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Research Paper

Porous poly(para-phenylene) scaffolds forload-bearing orthopedic applications

Amy L. DiRienzoa, Christopher M. Yakackib, Mareike Frensemeierc,Andreas S. Schneiderd, David L. Safranskie, Anthony J. Hoyta, Carl P. Fricka,n

aUniversity of Wyoming, Department of Mechanical Engineering, Laramie, WY 82071, USAbUniversity of Colorado Denver, Department of Mechanical Engineering, Denver, CO, USAcINM—Leibniz Institute for New Materials, Functional Surfaces Group, Saarbrücken, GermanydDevelopment and Plate-Design, AG der Dillinger Hüttenwerke, Dillingen, GermanyeMedShape Inc., Department of Research & Development, Atlanta, GA, USA

a r t i c l e i n f o

Article history:

Received 1 August 2013

Received in revised form

8 October 2013

Accepted 13 October 2013

Available online 25 October 2013

Keywords:

Poly(para-phenylenes)

Porous biomaterials

Biomedical devices

Mechanical behavior

Scanning electron microscopy

nt matter & 2013 Elsevie.1016/j.jmbbm.2013.10.012

hor. Tel.: þ1 307 766 4068frick@uwyo.edu (C.P. Fric

a b s t r a c t

The focus of this study was to fabricate and investigate the mechanical behavior of porous

poly(para-phenylene) (PPP) for potential use as a load-bearing orthopedic biomaterial. PPPs

are known to have exceptional mechanical properties due to their aromatic backbone;

however, the manufacturing and properties of PPP porous structures have not been

previously investigated. Tailored porous structures with either small (150–250 mm) or large

(420–500 mm) pore sizes were manufactured using a powder-sintering/salt-leaching tech-

nique. Porosities were systematically varied using 50 to 90 vol%. Micro-computed tomo-

graphy (mCT) and scanning electron microscopy (SEM) were used to verify an open-cell

structure and investigate pore morphology of the scaffolds. Uniaxial mechanical behavior

of solid and porous PPP samples was characterized through tensile and compressive

testing. Both modulus and strength decreased with increasing porosity and matched well

with foam theory. Porous scaffolds showed a significant decrease in strain-to-failure (o4%)

under tensile loading and experienced linear elasticity, plastic deformation, and densifica-

tion under compressive loading. Over the size ranges tested, pore size did not significantly

influence the mechanical behavior of the scaffolds on a consistent basis. These results are

discussed in regards to use of porous PPP for orthopedic applications and a prototype

porous interbody fusion cage is presented.

& 2013 Elsevier Ltd. All rights reserved.

r Ltd. All rights reserved.

.k).

1. Introduction

Poly(para-phenylenes) (PPPs) have been reported to have thehighest mechanical properties of any neat thermoplasticpolymer with modulus and tensile strength values as highas 8.7 GPa and 207 MPa, respectively (Morgan et al., 2006; Pei

and Friedrich, 2012; Vuorinen et al., 2008). Their backbonestructure consists of a direct linkage of repeating aromaticrings, which leads to exceptional mechanical strength andstiffness by providing strong anti-rotational biaryl bonds.Furthermore, the addition of side chains can help to limitthe chain mobility via steric hindrance as well as to increase

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molecular weight during processing. Other polymer systemscomposed of highly aromatic backbones, such as poly(phe-nylene sulfide) (PPS) or poly(etheretherketone) (PEEK), areconsidered as gold standards for applications where a com-bination of high strength, modulus, and stability at elevatedtemperatures is required; however, PPPs have higher strengthand modulus values compared to these systems. In addition,PPPs can be melt-processed at lower temperatures than PEEKand can be solution cast (Morgan et al., 2006).

Despite their outstanding material behavior, only a hand-ful of studies have investigated the potential use of PPPs as abiomaterial (Burstone et al., 2011; Vuorinen et al., 2008, 2011;Frick et al., in press). This is in large part because sufficientlyhigh molecular weight PPP suitable for mechanical loadinghas not been readily available until recently. These prelimin-ary studies have been promising in that they show noobvious cytotoxicity (Frick et al., in press) as well as stablemechanical behavior when submerged in an aqueous solu-tion (Vuorinen et al., 2011; Frick et al., in press). For load-bearing orthopedic applications, a polymer material withhigh strength, stiffness, and toughness to resist failure wouldbe advantageous over currently used materials. Metallicbiomaterials (e.g. cobalt–chrome, titanium, and stainlesssteel) have elastic modulus values roughly one to two ordersof magnitude higher than bone and as a result can create alocalized mechanical mismatch (i.e. stress shielding) (Bobynet al., 1992). Metallic devices are also incompatible withseveral standard imaging techniques (Shellock et al., 1992;Suh et al., 1998), and are difficult to remove if a revisionprocedure is necessary (Brand et al., 2000). Devices madefrom polymer materials are often radiolucent for X-rayimaging, compatible with magnetic resonance imaging(MRI), and can simply be drilled out by conventional equip-ment in a revision procedure if necessary. However, nearly allbio-polymers have strength/stiffness values well below PPP(Frick et al., in press). Furthermore, biodegradable materials,which are designed to degrade during the healing processand return the injury site back to its original anatomy,experience a fast and substantial decrease in mechanicalproperties due to the inherent nature of degradation (Roseand Oreffo, 2002). For example, it has been shown that thetensile load to failure of polyglyconate suture anchorsdecreases by 75% after 3 weeks in vivo (Demirhan et al., 2000).

In addition to appropriate mechanical properties, soft-tissue fixation devices must maintain a strong implant-bone interface; failure caused by loosing of this interface iscaused by three main factors: (a) bone resorption (osteolysis)caused by wear debris, (b) mechanical mismatch between theimplant and surrounding bone (stress-shielding), and (c)displacement between the implant and bone tissue (Slotenet al., 1998). Several studies have investigated porousscaffold materials for use in orthopedic applications includ-ing bone ingrowth scaffolds for permanent implant fixation(Agrawal and Ray, 2001; Causa et al., 2006; Converse et al.,2010; Converse et al., 2009; Hench, 1991; Karageorgiouand Kaplan, 2005; Kretlow and Mikos, 2007; Rezwan et al.,2006). The fundamental premise is that during healing theosteoblast cells will penetrate and proliferate into the open-cell porous scaffold, allowing for osteointegration (a physicalintermix of bone and porous scaffold upon healing).

Additionally, porous scaffolds could be used as a vehicle fortargeted delivery of cytokines such as bone morphogenicproteins (BMPs) (Chen et al., 2004; Wozney, 1989). Severalcommercialized bone ingrowth scaffolds are currently avail-able, such as sintered beads, wired meshes, and metallicfoams (Ryan et al., 2006). However, it is important to note thattraditional polymers are not widely used as load-bearingorthopedic scaffolds as they do not have sufficient mechan-ical properties.

In this study, we show that PPP can be used to create aporous, load-bearing biomaterial, which matches the stiff-ness of trabecular bone. Porous PPP samples were formed byhot-press powder sintering of a PPP/sodium–chloride mixturefollowed by particle leaching. This method was chosen for itsversatility in controlling porosity and pore size, via sodiumchloride volume fraction and size. Open-cell PPP scaffoldsranging in porosity from 50–90 vol% and pore sizes150–500 mm were chosen for the investigation to cover opti-mal ranges for osteointegration as suggested by previousstudies (Hulbert et al., 1970; Karageorgiou and Kaplan, 2005).Micro-computed tomography (mCT) imaging was used toconfirm an interconnected open-cell structure and that allsalt particles were successfully leached out. The morphologyof the scaffolds was characterized using scanning electronmicroscopy (SEM) imaging. PPP scaffolds were tested intension and compression to determine the effect of porosityand pore size on the mechanical behavior, and experimentalresults were compared to well-established foam theory. Theevolution of the porous structure post deformation wasinvestigated using SEM imaging to examine the progressionand extent of damage. To the authors’ knowledge, this studyis the first to manufacture and characterize porous PPP.

2. Experimental methods

2.1. Materials

PrimoSpire PR-250 powder was used as the PPP in this studyand was provided by Solvay Specialty Polymers, Inc (Alphar-etta, GA). The material was used in its as-received conditionwithout any further modification. The chemical structure hasbeen confirmed to be an aromatic backbone with aromaticside groups, shown in other work (Frick et al., in press).

2.2. Solid PPP Samples

Solid PPP samples used for tensile testing were formed bycompression molding using a hydraulic high-temperaturepress (Model DV-62-422, Pasadena Hydrualics, Inc.). Amountsranging from 4–6 g of PPP powder were pressed into soliddisks at a temperature of 250 1C, a force of 1334 N, and asintering time of 10 min. These disks were formed betweenflat aluminum molds wrapped in non-stick aluminum foilmeasuring approximately 12 cm in diameter and 0.5–0.8 mmin thickness. Dog-bone shaped samples for tensile testingwere cut from the flat disks described above to cross-sectionsof 2.9–3.2 mm in width and gauge lengths of 12–15 mm.Specimens were lightly sanded with 600-grit sandpaper toensure consistent sample surfaces. Compression samples

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were machined to 3.18�3.18�6.35 mm3 rectangles frominjection-molded plaques. Two sets of compressive sampleswere machined orthogonally from each other to verify therewas no directional dependence due to chain alignment fromthe manufacturing of the plaques.

2.3. Porous PPP samples

Porous PPP scaffolds were manufactured by powder sinteringand particle leaching. Sodium chloride crystals (Sigma Aldrich)were sifted into two different sizes: small (150–250 mm) andlarge (420–500 mm). Sodium chloride crystals and PPP powderwere massed and mixed thoroughly by hand to createmixtures containing 50–90 vol% salt crystals using largeand small crystals separately. Disks for tensile samplesmeasuring 10 cm in diameter and 0.7–1.8 mm in thicknesswere formed by pressing 12–20 g of salt/powder mixturesbetween flat aluminum molds covered in non-stick alumi-num foil using previously mentioned hydraulic high-temperature press at a temperature of 250 1C, a force of2669 N, and sintering times ranging from 60–120 min; sam-ples containing higher percentages of salt required the longersintering times. Samples for tensile testing were machined asdescribed above. Compression samples approximately15.5 mm in diameter and 15.3 mm high were formed in acylindrical aluminum mold by pressing the salt/powdermixtures in the hydraulic high-temperature press at a tem-perature of 250 1C, a force of 445 N, and sintering timesranging from 60–120 min depending on salt content. Com-pression samples were faced and turned down to an averagediameter of 9.5 mm on a lathe. The as machined sampleswere submerged in distilled water and agitated on a shakerplate at 60 rpm heated to 90 1C. The distilled water waschanged daily for 7–10 days. This time frame was chosen toensure that all sodium chloride was leached out of thesamples. The resulting porous samples were dried in avacuum oven at 90 1C for 24 h. Compression samples dimen-sions were measured prior to leaching and massed afterdrying. The actual porosity was determined by calculatingthe fraction of the scaffold composed of PPP:

Porosity¼ 1� ML

ρπr2Lð1Þ

where ML is the mass after leaching, ρ is the density of PPPtaken to be 1.19 g/cm3, r is radius and L is the length of thecylindrical sample.

2.4. Micro-computed tomography (μCT)

The μCT scans were performed using a Siemens Inveon X-raymicrotomography machine. Representative samples of eachporosity and pore size were scanned at a voxel resolution of21 mm. The 2-D tomograms were thresholded using Image Jand assembled into 3-D binarized images using SkyScan.Both 2-D tomograms and 3-D reconstructions were used tovalidate the open pore structure of the materials and com-plete leaching of the sodium chloride crystals. Morphometricparameters such as volume fraction, average strut thickness,average strut spacing, and degree of anisotropy were

determined using direct distance transformation methods(Hildebrand et al., 1999).

2.5. Mechanical testing

Uniaxial monotonic tensile and compression testing wasconducted on a MTS hydraulic load frame (858 Mini BionixII, MTS) equipped with a laser extensometer (LX 500, MTS) ata displacement rate of 0.01 mm/s. Reflective tape was placedon the gauge length of the tensile samples and the strain wasdefined as the ratio of the displacement from the laserextensometer to the initial gauge length. For compressivetests, reflective tape was placed on the compression platensand the width of the tape was subtracted from the displace-ment measured by the laser extensometer before calculatingthe strain. Stress was calculated by dividing the force by theinitial cross sectional area. Tensile tests were run to failure(defined by fracture) on dog-bone samples ranging from 50 to80 vol% using four different samples at least (n¼4) at eachporosity. Compression tests were run well into the third stageof compression (as fracture was not observed) for cylindricalsamples ranging from 50–90 vol% porosity using four to sixsamples at each porosity. Additionally, four samples of 50, 70,and 90 vol% porosity were tested to intermediate strains(1 within the first stage, 2 within the second stage, and1 within the third stage of compression). The modulus wasdetermined by finding the slope of the linear fit to the initialelastic region. Tensile strength was defined as the fracturestress in tension and the maximum stress at the first yieldpoint in compression. Compressive yield strength wasdefined as the intersection of the linear fit to the initialelastic region and the linear fit to the initial plastic region.

2.6. Scanning electron microscopy (SEM)

Scanning electron microscopy images were taken along thelength of the samples using a FEI Versa 3D™ DualBeam™

microscope. Specimens were investigated without any coat-ing in low vacuum mode at a base pressure of 100 Pa. Imageswere taken using a low vacuum secondary electron detectorat a voltage of 10 kV, current of 47 pA, and a working distanceof approximately 10 mm.

2.7. Statistical analysis

Statistically significant differences were determined by usinga Student's t-test with an alpha value of 0.05. The measuredphysical and mechanical properties were compared at eachporosity level between the two pore sizes to investigate theinfluence of pore size on the samples.

3. Results

3.1. Stress–strain behavior of solid PPP

Representative uniaxial mechanical behavior of solid PPP isshown in Fig. 1. Initial tensile loading demonstrated nearlinear elastic behavior with an average elastic modulus of5.0 GPa followed by plastic yielding at an average tensile

Fig. 2 – Representative lCT scan from the center section of a90% porous PPP scaffold with large (420–500 lm) pore size.The lCT imaging was used to confirm complete leaching ofsodium chloride particles and an interconnected open-cellstructure for all of the samples tested.

Fig. 1 – Tensile and compressive mechanical behavior ofsolid poly(para-phenylene) (PPP).

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strength of 141 MPa. Subsequent to yielding, the stressdropped approximately 20 MPa prior to reaching a flowstress plateau. Post-yielding softening is commonly observedfor thermoplastics and is associated with local segmentalrotation. With increasing strain, local polymer chains becomeincreasingly aligned which leads to the mild hardening slope.On average, the tensile strain-to-failure of PPP is approxi-mately 17% strain and fracture strength is 120 MPa. Incompression, PPP demonstrated identical linear elastic beha-vior but experienced yielding and flow at higher stress values.The compressive strength of PPP was measured at 168 MPa. Itis important to note that compressive samples shown did notfracture after 30% strain, but rather experienced shearingsuch that the top and bottom of the samples were notvertically aligned.

3.2. lCT scans

The μCT scans were performed on representative samples ateach porosity and pore size. The scans validated that thespecimens were comprised of an interconnected open-cellstructure. Density was also calculated after drying to ensurethat all sodium chloride was being leached from the scaf-folds. The μCT images confirmed that no salt crystalsremained present in the structure. A 3-D reconstruction ofthe internal cross-section of a 90% porous sample with largepore size is shown in Fig. 2. For this sample, μCT analysisverified the designated microstructure of the material. Por-osity was measured as 91.6%, while the average strut thick-ness and spacing was measured as 96 and 506 μm,respectively. Degree of anisotropy was 1.1, verifying that nodirectional dependence was created during the sintering–leaching process.

3.3. SEM imaging

The microarchitecture was also analyzed through SEM ima-ging. A comparison of the two pore sizes for each porosity isdemonstrated in Fig. 3. Images of these representative sam-ples reveal that the pores appear to be well distributed, whilethe size and shape of the pores are characteristic of the salt

crystals used. For both pore sizes, the porosity and the strutthickness of the respective sample can be qualitativelycharacterized. As porosity increases, the number of poresincrease and the separation between the pores decrease.

However, careful inspection of SEM images of the50–70 vol% porous samples reveal pores estimated to be overan order of magnitude smaller than the salt particles. Thesesmall-scale pores are likely places where the PPP did notpenetrate during sintering, such as a location where the saltcrystals were in direct contact. Comparing actual porositymeasurements to the initial mixtures, small-scale pores wereestimated to be between 0.2 and 0.8 vol%. No obvious correla-tion existed between small-scale pore volume fraction andvolume fraction of salt, salt crystal size, or sintering time. Thepresence of the small-scale pores is thought to aid in particleleaching (demonstrated by the complete leaching of 50 vol%porous samples). In 80 and 90 vol% porous samples, poresbegin to coalesce and disconnected struts are observedthroughout the structure, regardless of pore size.

3.4. Stress–strain behavior of porous PPP

Representative tensile strain-to-failure results for each por-osity and pore size are shown in Fig. 4 and average values ofthe mechanical behavior are summarized in Table 1. It shouldbe noted that the 90 vol% samples were not tested in tensiondue to the fragility of the cross-section over the gauge length.The tensile performance of the porous scaffolds does notfollow the same behavior as the solid material, as thesamples experienced little-to-no plasticity and low ductility.The modulus and tensile strength of the samples decreasedas porosity increased. When comparing the small (150–250 mm)

Fig. 3 – SEM images of PPP porous scaffolds. Small (150–250 lm) and large (420–500 lm) pore size ranges and increasingporosity (50–90 vol%) were compared.

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and large (420–500 mm) pore sizes at each porosity level, poresize did not consistently influence the mechanical propertiesof the samples, as the values were typically within onestandard deviation of each other. A statistically significantdifference was observed in modulus for the 50 and 80 vol%samples with different pore sizes, and in tensile strengthfor 50, 60, and 80 vol% samples with different pore sizes;

however, these differences were attributed to random batchvariation. There were no significant differences in strain-to-failure when investigating the influence of pore sizes at eachporosity.

Representative compressive stress–strain results for PPPsamples of each porosity and pore size are shown in Fig. 5and average values of the porosity and mechanical behavior

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are summarized in Table 2. The compressive performance ofthe porous scaffolds resembles foam theory, as the samplesexperience three stages: linear elasticity, plastic yielding, anddensification (Liu and Subhash, 2004). Similar to the tensilesamples investigated, modulus and strength decreased withincreasing porosity with no clear influence of pore size.

Fig. 4 – Tensile strain-to-failure behavior for PPP porousscaffolds as a function of porosity and pore size.

Table 1 – Elastic modulus, tensile strength, and strain-to-failuand one standard deviation (n designates statistical differencedistribution with poα¼0.05).

Sample n Tensile modulus (MPa)

50 vol%—150–250 mm 4 9707128n

50 vol%—420–500 mm 4 6417125n

60 vol%—150–250 mm 4 49272460 vol%—420–500 mm 4 411710570 vol%—150–250 mm 4 25575570 vol%—420–500 mm 4 28175880 vol%—150–250 mm 4 53715n

80 vol%—420–500 mm 5 122724n

90 vol%—150–250 mm – –

90 vol%—420–500 mm – –

Fig. 5 – Compressive stress-strain behavior of PPP porousscaffolds as a function of porosity and pore size.

The stress required to cause plastic deformation (2nd stage)and pore densification (3rd stage) decreased as porosityincreased (i.e. a much higher stress is needed to compress a50 vol% porous sample to 40% strain than for a 90 vol% poroussample). Specimens were tested until compressive behaviorclearly reached the 3rd stage. After testing all samples wereintact, with the exception of the 90 vol% porous samples withlarge pores, which crumbled when removed from the com-pression platens. The average measured porosity of thesamples matched within 3 vol% of the desired manufacturedvalue of porosity. It should be noted that the 90 vol% com-pressive samples were the only samples that demonstrated astatistical difference in porosity between pore sizes, whichhelps explain why the mechanical differences between thetwo pore sizes are so dramatic (i.e. the 90 vol%—150–250 mmsamples had average mechanical properties over two timeslarger than their 420–500 mm counterparts).

Fig. 6 follows the progression of damage in three 50 vol%porous large pore size samples loaded to various strains. Thestress–strain graphs shown in Fig. 6a illustrate that eachsample was loaded into a different stage of deformation. SEMimage in Fig. 6b was taken of a sample tested to 2% strain.This deformation falls within the 1st stage of compressionand corresponds to linear elastic behavior. No clear damage isobserved at this stage. Fig. 6c was tested to 9% strain. At 9%strain the material is in the 2nd stage of compression andexperienced plastic deformation. Strut buckling and fractureare observed in discrete locations along the sample surface.At 40% strain the sample has reached the third stage ofcompression (Fig. 6d). At this stage pores have readilycollapsed and cracks have formed between pores. Thisbehavior is closer to that of metal foams than traditionalpolymer foams that are elastomeric and experience highstrains before plastic deformation occurs.

Analogous to Fig. 6, Fig. 7 follows the damage progressionof 90 vol% porous large pore size samples. At this porosityonly a small linear elastic stage exists (o0.05% strain), andwas therefore not investigated. The SEM image shown inFig. 7b is of a sample tested to 23% strain. Even though thesample has been compressed well into the 2nd stage, it isdifficult to discern damage due to loading since the structurecontains disconnected struts prior to testing. After testing to61% strain (Fig. 7c), it is apparent that struts have buckled andthe material has entered the densification range. At 74%

re of porous scaffolds tested. Values represent the averagebetween samples of equal porosity but different pore-size

Tensile strength (MPa) Strain-to-failure (%)

1971.6n 2.870.41273.3n 3.170.31170.4n 3.770.57.271.5n 3.070.85.970.7 3.570.24.870.7 3.371.21.570.3n 3.870.62.470.6n 3.271.0– –

– –

Table 2 –Measured porosity, elastic modulus, and compressive strength for porous scaffolds tested in compression. Valuesrepresent the average and one standard deviation. (n designates statistical difference between samples of equal porositybut different pore-size distribution with poα¼0.05).

Sample n Measured porosity (vol%) Compressive modulus (MPa) Compressive strength (MPa)

50 vol%—150–250 mm 4 51.6%70.8% 1017776 22.475.350 vol%—420–500 mm 5 52.6%72.5% 964789 26.174.760 vol%—150–250 mm 4 59.4%72.3% 5847123 15.172.860 vol%—420–500 mm 4 59.7%74.1% 7307154 17.173.970 vol%—150–250 mm 6 71.0%70.6% 296727n 7.470.5n

70 vol%—420–500 mm 6 71.4%71.4% 374738n 9.370.9n

80 vol%—150–250 mm 4 80.4%70.6% 79719 1.770.4n

80 vol%—420–500 mm 4 81.0%70.3% 127732 2.770.4n

90 vol%—150–250 mm 6 87.9%70.2%n 2174.0n 0.4870.10n

90 vol%—420–500 mm 6 90.8%70.3%n 8.876.3n 0.2170.16n

Fig. 6 – Stress–strain behavior and SEM images illustrating the progression of damage in 50% porous 420–500 lm pore sizecompression samples at intermediate strains.

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strain (Fig. 7d), the material continues to accumulate damage.At this stage the sample is still whole and has not yet begunto crumble.

Fig. 8 compares the experimental modulus of all samplestested to the theoretical modulus predicted from rule ofmixtures, an analytical model of the mechanical behavior ofporous materials

E� ¼ Esð1�Vf Þ2 ð2Þ

where E� and Es are the elastic moduli of the porous scaffoldand the material without porosity, respectively, and Vf is thepore volume fraction ( Ji et al., 2006). Eq. (2) is similar to manyclassic analytical cellular models (e.g. Gibson and Ashby(1982), Ramakrishnan and Arunachalam (1993), and Waghet al. (1991)) and fit to experimental data taken from Gibsonand Ashby (1982), Lederman (1971), Ramakrishnan andArunachalam (1990). In general, for 50–70 vol% porosity,

experimental results closely matched theory. Compressivetests showedmore concentrated modulus values closer to thetheoretical modulus than tensile tests; however, experimen-tal results were always at or below the theoretical curve. Asporosity increased to 80 and 90 vol%, experimental resultsbegin to increasingly diverge from the theoretical curve.Foam theory cannot accurately predict the modulus of scaf-folds with porosities of 80 vol% and above since a traditionalfoam structure no longer exists.

The compressive yield strength was compared to Gibsonand Ashby's model for strength of open-cell foams (Arora et al.,1998; Gibson and Ashby 1988)

sn

y ¼ 0:23 syð1�Vf Þ3=2 1þ ð1�Vf Þ1=2h i

ð3Þ

where sny is the yield strength of the foam, sy is the yieldstrength of the solid material, ρn is the density of the foam and

Fig. 7 – Stress-strain behavior and SEM images focusing on the progression of damage in 90 vol% porous 420–500 lm pore sizecompression samples at intermediate strains.

Fig. 8 – Experimental modulus for all porous scaffolds testedin this study relative to theoretical modulus for PPPpredicted from rule of mixtures (Eq. (2)).

Fig. 9 – Experimental yield strength for porous scaffoldstested in compression relative to theoretical compressivestrength for open-cell foams predicted by Gibson and Ashby(Eq. (3)).

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ρs is the density of the solid material. It is important to notethat Eq. (3) is only applicable for 70%rVfr96%. For theprediction, the density of the porous material was assumedto be directly proportional to the volume percentage of PPP(i.e. a 70 vol% porous scaffold possesses a density of 30 vol% ofthat of the solid material). This prediction closely matches theexperimental yield strength behavior of the PPP porous scaf-folds as shown in Fig. 9.

4. Discussion

The goal of this study was to fabricate and mechanicallycharacterize PPP as a porous polymer scaffold for potential

use in orthopedic devices. We hypothesized that due to theinherent strength of PPP a porous polymer scaffold can becreated for load-bearing applications. Our results show that apowder-sintering/salt-leaching technique can be used tocreate porous PPP scaffolds with tailored porosity and poresize, which influence the mechanical properties of the scaf-folds. The sintering/leaching method showed consistent andrepeatable results, and has been used successfully in otheraromatic systems, such as PEEK (Converse et al., 2010).

The size, shape, and interconnectivity of porous scaffoldshave a significant influence on the potential osteointegrationof a material. Due to the wide range of potential uses andtypes of biomaterials, ideal pore volume fractions and sizes

Fig. 10 – A prototype PPP interbody fusion cage is comparedto a PEEK cage. The PPP cage wasmanufactured from 75 vol%porous material.

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are not precisely known; however, porous scaffolds shouldmatch the mechanical properties of the surrounding bone,and maximize porosity for enhanced osteogenesis and boneingrowth (Karageorgiou and Kaplan, 2005). Previous researchindicated that the minimum pore size should be 100 μm,(Hulbert et al., 1970), although subsequent studies demon-strated improved osteogenesis with pores sizes larger than300 μm due to higher vascularization and oxygenation(Karageorgiou and Kaplan, 2005). Interconnected porosity(i.e. an open-cell structure) is necessary for healthy cellularfunction (i.e., attachment, proliferation, and migration), diffu-sion of nutrients, and waste removal. While the optimalporosity and pore size for osteointegration may be debatable,the proposed manufacturing technique presented here showsthat PPP scaffolds can be created with a wide range ofmicrostructural parameters that can be tailored to possiblyoptimize osteointegration. Furthermore, mechanical testingof the PPP scaffolds (Figs. 4 and 5) showed that pore size hadrelatively little influence on mechanical behavior comparedto porosity.

There is extensive literature on the mechanical behaviorof porous structures both from an experimental (Brezny andGreen, 1989; Gupta et al., 2001; McIntyre and Anderton, 1979;San Marchi and Mortensen, 2001; Sun and Mark, 2002; Wanget al., 2007) and modeling (Di Prima et al., 2010; Gibson, 1989;Meinecke and Schwaber, 1970; Rusch, 1969, 1970; Sherwoodand Frost, 1992; Walter et al., 2009) standpoint. In compres-sion, porous materials experience three stages termed linearelasticity (i.e. bending), plastic yielding (i.e. buckling), anddensification. In the initial elastic regime cell wall “struts”deform by uniaxial loading or bending, although localizedplasticity has been shown to occur (Di Prima et al., 2007;Gibson, 2000). Further deformation causes the struts to buckleat discrete locations, and consequently the material under-goes plastic deformation or localized fracture resulting in aglobal stress plateau. Eventually, upon continued collapse ofthe cells, the structure is compacted and the stress–straincurve turns up toward higher stress at large strain. For theporous PPP tested in this study SEM imaging also verified thethree stages of compressive behavior in the PPP scaffoldsshown in Figs. 6 and 7. Notably, the porous structure had atendency to form cracks at strains beyond the yield stress. Incontrast to compression, tensile behavior of porous struc-tures had much lower failure-strength and -strain values dueto crack formation and propagation, associated with thepores themselves (Fowlkes, 1974; Kabir et al., 2006). Mechan-ical testing of the PPP scaffolds verifies that the scaffolds areless ductile with lower failure stresses when tested in tensioncompared to compression (Figs. 4 and 5).

The differences in tensile versus compressive behaviorshould be kept in mind when proposing porous materials forload-bearing applications. For example, in lumbar spinalfusions, stabilization rods are often used in conjunction withinterbody fusion cages to promote fusion (DiPaola andMolinari, 2008; Noshchenko et al., 2011a, 2011b). While thecages are placed in between the vertebrae and subjectedprimarily to compression, the rods are connected to pediclescrews and can experience tensile and shearing stresses. Theproposed porous PPP materials should likely be used in areassubjected primarily to compression, such as in the interbody

fusion cage, as the material has the ability to adapt to anoverload in stress through the densification process and helpavoid catastrophic failure. A prototype fusion cage made from75 vol% porous PPP was created to match the overall dimen-sions of a PEEK implant (Fig. 10). PPP is advantageous overPEEK from a manufacturing perspective, in that it has a lowermelting temperature, and its amorphous structure does notrequire careful processing to ensure a semi-crystalline struc-ture necessary for good mechanical behavior of PEEK. PorousPPP may unlock new design strategies for fusion cages, asthere is a need for better distribution of stresses across thevertebral endplates to reduce subsidence (Beutler andPeppelman, 2003), while still allowing the bone to fusethrough the implant. Future studies are needed to explorethe design and use of porous PPP implants.

The modulus of trabecular bone has been reported over awide range of 10 to 1000 MPa (Goulet et al., 1994; Liu et al.,2006; Poukalova et al., 2010). Factors such as type of bone(i.e. humerus vs. femur), age of patient, and history of disease(osteoporosis) can all contribute to the variability of bonemodulus; however, this study demonstrates that porous PPPcan be tailored from 50–90 vol% porosity to match the widerange of mechanical properties of trabecular bone. Comparedto traditional biopolymers, the initially high starting modulus(5 GPa) and yield strength (141–168 MPa) of solid PPP allows itto maintain a higher strength and stiffness with increasingporosity. A polymer with a more conventional solid modulusof 1 GPa can only be made 68 vol% porous before modulusdrops below 100 MPa according to the rule of mixtures. Bycomparison, PPP is able to maintain a modulus of 100 MPa at80 vol% porosity (Fig. 8). However, it is important to note thatthe modulus of porous PPP deviates significantly from foamtheory at porosities of 80 and 90 vol% due to disconnectedstruts observed in Fig. 3. This observation is an inherentlimitation of the particle leaching technique. For instance, thehighest theoretical packing density of spherical particles is74 vol%. While the cubic geometry of salt crystals allows for ahigher packing density, it results in many disconnected strutsabove 80 vol% porosity.

To date, no products made from PPP have been submittedfor approval by the FDA for biomedical use; however, PPP ischemically similar to PEEK, which has been successfully usedin numerous FDA approved biomedical devices (Kurtz andDevine, 2007). PPP has good chemical, thermal, and radiation

j o u r n a l o f t h e m e c h a n i c a l b e h a v i o r o f b i o m e d i c a l m a t e r i a l s 3 0 ( 2 0 1 4 ) 3 4 7 – 3 5 7356

stability, which is considered beneficial for sterilization andpost-irradiation aging for biomedical device use (Kurtz andDevine, 2007). Furthermore, PPP has been shown to be non-cytotoxic (Frick et al., in press). An additional critical issue inload-bearing polymers in vivo is their loss of mechanicalproperties due to water absorption (Drummond, 2008; Janand Grzegorz, 2005). PPP demonstrates only 0.7 wt% waterabsorption, and a minimal reduction in modulus (Frick et al.,in press). Low absorption is due to its bulky side groups,which act as diffusion barriers and prevent water moleculesfrom entering the network and bonding with hydrophiliccomponents (Barnes et al., 1988; Corkhill et al., 1987). How-ever, it is necessary to perform more tests to prove PPP isappropriate for orthopedic devices.

5. Conclusions

Tailored porous PPP scaffolds can be manufactured using asintering and leaching technique. Open-celled structureswere created as low as 50 vol% porosity, although many strutsbecame disconnected when porosity reached 80–90 vol%.The addition of porosity severely restricted the strain-to-failure of the PPP material in tension to less than 4%, whilethe compressive behavior of porous PPP experienced linearelasticity, plastic yielding, and densification. Experimentallymeasured modulus and strength values of the porous PPPscaffolds in both compression and tension matched wellwith theoretical predictions. Pore size did not consistentlyinfluence the mechanical behavior of the scaffolds, whichwas primarily influenced by porosity, over the ranges testedin this study.

Acknowledgements

The authors would like to thank Prof. Robert Guldberg, Prof.R. Dana Carpenter, and Angela Lin for their help regarding μCTimaging and analysis using direct transformation methods. Theauthors gratefully acknowledge assistance in porous samplefabrication and mechanical testing from Jonathan Schlotthauerand Dustin Bales. The authors would also like to thank EricJ. Losty and Jac Corless for their help in machining prototypeimplants. CPF and CMY acknowledge Solvay Specialty Poly-mers, LLC for their financial and material support of this work.

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