or 6.26 ppm. The small peaks observed

near the peaks i, j, and k in Figure 7B and 7C were attributed to the methacry-

late end-groups of the previously unreacted TMP arms, which are typical for

the low-molecular weight star-shaped polymers. The successful methacryla-

tion of the PEG-co-PDP oligomer was seen as a complete disappearance of

peak 1 in Figure 6C attributed to the last CH2 group preceding the hydroxyl

end-groups of the oligomer and an appearance of peaks i, j, and k in Figure 7D

attributed to the methacrylate groups. The dialyzed and freeze-dried

macromer was a light-brown powder showing great water-solubility at room

temperature.

Figure 7. 1H NMR spectra of the macromers in CDCl3: A) PCL1500,I B) PCL750,II C) PCL-co-

PDP,III and D) PEG-co-PDP.IV

The molecular weights (Mn) of the synthesized oligomers were calculated us-

ing 1H NMR integration and are listed in Table 1. For the PCL homopolymer

and the PCL-co-PDP copolymers, the average number of CL and MMD units

per oligomer was determined by comparing the peak integral of the TMP me-

thyl group (peak h) with the integral of the protons in the CL units (peak e)

and in the MMD units if any (peaks 1-2). The molecular weight of the PEG-co-

PDP oligomer was calculated by comparing the integrals of the protons in the

PEG macroinitiator (peak f ) and the protons in the PDP units (peak 3).

Table 1 also lists the thermal properties of the synthesized oligomers and

macromers. As the molecular weight of the PCL oligomers increased, the melt-

ing temperature (Tm) approached that of a linear high molecular weight PCL

(60 °C). The melting endotherms consisted of two peaks, which have been pre-

viously reported for both branched and linear polymers and has been ex-

plained by the distribution in length of the polymers, causing variation in the

size of the polymer crystallites. 235, 236 The molecular weight and thermal prop-

erties of the PCL750 polymers in different publications varied slightly due to

different synthesis conditions. For the low molecular weight PCL-co-PDP co-

polymers, no melting peaks were detected, indicating their amorphous charac-

ter. An increasing amount of MMD in copolymers increased the glass transi-

tion temperature (Tg), suggesting that the depsipeptide units made the struc-

tures of the copolymers more rigid when compared to the PCL homopolymer.

Unlike PCL macromers with a higher molecular weight, the PCL750 and PCL-

co-PDP macromers were liquid at room temperature, which was highly desired

because it allowed them to be processed in SLA without the use of heating or

solvents. Copolymerization of four-arm PEG with the depsipeptide decreased

the melting point of the oligomer when compared to the PEG homopolymer,

suggesting that the depsipeptide units interfered with the highly crystalline

structure of PEG.

Table 1. Molecular weights and thermal properties of the oligomers and macromers.

Oligomer Macromer

Oligomer name

Targeted Mn (g/mol)

1H NMR Mn

(g/mol)

DSC DSC Tg (°C)

Tm (°C)

ΔH (J/g)

Tg (°C)

Tm (°C)

ΔH (J/g)

State at RT

PCL750 I 750 810 -67 24 39 nd nd nd nd PCL1500 I 1,500 1,550 – 33 58 -62 36 54 Waxy PCL2250 I 2,250 2,680 – 44 61 -64 44 57 Solid PCL3000 I 3,000 4,200 – 49 65 -65 49 57 Solid PCL6000 I 6,000 6,030 – 54 64 – 54 60 Solid PCL750 II 750 702 nd nd nd -61 15 nd Liquid PCL750 III 750 741 -65 – – -64 – – Liquid PCL-PDP5 III 750 738 -60 – – -59 – – Liquid PCL-PDP10 III 750 731 -55 – – -56 – – Liquid PEG IV 10,000a 10,000a – 56 160 – 44 101 Solid PEG-PDP IV 11,000 10,850 – 45 81 – 43 77 Solid

nd = not determined; a as given by the chemical producer

In addition to 1H NMR analysis, copolymerization of both PCL and PEG with

depsipeptide units was monitored with FTIR. Figure 8A shows that for FTIR of

the PCL-co-PDP macromer, transmittance near 1540 cm-1 decreased with in-

creasing amount of MMD in the copolymerization. This decreased transmit-

tance was attributed to the peptide bonds (N-H bends) of the macromer, re-

vealing the successful introduction of depsipeptides into the PCL homopoly-

mer structure. The FTIR data were consistent with data from the 1H NMR

analysis. FTIR analysis also confirmed the copolymerization of PEG with the

depsipeptide units. Figure 8B shows the FTIR spectra of PEG before and after

copolymerization with MMD. ROP of PEG with MMD resulted in new peaks at

1750 cm-1 and 1670 cm-1 attributed to the C=O ester and C=O amide stretches,

respectively, and a new peak at 1540 cm-1 attributed to the N-H bend of the

peptide bonds in the copolymer. These FTIR data together with the 1H NMR

data confirmed that the intended PEG-co-PDP copolymer was successfully

synthesized.

Figure 8. A) FTIR curves attributed to the peptide bonds of the PCL-co-PDP macromersIII

and B) FTIR spectra of the PEG homopolymer and the PEG-co-PDP copolymer showing the new

ester and amide bonds of the copolymer.IV

4.2 Photocrosslinking characteristics

Before using the synthesized macromers in SLA, their photocrosslinking ca-

pacity was analyzed by mixing them with photoinitiators and photocrosslink-

ing polymer samples under exposure to UV or visible light. The crosslinking

density of photocrosslinked samples was analyzed by measuring the gel con-

tent and swelling degree of the crosslinked samples. Table 2 shows that the gel

content of the photocrosslinked PCL samples decreased from 99.8% to 89.8%

when the molecular weight was increased from 750 g/mol to 6000 g/mol,

while the swelling degree increased from 1.67 to 3.60. The gel content and

swelling degree values revealed that the short macromers formed a highly

crosslinked network, while the macromers with longer arms resulted in more

loosely crosslinked networks with increased solvent penetration to the poly-

mer network. Copolymerization of PCL with PDP only slightly decreased the

gel content when compared with the PCL homopolymer. Photocrosslinking of

PEG-co-PDP hydrogels (10% w/v) by SLA for extended time (10 min/layer)

revealed their maximum gel content to be lower than the gel content of the

PCL-based samples. This is due to the higher molecular weight and lower pol-

ymer concentration of the PEG-co-PDP samples compared with the PCL-based

samples. The swelling degree of PEG-co-PDP hydrogels in water was high due

to the highly hydrophilic nature of the crosslinked PEG-based polymers.

Table 2. Gel content and swelling degree of the photocrosslinked polymers.

Polymer Gel content (%) Swelling degree PCL1500 I 98.2 ± 0.5 2.04 ± 0.03 a PCL2250 I 95.4 ± 1.0 2.63 ± 0.03 a PCL3000 I 93.1 ± 0.6 2.97 ± 0.03 a PCL6000 I 89.8 ± 0.9 3.60 ± 0.04 a PCL750 III 99.8 ± 0.1 1.67 ± 0.01 a PCL-co-PDP5 III 99.2 ± 0.3 1.60 ± 0.01 a PCL-co-PDP10 III 98.3 ± 0.3 1.65 ± 0.01 a PEG-co-PDP c 89.4 ± 0.3 15.2 ± 0.3 b

a Swelling in chloroform; b swelling in water; c supports data in publication IV

Besides the gel content, which describes the ratio of macromers that are at-

tached to the polymer network with at least one arm and will not leach out of

the polymer, the double bond conversion of PCL-co-PDP samples was ana-

lyzed with FTIR. Figure 9 shows the FTIR peaks near 1640 cm-1 attributed to

the methacrylate double bonds of the PCL-co-PDP macromer. Before photo-

crosslinking, clear peaks were observed due to the unreacted double bonds.

After crosslinking, the peaks disappeared indicating that the double bonds

were successfully reacted. The FTIR data together with the high gel content of

the PCL-co-PDP samples confirmed that the resulting polymer networks were

highly crosslinked and the number of unreacted double bonds of the polymers

was very low after photocrosslinking.

Figure 9. FTIR transmittance curves attributed to the methacrylate double bonds of PCL-co-

PDP polymers before (Macromer) and after photocrosslinking (Network).III

To study the effect of molecular weight, and thereby crosslinking degree, on

the mechanical properties of photocrosslinked PCL films, the stress–strain

behavior of films was characterized by tensile testing. Despite the molecular

weight, all samples showed elastic deformation until failure of the specimen

without the occurrence of a yield point. Table 3 lists the average Young’s mod-

ulus, maximal tensile strength, and elongation at the break of the photocross-

linked PCL films. As the crosslinking density of the polymers decreased with

increasing molecular weight of the macromers, the Young’s modulus de-

creased and the elongation at break increased. Thus, increasing molecular

weight and the decreased crosslinking density made the photocrosslinked PCL

a less rigid and more elastic material.

Table 3. Tensile properties of photocrosslinked PCL films.I

Polymer Young's modulus (MPa)

Tensile strength (MPa)

Elongation at break (%)

PCL1500 15.4 ± 0.7 2.55 ± 0.12 19.3 ± 0.5 PCL2250 10.0 ± 0.2 2.49 ± 0.04 30.8 ± 0.1 PCL3000 7.1 ± 0.2 2.02 ± 0.13 38.2 ± 4.5 PCL6000 6.7 ± 0.4 2.87 ± 0.12 78.5 ± 6.7

4.3 In vitro biodegradation of photocrosslinked films

The degradation rate of photocrosslinked PCL and PCL-co-PDP copolymers

was evaluated by measuring the mass loss of the polymer films in PBS. Figure

10 shows that within six months in PBS, PCL and PCL-PDP5 samples lost ap-

proximately 8% of their initial mass, whereas PCL-PDP10 samples lost over

15%. The water contact angle measurements of the crosslinked films revealed

that PDP10 films were initially less hydrophobic (contact angle 75° ± 1°) com-

pared with PDP5 films (85° ± 1°) and PCL films (87° ± 1°). Therefore, the more

hydrophilic surface and lower gel content of PDP10 samples contributed to

their increased mass loss compared with PCL and PDP5 samples. The linear

mass loss and intact shape of the films indicated that the degradation occurred

via surface erosion, which was consistent with a previous study192 showing that

highly crosslinked polymer networks undergo surface erosion both in vitro

and in vivo. The dense hydrophobic polymer network restricted water diffu-

sion into the polymer matrix, making the polymer chains on the matrix surface

more susceptible to hydrolysis than the chains inside, thus resulting in surface

erosion.

Figure 10. In vitro degradation of photocrosslinked PCL and PCL-co-PDP films.III

4.4 Cytocompatibility of the polymers

The cytocompatibility of photocrosslinked polymers was evaluated by seeding

cells on a polymer surface and quantifying their metabolic activity using a col-

orimetric cell proliferation assay. Figure 11A shows that for the photocross-

linked PCL samples, the optical density, which is proportional to the metabolic

activity and number of living cells, increased after day 1 and was almost identi-

cal to the positive tissue culture polystyrene (TCPS) control sample at day 3.

Due to confluence, the seeded fibroblasts had no space for further growth, re-

sulting in decreased activity at day 7. For PCL-co-PDP films, Figure 11B shows

that the activity of HUVECs increased at every time point within 7 days of cell

culturing. At day 7, the optical density was around 10-fold when compared to

the value at day 1, indicating that the cells proliferated actively on the polymer

films.

To study in vitro cytocompatibility of the methacrylated PEG-co-PDP poly-

mer and compare it to the methacrylated PEG control polymer, PEG, PEG-co-

PDP, and PEG-co-PDP/RGDS hydrogels were photocrosslinked and HUVECs

(1×103) were seeded on the samples. Figure 11C shows that between days 4 and

14, metabolic activity significantly increased on PEG-co-PDP and PEG-co-

PDP/RGDS hydrogels, while no significant increase was seen on PEG gels. At

days 7 and 14, cell activity was significantly higher on PEG-co-PDP gels com-

pared to PEG gels, while the highest activity was detected on PEG-co-

PDP/RGDS hydrogels. Because the non-patterned PEG surface suppresses

non-specific protein adsorption and is anti-adhesive due to its lack of func-

tional groups, 237 improved cell proliferation on the hydrogel surface was sug-

gested to result from new functional groups introduced by the depsipeptide

units to the hydrogel polymer. RGDS adhesion peptides further improved cell

proliferation by providing integrin-specific attachment sites on the hydrogel

surface. The acrylated PEG with a molecular weight of 3,400 g/mol was suc-

cessfully used as a spacer between the hydrogel matrix and RGDS peptide,

coupling the peptide to the hydrogel and freeing the immobilized peptide from

steric hindrance. 238

Figure 11. Cytotoxicity studies of photocrosslinked polymers: A) MTS of NIH3T3 fibroblasts

on PCL1500 films,I B) MTT of HUVECs on PCL750 and PCL-co-PDP films,III C) AlamarBlue® of

HUVECs seeded on PEG, PEG-co-PDP, and PEG-co-PDP/RGDS hydrogels.IV

To test the suitability of photocrosslinked PCL and PCL-co-PDP polymers for

bone tissue engineering, mouse pluripotent C3H10T1/2 cells were seeded on

the polymer films, and cell proliferation was assessed by measuring the

amount of dsDNA within 14 days of cell culturing. Figure 12A shows that all

samples supported a greater than two-fold increase in the amount of dsDNA in

cells between days 3 and 14, indicating active cell proliferation on the photo-

crosslinked samples. It has been shown that C3H10T1/2 cells have the poten-

tial to differentiate into osteoblasts, which can be detected by measuring the

cell’s ALP activity. 239, 240 Here, Figure 12B shows that ALP activity was low

within 7 days of cell culturing and increased significantly before day 14. The

increased ALP activity suggested that the cells went through osteogenic differ-

entiation on the polymer films, indicating that the PCL-co-PDP polymers

would be particularly suitable to bone tissue engineering.

Figure 12. A) Cellular dsDNA content and B) ALP activity of C3H10T1/2 cells on photocross-

linked PCL and PCL-co-PDP films, and on a TCPS control.III

4.5 SLA fabricated scaffolds and hydrogels

4.5.1 Surface morphology

To fabricate porous mathematically defined TE scaffolds by SLA, a liquid resin

was needed. The resin consisting of the PCL1500 macromer was solid at room

temperature and required heating during scaffold fabrication. At 36 °C, just

above the melting point, the viscosity of the resin was 1130 cP and decreased to

660 cP at 46 °C. Scaffolds were built by SLA at 43–46 °C due to the suitable

viscosity of the resin in that temperature range. The PCL750 and PCL-co-PDP

resins used for scaffold fabrication were liquid at RT because of their low mo-

lecular weight, and therefore no heating or solvents were needed. The addition

of bioactive glass to the PCL resin still allowed for successful scaffold fabrica-

tion by SLA. A small amount of orange dye was added to the polymer resins to

prevent over-cure, thus improving scaffold resolution. Since no additional sol-

vents were used during scaffold fabrication, the final dimensions of the scaf-

folds were as designed and no material shrinkage was observed. PEG-co-PDP

hydrogels were 3D fabricated by SLA using a 10% w/v polymer solution and a

visible light sensitive photoinitiator. Hydrogels were photocrosslinked at room

temperature using three different light exposure times: 100 s, 120 s, and 160

s/layer. With all light exposure times, the resulting hydrogels closely resem-

bled the shape and dimensions designed in their CAD models.

Scanning electron microscopy (SEM) images in Figure 13A and 13B revealed

the smooth and uniform surface of photocrosslinked PCL scaffolds, while

PCL/BG scaffolds in Figure 13C and 13D demonstrated an uneven and rough

surface. BG particles were homogenously distributed throughout the PCL/BG

scaffold surface, where they remained uncovered and readily available for rap-

id ion release and interactions with cells. These findings highlighted the great

advantage of SLA over other scaffold fabrication techniques where ceramic

particles tend to be covered by a polymer layer on the scaffold surface. The

SEM image of a PCL-co-PDP scaffold in Figure 13E revealed that the gyroid

pore architecture of the scaffold closely resembled its mathematically defined

3D model. The cross-section of an air-dried PEG-co-PDP hydrogel in Figure

13F showed the hydrogel matrix to be virtually non-porous in macroporosity,

as designed in its CAD file.

Figure 13. A-B) PCL1500 scaffolds,I C-D) PCL-BG20 scaffolds,II E) a PCL-co-PDP scaffold,III

and F) a cross-section of an air-dried PEG-co-PDP hydrogel.

4.5.2 Pore size distribution and pore interconnectivity

μCT reconstruction of 3D fabricated PCL and PCL/BG scaffolds revealed an

open and interconnected pore structure of the scaffolds (Figure 14A-14D).

Pore size analysis revealed a narrow pore size distribution for both PCL1500

scaffolds (Figure 14E) and PCL750 and PCL/BG scaffolds (Figure 14G). In the

PCL1500 scaffolds, more than 90% of the pores ranged from 350 to 550 μm in

diameter. The average pore size was 465 μm, and the average porosity was 71

%, very close to the designed porosity of 70%. In the pCL750 and PCL/BG scaf-

folds, the mean pore sizes were 594 μm (BG0), 555 μm (BG5), 517 μm (BG10),

and 476 μm (BG20), revealing that increasing the amount of BG decreased the

average pore size of scaffolds. The total porosity of the scaffolds decreased

from 77 % (BG0) to 75 % (BG5), 70 % (BG10), and 63 % (BG20), which were

consistent with decreasing pore size. The lower porosity of composite scaffolds

compared with neat PCL scaffolds can be explained by the higher viscosity of

PCL/BG resins. Uncured resins with higher viscosity could not freely flow out

from the pores during scaffold fabrication, resulting in over-cure and lower

porosity. However, the mean pore size remained approximately 500 μm in all

scaffolds, which is considered optimal for bone formation. 22

The interconnectivity of pores was characterized by estimating the accessible

pore volume by μCT reconstruction. The accessible pore volume represents the

volume that can be reached from the outside of the scaffold by simulating the

permeation of a spherical particle. 120 Figure 14F and 14H show that there was

high accessible pore volume for 400 μm spheres in all scaffolds, indicating

great interconnectivity of the pore networks. High permeability of the scaffolds

is essential for the successful accommodation of cells and efficient transport of

nutrients and metabolites inside the 3D structure. The use of SLA allowed for

fabrication of highly permeable scaffolds, which can be considered a signifi-

cant advantage of the SLA technique when compared to conventional scaffold

fabrication methods where the pores typically remain unconnected.

Figure 14. A) A photograph and B) a μ-CT image of a PCL1500 scaffoldI; C) a μ-CT image of a

PCL750 scaffold and D) a PCL750-BG20 composite scaffold.II Below: Pore size distribution and

pore interconnectivity of E-F) a PCL1500 scaffoldI and G-H) a PCL750 and PCL/BG scaffolds.II

4.5.3 Mechanical properties of porous scaffolds

Besides pore properties, mechanical properties are also a key criterion in de-

signing a TE scaffold. The effect of BG on the mechanical strength of PCL/BG

scaffolds was evaluated by compression tests of dry and wet scaffolds. Figure

15A shows that the addition of 20 wt% of BG (BG20) in dry conditions in-

creased the compression modulus from 1.4 to 3.4 MPa when compared to PCL

(BG0) scaffolds. There was no statistically (p < 0.01) significant difference

between BG5 and BG10 scaffolds, while both were stronger than PCL scaffolds.

Wet BG20 scaffolds had an increased compression modulus from 1.1 to 2.5

MPa when compared to wet PCL scaffolds. There were no other statistically

significant differences between wet BG5, BG10, and BG20 scaffolds. For com-

parison, the compression modulus of nonporous PCL samples prepared by

SLA was 10.4 ± 0.8 MPa for dry samples and 10.3 ± 0.6 MPa for wet samples.

Water uptake of the scaffolds was 150 ± 7% (BG0), 160 ± 10% (BG5), 176 ± 8%

(BG10), and 190 ± 10% (BG20), while for nonporous PCL samples it was only

1.66 ± 0.02%. The low water uptake of nonporous PCL samples indicated their

highly hydrophobic nature. BG increased the water uptake of scaffolds, indi-

cating the more hydrophilic nature of the composite materials.

The effect of depsipeptide on the mechanical properties of PCL scaffolds was

also measured with compression tests. Like PCL/BG composite scaffolds, PCL-

co-PDP scaffolds showed a long linear stress–strain curve before yielding, in-

dicating the high elasticity of the materials. Figure 15B shows the load and

strain at the yield point, and the stiffness and compression modulus of the

scaffolds calculated using a strain range of 0–10%. No statistically significant

differences (p<0.05) were observed between properties of PCL and PDP5 scaf-

folds, while 10 mol% of depsipeptide in PDP10 scaffolds was sufficient to in-

crease these values significantly. PDP10 scaffolds had a three-fold higher load

and two-fold higher strain at the yield point when compared to PCL homopol-

ymer scaffolds. The compression modulus and stiffness were also significantly

higher for PDP10 scaffolds compared to other scaffolds. The increased me-

chanical strength of PDP10 scaffolds was attributed to L-alanine molecules in

the depsipeptide units forming strong intermolecular hydrogen bonds between

the amide groups, 87 enhancing the mechanical strength of the scaffolds. Ap-

parently, 5 mol% of depsipeptide was not sufficient to significantly change the

mechanical properties of PDP5 scaffolds when compared to PCL scaffolds.

Figure 15. Compression results of A) PCL/BG scaffoldsII and B) PCL-co-PDP scaffolds.III

4.5.4 Bioactivity of BG/PCL composite scaffolds

To evaluate the effect of BG on the bioactivity of PCL scaffolds, ion release

from the scaffolds and formation of precipitates were measured by monitoring

concentration changes of silica (Si), calcium (Ca), and phosphorus (P) in SBF

(Figure 16A). BG20 samples showed a strong decrease in Ca and P concentra-

tions, indicating that calcium and phosphorus were precipitating on the scaf-

folds, whereas BG10 scaffolds showed only a moderate decrease. The overall

Ca and P concentrations remained steady in BG0 and BG5 samples, indicating

that no precipitation occurred. No silica was detected in PCL samples, while all

PCL/BG samples had increased SBF Si levels, indicating that SiO2 from BG

was hydrolyzed and released from the scaffolds. Figure 16B shows SEM images

of PCL/BG scaffolds after 3 days, 10 days, and 21 days in SBF. Clear calcium

phosphate precipitates were seen on the surface of BG20 scaffolds after 3 days;

after 21 days, BG20 scaffolds were fully covered with calcium phosphate. Un-

like BG20 and BG10 scaffolds, BG5 and BG0 scaffolds did not exhibit precipi-

tates. The SEM-EDS elemental analysis of the BG20 scaffold surface after 3

days in SBF revealed that the molar ratio of calcium and phosphorus in the

precipitates was 1.64–1.86, which was close to the Ca/P ratio of the natural

hydroxyapatite Ca5(PO4)3(OH), whose ratio is 1.67. 241 Based on the key role of

the HA layer in interactions between bone tissue and the scaffold, the for-

mation of a HA-like calcium phosphate layer on composite scaffolds was con-

sidered an indication of their in vitro bioactivity. 242, 243 Although the SiO2 of

BG particles hydrolyzed in all PCL/BG scaffolds, it was concluded that BG con-

centrations below 10 wt% were not enough to make the scaffolds bioactive for

calcium phosphate precipitation.

Figure 16. A) Concentration of calcium, phosphorus, and silica within 21 days of scaffold

immersion in SBF and B) SEM images of scaffolds after immersion: a) BG20 after 3 days, b)

BG20 after 10 days, c) BG10 after 21 days, and d) BG20 after 21 days.II

4.5.5 Characteristics of 3D fabricated hydrogels

To study the effect of crosslinking time on the degradation and mechanical

properties of PEG-co-PDP hydrogels, hydrogel samples were fabricated with

SLA using three different light exposure times. The gel content of hydrogels

increased from 85% to 88% with increasing crosslinking time, while the swell-

ing degree decreased from 21 to 16. As more double bonds reacted within the

continued light exposure time, the crosslinking density and gel content of the

hydrogels increased, resulting in decreased mesh size and thus more restricted

water penetration into the hydrogel network. A similar decrease in swelling

capacity has been previously reported for PEG hydrogels, where the crosslink-

ing density was increased by increasing polymer concentration 32 or by de-

creasing the molecular weight of the polymer. 34 Here, the crosslinking density

was increased by extending the crosslinking time, and thus, there was no need

for changing the parameters of the hydrogel solution.

Table 4. Gel content and swelling degree of the PEG-co-PDP hydrogels prepared by SLA. IV

Light exposure time/layer (s)

Gel content (%)

Swelling degree

100 85 ± 2 21 ± 1 120 86 ± 1 18 ± 1 160 88 ± 1 16 ± 1

The effect of crosslinking time and crosslinking density on the hydrolytic

mass loss of photocrosslinked PEG-co-PDP hydrogels was monitored by incu-

bating the gels in PBS and weighing their remaining dry mass after predeter-

mined times. Figure 17A shows that the total mass loss of hydrogels decreased

with increasing crosslinking time. During 24 d in PBS, the 100 s samples lost

66% of their initial dry mass, while the 120 s samples lost 45% and the 160 s

samples lost 35% of their initial mass. The mass loss followed a zero order

curve trend, showing similar linear mass loss as previously reported for photo-

crosslinked PEG-co-PLA hydrogels. 244 During the degradation study, the hy-

drolysis of degradable bonds in the PEG-co-PDP chains resulted in broken

crosslinks, weakening the mechanical integrity of the hydrogel. When all the

crosslinks of the multi-arm copolymer were broken, the copolymer unit was

free to diffuse out of the hydrogel network, leading to hydrogel mass loss. The

difference in mass loss rates of PEG-co-PDP hydrogels can thus be explained

by differences in their crosslinking densities. When longer light exposure times

increased the crosslinking degree, more polymer arms were covalently bound

to the hydrogel network, making the release of polymer chains more difficult

and slowing down mass loss. The degradation of polymer networks was also

evidenced as an increase in the swelling degree of hydrogels during incubation

in PBS. Figure 17B shows that the swelling degree of hydrogels increased dur-

ing the 24-day degradation study from 21 to 58 for the 100 s samples, from 18

to 45 for the 120 s samples, and from 16 to 40 for the 160 s samples. The in-

crease in swelling degree was slow in the early stage of degradation, but accel-

erated in the later stage, indicating that the decrease in crosslinking density

occurred in a non-linear fashion. This result was consistent with a similar ex-

ponential increase in swelling degree of previously reported biodegradable

PEG-co-PLA hydrogels. 244

Figure 17. A) In vitro mass loss and B) swelling degree of the PEG-co-PDP hydrogels with

various crosslinking times during 24-day degradation study in PBS (pH 7.4).IV

To evaluate the effect of light exposure time and polymer degradation on the

local mechanical stiffness of 3D fabricated PEG-co-PDP hydrogels, AFM was

used to measure the Young’s modulus of wet samples right after SLA fabrica-

tion and after 7 d in PBS. Figure 18A shows that the stiffness increased with

crosslinking time. The initial Young’s modulus was 3.0 ± 1.0 kPa, 9.4 ± 1.2

kPa, and 37.9 ± 13.3 kPa for the 100 s, 120 s, and 160 s samples, respectively;

these were in the stiffness range of natural soft tissues regardless of the cross-

linking time. 245 Hydrogels composed of 10% w/v of PEG diacrylates with com-

parable chain lengths have shown bulk stiffness of 50 kPa 187 and 170 kPa, 183

indicating that our hydrogels were in the lower range of mechanical stiffness

for corresponding PEG hydrogels. Within 7 d in PBS (Figure 18C), Young’s

modulus decreased to 0.24 ± 0.02 kPa, 0.96 ± 0.50 kPa, and 1.52 ± 0.55 kPa,

for the 100 s, 120 s, and 160 s samples, respectively. Rheometric analysis of the

hydrogels revealed that the bulk storage modulus describing the elastic prop-

erties of gels increased with increasing crosslinking time. The storage modulus

was initially in the order of magnitude of 103−104 Pa (Figure 18B), while it de-

creased to 102−103 Pa (Figure 18D) after 7 d in PBS. The decrease of one order

of magnitude within the 7-day hydrolysis period was consistent with both the

AFM and rheometric studies. As the mechanical elasticity and strength of the

photocrosslinked hydrogel are caused by the retractive forces of the cross-

linked polymer chains, 244 the cleavage of biodegradable PEG-co-PDP chains

decreased the total forces, and together with the reduced polymer mass and

increased swelling of the gels, resulted in decreased mechanical stiffness of

degraded hydrogels.

Figure 18. Local Young’s modulus measured with AFM, and bulk storage and loss modulus

of the 3D hydrogels measured with a rheometric analysis A-B) right after SLA fabrication and C-

D) after 7 d in PBS.IV

4.5.6 Cell proliferation on scaffolds and within hydrogels

To further evaluate the in vitro bioactivity of PCL/BG scaffolds, fibroblasts

were seeded on scaffolds and their proliferation was measured for 14 days.

Figure 19A shows that cell activity increased between day 1 and 7 in each

group, and BG20 scaffolds showed the most active cell proliferation. At day 3

and 7, cell activity was significantly higher (p < 0.05) on BG10 and BG20 scaf-

folds when compared to BG0 and BG5 samples. At day 14, the most highly

populated BG20 scaffolds showed a decrease in cell activity, which was con-

cluded to result from cell confluence on the scaffold surface. In SEM and con-

focal fluorescence microscopy images (Figure 19B), typical fibroblast mor-

phology was observed after 14 days of cell culturing, and cells were growing

throughout the scaffold surface as well as inside the matrix. That indicated

high potential for the PCL/BG scaffolds as tissue engineering materials. Even

though the differences in cell attachment were not further studied, it was con-

cluded that the addition of BG to PCL scaffolds induced changes in the scaffold

surface chemistry and topography, which in turn enhanced cell attachment

and proliferation on the composite scaffolds when compared with neat PCL

polymer scaffolds.

Figure 19. A) Proliferation of human gingival fibroblast seeded on PCL and PCL/BG scaf-

folds measured by AlamarBlue® metabolic assay. B) SEM images (a-b) and fluorescence images

(c-d) of the PCL/BG scaffolds after 14 days of cell culturing II

To preliminarily study the cell attachment properties of PCL-co-PDP scaf-

folds, HUVECs were seeded on the scaffolds and cultured for 7 d. Due to the

highly hydrophobic character of the photocrosslinked polymer, cell seeding

was done by slowly absorbing the cell suspension into the scaffold. Figure 20

shows the cytoskeletal organization of cells after seven days of cell culturing

followed by F-actin/nucleus staining. At day 7, a dense cell layer was observed

on the scaffold surface indicating great cell attachment and proliferation of

cells on the scaffold. Since good capacity for cell attachment is a key property

of TE scaffolds, the great cell growth on photocrosslinked scaffolds in this

study indicated their high potential as biodegradable tissue engineering mate-

rials.

Figure 20. A-B) HUVECs on a PCL-co-PDP scaffold after 7 days of cell culturing. Rhodamine

phalloidin/DAPI staining shows F-actin (red) and cell nuclei (blue) of the cells.III

3D fabrication of cell-laden hydrogels with biomimetic complexity and en-

hanced functionality is particularly promising in tissue engineering and drug

screening; it provides more physiologically relevant data on cell behavior and

responses to stimuli than can be obtained by seeding cells on a surface of pre-

fabricated hydrogels. SLA was used to fabricate cell-laden PEG-co-PDP hydro-

gels using a filtered PEG-co-PDP/RGDS hydrogel solution mixed with HUVEC

suspension and by photocrosslinking samples in a layer-by-layer manner. The

effect of crosslinking time on cell proliferation was tested by fabricating the

cell-laden hydrogels with increasing light exposure time and by monitoring the

metabolic activity of encapsulated cells with the colorimetric AlamarBlue®

assay. Figure 21A shows that cells significantly increased their metabolic activ-

ity in all sample groups from day 1 to day 10. At each designated time point,

there was no significant difference in cell activity among the sample groups.

Representative fluorescence images of the 160 s samples show that the cells

were distributed homogenously in the gels after 1 d (Figure 21B) and 7 d of cell

culturing (Figure 21C). The initially non-porous PEG-co-PDP/RGDS hydrogels

supported proliferation of encapsulated HUVECs within the 10-day cell cultur-

ing window regardless of crosslinking time. A previous study elsewhere

showed that non-degradable, non-porous PEG hydrogels prepared by SLA had

no cell proliferation within 7-day cell culturing, while the addition of small

channels to the hydrogel improved cell proliferation. 36 Here, successful cell

proliferation was attributed to polymer degradation and subsequent mass loss,

and the decreased mechanical stiffness of the hydrogel matrix. This is con-

sistent with previous studies showing that the high stiffness of PEG hydrogels

tends to decrease the proliferation of encapsulated cells due to the physical

barrier caused by its non-degradable polymer network. 34, 181 By using biode-

gradable PEG-co-PDP macromer, hydrogels degraded simultaneously with cell

growth, so that the degrading gels continuously provided the cells with more

space and a softer environment to grow inside the hydrogel matrix.

Figure 21. A) Metabolic activity of encapsulated cells in 3D fabricated hydrogels within 10 d

of cell culturing measured with a colorimetric AlamarBlue® cell viability assay.IV Fluorescence

images of encapsulated HUVECs in 160 s hydrogels after B) 1 d and C) 7 d of cell culturing.

To demonstrate the capacity of the new PEG-co-PDP macromer in SLA-

based fabrication of tubular hydrogel constructs for vascular tissue engineer-

ing, ring-like cell-laden hydrogels and bifurcating vascular tubes were de-

signed and 3D fabricated by SLA. The fluorescence images in Figure 22A and

22B show that the cells were homogenously distributed in the hydrogel ring

after 1 d of cell culturing. Figure 22C shows the uniform and round shape of

the hydrogel ring. Figure 22D shows that the 3D fabricated bifurcating hydro-

gel constructs closely resembled their digital models and had a uniform and

smooth wall structure as required for functional vascular tissue engineering

grafts. The good cell encapsulation capacity and printability of the PEG-co-

PDP polymer demonstrated its high capacity for SLA-based fabrication of 3D

modeled vascular grafts, and more importantly, proved a significant addition

to the very limited group of biodegradable polymers suitable for SLA fabrica-

tion of cell-laden TE constructs.

Figure 22. A-B) Fluorescence images and C) a photograph of a ring-like hydrogel construct;

D) CAD models of bifurcating vascular tubes and the resulting hydrogel vessels visualized by

cross-section and top view. IV

5. Conclusions and Future Prospects

This thesis expanded the repertoire of biodegradable photocrosslinkable pol-

ymers that are suitable for 3D fabrication of highly defined TE scaffolds and

hydrogels by SLA. The emergence of new polymers is crucial to promote the

use of additive manufacturing and SLA in patient-specific, more effective and

safer healthcare. In this study, star-shaped PCL- and PEG-based oligomers

were synthesized through ring-opening polymerization, and photocrosslinka-

ble macromers were obtained by functionalizing the hydroxyl-ended oligomers

with methacrylate groups. The macromers were photocrosslinked using radi-

cal-forming photoinitiators that started free radical chain-growth polymeriza-

tion through the double bonds of the macromers. Photocrosslinking resulted

in polymer networks with high gel content. The porous mathematically defined

PCL-based scaffolds were built by visible light projection SLA using liquid

photocrosslinkable resins consisting of a macromer and a photoinitiator.

When the solid PCL macromer was used, it was heated above its melting tem-

perature to obtain a liquid resin for SLA without need for using additional sol-

vents. The porous scaffolds accurately resembled the structures modeled by

computer-aided design, and the high interconnectivity of the pores indicated

great potential for these scaffolds as TE constructs.

To enhance the bioactivity of PCL scaffolds, bioactive glass was added to the

PCL resin before fabrication of porous composite scaffolds by SLA. By using a

liquid low molecular weight PCL resin, the need for heating was avoided. The

use of SLA enabled homogeneous distribution of BG throughout the scaffolds,

leaving BG particles on the scaffold surface uncovered and ready for rapid in-

teractions with surrounding cells. Composite scaffolds showed precipitation of

calcium phosphate on their surface in simulated body fluid, indicating their in

vitro bioactivity. The improved bioactivity further enhanced the attachment

and proliferation of cells on the scaffold surface.

The degradation and mechanical properties of the PCL macromer were

tuned by synthesizing a new type of photocrosslinkable poly(ester amide)

macromer based on PCL and L-alanine-derived depsipeptide. The copolymeri-

zation of PCL with the depsipeptide increased its hydrophilicity, resulting in

faster hydrolytic degradation. The mechanical strength of SLA fabricated TE

scaffolds also increased with depsipeptide content, which was assumed to re-

sult from strong intermolecular hydrogen bonds between amide groups in the

depsipeptide units. In vitro cell studies revealed the excellent cytocompatibil-

ity of photocrosslinked PCL-co-PDP copolymers, suggesting their high poten-

tial as readily tunable polymers for SLA fabrication of biodegradable TE appli-

cations.

Since hydrophobic PCL scaffolds lack the dynamic water-rich environment

needed for cell encapsulation, a new hydrophilic, photocrosslinkable

poly(ether ester amide) macromer based on 4-arm PEG and the L-alanine dep-

sipeptide was synthesized. The resulting biodegradable, water-soluble PEG-co-

PDP macromer was used for 3D fabrication of cell-laden hydrogels by visible

light projection SLA. The gel content and mechanical stiffness of 3D fabricated

hydrogels increased with increasing crosslinking time, while the swelling and

mass loss of the gels in aqueous media decreased, indicating their increased

crosslinking density. The mechanical stiffness of the gels decreased by one

order magnitude within a 7-day incubation in PBS, indicating the degradation

of hydrolytically labile bonds of the hydrogel networks. The encapsulated cells

proliferated well during 10 days of culturing in the hydrogels regardless of

light exposure time. The advanced use of the new polymer in SLA was success-

fully demonstrated by fabricating bifurcating tubular structures as preliminary

vessel grafts. The new biodegradable, photocrosslinkable polymer was a signif-

icant addition to the very limited group of polymers currently available for SLA

fabrication of cell-laden TE hydrogels. In the future, more studies are needed to evaluate the in vivo biodegradation

and biocompatibility of the synthesized polymers. In addition, new approaches

to support the vascularization of scaffolds are highly interesting. Since it is

known that the diffusion limit of oxygen in mammalian tissue is 100 - 200 μm,

246-248 the cells forming the tissue need to be located within that distance of the

blood vessels to secure a sufficient supply of oxygen and nutrients. Therefore,

to grow tissues thicker than a few hundred micrometers, angiogenesis is need-

ed to form new blood vessels. As the spontaneous vascular ingrowth from the

host tissue to the implanted tissue structure is only several tenths of a mi-

crometer per day, 249 it can be hypothesized that the incorporation of vascular

structures to TE scaffolds in vitro may significantly speed up their in vivo vas-

cularization. This, in turn, would improve cell proliferation and tissue for-

mation inside the 3D tissue constructs. To address the challenge of incomplete

cell viability inside TE grafts, the next step will be to combine the 3D fabrica-

tion of rigid biodegradable TE scaffolds with the fabrication of cell-laden hy-

drogels to prepare prevascularized multi-material TE constructs.250 In that

approach, a mechanically stronger polymer scaffold is fabricated layer by layer

simultaneously with photocrosslinking a soft, cell-laden hydrogel inside the

scaffold. The stronger, hydrophobic polymer scaffold will then mechanically

support the softer hydrogel construct, and the cell-laden hydrogel will provide

the scaffold with cells and guide the cells to grow into a functional vascular

network inside the porous TE scaffold.

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Biodegradable polymers are increasingly used in biomedical applications. When implanted into a body, biodegradable material does not require removal surgery but will disappear through polymer degradation. In tissue engineering, porous biodegradable scaffolds are used to grow tissue structures that restore or improve natural tissue functions. Additive manufacturing techniques provide a feasible method to 3D fabricate patient-specific scaffolds that closely satisfy the requirements of their target tissue. In this dissertation, new biodegradable photocrosslinkable polymers were synthesized and used for stereolithography-based 3D fabrication of tissue engineering constructs. The development of new biodegradable polymers with great printing properties is a crucial step toward the successful use of additive manufacturing in more effective, personalized medicine.

Aalto-D

D 16

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015

9HSTFMG*agejbi+

ISBN 978-952-60-6491-8 (printed) ISBN 978-952-60-6492-5 (pdf) ISSN-L 1799-4934 ISSN 1799-4934 (printed) ISSN 1799-4942 (pdf) Aalto University School of Chemical Technology Department of Biotechnology and Chemical Technology www.aalto.fi

BUSINESS + ECONOMY ART + DESIGN + ARCHITECTURE SCIENCE + TECHNOLOGY CROSSOVER DOCTORAL DISSERTATIONS

Laura E

lomaa

Synthesis of biodegradable photocrosslinkable poly- mers for stereolithography-based 3D

fabrication of tissue engineering scaffolds and hydrogels

Aalto

Unive

rsity

2015

Department of Biotechnology and Chemical Technology

Synthesis of biodegradable photocrosslinkable poly- mers for stereolithography-based 3D fabrication of tissue engineering scaffolds and hydrogels

Laura Elomaa

DOCTORAL DISSERTATIONS