02 Classes of Materials - Chapter_Mow_Huiskes Book

A. El-Ghannam, P. Ducheyne Biomaterials In: Basic Orthopaedic Biomechanics and Mechanobiology Eds., Van C. Mow and Rik Huiskes, Lippincott-Williams and Wilkens, 2005

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LI005-11 Mow and Huskies August 14, 2004 17:34

11

BIOMATERIALSAHMED EL-GHANNAM

PAUL DUCHEYNE

1 THE MATERIALS SCIENCE TRIAD 00

2 CLASSES OF MATERIALS 002.1 Metals 002.2 Ceramics 002.3 Polymers 002.3 Composites 00

3 CRYSTAL STRUCTURE 003.1 Metals and Ceramics 003.2 Polymers 00

4 MECHANICAL BEHAVIOR 004.1 Mechanical Properties 004.2 Fracture Mechanics of Ceramics 004.3 Elastic Properties of Composites 00

5 WEAR RESISTANCE 005.1 Ultrahigh-molecular-weight

Polyethylene 005.2 Alumina 00

6 METALLIC BIOMATERIALS—CORROSION RESISTANCE AND RELATEDPROPERTIES 00

7 BONE BIOACTIVE CERAMICS 007.1 Calcium Phosphates—Structure,

Mechanical Properties, and Processing 007.2 Calcium Phosphates—Biological

Behavior 007.3 Calcium Phosphates—Coatings 007.4 Bioactive Glasses 007.5 Bioactive Glass Ceramics 00

8 TISSUE ENGINEERING AND RECENTDEVELOPMENTS 008.1 Bone Tissue Engineering 008.2 Engineering Surfaces with Peptides 00

1 THE MATERIALS SCIENCE TRIAD

Biomaterials including ceramics, metals, andpolymers have been widely used in surgery.Some biomaterials with relatively high mechani-cal strength such as the titanium alloy Ti-6Al-4V,the Co-Cr alloys, alumina, or zirconia have beenused in joint replacement devices. Others withbioactivity and bone bonding ability such as cal-cium phosphate ceramics and bioactive glasseshave been used as artificial bone grafts.

The principle that is of most value with re-spect to assessing properties of a material origi-nates from the internal structure of that material.At an ultramicroscopic and light microscopiclevel, materials are made up of atoms which

are associated with their neighbors in crystals,molecules, multiphase arrangements, and mi-crostructures. In this chapter we shall focus ex-tensively on structures and how they affect thebehavior of the biomaterial.

Biomaterials must be processed to meet therequirements of a specific biological applica-tion. The most familiar processing steps simplychange the shape of the materials by machiningor forging. However, processing commonly in-volves more than simply changing the shape ofthe material. For example, thermal processing isrequired to coat metallic implants with a calciumphosphate surface intended to enhance fixationand hence longevity of the implant. During coat-ing the internal structure of the material may be

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altered, in its turn resulting in a change of bi-ological properties. It is obvious then that theprocessing-induced structural changes must beunderstood, lest it be impossible to formulateappropriate processing steps.

The relationship among structure, properties,and processing is called the materials sciencetriad. This concept is all-important in bioma-terial science, no matter which material is con-sidered. Thus in this chapter we will always in-dicate properties as a function of structure andprocessing.

2 CLASSES OF MATERIALS

As is the case for materials, biomaterials can beconveniently grouped into three classes: metals,polymers, and ceramics. This scheme is basedprimarily on the chemical makeup and atomicstructure of the various materials. Composites,representing another group of materials, consistof combinations of two or more metals, poly-mers, or ceramics. An alternative classificationscheme relates to the biological properties of bio-materials, including bioinert or bioactive. In thischapter we do not use this method to organizematerial classes.

2.1 Metals

Metal atoms are held together in a crystal lat-tice by the interaction of the valence electrons(outermost electrons) with the positive metal-lic ions. These nonlocalized electrons are free tomove throughout the solid, because the valenceelectrons are not tightly bound to the metal ions.Such an arrangement is called the metallic bond.This distinguishes it from the covalent and ionicbonds present in polymers and ceramics, respec-tively. In these other configurations, electronsare not free to roam (Fig. 11-1). In what follows,Fig.11.1some illustrations of the importance of the typeof bond vis-a-vis biomaterials behavior are given.

The independent electrons in the metallicbonds can quickly transfer electric charge andthermal energy. As a result of the different bind-ing states of the electrons, there is a differencein thermal expansion coefficient between met-

FIGURE 11-1. Schematic representation of cova-lent, metallic, and ionic bonding.

als and ceramics. Since it is often desirable tocoat metallic implants with a bioactive ceramicfilm in order to improve implant fixation, the

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difference in thermal properties between metalsand ceramics is an important issue, as this differ-ence may be at the basis of destabilization of thecoating. Specifically, the difference in the ther-mal expansion coefficient between metals andceramics results in interfacial shear stresses thatcan create microcracks at the interface duringcooling subsequent to plasma-spraying a calciumphosphate coating onto a metallic implant. Af-ter implantation, preferential degradation of theceramic at the ceramic–metal interface can resultin removal of the coating layer and hence, poorfixation.

As the electrons can move easily in metals,neighboring atoms are weakly bonded, makingmetals easily deformable. This facilitates wear,except if special processing techniques are em-ployed with the intent to enhance bonding inthe metallic structure. Wear of a metallic im-plant can result from micromotion and frictionin the implantation bed against a counterpartthat has the same or higher hardness. Accumu-lation of wear debris has been related to asepticloosening and implant failure.

As the number of valence electrons increasesdown the periodic table, they become more



11. Biomaterials 3

localized and the metallic bonds become moredirectional, resulting in more brittle metals. Thisproperty is also related to the high melting tem-perature of such brittle metals as tungsten. Asthe mechanical properties (and also the chemicaland physical properties) of metals can be im-proved by alloying, most metals used in ortho-pedic surgery are alloyed. Obvious examples arethe alloys based either on titanium or on cobalt.

2.2 Ceramics

Ceramics are composed of compounds made ofmetallic and nonmetallic elements. With non-metallic elements typically oxygen, nitrogen, orcarbon, ceramics are frequently oxides, nitrides,or carbides. Ceramic biomaterials include cal-cium phosphate ceramics, bioactive glasses,bioactive glass ceramics, alumina, and zirconia.Each of these materials is relatively hard and brit-tle. In fact, hardness and brittleness are typicalproperties of ceramics, as are excellent resis-tance to high temperatures and corrosive envi-ronments. The basis of these characteristics isagain related to the type of atomic bonds. Themetallic elements release their outermost elec-trons to the nonmetallic atoms. This directionalbond produces electrons that are localized inthe structure. Thus, typical ceramic materials aregood insulators, both electrically and thermally.Considerable energy is usually required to sepa-rate the atoms that have lost electrons (cations)from the atoms that have gained electrons (an-ions). As a corollary, wear is low with ceramicimplant materials.

Glasses and carbon materials are commonlyregarded as ceramic materials, as they are com-

posed of ionically bonded components and dis-play the characteristic properties of ceramics,such as great hardness, lack of permanent de-formability, and excellent chemical inertness.These materials, however, do not have the long-range order typical of crystalline materials. Theiratomic arrangement is one of short order. Bioac-tive glass is a family of glass compositions withthe ability to bond to bone and enhance bone tis-sue formation. Collectively, bone bonding andbone tissue growth enhancement are termedbone bioactive behavior. Bioactive glasses arecomposed of silicon oxide, which is the networkformer in the glass structure, and calcium oxideand sodium oxide which are the network mod-ifiers. Figure 11-2 shows the schematic struc- Fig.11.2ture of a random glass network composed of anetwork former (SiO2) and network modifiers(Na2O and CaO). Some of the Si are bonded toeach other by bridging oxygen bonds and othersare coordinated with nonbridging oxygen bondsto network modifying ions in SiO2-Na2O-CaOglass. Other additives such as phosphorus pen-toxide (P2O5) can be added to modulate bonebioactivity reactions. With decreasing SiO2 con-centration, the instability and, hence, the bioac-tivity of the bioactive glasses increase. Bioactiveglass granules with a SiO2 concentration as lowas 45% by weight can be used as a resorbablebone graft material.

2.3 Polymers

Unlike metals, which have migrant electrons, thenonmetallic atoms of polymers are covalentlybonded. That is, elements such as C, N, andO have an affinity for attracting or sharing

FIGURE 11-2. Schematic structure of a randomglass network composed of a network former(SiO2) and network modifiers (Na2O and CaO).



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FIGURE 11-3. Various polymer chains. (From ParkJB. Biomaterials science and engineering. New York:Plenum Press, 1984.)

additional electrons. Thus, in covalently bondedelements, each electron is associated with a pairof atoms. Materials that contain only nonmetal-lic elements share electrons to build up largemolecules, often called macromolecules. Theselarge molecules contain many repeating units,or mers, from which comes the word polymers.The polymeric molecular arrangements can belinear, branched, or cross-linked (Fig. 11-3).The simplest example is polyethylene, whichFig.11.3is derived from ethylene (CH2 = CH2),in which the carbon atoms share electronswith two other hydrogen and carbon atoms!CH2(CH2!CH2)nCH2!, where n indicatesthe number of repeat units. To make a strongsolid, n should be well over 10,000, makingthe molecular weight of the polymer more than1 million grams/mole. Polymers bear a close re-semblance to natural tissue components such as

collagen. This makes it possible to incorporateother substances by direct bonding. An exampleis a heparin coating on polymer surfaces in aneffort to prevent blood clotting. During poly-mer processing additives such as antioxidants,antidiscoloring agents, and plasticizers are addedto facilitate the polymerization reaction and en-hance processability. Some of these additives aretoxic; therefore, pure medical-grade polymersare not readily available, except in very limitedcases.

2.4 Composites

Composite materials are solids that contain twoor more distinct constituent materials or phases,on a scale larger than the atomic. The term com-posite is usually reserved for those materials inwhich the distinct phases are separated on a scalelarger than the atomic and in which propertiessuch as the elastic modulus are significantly al-tered in comparison to those of a homogenousmaterial. Accordingly, reinforced plastics such asfiberglass as well as bone are viewed as compositematerials, but alloys are not. The properties of acomposite material depend on the shape, volumefraction, and interface between the constituentphases. The shape of the phases in a compos-ite material is classified into three categories asshown in Fig. 11-4. Most composite biomate- Fig.11.4rials have been developed to enhance mechani-cal and biocompatibility behavior. The strengthof composites derives in part from the fact thatfine fibers or whiskers of a material can be madenearly defect free for greatly increased strength,as is shown in Table 11-1. Carbon fibers have tab.11.1been incorporated in high-density polyethylenetibial plateaus of total knee replacements. Al-though strength was greatly enhanced, wear be-havior was fully unsatisfactory.

(A) (B) (C) FIGURE 11-4. Morphology of basic compositeinclusions: (A) platelet; (B) fiber; (C) particle.



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TABLE 11-1. PROPERTIES OF FIBERS USED FOR COMPOSITE MATERIALS

Class Material Tensile Strength (GPa) Young’s Modulus (GPa)

Whisker Graphite 20.7 675.7Al2O3 15.2 524.0Iron 12.4 193.1Si3N4 13.8 379.2SiC 20.7 689.5Boron — 441.3

Glass, ceramic Asbestos 5.9 186.2Drawn silica 5.9 72.4Boron glass 2.4 379.2

Polymer fibers High-tenacity nylon 66 0.8 4.8Metal wire Carbon steel 3.9 206.9

Molybdenum 2.1 365.4Tungsten 2.9 344.8

3 CRYSTAL STRUCTURE

3.1 Metals and Ceramics

Analyzing the atomic arrangements within asolid considerably facilitates studying the effectof bond type on resulting material properties.All metals, most ceramics, and some polymerscrystallize when they solidify. A crystalline ma-terial is characterized by long-range order andan infinitely repeating unit cell of atoms/ions.For example, metals can have their atoms ar-ranged in an orderly three-dimensional lattice,of which a two-dimensional section is shown inFig.11.5Fig. 11-5. In this two-dimensional lattice anysquare can be repeated by moving one latticespacing “a” in any direction. If this is extendedinto three dimensions, the spatial structure sodefined is a simple cube, and the crystal struc-ture is called a cubic crystal system. The unit cellis the smallest three-dimensional structure thatcan be repeated within the lattice. The unit cell

FIGURE 11-5. Two-dimensional section of a simplecubic lattice structure (a, lattice space).

of stainless steel is a cube with atoms located atthe corners and middle of each side. This crystalstructure is called a face-centered cubic (FCC)crystal. A cubic crystal with atoms located atcorners and in the center of the cube is called abody-centered cubic crystals lattice (BCC).

In addition to a cubic structure, there aresix other structures (Fig. 11-6). Apart from the Fig.11.6face-centered location of atoms, several otherso-called space lattices are possible; we indicatethese only as they specifically relate to the bioma-terials analysis. Figure 11-7 shows the arrange- Fig.11.7ments of the atoms in an FCC and a hexag-onal structure, specifically the hexagonal close-packed (HCP) structure. As this figure shows,these two lattice systems have the closest packingof atoms. In the FCC structure, the plane of clos-est packing runs through the three corners indi-cated by 1, 1", and 1"". Neighboring planes runthrough corners 2 or 3. Closest-packed planesof the HCP lattice are parallel with the base ofthe hexagonal prism. It is apparent that thereis considerable configurational symmetry inthese structures. This allows deformation of thestructure in a great variety of ways. Thus,since the cubic and hexagonal lattices are verycommon metal lattices, it further substantiatesthe easy deformability of metals. The crystalstructures, at room temperature, of some of thematerials used in hip surgery are summarized in tab.11.2Table 11-2.

Figure 11-8 shows the HCP structure of hy- Fig.11.8droxyapatite (HA), a bone bioactive ceramic.



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FIGURE 11-6. Geometric representation of the seven atomic lattices.

Comparison of the titanium HCP structure withthis structure reveals that the HA structure isconsiderably more complex, as a result of which,displacement of atoms within the lattice is diffi-cult. Thus the structure is resistant to deforma-tion and, when overloaded, fractures rather thandeforming permanently.

3.2 Polymers

Polymers are the result of chain reactions thatgive rise to linear, branched, or cross-linked poly-meric molecular arrangements (Fig. 11-3). Asthe extent of polymerization increases and themolecular chains become longer, the relative mo-bility of the chains in the structure decreases(Fig. 11-9). As a result, alignment of the chainsFig.11.9and formation of long-range order, i.e., crys-tallinity, is difficult. Polymers such as polyethy-lene and polyamide cannot be crystallized eas-ily, because the individual chains are long.Alignment is further rendered difficult by the

variability of the chain length. As the chain endsact as defects, this repeatedly interrupts the regu-lar continuity in the structure. Nevertheless, un-der optimal conditions, some crystallization ispossible: Various multimolecular structures areshown in Fig. 11-10. Fig.11.10

The chain mobility, which is directly propor-tional to the molecular weight, is related notonly to crystallinity, but also to mechanicalstrength. Thus, the molecular and microstruc-tural arrangements have a profound influenceon resulting properties. Factors affecting thestrength of polymers further include chemicalcomposition, side groups, cross-linking, copoly-merization, and blending. By way of example wecan consider the very strong aromatic polyamidfibers, better known by their trade name Kevlar.These fibers are made up of highly directionalmolecular chains, which confer on the fibersgreat strength. This structure, however, cannotbe obtained in thick sections. Thus, in order tomake optimal use of the excellent strength, these

FIGURE 11-7. FCC and HCP crystal lattices.



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TABLE 11-2. SOME MATERIALS ANDTHEIR REPRESENTATIVE EQUILIBRIUMCRYSTAL STRUCTURE AT BODYTEMPERATURE

Material Structure

Cobalt–chromium alloy FCCStainless steel AISI 316L FCCTitanium (Ti) HCP

Titanium alloy(Ti-6Al-4V annealed)

Two-phase material:!–HCP"–BCC

Tantalum (Ta) BCCNiobium (Nb) BCCGold (Au) FCCAlumina (Al2O3) HCPHydroxyapatite

[Ca10(PO4)6(OH)2]HCP

fibers are usually embedded in a matrix, i.e., usedin composites.

4 MECHANICAL BEHAVIOR

4.1 Mechanical Properties

Mechanical properties including elasticity andstrength are important properties to consider inthe selection of a material for a specific implantdesign. In this section we describe the variousmechanical properties relevant to biomaterials.

Many of the mechanical properties of a ma-terial take into account how it behaves underload. If a load, F, is applied to a material of crosssection A, it is subjected to a stress, # ,

# = F/A (1)

The units of # (N/m2 = Pa) show that stressand pressure are similar concepts. Under thisforce, there also will be a strain, $, due to changefrom the original length.

% = &l/l0 (2)

At low stress and low strain, stress is propor-tional to strain, i.e.,

# = E% (3)

E is constant for a material at low # . It iscalled Young’s modulus or modulus of elasticityof the material. E is a measure of stiffness (i.e.,resistance to strain) of a material, and is quitedistinct from strength. A small value of E meansthat a small stress gives a large extension (as inrubber); a large value of E indicates that the ma-terial is very stiff (as in diamond). The E valuesare 7 MPa and 1.2 # 106 MPa for rubber anddiamond, respectively.

Whereas equation 3 describes the mechani-cal behavior of a material at low stress, a moreextensive graphical representation of materialsdeformation is given in Fig. 11-11. This figure Fig.11.1shows the proportionality of stress to strain atlow stress, as well as the deviation from linearityat higher stress. The stress at which elongation isno longer reversible is called the limit of elastic-ity. For metallic and ceramic materials, the elasticregion shows a linear relationship between stressand strain, but for polymers the relationship canbe nonlinear. The yield strength is defined as thevalue of the stress when the strain is 0.2% morethan the elastic region will allow, i.e., the stress

FIGURE 11-8. The HCP structureof hydroxyapatite.



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FIGURE 11-9. Approximate rela-tion among molecular weight, Tg,Tm, and polymer properties.

level at which 0.2% permanent deformation oc-curs; this value is also referred to as the 0.2%proof stress. As it is very difficult to determinethe onset of permanent or plastic deformation,the yield strength is a more practical propertythan the limit of elasticity. Plastic deformation(plastic strain) is permanent strain and it occursbeyond the yield strength, as removal of stressleaves the material deformed. The ultimate ten-sile strength (UTS) is the maximum technicalstress (load divided by the original section) thematerial can sustain in tension. Ductility is the

strain at failure (usually expressed as percent).Fracture toughness is the measure of the energyrequired to break a material (the area under thestress-strain curve at failure). Fracture toughnessvaries with the material: it is 0.75 MPam1/2 forsilica glass, 4 MPam1/2 for diamond, and up to100 MPam1/2 for steel.

Hardness is another mechanical property. Itis the resistance of a material to the penetra-tion of its surface. The fatigue strength is themaximum stress a material can withstand for ei-ther an infinite or an arbitrary large number of

FIGURE 11-10. Two-dimensionalrepresentation of multimolecularpolymer structures: Extended chainfibers; folded-chain single crys-tals; glassy polymer amorphousstructure; and semicrystalline“fringed micelle” combinationstructure. (Reprint with permissionfrom Wunderlich, B. in Crystals oflinear macromolecules, ACS AudioCourse, American Chemical Society,Washington DC, 1973. Copyright (c)1973, American Chemical Society.)



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FIGURE 11-11. Stress deformation curve for a sim-ple tensile test. E, elastic modulus; se, proportion-ality limit; y.s., 0.02% offset yield stress; U.T.S.,ultimate tensile stress; df, total plastic deformationat fracture.

cycles. Usually the number of cycles selected is10 million. This represents about 10 years ofloading a hip prosthesis for a patient with lim-ited activity. As any device is subjected to cyclicloading, designing against fatigue is essential formost orthopaedic implants.

Typical mechanical property values of sometab.11.3tab.11.4 materials are given in Tables 11-3 and 11-4. Ta-tab.11.5 ble 11-5 summarizes Young’s modulus for some

of the commonly used materials.

TABLE 11-3. TYPICAL VALUES OF TENSILESTRENGTH FOR VARIOUS MATERIALS ATROOM TEMPERATURE

Tensile strengthMaterial (MPa)

Diamond 1.05 # 106

Kevlar 4,000High-strength carbon fiber 4,500High-tensile steel 2,000Superalloy 1,300Spider webs (drag line) 1,000Ti-6Al-4V 860CoCr alloy (F75) 655Aluminum 570Titanium (Grade 4) 550316L SS (F745, annealed) 485(Cold forged) 1,351Bone 200Nylon 100Rubber 100

4.2 Fracture Mechanicsof Ceramics

Biomaterials made of glass or ceramic (e.g.,bioactive glass or alumina) fail in a brittle man-ner. This is a process in which fracture occurswith little or no plastic deformation. Brittlefracture generally occurs by cleavage over par-ticular crystallographic planes (in crystallinestructures). Although not applicable here, it isuseful to know that at high temperatures, thecrystalline component of the ceramic can fail in-tergranularly. This occurs when grain-boundaryshearing takes place and cracks open up betweenthe grains, causing a local stress concentrationand ultimate fracture.

Fatigue fracture occurs in metals under cyclicstressing by the nucleation and extension of acrack within an extensively deformed area at ornear the surface. Such a phenomenon is unlikelyin ceramics. However, “static fatigue” or delayedfracture is common in ceramics. In this case,stress corrosion occurs preferentially at the tipof preexisting defects under the influence of astatic stress. This leads to subcritical crack pro-gression prorated to the time at which the partis stressed. Eventually fracture occurs some time AU: Preceding

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after the load was first applied. This kind of frac-ture, called fatigue in ceramics, is particularlysensitive to environmental conditions.

The static fatigue is dependent on the totaltime during which the load is applied, plus theloading rate and the total number of loading cy-cles. The relationship between the time to failuret and the fracture stress of a material can be de-scribed by the equation

log t = A/#f + B (4)

where A and B are constants. This relationshipis not valid at very short times but provides auseful description over a long time period.

The fatigue process has been associated withtwo phenomena: (1) a stress corrosion process,in which a sufficiently large stress enhances therate of corrosion at the crack tip relative to that atthe site, leading to a sharpening and deepeningof the crack and eventually resulting in failure,and (2) a lowering of the surface energy, ' , bythe adsorption of an active species. This leads to



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TABLE 11-4. FATIGUE PROPERTIES OF IMPLANT METALS

Fatigue Endurance LimitMaterial ASTM Designation Condition (At 107 Cycles, R = !1) (MPa)

Stainless steel F745 Annealed 221–280F55, F56, F138, F139 Annealed 241–276

30% cold worked 310–448Cold forged 820

Co-Cr alloys F75 As-cast/annealed 207–310P/M HIPa 725–950

F799 Hot forged 600–896F90 Annealed Not available

44% cold worked 586F562 Hot forged 500

Cold worked, aged 689–793 (axial tension R = 0.05, 30 Hz)

Ti alloys F67 30% cold-worked grade 300F136 Forged annealed 620

Forged, heat treated 620–689

aP/M HIP, Powder metallurgy produced, hot-isostatically pressed.

a decrease in the surface energy contribution tothe fracture surface work. However, as the esti-mated surface energies only represent about 30%of the total fracture surface work, changes in sur-face energy would not be sufficient to accountfor the observed long exposure to loads. For thisreason, as well as for its generally satisfactory de-scription of many experimental data, the stresscorrosion model is preferred. Although static fa-tigue is closely related to a corrosion process,the dissolution must be of a particular type thatleads to slow crack growth, increasing the stressconcentration at the crack tip until it grows to a

size that causes fracture. Such phenomena havebeen reported for alumina and zirconia headsof hip prostheses. Growth of subcritical crack-like defects till they reach a critical size can beprevented by designing for service stresses signif-icantly lower than those for which failure wouldoccur during the life of the device.

4.3 Elastic Propertiesof Composites

It is often useful to stiffen or harden a material,commonly a polymer, by the incorporation of

TABLE 11-5. ELASTIC PROPERTIES OF SOME TYPICAL MATERIALS

Materials Directionality Properties Modulus of Elasticity (MPa)

Cortical bone Anisotropic Longitudinal axis, 17,000Trabecular bone Anisotropic Longitudinal axis of femur

intertrochanteric 316 ± 293High-density polyethylene Isotropic 410–1,240PMMA Isotropic 3,000–10,000Stainless steel AISI 316L Isotropic 200,000Cobalt–chromium alloy Isotropic 220,000Titanium Isotropic 107,000Ti-6A1–4V Isotropic 110,000Carbon fiber–reinforced graphite

fibersAnisotropic Parallel to unidirectional

140,000Alumina Isotropic Single-crystal, 362,700

Polycrystal, 408,900



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particulate inclusions. The shape of the particleis important. In isotropic systems, stiff plateletinclusions are the most effective in creating astiff composite, followed by fiber inclusions. Theleast effective geometry is a spherical particle. Adilute concentration of spherical particulate in-clusions of stiffness Ei and volume fraction Vi, ina matrix, denoted by the subscript m, with a Pois-son’s ratio of 0.5, gives rise to a composite with astiffness E:

E = [5(E i ! Em)Vi/(3 + 2E i/Em)] +Em

(5)

Even if the particles are perfectly rigid comparedwith the matrix, their stiffening effect at low con-centrations is modest.

Fibers incorporated in a polymer matrix in-crease strength and fatigue life, in addition tothe stiffness (2). Unidirectional fiber compos-ites, when loaded along the fibers, can have astrength and stiffness comparable to that of steel,but with much less weight. However, if unidirec-tional fiber composites are loaded transverse tothe fibers, such a composite will be compliant,with stiffness not much greater than that of thematrix alone. If stiffness and strength is neededin all directions, the fibers must be oriented ran-domly. For such a three-dimensional isotropiccomposite,

E = (E iVi/6) +Em (6)

Thus, the stiffness is reduced by a factor ofabout 6 in comparison with an unidirectionalcomposite. However, if the fibers are aligned ran-domly in a plane, the reduction in stiffness is onlya factor of 3.

Conversely, when inclusions are used to re-duce the stiffness and the inclusions are more

compliant than the matrix, spherical ones re-duce the stiffness the least and platelets reduceit the most. In this case platelets resemble crack-like defects, and soft spherical inclusions are usedintentionally as crack stoppers to enhance thetoughness of polymers, with only a small sacri-fice in stiffness.

5 WEAR RESISTANCE

5.1 Ultrahigh-molecular-weightPolyethylene

Ultrahigh-molecular-weight polyethylene (UH-MWPE) (MW above 2 # 106 g/mol) has longbeen used for the acetabular cup of total hipprostheses and the tibial plateau and patellar sur-face of knee prostheses. The success of UHM-WPE in these applications is due to its favorableproperties, including abrasion resistance, im-pact strength, low coefficient of friction, chem-ical inertness, and resistance to stress cracking.Table 11-6 summarizes the mechanical and phy- tab.11.6sical properties of four different medical gradesof UHMWPE: compression molded sheet (1020CMS); ram extruded bar (1020 REB); partiallycrystallized 1050 CMS (Poly Hi); and partiallycrystallized 1050 REB (Poly Hi) (70).

In the majority of contemporary total jointreplacements, a metallic component articu-lates against UHMWPE. As a patient may beexpected to take an average of 1200 steps perday, the joint replacement is expected to with-stand millions of loading cycles during its ser-vice lifetime (148). In total hip replacements,where the articulating geometries consist of con-forming spherical surfaces, the wear occurs ata microscopic length scale (µm or less). Wearresistance has been related to its resistance to

TABLE 11-6. MECHANICAL AND PHYSICAL PROPERTIES OF MEDICAL GRADESOF UHMWPE (±± STANDARD DEVIATION)

1020 CMS 1020 REB 1050 CMS 1050 REB

Yield stress (MPa) 23.6 ± 0.1 23.6 ± 0.1 22.6 ± 0.1 22.3 ± 0.4UTS (MPa) 42.1 ± 2.6 37.2 ± 6.4 43.8 ± 3.5 39.9 ± 5.0% elongation 396 ± 20 376 ± 52 359 ± 20 354 ± 33Impact strength (kJ/m2) 161 ± 1.9 140 ± 1.4 93.7 ± 3.4 97.9 ± 2.9Density (g/cm3) 0.9366 ± 0.0001 0.9357 ± 0.0001 0.9308 ± 0.0003 0.9332 ± 0.0001



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multidirectional stresses (13,25,62,90,119,120,137,161–165). In total knee replacements,where the articulating surfaces consist of non-conforming cylindrical, toroidal or flat surfaces,the wear process includes various surface damagemechanisms ranging from pitting and delami-nation to burnishing and adhesive wear (5,161,172,174,175).

Focusing on acetabular components and theanalysis of retrieved devices, it is evident thatthere are three different mechanisms of wear. Ad-hesive wear is a process in which surface asper-ities of the polymer adhere to the metal surfaceand subsequently are torn off. As a result eithera polymer film is formed on the metal surfaceor polymer particles are released and entrappedin the joint. Such particles as well as remainingcement fragments can generate abrasive wear, asecond mode of wear. This process can also begenerated by asperities on the gliding metal part-ner. A third mechanism is fatigue wear; as a resultof creep or plastic flow, folds or cracks are formedthat cause small polymer particles to break off.

It is generally accepted that particulate debrisgenerated by mechanical wear (adhesion, abra-sion, and fatigue) of prosthetic componentsstimulates the generation of a pseudosynovialmembrane at the interface between implantand bone and the infiltration of fibrocytesand macrophages. In the presence of debris,these cells release various cytokines and medi-ators (such as IL-1", TNF!, collagenase, andprostaglandin E2) (69,85,89,155). These cy-tokines have been shown to be involved in boneresorption by activating osteoclasts (110). A re-cent study showed a positive correlation be-tween cytokine concentration in the looseningmembrane and the degree of underlying osteol-ysis (153). Many reports have investigated thesize of wear debris retrieved from failed totaljoint arthroplasty (TJA) and demonstrated thatphagocytosis and cell mortality increase withparticle size and concentration (18). The mech-anism that results in this cell death remains un-known, although a potential role for apoptosisin the pathogenesis of wear debris-associated os-teolysis has recently been suggested (19). Asep-tic loosening, which is the single most commoncause for long-term failure of TJA, is associated

with periprosthetic osteolysis with the incidenceof up to 25% of implant recipients (59,91,93).

The analysis of wear mechanisms in vivo read-ily points out the importance of the test condi-tions for the usefulness of wear data collectedin vitro. As loading controls the fatigue wear,both the loading cycle and the size and shapeof the contact area influence the wear behavior.Choosing the test solution is no less important,as this also affects the adhesive wear process. It iswidely recognized, however, that wear rates pro-duced in hip simulator studies are often lowerthan those generated in vivo (6,12,22), and thishas been explained by a lack of abrasive wear dueto roughened heads, third-body particles, andpolymer degradation. Although these parame-ters are clearly influential in wear debris produc-tion, the influence of alternative gait activitiesmust also be considered.

Hip simulator validation studies have corre-lated the mechanical behavior and surface mor-phology of polymeric biomaterials with the wearperformance (49,50,101). Because of the multi-directional motion prevalent at the articulatingsurface of total hip prostheses, the mechanicalbehavior of UHMWPE under multiaxial load-ing conditions has been shown to provide im-portant insight into the wear resistance (102).Based on equibiaxial tension testing performedon miniature 0.5 mm thick, 6.4 mm diame-ter disk specimens obtained from the surface ofwear-tested cups, work to failure has been signif-icantly correlated with the wear performance in ahip simulator for a wide range of virgin and cross-linked semicrystalline polymers (49,50). Fur-thermore, the surface morphology of the wear-tested acetabular components has been related tothe mechanical behavior of semicrystalline poly-mers, in that materials with greater toughness(work to failure) under multiaxial loading con-ditions were associated with lower surface rough-ness during hip simulator testing (101). Fig- Fig.11.12ure 11-12 shows an 8-station hip joint simulatorused in some of these studies.

Multiple factors can affect polyethylene wearand the production of wear debris in vivo af-ter joint arthroplasty. Such factors include theroughness and material of the femoral head,the method of polyethylene sterilization, and



11. Biomaterials 13

A

C

B

D E F

FIGURE 11-12. Photographs and schematics showing (A) MTS 8-station hip joint simulator, (B) physiolog-ical test setup, (C) fully constraining socket fixture, (D) partially constraining socket fixture, (E) location ofhorizontal torque cells, and (F) direction of torque measured.

the mechanical properties of the polyethyleneitself (149). The alterations of the mechani-cal properties of polyethylene associated withgamma irradiation in air and shelf life have beendemonstrated in the laboratory (24,26,169).Recent laboratory studies have confirmed thatthe wear resistance of ultrahigh-molecular-weight polyethylene can be significantly in-creased when applying additives or high-dose ir-radiation (21,68,105,121,125,176), with manystudies reporting extremely low quantities ofwear, typically as low as 2.0 mm3 per millioncycles for highly cross-linked polyethylenes. Astudy undertaken by Laurent et al. (105) re-ported no measurable wear after 20 millioncycles of normal walking and highlights thepotential of such materials. Radiographic mea-surements are also evidencing extremely lowin vivo wear rates of cross-linked polyethylene(176), with recent retrieval studies undertakenby Oonishi et al. (130) showing little evidenceof in vivo scratching and delamination on sock-ets irradiated at 100 Mrad. As these materialscontinue to maintain an acceptable biocompat-ibility, the overall success of polyethylene hiparthroplasty should greatly increase, thereby of-fering a highly cost-effective treatment.

5.2 Alumina

High-density, high-purity alumina is used inload-bearing hip prostheses because of its out-standing wear resistance and excellent corrosionresistance. It has a high Young’s modulus anda hardness second only to that of diamond.These properties have made the alumina-on-alumina couple for femoral heads and acetabularcups a materials combination of considerable im-portance. Most alumina devices are very finegrained polycrystalline !-Al2O3 produced bypressing and sintering at high temperature (1600to 1700$C). A very small amount of MgO(less than 0.5%) is used to aid sintering andlimit grain growth during sintering. Static andfatigue strength and fracture toughness are afunction of grain size and percentage of sin-tering aid (i.e., purity). Alumina with an aver-age grain size below 4 µm and a purity greaterthan 99.7% exhibits excellent flexural and com-pressive strength. Table 11-7 summarizes these tab.11.7properties.

Studies have reported excellent performanceof the alumina on alumina bearing in termsof low annual wear (<5 µm) (10,11,124,128,160,170). The long-term friction of an



14 11. Biomaterials

TABLE 11-7. PHYSICAL CHARACTERISTICS OF Al2O3 BIOCERAMIC

High-alumina Ceramics ISO Standard 6474

Alumina content (% by weight) >99.8 %99.50Density (g/cm3) >3.93 %3.90Average grain size (µm) 3–6 <7Surface roughness, Ra (µm) 0.02Hardness (Vickers hardness number, VHN) 2,300 >2,000Compressive strength (MPa) 4,500Bending strength (MPa) 550 400Young’s modulus (GPa) 380Fractures toughness 5–6(K1c) (MPa) 5–6

alumina-alumina joint prosthesis decreases withtime and approaches the value of a normaljoint. This lead to wear on alumina articulat-ing surfaces being nearly 10 times lower thanon polyethylene surfaces gliding against metallicFig.11.13heads (Fig. 11-13).

Careful operative technique with correct po-sitioning of the prosthesis is necessary to avoidexcessive wear (37,83,113). Impingement be-tween the femoral neck and the rim of the acetab-ular cup has been associated with massive wearand osteolysis (173) and high wear rates havebeen associated with “stripe” and severe wearof Mittelmeier prostheses (127). Nevertheless,given good surgical technique it is possible toachieve very low volumes of wear with ceramic-on-ceramic prostheses.

Histological studies of the pseudosynovial tis-sue obtained from around retrieved uncementedceramic–ceramic prostheses have identified nu-merous alumina ceramic particles with a meansize of 5 µm (100). In addition, Henssge et al.(79) observed particles up to 5 µm in diame-ter in the periprosthetic tissues from around ce-mented alumina–alumina prostheses. More re-cently, very small alumina wear particles in the

size range 5 to 90 nm were revealed by trans-mission electron microscopy (TEM), whereasSEM (lower resolution) revealed particles in the0.05 to 3.2 µm size range (72). The bimodalsize range of alumina ceramic wear debris over-lapped with the size ranges commonly observedwith metal particles (10 to 30 nm) and parti-cles of ultrahigh-molecular-weight polyethylene(0.1 to 1000 µm). It is possible that the two typesof ceramic wear debris are generated by two dif-ferent wear mechanisms in vivo; under normalarticulating conditions, relief polishing wear andvery small wear debris is produced, while underconditions of microseparation of the head andcup and rim contact, intergranular and intra-granular fracture and larger wear particles aregenerated.

The in vivo wear rate of the ceramic–ceramicjoint couple is reported to be up to 200 timesless than that of the metal–PE joint couple (37).Although ceramics are considered more bioin-ert than PE ceramic particles generated in vivo,when present in large concentrations, they caninitiate an inflammatory response in peripros-thetic tissue (107,108). Moreover, aluminaceramic particles could induce macrophage

FIGURE 11-13. Time dependence of(solid line) coefficient of frictionand (dashed line) index of wear ofalumina-alumina versus metal-PE hipjoint (in vitro testing). (Reprintedfrom Hench LL. Bioceramics: fromconcept to clinic. J Am Ceram Soc1991;74(7):1487–1510, by permissionof the American Ceramic Society.)



11. Biomaterials 15

TNF-! release (18,20) and apoptosis (19). Petitet al. (135) compared the macrophage responseto identically sized particles of alumina ceramic(Al2O3) and UHMWPE in terms of TNF-! re-lease and induction of apoptosis of J774 mousemacrophages. The stimulation of TNF-! re-lease was much greater (8 to 10 times higher)with UHMWPE than with Al2O3. However,the induction of apoptosis as measured by ac-tivation of caspase-3, PARP cleavage, and DNAladdering was different for these two particlecompositions, being faster and more importantwith Al2O3 than with UHMWPE. One of thehighlights of apoptotic cell death, as it appliesto periprosthetic osteolysis, is that the wholeprocess terminates in the elimination of deadmacrophages without the induction of a signif-icant inflammatory reaction (23). Therefore, itcould be possible that the ability of Al2O3 par-ticles to induce macrophage apoptosis may ex-plain the lower TNF-! release observed withthese particles and explain the differences seenin osteolysis patterns of ceramic–ceramic versusmetal–polyethylene articulations (107). The in-duction of macrophage apoptosis may thereforebe a desirable therapeutic endpoint.

6 METALLIC BIOMATERIALS—CORROSION RESISTANCE ANDRELATED PROPERTIES

Corrosion is the unwanted chemical reaction ofa metal with its environment, resulting in itscontinued degradation to oxides, hydroxides, orother compounds. Biological fluids in the hu-man body contains water, salt, dissolved oxy-gen, bacteria, proteins, and various ions suchas chloride and hydroxide. As a result, the hu-man body is a very aggressive environment formetals. Consequently, corrosion resistance of ametallic implant material is an important as-pect of its biocompatibility. Metallic biomate-rials are normally considered to be highly cor-rosion resistant because of the presence of anextremely thin passive oxide film that sponta-neously forms on their surfaces. These filmsserve as a barrier to corrosion processes in alloysystems that would otherwise experience very

high corrosion rates. That is, in the absenceof passive films, the driving force for corrosionfor typical implant alloys [e.g., titanium-based,cobalt chromium (CoCro)-based, and stainless-steel alloys] is very high, and corrosion rateswould also be high. The properties of these pas-sive oxide films depend to a large extent ontheir structure and chemistry, which are them-selves dependent on the substrate’s prior ther-mal, mechanical, and electrochemical history. Inthis section, the basis for the excellent corrosionresistance of currently used implant metals isexplained.

There are two CoCr alloys extensively usedin implant fabrications such as artificial joints,or stems of prostheses for heavily loaded jointssuch as knees and hips: the castable CoCrMoalloy and the CoNiCrMo alloy, which is usu-ally wrought by (hot) forging. The chromiumis a reactive element and is added to produce astable firmly adherent protective chromium ox-ide surface layer. Chromium also enhances thesolid solution strengthening of the alloy. Themolybdenum is added to produce finer grains,which results in higher strengths after castingor forging. As with stainless steel, molybdenumalso enhances the corrosion resistance of cobalt-chromium alloys (136). The inhomogeneousmicrostructure of the cast Co-Cr-Mo alloy ren-ders it more susceptible to corrosion than theforged alloy (36), presumably due to the presenceof chromium-depleted dendritic regions actingas the more anodic sites in a galvanic reaction.Wrought Co- Cr-Mo has lower carbon contentthan cast Co-Cr-Mo and, as a result a lowercorrosion resistance when tested in physiologicsolution (36,82). The metallic products releasedfrom the prosthesis because of wear, corrosion,and fretting may impair organs and local tissues,and moreover, some alloys with certain amountof Co can be toxic in the body. Low wear hasbeen recognized as an advantage of metal-on-metal hip articulations because of their hardnessand toughness.

Both c.p. Ti and Ti-6Al-4V possess excel-lent corrosion resistance for a full range of ti-tanium oxide states and pH levels. Titanium isa base metal in the context of the electrochem-ical series; however, it derives its resistance to



16 11. Biomaterials

corrosion by the formation of a solid oxide layerto a depth of 10 nm. Under in vivo conditionsthe oxide, TiO2 is a very stable reaction prod-uct. Corrosion currents in normal saline are verylow (10!8 Acm!2). The low dissolution rate andnear chemical inertness of titanium dissolutionproducts allow bone to thrive and therefore os-seointegrate with titanium.

Titanium implants remain virtually un-changed in appearance; however, even in theirpassive condition, metals, including titanium,are not inert. It is established beyond doubt, byboth in vitro and in vivo experiments (60,73,103,104,168), that there is a passive dissolution fromthe metal. Thus, linked to the issue of elec-trochemical behavior are several questions: (1)What material is released? (2) How much ma-terial is released under statistic conditions? (3)How is the release modified by wear conditions?(4) What subsequent reactions do the releaseproducts get involved in? (5) What percentageof the release products is excreted and what per-centage is retained? (6) Of the percentage that isretained, where does it accumulate? (7) What bi-ological response(s) will result from the retainedfraction (14,15,44,48,96,122,166)?

Stainless steel contains enough chromium toconfer corrosion resistance by passivity. The pas-sive layer (chromium oxide) is not as robust as inthe case of titanium or the cobalt–chromium al-loys. The relatively resistant varieties of stainlesssteel are the austinitic types 316, 316L, and 317,which contain molybdenum (2.5% to 3.5%).Even these types of stainless steels are vulnera-ble to pitting and to crevice corrosion aroundscrews, under certain circumstances such as in ahighly stressed and oxygen-depleted region (4).The corrosion resistance can be enhanced by in-creasing the thickness of the protective oxide us-ing concentrated nitric acid (“passivation”), byboiling in distilled water, or by electrochemicalmeans (anodization) (87). The reduction of car-bon to less than 0.03% has virtually eliminatedthe risk of intercrystalline corrosion, which canoccur when there is precipitation of chromiumcarbide at the grain boundary in stainless steelwith a carbon content above this value (133).Unfortunately, lowering the carbon content re-sults in lowering the ultimate tensile strength ofstainless steel.

Corrosion of an implant in the clinical set-ting can result in symptoms such as local painand swelling in the region of the implant, withno evidence of infection; only cracking or flak-ing of the implant (seen on roentgenograms),and excretion of excess metal ions. At surgery,gray or black discoloration of the surroundingtissues may be seen and flakes of metals may befound in the tissue. Corrosion also plays a role inthe mechanical failures of orthopaedic implants.Most of these failures are due to fatigue, andthe presence of a saline environment certainlyexacerbates fatigue. The extent to which corro-sion influences fatigue in the body is not pre-cisely known, although once an initial crack hasformed, crack propagation will be faster than inair or vacuum. Other mechanisms of corrosionsuch as fretting may also be involved at point ofcontact such as in the countersink of hip nails.

Different parts of the body undergo differ-ent types and rates of corrosion. Wounds andinfections can significantly change pH. In vivo,the equilibrium state between a metal and its re-action products (oxide, hydroxide, etc.), whichcauses passivation, may not occur if the reac-tion products are removed by the tissue fluidturnover. The replenishment of ions accumu-lated at the implant–tissue interface may causeadverse increase in the rate of ion release thatmay cause damage on the cellular level. Focus-ing on the major alloying elements in cobalt-based materials (Co, Cr, and Ni), Co and Niions bind to serum albumin, and Cr6+ binds tored blood cells (14). Chemical analyses of urinefrom animals subjected to metal salts indicatedthat most Co and Ni is rapidly excreted, whileless than 50% of the Cr is excreted, and this oc-curs at a slower rate than Co or Ni (15,123).Furthermore, organ levels of Co and Ni are notsignificantly elevated, whereas they are for Cr.Serum and urine analyses of patients with totaljoint replacements have also indicated a dose–response relationship (127). Koegel and Black(96), using a cast Co-Cr-Mo microsphere model,found dose related elevations in serum Co andCr, with peak concentrations achieved 3 days af-ter implantation. Scaling for the implant surfacearea to animal body weight ratio (300#), the Coand Cr elevations were 20 and 12, respectively. Ina related study (166), it was determined that the



11. Biomaterials 17

form of the released chromium was Cr6+, a morebiologically active form of Cr than Cr3+. Blacket al. (7) found that when a pyrolytic carboncoating was applied to cast Co-Cr-Mo, thecarbon-coated implants released more Co andCr than the uncoated implants. Relative mo-tion between the implant and tissue may havecaused the release of additional debris, possi-bly in the form of metal carbides. Ducheyneand Healy (44) determined that hydroxyapatitecoatings reduce the Ti and Al passive dissolutionrate from porous coated Ti-6Al-4V. Hydroxya-patite, however, did not produce a change in therelease kinetics of Co and Cr from Co-Cr al-loys. Other studies on Ti and Ti-6A1-4V haveshown that titanium is preferentially accumu-lated locally, with elevated levels of Ti detectedin adjacent soft tissue and bone (48). Healy andDucheyne (74) further determined that serumproteins increase the release rate kinetics of ti-tanium compared to solutions containing onlyserum electrolytes.

7 BONE BIOACTIVE CERAMICS

7.1 Calcium Phosphates—Structure, Mechanical Properties,and Processing

Bioactive materials including calcium phosphateceramics and bioactive glasses bond to bone andenhance bone tissue formation. The forms of cal-cium phosphate ceramics most widely used aretricalcium phosphate [Ca3(PO4)2, whitlockite],tetracalcium phosphate (Ca4P2O9), and hydrox-yapatite [Ca10(PO4)6(OH)2] (HA). There is awide variation of mechanical properties of syn-thetic calcium phosphates as given in Table 11-8.tab.11.8This variation of properties is the result of vari-ation of density and crystalline structure ofcalcium phosphates, in turn the result of dif-ferences in processing methods. Sintering of cal-cium phosphate ceramics is usually carried out inthe range of 1000 to 1500$C. Depending on thefinal firing conditions, the calcium phosphatecan be hydroxyapatite or "-whitlockite. In manyinstances however, both types of chemistries ex-ist in the same final product.

The phases formed at high temperature de-pend not only on the sintering temperature but

also on the partial pressure of water in the sin-tering atmosphere. This is because with waterpresent, HA can be formed and is a stable phaseup to 1360$C as shown in the phase equilibriumdiagram for CaO and P2O5 with 500 mm Hgpartial pressure of water (Fig. 11-14). The tem- Fig.11.14perature range of stability of HA increases withthe partial pressure of water, as does the rateof phase transition of tricalcium phosphate ortetracalcium phosphate to hydroxyapatite. Theideal Ca/P ratio of hydroxyapatite is 10:6 and thecalculated density is 3.219 g/cm3. Hydroxyap-atite has a high elastic modulus (40 to 117 GPa),whereas the elastic modulus of compact bone is12 to 18 GPa. Poisson’s ratio for hydroxyapatiteis about 0.27, which is close to that of bone(approximately 0.3).

At body temperature only two calcium phos-phates are stable when in contact with aqueousmedia such as body fluids. At pH less than 4.2,the stable phase is CaHPO4 · 2H2O (dicalciumphosphate or brushite), whereas at pH greaterthan 4.2 the stable phase is HA.

7.2 Calcium Phosphates—Biological Behavior

Daculsi et al. used transmission electron micros-copy to determine the in vivo degradation ofbiphasic implant materials which were variousmixtures of hydroxyapatite and tricalcium phos-phate (27). Implants were inserted in surgically

Au:Permission toreproduce?

created periodontal defects. At 6 months, a di-rect correlation was established between the rateof dissolution of the material and the quantityof newly formed microcrystals with Ca/P ratiossimilar to those of bone apatite. The transmis-sion electron microscopical studies carried out byJarcho et al. (88) and Tracy and Doremus (156),who each used dense, stoichiometric hydroxya-patite, showed direct apposition of bone crys-tals onto the synthetic apatite lattice. Whereasin these studies solid-solution ion exchange wasnot addressed, Davies et al. (32) documented ina more recent study a limited in vivo reactivityof this material.

A similar result of limited reactivity of dense,stoichiometric hydroxyapatite was also found bySchepers et al. (145), who implanted granules ofthis material in various bone tissue sites of the



TAB

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)

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4.5

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170

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52.

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8—

310–

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26.5

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200

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18



11. Biomaterials 19

FIGURE 11-14. Phase equilibrium diagram of cal-cium phosphates in a water atmosphere. Shadedarea is processing range to yield HA-containing im-plants.

beagle mandible. Bone tissue grew from the pre-existing bone along the particles that were closestto the defect wall. However, bone tissue did notgrow further than about 1 mm into the defect af-ter about 12 weeks. In addition, the presence ofmultinucleated cells was continuously observedfrom that time on, and the ceramic surface ex-hibited a moth-eaten appearance resulting fromthe local resorption of the material. In the areaswhere multinucleated cells were present, therewas no new bone tissue formation. Osteocon-duction and multinucleated cell-mediated re-sorption frequently had taken place on separatesides of a single particle. These phenomena areFig.11.15shown in Fig. 11-15a and b.

This experiment showed beyond doubt thatdense, stoichiometric hydroxyapatite is osteo-conductive in vivo. However, it also revealed thatthis material displays only a limited bone bioac-tivity. If a key element for bone bioactive behav-ior is the development of a carbonated apatitesurface, one might expect an excellent responseby using a carbonated apatite as implant mate-rial. However, implantation data do not un-equivocally support this hypothesis. The tissueresponse to a porous carbonated apatite mate-rial implanted in beagle mandibular bone wasvery similar to that with dense, stoichiometric

hydroxyapatite (145). Osteoconduction fromthe bone defect wall occurred over a limited dis-tance and bone tissue conduction into the centerof the defects was not observed. In addition, achronic, mild multinucleated cell resorption ofthe material started as early as 3 months afterimplantation. This limited osteoconductivity isattributed to the high stability of the hydroxya-patite ceramic in tissue fluids which results ina considerable induction time for precipitationof a fresh apatite layer from the interfacial tissuefluids onto the material surface. Another fac-tor which contributes significantly to the delay

A

BFIGURE 11-15. Histological section of dense, stoi-chiometric hydroxyapatite. (A) After 3 months, bonetissue has grown along the particle surface, but onlyover a distance of about 1 mm. (B) From 6 monthson, the particle surfaces have a moth-eaten appear-ance caused by multinucleated cell resorption. Theresorption is not paralleled by a bone formation pro-cess (Reprint from Schepers E, Declercq M, DucheyneP, Kempeneers R. Bioactive glass particulate materialas a filler for bone lesions. J Oral Rehab 1991;18:439–452, with permission.)



20 11. Biomaterials

in the formation of the bioactive apatite surfacelayer is protein adsorption.

7.3 Calcium Phosphates—Coatings

Coating metallic prostheses with a calcium phos-phate layer provides direct implant–bone bond-ing and better fixation and longevity. When aporous stainless-steel fiber network was coatedwith a slip-cast HA lining, a marked increasein bone ingrowth was observed in comparisonto the ingrowth in the same porous metal with-out HA lining. This effect was pronounced at2 and 4 weeks, but it disappeared at 12 weeks(45). This phenomenon provided the distincttherapeutic benefit of faster rehabilitation for pa-tients with HA-coated devices. However, struc-tural and compositional changes of the Ca-Pceramic occur as a result of processing of thecoatings. These variations of the material char-acteristics affect the rate of enhancement of thebioactive ceramic coating. Using an experimen-tal protocol that extensively controlled factorssuch as animal model, pore size, pore morphol-ogy, and properties of the metallic substrate, itwas found that there were statistically significantvariations in bonding strength in the first fewweeks after implantation arising from differencesin ceramic characteristics (40). Figure 11-16Fig.11.16shows that calcium-phosphate coatings signifi-cantly enhanced bone tissue growth fixation, butthat the effect differed among coatings. In addi-tion, these data show that the intensity of theeffect was consistent among the three early timepoints studied.

The clinical experience with HA coatings hasbeen excellent. Most trials have focused on theperformance of the femoral component. Therehas been variability in the implant design stud-ied, the method and characteristics of the ap-plied coatings, and the extent of the coating onthe surface of femoral stems (64). There havebeen numerous cohort studies but few random-ized clinical trials (29,30,65,66,143). Larger un-controlled series have reported valuable results(17,29,65). D’Antonio et al. reported a multi-center study of 316 hips (280 patients) with aproximally hydroxyapatite-coated titanium alloy

FIGURE 11-16. Shear strength as a function of im-plantation time for controls (porous-surfaced tita-nium) and porous specimens with various calciumphosphate coatings. Note that the greatest effecton bone tissue ingrowth fixation is associated withcoating CAP3, which is the coating with the high-est in vitro dissolution rate (Reprint from DucheyneP, Beight J, Cuckler J, Evans B, Radin S. Effect ofcalcium-phosphate coating characteristics on earlypost-operative bone tissue ingrowth. Biomaterials1990;11:531–540, with permission.)

stems that were implanted in a young and ac-tive patient population from 1987 through 1990(31). Clinically, these patients had early pain re-lief and rapid restoration of function. The resultsshow excellent lasting fixation of this taperedtitanium alloy stem coated proximally with athin, dense layer of hydroxyapatite. Another re-port by Hernandez et al. described the 11-yearfollow-up results of 52 unilateral primary hiparthroplasties performed with hydroxyapatite-coated stems (80). The femoral prosthesis usedwas a collarless titanium alloy implant, withproximal circumferential hydroxyapatite coatingand increased distal thickness to fit the proxi-mal diaphyseal region of the femur. At the endof the follow-up period, excellent clinical re-sults were recorded in 40 arthroplasties (77%).The 11-year survival rate was 92.3%. Geesinkand Hoefnagels (66) reported their 6-year resultsof 118 primary total hip arthroplasties using ametaphysis filling Ti stem that had a 50-µm-thick air plasma sprayed coating of hydroxyap-atite on the proximal third of its surface. Thestudy’s reported survival rate of the prosthesis at6 years was 100% for the 118 hydroxyapatitecoated stems and 99% for the hydroxyapatite-coated cups. Vidalain reported a 0.97 survivalrate of the Corail femoral stem at 10 years of



11. Biomaterials 21

implantation (158). This survival rate is consid-ered to be even higher than that of many seriesof Charnley prostheses, which remain the abso-lute standard in this respect. During that period(10 years), more than 6,700 implantations havebeen performed and analyzed by the surgeonsof the ARTRO Group, with a consistent annualrate of about 700 replacements. Functional re-sults have been rated as excellent using the PostelMerle d’Aubigne or Harris hip scores. Accordingto the PMA hip scale, the preoperative score was11.29; it was 17.01 at 1 year, 17.21 at 5 years,and 17.20 at 10 years. Actually, scores are verysimilar to those achieved with the best cementedprostheses, particularly for pain, with the ab-sence of significant thigh pain, which is typi-cal of the CORAIL hip prosthesis. Altogether,63% of patients are totally pain free and have re-covered normal motion and function. Based onthese clinical results it is concluded that osteo-conductive coatings have established an excellentperformance record in primary and revision hiparthroplasty.

In view of all the advantages of hydroxyap-atite coating, are there drawbacks or at least con-cerns or questions? It is now well established thathydroxyapatite is not toxic, does not produce in-flammatory or allergic reactions, and is not car-cinogenic. Still, four problems can be discussed:HA resorption and long-term stability, osteol-ysis that is occasionally observed, polyethylenewear and potential risk of granuloma, and thedifficulty of extracting an HA-coated stem, par-ticularly when it is perfectly ingrown.

Resorption of hydroxyapatite is more or lessinescapable, due to both chemical dissolutionand cell-mediated degradation. However, thekinetics of this degradation process is influ-enced by a number of factors which are still notclearly understood, and consequently, are not allwell known. Aebli et al. carried out a histologi-cal study of a proximally hydroxyapatite (HA)-coated femoral component, retrieved after 9.5years of good function (1). The HA coating hadcompletely degraded. Bone was in direct contactwith the titanium surface in all the areas whichhad been coated, with no interposing fibrous tis-sue. There were no signs of particles, third-bodywear, adverse tissue reactions, or osteolysis. Bone

remodeling was evident by the presence of re-sorption lacunae; tetracycline labeling showedbone laid down 6 years after implantation. Theloss of the HA coating had no negative effecton the osseointegration of the stem. It was con-cluded that the HA coating contributes to thefixation of the implant and that its degradationdoes not adversely affect the long-term fixation.

HA, when first used as a hip stem coat-ing more than 10 years ago, was consideredsufficiently insoluble at neutral pH not to bedegraded after implantation. However, it thenbecame apparent that the HA was decomposedduring plasma spraying. Several phases (!- or"-TCP, CaO, TTCP, oxyapatite, and amor-phous calcium phosphate) can be evidenced inthe HA coating after spraying. These differentphases occur in various amounts and at pref-erential locations in the coating depending onplasma-spray parameters and powder character-istics. CaO, for example, is found at the interfacewith the metal, whereas the contaminating cal-cium phosphate phases occur around the crys-talline phase of the coating grains. Most of thecontaminating phases are more soluble than HAat neutral pH and their presence in the HA coat-ing increases its degradability.

The coating environment is also determinantin coating degradation. The dielectric constantand pH of the extracellular fluids are key factorsand can be modified by cell activity. The pH can,for example, be modified by osteoclast activity,the pH in the resorption chamber of osteoclastsbeing as low as 4.5 to 5. Similar low-pH com-partments are found in cytoplasmic compart-ments such as lysosomes. These different factorsexplain the degradation of calcium phosphatecoatings evidenced in contact with or in the closevicinity of osteoclasts, macrophage, or giant cells.Degradation of the calcium phosphate coatingsconstituted by the stacking of calcium phosphategrains occurs preferentially at the interface be-tween the grains, where the most soluble phasesfacilitating the release of calcium phosphate par-ticles are located. Particle release by a materialimplanted in bone is a cause of concern dueto the osteolytic reactions that have been evi-denced in the vicinity of metal or polymeric par-ticles debris. Furthermore, Bloebaum et al. (8)



22 11. Biomaterials

suggested that the calcium particles emitted byacetabular coatings could migrate to the surfaceof polyethylene cups, thereby triggering third-body wear responsible for osteolysis and asepticloosening.

When HA-coated metallic implants havebeen retrieved from humans and animals, evi-dence has been found of loose HA particles andosteolysis (3,9,28,42,63). Bloebaum et al. re-ported that histologic evaluation of 14 retrievedimplants revealed HA particles, evidence of in-flammatory reactions and osteolysis (9). Fur-thermore, resorption of HA can create a voidbetween the metal implant and the surround-ing bone matrix (151), which may lead to me-chanical instability and eventual failure. Reikeraset al. reported the outcome of 191 acetabu-lar gritblasted titanium cups with a hemispher-ical design for press-fit insertion and coatedwith hydroxyapatite (139). The prosthesis wasmade of gritblasted titanium entirely coated withhydroxyapatite. The 155 patients, aged 15 to78 years, were operated on during the years 1991to 1993 and followed for 7 to 10 years. Duringthis period, 39 cups were revised because of me-chanical loosening, a further nine had radiolu-cent lines, and two had focal osteolysis. Noneof these 11 patients had clinical symptoms. Fail-ure was associated with age, wear, and radiolu-cency/osteolysis. At revision, it was found thatthe soft tissues were discolored, and that mostof the coating had disappeared. It has been con-cluded that this design of hydroxyapatite-coatedcups has a high rate of debonding and failure.Regardless of these reports, the overall clinicalbenefit of calcium phosphate coatings has beendecidedly positive. Basic science studies provedthat next-generation biomimetic hydroxyapatitecoatings will expand the application range of os-

teoconductive coatings to three-dimensional in-growth structures and combinations with antibi-otics, growth factors, or other achievements ofmodern biotechnology.

7.4 Bioactive Glasses

Bioactive glasses are silica-based glasses in whichsilica (SiO2) is the network former and alkalimetals, e.g., Na and K, or alkaline earth met-als, e.g., Ca and Mg, are the network modifier.The ratio of network former to network mod-ifier in the composition of bioactive glass de-termines its solubility in physiological solutionsand hence its bioactivity and resorbability. Ta-ble 11-9 summarizes the chemical composition tab.11.09of selected bioactive glasses and glass ceramicswith different degrees of bioactivity.

The amorphous structure of bioactive glassimpairs its mechanical strength and lowers itsfracture toughness. Tensile strength of most ofthe bioactive glass compositions in Table 11-9is in the range 40 to 60 MPa, which makesthem unsuitable for load-bearing applications.For some applications low strength is offset bythe glasses low modulus of elasticity of 30 to 35Gpa, which is close to that of cortical bone (7 to25 GPa).

A unique characteristic of bioactive glass isits ability to form a strong interfacial bondwith adjacent tissues in the host, and it hasbeen used successfully as a bone-substitute mate-rial in orthopedic and dental surgery (131,147,150,152,171). To understand the basic mecha-nisms of the formation of the bone–biomaterialsbond, two main approaches were consid-ered. One focused on studying the bone–biomaterials interface that developed in vivo.The examination of bonding zone revealed the

TABLE 11-9. COMPOSITION OF BIOACTIVE GLASSES AND GLASS CERAMICS(IN WEIGHT PERCENT)

Type Code SiO2 CaO Na2O P2O5 MgO K2O CaF2

Bioactive glass 45S5 45 24.5 24.5 6 — — —45S5F 45 12.5 24.5 6 — — 12.555S4.3 55 19.5 19.5 6 — — —

Bioactive glass ceramic Ceravital 40–50 30–35 5–10 10–15 2.5–5 0.5–3 —A-W GC 34.2 44.9 — 16.3 4.6 — 0.5



11. Biomaterials 23

consistent presence of an interfacial hydroxyap-atite layer (78,126). The other approach used invitro immersions in simulated physiological flu-ids or cell-containing media (99,77,53). Uponimmersion of bioactive glass in acellular, simu-lated body fluids (with ionic concentration andpH similar to the interstitial fluid) the followingreactions take place (77):

Leaching of Na+ (or K+) into solution (i.e.,exchange with H+ or H3O+ from solution) andformation of SiOH (silanol):

!Si!O!Na + H++ OH ' !Si!O!H+ Na+ (solution) + OH!

Loss of soluble silica (SiO2) and formation ofSiOH:

!Si!O!Si! + H2O ' !Si!OH!HO + Si!

Polycondensation of silanols to form hy-drated silica gel:

O!Si!OH + HO!Si!O 'O!Si!O!Si!O + H2O

Formation of amorphous calcium phos-phate layer: Ca2+ and PO4

3! migrate to thesurface through the SiO2-rich layer, forming aCaO–P2O5-rich layer on top of the SiO2gel layer(the CaO–P2O5 layer proceeds to grow by in-corporation of soluble Ca2+ and PO4

3! fromsolution).

In the absence of proteins and bone cells theamorphous calcium phosphate layer crystallizes

into apatite on surface [either hydroxycarbonateapatite (HCA) or fluorapatite—CO3

2! or F!

from solution].These reactions, however, does not occur as

such in vivo or in vitro in the presence of pro-teins, the HA layer does not form on the glasssurface. Only when osteoblasts are present on thebioactive glass surface does the HA surface formagain, evidently as a result of cellular activity(53–57).

There is a good correlation between the abil-ity of a biomaterial to develop a hydroxyapatitesurface layer in acellular SBF and its ability tobond to bone. Therefore, the degree of bioac-tivity of biomaterials is evaluated based on theirability to develop a hydroxyapatite surface layerin vitro. As already stated, this does not neces-sarily imply that the above dissolution–precipi-tation reactions take place at the material–tissuefluid interface in vivo. In addition to the majordifference in composition between tissue fluidand SBF as shown in Table 11-10, there is a tab.11.10considerable difference in the parametric condi-tions under which the above reactions take placein vivo and in vitro.

It is obvious from Table 11-10 that one can-not expect that the material behavior in vitro issimilar to that in vivo, especially in the first fewdays after implantation where the wound healingmechanism is in effect. Figure 11-17 schemati- fig.11.17cally shows a variety of events that were reportedto occur at the bioactive ceramic–tissue interface(41). The list does not imply a ranking in termsof time sequence or importance: (1) dissolutionfrom the ceramic (35,47,78,94,106,126,134);

TABLE 11-10. DIFFERENCES IN MEDIUM COMPOSITION AND PARAMETRICCONDITIONS UNDER WHICH THE BIOACTIVITY REACTIONS AT THE BIOMATERIALSURFACE TAKE PLACE IN VIVO AND IN VITRO

Parameter SBF (Simulated Body Fluid) Tissue Fluid

Cellularity Acellular Highly cellularProtein content No protein Protein-richReplenishment Non Fluid turnover (blood circulation)Variation in solution

composition during contactLimited to

dissolution–precipitationUnlimited, due to chemotaxis and

macrophage activitypH 7.4 Could reach 9 due to wound healing

mechanismTemperature 37$C Could reach 40$C due to wound healing

processes



24 11. Biomaterials

FIGURE 11-17. Schematic diagram representing the events which take place at the interface be-tween bioactive ceramics and the surrounding biological environment (Reprint from Ducheyne P,Bianco P, Radin S, Schepers E, Bioactive materials: mechanisms and bioengineering considerations.In: Ducheyne P, Kokubo T, van Blitterswijk CA, eds. Bone-bioactive biomaterials. Leiderdorp, TheNetherlands: Reed Healthcare Communications, 1993:1–12, with permission.)

(2) precipitation from solution onto the ceramic(27,34,106,126,134); (3) ion exchange andstructural rearrangement at the ceramic-tissueinterface (3,27,34,78,126,145); (4) interdiffu-sion from the surface boundary layer intothe ceramic (46); (5) solution-mediated effectson cellular activity (71,86,129,145,159); (6) de-position of either the mineral phase (a), or theorganic phase (b), without integration into theceramic surface (27,33,34,55,106); (7) depo-sition with integration into the ceramic (27,106,145); (8) chemotaxis to the ceramic sur-face (145); (9) cell attachment and prolifer-ation (33,71,86,118); (10) cell differentiation(145); and (11) extracellular matrix formation(33,55,118).

The observation of what transpires at the in-terface, however, does not represent a mech-anistic explanation for the effect which bio-ceramics have on bone tissue formation. Wecan focus on mechanisms by relying on an in-creasing body of evidence which suggests thatbone bonding and bone tissue ingrowth en-hancement are the result of multiple, parallel,and sequential reactions at the material–tissueinterface. These interactions either are related

to physicochemical phenomena that occur inthe presence or absence of cells, or are relatedto reactions affected by cellular activity. Animportant aspect of the overall reaction sequencebetween these materials and tissues is that, in theabsence of a biologically equivalent, calcium-deficient, carbonate-containing hydroxyapatite(HCA) surface upon implantation, dissolution,precipitation, and ion-exchange reactions leadto a biologically equivalent apatitic surface onthe implanted material. This reaction does notproceed by itself, but is accompanied by paral-lel reactions, such as adsorption and incorpora-tion of biological molecules and attachment ofsurrounding cells. Furthermore, cells that haveadhered to the ever-reacting material surface in-teract with the material and produce some of thesurface changes. In the inverse direction, i.e.,from material to environment, there is both asolution-mediated and a surface-controlled ef-fect on cellular activity, organic matrix depo-sition, and mineralization. The gradual changeof the ceramic surface to become a biologicallyequivalent HA with small crystal dimensions isa rate-determining step in the cascade of eventsunderlying bioactive behavior. All phenomena,



11. Biomaterials 25

collectively, lead to the gradual incorporationof the bioactive implant into developing bonetissue.

7.5 Bioactive Glass Ceramics

Approaches to achieving enhanced mechanicaland biochemical properties included transfor-mation of bioactive glass into glass-ceramic, andusing bioactive glass as a coating on a substrateor as a second phase in a composite combin-ing the characteristics of both into a new mate-rial (77). In these techniques, bioactive glass issubjected to thermal treatments which alter thematerial’s microstructure and limit its biologi-cal activity. Thermal treatment of bioactive glass(45S5) in the temperature range 550$ to 680$Cresulted in the formation of Na2Ca2Si3O9crystals (61). Raising the treatment tempera-ture to 800$C led to the development of cal-cium phosphate crystals similar to the structureof hydroxyapatite (78). Other investigatorsshowed that when bioactive glass in the systemMgO-CaO-SiO2-P2O5 was thermally treated,the glass transformed into apatite/wollastonite(A-W) glass ceramic (97). According to powderX-ray diffraction, the crystallized glass consistedof 38% apatite, 34% wollastonite, and 28 wt%residual glassy phase (97). The resulting bioac-tive A-W glass ceramic is characterized by a highmechanical strength (Table 11-11) and has beentab.11.11used successfully for load-bearing prostheses (75,98,177). However, in vivo studies showed thatthe improved mechanical strength of bioactiveglass ceramic is at the expense of the rate of bonebonding to the material (98,109).

An important change in the glass microstruc-ture, which usually precedes crystallization, is the

TABLE 11-11. PHYSICAL PROPERTIESOF A-W BIOACTIVE GLASS CERAMIC

Density (g/cm3) 3.07Bending strength (MPa) 215Compressive strength (MPa) 1,080Young’s modulus (GPa) 118Vickers hardness (HV) 680Fracture toughness (MPa 1/2) 2.0

glass-in-glass phase separation (112). After sep-aration, the glass no longer has a homogeneouscomposition, but rather, consists of two or moreimmiscible glassy phases of different chemicalconstituents. Unlike crystallization, phase sepa-ration in glass might not be visible by opticalmicroscopy and in most cases can be detectedonly by electron microscopy. Previous work re-vealed that thermal treatment of bioactive glassat 550$C resulted in glass-in-glass phase sep-aration and minor crystallization (58). Thesechanges in bioactive glass microstructure alteredthe corrosion behavior and led to a significantincrease in the negative ZP of the material. Inconjunction with the increase in negative ZPthere was a significant decrease in the amount ofserum protein adsorption onto the material sur-face. The inhibition of protein adsorption couldbe responsible for the slow rate of bone bond-ing to bioactive glass ceramic. It is also possiblethat conformation changes inhibited the activityof the protein adsorbed onto thermally treatedbioactive glass. Moreover, alteration of corrosionbehavior after bioactive glass crystallization de-layed the formation of the hydroxyapatite sur-face layer necessary for bone bonding.

8 TISSUE ENGINEERING ANDRECENT DEVELOPMENTS

8.1 Bone Tissue Engineering

The field of tissue engineering has recentlyemerged as a result of growing understanding ofthe mechanism of action of biomaterials in vitroand in vivo combined with latest advances incell isolation, culture procedures, molecular bi-ology, and biochemistry techniques. Tissue en-gineering offers a promising new approach torepair fractures that do not heal, major boneloss, and fractures associated with bone tumors.In this approach, a porous scaffold that seedsthe cells can serve as a vehicle to deliver cellsnecessary for tissue regeneration. With this ap-proach, it is possible to use the patient’s owncells so that immunosuppression is not needed.The design of a scaffolding material can signif-icantly affect the cell seeding and growth both



26 11. Biomaterials

in vitro and in vivo (111). Scaffolding materi-als should be biocompatible and biodegradable.The degradation products should be nontoxicand easily excreted by metabolic pathways. Thescaffold material should be mechanically strongto maintain structural integrity during culture.The ideal scaffolding materials also should beeasy to fabricate into a desired shape and have acontrolled porous architecture to allow for cellgrowth, tissue regeneration, and vascularization.The scaffolding materials should be osteocon-ductive so that osteoblasts and osteoprogenitorcells can adhere, migrate on the scaffolds, dif-ferentiate, and synthesize new bone matrix. Thenew tissue should also have the potential to growand remodel, which is especially important forpediatric patients.

Use of degradable scaffolds to support cellproliferation and extracellular matrix synthesisimplies a transient microenvironment that mayinfluence cell behavior. Poly(!-hydroxyl acids),such as poly(L-lactic acid) (PLLA), poly(glycolicacid) (PGA), and poly(D,L-lactic acid-co-glycolicacid) (PLGA), satisfy many of these material re-quirements and already have been fabricated intoscaffolds for cell transplantation and tissue engi-neering (81,84,178,179). One of the disadvan-tages of these materials is that the degradationproducts reduce the local pH value, which, inturn, may accelerate the polyesters’ degradationrates (157) and induce an inflammatory reac-tion. Another disadvantage is that the mechan-ical properties of the highly porous scaffoldsmade from poly(!-hydroxyl acids) are relativelyweak, which limits their use for bone-tissue re-generation, especially in the in vivo implant site.Incorporation of HA with PLLA resulted in acompact composite with little porosity and im-proved mechanical strength. By monitoring for24 weeks the pH value variations during theirincubation in phosphate-buffered saline (PBS),it has been found that the pH value of PLLA/HAcomposite is more stable than that of pure PLLAand pure HA (142,167).

Hydroxyapatite ceramic and its compositesare composed of the same ions as bone andare considered to be suitable candidate mate-rials for the synthesis of bone tissue engineer-ing scaffold. However, hydroxyapatite ceramic,

which is stoichiometric hydroxyapatite, resorbsvery slowly or remains unchanged in bone sites(95,114,115). Bioactive glass (45S5) is an at-tractive scaffold material because it stimulatesosteoprogenitor cell function and possesses con-trolled resorbability (39,43,52,55,76,146). Thegoal is to use the in vitro synthesized bone asa graft material to treat major bone loss in or-thopedic or maxillofacial surgery. Modificationof the bioactive glass surface before cell seedingserves two goals: (1) to minimize the adverse ef-fect on medium pH due to the accumulationof the bioactive glass degradation products and(2) to create the bioactive hydroxyapatite sur-face layer, which otherwise will not form in thepresence of serum. It was discovered that the cal-cium phosphate layer that forms on the surface ofbioactive glass is highly selective toward adsorp-tion of serum fibronectin, an attachment pro-tein known to enhance osteoblast adhesion andactivity (53,54). Other bioactive materials suchas hydroxyapatite ceramic or unmodified bioac-tive glasses do not have any selectivity toward fi-bronectin. Thus, when bone cells were culturedon bioactive glass pretreated to be covered witha dual layer of calcium phosphate and serumprotein, abundant and expeditious bone tissuewas formed on and within the porous template.Figure 11-18 is a scanning electron micrograph Fig.11.18of the cross section of porous BG showing theformation of extracellular matrix and boneliketissue after 7 days in culture. The cells also exhib-ited a high alkaline phosphatase activity and syn-thesized collagen type I and osteocalcin (57).

8.2 Engineering Surfaceswith Peptides

Adherence of osteoblasts to a biomaterial sur-face is dependent on topography, chemistry, andsurface energy. Cell attachment activity is also re-lated to a tripeptide sequence, arginine-glycine-aspartic acid (RGD), located in the cell-bindingdomain of many adhesion molecules (144). Fol-lowing attachment, there is reorganization of cy-toskeletal proteins, resulting in the flattening andspreading of the cell and formation of focal con-tacts, replete with clustered integrins. This formof cytoskeletal organization participates in cell



11. Biomaterials 27

FIGURE 11-18. SEM micrograph of thefractured surface of the porous modi-fied bioactive glass template seeded withneonatal rat calvaria osteoblasts for 7days. Note the uniform formation ofextracellular matrix and bonelike tissuethroughout the porous disk. (With per-mission from El-Ghannom A, Ducheyne P,Shapiro IM. Bioactive material templatefor in vitro systhesis of bone. J BiomedMater Res 1995;29:359–370.)

signaling events and serves to regulate cell be-havior (67). However, the adsorption of attach-ment proteins may not be successful in promot-ing long-term cell attachment to biomaterialsin vivo because of several limitations. These in-clude protein desorption and/or exchange anddenaturation, the removal of the oligopeptidesequences due to proteolysis, difficulties in con-trolling the sequence presentation (i.e., C- orN-terminus binding), and specifically arranginga sequence or a series of sequences on the surface(117,138).

In order to derivatize in a stable manner thesurfaces with RGD-sequence peptides, covalentimmobilization of the active peptide is neces-sary. The use of covalently coupled small pep-tide mimics allows control over the density andorientation of ligand attachment. Many investi-gators have bound RGD peptides to polymericsurfaces and evaluated cell activity (116,132).These studies have shown that the chemistryof the surface layer can modify the conforma-tion of the peptide and thereby change its in-teraction with cells. For example, Drumhellerand Hubbell showed that complete cell spread-ing on poly(ethylene glycol) networks requireda peptide density three orders of magnitudegreater than that required for silanated glass(38). A 15-mer peptide containing the RGD

motif grafted onto a quartz surface enhancedthe strength of bone cell attachment and pro-moted spreading and focal contact formation(141). In a parallel investigation, the surface ofquartz was engineered to contain different ra-tios of cell binding (–RGD–) and heparin bind-ing (–FHRRIKA-) domains (140). It was foundthat heterogeneous peptides containing –RGD–and -FHRRIKA– in the ratio 75:25 or 50:50enhanced cell spreading and mineralizationcompared to homogenous surfaces containingeither –RGD– or –FHRRIKA–. Results of thestudies discussed above provide evidence that theimmobilized peptides containing the RGD se-quence serve to enhance cell attachment to poly-meric and ceramic surfaces. In most cases, thebiomaterial exhibits a new topography, surfacechemistry, and functionality depending on themodification strategy. Even the method chosento activate the surface (e.g., wet versus dry chem-istry) can alter the biomaterial surface properties.

Cell behavior on surfaces grafted with theminimal tripeptide RGD devoid of any flank-ing groups or spacers was investigated. Usingan atomically smooth silicon substrate coveredwith a reproducible oxide layer, the oxide sur-face was activated using 3-aminopropyl tri-ethoxysilane (APTS) (51). This layer exhib-ited nanoscale roughness and presented amine



28 11. Biomaterials

functional groups for grafting the RGD tripep-tide. The attachment of RGD to the APTS sur-face was confirmed by time-of-flight secondaryion mass spectrometry analyses. Contact anglemeasurements indicated that the hydrophobic-ity of the APTS surface was significantly lowerthan that of the surface with grafted RGD(RGD-APTS). Atomic force microscopy rough-ness measurement values showed that surfacescovered with RGD-APTS were smoother (Ra =0.71 nm) than those covered with APTS alone(Ra = 1.59 nm). MC3T3-E1 bone cells wereseeded on these various substrates and incu-bated in a serum-free medium for different pe-riods of time. Whereas attachment was simi-lar on both APTS and RGD-APTS surfaces,cell spreading and cytoskeletal organization wereclearly favored on the relatively smooth graftedpeptide surface. Thus, neither surface roughnessnor the amine group alone of the APTS layeraccount for improved cell–biomaterial interac-tion. These data suggest that the combination ofsurface roughness and RGD modification pro-vides the most favorable surface for cell spread-ing and cytoskeletal organization. These eventssignificantly enhanced downstream maturation-related activities such as those linked to cytodif-ferentiation and mineralization of the extracel-lular matrix.

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