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University of Groningen
Biomedical polyurethane networksBruin, Peter
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Rijksuniversiteit Groningen
BIOMEDICAL POLYURETHANE NETWORKS
ter verkrijging van het doctoraat in de
Wiskunde en Natuurwetenschappen
aan de Rijksuniversiteit Groningen
op gezag van de
Rector Magnificus Dr. S.K. Kuipers
in het openbaar te verdedigen op
vrijdag 6 november 1992
des namiddags te 4.00 uur
door
PETER BRUIN
geboren op 22 maart 1963
te Hoorn
Promotor: Prof. Dr. A.J. Pennings
Dear Sir or Madam will you read my book
it took me years to write will you take a look
(from "Paperback writer" by the Beatles)
Dankwoord
Iedereen die op enigerlei wijze heeft bijgedragen aan de totstandkoming
van dit proefschrift wil ik hiervoor bedanken, met name:
mijn promotor prof. dr. A.J. Pennings voor de geboden mogelijkheid om
onder zijn deskundige toezicht dit promotieonderzoek uit te voeren,
prof. dr. G. Challa, prof. dr. J.H. Teuben en prof. dr. B. Witholt voor de
bereidwilligheid om zitting te willen nemen in mijn promotiecommissie en
het manuscript te beoordelen,
Annemarie Brummelhuis, Andries Hanzen, Henk Hoppen, Koen Knol, Hendrik
Luttikhedde, Edwin Meeuwsen, Henk-Jaap Meijer, Joke Smedinga, Gert-Jan
Veenstra, Norbert Wolberink, Geartsje Zondervan voor de waardevolle
bijdragen aan het experimentele werk,
Adams Verweij, Harry Nijland (electronen microscopie en fotografie), Anne
Appeldoorn (instrumentmakerij), Henk Knol (glasblazerij), Harm Draaijer &
Jan Ebels (microanalyse afdeling) voor de practische assistentie en
technische ondersteuning,
dr. Jan Willem Leenslag zonder wiens werk dit proefschrift er heel anders
uitgezien zou hebben,
dr. Marcel Jonkman en drs. Jean Coenen voor de prettige samenwerking op
het terrein van de wondbedekking,
dr. Berend van der Lei voor zijn bijdrage aan het vaatprothese onderzoek,
drs. Pek van Andel, drs. Gerard van der Veen en prof. dr. J.G.F. Worst
voor het enthousiasmeren voor de oogheelkundige toepassing van polymere
materialen,
en verder (zonder anderen te kort te willen doen) Machiel Bos, Jacqueline
de Groot, Atze Nijenhuis, Jan "Les" Paul Penning , alle "anonieme"
collega's, studenten, secretaresses, andere medewerkers en staf van de
vakgroep polymeerchemie voor de hulp, discussies en sfeer tijdens mijn
verblijf in jullie midden.
CONTENTS
Chapter 1 Introduction 1
Chapter 2 Design and synthesis of biodegradable 17
poly(ester-urethane) elastomer networks composed of
non-toxic building blocks
Chapter 3 Biodegradable lysine diisocyanate-based
poly(glyco1ide-co-E-capro1actone)urethane network
in artificial skin
Chapter 4 A new porous polyetherurethane wound covering
Chapter 5 A two-ply artificial blood vessel of polyurethane and 53
poly(L-lactide)
Chapter 6 Autoclavable highly cross-linked polyurethane networks 77
in ophtalmology
Summary
Samenvatting
Chapter 1
Introduction
Polymer networks
Real cross-linked polymer networks always deviate from ideal, perfect
networks, which are defined as random, homogeneous collections of
(Gaussian) chains between multifunctional junction points (cross-links)
under the condition that all functionalities of the junction points have
reacted with the ends of all and different chains (1,2). In other words,
the ideal network entirely consists of elastically effective chains,
meaning chains connecting two neighbouring cross-links in the network,
able to transfer a retractive force throughout the material if subjected
to an elastic deformation.
Chompff stated that real polymer networks consist of inhomogeneous
structures because if polymers form homogeneous continua their strengths
would theoretically be about one hundred times higher than is observed
experimentally ( 3 ) . The degree of inhomogeneity of polymer networks
depends on the way in which the network has been formed and the ultimate
strength, for example, is sensitively dependent on such defects (4).
Besides inhomogeneity in cross-link distribution there are other network
imperfections which may be introduced upon network formation: network
defects, like dangling or pendant chain, i.e., a chain attached to the
network at only one of its ends (unreacted functional endgroups, for
instance), elastically inactive loops (as a result of intramolecular
cross-linking), chain entanglements (51, and heterogeneity due to phase
separation ( 1 1.
Polymeric, chemically cross-linked networks can be formed in three
different ways:
-cross-linking (vulcanization) of existing linear polymers.
-chain cross-linking (co)polymerization.
-stepreaction of small molecules (multifunctional monomers and/or
prepolymers), all of which are reactive at the same time.
Polymer networks need not to be formed exclusively by chemical pathways
leading to permanent networks. Cross-linking by physical aggregation of
polymer chains also results in network structures. Examples of polymer
gels containing physical cross-links include microcrystalline polymers,
ionomers, chelation polymers, blockcopolymers (thermoplastic elastomers,
like polystyrene-butadiene triblockcopolymers and segmented
polyurethanes), stereocomplexes. In the case of physically cross-linked
polymer networks the cross-links are not permanent; and the physical
aggregation is often (thermo)reversible (2.10).
Cross-linking of linear polymers
Polymer chains can be cross-linked by chemically joining different
primary, linear polymer molecules. The techniques generally used to
introduce cross-links are peroxide decomposition, high-energy irradiation
and sulfur addition to skeletal or side-chain double bonds (2.6.7). All of
these cross-link methods are statistical processes. Cross-links are
introduced in a highly random manner resulting in polymeric networks
having a rather undefined network topology (2,s).
The minimum number of dangling ends is inversely proportional to the
number-average molecular weight of the starting polymer ( 7 , s ) . Due to the
finite molecular weight, dangling ends will always be present in the final
network. Networks cross-linked by means of high-energy radiation may
contain even higher concentrations of dangling chain ends arising from
chain-scission occurring during the cross-linking process.
In the case of peroxide curing, especially when carried out in solution.
peroxide radical fragments may chemically contaminate the initial polymer
chains ( 5 ,9 ) . Networks cross-linked in solution usually contain a lower
concentration of trapped entanglements (compared to cross-linking in the
melt), but here loop formation becomes important (2,10,111.
Chain cross-linking (co)polymerization
The free radical initiated chain (co)polymerization of monovinyl and
polyvinyl monomers leads to the formation of polymeric networks. Typical
examples are the bulk photopolymerization of diacrylates resulting in
glassy, densely cross-linked networks (12,131; the copolymerization of
styrene and divinylbenzene in an inert solvent leading to phase separated,
microporous polystyrene gels (14); and the copolymerization of acrylarnide
and N,N'-methylene bisacrylamide in aqueous solution leading to highly
swollen polyacrylamide gels (151.
Boots used the kinetic gelation model for the simulation of the free
radical chain cross-linking polymerization and showed that network
formation by chain reactions is an intrinsically inhomogeneous process, in
contrast to network formation by stepreactions (16-20). Snapshots taken
during a simulation of polymerization showed an inhomogeneous spatial
distribution of polymer. The homogeneity of the network created could be
improved by decreasing the kinetic chain lengths (In practice this means
increasing the initiator concentration or the temperature). Only at
unrealistically high initiation rates the network formed by the chain
reaction process becomes as homogeneous as the one formed by a step
mechanism. It should be mentioned that for chain cross-linking
polymerizations in the bulk the model implies homogenization towards the
end of the reaction. Complete monomer conversion in the bulk
polymerization is an unlikely event as a result of premature
vitrification.
According to DuSek cyclization plays a dominant role in the chain reaction
network formation. Primarily internally cross-linked microgel particles
are formed, which are linked through peripheral double bonds to form a
heterogeneous gel (21,221. In an inert solvent this leads to the formation
of macroporous, inhomogeneous structures; in bulk reactions
re-homogenization seems likely since polymers cross-linked to high
conversion appear, optically for example, to be quite homogeneous in
agreement with the model described by Boots.
Experimental evidence for the heterogeneity of polyadrylamide gels,
representative for chain cross-linked polymer networks, comes from several
independent studies, such as swelling measurements (231, scattering
studies (41, kinetics (241, electron microscopy (251, and permeability
studies (26-29).
The latter will be discussed in some more detail. Silberberg and coworkers
measured the permeability of polyacrylamide gels (made by the free radical
copolymerization of acrylamide and N,N'-methylene bisacrylamide in water)
and aqueous solutions of linear polyacrylamide. First of all, they
observed that when the total monomer concentration was kept constant, but
the percentage of the cross-linking monomer was increased, the
permeability of the gel rose markedly. It was also found that, at a
comparable concentration, a system of uncross-linked polymer possessed the
lowest permeability. This could only be accounted for by an inhomogeneous
(nonuniform) distribution of the monomer in the gel. In the gel two
regions are formed: one, containing a high proportion of the gel substance
(having a high cross-link density), is essentially non-draining and the
other, containing a very low fraction of the gel substance, is "freely"
draining. Microgel particles, formed in the early stage of the
polymerization, are weakly linked to form a macroscopic gel. The effective
concentration of polymer in the freely draining region is well below the
total monomer concentration as a result of the clustering. Fluid can thus
move faster through the heterogeneous gel than through a polymer solution
of corresponding overall, but uniform concentration. The permeability is
increased when more gel substance is incorporated into the non-draining
region, which happens when the cross-linker concentration is increased.
Another salient aspect is that the faster the initiation, the less
permeable a gel becomes (291. In other words, a high rate of initiation
leads to a more homogeneous network, which again is in full agreement with
the results of the computer simulations done by Boots.
Stepgrowth polymerization
Network structures may also be formed by the random stepgrowth
polymerization of monomers (or prepolymers), at least one type of which
has a functionality of 3 or greater. A classic example is the
polycondensation of dicarboxylic acids with glycerol leading to a
cross-linked polyester (7). Other polymer networks synthesized by
stepreactions include polyurethanes, epoxies, formaldehyde-based
thermosets (Bakelite. Novolac, etc. 1 , polydimethylsiloxane elastomers
(30,311.
By using stepreactions for the formation of polymer networks the topology
of the resulting polymeric network is well defined. Networks having known
values of cross-link functionality and the molecular weight between
cross-links (with known distibution of Mc) are called model polymer
networks (2,6,31-35). These are usually prepared by selectively
end-linking bifunctionally terminated chains (telechelic prepolymers)
with a multifunctional cross-linker. The resulting networks thus have
cross-links of functionality equal to that of the cross-linking agent, and
values of Mc corresponding to the values of the number-average molecular
weight (and its distribution) of the prepolymer prior to cross-linking.
Model elastomeric polydimethylsiloxane (PDMS) networks were prepared by
reaction of hydroxylterminated PDMS chains with tetraethylorthosilicate;
polyurethane elastomers were synthesized by end-linking of
dihydroxylterminated PPO, PEO or polycaprolactone prepolymers with a
triisocyanate, or a polytriol with a diisocyanate (38). These model
elastomeric networks were used to gain some more insight into the
relationships between the structure of a network and its (ultimate)
properties.
One of the imperfections known to be present in network structures is the
dangling chain. The incidence of dangling-end network imperfections in
model networks is very small when the end-linking reaction is carried out
stoichiometrically and to high conversion of functional groups. By
unbalancing the stoichiometry in the end-linking reaction or by
incorporation of monofunctional prepolymers known numbers of dangling
chains were introduced into the network structure (8). As expected an
increase in the concentration of dangling chains had a negative effect on
the ultimate properties, both ultimate strength and maximum extensibility,
of the elastomer.
Another interesting result obtained from these model networks concerns the
network chain length distribution. Properties of bimodal networks,
consisting of very short chains (Mc around a few hundred gmol-i) and
relatively long chains (Mc around 20000 gmol-l), were compared with
unimodal elastomeric networks. It was found that increasing the number of
very short chains in bimodal networks did not show any decreases in
ultimate properties. The strain, even at high elongations, is apparently
reapportioned (nonaffinely) within the network so as to ignore as long as
possible the difficultly deformable short chains (35,361. According to the
"weakest-link" theory rupture of an elastomer is caused by failure of the
shortest network chains (37). This theory is based on the assumption of
affine deformation, which does not seem to be correct. Bimodal networks
containing very large concentrations of short chains (90-95 mol %) were
found to have both high ultimate strength and high maximum extensibility
and as a result of that high toughness, in contrast to both unimodal
networks containing either only short chains or only long chains. This
result is rather surprising, since usually an elastomer will have good
ultimate properties only when reinforced with some mineral filler or hard,
glassy domains in the case of a multiphase polymer, or when it can
generate its own reinforcement through strain-induced crystallization
(39). Since the bimodal elastomeric (PDMS) networks studied have a low T 0
(-40 C) they can not crystallize upon stretching at room temperature.
These bimodal networks show upturns in the modulus at high elongations
that are diminished by neither increase in temperature nor swelling,
unlike networks that can undergo strain-induced crystallization. Any
intermolecular reinforcing effects can thus be ignored. Due to their
limited extensibility of the short chains in the bimodal network the
modulus and ultimate strength are high. The long chains present in the
network somehow delay the growth of the rupture nuclei required for
catastrophic failure. Beyond 95 mol % short chains properties decline,
because of increasing brittleness (21. The unusual ultimate properties of
bimodal networks are due to limited chain extensibilities (non-Gaussian
behaviour) rather than to reinforcing effects (40).
The extent of cyclization in model networks, prepared in the absence of a
diluent, is rather weak when compared to networks formed by chain
cross-linking polymerization (22,411. Dilution may considerably increase
the cyclization and may cause the formation of inhomogeneities, or even
phase separation (11,41). The extent of cyclization depends predominantly
on the polymerization mechanism (22). Stepto stated that the formation of
a perfect network from an end-linking polymerization requires that pre-gel
intramolecular reaction is negligible and that post-gel intramolecular
reaction always leads to elastically ac-tive chains. These requirements are
unlikely to be met. He showed that inelastic loops will arise from both
pre-gel and post-gel intramolecular reactions, which will always occur in
non-linear stepgrowth polymerizations (42-45).
Summarizing we may state that in principal network formation by
stepreactions is a homogeneous process as shown by Boots (18,191 and model
networks with a known structure can be obtained, having small numbers of
dangling-end network imperfections. However, perfect networks free of any
imperfections are not accessible experimentally (44,461.
Polymer networks in medicine
Polymer networks, especially elastomers, have been used in biomedical
applications for several decades (47). Silicone rubber (medical grade) for
example is used in numerous applications including mammary and facial
prostheses, contact and intraocular lenses (55-571, prosthetic heart
valves (48-50). Silicone elastomers are (usually peroxide) cross-linked
polysiloxanes, predominantly PDMS, that are mostly reinforced with
ultrafine particle silica filler to improve the mechanical properties.
Silicone rubbers are highly hydrophobic materials known for their
inertness, high oxygen permeability and relatively good
bloodcompatibility. Thermoplastic polyurethane elastomers, which will be
discussed in the next section, have also been used extensively for medical
applications.
Another class of biomaterials concerns the hydrogels, water swollen
polymers, usually polymer networks (51-53). Hydrogels are very interesting
materials since they superficially resemble living soft tissue in their
physical properties. Hydrogels have relatively high water contents and a
soft, rubbery consistency, causing no mechanical, frictional irritation to
the surrounding tissue, and allow the permeation and diffusion of small
molecules, metabolites just like living tissue. Hydrogels have a low
interfacial free energy in aqueous surroundings. The higher the water
content of the gel, the lower the interfacial free energy becomes. This
low interfacial tension should reduce the tendency of proteins in body
fluids to adsorb and to unfold upon adsorption and also minimize the
driving force for cell adhesion. Minimal protein interaction may be
important for the acceptance of any material when implanted. The
denaturation of proteins by surfaces of implant materials may serve as a
trigger mechanism for the initlation of thrombosis or for other biological
rejection mechanisms. Hydrogels are considered (soft tissue) biocompatible
and bloodcompatible, but suffer from poor mechanical properties. The
higher the water content of the gel, the poorer the mechanical properties
become. The inferior mechanical properties severely limit the potential
applicability of hydrogels. Nevertheless, hydrogels have found many
biomedical applications, like soft contact lenses (57). wound coverings
(541, drug delivery systems, foldable intraocular lenses (55,561,
encapsulation of living cells (58). There are ways to overcome these
mechanical limitations: surface grafting of hydrophilic polymers onto
mechanically strong (hydrophobic) polymers, for example grafted
polyacrylamide or polyethyleneoxide onto polymers for vascular prostheses
to improve the bloodcompatibility (59,601, copolymerization of a
hydrophilic monomer with a more hydrophobic one to form a linear
blockcopolymer, or formation of interpenetrating polymer networks (52).
Wichterle and Lim developed the concept of synthetic, polymeric hydrogels
designed for biomedical applications (53). They described the synthesis of
,lightly cross-linked poly(2-hydroxyethyl methacrylate) (PHEMA) gels. Other
types of hydrogels are prepared from (meth)acrylamide, vinylalcohol and
N-vinylpyrrolidone monomers (51,521. Also hydrophilic networks with
polyethyleneoxide have been synthesized by cross-linking (high molecular
weight) PEO by g-irradiation or by peroxide curing (60.61). Better defined
PEO containing polyurethane networks were obtained by reaction of low
molecular weight PEO diols with triisocyanates or by mixing diisocyanates
with di- and triols (62,631.
Finally, dental restoration materials based on photocurable (aromatic)
dimethacrylates should be mentioned. Concerns over the toxicity of dental
amalgam and an increased emphasis on aesthetics have popularized the
development and clinical use of dental composite restorative materials.
These dental composite resins are composed of a photopolymerizable
dimethacrylate matrix filled with fine (treated) silica particles to
increase the hardness and to lower the overall shrinkage upon
polymerization of the composite material (64).
Polyurethanes
Polyurethanes are a class of polymers having only in common the presence
of urethane bonds somewhere in their chains. The name polyurethane given
to a polymeric material does not tell anything about its chemical and
physical characteristics. Polyurethanes may be lightly or highly
cross-linked or uncross-linked and be highly crystalline, elastomeric or
amorphous and glassy. In the biomedical field polyurethane usually stands
for thermoplastic polyurethane elastomer. Thermoplastic polyurethanes,
also named segmented polyurethanes, are linear blockcopolymers composed of
chainextended diisocyanate hard segments dispersed in a soft segment
polyol matrix (65,661. All commercially available biomedical
polyurethanes, like Biomer, Estane, Cardiothane, Pellethane, Tecoflex, are
composed of an aromatic diisocyanate MDI (4.4' -methylenediphenyl
diisocyanate), except for Tecoflex which contains hydrogenated MDI, a
cycloaliphatic diisocyanate. The soft segment in commercial polyurethanes
is mostly a polyethermacrodiol (polytetramethyleneoxide) having a
molecular weight of 1000-2000 gmol-'. The chainextender is either
ethylenediamine (in the case of a polyurethaneurea) or l,4-butanediol
(formation of a polyurethane). Linear segmented polyurethanes are
preferrably synthesized in a two-step process, where diisocyanate and
polyol are reacted first, and then chainextended with a diol or diamine.
This method of preparation, in contrast to the one-step process where all
reactants are mixed together simultaneously, leads to blockcopolymers with
better defined structures and better properties. These polyurethane
blockcopolymers exhibit microphase separation due to the incompatibility
of the hard and soft chain segments. Elastomeric behaviour is observed
because the hard domains, having a high glass transition temperature or a
high melting temperature, act as multifunctional physical (hydrogen
bonding) cross-links and as a reinforcing filler in the soft (low T
matrix.
Since these polyurethane elastomers are rnultiphase elastomers they show
very good ultimate mechanical properties compared to one-phase
non-crystallizable elastomeric polymer networks (39,671. In such one-phase
elastomers microcracks, once formed, encounter little resistance to growth
because the network chains are highly mobile. High strength and toughness
result from mechanisms that impede crack growth. An elastomeric material
can only exhibit good ultimate mechanical properties when it consists of
two phases. These two phases normally consist of a rubbery amorphous
matrix containing glassy or crystalline domains or reinforcing filler
particles. In the case of crystallizable elastomers, the second phase is
generated upon stretching. Crystallites formed act as reinforcing domains
in the network and thus increase the ultimate properties. Strain-induced
crystallization provides plastic domains which block, retard or deflect
growing cracks. Filler particles or plastic, hard domains (in
blockcopolymers) may also deform plasticly in high-stress regions, thereby
relieving stress concentrations and dissipating energy.
Elastomeric polyurethanes due to their good mechanical properties (high
tensile strength, good flex life, good tear strength, high toughness),
reasonable bloodcompatibility and biocompatibility, have been used in many
medical applications, such as total artificial heart, heart valves,
vascular prostheses, wound dressings etc. (65).
By mixing segmented polyurethanes with (5-20 wt%) high molecular weight
poly(L-lactide) (PLLA), Gogolewski, Leenslag and Pennings developed
elastomeric, biodegradable mixtures with remarkable in vivo performance.
Quenched physical polyurethane/PLLA mixtures, in porous form, were
successfully applied as a small-caliber vascular prosthesis, artificial
skin, meniscus lesion repair material, nerve guide (68-74).
Since these polyurethanes are not cross-linked through covalent chemical
bonds, but physically through hydrogen bonding, they show stress softening
(stress hysteresis) when subjected to multiple stretching, which is a
serious limitation of these elastomers. This phenomenon is attributed to a
disruption of hard segments with strain, leading to a decrease in their
ability to reinforce the rubbery phase upon strain cycling. This problem
can be overcome by chemically cross-linking the linear polyurethane chains
(75). Another disadvantage of commercial biomedical polyurethanes concerns
their composition. As mentioned before, most polyurethanes are composed of
the aromatic diisocyanate MDI since aromatic polyurethanes show better
microphase separation than those based on aliphatic diisocyanates,
resulting in polyurethanes with superior mechanical properties (65.66).
Degradation (through hydrolysis) of the polymer may result in the
formation of the toxic, carcinogenic, mutagenic MDA, methylenedianiline.
Although it has not been shown unambiguously that MDI-based polyurethanes
induce the formation of cancer, it would be more elegant and safer to seek
for a replacement for this component in the polyurethane formulation. The
use of cycloaliphatic diisocyanates, for instance hydrogenated MDI
(Tecoflex) or 1,4-trans cyclohexanediisocyanate, also leads to segmented
polyurethanes having good ultimate properties. Aliphatic diisocyanates,
which are not that suited for the synthesis of thermoplastic polyurethane
elastomers, may be used for the formation of chemically cross-linked
polyurethanes. Especially aliphatic diisocyanates, producing non-toxic
diarnines (eg. lysine, 1,4-diarninobutane ) after eventual degradation, seem
the ultimate choice for the synthesis of biomedical polyurethanes.
Aim and survey of this thesis
As its title implies, the aim of this thesis is to investigate the
possibilities of polyurethanes, especially cross-linked ones, as
biomaterials. The present thesis can be considered an extension of
previous work from our laboratories on elastomeric
polyurethane/poly(L-lactide) mixtures as biomaterials, with the emphasis
here on polyurethanes. In this thesis the preparation, properties and
medical applications of some new polyurethane networks will be discussed.
In chapter 2 the design and synthesis of biodegradable lysine
diisocyanate-based elastomeric polyurethane networks are discussed. The
polyurethane networks, designed to release only non-toxic degradation
products, are prepared from hexafunctional hydroxy terminated starshaped
copolyesters, synthesized by ring-opening copolymerization of L-lactide or
glycolide and &-caprolactone, initiated by myo-inositol; and cross-linked
with ethyl 2,6-diisocyanatohexanoate (lysine diisocyanatel (761.
Chapter 3 is concerned with the evaluation of a lysine diisocyanate-based
polyurethane network (described in previous chapter) as a material for the
construction of a macroporous bottom-layer (dermal analogue) in a
two-layer artificial skin. In vitro and in vivo degradation studies are
presented (77 1.
Chapter 4 deals with the preparation and evaluation of a microporous
polyetherurethane wound covering having a high water vapour permeability.
The elastic, very thin (15-20 pm) wound covering, prepared by means of a
phase inversion process, has been tested on partial-thickness wounds in
guinea pigs and on donor sites in the clinical situation as well. It is
shown that an accelerated wound healing (enhanced reepithelialization)
under this polyurethane membrane is very likely caused by its high water
vapour permeability (78-85).
Chapter 5 describes a two-ply biodegradable microporous small-caliber
vascular prosthesis composed of polyurethane and poly(L-lactide). The
microporous innerlayer, which is supposed to be highly antithrombogenic,
has been made by cross-linking of a mixture of linoleic acid and a
cycloaliphatic polyetherurethane with dicumylperoxide. The outer ply,
containing much larger pores, has been constructed from
polyurethane/poly(L-lactide) mixtures. The biological performance of the
artificial blood vessels and the effects of cross-linking the innerlayer
with peroxides in the presence of linoleic acid on the antithrornbogenicity
of the prosthesis, the prevention of aneurysm formation and the rate of
degradation are discussed (75).
In chapter 6 the synthesis and properties of densely cross-linked
polyurethane networks, and their potential use in ophtalmology
(intraocular lens, keratoprosthesis) are described. Glassy polyurethane
networks are obtained from the bulk reaction of low molecular weight
polyols and (cyc1o)aliphatic diisocyanates. It is shown that these
optically transparent materials, which can be sterilized by autoclaving,
are rather well tolerated in rabbit eyes (86).
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pub1 ished
Chapter 2
Design and synthesis of biodegradable poly(ester-urethane) elastomer
networks composed of non-toxic building blocks
Summary
Biodegradable poly(ester-urethane) networks, designed to release only
non-toxic degradation products, were prepared from hydroxy terminated
starshaped prepolymers, synthesized by ring-opening copolymerization of
L-lactide or glycolide and E-caprolactone initiated by myo-inositol, and
ethyl 2,6-diisocyanatohexanoate. The poly(ester-urethane) networks, having
T s in the range 0-10 OC and gel contents up to 95 %, showed rubber-like 4 behaviour and after extraction relatively high. tensile strength (30-40
MPa) .
Introduction
Polyurethanes are considered "excellent" biomedical materials possessing
good mechanical and physical properties and showing relatively good bio-
and bloodcompatibility (1). For these reasons segmented polyurethane
elastomers have been used in biodegradable polyurethane/poly(L-lactide)
(PLLA) mixtures for application as vascular prosthesis (2,6), meniscus
prosthesis (31, artificial skin (4) and nerve guide (5). which have been
developed in our laboratory.
The in vivo rate of degradation, after initially observed fragmentation,
of the polyurethane/poly(L-lactide) mixtures appeared to be very low (6).
A further complication was the observation of creep failure upon dynamical
(cyclic) loading as a consequence of stress softening, always associated
with thermoplastic elastomers, which led to aneurysms of the artificial
blood vessels (6).
However, the major drawback of so called biomedical grade polyurethane
elastomers, like Biomer or Estane, used for biodegradable applications, is
their chemical composition. These polyurethanes contain the aromatic
diisocyanate 4.4'-methylenediphenyl diisocyanate (MDI), which is converted
to the toxic, mutagenic, carcinogenic diamine 4,4'-methylenedianiline
(MDA) after degradation ( 7 , s ) . This problem has been overcome by using
Stashaped polyester prepolymer
HO OH
Llnear plyester prepd ymer
(cyc1o)aliphatic diisocyanates like hydrogenated MDI, 4,4'-methylene-
dicyclohexyl diisocyanate HlzMDI in Tecoflex (71 or hexamethylene
diisocyanate (HDI) (91, but the corresponding diamines are still more or
less toxic.
Therefore, it is more elegant to use L-lysine based di-(or tri)isocyanates
for the synthesis of biodegradable polyurethanes. In scheme 1 two
approaches to the design of such degradable poly(ester-urethane) elastomer
networks, which are designed to release only non-toxic degradation
products, are depicted.
Schindler et al. have already reported on the alcohol initiated
ring-opening polymerization of &-caprolactone. In this manner starshaped
polycaprolactone polymers wcre obtained by using sugar alcohols, like
sorbitol, xylitol or ribitol (10). Lately, Pitt et al. have reported on
the synthesis of biodegradable polyurethanes composed of trihydroxy
terminated prepolymers, made by glycerol initiated ring-opening
copolymerization of a 1: 1 mixture of 6-valerolactone and E-caprolactone,
cross-linked with HDI (11). Lipatova et al. have investigated the
(enzymatic) hydrolysis of poly(ester-urethane) networks, containing 20 wtX
sugar as a filler (12). From the literature, degradable copolyesters of
L-lactide or glycolide and &-caprolactone are well known (13, 14).
The polyurethane networks in scheme 1 ( A ) are built up from hexafunctional
hydroxy terminated starshaped polyester prepolymers, synthesized by
ring-opening copolymerization of L-lactide or glycolide and c-caprolactone
initiated by myo-inositol, which can be cross-linked with ethyl
2,6-diisocyanatohexanoate (i.e. lysine diisocyanate). The degradation
products, myo-inositol, a vitamin widely distributed in the human body
(151, L-lactic acid or glycolic acid, 6-hydroxyhexanoic acid, L-lysine and
ethanol, which are set free upon biodegradation of this polyurethane
network, are all non-toxic which is very essential for the use as a
degradable biomedical material. This chapter reports in more detail on
these networks. The other polyurethane network (scheme 1 (B)) consists of
linear polyester prepolymers, dihydroxy terminated copolyesters of
L-lactide or glycolide and e-caprolactone, using 2-isocyanatoethyl
2,6-diisocyanatohexanoate as a cross-linking agent (16).
Experimental part
Prepolymer synthesis
L-lactide (C. C. A. Gorinchem, The Netherlands; recrystallized from dry
toluene) or glycolide (DuPont) and e-caprolactone (Janssen Chemical,
Belgium; distilled) and myo-inositol (Merck) were dissolved in dry DMF at 0 140 C. Stannous octoate (Sigma Chem. Corp. USA; 0,5 wt%) was added and
0 polymerization was carried out for 20 h at 120-130 C under nitrogen
atmosphere. After removal of the solvent i. vac. a tacky, yellowish
prepolymer resulted, which could be precipitated in ethanol (-70 O C ) from
chloroform solution and subsequently dried i. vac. at ambient temperature.
Ethyl 2,6-diisocyanatohexanoate (17)
L-lysine monohydrochloride (Janssen Chemical, Belgium) was first converted
to L-lysine ethyl ester dihydrochloride by refluxing in ethanol while
passing HC1 gas through the solution. The dihydrochloride was phosgenated
in o-dichlorobenzene or dioxane at 100-110 OC for ca. 8 h. The crude
diisocyanate was purified by vacuum distillation (bp. 125 OC/O,I mmHg).
Network formation
Prepolymers poly(L-lactide-co-E-caprolactone)~ were cross-linked by
treatment with ethyl 2,6-diisocyanatohexanoate ([OHI/[NCOl = 1) in
toluene, and poly(glyco1ide-co-E-caprolactone) prepolymers in CH C1 . Thin 2 2
films were obtained by reaction in a Petri-dish at room temperature under
nitrogen for one day and post-curing at 100-110 OC for 3 h. The elastic, 0 transparent films were dried at 50 C i. vac. Porous materials were
synthesized by curing a viscous slurry of the prepolymer, diisocyanate,
solvent and an amount of dried NaCl powder of variable particle size by
the method described previously. Afterwards the salt was removed by
washing the NaC1/ polymer mixture with water.
Characterization
Gel contents (in wt%) were determined by extraction of the networks with
chloroform. The extracted networks were first carefully air-dried and then
dried i. vac. at 50 OC to constant weight.
Swelling measurements were carried out on extracted networks in chloroform
at room temperature. The degree of swelling was calculated from the weight
increase after swelling using the densities of chloroform (p=1,48 g/cm3)
and the dry extracted networks (p=0,90-0,95 g/crn3).
Thermal analysis of the networks was performed by means of a Perkin-Elmer
DSC-7, calibrated with I. C. T. A. (International Confederation of Thermal
Analysis) certified reference materials and operated at a scan speed of 10
'chin.
Mechanical properties were determined at room temperature using an Instron
(4301) tensile tester equipped with a 10 N load-cell, at a cross-head
speed of 12 mm/min. Specimens (15 x ca. 0,75 x ca. 0,25 mm) were cut from
(unlextracted thin films.
An I. S. I.-DS 130 scanning electron microscope was used to study the
microstructure of the porous materials.
Results and discussion
Poly(ester-urethane) networks were formed by treatment of hexahydroxy
terminated starshaped prepolymers with ethyl 2.6-diisocyanatohexanoate.
Prepolymers were synthesized by ring-opening copolymerization of L-lactide
or glycolide with c-caprolactone in a 1: 1 mole ratio, initiated by
myo-inositol (hexahydroxycyclohexane) using stannous octoate as a
catalyst.
A rather unique aspect of these polyurethane networks is the use of ethyl
2,6-diisocyanatohexanoate, which has not been reported that of ten in the
literature (21-23). Besides the fact that L-lysine is the degradation
product of the incorporated diisocyanate, hydrolysis first of the ethyl
ester results in the introduction of a carboxylic group into the network.
The choice of prepolymers poly(L-lactide(or glyco1ide)-co-c-caprolactone)~
was based on the idea to obtain elastomeric polyurethanes exhibiting a
high rate of degradation (see next chapter). Polyurethane networks
composed of prepolymers poly(L-lactide (or glyco1ide)-co-myo-inositol)~
have glass transition temperatures above room temperature. Therefore,
copolyester prepolymers containing L-lactide (or glycolide) and
E-caprolactone in a 1:l mole ratio were used in order to obtain
poly(ester-urethane) elastomer networks, having T values far below room 9
temperature. The other extreme, polyurethane networks composed of only
poly(&-caprolactone) prepolymers are expected to degrade more slowly than
the co-poly(ester-urethane) networks. For some applications, however, a
low rate of biodegradation seems desirable ( 3 ) .
In table 1 some relevant data of the poly(ester-urethane) networks are 0 collected. The T values of the networks were in the range of 0-10 C,
g
depending on the branch length. The T can be lowered by increasing the B
branch length of the copolyesterprepolymers. The T of a network was '3
raised after extraction with chloroform. Remnants of unreacted monomers
(and oligomers), which were removed by the extraction procedure had a
plasticization effect on the networks. These remnants also had to be held
partially responsible for the observed gel content. Polyurethane networks
with the highest gel contents (95%) were obtained by using precipitated
prepolymers. Even higher gel contents can be expected by using a slight
excess of isocyanate groups ([OHI/[NCOI < 1). Besides urethane bond
formation excess cross-linking can take place through formation of
allophanate groups (18).
Table 1. Poly(ester-urethane) network data
~ 0 1 ~ - prepolymer T gel elongation tensile degree g
urethane branch content at break strength of
networka) length b, (OC) ( % I ( % I (MPa) swelling c
a#l=~repolymer poly(myo-inositol-co-glycolide-co-c-caprolactone + ethyl 2,6-diisocyanatohexanoate; #2=extracted network 1; #3=precipitated
prepolymer poly(myo-inositol-co-glycolide-co-e-caprolactone + ethyl
2,6-diisocyanatohexanoate; #4=extracted network 3; #5=prepolymer
poly(myo-inositol-co-L-lactide-co-c-capro1actone) + ethyl 2.6-diiso-
cyanatohexanoate; #6=extracted network 5.
b)~ranch length: number of lactone molecules (L-lactide, glycolide, c-
caprolactone) per OH group of myo-inositol, calculated from the initial
proportions of starting materials employed.
"1n chloroform at 20 OC.
Fig 1 . Stress-strain curves of poly~glycolide-co-~-capr01actone~-
urethane networks before (-----I and a i t c r (- 1 ~xlraction with
chloroform (respectively networks 3 and 4 i n t a b l e 11.
Fig. 2 . Scannrng t : l t . c t r u n mzcruzraph L,: J F ~ I c u s p . ~ l y ( e s l p r - ~ r e thane
matrix.
The networks were also characterized by their degree of swelling in
chloroform, which ranged from ca. 3,O for the glycolide based networks to
4,75 for the L-lactide based networks.
Fig. 1 shows typical stress-strain curves of poly(glyco1ide-co-c-
caprolactone) networks before and after extraction with chloroform. All
the polyurethane networks showed rubber-like behaviour, but from table 1
and fig. 1 it is clear that the extracted polyurethane networks exhibit
better tensile properties, increased elongation at break and higher
tensile strength (30-40 MPa). Only the extracted networks exhibit
pronounced strain-induced crystallization. Crystallites thus formed have a
reinforcing effect within the network, and thus increase its ultimate
strength and maximum extensibility. The presence of diluent (plasticizer)
suppresses the strain-induced crystallization and thus diminishes the
ultimate properties (20).
Fig. 2 shows a scanning electron micrograph of a porous
poly(ester-urethane) matrix, which was obtained by curing of prepolymer
poly(L-lactide-co-e-caprolactone) with ethyl 2,6-diisocyanatohexanoate in
the presence of an amount of salt (pore volume ca. 85%). In a very
straight forward way (salt casting method) porous materials of these
polyurethane networks for degradable biomedical applications can be
constructed. Preliminary experiments in guinea pigs have shown that the
poly(ester-urethane) networks biodegrade when implanted subcutaneously
(19). Concluding, we state that degradable poly(ester-urethane) networks,
designed to produce only non-toxic degradation products, as described
here, are very promising biodegradable materials. Further work and
especially in vitro and in vivo degradation studies are in progress (19).
References
1. M.D. Lelah, S. L. Cooper, "Polyurethanes in Medicine", CRC Press,
Boca Raton, Florida. 1986
2. S. Gogolewski, A. J. Pennings, Makromol. Chem., Rapid Commun., 3, 839
( 1982)
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Nielsen, H. W. B. Jansen, Proc. PIMS V, Noordwi jkerhout, The
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( 1983
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Pennings, Biomaterials, 11, 286 (1990)
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Chapter 3
Biodegradable lysine diisocyanate-based poly(glyco1ide-co-c-capro1actone)-
urethane network in artificial skin
s-Y A biodegradable lysine diisocyanate-based poly(glyco1ide-co-c-
capro1actone)urethane network has been evaluated as a material for the
construction of a macroporous bottom-layer (dermal analogue) in a
two-layer artificial skin.
High rates of in vitro degradation were observed; degradation of the
porous poly(glyco1ide-co-E-capro1actone)urethane networks was faster in
vivo than in vitro.
Subcutaneous implantation in guinea pigs showed that the porous
polyurethane networks allowed rapid cell ingrowth, degraded almost
completely between 4 and 8 weeks after implantation and evoked no adverse
tissue reaction.
Introduction
Recently we showed that epidermal wound healing of partial-thickness
wounds was accelerated when covered with a microporous polyetherurethane
membrane of high water vapour permeability 1 2 The healing of
full-thickness wounds is much more complicated because there are almost no
epidermal islands left in the wound-bed where skin regeneration may
commence. These wounds heal primarily by wound contraction, resulting in
scar formation ( 3 ) . Therefore a two-layer artificial skin is needed,
comprising a macroporous, biodegradable bottom-layer functioning as a
scaffold for skin regeneration, which enables fibrovascular ingrowth and
which should be resorbed when cell ingrowth is complete; and a
non-degradable top-layer, providing a barrier against infection and
optimal water vapour permeability, which can be peeled off the wound after
healing. This may be combined with seeding epidermal cells in the
bottom-layer (stage 2 artificial skin). This concept for covering
full-thickness wounds was originated by Yannas and Burke ( 4 , s ) . It has
also been described by Gogolewski and Pennings who used polyurethane/
poly(L-lactide) mixtures to construct a two-layer biodegradable artificial
skin ( 6 ) . However, the elastomeric polyurethanes used do not seem to be
ideal for biodegradable applications for two reasons. First, the rate of
degradation is too low, which is especially a problem in case of
applications like biodegradable artificial skin when a high rate of
degradation is desirable. Second, the segmented elastomeric polyurethane
is capable of releasing the toxic, carcinogenic methylenedianiline upon
degradation, as a result of the incorporated aromatic diisocyanate MDI
(7.81.
To overcome these problems we have developed new lysine diisocyanate-based
polyesterurethane elastomer networks, designed to degrade rapidly, thereby
releasing only non-toxic degradation products as outlined before ( 9 ) .
OH
HO
OH
Glycolide + c-caprolactone + myo-inositol ----+
Poly(glyco1ide-co-c-caprolactone) prepolymer
1 HO OH
0 + QCN-FH-$ lysine diisocyanate I$H21r OEt NCO
Poly(glyco1ide-co-s-caprolactonelurethane network
Figure 1. Synthesis of polyesterurethane networks.
Figure 1 shows how these polyesterurethane networks are built up. In
short, hexahydroxyterminated starshaped poly(g1ycolide-co-E-caprolactone)
prepolymers are synthesized by ring-opening copolymerization of glycolide
and E-caprolactone initiated by myo-inositol using stannous octoate as a
transesterification catalyst (9,101. These prepolymers are cross-linked
with 2,6-diisocyanato ethylhexanoate (referred to here as lysine
diisocyanate) to form poly(glyco1ide-co-E-capro1actone)urethane networks.
This chapter reports on the preparation, the physical characteristics and
biological performance after subcutaneous implantation of a porous
bottom-layer of a two-layer artificial skin, composed of this lysine
diisocyanate-based polyesterurethane network; and also how this can be
combined with the previously described polyetherurethane top-layer having
high water vapour permeability to form a stage 1 artificial skin (4.5).
Experimental part
Synthesis of prepolymer poly(glyco1ide-co-&-caprolactone)
Poly(glyco1ide-co-E-caprolactone) prepolymers were synthesized as
described elsewhere (9) .
Two prepolymers differing in glyco1ide:c-caprolactone ratio were
synthesized. Prepolymer A contained glycolide and z-caprolactone in a 1:l
mole ratio, with a calculated branch length of 6 lactone units (i.e.
glycolide or E-caprolactone) per OH group of myo-inositol. Prepolymer B
was synthesized from a 1:1.7 glyco1ide:c-caprolactone feed mole ratio. The
calculated branch length was 7.6 lactones per OH group of myo-inositol.
Porous poly(glyco1ide-co-E-capro1actone)urethane network (bottom-layer)
Poly(glyco1ide-co-E-caprolactone) prepolymer was dissolved in
dichloromethane and freshly distilled lysine diisocyanate (synthesized as
described elsewhere (9)) was added ([OHI/[NCOI = 1). This solution was
mixed with an amount of dry NaCl particles, resulting in a very viscous
slurry. The volatile solvent was allowed to evaporate during this process.
The mixture was then poured into a Petri-dish and extra salt was sprinkled
on top to avoid skin formation. Cross-linking reaction was carried out at
room temperature for one day while all of the solvent was allowed to
0 evaporate and post-curing at 100-120 C for at least 5 h. under nitrogen
atmosphere. Afterwards, the salt was leached out with water and a porous,
sponge-like sheet resulted after subsequent air drying. The porous
networks were extracted with chloroform and dried carefully to constant
weight, from which gel contents (in wt%) were calculated. The porevolume
of the porous materials was calculated from the weight ratio of the
(prepolymer + diisocyanate) and salt.
Two-layer artificial skin
The method for the construction of the polyetherurethane top-layer has
been described earlier (1). This porous PEU top-layer could be glued to
the porous bottom-layer (thickness ca. 2 mm) by using a viscous PEU
solution in THF which was cast in a thin layer onto the top-layer and
subsequently glued to the bottom-layer, dried and placed in water. Thus a
two-layer membrane was constructed.
In vitro degradation
Porous, extracted poly(glyco1ide-co-E-capro1actone)urethane network
samples, pore size 90-250 pm, porevolume 80 % (1 x 1 cm x 2 mm) were
subjected to degradation at 37 + 1 OC in phosphate buffer, pH=6.9.
Degradation was monitored by determination of the weight change.
An I.S.1.-DS 130 scanning electron microscope was used to study the
structure of the porous materials.
In vivo degradation and cell ingrowth
Strips (2 x 2 x 10 mm) of porous lysine diisocyanate-based
poly(glyco1ide-co-e-capro1actone)urethane (pore size 90-250 pm, pore
volume 80 %) were subcutaneously implanted in the dorsum of guinea pigs
(n=4), weighing between 300 and 400 grams. Each animal received six
strip-implants, three based on prepolymer A and three based on prepolymer
B. Every strip was implanted via a separate incision in a surgically
created pocket underneath the panniculus carnosus. The cutaneous incisions
were closed by interrupted 6-0 polyglycolic acid sutures.
The animals were sacrificed 2, 4 , 8 and 12 weeks after implantation and
the location of the implants was identified by blunt dissection of the
complete dorsal skin from the underlying fascia of the paravertebral
muscles. The implants were harvested by wide excision with scalpel and
immersion fixed in 10 % formalin.
The specimens were histologically processed as described previously (17).
Briefly, the specimens were embedded in glycol methacrylate resin, cut
perpendicularly to the axis of the strip-implant in 2 pm thick sections
(thus visualizing the initial 2 x 2 mm cross-surface area), and stained
with Sudan black B and hematoxylin. The Sudan black B stains the polymer
material dark green. Sections were photographed with a Zeiss
Photomicroscope 111.
Results and discussion
All lysine diisocyanate-based poly(glyco1ide-co-c-capro1actone)urethane
networks synthesized according to figure 1, having gel contents 92-95 %,
were extracted with chloroform to remove unreacted monomers (and
unreactive oligomers), which function as swelling agents. The extracted
networks show significantly improved ultimate mechanical properties
(tensile strength, elongation at break) in comparison with the unextracted
networks. Strain-hardening only displayed by the extracted networks is
apparently due to strain-induced crystallization, which is hindered by
swelling agents, even if present in small amounts (the sol fraction
comprises only a few percentages) (9,181. Another reason for extracting
the networks, besides improvement of the mechanical properties, is the
removal of the residual monomers like c-caprolactone, which might give
rise to undesired tissue reactions when implanted.
Porous, sponge-like materials, in sheet form, were obtained by "in-situ"
cross-linking of the prepolymers with lysine diisocyanate in the presence
of an amount of NaCl particles (saltcasting method1 and afterwards
leaching out the salt with water. Fig. 2 shows a scanning electron
micrograph of a porous poly(glyco1ide-co-c-capro1actone)urethane network,
with a mean pore size of 90-250 pm and a pore volume of 80 %, prepared by
saltcasting. It can be seen that by using this saltcasting method an open
porous structure was obtained. As stated in the introductory part, this
porous poly(glyco1ide-co-E-capro1actone)urethane network should function
as a bottom-layer of a two-layer artificial skin allowing fibrovascular
ingrowth and thus has to exhibit an open pore structure.
Another important characteristic of such a bottom-layer is its
degradability. Once the cell ingrowth is complete the porous scaffold has
no function anymore and ideally be resorbed from this moment on, i.e.
after ca. 3-4 weeks. So the material used, should exhibit a high rate of
degradation. For this reason a prepolymer composed of glycolide building
blocks was chosen for the formation of a polyesterurethane network, since
it is known that polyglycolide and its copolymers show a high rate of
degradation (11,12) as compared with poly(L-lactide), for instance.
Hydrolysis of semicrystalline polyesters first takes place in the
amorphous regions, followed by degradation in the crystalline phase (12).
Therefore, it is concluded that purely amorphous polyesters will show a
high rate of degradation as confirmed by the work of Gilding and Reed on
amorphous poly(L-lactide-co-glycolide) (11) and Schindler and Pitt on
amorphous, cross-linked elastomeric poly(valero1actone-co-E-caprolactone)
(13.14,15).
To obtain elastomeric, amorphous polyesterurethane networks, prepolymers
had to be built up from glycolide and E-caprolactone, since
polyesterurethanes from only polyglycolide prepolymers have a too high T .
From the literature, linear copolyesters of glycolide and c-caprolactone
are known (16). Incorporation of &-caprolactone into the branches of the
starshaped prepolymers lowered the T as compared with pure 4
polyglycolide-based branches, so that elastomeric polyurethane networks
with T far below roomtemperature resulted. Besides lowering the '3
glasstransition temperature, crystallization of polyglycolide (which may
happen when the branchlength is long) will be suppressed. All
poly(glyco1ide-co-c-capro1actone)urethane networks were amorphous as
observed by DSC.
Figure 2 . Scanning eIectron micrn~raph n f a porous lysine
diisocyanatc-baecd po!y(glycolide-co-r-caprolsctone)tlretha~~c n~twcrk.
wi tll a mean pore size of 9C-250 pm and a pore volume of 80 %,
prepared by salt-castlng.
Figure 3. Scarrnirlg r , lecL~ cj11 [ I I ~ C I ugt dph oi t h ~ . po rous polyurethane
network ( A ) depxctcd in f i g u r e 2 degraded i n v l t r o f o r 4 wk.
Table 1. In vitro degradation of porous lysine diisocyanate-based poly-
(glycolide-co-E-capro1actone)urethane networks A and B
Time (weeks) X weight loss
Table 1 summarizes the weight loss observed under in vitro conditions for
two elastomeric, porous lysine diisocyanate-based poly(glyco1ide-co-
c-capro1actone)urethane networks A and B with prepolymer
glyco1ide:c-caprolactone feed mole ratio of 1: 1 (prepolymer A) and 1: 1 . 7
(prepolymer B ) , respectively. Significant weight loss occurred after 2
weeks already for A, whereas B showed a comparatively delayed degradation
pattern, because B was built up from the E-caprolactone-richer prepolymer.
Thus by varying the feed mole ratio of glyco1ide:c-caprolactone in the
prepolymer, the rate of degradation of the resulting polyurethane network
can be controlled, due to the greater hydrophobicity of c-caprolactone
units compared with the relatively hydrophilic glycolide units.
Figure 3 shows a scanning electron micrograph of the same porous network A
as in fig. 2 degraded in vitro for 4 weeks. Besides 15 % weight loss (see
table 11, the porous network had also degraded visually. The porous
structure had partially collapsed and the sharp edges had been smoothed.
Two weeks later the porous structure had completely collapsed and the
network had turned into a tacky, chloroform-soluble polymer. Again,
network B showed a delayed degradation in comparison with A. High rates of
degradation in vitro were observed, owing to the amorphous nature of the
polyesterurethane networks (13). Random hydrolytic chain cleavage
apparently caused immediate weight loss, because of the formation of
water-soluble degradation products.
In contrast to the in vitro degradation results, the in vivo results
showed no difference in rate of biodegradation between samples made from
prepolymer A and prepolymer B. All porous implants showed signs of
degradation after four weeks implantation, reflected by distortion of the
outer dimensions, erosion and "foaming" of interporous walls (Figure
4a,b). After eight weeks the implants were almost completely degraded
(Figure 4c,d). At that time polymer remnants had lost their affinity for
Sudan black B and appeared as transparent particles which had been
engulfed by multinuclear giant cells. No polymer particle could be
detected after twelve weeks implantation, nor could the implantation sites
be identified. Histological evaluation of "blindly" taken twelve-week
tissue samples did not show polymer material or scar tissue.
Cell ingrowth was already seen in the two-week samples. Cells had filled
the complete labyrinth of micropores and consisted of macrophages,
epithelioid cells, fibroblasts and endothelial cells. The endothelial
cells had a lumen by that time, thus forming capillaries deep in the pores
of the implant. These capillaries had grown to 40 pm wide vascular
structures by the eigth week (figure 4d). In the course of weeks,
epithelioid cells predominated the infiltrate fusing into multinuclear
giant cells.
The poly(glyco1ide-co-c-capro1actone)urethane material can be considered
biocompatible, since no adverse tissue reactions developed. The implants
did not evoke any granulocyte-mediated inflammatory reaction. A thin
fibroblast layer initially encapsulated the polymer strip, but merged in
the surrounding loose connective tissue by the eigth week (figure 4c).
Connective tissue fibers, identified as being type I11 collagen using
Herovici's connective tissue stain, had been deposited, probably by
fibroblast, into the pores of the implant by the fourth week.
Figure 4.
Histology of cross-section of porous biodegradable lysine
diisocyanate-based poly(glyco1ide-co-e-capro1actone)urethane networks, 4
(a,b) and 8 (c,d) weeks after subcutaneous implantation in the guinea pig.
The implant was originally 2 x 2 mm in cross-section. (Sudan black B and
hemotoxylin; skin is to the top; bars represent 100 pm).
a. Four weeks after implantation: the polymer strip has colllapsed to 0.6
x 1.9 mm and is encapsulated by a thin fibroblast capsule. Polymer
fragments of interporous walls (PI are separated by ingrowing cells and
blood vessels (BV).
b. High power view of a. Epithelioid cells (E) and multinuclear giant
cells (MNG) engulf the polymer particles. The two polymer fragments
present at the bottom of the figure ( P I show "foaming" as a result of
resorption.
c. Eight weeks after implantation (note the same magnification as in a. 1:
The polymer strip has now collapsed to 0.09 x 1.9 mm and is penetrated by
numerous blood vessels (BV). The fibrous capsule has been absorbed by the
surrounding loose connective tissue.
d. High power view of c. Polymer particles (PI have been engulfed or
phagocytosed by multinuclear giant cells (MNG) and do not stain anymore
with Sudan black B. Note the numerous blood vessels (BV) involved in the
process of degradation.
Degradation of the porous poly(glyco1ide-co-E-capro1actone)urethane
samples was faster in vivo than in vitro. This faster biodegradation might
well be explained by the mechanical strain of ingrowing cells, the
additional effect of biologically available enzymes, and the subsequent
intracellular degradation of small polymer particles.
In conclusion, this new biodegradable poly(glyco1ide-co-c-caprolactone)
urethane seems promising as a material for the construction of a
macroporous bottom-layer, with a mean pore size of 90-250 pm (dermal
analogue) in a two-layer artifical skin, since it evokes no adverse tissue
reactions, allows rapid cell ingrowth and degrades almost competely
between 4 and 8 weeks after implantation. Futher studies are planned to
examine the efficacy of the two-layer artificial skin in a full-thickness
wound model in the Yorkshire pig.
References
1. P. Bruin, M. F. Jonkman, H. J. Meijer, A. J. Pennings, J. Biomed. Mater.
Res., 24, 217 (1990)
2. M.F. Jonkman, P. Bruin, E . A . Hoeksrna, P. Nieuwenhuis, H. J. Klasen,
A. J. Pennings, I. Molenaar, Surgery, 104, 537 (1988)
3. W. van Winkle, Surg. Gynec. Obstet., 124, 369 (1967)
4. I.V. Yannas, J.F. Burke, J. Biomed. Mater. Res., 14, 65 (1980)
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and A. S. Hoffman), Ni jhoff Publishers, The Netherlands (19861, p. 221
6. S. Gogolewski, A. J. Pennings, Makromol. Chem., Rapid Commun., 4, 675
(1983)
7. M. Szycher, V.L. Poirier, D. J. Dempsey, Elastomers and Plastics, 15,
81 (1983)
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( 1987 )
9. P. Bruin, G. J. Veenstra, A. J. Nijenhuis, A. J. Pennings, Makromol.
Chem. , Rapid Commun. , 9, 589 (1988)
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(1982)
11. A.M. Reed, D.K. Gilding, Polymer, 22, 494 (1981)
12. D.F. Williams, J. Mater. Sci., 17, 1233 (1982)
13. C.G. Pitt, R. W. Hendren, A. Schindler, S. C. Woodward, J. Contr.
Release, I, 3 (1984)
14. A. Schindler, C.G. Pitt, Polymer Preprints, 23(2), 111 (1982)
15. C. G. Pitt, A. E. Schindler, "Biodegradable polymers of lactones". U. S.
Patent 4.379.138.
16. H. R. Kricheldorf, T. Mang, J. M. Jonte, Macromolecules, 17, 2173 (1984)
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Chapter 4
A new porous polyetherurethane wound covering
Summary
A polyetherurethane (PEU) wound covering with non-interconnected
micropores up to approximately 5 pm has been prepared by means of a phase
inversion process. This highly elastic, very thin (15-20 pml, pliable
wound covering showed good, immediate adherence to wet wound surfaces and
high water vapour permeability, but was impermeable to bacteria.
In guinea pigs epidermal wound healing of partial-thickness wounds under
PEU wound coverings was accelerated as compared with uncovered controls
and an occlusive wound covering, OpSite. Water in liquid form or wound
exudate could not leak through the PEU covering, but its high water vapour
permeability induced concentration of the wound exudate into a jellylike
clot layer, which apparently accelerated reepithelialization.
The main conclusion from a clinical study on 20 donor sites was that the
use of the PEU covering reduces pain, besides prevention of fluid
retention and enhanced reepithelialization.
Introduction
The concept of a two-layer artificial skin for covering full-thickness
(burn) wounds, consisting of a biodegradable, porous bottom-layer
functioning as a temporary template for skin regeneration and a protective
non-degradable top-layer was first described by Yannas and Burke (1).
Gogolewski and Pennings constructed an artificial skin based on
polyurethane/poly(L-lactide) mixtures using the above concept (2).
Our current concept of the synthetic skin substitute consists of a
microporous vapour permeable polyetherurethane (PEU) top-layer and a
separate bottom-layer, composed of a biodegradable polyesterurethane
elastomer network, which is designed to degrade rapidly to non-toxic
degradation products, not having the disadvantages like low rate of
degradation and release of the toxic, carcinogenic, aromatic diamine
4.4'-methylenedianiline upon degradation of the segmented polyurethane,
associated with the PU/PLLA mixture (3).
The PEU top-layer in itself can also serve as a wound covering, acting as
a temporary covering of donor sites and second degree burns
(partial-thickness wounds).
It is commonly accepted that such a wound covering should be adherent,
elastic, pliable, impermeable to bacteria, easy to handle, non-toxic,
hemostatic and also allow the proper water vapour transport through the
covering (4,5). However, no data concerning the optimal water vapour
permeability of wound coverings, to prevent desiccation of the wound and
avoid simultaneously fluid accumulation under the covering, are available
(5.6). Despite the fact that many wound coverings are commercially
available (e.g., OpSite, Biobrane, Omiderm, Epigard) many surgeons still
prefer treatment of donor sites and second degree burns with conventional
methods like tulle gras dressing, which may indicate that the ideal.
synthetic wound covering has yet to be developed (5.8). This chapter
reports on the preparation and characteristics of this new
polyetherurethane wound covering.
Materials and Methods
PEU wound covering
In this study Tecoflex EG-80A (9) (Thermedics Inc. a so called second
generation medical-grade segmented cycloaliphatic elastomeric
polyetherurethane (PEU) was used as supplied.
A 5% (w/w) PEU solution in tetrahydrofuran (THF), containing 1% (w/w)
lithiumchloride (LiC1) was refluxed for 2 h under nitrogen atmosphere and
subsequently stirred for 2 h at room temperature. A film of this polymer
solution was cast on a glass plate, using a 500 pm cast-iron. The glass
plate was placed just above a water layer in an open box. Within some
minutes the clear film turned white. Half an hour later on, when most of
the solvent had evaporated, the non-translucent film was dried in vacuo at
50 OC for 2 h and subsequently placed in water to detach the film from the glass plate and to remove the salt. Finally the white elastic film was
dried. The PEU covering was sterilized by means of gamma radiation (25
kGy) .
Porous PU membrane
A Petri-dish containing a 7% (w/w) polyurethane (Estane, Goodrich U.S.A. )
solution in 1,4-dioxane was placed above a 1:l v/v water/l,4-dioxane
mixture in a closed box. After 2 days the porous membrane was placed in
water and subsequently dried.
Water vapour permeability
The wound covering was stretched in a screwed open cap onto a glass cup,
which was partially filled with water and inverted so that the wound
covering was in direct contact with water.
A ServoMed evaporimeter (Model EP-lC, ServoMed AB, Vallingby, Sweden) was
used for measuring the water vapour transmission rate (WVTR in gm-2h-1) at
various water vapour pressure differences across the wound covering. The
measurements were performed in a closed cabinet to prevent disruptive air
currents as described elsewhere in detail by Erasmus and Jonkman (6,7).
Stress-strain measurements were performed on cut specimens ( 5 x 20 mm)
from the above described dry PEU film and the commercial wound covering
OpSite (10) (thickness 28 pm, Smith & Nephew Ltd., Hull, U.K. at room
temperature using an Instron (4301) tensile tester equipped with a 10 N
load-cell, at a cross-head speed of 50 mm/min.
An I.S.1.-DS 130 scanning electron microscope was used to study the
microstructure of the membranes.
Animal and clinical studies
Under sterile conditions, at the back of each of 61 guinea pigs two
partial-thickness wounds (2 x 2 cm) with a mean depth of 0,32 mm were made
with a dermatome. The wounds were either covered with a dry, sterilized
PEU membrane or with OpSite or were left uncovered. The coverings were not
changed during the experiment. All wounds were evaluated histologically 1
to 14 days after excision as described in detail previously (12)
Twenty adult burn wound patients undergoing split-skin graft procedures at
the Burns Centre of the Roman Catholic Hospital of Groningen were
candidates for the clinical study. Skin grafts with a mean thickness of
0,30 mm were taken with a dermatome. Half pf the donor site was covered
w i t h the PZU cuverinq, tlre crther h a l f w l t t ~ = s l n p l e laycr of tuIle eras
drrs9lne (paraffin gauze) . The complete d a ! l ~ r rile w , ~ s cover~d w i t h f o u r
absorptive co t t on pads, comprising 32 s Lng!r l a y ~ r s f i n r rnrmsh ~ a i ~ p , hn'd
i n place with a crepe bandage The c g t t o n ~ d d s a-d crepe handage above t h e
PEU cover in^ uere cut abay hr twccn 1 and 5 days a f t e r npet-ation, thus
allowing free ventilation. T h e FEU cuvur i r ) ~ w , l s peeled o f f th i l bound
between 5 and 21 d a y s a f t e r operatlun '[ tw crepe bsndnjie and Lau7t.s or. t o g
of t h e t u l l e eras were removcd b v t w ~ ~ r , 7 a n d 10 d a y s a f t c t nperat i n n
B i n p s l e s were t a k ~ n between 3 days to 3 m o n t h s af t u r o p c r z l i u r ~ and sl~died
histologically as drscritcd e l s c h h e r e in dctall (15)
Results and discussion
Figure 1 presents a s c a n n i n e electrotj micrograph of a PtJ membrane,
possessing a w r y rc.gular pore q t s t i c t u r ~ I t car1 be s e m t h ~ t t h e large
pores are interconnected with c m a l l e r o n e s . l h i s pornus mrmbrane was
obtained by s l o ~ cvapuratiurt u f k l i c - sc7!vur~? (1,4-dioxane] and
slmu1:aneously s l o w diffuszon n f water vapou1- ( r w n s r> lven t ) i n t o t h e
polymer solution.
F i w r e 1. Scann lne e l e c t r o n rnIcrr,gi,?ptl o: a PLI rnerrbranr prt-pared f rom
a 7 utX polymer solutie? i n 1.4-dinxzne us in€ a I : 1 v/v 1.4-dinvan?/
w a t ~ r snlvent/nonsol v ~ n t mixtur-e. N o t e t h r I - P R U ~ ~ I - pure s t r u c t u r e .
I ' lgure 2. Scannine e l e c t r o n rnicrazral>h of a PEU mr:mbr-arw pr cparrzd by
c a s t i n g a fiPn from, a 5 w t % polymer solution i n TIIF, containing i w t X
LIC1. The rough topslde of t l e r n ~ m b r a n ~ i s t h e bo t to rns idc o f the FEU
wound covering, which faces the wound.
Figure 3 . Scanning electron micrograph u f a crcss-section of t h e PFU
wound covering.
By using a more volatile solvent (THF) combined with the presence of the
hygroscopic lithiumchloride in the polymer solution, which might attract
water vapour, the process of phase inversion was accelerated, which
resulted in a different, less regular membrane structure.
Figure 2, 3 and 4 show scanning electron micrographs of the dry PEU wound
covering, prepared by casting a film of a 5 wt% polymer solution in THF,
in the presence of 1 wt% LiCl. The topside of the PEU membrane, which
actually is the bottomside of the wound covering facing the wound, showed
a rough surface, composed of pits and small pores up to approximately 5 pm
(fig. 2). The membrane of thickness 15-20 pm contained micropores up to
ca. 5 pm, which were not interconnected (fig. 3 ) . The bottomface of the
membrane (i.e. the topside of the wound covering), which was originally
stuck to the glass plate, showed a smoother surface with pores up to ca. 5
Mm (fig. 4). As a result of this structure water in liquid form (or wound
exudate) could not leak through the PEU membrane, but water vapour
diffused at a high rate (see below). Due to the porous structure and its
hydrophilic characteristics the PEU membrane showed good adherence to wet
surfaces. The dry PEU wound covering immediately adhered well to
partial-thickness wound beds, thereby sealing the wound and forming a
barrier against infection, since it was shown that the covering was
impermeable to bacteria (12).
Figure 5 presents the water vapour transmission rate (WVTR) vs. the water
vapour pressure difference across two wound coverings, namely the PEU
wound covering and OpSite, measured in vitro, "upside down" (6.7) . The
slopes of the lines empirically found are designated water vapour
permeance (WVP) . The water vapour permeability of the PEU wound covering, expressed as
water vapour permeance was 20,l gm-2h-'k~a-', which was much higher than
the corresponding value found for OpSite (5.3 gm-2h-1k~a-'), which is
considered an occlusive wound covering ( l o ) , but less than Omiderm (24.6 gm-2h-1k~a-' ) , a commercially available polyurethane wound covering which has been grafted with polyacrylamide and which is known to be a high
vapour permeable wound covering (14). It is known that the use of
occlusive wound coverings in partial-thickness wounds accelerates
reepithelialization. stimulates wound healing, prevents wound desiccation
P N
ca.
n t I 3 L 5 r; Wnler vopor pressure d,Hermce (k%l
F l y r e 5. Wntar vapour p 6 r t r o a b i l i t y o f wound coverin~s: Water vapokir
transmission rate (WVTR) a s a f u n c t i ~ n of the water vapour pressure
diffcrcnce across t hc wound covering. The slopes of the individual -2 -1
lines are designated water vapour per rrlt,ance Igln 11 k ~ a - I ) .
and body heat loss, but has the disadvantage of accumulation of wound
exudate under the covering, which may lead to infection, especially in
case of occlusive film dressings like OpSite (10,181. It is suggested that
a more water vapour permeable wound covering, like the constructed PEU
membrane, will avoid the latter.
Figure 6 shows typical force-strain curves of the dry PEU wound covering
and for comparison OpSite, a commercial, clinically used wound covering.
It is clear that the dry PEU wound covering was far more elastic than
OpSite, having a lower modulus, but exhibiting lower strength at break.
A wound covering should be elastic if it has to be applied over bending
surfaces like over joints, to facilitate an intimate cover of the wound.
The thickness of the PEW membrane (15-20 pm) together with the elasticity
made the PEU membrane very pliable, a property which is needed to enclose
the wound surf ace very near (conf ormabi 1 i ty ( 16 1 1. The PEU covering is
elastic both in the dry and wet state, in contrast to some hydrogel-based
wound coverings, like Omiderm for instance, which are only elastic in the
wet state and contract the wound when losing water (vapour) during the
healing process (17).
The force-strain behaviour of a PEU membrane determined directly after
preparation differed from an aged PEU membrane as can be seen in fig. 6.
The P N membrane was prepared by casting of a LiCl containing refluxed
polymer solution. Heat treatment of a polyurethane solution will result in
complete dissolving of the physically cross-linked PU chains by breaking
up the existing supermolecular structure in solution. The LiCl keeps the
chains from aggregation upon cooling down to room temperature by
complexation to the urethane bonds, thus enhancing the solubility of the
PU and leading to the formation of a new super molecular structure (13).
A membrane cast from this PEU solution showed relatively high elasticity
(low modulus). However, after 1 to 3 months the elasticity had decreased
slightly. The at room temperature aged PEU membranes also exhibited
decreased elongation at break.
Preliminary studies on wound healing of partial-thickness wounds in rats,
using the PEU wound covering were encouraging (11 1. Extensive studies in
guinea pigs (121, in which wound healing of 122 partial-thickness wounds
under PEU coverings was compared with wounds covered with OpSite, an
Fieure 5 . Force - s t r - a in bchaviour of. ( a ) PEll wuur~d covcrine 3 mcnths
after preparation, IkE PEL1 woirrid covrring r l i r r r t l y aft-r prppmration
and CpSite
-
F i ~ i i r e 7 . F r e s h split-thickn~ss d r ~ n c s s i t e t~:>i!- dres srd wi t.h t.he PW
membrane Iriehtl and hali with a pararrin gauze ( l e f t ) . N u t e t h a t the
PEIJ membrane hecunles I r k ~ ~ s ~ i s [ t . r ~ l wt!e11 v~ic-ker i t.rr t he woll lrt lberl .
occlusive film dressing, and uncovered controls, showed that
reepithelialization and keratinization were enhanced in wounds covered
with the high water vapour permeable, porous PEU membrane compared with
wounds covered with OpSite or uncovered air exposed controls. The
percentage of reepithelialization on day 2 after operation was 85 % in
wounds covered with the PEU membrane, whereas it was 66 % and 35 %,
respectively in wounds covered with OpSite or exposed to air. In PEU
covered wounds 100 % reepithelialization was attained by day 3, one day
earlier than in the other wounds. Under the PEU covering the wound exudate
had turned into a jellylike clot layer by day 1 as a result of the high
water vapour permeablity of the porous covering. The high water vapour
permeability of the PEU covering prevented fluid retention (as was
observed in the case of OpSite covering) as well as complete wound
desiccation (uncovered controls). The jellified clot layer underneath the
PEU covering apparently provided an ideal matrix for epidermal wound
healing.
In order to evaluate the clinical efficacy of the PEU wound covering, it
was compared with the conventional treatment of tulle gras dressing (see
fig. 7) plus absorptive gauzes and crepe bandage on split-thickness skin
graft donor sites of 20 burn wound patients (15). The, initially fluid,
wound exudate under PEU coverings concentrated into a jellylike clot layer
after the extra gauzes had been cut away, which is needed to allow free
ventilation. After 5 days already the PEU covering could be peeled off the
wound without pain or epithelial damage. Clinically and histologically no
significant difference was observed in the rate of healing between the PEU
and tulle gras covered wounds, which may be explained by the fact that
tulle gras packed in a thick layer of gauzes and bandage prevented wound
desiccation, as a result of a sultry effect caused by the gauzes. Both
treatments enhanced reepithelialization at a similar rate. Further it was
observed that the use of the PEU wound covering reduces pain completely
compared with the rather painful treatment with tulle gras. One point to
note, finally, is that the PEU wound covering is not hemostatic in itself,
which, however did not turn out to be a problem in the clinical situation.
Acknowledgements
The authors would like to thank Dr. J.W. Leenslag for his contribution to
this work.
References
1. I.V. Yannas, J .F . Burke, J. Biomed. Mater. Res., 14, 65 (1980)
2. S. Gogolewski, A. J. Pennings, Makromol. Chem., Rapid Commun., 4, 675
(1983)
3. P. Bruin, G. J. Veenstra, A. J. Nijenhuis. A. J. Pennings, Makromol.
Chem. . Rapid Commun. , 9 . 589 (1988) 4. M. J. Tavis, J. Thronton, R. Danet, R . H . Barlett, Surg. Clin. North
Am., 58, 1233 (1978)
5. K. J. Qu~M, J.M. Courtney, J . H . Evans, J. D.S. Gaylor, W.H. Read,
Biomaterials, 6, 369 (19851
6. M.F. Jonkman. I. Molenaar, P. Nieuwenhuis, P. Bruin, A. J. Pennings,
Biomaterials, 9, 263 (1988)
7. M. E. Erasmus, M. F. Jonkman, Burns, 15, 371 (1989 )
8. C. P. Artz, D. R. Yarbrough, in: "Textbook of Surgery ed. ll",
W.B. Saunders Company, Philadelphia, 295 (1977)
9. M. Szycher, V.L. Poirier, D. J. Dempsey, J. Elastomers and Plastics,
15, 81 (1983)
10. S. R. May, in: "Burn Wound Coverings" (Ed. D. L. Wise), CRC Press, Boca
Raton, Florida, 53 (1984)
11. M. F. Jonkman, H. J. Mei jer, J. W. Leenslag, A. J. Pennings, P.
Nieuwenhuis, I. Molenaar, in: "Biomaterials and Clinical Appl.",
Elsevier Science Publishers, Amsterdam, 361 (1987)
12. M.F. Jonkman, P. Bruin, E.A. Hoeksma, P. Nieuwenhuis, H. J. Klasen,
A. J. Pennings, I. Molenaar, Surgery, 104, 537 (1988)
13. A.G. Zhigotskii, Z.N. Pazenko, T. I. Zhila, A.A. Panasevich, A.G.
Yakovenko, International Polymer Science and Technology, 3, 28 (19761
14. D. Behar, M. Jaszynshi, N. Ben Hur, J. Golan, A. Eldad, Y. Tuchman, N.
Stevenberg, B. Rudensky, J. Biomed. Mater.Res., 20, 731 (1986)
15. M.F. Jonkman, J.M. F. H. Coenen, P. Bruin, A. J. Pennings, H. J. Klasen,
Burns, 15, 211 (1989)
16. D. Queen. J.D.S. Gaylor. J.H. Evans, J.M. Courtney, W.H. Reid,
Biomaterials, 8, 372 (1987)
17. C. Cristofoli, M. Lorenzini, S. Furlan, Burns, 12, 587 (1986)
18. V. Falanga, Arch. Dermatol., 124, 872 (1988)
Chapter 5
A two-ply artificial blood vessel of polyurethane and poly(L-lactidel
S-Y
A biodegradable microporous small-caliber vascular prosthesis has been
developed that consists of two layers. The inner layer has been made
highly antithrombogenic by cross-linking of a mixture of linoleic acid and
a cycloaliphatic polyetherurethane with dicumylperoxide. Microporosity was
introduced by adding sodiumfluoride crystals of about 5 pm in diameter
prior to cross-linking and leaching them out afterwards. The outer ply has
been constructed by precipitating a (95/5) physical mixture of
polyesterurethane and poly(L-lactide) from solution in the presence of
sugar crystals with dimensions in the range 30-90 pm which were removed by
exposing the prosthesis to water.
The two-ply prostheses were tested in vivo by replacing 1 cm of the
abdominal aorta of rats. All the prostheses remained patent at least up to
1 year and did not exhibit any aneurysmal formation. The inner layer was
covered with endothelial cells and several layers of smooth muscle cells.
Introduction
Since the majority of deaths in the western countries are caused by
malfunctioning of diseased arteries, there has been a virtually unlimited
clinical need for arterial prostheses. This has brought about a
considerable research effort at constructing polymeric conduits for blood
and testing their biological performance. Especially small-caliber
artificial blood vessels with inner diameters of less than 6 mm seem to be
difficult to design and often exhibit poor patency as a result of blood
clotting.
Our approach to solving this problem has been to fabricate compliant,
microporous tubes of biodegradable polymer mixtures of polyurethanes and
poly(L-lactide), which act as temporary scaffolds for the ingrowth and
overgrowth of cells so that a neo-artery could develop 11.21. It was
indeed found that our vascular prosthesis when implanted in rats to
replace 1-cm length of the abdominal aorta induced the formation of a new
arterial wall consisting of an inner endothelial lining (neo-intima),
several layers of smooth muscle cells connected by elastin and collagen
(neo-media) and an outerlayer of fibrohistiocytic tissue constituting a
neo-adventitia [31. At one year after implantation some neo-arteries were
found to have newly formed smooth muscle cells that were
circumferentially arranged and elastic fibers formed concentric layers as
in the natural aorta. The development of the new artery wall is to be
attributed to a very well-matched rate of desintegration of the vascular
prosthesis and the rate of tissue ingrowth and the mechanical stimulation
on cell growth and orientation by the arterial pulsation of the blood. The
rate of fragmentation of this compliant prosthesis has been enhanced by
the rapid precipitation of the physical mixture of the polyesterurethane
and the high molecular weight poly(L-lactide) as used in the applied
dip-coating technique [4,181. The polymers solidify far from equilibrium
conditions thereby generating a substantial amount of residual stress in
the porous material. These residual stresses, as well as those arising
from the blood pulses and the large internal surface make the structural
breakdown by the hydrolytic and enzymatic environment quite unpredictable
[5-71. This has also led to the formation of aneurysms, a catastrophic
dilation of the artery wall, in several cases.
If one wishes to ever apply these artificial blood vessels to patients for
the replacement of damaged veins and arteries due to atherosclerosis it
obviously is a basic requirement that the rate of fragmentation of the
implants does not exceed the rate of tissue formation. But the major
problem arises if the prosthesis measures more than one cm in length
because the smooth muscle cells and endothelial lining have to grow mainly
from the anastornotic sites and require much more time to cover the entire
luminal surface of longer prostheses 181. This difficulty might be
circumvented, for instance by seeding of endothelial cells [91 and smooth
muscle cells all along the inner wall of the prosthesis as investigated by
Wildevuur et al. [lo]. But also improving the implant construction may be
sufficient as demonstrated by Gogolewski et al. [ I l l . They implanted
prostheses that measured 6 cm in length in healthy pigs and observed the
development of a neoartery which also grew in length reaching 12 cm after
one year. The key feature is that blood should not coagulate when it comes
in contact with the polymer surface in vivo which is initiated by the
adsorption of platelets. This may be achieved by special polyurethanes
which preferentially adsorb albumin and therefore almost no platelets, as
developed by Lyman et al. 1121. Ikada [I31 improved the
antithrombogenicity by grafting polyacrylamide on polymer surfaces. Bots,
van der Does and Bantjes 1141 made use of polyethers as biomaterial for
vascular prostheses and Klopper 1151 made them of collagen cross-linked
with glutaraldehyde. Bamford I161 has been able to prevent platelet
aggregation by attaching a synthetic analogue of prostacyclin to a variety
of polymer surfaces. Grafting of heparin [17,39,401 onto the luminal side
of our prosthesis decreased the platelet deposition but did not improve
the patency rate [171.
In the present study an attempt has been made to increase the
antithrombogenicity and lifetime of the vascular prosthesis by
cross-linking a thin inner layer of microporous polyurethane in the
presence of linoleic acid with dicumylperoxide. The negative charges [211
of the carboxyl groups [221 on the inner wall repel the platelets which
are also negatively charged. In addition, carboxyl groups appear to favour
adherence and spreading of cells 119-201 which therefore may promote the
attachment of layers of smooth muscle cells and endothelial cells. This
thin antithrombic inner layer was covered with a compliant microporous
material composed of a physical mixture of polyurethane and high molecular
weight poly(L-lactide) that was used successfully in previous animal
experiments.
This chapter describes the preparation of the two-ply artificial blood
vessel and its biological behaviour in rats. The prostheses that were
implanted were all patent for at least one year and did not exhibit any
aneurysm formation.
Experimental
Materials
The poly(L-lactide) (Mv=9.6. lo5) used throughout this study was
synthesized by ring-opening polymerization of L-lactide that was catalyzed
by Sn(I1)-2-ethylhexanoate. A medical grade segmented polyesterurethane
(Estane 5701 F1, Goodrich USA) was used after being purified by
precipitation from solution in N,N-dimethylformamide (DMF). The solution
was poured into ice-water. The polymer was washed with ethanol (96%) and
ether and dried in vacuo at 40'~. A second medical grade polyetherurethane
(Tecoflex EG 80A, Thermedics Inc., USA) was used as supplied.
The solvents, DMF, tetrahydrofuran (THF) and 1,4-dioxane (Janssen Chimica,
Belgium) were purified prior to use according to standard procedures.
Chloroform, linoleic acid (9,12-octadecadienoic acid, cis, cis) (no.5353
Merck, Germany), sugar crystals (Suiker Unie, The Netherlands) and NaF
(Baker, UK) were used as supplied. Dicumylperoxide (Dicup R, Hercules,
USA) was purified by repeated recrystallization from methanol.
Prostheses preparation
The inner layer was made of a suspension that was prepared by mixing 0.05
g of linoleic acid and 2 g of small NaF crystals (size 5-15 pin) and 0.033
g of dicumylperoxide with 6.5 g of a solution of Tecoflex (3.4 wt %) in
chloroform. A glass mandrel (diameter 1.4 mm) was coated with the inner
layer by immersing it in the suspension and pulling it out slowly in order
to obtain a film of uniform thickness. The coated mandrel was put in an
oven at a temperature of 150 OC under a dry nitrogen atmosphere for 30
minutes.
For the preparation of the outer layer a 32.5 wt % solution was made of
Estane 570l/poly(L-lactide) (95/5(w/wl) in a mixture of 1.4-dioxane/DMF
(1/3(v/v)). At first the high molecular weight poly(L-lactide) (PLLA) was
dissolved in the refluxing solvent mixture. After cooling the solution
down to room temperature, the purified polyurethane was added and the
temperature was raised to 150 OC. Subsequently, the solution was quickly
cooled down to room temperature and 2 g of sugar crystals (size 30-90 wm)
were added to 3 g of this solution to obtain a porous outer layer.
This viscous suspension was put on the mandrel already covered with the
inner layer. Next the mandrel was rolled over a glass plate that was
covered with a thin layer of small NaF crystals in order to prevent the
formation of a skin and to obtain a pore structure on the outside of the
prosthesis. Two small rings on the mandrel, at a distance of 5 cm, took
care of achieving a wall thickness of 0.7 mm. The polymer mixture was
precipitated in a coagulation bath containing a mixture of ethanol and
water (8/l(v/v) at 20'~. After 30 minutes the prosthesis was stripped off
the mandrel and submerged in water at room temperature in order to leach
out the sugar and NaF crystals. The prostheses were left for several days
in the distilled water which was renewed several times before
sterilization and implantation.
Characterization of the prostheses
A scanning electron microscope (ISI-DS-130) was used for examining the
microstructure of the prostheses prior to, as well as after implantation.
Stress-strain measurements were carried out using an Instron 4301 tensile
tester, equipped with a 1 N load cell at a cross-head speed of 12 mm/min.
Cross-sections of the prostheses were determined using a Profile Projector
(Nikon, model 6C, Japan) with an accuracy of 0.005 mm.
Wet samples were subjected to mechanical deformation in their longitudinal
direction and the distance between the clamps was 15 mm. Extraction of the
sol fraction was performed in an excess of tetrahydrofuran at room
temperature for 72 hours. The porosity of the prostheses was determined by
density measurements. The radial strength of the prostheses was tested by
clamping one side and filling it with water to which a pressure of 2 m
water was applied. In animal studies only those prostheses were used that
could withstand this pressure and did not exhibit leakage as a result of
structural defects.
Implantation
22 prostheses were gas-sterilized with ethylene oxide according to
standard procedures. The in vivo experiments were performed by resecting 1
cm of the abdominal aorta of rats and replacing it by the two-ply
artificial blood vessels.
Results and discussion
General design considerations
This section describes our approach to the evaluation of the basic design
parameters for an artificial blood vessel. The development of such a
prosthesis is primarily an interfacial problem [23] because the luminal
side of the conduit is in contact with streaming blood and the outside
interferes with surrounding tissue which initiates vascularization and
penetration of cells as well as biofragmentation. Hench and Ethridge [231
developed the theory that the properties of an implant should match as
closely as possible those of the organ or tissue that has to be replaced
in the body of human beings or animals. This "natural approach" implies in
the case of an artificial blood vessel that their mechanical properties
should correspond to those of the arteries to be substituted and that the
inside should not cause any thrombus formation. In the natural artery this
is ideally effected by the endothelial lining at the bloodstream
interface. If the vascular prosthesis is not completely covered by
endothelial cells, the risk remains of thrombus formation with ultimate
occlusion. Endothelial cells grow from the anastomotic sites and are
deposited from the flowing blood if sufficient sites for anchoring are
available 181. Particularly in the case of human beings whose diseased
blood vessels measure more than 1 cm in length, quite some time will be
required before the entire luminal side is covered by smooth muscle cells
and endothelium and additional precaution has to be taken concerning
antithrombogenicity and cell adhesion. No material is completely
thromboresistant and also the mechanism of thrombosis prevention is
unknown. We have chosen for introducing COOH groups by cross-linking
polyurethanes with dicumylperoxide in the presence of linoleic acid
because negative charges in general increase the thromboresistance due to
the fact that blood platelets are also negatively charged. Another
advantage of the COOH groups on the luminal side of the prosthesis is that
they strongly promote the adherence and spreading of cells. The
antithrombogenicity may be further enhanced by adsorption of albumin
133,341, the protein that occurs so abundantly in blood. This protein is
known to bind strongly to fatty acids by interaction of the COOH groups
with the free NH groups of the lysine residues. Evidence for this 2
interaction has been obtained from studies of NMR relaxation times [291,
although it is believed that hydrophobic interactions at certain binding
sites may be dominating factors 1301.
Another advantage of the carboxyl groups on the luminal side is that they
promote strongly the adherence of cells which often only grow and
proliferate when attached to a suitable surface [26,271. This is the
reason why this cross-linked inner layer has also been made microporous,
thereby increasing the adherence and spreading of cells as well as
permitting transport of nutrients.
A second reason for putting chemical cross-links into polyurethanes is to
eliminate the serious limitation of the stress softening or hysteresis
which occurs especially in cyclic loading. It is to be attributed to the
deformation and rupture of the hard segment domains. This cyclic
creep-failure, invoked by the arterial pulsation of the blood, may lead to
the formation of aneurysms, the catastrophic dilation of the prosthesis.
Chemical cross-links also impede the structurization of the hard segment
domains on which the platelets seem to adsorb preferentially L24.251,
thereby inducing blood clotting. Furthermore, the cross-links improve the
fatigue life of the prosthesis and diminish the possibility of the
occurrence of environmental stress cracking.
A third reason for selecting the chemical network structure of the inner
layer is that the biofragmentation might be retarded. The implantation of
a foreign device as well as cellular damage will elicit an inflammatory
response and the release of hydrolases which are enzymes from lysosomes
that cause proteins, lipids and other biomolecules to fragment [281. These
enzymes may be actively involved in the phagocytosis of the artificial
blood vessel. But acids are also highly efficient catalysts in cleaving
the urethane linkages t5-7,35, 411.
In summary, the microporous inner layer of our proshesis is supposed to be
highly antithrombogenic and has an improved creep resistance that is aimed
atpreventing aneurysmal dilation, thereby resembling the intima and media
of a natural artery. A second outer layer is made of
polyurethane/poly[L-lactide) mixtures with rather large pores for
vascularization and cell ingrowth so that fibrohistiocytic tissue may
develop, as in the adventitia of natural blood vessels. This two-ply
prosthesis will have better anisotropic elastic properties matching closer
those of arteries in vivo.
Construction and properties of the inner layer
The inner layer of the prosthesis was prepared by immersing a smooth glass
mancirel i n a 7 . 4 %. s o l u l i on r l f 1 he cycl.o;il i phdt i c 1 yu[ c l l l ; l r~c , 1-<!cot i cx ,
I n chlnrofor-m c c ~ n : a l n i n ~ 30 % s l , , l l l NaF crystals I*l; ivlny, ;lirnrnsior.s in tbc
Irsngc 5-15 p m l , dic~irnylp(+t-ozlidf, 2nd 1 i ~ l o l e i c ;#:id.
T l w l a y e r thickr!css and u n i f o r ; r ~ i t v was a c h l e v c d by pulline th*= mandrel
slow1 y out n f t h ~ pfl l ymeric sulpr 'nsior l arir l t l l r - r1ilorr)i { r r fn W;~S 1 l r v w u ~ ~ l t c )
e v d [ ~ u i d t ~ Ily p u t t i l l g the ~ u v ~ r t - r I rnantlrcr i n an o ~ c f i a t 1 5 0 " ~ f o r 313
r n i n u t ~ s . I n t h i s w a y th? polyrrrr ct.3.n:: w c r : c~-r~~-.s-l i1tkc.c i n ;I staye o f
r'rlndomnrss, i . c=. . wi thuut the formation o t I.hr d r m ~ i r ~ c ; n f the hard
s~mgrnen ts .
A f t c r the synt .hcs i s of thc chcrniual n ~ t u c l r k I lhe s11;lll FIaF c r y s t a : ~ wprp
l ~ a r h r r l n11t w i t h H O and a micrclmrous s t r - u c t r ~ r ? r*mll i !br.d a s is shown by
the s c a n n i r . ~ electr-an micru,,l'ny~li of' 1 i g . 1 [-?l ,JI l 1111s microcraph a l s n
reveals C h d L L t i ~ ? surracc i s n o t rnmplrl = ! y : < m o o t h t n ~ t ha: a certain
r o u ~ h n ~ s s w t ? i r l r h a s a i a v o u r a b l n pffpr t rmii t h~ rt'l l ; I ~ I I I I T i o r l and possibly
also or1 t t ~ c flow beh~viuur of t h ~ uIood
Fie: I SFM mi crop rapt^ o f t l ~ r ml I:!-aporous skrur tun; of the inner layer of
t n ~ two-ply v n s c u l a r prozt hrs:,; p r r p a r r d 1 1 om ,I 3 . 4 X solut ~ n n of
cyclnal ipha t ic. pol y~lrr t l i : r lh? Tet <milex i n c : ~ 1t11.ofc)1 411 an11 30% ?ma 1 1 NaF
C I - y s t a l s whj ch were Ieiir1i.zd o l ~ : w l t ! ~ w n t r r
I
Fig. 2. Stress-strain dependance 6 0 0 -
of the cycloaliphatic
polyurethane Tecoflex (A) and of
the microporou~ polymer (porosity
35%) (B). Sample (C1 has a
porosity of 55%.
0 100 2 0 0 300 LOO
straln E l o l o I
0 0 200 LOO 500 800 1000
Fig. 3. Stress-strain dependance of the microporous ( 5 5 % ) aliphatic
polyurethane Tecoflex (A) illustrating the effect of cross-linking with
dicumylperoxide (DCP) at 150 OC for 30 minutes (B) cross-linked with 7.5
W/W % (DCP); (C) cross-linked with 15.0 w/w% DCP; (D) cross-linked with
20.0 w/w % DCP.
Fig. 4. The gel content of
the microporous (55%)
polyurethane Tecoflex
cross-linked with 15 w/w %
dicumylperoxide as a
function of the linoleic
acid concentration.
These small protuberances may just suppress the formation of eddying
vortices that swirl up from the surface in areas of turbulent flow.
The overall porosity of the inner layer is of the order of 50% and
diameter of the interconnected pores are indeed in the range 5-15 pm. As a
result of the pores there is a considerable decrease in modulus of the
polyurethane as illustrated by the stress-strain curves for specimen with
different degrees of porosity in fig. 2. The elastic modulus scales +2
approximately as 4 where 4 is the volume fraction of polymer [36-381. P
Since the polyurethanes and especially the ones based on aliphatic
diisocyanates exhibit a pronounced stress softening, which is bound to
lead to the development of aneurysms, we have attempted to diminish this
disadvantageous property by cross-linking with dicumylperoxide for 30
minutes at 150 OC. This reaction period is six-times the half-life of the
peroxide at this temperature which could not be increased because of the
thermal degradation of the polyurethane. The effect ofcross-linking of the
porous polyurethane is a pronounced reduction in Young's modulus as well
as in tensile strength and elongation at break, as manifested by the
stress-strain curves of fig. 3. Chemical cross-linking of randomized
polyurethane molecules is likely to lower the number of hydrogen bonds
that can be formed and will impede structural organization as hard
segments domains which will therefore lead to a lowering of the Young's
modulus. The decrease in tensile strength may be due to network
inhomogeneities but also degradation may have contributed to this marked
fall. In order to verify whether chain scissioning did occur, sol-gel
analysis were performed by swelling the networks in tetrahydrofuran. For
the polyurethane samples that were cross-linked with 15% by weight of
dicumylperoxide a gel content of 95% was attained, suggesting that, in
view of the rather broad molecular weight distribution, molecular
degradation does not seem to be a predominant factor. The cross-linking of
the polymer chains has been found to take place by recombination of free
radicals on the nitrogen and the carbon atom adjacent to the chain oxygen
l43.451. Both are losing a hydrogen atom by the dicumylperoxide fragments.
The cross-linking efficiency could be enhanced by adding linoleic acid
which contains two double bonds and mainly served to introduce carboxyl
groups for improving the antithrombogenicity. Fig. 4 shows that adding the
linoleic acid resulted in an increase of the gel content and at 20 %
linoleic acid a gel content of 100% was achieved, indicating that chain
scissioning did not occur. Supplying linoleic acid also had a favourable 2
effect on the stress at a strain of 300 % and increased from 50 N/cm to 2 70 N/cm as shown in fig. 5. In order to acquire some indication of the
creep and fatigue resistance the uncross-linked porous polyurethane sample
and one that was cross-linked under optimum conditions with 15%
dicumylperoxide and 20% linoleic acid, the testing-samples were stretched
20-times to a strain of 50% and allowed to recover for 2 hours.
Subsequently the stress-strain measurements on both samples showed a
permanent deformation of 16% for the porous polyurethane (fig. 61, whereas
the network exhibited still some hysteresis but no creep at all (fig. 7).
The most pertinent result of the cross-linking of the polyurethane in the
presence of linoleic acid is the marked increase in tensile strength which
is equal to that of the uncross-linked porous polyurethane as illustrated
by the stress-strain curves of fig. 8. Another salient feature is that the
stress-strain curve of the cross-linked porous polyurethane is concave and
therefore much closer approaches that of the aorta than the pure
polyurethane.
Fig.5. The stress at a
strdin of 300% for the
microporous (55%) polyurethane
Tecoflex cross-linked
with 15 w/w % dicumylperoxide as
a function of the linoleic acid
concentration.
. z -
J
Fig. 6. Stress-strain curves
for the uncross-linked
microporous (55%)
cycloaliphatic polyurethane
Tecoflex illustrating a permanent
deformation of 16% after being
loaded 20-times up to a strain of
50%.
0 2 0 40 6 0
strain E ( ' 1 0 )
I 12 -
I 1 Fig. 7 . Stress-strain curves for
the cross-linked microporous (55%)
1 cycloaliphatic polyurethane
Tecoflex after being loaded
20-times up to a strain of 50%.
There is no permanent deformation
and only a little hysteresis.
0 20 40 6 0
s t r a ~ n E t o l o )
Fig. 8. Stress-strain behaviour of
the microporous (55%) polyurethane
Tecoflex cross-linked with 15% DCP
and 20% linoleic acid ( A ) compared
with that of the uncross-linked
microporous polymer (B).
0 200 4 00
strain E ( ' l o )
Fig. 9. S M rnic:-ograpll ol t , lc . twcl-ply a r t i T i r i - + I L > I o , d w ? s ~ c I cornposed of a
cross-linked m i u r u p u r - o u s pulyui -et l~ane i n n e r layer and a- ouLer layer- wi t l ~
larger pores made af a physical mIxtur c o f a polyesturethane and
polytl-lactide) I95/51.
F i g . 10. SEM I : ~ ~ L I U ~ I clpli uf d c ush-tre~ - i o ~ i UI t 1 l r dl t i f i l i a 1 1,lood vessel
showing small pores 111 t h e polyurethane/poly(L-lactide) mixture I n
addi t i n n t o l a r g e p o r t s which o r i g i n a t e d irvrn leaching o u t the suRar
crystals.
Flg. 11. StM micrograph of t ~ c a r t l f l c la l blood vr:ssel I 1 l u s t r a t ing t h e
firm connection between the IOU-pm-thick i n n e r l a v c I- and t h r o u t e r layer
uith the large pcres.
50 - F i g . 12. " t r e s s - 5 t r a i n behnv i v u r uT
t l ~ e u u l e r l a y r - : . The cor,cavr cu rve A - LO - repses-nts t h c hehavicur 01' t h r
physical mixture of pn 1 y u r ~ t h a n r and
pnly(L-lactide) (q5/5) dissolved i n a
dioxane/DMF mixturc at 150 "C and
guencf i~d t o rocn I c m p e r i i t l l ~ ~ . C u r v y &
describes the strr%s+straln f o r t h ~
sanr polvmer m l x l u r e dissulvcd A ? 70'~
0 20 i D 6 0
5tra1n E ( ' 1 0 )
Construction of the outer layer
The optimal dimensions of the pores in the outer layer should be in the
range of approximately 25 to 150 pm. Previous studies have shown that a
homogeneous pore structure of this size facilitates cell ingrowth and
vascularization. Instead of applying a dip-coating technique which had to
be repeated many times in order to acquire a layer of sufficient thickness
a salt leaching method was developed. At first a rather viscous solution
was prepared of polyurethane and poly(L-lactide) in a mixture of
1,4-dioxane and dimethylformamide (1/3 ratio) to which sugar crystals
(30-90 pm) were added. The mandrel covered with the under layer was rolled
in this suspension and subsequently the polymer mixture was precipitated
in a coagulation bath of ethanol /water (8/1). The sugar crystals were
dissolved in water leaving the required pore structure. Fig. 9 presents a
scanning electron micrograph (SEM) of the two-ply artificial blood vessel
and fig. 10 gives a scanning electron micrograph at higher magnification,
which reveals that in addition to the large pores originated from the
leaching of the sugar crystals the polymeric material also contains very
small pores which are of importance for the fragmentation and gradual
degradation of the prosthesis. The connectedness between the inner layer
and the outer one is clearly demonstrated by the scanning electron
micrograph of fig. 11. The significance of the 5% of high molecular weight
poly(L-lactide) in the mixture with the polyurethane is essentially that
both polymers are highly entangled in these concentrated solution which
hinders the formation of hard segment domains and consequently increases
the rate at which the material can fragmentate. This effect is also
manifested in the stress-strain curve which turns out to be somewhat
concave for the polymer mixture dissolved in dioxane/DMF effectively at
150 OC and subsequently quenched at room temperature, whereas the curve
for the polymers dissolved at room temperature is convex (fig. 12).
Therefore, it is very essential that the polymers are dissolved at high
temperature in order to make the urethane linkages more accessible for
hydrolytic and enzymatic fragmentation which basically determines the rate
of cell ingrowth and development of the fibrohistiocytic tissue.
Biological ~valuat icra
TIE a r t i f i c::ll ; I T - l PI-ier; wcrc i r r . ~ l 2 n t u d i n 7 2 I-at ,s \ ) y rrhquvt-t in^ 1 cm of t h e
abdominal a c r t x artd - i : i l t y t t ~ t i v hy weans of
m i c r o s u r g l c n l t r>r -hnlques Up t o s>np yr:br ~ 1 - thc ?.I t i f i c i a l blood vvssels
wcrr patcnt [ I , ' ] .
R p ~ a f - e n t l y , t h e i nt:nrpi,raL i o n of t h e ! !n:>l P i I- a c i d 1nt.c thr i r m ~ 1 l ~ l l ayer
p r o v i d ~ d good a n t u +hi o n i ~ n g e n i c i t y . Two J ) I - { I ? ~ h ~ s e s w h i c h were ~ n a r ~ u f ' a c t u r ~ d
w1l.t) a lhinner o u t e r !dyrt 'xV. i lr i tc~d 6 slicht dilation as biell as nncu!-ysm
for-nl:~t ion af tcr 4 ucck:;. Ttic s:rc:%-7t.r ,!. 11 cu t ves oi' I t , rsr? t w r ~ p r u ~ l llebes
pr i c l r to implantntion a r e ccm~arcd with '.h,qt. cnf Lhu dhriominal aorta nt~ri of
a g~ a f t , wi t h Ills? r o r r ? c t lzyr-l- l . l i i c k t ~ ~ c ~ in I - i ~ . 1 3 . Arlci:rysm format 1011
can be avoided if the s t ress of ttu- als2arnirt;il L I I J I ~ L I illld Ih:tt the
proslhesis a r e ~ r q 1 1 ; . + 1 at. a s t , r ; . l~~ u i ?,(:'%.. TI , < h n ~ j l d tbc re:narkt?rl t 1 1 ; l i t h e s e
arleut ysrr.;~ 1 r j i l a l l o n s wcrc wlly o t r e r v c d v i ::lual l y and :!LC l,t!mn i n i r l ~
prosthesr:s wcrc found t o 1 ~ v s Lhc original rilan~el.nr- ii!; is i . l l ( ~ s t ~ - a t e d by
the ~ h n l u r n i c ~ i l p r ~ p h of t h ~ two-p ly val;c;llar ~ r n s t . h ~ s i s u n e year a r t e r
implantation ( f i g . 1 4 ) .
F i g . 15. SEM 1nicrograpl.1 of- tl;t-: l~rntlrial su~-f:ict- somr:wi~vr-c j n thr! m i ~ l r j l e of
the two-ply vascular prosthesis a t I n c yrar a l t e r implantatlon i n t n I h r
abdominal x o r t a Q T a ra i . l l . 1 1 ~ the ~ m n o t h r l > d o t t 1 ~ 1 i a l J i n ing .
M~vnification 1500 X I .
The luminal surface oi the p r o s t l ~ c s ~ s t u rned o u t t o be cornplrtcly covered
with endothelial cells, as demonstrated by t h e scanning e lect rnn
micrograph prescntrd i n Tig. 15.
I ~ i s t o l o g i c a l s t u d l e s of longitudinal and transverse cross-sections of the
prosthesis 1 year a f t e r i r n p l a n i a t i v ~ ~ rrvcalcd t h a t i r ~ a ~ l d i l i r j n t v the
complete endothelial l i n i n g making u~ a neo-int ima also smooth muscle
cells and elastine were formed. r a t h e r firmly attached t o the inncr l a y e r
nf t h e prosthesis [ f i g . 16) .
Thc light micrograph of a transverse cross-section presen ted i n rig. 17
~110~5, i n addition t o the neo-intirna 2nd neo-media, a l s o the gen~ratlon o f
flbrohystiocytlc t issue in the oute r l a y e r . Furthermore. it discloses the
presence of cap i l l a r ips and f r a ~ n e n C a l i o n of the polymeric prosthesis
material. T h i s ]pads to the conc l i l s i on t h a t one year after implantation
the deve loprn~nt of a neo-adventitia hes a l s o s tar ted .
F i g . 16. LighC micrograph of a l o n g i t u d i n a l scction D T t h r two-ply vascular
prosthesis at one year a i t r r implantatiun i n t u the ai~rlurnilbal a u r t d o f a
r a t . The inner d e n s e Xaycr, P i , artd bht. n u t e r loose layer, Po, have been
made v i s i b l e by staining u i t h Sudan black B 1441. Arrow marks weak spot
in the inncr layer The prosthesis is lined w i t h a neo-intima 111 and
neo-media (MI magnification 100 x.
F i c . 1 7 . L i c b t m l c r o ~ r d p i ui .j I r a r ~ 5 v p r ~ ; r r ~ n = ; - ~ : ~ c ' i o n n f the two-ply
vascular p r o s t h x i ~ a t OIIC j c l r a f t r r i~pl?ntltiun i n t o t h e ahdomlnal
a o r t a of a r a t . The p r o s t h e s ~ ? ha7 I hv i u r m clt a r1r.u-..r tcr y i n which t h r e e
s e p ~ r a t e layprs can h~ rrcngnizrd. At7 I t i t ~ r r I l l 1 layer, I .
[ l ieu- int i m a Subintima 1 l ay ; r? of r m o o l h nusclp cel 1s. M, In~o-mccl i a l
c 0 n t a i n i r . g r l n s t i n , F Fi Prnhyst iocytic t iss-1s l k u ~ h a s orrani r ~ d t h r
Aeslntegrating pros t,hesis, A , (neo-advent 1 t l a). Nrl t r - t i l ~ ~ T C T P T I L C ~ l f
cnpillarirs. C . anrl ~ ~ r n ~ t l ~ r t j r matcrial, I \ i i r l t l c neo-advcntltla
Hn~nificat ton 7(J r ) x
Conclusions
The two-ply microporous, biodegradable, and compliant artificial blood
vessel described in this chapter has proved to function adequately as a
scaffold for the generation of a neo-artery in rats. The cross-linking of
the microporous polyurethane inner layer with dicumylperoxide in the
presence of linoleic acid has substantially improved the
antithrombogenicity of the prosthesis. This may be due to the fact that
the fatty acid moieties extending into the lumen may tightly bind to
albumin which prevents the platelets from being adsorbed.
Carboxyl groups also promote the adherence of cells which resulted in a
firm connection between the smooth muscle cells of the neo-media and the
inner layer of the prosthesis.
Another salient feature is the pronounced increase in tensile strength of
the porous polyurethane when the cross-linking efficiency is enhanced by
adding linoleic acid. This strength-increasing factor has substantially
diminished the risk of ameurysm formation.
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Chapter 6
Autoclavable highly cross-linked polyurethane networks in ophtalmology
Summary
Highly cross-linked polyurethane networks have been prepared by the bulk
stepreaction of various low molecular weight polyols and (cyc1o)aliphatic
diisocyanates. All these polyurethane networks were optically transparent,
colourless, amorphous glassy thermosets. The properties of the glassy
polyurethane, obtained from the bulk reaction of a tetrafunctional
secondary aminoalcohol tetrakis(2-hydroxypropy1)ethylenediamine or Quadrol
(containing an internal tertiary amino group, that can catalyze the
urethane reaction) and hexamethylenediisocyanate (HDI) in stoichiometric
proportions, have been investigated in more detail. This glassy
polyurethane, with an ultimate glass transition temperature of 85 OC, and
a very low degree of swelling in chloroform (1,271, exhibited good
ultimate mechanical properties (tensile strength 80-85 MPa, elongation at
break ca. 15 %, modulus ca. 1,s GPal. Infra-red spectra of these
hydrophobic polyurethane networks (water uptake ca. 1 %) revealed the
absence of an isocyanate absorption, indicating that all isocyanates,
apparently, had reacted during the cross-linking reaction. Preliminary
experiments and suggestions to increase the hydrophilicity of the networks
have also been described.
In contrast to poly(methylmethacry1atel [PMMA), which has been used
successfully as an intraocular lens material the last 15 years, these
transparent cross-linked polyurethanes can be sterilized simply by
autoclaving. The possiblity of an autoclavable lens is especially
interesting for use in eye surgery in the developing world where the
majority of the blind people live. These highly cross-linked
Quadrol/HDI-based networks, after being autoclaved, were implanted in
rabbit eyes, either in the form of small circular disks or in the form of
a keratoprosthesis (artificial cornea). It was shown that the material was
well tolerated by the rabbit eyes. A serious opacification of the cornea,
a direct result of an adverse reaction to the implant, was never seen.
Even one year after implantation of a polyurethane keratoprosthesis the
eye was still "quiet" and these findings were comparable with the ones
obtained from implantations of keratoprostheses made of PMMA or stainless
steel/glass. These results show that the transparent highly cross-linked
polyurethane network seems suited for use in ophtalmic applications, like
autoclavable intraocular lenses or keratoprostheses.
Introduction
It is estimated that cataract, i.e. an opacification of the crystalline
lens of the eye and the main cause of blindness, is responsible for
approximately 20 million blind people world-wide; most of them live in
developing countries in Asia or Africa (1-4). It is now recognized that
cataracts are not only another sad consequence of ageing. Other possible
risk factors include malnutrition, sunlight exposure, smoking. This has
led to the hypothesis that oxidative damage plays a major role in
cataractogenesis (5,6). Surgical removal of the cataractous natural lens
is the only medical treatment available for cataract patients. Spectacles
or contact lenses used to be the conventional way of replacing the natural
lens, but turned out to be far from ideal. However, implantation of an
intraocular lens in the place of the removed cataractous lens is presently
the best way to correct aphakia and to visually rehabilitate the cataract
patient (7.8). Modern intraocular lens implantation, using artificial
lenses made of poly(methy1methacrylate) (PMMA), had its start in 1949
after the Second World War during which Ridley, an English eye-surgeon who
performed the first PMMA lens implantation, noticed that Perspex splinters
of canopies of airplanes caused no irritations in the eyes of pilots.
Since then many intraocular lens designs have been developed and implanted
in cataract patients (8,9,12). The number of PMMA lens implantations has
increased enormously during the last 15 years and nowadays lens
implantation is considered a routine operation (7,101.
It is generally accepted that the artificial lens should be optically
satisfactory, inert, non-toxic, biocompatible, lightweight (glass, for
example, has always been considered as a possible, autoclavable lens
material, but weight problems have limited its use (1211, structurally
sound, durable, ultraviolet light absorbing, resistant to laser treatment
(in case of secondary cataract formation), easily implanted and securely
fixated and of course be sterilized safely (8,111. PMMA meets almost all
of the requirements listed, but does have some disadvantages as a lens
implant material. Although non-toxic, PMMA is extremely damaging to the
corneal endothelium if in contact during implantation surgery (13,18,19).
Endothelial cells adhere to the hydrophobic PMMA and may be stripped off.
Surgical skill is needed to minimize contact adhesion. Due to its rigidity
PMMA may cause mechanical irritation of uveal tissue. Another serious
problem with PMMA lenses has been with regard to sterilization. PMMA can
not be sterilized simply by autoclaving due to its relatively low T (100
OC), and therefore has to be sterilized either with the toxic ethylene
oxide or by sodium hydroxide sterilization. The last method has been
prohibited in the United States by the FDA, but is still used in Europe
(7.9). Both sterilization methods are, unlike autoclaving, not without a
certain risk. According to eye-surgeon Worst there is a need for an
autoclavable intraocular lens, especially for cataract surgery in
developing countries (14). Research in this field has led to the
development of other lens materials that are autoclavable. On the one hand
amorphous, aromatic thermoplastics with a very high glasstransition
temperature, like polycarbonate, poly(ether)sulphone, polyimide, fulfil
this requirement (7,11,15,16). On the other hand, polymeric networks can
be used as autoclavable intraocular lens materials. Hydrophobic silicone
elastomers and hydrogels (polyHEMA), which have been used in other medical
applications, e.g. soft contact lenses (171, for a long time, are examples
of polymeric networks which have been considered and evaluated as
potential, clinical intraocular lens materials (11,18-23). Both materials
are elastomeric and can be folded, which means that artificial lenses made
of these materials can be inserted via a smaller incision into the eye
than in the case of glassy polymers. Hydrogels are hydrophilic materials
having a soft consistency and are known for their soft tissue
biocompatibility (24 ) . So, these materials are expected to be less
damaging to the eye, especially to the corneal endothelium.
Besides opacification of the eye-lens, the cornea may become
non-transparent, leading to so-called corneal blindness. Throughout the
world about ten million people, mainly living in the developing world,
suffer from corneal blindness, for example from trachoma (2). If a corneal
graft is not available, due to a lack of donor eyes, or not indicated,
implantation of a keratoprosthesis is the only possibility to restore
sight (25). A keratoprosthesis (an artificial corneal is usually made of
the same materials that are used for intraocular lenses. The requirements
for an intraocular lens material and a keratoprosthesis material are
virtually the same.
In this chapter we will describe the synthesis and properties of glassy,
highly cross-linked polyurethanes and their potential application in
ophtalrnology as an autoclavable ocular implant material. Polyurethanes,
which in general are relatively biocompatible and used in many medical
applications (261, have never been considered as materials that might be
used in ophtalmic applications. All this motivated us to synthesize a
series of new, densely cross-linked polyurethane networks by stepgrowth
polymerization of low molecular weight polyols and (cyc1o)aliphatic
diisocyanates.
Experimental
The polyols used in this study were:
tetrakis(2-hydroxypropyl)ethylenediamine (Quadroll,
triisopropanolamine (TIPA, mp. 48-52 O C ) ,
triethanolamine (TEA),
tetrakis(2-hydroxyethy1)ethylenediamine.
bis-N,N-(2-hydroxyethyl)isopropanolamine (BHEIPA),
tetrakis(2-hydroxyethyl)methylaminomethylmethane,
octakis(2-hydroxypropyl)pentaerythrityltetraamine ("octaol"),
trimethylolpropane (TMP, mp. 60-62 O C ) ,
pentaerythritol (mp. 260 OC),
glycerol,
2.2-bis(hydroxymethy1)-2,2', 2"-nitriloethanol (BIS-TRIS, mp. 104 OC).
The chemical structures of these polyols are shown in figure 1. All
polyols were liquids at room temperature, unless a melting point is
mentioned in brackets, and were purified, if possible, by distillation
under reduced pressure. Most polyols were commercially available. Two new
polyols were synthesized as described below.
Tetrakis(2-hydroxyethyl)methylaminornethylmethane
A stirred mixture of 1 eq. pentaerythrityltetrachloride and 9 eqs.
N-methyl-ethanolamine was refluxed for 8 days under a nitrogen atmosphere
at ca. 165'~. Excess N-methylethanolamine then was distilled off and to
the residue ethanol and 4 eqs. powdered potassium hydroxide were added.
After stirring, ethanol was removed and the residue was extracted with
chloroform. After removal of the solvent, fractional vacuum distillation
yielded the product, a viscous, colourless liquid, bp. 148-155 O~/0,007
mbar. Analysis calculated for C H N 0 C 56.04, H 10,99, N 15,38. 17 40 4 4'
Found: C 56,27, H 10,76, N 15,49.
Octakis(2-hydroxypropyl)pentaerythrityltetraamine was prepared in the same
way as the previous compound from pentaerythrityltetrachloride and
diisopropanolamine. The octafunctional polyol was isolated, in poor yield,
as a yellowish, very viscous liquid, which could be decolorized by using
activated carbon. Bp. 200 '~/0,005 mbar. Anal. calcd, for C H N 0 : C 29 64 4 8
58,39, H 10,74, N 9,39. Found: C 58,24, H 10,46, N 9,33.
The polyfunctional amines used were pentaerythri tyltetraamine (also known
as tetrakisaminomethylmethane). tetrakis(N-propylaminomethyl)methane,
tris(2-aminoethy1)amine. In figure 1 the structural formulas are depicted.
The syntheses of the two tetraamines are described below.
Pentaerythrityltetraamine was synthesized from pentaerythritol in 5 steps.
In the first step pentaerythritol was converted to the
tetrabenzenesulfonate according to a published method (27). Following a
patented method (281, the tetrabenzenesulfonate was reacted with
sodium-p-tosylamide in N-methylpyrrolidone solution at 200 OC for 20 hours
to yield the tetratosyl-amide, a compound also described earlier by
Litherland and Mann, who started from pentaerythrityltetrabromide (29,301.
In the third step the tetratosylamide was hydrolyzed with 80% sulfuric
acid resulting in the formation of pentaerythrityltetraamine disulphate.
The next step was the continuous extraction of the disulphate with sodium
hydroxide in benzene to give the tetrahydrate of
pentaerythrityltetraamine. Both last steps were described in the
literature (29,311. Finally, the tetrahydrate was converted to the pure,
hygroscopic tetraamine by azeotropic distillation with benzene. The
overall yield was ca. 60%. Anal. calcd. for C H N : C 45,40, H 12.19, N 5 16 4
O C N + C H ~ ~ ~ N C O
Hexamethylenedllsocyanate (HDI)
OCN CH,
OCN NNC0 trans 1.4qclohexanedtisayanate (I-CHDI)
OCN-CH-CyO (Ial \"*F*.
NCO
Lysine dilrocy~me (LDI)
F", FH, HO-CH-CH
\2 ,CHZ-CH-OH
N-CH2-CH2-N
no-CH-cn2 \cH2-cn-oH I I CHI CH,
FH, CH,-CH-OH
HO-CH-CHz-N
\cH,-cH-oH
CHI
CH.-OH I
HC-OH I cn2-OH
Glycerol
42.37. Found: C 45,37, H 12,19, N 42,17.
TetrakidN-propylaminomethyl)methane
1 eq. pentaerythrityltetrabenzenesulfonate and 8 eqs. N-propylamine were
refluxed in N-methylpyrrolidone during 80 hours. The solvent was then
removed and 8 eqs. powdered KOH and ethanol were added to the residue.
After stirring for some time, the ethanol was distilled off and the crude
product was extracted with diethylether, which was removed subsequently.
The product, a colourless liquid with a strong odour, was obtained by
fractional vacuum distillation, bp. 130~~/0,028 mbar. Anal. calcd. for
C17H40N4: C 68,0, H 13.33, N 18,66. Found: C 67,90, H 13,37, N 18,37.
The diisocyanates hexamethylenediisocyanate (HDI), isophoronediisocyanate
(IPDI), trans 1,4-cyclohexanediisocyanate (tCHDI) were commercially
available, and ethyl 2,6-diisocyanatohexanoate (lysine diisocyanate, LDI)
was synthesized according to a previously described procedure (32). In
fig. 1 the structural formulas are shown. All diisocyanates were vacuum
distilled prior to use.
Polyurethane network preparation
A polyol and a diisocyanate were thoroughly mixed in stoichiometric
proportions ([NCOI/[OH]=l) at roomtemperature under a nitrogen atmosphere.
The homogeneous, colourless mixture was degassed repeatedly and allowed to
gelate at room temperature. Post-curing at a temperature above T (i.e. gm
the glass transition temperature of the fully cured sample) yielded
optically transparent, glassy thermosets. Cure of the polyurethane
thermosets could also be achieved isothermally at a temperature above T 9'
Alternatively, polyurethane networks were formed very fast when the
mixture was polymerized by means of microwave heating. Within minutes
gelation was achieved. Microwave cure was conducted in an ordinary
domestic microwave oven (microwaves with 2,45 GHz frequency).
All polyaminoalcohols, or poly(hydroxyalkyl)amines, from figure 1 were
colourless, viscous liquids at room temperature, except TIPA which
crystallized very slowly to a waxy solid after melting or distillation.
Prior to mixing this trio1 had to be melted, just like TMP. The latter
compound was mixed with diisocyanates at temperatures 80-110 OC, and also
allowed to gelate at these temperatures. Gelation in these formulations
was rather fast (ca. 0,5 hr. 1 . All diisocyanates from fig. 1 were
colourless liquids at room temperature, except tCHDI, melting at 60 OC.
This diisocyanate had to be mixed with polyols above its melting point.
Gelation at these temperatures is very rapid. Triethanolamine was miscible
with liquid diisocyanates only at temperatures higher than 40-50 'c.
Polyisocyanurate network formation
Hexarnethylenediisocyanate. containing 0,25 wt.-% stannous octoate as a
catalyst, was allowed to polytrimerize at 125 OC for some days. Gelation
at this temperature took about three hours. The glassy, transparent solid
was post-cured at 180 OC.
Characterization
Glass transition temperatures of the polyurethane networks were determined
using a Perkin-Elmer DSC-7, calibrated with ICTA (International
Confederation for Thermal Analysis) certified reference materials, and
operated at a scan-speed of 10 'chin.
Tensile testing was performed on rectangular-shaped specimens (ca. 50x6~2
mm), machined from glassy polyurethane samples, at room temperature using
an Instron (4301) tensile tester, equipped with a 5 kN load-cell, at a
cross-head speed of 10 mm/min. The gauge length was 25 mm. The reported
tensile data are the mean values from at least six tests. Failure was
nearly always initiated at the clamps as would be expected for rectangular
shaped specimens.
Swelling measurements were carried out on polyurethane samples weighing
less than ca. 0,5 g. that were immersed in chloroform at room temperature
for 2 days. The volume degree of swelling was calculated from the weight 3 increase, using the densities of chloroform (p=1,48 g/cm and the
polyurethane networks. Both the Quadrol/HDI-based network and the
TIPA/HDI-based network had a density of 1,135 g/cm3, which were determined
by weighing samples in air and by weighing them submersed in water.
Infra-red measurements were carried out on ultra thin (ca. 10 pm)
polyurethane films with a Bruker IFS-88 FT-IR spectrometer.
UV/VIS transmission spectra of samples having 1 mm thickness were recorded
on a Pye Unicam SP 8-200 UV/VIS spectrophotometer.
The refractive index of a Quadrol/HDI-based polyurethane network was
measured on a thin film wetted with dichlorobenzene (having a higher
refractive index than the polymeric material) with an Abb6 refractometer
( ATAGO .
Implantations
Circular disks with smooth edges having 6 mm diameter and 1 mm thickness,
containing one big central hole and three smaller peripheral holes for
fixation, were constructed from a fully cured Quadrol/HDI-based
polyurethane network sample. After a cleaning procedure, the disks were
sterilized by autoclaving at 120 OC for 20 minutes. Three disks were
inserted into the anterior chamber of rabbit (Chinchilla) eyes through a
corneal incision. Prior to implantation, the eyes had been made aphakic
(lens less). The disks were hung up (fixated) in front of the pupil using
two stainless steel 70 pm wires, led through the peripheral holes. The 70
pm wires were knotted together on the sclera.
One "mushroom"-shaped keratoprosthesis (artificial cornea) was cut on a
lathe from a fully cured Quadrol/HDI-based polyurethane network sample
(see figure 2). After polishing, cleaning, and autoclaving the
keratoprosthesis for 20 minutes at 120 OC, it was implanted by Van Andel
as a "porthole" in a cornea of an aphakic (Chinchilla) rabbit eye and
fixed on the eye like a "champagne cork" on a bottle. The central "column"
of the keratoprosthesis (the "leg" of the "mushroom") perforates the
centre of the cornea through a 3 mm trephined hole. The "hat" of the
"mushroom" (with a diameter of 6 mm) lies on the cornea and is "anchored"
with two permanent 70 pm thin soft stainless steel (type: AINSI 316) wires
around the whole eyeball, in two planes perpendicular to each other. The
keratoprosthesis works now as a valve: the peribular fixation keeps the
valve on the trephined hole and the internal pressure of the eye pushes
the corneal rim around the trephined hole against the back of the "hat" of
the "mushroom". The resulting pressure on the interface between
keratoprosthesis and cornea prevents the leaking of aqueous humour, the
melting away of corneal tissue, and the epithelial downgrowth
(fistula-formation) (33 ) .
Figure 2. A "mushroom"-shaped keratoprosthesis.
Results and discussion
As outlined in the introduction, densely cross-linked polyurethanes might
be interesting materials with potential ophtalmic applications, like
autoclavable intraocular lenses, and keratoprostheses.
First, it will be described how such polyurethane networks can be
synthesized. From the literature numerous examples of elastomeric model
polyurethane networks, formed by the endlinking of hydroxylterminated
prepolymers with polyisocyanates, are known (34-40). In contrast to
elastomeric polyurethane networks, densely cross-linked polyurethane
networks are relatively unknown (41,421. These model networks are formed
by stepreactions, which according to Boots is an intrinsically homogeneous
process, unlike network formation by chain reactions (43.44). In this way
polymer networks with a well-defined topology and a minimal number of
dangling-end network imperfections can be obtained when the cross-linking
reaction is carried out stoichiometrically and to very high conversion of
the functional groups (45). DuSek and Stepto pointed out that the extent
of cyclization in model networks prepared in the absence of a diluent is
rather small, but never negligible. Dilution may considerably increase
intramolecular reactions, leading to elastically inactive loop structures
and may cause the formation of inhomogeneities, or even phase separation
(micro-, macrosyneresis) (39,461.
All this has led to the idea to prepare polyurethane networks from pure,
low molecular weight polyols and diisocyanafes in the bulk and to obtain
highly cross-linked networks that are in principle homogeneous and contain
a minimal concentration of network imperfections, which should result in
materials with good (ultimate) mechanical and optical properties.
It is obvious that a first requirement for a network-forming reaction, in
which more than one component participates, is that, in the absence of a
solvent, the reactive components have to be miscible. Applying this to the
polyurethane network formation this means that the polyol and the
diisocyanate have to miscible. It appeared that not all low molecular
weight polyols from figure 1 were miscible with the (cyc1o)aliphatic
diisocyanates listed. For instance, pentaerythritol or glycerol were not
miscible with any of the diisocyanates at any temperature. No sign of
reaction could be observed. However, in the presence of a solvent (e.g.,
DMF) at room temperature, transparent swollen gels were obtained.
A second requirement is that in the case of a miscible formulation the
reactivity of the components should be low, low enough to achieve complete
miscibility before the polymerization reaction takes place in a
controllable manner. It turned out that some polyols,
tetrakis(2-hydroxyethy1)ethylenediamine. tetrakis(2-hydroxyethy1)methyl-
aminomethylmethane, BIS-TRIS were too reactive. During mixing these
polyols with diisocyanates noticeable polymerization (gelation) took
already place, resulting in a very macroscopically heterogeneous network
formation. It was also seen that low molecular weight polyfunctional
amines, primary or secondary amines, tri- or tetrafunctional (see fig. 11,
which are known to be highly reactive towards isocyanate groups (resulting
in urea bonds instead of urethane bonds in the case of the reaction of an
alcohol with an isocyanate), reacted instantaneously with diisocyanates,
both in the absence and in the presence of a solvent (for instance, DMF,
toluene). So, macroscopically homogeneous polyurea networks could not be
obtained by the direct addition of polyfunctional amines and
diisocyanates, due to the fact that the very fast polymerization reaction
interfered with the molecular mixing. Any attempt to slow down the
reactivity of the amine group was unsuccessful.
The capping or blocking of isocyanates is a way to prevent premature
reaction of the polyurethane (or polyurea) components. Isocyanate groups
are reacted with compounds which form a thermally weak bond. The reactive
isocyanate can be liberated at elevated temperatures [ca. 150 OC or
higher). Examples of compounds used for the blocking of isocyanates are:
phenols, caprolactam, oximes and 6-dicarbonyl compounds (ethyl malonate)
(47-49 ) . Fortunately, there were several low molecular weight polyols, having a
relatively low reactivity, that were miscible with diisocyanates:
tetrakis(2-hydroxypropy1)ethylenediamine (Quadrol), triisopropanolamine
(TIPA), triethanolamine (TEA), octakis(2-hydroxyethy1)pentaerythrityl-
tetraamine ("octaol"), bis(2-hydroxyethyl)isopropanolamine (BHEIPA),
trimethylolpropane (TMP). At first sight it looks like these polyols have
nothing in common, but this is an overstatement. First of all, one can see
that the polyols containing only secondary alcohol groups (TIPA, Quadrol,
octaol) are miscible with diisocyanates. All three compounds are so-called
secondary polyaminoalcohols, or poly(hydroxya1kyl)amines which can be
regarded as the reaction products of the corresponding amines and
propylene oxide. It is noteworthy that triethanolamine can be mixed with
diisocyanates whereas the analogous tetrakis(2-hydroxyethy1)-
ethylenediamine was considered too reactive to be miscible with
diisocyanates to form a homogeneous mixture that can be polymerized
controllably. This may be explained by the fact that the gelation of the
trifunctional primary polyol TEA/diisocyanate formulation is much slower
than the gelation of the tetrafunctional primary polyol/diisocyanate
mixture. The same was observed in the corresponding secondary polyol
systems: TIPA/HDI gelates at room temperature in ca. two days, Quadrol/HDI
gelates at room temperature in ca. 6-10 hours. So in the case of the trio1
the time available for mixing (before gelation sets in) is much longer
than in the tetra01 case. Due to the fact that primary alcoholgroups are
more reactive towards NCO groups than secondary alcohol groups, the
gelation time becomes shorter when in the formulation TIPA/HDI TIPA is
replaced with TEA, or in the case Quadrol/HDI Quadrol is replaced with the
analogous primary tetraol. In the latter case the polymerization
(gelation) interferes with the mixing. The gel time is too short to permit
sufficient mixing. The compound BHEIPA can be considered a polyol with
properties lying between those of TEA and TIPA. Trimethylolpropane,
finally, which also contains three primary OH groups, just like TEA, was
also miscible with diisocyanates above its melting point (60'~) (From the
literature aromatic, highly cross-linked polyurethane networks made from
TMP and MDI are known (41)). In this context it is strange that glycerol,
containing two primary and one secondary hydroxyl group, can not be mixed
with any diisocyanate. From the above, it may be clear that it is rather
unpredictable whether or not a certain pair of polyol/diisocyanate is
miscible.
Figure 3. The mechanism of the tertiary arnine catalyzed urethane reaction.
Another interesting thing that the above mentioned polyols have in common,
except TMP, is the presence of tertiary amino groups in the molecular
structure. Tertiary amines are known catalysts for the urethane reaction
(26,47-49). The mechanism for tertiary amine catalysis involves the
formation of an intermediate active complex of the isocyanate and the
tertiary amine. After the OH group is added, the intermediate complex
rapidly rearranges to the urethane linkage (figure 3 ) . So, these
polyaminoalcohols contain an internal, built-in catalyst (48). No external
catalyst is needed for the polyurethane reaction, which is interesting
with regard to the potential biomedical application of the polyurethane
networks. Added catalysts, for instance toxic tin salts, may be
responsible for complications after implantation.
The whole problem of miscibilty can be avoided by using only one compound
in the formation of a polymer network. An example is the polytrimerization
of diisocyanates in the presence of a suitable catalyst (e.g. stannous
octoate), resulting in polyisocyanurate networks. At temperatures higher
than 120 OC HDI formed a glassy, transparent, densely cross-linked network
with a high ultimate glass transition temperature of 150 OC. This high T 4m
is likely due to the presence of rigid 6-membered isocyanurate ring
structures in the network.
Polyurethane networks were prepared by firstly mixing the polyol component
and the diisocyanate under a nitrogen atmosphere, to avoid side reactions
of the isocyanate group with water which gives rise to undesired bubble
formation (or foaming). After mixing, the homogeneous, colourless viscous
mixture was allowed to gelate (usually at temperatures below Tgm).
Directly after gelation, the glass transition temperature of the network
is equal to the gelation temperature. Due to vitrification the chemical
reactions in principle are quenched. The mobility of the functional groups
in the glassy network that have not reacted yet is very low, which means
that the network formation takes place very slowly. After vitrification,
the cross-linking reaction becomes diffusion-controlled ( 5 0 ) . The samples
were fully cured at temperatures above the ultimate glass transition
temperature of the polyurethane networks T In order to reach a maximum 4m'
conversion of the functional groups, the cure temperature has to be above
the T of the system. The glass transition temperature is a sensitive gm
measure of the functional group conversion (51).
Instead of the conventional thermal curing of the samples, the
polyurethane resins may be cured by means of electromagnetic waves in the
microwave frequency range (2,45 GHz 1. In contrast to thermal heating,
which involves heat conduction and thermal lag associated with it,
microwaves can generate heat directly within the sample and thus offer the
possibility of very fast, uniform curing which should result in improved
physical/mechanical properties of the final material (52). Microwave
heating has been used for the cure of epoxy resins and polyurethane foams
(53).
All polyurethane networks were optically transparent, colourless,
amorphous, glassy materials. The refractive index of a fully cured
Quadrol/HDI-based polyurethane network is 1,50, slightly higher than the
value (1,491 reported for PMMA 7 9 1 1 1 5 1 The networks were
macroscopically homogeneous; apparently no phase separation had occurred
during the process of network formation. In table 1 the ultimate T 's of 9
the densely cross-linked networks, obtained from the bulk reaction of a
polyol and a (cyc1o)aliphatic diisocyanate in stoichiometric proportions,
are collected.
Table 1. Ultimate glass transition temperatures (in OC) of polyurethane
networks obtained from the bulk polymerization of a polyol and a
diisocyanate in stoichiometric proportions
TI PA Quadrol octaol TEA BHEI PA TMP
HD I 75 85 106 33 45 83
LD I - 72 - 34 - - IPDI 160 165 - 133 - 184
tCHDI 178 190 - 125 - 227
The ultimate glass transition temperature of a HDI-based polyurethane, as
can be seen in table 1, increases when the functionality of the polyol
increases (going from TIPA to Quadrol to the octaol), which is due to an
increase in the cross-link density. The T is also markedly raised when gm
an aliphatic diisocyanate (HDI, LDI) in a particular formulation is
replaced with a cycloaliphatic one (t-CHDI, IPDI), which is a consequence
of the higher rigidity in the case of a cycloaliphatic polyurethane
network.
Since all these polyols, and all diisocyanates as well, are miscible, the
TP of a polyurethane network can be varied endlessly. The T of a
4m multi-component system then lies between the T 's of the networks
4m resulting from the individual pair of reactants (two-component system). An
interesting example is the polyurethane network made from a polyol mixture
of TEA and TIPA in a 2 : l mole ratio, and HDI in stoichiometric
proportions. This polyurethane had a T (44 OC) that was virtually equal '3m
to the one of the BHEIPA/HDI-based polyurethane ,as could be expected, and
inbetween the T values of the TEA/HDI-based and the TIPA/HDI-based gm
polyurethane network.
Another way of influencing the T is to carry out the network-forming gm
reaction using off-stoichiometric proportions of reactive groups. For
example, when 1 eq. Quadrol is reacted with 1,5 eqs. HDI, instead of the
stoichiometric 2 eqs, the T of the final network drops ca. 20 OC. In 4m
this case a considerable amount of unreacted alcohol groups (dangling
ends) are present in the network, which are responsible for this lowering
of the T One can also say that the cross-link density is much lower gm'
than in the case of a stoichiometric network formation. Another example is
the drop in T when the octaol is reacted with 2 eqs. HDI instead of the 4m
stoichiometric 4 eqs.. The ultimate glass transition temperature of the
network now drops from 106 OC to 54 OC.
From swelling measurements using chloroform, it could be established that
the polyurethane networks were definitely highly cross-linked, i.e. having
a low degree of swelling. The volume degrees of swelling for the
Quadrol/HDI-based and the TIPA/HDI-based polyurethane network were 1.27
and 1.60, respectively. The higher degree of swelling of the trifunctional
aliphatic polyurethane network compared to the corresponding
tetrafunctional network is ascribed to the lower cross-link density in the
trifunctional network. It should be mentioned that in all cases it was
found that the gel content was 100 % (The drying of the swollen networks
to constant weight took about 3 weeks in a vacuum oven at 100 OC). No sol
fraction could be detected in any cross-linked polyurethane network. It
has been said before that by extraction of tight polymer networks, and
subsequently drying, the gel content can not be determined accurately,
since the actual sol fraction may not be able to diffuse out of the dense
network.
Infra-red spectra of fully cured Quadrol/HDI-based or TIPA/HDI-based
polyurethane networks, both spectra were virtually identical, revealed the
absence of an isocyanate absorption at ca. 2250 cm-l, indicating the
completeness of the cross-linking reaction. Apparently, all isocyanate
groups have reacted, either with the hydroxyl groups (urethane bond
formation) or with urethane bonds to form allophanate linkages. The latter
should not be excluded since the curing was carried out at ca. 100-110 OC,
at temperatures where the allophanate formation becomes competitive with
the urethane formation (54). The isocyanate absorption at 2260 cm-',
however, was not absent in the infra-red spectrum of the HDI-based
polyisocyanurate network. This indicates that in the case of a
polytrimerization, where three NCO groups have to react simultaneously to
form ring-like isocyanurate structures, residual isocyanate groups can be
detected in the final network. Apparently, a complete conversion of the
reactive groups can not be attained.
The Quadrol/HDI-based and TIPA/HDI-based polyurethane networks were
subjected to tensile testing. Both glassy materials showed virtually
identical stress-strain behaviour. Figure 4 shows a typical stress-strain
curve of a Quadrol/HDI-based polyurethane network. The tensile strength of
the two aliphatic polyurethane thermosets was in the range 80-85 MPa. Both
formulations had comparable moduli (1,5 GPa) and comparable strains at
break (ca. 15 % I . The polyurethane glasses usually yielded prior to break.
a phenomenon also displayed by other glassy cross-linked networks, for
example epoxies (55-571, and by glassy thermoplastic polymers, like PMMA,
as well. The plastic deformation of glassy thermosets below T is similar
to that of amorphous glassy thermoplastics (58) . The mechanical properties
of these densely cross-linked polyurethane networks will be discussed in
more detail in a forthcoming paper (591, in which also results of dynamic
mechanical measurements will be presented.
5 10 15 20
Strain (%)
Figure 4. Stress-strain behaviour of a glassy polyurethane network
obtained from the bulk polymerization of Quadrol and HDI in stoichiometric
proportions.
Some densely cross-linked polyurethanes were immersed in water for about
two months to measure their water uptake by weighing. The water uptake of
Quadrol/HDI-based, octaol/HDI-based, TEA/HDI-based polyurethanes were 1.2
%, 1%, and 8 %, respectively. The water uptake of the Quadrol/HDI-based
polyurethane was increased when the cross-linking reaction was carried out
off-stoichiometrically. The residual, polar hydroxyl groups in the
network, obtained from the polymerization of Quadrol with 1,s eqs. HDI,
were responsible for a higher water uptake (3.1 %) of this network. The
same was observed for the octaol/HDI-based network. The water uptake was
increased to ca. 8 % when the octaol was reacted with 2 eqs. HDI instead
of the stoichiometric 4 eqs. In any case it was seen that the uptake of
water dramatically lowered the T of the network, due to a plasticizing g
effect. The T of the stoichiometric Quadrol/HDI network dropped from 85 g
to 48 OC, and the T of the TEA/HDI network dropped from 33 to -20 OC
(resulting in a rubberlike polymer). It is well-known that the presence of
water in polymer networks positively contributes to their
biocompatibility. Hydrogels, usually water swollen polymer networks, are
said to be biocompatible due to a low interfacial tension which may be
exhibited between the hydrogel surface and an aqueous solution (as in
living tissue). An attempt to make the polyurethane networks more
hydrophilic, so that the water uptake would increase, was the
incorporation of low molecular weight polyols containing besides OH
groups, carboxyl groups in their molecular structure:
N.N-bis(2-hydroxyethy1)glycine (bicine), bis(hydroxymethy1)propanoic acid.
These two compounds could be "dissolved" in the polyaminoalcohols from
figure 1 only at high temperatures (ca. 100-150~~). The resulting clear
polyol mixture was then reacted with the stoichiometric amount of HDI, but
the water uptake of the resulting network was not noticeably higher than
in the case of a normal polyurethane formulation. Apparently, the carboxyl
groups had reacted during the course of the experiments, either with the
hydroxyl groups (esterification) or with the NCO groups.
Besides low molecular compounds, linear polymers may be incorporated into
the polymer network. Thus, so-called semi-interpenetrating networks are
formed. For example, poly(N-vinylpyrollidone), a water-soluble polymer,
can be dissolved in Quadrol (at elevated temperatures), and the resulting
polymer solution can be cross-linked with a diisocyanate to form a glassy
thermoset.
PMMA intraocular lenses have been surface modified with covalently linked
heparin to increase the biocompatibility (60,611. A (PMMA) lens surrounded
by a solution of hydrophilic monomers can be 7-irradiated, resulting in a
thin coating of grafted hydrophilic polymer (62). Another method used to
modify the surface of a polymeric material is the glow-discharge technique
(53,631. In this way surfaces can be made more hydrophilic, which might
also work out nicely for the polyurethane networks described here.
Alternatively, the surface of the polyurethanes might be modified to
increase the hydrophilicity by usual grafting techniques or by the two
following methods. Since NCO groups can also react (slowly) with amide
groups, the liquid polyol/diisocyanate mixture may be poured onto a dry
polyacrylamide film, and subsequently cured at high temperatures (ca.
1 1 0 ~ ~ ) . In this very straight-forward manner polyacrylamide may be grafted
chemically to the polyurethane. Another possibility is the "surface
etching" of the polyurethane network with a dicarboxylic acidchloride
(e.g. succinyl chloride) for a short period of time, and then placing the
whole sample in water in order to create carboxyl groups at the surface.
This last method is in slight analogy with the sodium hydroxide
sterilization technique used for the sterilization of PMMA, for instance
(7,9). Local hydrolysis of the methyl ester at the surface through the
action of sodium hydroxide results in the formation of carboxyl groups at
the surface, making the surface more hydrophilic. Although not intended as
a surface modification method, this sodium hydroxide sterilization might
act this way in case of PMMA.
As a pilot experiment, three disks made of the Quadrol/HDI-based
polyurethane network were implanted in rabbit eyes to see how the material
was tolerated in the eye. All three implantations, basically, gave the
same results. A serious opacification of the whole cornea, indicating an
adverse reaction to the implant material, was never seen. The cornea was
nearly completely clear two weeks after the surgical procedure. Only near
the site of the corneal incision, the cornea was a little hazy, and
starting from that site some vascular ingrowth, which was likely due to
some surgical trauma, could be seen. Figure 5 shows a photo of a rabbit
eye c0ntainin.q a c i r cu l a r d l s k m ~ d e o r b h ~ 1 1 3 ~ h l y rross-linked
polyurethanp tun u c c k s af t r r ~ m p l q n t a t i o n . From t h e ~ c p r c l l r n i n a r y
experiments i t may LC voncl~d-d t h t t h e !~ighly cr uss-1 inked pwlyur e t h a n e
material was rattler well tolerated i r k [lie r a b b i t eye. The material d i d n o t
t u rn out Lo be acute toxic and seems su l led for use i n ophtalmic
applications. These encouraging results have l e d t o t h r addi l iona l
implantation nf a k~ratoprostheris l n a r a b b f t e y e .
The preliminary experiment with onc lathe-cut keratoprosthesis uT t h e
prradro:/IIDI-based p o l y u r e t h a n n nclwurk implantrd i n on? Chinrhllla-rabbit
eye also showed Chat t i le materlal was t o l ~ r a t c d very w ~ l l by t h e hcal thy
rabbit eyr. I n e cornea stayed clear, indicating t h a t t h ~ ~ n d n t h ~ l i u m ,
stroma and epitheljum n f t h ~ cnrnp: and t h p antericr cycrhamber d i d not
Figure 5 . Photograph of a rahbi t ~ y c with a clrrular Implant made of a
highly cross-linked Quadrol/lIDI-bnscd po:yure thane nclwork t w o weeks aC ter
implantation.
Figure 6 . Pho tugrdp i u f a I dbLlL cys5 W L t l r a kera toprosthesis made o f a
h i p ; h l y crnss-l i nk rd Quadrol / l lnI-based po lyure t l l ane network o n e year ar ter
i m p l a n t a t i u r ~ .
Figure i. Ha.bb i r e y e 3 : L I; . i n i sla.inL~!:;c. I. 'J/YI,+SS
keratopros t n e s i s
react on the implant and/or its material. Even one year after implantation
(see figure 61 the eye was still "clear" as it is called in the clinical
terms. The opacification seen around the keratoprosthesis and the
incision, made to extract the lens, are the result of
"Kammerwasserausbruch": a reaction of the stromal tissue to the leakage of
aqueous humour into the stroma of the sclera, unavoidably caused by the
surgical trauma. The overgrowth of the keratoprosthesis by the stroma and
epithelium of the sclera is not a reaction to the material as such. This
also happens after implanting a keratoprosthesis made of inert PMMA CQ
(clinical quality) or inert stainless steel/glass (see figure 7 )
(33,64,65). These results show that the transparent highly cross-linked
polyurethane network is suited to make inert, autoclavable
keratoprostheses.
The crystalline lens and cornea filter most of the solar ultraviolet
radiation, having wavelengths from approximately 285 to 400 nm and which
itself is damaging to the retina, to which the eye is exposed. Quanta with
wavelengths below 300 nm are almost completely absorbed by the cornea.
Ultraviolet radiation ranging from 300 to 400 nm is normally absorbed by
chromophores in the crystalline lens. After cataract extraction, the
posterior segment of the eye is therefore exposed to UV radiation not
normally encountered in the phakic state. The ideal intraocular lens
should be ultraviolet light absorbing, i.e. absorb UV light with
wavelengths below 400 nm. Nowadays nearly all intraocular lenses contain a
UV-absorbing chromophore, either in the form of a low molecular additive
or polymer bound (66-681. It appeared that Coumarin 102 (see figure 81, a
so-called laser dye, was soluble in the polyol component of the
polyurethane formulation. This compound could withstand thecuring process
without losing its UV light absorbing activity, unlike the conventional
hydroxyl containing UV absorbers, like hydroxybenzophenones or
hydroxybenzotriazoles (69) (these compounds contain hydroxyl groups that
can react with the diisocyanate used and consequently losing their
UV-absorbing ability). In figure 8 UV/VIS transmission spectra are shown
of a Quadrol/HDI-based polyurethane sample ( 1 mm thickness) with and
without the additive Coumarin 102, in low concentration. As can be seen
the polyurethane thermoset without the additive transmits light with
wavelengths above 275 nm (ca. 90 % transmission). A few promille Coumarin
102, which is a fluorescent compound with its absorption maximum at h=390
nm and its fluorescence maximum at h=468 nm, is enough to completely
absorb the UV light with wavelengths below 450 nrn.
Coumarin 102 FH'
200 300 400 500 600 700
Wavelength (nm)
Figure 8. UV/VIS transmission spectra of the additive-free
Quadrol/HDI-based polyurethane network ( A ) , and of the same material
containing 0,08% Coumarin 102 ( B ) , a UV-absorbing chromophore (sample
thickness 1 mm) .
Finally, there may be some concern about the long-term stability of the
highly cross-linked polyurethanes. As known from the literature (26,701,
polyurethanes in general do (bioldegrade, although nearly all
polyurethanes investigated are (uncross-linked) thermoplastic elastomers.
In vivo degradation is usually hydrolytic, although tissue enzymes may
also participate in the degradation process. Polyurethanes can be made
deliberately degradable (see ref. 32), but also hydrolytically stable.
Since the polyurethane networks described here are rather hydrophobic and
very densely cross-linked, pronounced degradation is not to be expected.
An indication for this is that the T of a Quadrol/HDI-based sample g
immersed in water for 20 months did not differ from the value of a sample
of the same material immersed in water for 2 months ( T =48 O C ) . Another 4
noticeable feature is that the rabbit eye with the implanted polyurethane
keratoprosthesis was quiet one year after the surgical procedure.
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Summary
This thesis is concerned with the synthesis, physical properties and
potential biomedical applications of some polyurethanes, especially
cross-linked ones. Polyurethanes are a class of polymers having only in
common the presence of urethane bonds somewhere in their chains. The name
polyurethane given to a polymeric material does not tell anything about
its chemical and physical characteristics. Polyurethanes may be lightly or
highly cross-linked or uncross-linked and be highly crystalline,
elastomeric or amorphous and glassy. In the biomedical field polyurethane
usually stands for thermoplastic polyurethane elastomer. Thermoplastic
polyurethanes, also named segmented polyurethanes, are linear
blockcopolymers composed of chainextended diisocyanate hard segments
dispersed in a soft segment polyol matrix. Due to their good mechanical
properties (high tensile strength, good tear strength, high toughness,
good flex life), reasonable bloodcompatibility and biocompatibility,
elastomeric polyurethanes have been used in many medical applications,
like total artificial heart, heart valves, vascular prostheses, wound
dressings etc. By mixing segmented polyurethanes with (5-20 wt.%) high
molecular weight poly(L-lactide) (PLLA), Gogolewski, Leenslag and Pennings
in Groningen developed elastomeric, biodegradable mixtures with remarkable
in vivo performance. Quenched physical polyurethane/PLLA mixtures, in
porous form, were succesfully applied as a small-caliber vascular
prosthesis, artificial skin, meniscus lesion repair material, nerve guide.
Since these polyurethanes are not chemically cross-linked (formation of
permanent cross-links), they show stress softening (stress hysteresis)
when subjected to cyclic deformation. This problem can be overcome by
chemically cross-linking the linear polyurethane chains, for instance,
with peroxides (see chapter 5 ) . Another drawback of commercial biomedical
polyurethanes concerns their chemical composition. Nearly all these
polyurethanes (Biomer, Estane, Pellethane, etc.) are composed of an
aromatic diisocyanate MDI (4,4'-methylene diphenyl diisocyanate).
Degradation (through hydrolysis) of the polymer may result in the
formation of the toxic, carcinogenic, mutagenic MDA.
4,4'-methylenedianiline, the degradation product of the incorporated
aromatic diisocyanate. Although it has not been shown unambiguously that
MDI-based polyurethanes induce the formation of cancer, it would be more
elegant and safer to seek for a replacement for this component in the
polyurethane formulation. The use of cycloaliphatic diisocyanates, for
instance, hydrogenated MDI in Tecoflex, also leads to segmented
polyurethanes with good ultimate properties. Aliphatic diisocyanates,
which are not particularly suited for the synthesis of thermoplastic
polyurethane elastomers, may be used for the formation of chemically
cross-linked polyurethanes. Especially aliphatic diisocyanates, producing
non-toxic diamines (e.g., lysine, 1,4-diaminobutane) after eventual
degradation, seem the ultimate choice for the synthesis of biomedical
polyurethanes.
In chapter 2 such lysine diisocyanate-based elastomeric polyurethane
networks are described. These polyurethane networks, designed to release
only non-toxic degradation products, were prepared by cross-linking
hexafunctional starshaped prepolymers with ethyl 2,6-diisocyanatohexanoate
(i.e., lysine diisocyanate). The hydroxy terminated prepolymers were
synthesized by the ring-opening copolymerization of L-lactide or glycolide
and E-caprolactone initiated by myo-inositol, a vitamin. The
polyesterurethane networks, having T 's in the range 0-10 OC and gel
contents of 90-95 %, showed rubber-like behaviour. It is noteworthy that
the chloroform-extracted networks exhibited much better tensile properties
(tensile strength ca. 30-40 MPa) than the unextracted networks (tensile
strength ca. 10 MPa). Only the extracted networks exhibited pronounced
strain-induced crystallization. The presence of plasticizer (sol fraction)
suppressed the strain-induced crystallization.
In chapter 3 the polyurethane networks described in the previous chapter
are evaluated as potential materials for the construction of a macroporous
bottom-layer (dermal analogue) in a multi-layer artificial skin. An
amorphous, elastomeric, porous lysine diisocyanate-based poly(glyco1ide-
co-E-capro1actone)urethane network degraded fast in vitro. In vivo the
same material was degraded even faster. Subcutaneous implantation in
guinea pigs showed that the porous polyurethane networks degraded almost
completely between 4 and 8 weeks after implantation, allowed rapid cell
ingrowth and evoked no adverse tissue reaction. The lysine
diisocyanate-based polyurethane networks can be considered biocompatible.
Chapter 4 is concerned with the preparation and characteristics of a
porous polyurethane wound covering. A very thin, porous membrane (15-20
pm) was prepared by means of a phase inversion process. This elastic film,
made of a cycloaliphatic polyetherurethane (Tecoflex), contained
micropores up to approximately 5 pm. The porous wound covering was
impermeable to bacteria. The polyurethane membrane appeared to be very
permeable to water vapour, whereas water in liquid form or wound exudate
could not leak through the membrane. In guinea pigs epidermal wound
healing of partial-thickness wounds under polyurethane wound coverings was
accelerated as compared with uncovered controls and an occlusive wound
covering (Op-Site). The high water vapour permeability of the polyurethane
wound covering induced concentration of the wound exudate into a jellylike
clot layer, which apparently accelerated reepithelialization. The main
conclusion from a clinical study on 20 split-skin donor sites was that the
use of the polyurethane covering reduces pain (as compared with the
conventional treatment of tulle gras dressing), besides prevention of
fluid retention and enhanced reepithelialization.
Chapter 5 describes a two-ply biodegradable artificial blood vessel made
of polyurethane and poly(L-lactide). The microporous innerlayer of the
small-caliber vascular prosthesis was constructed from a cycloaliphatic
segmented polyurethane (Tecoflex) cross-linked with dicumylperoxide in the
presence of linoleic acid. The reason for introducing chemical (permanent)
cross-links into the polyurethanes is to eliminate the serious limitation
of stress softening which occurs especially in cyclic loading. Cyclic
creep-failure, resulting from the arterial pulsation of the blood, may
lead to the formation of aneurysms (catastrophic dilation of blood
vessels). Furthermore, carboxyl groups of the linoleic acid on the luminal
side of the prosthesis contribute positively to the antithrombogenicity of
the artificial blood vessel. It appeared that adding linoleic acid during
the peroxide vulcanization led to a maintenance of the tensile strength of
the prostheses. The outer ply was constructed by precipitating a (95/5)
physical mixture of a polyesterurethane and poly(L-lactide) from solution
in the presence of sugar crystals in the range 30-90 pm which were removed
by exposing the prosthesis to water. The two-ply vascular prostheses were
tested in vivo by replacing 1 cm of the abdominal aorta of rats. All the
prostheses remained patent at least up to one year and did not exhibit any
aneurysmal formation, which has to be ascribed to the improved
creep-resistance of the prosthesis as a result of the cross-linking. The
inner layer of the prosthesis was covered with endothelial cells and
several layers of smooth muscle cells. Essential components of a
neo-artery were regenerated.
Chapter 6 deals with the synthesis, properties and potential ophtalmic
applications (intraocular lenses, keratoprostheses) of highly cross-linked
polyurethane networks. Such polyurethanes were prepared by the bulk
stepreaction of various low molecular weight polyols and (cyc1o)aliphatic
diisocyanates. All these polyurethane networks were optically transparent,
colourless, amorphous glassy thermosets. The properties of one particular
glassy polyurethane, obtained from the bulk reaction of a tetrafunctional
aminoalcohol tetrakis(2-hydroxypropy1)ethylenediamine (Quadroll and
hexamethylenediisocyanate (HDI) in stoichiometric proportions, were
investigated in more detail. This glassy polyurethane, with an ultimate
glass transition temperature of 85 OC, and a very low degree of swelling
in chloroform (1,271, exhibited good ultimate mechanical properties
(tensile strength 80-85 MPa, elongation at break ca. 15 X , modulus ca. 1,5
GPa). Infra-red spectra of these polyurethane networks revealed the
absence of an isocyanate absorption, indicating that all isocyanates.
apparently, had reacted during the cross-linking reaction.
These transparent cross-linked polyurethanes can be sterilized simply by
autoclaving, in contrast to polymethylmethacrylate (PMMA) which has been
used successfully as an intraocular lens material the last 15 years. The
possibility of an autoclavable lens is especially interesting with respect
to eye surgery in the developing world where the majority of the blind
people live. These highly cross-linked Quadrol/HDI-based networks, after
being autoclaved, were implanted in rabbit eyes, either in the form of
small circular disks or in the form of a keratoprosthesis (artificial
corneal. It was shown that the material was well tolerated by the rabbit
eyes. A serious opacification of the cornea, indicating an adverse
reaction to the implant, was never seen. One year after implantation of a
polyurethane keratoprosthesis the eye was still "quiet". These results
show that the transparent highly cross-linked polyurethane network seems
suited for use in ophtalmic applications.
Samenvatting
Dit proefschrift beschrijft een onderzoek naar de synthese, eigenschappen
en mogelijke biomedische toepassingen van een aantal (hoofdzakelijk
vernette) polyurethanen. De term "polyurethaan" heeft betrekking op een
klasse polymeren die slechts de aanwezigheid van urethaan bindingen ergens
in de polymere ketens gemeen hebben. De naam polyurethaan gegeven aan een
polymeer materiaal zegt weinig over de chemische struktuur en de fysische
eigenschappen ervan. Polyurethanen kunnen licht of sterk vernet of
onvernet zijn, hoog kristallijn, rubberachtig of amorf en glasachtig zijn.
In de biomedische wereld heeft de naam polyurethaan doorgaans betrekking
op thermoplastische polyurethaan elastomeren. Deze zogenaamde
gesegmenteerde polyurethanen zijn lineaire blokcopolymeren bestaande uit
harde segmenten ("chainextended" diisocyanaten) die gedispergeerd zijn in
een matrix gevormd door zachte polyol segmenten. Door hun goede
mechanische eigenschappen (hoge treksterkte, taaiheid, goede
scheursterkte) en acceptabele bloed- en biocompatibiliteit, hebben
multifase elastomere polyurethanen succesvolle medische toepassing
gevonden, bijvoorbeeld als kunsthart, hartklep, kunstaders, wondbedekking.
In Groningen werden dergelijke polyurethanen door Gogolewski, Leenslag en
Pennings gemengd met 5-20 X hoog moleculair poly(L-melkzuur) (PLLA). De zo
verkregen elastomere, biodegradeerbare mengsels voldeden opmerkelijk goed
in een aantal degradeerbare biomedische toepassingen.
Aangezien deze elastomere polyurethanen niet chemisch vernet zijn, maar te
beschouwen zijn als fysische netwerken, vertonen ze zogenaamde "stress
softening" bij cyclische belasting. Dit probleem kan verholpen worden door
de lineaire ketens alsnog chemisch te verknopen, bijvoorbeeld met
peroxides (zie hoofdstuk 5). Een ander nadeel van de commercieel
verkrijgbare biomedische polyurethanen heeft betrekking op de chemische
samenstelling. Bijna a1 deze polyurethanen (Biomer, Estane, Pellethane,
enz.) zijn opgebouwd uit een aromatisch diisocyanaat 4,4'-methyleendifenyl
diisocyanaat (MDI). Degradatie, door hydrolyse, van het polymeer kan
uiteindelijk leiden tot de vorming van het toxische, carcinogene, mutagene
4,4' -methyleendianiline (MDA) , het degradatieprodukt van het
gePncorporeerde aromatische diisocyanaat. Alhoewel het gezegd dient te
worden dat de implantatie van een polyurethaan op basis van MDI niet
ondubbelzinning aantoonbaar geleid heeft tot de vorming van een vorm van
kanker, zou het toch eleganter en veiliger zijn om te zoeken naar een
vervanger voor MDI in de polyurethaan samenstelling. In plaats van
aromatische diisocyanaten zijn cycloalifatische diisocyanaten gebruikt
(bijvoorbeeld in Tecoflex), hetgeen ook gesegrnenteerde polyurethanen met
goede eigenschappen oplevert. Vooral alifatische diisocyanaten, die er
voor zorgen dat niet toxische diamines (lysine, l,4-diaminobutaan) als
degradatieprodukten kunnen worden gevormd, lijken een goede keuze voor de
synthese van biomedische (vernette) polyurethanen.
In hoofdstuk 2 worden dergelijke rubberachtige polyurethaan netwerken op
basis van ethyl 2,6-diisocyanaathexanoaat ("lysine diisocyanaat")
beschreven. Deze polyurethaan netwerken, die zodanig ontworpen zijn dat
bij degradatie uitsluitend verbindingen worden gevormd die niet toxisch
zijn, worden verkregen door hexafunctionele stervormige prepolyrneren te
vernetten met het diisocyanaat op basis van lysine, een aminozuur. De
prepolyrneren op hun beurt worden gesynthetiseerd via de
ringopeningscopolymerisatie van hetzij L-lactide of glycolide, en
c-caprolacton, geYnitieerd door myo-inositol, een vitamine. De zo gevormde
biodegradeerbare polyesterurethaan netwerken zijn rubberachtig (T 's in de 'J
orde van 0-10 OC). Opmerkelijk is dat de rubber netwerken na extractie met
chloroform (de netwerken hebben gel percentages in de orde van ca. 90-95
% I zeer goede mechanische eigenschappen bezitten in vergelijking met de
niet geextraheerde netwerken, hetgeen wordt toegeschreven aan
"strain-induced crystallization". Na extractie van de sol fractie, die als
weekmaker fungeert en daardoor strain-induced crystallization verhindert,
stijgt de treksterkte van deze rubber netwerken van ca. 10 MPa naar 30-40
MPa. In hoofdstuk 3 wordt de biomedische toepasbaarheid van de
polyurethaan netwerken zoals beschreven in het voorgaande hoofdstuk
geevalueerd, en dan met name de toepasbaarheid van deze materialen als
biodegradeerbare, poreuze onderlaag in een meerlagen kunsthuid. Een amorf,
rubberachtig, poreus, poly(glyco1ide-co-E-capro1acton)urethaan netwerk op
basis van lysine diisocyanaat bleek in vitro snel te degraderen. Hetzelfde
rnateriaal bleek in vivo (subcutane implantaties bi j cavia' s) nog sneller
afgebroken te worden. De poreuze implantaten waren na 4-8 weken bijna
volledig gedegradeerd. Ondertussen was in de poreuze materialen duidelijke
celingroei waarneembaar. Het polyurethaan op basis van lysine diisocyanaat
bleek geen ongunstige weefsel/ontstekingsreactie te veroorzaken en zou als
biocompatibel bestempeld mogen worden.
Hoofdstuk 4 beschrijft de vervaardiging en eigenschappen van een poreuze
polyurethaan wondbedekking. Door middel van een fase inversie proces werd
een zeer dunne, poreuze membraan (15-20 pm) verkregen. Deze elastische
film, gemaakt van een cycloalifatisch polyetherurethaan (Tecoflex), bevat
een groot aantal microgaatjes ter grootte van ca. 5 pm. De poreuze
polyurethaan wondbedekking bleek ondoorlaatbaar voor bacterien (en voor
water in vloeibare vorm, b.v. wondvocht), maar zeer permeabel voor
waterdamp. Bij cavia's bleek dat de epidermale wondgenezing van 0.3 mm
diepe schaafwonden bedekt met de polyurethaan wondbedekking sneller
verliep dan die van onbedekte wonden en wonden bedekt met een occlusieve
wondbedekking (Op-Site). Deze stimulerende werking op het
wondgenezingsproces komt waarschijnlijk door de hoge
waterdampdoorlaatbaarheid van de poreuze polyurethaan film die het
wondvocht concentreert tot een soort gelachtige substantie. De
belangrijkste conclusie van het klinisch gebruik op 20 split-skin
donorplaatsen was dat de polyurethaan wondbedekking de pijn duidelijk
vermindert, in vergelijking met een conventionele behandeling met
ingezwachteld paraffine gaas (tulle gras). De polyurethaan wondbedekking
zorgde ervoor dat wondvochtretentie werd voorkomen, maar de epithelisatie
niet meer werd versneld dan bij gebruik van paraffine gaas.
In hoofdstuk 5 wordt een tweelagen biodegradeerbare kunstader, gemaakt van
polyurethaan en poly(L-melkzuur) beschreven. De microporeuze binnenlaag
van de kleine-diameter vaatprothese werd vervaardigd van een
cycloalifatisch gesegmenteerd polyurethaan (Tecoflex), dat vernet werd met
behulp van dicumylperoxide in de aanwezigheid van linolzuur. Op deze wijze
werd de kruipweerstand van het polyurethaan sterk verbeterd. Bovendien
zorgden de carbonzuur groepen van het linolzuur voor een niet trombogeen
binnenoppervlak van de vaatprothese. Ten derde zorgt de toevoeging van
linolzuur bij het vernetten met peroxides dat de sterkte als gevolg van de
vernetting niet afneemt. De macroporeuze buitenlaag van de prothese werd
gemaakt van een snel uit oplossing neergeslagen polyurethaardPLLA (95/5)
fysisch mengsel. De tweelagen kunstaders bleken in vivo (vervanging van 1
cm van de buikslagader in ratten) goed te functioneren. Alle protheses
waren na 1 jaar nog functioneel en lieten geen aneurysma' s (catastrofale
verwijding van de ader) zien, hetgeen toegeschreven dient te worden aan de
sterk verbeterde weerstand tegen kruip van de protheses als gevolg van de
chemische vernetting. De essentiele componenten van een natuurlijke
vaatwand bleken te zijn geregenereerd: de binnenlaag van de protheses was
volledig bedekt met endotheelcellen en enkele lagen gladde spiercellen.
Hoofdstuk 6 beschrijft de synthese en eigenschappen van zeer dicht
vernette polyurethanen, alsmede de potentiele toepassing van deze
materialen in de oogheelkunde, bijvoorbeeld als intraoculaire lens (ter
vervanging van de door staar troebel geworden natuurlijke ooglens) of als
keratoprothese (kunsthoornvlies). Deze sterk vernette polyurethanen werden
gesynthetiseerd door verschillende laag moleculaire polyolen te laten
reageren met (cyc1o)alifatische diisocyanaten in de afwezigheid van een
oplosmiddel. De zo verkregen polyurethaan netwerken ziJn optisch
transparante, kleurloze en arnorfe, glasachtige materialen. De
eigenschappen van een bepaald polyurethaan thermoharder, representatief
voor de hier gesynthetiseerde netwerken, verkregen door de stapreactie in
de bulk van een tetrafunctioneel aminoalcohol tetrakis(2-hydroxypropy1)-
ethyleendiamine (Quadrol) en hexamethyleendiisocyanaat (HDI) in
stoichiornetrische hoeveelheden, werden nader onderzocht. Bij volledige
uitharding heeft dit glasachtig polyurethaan netwerk een maxirnale T van g
85 OC, een zeer lage zwelgraad in chloroform (1,271 en goede mechanische
eigenschappen (treksterkte 80-85 MPa, rek biJ breuk ca. 15 %, Young's
modulus 1,5 GPa). Uit het infrarood spectrum bleek dat alle (toxische)
isocyanaat groepen, bij volledige uitharding, weggereageerd waren.
In tegenstelling tot polymethylmethacrylaat (PMMA), dat de laatste 15 jaar
met veel succes als intraoculaire lens materiaal gebruikt is, kunnen deze
vernette polyurethanen gesteriliseerd worden door autoclavering. De
mogelijkheid om een kunstlens te kunnen autoclaveren is met name
interessant voor de oogchirurgie in ontwikkelingslanden, waar de
meerderheid van alle blinde mensen leven. Implantaties van het
polyurethaan netwerk op basis van Quadrol en HDI in konijneogen lieten
zien dat het materiaal goed getolereerd werd. Een jaar na de implantatie
van een polyurethaan keratoprothese was het oog nog steeds rustig. De
resultaten van de polyurethaan keratoprothese waren vergelijkbaar met
implantaties van PMMA of metaal/glas keratoprotheses. A1 deze resultaten
laten zien dat het dicht vernette polyurethaan netwerk geschikt lijkt voor
toepassing in de oogheelkunde.