SINTERING TEMPERATURE EFFECTS ON THE MECHANICAL PROPERTIES OF

POROUS-COATED TI-6AL-4V €LI ALLOY

Jeff Archbold

A thesis submitted in conformity with the requirements for the degree of Master of Applied Scierice

Graduate Depart ment of Metallurgy and Material Science University of Toronto

Copyright by Jeff Archbold 1 999

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Cementless fixation of orthopaedic implants through bone ingrowth allows long-term implant anchorage in bone. Sintering of metallic particles onto the substrate of an orthopaedic implant results in a porous surface structure with a 3-dimensional interconnected pore geometry into which bone tissue can attach. Shedding of sinter-bonded particles from the substrate due to loading of the implant can resuit in unacceptable third-body Wear and loosening of the implant. This study examined the shear strength of the particlelsu bstrate interface for specimens sintered at different temperatures. Fatigue properties of sintered porous coated Ti-6Al-4V EL1 ratating beam fatigue specimens and notched, as-sintered Ti-6AI-4V EL1 axial fatigue specimens were studied as a function of sintering temperature to ensure that increased sintering temperature did not adversely affect the fatigue endurance strength.

The alloy studied was Ti-6AI-4V €LI. The particles used were atomized spherical particles approximately 150pm to 250pm in diameter (-60/+8O rnesh). Sintering, and heat treatment to simulate sintenng, was dune at 1 2 5 0 * ~ , 1 275'~, 1 ~OO'C, and 131 5-c in a vacuum ( 4 0 - 5 Torr ) furnace.

The geometry of the particlelsubstrate interface region was examined using scanning electron microscopy to determine sinter neck contact diameters. Higher sintering temperatures were observed to increase the average sinter neck diameter, and hence sinter neck area, with a dramatic increase occurring between 1300'~ and 131 5 ' ~ .

The average shear force required to shear the particles from the substrate varied from 960N for the specimens sintered at 1250'~. up to 1880N for the specimens sintered at 131 5-C.

Fatigue strength was determined using both ASTM axial fatigue test specimens at a fatigue stress ratio of R=0.1, and R. R. Moore rotating beam test specimens at R=-1. The rotating beam fatigue specimens were coated with 150pm to 250prn diameter (-60/+80 rnesh) sphencal particles and sintered at the prescribed temperatures.

There was no significant difference in fatigue strengths between neither the notched as-sintered axial fatigue specirnens heat-treated at 1275 '~ .

1 ~ O O ' C , and 131 5 ' ~ , nor for the porous-coated rotating beam fatigue specimens sintered at 1250'~ and 131 5 ' ~ .

Therefore, this study suggests that by increasing the sintering temperature for porous coated implants, the interface shear strength of the particle/substrate interface can be increased without a significant reduction of the fatigue strength, at least within the temperature range studied. The machined notched specimens appeared to model the fatigue behavior of the porous coated specimens closely.

Acknowledgments

Many people have helped me with this research. I have received support in everything from data collection to brainstorming to a sympathetic ear when things did not go as onginally planned. I wish ta thank al1 of you: David Abdulla, Keith Carter & Innova, Robert Chemecky, Mark Fitiaggi, Chris Pereira, Robert Piltiar, and Jeff Wetls.

TABLE OF CONTENTS

PAGE

ABSTRACT Acknowledgrnents Table of Contents lndex of Tables lndex of Figures

ii

iv v viii ix

INTRODUCTION 1 Sintering 4

The Stages of Sintering 4

The Mechanisms for Material Transport During Sintering 6 Driving Forces for Material Transport During Sintering 10

Sintering for Orthopaedic and Dental Implants 12

Titanium-6Aluminum-4Vanadium EL1 13 Ti-6AI-4V EL1 for Orthopaedic Implants 14

Fatigue 15 Fatigue Crack Initiation 16 Fatigue Crack Growth 17

Fatigue of Titanium Alloy 19

Fatigue of Porous-Coated Titaniurn Alloy 19

2 OBJECTIVE 21

3 METHOOS AND MATERIALS 3.1 Specimen Preparation 3.1 .1 Sintering 3.1.2 Interfacial Shear Strength Specimens 3.1.3 Fatigue Specimens 3.1.3.1 Axial Fatigue Specirnens 3.1 -3.1 .1 U n-Notched Axial Fatigue Specimens 3.1 -3.1.2 WireEDM-Notched MA Axial Fatigue Specimens 3.1.3.1.3 WireEDM-Notched BA Axial Fatigue Specimens 3.1 -3.1.4 Surface Finishing of Axial Fatigue Specimens 3.1.3.1 -5 Measurement of Cross-Sectional Area

of Axial Fatigue Specimens Rotating Beam Fatigue Specimens Metallographic Characteriration Hardness Testing of Notched Axial Fatigue Specimens Scanning Electron M icroscopy Sinter Neck Measurements Particle-Center-Approach Measurements Finite Elernent Analysis Mechanical Testing Testing of Shear Specimens Testing of Axial Fatigue Specimens Testing of Rotating Beam Fatigue Specimens

RESULTS Sinter Neck DiametedParticle Diarneter Ratios Shear Force Microstructure and Grain Size Hardness FEM Analysis of Notch of Axial Fatigue Specimens Fatigue Strength Axial Specimen Fatigue Strength Rotating Bearn Specimen Fatigue Strength

5 DISCUSSION 5.1 Sintering 5.2 Interfacial Shear StrengtMFracture Characteristics 58 5.3 Hardness 5.4 Fatigue Strength 5.4.1 AxialSpecirnenFatigue 5.4.2 Rotating Beam Specimen Fatigue

6 Conclusions

REFERENCES

APPENDICES

vii

INDEX OF TABLES

Chemical Composition of Materials Sinter Neck Ratio and Height Summary ANOVA and Paired T-Tests for Contact Ratio ANOVA and Paired T-Tests for Sinter Neck Height Shear Force Summary ANOVA and Paired T-Tests for Shear Force Grain Size Summary Axial Specimen Fatigue Results Summry Rotating Beam Fatigue Results Summary

PAGE

viii

INDEX OF FIGURES

Sinter Particle and Sinter Neck Sintering Diagram for Ag Spherical Particles Sintering Diagram for fi Spherical Particles 1 nterfacial S hear Strength Specimen Shear Tests Axial Fatigue Specimen WireEDM Notches Rotating Beam Specimen Porous Coated Rotating Beam Specimen SEM Micrograph of Hardness Test of Failed

Axial Fatigue Specimen Single Notch Specimen FEM Mesh Contact Ratio vs. Sintering Temperature SEM of Sintered, Porous Coated Rotating Beam

Fatigue Specimen, 1 2 5 0 ' ~ SEM of Sintered, Porous Coated Rotating Beam

Fatigue Specimen, 1 31 S'C Shear Force vs. Displacement Graph 131 5C-b Light Micrograph of Microstructure of Rotating Beam

S~ecimen Sintered at 1250-c Light Micrograph of Microstructure of Rotating Beam

Specimen Sintered at 131 S'C Hardness for 1275 Specimen FEM Analysis of Notch

PAGE

5 9 9 24 24

28 28 31 31

34 36 43

44

45 46

48

49 51 52

SEM of Surface of Sheared Specimen Sintered at 1 3 1 5 ' ~ 61 SEM of Fracture Surface of MA

Axial Fatigue Specimen 64 SEM of Fracture Surface of BA (1 275C)

Axial Fatiaue S~ecimen 64

INTRODUCTION

Total hip arthroplasty, while an extremely invasive procedure, has become a common surgical procedure'. For older, or less active implant recipients, cemented femoral prostheses are generally preferred due to their lower cost and the ability for patients to safely apply load to the implant soon after surgery. This avoids potentially debilitating periods of inactivity. For younger, o r more active patients, for whom acrylic bone cernent may not provide reliable long-term fixation due to the mechanical property limitations of the c~ rnen t , cementless implants, with their special surfaces ont0 w l i i ~ i t bone tissue can attach, are generally preferredz. Bone tissue is constantly being resorbed and produced by the body. This rneans that the mechanical bond between a cementless implant and surrounding bone tissue wi l l be constantly renewed and maintained ove? the lifetime of the implant recipient, under ideal conditions.

A variety of cementless implant designs using different rnethods of preparation are currently available, including variations of plasma-sprayed coatings and porous surfaces. Plasma-spraying of implant substrates with nydroxylapatite (HA) ceramics creates a surface on the implant with a chernical composition sirnilar to bone tissue. These surfaces have been found to improve the apposition of bone tissue on to an implant during healing. Plasma sprayed HA coatings tend to be resorbed by the human body over prolonged periods of implantation due to amorphous regions with high solubility, diminishing their long-term osseointegra!ion properties. Titanium plasma-sprayed coatings exhibit a 'roughened" surface morphology. which improves osseointegration with respect to smooth, press-fit implants. There is, however, no three- di mensional interconnected porosity with a plasma-sprayed coating into which bone tissue can mechanically interlock. Surfaces with three-dimensional interconnected porosity are preferred for good long-term fixation of the implant within bone'. A sintered porous coating can provide this surface, with the porous zone being formed by sintering metal powders or wires. For multi-Iayered

particle designs, a surface with a 3-dimensional interconnected series of pores is created. The s i re and shape of the pores can be varied through the use of different sized metallic powders or other particle forms, including fibres, wires, and mesh,

While plasma-spraying and sintering metal particles can create surfaces which promote long-term fixation of bone to the implant, there i s an associated negative effect on the mechanical properties of the implant that should be considered prior to use.

Due to the forces experienced by an implant, and in particular femoral implant components of total hip joint replacements, debonding of sintered particles (pawder or wires) from one another and the substrate i s a concern. Sintered particles (especially the sp herical powder particles used by many implant manufactu rers) which break away from the surface of the impfant compromise bone/implant fixation3. This can necessitate implant revision due to loosening of the implant i n the femur. In addition, loosened particles can migrate along the femoral implant to the acetabular cup region of the implant. If the particles become entrapped between the metallic head of the femoral component and the ultra- hig h-rnolocular-weight-polyethylene (UHMWPE) acetabular cup, increased Wear of the UHMWPE cup due to the third-body contamination can occur. The UHMWPE Wear particles are believed to be responsible for the massive endosteal osteolysis which is often observed radiographically in failing implants4. In these cases, not only does the implant become loose, but the severe loss of bon8 tissue around the implant makes further revision difficult.

The sinter neck region, that i s the region of bonding between particles and substrate, acts as a stress riset, which has been shown by others ta dramatically diminish the the fatigue strength of titanium a l l o y ~ ~ ~ ' * ~ . In addition, the sintering treâtment causes a change in microstructure of Ti-6AI-4V alloys from Mil l Annealed (MA) condition (fine equiaxed a-grains with uniform distribution of P-phase particles at the grain boundaries) to a Beta Annealed (BA) structure (lamellar a-phase and pphase regions). This has been found to reduce fatigue strengths of unnotched Ti-6AI-4V alloys by

approximately 10 to 20%'. Higher sintering temperatures should result in increased grain size. Larger grain sizes generally lower the yield strength of an alloy and can reduce resistance to crack initiation, which may decrease its fatigue endurance strengtha?

The resistance to fatigue fracture represents a primary requirement of load-bearing implants. Fortunately, failures of femoral hip implant components due to fatigue are rare3 because implants are made sufficiently large in cross-sectional area to ensure stresses acting during normal use are below the fatigue lirnit or the endurance strength corresponding to the number of cycles expected to be experienced in the Iifetime of an implant. For convenience, the fatigue endurance strength corresponding to 1 O' cycles is used. lncreasing the cross-sectional area of a well- bonded implant results in a greater portion of applied force being borne by the higher stiffness implant and a corresponding lower force acting on the host bone. Due to the nature of bone remodelling, this can result in the reduction in the quantity of bone tissue around an implant over time. The loss of bone tissue around the implant makes bon8 fracture more likely and implant revision more difficu lt. Additionally, for implants made of alloys displaying high notch fatigue sensitivity (e.g. Ti-6AI-4V). zones of porous- coating, with their associated stress concentration regions, are not placed over areas of the implant experiencing high tensile stresses, Le. lateral aspects of femoral stems. However, this introduces other problems such as intrinsic tracts for polyethyfene Wear debris migration (from implanted prosthetic acetabular cups) to more distal implant zones, the reduction of osseointegration zones where bone is least likely to be resorbed over extended periods of implantation due to the high transfer of loading force from implant to bone tissue, and weaker fixation of the implant in the bond0.

This study examines one approach to improving the sinterbo nding of sp herical powder particles to the substrate. Further, the effect of implementing this approach on the fatigue properties of notched and porous coated Ti-6AI-4V €LI specimens i s examined.

1 -1 Sintering

Sintering is a process by which bonding between powder/particulate matter i s achieved as a result of atomic diffusion at elevated temperatures in appropriate furnace conditions. The application of pressure during sintering will also promote atomic diffusion processes resulting in increased degrees of powder/particulate bonding. Localized melting may occur (liquid phase sintering), although sintering is possible without melting. The sinte ring process decreases porosity. and hence results in densification of structures. All sintering, however, involves thermally-induced material transport.

Sinterbonds may be formed between particles and particles, particles and solid substrates, fibres and fibres, and fibres and solid substrates. The common factor in these examples is that there i s a least one component possessing a high surface area-to- volume ratio. This is important as sintering promotes a reduction in the free energy of a system, and the diminution of surface area and surface free energy is a prime driving force in the creation and growth of bond regions between components.

1.1.1 The Stages of Sintering

The preparation of cornponents or regions of components for sintering involves initially the preparation of a 'green' state structure characterized by particlelparticle contact without interparticle bonding. At this point there is only adhesion between particles, and the volume density can be approximately 213 of full density (depending on the particle size distribution and shape). This is termed the 0th Stage of sintering.

Once heat is applied to the 'green' structure the first stage of si ntering begins. The contact points between powder particles become bonded as material transport (See 1.1.2) creates 'sinter necks' between particles, see Figure 1.1. With increased sintering time/temperature/pressure the sinter necks continue to increase in diameter, up to a sinter neck diameter/particle diameter of

approximately 0.4. There i s very litt le densification of components at this point. The particles remain discreet and clearly defined, al though grain boundaries do form between part ic lesl 1.

Individual particles, characterized by sinter neck contact points i n the first stage of sintering, begin to coalesce in the second stage of sintering. Most of the densification occurs in the second stage of sintering. The pores tend to a cylindrical shape between particles. Grain growth is significant.

The third. and final, stage of sintering leads to densities of 90% - 95% of the theoretical full density. Pores, which had been interconnected i n the first stage, become increasingfy isolated and spheroidised. These sphericai pores shrink slowly in the cases where there are no entrapped gases (vacuum sintering) or the entrapped gases are soluble i n the matrix. Pores containing insoluble gases shrink unt i l the pressure of the entrapped gas is i n equi l ibr ium with the sintering-induced contracting pressure of the pore.

FIGURE 1.1: Sinter Particle and Sinter Neck

1.1.2 Mechanlsms for Material Transport Duting Slnterlng

There are four different mechanisms for material transport i n the sinte ring process of crystalline materials; surface diffusion, grain boundary diffusion, volume diffusion, and evaporation & recondensation. For amorphous materials, viscous flow can cause significant material transport.

Evapo ration & recondensation i s rarely a significant contributor to material transport i n the sintering process due to the (usual ly) minute partial pressures of metals and cerarnics. Two exceptions are uranium oxide (UOz) and T-iront2. When evaporation and recondensation does take place, the distances between the centers of adjacent particles do not change. There i s no 'particle-center-approach'.

Of the three rernaining mechanisms of material transport for crystalli ne structures, surface diffusion normally has lower activation energy than either grain-boundary diffusion or volume diffusion. For this reason, surface diffusion is the mechanism which most contributes to the formation and early growth of sinter necks and why surface diffusion is usually the predorninant mechanism during first stage sintering". The mobilty of surface ato ms, driven by thermodynamics and surface energy differences, leads to material transport from highly convex surfaces (hig h energy zones) to highly concave surfaces (lower energy zones). Sinter neck zones are characterized by large differences in free energy between sharply concave sinter neck zones and convex particle macrosurfaces. Once a certain sinter neck diameter/particle diameter ratio i s attained the differences between concave and convex zones diminishes. The effect of surface diffusion then ceases to dominate. As with evaporation & recondensation. there is little particle-center-approach with surface diffusion.

Grain boundary diffusion and volume diffusion are the mechanisms which cause the greatest densification of particles

during sintering. Grain boundary diffusion i s associated with a lower activation energy than volume diffusion, but i s not always the mechanism which dominates material transport du ring the second stage of sintering. The example of sintering of copper wires with and without grain boundaries at the surface of the wires illustrates this13. Copper wires with grain boundaries a t the surface exhibit dominance by grain boundary diffusion for sintering conditions generalized as temperatures less than T/T,=0.75 (T in 'K). Copper wires without grain boundaries at the surface do not exhibit any dominance by grain boundary diffusion. Surface diffusion dominates the entire first stage of sintering, followed by volume diffusion. Both volume diffusion and grain-boundary diffusion mechanisms cause particle-center-approach and as such they are the mechanisms that lead to densification i n stages two and three of sintering. Depending on the sintering temperature, the time of sintering, or the phase of the metallmaterial, volume diffusion or grain boundary diffusion can dominate.

Pioneering work by Ashby and CO-workers13 concluded that one of the four mechanisms for material transport during sintering predominates for given temperatures and sinter neck diameterlparticle diameter ratios. Ashby proposed the construction o f sintering diagrams to describe the the dependance of sintering mechanism on temperature and sinter neck diameter/particle diarneter ratios.

Sintering diagrams map the dominating mechanism for material transport for a given condition (Figure 1.2). lsochronic lines depict time of sintering. All mechanisms contribute at al1 times, but outside specifically defined areas of dominance in a sintering diagram, the contribution of other mechanisms is small or negligeable. The lines on the sintering diagrams denoting boundaries between dominance of mechanisms shift according to impurities, particle sire, applied pressure, and sintering atrnosphere. Some surface oxides promote sintering, as volatile oxides evaporate and leave a very pure surface on the particlesl1. Thick oxide layers with melting points greater than the metal substrate inhibit sintering, as the oxide layer forms a protective

shield which prevents the transport of metallic atoms. The rate of sintering for smaller diameter particles i s greater than that for larger diameter particles, with al1 other parameters equal, due to the influence of particle s i d 3 on sinter-rate equations (see 1.1 -3).

The application of pressure du ring sinte ring increases the rate of the sintering processi'. Volume diffusion from dislocation sources (compared to volume diffusion from surface and grain boundary sources) has a negligeable contribution to material transport when there is no externally applied pressure to sinter particles. Upon application of load, however, volume diffusion from dislocation sources has a significant contribution to sintering, and thus increases the sintering rate.

FIGURE 1.2: Sintering Diagram for Ag Spherical Particles13

FIGURE 1.3: Sintering Diagram for Ti Spherical ParticlesY2

T['C 1

200 WO , [O00 1400

25 bdrydiff from bdry ~ v o l l i f f t rom bdry

bdry diffusio n v o l urne @ff u si0 n from bdry t r a m b a r y

1.1 -3 Driving Forces for Material Transpoit Durlng Sintering

When sintering is performed with negligible pressure, the driving force that leads to the creation and growth of sinter necks, and eventually to densification, is the difference in surface energy due to differences in curvature between particles and sinter n e c k ~ l p o r e s ~ ~ . From Figure 1.1 the curvature difference K at the sinter neck is:

p is the radius of curvature for the sinter neck r i s the radius of the sinter neck

Surface diffusion, with the influence of the curvature of the particles alsa contributing, has:

K = (1 /p - 11 r) + 2lR for contact between two spherical particles

K = (1 l p - 1 / r) + 1 I R for contact between a flat substrate surface and a spherical particle, where R is the radius of the sinter particle

This expression does not tend to zero as r -. R. As the curvature difference must approach zero as the sinter neck radius approaches the particle radius, the expression is multiplied b y (1 - r/R) to give:

KI = (1/p - 1lr + 21R ) (1 - r/R) for surface diffusion

For transport of material via boundary diffusion and volume diffusion during Stage One sintering the curvature of the sinter

particles does not influence K. The curvature difference is labelled KL and reverts to K, = ( I l p - I l r )

For Stage Two and Stage Three sintering, where pores are cylindrical or spherical, and particles are no longer discrete. the curvature difference is reduced ta that of the pore:

p is the radius of curvature for the pore (which originally vas the radius of curvature of the sinter neck)

The tensile stress a at the sinter neck region becomes:

o = 9 K . cp i s the surface tension

This becomes the driving force for the increase in sinter neck diameter.

The growth of the sinter necks is given empirically as:

r = sinter neck radius t = time of sintering F(s) = function dependent on temperature, particle s i n ,

geometry, etc. n = f(materia1 transport mechanism)

For surface diffusion, 'n' has values reported from 5 to 7.5, and mainly between 6.5 and 7.5, depending on the sintering conditions. The value of 'n' for grain boundary diffusion is approximately 6. For volume diffusion, 'n' i s accepted as having a value of 4-5.

The rate of sinter neck growth due to surface diffusion is given byq3:

0, i s the surface diffusion coefficient 6, i s the effective surface thickness from which atoms can

diffuse y, i s the surface free energy R i s the atomic volume T is the temperature in 'K

The rate of sinter neck growth i s proportional to the third power of K,. KI increases with diminishing particle size (smaller particles are more convex than larger particles), and this i s why smaller particles sinter more quickly than larger particles. The difference in sintering rates between larger and smaller particles affects the time required to achieve certain r/R values. meaning smaller particles begin Stage Two sintering before larger particles. This could lead to the case where particles of different sizes, adjacent to one another, could exhibit different mechanisms for material transport during the sintering process.

1 .1.4 Sintering for Orthopaedic and Dental Implants

Sintering is used to cisate porous surface zones for bone ingrowth for bone-interlacing, permanent implants. By selecting conditions that limit sintering primarily to Stage One. three- dimensionally interconnected porosity, into which bone tissue may grow and anchor, results. The size and shape of the pores i s dictated by the starting size of the spherical particles that are sintered to the surface of the implant. Particles are placed as a coating onto a preformed substrate, with the sintering resulting in consolidation of this coating ta form an integral porous structure. Orthopaedic implants require a sintered coating on the surface of a

milled or forged substrate (in the case of titaniurn alloy implants). or cast substrate (in the case of Co-Cr implants). The shape of the implant substrate i s usuatly irregufar. This precludes using pressure ta hold particles on the surface of the implant prior to sinteri ng, with certain exceptionst4. A sacrificial binderfglue is used ta initially hold the sinter particles on the surface of the su bstrate. The binder i s an organic compound. i s sacrificial and formulated to burn off the implant at low temperatures ( < 5 0 0 ' ~ ) du ring sintering.

Titanium (and i ts alloys) i s classified as one of the light alloys. It has a density 43% less than that of steel, yet with comparable strength. The Young's modulus for titanium is 120GPa, versus 21 0GPa for most steels and Co-based alloys. It has excellent corrosion resistance, forming a very stable layer of titanium oxide when exposed ta air.

Commercially pure titanium (cpTi) has a hexagonal close packed (HCP) crystal structure below 8 8 2 ' ~ (a -phase). Above this t e mperatu re, the crystal structure transforms to body centered CU bic (BCC) atomic packing (P -phase). The addition of aluminurn as an alloying agent stabilizes the a -phase. Vanadium stabilizes the B - phase.

Titanium alloy with 6% aluminum and 4% vanadium i s the uworkhorse" alloy when high strength i s required. Tensile strengths of over 1000MPa are attained with this a l l ~ y ' ~ . The addition of the alloying elements (6% Al and 4% V) raises the transformation te mperature for Ti-6AI-4V to approximately 9 9 2 ' ~ " . At temperatures below 9 9 2 ' ~ , a two-phase a + p structure forms.

Ti-6AI-4V alloy i s usually supplied in the miIl annealed (MA) condition. It i s produced by mechanical deformation just below 9 9 2 * ~ , followed by a heat treatrnent at approximately 8 0 0 ' ~ . which i s in the a+p field". This comprises a mixture of both a phase and p phase in a fine grained, two-phase alloy. Equiaxed a-phase makes up the bulk of the alloy, with about 15% by volume of small

$-phase particles, located primarily at the grain boundaries and triple points.

Extra low interstitial, or ELI, is the ASTM F136-84 designation for Ti-6AI-4V alloys with oxygen. nitrogen, and hydrogen percentage content below 0.1 3, 0.05. and 0.012 weight percent respectively. If the interstitial oxygen. nitrogen, and hydrogen content i s significantly higher than this, embrittlement becomes a factor. In cyclic loaded applications, €LI i s desirable to improve fatigue performance and is required for most orthopaedic biornaterial applications.

As noted, above 9 9 2 ' ~ Ti-6AI-4V transforms to BCC. The rate of cooling from above this temperature determines the microstructure of the two-phase, a+p structure at iower temperatures. Water quenching can result in a fine lamellar a+p

structure. while slow cooling (l'C/min) from temperatures well above 9 9 2 ' ~ resufts in a coarse lamellar structure termed beta- annealed (BA). It is alsa possible to form other variants of this microstructu te designated as basket-weave structures. One way to create a basket-weave microstructure involves heating the titanium alloy above the a - p allotropie transformation temperature and then quenching. The titanium alloy is then held in the a range, usually at BOO'C for 24 hours. This treatment creates a broken-up- structure (BUS), a base microstructure of p-phase grains with interspersed a-phase1'. The BUS microstructure exhibits improved fatigue characteristics versus notched MA microstructures and bath un-notched and notched BA rnicrostructures17*1?

1.2.1 Ti-6AI-4V EL1 foi Orthopaedic Implants

Titanium alloy is the material of choice for many orthopaedic implants for a variety of reasons.

Surgical implants are exposed to body fluids ranging from acidic to alkaline. Titanium and its alloys are very corrosion resistant. This reduces the concern of long-term degradation of the implant itself as metallic ion release into the body is minirnized.

The oxides formed on titanium alloys appear to have a beneficial effect on the bone-interfacing (osseointegration) characteristics of titanium implants. especially compared to Co-Cr and steel alloy implants. The (predorninantly) TiOs layer has been

reported to be a preferied bonding site for hyaluronic acid r no lec~ les '~ . According to Albrektsson et al'', hyaluronic acid forms a base ont0 which proteoglycans attach. The proteoglycans form a cernent l ine of 'gluew ont0 which bone tissue attaches. Titanium thus is hypothesized to have a greater affinity for bone apposition than the Co-Cr and steel implant alloys.

The shear strength of Ti-6AI-4V is approximately 600MPa20. As noted, the Young's Modulus of titanium alloy i s 120GPa. while for Co-Cr and steel alloys the Young's Modulus is over 2OOGPa. The Young's Modulus for cortical bone is only 20GPa. The lower modulus of elasticity for titanium alloys versus steel alloys aids in minimizing bone tissue resorption around an implant. This i s a second principle reason why titanium alloy implants are preferred by many surgeons over Co/Cr implants.

1.3 Fatigue

Fatigue failure results from the repetitive application of force below the ultimate strength of an object/specimen. It is the result of the initiation and growth of cracks due to the repetitive loading unti l ultimate failure.

Fatigue testing i s usually perforrned by applying a cyclic force which varies from a minimum stress to a maximum stress. A simple sinusoidal force pattern i s commonly used. The difference in the applied stress gives the stress intensity factor range AK. K is a measure of the elastic stress at a crack tip. K is given by2':

Where a i s the average stress over the entire specimen, a is the length of the existing crack at the surface of the specimen (for cracks contained entirely within the substrate of a specimen, that i s

cracks that are not present at the surface of a specimen, the length of the crack is defined as 20.) and f(oAN) is a dimensionless parameter dependant on the specimen geornetry, crack geometry. and applied stress.

The ratio of minimum stress to maximum stress in cyclic loading is called the stress ratio R, (R = Gminlamm). The stress ratio affects fatigue properties due to the phenornenon of crack cfosure. Stress ratios which are negative indicate that the test specimen undergoes compressive stresses. This compressive loading acts to close the existing crack tip, retarding crack growth. Stress ratios which are positive have no compressive loading (except in the cases where R is only slightly greater than zero, and residual stresses act to compress the crack tip) and crack closure is inhibited.

There are three stages that describe the lifetime of an objectkpecimen resulting in fatigue failure. The first stage i s the initiation of cracks on slip planes due to dislocation interactions. The second stage is stable crack growth, where there is a linear log-log relation between the increase in crack length per cycle, d d d n , and the difference in stress intensity over the cyclic loading range, AK. Striations on the fracture surface are evident for the second stage of crack growth. Once the crack growth becornes unstable, i.e. log da/dn is non-linear with respect to log AK, then the third stage of crack growth has been achieved. Crack growth instability occurrs resulting in ultimate failure.

If a material or specimen exhibits an infinite life under cyclic loading below a certain threshold, this threshold is terrned the fatigue limit. Materials which do not exhibit a fatigue limit have a fatigue endurance strength defined as the stress level at which failure has not occurred after (typically) 10' cycles.

1 -3.1 Fatigue Crack Initiation

The first stage of fatigue is the initiation of fatigue cracks. All materials (except rnonocrystalline structures at O'K) have interstitial o r substitutional impurities, vacancies, and dislocations. Grain

boundaries are prime locations for these lattice imperfections. During cyclic loading. even at mean stresses below the yield stress, these imperfections result i n localized slip and the development of slip-planes (even perfect crystalline structures wi l l develop slip planes under sufficient loading. although at higher stresses than required when imperfections exist). These slip planes eventually emerge at the surface of a specimen/object. creating a surface 'slip line". The slip iine acts as a stress riser from which a microcrack can develop under continued cyclic loading. Following crack initiation, crack propagation rate then determines the remaining Iife of a specirnen.

Crack initiation may consume a lot of the fatigue life of a specimen/object. For this reason, grinding and polishing of the surface, or any other process to reduce surface defects, has a much more significant effect on fatigue life compared with improving crack propagation characteristics of a specimen.

1.3.2 Fatigue Crack Growth

Once a crack has formed, cyclic loading will cause the crack to increase in length, provided the stress i s of sufficient magnitude. The application of cyclic stress of sufficient magnitude causes stable crack growth at first, followed by accelerating crack growth unti l the crack length reaches the critical length. resulting i n f inal fracture due to catastrophic crack growth under the applied load conditions.

Crack growth from the surface of a specimen/object can be described using Linear Elastic Fracture Mechanics (LEFM) and the stress intensity factor range AK. The Paris equation for crack growth describes the relationship between crack length growth per applied cycle and AK for the second stage of fatigue life. the stable crack growth stage, i.e.,

C and m are empirical constants

A second formula, the Forman equation, describes both the second stage and third stage of fatigue life. where crack growth accelerates until K,. the critical stress intensity factor. is reached.

dadn = C(AK)' /[(1-R)K, - A K ]

LEFM can describe much of the fatigue life of a specimenlobject, however at the tip of a crack there is a plastic zone which influences the behavior when varying the stress ratio R and when cracks emanate from notches. For these conditions Elastic Plastic Fracture Mechanics (EPFM) bette? descri be crack growth and crack closure.

Because the crack tip geornetry predicts a crack tip of an almost infinitesimally small radius, theoretically the stress concentraion at the tip would approach an infinite value. Since there is a yield strength which limits the stress that can be sustained by the bulk material at the crack tip. plastic deformation about the crack tip results. Residual stresses generated by the plastic deformation act to close cracks and hence reduce crack growth. The stress ratio has a pronounced effect on crack closure, with greater R-values, for the same AK, reducing crack closure and lowering fatigue life.

Cracks originating at notches are affected by the plastic deformation zone normally created atound the notch tip. The crack growth rates for these cracks are higher than those for cracks in the LEFM bulk region of a specimen/object.

The growth of cracks in plastic zones under cyclic loading depends on the size of the plastic zone. The actual crack length increment is given by A a i .

(dddn),. is the constant amplitude crack growth rate for the given AK

is a variable dependant on the size of the plastic zone

1.3.3 Fatlguo of Titanium Alloy

Titanium alloys do not exhibit infinite life, and as such are described by a fatigue endurance strength. In the MA condition, smooth Ti-6AI-4V alloy specimens exhibit fatigue strengths of about 625-650MPa5*7*22. As-sintered smooth specimens exhibit different fatigue endurance strengths depending on whether machining of specimens occurs before or after the as-sintered heat treatment. BA specimens machined and ground after heat treatment have fatigue strengths around 5 0 0 ~ ~ a ~ * ' ~ . BA specimens machined and g round (i.e. smooth surface) before heat treatment have fatigue strengths reduced to around 375MPa5*7v"*22. Defects on the surface of the specimens that arise from t h e heat treatment (thermal etch lines) decrease the time required to initiate cracks at the surfaces of specimens heat treated after the machining and grinding steps. This lowers the fatigue limit versus specimens which have these defects machined and ground from the surface after heat treatment.

Grain s i te has been found to have a significant effect on fatigue endurance strength. The BA microstructure for Ti-6AI-4V. with relatively large grains, reduces long crack propagation rates versus finer grained MA Ti-6AI-4V, however the crack initiation and short crack growth of BA Ti-6AI-4V is inferiora*'.

Notched titanium alloy specimens have drastically reduced fatigue strengths. Reductions in fatigue strengths to befow 200MPa is quite common5*6*7J8*22 . High strength MA and BA titanium alloys have low fracture toughness and crack retardation c h a r a c t e r i s t i ~ s ~ ~ - * ~ . This rnakes many titanium alloys very susceptible to notch effects.

1.3.4 Fatigue of Porous Coated Titanium Alloy

An implanted femoral component of a hip implant experiences load with every step that the recipient takes. Assuming the average person walks 5km every day, over a year each hip is stressed approxirnately 10' cycles/year, Due to the high number of cycles

that would be experienced over the expected period of implant function. fatigue of the implant must be considered.

From 1.3.3 it is noted that there i s a drastic reduction in the fatigue strength of titanium alloy due to notching of the surface of fatigue test specimens. The porous coating on orthopaedic implants acts as a notch on the surface. The sinter neck, with a radius of CU Nature of approximately 3-5pm. concentrates stresses experienced by the bulk of the specimen/implant. Two-dimensional FEM analysis of sinter necks using varying sinter neck diameters and particte diameters found the stress concentration factor at the sinter neck to be approximately 2.5-35*23. A plastic zone around the sinter neck notch increases the ease with which cracks propagate (and the growth of cracks in porous-coated implants are best described by EPFM). This reduces the fatigue limit of rotating bearn fatigue specimens from above 625MPa for smooth MA specimens, and up to 500MPa for as-sintered smooth fatigue test specimens, to around 200MPa. The fatigue limit for porous coated specimens decreases with increasing spherical particle powder diameterlao2? Larger particles. however, provide larger pores, and as animal studies have shown, pore sizes above a certain minimum value appear necessary for allowing bone ingrowth to occur during the shortest tirne periods'sZ4. This limits the possibility of increasing fatigue life throug h geometrical means. Attempts to improve fatigue life thraug h microstructu ral changes have resulted in the use of the broken-up-structure (BUS) which has been found to increase fatigue strengths by about 50MPat0. This procedure i s expensive, and increases the initial cost of any implants which might use this heat treatment to improve fatigue life. Furthermore, it introduces concerns related ta possible contamination of the porous surface region.

Sintered porous coatings on orthopaedic implants made using powders, fibres, or wires, provide a good anchor for bone tissue ingrowth which usually results i n long-term fixation of the load-bearing implant within the body. There are, however, areas for possible improvement. lncreasing the bond strength betwee n the porous surface construct and the substrate as wel l as within the porous zone itsetf (Le. interparticle bonding) wauld reduce the likelihood of failure of the porous-surfaced region of the implant. This would ameliorate the problems of third body Wear and implant looseni ng due to substrate/particulate or interparticulate bond breakdown, which results in particles separating from the implant. Altering the sintering conditions may improve bonding of the sintered particles (or fibres) to each other and to the implant substrate. However, i f improved sinter bonding occurs by increased densification of the porous surface region, and hence results in loss of porosity required for effective fixation by bone ingrowth, or i f the process required to cause the improved bonding results in a significant diminuation in the fatigue endurance limit of the implant, altering the sintering conditions may not be practical.

Previous studies on the mechanical properties of porous- coaled implants focussed on either the sinter bonding betweer: particle and substrate or the fatigue strengths of porous-coated versus non-coated specimens. This study examines the effects on mechanical properties of porous-surfaced specimens by varying si ntering temperatures. Differences in sintering temperature should alter the microstructure of specirnens, as well as the diameter and geometry of the stress-concentrating sinter necks. This may affect both the shear strength of the sinterbonds and the fatigue endurance strength of the specimens.

Using Ti-6AI-4V metal powder for formation of the porous surface region, this study aims to;

1. determine the dependance of sinter neck diameter/particle diameter ratios & sinter neck geometry on sintering temperature, al1 other parameters being kept constant.

2. f ind the relationship between sintering temperature, sinter neck diameter/particle diameter ratios, sinter neck geometry, and the shear strength of the particle/substrate interface.

3. ascertain the effect of sintering temperature on the fatigue endurance strength of porous-coated fatigue specimens for the specific system studied (1 50pm-250pm-si zed Ti-6AI-4V powder particles sintersd to Ti-6AI-4V substrate). For control, the fatigue endurance strengths of heat-treated (as-sintered) notched fatigue specimens were studied.

This study i s intended to provide new information for the design of porous-coated implants such that the best compromise between shear strength of sintered particles, fatigue endurance strength of implants, and implant surface porosity for tissue ingrowth can be obtained.

3 METHODS AND MATERIALS

3.1 Specimen Preparation

3.1 .i Sintering

S inte ring was performed using a hig h temperature vacuum furnace, (Centorr hc.. Suncook. New Hampshire). The specimen heat treatment procedure began with a rough vacuum (5x1 Torr) of the furnace chamber. High purity hetium gas was used to backfill the furnace chamber, and then rough vacuum was again attained. A high vacuum (7x1 Torr at room temperature) was then achieved and held for 30min before heating of the specimens. The furnace took 20min to reach (T.-40)'~. with T. being the desired sintering temperature. It took a further 5min to attain the selected sintering temperature. This method of heat-up prevented temperature overshoot. The sintering temperature was held for two hours. Temperature was controlled by a Microstar Programmer, (Microstar Inc., Minneapolis, Minnesota). The in-furnace-cool (IFC) required approximately 2hrs to return to roorn temperature.

3.1 .2 Interfacial Shear Strength Specimens

The interfacial shear strength test specimens were designed and test methods developed to compare the interfacial bond strength between particles and substrat8 for the different sintering conditions used in this study. Figure 3.1 illustrates the test specimen geometry. This specimen design was developed at the Centre for Biomaterials for assessing the shear strength of the particle/substrate sinterbond interface.

The Ti-6AI-4V ELl specimen substrates were machined from 15mm (0.588") diameter Ti-6AI-4V EL1 rod (Dynamet Inc., Washington. PA Chemical Composition: Table 3.1) supplied in the miIl annealed (MA) condition. The substrates were prepared for coating with particles by first degreasing using acetone. The 15mm

iFIGURE 3.2: Shuw Torta

Not to Scale

Existing Bar Stock 14-950mm DIA (0.588' DI A)

Pororu Coaliflg

nxed End

Shearing Coating Ring

Shear Specimen

end of the specimens was masked with masking tape placed circumferentially approximately 3mm from the end of the specimen. A 3mm band of sintered particles was placed in this region, sufficient to ensure that a maximum force was recorded during testing before the entire coating was sheared from the substrate, but also narrow enough to prevent build-up of debonded particles at the shear collet resulting in excessive frictional resistance developing during testing. The end of each specimen was coated with an organic binder of 1 :10 parts acryloid:toluene. The specimen was then dipped i n titaniurn alloy powder. consisting of 150pm to 250pm particle diameters, (-60/+80 mesh particle size. Nuclear Metals Inc., Concord, Mass.). The chemical composition of this powder i s presented in Table 3.1. After particle coating. the binder was allowed to dry, resulting in particles adhering to the substrate via the organic binder. The interfacial shear strength specimens were sintered in the high temperature vacuum furnace at 1 2 5 0 ' ~ , 1 2 7 5 ' ~ . 1 3 0 0 ' ~ . or 131 5 ' ~ . There were 11 specimens sintered at 1 2 5 0 ' ~ . 10 at 1 2 7 5 ' ~ . 9 at 1 3 0 0 ' ~ . and 17 sintered at 131 5'~.

3.1.3 Fatigue Specimens

Two types of fatigue test specimens were tested. Notched axial fatigue specimens, heat treated at different temperatures, were studied to isolate the effect of temperature on fatigue endurance strength. Porous-coated rotating beam specimens. sintered at different temperatures, were studied to observe the effects on fatigue endurance strength by sinter neck geometry and multiple crack initiation sites (sinter necks), in addition to sintering temperature.

3.1 -3.1 Axial Fatigue Speci mens

3.1 -3.1 .1 Un-notched MA Axial Fatigue Specimens

Un-notched MA axial fatigue test specimens were machined frorn blocks of Ti-6AI-4V EU. (Figure 3.3). The Ti-6AI-4V €LI was supplied as 1.25"x3.00' (31.7mm x 76.2mm) bar stock in the MA CO ndition, (Teledyne Allvac. Monroe, NC. Chemical Composition: Tab le 3.1 ). Some warpage or bowing of the specimens occurred due to rnachining. To rninimize the bowing. test specimens were cold pressed until bowing at the center of the specimens was less than 1 OOpm, relative to the ends of the specimens. Bowing up to 100pm would be eliminated during the final grinding and polishing of the test specimen surfaces in preparation for fatigue testing.

To eliminate the added interna1 stresses generated by the coid setting of warped specimens, a stress relieving heat treatment was performed on al1 un-notched MA axial fatigue test specirnens. This treatment consisted of a ~ O O ' C , 15min anneal in vacuum with an IFC. The surfaces of the un-notched MA axial fatigue test specimens were then finished by the process described in 3.1 -3.1.4.

3 -1 -3.1 -2 WireEDM-Notched MA Axial Fatigue Specimens

Stress concentrating notches, which had been found in preliminary studies to simu fate the notch effect associated with the sinter neck regions of porous coatings on fatigue endurance strengths of titanium alloy, were machined into MA axial fatigue test specimens after the stress relieving heat treatment described in 3.1 -3.1 -1 .

The notches were machined into the sides of the axial fatigue specimens that had been prepared with a nominal thickness of 0.1 38" (3.50mm). The notches were machined by wire electrodischarge machining (WireEDM) using wire rated at 75pm (0.003") diameter, (Strite Industries. Cambridge, Ontario).

The wires were programmed to proceed 200prn (0.008') into the surface of the specimens once electrodischarge cornmenced (Figure 3 -4).

The surfaces of the WireEDM-notched MA axial fatigue test specimens were then finished by the process described in 3.1 -3.1 - 4 ,

FIGURE 3.3: Axial Fatigue Spodman

FIGURE 3.4: WireEDM Notches

t 0.003" MAX' (75um MAX)

3.1.3.1.3 WlreEDM-Notched BA Axial Fatigue Specimens

A number of axial fatigue test specimens were heat treated at 1 2 7 5 ' ~ , 1 3 0 0 ' ~ . or 131 S'C, creating a beta-annealed (BA) microstructure.

Prior to machining, a block of the MA Ti-6AI-4V was heat- treated for one hour at the desired temperature in the high temperature vacuum furnace as described in 3.1.1. The axial fatigue specimens were machined from this block. Warpage or bowing due to the machining was corrected to within 100pm and the specimens were then notched by WireEDM as described in 3.1 -3.1 -2 . Since WireEDM involves localized melting immediately followed by a water quench. the microstructure in the zone around the notch was not the same as the bulk of the specimen. To ensure a uniform microstructure throughout the test specimens, the specimens were heat treated a second time at the desired temperature for one hour. following the protocol described in 3 -1.1 . The total time spent at the desired temperature was 2hrs. consistent with the time for sintering.

After the second high temperature heat treatment the specimens were again cold pressed to minimize warpage and bowing. The stress relieving heat treatment described in 3.1 .3.1.1 followed.

3.1 .3.1.4 Surface Finlshing of Axial Fatigue Specimens

The axial fatigue specimens were ground and polished using an automatic polisher (Buehler Ecomet S. Buehler Inc., Lake Bluff, III.) using a method developed specifically for titanium alloy. (Dialog Method, Buehler Inc.). Two specimens at a time were mounted on a flat base fixture for attachment to the rotating machine head (Buehler Automet 2). The first step in grinding used 120 grit S i c grinding paper with water lubrication. This step was repeated until the warpage or bowing of up to 100pm was eliminated. The second step used 240 grit S i c grinding paper with water lubrication. Subsequent steps used 6pm diamond powder in

oi l suspension (Buehler #40-6542) and a platen, and with 1pm alurni na paste added approximately every 30sec (Buehler Alpha Micropolish t40-6361. diluted 3:1 with DH20) ont0 a polishing pad.

The above steps were repeated for the second side of the specimens.

3.1.3.1.5 Measutement of Cross-Sectional Area of Axial Fatigue Specimens

The cross-sectional area of the axial fatigue specimens was required to set up the loads for the fatigue tests.

The width measurement used to determine the stresses experienced by the test specimens was calculated as the measured width of the specimens minus 400pm, which was twice the nominal depth of each notch.

3.1.3.2 Rotating Beam Fatigua Specimens

The rotating beam fatigue specimens were machined from 1 5mm diameter MA Ti-6AI-4V EL1 rod stock, (Dynamet lnc., Washington, PA), (F igure 3.5.1). Testing was conducted using an R. R. Moore Rotating Beam test machine. The R. R. Moore rotating beam test i s an industrial test that has been used previously to test fatigue strengths of porous-coated biomaterials. The specimens were wet ground with S i c metallographic papers down to 600 grit, with the grinding direction being kept parallel to the longitudinal axis of the specimens. prior to coating with the 150pm-250prn diameter (-60/+80 mesh) Ti6-AI-4V €LI particles, (F igu re 3.5.2).

The ground specimens were degreased using acetone. Part of the waisted region was masked with masking tape leaving an exposed zone approximately 10mm long in the mid-length of the waisted region for coating with the Ti-6AI-4V EL1 powder. The rotating beam fatigue specimens were sintered in the high

FGURE 35-1 : wLri.ri

114- DIA 0.25" (6.35mm DIA) (6.35m)

0 0 . 5 0 ~ Nominal

f '-' b

2-50" RAD (63.5mm RAD)

# # . I I - 4 +-

1 Omrn (0.4")

temperature vacuum furnace for two hours at 1 2 5 0 ' ~ or 131 5 ' ~ . followed by an IFC.

After sintering, the ends o f the specimens were remachined to ensure specirnen concentricity. A special collet was fabricated to hold the specimens within the waisted region without disturbing the sintered porous-coated reg ion du ring machining.

3.1.4 Metallographlc Cheracterlzatlon

After fatigue testing. sections were cut from the end portions of the test specimens and prepared for metallographic examination. Four metallographic samples were prepared from each of the 1 2 5 0 ' ~ and 131 5 ' ~ rotating beam fatigue specimens, and four sarnples were prepared from the axial fatigue test specimens for each of 1 2 7 5 * ~ , 1 3 0 0 ' ~ . and 131 5 ' ~ heat treatment conditions. These sarnples were mounted i n thermosetting bakelite mount and then ground and polished using the protocol of 3.1.3.1.4, with the addition of a fifth step to remove the lprn scratches. This was a chemical polish using an alkali solution (Buehler Mastermet Y40- 6370) and a special polishing pad, (Buehler Chernomet 1 #40- 791 8 ) . The surface of the titanium alloy samples was etched using Kroll's Reagent (1-3 mL HF, 2-6mL HNOÎ, in 100mL H20) for about

30 seconds. Grain s i re was determined using a planimetric method according to ASTM €01 12.

3.1 .5 Hardness Testing of Notched Axial Fatigue Specimens

Prelirninary studies had shown that heat treatment of the WireEDM notched axial fatigue specimens after the machining of the notches improved the fatigue endurance strengths when compared with axial fatigue specimens heat treated before WireEDM notching. The single-notch, BA axial fatigue specimens from 3.1.3.1.3 were heat treated in two steps to minimire the microstructural differences between the bulk of the specimen alloy and the region of the specimen that was melted and quenched by WireEDM. Hardness testing was performed to determine i f fatigue

endurance strengths may have been affected by an embrittled zone around the notches. Sections were prepared from failed specimens of the notched MA and the three BA specimens. The sections of the specimens were approximately 15mm long and included the fracture surface, which originated at the WireEDM notch. These sections were mounted i n thermosetting bakelite. Polished surfaces were prepared by grinding with 120 grit S i c metallographic grinding paper. At least 200prn of the cut surface was removed. Final grinding and polishing was done using the protocol described in 3.1.3.1.4, and the chernical polish of 3.1.4.

Hardness testing was performed using a Leitz microhardness tester with tests performed using both the notch-edge of the specimen and the machined notch as Opm data reference points. See Figure 3.6. Indentations were made starting as close to the Opm data reference points as possible, typically 1 5pm. Indentations were repeated approximately every 1 Opm. although not in a direct radiating l ine from the Opm data reference points as this would not allow adequate spacing between indentations. Measurements were made up to a distance of approximately 75pm from Opm data reference points. Additional measurements were then made every 25pm (approximately) up to a distance of 200pm from O p r n data reference points. Hardness readings were recorded using the Vicker's scale.

FIGURE 3.6: SEM Micrograph of Hardness Test of Failed Axial Fatigue Specimen

3.2 Scanning Electron Microscopy

3.2.1 Sinter Neck Measurements

The sinter neck diameters between the 150pm-250prn Ti-GAI-

4 V EL1 particles and the substrate were determined using both sinter porous-coated disc-shaped specimens cut from the same rod as the shear specimens and the ends of shear specimens where undisturbed particlelsubstrate interfaces remained. The specimens were mounted in a scanning electron microscope (Hitachi S-2500. Mito City, Japan) with an integral Micro Scale Unit to measure

distances on the view screen. The sinter neck diameter was taken as the smallest diameter region between particle and substrate. The particles were not perfectly spherical so the diameter of the particles was taken as the greatest width of the particle in a plane parallel to the plane of the sinter neck.

3.2.2 Particle-Center-Appiaech Measurements

To examine the effects of sintering temperature on particle- ce nter-approach, measurements were taken of particles that CO ntacted the su bstrate of sintered specimens. Particles were chosen for measurement i f they appeared ta be spherical to the naked eye as observed on the monitor of the scanning electron microscope. One measurement was made of the width/diameter of the particle in the plane parallel to the substrate and a second measurement was made of the height of the particle perpendicular to the surface of the substrate. The ratio of 'height' divided by 'widthldiameter' gave a measure of the degree of particle-center- approach.

3.3 Finite Element Analysls

The single notch axial fatigue test specimen was modeled on the ANSYS CAD program in order to determine the stress concentration factor of the WireEDM notches. The specimen was rnodeled as a block of metal with the same Young's Modulus and Poisson's Ratio as titanium, with the same thickness as requested on design drawings (3.50mm or 0.1 38') and with a height one half that of the actual sample (3.57mrn or 0.1 40"). A single notch with a width of 75pm (0.003") and a depth of cut into the specirnen of 200pm (0.008"). including the idealized tip with a radius of 37.5pm (0.001 5")' was cut from the modeled block. Boundary conditions created the necessary symmetry to duplicate the effects of a full height specimen with a notch on each side.

The model was meshed by the software prompts from the operator. Element s i re was

using command deemed sufficiently

srnall when a 10% increase in the number of elements incfeased the determined stress concentration factor by less than 1%.

The applied force acting on the model was chosen to give a normal tensile stress of 300MPa within the specimen, calculated a s s u m i n g no stress concentrators, (i.e. 300MPa = FIA, A = cross- sectional area of an unnotched specimen). This figure was chosen as it i s an upper limit for fatigue endurance strengths reported by others for porous coated Ti-6AI-4V fatigue s p e c i r n e n ~ ~ * ~ ~ ' ~ ~ ~ ~

FIGURE 3.7: Single Notch Specimen FEM Mesh

Nodes along the bottorn of the mode1 are constrained vertically. Nodes dong the left-hand side of the model are constrained horizontally. The tensile force is applied horizontally to the right-hand side of the model.

3.4 Mechanical Testing

3.4.1 Testing of Shear Specimens

The shear specimens were tested on a Uniaxial S ervohydraulic lnstron 850 1 Universal Testing Machine, (Instron Inc., Canton, MA). One end of the shear specimen was placed in one actuator grip of the machine while a ring designed to shear off the sintered particles was attached to the opposite grip. The grips were separated at a rate of 1.25mm/min and the force recorded using Control High Strain Rate Testing software, (Chris Pereira, University of Toronto, Figure 3.2). The shear force versus displacement was plotted for each specimen. The shear force was determined by subtracting the average force acting on the load cell before the sharp rise in force (considered to represent the frictional force due to the ringkpecimen-substrate interaction) from the maximum force attained during shearing of the sintered particles from the solid substrate.

3.4.2 Testing of Axial Fatigue Specimens

The axial fatigue specimens were tested using the lnstron 8501 universal test machine. The tests were run using a sinusoidal wave function with R=0.1 (R = q,,i,/a,,,) at 50Hz in air at

room temperature. Loads were selected for the first specimen of each test group to give maximum nominal stresses below the expected fatigue endurance strength for the Ti-6AI-4V specimens. If runout occurred, typically for 1OE6 cycles, the load was increased for the subsequent test on the same specimen. The load was increased in a step-wise manner until specimen failure occurred within 10E6 cycles. Once the first specimen failed, the applied maximum load and resulting stress level for the next specimen was chosen to allow efficient definition of an S/N curve (at least in the region of the endurance strength). For the notched specimens, the fatigue endurance strength was determined in this mannet to within +/- 5MPa.

For the MA condition, eig ht specimens were used for 13 tests. For the specimens treated at 1 2 7 5 ' ~ . eight specimens were used for 1 7 tests. For the specimens treated at 1 300-C, seven specirnens were used for 10 tests. For the specimens treated at 131 S'C, seven specimens were used for 21 tests.

3.4.3 Testing of Rotating Beam Fatigue Specimens

Initially. axial fatigue testing of notched specimens was undertaken. When it was observed that there was no difference in fatigue strengths between specimens heat treated at 1 2 7 5 ' ~ . 1 ~ O O ' C , and 131 5 ' ~ . it was decided to perform rotating beam fatigue tests with sintered. porous-surfaced specimens to determine if the axial fatigue tests were truly representative of the notch fatigue properties as indicated by previous studies of porous- su rfaced Ti-6A1-4V specimens. lnclusion of 1 2 5 0 ' ~ as the lower sintering temperature for th8 rotating beam fatigue specimens was selected to provide a greater extrerne in both sinter neck contact ratios and microstructure. This temperature also corrosponded to reported sintering temperatures used for preparing porous-surfaced Ti-6AI-4V dental implants, albeit with different particle sizes.

The rotating beam fatigue specimens were tested in air using an R. R. Moore rotating beam test bench a t a test frequency of approximately 100Hz. The applied load and resulting maximum stress for the first specimen was chosen to be significantly below the expected fatigue endurance strength to produce a runout of 10E6 cycles. Applied stresses were increased for subsequent tests on the first specimen unti l failure occurred. Once failure of the first specimen occurred, the stress level for the next specimen was chosen to best describe the S/N curve (at least in the region of the endurance strength). Stress levels were increased i n 20MPa increments.

TABLE 3.1 : Chernical Composition of Materials

(AI1 vaiues are in w e i g M , baiance Ti.)

A l l o y A l V h C O N H Y

Ti-6AJ4V €LI MAX, 6.5 4.5 0.25 0.08 0.13 0.05 0.012

(ASTM F 136-84) M N 5.5 3-5 - - - - - -

n-64-4v ELI

(Axial Fatigue Test Spedrnens) 6.1 2 4.05 0.1 1 0.039 0.09 0.005 0 . W 0.001

(Supplied by Teledyne Alivac)

Ti-ôAI4V EL1

(Rotating Beam Test Spedmens) 6.1 5 4.2 0.20 0.011 0.13 0.012 0.001 0.001

(Supplied by Dynamet Inc.)

Ti-6AI4V EL1 Partides 6.1 3.9 0-15 0.125 0.02 0.014 0.007 0.005 (Supplied by Nudear Metals)

4.1 Sinter Neck DiameterIParticle Diameter Ratios

Sinter neck measurements were recorded to determine how altering the sintering temperature for a fixed sintering time (2 hours) affected the geometry of the particle-substrate junction. Sinter neck diameters were normalized with particle diameters to get a sinter neck diameter/particle diameter ratio, r /R This ratio i s referred to as the 'contact ratio'. The number of sinter necks exarnined totalled 60 from tliree specimens for the 1250*~ sintering condition, 85 from three specimens for the 1275'~ sintering condition, 80 from three specimens for the 1300'~ sintering condition, and 80 from four specimens for the 1315'~ sintering condition.

With the exception of particles with diameters of less than about 1 OOpm, there was little evidence of particle-center-approach. For both the 1250'~ specimens and the 131 5 ' ~ specimens. examining particles with diameters of at least 150pm (with the rare exception), the particle 'height' to particle 'width/diameter' ratio was 0.94. (1 2 5 0 ' ~ : 0.94 +/- 0.05; 131 5 ' ~ : 0.94 +/- 0.04) Surface diffusion causes material ta be transported from the 'top' of the sinter particle to the sides, and eventually to the sinter neck. Because of this, even particles affected only by surface diffusion wil l exhibit a particle 'height' to particle 'width/diameter' ratio less than 1 .O.

Specimens examined demonstrated that increasing the sintering temperature increased the sinter neck contact diameters, and hence the contact ratios. See Figure 4.1. ANOVA tests for the contact ratios of the four groups of specimens indicated significance. The difference between the Contact ratio for the 131 5 ' ~ specimens and the 1250'~. 1 2 7 5 ' ~ and 1300'~ specimen contact ratios i s significant to <0.0001 based on paired T-tests. Using Pe0.05 for paired T-tests to compare the contact ratios

arnong the three lower sintering temperatures, only between 1 2 5 0 ' ~ and 1 3 0 0 ' ~ i s there a significant difference (Table 4.2).

Observations made while photographing failed R.R. Moore waisted specimens showed that there was a difference in the size of the "sinter zone". At higher sintering temperatures there appeared to be greater heights to the sinter necks. See Figure 4 . 2 A and Figure 4 - 2 8 . Subsequent measurements of the height of the sinter neck were made. Measurements showed an increase in si nter neck heig ht with increasing sintering temperatu re.

TABLE 4.1: Sinter Neek Ratio and Height Summary

TABLE 4.2: ANOVA and Paired T-Tests for Contact Ratio

ANOVA Table for Nock R d i o OF Surn of Sauares Mean Saare F-Value P-Value

Fishefr PLSD for Nock Ratk Effstt: TsmpGroup Slgnifkance W m l : 5 %

Mean Diff- Crit- Diff P-Value

Temp.Group Residual

TABLE 4.3: ANOVA and Paired T-Tests for Sinter Neck Helght

ANOVA Table for Slntor Hdght OF Surn of Sauares Mean Sauare F-Value P-Value

Model II estimate of beniurren component variance: 28-71 9

3 301

- - - -- - -

Ternp. Group 1 3 ( 20.528 1 6.843 1 11.412 c.0001 ] -- -

Model II estirnate of between component variance: .208

7

6683.288 14904.660

Fisher's PLSD for Sinter HeIgM Effect: Temp. Group Signlficrince Level: 5 96

- - - pp

2227.763 49.51 7

Mean Diff. Crit. Diff P-Value

- -

44.990 ~ 0 0 0 1

S S S

S S

T 1250C, T 127% T 1250C, T 1300C T 1250C. T 131 5C T 1275C. T 1300C T 1275C, T 131 SC

T 1300C. T 1315C

-.607 --71 O

-1.1 60 -. 1 03

-.553 -.450

.396

-396 .396

.396 -396 -396

.O030

-0006 <.O001

.6063 ,0066

.O263

Figure 4.1 : Contact Ratlo vs. Sintering Temperature

0.5 ,

0.45 : 0

I 1 L

r Contact

0.05 : - Ratio O, I 1

1 250 1275 1300 1315

Sintering Temperature (C)

4 .2 Shear Force

An example of a Shear Force vs. Displacement Graph is given in Figure 4.3.

lncreasing the sintering temperature generally increases the force required to shear the ring of particles from the specimen substrate, as shown in Table 4.4. The exception to this was the near identical shear forces recorded for the 1275.~ specimens and the 1 3 0 0 ' ~ specimens. The specimens sintered a1 1250'~ sustained shear forces moderately Iower (by approximately 17%) than the specimens sintered at 1275'~ and 1300'~. There was a large (approximately 45%) increase in shear force recorded for the 131 S'C specimens compared to the 1275'~ and 1300'~ specimens. ANOVA tests showed significance, and using Pt0.05 for paired T- tests, only the results between the specimens at 1315'~ and al1 of I ~ s o ' c , 1 2 7 5 ' ~ ~ and 1300'~ were significant (Table 4.5).

FIGURE 4.2a: SEM of Sintered, Porous-Coated Rotating Beam Fatigue Specimen, 1 2 5 0 ~ ~

FIGURE 4.2b: SEM of Sintered, Porous-Coated Rotating Beam Fatigue Specimen, 131 5 ' ~

FIGURE 4.3: Shear Force vs. Displacement Graph

Force (N)

0.2 0.4 0.6 0.8

Displacement (mm)

TABLE 4.4: Shear Force Summary

TABLE 4.5: ANOVA and Paired T-Tests for Shear Force

Fisher's PLSD for Show F o m Effsct: Tarnp.Group SignifIcance Level: 5 %

ANOVA Table for Show Force DF Sum of Squares Mean Square F-Value P-Value

T q G i a i p Residual

4.3 Microstructure and Grain Size

Mean Diff. Crit. Diff

Both the bar stock used for axial fatigue specimens and the rod used for the shear force specimens, rotating beam specimens, and addit ional sinter neck measurement specimens were supplied i n the Mi Il Annealed condit ion. The heat treatments, including sintering, at 1 2 5 0 ' ~ , 1 2 7 5 ' ~ , 1 3 0 0 ' ~ . and 131 5 * ~ , created a beta- annealed (BA) microstructure. Analysis of the grain sizes for the BA specimens showed increasing grain sizes with increasing sintering temperature. (see Table 4.6 and Figure 4.4a & Figure 4 .4b )

Model I I estirnate of between comportent variance: 136624.227

3

43

T 1250C, T l275C

T 1250C, T 1300C

T 1250C, T 131%

T 1275C, T 1300C

T 1275C, T 1315C

T 1300C, T 131 SC

503241 3.386 1 l67747l.l29 471 6407.891 1 109883.904

-222.545 -218.990 -797.487

3.556 -574.941 -578.497

291 -826

300.198

258.446 306.879

266.176

275.329

15.294 <.O001

TABLE 4.6: Grain S ize Summary

Specimenlcondition Average Grain Diameter

R. R. ~ o o r e / l 2 5 0 ' ~ R.R. ~ o o r e / l 3 1 S'C

ASTM Axial11 2 7 5 ' ~ ASTM A x i a I f 1 3 0 0 ' ~ ASTM Axia1 /1315 '~

FIGURE 4.4s: Light Micrograph of Microstructure of Rotating Beam Specimen Sintered at 1 2 5 0 ' ~ (1 1.4mm diameter)

FIGURE 4.4b: Microstructure ot Rotating Beam Specimen Sintered at 181 S'C (1 1.4mm diameter)

4.4 Hardness

The machining of notches into the axial fatigue test specimens was done by wire electrodischarge machining. This created a zone around the machined notch which had undergone melting immediately followed by a water quench. As per 3.1.3.1 -3, the heat treatment procedure was divided into two treatments of one hour each, with the second heat treatment after the machining of the notches. It was hoped that the effects of melting and quenching on hardness, and in turn fatigue properties, could be minimized by the post-rnachining heat treatrnent causing any accumulation of oxygen, nitrogen, and hydrogen interstitials in the machining zone to diffuse through the titanium alloy due to their

high solubility. Testing was performed to measure any change in the hardness around the notches which may have affected fatigue properties.

The hardness tests, while exhibiting significant scatter, indicated no effect due to the EDM machining technique when followed by the heat treatment. All specimens generally had a Vicker's hardness number of approximately 400-450 both close to the notch ( (50pm ) and at distances up to 200pm from the notch. The amount of scatter in the hardness results rnay, however, mask an effect extremely close to the machining zone (4 5pm from the notch) which may be discernable with a more precise hardness tester. Figure 4.5 is an example for one of the 1275C axial fatigue test specimens. All graphs are recorded in Appendix 7.5.

4.5 FEM Analysis of Notch of Axial Fatigue Specimens

The computer finite element method analysis of the effect of the notch on the stresses in the axial fatigue test specimen showed a stress concentration factor of 6.03 = 6.0. As expected for a

simple geometry under straight tension, the maximum stress was located at the end of the radiused tip of the notch. See Figure 4.6.

FIGURE 4.5

Hardness for 1275 Smimen

O 20 40 60 80 106120140160180200 Distance from Notch

(um)

FEM Analysis of Notch

4.6 Fatigue Strength

The fatigue strength for this study was expressed as the stress level at which approximately 50% of specimens failed and 50% experienced runout.

4.6.1 Axial Specirnen Fatigue Strength

Initial fatigue tests were done on axial fatigue specimens that were un-notched and in the MA condition. However, due to the high strength of the un-notched specimens, many failures occurred in the grip region of the specimens. Based on the limitted number of successful tests (two failures in the test region of the specimens at 700MPa, and one runout at 650MPa & one failure in the grip end after 1.2E6 cycles at 650Mpa) the fatigue strength was estimated to be 650MPa+/- SOMpa, consistent with literature values for Ti-GAI- 4v.

Notched specimens had a nominal width of 7.1 mm before notching. The depth of the notches was approximately 0.2mm per notch, leaving a notched width of 6.7mm. This creates nominal stresses experienced at the notched region of the specimens 6% g reater than unnotched specimens. Stresses Iisted in Table 4.7 incorporate the 6% increase due to the reduction in cross-sectional area due to notching.

The MA specirnens, with their fine-grained microstructure, exhibited very little scatter in fatigue strength. There were no runouts above 143MPa, and no failures below 143MPa. The BA specimens, with their much greater grain sizes, demonstrated some scatter in fatigue strengths. The group of specimens heat treated at 1 2 7 5 ' ~ had a failure at 170MPa, and a runout at 185MPa. The 1 3 0 0 ' ~ group had a runout at 19OMPa. The 131 5 ' ~ group had failures at 170MPa and runouts up to 195MPa. All specimens failed at the notch.

TABLE 4.7: Axial Specimen Fatigue Results Summary

Mill Annealed 1 2 7 5 ' ~ 1300'~ 131 S'C

4.6.2 Rotating Beam Specimen Fatigue Strength

The rotating beam fatigue tests were less precise that those for axial fatigue (Table 4.8). The inherent randomness of coating the specimens with particles creates a variable. The installation of the specimens into the rotating grips required numerous attempts to ensure eccentricity was within acceptable limits. The application of the load to the specimen was accomplished with a slider on a graduated bar; the applied loads were not exact.

Very few specimens failed at the center of the waisted section. The porous-coated region comprised only the middle 1 Omm of the specimen. The application of the particles resulted in an increase in the cross-sectional area over which stress was distributed, but with the outer porous-coated zone having a lower effective Young's Modulus by virtue of its lower density. The region which experienced the greatest stress, allowing for stress risers resulting from the transition from a fully dense region to a lower density region, may in fact have occurred near the edge of the particle coated region of the specimen. This could explain the rnajority of the failures occurring nearer the edge of the band of porous-coating on the specimens.

The group of specimens sintered at 1 2 5 0 ' ~ had three failures at 180MPa, tnree failures at 200MPa, and two runouts at 200MPa. The group of specimens sintered at 131 S'C had four runouts at 180MPa, and three failures at 180MPa.

TABLE 4.8: Rotating Beam Fatigue Rmults Summciry

Sintering Temperature

Endurance Limit (RA Pa +/- 20MPa)

5.1 Sintering

The average sinter neck diameter/particle diameter contact ratio increased with increasing temperature. Between the specimens sintered a t 1 2 5 0 ' ~ and 1 ~ O O ' C , the contact ratio increased linearly from 0.34 to 0.38. The specimens sintered at 131 5 ' ~ had a contact ratio of 0.46. The modest 1 S'C increase in sintering tempe rature (approximateiy 1% increase in absolute temperature) between 1 3 0 0 ' ~ and 1 31 S'C resulted in an increase in the contact ratio of 20%. This corresponds to a 45% increase in the contact area between sintered particles and the substrate. The sharp increase in the sinter neck diameter/particle diameter ratio might be explained by a shift i n dominant material transport mechanisrn due to the increase in sintering temperature from 1300'~ to 131 5 ' ~ .

The sintering diagram (Figure 1.3) for titanium developed by Uskokovic et a1.12 shows the dominance of surface diffusion early in the sintering process, bath above and below the f'3 -transition temperature. Below the B -transition temperature (approximately 9 9 2 ' ~ for Ti-6AI-4V ELI), the dominant mechanism for material transport can shift from surface diffusion ta grain boundary diffusion, once a certain sinter neck diameter has been attained. Above the $ -transition temperature, the dominant mechanism for material transport can shift from surface diffusion to volume diffusion.

Recall the empirical formula for sinter neck growth:

r = sinterneck radius t = tirne of sintering F(s) = function dependent on temperature, particle size,

geometry n = f(materia1 transport mechanism)

For surface diffusion, 'n' i s accepted as being 7. For volume diffusion, 'n' i s accepted as having a value of 4". The transition from surface diffusion to volume diffusion would result in a su bstantial increase in the sinter neck growth rate. This would account for the large increase in sinter neck diameters observed for specimens sintered at 131 5 ' ~ compared with those specimens sintered at 1 3 0 0 ' ~ .

From the work of Uskokovic et al.'*, volume diffusion can become the dominant mechanism for material transport during sintering of titanium (taking over from surface diffusion) while sti l l in the Stage One region of the sintering diagram. From the sintering diagram constructed for pure titanium, the transition from surface diffusion to volume diffusion occurs , depending on the sintering temperature, at a contact ratio of between approximately 0.38 to 0.42. If Ti-6AI-4V alloy closely follows the diagrarn for pure titanium, then volume diffusion occurs when the specimens are sintered at 131 S*C with the 150pm-250pm particles. Further evidence for a shift from surface diffusion to volume diffusion during sintering is found i f Stage Two sintering has begun. From the sintering diagram for titanium, volume diffusion is definitely the dominant mechanism for material transport in Stage Two sintering. Ashby denoted the transition between Stage One sintering and Stage Two sintering as the point at which the driving force for mate rial transport due to curvatu re differences within pores approached zero? Ashby chose dR = 0.6 to represent this point of transition. Like Ashby, Mohan et alz5 considered the point at which pore shrinkage due to curvature of the pore surface dominates material transport as the transition between Stage One and Stage Two. Mohan calculated mathematically that this occurs at approxirnately dR = 0.38 or r/R = 0.39, depending on the sintering conditions and the sinter particles. Computer simulations by Mohan et al for U 0 2 particles gave transition from Stage One to Stage Two

at r/R = 0.4 ta r/R = 0.42. The sinter neck contact diameter ratio for the 1 2 5 0 ' ~ sintered specimens was 0.34, for the 1 2 7 5 ' ~ sintered specirnens was 0.36, for the 1 3 0 0 ' ~ sintered specimens was 0.38, and that for the 1 3 1 5 ' ~ sintered specimens the ratio was 0.46. A

transition from Stage One to Stage Two sintering at r/R= 0.39 for Ti-6AI-4V ELI would appear to correspond to the significant occurrence of volume diffusion, resulting in the associated increase in the neck growth rate. A shift to Stage Two sintering mandates that a transition in dominant mechanism for material transport to volume diffusion, from surface diffusion has occurred. This transition between 1 3 0 0 ' ~ and 131 S'C for these specirnens explains the sharp increase in the sinter neck contact diameters of the 131 5 ' ~ specimens. The fact that there is no more particle- center-approach in the specimens sintered at 131 5 ' ~ than with the specimens sintered at 1 2 5 0 ' ~ . indicates that volume diffusion is not the primary mechanism for material transport in the 131 5 ' ~ specimens. but that it i s only a factor at the end of the sintering process for these specimens under these sintering conditions.

The sintering conditions studied for the 150pm-250pm particle range d id not cause significant densification, as demonstrated by the fact that there was very Iittle change in interparticle-distance during sintering for al1 temperatures studied. The exceptions occurred when there was sintering involving small (<100wrn diameter) particles, and even in these cases the change in interparticle-distance was moderate. The sintering rate for smaller particles i s greater than that for larger particles, as shown in 1.1.3. The increase in particle-center-approach for the smaller particles is explained by the fact that smaller particles undergo the transition from Stage One to Stage Two sintering before larger particles. assuming identical sintering temperatures. Smaller particles would experience greater amounts of material transport by volume diffusion than larger particles, under the same sintering conditions.

5.2 Interfacial Shear StrengthIFracture Charateristics

The sinter neck measurements of section 4.1 indicated that the sinter neck diametWcontact diameter ratio increased with sintering temperature. Examination of the interfacial shear strength specimen substrat0 after shear strength tests revealed 'craters"

(Figure 5.1) where each particle had previously formed a sinter neck reg ion. This indicates good metallurgical bonding of particles to the substrate in these regions resulting in cohesive (as opposed to adhesive) failure on shear testing. As the shear strength tests examine the forces required to shear the particles from the substrate, greater sinterbond area should require greater shear forces for particle removal from the substrate. From section 5.1, it

was observed that there was an increase in the contact area of 45% from specimens sintered at 1 3 0 0 ' ~ to those sintered at 131 5 ' ~ . Comparing the shear forces sustained by the 1300k specimens and the 13 1 5 ' ~ specimens demonstrated the same 45% increase. The same results are observed when the data from specimens sintered at 1 2 5 0 ' ~ and 131 S'C are compared. The sinter neck diameter/particle diameter ratio was 0.34 for the 1 2 5 0 ' ~ specimens versus 0.45 for the 131 5 ' ~ specimens. This is equivalent to a 75% increase in contact area between particles and substrate. The average force required to shear the sintered particles from the substrate for the 131 5-c specimens was 1880N. This is 75% greater than the average force of 1080N for the 1 2 5 0 ' ~ specimens. The shift to volume diffusion from surface diffusion as the prirnary mechanism for material transport during the sintering of the particles does not affect the shear failure stress. A greater contribution from volume diffusion only increases the area of the sinter neck regions.

From the plots of shear force vs. displacement, for example Figure 4.3, i t was found that from the point at which the shearing collet contacted sintered particles over the entire circumference of the collet, to the point on the F vs. x graph when maximum shear force is attained, the displacement x was approximately 300pm.

Visual examination of the surface of the shear specimens after testing indicated that the sinter necks had occupied an estimated 15% of the total area of the substrate surface. Using a shear force of 1500N to shear particles from the substrate, a displacement x of

300pm over which the l5OON was sustained, and 15% contact area of the substrate by particles, the strength of the sinterbonds can be approximated by:

Area of sinterbonds = D x ~r x x ~ 1 5 % D = 15mm D = diameter of

specimen

Area of sinterbonds = 15 x n x 0.300 x 0.15

Area of sinterbonds = 2.1 mm2

S hear force = 1500N, Area = 2.1 mm2

Strength of bonds = 1500N 12.1 mm2 = 71OMPa (Approximate)

The generally accepted shear strength for Ti-6AI-4V alloys is 600MPaZ0, supporting the observation that there is a true cohesive bond of particfe and substrate due to the sintering process.

FIGURE 5.1 : SEM of Surface of S heared Specimen Sintered a t 131 5 ' ~

5.3 Hardness

Fatigue strengths are dependent on many factors. For the notched fatigue specimens, the possibility that the WireEDM machining affected the fatigue strength was considered. Despite

the fact that the notches were created by a process which involved localized melting and rapid resolidification. the hardness tests showed no significant difference in hardness between regions close to the notches (>50pm) and regions greater than 100pm from the notch. This indicated that the second heat treatment, performed after the notching and resulting in a total time at temperature of two hou rs, minimized microstructural effects that may have negatively affected the fatigue strength of the axial fatigue test specimens.

5 -4 Fatigue Strength

5.4.1 Axial Specimen Fatigue

The fatigue strength of 650MPa for the unnotched MA axial fatigue specimens is slightly higher than the average of approximately 625MPa given for waisted rotating beam specimens i n previous studies, but well within the range reported in the Iiterature.

The fatigue tests of the MA axial fatigue test specimens demonstrate the high sensitivity of Ti-6AI-4V to notches. From FEM analysis, the notch created a stress concentration of 6.0 at the tip of the radius. The reduction in the fatigue strength from 650MPa for the unnotched MA specimens to 143MPa for the notched MA specimens is almost an 80% decrease. The fracture surface of the MA specimens was quite smooth, indicating that the fine-grained MA microstructure is not especially resistant to crack propagation. The as-sintered axial fatigue specimens, heat-treated at 1 2 7 5 ' ~ . 1 300'~. and 1 31 5 ' ~ . al1 demonstrated fatigue strengths of 175MPa. The fracture surfaces were not as smooth as those of the MA fatigue specimens. The irregular fracture surface is evidence of the greater resistance to crack propagation of BA microstructures versus MA microstructures. As the notch acts as a pre-existing crack. the threshold stress intensity factor, AKthr, w hich measures the ability of a material to arrest unstable crack propagation, becomes important. For MA Ti-6AI-4V alloy, AKthr i s reported to be 6MPa dm, white for BA ~ i - ~ A ' I - 4 ~ alloy, BKthr i s reported to be

1 2MPa dm26927. The hig her fatigue strength of the pre-notched BA specimens versus the pre-notched MA specimens confirms this.

Despite the significant coarsening of the microstructure of the specimens with increasing heat treatment temperature (average grain diameters of 2.7mm for the 1 2 7 5 . ~ specimens to 6.2mm for the 131 S'C specirnens), there was no difference i n fatigue strengths. The pre-existing notches act as a pre-existing crack, which minimize the superior crack initiation retardation characteristics of finer microstructures. There did, however, appear to be an increass in scatter in the S/N cuwe with larger grain sire. The studies confirmed the susceptibility of titanium alloys to notch defects with respect ta fatigue strength.

FIGURE 5.2a: MA Axial Fatigue Specimen Fracture Surface

FIGURE 5.2b: BA (127S°C) Axial Fatigue Specimen Fracture Surface

5.4.2 Rotating Beam Specimen Fatigue

The rotating beam fatigue tests mirrored the axial fatigue tests in that there was no significant change in fatigue endurance strength for specimens sintered between 1 2 5 0 ' ~ and 131 5 k . Though the grain size almost doubled for the fi-annealed

specimens. which can reduce resistance to crack initiation and short crack growth. the notch effect associated with the sinter neck regions had an influence on fatigue strength that overshadowed any effect due to coarsening of the microstructure. Extrapolation and interpolation of prior FEM analysis of the stress concentration factor for various sinter neck geometries give. for the sinter necks on the rotating beam specimens of this study. a stress CO ncentration factor of approximately 3'e2? This value changes depending on the radius of the sinter neck and the sinter neck contact diameter ratio- An increase in the sinter neck contact diameter ratio tends to increase the stress concentration factor, while a greater sinter neck radius decreases the stress concentration factor. The value of '3' i s about half that determined in the present study by FEM for the axial fatigue test specimens.

The sinter neck contact ratio for the 131 S'C specimens was 0.46, and for the 1250'~ specimens it was 0.34. The computer FEM of sinter necks with varying contact diameter ratios reported by Ducheyne et al2' predict lower fatigue strengths for the higher contact ratio 131 5 ' ~ specirnens. The explanation for this was suggested to be due to stress contouring, whereby greater sinter neck contact allowed more stress to be carried into the particle and hence around the sinter neck. This shoufd have resulted in a lower fatigue endurance limit compared with the 1250'~ f3-annealed specime ns. The computer simulations, however. considered only single particles and rows of particles. The growth of sub-critical cracks and the influence of a complete coating in three dimensions was not included in the computer simulation study. The possible interaction of a series of sub-critical cracks, which would occur on a porous-coated surface, was not considered.

Examinations of photomicrographs of the sinter neck contact regions of the rotating beam specimens showed a difference between the specimens sintered at 1250'~ and 131 S'C. For the specirnens sintered at 131 5 ' ~ (Figure 4.2b). the sinter necks appeared larger in diameter. greater in height. and had a much 'smoother' appearance that the sinter necks of the 1 2 5 0 ' ~ specirnens, (Figure 4.2a). The minute roughness of the sinter

necks of the 1 2 5 0 ' ~ specimens could promote crack initiation to a greater extent than for the sinter necks of the 131 S'C $-annealed specimens. The inclusion of volume diffusion as a significant mechanism for material transport for the 131 5 ' ~ specimens not only increased the sinter neck contact diameters. but also appeared to increase the radius of the sinter neck to the extent that fatigue strength was not reduced due to the coarser microstructure and the influence of stress contouring on the larger diameter sinter necks.

The axial fatigue specimens were tested at R=0.1, with the entire cross-section of the specimen experiencing this load. Rotating beam specimens are tested at R=-1 . Only the outer surface of the rotating beam bending specimen experiences maximum stress; the stresses in the specimen decrease from a maximum to zero along the line from the surface to the axis of rotation. Material a i the neutral axis of the rotating beam specimen experiences no load until cracks have propagated sufficiently to alter the location of the original neutral axis of the rotating beam specimen. At R=-1 , the stress amplitude is greatest, which tends to initiate cracks to a greater degree than tests at R10.1. There is. however, the largest amount of crack closure at R=-1. Fatigue tests at R=-1 will tend to propagate cracks slower than at R=0.1, for the same AK, as demonstrated below.

To account for crack closure. a AKeff was p r ~ p o s e d ~ l * ~ ' . Two empirical formulas for AKeff are given below.

For 1 ., RI-1, U = 0.1 R10.1, U = 0.54 = 0.55

For 2., R=-1 , U = 0.3 Rz0.1, U =0.59 =0.6

Using a maximum stress of 200MPa, at R=-1, AK=400MPa.

(1 .) AKeff = 400MPa x 0.1 = 40MPa. (2.) AKeff = 400MPa x 0.3 = 120MPa.

At R=O.l , using a maximum stress of 2OOmpa. AK =180MPa. (1 .) AKeff = 180MPa x 0.55 - 100MPa. (2.) AKeff = 180MPa x 0.6 = 1 iOMPa.

Th8 more complex formula (2.) gives comparable effective stress intensity factors for both the Rs0.1 and R=-1 fatigue tests. This explains the sirnilar fatigue endurance strengths recorded for both the notched, axial fatigue specimens and the sintered, porous- coated rotating beam fatigue specimens.

1. For sintering of 150pm-250vm (-60/+80 mesh) Ti-6AI-4V EL1 spherical particles to Ti-6AI-4V EL1 substrate, with 2 hours at the prescribed temperature, there i s a change in the dominant mechanism for material transport between 1 3 0 0 ' ~ and 131 S'C. The transition from surface diffusion to volume diffusion, under the described conditions, does not lead to significant densification. Volume diffusion significantly increases the sinter neck contact diameter ratio under the conditions of these experiments.

2. The increase in the sinter neck contact diameter ratio leads to a corresponding increase in the force required to shear the particles from the substrate. The shear forces resisted by the particles sintered at 131 S'C are 45% greater than those resisted by the particles sintered at I~OO'C , and are 75% greater than those resisted by the particles sintered at 1 2 5 0 ' ~ .

3. There is no significant change in fatigue strengths between notched axial fatigue specimens heat treated for two hours at 1 2 7 5 ' ~ , 1 3 0 0 ' ~ . and 1 31 S'C respectively, despite the change in microstructure.

4. There is no significant change in fatigue strengths between rotating beam specimens sintered for two hours at 1 2 5 0 ' ~ and 1 31 5 ' ~ respectively using 150pm-250pm (-60/+80 mesh) sp herical powder porous-coating. despite the change in microstructure and particle/su bstrate interface geometry.

5. When forming porous coatings on Ti-6AI-4V EL1 orthopaedic implants using 150prn to 250pm (-60/+80 mesh) Ti-6Al- 4V EL1 spherical particles, sintering at 1 3 1 5 ' ~ for 2 hours will increase the force required to shear the particles from the substrate versus sintering at lower temperatures. There is no decrease in fatigue strength by sintering at the higher temperature. Any

densification of the porous surface at the higher temperature is insignificant with respect to bone ingrowth ability.

6 . For sinter particles of diameters different from 150pm to 250pm (-60/+80 mesh), there is the likelihood that there exists a sintering temperature which will, for a two hour sinter, achieve the same significant increase in contact ratio (and shear force) over slightly lower sintering temperatures, while rnaintaining porosity and fatigue endurance strength, as observed with the 1SOpm to 250pm (-60/+80 mesh) diameter sinter particles used in this study. For smaller particles, the ideal sinter temperature wil l be lower than 131 5 ' ~ , and for larger particles the ideal sinter temperature wil l be greater than 131 S'C.

RECOMMENDATIONS FOR FUTURE STUDY

From the above conclusions, it is recommended that a study be conducted to determine the optimum sintering temperature and/or sinte ring time for spherical particle porous coated systems of titaniurn alloys. The optimum sintering temperature and/or sintering time would be determined when the contact ratio of the sinter particles approaches 0.45 using the the most economical compromise between higher sintering temperature and longer sintering time. From these results for various sinter particfe diameters, manufactu rers of porous coated systems for orthopaedic and dental applications could optimize porous coatings for shear strength, fatigue strength, and porosity.

REFERENCES

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J. Robbins, Medical materials. Materials Edge, 10, MarchIApriI, 25-43, 1989

S. D. Cook, K. A. Thomas, Fatigue failure of noncemented porous-coated implants, J Bone Joint Surg, 73, 20-24, 1991

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G. R. Yoder, L. A. Cooley, T. W. Crooker, A cornparison of microstructural effects on fatigue-crack initiation and propagation in Ti-6AI-4V, 23rd S, SD, &MTL Conf, May, 1984

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R. M. Pilliar, D. A. Deporter, P. A. Watson, N. Valiquette, Dental implant design - effect on bone rernodeling, J Biorned Mater Res, 25, 467-483, 1 991

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M. F. Ashby, A first report on sintering diagrams, Acta Met, 22, 275-288, 1974

P. Ducheyne, M. Martens, Orderly oriented wire meshes as porous coatings on orthopaedic implants II: the pore size, interfacial bonding and microstructure after pressure sintering of titanium OOWM, Clinical Materials, 1, 91 -98, 1986

Metals Handbook, AS M, 1985

E. W. Collings, The Physical Metallurgy o f Titanium Alloys, Am Soc Metals, Metals Park, Ohio, 1984

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S. B. Young, Fatigue of porous coated titanium implant alloy, MASc Thesis, University of Toronto, 1989

T. Albrektsson et al., The interface zone of inorganic implants in vivo: titanium implants in bone, Annals Biom Eng. 1 1, 1 - 27,1983

Materials Properties Handbook Titanium Alloys, AS M. 1 994

S. A. Meguid, Engineering Fracture Mechanics. Elsevier Science Publishers, 1989

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D. Wolfarth, M. Filiaggi, P. Dusheyne, Parametric Analysis of Interfacial Stress Concentrations in Porous Coated Implants, J Applied Biomat, 1, 3-1 2, 1990

J. D. Bobyn, R. M. Pilliar, The optimum pore s i re for the fixation of porous-surfacéd rnetal implants by the ingrowth of bone, Clin Orthop, 150. 263-270, 1980

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G. W. Kuhlman, A. K. Chakrabarti, Room temperature crack propagation in beta-titanium alloys, Microstructure Fracture Toughness and Fatigue Crack ~ r o wth Rate in fitanium Alloys, proc. TMS-AIME Annual Symposia, Denver, February 1987, eds. A. K. Chakrabarti, J. C. Chesnutt, The Metallurgical Society, Warrendale, Pennsylvania, 3-1 5, 1987

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28- H. L. Ewalds, R. J. H. Wanhill, Fracture Mechanics, Edward Arnold, 1984

7 APPENDIX

7.1 Contact Ratios

Bead Diameter Contact Diameter (um) (um)

Contact %

Average Average Average

Bead Diameter (um)

Contact Diameter (um)

Contact %

Average Average Average

Bead Diameter Contact Diameter Contact % (um) (u m)

Average Average Average

1315C Specimens

Bead Diameter Contact Diameter (um) (um)

Contact %

Average Average Average

7.2 Sinterneck Helghts

7.2.1 1250C Specimens

1250C Sample Heig ht (um)

Average 3.6

Std. Dev. 1.1

7.2.2 1275C Specimens

12756 Sarnple Heig ht (um)

Average 4.2

Std. Dev. .5

l3OOC Specimens

l3OOC Sample Heig ht (um)

Average 4.3

Std. Dev. .6

7.2.4 131 SC Specimens

1315C Sample Height (um)

Average 4.8

Std. Dev. -7

7 .3 Particle-Center-Approach

Bead Width Bead Height (um) (um)

Average 94

Standard Dev. 4.805424509

7.3.2 1315C Specimens

Average 94

Standard Dev. 4.1 1 3002643

7.4 Shear Force Specimens Data

1250C 1 275C 1 3 0 a

Average Average Average Average 1085.45455 1308 1304.44444 1883.52941

STD DEV STD DEV STD DEV STD DEV 152--339334 150.908361 167.56425 504.491 59

7.5 Hardness

7.5.1 -1 Hardness Plot: 127SC, Surface as Reference

7.5.1 -2 Hardness Plot: 1275C. Notch as Reterence --

O 20 40 60 80 100120140160180200 Distance from Notch

(um)

7.5.2.1 Hardness Plot: 1300C, Surface as Reference

7.5.2.2 Hardness Plot: 1300C, Notch as Reference

O 20 40 60 80 100120140160180200 Distance from Notch

(um)

7.5.3.1 Hardness Plot: 1315C, Surface as Reference

500

450

400

3 350 z -r, 300 a x 250 L aa 200

0 15J 100

50

O O 20 40 60 80 100120146160180200

Distance from Surfa- (ml

7.5.3.2 Hardness Plot: 1315C, Notch a i Reference

O 20 40 60 80 iooi20i4oi66i80200 Distance from Notch

(um)

7.6 Fatigue

7.6.1 Axial Fatigue

7.6.1.1 Mill Annealed Specimens Fatigue Data

Fatigue Results MA Single Notch

Cycles Failures Runouts

7.6.1 -2 12tSC Specimens Data

Fatigue Results

Cycles

BA, l27SC Single Notch

Failures Runouts

7.6.1.3 1300C Specimens Data

Fatigue Resuits BA, l3OOC Single Notch

Cycles Failures Runouts

7.6.1 -4 131 5C Spaclmens Data

Fatigue Results BA, 131 5C Single Notch

Cycles

302000 10000000 10000000 10000000 10000000 4000000 400GOOO 849000 193000 337000 10000000 4200000 945000 10000000 4500000 1 100000 6000000 4200000 8950000 786000 10000000

7.6.2

7.6.2.1

Failu ?es Runouts

Rotatlng Beam Fatigue

125OC Porous Coated Speclmens Data

Fatigue ResuRs

Cycles

PC, 1250C RRMoore

Fai lures Runouts

150 1 75

200 1 75

175 1 75

7.6.2.2 131 SC Porous Coated Specimens Date

Fatigue Resuits PC. 1 31 SC

Cycles Failures

RRMoore

Runouts

160 1 80

160 180

160

160 180

160

160