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Biodegradation of polyurethane under fatigue loading

Michael J. Wiggins,1,2 James M. Anderson,2,3 Anne Hiltner1,21Center for Applied Polymer Research, Case Western Reserve University, Cleveland, Ohio 441062Department of Macromolecular Science, Case Western Reserve University, Cleveland, Ohio 441063Department of Biomedical Engineering and Institute of Pathology, Case Western Reserve University,Cleveland, Ohio 44106

Received 27 June 2002; revised 6 August 2002; accepted 14 August 2002

Abstract: A method utilizing expansion of a diaphragm-type film specimen was developed to study in vitro biodeg-

radation of poly(etherurethane urea) (PEUU) under condi-tions of dynamic loading (fatigue). A finite element modelwas used to describe the strain state, which ranged fromuniaxial at the edges of the film to balanced biaxial tensilestrain at the center. During testing, the film was exposed toa H2O2/CoCl2 solution, which simulated in vivo oxidative

biodegradation of PEUU. The extent of chemical degrada-tion was determined by infrared analysis. Physical damageof the film surface was characterized by optical microscopyand scanning electron microscopy. Dynamic loading did notaffect the rate of degradation relative to unstressed andconstant stress (creep) controls in regions of the film thatexperienced primarily uniaxial fatigue; however, degrada-

tion was accelerated in regions that experienced balancedbiaxial or almost balanced biaxial fatigue. It was concluded

that the combination of dynamic loading and biaxial tensilestrain accelerated oxidative degradation in this system.Chemical degradation produced a brittle surface layer thatwas marked by numerous pits and dimples. Physical dam-age of the surface in the form of cracking occurred only infatigue experiments. Cracking was not observed in un-stressed or creep tests. Cracks initiated at the dimples pro-duced by chemical degradation, and propagated in a direc-tion that was determined by the strain state. 2003 WileyPeriodicals, Inc. J Biomed Mater Res 65A: 524-535, 2003

Key words: polyurethane; biodegradation; fatigue; creep

INTRODUCTION

Polyether-based polyurethane (PEU) has been uti-lized in a number of implantable biomedical devices.1

One of the first biomedical applications of PEU uti-lized the material as the outer insulating sheath sur-rounding cardiac pacing leads.2 From early experiencewith pacing leads, PEU was recognized as biocompat-ible and was thought to be biostable. It was not untilnearly a decade after PEU was first used in pacingleads that it was recognized that a small percentage ofleads were failing.3 Lead failure occurred when phys-ical damage to the PEU, in the form of cracks andtears, compromised its ability to act as an insulator.Analyses of failed leads led to the understanding thatphysical damage could occur in response to residualstress or strain. This type of physical damage waslabeled as environmental stress cracking (ESC).4 An-nealing the PEU to reduce residual stresses, e.g., hoop

stresses due to interference fits between the outer PEUinsulating sheath and the conductor coil, and aban-doning lead designs that caused large bending strainson the PEU, e.g., in the J-type leads, partially resolvedthe problem of ESC.3,5,6 It was also recognized thatchemical degradation of the PEU macromolecule wasinvolved in lead failure and that biostability was likelya function of interrelated phenomena of physical dam-age and chemical degradation.7 It is now understoodthat the mechanism of chemical degradation involvesoxidative attack on the polyether soft segment, start-ing with abstraction of a proton from the -methyleneposition and leading to chain scission and/or chemicalcrosslinking.8,9

Based on early experience with pacing leads, a num-ber of in vivo and in vitro experiments were devised tofurther understand the influence of mechanical stresson chemical and physical degradation. It was shownthat constant uniaxial stress (creep mode) affected thein vitro biostability of a poly(etherurethane urea)(PEUU) in two ways.10,11 It inhibited chemical degra-dation by causing stress-induced crystallization of theether soft segment. On the other hand, stress enhanced

physical damage. Stress concentration at surface

Correspondence to: A. Hiltner; e-mail: [email protected] grant sponsor: National Institutes of Health; con-

tract grant number: HL-25239

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flaws, created by localized chain scission, caused theflaws to enlarge first into small pits and then intolarger voids as the pits coalesced. Similar results wereobserved for in vitro degradation of the same PEUUunder constant uniaxial strain (stress relaxationmode).12 Increasing uniaxial strain caused progres-

sively more orientation of soft segment chains, whichcorresponded to a progressive decrease in the amountof chemical degradation. Physical damage was en-hanced by constant uniaxial strain, but only at veryhigh strains.

In addition to uniaxial strain, the affect of constant biaxial strain on biostability has also been tested.12

Various levels of unbalanced biaxial strain weretested. Biaxiality inhibited orientation of the ether softsegment. Orientation of the soft segment in the direc-tion of highest strain did occur, but only at higherstrains, depending upon the ratio of the level of strainin each direction. As would be expected, chemicaldegradation was inhibited when orientation did oc-cur. When no orientation was detected, biaxial strainhad no effect on chemical degradation. Chemical dam-age, particularly at high levels of biaxial strain, en-hanced physical damage. Biostability in these exam-ples can be explained by competing factors:orientation inhibited chemical degradation whereasthe applied load enhanced physical damage.

Although understanding the effects of constantstress or strain on biostability helps considerably inexplaining the phenomenon of ESC in pacing leadinsulation, little research has been devoted to the ef-

fects of dynamic loading on biostability. The potentialimportance of dynamic loading can be appreciated ifone considers that pumping diaphragms, present inventricular assist devices and total artificial hearts,require a flexible membrane that exhibits long-termbiostability under dynamic stress. Potential problemswith using PEU as pumping diaphragms could beanticipated by studying the effect of dynamic stress onPEU biostability. It is reasonable to anticipate a signif-icant effect of dynamic strain on the working lifetimeof a PEU pumping diaphragm, considering the dem-onstrated effects of constant stress and strain on bio-

stability, and well-documented observations that dy-namic stresses promote chain scission, crack initiationand propagation in other environmentally challengedsynthetic elastomers.1316

Various researchers have investigated the effects ofdynamic stresses on PEUs,1722but most of the studiesdid not directly address biostability. Furthermore,most fatigue experiments were performed in uniaxialtension, whereas the diaphragm operates under mul-tiaxial stresses. In the present study, a method wasdeveloped to study the effect of dynamic mechanicalstress (fatigue) on the biostability of PEUU. Thismethod allowed fatigue experiments to be performed

in an environment that has been shown to reproduce

the oxidation of PEU observed in vivo. Moreover, theexperimental design permitted uniaxial and balancedbiaxial fatigue to be performed simultaneously. Forcomparison, constant stress experiments and experi-ments with no applied stress were also performed.During fatigue, the PEUU was exposed to a H2O2/

CoCl2 environment that simulates in vivo degradationof PEUU.23 Chemical degradation and physical dam-age were evaluated as measures of biostability.

MATERIALS AND METHODS

Materials

The poly(etherurethane urea) (PEUU) was described pre-viously.24 It had a poly(tetramethyleneglycol) soft segment(Mn 2000) and a di(p-phenyl isocyanate) (MDI) hard seg-ment chain extended with diamines. To prevent prematurefailure of the PEUU at the clamped edges of the fatiguespecimen, the tested films were bilayers of PEUU and apolycarbonate polyurethane that was stable to oxidativedegradation.25 The polycarbonate polyurethane (PCU) had ahexanediol carbonate soft segment (Mn 1939) and an MDIhard segment chain extended with diamines. Both polyure-thanes were in a solution of dimethyl acetamide, 20% solids

by weight.To fabricate the bilayer film, the PCU solution was cast

onto a glass substrate to a uniform thickness using a doctorblade. The solvent was removed in vacuo at ambient tem-perature. The PEUU was then cast over the PCU layer with

a doctor blade and dried in vacuo. The PCU base layer andthe PEUU top layer were both approximately 105110 mthick, giving a total film thickness of 210220 m.

In vitro fatigue and creep experiments

Testing was carried out with a custom-designed filmholder coupled to a commercial Flexercell Strain Unit (Flex-cell International Corp., McKeesport, PA). The design of thefilm holder allowed deformation of a circular section of thefilm by application of a vacuum to the bottom of the film

(PCU side), whereas the top of the film (PEUU side) wasexposed to an in vitro testing solution. The film holderconsisted of two blocks, an upper block constructed ofpoly(oxymethylene) (POM) and a lower block constructed ofaluminum. A single continuous film was inserted betweenthe blocks, which were secured with a series of screws. The

blocks contained eight specimen wells. One well is shown inFigure 1. The space above the specimen was filled with theappropriate test solution, and the film was deformed byapplication of a vacuum from below. The magnitude andfrequency of the applied vacuum were controlled with theFlexercell Strain Unit. During a typical test, four of the wellswere subjected to loading, either fatigue or creep, and theremaining four wells were not loaded and served as con-

trols. A typical experiment lasted for 12 days. At intermedi-

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ate time points, 3, 6, and 9 days, a section of the film from aloaded well and from an unloaded well were removed foranalysis.

To reproduce the mechanism of in vivo PEUU degradationat an accelerated rate,12 the wells were filled with a solutionof 20% H2O2/0.1M CoCl2. As controls, wells were filled witha solution of 20% H2O2 or with distilled water. All solutionswere refreshed every 3 days, and all testing was done at37C.

Fatigue testing was carried out at a frequency of 1 Hz witha maximum load of 50 kPa. The stress as a function of timewas measured during fatigue testing with a pressure trans-ducer. The fatigue waveform as measured in this way isshown in Figure 2. The mean load applied during fatiguetesting was 25 kPa. Creep experiments were performedat a constant applied load of either 50 or 25 kPa.

Characterization of strain state

Deformation varied across the film from essentially uni-axial strain at the edges to balanced biaxial strain at thecenter. The strain was resolved into radial and tangentialcomponents, r and t, respectively, as shown in Figure 1(b).The deflection at the center of the film (d) was measuredfrom the vertical displacement of a pipette. A separate de-termination of the strain at the center of the film () wasobtained by depositing gold squares on a film and measur-ing the changes in dimension of the squares at maximum

deflection with an optical microscope (OM). The pattern was

deposited by placing a flat metallic mesh on the outer PEUUsurface of the film (40 lines per inch) and sputter coating thefilm with 200 of gold in a Hummer 6.2 sputtering system(Anatech Ltd., Alexandria, VA). The mesh was removed,leaving gold squares on the film surface. It was only possibleto measure the strain at the center of the film in this way

because of curvature during deformation.

After direct measurement of d and a silicone cast of thedeformed film was made for determining the strain distri-

bution across the film. An RTV11 silicone base compoundand dibutyl tin dilaurate curing agent was used according tothe manufacturers recommendation to make the cast (GESilicones, Waterford, NY). The silicone base compound andcuring agent were mixed and poured into the well. Aftercuring for 24 h, the film and the silicone cast were removedfrom the loading cell. The film remained adhered to thesilicone cast upon removal. Because the film was clear, thegold squares were visible through the film. With the filmadhered to the silicone cast, the dimensions of the goldsquares were measured by optical microscopy keeping the

line of sight perpendicular to the surface of the silicone castto ensure accurate measurements. The radial and tangentialstrains as a function of distance from the center of the filmwere measured in this way. The deflection was also mea-sured from the silicone cast. To obtain the strain in thethickness direction, the silicone cast was bisected throughthe diameter. The thickness of the film as a function ofdistance from the center was measured from micrographsrecorded with a scanning electron microscope (SEM; JEOL840-A, Peabody, MA). Before examination in the SEM thesilicone cast and the adherent film were coated with 100 ofgold.

The deflection was measured every 3 days during theexperiment. Because significant creep occurred during the

course of an experiment, a finite element model (FEM) wasdeveloped to predict the strains on the film surface based onthe measured deflection. The FEM also served to confirm thestrain distribution that was measured with the silicone cast.Finite element modeling was done with the software pack-age ABAQUS/Standard, version 5.6 (Hibbitt, Karlsson &Sorensen, Inc., Pawtucket, RI). The film was modeled usinga total of 200 two-node linear elastic axisymmetric elements.The model was defined to be 2.5 cm in diameter and 220 m

Figure 2. Load versus time curve for two cycles of thefatigue test.

Figure 1. Schematic diagram of the test cell used for fa-tigue and creep experiments: (a) cross-sectional view and (b)top view.

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thick. Nodes around the boundary of the model were con-sidered pinned. Loading was distributed over the bottomsurface of the elements and was applied incrementally infour steps. The software provided the positional displace-ment and the tangential and radial strain at each node.

Surface characterization

Before and after treatment, each film was analyzed byattenuated total reflectance Fourier transform infrared mi-crospectroscopy (FTIR). Data were collected with a Nexus870 FT-IR bench coupled to a Continuum microscope(Thermo Nicolet, Madison, WI). The microscope utilized anattenuated total reflectance (ATR) slide-on attachment witha germanium crystal to allow micro-ATR analysis. Spectrawere collected at a resolution of 2 cm1 for 32 scans. Spectrawere collected every 1 mm along a diameter over the filmsurface. Each spectrum sampled an area 150 150 m.

Chemical degradation was quantified as the percent etherremaining after treatment compared with the untreated film.Ether content was determined from the absorbance of the1110 cm1 peak (asymmetric stretch of soft segment ether) inthe FTIR spectra relative to the internal reference peak at1595 cm1 (stretching of the carbon-carbon double bond inthe aromatic hard segment). The amount of crosslinking wasquantified from the absorbance of the 1174 cm1 peak(stretch of soft segment branched ether) relative to the in-ternal reference peak. Physical damage was characterized byOM and SEM. Representative SEM images were recorded atthe center of treated films and at 4, 8, and 12 mm from thecenter. All SEM images are of surfaces tilted by 60 to show

surface morphology.

RESULTS AND DISCUSSION

Finite element model

The finite element model (FEM) was tested by com-paring the displacement at the center (d) and the strainat the center of film ( t r) at maximum loadingto the FEM prediction. The measured values of d and

were 6.22 mm and 17.2%, respectively, at a loadingof 50 kPa. Initially, the tensile modulus from uniax-ial stress-strain measurements, E 12.1 MPa, and atypical Poissons ratio for an elastomer, 0.49, wereused in the FEM. These values, however, predicted dand to be much larger than the measured values(Table I). The FEM accurately predicted the measuredvalues of d and with E 14.1 MPa and 0.3(Table I).

From the silicone cast, d was measured to be 6.61mm and was measured to be 19.0%. These valueswere higher than those directly measured from thefilm. This was attributed to creep of the film during

the 24-h period required for complete curing of the

silicone resin. Accommodating the increase in d and required a decrease in E from 14.1 to 12.4 MPa in theFEM (Table I).

To further test the FEM, the strain distribution onthe film measured from the gold grid on the siliconecast was compared with the strain distribution fromthe FEM. Despite scatter in the measured r and tdata, which was attributed to imperfect alignment ofthe silicone cast surface in the optical microscope, theFEM with E 14.1 MPa and 0.3 accurately pre-dicted the distribution in radial strain (r), tangentialstrain (t), and strain in the thickness direction (z) asmeasured from the silicone cast (Fig. 3). Representa-tive data is plotted in Figure 3. Radial and tangentialstrains were maximum and equal at the center (bal-anced biaxial strain state). Radial strain decreasedslightly toward the edge, whereas tangential straindecreased to zero at the edge (uniaxial strain state).

The tensile strains were accompanied by a decrease inthe thickness of the film.In formulating the FEM it was assumed that the

strains were small enough that the material behavedas a linear elastic solid. This assumption was reason-able based on the fact that the strains involved were20% and the uniaxial stress strain curve was approx-imately linear up to 50% strain. It was also assumedthat the modulus measured in uniaxial tension accu-rately described the response to biaxial deformation.This assumption may not have been valid. Some elas-tomers are reportedly stiffer in balanced biaxial ten-

sion than in uniaxial tension.26

This could explain whythe FEM required a higher modulus than the mea-sured uniaxial tensile modulus to produce the mea-sured values of d and .

An increase in volume during deformation was in-dicated by the strain measurements from the siliconecast and was reflected by the Poissons ratio of 0.3 inthe FEM. The Poissons ratio of polyurethanes similarto the one in this study has been measured to be between 0.45 and 0.5.27 Although using a Poissonsratio 0.5 was counterintuitive, a value of 0.3 accu-rately predicted the strain distribution in the de-formed PEUU film and therefore was used in subse-

quent calculations.

TABLE IComparison of Measured and Predicted Strains

Method d (mm) (%)

FEM Parameters

E (MPa) v

Direct measure 6.22 0.08a 17.2

FEM 7.45 24.9 12.1 0.49FEM 6.22 17.4 14.1 0.3Silicone cast 6.60 0.06a 19.0

FEM 6.61 19.4 12.4 0.3aAverage SD; n 4.



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Strain state during deformation

The time dependence of the maximum displace-ment was measured over the course of fatigue, andcreep experiments were carried out in different envi-

ronments. Figure 4 is representative data showing thechange in the deflection (d), defined as the differencebetween the maximum deflection measured at sometime point and the maximum deflection measuredimmediately after loading commenced. In Figure 4(a),d is plotted for creep loading at 50 kPa for filmstested in water, H2O2, or H2O2/CoCl2. In all threetreatments, an initial rapid increase in the deflectionwas observed between day 0 and day 3. After day 3,there appeared to be little or no further increase in themaximum deflection. Films tested in H2O2 and H2O2/CoCl2 deflected an additional 2.3 mm over the initial

displacement, whereas films tested in H2O deflectedonly an additional 1.6 mm. It is known that swelling ofPEUU by either H2O or H2O2 increases the compliancedue to plasticizing effects, and the effect is more pro-nounced with H2O2, which swells PEUU by 6.7% byweight, whereas H2O swells PEUU by 1%.

11 The in-crease in deflection between day 0 and day 3 wasattributed primarily to the increase in compliance dueto swelling. This behavior has been observed in pre-vious studies of PEUU creep behavior.10 One wouldexpect d to continue increasing after day 3 because ofthe viscoelastic nature of the PEUU. Indeed, this be-havior was observed previously and was termed pri-

mary creep.10 However, the magnitude of the effect

was small and would have been within experimentalerror of the d measurement.

In addition to primary creep, secondary creep waspreviously observed in PEUU treated in H2O2/CoCl2.

10 Secondary creep was characterized by an in-

crease in slope of the logarithmic creep response atlonger time points. It arose from the effect of chemicaldegradation on the bulk compliance of the film. Sec-ondary creep was not observed within the time scaleof the present experiments. Although chemical degra-dation responsible for secondary creep certainly oc-curred in the present experiments, construction of thefilm with an oxidation-resistant PCU base layer anddesign of the experiment with only the PEUU sideexposed to the in vitro solution would have minimizedthe effect of chemical degradation on bulk propertiesof the film.

The trend in d observed with creep loading at 50kPa was also observed in fatigue loading [Fig. 4(b)].There was an initial rapid increase in the deflectionbetween day 0 and day 3, with no further increase atlonger times. Similarly, a larger deflection was ob-served after day 3 for H2O2 treatments than for H2Otreatment. Comparing the magnitude of d betweencreep loading at 50 kPa and fatigue loading at 50kPa in H2O2/CoCl2 in Figure 4(c), it is seen that dafter day 3 was about twice as great in creep loadingas in fatigue loading, 2.3 and 1.2 mm, respectively. Thedeflection during creep loading at 25 kPa was 1.2

mm, the same as in fatigue at 50 kPa. This meant that

Figure 3. Measured strain distribution on the film at maximum displacement for 50 kPa loading (data points) and the FEMprediction (solid lines).

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the mean load applied to the film controlled the in-crease in deflection between day 0 and day 3.

The FEM was used to obtain the strain distributionon the film after day 3 of fatigue and creep testing. Themodulus used in the model was decreased to obtain amatch with the measured deflection. An initial loading

of 50 kPa produced an average displacement of 6.22mm. Between day 3 and day 12 of fatigue in H2O2/CoCl2, the deflection increased to an average value of7.34 mm (Table II). The modulus in the FEM decreasedfrom 14.1 to 10.1 MPa to match the new deflection andthereby accommodate the increase in effective compli-ance of the film. This corresponded to 23.3%. Thiswas significantly higher than the value measured im-mediately after loading, 17.2%. The average de-flection between day 3 and day 12 for creep loading at50 kPa was 8.56 mm. Decreasing the modulus in theFEM to 7.7 MPa to match the deflection gave

30.3%. These results are summarized in Table II.Subsequently r and t were obtained from the FEMusing the modulus values in Table II. In Figure 5(a), rand t are plotted against the distance from the centerof the film for fatigue loading. The strains were max-imum and equal at the center of the film. Toward theedge, t quickly decreased to zero, whereas r slowlydecreased to 70% of its maximum value. Region Iwas defined as the area where r and t were bothgreater than 85% of their maximum value. This regionof primarily biaxial strain was located within 5 mm ofthe center of the film. Region II was the area between

5 mm from the center and the edge of the film wherethe strain state gradually changed from almost bal-anced biaxial to uniaxial. Although the magnitude ofthe strain in creep loading at 50 kPa was greater thanin fatigue, the shapes of the r and t distributionswere the same.

The ratio t/r was a measure of strain biaxiality andis plotted against the distance from the center of thefilm in Figure 5(b) for fatigue loading after day 3. Inregion I the strain was nearly balanced biaxial witht/r 0.92. In region II, the tangential strain rapidlydecreased toward the edge, producing nearly uniaxialstrain with t/r 0.

The corresponding nominal radial and tangential

Figure 4. Change in maximum deflection (d) over timefor several experiments: (a) creep at 50 kPa, (b) fatigue,and (c) creep and fatigue.

TABLE IIPrediction of Based on Measurement of d in Fatigue and Creep Loading During Equilibrium

Method d (mm) (%)Effective modulus,

E (MPa)

Direct measure immediately after 50 kPa loading 6.22 0.08a 17.2 0.09a

Direct measure from fatigue loading and creep loading at 25 kPa 7.34 0.05a,b

FEM prediction 7.33 23.3 10.1Direct measure from creep loading at 50 kPa 8.56 0.08a,b

FEM prediction 8.55 30.3 7.7aAverage SD; n 4.bAverage deflection measured between day 3 and 12 of testing.



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strain rates ( r and t, respectively) for fatigue loadingare plotted as a function of distance from the center ofthe film in Figure 5(c). The radial and tangential strainrates were obtained by dividing the maximum strainby one-half of the loading period. Both r and t werehighest at the center of the film; r remained within85% of the maximum value in region I, whereas tdecreased to zero at the edge.

Chemical degradation

The extent of chemical degradation was measured

by the change in intensity of the normalized ether

band at 1100 cm1. The intensity of the 1100 cm1

band for the unstressed control did not vary across thefilm, and the average intensity from several controlexperiments is indicated on each plot as a solid line.Consistent with previous observations,9,12,23 the inten-sity decreased progressively from 9485 to 7555%

after 3, 6, 9, and 12 days exposure to the oxidativeH2O2/CoCl2 solution.The relative intensity of the 1100 cm1 band as

function of position on the film for creep loading at50 kPa at 3, 6, 9, and 12 days is plotted in Figure 6 forrepresentative creep experiments. At all time pointsthe change in ether intensity was constant across thefilm and showed approximately the same amount ofchemical degradation as the unstressed control. Scat-ter in the data was always present, especially at latertime points. Similar scatter was observed in the con-trol data. This scatter was attributed to inhomoge-neous degradation of the surface of the film. Regionsof higher degradation associated with formation ofsmall pits or dimples were typical of in vivo and invitro degraded PEUU films.9,12 Creep loading at 25kPa produced the same results as creep loading at 50kPa in that the ether loss was constant across the filmand was comparable in magnitude to that of the un-stressed control.

Under the creep loading conditions used in thepresent study, neither uniaxial creep near the edge norbalanced biaxial creep at the center of the film affectedthe rate of biodegradation. Although previous studieshave reported that uniaxial creep inhibits in vitro oxi-

dation of PEUU due to strain-induced crystallizationof the soft segments,10,11 the strains used were muchhigher, between 150 and 750%. In the present creepexperiments, the edge of the film experienced an ini-tial uniaxial strain of only 10%. This was much lowerthan the 150% required for the onset of strain-inducedcrystallization. Soft segment orientation at strains aslow as 50% was observed to inhibit degradation inuniaxial stress relaxation, but not in biaxial stress re-laxation.12 If orientation occurred in the fatigue andcreep experiments near the edge, it was not highenough to significantly influence degradation.

The ether intensity for a representative fatigue ex-periment is compared with the unstressed control at 3,6, 9, and 12 days in Figure 7. After 3 days of fatigue,the ether intensity was constant across the film andshowed approximately the same amount of ether lossas the unstressed control. After 6 days, the center ofthe film (region I as defined in Fig. 5) showed moreether loss than the rest of the film and showed moreether loss than the unstressed control. The edge of thefatigued film (region II in Fig. 5) showed about thesame or even less ether than the unstressed control.The difference in ether intensity between the centerand edge of the fatigued film became more distinct at

9 and 12 days.

Figure 5. FEM prediction of the strain state on the filmsurface at maximum displacement for 50 kPa loading: (a)radial (r) and tangential (t) strains, (b) biaxiality (t/r), and(c) strain rate.

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Figure 8 compares the rate of ether loss at region IIand at the center of the film in one fatigue experimentto the rate of ether loss of the unstressed control. In thecontrol, ether loss was gradual and fell to about 56% ofits initial value at day 12. In region II of the fatiguedfilm, ether loss was similar to the control. Considering

the experimental scatter, ether loss in the control andin region II of the fatigued film were not considered tobe significantly different. In the center of the fatiguefilm, ether loss was significantly greater than in thecontrol. The ether content fell to 35% of its initialvalue at day 12.

Figure 6. Percent ether remaining for unstressed controls (solid lines) and creep at 50 kPa (data points) at (a) 3 days, (b)6 days, (c) 9 days, and (d) 12 days.

Figure 7. Percent ether remaining for unstressed controls (solid lines) and fatigue tests (data points) at (a) 3 days, (b) 6 days,(c) 9 days, and (d) 12 days.



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Previous studies demonstrated the effect of soft seg-ment orientation and crystallization on the degrada-tion rate of PEUU films mechanically deformed eitherunder conditions of creep or stress relaxation. Appar-ently the imposed strains were too small for these

effects to be observed in the present creep experi-ments. However fatigue loading, especially biaxial fa-tigue, did accelerate degradation although the im-posed strains were actually smaller than in creep. Thedifference between creep and fatigue results must beattributed, at least in part, to the effect of strain rate.

The dependence of chemical degradation on strainrate can be tested easily by varying the frequency ofthe fatigue experiment. This is the subject of a futurepublication.

Physical damage

Surfaces of films exposed to constant stress (creep)or no stress (unstressed controls) were featurelesswhen viewed in the optical microscope. Films fatiguedup to 6 days were also featureless [Fig. 9(a)]. However,

cracks appeared on the surface of fatigued films atlater time points. Cracking was observed as early asday 9, although it usually did not appear until day 12or even later. The example in Figure 9(b) shows a filmthat had started to crack after 9 days of fatigue. Ini-

Figure 8. Comparison of degradation rate for fatigue testsand unstressed controls, average SD, n 4.

Figure 9. Optical micrographs of films after (a) 3 days, (b) 9 days, (c) 12 days, or (d) 18 days of fatigue in H 2O2/CoCl2.

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tially, cracking was confined to a small region in thecenter of the film. After 12 to 15 days, cracks weremore numerous, and the cracked area had spreadoutward from the center [Fig. 9(c)]. At later timepoints, 18 21 days, cracking appeared over the entiresurface of the film [Fig. 9(d)]. The crack direction in thecenter tended to show no orientation, whereas cracksnear the edge tended to show tangential orientation.

Examination in the SEM confirmed that the initiallyfeatureless surfaces remained featureless at the micronscale at day 3 [Fig. 10(a)]. However after 6 days all thesurfaces exposed to H2O2/CoCl2, including un-

stressed controls, creep and fatigue films, had devel-oped small dimples that became more numerous atday 9 and day 12 [Fig. 10(b)].

The type of loading had no effect on the size ordensity of the dimples, nor was there any difference indimpling between the center and edge of the film (Fig.11). Dimpling preceded cracking in the fatigue exper-iments. The cracks appeared to initiate at the dimples,as indicated by the arrows in Figure 11(e,f). The direc-tion of crack propagation correlated with the strainstate of the fatigued film. Nearly balanced biaxialstrain in the center (region I) caused the direction of

crack propagation to be nonpreferential. Closer to theedge (region II) where the strain state approacheduniaxial, cracks propagated in the direction perpen-dicular to the highest strain, i.e., tangentially.

Oxidative degradation of PEUU in vitro and in vivoresults in soft segment crosslinking in addition tochain scission.23 The combined effects of crosslinkingand chain scission produce a brittle surface layer onthe PEUU film. The brittle layer was previously esti-mated to be on the order of 10 m thick and crackedeasily if the degraded film was bent or stretched.9 Inthe present study, crosslinking of all the films exposedto H2O2/CoCl2, whether in creep, fatigue, or un-

stressed, was confirmed by the presence of a charac-

teristic band at 1174 cm1 in the infrared spectrum. Itis speculated that the brittle surface layer crackedunder dynamic loading conditions encountered in fa-tigue. The importance of dynamic loading is indicated by the fact that no cracks were observed in creepexperiments. Dimples facilitated crack initiation byacting as sites of stress concentration with high localstrain rate.

CONCLUSIONS

A method utilizing expansion of a diaphragm-typefilm specimen was developed to study biodegradationof PEUU under conditions of dynamic loading (fa-tigue). Experiments were also carried out withoutloading and with a constant load (creep). The strainwas not uniform over the film. A finite element modelwas used to describe the strain state, which rangedfrom uniaxial at the edges to balanced biaxial tensilestrain at the center. In the range of strains tested, creephad no effect on the rate of in vitro chemical degrada-tion of PEUU, compared with unstressed controls,

regardless of the strain state. Dynamic loading did notaffect the rate of degradation in regions of the film thatexperienced primarily uniaxial fatigue; however, deg-radation was accelerated in regions that experiencedalmost balanced biaxial fatigue. From these observa-tions, it was apparent that accelerated oxidation re-sulted from the combination of dynamic loading andbiaxial tensile strain in this system.

Chemical degradation produced a brittle surfacelayer that was marked by numerous pits and dimples.Physical damage of the surface in the form of crackingoccurred only in fatigue experiments. Cracking wasnot observed in unstressed or creep experiments.

Cracking occurred first in biaxially loaded areas of the

Figure 10. Scanning electron micrographs of (a) an untreated control film and (b) an unstressed film treated for 12 days inH2O2/CoCl2.



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specimen that experienced accelerated chemical deg-radation. Cracks initiated at dimples produced bychemical degradation and propagated in a directionthat was determined by the strain state. In biaxiallyloaded areas, there was no preferential direction ofcrack propagation. In primarily uniaxially loaded ar-eas, the cracks propagated perpendicular to the direc-

tion of maximum tensile strain.

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Figure 11. Representative scanning electron micrographs from three regions of films loaded in either creep at 50 kPa orfatigue, all after 12 days of treatment in H2O2/CoCl2.

534 WIGGINS ET AL.


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