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Degradation of poly-L-lactide. Part 1: in vitro and
in vivo physiological temperature degradation
N A Weir1, F J Buchanan1*, J F Orr1, and G R Dickson2
1School of Mechanical and Manufacturing Engineering, Queen’s University Belfast, Belfast, UK
2Department of Trauma and Orthopaedic Surgery, Queen’s University Belfast, Belfast, UK
Abstract: Poly-L-lactide (PLLA) is one of the most significant members of a group of polymers
regarded as bioresorbable. The degradation of PLLA proceeds through hydrolysis of the ester linkage
in the polymer’s backbone and is influenced by the polymer’s initial molecular weight and degree of
crystallinity. To evaluate its degradation PLLA pellets were processed by compression moulding into
tensile test specimens and by extrusion into 2 mm diameter lengths of rod, prior to being sterilized
by ethylene oxide gas (EtO) and degraded in both in vitro and in vivo environments. On retrieval at
predetermined time intervals, procedures were used to evaluate the material’s molecular weight,
crystallinity, mechanical strength, and thermal properties. Additionally, the in vivo host tissue’s bio-
logical response was analysed. The results from this study suggest that in both the in vitro and in
vivo environments, degradation proceeded at the same rate and followed the general sequence of
aliphatic polyester degradation, ruling out enzymes contributing and accelerating the degradation
rate in vivo. Additionally, the absence of cells marking an inflammatory response suggests that the
PLLA rods investigated in vivo were biocompatible throughout the 44 weeks duration of the study,
before any mass loss was observed.
Keywords: bioresorbable, poly-L-lactide, degradation, molecular weight, crystallinity
NOTATION 1 INTRODUCTION
Over the past 40 years polymer scientists, workingDSC differential-scanning calorimetry
closely with those in the medical device and clinical fields,EtO ethylene oxide gas
have made tremendous advances in understanding andGPC gel-permeation chromatography
applying bioresorbable polymers to more sophisticatedLM light microscopy
applications. Currently, applications for these promisingMn number average molecular weight
biomaterials cover a broad range of clinical and scientificm
0
initial mass
disciplines, including sutures and fracture-fixation devicesm
t
mass at time, t
to support tissue regeneration [1], as drug deliveryPDLA poly-D-lactide
systems in the pharmaceutical industry [2], and, withPDLLA poly-DL-lactide
recent scientific advances, as scaffolds in the field ofPGA polyglycolide
tissue engineering [3].PLA polylactide
Bioresorbable polymers belonging to the aliphaticPLLA poly-L-lactide
polyester family currently represent the most attractivePI polydispersity index
group of polymers that meet the various medical andSEM scanning electron microscopy
physical demands for safe clinical applications. ThisTEM transmission electron microscopy
is mainly due to their high level of biocompatibility,Tg glass transition temperature
acceptable degradation rates, and versatility regardingTm melting temperature
physical and chemical properties [4]. Undoubtedly, two
of the most significant members of the aliphatic poly-DH
melt
enthalpy of fusion
ester family are the poly(a-hydroxy acids), polyglycolide
The MS was received on 2 February 2004 and was accepted after revision (PGA), and polylactide (PLA) [5]. High molecular
for publication on 17 June 2004. weight PGA and PLA were first introduced in the 1950s
* Corresponding author: School of Mechanical and Manufacturing
[6] and [7]; however, they were initially discarded andEngineering, Queen’s University Belfast, Ashby Building, Stranmillis
Road, Belfast, BT9 5AH, UK. email: f.buchanan@qub.ac.uk research was abandoned into the polymerization of other
H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
308 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON
a-hydroxy acids because of their poor thermal and hydro- molecular weight, mechanical strength, crystallinity, and
mass change at predetermined time intervals throughoutlytic stabilities, which did not allow them to be used as
regular plastics for long-term industrial use [8]. Never- degradation, and is a follow up to a previous study that
examined the processing, annealing, and sterilization oftheless, this very instability has proven to be of immense
importance, enabling PGA, PLA, and other aliphatic PLLA [25]. Additionally, this study examines the in vivo
host tissue’s biological response, to determine the bio-polyesters to be developed as synthetic bioresorbable
polymers for medical use. compatibility of the processed material. Test techniques
utilized include tensile and shear testing for mechanicalPLA was first investigated for medical use by Kulkarni
et al. [9], who reasoned that it would be useful for strength, differential-scanning calorimetry (DSC) to deter-
mine thermal properties and crystallinity, gel-permeationsurgical implants since hydrolysis would yield lactic acid,
a normal intermediate of carbohydrate metabolism and a chromatography (GPC) to measure molecular weight,
and conventional histological analysis techniques forproduct of muscle contraction [10]. Due to the asymmetry
of the PLA molecule it exists as two optical isomers, examining any inflammatory response to the implants
in vivo. The aim of the study is to develop a completepoly-D-lactide (PDLA) and poly-L-lactide (PLLA),
with a synthetic blend of D-lactide and L-lactide yielding understanding of the processed PLLA’s degradation
profile, and to provide the necessary control data forpoly-DL-lactide (PDLLA) [11]. PDLLA is an amorphous
polymer, while PLLA, the naturally occurring isomer, investigation of methods to accelerate the degradation
of PLLA relative to its physiological degradation rate.is semicrystalline and the most common bioresorbable
polymer used for orthopaedic devices [12], due to its
relatively high tensile strength and low elongation. PLLA,
2 MATERIALS AND METHODSlike all aliphatic polyesters, degrades in vivo through
simple hydrolysis of the hydrolytically unstable ester
2.1 Materialslinkage in the polymer’s backbone, with the degradation
products ultimately metabolized to carbon dioxide and The polymer studied in this investigation, poly-L-lactide
water, and eliminated from the body [13]. (PLLA) ResomerA L (batch number 26033), was sup-
For semicrystalline aliphatic polyesters like PLLA plied in a sealed moisture-proof container by Boehringer
there is a general concensus that degradation proceeds Ingelheim (Ingelheim, Germany) in pellet form.
via random bulk hydrolysis in two distinct stages [14].
The first stage is characterized by the preferential attack
2.2 Methodsof the ester linkages in the more accessible amorphous
regions [15], with the second stage characterized by the 2.2.1 Processing
onset of mass loss and attack of the less accessible
The PLLA was processed by compression mouldingcrystalline regions [16–18]. In conjunction with this
into plates 0.8 mm thick and by extrusion into a 2 mmtwo-stage mechanism, work by Li et al. [19], examining
diameter rod using techniques detailed previously bythe degradation of PDLLA, showed conclusively that
Weir et al. [25]. ASTM D638-99 type-V tensile samplesdegradation occurs more rapidly in the centre than at
were then cut from the compression-moulded plates andthe surface. This is known as heterogeneous degradation
30 mm lengths cut from the extruded rod. The tensileand is widely accepted as occurring in both lactide and
and extruded rod samples were then annealed at 120 °Cglycolide polymers [20] and [21]. This heterogeneous or
for a period of four hours in a preheated air-circulatingautocatalytic degradation mechanism results from the
oven (see Fig. 1) prior to being sterilized using ethylenehydrolytic cleavage of the ester bonds forming new acidic
oxide gas (EtO) by Griffith Microscience (Derbyshire,carboxyl end groups. As degradation proceeds the soluble
UK) on their standard EtO cycle for medical polymers,oligomers produced close to the surface can escape, while
i.e. ‘Cycle 33’ [25].those in the centre cannot diffuse out of the polymer.
This results in a higher internal acidity, with the carboxyl
2.2.2 In vitro degradation
end groups catalysing the ester hydrolysis reaction,
and a differentiation between the surface and interior The initial mass, m
0
, of each of the tensile and extruded
rod samples were recorded. Individual specimens weredegradation rates. The degradation rate of PLLA, and
aliphatic polyesters in general, is strongly related to their then placed in 28 ml screw-top glass bottles, fully
immersed in a pH 7.4 phosphate-buffered solution, inmaterial properties with crystallinity, molecular weight
and distribution, orientation, unreacted monomer, and accordance with ISO 15814:1999, and placed in an air-
circulating oven, maintaining the temperature at a con-the presence of impurities all playing significant roles
[1]. This is illustrated with reports on the time for the stant 37 °C. The pH of the solutions was periodically
monitored throughout the duration of each of the studies,complete resorption of PLLA ranging from 80 weeks
[22] to 5.7 years [23] and [24]. although in each case the ratio of the phosphate-buffered
solution in millimetres to the polymer’s mass in gramsThis study investigates the in vitro and in vivo
degradation of PLLA by evaluating its properties of was greater than 30:1.
H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
309DEGRADATION OF POLY-L-LACTIDE. PART 1
Eight PLLA extruded rod samples were removed at strengths of the retrieved extruded rod samples were
measured as an indication of mechanical strength. Theeach follow-up time (five mechanical test samples and
three mass-change samples), while six PLLA tensile shear test employed was adapted from BS 2782: Part 3:
Method 340B: 1978, Determination of Shear Strength ofsamples were removed at each follow-up time (three
mechanical test samples and three mass change samples). Sheet Material, and was similar to the method employed
by Suuronen et al. [26]. The retrieved PLLA rod samplesFollow-up times for the in vitro studies are given in
Table 1. were slotted into a hole aligned between two halves of
a shear bracket and sheared simultaneously at both ends
2.2.3 In vivo degradation under a constant strain rate of 5 mm/min. Shear strength
in MPa was calculated from equation (1)The Sprague Dawley rat was chosen as the animal model,
with the 30 mm lengths of 2 mm diameter extruded shear strength (MPa)=F/2A=F/2pr2 (1)
PLLA rod solely investigated. The samples were pre-
where F=load at maximum (N) and r=average radius
pared similarly to the rods for the in vitro studies, how-
of sample (mm). The overall shear strength was divided
ever, before sterilization the ends of each of the rods were
by two in equation (1) to account for the double shearing
blunted by grinding on a Struers (Rødovre, Denmark)
action taking place.
rotating grinding machine using Struers silicon carbide
In accordance with ISO 15814:1999, the retrieved
paper (grit 1200) and the initial mass (m
0
) of each of
material underwent mechanical testing while ‘wet’, with
the samples was recorded. Three PLLA rod samples were
testing conducted within three hours of retrieval from
implanted per rat, one each for mechanical testing, mass
both the in vitro buffered solution and in vivo animal
loss, and histological analysis to examine the host tissue’s
model.
response. A total of 12 male-weight-matched (#350 g)
Sprague Dawley rats were used for the in vivo studies, Mass change. On retrieval, the in vitro and in vivo
with the material implanted subcutaneously in the rats’ samples were dried immediately with a paper towel to
dorsum, with the three samples placed about the dorsum remove any surface moisture before being weighed using
midline. an electronic balance (Mettler Toledo, Fisher Scientific,
At predetermined time intervals (Table 1), the PLLA UK) to determine the percentage swelling of the polymer
rod samples intended for mechanical strength and mass- and water uptake. The samples were then dried in a
change analysis were surgically removed from the rats’ vacuum oven (Townson+Mercer, Altrincham, UK) at
dorsal subcutaneous tissue and separated from any approximately 30 °C for 48 hours at a vacuum of 0.68 bar
adhering tissue. The rod samples intended for histological and reweighed to obtain their mass at time, t(m
t
). The
analysis were also removed, together with surrounding overall percentage mass change after drying was then
adherent tissue, and placed in a fixative solution and calculated from equation (2)
stored in a refrigerator at 4 °C to preserve the tissue until
it was ready to be processed for light microscopy (LM),
percentage mass change=
m
t
−m
0
m
0
×100 per cent
scanning electron microscopy (SEM) and transmission
electron microscopy (TEM). Three rats were sacrificed
(2)at each time interval by administering an increasing dose
of CO
2
gas. All experimental procedures were carried
Molecular weight and thermal properties. Following
out under approved Home Office Project and Personal
mass change measurements the dried PLLA samples
Licence cover.
were reused for gel-permeation chromatography (GPC)
analysis to determine their weight and number average2.2.4 Characterization of retrieved in vitro and in vivo
molecular weights (Mr) throughout degradation and alsomaterial
for differential-scanning calorimetry (DSC) to determine
Mechanical properties. The mechanical properties of their thermal properties and percentage crystallinity.
the PLLA tensile samples were determined using a
JJ Lloyd EZ 50 tensile testing machine (Hampshire, Molecular weight. The GPC analysis was conducted
by Rapra Technology Ltd (Shropshire, UK). SamplesUK), equipped with a 1 kN load cell and tested at a
constant strain rate of 10 mm/min. Young’s modulus, were prepared by adding 10 ml of choloform solvent to
20 mg of sample taken through a cross-section of thetensile strength, and extension at break were calculated
from each of the load versus extension curves. The shear material. A Pl gel-mixed bed column with a refractive
Table 1 In vitro and in vivo follow-up times
Follow up Weeks
In vitro 4 10 20 26 32 38 44 50 57 65
In vivo 10 26 38 44
H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
310 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON
index response detector was used. The GPC system was at a heating rate of 10 °C/min., providing measurements
of glass transition temperature, Tg, melting point, Tm, andcalibrated with polystyrene and all results were expressed
as ‘polystyrene equivalent’ molecular weights. It should enthalpy of fusion, DH
melt
in J/g. The DSC results were
derived from this single heating cycle to provide a truebe noted that this is a relative technique, rather than an
absolute technique, for determining molecular weight. indication of changes in the polymers’ thermal properties
and morphology as a direct result of degradation. TheMolecular weight is considered to be the most
important and sensitive parameter for modelling bio- enthalpy of fusion, DH
melt
was then used to calculate
the polymers’ percentage crystallinities relative to theresorbable polymer degradation [27–29]. In particular,
with Mn directly related to the scission of the polymers enthalpy of fusion of a 100 per cent crystalline sample
of PLLA, reported to be 93 J/g [32]chains, a number of relationships have been derived
relating the changes in Mn with time to the hydrolysis
percentage crystallinity=(DH
melt
/93)×100 per cent
rate of the unstable ester linkages. Anderson [30] and
Chu [17] reported a statistical method for relating (7)
molecular weight to hydrolysis rate; assuming that the
extent of degradation was not large, they reported the Investigation of the biological host tissue’s response. The
following kinetic relationship based on the polymers Mn retrieved PLLA extruded rod samples from the in vivo
animal model were prepared and investigated using LM,1/Mn
t
=1/Mn
0
+kt (3)
TEM, and SEM, to observe structure, ultrastructure, and
where Mn
t
=Mn at time t; Mn
0
=Mn at t=0; k=rate surface topography of the tissue-implant environment.
constant; and t=time. If the theory holds true, a linear Conventional techniques were used in the preparation
relationship should exist between 1/Mn versus time, up of each of the samples [33] and [34]. Specimens for
until the point of mass loss. LM were fixed in 10 per cent buffered formaldehyde,
However, a disadvantage of this statistical approach while those for SEM and TEM were preserved in 3 per
is that it does not account for the possibility of auto- cent glutaraldehyde in 0.1 M sodium cacodylate buffer,
catalysis accelerating the polymer’s degradation rate. Pitt pH 7.2–7.4. The local host biological tissue’s response
and Gu [31] derived a relationship based on the kinetics to the implant was analysed after 36 and 44 weeks.
of the ester-hydrolysis reaction, accounting for auto-
catalysis by the generated carboxylic acid end groups,
described by the rate equation 3 RESULTS
d(E)/dt=−d(COOH)/dt=−k(COOH)(H
2
O)(E)
3.1 Visual examination
(4)
Initially, at 0 weeks the annealed PLLA tensile and
where (COOH), (H
2
O), and (E) represent the con- extruded rod specimens were opaque and off-white
centrations of carboxyl end groups, water, and esters in colour (Fig. 1). At 32 weeks small areas of the
respectively. tensile specimens became more intensely white and as
On further analysis of equation (4) and assuming that degradation time increased more white areas became
the ester and water concentrations remain constant and visible (Fig. 2). Both the in vitro and in vivo extruded
the concentration of acid end groups is equal to 1/Mn, rod specimens remained opaque and off-white in colour
it can be shown that throughout the duration of the study (Fig. 1), with no
whiter areas becoming visible. The sizes and shapes ofMn
t
=Mn
0
e−kt (5)
both the tensile and extruded rod specimens did not
If this relationship holds true, a linear relationship change during ageing.
should exist between the ln Mn versus time up until the
point of mass loss
ln Mn
t
=−kt+ln Mn
0
(6)
An advantage of using GPC to measure molecular
weight is that it also provides information on the
samples’ molecular weight distribution, through con-
sideration of the whole GPC chromatogram, providing
further insights into the complex nature of bioresorbable
polymer degradation.
Thermal properties. The thermal properties of the dried
retrieved PLLA samples were analysed using a Perkin
Elmer DSC 6 (Beaconsfield, Buckinghamshire, UK) test-
Fig. 1 PLLA tensile and extruded rod test samplesing machine, over a temperature range of 40 °C to 200 °C
H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
311DEGRADATION OF POLY-L-LACTIDE. PART 1
Table 2 In vitro and in vivo molecular weight results versus time results (PI=polydispersity index)
Tensile samples Extruded rod samples
Degradation In vitro In vitro In vivo
time
(weeks) Mw Mn PI Mw Mn PI Mw Mn PI
0 424 000 158 500 2.67 339 000 155 000 2.18 399 000 155 000 2.57
4 339 000 143 000 2.37 312 500 129 000 2.42
10 309 000 120 000 2.58 249 000 85 250 2.9 226 500 87 750 2.56
20 199 000 72 500 2.74 176 000 82 900 2.12
26 199 000 93 850 2.11 149 000 70 150 2.13 153 500 70 050 2.23
32 159 000 65 800 2.4 127 000 60 800 2.08
38 133 500 53 050 2.51 89 100 40 850 2.18 105 500 54 250 1.94
44 74 900 22 500 3.33 76 400 37 850 2.02 69 800 32 600 2.15
Table 3 Thermal properties of extruded PLLA rod in vitro
Degradation time
(weeks) % crystallinity Tm, °C Tg onset, °C
0 40.7 184.2 68.9
10 45.5 184.6 70.7
26 44.9 184.8 67.2
32 40.8 183.8 65.5
38 53.4 183.6 63.6
44 53.6 182.4 58.7
Considering the data presented in Table 2 in con-
junction with equations (3) and (6), Fig. 4 presents
plots of 1/Mn and ln Mn versus time, modelling the
uncatalysed and autocatalysed degradation models, with
linear trendlines fitted and R2 correlation coefficients
displayed. In both cases a higher correlation coefficient
was achieved for the ln Mn versus time relationship,Fig. 2 Degraded PLLA tensile samples at 37 °C in vitro
describing the autocatalytic degradation mechanism.
The GPC chromatograms for the PLLA tensile3.2 Molecular weight versus time
samples remained monomodal throughout successive
The molecular weights of the PLLA tensile and extruded weeks of degradation (Fig. 5) and shifted towards lower
rod samples in vitro and in vivo decreased with time molecular weights as degradation time increased. A
(Table 2). After 44 weeks the Mw of the extruded rod similar trend was observed for the extruded rod samples
samples decreased by approximately 80 per cent, while in vitro and in vivo.
the Mw for the tensile samples in vitro and extruded rod
in vivo decreased by approximately 82 per cent. A similar
decreasing trend was also observed for the extruded rod 3.3 Mass change versus time
and tensile samples’ Mn in vitro and in vivo (Fig. 3).
Before drying, a similar pattern for the percentage massHowever, no obvious pattern in polydispersity index
change of both the tensile and extruded rod samples(PI) (Table 3), the ratio of Mw/Mn, with time could be
in vitro was observed (Fig. 6). After four weeks the massderived from the molecular weight data.
of both sets of samples had increased by approximately
0.6 per cent, this increase remained relatively constant
up until about week 44, when the polymer’s mass
increased to approximately 1 per cent. Further increases
in mass before drying were observed throughout sub-
sequent weeks up to approximately 2.5 per cent for both
the tensile and extruded rod samples at 65 weeks. After
drying, a similar pattern was again observed for both
sets of samples, with minimal mass loss observed at 57
weeks for both the tensile and extruded rod samples.
However, at 65 weeks the tensile samples’ mass had
Fig. 3 Comparison between in vitro and in vivo samples’ Mn decreased by approximately 1.2 per cent with a 0.9 per
H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
312 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON
Fig. 4 Uncatalysed and autocatalysed models for tensile and extruded rod samples
cent mass loss observed for the extruded rod samples. A
similar pattern was also observed for the extruded rod
samples mass change in vivo, however, the increases in
mass before drying were not as large, peaking at approxi-
mately 0.8 per cent after 44 weeks, with no significant
mass loss observed after drying.
3.4 DSC analysis versus time
For both the tensile and extruded rod samples in vitro,
a general trend was observed of increasing crystallinity
and decreasing Tg onset temperature with degradation
time (Tables 3 and 4). Additionally, a slight but signifi-Fig. 5 PLLA tensile samples molecular weight distributions
cant decrease in both the tensile and extruded rodat 0, 10, 32, and 44 weeks in vitro
samples melting point, Tm, was also observed after 44
weeks. The results of the thermal analysis conducted on
Table 4 Thermal properties of PLLA tensile samples in vitro
Degradation time
(weeks) % crystallinity Tm, °C Tg onset, °C
0 44.8 182.1 66.9
4 45.5 181.4 68.8
10 41.6 181.6 67.7
20 42.1 182.1 66.8
26 47.0 181.4 65.7
32 51.9 180.3 61.9
38 54.8 181.0 63.2
44 58.7 179.1 57.3
Fig. 6 In vitro mass change analysis
H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
313DEGRADATION OF POLY-L-LACTIDE. PART 1
the extruded rod in vivo also followed this general trend, time, with degradation proceeding further, the main
melting peak began to shift to lower temperatures, untilwith the in vitro and in vivo extruded rod results proving
to be very similar. eventually the two peaks merged and the smaller peak
appeared as a shoulder on the larger peak (Fig. 7(f )).The DSC thermograms for the tensile samples at 0, 4,
10, 20, 38, and 44 weeks are shown in Fig. 7. At 0 weeks
(Fig. 7(a)), a small endothermic peak commencing at
approximately 67 °C was observed, relating to stress 3.5 Mechanical strength versus time
relaxation at the polymer’s Tg [35]. As the temperature
The mechanical strengths of both the tensile and
increased further a small endothermic dip was observed
extruded rod samples deteriorated with time, with the
just before melting commenced, followed by the main
tensile strength of the compression-moulded samples
melting peak. It is suggested that the dip before melting
reduced to approximately zero in 44 weeks (Table 5).
was caused by some crystallization of the polymer.
After 44 weeks the samples were very brittle and could
Although the polymer was annealed prior to degradation,
not be gripped in the tensile test grips without fracturing.
with the aim of limiting crystallization throughout the
A similar trend was observed initially for the loss of
study, close to the polymer’s melting point the chain
shear strength for the extruded rods in vitro and in vivo
mobility would have increased, allowing some of the
(Fig. 8). However, after 44 weeks in vitro, the rod
amorphous segments to order themselves into a more
samples had lost approximately 52 per cent of their
crystalline structure.
original strength, compared to only 26.3 per cent for the
At ten weeks (Fig. 7(c)), as degradation increased,
samples in vivo.
presumably in the amorphous regions, the initial dip
before melting observed at 0 and 4 weeks was reduced,
with less amorphous regions remaining capable of
3.6 Biological host tissue’s response
crystallization.
At 20 weeks (Fig. 7(d)) the endothermic dip before The combination of LM, TEM, and SEM proved useful
in determining the relationship between the PLLA rodmelting had disappeared and a small peak appeared to
form in its place. It is speculated that this new peak implants and surrounding biological tissue. After 36
weeks the PLLA implant appeared to stimulate the pro-represented the melting of new crystallites formed by the
crystallization of internal degradation by-products. The duction of a fibrous tissue capsule (Figs 9(a) and (b))
in which type-1 collagen fibre production was extensivereduction of the amorphous regions and crystallization
of the degradation by-products resulted in the polymer’s (Fig. 9(c)). The TEM image of the fibrous capsule
surrounding the PLLA implant showed the presence ofoverall crystallinity increasing throughout degradation
(Tables 3 and 4).
At weeks 38 and 44 (Figs 7(e) and (f )), the newly Table 5 Deterioration in PLLA’s tensile properties
formed peak appeared to have grown and moved to a throughout degradation
higher temperature, evidence that the newly formed
Degradation Young’s Tensile Extensioncrystallites’ size may have been increasing. At the same
time modulus strength at break
(weeks) (MPa) (MPa) (mm)
0 668.4 64.3 1.6
4 618.4 53.8 1.6
10 625.2 60.3 1.5
20 433.0 23.7 0.8
26 528.7 36.9 1.0
32 226.5 9.9 0.4
38 284.2 8.2 0.3
44 – 1.0 0.4
Fig. 7 PLLA tensile samples’ DSC thermograms at 0, 4, 10, Fig. 8 Shear strength comparison of PLLA rod versus time
in vitro and in vivo20, 38, and 44 weeks
H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
314 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON
the absence of cells marking an inflammatory response
at 36 and 44 weeks would suggest that the PLLA rod
investigated was biocompatible throughout the 44 weeks
duration of the study.
4 DISCUSSION
4.1 Degradation mechanisms
The results presented show that the in vitro and in vivo
degradation of PLLA commences almost immediately,
with the in vitro tensile and extruded rod samples losing
approximately 20 per cent of their initial molecular
weight at four weeks, the first time point analysed.
Autocatalysis. The higher R2 correlation coefficients
(Fig. 4) achieved for the ln Mn versus time relationship
shows a closer approximation of the autocatalysed model
(equation (6)) to the experimental data, compared to
the uncatalysed model, with degradation accelerated by
the newly formed carboxylic acid end groups generated
by the continual ester hydrolysis reaction. However, it
cannot be concluded that the mechanism is exclusively
autocatalytic. Investigating the in vitro degradation of
SR-PLLA at 37 °C, Pohjonen and To¨rma¨la¨ [36] observed
a similar trend. They reported correlation coefficients of
0.989 for the autocatalysed model and 0.910 for the
uncatalysed model, confirming the findings of the present
study and at least, according to theory, suggesting an
autocatalytic degradation mechanism. In contrast, in a
comparative study investigating the molecular weight
versus time data available in literature for semicrystalline
aliphatic polyesters such as PLLA and amorphous poly-
mers such as PDLLA, Anderson [30] reported that no
clear distinction could be derived between the uncatalysed
1/Mn and autocatalysed ln Mn plots versus time for semi-
crystalline polymers. However, for amorphous polymers,
the results were reported to be much more consistent
with an autocatalytic mechanism, with higher correlation
coefficients achieved for plots of ln Mn versus time.
Although, the correlation coefficients for each study were
not given, making comparisons to the present study
difficult. As a result of these studies Anderson [30] con-
cluded that the hydrolytic degradation of semicrystalline
polyesters may not proceed exclusively by non-catalytic
or autocatalytic mechanisms, speculating that both may
contribute to the rate of chain scission.
Relationship between molecular weight distribution and
degradation. Considering the GPC curves for the tensileFig. 9 Images of PLLA biological tissue after 36 weeks
implantation samples (Fig. 5) it is interesting to note that they remained
monomodal throughout successive weeks of degradation.
In contrast, many researchers have reported that asfibroblasts (Fig. 9(c)) in an extracellular matrix (inset)
composed extensively of type-1 collagen fibre bundles the degradation of PLA and PGA aliphatic polyesters
proceeds the initially monomodal GPC curve becomesrunning in different orientations. While tissue disturbance
during surgery produces an initial inflammatory response, bimodal and even multimodal in nature [14] and [37].
H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
315DEGRADATION OF POLY-L-LACTIDE. PART 1
The bimodal nature of these GPC curves was originally butions were the result of preferential degradation of the
amorphous regions, a view supported by Fischer et al.assigned to the difference in degradation rates in the
[32]. It would be anticipated that due to the semicrystallineamorphous and crystalline regions [14], [32] and [37].
nature of the compression-moulded and extruded PLLAHowever, with molecular weight usually determined by
investigated in the present study [25], the GPC curvestaking samples from the bulk of the polymer, comprising
of both would become bimodal and even multimodalthe interior of a lower molecular weight than the surface
in nature due to the preferential degradation of the[37], Li et al. [19] were the first to assign this bimodal
amorphous regions. However, Fig. 5 shows that this isbehaviour to the autocatalytic effect and faster internal
not the case, with the molecular weight distributionsdegradation. It is currently understood that for PLA and
remaining monomodal throughout the 44 weeks durationPGA polymers and their copolymers, this bimodal nature
of the study, contradicting the findings of Li et al. [20]can be accounted for by three different mechanisms
and Pistner et al. [39]. It must also be assumed that withrelated to the polymers morphology [14]. First, by faster
the samples intended for molecular weight analysis takeninternal degradation; however, this mechanism is most
through a cross-section of the material, the suspectedcommonly observed for initially amorphous polymers,
autocatalytic mechanism did not result in a large enoughwhich are not believed to be capable of crystallization
surface-interior differentiation to yield curves containingeven throughout degradation [37], for example, a 50:50
two distinct molecular weight species.copolymer of PLA and PDLA [19]. Second, for semi-
crystalline polymers the bimodal nature has been attri-
buted to selective degradation of the amorphous regions, Bulk degradation. The time delay before mass loss
with the surface-interior differentiation reported not to observed in this study (Fig. 6) is in agreement with
be large enough to yield bimodal GPC chromatograms the reported general sequence of aliphatic polyester
[14], supporting Anderson’s theory [30] that the hydro- degradation, which suggests molecular weight loss is
lytic degradation of semicrystalline polyesters does not observed first, before loss of mechanical strength and
proceed exclusively by non-catalytic or autocatalytic before any physical mass loss is observed [8]. This is
mechanisms. Finally, the bimodal nature of the GPC accounted for by the fact that water diffusion into the
chromatograms has been attributed to the crystallization polymer is faster than the hydrolytic degradation of
of low molecular weight degradation by-products in the polymer’s ester linkage, suggesting that ester-bond
initially amorphous polymers, for example, amorphous cleavage is the rate-limiting step in the degradation of
PLLA and a 75:25 PLA/PGA copolymer [38], which are aliphatic polyesters [40]. This results in degradation pro-
capable of crystallizing throughout degradation. Once ceeding in the bulk of the polymer, resulting in a time-
lag before any mass loss is observed, as the polymer’sthe low molecular by-products crystallize they become
molecular weight has to be reduced to a critical valueresistant to degradation and appear as a low molecular
before soluble oligomers can be released. In contrast, forweight peak on the GPC curve.
bioresorbable polymers regarded as surface eroding,The monomodal nature of the GPC curves obtained
such as those belonging to the polyanhydride and poly-for the semicrystalline PLLA prepared by annealing
orthoester families [41], mass loss is observed almostand investigated in this study appears to contradict
immediately, as the chain scission of their more reactivethe findings of other researchers investigating similar
unstable linkages, in comparison to the ester linkage insemicrystalline PLLA. Li et al. [20], investigating the
aliphatic polyesters, is faster than the diffusion of waterdegradation of semicrystalline PLLA prepared by anneal-
molecules into the polymer [40].ing at 130 °C for two hours, with an initial crystallinity
of 72 per cent deduced from XRD measurements,
observed that the initial monomodal molecular weight Polymer morphology and degradation. The results of the
distribution became multimodal after 18 weeks. After 50 DSC analysis (Fig. 7) appear to provide evidence that
weeks Li et al. [20] observed that the GPC curve became the low molecular weight degradation by-products are
bimodal, with the peak corresponding to high molecular capable of crystallizing, due to their greater mobility,
weight being more prominant for the surface than for and contribute to the samples’ increasing crystallinity.
the centre, suggesting autocatalysis. At 90 weeks the This is evident by the emergence of a small peak forming
GPC chromatogram then became almost monomodal and eventually merging with the main melting peak. The
and was composed of a single low molecular weight crystallization of these internal degradation by-products
peak. Pistner et al. [39] observed a similar profile for the resulted in the polymer maintaining its structural integrity
GPC chromatograms of semicrystalline PLLA with an throughout the duration of the study. In contrast, hollow
initial crystallinity of 73 per cent measured by DSC, with structures have been reported for intrinsically amorphous
a low molecular weight shoulder observed after eight polymers since their degradation products are not
weeks, becoming more important as degradation time believed to be capable of crystallizing, for example, in
proceeded. Both Li et al. [20] and Pistner et al. [39] the case of a 50:50 copolymer of PLLA/PDLLA [19].
The decreasing peak melting temperature, observed mostconcluded that the multimodal molecular weight distri-
H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
316 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON
significantly at 44 weeks (Tables 3 and 4) and deter- In vitro and in vivo degradation rates. Literature regard-
ing the role of enzymes on the degradation of aliphaticmined from a single heating cycle, is consistent with the
hypothesis that the initially crystalline regions are resistant polyesters is often contradictory. Many authors have
reported that enzymes may be involved in the latter stagesto degradation, resulting in a two-stage degradation mech-
anism with the amorphous regions being preferentially of degradation, when the polymer has fragmented and the
molecular weight is sufficiently small [47–49]. However,attacked [14]. However, once the amorphous regions have
been exhausted, the less accessible crystalline regions the role of enzymes during hydrolysis of the polymer
bulk remains unclear. In comparative in vitro and in vivoare then solely attacked and disrupted, resulting in a
decreased size of the initially present crystallites and studies, Vasenius et al. [50] have reported significantly
faster degradation of PGA rods in vivo, with Matsusuehence a reduced melting point [42]. Although the melt-
ing point of bioresorbable polymers is also known to be et al. [51] also reporting faster in vivo degradation
of PLLA. In each case the faster in vivo degradation ratedependent on molecular weight, the extent of this is most
readily determined by considering a reheat DSC run. was attributed, in some part, to the action of enzymes.
However, comparative studies by Hooper et al. [52] andSince the fusion of the first run destroys the polymer’s
initial crystalline structure, crystallization on cooling Pitt et al. [53] have reported no significant differences in
the degradation rates of poly(a-hydroxy acids) in vivo,involves the degraded chains only [36], confirming that
in the present study the decrease in melting temperature, with Cam et al. [54] reporting their degradation to
be practically independent of enzymes. The similaritydetermined from a single heating cycle, was most likely a
result of a reduction in crystallite size and not decreasing between the results of the molecular weight (Fig. 3) and
DSC analysis for the PLLA rods investigated in vitromolecular weight. It is speculated that the decreasing Tg
observed as degradation time increased (Tables 3 and 4) and in vivo in the present study suggests that the
degradation of PLLA is independent of enzymes, and inis related to the reduction in molecular weight of the
polymer’s chains in the amorphous regions, with a similar agreement with Timmins and Lenz [55], who reported
that enzymes capable of catalysing hydrolysis are them-trend also observed by Li [14], Duek et al. [43], Joukainen
et al. [44], and Kelloma¨ki et al. [45]. Interestingly, a selves macromolecules, unable to penetrate into the poly-
mer bulk. Therefore, any enzyme-contributed reactionsmall amount of water within a polymer is also known
to have a marked plasticizing effect, causing a reduction would be heterogeneous and confined to the surface of
the polymer, with a reduction in mass observed, but littlein the polymer’s Tg. A study by Siemann [46], investi-
gating the influence of water on the glass transition of change in the polymers overall molecular weight [4].
The significant loss of molecular weight (Table 2) andpoly(dl-lactic acid) by DSC, reported a 12 K decrease
in Tg after samples were exposed to water for six hours negligible mass loss (Fig. 6) observed for the PLLA rod
investigated in vivo in the present study would suggestprior to testing. However, in a further study, investi-
gating samples exposed to water and then dried to a that degradation proceeded predominantly in the bulk
of the polymer by non-enzymatic hydrolysis, similar toconstant mass before testing, the Tg remained the same
as the untreated samples. In the present study the the mechanism observed in vitro. However, this does
not rule out the influence of enzymes at later stages ofsamples were dried to constant mass before DSC testing
was conducted, ruling out water acting as a plasticizer, the degradation process, particularly when mass loss
becomes significant.and confirming a reduction in molecular weight as the
most probable cause for the decreasing trend in Tg.
However, this underlines the problem that for accurate
Tg measurements, representative of the polymer’s con-
4.2 Biological response
dition in service, test regimes need to be developed that
can accurately monitor the polymer’s Tg while the The production of a fibrous capsule around bio-
resorbable implants has been observed previously [56–58]samples remain ‘wet’.
and is regarded as part of the body’s natural response
to implants made of diverse materials [49]. SurgicalMechanical strength. Since degradation predominantly
occurred in the amorphous regions, disrupting the tie intervention, such as the implantation procedure under-
taken in this study, would initiate inflammation as achains holding the crystallites together, coupled with the
decreasing molecular weight and increasing crystallinity, response to injury. However, the absence of inflam-
matory cells at 36 and 44 weeks suggests that PLLAit is not surprising that the mechanical properties of the
PLLA investigated decreased so rapidly. However, con- is biocompatible throughout the early stages of its
degradation. It is understood that the onset of mass loss,sidering the similarities between the molecular weight
loss and the results derived from the DSC analysis for particularly in fast degrading aliphatic polyesters such
as PGA, can result in an inflammatory reaction due tothe extruded PLLA rod in vitro and in vivo, it is difficult
to speculate at this stage why the samples in vitro the sudden release of acidic degradation by-products,
causing a large change in pH of the surrounding mediaappeared to lose their strength more rapidly than those
in vivo. [59]. In the present investigation the PLLA degradation
H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
317DEGRADATION OF POLY-L-LACTIDE. PART 1
scaffolds and cells. In Synthetic Biodegradable Polymerstudy in vivo was terminated before any mass loss was
Scaffolds (Eds A. Atala and D. J. Mooney) 1997, pp. 1–14observed, although it is speculated that any inflam-
(Birkha¨user, Boston, MA, USA).matory response observed as a direct result of the onset
4 Li, S. and Vert, M. Biodegradation of aliphatic poly-of polymer mass loss would be mild. In comparison to
esters. In Degradable Polymers Principles & Applications
fast degrading PGA implants, it is anticipated that the
(Eds G. Scott and D. Gilead) 1995, pp. 43–87 (Chapman
release of acidic degradation products from the slower & Hall, London).
degrading PLLA would be less intense. This would 5 Chu, C. C. Biodegradable polymeric biomaterials: an
result in the surrounding tissue being more capable of overview. In The Biomedical Engineering Handbook
eliminating any such debris more efficiently, reducing the (Ed. J. D. Bronzino) 1995, pp. 611–626 (CRC Press, Boca
Raton, FL, USA).risk of a severe inflammatory reaction developing that
6 Higgins, N. A. Condensation of Polymers of Hydroxyaceticwould require further surgical intervention.
Acid. US Patent 2 676 945, 1954.
7 Schneider, A. K. Polymers of High Melting Lactide. US
Patent 2 703 316, 1955.5 CONCLUSIONS
8 Middleton, J. C. and Tipton, A. J. Synthetic biodegradable
polymers as orthopedic devices, Biomaterials, 2000, 21,
The results of the analytical characterization studies 2335–2346.
conducted on the retrieved PLLA samples in vitro and 9 Kulkarni, R. K., Pani, K. C., Neuman, C. and Leonard, F.
in vivo provides strong evidence to support the findings Polylactic acid for surgical implants, Arch. Surg., 1966,
of other researchers investigating similar bioresorbable 93, 839–843.
polymers. Additionally, the results from the in vivo 10 Hofmann, G. O. Biodegradable implants in orthopaedic
surgery—a review on the state-of-the art, Clin. Mater.,studies would suggest that throughout the first stage of
1992, 10, 75–80.degradation, before mass loss is observed, PLLA is bio-
11 Ciccone, W. J., Motz, C., Bentley, C. and Tasto, J. P.compatible and degrades at the same rate in vitro and
Bioabsorbable implants in orthopaedics: new developmentsin vivo. However, the results of the present studies do
and clinical applications. J. Am. Acad. Orthop. Surg., 2001,
appear to indicate that for semicrystalline polymers, like
9, 280–288.
the PLLA investigated, no clear differentiation between 12 Barber, F. A. Resorbable materials for arthroscopic
surface and interior degradation could be observed that fixation: a product guide. Orthopedic Special Edn, 2002,
would clearly point to an autocatalytic degradation 8, 29–37.
mechanism. As a result it is speculated that as poly- 13 Hayashi, T. Biodegradable polymers for biomedical uses,
Prog. Polym. Sci., 1994, 19, 663–702.mer crystallinity increases, the importance of the auto-
14 Li, S. Hydrolytic degradation characteristics of aliphaticcatalysis degradation mechanism may become less
polyesters derived from lactic and glycolic acids, J. Biomed.significant.
Mater. Res. (Appl. Biomater.), 1999, 48, 342–353.
15 Vert, M., Li, S. and Garreau, H. New insights on the
degradation of bioresorbable polymeric devices based on
ACKNOWLEDGEMENTS
lactic and glycolic acids, Clin. Mater., 1992, 10, 3–8.
16 Ali, S., Doherty, P. J. and Williams, D. F. Mechanisms
The authors would like to thank Mr David Farrar of polymer degradation in implantable devices. 2.
at Smith & Nephew Group Research Centre (York, Poly(DL-lactic acid), J. Biomed. Mater. Res., 1993, 27,
UK), Boehringer Ingelheim (Ingelheim, Germany) for 1409–1418.
17 Chu, C. C. Degradation and biocompatibility of syn-supplying the PLLA, Griffith Microscience (Derbyshire,
thetic absorbable suture materials: general biodegradationUK) for the ethylene oxide sterilization, and Rapra
phenomena and some factors affecting biodegradation.Technology Limited (Shropshire, UK) for the molecular
In Biomedical Applications of Synthetic Biodegradableweight characterization. Finally, the EPSRC (Swindon,
Polymers (Ed. J. O. Hollinger) 1995, pp. 103–128 (CRC
UK) for financial assistance.
Press, Boca Raton, FL, USA).
18 Mainil-Varlet, P., Curtis, R. and Gogolewski, S. Effect of
in vivo and in vitro degradation on molecular and mech-
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Degradation of poly-L-lactide. Part 1, IMechE, 2004

  • 1. 307 Degradation of poly-L-lactide. Part 1: in vitro and in vivo physiological temperature degradation N A Weir1, F J Buchanan1*, J F Orr1, and G R Dickson2 1School of Mechanical and Manufacturing Engineering, Queen’s University Belfast, Belfast, UK 2Department of Trauma and Orthopaedic Surgery, Queen’s University Belfast, Belfast, UK Abstract: Poly-L-lactide (PLLA) is one of the most significant members of a group of polymers regarded as bioresorbable. The degradation of PLLA proceeds through hydrolysis of the ester linkage in the polymer’s backbone and is influenced by the polymer’s initial molecular weight and degree of crystallinity. To evaluate its degradation PLLA pellets were processed by compression moulding into tensile test specimens and by extrusion into 2 mm diameter lengths of rod, prior to being sterilized by ethylene oxide gas (EtO) and degraded in both in vitro and in vivo environments. On retrieval at predetermined time intervals, procedures were used to evaluate the material’s molecular weight, crystallinity, mechanical strength, and thermal properties. Additionally, the in vivo host tissue’s bio- logical response was analysed. The results from this study suggest that in both the in vitro and in vivo environments, degradation proceeded at the same rate and followed the general sequence of aliphatic polyester degradation, ruling out enzymes contributing and accelerating the degradation rate in vivo. Additionally, the absence of cells marking an inflammatory response suggests that the PLLA rods investigated in vivo were biocompatible throughout the 44 weeks duration of the study, before any mass loss was observed. Keywords: bioresorbable, poly-L-lactide, degradation, molecular weight, crystallinity NOTATION 1 INTRODUCTION Over the past 40 years polymer scientists, workingDSC differential-scanning calorimetry closely with those in the medical device and clinical fields,EtO ethylene oxide gas have made tremendous advances in understanding andGPC gel-permeation chromatography applying bioresorbable polymers to more sophisticatedLM light microscopy applications. Currently, applications for these promisingMn number average molecular weight biomaterials cover a broad range of clinical and scientificm 0 initial mass disciplines, including sutures and fracture-fixation devicesm t mass at time, t to support tissue regeneration [1], as drug deliveryPDLA poly-D-lactide systems in the pharmaceutical industry [2], and, withPDLLA poly-DL-lactide recent scientific advances, as scaffolds in the field ofPGA polyglycolide tissue engineering [3].PLA polylactide Bioresorbable polymers belonging to the aliphaticPLLA poly-L-lactide polyester family currently represent the most attractivePI polydispersity index group of polymers that meet the various medical andSEM scanning electron microscopy physical demands for safe clinical applications. ThisTEM transmission electron microscopy is mainly due to their high level of biocompatibility,Tg glass transition temperature acceptable degradation rates, and versatility regardingTm melting temperature physical and chemical properties [4]. Undoubtedly, two of the most significant members of the aliphatic poly-DH melt enthalpy of fusion ester family are the poly(a-hydroxy acids), polyglycolide The MS was received on 2 February 2004 and was accepted after revision (PGA), and polylactide (PLA) [5]. High molecular for publication on 17 June 2004. weight PGA and PLA were first introduced in the 1950s * Corresponding author: School of Mechanical and Manufacturing [6] and [7]; however, they were initially discarded andEngineering, Queen’s University Belfast, Ashby Building, Stranmillis Road, Belfast, BT9 5AH, UK. email: f.buchanan@qub.ac.uk research was abandoned into the polymerization of other H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 2. 308 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON a-hydroxy acids because of their poor thermal and hydro- molecular weight, mechanical strength, crystallinity, and mass change at predetermined time intervals throughoutlytic stabilities, which did not allow them to be used as regular plastics for long-term industrial use [8]. Never- degradation, and is a follow up to a previous study that examined the processing, annealing, and sterilization oftheless, this very instability has proven to be of immense importance, enabling PGA, PLA, and other aliphatic PLLA [25]. Additionally, this study examines the in vivo host tissue’s biological response, to determine the bio-polyesters to be developed as synthetic bioresorbable polymers for medical use. compatibility of the processed material. Test techniques utilized include tensile and shear testing for mechanicalPLA was first investigated for medical use by Kulkarni et al. [9], who reasoned that it would be useful for strength, differential-scanning calorimetry (DSC) to deter- mine thermal properties and crystallinity, gel-permeationsurgical implants since hydrolysis would yield lactic acid, a normal intermediate of carbohydrate metabolism and a chromatography (GPC) to measure molecular weight, and conventional histological analysis techniques forproduct of muscle contraction [10]. Due to the asymmetry of the PLA molecule it exists as two optical isomers, examining any inflammatory response to the implants in vivo. The aim of the study is to develop a completepoly-D-lactide (PDLA) and poly-L-lactide (PLLA), with a synthetic blend of D-lactide and L-lactide yielding understanding of the processed PLLA’s degradation profile, and to provide the necessary control data forpoly-DL-lactide (PDLLA) [11]. PDLLA is an amorphous polymer, while PLLA, the naturally occurring isomer, investigation of methods to accelerate the degradation of PLLA relative to its physiological degradation rate.is semicrystalline and the most common bioresorbable polymer used for orthopaedic devices [12], due to its relatively high tensile strength and low elongation. PLLA, 2 MATERIALS AND METHODSlike all aliphatic polyesters, degrades in vivo through simple hydrolysis of the hydrolytically unstable ester 2.1 Materialslinkage in the polymer’s backbone, with the degradation products ultimately metabolized to carbon dioxide and The polymer studied in this investigation, poly-L-lactide water, and eliminated from the body [13]. (PLLA) ResomerA L (batch number 26033), was sup- For semicrystalline aliphatic polyesters like PLLA plied in a sealed moisture-proof container by Boehringer there is a general concensus that degradation proceeds Ingelheim (Ingelheim, Germany) in pellet form. via random bulk hydrolysis in two distinct stages [14]. The first stage is characterized by the preferential attack 2.2 Methodsof the ester linkages in the more accessible amorphous regions [15], with the second stage characterized by the 2.2.1 Processing onset of mass loss and attack of the less accessible The PLLA was processed by compression mouldingcrystalline regions [16–18]. In conjunction with this into plates 0.8 mm thick and by extrusion into a 2 mmtwo-stage mechanism, work by Li et al. [19], examining diameter rod using techniques detailed previously bythe degradation of PDLLA, showed conclusively that Weir et al. [25]. ASTM D638-99 type-V tensile samplesdegradation occurs more rapidly in the centre than at were then cut from the compression-moulded plates andthe surface. This is known as heterogeneous degradation 30 mm lengths cut from the extruded rod. The tensileand is widely accepted as occurring in both lactide and and extruded rod samples were then annealed at 120 °Cglycolide polymers [20] and [21]. This heterogeneous or for a period of four hours in a preheated air-circulatingautocatalytic degradation mechanism results from the oven (see Fig. 1) prior to being sterilized using ethylenehydrolytic cleavage of the ester bonds forming new acidic oxide gas (EtO) by Griffith Microscience (Derbyshire,carboxyl end groups. As degradation proceeds the soluble UK) on their standard EtO cycle for medical polymers,oligomers produced close to the surface can escape, while i.e. ‘Cycle 33’ [25].those in the centre cannot diffuse out of the polymer. This results in a higher internal acidity, with the carboxyl 2.2.2 In vitro degradation end groups catalysing the ester hydrolysis reaction, and a differentiation between the surface and interior The initial mass, m 0 , of each of the tensile and extruded rod samples were recorded. Individual specimens weredegradation rates. The degradation rate of PLLA, and aliphatic polyesters in general, is strongly related to their then placed in 28 ml screw-top glass bottles, fully immersed in a pH 7.4 phosphate-buffered solution, inmaterial properties with crystallinity, molecular weight and distribution, orientation, unreacted monomer, and accordance with ISO 15814:1999, and placed in an air- circulating oven, maintaining the temperature at a con-the presence of impurities all playing significant roles [1]. This is illustrated with reports on the time for the stant 37 °C. The pH of the solutions was periodically monitored throughout the duration of each of the studies,complete resorption of PLLA ranging from 80 weeks [22] to 5.7 years [23] and [24]. although in each case the ratio of the phosphate-buffered solution in millimetres to the polymer’s mass in gramsThis study investigates the in vitro and in vivo degradation of PLLA by evaluating its properties of was greater than 30:1. H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 3. 309DEGRADATION OF POLY-L-LACTIDE. PART 1 Eight PLLA extruded rod samples were removed at strengths of the retrieved extruded rod samples were measured as an indication of mechanical strength. Theeach follow-up time (five mechanical test samples and three mass-change samples), while six PLLA tensile shear test employed was adapted from BS 2782: Part 3: Method 340B: 1978, Determination of Shear Strength ofsamples were removed at each follow-up time (three mechanical test samples and three mass change samples). Sheet Material, and was similar to the method employed by Suuronen et al. [26]. The retrieved PLLA rod samplesFollow-up times for the in vitro studies are given in Table 1. were slotted into a hole aligned between two halves of a shear bracket and sheared simultaneously at both ends 2.2.3 In vivo degradation under a constant strain rate of 5 mm/min. Shear strength in MPa was calculated from equation (1)The Sprague Dawley rat was chosen as the animal model, with the 30 mm lengths of 2 mm diameter extruded shear strength (MPa)=F/2A=F/2pr2 (1) PLLA rod solely investigated. The samples were pre- where F=load at maximum (N) and r=average radius pared similarly to the rods for the in vitro studies, how- of sample (mm). The overall shear strength was divided ever, before sterilization the ends of each of the rods were by two in equation (1) to account for the double shearing blunted by grinding on a Struers (Rødovre, Denmark) action taking place. rotating grinding machine using Struers silicon carbide In accordance with ISO 15814:1999, the retrieved paper (grit 1200) and the initial mass (m 0 ) of each of material underwent mechanical testing while ‘wet’, with the samples was recorded. Three PLLA rod samples were testing conducted within three hours of retrieval from implanted per rat, one each for mechanical testing, mass both the in vitro buffered solution and in vivo animal loss, and histological analysis to examine the host tissue’s model. response. A total of 12 male-weight-matched (#350 g) Sprague Dawley rats were used for the in vivo studies, Mass change. On retrieval, the in vitro and in vivo with the material implanted subcutaneously in the rats’ samples were dried immediately with a paper towel to dorsum, with the three samples placed about the dorsum remove any surface moisture before being weighed using midline. an electronic balance (Mettler Toledo, Fisher Scientific, At predetermined time intervals (Table 1), the PLLA UK) to determine the percentage swelling of the polymer rod samples intended for mechanical strength and mass- and water uptake. The samples were then dried in a change analysis were surgically removed from the rats’ vacuum oven (Townson+Mercer, Altrincham, UK) at dorsal subcutaneous tissue and separated from any approximately 30 °C for 48 hours at a vacuum of 0.68 bar adhering tissue. The rod samples intended for histological and reweighed to obtain their mass at time, t(m t ). The analysis were also removed, together with surrounding overall percentage mass change after drying was then adherent tissue, and placed in a fixative solution and calculated from equation (2) stored in a refrigerator at 4 °C to preserve the tissue until it was ready to be processed for light microscopy (LM), percentage mass change= m t −m 0 m 0 ×100 per cent scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Three rats were sacrificed (2)at each time interval by administering an increasing dose of CO 2 gas. All experimental procedures were carried Molecular weight and thermal properties. Following out under approved Home Office Project and Personal mass change measurements the dried PLLA samples Licence cover. were reused for gel-permeation chromatography (GPC) analysis to determine their weight and number average2.2.4 Characterization of retrieved in vitro and in vivo molecular weights (Mr) throughout degradation and alsomaterial for differential-scanning calorimetry (DSC) to determine Mechanical properties. The mechanical properties of their thermal properties and percentage crystallinity. the PLLA tensile samples were determined using a JJ Lloyd EZ 50 tensile testing machine (Hampshire, Molecular weight. The GPC analysis was conducted by Rapra Technology Ltd (Shropshire, UK). SamplesUK), equipped with a 1 kN load cell and tested at a constant strain rate of 10 mm/min. Young’s modulus, were prepared by adding 10 ml of choloform solvent to 20 mg of sample taken through a cross-section of thetensile strength, and extension at break were calculated from each of the load versus extension curves. The shear material. A Pl gel-mixed bed column with a refractive Table 1 In vitro and in vivo follow-up times Follow up Weeks In vitro 4 10 20 26 32 38 44 50 57 65 In vivo 10 26 38 44 H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 4. 310 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON index response detector was used. The GPC system was at a heating rate of 10 °C/min., providing measurements of glass transition temperature, Tg, melting point, Tm, andcalibrated with polystyrene and all results were expressed as ‘polystyrene equivalent’ molecular weights. It should enthalpy of fusion, DH melt in J/g. The DSC results were derived from this single heating cycle to provide a truebe noted that this is a relative technique, rather than an absolute technique, for determining molecular weight. indication of changes in the polymers’ thermal properties and morphology as a direct result of degradation. TheMolecular weight is considered to be the most important and sensitive parameter for modelling bio- enthalpy of fusion, DH melt was then used to calculate the polymers’ percentage crystallinities relative to theresorbable polymer degradation [27–29]. In particular, with Mn directly related to the scission of the polymers enthalpy of fusion of a 100 per cent crystalline sample of PLLA, reported to be 93 J/g [32]chains, a number of relationships have been derived relating the changes in Mn with time to the hydrolysis percentage crystallinity=(DH melt /93)×100 per cent rate of the unstable ester linkages. Anderson [30] and Chu [17] reported a statistical method for relating (7) molecular weight to hydrolysis rate; assuming that the extent of degradation was not large, they reported the Investigation of the biological host tissue’s response. The following kinetic relationship based on the polymers Mn retrieved PLLA extruded rod samples from the in vivo animal model were prepared and investigated using LM,1/Mn t =1/Mn 0 +kt (3) TEM, and SEM, to observe structure, ultrastructure, and where Mn t =Mn at time t; Mn 0 =Mn at t=0; k=rate surface topography of the tissue-implant environment. constant; and t=time. If the theory holds true, a linear Conventional techniques were used in the preparation relationship should exist between 1/Mn versus time, up of each of the samples [33] and [34]. Specimens for until the point of mass loss. LM were fixed in 10 per cent buffered formaldehyde, However, a disadvantage of this statistical approach while those for SEM and TEM were preserved in 3 per is that it does not account for the possibility of auto- cent glutaraldehyde in 0.1 M sodium cacodylate buffer, catalysis accelerating the polymer’s degradation rate. Pitt pH 7.2–7.4. The local host biological tissue’s response and Gu [31] derived a relationship based on the kinetics to the implant was analysed after 36 and 44 weeks. of the ester-hydrolysis reaction, accounting for auto- catalysis by the generated carboxylic acid end groups, described by the rate equation 3 RESULTS d(E)/dt=−d(COOH)/dt=−k(COOH)(H 2 O)(E) 3.1 Visual examination (4) Initially, at 0 weeks the annealed PLLA tensile and where (COOH), (H 2 O), and (E) represent the con- extruded rod specimens were opaque and off-white centrations of carboxyl end groups, water, and esters in colour (Fig. 1). At 32 weeks small areas of the respectively. tensile specimens became more intensely white and as On further analysis of equation (4) and assuming that degradation time increased more white areas became the ester and water concentrations remain constant and visible (Fig. 2). Both the in vitro and in vivo extruded the concentration of acid end groups is equal to 1/Mn, rod specimens remained opaque and off-white in colour it can be shown that throughout the duration of the study (Fig. 1), with no whiter areas becoming visible. The sizes and shapes ofMn t =Mn 0 e−kt (5) both the tensile and extruded rod specimens did not If this relationship holds true, a linear relationship change during ageing. should exist between the ln Mn versus time up until the point of mass loss ln Mn t =−kt+ln Mn 0 (6) An advantage of using GPC to measure molecular weight is that it also provides information on the samples’ molecular weight distribution, through con- sideration of the whole GPC chromatogram, providing further insights into the complex nature of bioresorbable polymer degradation. Thermal properties. The thermal properties of the dried retrieved PLLA samples were analysed using a Perkin Elmer DSC 6 (Beaconsfield, Buckinghamshire, UK) test- Fig. 1 PLLA tensile and extruded rod test samplesing machine, over a temperature range of 40 °C to 200 °C H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 5. 311DEGRADATION OF POLY-L-LACTIDE. PART 1 Table 2 In vitro and in vivo molecular weight results versus time results (PI=polydispersity index) Tensile samples Extruded rod samples Degradation In vitro In vitro In vivo time (weeks) Mw Mn PI Mw Mn PI Mw Mn PI 0 424 000 158 500 2.67 339 000 155 000 2.18 399 000 155 000 2.57 4 339 000 143 000 2.37 312 500 129 000 2.42 10 309 000 120 000 2.58 249 000 85 250 2.9 226 500 87 750 2.56 20 199 000 72 500 2.74 176 000 82 900 2.12 26 199 000 93 850 2.11 149 000 70 150 2.13 153 500 70 050 2.23 32 159 000 65 800 2.4 127 000 60 800 2.08 38 133 500 53 050 2.51 89 100 40 850 2.18 105 500 54 250 1.94 44 74 900 22 500 3.33 76 400 37 850 2.02 69 800 32 600 2.15 Table 3 Thermal properties of extruded PLLA rod in vitro Degradation time (weeks) % crystallinity Tm, °C Tg onset, °C 0 40.7 184.2 68.9 10 45.5 184.6 70.7 26 44.9 184.8 67.2 32 40.8 183.8 65.5 38 53.4 183.6 63.6 44 53.6 182.4 58.7 Considering the data presented in Table 2 in con- junction with equations (3) and (6), Fig. 4 presents plots of 1/Mn and ln Mn versus time, modelling the uncatalysed and autocatalysed degradation models, with linear trendlines fitted and R2 correlation coefficients displayed. In both cases a higher correlation coefficient was achieved for the ln Mn versus time relationship,Fig. 2 Degraded PLLA tensile samples at 37 °C in vitro describing the autocatalytic degradation mechanism. The GPC chromatograms for the PLLA tensile3.2 Molecular weight versus time samples remained monomodal throughout successive The molecular weights of the PLLA tensile and extruded weeks of degradation (Fig. 5) and shifted towards lower rod samples in vitro and in vivo decreased with time molecular weights as degradation time increased. A (Table 2). After 44 weeks the Mw of the extruded rod similar trend was observed for the extruded rod samples samples decreased by approximately 80 per cent, while in vitro and in vivo. the Mw for the tensile samples in vitro and extruded rod in vivo decreased by approximately 82 per cent. A similar decreasing trend was also observed for the extruded rod 3.3 Mass change versus time and tensile samples’ Mn in vitro and in vivo (Fig. 3). Before drying, a similar pattern for the percentage massHowever, no obvious pattern in polydispersity index change of both the tensile and extruded rod samples(PI) (Table 3), the ratio of Mw/Mn, with time could be in vitro was observed (Fig. 6). After four weeks the massderived from the molecular weight data. of both sets of samples had increased by approximately 0.6 per cent, this increase remained relatively constant up until about week 44, when the polymer’s mass increased to approximately 1 per cent. Further increases in mass before drying were observed throughout sub- sequent weeks up to approximately 2.5 per cent for both the tensile and extruded rod samples at 65 weeks. After drying, a similar pattern was again observed for both sets of samples, with minimal mass loss observed at 57 weeks for both the tensile and extruded rod samples. However, at 65 weeks the tensile samples’ mass had Fig. 3 Comparison between in vitro and in vivo samples’ Mn decreased by approximately 1.2 per cent with a 0.9 per H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 6. 312 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON Fig. 4 Uncatalysed and autocatalysed models for tensile and extruded rod samples cent mass loss observed for the extruded rod samples. A similar pattern was also observed for the extruded rod samples mass change in vivo, however, the increases in mass before drying were not as large, peaking at approxi- mately 0.8 per cent after 44 weeks, with no significant mass loss observed after drying. 3.4 DSC analysis versus time For both the tensile and extruded rod samples in vitro, a general trend was observed of increasing crystallinity and decreasing Tg onset temperature with degradation time (Tables 3 and 4). Additionally, a slight but signifi-Fig. 5 PLLA tensile samples molecular weight distributions cant decrease in both the tensile and extruded rodat 0, 10, 32, and 44 weeks in vitro samples melting point, Tm, was also observed after 44 weeks. The results of the thermal analysis conducted on Table 4 Thermal properties of PLLA tensile samples in vitro Degradation time (weeks) % crystallinity Tm, °C Tg onset, °C 0 44.8 182.1 66.9 4 45.5 181.4 68.8 10 41.6 181.6 67.7 20 42.1 182.1 66.8 26 47.0 181.4 65.7 32 51.9 180.3 61.9 38 54.8 181.0 63.2 44 58.7 179.1 57.3 Fig. 6 In vitro mass change analysis H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 7. 313DEGRADATION OF POLY-L-LACTIDE. PART 1 the extruded rod in vivo also followed this general trend, time, with degradation proceeding further, the main melting peak began to shift to lower temperatures, untilwith the in vitro and in vivo extruded rod results proving to be very similar. eventually the two peaks merged and the smaller peak appeared as a shoulder on the larger peak (Fig. 7(f )).The DSC thermograms for the tensile samples at 0, 4, 10, 20, 38, and 44 weeks are shown in Fig. 7. At 0 weeks (Fig. 7(a)), a small endothermic peak commencing at approximately 67 °C was observed, relating to stress 3.5 Mechanical strength versus time relaxation at the polymer’s Tg [35]. As the temperature The mechanical strengths of both the tensile and increased further a small endothermic dip was observed extruded rod samples deteriorated with time, with the just before melting commenced, followed by the main tensile strength of the compression-moulded samples melting peak. It is suggested that the dip before melting reduced to approximately zero in 44 weeks (Table 5). was caused by some crystallization of the polymer. After 44 weeks the samples were very brittle and could Although the polymer was annealed prior to degradation, not be gripped in the tensile test grips without fracturing. with the aim of limiting crystallization throughout the A similar trend was observed initially for the loss of study, close to the polymer’s melting point the chain shear strength for the extruded rods in vitro and in vivo mobility would have increased, allowing some of the (Fig. 8). However, after 44 weeks in vitro, the rod amorphous segments to order themselves into a more samples had lost approximately 52 per cent of their crystalline structure. original strength, compared to only 26.3 per cent for the At ten weeks (Fig. 7(c)), as degradation increased, samples in vivo. presumably in the amorphous regions, the initial dip before melting observed at 0 and 4 weeks was reduced, with less amorphous regions remaining capable of 3.6 Biological host tissue’s response crystallization. At 20 weeks (Fig. 7(d)) the endothermic dip before The combination of LM, TEM, and SEM proved useful in determining the relationship between the PLLA rodmelting had disappeared and a small peak appeared to form in its place. It is speculated that this new peak implants and surrounding biological tissue. After 36 weeks the PLLA implant appeared to stimulate the pro-represented the melting of new crystallites formed by the crystallization of internal degradation by-products. The duction of a fibrous tissue capsule (Figs 9(a) and (b)) in which type-1 collagen fibre production was extensivereduction of the amorphous regions and crystallization of the degradation by-products resulted in the polymer’s (Fig. 9(c)). The TEM image of the fibrous capsule surrounding the PLLA implant showed the presence ofoverall crystallinity increasing throughout degradation (Tables 3 and 4). At weeks 38 and 44 (Figs 7(e) and (f )), the newly Table 5 Deterioration in PLLA’s tensile properties formed peak appeared to have grown and moved to a throughout degradation higher temperature, evidence that the newly formed Degradation Young’s Tensile Extensioncrystallites’ size may have been increasing. At the same time modulus strength at break (weeks) (MPa) (MPa) (mm) 0 668.4 64.3 1.6 4 618.4 53.8 1.6 10 625.2 60.3 1.5 20 433.0 23.7 0.8 26 528.7 36.9 1.0 32 226.5 9.9 0.4 38 284.2 8.2 0.3 44 – 1.0 0.4 Fig. 7 PLLA tensile samples’ DSC thermograms at 0, 4, 10, Fig. 8 Shear strength comparison of PLLA rod versus time in vitro and in vivo20, 38, and 44 weeks H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 8. 314 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON the absence of cells marking an inflammatory response at 36 and 44 weeks would suggest that the PLLA rod investigated was biocompatible throughout the 44 weeks duration of the study. 4 DISCUSSION 4.1 Degradation mechanisms The results presented show that the in vitro and in vivo degradation of PLLA commences almost immediately, with the in vitro tensile and extruded rod samples losing approximately 20 per cent of their initial molecular weight at four weeks, the first time point analysed. Autocatalysis. The higher R2 correlation coefficients (Fig. 4) achieved for the ln Mn versus time relationship shows a closer approximation of the autocatalysed model (equation (6)) to the experimental data, compared to the uncatalysed model, with degradation accelerated by the newly formed carboxylic acid end groups generated by the continual ester hydrolysis reaction. However, it cannot be concluded that the mechanism is exclusively autocatalytic. Investigating the in vitro degradation of SR-PLLA at 37 °C, Pohjonen and To¨rma¨la¨ [36] observed a similar trend. They reported correlation coefficients of 0.989 for the autocatalysed model and 0.910 for the uncatalysed model, confirming the findings of the present study and at least, according to theory, suggesting an autocatalytic degradation mechanism. In contrast, in a comparative study investigating the molecular weight versus time data available in literature for semicrystalline aliphatic polyesters such as PLLA and amorphous poly- mers such as PDLLA, Anderson [30] reported that no clear distinction could be derived between the uncatalysed 1/Mn and autocatalysed ln Mn plots versus time for semi- crystalline polymers. However, for amorphous polymers, the results were reported to be much more consistent with an autocatalytic mechanism, with higher correlation coefficients achieved for plots of ln Mn versus time. Although, the correlation coefficients for each study were not given, making comparisons to the present study difficult. As a result of these studies Anderson [30] con- cluded that the hydrolytic degradation of semicrystalline polyesters may not proceed exclusively by non-catalytic or autocatalytic mechanisms, speculating that both may contribute to the rate of chain scission. Relationship between molecular weight distribution and degradation. Considering the GPC curves for the tensileFig. 9 Images of PLLA biological tissue after 36 weeks implantation samples (Fig. 5) it is interesting to note that they remained monomodal throughout successive weeks of degradation. In contrast, many researchers have reported that asfibroblasts (Fig. 9(c)) in an extracellular matrix (inset) composed extensively of type-1 collagen fibre bundles the degradation of PLA and PGA aliphatic polyesters proceeds the initially monomodal GPC curve becomesrunning in different orientations. While tissue disturbance during surgery produces an initial inflammatory response, bimodal and even multimodal in nature [14] and [37]. H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 9. 315DEGRADATION OF POLY-L-LACTIDE. PART 1 The bimodal nature of these GPC curves was originally butions were the result of preferential degradation of the amorphous regions, a view supported by Fischer et al.assigned to the difference in degradation rates in the [32]. It would be anticipated that due to the semicrystallineamorphous and crystalline regions [14], [32] and [37]. nature of the compression-moulded and extruded PLLAHowever, with molecular weight usually determined by investigated in the present study [25], the GPC curvestaking samples from the bulk of the polymer, comprising of both would become bimodal and even multimodalthe interior of a lower molecular weight than the surface in nature due to the preferential degradation of the[37], Li et al. [19] were the first to assign this bimodal amorphous regions. However, Fig. 5 shows that this isbehaviour to the autocatalytic effect and faster internal not the case, with the molecular weight distributionsdegradation. It is currently understood that for PLA and remaining monomodal throughout the 44 weeks durationPGA polymers and their copolymers, this bimodal nature of the study, contradicting the findings of Li et al. [20]can be accounted for by three different mechanisms and Pistner et al. [39]. It must also be assumed that withrelated to the polymers morphology [14]. First, by faster the samples intended for molecular weight analysis takeninternal degradation; however, this mechanism is most through a cross-section of the material, the suspectedcommonly observed for initially amorphous polymers, autocatalytic mechanism did not result in a large enoughwhich are not believed to be capable of crystallization surface-interior differentiation to yield curves containingeven throughout degradation [37], for example, a 50:50 two distinct molecular weight species.copolymer of PLA and PDLA [19]. Second, for semi- crystalline polymers the bimodal nature has been attri- buted to selective degradation of the amorphous regions, Bulk degradation. The time delay before mass loss with the surface-interior differentiation reported not to observed in this study (Fig. 6) is in agreement with be large enough to yield bimodal GPC chromatograms the reported general sequence of aliphatic polyester [14], supporting Anderson’s theory [30] that the hydro- degradation, which suggests molecular weight loss is lytic degradation of semicrystalline polyesters does not observed first, before loss of mechanical strength and proceed exclusively by non-catalytic or autocatalytic before any physical mass loss is observed [8]. This is mechanisms. Finally, the bimodal nature of the GPC accounted for by the fact that water diffusion into the chromatograms has been attributed to the crystallization polymer is faster than the hydrolytic degradation of of low molecular weight degradation by-products in the polymer’s ester linkage, suggesting that ester-bond initially amorphous polymers, for example, amorphous cleavage is the rate-limiting step in the degradation of PLLA and a 75:25 PLA/PGA copolymer [38], which are aliphatic polyesters [40]. This results in degradation pro- capable of crystallizing throughout degradation. Once ceeding in the bulk of the polymer, resulting in a time- lag before any mass loss is observed, as the polymer’sthe low molecular by-products crystallize they become molecular weight has to be reduced to a critical valueresistant to degradation and appear as a low molecular before soluble oligomers can be released. In contrast, forweight peak on the GPC curve. bioresorbable polymers regarded as surface eroding,The monomodal nature of the GPC curves obtained such as those belonging to the polyanhydride and poly-for the semicrystalline PLLA prepared by annealing orthoester families [41], mass loss is observed almostand investigated in this study appears to contradict immediately, as the chain scission of their more reactivethe findings of other researchers investigating similar unstable linkages, in comparison to the ester linkage insemicrystalline PLLA. Li et al. [20], investigating the aliphatic polyesters, is faster than the diffusion of waterdegradation of semicrystalline PLLA prepared by anneal- molecules into the polymer [40].ing at 130 °C for two hours, with an initial crystallinity of 72 per cent deduced from XRD measurements, observed that the initial monomodal molecular weight Polymer morphology and degradation. The results of the distribution became multimodal after 18 weeks. After 50 DSC analysis (Fig. 7) appear to provide evidence that weeks Li et al. [20] observed that the GPC curve became the low molecular weight degradation by-products are bimodal, with the peak corresponding to high molecular capable of crystallizing, due to their greater mobility, weight being more prominant for the surface than for and contribute to the samples’ increasing crystallinity. the centre, suggesting autocatalysis. At 90 weeks the This is evident by the emergence of a small peak forming GPC chromatogram then became almost monomodal and eventually merging with the main melting peak. The and was composed of a single low molecular weight crystallization of these internal degradation by-products peak. Pistner et al. [39] observed a similar profile for the resulted in the polymer maintaining its structural integrity GPC chromatograms of semicrystalline PLLA with an throughout the duration of the study. In contrast, hollow initial crystallinity of 73 per cent measured by DSC, with structures have been reported for intrinsically amorphous a low molecular weight shoulder observed after eight polymers since their degradation products are not weeks, becoming more important as degradation time believed to be capable of crystallizing, for example, in proceeded. Both Li et al. [20] and Pistner et al. [39] the case of a 50:50 copolymer of PLLA/PDLLA [19]. The decreasing peak melting temperature, observed mostconcluded that the multimodal molecular weight distri- H01004 © IMechE 2004 Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 10. 316 N A WEIR, F J BUCHANAN, J F ORR AND G R DICKSON significantly at 44 weeks (Tables 3 and 4) and deter- In vitro and in vivo degradation rates. Literature regard- ing the role of enzymes on the degradation of aliphaticmined from a single heating cycle, is consistent with the hypothesis that the initially crystalline regions are resistant polyesters is often contradictory. Many authors have reported that enzymes may be involved in the latter stagesto degradation, resulting in a two-stage degradation mech- anism with the amorphous regions being preferentially of degradation, when the polymer has fragmented and the molecular weight is sufficiently small [47–49]. However,attacked [14]. However, once the amorphous regions have been exhausted, the less accessible crystalline regions the role of enzymes during hydrolysis of the polymer bulk remains unclear. In comparative in vitro and in vivoare then solely attacked and disrupted, resulting in a decreased size of the initially present crystallites and studies, Vasenius et al. [50] have reported significantly faster degradation of PGA rods in vivo, with Matsusuehence a reduced melting point [42]. Although the melt- ing point of bioresorbable polymers is also known to be et al. [51] also reporting faster in vivo degradation of PLLA. In each case the faster in vivo degradation ratedependent on molecular weight, the extent of this is most readily determined by considering a reheat DSC run. was attributed, in some part, to the action of enzymes. However, comparative studies by Hooper et al. [52] andSince the fusion of the first run destroys the polymer’s initial crystalline structure, crystallization on cooling Pitt et al. [53] have reported no significant differences in the degradation rates of poly(a-hydroxy acids) in vivo,involves the degraded chains only [36], confirming that in the present study the decrease in melting temperature, with Cam et al. [54] reporting their degradation to be practically independent of enzymes. The similaritydetermined from a single heating cycle, was most likely a result of a reduction in crystallite size and not decreasing between the results of the molecular weight (Fig. 3) and DSC analysis for the PLLA rods investigated in vitromolecular weight. It is speculated that the decreasing Tg observed as degradation time increased (Tables 3 and 4) and in vivo in the present study suggests that the degradation of PLLA is independent of enzymes, and inis related to the reduction in molecular weight of the polymer’s chains in the amorphous regions, with a similar agreement with Timmins and Lenz [55], who reported that enzymes capable of catalysing hydrolysis are them-trend also observed by Li [14], Duek et al. [43], Joukainen et al. [44], and Kelloma¨ki et al. [45]. Interestingly, a selves macromolecules, unable to penetrate into the poly- mer bulk. Therefore, any enzyme-contributed reactionsmall amount of water within a polymer is also known to have a marked plasticizing effect, causing a reduction would be heterogeneous and confined to the surface of the polymer, with a reduction in mass observed, but littlein the polymer’s Tg. A study by Siemann [46], investi- gating the influence of water on the glass transition of change in the polymers overall molecular weight [4]. The significant loss of molecular weight (Table 2) andpoly(dl-lactic acid) by DSC, reported a 12 K decrease in Tg after samples were exposed to water for six hours negligible mass loss (Fig. 6) observed for the PLLA rod investigated in vivo in the present study would suggestprior to testing. However, in a further study, investi- gating samples exposed to water and then dried to a that degradation proceeded predominantly in the bulk of the polymer by non-enzymatic hydrolysis, similar toconstant mass before testing, the Tg remained the same as the untreated samples. In the present study the the mechanism observed in vitro. However, this does not rule out the influence of enzymes at later stages ofsamples were dried to constant mass before DSC testing was conducted, ruling out water acting as a plasticizer, the degradation process, particularly when mass loss becomes significant.and confirming a reduction in molecular weight as the most probable cause for the decreasing trend in Tg. However, this underlines the problem that for accurate Tg measurements, representative of the polymer’s con- 4.2 Biological response dition in service, test regimes need to be developed that can accurately monitor the polymer’s Tg while the The production of a fibrous capsule around bio- resorbable implants has been observed previously [56–58]samples remain ‘wet’. and is regarded as part of the body’s natural response to implants made of diverse materials [49]. SurgicalMechanical strength. Since degradation predominantly occurred in the amorphous regions, disrupting the tie intervention, such as the implantation procedure under- taken in this study, would initiate inflammation as achains holding the crystallites together, coupled with the decreasing molecular weight and increasing crystallinity, response to injury. However, the absence of inflam- matory cells at 36 and 44 weeks suggests that PLLAit is not surprising that the mechanical properties of the PLLA investigated decreased so rapidly. However, con- is biocompatible throughout the early stages of its degradation. It is understood that the onset of mass loss,sidering the similarities between the molecular weight loss and the results derived from the DSC analysis for particularly in fast degrading aliphatic polyesters such as PGA, can result in an inflammatory reaction due tothe extruded PLLA rod in vitro and in vivo, it is difficult to speculate at this stage why the samples in vitro the sudden release of acidic degradation by-products, causing a large change in pH of the surrounding mediaappeared to lose their strength more rapidly than those in vivo. [59]. In the present investigation the PLLA degradation H01004 © IMechE 2004Proc. Instn Mech. Engrs Vol. 218 Part H: J. Engineering in Medicine
  • 11. 317DEGRADATION OF POLY-L-LACTIDE. PART 1 scaffolds and cells. In Synthetic Biodegradable Polymerstudy in vivo was terminated before any mass loss was Scaffolds (Eds A. Atala and D. J. Mooney) 1997, pp. 1–14observed, although it is speculated that any inflam- (Birkha¨user, Boston, MA, USA).matory response observed as a direct result of the onset 4 Li, S. and Vert, M. Biodegradation of aliphatic poly-of polymer mass loss would be mild. In comparison to esters. In Degradable Polymers Principles & Applications fast degrading PGA implants, it is anticipated that the (Eds G. Scott and D. Gilead) 1995, pp. 43–87 (Chapman release of acidic degradation products from the slower & Hall, London). degrading PLLA would be less intense. This would 5 Chu, C. C. Biodegradable polymeric biomaterials: an result in the surrounding tissue being more capable of overview. In The Biomedical Engineering Handbook eliminating any such debris more efficiently, reducing the (Ed. J. D. Bronzino) 1995, pp. 611–626 (CRC Press, Boca Raton, FL, USA).risk of a severe inflammatory reaction developing that 6 Higgins, N. A. Condensation of Polymers of Hydroxyaceticwould require further surgical intervention. Acid. US Patent 2 676 945, 1954. 7 Schneider, A. K. Polymers of High Melting Lactide. US Patent 2 703 316, 1955.5 CONCLUSIONS 8 Middleton, J. C. and Tipton, A. J. Synthetic biodegradable polymers as orthopedic devices, Biomaterials, 2000, 21, The results of the analytical characterization studies 2335–2346. conducted on the retrieved PLLA samples in vitro and 9 Kulkarni, R. K., Pani, K. C., Neuman, C. and Leonard, F. in vivo provides strong evidence to support the findings Polylactic acid for surgical implants, Arch. Surg., 1966, of other researchers investigating similar bioresorbable 93, 839–843. polymers. Additionally, the results from the in vivo 10 Hofmann, G. O. Biodegradable implants in orthopaedic surgery—a review on the state-of-the art, Clin. Mater.,studies would suggest that throughout the first stage of 1992, 10, 75–80.degradation, before mass loss is observed, PLLA is bio- 11 Ciccone, W. J., Motz, C., Bentley, C. and Tasto, J. P.compatible and degrades at the same rate in vitro and Bioabsorbable implants in orthopaedics: new developmentsin vivo. However, the results of the present studies do and clinical applications. J. Am. Acad. Orthop. Surg., 2001, appear to indicate that for semicrystalline polymers, like 9, 280–288. the PLLA investigated, no clear differentiation between 12 Barber, F. A. Resorbable materials for arthroscopic surface and interior degradation could be observed that fixation: a product guide. Orthopedic Special Edn, 2002, would clearly point to an autocatalytic degradation 8, 29–37. mechanism. As a result it is speculated that as poly- 13 Hayashi, T. Biodegradable polymers for biomedical uses, Prog. Polym. Sci., 1994, 19, 663–702.mer crystallinity increases, the importance of the auto- 14 Li, S. Hydrolytic degradation characteristics of aliphaticcatalysis degradation mechanism may become less polyesters derived from lactic and glycolic acids, J. Biomed.significant. Mater. Res. (Appl. Biomater.), 1999, 48, 342–353. 15 Vert, M., Li, S. and Garreau, H. New insights on the degradation of bioresorbable polymeric devices based on ACKNOWLEDGEMENTS lactic and glycolic acids, Clin. Mater., 1992, 10, 3–8. 16 Ali, S., Doherty, P. J. and Williams, D. F. Mechanisms The authors would like to thank Mr David Farrar of polymer degradation in implantable devices. 2. at Smith & Nephew Group Research Centre (York, Poly(DL-lactic acid), J. Biomed. Mater. Res., 1993, 27, UK), Boehringer Ingelheim (Ingelheim, Germany) for 1409–1418. 17 Chu, C. C. Degradation and biocompatibility of syn-supplying the PLLA, Griffith Microscience (Derbyshire, thetic absorbable suture materials: general biodegradationUK) for the ethylene oxide sterilization, and Rapra phenomena and some factors affecting biodegradation.Technology Limited (Shropshire, UK) for the molecular In Biomedical Applications of Synthetic Biodegradableweight characterization. Finally, the EPSRC (Swindon, Polymers (Ed. J. O. Hollinger) 1995, pp. 103–128 (CRC UK) for financial assistance. Press, Boca Raton, FL, USA). 18 Mainil-Varlet, P., Curtis, R. and Gogolewski, S. Effect of in vivo and in vitro degradation on molecular and mech- REFERENCES anical properties of various low-molecular-weight poly- lactides, J. Biomed. Mater. Res., 1997, 36, 360–380. 1 To¨rma¨la¨, P., Pohjonen, T. and Rokkanen, P. Bioabsorbable 19 Li, S. M., Garreau, H. and Vert, M. 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