76-9939 BOWERS, Davtd Francis, 1937- CORROSION FATIGUE: TYPE 304 STAINLESS STEEL IN ACID-CHLORIDE AND IMPLANT METALS IN BIOLOGICAL FLUID. The Ohio State University, Ph.D., 1975 Engineering, metallurgy
Xerox University Microfilms, Ann Arbor, Michigan 48ioe CORROSION FATIGUE: TYPE 304 STAINLESS STEEL IN ACID-CHLORIDE AND IMPLANT METALS IN BIOLOGICAL FLUID
Dissertation
Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University
by
David Francis Bowers, B.S., M.S. Metallurgical Engineering
The Ohio State University 1975
Approved by
(dviser Department of Metallurgical Engineering
Reading Committee: Dr. Mars G. Fontana Dr. Frank Beck Dr. John P. Hirth To My Wife and Daughter TABLE OF CONTENTS
Page
ACKNOWLEDGEMENTS...... iv
VITA...... vi
LIST OF TABLES...... viii
LIST OF FIGURES...... x
INTRODUCTION...... 1
I. LITERATURE SURVEY...... 4
1.1 Concepts and Theory of Fatigue Crack Growth.. 4 1.2 Fracture Mechanics Approach to Crack Growth in Engineering Structures...... 18 1.3 Effect of Microstructure on Fatigue Cracking Growth...... 24 1.4 Mechanical Variables in Corrosion Fatigue.... 41 1.5 Environmental Interaction with Fatigue Crack Growth...... 59
II. EXPERIMENTAL DETAILS...... 92
2.1 Test Apparatus and Specimen Design...... 93 2.2 Test Material...... 109 2.3 Test Procedure...... 112
III. EXPERIMENTAL RESULTS AND DISCUSSION...... 118
3.1 Polarization Properties of Type 304 Stain less Steel in 5NH2SO4 and IN NaCl...... 119 3.2 Corrosion Fatigue Crack Propagation in Air and at Open Circuit Potential...... 122 3.3 Corrosion Fatigue Crack Growth at Applied Potential...... 128
ii Table of Contents (Continued)
Page
III. 3.4 Metallography and Fractography...... 131 3.5 Environmental Mechanisms Involved in the Corrosion Fatigue Behavior...... 138
IV. CONCLUSIONS . . 144
APPENDICES
A. A Review of the Literature Concerning Metallurg ical Invesitgations and Related Studies of Surgical Implant Materials...... 147
B. The Corrosion Fatigue Behavior of Biomedical Alloys...... 306
BIBLIOGRAPHY 379 ACKNOWLEDGEMENTS
A very special thanks is rendered to my employer,
Dr. R.W. Staehle whose constant efforts provided the research grants to financially support my education and thesis work during the past five years. His dedication to professional duties will never be forgotten as this example provides the necessary impteus for me and fellow co-workers to always strive for professional excellence both academically and in research. I express my best wishes for his continued success in the future. In addition, I offer my thanks to
Professors Fontana, Beck and Hirth for their suggestions and comments.
The following persons are gratefully acknowledged for their prompt assistance: Mr. R. Farrar for the photography and S.E.M. fractography; Mr. R. Justus and his assistants for machine shop work; Mrs. E. Orwig for typing this manuscript.
My grateful thanks is offered to the National Science
Foundation and the Electric Power Research Institute who provided the financial support for my thesis work.
iv No words can express the gratitude in my heart for the great personal sacrifice my wife and daughter sustained during these past five years. I sincerely hope they will receive the rewards so richly deserved in the years to come.
v VITA
May 8 , 1937...... Born-Birmingham, Alabama
1959...... B.S. Metallurgical Engr'g, University of Alabama,
1959-1965...... Project Engineer and Supervisor, Ross Meehan Foundry, Chattanoaga, Tenn.
1959-1967...... U.S. Army Reserve; Platoon Leader-169th Engr Bn. Co. Executive Officer- HQ, 101st Engr. Combat Battelion.
1965-196 6 ...... Research Assistant, Dept. of Metal. Eng., Univ. of Wisconsin, Madison, Wise.
1966...... M.S., Metal En., University of Wisconsin, Madison, Wis.
1966-1970...... Plant Metallurgist, Dayton Malleable Iron Co., Columbus, Ohio.
1970-1975...... Research Assistant, Dept. of Metal En., Ohio State University. PUBLICATIONS
"The Corrosion Fatigue Behavior of Biomedical Alloys, be published 1975, Journal of Biomedical Research
FIELDS OF STUDY
Major Field: Metallurgical Engineering
Studies in Electrochemistry and Corrosion: Professors M. G. Fontana, R. W. Staehle
Studies in Thermodynamics, Kinetics and Process Metallurgy Professors G. R. St. Pierre, R. A. Rapp
Studies in Physical Metallurgy Professors G. W. Powell, J. P. Hirth J. W. Spretnak
vii LIST OF TABLES
Table I - Procedures used in photo-fabrication technique.
Table II - Test material chemical composition and mechanical properties.
Table III - Geometrical multiplying factors for each (—) ratio of 5 mil crack length increments.
APPENDIX A
Table IV - Classification of Metals to Implant Locations and Functions.
Table V - A.S.T.M. Chemical Composition and Mechanical Properties of Implant Metals.
Table VI- Standard E.M.F Series of Metals.
Table VII - Galvanic Series of Some Commercial Metals and Alloys in Seawater.
Table VIII-Anodic Back E.M.F. (A.B.E.) of Metals in Equine Serum.
Table IX - Predicted Electrolytic Crevice Corrosion Resist-, ance
Table X - Relative Crevice Corrosion Resistance of Metals and Alloys in Seawater.
Table XI - Comparison of Bone Strength to Intramedullary Fixation Devices.
Table XII- Effect of Alloy Composition on SCC Resistance as Measured by Fracture Toughness.
viii Table XHE-Examples of Test Results in Measuring Material Resistance to SCC.
Table XIV - Results of Visual Examination and Weight Measurements For SCC Specimens.
Table XV r Results of Fatigue and Stress Corrosion Tests for Pacemaker Electrode Materials.
Table XVI - Strength Reduction Factors and Fretting Fatigue Limit.
Table XVII - Statistical Evaluation of Metal Performance in Implant Applications,
APPENDIX B
Table XVHI— Composition of Lactated Ringer's Solution.
Table XIX-Fatigue Specimen Processing Procedure.
Table XX- Chemical Composition and A.S.T.M. Specification for Test Materials.
Table XXI-Mechanical Properties and A.S.T.M. Specifications for Test Material.
Table XXn-Heat Treatment and Processing of Fatigue Specimens. LIST OF FIGURES
Figure 1 Sequence of crack opening and closing under the repeated loading cycle 0 to or.
Figure 2 Notation for crack growth theory.
Figure 3 Comparison of crack growth theory parameters to those of LEFM.
Figure 4 Kj values for various crack geometries.
Figure 5 Effect of thermal aging on fatigue crack growth of type 304 and 316 stainless steel.
Figure 6 Effect of grain size on fatigue crack growth of 309 stainless steel.
Figure 7 Crack growth rates for marazing and stainless steels.
Figure 8 . Comparison of fatigue crack growth in weld and parent metal of type 304 composition.
Figure 9 Comparison of fatigue crack growth of type 308 weld and type 304 basemetal at 800°F.
Figure 10 Influence of cold work in fatigue crack growth of type 304 at 800°F.
Figure 11 Fatigue crack morphology in cast (CF-8 ) stainless steel.
Figure 12 The effect of frequency on fatigue crack growth of type 304 at 1100°F.
Figure 13 Idealized crack growth of type 304 at 1000°F.
x Figure 14 Design curve for the effect of cyclic frequency on type 304 stainless steel at 1000°F.
Figure 15 Fatigue-crack propagation behavior of type 304 stainless steel at 75°F over a frequency range of 0.067 to 6.7 H2 *
Figure 1.6 S.E.M. of fracture surfaces of specimens exposed to different hold times in loading.
Figure 17 Logarithmic plot of fatigue crack length vs strain cycles for type 304 stainless steel.
Figure 18 Effect of cyclic stress ratio on fatigue crack growth of 304 at 1000°F.
Figure 19 Same as Figure 17 but plotted as a function of effective stress intensity.
Figure 20 AK vs da/dN for 18/8 austenitic steel using Frost and Pook's analyzing crack growth.
Figure 21 Types of fatigue crack growth behavior.
Figure 22 Effect of humidity on crack growth rates in type 304 stainless steel at room temperature.
Figure 23 Potential-pH diagram of iron-water system based on the oxides of iron.
Figure 24 Corrosion fatigue crack growth rate versus frequency for nickel-base superalloy.
Figure 25 Growth rates of fatigue cracks in 12% chromium steel as influenced by frequency, environment and heat treatment.
Figure 26 A schematic diagram of the superposition model.
Figure 2 7 Superposition model prediction for a slow frequency. Figure 28 Metallographic photos of fatigue cracking of type 304 in active and passive range MgCl-2 environment.
Figure 29 S.E.M. photographs of the same specimens shown in Figure 27.
Figure 30 Photomicrographs and S.E.M. photos of specimen exposed to SCC conditions in MgCl2 .
Figure 31 Photomicrograph and SEM photos of specimens exposed to superimposed cyclic loading on SCC conditions in MsCl2 *
Figure 32- Same as Figure 30 but cyclic loading imposed at higher frequency.
Figure 33 S.E.M. surface area photograph of the specimen shown in Figure 31. Slip line impingement at grain boundaries is shown.
Figure 34 Effect of applied current on corrosion fatigue life of type 304 in 10% oxolic acid.
Figure 35 Effect of applied potential on corrosion fatigue life of type 304 in 10% oxolic acid.
Figure 36 Effect of potential on fatigue crack growth of high strength aluminum in aqueous halide solutions.
Figure 37 Effect of potential on the growth rate of SCC in a high strength aluminum exposed to aqueous halide solutions.
Figure 38 Effect of frequency on potential controlled fatigue behavior of type 304 in acid-chloride medium.
Figure 3 9 Effect of stress ratio on potential controlled fatigue behavior of type 304 in acid-chloride medium. Figure 40 Direct stress fatigue machine (Wiese Model No. 101) used in crack growth study.
Figure 41 Calibration curve for load (pounds) vs strain ( in/in) in the direct stress fatigue machine.
Figure 42 Calibration curve for minimum load to pressure gauge reading for the air-hydraulic load maintenance.
Figure 43 Calibration curve-C.O.D. voltage to crack length from notch root.
Figure 44 Compact tension specimen.
Figure 45 DCB specimen configuration.
Figure 46 Photograph of DCB specimen before and after test
Figure 47 Photograph of wrap around teflon cell used in corrosion fatigue tests.
Figure 48 Dimensions of corrosion cell.
Figure 49 Photograph of test equipment used in corrosion fatigue tests.
Figure 50 Crack growth data for type 304 in air.
Figure 51 Comparison of crack growth data in air to vacuum
Figure .52 Polarization behavior of type 304 in 5N t^SO^ and IN NaCl solution.
Figure 5 3 Comparison plot of polarization data for type 304 in acid-chloride media.
Figure 54 Crack growth data at open circuit and applied potential in comparison to the same behavior in air.
xiii Figure 55 Transgranular crack morphology in type 304 fatigue specimen.
Figure 56 Fractograph showing the secondary cracking behavior of type 304.
Figure 57 Fractograph showing the pitting and crack tunneling observed for corrosion fatigue behavior of type 304.
xiv LIST OF FIGURES (continued)
Appendix A
Figure 58 An illustration of vitallium hip prosthesis, in situ, and the various designs available for this purpose.
Figure 59 Illustration of the mandible jaw prostheses-- stainless steel acrylic composite.
Figure 60 Illustrated application of type 316 stainless steel in a shoulder prosthetic device.
Figure 61 Type 316 in the Flatt finger prosthesis.
Figure 62 Cast vitallium elbow prosthesis and in in situ.
Figure 63 Titanium knee prosthesis in situ and general design.
Figure 64 Illustration of stainless steel wire fixation.
Figure 65 Illustration of the use of type 316 stainless steel in various bone plate and screw designs.
Figure 66 Illustration of type 316 stainless steel hip nail design and x-ray showing device in situ.
Figure 66 Illustration of type 316 stainless steel Harring ton Rods.
Figure 67 X-ray of Harrington Rod in situ.
Figure 68 Illustration of the use of tantalum and methyl methacrylate in cranioplasty.
Figure 69 X-ray of stainless steel intramedullory nail for internal fixation of a multiple fracture of the femur. xv Figure 70 X-ray of stainless steel (Kirschner) wire fixation for a fracture of the neck of the radius in a child.
Figure 71 . Corrosion process showing the electron flow between active and inert electrodes.
Figure 72 A type 301 stainless steel slip coupled with an Olivercrona silver clip in 3% saline-agar solution.
Figure 73 A type 301 stainless steel clip coupled with a type 420 stainless clip immersed in 3% saline- agar solution.
Figure 74 Typical example of electrode kinetic behavior of pure zinc in an acid solution.
Figure 75 Polarization studies for an active-passive system.
Figure 76 Polarization behavior of Inconel 600 annealed and cold worked type 316 L stainless steel, titanium and vitallium.
Figure 77 Circuit diagrams for conducting linear polariza tion measurements.
Figure 78 Comparison of in-vitro corrosion rates to those in-vivo for type 304 stainless and titanium.
Figure 79 Schematic potential time measurements showing three types of behavior.
Figure 80- Potential time measurements for various alloy systems in simulated body solution.
Figure 81 Example of crevice corrosion in type 316 stainless Thorston nail and plate.
Figure 82 Schematic representation of the mechanism of crevice corrosion in two stages.
xv i Figure 83 The effect of CO2 and 02 partial pressure on anodic polarization of 316L stainless steel.
Figure 84 Schematic representation of crevice corrosion of polarization behavior of CW and annealed stain less steel.
Figure 85 Potential-time tests for scratched implant alloys in Hank's solution.
Figure 86 Effect of scratch on the polarization behavior of an active passive metal.
Figure .87 Microstructure of cast vitallium.
Figure 88 Effect of cystine on the polarization behavior of copper in Ringer's solution.
Figure 89 Effect of cystine on the polarization behavior of nickel in Ringer's solution.
Figure 90 Harrington spine instrumentation.
Figure 91 Strain gauge instrumentation in application of Harrington Rod.
Figure 92 Combined effect of pitting and fatigue.
Figure 93 Effect of corrosion rate and endurance of 0.18%C steel in 370 NaCl solution.
Figure 94 Crack growth rate vs stress intensity range for Ni-Mo-V steel.
Figure 95 Types of fatigue crack growth.
Figure 96 Effect of loading frequency on crack propagation of A-533B (HSST) steel.
Figure 97 Load-time profile for vertical hip reactions during level walking.
xvii Figure 98 In vitro testing Harrington Rods in Ringer's solution.
Figure 99 Effect of body environment on fatigue crack growth of type 316 stainless steel.
Figure 100 Schematic representation of stress intensity on crack growth in corrosion fatigue.
Figure 101 Effect of fretting on fatigue limit of 1090 steel.
Appendix B
Figure 102 Schematic diagram of corrosion fatigue testing cell.
Figure 103 Plan view photographs of Type I (top) and Type II specimens used in corrosion fatigue studies.
Figure 104 Metallographic cross-sections of alloys used in corrosion fatigue studies.
Figure 105 Alternating stress (based on initial deflection) vs number of cycles for corrosion fatigue. Type 316 LVM and titanium exposed at two cyclic frequenceis to air and lactated Ringer's solution
Figure 106 Effect of specimen design on the corrosion fatigue endurance of Type 316 LVM stainless steel and titanium in lactated Ringer's solution, pH7, 37°C at 50 cpm.
Figure 107 Corrosion fatigue endurance of biomedical alloys in lactated Ringer's solution 37°C, 50 cpm. Solid symbols -pH2; open symbols -pH7.
Figure 108 Effect of pH on corrosion fatigue endurance of vitallium and 316 LVM, 50 cpm in Ringer's solution 37 °C.
xviii Figure 109 Effect of pH on corrosion fatigue endurance of Ti and Ti-6A1-4V at 50 cpm in Ringer's solution, 37°C.
Figure 110 Effect of prior stress relief on corrosion fatigue endurance on 316 LVM.
Figure 113 Open circuit potential-time plot for test specimen of each metal at rest in Ringer's solution (pH7 and 2 ).
Figure 114 Open circuit potential (MVH ) vs time (HR) for Type 316 LVM in corrosion fatigue.
Figure 115 Open circuit potential (MV^) vs time (HR) for MP-35N during corrosion fatigue.
Figure 116 Open circuit potential (MCH ) vs time (HR) for Ti-6A1-4V during corrosion fatigue.
Figure 117 Open circuit potential (MVH) vs time (HR) for titanium in corrosion fatigue.
Figure 118 Open circuit potential (MVjj) vs time (HR) for vitallium in corrosion fatigue.
Figure 119 A common transgranual mode of crack propagation was noted in all implant metal corrosion fatigue failures. Typical examples are shown above: (a) Vitallium (b) Ti-6A1-4V (c) Type 316 LVM stainless steel
Figure 120 Plane surface characterization after exposure to corrosion fatigue conditions. Note the difference in the tendency toward surface rupture or second ary cracking.
xix Figure 121 The effect of changing the environmental pH on the corrosion fatigue fractures of Type 316 LVM stainless steel and MP-35N, Top 316 LVM, left-pH 7, 60 ksi, 72,665 cycles; right-pH 2, 75 ksi, 36,401 cycles* Bottom-MP-35N, left-pH 7, 90 ksi, 98,883 cycles; right-pH 2, 80 ksi, 71,225 cycles.
Figure 122 Same as Figure 18a. Top Vitallium, left-pH 7, 55 ksi, 147,378 cycles; right-pH 2, 80 ksi, 35,867 cycles. Bottom--Titanium, left-pH 7, 80 ksi, 73,115 cycles; right-pH 2, 80 ksi, 27,006 cycles.
Figure 123. Corrosion fatigue fracture surface of Ti-6A1-4V. (a) 110 ksi, pH 7, 92,683 cycles (b) 80 kis, pH 2, 190,401 cycles
Figure 124 Open circuit potential vs time plots for Type 316 LVM observed during crack initiation study.
Figure 125 Open circuit potential vs time plot for Titanium observed during crack initiation study.
Figure 126 Open circuit potential vs time plot for vitallium observed during crack initiation study.
Figure 127 Open circuit potential vs time plot for Ti-6A1-4V observed during crack initiation study.
Figure 128 S.E.M. Montage of the fracture surface of 316 LVM stainless steel showing three distinct regions of crack propagation.
Figure 129 S.E.M. Montage of the fracture surface of Vitallium and MP-35N depicting extensive region of slow crack growth (from plane surface to center) where both stress and corrosion participate in total damage to failure.
xx Figure 130 Typical areas on the fracture surface of titanium and Ti-6A1-4V. The fracture morphology in these metals did not change across the fracture plane.
Figure 131 Effect of applied potential on the corrosion fatigue endurance of titanium and Ti-6A1-4V in Ringer's solution, 37°C, aerated, at 50 cpm. Applied stress = 90% Y.S.
Figure 132 Current density (-) vs cycles in corrosion fatigue of titanium and Ti-6A1-4V under applied cathodic potentials, 50 cpm, 3 7 °C Ringer's, stress = 90% Y.S.
Figure 133 Current density (+) vs number of cycles corrosion fatigue on Ti-6A1-4V under applied anodic potentials 50 cpm, 37°C Ringer's solution applied stress = 90%, Y.S.
Figure 13^ Applied potential (MVjj) vs average current density ( m A / c m 2 ) for Ti and Ti-6A1-4V in Ringer's solu tion, 37°C aerated, pH 7, 50 cpm. Applied stress = 90% Y.S.
Figure 135 The effect of applied potential on the corrosion fatigue fracture of Ti-6A1-4V. Note the increas ed corrosion damage (pitting) that occurred when anodic potentials were used in these tests. The applied stress was 90% of the yield strength.
Figure 13.6 Fractography of titanium specimens used in potentio- static tests. Metal dissolution in the form of deep micro-cracking occurred at high anodic ap plied potentials (right).
xx i LIST OF SYMBOLS
CF Corrosion Fatigue
SCC Stress Corrosion Cracking
S.E.M Scanning Electron Microscope
L.E.M ,F. Linear Elastic Fracture Mechanics
C.O.D Crack Opening Displacement
Stage I Crack Initiation
Stage II Crack Propagation
W.O.L Wedge Open Loading--a type of test specimen
DCB Double Cantilever Beam--a type of test specimen
KI Stress Intensity Factor for Mode I Crack Opening in psi^in of ksi'/Tn or MN-m“3/2
Fracture toughness, the Kj value for unstable K IC crack growth. Same units as Kj.
AK Stress intensity factor range for fatigue loading=Kmax(l-R). Same units as Kj.
P Load in pounds or kilograms
B Specimen thickness in inches or cm.
W Specimen width in inches or cm.
xxii Load configuration--Stress applied normal to width or transverse direction (W) and crack path is directed in the rolling (R) or longitudinal direction.
Hertz-cyclic frequency in cycles per second.
Half crack length in a centrally notched plate or edge crack length in an edge notched plate in inches or cm.
Stress ratio - ^-.stress = max K mm. stress mm. K Nominal applied stress in ksi (thousand pounds per square inch) or Mega newtons per square meter.
Mean stress = 1/2(°max+ ^in^ ksi or kg per sq. mm
Alternating stress = 1/2 cr . ) in ksi , max m m or kg per sq. mm
Yield stress in ksi or kg per mm .
Modulus of elasticity or Young's modulus in ksi or kg per mm^.
Displacement in inches or centimeters.
Half specimen width for DCB specimens in inches or cm
Initial crack length in inches or cm
Cyclic frequency in Hz
Mass transfer rate, in moles/cm^-sec
Velocity in cm/sec
xxiii Absolute temperature in degrees kelvin
Faraday's constant=96,500 coulombs/equivalent
Concentration of specie, i, in moles/cm^
Net number of electrons absorved or released in the stoichiometric reaction.
Current density in amperes km^
Exchange current density in amperes/cm^
Current in amperes.
Equilibrium potential, volts.
Standard equilibrium potential, volts.
Corrosion potential, volts.
Electric field potential, volts.
Corrosion current, amperes.
Number of ions of specie, i.
Mass of material in gms.
xx iv INTRODUCTION
Corrosion fatigue may be defined as a failure process in materials subjected to cyclic loading in an aggressive environment. In metals the process involves the initia tion of a crack from some pre-existing flaw and its propagation to final fracture. The environment serves to aide both crack initiation and propagation by corrosive action. The instances where corrosion fatigue was identified as the mode of failure in metal structures are numerous but this thesis will concern studies of the fatigue response of type 304 stainless steel in acid- chloride media and implant metals in human body-like environments.
The research work accomplished in a study of the fatigue crack propagation in type 304 stainless steel is discussed in the first part of the thesis. In addition, a review of the literature concerning related work in this field is also presented. The literature survey was intended to cover only information concerning the important parameters related to corrosion fatigue crack growth in type 304 stainless steel. However, other metal- environment systems were used when it was necessary to 2. illustrate metallurgical considerations which are important to corrosion fatigue.
The crack growth rate for type 304 in air and
5N H^SO^-IN NaCl was determined in direct tension-tension type fatigue tests at a load ratio, R, equal to 0.05 and a cyclic frequency of 3Hz. The tests conducted in acid- chloride were done at open circuit as well as applied potentials. The graphical representation of the results show three regions of crack growth dependence on applied stress intensity. Crack tunneling or branching beneath the main crack front occurred only in region II of crack growth. Anodic applied potentials accelerated crack growth while cathodic potentials reduced it.
The observed behavior is explained in terms of the environmental mechanisms operating during the three re gions of crack growth. Corrosion of the metal in the active state produced a dissolution rate which was a significant part of the crack growth rate during region
II. The possibility of corrosion product wedging is discussed for region III while in region I metal dissolution processes predominate. In all three regions there is the possibility of hydrogen entry at the crack tip and martensite formation within.the plastic zone.
A study of the corrosion fatigue crack initiation 3. process was also done as a part of this research for thesis work. The discussion of this work is presented in Appendix A and B. A literature survey concerning the predominant forms of corrosion reported for metal implants as well as related metallurgical investigations in the bio-medical field is presented in the first appendix.
The fatigue behavior of five biomedical alloys was determined by cantilever, bending fatigue tests in lactated Ringer's solution. The metals were type 316
LVM stainless steel, HS25 (a cobalt chromium alloy), commercially pure titanium, Ti-6 A1-4V, and multiphase alloy MP-35N. The tests in solution were conducted at
37°C (human body temperature) and a cyclic loading rate of 50 CPM, Air fatigue tests at the same cyclic fre quency as well as at 1700 CPM. provided a baseline to show the effect of environment in the corrosion fatigue process. The fatigue endurance was studied as a function of applied stress, solution pH and applied potential
(Ti. and Ti.-6 A1-4V only). The results of this work may be found in Appendix B. I. LITERATURE SURVEY
1.1 Introduction to the Concepts and Theory of Fatigue Crack Growth
In order to understand the mechanism of fatigue
crack growth from a scientific point of view, it is
important to start with the theories and concepts that are
used to analyze this process. The following section will
present the theory of fatigue crack growth when the material
is treated as an elastic continuum. It will show the
correction necessary for plasticity at the crack tip. Finally
a discussion of some of the factors considered important when corrosion is active at the crack tip.
1.1.1 Continuum Elasticity Treatment
The theory proposed by Frost and Dixon is
considered to contain all the essential elements necessary to
evaluate crack growth under cyclic loading conditions.^ Con
sider an infinitely wide thin sheet of homogeneous inert
material in which a central transverse crack of uniform length
is growing under a uniaxial zero to maximum loading cycle.
(Figure 1). Also consider the materials elastic, i.e., no
gross yielding plastically and the crack grows in a direction 5. perpendicular to the direction of loading.
Changes in the profile of the crack are assumed to follow the pattern given below:
NOMINAL STOSS (a) O 2 /
(b)
(c) do 2/(i On
to)
(e) o
THE SEQUENCE (a) TO(e) IS REPEATED FOR EACH SUCCESSIVE CYCLE
Figure 1 - Sequence of crack opening and closing under the repeated loading cycle 0 to e r . (Ref. 1) (1) under no load the crack is closed;
(2 ) as the load is increased to the maximum stress
the crack tip opens as a hinge and then blunts, the crack as a whole having an elliptic profile;
(3) as the load is reduced the blunted end resharpens and at zero load the crack closes completely.
When blunting occurs it can do so both elastically and
plastically but in the latter case for 90° growth, it does
so by slip out of the plane of the crack and gives rise
to striation markings on the fracture surface. Also
blunting is allowed by unbonding of the atoms at the crack
tip, that is, microcleavage. The fracture mechanism
assumed is that part of the blunted crack tip profile which
is subjected to tensile stresses retains this length when
the load is removed and then increases the initial crack
length. The crack is assumed to open up into an ellipse
under the normal stress, tion is given by Inglis1 elastic solution. If the crack is assumed to be the ellipse shown in Figure 2, it becomes necessary to calculate the advance of the tip, Al, for each cycle of load, AN. This is done by first finding an expression for S^g along the perimeter of 7. the ellipse,. It requires the transformation from ellipitic to cartesian co-ordinates, convential series approximations, and elliptic integral approximations. It is not necessary to show the details of these steps, the interested reader may follow them in the referenced text. Recognizing that the ellipse is defined by a=aQ when the normal stress has opened it up and the overall crack at zero load is 21 then a, b, and 1 are given approximately by: a = 1(1 - (£)) b = i s ± E The resulting expression for is given by: SA B - ^ + 4 £ - 3)1 Next the peripheral length of the original crack that deforms into the length on the opened-up ellipse is calculated in a similar manner. It is given by: SA ’B ' = ^ E Now the increase in crack length, Al of that part of a quadrant of the loaded crack in tension is given by: NOMINAL STRESS Figure 2. - Notation for crack growth theory (ref. 1 ) 2 Al = S^-g - i g i = 2 (In ^=- - 1 ) (3) G Since growth occurs under each cycle of loading, the desired expression for fatigue crack growth is therefore (In ^ - 1 ) (4) dN e2 ° In actual tests to confirm this theory the author found that 3 crack growth is more nearly proportional to c r 1. However the use of equation (4) can be shown to yield quite good agreement with growth laws described by other investigators. An estimate of the increase in crack tip profile due to plastic deformation is based on an experimental relationship estimating the total elastic-plastic strains around a crack in a non-workhardening material, namely, ( / ) 1 / 2 e E s t e where is the total strain, is the elastic strain, E is Young's modulus and E is the secant modulus corresponding o to the strain, € . The average tangential strain over SAB is given by elastic theory as = SAB~SA'B' = Al AB SA .B * SA 'B' and the plastic strain is E 1/2 ^ A B ^ p ~ t ^ a v . ^ a b = /E s 1/2 Al EsT7 av* S . A . B , F 1/2 F 1/2 where (~r-)ay is the value of (^— ) ' corresponding to a s s strain (^AB^p* ^ p *‘S t*ie increase in length (Sajj-Sa'B 1 ) under elastic-plastic conditions then 4lp = <£AB>P SA'B' = (fj) ^ 2 Al (5) Let Y be the flow stress for this material and then plastic 10, deformation will only modify the crack tip profile if € > 1 “ . Thus for | (f-)av = f(€AB)p and from equation (5) s (E ) * Ai . . V av Y Es av gA ,B. rE ^ / 2 = EAl or (— )„.7 = substituting again E 0 av YSa5A'B' * in equation (5)gives A 1 = E but SA'B' = U ~ from E 2 and Al = (In — ■ - 1) from (3) E2 3 2 therefore Al = (In ^ - l)2 (1 + % + 2- + ...) P 2 YE ^ E since ^ 1 . then E di = crjCTl 1 ((in i n M _ n 2 (6) dN v ^ 7 v 2E2Y As previously stated, experimental results have confirmed the 3 proportionality of growth rate to cr 1 and further tests 11. showed that ln(^£) does not vary far from a value of 9 . Therefore the final expression for fatigue crack growth in a non-workhardening material with crack tip plasticity correction becomes ^ 2 5 6 — = ------where the (7 ) dN e ^Y loading cycle 0-C is expressed &a . 1.1.2 LEMF Approach Other models for crack growth have used the above theory to express growth rate in terms of stress intensity at the crack tip. These analyses utilize the principles of linear elastic fracture mechanics to analyze the stress distribution at the tip of a crack geometry similar to that given abcve. For a nominal stress, <7", perpendicular to the crack direction the stress intensity factor, K, at the tip of an edge crack in a semi-infinite plate is given by: Kj = 1.1 < r (tra)^^ where a is the half crack length (equivalent to 1). Thus as the crack length 2a or the stress, 9 increases a corresponding increase in stress intensity occurs. Therefore the stress intensity factor provides a description of the effects of the nominal stress and crack length on the stress field at a crack tip. Under cyclic loading, a stress intensity factor range is defined by AK = Kmax - Km £n * Thus for a given notched specimen configuration, the notch tip defines a starting crack geometry and the stress intensity factor maybe calculated. This provides a convenient laboratory method of studying crack growth response of materials in any type of environment. The geometrical multiplying factors for a number of different specimen geometries are readily available in the literature. In addition this method has utility in studying flaw growth in engineering structure (see next section). At this time the utilization of the fracture mechanics approach is by far the most popular method of evaluating 2 crack growth. Paris and Erdogan have shown that fatigue crack growth in a large number of materials is given by: & = B(AK)m (8 ) dN where B is a material constant and the value of the exponent m varies between 2 and 4 dependent upon the particular material. This expression is known as the power law relationship and it will be used throughout the remainder of this text. 13. Pook compared the fatigue crack growth of a variety of materials by utilizing both expressions (7) and (8 ). In order to evaluate the constant in expression (7) the integrated form becomes, In _£ = A a r ■ . a 0 a and the constant A is then given by 1 a f A = — n— In ~ < x \ ao where a Q = the initial crack length. af = the final crack length at the end of the test. N = number of stress cycles the crack takes in growing from aQ to af. Typical crack growth data plotted for low alloy and maraging steel using both expressions is shown in Figure 3. The good correspondance between the two methods is readily apparent. 1.1.3 Corrosion Fatigue Considerations While both methods of analysis have equal merit, the linear elastic fracture mechanics approach is the most popular and it is used in this thesis work as well as the remainder of the literature review that follows. Since this work concerns fatigue growth in aqueous corrosive environments it is important to analyze the additional considerations for M N D FROM CRACK GROWTH IO TESTS'111 ON MILD STEEL R = OT4 * 0 54 9310 STHL REF 20 ->»M e .o ■b LATTICE SPACINC/CYCLE A MARAGING STEEL R = O 64 IO' dc/d'V iin cycle l______» — * I O '8 I O '7 IO '6 in/cycle Figure . - The solid line through the scatter band, indicated 9310 steel ref. 20, is from Paris' work. M 4> 15. fatigue crack growth when corrosive processes are active at the crack tip. Speidel, et al., have reported an outline of an electrochemical mass transport kinetic model for this 3 purpose. The conditions at the tip of a crack differ from the bulk environment in that the ions are consumed and formed in the crack by (a) dissolution of metal; (b) hydrolysis of metal ions in solution; (c) formation of oxide film and (d) hydrogen ion discharge. In order to calculate corrosion fatigue crack growth rate, the concentration and potential at the crack tip must be determined. This is possible with knowledge of the kinetics and equilibria of crack tip- media reactions, transport properties of the ions in solution and the crack geometry. Unfortunately much of this informa tion is unavailable at the present time but the principals necessary in such analyzes can be outlined as follows. The mass transport of a particular ionic species, j^, is given by: dc-j j. = -D. --- - ^ + C.:v (9) J i i dy 1 RT 1 1 dy 1 when the first term is Fick's law of diffusion, the second is migration of charged ions in an electrical field, and the third term is convection by mass flow of the electro- lyte. The model used to calculate mass transport is shown in Figure 3a . If the ionic fluxes along the wall of the crack, jw £, are determined by electrochemical kinetics then they can be related to mass transport fluxes, M^, along the axis of a small angle crack by a mass balance. (10) and a mass balance at y, gives: (ID The ionic fluxes along the wall of the crack are calculated by the electroneutrality condition, Zi c± = 0 (12) and for hydrolysis of metal ions in solution or metal dissolution is given by, (13) The difficulty arises in the development of kinetic equations which describe the nucleation and growth of oxide films on freshly created surfaces„ Experiments performed on titanium in 12m HCl show that new surfaces formed by Figure 3a - Schematic representation of a crack tip moving with velocity V illustrating the (moving) coordinate system used in the analysis. (Crack angle-y, mass transport rate-M-^, mass transport flux=jw ^). 18* fracture are much more active than steady state surfaces. For example, the rate of hydrogen ion reduction was a +3 factor of 10 higher and the dissolution rate to ions was 105 higher. Furthermore the measured current transients for titanium surfaces indicate very rapid passivation (of the order of milliseconds. Though current measurements can be related to flux by H = ziFj tA <14> an acceptable solution of the equations is not available to date. 1.2 Fracture Mechanics Approach to Crack Growth in Engineering Structures A review of the literature reveals that various techni ques are used to evaluate the SCC or CF crack growth behavior of a particular material environment system. In the past most studies concerned both the initiation and propagation stage of crack growth and the results were reported as total time to failure in SCC or fatigue endurance limits. One such study appears in the appendix of this text. In recent years the most common way of evaluating stress corrosion or 19. corrosion fatigue crack growth has been through the methods of linear elastic fracture mechanics. The popularity of this technique is the direct correlation of lab test results to structural design parameters. With some knowledge of a material's fracture toughness together with the flaw size and applied stress in a structure, a designer may use these test results in a predictive manner to avoid catastrophic failure. The following review will outline the usefulness and application of linear elastic fracture mechanics in design and relate the important variables used to evaluate fatigue or stress corrosion crack propagation of 304 stain less steel in various environments (with emphasis on chloride containing media). The details of the stress analysis used to characterize the fracture mechanics approach to crack growth is adequately described in other reviews,^>^_10 The purpose here is to show the value of this approach in design and its usefulness in evaluating structures under operating conditions. The principal feature of fracture mechanics is that the stress field ahead of a sharp crack can be described in terms of a single parameter, K, the stress intensity factor. This parameter is a function of stress level, cr? and flaw size, 20. a. When the deformation is such that the applied stress is normal to the crack surface, it is designated Mode I and the stress intensity factor, Kj. The mathematical express ions for Ki in various crack geometries is shown in Figure 4 . ^ At maximum constraint, the critical stress intensity for static loading and plane strain conditions is designated Kj^, the critical stress intensity factor. 3 If the Kc of a particular structural steel is 66 MN/m? (60 ksi in) by lab analysis and its yield strength is 552 MN-m-L/2 (80 ksi) at the service temperature, load rate, and plate thickness, then the maximum flaw size tolerable is 145mm (5.7 in) for a design stress of 138 MN-m^/^ (20 ksi). When the design stress is 310 MN-m^/^ (45 ksi) the same 1 1 material would tolerate a flaw size of only 28mm (1.1 in). Thus if structural design requires yield stress (CTy) loading, the maximum allowable flaw size in a component is given by: a ' ^K I r 2 K lc and the ratio of (— —) becomes a useful parameter for measuring toughness of structural materials. When structures are subjected to fatigue or stress corrosion cracking conditions the initial undetected flow 21. can grow to a critical size. The fracture mechanics approach again provides an available means of accessing sub-critical crack growth.12-14 gy assuming a flaw size commensurate with quality of fabrication, aQ, the number of cycles necessary for crack growth to a size critical, a^ , for brittle fracture can be calculated. This is possible from known fatigue crack growth laws,eg.,: Martensitic Steels: da/dN = 0.66x10"® Ferrite-Pearlite Steels: da/dN = 3 <,6x 10-10 (AKj)^ where da/dN = fatigue crack growth per cycle of loading (mm/ eye.)(in/eye). AKg=stress intensity factor range, MN/m® ^2 (ksi\TTn). From Figure 4, for an edge crack in tension AK = 1.lZ'fifAa>laavg for a fatigue load stress range, Ac. The aavg is the average value between the assumed flaw size, aQ, and the critical value calculated from known fracture toughness, Kic, and design stress, O ^ x * data (Figure. 4, part 3). Finally for a martensitic steel da/dN is replaced by Aa/AN and AN for each increment of crack growth is given by: Aa AN = ------. 66x 10-® [ 1 . l i f t Acr/a^ g ]2 •25 James shows the usefulness of this technique in estimat ing crack extension in a piping elbow located at the ins 1 1 K’.f J lJ C iH IIIICKHI.j S C R A C K K. o /•' o rrrr L l j - L b 1 MJRI'ACC C RACK i.i .ir• (i o v/iti.Kt; o iin,'Ac, u ) TTTTT, CDGE C RACK cLLLL" crm Figure 4 - Kj- values for various crack geometries, (ref 11). 23. inlet to a liquid-metal fast-breeder reactor.Although the assumptions and approximations required to analyze this structure were necessarily complex, the direct application of the above principles were illustrated. Other investigators have shown the effectiveness of this technique in other structures ’ while Wei relates the philosophy of the 18 fracture mechanics approach in design. Thus the only required information in the fracture mechanics approach to structural design is accurately 19 summarized by Luttrell, et al,, as follows: 1. Knowledge of the inherent fracture toughness (Kj c ) of the material at the temperature and loading rate coincident with the intended application. 2. Fatigue crack growth data expressed in terms of the stress intensity factor for the material-environment system at the operating temperature. 3. The estimated flaw size, shape and orientation of the largest flaw presumed in the structure. 4. A stress analysis of the structure to predict nominal stresses. 5. The appropriate stress-intensity expression for the geometry relating material property data to nominal applied stress and defect. 24. 1.3 EFFECT OF MICROSTRUCTURE ON FATIGUE CRACK GROWTH 1.3.1 Thermal Aging Thermal aging of type 304 stainless steel structural components may occur under normal operating conditions in any high temperature application or in the heat affected zone of welded parts. The aging process occurs in the temperature range of 700 to 1200°F and results in carbide precipitation within the austenite matrix. The precipitation is preferential to grain boundaries and the steel is said to be in the "sensitized state". Michel and Smith found that thermal aging (5000 hours at 1100°F) produced only a slight increase in fatigue crack growth resistance of solution annealed 304 or 316 stainless 20 steel compared to unaged specimens. However for cold worked specimens, the thermal aging decreased room temperature fatigue crack growth resistance for both steels compared to unaged material (Figure 5). The authors attributed the observed behavior to enhanced precipitation of carbide and intermetallic phases in the cold worked material. In addition, the increased availability of intergranular pre cipitation sites (dislocations and stacking faults) in the FATIGUE CRACK GROWTH RATE (IN /CYCLE I i-3 020 10 TES NEST FCO RNE (KSI^iwT) RANGE FACTOR INTENSITY STRESS 0 * 3 9 5 ( ET A HEAT 1100 *F SOLUTION - ANNEALED - SOLUTION YE 0 SANES STEEL STAINLESS 304 TYPE UNAGEO 20 iue '- fet fteml gn o fatigue on aging thermal of 5' Effect -Figure tils sel (ref. 316 steel, and 304 Type instainless propagation crack (ml ) l£ m /( N ( M »/, 40 GD 00 HR AGED 5000 T O* (593*0 * 3 9 5 ( MOO*FAT 60 60 0 100 60 00 100 > j i j io' 10 0 2 TES NEST FCO RNE KSll ) l/lN S (K RANGE FACTOR INTENSITY STRESS NGDAE 5000 HR 0 0 0 5 AGED UNAGED (593* C) (593* SOLUTION-ANNEALED YE 1 SANES STEEL STAINLESS 316 TYPE MOO* F MOO* ), 20 20 MN( )5'z) N/(m 0 100 60 25 FATIGUE CRACK GROWTH RATE (IN /CYCLE) 10 * 0 0 0 0 BO 60 40 20 10 T MOO°F (593*C) AT TES NEST FCO RNE /l ) lN l/ S K ( RANGE FACTOR INTENSITY STRESS 5000 HR 5000 AGED YE 0 SANES STEEL STAINLESS 304 TYPE ET A HEAT Figure 5 - (continued) Effect of thermal of Effect (continued) - 5 Figure 593*C) ( 1100* 1100* gn o ftge rc rpgto i Type in propagation crack fatigue on aging 0 ad 1 sanes te, rf 20). (ref. steel, stainless 316 and 304 F COLO-WORK 2?>% I_ J i j -i— (M N / ( m ) \ ) \ ) m ( / (M N 40 UNAGEO u. < v U 60 GO 26. UO 100 10 ' 27. cold worked material provided a major microstruetural contribution to the difference in crack growth observed for the annealed material. James found similar benefits from thermal aging similar materials in the annealed condition but only in tests at high temperature (1000°F) and only in the lower AK r a n g e s . 2 1 in contrast to the previous investigators, no such benefit was found in room temperature tests, in fact, in the higher AK range the aged material had higher crack growth rates. In addition, this study included aging effects from welding annealed type 304 with 308 stainless steel. Again, slower crack propagation rates were found for aged 304-308 weldments in comparison to unaged 304 basemetal. The author explains the lower crack growth in aged material at low AK's due to crack tip blunting. At higher AK, the second phase particles cannot sustain the large plastic strain and blunting is less effective. The improvement in crack growth resistance at 1000°F in the aged materail was accounted for by better compatibility (matrix to second phase particle) with respect to strain. Also crack growth rates at room temperature maybe influenced by deformation induced martensite but in high temperature tests this effect is Fatigueci;iek propagation k.si v'm 2 0 4 0 100 o o 0 20 Q 23 mm /c y c le .n-4 In / cycle V-5 20 JO 100 ,3/2 &K, M N /m ; Figure 6 - Fatigue crack growth rate da/dN for two grain sizes as a function of stress intensity range AK in 309S stainless steel. The line is a computed least-squares fit. (ref. 22). 29. eliminated. 1.3.2 Grain Size Another microstruetural property which has a significant effect on strength is grain size. Thompson studied grain size effects on fatigue crack growth of Type 22 309S stainless steel in air. The chemical composition of this steel differs from that of Type 304 in nickel content in order to attain a stable austenitic matrix and preclude martensite formation. Thompson found fatigue crack growth to be independent of grain size over 4 decades of growth rate. The grain size investigated was 45 and 480/*m. (Figure 6 ) In contrast, Hertzberg, et al., showed some dependence of fatigue crack growth on grain size, in AISI Type 305 2 o stainless steel. J Again the type 305 stainless steel was selected because it does not undergo a strain induced martensite transformation. According to the relationship: = A(AK)m , Hertzberg found the value of m to increase from the coarse grain size, 1.1 mm, (m=2.9) to that of the fine grain size, .04mm (m=3.8). However, as Thompson points out, the coarse grain size coupled with the specimen geometry used in this 30. study does not ensure polycrystallinity. Thus grain size effects in the type 305 fatigue crack propagation are questionable. 1.3.3 Martensitic Formation Since the fatigue crack growth rate in martensitic and austenitic steels is not the same, the influence of strain induced martensite in unstable austenitic steels, e.g., 24,25 Type 304 stainless steel, is an important microstruetural property. Pelloux, et al., found the fatigue crack growth rate of metastable stainless steels (Fe-31.570Ni, Fe-16Cr-13Ni, and Type 301 stainless steel) were strongly temperature 0/ dependent in the Mg-M^ temperature range. At a given AK level, the growth rate decreased as the volume fraction of strain induced y+ -a phase transformation at the crack tip increases with decreasing temperature. In another study, the same author compared fatigue crack propagation in martensitic and austenitic steels.^(Fig.7) Again at a given AK level, (AK<30 ksi in) the austenitic steels show excellent resistance to crack growth as compared to quenched and tempered martensitic steels. The difference was attributed to the deformation induced phase transformation and to the low stacking faults of the austenitic alloys. 100 TTTTT i— r—nssfrpTT d I d N 7i I0"6 inches /c r*210 < i m= 2.5 >- 11 • 304 o. O 16-13 x 24-20 1 In air 2i! □ M 9H>- A MA 4 $ v MA n argon - i s'■ k*r'* $ m= 6.5 r AKj Ksi\/inches 0.1. J_L _L J I 1 I Ll_l_ _L _L 10 100 STRESS INTENSITY FACTOR RANGE Figure 7 - Crack growth rates vs AK for maraging and stainless steels, (ref. 25) 3 2 - It is interesting to note that when all these results are compared to those reported previously for grain size effects in type 309 stainless, they fall within a scatter band less than a factor of 2 in AK. This suggests that differences between stable and unstable austenitic stainless 22 steel are not great. 1.3.4 Operating History and Processing The use of type 304 stainless steel as a structural material in nuclear reactor vessels promoted a study of radiation induced changes in microstructure and its effect 2 (i on fatigue crack propagation. Both unirradiated and irradiated and aged type 304 and 316 stainless steel were 21 7 investigated. A fast neutron irradiation of 1.3x10 n/ctir' was used but crack growth was essentially the same for both irradiated and unirradiated specimens. The authors expect greater difference in fatigue response due to radiation induced strengthening at higher fluence levels. In another study neutron radiation of A533-B, a pressure vessel steel, produced greater fatigue crack growth rates at low AK levels in tests in 27 room air. These results are in accord with other investi gators showing radiation induced void formation.23 In many applications it is common to weld type 304 33. stainless steel structures with type 308 filler material. Since it is not feasible to post heat treat the welded components in most cases, it becomes necessary to evaluate the problems created by the welding process. Two important considerations are the strength of the weld and low tempera ture sensitization of the heat affected zone. The effect of sensitization on fatigue crack growth was reported under thermal aging. It was not found to be significant at low stress intensity particularly at high operating temperatures. However it can become a real problem when conditions are favorable for SCC. Smith, et al., studied room temperature fatigue crack growth in type 308 weld deposits and the results were essentially the same as type 304.^(Figure 8) At higher temperature (800°F) the type 304 material exhibited improved crack resistance in the low (<30 ksi'Tin) stress intensity range but the important conclusion in this study was that 308 weldments have adequate strength from a fatigue standpoint (Figure 9). The effect of crack orientation relative to the weld and thermal aging of the type 308 weldments on fatigue crack o n growth in type 304 has also been reported. v The results show some benefit in crack resistance for aged 304-308 weld- 3X 10.-4 TYPE 304 STAINLESS STEEL TYPE 308 WELD DEPOSIT 7 7 F (2 5 C ) *4 u i u WELD o UJ o PLATE o WELO 2 0 -10 6 0 0 0 100 STRESS INTENSITY FACTOR RANGE Figure 8 - Comparison of fatigue crack growth behavior of Type 308 weld deposit with that of Type 304 plate at room temperature* (ref. 27) Figure Figure growth behavior of Type 308 weld with weld 308 Type of behavior growth ht fTp 34 lt a 80 F.(ref. 800 at plate 304 Type of that FATIGUE CRACK GROWTH RATE - (IN ./C YC LE) 5XIO *4 - g TES NEST FCO RANGE-KSI ) I^ S -IK E G N A R FACTOR INTENSITY STRESS YE TILS STEEL STAINLESS 4 0 3 TYPE oprsn f aiu crack fatigue of Comparison 8 0 0 F (4 2 7 C) 7 2 (4 F DEPOSIT 0 0 WELD 8 8 0 3 TYPE 0 WO 80 0 6 4 0 0 2 PLATE o WELD a PLATE 7 7F (25 C) (25 7F 7 54.8 m WELD 36. merits in comparison to unaged basemetal. There was little or no benefit between aged 304 and unaged 304-308 welds. Similar crack growth data was shown for cracks oriented along the direction of the weldment and transverse to the weld. However severe crack branching negated acceptable crack growth measurements when both weld and parent metal were in the crack path. These results indicate that as long as the crack is propagating entirely in weld metal, the fatigue strength is comparable to that of the basemetal. Only when the heat affected zone lies in the crack path is there great cause of concern. James examined the system of basemetal, heat affected 2' zone and weld deposit for crack growth in the 304-308 couple. Using tapered DCB specimens, the 308 filler metal was placed at mid-length and normal to the path of crack propagation. In almost all cases, crack growth decreased as the crack approached the weld, i.e., in the heat affected zone, and then increased as the crack grew through the weld. For this configuration, the crack growth behavior is attributed to residual compressive stresses in the heat affected zone and tensile stresses in the Weldment. When the weld deposit was placed parallel to the crack path, crack growth was the same or less than that in the basemetal. 37. 1.3.5 Texture and Presence of Second Phases James studied microstructural changes induced by cold work on fatigue crack growth in type 316 stainless 30 steel over a range of temperatures. A comparison of these results to crack growth in the annealed condition showed a greater resistance to crack propagation in the cold worked state. Shahinian, et al., confirmed this finding for cold 27 worked type 304 stainless steel. Cold work was paricularly effective for improved crack resistance in high AK range at 800°F. (Figure 10) Hertzberg, et al., found similar response to fatigue crack growth in cold worked (50%) and annealed type 305 23 stainless steel. The exponent in the power law relation ship was 2.8 and 2.9 respectively. However the same material in a recrystallized state exhibited more sensitivity to the stress intensity factor as its m value was 3.8. Furthermore fatigue crack propagation was substantially affected by load orientation in the cold worked and recrystallized state. The greater sensitivity was found for loading transverse to specimen grain orientation. These results indicate that it is important to characterize microstructural texture in evaluating fatigue crack growth in stable austenitic alloys. 2 X I0 '4 TYPE 3 0 4 STAINLESS STEEL 800 F (427 C) . - 4 ANNEALED-. ui o COLD WORKED 25 % m=2.3 . - 5 o UJ O u . »"6 20 40 60 80 100 STRESS INTENSITY FACTOR RANGE- (KSIvlN.) Figure 10 - Influence of 25 percent cold work on crack growth characteristics of Type 304 stainless steel at 800°F (427°C). (ref. 27) The effect of ferrite within the austenite matrix of SA-351 Grade CF-8 was evaluated in a study of fatigue 31 crack propagation in cast stainless steels. The essential difference between the results for cast and wrought grades was that extensive crack branching occurred in the cast specimens. (Figure H ) This had the effect of lowering crack growth rates in cast specimens relative to wrought specimens and it was persistant over test tempera tures of 75, 600, 800 and 1000°F. These results were similar to that reported previously concerning weld deposits and it is to be expected since a weldment represents a cast material. Still further evidence of improved fatigue strength 3 2 in two phase alloys exist in the literature and the thesis research reported in Appendix B . 40. . .. o 4 ^ - M ^ , J fc -t'tfUl'‘$l- Figure 11 - Crack extension through the austenite/ferrite microstructure of cast stainless steel-Type CF-8 . (ref. 31) 1.4, Mechnical Variables in Corrosion Fatigue The mechanical variables which influence the crack growth response in fatigue are mean stress (R ratio), load profile (tensile and/or compressive hold times), and frequency. In practice these variables maybe more difficult to define than in the laboratory because of complexities in structural design geometry or fluctuations in the operations of the process that the structure supports. However by careful analysis of the worst possible operating conditions and the employment of reasonable factors of safety, these variables have been successfully identified and correlated with lab tests. The following review describes the influence of the mechanical parameters on crack growth in type 304 stainless steel and other realted material-environ- ment systems. 1.4.1 Effect of Cyclic Frequency When all other variables are held constant, a lower frequency allows greater time for environmental effects on each cycle of fatigue crack growth. James studied fatigue iue 2 Te fet f rqec n the on frequency of effect The - 12 Figure FATIGUE-CRACK GROWTH RATE da/dN. inch/cycle ye 0 sanes te t 00.rf 33) 1000F.(ref. at steel stainless 304 Type rc rwhbhvo f ouin annealed solution of behavior growth crack 'b 0 1 0 5 10‘ 10'4 TES NTNST FCO RNE l lnl3,Z /lin lb , K A RANGE FACTOR SITY TEN IN STRESS TES NTNST FCO RNE kg/(mm>3/2 . K A RANGE FACTOR SITY TEN IN STRESS 7 O O 5 40 cpm 40 6 A 5 O 4 O 3 □ 2 V 1 + SPECIMEN NUMBER 4000 cpm 4000 . cpm Q.4 / + 8 cpm 083 G 400 cpm 400 FREOUENCY 4000 cpm 4000 400 cpm 400 cpm 4 cpm 4 0.4 cpm 0.4 cpm 3 08 1 CYCLIC 10 10 10 - ' -i IAIICUE CRACK GRIAVIH RATE da/dN. .'.m /cycle Figure 13 - Idealized crack growth behavior growth crack Idealized - 13 Figure fTp 34 tils sela 10F (ref.33) 1000F. at steel stainless 304 Type of FATIGUE-CRACK GROWTH RATE da/dN. in ch/cycle 104 TES NTENSI ATR AG AK. l ‘ r m /lm g k . K A RANGE FACTOR Y IT S N E T IN STRESS TES NT T ATR AG AK. b/ ' n /u lb . K A RANGE FACTOR ITY S N TE IN STRESS 7.63 4000 cpm 4000 .8 cpm 0.083 40 cpm 40 cpm 4 0.4 cpm 0.4 3/2 3/7 ; ; I CUE -CRACK GROWTH RATE da/dN. mm/cycle 44. crack propagation in solution annealed type 304 stainless steel over a frequency range of 0.013 to 67 Hz at 1000°F.^ The results showed two regions where growth is either frequ ency dependent or frequency independent (Figs. 12 & 13). The frequency dependent region followed the power law behavior, i.e., = A(f)AKn dN and by assuming n as a constant, he determined A(F) at each different frequency. A plot of A(f) vs frequency demonstrat ed that A(f) approaches a constant at higher frequencies and increases at lower cyclic rates (Figure 14). These results were interpreted as being the influence of environmental oxidation rather than a creep-fatigue interaction. It is interesting to note that in another study by the same author, crack growth was independent of frequency over a cyclic range of 0.067 to 6.7Hz (Fig. 15). These test results were obtained with the same material and load condi tions reported previously except the tests were conducted in air at room temperature. Thus environmental influence at low frequency appears to be absent in the fatigue response of type 304 in room air. In contrast, this writer always found lower fatigue endurance in type 316 stainless steel when the test frequency was decreased from 30 to 1Hz in room air. AI SI TVP[ Joa = '> '■UEt c iiitto in « iJ rvciiC ;afO-tNO ’ ivdevminule Figure 14 - Design curve for the effect of cyclic frequency on Type 304 stainless steel at 1000F. (ref. 33) (see Appendix B ) ° Other investigations confirm low frequency environmental effects on crack growth in a variety of material-environment systems. For example, Barsom shows a substantial increase in growth rate from air to 3 per cent NaCl solution as the 34 frequency is changed from 10 to 0.1 Hz. These test results on a 12 Ni-5 Cr-3 Mo steel were essentially the same in air ,3.365 UJ o > o X o •o © reo CJ re-C V © o SPECIMEN NO. 1. RW SPECIMEN MO. 2. RW SPECIMEN NO. 3. RW SPECIMEN NO. A, WR SPECIMEN NO, 5, WR 1 0,5' STRESS INTENSITY FACTOR RANGE AK. p s i / l N . Figure 15 - Fatigue-crack propagation behavior of type 304 stainless steel at 75°F over a frequency range of 0.067 to 6.7 Hz. (ref. 33) 47. or solution at 10 Hz. At lower frequencies, the difference in crack extension per cycle of fatigue for air and solution was much greater. The environmental contribution was a maximum between 0 and 0.0167 Hz. Speidel shows a similar frequency effect for Nimonic 105 35 in distilled water. In this case a critical frequency below which crack rates in water increase markedly was established at 0.1 Hz. The critical frequency for type 4340 steel water system was approximately 5 Hz and the value for Inconel 600-NaOH system was between 0.1 and 1 Hz. Finally, endurance test results for type 304 stainless steel-acid-chloride system obtained by investigators in this laboratory always show a substantial decrease in fatigue 36 life as the frequency is decreased from 10 to 0.01 Hz. Thus for the system used in this investigation as well as that for many other material-environment couples, the delet erious effect of low frequency is accentuated as the environment becomes more innocuous. 1.4.2 Effect of Load Profile Information on the effect of load profile on fatigue crack growth in type 304 stainless steel could only be found in one study. Cheng, et al., investigated the effects of hold time on the low cycle fatigue behavior of 304 37 stainless steel at 593°C. The load profiles were fully reversed cycling with no hold time, reversed cycling with tension hold only, reversed cycling with compression hold only and equal alternation of tension and compression holds. In S-N type plots, the tension hold only cycling gave shorter fatigue life in the same strain range compared with other profiles. Furthermore the fractography of tension hold specimens exhibited intergranular fracture patterns with no apparent striations while specimen in all other profiles were transgranular-ductile fractures with striations. (Fig. 1 5 ) 38 39 Since other investigators * have reported similar changes in crack morphology and fatigue response by environmental changes (air to vacuum, for example), both load profile and test media are important parameters in evaluating fracture behavior. Cheng and co-workers also took striation counts on fracture specimens to evaluate crack growth rates for all but the tension-hold load profiles. In this case the symmetrical and compression hold time data were similar, indicating comparable crack propagation rates but the zero hold time data had a lower rate. (Figure 1 7 ) Barsom studied the effect of 5 different load profiles NO HOLD TIME 5Qu ONE MINUTE TENSION HOLD §? ONE MINUTE COMPRESSION HOLD ONE MINUTE TENSION AND COMPRESSION HOLD Figure 16 - Scanning electron micrographs of the fracture surfaces of specimens tested under various loading conditions (ref. 37) iue 7 Lgrtmc lt f aiu crack fatigue of plot Logarithmic - 17 Figure eghv sri cce fr 0 sanes steel.(ref.37) stainless 304 for cycles strain vs length CRACK LENGTH, a, mm 10.0 0.2 .4 0 6.0 0.6 0.6 2.0 .0 4 e.o o.e , 1000 2000 3000 4000 0 0 4 0 0 0 3 0 0 0 2 0 0 0 1 0,1 1.0 n I F YE TILS STEEL STAINLESS 4 0 3 TYPE TAN AE*4 O' sec'1 e s '3 IO x 4 * RATE STRAIN EMPERAT "593*C E R TU A R E P M TE OD I " mln I E" TIM HOLD . V I • RANGE STRAIN TOTAL OPESO HOLD COMPRESSION A EO OD TIME HOLD ZERO O YMTIA HOLD SYMMETRICAL 0 r . -rt c pl of a FR. -1. . FiR f o ta a d f o t lo p ic m .-irith R o —L 5 TAN YLS N CYCLES, STRAIN 1 ------r $ i f (9 < k A £ o Ao o o oA c £o° O k o o—i o — 51. on fatigue crack growth in a 12Ni-5Cr-3Mo maraging steel in both air and NaCl environments.^0 There was no effect of load waveform in tests conducted in air. However the corrosion fatigue data in NaCl media demonstrated a signifi cant increase in crack growth rates for sinusoidal or triangular load profiles. This means that environmental influence is enhanced by slow rates of cyclic loading (triangular positive sawtooth sinusoidal) when compared to faster rates (negative sawtooth, square). The results in both of the above studies as well as those concerning cyclic frequency effects in other material-environment systems indicate that rates of load application and hold times are most significant to fatigue behavior in the presence of a innocuous environment or at high temperatures. Apparently for some material-environment systems it is only significant under these conditions. 1.4.3 Effect of Mean Stress Another important mechanical variable in fatigue behavior is mean load or stress. The two most common ways of expressing mean stress are in terms of either stress ratio or in the form of maximum and minimum applied stress. The forms are easily converted by means of the following 52. expressions: (7 max <7min (15) ( 7 max tf'min (16) 2 ^jnin < f" n (17) max _ ^Tm Ca (18) James studied the effect of stress ratio on fatigue crack growth in type 304 stainless steel at elevated temperature (1000°F).^ Rvalues ranged from -0.15 to 0.75 and crack growth increased from minimum to maximum R ratios at any given value of AK. This effect was more propounced at low stress intensity and slow crack growth ranges (Fig ]8). It could be accounted for by increased environmental effects during slow crack growth at high vs low R ratios. However the author surmises that closure of crack surfaces under compressive stresses terminates crack growth during that portion of the load cycle at low R ratios. It was possible to normalize all the data at different R ratios by using the following equation: STRESS INTENSITY FACTOR RANGE AK. Wj/lmml3'2 - CJ c c3 c< c I o c c o V? o nc< V : ft. ISO (J5 1(W 115. type .iw siM Nirss sun iso!irnoN annlaud) CYCUD IN AIR A! mf\ IM R X I AOOcpm. SAWtOOlH WAVEfORM STRESS INTENSITY EACIOR RANGE AK. IbMin.l Figure 18 - Effect of cyclic stress ratio on the fatigue-crack growth of Type 304 stainless steel at 1000°F: crack growth rate as a function of stress intensity factor range, (ref. 41) 54. § - c Kmai1^)”3 ” (19> For n = 1, various values of m were used to correlate the data. A value of m=0.5 gave the best correlation. (Figure 19) The slope transition at AK = 11 ksi /In had been noted in a previous study of the same material-environment system and remains unaccounted for at this time. The importance of equation (19) and its usefulness in structural design deserves further attention. Several modifications to the power law relationship developed by 2 Paris and co-workers, namely, . C has been attempted by various investigators as a means of correlating data at different stress ratios. For example, 42 a popular expression by Forman, et. al., is as follows da _ C(AK) (20) dN (l-K)Kc-AK This expression requires knowledge of material fracture toughness, Kc. In the case of type 304 and other low strength austenitic materials this information is not defined and the usefulness of equation (19)is readily apparent. In 3/2 ElFECTIVE STRtSS INTENSITY. kg/lmml n R -- o. 7?n TYPE 304 S IA IN I.IS S M i l l IM IU H IO N ANNFAIEDI . CYCIID IM AIR Al lUllPI W 'C I <1(X) cpm. SAWIOCJlllVVAVIIORM EFFECTIVE STRESS INTENSITY, Kmax[*-R)° ‘5- ^ H M 112 Figure 19 - Effect of cyclic stress ratio on the fatigue-crack growth of Type 304 stainless steel at 1000°F: crack growth rate as a function of the effective stress intensity, (ref. 41) 56. addition when James replots the data obtained from equation (20) on a material of known fracture toughness, a better correlation was obtained with his formulation. Frost and co-workers used a different approach in studying the effect of mean stress on fatigue crack growth air room temperature air for a variety of materials 43 including type 304 stainless steel. By varying the mean and alternating stress, (Note for - 1, then R = 0, and (T ^ J c T ^ - 1, R = 0, from equation Iff) . The results of a plot of CM vs 304 stainless steel composition showed a linear increase in the value of CM from £Tm 4ra ratios beginning at 3 down to 1. This corresponds to R ratios from 0.5 to 0. This confirms the previous study wherein the threshold for crack growth was greater at increased R ratios. Pook used Frost's crack growth data and his own results to establish fatigue crack propagation in terms of linear elastic fracture mechanics.^ He established a threshold stress intensity range, AK for crack growth as well as AK values during crack growth using the following expressions: K r = 1.1 The threshold values for crack growth in the type 304 material at different R values were as follows: R 103 LB-FT/in3 ^ 2 -1 6.0 5.5 0 6.0 5.5 0.33 5.9 5.4 0.62 4.6 5.2 0.74 4.1 3.7 This confirms the previous reported fact that crack growth is accentuated at higher mean stress in the low range of stress intensity. Further agreement to previous findings concerning crack growth above the threshold stress intensity was also demonstrated. (Figure 2.0) While the range of AK covered in this analysis was much smaller, the effect of increased crack growth rate at higher R ratios is shown in Pook's results. 58. R = O 33 R = O 62 R = O 74 •o+ BAND FROM CRACK GROWTH TESTS I LATTICE SPACING CYCLE ia/ i N mr»/cycl« I______I IC C 8 I O ' T in /cyclt Figure 20 - AK against da/dN for 18/8 austenitic steel. (ref. 44) 1.5 Environmental Interactions With Fatigue Crack Growth 1.5.1 Phenomenological Behavior The contribution of environment to fatigue crack growth is always dependent upon the magnitude of applied stress and the severity of the damaging species to the particular material. In the research work of this thesis, it concerns the influence of an acid-chloride media on type 304 stainless steel. Therefore, it is important to review what others have found concerning, the effectofaqueous chloride corrosion processes on the fatigue behavior of the type 304 material. The fundamental way of depicting the role of the environment in corrosion fatigue is the commonly accepted 45 approach presented by McEvilly and Wei. The types of crack growth behavior of high strength steels in aggressive environments is shown in Figure 21. Type A represents those material-environment systems where the environmental role demonstrates a symergistic action of corrosion and fatigue damage. The type B behavior represents systems where there is substantial contribution of environment enhanced sustained load component, i.e., SCC. Thus environmental 60. effects enter at some threshold value of stress intensity, namely, K^-g^. Type C behavior is typlified by systems which have type A behavior below K^g^g and type B behavior A ggrstsrvr Aggressive Aggressive "Kiscc" KfcO'Kc *°g Kmox Typs C Figure 21 - Types of fatigue crack growth behavior, (ref. 45) above it. Various material-environment systems have been 46 47 48 demonstrated to follow one of these behavioral types ’ ’ but very little information is available for the corrosion fatigue behavior of type 304 stainless steel. The fatigue crack growth behavior of type 316 stainless steel in human body environments depicts a type A behavioral Figure 22 - Effect of humidity on crack growth crack on humidity of Effect - 22 Figure ae i ye 0 sanes te a room at steel stainless 304 Type in rates eprtr, rf 27) (ref. temperature, I j FATIGUE CRACK GROWTH RAT E — {IN./-CTCI_E) 2 X 10 ' TES NEST FCO RNE-( vN.) IvlN S (K - RANGE FACTOR INTENSITY STRESS ( C) 5 (2 F 7 7 STEEL STAINLESS YE 4 0 3 TYPE 0 4 0 2 UI AIR HUMID R AIR DRY m= 3.4 m= AIR ROOM GO 80 100 62. pattern. (Figure 104;, Appendix A). This means the solution makes a significant contribution to fatigue crack growth at low and intermediate stress intensities. The severe corrosion attack and corrosion products in the vicinity of the crack which was observed in this study provides further proof of this phenomenon. James found the crack growth behavior of type 304 in liquid sodium environments followed the response of materials in an inert environment.^ The improved resistance to fatigue crack growth over that oberved in air was attributed to the lack of oxygen at the crack tip. Thus as pointed out above the lack of the damaging species within the environ ment greatly improves resistance to crack growth. In a separate study by Smith^O the aggressive specie was definite ly identified to be the oxygen content in an environment- material system similar to that used by James. The fatigue crack growth of type 304 in humid air (relative humidity 90%,) again resembles type A behavior (Figure 22).^7 The comparison between crack growth rates in dry and humid air shows a substantial increase in growth rate at AK levels less than 30 ksi/Tn. In the intermediate stress intensity level the increase is not as great (50% 63. increase) while at high levels of AK ( 35 ksi in) the rates are the same. Thus for this system, synergistic actions of fatigue and corrosion are manifested in the fatigue crack propagation. 1*5.2 Interaction of Stress Corrosion Cracking and Corrosion Fatigue One of the first studies to establish the interface between stress corrosion cracking and corrosion fatigue was done by Barsom.^l The system of 12Ni-5Cr-3Mo maraging steel and 3 1/2 per cent NaCl solution was evaluated in both SCC and CF tests. Crack growth rate was frequency dependent in the chloride solution and could be expressed in terms of the power law by: § - ® o < ‘ > A k 2 where DQ(t) is a constant for a given frequency and is a function of time for various cyclic periods. The factor DQ (t) was a constant for a given frequency and reached its maximum value between 0 and 0.016 Hz. At very low frequencies fatigue crack growth occurred below the threshold stress intensity for SCC in the 3 1/2 per cent NaCl environment. No frequency dependance was found for crack growth when the solution pH was changed from pH 7 to pH 13.5. 64. The implications of this work are that two forces exist in this system--one which accelerates growth by hydrogen embrittlement, SCC and CF at low frequencies in pH 7 environments and a second which retards growth by oxide film stability at any frequency in the pH 13.5 environment. A further insight into the corrosion processes operating in this system is evident when one examines the Pourbaix diagram for iron. (Figure 23) Note that corrosion is minimized by 0H“ ion passivation when the pH is changed from 7 to 13.5. Furthermore a quite different approach to reach immunity would have been to lower the electrochemical potential. Speidel has studied other systems which are susceptible to SCC and subjected to fatigue loading.^ His description of material behavior under these conditions is termed cyclic dependent crack growth. A typical example of this work is shown by the plot of crack growth rate vs test frequency for Nimonic 105 in water (Figure 24 ). At frequencies of 1Hz and greater the fracture is frequency independent (trans- granular) and is termed "cyclic dependent crack growth". Below 0.1 Hz, the growth becomes frequency or "time dependent", i.e., it is the time spent in each load cycle (at 65. E(v) -/ 0,8 0,6 fc(0HV 0,4 0,2 - 0,2 - 0,2 -0,4 -0,5 - 0,8 - 0,8 Fe - 1,2 -1*4 - 1,6 .“ 1,8 Figure 23„ Potential pi I Diagram of iron-watar system based cm the oxides of iron, (adapted from Pourbaix.. (ref. 52) 10 Nickel - base superalloy Nimonic 105 predicted from see crack growth rate 10 • exclusively intergranular 10 superposition partly transgranular i 10 cycle dependent .-6 crack growth 10 measured \ time dependent \ crack growth \ - 7 10 ( % stress intensity range , a K 53 M N m -8 environment : water 10 temperature : 23*C sinusoidal load wave R = 0 ,-9 10 -i 10 10 10 10 frequency , Hz Figure 24 - Corrosion-fatigue crack growth rate versus frequency for nickel-base superalloy. 67. Stresses high enough for SCC to occur) that controls overall crack growth. It is interesting to see how this interplay of SCC and CF changes this time or cyclic dependency behavior for other systems. For example the plot for a 12% steel in vacuum or water environment for two conditions of heat treatment shows that when this steel is in the proper temper condition, i.e., SCC resistant, there is no frequency dependence of crack growth but the growth rate increase at low frequency in the untempered condition definitely indicates a SCC process (Figure 25). Thus environments can enhance the fatigue crack growth rate in two different ways: by "stress corrosion under cyclic loads" (strongly frequency dependent) and by "true corrosion fatigue" for which frequency dependance is insignificant. Like Barsom, Speidel notes that the frequency effects are found at AK's below ana states that perhaps a different and smaller Kig^Q should be defined for SCC under cyclic loading if this wasn't a con tradiction in itself. Even though the interplay of SCC and CF has been observed for a number of other systems, the problem remains one of correlating the fatigue behavior with SCC data to yield a resultant environment enhanced fatigue behavior. One solution demonstrated in the literature is that of super- fatigue crack growth rate , A a /A N , [m /c yc le] 10 -i iue 5 Got ae o ftge cracks fatigue of rates Growth - 25 Figure rqec, niomn ad et treatment heat and environment frequency, in rf 35) (ref. 12 , h 0,5 en specim DCB thick 1cm = ts tmpeaur 2 C 23 re eratu p tem test , 0 = R ap pu 1 710 C/hir 10 h plus sarpe 12 et rament treatm heat % 63 MN m N M 3 6 = K A % chromium steel as influenced by influenced as steel chromium % chrom ium s t e e l, 0,2 r; C r; 0,2 l, e e t s ium chrom ylc rqec , 11/ | , frequency cyclic 0 6 0 1 i “ ------C, i cool air , C 1 ------% h environment • ■ 2 o r vacuum A - • V . 8 6 69. position models. Various investigators from Lehigh University have used this model and typical among these is 53 the one shown in Figure 53 . Here it is assumed there are no mechanical effects of load profile, load ratio and fre quency in the referenced environment. Using sustained load stress corrosion results (B in the figure) and a knowledge of K variation in time (C plot) for one cycle of loading a crack velocity Aa/At vs time plot is obtained, part D of this figure. The area under this curve is obtained and it represents the environmental component of crack growth resulting from one cycle of loading at the specific K levels. Superposition of this component upon base line fatigue data, i.e., fatigue data from an inert environment, gives the resultant environment enhanced fatigue behavior for the specific AK. A point by point calculation for a number of different AK values produces the resultant environ ment enhanced fatigue behavior shown. Such a model predicts increased cyclic growth rate for loading profiles which spend longer periods of time per cycle at elevated K levels. Thus the effects of lower frequency or higher mean load is well demonstrated by this model. One major shortcoming in this model is its inability A. Fatigue Behavior In Both 0. Sustained Loodlng Crack Velocity Reference and Agresslve vs. Time. Environments. Resultant— ■ Environment Sustained Loading Enhanced Component for Area ■ Sustained Behavior 'Specified AK Loading Component o for Specified AK <3 Reference Envlroment TIME MIN TIME B. Sustolned Loading Stress Corrosion C. K vs. Time for One Cycle of Loading Behavior. 70. Figure 26 - A Schematic Diagram of the Superposition Model (ref. 53) iue 7 SproiinMdl Prediction Model Superposition “ 27 Figure Ao/AN ~ IN./CYC. o aSo rqec, rf 53) (ref. Frequency, Slow a for to'*h- 0 1 10 10 10 -4 -> I*® “ - od rfl Sn Wave ,Slne Profile Load CPS «V2 f , 5 R ■ .0 0 Temp. Room ot Fatigue 0.27" • Anneol Mill Thlckneu I I- - 8 I- T 5% NoCI % .5 3 _ _ _ _ — cul Reiulte Actual ■— Superpoiltlon Curve *B" Curve Model,Uslng ActualReeult* AK« ■ ksi 0 3 0 2 /or. 60 80 100 0 8 0 6 0 5 72. to account for environment enhanced fatigue effects below Kxscc* T^e authors used an extrapolation of the data to obtain values below . Although an optimum frequency effect caused two such extrapolations to be made in the cited work, the agreement between model prediction and actual results was quite good as shown in Figure 27* As noted the model predictions are either very close or conservative to actual results. Thus far, this review has included studies of the interplay of SCC and CF mechanisms through lab experiments involving principles of LEFM. These as well as all other studies using this method have shown these principles to be a proven technique of accessing stage II crack growth in recent years. On the other hand some consideration should also be given to the microstructural and electrochemical observations which may help to delineate areas of SCC and CF in a given material-environment system. In this regard the work of Hirth and co-workers using the 18-8 type stainless 54 steel in a MgC^ environment is interesting. Tests were conducted at various temperatures to vary the pH^andQ.C. potential measurements were monitored during the course of each test. In addition to sustained load tests, sustained 73. W 1560.73 Figure 28 - Metallographic section cut longitudinally through the specimen. a) Crack path of the specimen passive range. b) Crack path of the specimen active range 100X (ref. 54) load with superimposed alternating load and circular bend fatigue tests were performed. The importance of analyzing all these parameters to evaluate SCC and CF processes in both stage I and II crack growth is shown in the following sections. The microstructure of fatigue test specimens is shown in Figure 2 8 . The left portion shows fatigue cracking at room temperature and 60°C and electrochemical measurements indicate a passive condition--note that it appears as an isolated crack and endurance values indicate it propagates at a high velocity; the right side of this photograph shows the cracking obtained boiling M gC^ where the metal is in the active condition--here there is a multitude of cracks and propagation of individual cracks results in a lower over all crack growth rate than at lower temperatures. Figure 2g shows the fractography obtained on specimens at various temperatures and supports the previous microstructural observation. At the top (a and b) there is very little if any dissolution effect at low temperature (in the passive range) and clear striations are noted while the lower right (d) shows active dissolution at the boiling temperature so as to obliterate striations. The microstructural and fractographic appearances of SCC test specimens are equally interesting. Figure 30 shows these features for a sustained load test specimen. The photomicrograph at the left (a) shows a transgranular mode of propagation with branching and the fractography at the right (b and c) shows a mixed mode of initiation and dimpled final fracture area. Figure 31 shows these features for similar SCC conditions but this time under cyclic loading at very low frequencies - '■ * 'jctS Figure 2 9 - SEM photographs of the fracture formation shown in Figure 27. a) Specimen stressed at 16 kp/mm^ at room temperature. 500X; b) Section cut from (a). 1700X; c) Remaining fracture region of this specimen. 500X; d) Section cut from the fracture surface of a ground specimen loaded at 26 kp/mm^ at boiling temp, which sustained 174,000 load cycles, (ref. 54) IW 1560.6a-cl''.— r—I rf Figure 30 - Specimen from the steel X2CrNil810 which was exposed in boiling 357, M g C ^ solution at a growing tensile stress with low load frequency (1 min 25 kp/mm^, then 10 min, 5 kp/mm^); Lifitime 796n corresponding to a number or load cycles of 3466. a) Longitudinal cut (caustic V2A Etch, 200X) b) SEM photograph of the fracture image in the region of high crack propagation velocity (42OX). c) Section from b at 1400X. d) SEM photograph of the fractures images in the region of fracture initiation. (420X). (ref. 54) Figure 31 - Specimen from the steel X 2CrNil810 which was exposed in boiling 357c MgCr2 solution at a growing tensile stress with low load frequency (1 min. 25 kp/mm2 , then 10 min., 5 kp/mm 2); Life time 796 h. corresponding to a number of load cycles of 3466. a) Longitudinal cut (contain V2A Etch. 200X) b) SEM photograph of the fracture image in the region of high crack propagation velocity (420X); c) Section from (b) at 1400X; d) SEM photograph of the fracture images in the region of fracture initia tion (420X). (ref. 54) 78. (.0016 Hz), note transgranular cracking persists but now definite cleavage area with high crack propagation velocity appear as well as quasi intercrystalline cracks. At higher frequencies of sustained load cyclic testing, namely, SOHz^ there is similar intercrystalline damage but here it is restricted to the initiation area and gives way to the striations as shown in Figure 32. The authors conclude that the interaction of SCC and CF appear in regions of low crack velocity and a final fractograph of this region at high magnification, (Figure 33), indicates that it is the engagement of slip lines at the grain boundaries causing an energy increase that produces an intercrystalline or quasi-intercrystallins damage. 1.5.3 Effect of Applied Current or Potential in Corrosion Fatigue The importance of electrochemical variables such as pH, applied potential or current, is well established in corrosion processes yet the application of this technology to corrosion fatigue is still in the forefront. In particular very little information is available concerning electrochemical control in the sytem under investigation in this thesis work. The effect of galvanic static and potentiostatic control Figure 32 - SEM photograph of the fracture image of a specimen exposed in boiling 357, Mo, Cr2 solution at growing tensile stress (+12, +10 kp/mm^). a) Region of fracture initiation in the zone of low crack propagation velocity, 66X b) Section from a, 400X. c) Overview of conversion to the region of high crack propagation velocity, 400X. d) Region of high crack propagation velocity, 1320X. (ref. 54) Figure 33 - SEM surface area photograph of the specimen shown in Figure 32 near the fracture, 360X. (ref. 54) in the corrosion fatigue of type 304 in 10 per cent oxalic acid produces dramatic changes in endurance.^ Under galvanostatic control, the .corrosion fatigue life decreased substantially with increased applied current. (Figure 34) The steel was in a sensitized condition and fatigue cracking initiated intergranularly. At a shallow depth (20jam ) below the initiation point the crack path was transgranular and it was at this point that the corrosion fatigue process is believed to be operating. The oxalic acid is such a good electrolyte that at high current densities in corrosion fatigue the corrosion fate is so great that the grains are corroded away before there is any crack. In potentiostatic control tests both sensitized and unsensitized 304 stainless steel was used in the oxalic acid medium and both material conditions exhibited the same behavior. In the passive potential range of +50 to +1100mV a slight increase in endurance life was obtained relative to open circuit (±0)potential. (Figure 35) Even greater life is shown for specimens controlled at high negative values of potential, i.e. cathodic protection. In fact, cathodic protection afforded an 11.5% increase in life relative to air. When high anodic potentials were applied rthout ext. cu rt variables: _ galvanostaticly applied curt density mA/cm? cycles Figure 34 - Corrosion fatigue of sensitized steel X 12 CrNi 18-8 in 10% oxalic acid (22 C) under different electrochemical conditions, (ref. 55) I 1 • -____ -JC ? | s j r r — T ‘ C f y 2 ■ V £V) 4 +sri .r 'varii jfa/es: ^± o m- ~ polenfiost. op,died potentiolmV 107 cycles Figure 3.5 - Corrosion fatigue of sensitized steel X 12 CrNi 18-8 in 10% oxalic acid (22 C) under potentiostatic conditions, (ref. 5 5 ) 84. to specimens ( -t-llOGmv) intergranular corrosion occurred and endurance life was drastically reduced. In addition corrosion fatigue results on a Niobium stablized (type 347) stainless steel in FeSO^-NaCl or 3 per cent NaCl solution show similar behavior to the above in the active, passive and transpassive state. Thus the use of electrochemical measurements and control during the corrosion fatigue process provide a more complete evaluation of the particular material-environment system. In fact further use of these methods in studying the corrosion fatigue behavior of a varity of materials in chloride environments was a part of this thesis work and is presented in Appendix B . The utilization of electrochemical control is best evaluated when the initiation and propagation stages of fatigue crack growth are separated. For this case one must turn to other material-environment systems. One such study should illustrate the important behavioral trends that are observed in stage II crack growth under applied potential. The effect of applied potential on the corrosion fatigue crack growth rate of an aluminum alloy in aqueous halide 3 environments is shown in Figure 36 . A substantial decrease ALUMINUM ALLOY 7079-T651 2.5 cm THICK PLATE CRACK ORIENTATION: TL , 5m AqUEOUS K l SOLUTION TEMPERATURE 23° C PN = ® -i/o ID'' A K-20 Kg mm SIN U S O ID A L LOAD WAVE R - 0 4 cps 10 PITTING 10-6 I -1.4 -1.2 -1.0 -0.8 -a 6 -0.4 -0.2 "p o t e n t ia l (Volts versus E ^ f H4I Figure 36 - Effect of potential on the growth rate of corrosion fatigue cracks in a high strength aluminum alio exposed to an aqueous halide solution. 85. in growth rate is noted under cathodic potentials while anodic polarization accelerated it. At both very anodic or cathodic potentials the growth rate vanishes.. It is interesting to note a similar behavior in the same system for conditions of SCC. (Figure 37) The major difference is that cathodic protection retards crack growth in SCC more than in CF. The importance of these findings in controlling the damage of corrosion fatigue in structural application is obvious but much more work is needed to establish data on other material-environment systems. Recent work in this laboratory has shown the effect of electrochemical control in the corrosion fatigue life of type 304 stainless steel in a variety of sulfate-chloride 36 environments. Since the present thesis work concerns the same material and one of those solutions, namely, 5NH2SO4 - 1N NaCl, it is important to state the important findings of that study regarding this particular material-environment sytem. The important results are as follows: 1) Effect of frequency. The endurance life is lowered with decreasing frequency regardless of applied potential. Figure 38 ; Some cathodic protection is observed at the lower frequency. The crack morphology changes from trans- ALUMINUM ALLOY 7079-7651 2.5 cm THICK PLATE CRACK ORIENTATION fL 5 m AQUEOUS Kl SOLUT ION TEMPERATURE: 23°C -3/2 STRESS INTENSITY K - 46 kg.m m z A -2 rx i c i - o.c, -1.0 -0.8 -a 6 -0.4 POTENTIAL (volts versus E, . +1 H2/H ' Figure 3 j - Effect of potential on the growth rate of stress corrosion cracks in a high strength aluminum alloy exposed to an aqueous halide solution. mV SCE - 200 200 0.0 i iue . - fet f rqec o oeta controlled potential on frequency of 3.8 Effect - Figure aiu bhvo, (e. 36) (ref. behavior,, fatigue Type Type 5 NH requency F 0 0 © © O O 2 SO ./ Hz 0.0/ 0 SS 304 4 O.l Hz O.l 0 Hz 10 /IN NaCl /IN No. of Cycles to Failure Failure to Cycles of No. ------> 10 s 8 8 . granular at 0.1 Hz. 2) Effect of R ratio. The fatigue life decreased as the R ratio decreased from 0.5 to -1 when the applied poten tial was slightly more noble or more active than the corrosion potential. Some overlapping of this trend occurred at high anodic potentials (Figure 3 9). The most damaging condition, i.e., shortest life, at any R ratio occurred approximately 250 mv more noble than the corrosion potential i.e., in the passive condition. 3) The crack and corrosion damage observed at R=-l and applied stress=yield stress at different applied potentials are given below. These observations were used in an effort to evaluate the crack initiation stage of corrosion fatigue. Applied Potential Observation -750mv No cracks up to 50% of life at 60%, cathodic cracks occurred intragran. at sur range face and mixed just below surface. No cracks from pits. -350mv Cracks at 45% of life intragranular. cathodic at surface along g.b. as well as range across grains beneath surface. Pits observed but no cracks from these attributed to higher reactivity or sensitivity to chem. attack at g.b. to start crack-then cycling cracks created a notch effect causing trans- crystalline development. Applied Potential mV, SCE iue_9 fet fsrs rtoR n oeta controlled potential on R ratio stress of Effect _39 Figure aiu eair (e. 36) (ref. behavior. fatigue -250 - 500 — 250 I 3 0 5I\ H ye 0 Sanes Steel Stainless 304 Type E D r B A in M x a M i i i i i i i i i i 9 - - SO - R = R = R = R = R = R = R 4 i s p 0 0 0 5 3 /IN Ni/IN i s p 0 5 7 8 + 0. 5 .7 -0 25 .2 0 - L0 -L rmcx m cr cr mm. cr 25 .2 0 .5 0 to i s p 0 0 0 5 3 ,a'Jl 10 ' C y c l e st oF a i l u r e l —: i i i i :— ;— ~l— 10 - rC rD i i r i i i i I T 10 * vo 0 0 90. Applied Potential Observation -250mv Active dissolution and protrusion active range grooves at 20% of life, these sites sharpened into cracks, at 60%, of life crack propagation transgran- ular, 80% of life grooving starts at g.b.; crack initiation was great er part of life, thus crack propa gation-short. Total life shortest than at any potential tested. + 100mv Slip steps and pits early, at 25%, of life G.B. attack, slip steps broaden and intergran. cracks init iated, at these sites-50% of life. Slip line attack only in region of high or max. cr. At 75% of life both hemispherical and irreg. pits formed with cracks emanating from bottom. +250mv Pits formed at 25% life and strips passive range of metal exuded on cavity surface. Cracks nucleated, at base of pit 25% of life and connected between pits-37%, of life. Propagation from these sites was transgran., at 50% life No cracks in unpitted areas, +500mv Pits and general dissolutions at SCE 25% life,g.b. attack and slip line attack. At 50%, life still no cracks at longer times more dissolution-no cracks failure by dissolution. Quite apart from the qualitative importance of the corrosion fatigue crack initiation process for type 304 in acid chloride solution (given above) it is quantitatively more important from design considerations to evaluate corrosion fatigue crack propagation in this sytem. It is the purpose of the following chapters to describe how this work was accomplished and then its results. 2. EXPERIMENTAL DETAILS Introduction The purpose of this chapter is to describe the test apparatus, material and procedures used in accessing fatigue crack propagation. The experiments were conducted in order to define the influence of acid-chloride environments on stage 2 crack growth in type 304 stainless steel. The effect of this environment on both stage 1 (crack initiation) and stage 2 crack growth processes has been determined by a 36 previous investigator in this .laboratory. However, it is of greater practical interest to design when the processes are separated and a LEFM approach is used to analyze fatigue crack growth. This chapter will present the analytical tools used in this thesis study in order to utilize a fracture mechanics approach to crack growth. 93. 2.1 Test Apparatus 2.1.1 Fatigue Machine Design All fatigue tests were conducted on a direct stress machine manufactured by Wiese Precision Lab (Model W-101) (Figure 40). The initial 10 to 30 Hz frequency capability was modified prior to installation. This was accomplished by changing the motor V-belt direct drive to a system of sprockets and speed reducers. A positive drive at low speeds was attained through an arrangement of roller chains and sprockets. The machine is capable of producing applied stress in bending as well as direct tension. It has a load capability of 4 0 0 0 pounds in direct stress and 500 pounds in bending. The direct stress system was calibrated by attaching a proving ring to the load train for load readings while corresponding strain readings were taken through the load cell. The calibration curves load to strain and load to pressure gage readings is shown in Figures 41 and 42 respect ively. The load maintainer controls the minimum load in direct stress and it is operated by an air-hydraulic system. Cyclic load or mean stress is applied by adjusting a variable Figure 40 - Direct stress fatigue machine (Wiese Model No. 101) used in crack growth study. t Load (lb.) 0 0 0 4 0 0 4 4 0 0 8 4 0 0 4 2 0 0 6 3 2000 0 0 8 2 0 0 2 3 1200 1600 0 0 4 0 0 8 0 iue41 Clbain uv fr load for curve Calibration - 1 4 Figure 100 stress fatigue machine. fatigue stress (pounds) vs strain strain vs (pounds) airto of /I oa System (pin./in.) Total Strain to Ob.) -/OI W Load f o Calibration 0 0 7 0 0 6 0 0 4 0 0 3 0 0 2 e i f j . i n . / i n . ) - * ' in /i n i ^ y ( 0 0 5 i te direct the in ) 0 0 9 0 0 8 95. 4 5 0 0 Figure 42 - The calibration curve for minimum for curve calibration The - 42 Figure hydraulic load maintainer. load hydraulic od o rsuegg raig o te air- the for reading gage pressure to load O O ro o o in ro O O O M C O in O -e— (qi) PD01 O O o o in o O O O o ro Pressure Gauge Reading (psig) 96. 97. throw crank. Load measurement is accomplished through a load cell connected to a strain indicator and cyclic load while running is monitored continuously by connecting the load cell to an oscilloscope. The obvious advantages of this machine are its simplicity in design and direct personal control during operation. However two operational problems continuously persist during testing and these represent sources of error in test results. The load cell operates off of flexture plates directly connected to the load train. Any substantial change in load train deflection is reflected by the load cell and shows up on the oscilloscope as a change in load ratio. Thus as crack growth occurs in the specimen, the crack opening displacement changes and load train deflection changes. In most cases, this problem is overcome by adjust ments to the load maintainer while running or by a shutdown to re-set proper load control. In any case, it requires continual surveillance and appropriate action. The degree of error in this case was never allowed to exceed 40 pounds in minimum or maximum load applied. The second problem was that load cell readout on the 98. oscilloscope is accurate to its least scale division in this case 20 microinches/in of strain. This means that variations of 10 microinches/in. strain in minimum and maxi mum load go undetected during operation. The degree of this error is again about 40 pounds so that total error from both sources should not exceed 100 pounds in loading. 2.1.2 C.O.D. Instrumentation Crack growth measurements in corrosion fatigue were made by an M.T.S. crack opening displacement gage (Type 632.02B22). The clip on gage was attached to knife edge holding clamps spanning the specimen notch. Displacement readings were monitored through an M.T.S. receiving unit and continuously recorded on a linear recorder. The calibration of displacement readings to crack growth was done optically. A system of grid lines was photographically etched on the specimen surface so that optical readings of crack growth could be taken with the aide c /: of 10X power magnifying glass. This technique has been reported by others and does not affect crack growth rates. ^ The grid line spacing was 0.020 inches and readings were easily taken to the nearest 0.010 inches. The details of 99. the photo-etch method are given in Table I. Four specimens were tested to establish the relationship between optical and C.O.D. measurements for crack growth. The calibration curve for this relationship is shown in Figure 43. 2.2 Test Specimen Design Compact tension specimens were used for fatigue testing. The dimensions, shown in Figure 44, were made according to A.S.T.M. specification E399-72 except the nominal thickness tr o was 0.5 inches. Specimen grips and pins were machined to the dimension specified by A.S.T.M. E99 for a tension testing clevis. For corrosion fatigue tests the pins were insulated with Flourglas pressure sensitive tape (No. 2355-2). Wedge opening loading type test specimens were used in stress corrosion cracking experiments.59 Several one inch thick specimenswere tested to establish the stress intensity below which SCC does not occur, i.e., However the high ductility of this metal prevented cracking at the root of the notch. Also in this particular design the overall geo metry and notch tip radius was not considered appropriate for SCC of low strength, highly ductile materials. The double cantiliver beam test specimen was then selected as an 100. Table I Photo-Fabrication Technique Step Procedure 1 Specimen surface hand ground through 600 grit abrasive paper and polish with 1 micron dismound. 2 Abonox wash and rinsed in hot water--final rinse with alcohol--still air dry. 3 Mask off all areas of the specimen which are to remain unetched. The following steps are done in a darkroom with only gold (anilier) safety light. Precaution is also taken to protect hands and eyes while handling the photo-resist chemicals. Also avoid breathing the fumes given off by the chemicals--cover all containers immediately upon completion of each step. 4 Apply Kodak (KTFR) photo-resist by dipping several times. Air dry on hangers in dust free area. 5 Post bake 10 minutes at 90°C then air cool to room temperature. 6 Place specimen in specimen holder and position film negative on the locating pins. Position a glass plate over the film so that intiniate contact is established between film and specimen surface. Large rubber bands are used to hold the glass plate in tact. 7 Expose for 2 minutes to ultra-violet light at a distance of 4 to 5 inches. 8 Remove specimen from holder and reverse sides--then repeat steps 6 and 7. 9 Dip and agitate specimen in developing solution for approximately 2 minutes. 10 Air dry in dust-free area. 101. Table i (Continued) Step Procedure 11 Post bake 20 minutes at 100 to 110°C. The following steps maybe done in white light. 12 Dip and agitate specimen in 42° Baume ferric chloride etching solution for approximately 5 minute or until the grid lines are clearly visible. 13 Rinse in hot water and air dry. 14 Remove undissolved resist chemicals from the specimen surface in chloroethylene stripping solution. The specimen surface must be scrubbed with a brush for effective stripping. 15 Hot water rinse and air dry. COD Voltage (V.) - 0.7 -0 2.7 -2 -2.3 -4.3 -3.9 -3.5 -3. 0.3 -5.0 iue4 - airto cre ewe optical between curve Calibration - 43 Figure -5.0 —♦ Mnmm otg Setting Voltage Minimum ♦ o— Seie 2 Specimen ■ airto- rc Lnt,, I. t C.O.D. (V) to Calibration-- (IN.) Length,a, Crack pcmn 4 Specimen • 3 Specimen 0 5.0 0 -4.82 j _ * .0 .0 .0 .0 1.70 1.60 1.50 1.40 1.30 ______4. 3 .7 -4 4.60 -4 n ... measurements. C.O.D. and rc Lnt,a (in.) a, Length, Crack i ______ -4.62 i ______i ______102. -0.625 (±.0125) CM 0.0937 D a T 0.25 0.1875 0.6875 .375 0.2718; ;±.0I25) WO' 0.2718; -1.125 m 0.25~i CM 0.6875 (±.0125) 0.1562 0.25 m 2.5(±.0125) 0.500 3.I25(±.025)------(±.025) Figure 44 - Compact tension specimen, a l l dimensions in inches. 104. alternate design. This specimen geometry has been used ft 1 quite successfully for aluminum and titanium alloys. The specimen dimension are shown in Figure 45 and photographs of the specimen before and after testing in Figure 46. Stress intensity values are calculated from the total de flection, 8 , at the load line after application of the load. f\ 1 The mathematical relationship is as follows: K = Eht3h(a+0.6h)2+h3 ]1/2 4[ (a+0 .6h)3+h^a] The same optical grid technique (reported under COD instrumentation) was employed for crack growth measurements. 2.1.3 Test Cell Design and Insulation A wrap around test cell was used for corrosion fatigue experiments. The cell was constructed from Teflon. Impolene tubing was used to conduct the acid-chloride solutions. The overall size and dimensions are shown in Figs. 47 and 48 res pectively. The flow rate control is accomplished by Teflon stop cocks both up and down stream in the system. The cell is 105. □5P ■p ■ r*p -a J J m BOLT V GROOVE SCC CRACK see f *« a c tu re CRACK FRONT m e c h a n ic a l./// I A SECTION A A Figure 45 - DCB specimen used to determine stress-corrosion crack velocity as a function of applied stress intensity Kj (after Speidel61). 106. Figure 46 - Photograph of DCB specimens before and after SCC test. Figure 4.7 - Photograph of wrap around teflon cell used in corrosion fatigue tests. 108. 5/16 Drill 3-1/8 5/8 5/16—1 1-1/4 — Slot 3/32x9/16 O-Ring Groove 3/16 Drill 3 Holes Hollow Out I x 2-1/4 Long y |*-5/8-*| I x 1-1/4 in. Section A~A Glass Window Drill S Top 3/4-20 Figure 48 - Corrosion Cell-Teflon Construction 109. o-ring sealed to the specimen surface and all specimen to cell interfaces are sealed with a silicone sealant. A photograph of this arrangement is shown in Figure k l . 2.1.4 Electrochemical Apparatus A Wenking potentiostat (Model 6839-61TRS) was used in controlled potential tests and a rotary resistor switch and recorder were connected in series with the potentiostat to monitor corrosion current continuously. Potential measurements were made with reference to a standard calomel electrode while counter electrodes were constructed of Ti- 0.2%, Pd alloy. In addition to the above apparatus, a motorpoteniometer (Model MP165) and Heathhit electrometer were used for polarization and open circuit potential measure ments. A photograph of this equipment is shown in Figure- 49. 2.2.1 Test Material and Environment The material used in this study was solution annealed type 304 stainless steel. The chemical composition and mechanical properties are shown in Table II. The test solution was 5N H2SO4 -IN NaCl. The solution was prepared from reagent grade 96% sulfuric acid, NaCl and double distilled water. 110. Figure 49 - Photograph of test equipment used in corrosion fatigue tests. Table II Certified Chemical Analysis of ASTM A240-72B Type 304 Stainless Steel United States Steel Corporation Heat No. 2P6901 C Mn P S Si Cu Ni Cr Mo B Co Pg .052 1.66 |.024 0.019 .51 .20 9.85 18.65 0.20 0.009 0.12 0.001 Certified Mechanical Properties Yield Pt* Tensile Strength Elongation Hardness (psi) (psi) (%) BHN (3000 kg) 43,800 80,100 68.0 170 *Yield point evaluated at 0.0050 inch extension 2.3 Experimental Procedure 2.3.1 Fatigue Testing Test specimens subjected to photofabrication of the grid reference lines were prepared according to the procedure outlined previously (Section 2.1.2,C.O.D. Instrumentation). All other specimens were ground and polished through 1 micron diamond along the direction of crack propagation. The specimen was placed in the alignomatic grips and all loads removed in the linkage of the load train. Output leads from the load cell were connected to strain gage instrumenta tion for zero strain readings and oscilloscope calibration. Then all load train linage was tightened and the desired strain was imposed on the specimen by activating the hydraulic system for minimum strain and adjusting the variable throw crank for mean strain. Strain readings are converted to load by the calibration curves given previously. (Section 2.1.1 Fatigue Machine Design). The system was checked dynamically for correct minimum and maximum load and proper load profile by using the oscilloscope. All fatigue test specimens were precracked a distance of 0.08 to 0.1 inches prior to testing. The crack distance was 113. measured with a Gaertner traveling microscope (least dial division 0.0001 in.) at a magnification of 50X. The maximum stress intensity factor used in precracking never exceeded 60 per cent of that used in the fatigue tests. The precracking was done in 20 mil increments such that each successive increment of crack growth from notch root was cracked at lower loads and the final or lowest load was that needed to produce the desired stress intensity for the slowest crack growth measurements. A sine wave load profile was used throughout the complete test. This form of loading has been shown to exert the full environmental effect on crack 40 growth in corrosion fatigue. The test frequency was 3 Hz. After precracking, the test cell was attached to the specimen, and the loading procedure repeated. The appropriate electrodes were inserted in the test cell after ensuring proper flow control of the acid-chloride media. C.O.D. instrumentation was used to measure crack growth throughout the balance of the test. The crack growth rate was calculated for each increment (10 to 20 mils) of crack extension and the stress intensity factor range was based on the crack length for that increment using the following equations: 114. K = — --- [29.6 (— ) 0 *5 -185.5 (— )1 * 5+655.7 (fb 2 *5 -1017 (£) 3 * b w ° . 5 w w w w +638.9(~)4 *5] w p A complete range of multiplying factors, (—) ratios, for w each 5 mil increment of crack length was calculated and is shown in Table 3 . Severe plastic deformation in the form of necking was observed in air tests at crack lengths corresponding to 0.65(^). All tests were terminated at that point. Thus crack length measurements were restricted to the range of 0.482 > (|) < 0.650. Stress Corrosion Cracking Tests In the stress corrosion cracking tests the procedure consisted of loading the specimens by turning the bolts to a predetermined deflection at the load line. If no crack developed at this point, the bolts were turned to yield greater deflections. In both WOL and DCB specimens, the material either returned elastically or yielded plastically without crack formation. Thus no data on a SCC threshold 115. Table ILL- Multiplying Factors f(— ) for Increments of Crack Growth w a a/w f (a/w) a a/w f (a/w) .125 .450 8.34 1.335 .534 10.68 .135 .454 8.44 1.340 .536 10.74 .145 .458 8.53 1.345 .538 10.81 .150 .460 8.57 1.350 .540 10.89 .155 .462 8.61 1.355 .542 10.97 .160 .464 8.67 1.360 .544 11.03 .165 .466 8.71 1.365 .546 11.12 .170 .468 8.75 1.370 .548 11.20 175 .470 8.81 1.375 .550 11.26 ; iso .472 8.86 1.380 .55 2 11.34 .185 .474 8.90 1.385 .554 11.41 .190 .476 8.97 1.390 .556 11.51 .195 .478 9.02 1.395 .558 11.59, .200 .480 9.06 1.400 .5.60 11.67 .205 .482 9.11 1.405 .562 11.75 .210 .484 9.17 1.410 .564 11.82 .215 .486 9.22 1.415 .566 11.92 .220 .488 9.27 1.420 .568 12.01 .225 .490 9.33 1.425 .570 12.09 .230 .492 9.38 1.430 .572 12.17 .235 .494 9.44 1.435 .574 12.27 .240 .496 9.49 1.440 .576 12.36 .245 .498 9.57 1.445 .578 12.45 .250 .5 00 9.60 1.450 .580 12.54 .255 .502 9.67 1.455 .582 12.63 .260 .504 9.72 1.460 .584 12.74 .265 .506 9.78 1.465 .586 12.83 .270 .508 9.85 1.470 .588 12.93 .275 .510 9.90 1.475 .590 13.04 .280 .512 9.96 1.480 .592 13.13 .285 .514 10.02 1.485 .594 13.22 .290 .516 10.08 1.490 .596 13.34 .295 .518 10.16 1.495 .598 13.45 .300 .520 10.21 1.500 . 600 13. 62 .305 .522 10.28 1.505 .602 13. 66 .310 .524 10.34 1.5,10 .604 13.77 .315 .526 10.42 1.515 . 606 13.88 .320 .528 10.48 1.520 .608 14. 00 .325 .530 10.54 1.525 .610 14.10 .330 .532 10.61 1.530 .612 14. 22 116. Table III (Continued) a a/w f(a/w) 1.535 .614 14.34 1.540 . 616 14.46 1.545 .618 14.58 1.550 .620 14.70 1.555 .622 14.84 1.560 .624 14.96 1.565 . 626 15.09 1.570 .628 15.23 1.575 .630 15.36 1.580 .632 15.48 1.585 .634 15.62 1.590 . 636 15.66 1.595 .638 15.91 1.600 .640 16.04 1.610 .644 16.33 1.620 .648 16.64 1.625 .650 16.85 1.630 .652 16.94 1.640 . 656 17.26 1.650 . 660 17.59 1.660 .664 17.94 1.670 . 668 18.29 1.675 .670 18.46 1.680 .672 . 18.64 1.690 . 676 19.00 1.700 .680 19.38 1.710 .684 19.77 1.720 .688 20.17 1.725 .690 20.38 1.730 .692 20.59 1.740 .696 21.00 1.750 .700 21.57 117. stress intensity or crack growth rate could be determined in this study. Electrochemical Polarization Tests The standard l/4xl/2 inch cylindrical test specimen was used in obtaining the polarization behavior of the material in the acid chloride environment. The potentiodynamic method was used at a scan rate of 25 millivolts per minute in both anodic and cathodic directions. Several of these tests were performed to ensure reproducibility of polariza tion data with different counter electrodes and solution volumes. Titanium-0.2%Pd material was finally selected as adequate counter electrodes for corrosion fatigue tests at applied potentials. 3. DISCUSSION OF RESULTS The following sections present the experimental results obtained in the study of fatigue crack growth of type 304 stainless steel in air and acid-chloride environments. The remainder of the thesis work concerning the corrosion fatigue properties of biomedical alloys is given in Appendix B. The environmental mechanisms which influence the controlling factors in fatigue crack propagation are used to evaluate the crack growth rates observed in this study. The major sections of this chapter are as follows: 3.1 Polarization Properties of Type 304 Stainless Steel in 5NH2SO4 and 1 N NaCl. 3.2 Corrosion Fatigue Crack Propagation in Air and at Open Circuit Potential. 3.3 Corrosion Fatigue Crack Growth at Applied Potential 3.4 Metallography and Fractography 3.5 Environmental Mechanisms Involved in the Corrosion Fatigue Behavior. 118 119. 3.1 Polarization Properties of Type 304 Stainless Steel in 5N H 9SO4. and IN NaCl Before evaluating the corrosion fatigue crack growth as a function of applied potential, it is important to establish the active-passive-transpassive behavior of this material in the reference environment. The sulfuric acid is a strong passivating medium for 304 stainless steel and the innocuous specie is the chloride ion. The polarization behavior is shown in Figure 50. The corrosion potential, i.e., freely corroding potential was observed to occur at -430 to -440 mV S.C.E. An active peak occurred at -50 mV SCE and the active to passive transition occurred in the range of -40 to +50 mV SCE. The passive range extended from this point to approximately +500 mV SCE. The passive film breakdown with subsequent pitting occurred at a transpassive potential range of +600 to +800 mV SCE. The results of this work maybe compared to that of a 36 previous investigator in Figure 51. Generally, in the referenced work there is a shift of all ranges in the noble direction relative to open circuit potential. This undoubtedly is due to the difference in solution concentration and pH. In the referenced work the solution is at lower pH and thus reaction kinetics for H+ ion reduction occur at Potential (mV, SCE) 0 0 4 - -200 300 0 -3 500 0 -5 800 0 0 7 0 0 4 0 0 6 0 0 5 -100 0 0 3 200 0 0 1 -C -6 iue5 - oaiain eair fTp 304 Type of Behavior Polarization - 50 Figure 0.1 stainless steel in 5N H 5N in steel stainless NNC. cn aews 25mV/min. was rate Scan NaCl. IN urn Dniy MA/ 2) M /C A (M Density Current 1.0 HS4I al 25°C ° 5 2 NaCl, H2S04-IN N 5 ye Stainless Steel 4 0 3 Type 10 2 SO 4 and 100 120. 1000 121. '- 4 * 1 LiJ O CO in ir> OJ VJ lOo o c: n) o CL ( un/Avujv A11siLaa_.1uo.ijnn Figure 51 - Polarization behavior of 304 stain less steel in ION H2SO4 and ION H 2SO4 /O.IN NaCl.3^ 122. more noble potentials. The results of this study established the regions of best interest to evaluate corrosion processes during fatigue tests. 3.2 Corrosion Fatigue Crack Propagation in Air and at Open Circuit Potential The results for tests conducted in air and acid- solution are shown in Figure 52. The scatter in data points at specific values of stress intensity is caused by the irregular crack path of this material. Furthermore valid data points were restricted to values of stress intensity less than approximately 27.3 MN-m"3/2 ^30 ksi \PIn) as crack tip necking occurred above this range. Open circuit corrosion fatigue crack growth in 5N H 2S04-lN-NaCl is substantially higher than that shown for results in air. The general pattern of these results follow the trend for corrosion fatigue crack growth shown by others as demonstrated schematically in Figure 1 0 4 . (Appendix A) The three regions of crack growth depicted in this diagram are as follow: Region I - at low AK values, the crack growth is extremely dependent on stress intensity. Crack Growth Rate, AN ' CYCLE 0 iue 2 Croin aiu cak propa crack fatigue Corrosion - 52 Figure orso Ftge rc Growth Crack Corrosion Fatigue ye 0 Stainless Steel, Type 304 5N H N 5 ▲ 2/=3Hez R=0.05, 1 1 2 2 3 3 40 35 30 25 20 15 10 5 5 2 SO 10 4 N al t pn ici (O.C.) Circuit Open at NaCl IN ala oe crut 25°C. 5 2 circuit, open at NaCl stainless 304 type of gation steel in air and 5N H 5N and air in steel MN- 3/ ) /2 "3 -m N (M K A (Ksi-ZTn) K A 15 20 ■■ 2 SO 4 3025 -IN ^6 -r s 123. AN '•CYCLE 124. Region II - at intermediate values of stress intensity, crack growth is strongly influenced by environment and its stress dependance is reduced. Region III - the three types of behavior depicted in this region depend on the particular metal-environment couple but in this study the type A behavior is observed, i.e., type 304 stainless steel in acid-chloride solutions. This behavior indicates further crack acceleration by an environmentally enhanced stress dependance. The three regions of crack growth corresponds to slope transitions in a log-log plot of crack growth versus stress intensity factor range. (Figure 52). The transition from region I to II and II to III occur at approximately -3 /2 20 and 25 MN-m respectively. Changes in slope represent different values of the exponent, n, in the power law relationship: M = C AKn AN Different values of n are not uncommon for fatigue behavior in austenitic steels. Values have been reported to range from 3 to 10 dependent on the particular material-environ- ment s y s t e m . 24 Slope transitions in fatigue crack growth AK (ksi VTn) 10 20 30 40 50 60 80 100 ------!------1----- 1---- 1-- 1— I I I '|------Corrosion Fatigue Crack Growth Type 3 0 4 Stainless Steel R=0.05, v=3Hz O Air r 6 10 A 5 N H 2 SO4 IN NaCI at Open Circuit (O.C.) 6.3 2.6 6.8 r 7 10 J I L 10 20 30 40 50 60 80 100 AK (MN-nrf3/2) Figure 53 - Corrosion fatigue crack growth of type 304 in air and acid-chloride in logarithmic plot. 126. results for 18-8 type stainless steel have been reported in air, elevated temperature air, and aqueous chloride environments.^’ The reasons for this behavior have either remained unclear or they were accounted for by plasticity considerations and phase transformations at the crack tip. The explanation of this behavior in the present system will be discussed later however it is interesting to note that these transitions were also evident in open circuit potential measurements. (Table 4). As crack growth changes from region I to II, the potential changes from values near the corrosion potential to those correspond ing to the active range. The potential measurements in region III of crack growth correspond to the passive range. 127. Table IV Open Circuit Potential Measurements during C.F. Crack Growth O.C. Potential No. of Polarization Region (mV., SCE) Cycles Behavior I -395 0 Corrosion -385 10000 Potential -360 22000 Range II -250 35000 Active -200 37000 Range - 50 40000 Passive III 0 42000 Range +100 to +125 to Failure Table V Current Measurements during C.F. at Anodic Applied Potentials Crack Growth Current, i, No. of Cycles Region (ma) (N) I +180 to 200 0 +240 5400 +260 12000 . +320 24000 II +315 39000 +300 56000 III +290 62000 128. 3.3 Corrosion Fatigue Crack at Applied Potentials Fatigue crack growth for type 304 in 5N H2SO4 -IN NaCl at applied potential is shown in Figure 54. The results at anodic applied potentials are substantially higher than air results and tend to parallel open circuit behavior except the growth rate is somewhat higher in all three regions. At cathodic applied potentials fatigue crack growth is suppressed but not eliminated. This slow rate of crack growth restricted observations to region I where a growth rate similar to that for air is depicted. The linear relationship between crack growth rate and stress intensity range is shown in Figure 55. A somewhat greater stress intensity dependance is demonstrated in region I and III at anodic applied potentials relative to open circuit results. In region II, the dependence is identical. The growth rate in all three regions is shifted to higher values at anodic applied potentials. At cathodic potentials the stress intensity dependance is similar to that for air and the growth rates are nearly the same. The current measurements taken during corrosion fatigue at applied anodic potentials are shown in Table V. Again the transition in crack growth is indicated by these Crack Growth Rate, AN 'CYCLE Figure 54 - Corrosion fatigue crack growth crack fatigue Corrosion - 54 Figure f v-7 6 0 0 5 0 4 5 3 0 3 5 2 0 2 15 10 5 0 Corrosion Crack Fatigue Growth =.5 is=3Hz R=0.05, Air O cd hoie t C-200 M (SCE) MV 0 0 2 - .C O at Chloride Acid" ■ y e 0 Stainless Steel, 304 Type Ai Clrd a Oe Crut (O.C.) Circuit Open at "Chloride Acid ▲ 0 Acid-Chloride at 0 .C .+ 2 0 0 MV (SCE) MV 0 0 2 .+ .C 0 at Acid-Chloride 0 Region 5 behavior of type 304 in air and acid- and air in 304 type of behavior hoie t ple potentials. applied at chloride 10 K (ksiVTn) AK MN- 3/ ) /2 '3 -m N (M K A 15 20 53 35, 30 25 O.C. -Air 10 Crack Growth Rate, AN 'CYCLE I0"6 h c 7h icr7 Figure 55 - Logarithmic plot of crack growth crack of plot Logarithmic - 55 Figure 10 s.5 v- z H -3 v Rs0.05, ye 0 Sanes Steel Stainless 304 Type orso Ftge rc Growth Crack Fatigue Corrosion 0 0 100 0 8 0 6 50 0 4 30 0 2 s tes nest ag fr all for range intensity stress vs orso ftge et results. test fatigue corrosion _1 5 H 5N ■ O A 5N H 5N A 0 _____ Air K MN-n3/ ) /2 3 -rn N (M AK 5N H 5N AK AK (ksi 1 1 1 1 2 SO — RegionI *— 2 2 4 S04 -IN NaCI (0.C.+200 MV. SCE) (0.C.+200 NaCI -IN S04 0 I (O.C.) NaCI -IN S04 I aI(..20 V SCE) MV (0.C.-200 NaCI -IN I I II I T| Region 11 Region E yTn) __ LEGEND 1 1 1 1 1 100 ______ 10' 10 I0‘ \-5 130. AN 'CYCLE 131. electrochemical measurements. More active dissolution is denoted during region II of crack growth by higher current measurements. 3.4 Metallography and Fractography The fatigue fracture surfaces were sectioned for fractography in such a way that each region of crack growth could be scanned to detect any changes in crack morphology. In addition, these sections were cut longitudinally for metallographic observations. The crack morphology in region I and III was similar in both air and acid chloride environments. The crack propagated transgranularly in an irregular manner as shown by the wavy nature of the striated surface. (Figure 56) The preferential dissolution at sites of plastic deformation and grain boundaries in the acid chloride environment are shown by the fractography in Figure 5 7 and 58. This dissolution occurred after the crack had passed and simply shows the action of corrosion processes on the fracture surface. No strong environmental influence in the fracture mechanism could be detected in these crack growth regions. Environmental interactions with fatigue is clearly demonstrated during crack growth in region II. Crack 132. Optical photomicrograph (region I & III) SEM (region I and III) Figure 56 - Crack morphology of type 304 in region I and III of crack growth in air. Above-typical metallo- graphic appearance in region I or III. Below-scanning electron micrograph of the same regions. Optical photomicrograph (region I and III) SEM (region I and III) Figure 5 7 - Crack morphology of type 304 in 5N H^SO^- 1N NaCI during region I and III of crack growth. Above- typical metallographic appearance in region I or III. Below-scanning electron micrograph of the same region. 134. Optical photomicrogaph (region I and III) fpllllllll .. * I V i l l & T l i T k k . fcp Anodic Cathodic SEH (region I and III Figure 58 - Crack morphology of type 304 after CF in acid chloride at applied potentials during region I and III of crack growth. Above-typical metallographic appearance; Below-scanning electron micrographs of this region. 135. tunnelling beneath the main transgranular crack is clearly evident in the photomicrographs of this region particularly in the acid-chloride media. (Figure 58 and 59). The tunnelling cracks show no preferences for grain boundaries but propagate in a transgranular manner similar to the main crack. The fractography confirmed these findings and the tunnelling cracks appear as deep fissures in the fracture surface. No dimples on striated steps were observed in this region because of the extensive dissolution that occurred in the form of pits and trenches associated with these areas as well as grain boundaries. 136 Optical photomicrograph (region II) j u m Air Open Circuit SEM (region II) Figure 59 - Crack morphology of type 304 after CF in air and acid chloride (open circuit potential) during region II crack growth. Above-typical metallographic appearance (note:crack tunnelling below main crack fror\t): Below-scanning electron micrograph of the same regions. 137. Optical photomicrograph (region II) SEM (region II Figure .60 - Crack morphology of type 304 after CF in acid chloride at anodic applied potentials during region II crack growth. Above-typical metallographic appearance; Below- scanning electron micrograph of the same region. 13 8. 3.5 Environmental Mechanisms Involved in the Corrosion Fatigue Behavior The corrosion fatigue behavior of type 304 stainless steel in 5N H2SO4 -IN NaCI follows the pattern of fatigue crack growth wherein there are three regions of stress dependance. The results at open circuit or applied potential indicate that the environmental influence in region I and III acts in such a way as to accelerate crack growth but the stress dependance for crack growth still dominates. The role of the environment then is manifested by its influence on fracture mechanisms at the crack tip. Open circuit potential measurements during region I crack growth indicated that oxide film formation does not occur in the area around the crack, though conditions at the crack tip maybe quite different. In region III, cracking occurs in the passive state and film formation is the important environmental mechanism. The accelerated crack growth above that produced in air for region III growth could be associated with corrosion product wedging at the crack tip. Nielsen demonstrated that corrosion products produced by advancing stress corrosion cracks in stainless steel exposed to aqueous chloride environments can exert a measureable stress.^3 Pickering, et al., confirmed this 139. fact finding that pressures of 4000 to 7000 psi could be exerted by corrosion product buildup in a constricted r e g i o n . Oxide film formation at the crack tip in this material-environment system would indicate a local increase in pH. Staehle has explained that the wedging would be aggravated in solutions of higher pH where Fe(0H)g is more insoluble.^ In addition to wedging effects oxide formation could affect plasticity of the substrate in an uncertain way so as to accelerate the formation of new crack surfaces. This could occur by reducing the reversability of slip and micro-twinning which is characteristic of the planar slip behavior of this material. It could also depend on the ductility of the film as crack propagation through brittle films is expected to exceed that for ductile films.95 a final contribution of film formation maybe the prevention of rewelding freshly created crack faces during crack closure. Region II crack growth was observed to occur in the absence of films and where dissolution processes were the most active. This was manifested by a reduction in stress dependance and dissolution trenching or tunneling beneath the 140. main crack front. The importance of dissolution in this region of crack growth can be shown best by a comparison of penetration to crack growth rate. The crack growth rate in this region is in the range of 10"^ in./cycles or 7.5x10"-* cm/sec at the test frequency. The penetration rate is given by: P.R. (_£“ ) = sec nF where ry j = corrosion current density, amp/cmz for active dissolution =0.1 amp/cm^ (Figure 51). A = gram molecular weight 56 for Fe. n = no. of equivalents/mole = 2 . F = Faradays constant = 96,500 coulomb/equivalent = density (for iron = 7.86 g/cc) /.P.R. = 0.1(56)/2(96,500)(7.86) = 0.36 x 10"5 cm/sec. Thus crack tip dissolution makes a significant contribution to region II crack growth. The same calculation for region I and III show penetration rates are two and four orders of magnitude less than the above, respectively. In all three regions of crack growth there are some additional considerations which are important to the overall corrosion fatigue behavior of this system. Dissolution processes in this low pH environment could produce hydrogen locally at the crack. This could enhance fracture through embrittlement by induced lattice strain. In addition, hydrogen entry at the crack tip could alter alloy chemistry and the corresponding Mg and MD temperature range for mastensite formation. In fact the high work hardening coefficient of this alloy is related to a austenite-martensite phase transformation resulting from plastic deformation. The presence of both a and £ martensite along the crack front has been reported for type 304 stainless steel in both corrosion fatigue and stress corrosion tests^^ The detection of martensite requires systemic and sophisticated techniques, eg., x-ray diffraction, transmission electron microscopy, selective etching methods, which were not a part of this study. However, the corrosion processes at the crack tip in the acid-chloride environment make hydrogen readily available and martensite formation by austenite saturation with hydrogen could occur in the plastic zone ahead of the crack tip. In addition to the enhancement of phase transformation 142. by environmentally induced hydrogen effects, there is the possibility that the presence of hydrogen could aide fracture by void coalescence. The dissolution occurring on the fracture surfaces in the acid-chloride solution obliterated any evidence of dimpled striations. However no such evidence could be found on striated faces in air or fatigue fracture of comperable alloys in less reactive environments (see Appendix B ) . In addition to hydrogen effects, the environment could assist crack propagation by weakening of atomic bonds. For example, Uhlig's stress adsorption model for stress corrosion cracking supports this mechanism.^ Accordingly cracking would proceed by weakening of already strained atom bonds through adsorption of the environment or its constituents. The rather wide and wavy nature of the crack front in this ductile material makes the possibility of bond breakage doubtful. Furthermore dissolution in this environment offers greater energy reduction than that caused by adsorption. Finally environmental mechanisms involving the reduction of surface energy (Rebinder effect), dissolution of piled-up dislocations which impede plastic flow in the straining metal and the enhancement of plastic flow through 143. corrosion processes which inject dislocations at the slip planes are all possible in corrosion fatigue and undoubtedly related to the crack growth observed in this study. Their specific contribution to corrosion fatigue crack propagation is not clear at this time and much more work is needed in the physio-chemical principles which delineate their precise role. 144. 4.0 CONCLUSIONS 1. The corrosion fatigue crack propagation rate of type 304 stainless steel in 5N H2SO4 - IN NaCI is substant ially higher than air. This behavior was found at a stress ratio of 0.05 and test frequency of 3Hz. 2. The crack growth behavior in the acid-chloride medium can be described by three regions of crack growth dependance on stress intensity. 3. Region I crack growth is strongly dependent on stress intensity and the role of the environment is to intensify this dependance by dissolution. 4. Region II crack growth is strongly influenced by the environment and the stress dependance is reduced. Corrosion occurs on the metal in the active state and crack tunneling beneath the main crack is observed only in this region. 5. Region III crack growth shows a strong dependance on stress intensity and the environment acts in such a way as to accelerate cracking rate. The environmental mechanism could involve corrosion product wedging effects because crack 138 145. growth in this region occurred while the metal was in the passive (film forming) state. 6 . Corrosion fatigue crack growth of type 304 in 5N H 2SO4 - IN NaCI is affected by applied electrochemical potentials. Anodic potentials accelerate it and cathodic potentials reduce it relative to open circuit behavior. 7. The crack path was transgranular in either environment during all three regions of crack growth. This was true at open circuit or applied potentials. 8 . A linear relationship between crack propagation rate and stress intensity for all three regions of crack growth shows the value of the exponent, n, in the power law relationship, ^ = CAKn, varies in the following way: n = 4 in air (region I). n = 5.12 in acid-chloride at cathodic potentials (region I). n = 6 .8 , 7.0 in acid-chloride at open circuit and anodic applied potentials respectively (region I). n = 2.6, 1.4, 1.4 in air and acid-chloride at anodic and O.C. (region II). n = 6.3 and 8 in acid-chloride at open circuit and anodic applied potentials respectively, (region III). 9. Attempts to determine a threshold stress intensity for corrosion fatigue or stress corrosion cracking were unsuccessful for this material-environment couple. Appendix A A REVIEW OF THE LITERATURE CONCERNING METALLURGICAL INVESTIGATIONS AND RELATED STUDIES OF SURGICAL IMPLANT MATERIALS. D. F. Bowers, July, 1972 Submitted as part of the annual report, "Development of Advanced Metallic Materials for Implants and Prostethic Devices," sponsored by the National Science Foundation, 1971-72, 148. 1.0 INTRODUCTION The purpose of this monograph is to present an up- to-date account of the application of metals as repair or replacment parts in the human body. While the historical development of metals in medical applications is an important record, (adequately covered elsewhere..) this review outlines the more recent research conclusions in both medical and metallurgical studies, with emphasis on the latter. The report has been written in a style that seeks to be informative to those unfamiliar with metal technology as well as practicing metallurgists who are unacquainted with the biomedical metal applications. The subject is discussed along five metallurgical points of interest which have been most extensively reported in the literature. First, the corrosion properties of implant metals, a deciding factor for human body compatibility, has received widest attention in the biomedical field. The types of corrosion associated with present day implants are pitting, galvanic, crevice, intergranular and surface abrasion combined, with fretting corrosion. Each type is defined and its related identification in implant application reported. Second, the mechanical properties considered 149. important for implant design are reviewed. The significance of physical metallurgical principles is emphasized in evaluating these properties. Although much of this research interest in the combined effects of corrosion and stress on implant metals is current, those studies recorded are reviewed with particular emphasis on the property of corrosion fatigue. Next, the number and cause of metal failures in implant applications is surveyed from worldwide reports. Finally, a confrontation with clinical problems old and new is used to establish optimum metal properties compatible with the acceptance of metals in the body. 1•1 Illustrated Medical Applications of Metals The current use of metals for the human body maybe classified as that required for surgical implants or externals devices. External devices, eg., artificial limbs, braces, traction units, etc., are not within the scope of this research effort but excellent reviews of metal performance in this regard exist elsewhere. Surgical implants are devices designed from suitable materials for extensive use in reconstructing internal o r g a n s . 73 when the devices replace a diseased or otherwise defective part of the anatomy, it is called a prosthesis. The most common type, the hip prosthesis, (Charnley) design is shown in situ in Figure 60 (bottom of 60) as well as the design diversity in this device (also in Figure 60).^“* Other parts of the body where metals are used in prosthetic devices, eg., mandible (jaw), humerus (upper arm), finger, elbow and knee are illustrated in Figures 61-65.^“^ Internal fixation devices are implants designed for the immobilization of fractured bones during the process of bone healing. For simple fracture location these devices maybe constructed of wire or plate and screw combinations as illustrated in Figures 6 6 , 67 and 73 or they maybe of special fabrication eg., intramedullary nails, steinman pins, hip nails, for more complicated fracture of long bones (in tramedullary nails), small bone fracture (pins) or fracture in bones located near the joints (hip nails) is illustrated in Figures 68 and 80. These illustrations demonstrate a variety of designs are available to the surgeon for properly handling a particular fracture. Some typical examples of this utility for a variety of bone fracture locations is shown by the x-rays in Figures 66-6 8 . ^ Still another metal implant device of the internal fixation type is that used for correction of skeletal 151- deformities, eg., Harrington Rods in the treatment of scoliosis, metal fasteners (clips and stables) for binding living tissue. These short term theraputic devices are illustrated in form in Figure 69 and in practice by the x-rays in Figure 7 0 . ^ The types of metal currently used for surgical implants are shown with respect to body anatomy location as well as particular application in Table VI. It appears the most widely used metals are Type 316 stainless steel in its various forms, i.e., standard (316), extra-low carbon (316L) or extra-low vacuum melted (316 LVM), wrought and cast cobalt-chromium alloys (called Vitallium) and titanium. The alloys of titanium, stainless steel, and cobalt-base metals are currently under clinical investigation but documented proof of success is yet to be reported. The noble metals, gold, platinum, or silver, are limited to dental use or applications where strength is not important, however, some platinum alloys, eg., Paliney 7 a platinum paladium alloy with substantial amounts of gold and silver, and Pt-407> Iridium demonstrate adequate strength with necessary electrical properties for pacemaker wires.^3 There are currently 30 A.S.T.M specifications covering 152. surgical implants with respect to material and product quality.^ All of these pertain to the three predominant metals used for implant devices, namely, Type 316 stainless steel ( 3 types), vitallium (wrought and cast forms) and titanium (unalloyed). The specified composition and strength requirements for these alloys are shown in Table VII. As noted from Table VII, there are certain advantages and disadvantages derived from the material properties of the different metals and this will become even more apparent as the technological reports describing each metal's performance is reviewed in this monograph. In summary, surgical implants maybe classified as internal prostheses, fixation or corrective devices. The variety of designs available provide the surgeon with necessary selective means to treat a particular fracture or bone replacement location. Type 316 stainless steel, vitall ium and titanium are used extensively in the manufacture of these devices. The following review will define the observed physical and chemical reactions involved with these metals and the body: second, it will discuss the research efforts and conclusions of clinical and metallurgical studies con ducted in this field and related areas and finally in success 153. and failure of these metals in implant applications is evaluated with the object of identifying the necessary improvements required in metals properties and the research efforts needed to appraise promising prospective metal systems. • 154. 2.0 Implant Metal Corrosion Corrosion is an electrochemical reaction involving two partial reactions, namely, the anodic or oxidation (loss of electrons by the metal) reaction and the cathodic or reduction (acceptance of electrons by ions in solution) 85 reaction as shown by the examples in Figure 75. These partial reactions occur simultaneously and at the same rate and usually at different locations on the metal surface. The readiness of a metal to lose its electrons in an aqueous solution at 25°C is measured by the potential it develops with respect to some standard or reference electrode, eg., hydrogen. Studies of this nature have led to the electromotive force (emf) series for pure metals (see Table VIII) where it is noted a metal’s readiness to lose electrons increases in descending the list from gold to potassium.^ Thus the less reactive metals (more corrosion resistant) are located at the top of the list with more positive potentials„ Futhermore it has been found that if any two metals are placed in contact in solution the poten tial developed is the absolute difference between their respective potentials in the series. Therefore, metals widely separated in the list develop the highest potentials 155. and maybe expected to participate in active corrosion. This is usually the case (galvanic corrosion) but not always because of other retarding factors, eg., polarization, which will become apparent later. The emf series was established for pure metals in a particular environment but this is seldom found in practice, for example, both stainless steel and vitallium are alloys composed of at least three metals and they are immersed in body fluids. Various investigators have applied the same (emf) principles to systems involving commercial alloys and QfL 0"7 more practical environments. ’ ' These results are arranged in similar manner, called the galvanic series, eg., metals in seawater (Table IX). The same interpretation applies to Table IX as that used for the emf series (Table VIII). Clarke and Hickman established a similar series for metals in biological environments (Table X).^ The table shows the potentials as anodic back emf's because the results account for the internal resistance of the solution as well as the fact that the measurements were made with respect to calomel standard electrode rather than hydrogen. Again the same interpretation is used, i.e., the inert or less active (more corrosion resistant) metals (exposed 156. independently to simulated body environments) are located at the top of the list with more positive potentials. Also the coupling of two metals widely separated in the series predicts active corrosion. As before, the more active metal (lower position in the series) is anodic to the other more noble metal (higher position in the series) and usually suffers the greatest corrosion damage. 2.1 Galvanic Corrosion The above discussion leads to the first type of corrosion frequently observed with metal implant devices, namely, galvanic corrosion. Galvanic corrosion occurs when two metals of differing electrochemical nobilities are in electrical contact and exposed to an aqueous environment. When this coupling results in such a way that the corrosion rate of the less noble metal is greatly accelerated beyond that it would experience by itself (non-coupled), it is QQ called galvanic corrosion.00 Metal nobility is given by Table IX and X explained previously. The most likely occur- rance in surgical implants would be (for bone plate and screw) combinations of different metals*^ 990 or implant surface contamination with a dissimilar metals eg., contact with surgical tools during implantation.^ 157. Today , the danger of using implant components of different metals has been largely eliminated due to earlier recognition of galvanic a ctio n ^ >93,94 ancj aiso the possible combinations in orthopedic surgery (see Table VI) are now restricted to that between Type 316 stainless steel, vitallium and titanium. While the close proximity of their position in the galvanic series predicts good compatibility as couples, changes in the environment eg., healing processes, impart unexpected behavior and uniformity of metal for 81 95 parts of a multi-component device, is always recommended. ’ When other metals exist as implants, eg., neurosurgery (Table VI), the occurrance of galvanic corrosion is again possible. Recent evidence of this phenomenon was confirmed by McFadden in colorimetric and chemical tests on neurosurg- 90 ical implants (Figures 77 and 78). In this case, there was a possibility for galvanic corrosion when surgical clips of silver and stainless steel were used together or when coupling two different types of stainless steel. Another sources of galvanic corrosion reported for implants in caused by metal contaminations. The relatively soft and highly polished surface of annealed implant devices can be contaminated with deposits of tool steel particles 158. during surgery eg., orthopedic fixation of a fracture. 96 The tool steel drills, screwdrivers, clamping devices, etc., may have compositions which rank them far below that of the titanium, stainless steel or vitallium in the galvanic series. Therefore any tool steel deposit on or near the device during insertion acts as an anode and results in cathodic attack of the implant.^ Furthermore this possibil ity, resulting in large cathode to small anode area, is known to be a most undersirable effect for galvanic corros- O C ion. This form of corrosion is even further enhanced when broken drill points or other metallic foreign bodies eg., (bullets or shrapnel) are allowed to remain near an implant as pointed out by Laing and co-workers. 96 In various 91 96 97 studies by these investigators » * the use of vitallium screwdrivers for vitallium screws and Type 420 stainless steel screwdrivers for titanium or 316 stainless screws was recommended. 2.2 Pitting Corrosion Another form of corrosion occurring with implant devices is that of pitting. Pitting is an intense localized attack that results in holes in the metal.The pitting of corros ion resistant metal surfaces is usually associated with the 159. breakdown of passivity by an aggressive environment. Body fluids are typical of such media because of the presence of a high concentration of chloride ions (a well known pitting agent). Passivity is a generic term ascribed to those metals whose corrosion resistance is attributed to the presence of a thin oxide or hydrate surface film.The orthopedic metals, titanium, type 316 stainless steel, and vitallium are examples of metals with passive oxide films. Thus the success of their application as implants depends on the stability of these films. The oxide film on stainless steel is that due to its chromium content while titanium has its own oxide and vitallium is probably a complex cobalt chromium oxide, A convenient means of evaluating film stability (or passivity effects) is attained by conducting polarization studies. There are many excellent metallurgical texts con cerning this subject^,88»99,100 ^ut ^Q r pUrp0Se Qf this review, the following is sufficient. In order to understand the significance of polarization studies, it is necessary first to recall from the previous discussion, the standard electrode potential (emf series, Table VII) associated in ranking metal nobility. These are 160. referred to as equilibrium potentials because their measure ment is based on the standard free energy change for that particular metal-solution reaction in comparison to hydrogen. The potential represents a measure of the anodic reaction, i.e., loss of electrons by the metals, as is balanced by the cathodic (reduction) reaction, i.e, acceptance of the metal electrons by ions in solution. Measurement of the kinet .s of these partial reactions (thus the corrosion process) is greatly facilitated when the particular metal in the series (called the working electrode) is short circuited with another non-participating metal (called the counter-electrode). The overall reaction is still the same but, the processes (oxidation-reduction) have been separated as shown in Figure 76.®-* If additional circuitry incorporating the counter electrode is now used to measure the potential developed at the working electrode surface, the measurement is found QQ to deviate from the equilibrium potential. ^ This deviation is referred to as polarization. Polarization can be defined as the displacement of electrode potential resulting from a net current.®-* The magnitude of polarization is expressed by overvoltage which maybe positive or negative 161. with respect to the equilibrium potential. For example, (Figure 76) in the case of zinc coupled with platinum (counter electrode), the potential developed is -0 .66v and its (equilibrium potential Table VTI)is -,76v so the over voltage is +0.1v, i.e., (-0.76 + 0.1 = -0.66v). When the magnitude of the overvoltage is greater than ±50mv along with other considerations e.g., activation polarization, the reaction rate for a particular metal- solution system is given by the Tafel equation, (Figure 79). Here the zinc and hydrogen equilibrium potentials in an acid solution are associated with a particular measured current density, iG (exchange current density) determined in separate studies and the overvoltage deviations are shown by lines construeted.with the proper Tafel slope.These deviations are drawn as a function of the measured anodic current. In the illustration, (Figure 79), the point of intersection between the lines representing the anodic (metal oxidation) and cathodic (hydrogen ion reduction) reaction determine the potential and current density where the total rates of oxidation and reduction are equal. Thus it defines the corrosion process and is called the corrosion potential, ECOrr.j an(* the corrosion current density, icorr. 162. In summary, polarization techniques provide a means of studying the kinetics of corrosion reactions in different environments by current and potential measurements. Although many details of corrosion kinetics have been necessarily omitted in order to conserve space, the above discussion defines some of the terminology needed in the following review. The implant metals (vitallium, stainless steels and titanium) react somewhat differently to polarization studies. These alloys exhibit active-passive transition as shown in Figure 80. As before the metal oxidation reaction (M-> M+ e) proceeds in the direction of positive potential away from its equilibrium potential at increasing current density. However at some point, designated Epp (the primary passive potential) there is a abrupt change to a lower dissolution rate (current density ip) which remains constant Q C over a considerable range of potential.OJ This is the passive region and its presence is due to surface film formation. Increasing the potentials still further, namely, at potentials more noble than Et, the dissolution rate (active corrosion) again increases and this region is called the transpassive region. The transpassive region 163. represents corrosion conditions where the film is no longer stable (breakdown of passivity) and pitting occurs. The potential designated E-p, is commonly called the pitting potential, Ep. When chloride ions are present, the pitting potential is lower and the range over which the metal is passive has been reduced (Figure 80-89). From a corrosion standpoint, the study is not complete without consideration of the cathodic (reduction) reactions. Figure 76 shows three examples of possible reduction reactions in an active-passive system. It is important to note that there is a marked difference in the corrosion potential and associated corrosion current density for the two cases iQlj 1q2 anc* -*-03* reduction kinetics for case results in a much greater dissolution rate (increased icorr) than for case iQ3 * Thus the type and degree of reduction kinetics are equally important in describing the corrosion process for a given active-passive metal solution environment. Experiments involving polarization techniques can be performed on a variety of metals or alloys in any desired type of solution; In particular, the adaptation of this procedure to evaluate the breakdown of passivity and onset 164. of pitting for surgical alloys in body fluids, is widely accepted in the biomedical field. For example, Figure 81 shows the results of recent study conducted by Mueller and Greener on current orthopedic alloys and two other 101 companion metals in Ringer’s solution (simulated body fluid). The polarization behavior shows that breakdown and onset of pitting corrosion occurred at approximately 600mV for both Type 316 and 316L stainless steel in the annealed condition. Vitallium (in the cast form) exhibited passivity breakdown by dissolution or breaking away of a brown corrosion layer at approximately 700mV. This type of transpassive dissolu tion was associated with a smaller rate relative to that of 316 or 316L stainless steel (note current density in the range). Pure titanium did not show breakdown and remained passive throughout the range of potentials studied. The passive breakdown potential for type 316 stainless steel in the cold worked condition was approximately 250mV less noble than that of the annealed state.(Figure 81). Thus for the same reduction reactions, the associated corrosion potential and current density in the passive range of annealed state (point A, Figure 81) would correspond to the transpassive region of the cold worked state (point B) and 165. film breakdown with onset of pitting would be predicted for the cold worked metal. This study also showed the more active potential of polarized Inconel 600. So the important conclusion of this analysis with respect to the form of implant corrosion is that the pitting attack observed on stainless steel was not present on cast vitallium or titanium. The pitting of orthopedic stainless steels becomes significant from a clinical standpoint when the presence of corrosion products become toxic to the local tissue or bone area. If small amounts of corrosion products (as characteristic of pitting corrosion) are detrimental to local body processes, then determination of low corrosion rates is a necessary requirement for clincial evaluation 102,103 of implant acceptance by the body. Greene and co-workers adopted the polarization method to accomplish this objective. The circuit diagram and a proposed biomedical design (Figure 82) was used in performing linear polarization measurements. This method employs a sequence of constant current applications to the corroding electrode (WE) and measuring the resulting potentials.^7 By plotting the overvoltage as a function of applied current density, a linear relation is obtained for 166. 102 values up to lOmV, and the slope has units of resistance. The resistance can be related to the corrosion rate (via Stern and Weisert-^^) and when this is associated to the time intervals of current applications a complete corrosion rate-time record is established. The simplicity of the circuit lends itself well to studying the in-vivo corrosion behavior of implant metals (as shown in Figure 82b) and comparing these results to those obtained in-vivo tests. Greene, et al., made such a 103 comparison for type 304 stainless steel and titanium. In vivo tests, conducted in dogs and rabbits, were of necessari ly short duration because the sensitivity of such measure ments to muscle movement required the experiments to be performed on anesthetized animals. The isotonic saline solution used for in-vitro tests correlated well with in-vivo results for 304 stainless but proved more damaging to titan ium for the same comparison. Steam sterilized samples of either metal showed a lower corrosion rate than unsterilized ones in both types of tests. In addition, the corrosion rate of both metals (sterilized) was the same in living tissue for this short period of implantation, (6 hrs dura tion, Figure 83). 167. So polarization techniques provide a means for evaluating passivity characteristics and pitting potential of metals in simulated body solutions. But it is important to recognize that the nature of pitting in this chloride containing environment is normally rapid only after prolonged exposure. Furthermore, pitting occurs by transient film breakdown, indicating that metal film repair as measured by its rate of oxide formations is an important consideration for applications involving long periods of time as in the case of implants. In an excellent study to find improved alloys for implants Hoar and Mears examined the rate of metal conversion into oxide for alloys in the passive state in both vivo and vitro tests.The characteristic behavior for film breakdown with general or pitting corrosion as well as that for a film that repairs itself and remains in tact is shown by their potential-time measurements (Figure 84). Only those metals exhibiting the latter behavior would be considered successful for long exposure periods in the body. The experiments in vitro were conducted on a number of metal alloys systems in Hank's solution (simulated body fluid) and a 0.17M sodium chloride solution, (Figure 85). The 168. results show that stainless steel and nickel alloys undergo film breakdown and pitting while titanium, vitallium and pure tantalum demonstrated the desired oxide film conditions. The authors verified the vitro results by similar vivo measurements in goats (tibial plate and screw appliance) and humans (finger implants) for exposure periods as long as 90 days. In another part of this investigation, the pitting potential for the various alloys was determined in deaerated 0.17M sodium chloride solution and verified in human blood. Increasing the chloride content or the anodic potential, in the rapidly moving solutions, resulted in the appearance of discrete, nearly hemispherical, brightened pits on the surface of all alloys. The very high breakdown potential for vitallium and titanium (greater than that where oxygen reduction occurs, thus validating the use of deaerated solution) indicates film stability and good resist ance to pitting attack relative to the nickel alloys and stainless steel. The implications of the above, relative to implant properties as well as other conclusions made in this study will be reserved until later in this report (see Future 169. Technological Improvements for Implant Metals). In summary, the corrosion process and its kinetics are adequately predicted by the type of anodic-cathodic reactions associated with the given system. Metal dissolution in the form of pitting occurs by a localized reactions at the metal surface that is associated with oxide film stability. Polarization techniques provide an accessible method for studying film breakdown on active-passive alloy systems, in particular for implant metals. The occurrance of this form of corrosion for surgical implants has been reported by many investigators^’^ ’ 103-108 particularly for the stainless steels. 2.3 Crevice Corrosion The third form of corrosion found in implant applica tions is that of crevice corrosion. Crevice corrosion is another form of localized attack that occurs at shielded 113 areas on metal surfaces in certain environments. Typical examples of such shielded areas in implants would be between bone plate and the supporting screw or at junction points between bone and a single holding device, such as, bone screw, wire, pin, or nail. Figure 86 shows a case of crevice attack on a Thornton nail and plate assembly. The mechanism for crevice corrosion in chloride containing environments, e.g., body fluids, maybe represented 85 by a two stage process, Figure 87. Initially, the oxygen reduction process proceeds at the same rate on metal surfaces both inside and outside the crevice. When the available oxygen inside the crevice is depleted, the external reduction kinetics is still sufficient to balance the metal dissolution rate and stage 1 of the mechanism continues (Figure 87). At this point, there is an excess of metal ions (positive charge ions) created in the crevice and chloride ions begin to migrate into this area to maintain charge neutrality. When the resulting metal chloride concentration is sufficient, hydrolysis occurs and the chloride dissociates into an insoluble hydroxide and a free acid. Thus the fluid within the crevice becomes highly concentrated with chloride and hydrogen ions, insoluble hydroxide and acid. The presence of these elements in the crevice accelerates metal dissolution which in turn must be balanced by increased oxygen reduction on the external sur faces and stage II of the mechanism begins. In stage II, the metal is rapidly attacked within the crevice and the external surfaces are cathodically protected by the increased m . O C oxygen reduction. The damaging effect of the chloride ion on film stability (Figure 80) makes passive metal systems such as the implant alloys, particularly vulnerable to crevice corrosion. For example, the occurrance of this form of corrosion on an isolated bone screw and its interaction with the bone has been shown to result in a bony erosion (bone resorption) problem that continues to plague orthopedic therapy.Apparently the crevice fluid and electro chemical processes described above, attack the bone as well, and the introduction of mechanical strains by muscle loading accelerates this attack at both ends of the screw. Other authors have also reported a zone of bone resorption 81 surrounding isolated screws embedded in bone. The bone reaction mechanisms will be reviewed elsewhere but the documented evidence of crevice corrosion causing implant removal is well supported in the literature. The crevice attack on implant metals requires break down of the protective oxide film and as noted previously (see 2.2 Pitting Corrosion) the susceptibility of this occurrance in body fluids is far greater with the stainless steels. Thus, with regard to orthopedic applications consideration should be given to type 316, 316L or 316LVM stainless steel. Electrochemical measurement of potential and current density in polarization techniques on Type 316L, annealed and cold worked, was recently performed by Tennese in studying crevice corrosion of orthopedic implants.In vitro-tests in Ringer's physiological solution of pH 7.2-7.4 with varying concentrations of O2 and CO2 showed that CO2 is greater than C>2 in its effect on the protective metal surface film, (Figure 88). The general effect of high CO^ was to prevent the expected difference in breakdown potential between the annealed and cold worked metals. It is inter esting to note (Figure 88 ) for the high CO2 condition that the current density in the passive range for annealed type 316 was at least one order of magnitude less than that observed for the cold worked condition. This same reduced corrosion rate was also observed for high CO2 conditions versus low CO2 when the, metal was in the same (annealed) heat treated state. Thus, the influence of CO2 concentration on the corrosion rate in the passive range was considered equal in importance to that of oxygen. Tennese proposes a crevice corrosion mechanism for conditions when both (annealed and cold worked, 316 stainless) 173. are coupled in a surgical implant, eg., the case of a screw (cold worked) and bone plate (annealed). The mechanism states that the passage of the higher anodic current in the cold worked metal (point C, Figure 89) results in film breakdown and crevice attack in the form of pitting on the annealed metal (see point A, Figure 89). The author concludes that the combined use of stainless steel of varying hardness favors the occurrence of crevice corros ion in shielded areas and the most inert device would be one fabricated completely of annealed material. Unfortunately additional research in the biomedical field concerning the improvements required in material or design to overcome the problems with crevice corrosion have not been reported, the superiority of Hastelloy C to stainless steel was determined by comparison of short time polarization measurements to predetermine (4 1/2 year) 116 corrosion rates of these metals in seawater, Table XII. Furthermore, evaluations of crevice resistance in quiet (unmoving) seawater environments rank titanium and Hastelloy C as inert materials (Table X I ) . A similar comparison in H 2 saturated sulfuric acid solution showed the same good position of these metals and included Ti-6A1-4V alloy as 174. 1 I O even more resistant to crevice attack. In summary, the two stage mechanism of crevice corrosion requires a substantial incubation time then it accelerates rapidly and frequently results in bone resorption. Many times this causes premature implant removal. The crevice form of corrosion is most common in multicomponent devices of stainless steel composition. For type 316L stainless steel couples composed of the annealed and cold worked forms, the influence of CO2 and O2 concentration is important but the mechanism appears to be the effect of increased passivation current of the cold worked metal when conducted through the annealed metal. Titanium, vitallium and a wrought version of cast vitallium (similar to Wulff's IIO alloy) are superior to Type 316 stainless steel in resistance to crevice attack while Hastelloy C and T106A1-4V show even more promise in this regard. 2.4 Surface Abrasion and Fretting Corrosion The corrosion of implant alloys is fundamentally related to the integrity of its surface film. The disruption or breakdown of this protective film by electrochemical action of the environment has attested to the different forms of corrosion described above. Implant metal surfaces may 175. also suffer damage when contacted with surgical tools or when held together, as parts of a multicomponent devices, eg., intimate contact of screw and bone plate. Surface abrasion occurs in the former case and it is continuous (fretting corrosion) throughout the period of healing in the latter. When the contact surface is under load and subjected to repeated motion so that slip or deformation occurs, the process is called fretting corrosion. This form of corrosion is particularly detrimental in the body because it produces corrosion products (metal oxide particles) which may be very toxic to body t i s s u e s . Furthermore the pits and grooves associated with fretting act as points of stress concentration for fatigue crack initiation. There are at least two possibilities for the formation of corrosion products due to surface abrasion by surgical tools during implantation. The first is a terminal effect in that once the point of film rupture heals the production of corrosion products stops. It then becomes a matter of the ability of the metal film to reform or heal itself. This is termed repassivation and it is associated with a particular dissolution rate and time of repair. Comparative dissolution rates for implant alloys as given by the passivation current density has been previously reported^^ and the repair time for a surface scratch was demonstrated by potential-time curves in Hoar and Mears work^^ (Figure 91). The latter investigation shows instantaneous repair for titanium or titanium-16 per cent molybdenum alloy but repair times of 4 and 12 hour duration was found for vitallium and type 316 stainless steel respectively. Thus while the abrasion effects maybe terminal,the production of corrosion products could continue for long times after implantation. The second possibility for abrasion induced corrosion processes occur when the scratched surface becomes a part of a crevice between surface of mating components, eg., screw and bone plate. Repassivation of the scratched surface is not possible when it lies in a crevice because of the large ohmic drop reportedly occurring along the long narrow electrolytic path between it and the cathodic surface.In this case, the reduction curve accounting for the ohmic drop (called load-line curve, CQ ^ , Figure 92) intersects the polarization curve in the active region. Within a operating crevice the magnitude of this active 177. current is sufficient to cause serious pitting of stainless steel in a few months. The condition is avoided with metals of more negative passivation currents densities such as titanium or its alloys. The other form of corrosion associated with abrasion of implant surfaces was that of fretting. Fretting corrosion produces abraded particles which could adhere to implant surfaces and thereby lead to other forms of corrosion. If the particles differ in composition or metallurgical structure to that of the surface, galvanic corrosion is ... 81,96,97 possible. Alternatively, regardless of condition, the deposit shields a part of the surface creating the necessary 85 criteria for crevice attack. Metal susceptability to fretting corrosion is directly related to its wear properties (reviewed elsewhere in this report) and vitallium is 120 particularly vulnerable from this standpoint. In fact, when vitallium is used in prosthetic devices, (knee prostheses) the fretting may yield damaging results to the body. Type 316 stainless steel seems to be less of a problem under fretting conditions in multicomponent devices, (screw-bone plate) than titanium or vitallium. 178. Summarizing, surface abrasion and fretting corrosion represent the most perplexing forms of implant corrosion because of the inherent design requirements of multicomponent devices and the surgical techniques used in bone fixation. Furthermore all current metals used in orthopedic surgery are highly subject to this form of corrosion either from abrading the surface film during surgery or fretting during theraputic applications. 2.5 Intergranular Corrosion The conditions for intergranular corrosion distinctly belong to metallurgical processing techniques or high temperature metal applications. Therefore the only instances reporting its occurrance in the biomedical field are consequences of manufacturing errors. Intergranular corrosion is the preferential attack of grain boundary areas originating from the presence of elements of lower corrosion resistance. These elements maybe impurties such as inclusions or metal phase depleted 85 of one of the alloying compotents. The most susceptible implant alloy to this form of corrosion is sensitized type 316 stainless steel. A sensitized condition is manifested by a chromium depleted zone adjacent to the grain boundary 179. which is much less corrosion resistant that the matrix grains. The carbide precipatation causing chromium depletion occurs for carbon contents greater than .02 per 85 cent in a temperature range of 950 to 1450°F. Proper heat treating techniques and/or avoiding this sensitizing temperature range during application will prevent inter granular corrosion of type 316 stainless steels. Such heat treatments are adequately specified for stainless steels 84 in surgical implant applications. In addition recent manufacturing processes provide vacuum melted, extra low carbon grades of type 316 stainless steel (316 LVM) for virtually all implant devices (see Table VI) preventing sensitizing possibility. Furthermore available specifications covering the metal cleanliness from an inclusion or impurity standpoint should preclude instances of inter- 123 granular attack in this regard. Although there are no reported problems with preferent ial attack in cast vitallium alloys the possibility of such occurrance appears ever present, (Figure 93). In this case the line of spherical carbides may prove a site of 81 weakness. In conclusion, the presence of intergranular corrosion 180. in current implant applications has been reported for stainless steels devices but is simply the result of unacceptable metallurgical processing. The danger of this occurrance is avoided by appropriate choice of reputable manufacturers and/or specifying available corrosion tests to check heat treatment of the higher carbon grades or 123 specifying the lower carbon, vacuum melted grade. Finally, pure titanium is immune to this form of corrosion, but the presence and orientation of carbides could make vitallium susceptible. 2.6 Summary of Implant Corrosion The electrochemical nature of the corrosion process has been defined only in enough detail to describe the five most reported forms of corrosion found in surgical implant applications. These forms were galvanic, pitting, crevice, surface abrasion and fretting, and intergranular corrosion. Current applications where these forms of corrosion are found is limited to the three most widely used metals, namely, 316 stainless steel, Vitallium, and titanium. The fundamental aspects of corrosion involve simultane ous oxidation (anodic) and reduction (cathodic) reactions of a given metal solution environment. The oxidation reaction 181. was described by the readiness of a given metal to lose electrons and when this was measured by electromotive force in a standard cell, the emf series was established. Reduction reactions were possible in five different ways (Figure 76) and thus pose no problem in completing the total corrosion process. In more practical situations the standard emf series was shown to be inadequate and the galvanic series was introduced. Finally polarization studies were used to describe the techniques used to study the kinetics of active or active-passive systems. Galvanic corrosion was recognized early in biomedical metal applications and today its occurrance in orthopedics is reported mainly as surface contamination from surgical tools or clamping devices. Furthermore there are only three metals widely used in this field and all have relative close positions in the galvanic series. This limits the possibility of mixed metals as well as accelerated attack in cases of mixed couples. Such success was not found in the field of neurosurgery where the possibility of more active metals, eg., silver plus Type 301 stainless steel, may still exist as galvanic couples. Pitting corrosion of implant metals was shown to be 182. dependent on the stability of a protective oxide film. The use of polarization techniques was used to describe the breakdown potential (beginning of transpassive region) and onset of pitting for the three most common implant metals. Furthermore the type and degree of the reduction process described the expected behavior with regard to pitting for a particular metal-solution environment. The damaging effect of the chloride ion (as exists in body fluids) presented a more vulnerable anodic position for these reduction processes to operate and thus induce pitting. Various studies including linear polarization and anodic polarization methods showed the superiority of titanium to vitallium or stainless steel for pitting resistance in simulated body fluids. Crevice corrosion of Type 316L stainless steel in biomedical environments was the only reported study found for implant metals though it represents a major problem in this field due to the multitude of possible sites for its occurrance. The study of crevice attack on type 316L annealed and cold worked stainless steel in Ringers solution at pH 7.2 to 7.4 was conducted with varying oxygen and carbon 183. dioxide contents. The presence of high CO2 contents proved equally as important as 0£ in cathodic polarization as demonstrated by the prevention of expected differences in breakdown potentials of annealed and cold worked metal. The high CO2 condition also had the effect of lowering the passive current density of the annealed metal relative to the same condition at low CO2 contents. Crevices formed by the union of cold worked and annealed 316L stainless steel were shown to be active sites for crevice corrosion and the use of 316L in the annealed condition was recommended for all components of a multicomponent device. Nonbiomedical but related studies in chloride containing environments demonstrated a superior rating for titanium and its alloys as well as Hastelloy C in resistance to crevice corrosion. Surface abrasion and fretting corrosion were found to be the result of implant surface contact by surgical tools or a holding device used as part of the implant. The formation of corrosion products by an abraded surface is a function of repassivation current and film repair time. Repair times can be quite long for some implant alloys and although the metal dissolution process is low, the products maybe toxic. If the abraded surface lies in a crevice 184. formed by two parts of a implant device, the action of the crevice accelerates the scratch beyond repassivation and pitting will result. Metal wear characteristics become important in fretting corrosion and all current orthopedic alloys render problems in this respect. Intergranular corrosion is the result of improper metallurgical processing. In the case of type 316 stainless steel, it is a preferential attack that dissolves chromium depleted metal near the grain boundaries. The zone of chromium depletion is avoided with proper heat-treating or selection of the extra low carbon grades of this stainless alloy, ie., type 316L. These conditions are assured by appropriate A.S.T.M specifications which specify corrosion tests to insure proper heat-treatments and/or chemical compositions complying with the low carbon grade. The metal quality required cannot be accomplished by specifications alone--selection of reputable manufacturer is also necessary. Currently, it is encouraging to note the low incidence of this form of attack for type 316 stainless implants but the danger of its occurrance is also present and should be recognized. Intergranular corrosion was not a problem for applications of pure titanium but the carbide morphology of 185. cast vitallium was suspected as a path for this preferential attack. In conclusion, the subject of implant corrosion was reviewed from the standpoint of those metals currently used and biologically acceptable to the human body. The technical aspects of the mechanisms presented for the possible forms of corrosion were reluctantly limited, particularly with regard to electrochemistry: nevertheless from the above discussion it should be recognized that many of the corrosion forms have similar physical appearances and maybe combined so that evaluations require close attention to the particular surgical technique and applica tion as well as proper metallurgical interpretation. Finally, corrosion resistance is only one of the important propert ies of metals insuring a successful implant; the total metallurgical success is also dependent upon its mechanical properties and its response to combined corrosion-mechanical processes. 186. 3.0 Introduction to the Mechanical Properties of Implant Metals One criteria in the proper design of implant devices from a structural standpoint is knowledge of the important engineering properties of construction materials. Equally important is the knowledge of the required strength for proper functioning in the body. The following sections will define the more important engineering properties of metals by emphasizing the metallurgical mechanisms which constitute their origin. Also included is a comparison of the magnitude and meaning of these properties among present day implant metals. A brief description of the studies relating the classification of the types of loads to anatomical function will be used to introduce the recent analyses of various implant designs. 3.1 Relation of the Mechanical Properties of Metals to Implant Strength The conventional tensile test establishes most of the standard metal properties used in structural designs. This is a static test which measures the ability of the metal to withstand a steadily increasing tensile load. Representa tive results from the test are indicated by the minimum 187. allowed mechanical properties of implant alloys shown in Table VI. The yield strength approximates that a stress be yond which metal behavior is no longer elastic and permanent deformation will occur. For this reason, design safety factors are based on this property, eg., a safety factor of 5 (common for implant designs) for Grade 4 titanium would mean the device is never expected to be loaded more than approximately 14,000 psi. When metal is subjected to tension it elongates in the longitudinal direction while contracting in the transverse directions. The ratio of strain in the transverse direction to that in the longitudinal direction is called Poisson's ratio. The effect of deformation in both directions as measured by this ratio is important in design-calculations (see implant design studies); it is accepted to be 0.33 for most metals.1^4 In body locomotion a more realistic evaluation of implant metal yielding to deformation is called the dynamic 1°5 yield strength. This property is determined from a cyclic loaded testing system and represents the point where stable hysterysis is established. It is similar to the static system in that a constantly increasing load is used but now 188. the load is applied in both tension and compression as the metal specimen undergoes a 4 stage load cycle, i.e., 1) tensile load; 2 ) unload; 3) load in the opposite direction (compression); and 4) unload again. With subsequent cycles the load required to produce the same 1st cycle deformation is found to either increase or decrease until finally a stable system (or hysterysis loop) is reached where all subsequent cylces require the same stress (the dynamic yield point) to obtain the value of strain achieved in the 1st cycle. Some typical results for various metals are given in the literature.The significance of these results from a dynamic design standpoint is that the yield points are reduced from those determined statically particularly for cold worked material. Unfortunately no studies of this type have been performed and/or reported for the implant metals. Another useful design property related to the elastic range i.e., stresses and strain below the yield point, is called the modulus of elasticity. It also represents the ability of the material to withstand strain without permanent deformation. In this regard, the modulus value for stainless steel and vitallium is approximately twice that 189. of titanium (Table VII). The usefulness of this property is described later. Since the type of loading varies with application or a particular anatomical location, it is necessary to have knowledge of metal properties where loading is other then direct tension or compression. It can be shown that if twisting or torsional loads are applicable, the shear stress is one half the tensile yield stress when the load is applied uniaxially.1^7 Also the shear stress is a direct function of the applied torgue, the specimen length, the angle of twist and the polar moment of inertia, all measure- able or calculated quantities. Thus the applied torgue to a given size bone screw related to the shear strength attained by the holding device as well as the state of stress in the screw. Within the elastic range the shear stress in considered proportional to the shear strain and the constant of proportionality is called the modulus of rigidity 1 26 or modulus of elasticity in shear. This design modulus is also easily calculated from the shear stress parameters mentioned above. There are many instances in biomedical applications of metals where bending loads are important, e.g., hip or 190. knee prosthesis hip nails, intramedullary nails, Steinmann pins. The bending strength at a particular location is called the bending moment which is calculated simply by multiplying the load by the distance from that point to the point of load application (the distance is called the moment arm). The bending modulus is a function of the modulus of elasticity and the area moment of inertia (calcula ted from the geometrical design dimensions). For example, the bending moment at the distal end of a hip prosthesis with a 6 inch stem length maybe 2880 in lb for a 160 lb body weight (assuming the full load of 3 times body weight is applied at the top of the stem). The ultimate bending moment of a given structural item represents its maximum bending strength and it is directly related to the product of the area moment of inertia, I, and ultimate tensile strength, crujg, a°d indirectly to the distance, Ymax, from the neutral axis to the outermost fibers, i.e., (M ^t= 128 l / Y ma x ^UTS^* Examples of typical values for these parameters in implant alloys will be given presently. In conclusion, there is one important property for structural design where applications involve repeated or varying forces, namely, fatigue strength. The fatigue strength of a material is represented by the endurance limit: that stress at which the material will endure an infinite number or a given number of repeated loadings. The endurance limit maybe established under various types of loads such as those mentioned above, e.g., cantiliver (plane) bending, rotating beam, reversed axial stresser or torsion. Regardless of the type of loading, the preservation of surface condition is equally important to that observed for good corrosion resistance because virtually all fatigue failures are initiated at the surface. In addition, there are a number of other variables, e.g., stress concentration, corrosion, temperature, overload, metallurgical structure, residual stresses and combined stresses, which tend to 127 alter the endurance limit test results. Thus, the proper characterization of metal surface, type and frequency of loading and environmental conditions are all important for meaningful information of metal fatigue in a given application. As a relative measure of the fatigue strength of implant alloys and similar metals in air at room temperature it can be shown that their endurance limit represents only 35 to 40 per cent of the ultimate tensile strength (Table VII). It is also noted that these are estimated endurance limits, 192. apparently actual test results are either unfounded or not published. 3.2 Metallurgical Significance of Mechanical Properties The mechanical properties, previously described, are primarily dependent on the metallurgical condition of the metal. Excellent reviews of the metallurgical variables i o n i o/* involved are present in the literature but for the purpose of this review the influence of work hardening, grain size, and precitation hardening are deemed sufficient. Metal crystals are composed of a three dimensional array of atoms wherein the position of the atoms are fixed in stable equilibrium by the balance of attraction and repulsive forces associated with their nuclear charge. When the atoms are displaced from their positions, crystal imperfections or dislocations are produced. Dislocations may exist as a result of crystal growth during solidification or they are produced by applied stress as with metal forming to produce wrought parts. For example, the dislocation density of annealed polycrystalline metals maybe 10^ or 10^ lines/cm2 while the cold worked condition it is 1 0 ^ or 1 0^ lines/cm.66 -jhe distortion necessary to produce the large 193. density of dislocation (cold worked condition) requires applied stresses beyond the elastic range of the metal yield point), ie., plastic deformation. The increase in stress called work hardening is related to interactions between numerous dislocations. ^ 1 por example, the substantial difference between the annealed and cold worked strength properties of type 316L stainless steel (Table VII) is due in large part to the high work hardening capacity (numerous dislocation interactions) of this metal. When second phase particles, eg., carbides or alloy phases are present, the motion of dislocation is impeded and again increased stress is required for plastic deforma tion. This strengthening mechanism, called precipitation hardening, is manifested in the high strength properties of vitallium and titanium alloys. Grain refinement in metals is enhanced by cold working processes or controlling heat-treating conditions for the annealed state. The increased number of grain boundaries found in such fine grain metals act as additional barriers to dislocation motion and thus greater resistance to deforma tion. The size of grain as prescribed in A.S.T.M test methods results in a rating from 0 to 8 with the lower number 194. indicating a more coarse grain size.^^ Recognition of the strengthening attained by grain refinement has resulted in the use of the fine grained condition for most implant devices. In summary, the mechanical properties of metals that are important from a design standpoint are determined by static tensile or compressive, tests or dynamic (fatigue) tests. Some of the metallurgical mechanisms associated with the attainment of these properties are defined as well as the significance of metal-working and heat-treating operations. 3.3 Implant Design Studies The evaluation of various implant designs from an engineering standpoint is a necessary requirement for proper functioning during implantation. In performance studies of hip prostheses, Patrick; et al. subjected both intact (human) femurs and hip prostheses (Moore, Thompson and Eicher 137 type) to static and modified fatigue tests. The modified fatigue test on the combined femur-prothetic device or the device alone showed the Eicher design strongest due to larger bearing areas at the medial and proximal end. Static compression tests on femurs gave ultimate loads of 2030 lb 195. while that combined with prosthetic failed at 1030 lb. In another part of the same study, failures bj' impact on patella, femur and pelvis were recorded at loads of 1500 to 3800 pounds. Further comprehension of the magnitude and types of expected hip loads is found in other studies. For example, in Koch's early study of bone architecture, he calculated that the load to failure on the femoral head is 1820 pounds while the maximum strength developed by the action of all thigh muscles was l/7th that of the femoral shaft.In a more current study, shear failure occurred at the head of the femur at 1930 pounds in compression but failures of the femoral shaft in transverse bending was reported to appear at 1106 pounds. The effect of an implanted hip prosthesis made of stainless steel, vitallium, and titanium was to lower the compressive failure loads to 520, 1870 and 1720 pounds, respectively. It is also interesting to note that the success of fixation by other means, eg., hip nails, blade plates, plates with screws, depend on loads below that causing withdrawal or bending failure. This report shows maximum withdrawal and bending loads of 115 to 160 and 140 to 160 pounds, respectively, for individual screw fixation 196. in cortical bone. It would be helpful from an engineering standpoint to analyze the strength properties of bones under all the different types of load test parameters established for structural materials. One such study by Frankel, et al., (Table XIII) has shown the bending modulus of bone 5 to 10 times more rigid then the best intramedullary nail and 100 times better in torsion.The significance of implant design geometry in improving the observed strength properties was also emphasized. For example, a change in the working length* of the nail was shown to be important because the mathematical relationship describing stiffness in bending is inversely proportional to the square of this length and in torsion it is inversely proportional: furthermore the advantage of nail shape and placement was also demonstrated by similar calculations. Related studies on cloverleaf nails have shown the importance of nail diameter to bending 140 resistance as well as nail axis orientation (Table XIII). Various methods have been used to evaluate the strength *The working length is defined as that portion of the nail crossing the fracture site between the two area of contact in proximal and distal fracture fragments. 197. of bone plate and screw combinations. The most prominent study, fashioned after that used by the aircraft industry, determines the holding power of screws as measured by the torque used in tightening screw to bone. This torque was found to vary with age or bone location from 5 to 25 inch/pounds, but the clinical significance of this technique (prescribed for osteoporosis) was to determine the stripping torque required in removing a screw from a drilled hole near the fracture area, then utilize no more then 75 to 80 per cent of this stripping torque in applying the screws to hold the internal fixation device. As part of this same report the shape (design of threads) and size were qualitat ively evaluated by "pull-out" tests using bone plate counterpressure and it was concluded that the current design and small size available were most efficient. Quantitative analysis from a previous me tioned investigator disclosed that the average bone screw fails at 30 to 35 inch-lb of torque but that four inch-pounds of torque in a 20 thread per inch screw provides 811 psi compression which is deemed 139 sufficient for successful fixation. Very few studies have been reported on the dynamic testing of bone, implant device or their combination, yet 198. this is an important design feature. Recently Kraus presented an excellent review of the mechanical properties of human compact bone in which he reported (from other studies) the average life of tibial, femoral and fibular samples subjected to cyclic applications and removals of 5000 psi were 2.378, 1.188 and 2.841 million cycles respectively. The endurance limit for equal tensile and compressive loads on wet samples of femoral bone was given as 4000 psi. However, Kraus notes that the fatigue test of dead bone is of doubtful importance due to the self healing nature of bone in-vivo. A mathematical model to account for some of these vivo conditions, eg., aging, disease, self healing nature, was used to describe a proposed failure time for fatigue type loading but the magnitude of the variables involved requires further definition. Fatigue studies of the implant metals will be reserved until later and those concerning the actual devices seem to be neglected. In summary, the mechanical properties important to the design of any engineering structure have been described by reviewing more current studies that emphasize stress- strain analysis of bone, implant or their combination. The types of loads considered important in the body are 199. compression, torsional and bending loads. From studies of compression loads, vitallium provided superior strength performance in hip prosthesis in comparison to type 316 stainless steel or titanium, however all were inferior to bone compression strength. The results for intramedullary nails as well as bone screw (apparently of type 316 stainless steels though not always indicated) show the importance of using established engineering principles and mathematical formulation, e.g., bending, torque, to evaluate design geometries. The infonnation regarding fatigue strength of implanted structures and bone was given the same brevity as exists in the literature, however, the endurance of hip prosthesis alone or combined with bone was found to be far less than that of human compact bone. In- vivo or in-vitro measurements for this type loading on the implant metals are more aptly concerned with corrosion fatigue which will be discussed presently. 200. 4.0 Effect of Combined Processes on Implant Metals There are three combined processes considered important to this review, namely, stress corrosion, corrosion fatigue and corrosion wear. Stress corrosion is a failure phenomenon observed for metals that crack when stressed in a particular environment but remain intact when unstressed in the same environment. Cracks, can progress through the metal intergranularly or transgranularly at velocities varying from inches per minute to inches per year.^ The human body environment combined with the stress associated with implant application represent conditions highly susceptible to stress corrosion cracking yet its occurrence has not been reported. Consequently the vitro studies of stress corrosion among orthopedic alloys discussed in this review will be augmented by related studies which may benefit its future evaluation in biomedical applications. Likewise, the phenomenon of corrosion fatigue is likely to occur with metals in the body yet failure rates by this 129 combined process are reported to be less than one per cent. Corrosion fatigue maybe defined as a phenomenon where material damage caused by repetitive loads is aided by the corrosion process. Unless a structure is enclosed in a 201. complete vacuum, all materials, that undergo repeated loading are subject to corrosion fatigue. Thus the fatigue strength of metals is progressively reduced in going from vacuum to air to more corrosive environments. In particular, the fatigue strength of metals in the body maybe as low as 20 per cent of its ultimate tensile strength. The fatigue and/or corrosion fatigue process is usually described by its failure node, i.e., two stages of crack development, namely, initiation (1st stage) and propagation (2nd stage). By far the most extensive studies have concerned stage II (propagation) because of limited observation or measuring techniques inherent to initiation. Also this direction of attention stems from the established technology concerning the fracture toughness of structural materials which renders great utility to crack propagation studies. The adoption of these techniques to in-vivo or vitro implant metal studies is limited, so supplementary information from related fields must be offered for a more complete evaluation. Furthermore the influence of environ ment, metallurgical processing variables, surface condition and design geometry must be characterized. The combined process of corrosion-wear in implant applications is most manifest in device undergoing fretting 202. corrosion as noted by the high incidence of its occurrence among multicomponent devices as well as its influence on the corrosion fatigue process. Therefore, in this section, discussion of this dual process is limited to the subject of fretting corrosion or related to corrosion fatigue but a more detailed analysis of all mechanisms is presented elsewhere (see section on Wear). 4.1 Stress Corrosion Cracking (SCC) Since the conditions of tensile stress and corrosive environment exist in biomedical application of metals, the possible occurrence of stress corrosion cracking in implant metals also exists. This necessatates a knowledge of (SCC) mechanisms and how they are measured for at least qualitative implant material evaluations. While the information concerning implant metals is meager, technology in related fields is extensive. Stress corrosion cracking involves the dual nature of the stress and environment to accelerate concentrated metal dissolution via, a propagating crack, that ultimately causes failure. The nature of the stress is to help break down the protective oxide surface film and create points of high stress concentration both at the surface and metal 203. interior. The role of the environment is to preferentially attack areas of high surface energy by adsorption and diffusion processes. One theory reported by Uhlig^^ proposes that surface imperfections, (stacking faults--or aggregates of dislocations) caused by residual or applied tensile stress may speed the diffusion of impurity alloy atoms to the imperfection arrays and electrochemical dissolu tion occurs as a result of this localized segregation. This was called the electrochemical mechanism and it includes oxide film rupture and/or corrosion pitting type dissolution process at the crack tip to enhance crack growth. The other theory for the SCC mechanism by the same author is called "stress-sorption cracking" where adsorption of environ mental species causes further weaking of atom bonds which are already strained from applied or residual tensile stress es. The resulting reduction of surface energy causes pre ferential crack growth and failure. Environmental composition plays a major role in this mechanism and the presence of known innocuous anions, e.g., chloride, maybe expected to take part in the reaction by adsorption processes. The role of chloride or other dissolv ed elements in body fluids provides such a environment. For example, in a excellent study by Svare, et al., the passivation of nickel and copper in Ringer's solution containing organic substances (of average human body concentration) was determined while simultaneously measuring organic adsorption. The substances alanine and bovine plasma albumin had little effect on the rate of anodic dissolution of copper but the rate was decreased by two orders of magnitude in the presence of cystine (Figure 93). Cystine is a crystalline amino acid produced in the digestion of proteins.. Similar experiments with nickel show the same effect for the former organics but an opposite effect for cystine in that the rate of dissolution continu ally increases without evidence of passivation (Figure 94). From the observed potentials for cystine adsorption and copper oxide film formation the authors proposed that the improved passivation of copper could be due to a doping effect of cystine on the semiconductive properties of copper oxide. In the case of nickel, the charge donation by cystine could increase the electrical resistivity of the nickel carbonate thereby decreasing passivation capacity. Thus, the effect of dissolved substances in body fluids has a marked effect on passive metal film stability for nickel 205. and copper which not only lends credence to the role of adsorption in the stress sorption mechanism but also makes passive alloy systems in body environments highly suspect to stress corrosion cracking. Unfortuantely, no information was found regarding the effect of these organics (particularly cystine) on implant metals. Before going on to the more important studies evaluating a metal's susceptability to SCC, it is necessary to review one of the adopted measuring techniques which will be important also in the review of corrosion fatigue. The concept used to describe a material's resistance to fracture is based on linear elastic fracture mechanics and this analysis has defined a material property called fracture toughness. Fracture toughness of a material is described by a crack tip stress intensity parameter (K). By analyzing the observed behavior of an advancing crack, it has been found that the opening mode, called mode I, can be described mathematically by a series of equations where the stress in the vicinity of the crack tip is a function of the radial and angular displacement from the crack tip, (r,0), and the stress intensity factor, Kj. The Kj factors depend on loading, specimen or structure geometry and crack size, 206. but these factors have been worked out for a number of conditions and are available from the literature. A critical crack extension can then be defined by subjecting a precracked specimen to increasing loads and the value corresponding to the load at which unstable crack extension occurs is called Kj , the plane strain crack toughness.^^ Thus the Kj value is a material property that maybe used in design in a way similar to that described previously for yield strength, modulus of elasticity, etc. However, it must be emphasized that the parameter is valid only when there is constraint to plastic action at the crack tip. Nevertheless, many good estimates of the load to cause fracture under certain structural configurations have been accomplished using this parameter. Routine material evaluation for stress corrosion cracking can be accomplished by recording the time to failure of precracked spcimens in a corrosive environment subjected to different loads. By progressively reducing the load, a value of K-j- is reached below which no failure can occur under statically applied loads. This level of Kj has been designated Klg^Q and it is more often defined as that point where no failure is observed after some interval of time.^^ 207. Comparitive values of K-j- for materials exposed to SCC various environments are established in the literature 147 (e.g., Table XIV). Among the other types of tests used for evaluating a material's resistance to stress-corrosion cracking are the U-bend test and the conventional time to failure test. In the time to failure tests various devices can be used to apply a constant load that insures longitudinal tensile stress in the specimen. A simple arrangement consists of a wire specimen loaded by suspended dead weights. By enclosing the specimen in a cell of appropriate size to contain the environmetnal media, and electrochemical monitor ing devices, it is possible to determine potential and current time measurements that allow evaluation of the metal dissolution rate and passivation characteristics while undergoing conditions of stress corrosion cracking. Numerous tests of this kind have been performed to evaluate SCC resistance for various materials in a particular environment, e.g., (Table XV).149 The U-bend test is a constant strain or deformation test wherein a metal strip is bent in a U-shape so that there is a maximum tensile stress associate with outer fibers of 208. o c the closed end. This quantitative test is then complete by exposing it to the environment and examining the specimen surface for cracks. Thus the simplicity of this qualification procedure, is evident, yet the use for this purpose is difficult to find. One type of U-bend test has been reported for Type 316L stainless-Schneider pin stack (8 and 11mm) in a study to determine the proper stress 84- relief heat treatment of type 316L stainless steel. The metal specimens were tested in the cold worked condition as well as 4 different stages of stress relief. The tests consisted of exposing these specimens to five 48 hour cycles in boiling 42 per cent magnesium chloride. This system is known to be damaging to stainless steels by promoting the development of stress corrosion cracking. The only metal condition found susceptible was that of the cold rolled state in the 11mm specimens (Table XVI) . Another example where conditions for SCC or corrosion fatigue in biomedical applications exists is found in the treatment of scoliosis (curvature of the spine) by Type 316 stainless steel Harrington rod devices. In probably the first, vivo study of actually measuring human body forces, Nachemson, et al, used implanted wireless telemetry to measure post-operative distracting forces on the modified Harrington R o d . 151 The restraint placed on the rachet rods during and after operation is interpreted as being representative of a static tensile force. This distracting force acts on the spinal column so as to restore an upright position during healing (Figure 95). As healing progresses the restraint or distracting force is lessened and the rod distracts. The vivo test results taken from four patients established distracting forces of 20 to 40 kiloponds (44 to 88 pounds) during operation and decreased with time post-operatively until after 10 days the force was 1/3 the maximum found during operation. Confirmation of these results was found in a different study using strain gages instrumentation (Figure 96). 157 Here the forces during operation were 20-30 kiloponds (44-66 pounds) and static tests on the autopsy rod showed failure to occur in the Gruen hook at 55 ± 10 kiloponds (121±22 pounds). The authors believe the latter tests demonstrate adequate design strength for this application but this is questionable unless the vitro tests were conducted in the proper environment. The above studies then serve as vivo examples of static tensile loads in a human body environment. 210. In summary, the combined process of static tensile stress and corrosive environment defines the conditions for stress corrosion cracking in metals. Two of the mechanisms proposed for stress corrosion cracking were described on the basis of electrochemical processes and stress-sorption cracking. These mechanisms are necessarily complicated by the reactivity of the various constituents in the environment; this was demonstrated for nickel and copper by the role of dissolved organics in concentrations compatiable to that found in body fluids. The concept of fracture toughness, Kj , was adopted to represent a means c of evaluating design loads on structures in corrosive environments by defining a related stress intensity factor, Ki . Fracture toughness was shown to be a material uliL property which has found extensive use in relation to crack propagation in many metals (particularly those of high strength) in a corrosive environment. Further use of this property will be discussed in the section on corrosion fatigue. The treatment of scoliosis via stainless steel Harring ton Rods in an immobilized patient was considerd a good example for SCC conditions, though no failures by this 211. process have been reported. Force measurements on the rods ranged from a post-operative low of 22 pounds to a high of 66 pounds (average). A 60 pound load on the type of design mentioned would represent a 1200 psi tensile stress on the rod in the vivo environment. Simplified qualitative tests to measure material reistance to SCC are represented by description of constant load and constant strain tests. One of these, the U-bend tests on a type 316L stainless implant device in the cold rolled state exhibited slight cracking in conditions that promote stress corrosion cracking. Utilization of these or other tests to evaluate the susceptibility of implant alloys to stress corrosion cracking have not been found. 4.2 Corrosion Fatigue and Implant Alloys An in depth study of the combined process of corrosion and fatigue in material applications would involve an extensive number of considerations. Therefore this review will limit discussion to only those considerations which place the subject in proper perspective. The number of variables associated with metal fatigue has been alluded to in section on mechanical properties. Among these variables, the effect of corrosion type on fatigue crack initiation 212. and corrosion rate on the propagation of these cracks is of fundamental importance. By using electrochemical methods a minimum corrosion rate can be established for onset of corrosion-fatigue cracking. Furthermore the relationship of pitting, intergranular, and crevice cor osion to the \ initiation process can be demonstrated. The effect of surface condition on both the macro scopic and microscopic scale deserves emphasis because the former directly affects fatigue endurance and the latter influences the dynamic stability of the metal in a corrosive environment. Specific surface treatments, e.g., shot- peening, improve endurance by introducing surface compressive stresses that restrain crack growth at surface stress con centrations and notches. In some metals other treatments, e.g., mechanical polishing, remove surface roughness and produce a soft surface that is subsequently work hardened, thus providing better resistance to crack initiation. The mechanical breakdown of metal surface films in a corrosive atmosphere by dynamic loading results in accelerated corrosion and rapid fatigue crack initiation and propagation. Therefore the mechanical properties of films or diffusion layers deserves some consideration. 213. The relation of fracture toughness to crack progaga- tion previously shown to represent a current design parameter in the evaluation of materials in stress corrosion cracking is also applicable to corrosion fatigue. Many current studies utilize this parameter to examine fatigue crack velocity in various materials as a function of environment, cyclic frequency and corrosion rate. Therefore it seems appropriate to review this work as an introduction to the investigations involving corrosion fatigue of implant metals. The vivo failure of implant metals by fatigue has been reported to be less than one per cent of all premature 64 implant removals. However, accurate statistics of all implant failures due to any cause are difficult to find and no matter how few the number may be the theraputic complications of additional surgery for removal may defeat the original purpose of implantation. Inspite of this fact, there are very few studies evaluating the corrosion fatigue properties of implant metals in vivo or vitro environments. Therefore, this review will additionally include results of corrosion fatigue for current or perspect ive implant metals in related environment. 214. 4.2.1 The Role of Corrosion in Dynamic Loading Applications The fracture of metals subjected to repeated loads is associated with points of stress concentration usually located on its surface. Stress concentrations exist in design, e.g., sharp corners or fillets, bolt or screw holes, slots and at surface discontinuities. The disruption of the surface by corrosion processes create discontinuities that facilitate the formation and propagation of cracks under the action of repeated stress. Any of the five forms of corrosion previously described for implant applications can produce this effect and most notable among these are pitting or crevice attack and surface abrasion and/or fretting. If intergranular corrosion occurs, it too can lead to reduced endurance. Pitting and/or crevice attack appears at localized areas on the metal surface as small holes or depressions which have length to diameter ratios near unity. The damaging effect of such notches on metal fatigue, was first observed by McAdam in studies of low alloy and mild steel in three different aqueous environments. By perform ing fatigue tests on as machined specimens and previously corroded specimens as well as those corroding under stress, 215. he concluded the combined action of corrosion and stress always led to shorter endurance than that obtained for the other conditions. The most damaging environment in the combined tests was that of river water equal to 1/3 to 1/6 the salinity of seawater (note the similarity in chloride concentration to that of body fluids). Surface film breakdown and onset of pitting is greatly enhanced in solutions containing chloride. The corrosion fatigue cracks leading to failure were observed to emanate from these surface pits (Figure 97). The observed damaging effect of pitting on fatigue endurance may be related to the shape of the pit. Laird reports from other studies that endurance limits for steel were equal in either air or 370 NaCl + NaOH (ph 12) even though a few randomly distributed pits were present in the latter environment. 1^5 This result was attributed to the low stress intensity associated with the hemispherical pits produced in the NaCl solution. The author believes the damage occurs by crack initiation in the normal way, i.e., due to mechanical stress which is subsequently accelerated via, the dissolution action associated with pitting. Using microscopic examination of the surface of a 216. mild steel specimen shortly after immersing it in a 3% sodium chloride solution and subjecting it to dynamic loading, Uhlig, et al., discoverd mechanical surface damage in the absence of pitting. This damage took the form of surface disruptions called extrusions and intrusions which are commonly found in specimens subjected to repeated stresses without exposure to corrosive environments. Once intrusions (micro-cracks) were formed, elongated pits grew from these sites in a manner characteristic of dissolu tion by crevice or pitting mechanism. Thus for these conditions the crack preceded the pit.^^ Since corrosion is a time dependent process the above facts are not too surprising rather it is important to recognize the terminating effect on time or incubation periods when repeated stresses are present. Speckhardt used electrochemical methods (potential and current density time measurements) to show the incubation period for pitting can be shortened considerably by mechanical stress even in solutions that are not very corrosive. By controlling the potential in corrosion fatigue tests on stainless steel (18Cr, 8Ni) in 3% sodium chloride solution so that the passive state was maintained, no pitting corrosion occurred 111. n and a relative high endurance limit of 20.5 kp/mm (approx. 30,000 psi) was established. However when the potential was increased beyond the passive into the transpassive range, the specimen surface was rapidly attacked in the form of pits changing to cracks and low endurance was observed even at repeated stresses as low as ±4 kp/mm (approx. 5700 psi). Other workers have demonstrated that this benefitial effect of polarization on stainless alloys is not afforded to specimens exposed to more acid solutions ( in the passive range for the acid solutions also affects the incubation period for pit development. The decreased endurance observed in the presence of corrosive conditions can then be related to the severity of pitting attack for different metal-solution system. In fact Uhlig and coworkers suggest that a critical corrosion rate must exist below which the damaging effects of the environment are not observed.156 By anodically polarizing mild steel specimens to a fixed current density (corrosion rate) during simultaneous corrosion fatigue tests in deaerated 3% sodium chloride solution, a critical current density was established (2 A/cm^) below which maximum test 218. life was observed (Figure 98). Higher corrosion rates at the same stress level substantially reduce endurance. Although the critical corrosion rate corresponded to small metal loss (1.9x10”^ mg/cm^/cycle), its significance lies in the reduced initiation time of stage I crack development. For example, it was reported from other studies on low alloy steels that while increased applied stress reduced endurance life, there was little change in the critical corrosion rate for initiation of fatigue cracks. Thus the effect of increasing stress at low corrosion rates was to accelerate crack propagation. The critical corrosion rate has not been found to be the same for all metals, and this is probably due to difference in the stability of their surface films in the presence of electrochemical and/or mechanical action. 4.2.2 Influence of Surface Condition on Corrosion Fatigue The surface condition associated with a given metal part is dependent upon the nature of the surface response to final manufacturing processes and subsequent exposure to the environment. Since implant metals possess work hardening characteristics during processing and a stable oxide film, 219. it becomes important to look at the effect of corrosion fatigue on these properties. Generally speaking, the work-hardening of metal surfaces during processing e.g., grinding and mechanical polishing, results in increased resistance to repeated stresses by blocking the motion of dislocations. Under applied stress the motion of disloca tions is more easily accomodated along certain crystallo- graphic planes called slip planes. The slip associated with cyclic stress has been found to be even more preferential, occurring in so called "persistent slip bands". The surface appearance of wavey slip lines and microscopic observation of persistent slip bands represents the physical evidence of the fatigue mechanisms. Slip occurring in the persistent slip bands on the surface produces localized surface dis ruptions called extrusions and intrusions which represent crack nucleation. Once such cracks have nucleated, they first grow in the slip band in shear (stage I crack growth) and the grow normal to the principle tensile stress direction (stage II crack growth) . Processes which tend to retard slipping increase fatigue resistance and improve strength properties. For example, Speckhardt depicted a 6.57o improvement in endurance 220. to work hardening of a ground surface of 18-8 type stainless steel in comparing the air fatigue limit of ground surface to those electrapolished.Another example of this effect is found in Kurisumi’s test on type 305 stainless steel which showed a reduction in fatigue limit from 26.3 to 19.9 kg/mm^ in going from a machined surface condition to one electropolished. In addition, the electropolished condition was extremely notch sensitive to cyclic stress while the machined condition was non-sensitive to this effect.^0 The significance of their results in regard to implant metals can be recognized when it is recalled that implants are electropolished metals subject to pitting corrosion in the body and that pits represent stress concentration points as notches on metal surfaces subjected to cyclic loading. Improved resistance to fatigue crack initiation also depends on the stability of the surface oxide film. When the film remains intact under the action of a corrosive environment, resistances to corrosion fatigue is solely dependent on film strength. Grosskreutz explains an effective parameter in evaluating this intact film condition is the ratio of the shear modulus film to substrate, / , i.e., r = Gc/Gg where Gc and Gs are the shear modulus of the coating or film and substrate respectively. A favorable condition for retarding the action of surface slip is reported for values of 'P. i. For example, for aluminum the values of F* are 4.8 and 0.29 for vacuum and humid air respectively and microscopic surface examination of samples fatigued at low strain amplitudes in these two environments showed a drastic reduction of surface slip under vacuum 161 conditions in the early stages of fatigue. Therefore the evaluating of materials with intact films in the long-life region (low stress amplitude) of a given corrosion fatigue system should include the determination of the ratio , /~ . When higher stresses are applied, surface film fracture occurs when the strain amplitude in the substrate exceeds the fracture strain of the film. But plastic deformation or slip in the substrate during fatigue loading will accumulate at relatively low stresses and the emergying slip steps at the surface create strain concentrations which will eventually rupture the film and environmental protection is lost. In this regard thicker films offer some what more protection than thin ones but the main consideration in resisting crack initiation seems to depend on the adherance strength of the film and its ability to suppress plastic deformation. 222. Unfortunately, there is very little quantiative information given for evaluating corrosion fatigue crack initiation particularly for implant metals in-vivo environments. Perhaps further study using the above considerations will unlock the complications regarding the effect of corrosion and surface condition on the initiation process. 4.2.3 Fracture Mechanics and Corrosion Fatigue The complicated nature of crack initiation in metals subjected to cyclic stresses has advanced the study of crack growth or stage II cracking. Since the existance of surface stress concentrations in virtually all structural designs eliminates the need to develop preferred sites for crack nucleation, fatigue endurance thus becomes a direct function of crack propagation rate. Linear elastic fracture mechanics technology has been adopted to establish fracture instability in the presence of a crack. It was shown previously that the distribution of stresses in the vicinity of a crack tip depend only on the distance and angular displacement from the tip and the magnitude of stress at a given location is described by the stress intensity factor, Kj. The value of Kj where instability or rapid crack growth 223. to failure, occurs is designated KT . When alternating c stress is applied, Paris, et al., proposed the rate of fatigue crack growth is realted to an alternating crack tip stress intensity factor AK. The controlling influence of AK on crack growth, (da/dN) was clearly demonstrated in fatigue tests of low alloy steels (Figure 99). Similar conclusions for a wide variety of structural alloys led to the generalized expression, da/dN = CoAKn where n is the slope of log da/dN vs logAK relationship and Co is an empirical constant which depends upon material properties, frequency, mean load, and other secondary variables. Although this analysis is applicable only when constraint to plastic deformation at the crack tip exists, i.e., elastic behavior, the stress intensity factor AK, a material property, is directly related to crack growth. Often the influence of other variables, e.g., environment, cyclic rate, on fatigue endurance can be evaluated by their effect on crack propagation for a given AK range. Fatigue crack growth in an inert environment is most frequently observed to obey the above equation for high strength steels when the plastic zone at the crack tip is relatively small and in this instance the rate of crack 224. growth is dependent on the square or fourth power of the I stress intensity range (AK) . The corrosion fatigue crack growth response has been broadly classified by McEvily in terms of three general behavioral patterns (Figure 100).^“* The first type of material-environment system is characterized by a reduction in threshold for crack growth and increases in the rate of crack growth occur for given K levels.* As K values approach KT environmental c influences diminish due to reported mechanical chemical interactions or the rate limiting nature of transport processes. In the second type of behavior, the environmental effects are most significant above a threshold for stress corrosion cracking Kj and negligible below it. The bUij environmental effects are again negligible at values of K approaching Kxc * A typical system for this behavior is steel in a hydrogen environment. The third pattern exhibits behavior that falls between the first two types, with environmentally enhanced crack growth below threshold and accelerated growth above threshold. Many environment-material systems have been found to obey the third behavior. An excellent example of corrosion *The value of K is directly related to AK so no definition is lost by using K instead of AK. 225. fatigue studies in such a system was conducted by Barsom in his analysis of the effect of wave form of cyclic stressing on rates of crack growth below i-n maraging steels. ^ Maraging steel is an ultra high strength iron alloy composed of 12Ni 5Cr and 3Mo. The tests were conducted at various frequencies in 3 per cent sodium chloride solution under sinusoidal (zero mean stress or equal tension-compress- ion alternating stress) and various programed (unequal tension-compression) loads. The results show that environ mental effects in corrosion fatigue below Kj depend primarily on the time spent in periods where plastic deforma tion is occurring. For example, the crack growth at 6cpm under a square load pattern (tension-tension) was found equal to the growth rate per cycle at 600 cpm under sinusoidal loading. Furthermore environment enhanced crack growth below Kj occurred only during increasing tensile stresses. \jL Thus there is no environment effect on corrosion fatigue below under the compression side or the constant load portion of the loading program. The significance of environmental and frequency effects at a given value of cyclic stress intensity factor, AKj on implant alloys can be appreciated when recalling 226. the high strain-low cycle conditions in biomedical applications. Although these variables have not been extensively investigated for implant alloys, a more recent study of these effects on low alloy steels in high temperature aqueous environments is a case in point. The load controlled tension-tension test results showed a strong environmental influence on crack growth ar constant 164 AKj, when frequency is lowered below 20 cpm (Figure 101). The crack growth rate under these conditions may be as much as 100 times greater than that found in ambient air. In summary, the current trend in application of linear elastic fracture mechanics was adopted to study the rate of crack growth in metal-system environments. This technology has established the fundamental rate controlling parameter AK which shows a linear relation to crack growth, da/dN. The influence of some of the variables common to fatigue and/or corrosion fatigue conditions on crack growth has been demonstrated by the numerous studies of these parameters in the literature. 4.2.4 Studies of Corrosion Fatigue in Vitro or Vivo The previous discussion presented some of the more 227. important considerations necessary to evaluate the corrosion fatigue behavior of metal-environments systems. One of the earliest among these was the endurance limit or more importantly the corrosion fatigue endurance limit. Many authors do not ascribe much importance to this property, yet it is still used by designers in evaluating materials of construction. Except for isolated cases, e.g., Morral^^ reports an air endurance limit of cast vitallium in bending at 27000 psi, the establishment of this property for the orthopedic alloys in vivo-type environments seems unfounded or unpublished. In order to conduct significient corrosion fatigue tests in vitro the magnitude and type of loading, its wave form . and the cyclic frequency of the implant in vivo must be established. Since muscle and bone action varies with anatomical position the type of loading depends on implant location. In addition, in most all of the anatomical functions reported, there is a combined load present so that direct association with the available types of fatigue tests does not exist at present. The magnitude of loading also depends on degree of locomotion and body function at the site of implantation and the most often reported information in this regard is in 228. the hip location.73,129,150,166 Loads ranging from a minimum of 2.5 to 3 times body weight to a maximum of 6 times body weight are reported. It should also be noted that these loads are based on a condition of full weight bearing on one leg and their magnitude has been confirmed 73 in-vivo. Wheeler and James report from other studies that vivo measurements on a Austin Moore hip prosthesis showed the maximum at 4.33 times body weight when the subject was running, and 3.3 times body weight in level walking. The vivo evaluation of cyclic waveform and frequency has also been restricted to reports for the hip. Estimated frequencies vary from 3 to 51 cycles per minute (cpm) based on level walking but no estimates are available from body at rest or such conditions as ro ling or turning during bed rest.129 The importance of these frequency discrepancies can be realized when recalling the effect of frequency on corrosion fatigue behavior i.e., stage II crack growth. The frequency of 51cpm was reported from a study of an instrument ed human leg during normal level walking. This same investigation considered the cyclic waveform of vertical hip reactions, closely approximate a square pattern (Figure 229. 102).166 Various studies on other parts of the ai'-atomy applic able to implants devices have failed to establish the exact nature of the response to inflicted loads by the body. However, it appears in these cases that loads are axial (tension or compression) combined with either bending or torsion.167,168 Much more research is required via, vivo strain measurements, to establish frequency and cyclic waveform in different body locations. The importance of surface finish and corrosion rate on the corrosion fatigue properties of metals was previously demonstrated by several reviews from the literature. In this regard the significant considerations for evaluation of implant materials include, the composition and mechanical stability (strength) of oxide films versus that for diffusion coatings, shot-peening, or other surface treatments, the minimum corrosion rate below which infinite endurance is attained, and the determination of stage I (initiation) cracking mechanisms for different implant alloys. There have been no studies found in vivo to establish these parameters, in fact, very few reports concerning in vitro corrosion fatigue even characterize the surface of the test 230. specimens let alone microscopic and electrochemical examination of crack nucleation. Perhaps the best type of corrosion fatigue tests are those conducted on actual implant devices. For example, in vitro testing of pacemaker electrode wire of five different compositions was recently conducted by van Heeckeren et al., in a search for better electrode materials. ^ 9 The cold drawn and stress relieved wire specimen were of the following metal compositions Elgiloy (40Co, 20Cr, 15Ni, 7Mo, 2Mn, 15He, 15.8Fe), Paliney 7(35Pd, 30Ag, 14Cu, loPt, lOAu, lZn), Type 304 stainless steel (19Cr, 9Ni, Fe base), Pt 10 per cent Iridium and Pt 40 per cent Iridium. Rotating beam fatigue tests at a cyclic frequency of 360 rpm, and a stress equal to approximately 1/3 ultimate tensile strength were conducted on wire loops suspended in air, distilled water and 0 .97. sodium chloride solution (simulated body fluid). The endurance under corrosion fatigue conditions, calculated as per cent of air endurance, for the metals listed in the order given above was 11.1, 37.8, 0.2, 94.4* and 45.8 per *Although this material had the best corrosion fatigue behavior reported, it was tested at a much lower strength level than the other materials and thus was not considered "ideal" implant application. 231. cent respectively (Table XVII). It is interesting to note that this fatigue analysis properly conducted in a vivo type environment, disproved the previous favorable results found for Elgiloy in accelerated fatigue testing in a i r . 170 The accelerated tests subjected insulated electrode wire specimen i.e., teflon and siloastic ocated, to flexion, rotation and elongation motions which are more indicative 169 of cardiac conditions than those of van Heeckeren. Therefore based on these two studies it appears that selection of the electrode material that is best from a fatigue standpoint depends on whether the metal is exposed to body fluids or insulated from them and which loading system was correct. The vivo measurements of loads applied during surgery and post-operative on Harrington Rods have been shown to vary from 20 to 50 kiloponds (44 to 100 lb). Re ported fatigue failure of these rods recently prompted 168 additional intravital measurements and in vitro testing. By subjecting the critical failure region of the rod devices to bending fatigue (constant deflection) in Ringer's solution (37°C) at an undisclosed frequency, fatigue failure occurred at 50kp in 100,000 cycles. Some of the rods were shot 232. peened prior to testing but the amount of improvement was not considered significant (Figure 103). The author concludes that the design has adequate corrosion fatigue strength based on previously reported vivo load measurements but noted these measurements must not be indicative of true load variations. As a consequence, restriction on theraputic activity was recommended. Wheeler and James applied fracture mechanics to in vitro and air fatigue tests on type 316 stainless steel in 166 Ringer's solution at 98°F . The waveform and frequency, described earlier, simulated a vertical hip reaction. A definite increase in fatigue crack growth occurred for specimen tested in Ringer's solution (Figure 104). General corrosion was observed on the specimen surface and severe attack in the vicinity of the crack was caused by the corrosion products present. While no direct correlation with the results given in the latter report is possible from related studies (because of load variations, specimen and/or environment), it is fundamentally important to pursue the fracture mechanics approach to corrosion fatigue for evaluation of other possible implant materials. Piper et al., have studied fatigue 233. crack growth as a function of stress intensity factor below Klcrv-. on titanium alloys at a cyclic frequency of 120cpm 171 in 3.5% NaCl solution. Their findings for these conditions in terms of crack length versus cyclic life listed Ti-4Al-3Mo-lV (duplex annealed), Ti-6A1-4V (duplex annealed) and Ti-6A1-4V (in the mill annealed or beta stabilized, 1250 temper) in increasing order of crack growth per cycle, i.e., crack growth is faster in Ti-6A1-4V mill annealed- 1250 temper then the same alloy duplex annealed or Ti-4A1- 3Mo-lV. Spiedel used fracture mechanics to study the stress corrosion and corrosion fatigue properties of Ti-6A1-4V alloy in two different metallurgical conditions. Values of K^scc ^ ksi/in as compared to a value of 45vlksi in were reported for an alloy containing 0.183 oxygen while the respective value for 'Kj of 93 and 95^1 ksi in in another alloy of similar composition but with 0.116 oxygen. This shows the effect of oxygen content on the tendency for stress corrosion cracking. The effect of the higher oxygen content also extended the range over which SCC occurred. Spiedel also depicts the correlation between stress corrosion cracking and corrosion fatigue previously reported (see McEvily, Type C metal environment behavior) by dividing the plot of crack growth vs stress intensity into three regions (Figure 105). In region I (low AK values) a threshold stress intensity can be established at a point where the damaging effect of a mildly corrosive environ ment. Region II depicts true corrosion fatigue effects as represented by highly accelerated crack growth in comparison to inert environment. In region III, the high stress intensity range, the behavior may be dominated by time dependent stress corrosion plot A (Figure 105) (low frequency fatigue condition) or corrosion fatigue dominated, plot C, (where fatigue crack growth exceeds SCC crack growth in inert environments) or some behavior in between these two, plot B. In summary the above discussion has stated the concurrance or absence of some of the conditions considered important to the evaluation of corrosion fatigue resistance in-vivo or in-vitro tests. The effect of the corrosion form and its rate on the corrosion fatigue of metals has been reported. Likewise, proper characterization of metal surface condition was mentioned. Implant metal application involve high strain-low frequency conditions in various 235. waveforms dependent on the type of combined load present. One study showed that a square waveform properly simulates vertical hip reactions and the magnitude of these forces was reported to be approximately 3 to 5 times body weight. A suitable vitro environment used in implant corrosion fatigue studies was Ringer's solution. Corrosion fatigue tests have been reported for type 316 stainless steel and pacemaker electrode material. 4.2.5 Combined Processes - Corrosion - Wear The combined process of corrosion and wear has been discussed from the standpoint of its corrosion form commonly found among implant applications, namely, fretting corrosion. This insidious form of corrosion has been linked with fatigue failure for many years in a variety of fastner applications, e.g., rivets, bolted junctions, wheel hub and axle. The technology developed to access these failures should therefore be considered in evaluating biomedical materials used in any multicomponent implant where there is some motion between moving surfaces. The amplitude of relative motion need not be large, in fact, damage can result from slip of as little as 10“^.^72 Furthermore the contact surfaces do not have to be both metal, e.g., bone rubbing 236. on metal, so that susceptibility of fretting damage leading to corrosion fatigue of relevant implant devices is always imminent. A most common test used in evaluating the effect of fretting on a material's fatigue resistance is the same as that associated with fatigue testing to establish an endurance limit. In this case clamping devices apply pressure to the surface of specimens during the fatigue test. The test results show that the endurance limit is substantially reduced from that observed without fretting (Figure 106). The fretting fatigue limit is defined as the critical number of cycles to initiate a crack and it is found to be a function of the material, applied stress, and cyclic frequency. For a given material a higher stress and/or lower frequency reduces the fretting fatigue limit. Waterhouse has reported tests on variety of materials to establish the strength reduction factors and fretting fatigue limit. In this case greater reductions in strength or lower fretting fatigue limits were observed when the contact- i to ing surfaces were of different hardness (Table XVIII). As part of this same study two stage experiments were conducted in which the specimen surface is fretted first 237. and then fatigued (stage I), and tests where fretting and fatigue occur simultaneously (stage II). Stage I test results show strength reduction resembling that associated with corrosion fatigue but the co-joint action of fretting and fatigue in stage II tests gave by far the most damaging results. All of these results were from tests conducted in air. When an aggressive environment has access to the fretted area of less corrosion resistant materials the corrosion process dominates the nucleation of fatigue cracks. For example, preliminary results for fretting tests on aluminum in sodium chloride solutions show the reduction in fatigue strength follows the same behavior as that of 173 corrosion fatigue without fretting conditions. ' This is not to imply any benefit because in this case the corrosion fatigue strength is much less then that obtained for the fretted condition in air. Apparently the fretting action aides the corrosion process so that fatigue failure is more rapid. It is interesting to note Cohen's results in this regard. In a study that subjected standard bone plate and screw combinations into a 10 pound cyclic (50cpm) load, he determin ed the corrosion rate by weight loss measurements. The combination of screw and plate assemblages with two somewhat different compositions of type 316 stainless steel showed much less weight loss than when either was coupled with vitallium as their counterpart. This was explained by the abrasive action of vitallium, i.e., greater fretting damage, on the stainless steel surfaces. Control tests conducted statically (no load) on similar combination of bone plate and screws for periods of 8 times the duration of cyclic tests showed negligible weight loss. Thus the accelerated corrosion rate in the presence of fretting due to cyclic loading is clearly demonstrated. The combined process of corrosion and wear can also lead to clinical failure of the implant devices. In a very recent metallurgical study of a Smith-Peterson hip nail, removed due to pain and tenderness near the incision area, crevice attack in the form of pitting was observed i 7c near abraded regions of plate-screw junctions. ' Crevice attack occurring in shielded areas between bone and plate or plate and screw is greatly enhanced in the presence of fretting because the sliding contact causes breakdown of the protective oxide film. The increased metal dissolution, 239. apparently produces a greater requirement on body tolerance to the type of corrosion products generated. In this case the excessive vitallium corrosion products caused localized tissue reactions and implant removal. In summary the effect of fretting on the fatigue behavior of metals has revealed a marked reduction in associated endurance limits. Tests used to establish this effect in corrosive environments demonstrated that when the agents were aggressive to the metal, the fatigue damage is dominated by the corrosion process and the endurance resulting is less than that obtained fro fretting alone. The generation of fretting by cyclic loading both in vitro and in vivo has been shown to increase the corrosion rate of rate of two current orthopedic alloys, type 316 stainless steel and vitallium. The severity of such damage in the latter case has led to the clinical failure of a internal fixation device. 240. 5.0 Statistics of Implant Metal Failures The purpose of this section is to provide an insight as to the degree of involvement all the factors previously mentioned in Section 1-4. By combining all available reports on implant failure statistics and their causes , the predominant damage effects to each type of implant metal is its respective application may be depicted. The results are shown in Table XIX. There are several points persuant to interpretation of this compilation. First, since most of the reports reviewed concenned only orthopedic implant metals, the coverage is restricted to stainless steel (mostly types 316 and 316L), vitallium (both cast and wrought) and titanium. Next, the data includes performances or quality inspections as well as implant failures so that attention is directed to the column B of each item before interpretation. Third, the use of parentheses in column C denotes numbers of a particular type of device in the survey while the same symbol in column E and H represents percentage of a particular type metal device damaged and fractured respectively. Last, where necessary some interpre tation was used in classifying the metallurgical causes given in these reports, e.g., if extreme wear was listed for 241. a screw and plate device, it was classified as fretting. In considering type metal (column D) and the associated damage reported (column E), it appears all are subject to corrosion in any one or more of the five corrosion forms mentioned in Section 2. Titanium is the most corrosion resistant but lacks strength (item 9 and 11) while stainless steel and vitallium appear particularly susceptable to crevice and fretting attack. As with all engineering structural applications, design is critical for successful load bearing applications and the corrosion occurring at stress concentration subsequent to failure was a common cause of implant damage (column F). Another cause was the corrosion susceptability of the metal itself but based on the previous review (Section 2) the actual results of metal performance in the body should not be surprising. Although surgical technique was a major cause of removal in these and other studies, these results are intended to emphasize metallurgical factors. Finally, a very surprising cause of implant damage, frequently demonstrated in these surveys (item 1, 3, 4, 10, 11, column F), was the poor quality obtained from the manufacturers. The incidence of pitting (due to lack of control in electropolishing 242. technique), surface abrasion, improper heat treatment and even off-analysis metal can be linked to enhanced corrosion damage. The incidence of implant fracture was low in comparison to that of corrosion damage and the most common fracture mode was fatigue. Here stainless steel out performed vitallium or titanium, e.g., item 9, column H. The detrimental influence of stress concentrations in design or defects in manufactured surface conditions were listed as initiating mechanisms with corrosion as an accelerating factor. In conclusion, the available information concerning implant performance in the human body has shown that the predominant metal damage is corrosion which may be combined with wear and/or fatigue in load hearing applications. Stainless steel (type 316 or 316L) proved least corrosion resistant with pitting and crevice forms most prevalent. The most common forms of corrosion found for vitallium were fretting and crevice attack. Though titanium was also subject to fretting damage, it was by far the most corrosion resistant. Based on the total number of implants considered in this survey, all metals performed well with respect to 243. failure by fracture. However corrosion fatigue seems most susceptible due to design complexity of short terms devices or metal strength in long term implai?ts. Stainless steel proved the better material in this regard but information is meager. 5.1 Technological Improvements for Implant Metals This report has reviewed the current literature concerned with the placement arid performance of metals in the body. The discussion dealt primarily with the more recent metallurgical investigations concerned with the response of metals to human body environments and it had avoided the equally important but opposite effect i.e., bone and tissue reactions to metal. As an appropriate conclusion to this report therefore, it is noteworthy to mention some technological improvements applicable to the established problems areas. Of course, any technical proposal must be accompanied by medical approval. With regard to corrosion it appears that a substantial improvement could be attained by comparing various materials to in vitro tests that simulate conditions where crevice, fretting and pitting attack are likely to occur. The crevice test conditions for the respective materials should 244. duplicate both small (as with bone plates and screw) and large (as between bone and plate) shielded areas and incorporate the previously mentioned electrochemicr1 measure ments and microscopic evaluations. Also by testing within the expected pH range and temperature associated with vivo measurements, it is possible to evaluate the resistance of the same materials to pitting attack. Titanium and its alloys appear to be most resistant from the information provided in this literature review but additional qualifica tion is necessary and other materials of proven stability under these conditions in other fields have yet to be tested. The evaluation of material resistance to fretting could involve testing apparatus used in identifying fretting endurance limits for more rapid primary qualification and followed up with tests on the most resistant of these 1 “71 bio-acceptable metals similar to those conducted by Cohen. Furthermore much more developmental work in fastener design could be a accomplished to overcome the fretting problems 5 6 as evidenced by the recent Swiss investigations. All of these methods for evaluating corrosion resistance could be conducted in conjunction with appropriate vivo tests in animals so that information regarding metal interaction in 245. bone resorption and tissue reaction processes can be more properly identified. Considering the performance of implant metals in load bearing applications, it appears necessary to establish metal strength in body flu id environments under dynamic loading conditions. Although the magnitude and type of combined loads related to body-implant functioning is not accurately identified at the present time, standard fatigue tests modified to include the corrosive environment could be used to obtain comparitive corrosion fatigue resistance of the present orthopedic alloys as w ell as other prospective m aterials. By u tilizin g previously mentioned electrochemical and m etallurgical techniques in these tests, the initia tio n v and propagation of metal fatigue damage can be studies in an effort to find improved m aterial performance. Next combined loading techniques could be conducted on the best of the above tested materials and axial loading combined with torsion appeared to best in this regard. In conclusion, a m etallurgical review of the literature concerning metals used for im plantation has demonstrated some of the important considerations required in evaluating materials for implant performance. One of these metal properties namely, corrosion fatigue behavior is currently under investigation and prelim inary results are discussed in a subsequent section of the report. 247. 248. Table V I A Classification of Metals to Implant Locations and Functions Anatomy Location Implant Name Implant Type M e ta l No. Name F u n c tio n 1 S k u ll Skull Plates, C ra n io - Vitallium , Screws, Wire p la s t y T ita n iu m , Mesh T a n ta lu m Type 316 SS 1 A r t e r i e s C lip s T r e a t G o ld , m ent o f S i l v e r A neurysm M a n d ib le Mandibular R e sco n - V i t a l l i u m Prosthesis, s t r u c t i v e Bone Plate Appliances and W ire f o r th e Mesh j aw 2 V e r te b r a H a r r in g to n T r e a t - Type 316 SS Rods m ent o f S c o l io s i s S p in a l S p in a l Type 316 Plates and F u s io n SS W ire s 3 Clavicle Clavicular F i x a t i o n Type 316 N a i l s , o f c l a SS Screws and v i c l e d i s P in s l o c a t i o n o r f r a c t u r e 4 Scapula Carpal R e d u c tio n Type 316 S ca p h o id o f f r a c SS Screws t u r e s o f 1 th e c a r i p a l s c a p h o id bone 249. Anatomy Location Implant Name Implant Type Metal No Name F u n c tio n 5 C hest Pacem aker H e a r t- T ita n iu m , a s s is t P la tin u m - d e v ic e s I r i d iu m , N ic k e l A l l o y s , E l g i l o y 6 S h o u ld e r S h o u ld e r P ro x im a l V i t a l l i u m P r o s th e s is H um eral R e p la c e m ent Jewett Nail F i x a t i o n Type 316 and P la te o f P r o x i SS and S ta p le s m al End V i t a l l i u m o f Humerus 7 Humerus Stevens-Street R escon- T i-6 A 1 -4 V Elbow Pros s t r u c t i o n t h e s is Device for th e E lbow Mechanical R e p la c e V i t a l l i u m Elbow Joint ment P ro s t h e s is K u n ts c h e r H um eral Type 316 Humerus N ail F r a c tu r e SS F i x a t i o n 8 U ln a Proximal or R e p la ce - V i t a l l i u m Distal Ulna m ent o f P r o s th e s is P ro x im a l and D istal Ends o f U ln a Kuntscher V- F r a c tu r e Type 316 type, Vesely- F i x a t i o n SS Street Split o f U ln a ty p e , and Schneider Self- B ro a c h in g Intram edullary N a ils 250. Anatomy Location Implant Name Implant Type Metal N o . • Name Function 9 R a d iu s Radius Cap or H e lp Re V i t a l l i u m Head P r o s s to r e t h e s is F u n c tio n o f E lbow J o in t Y Plates and R a d iu s Type 316 Kuntscher V- F r a c tu r e SS type Radius F i x a t i o n N a il 10 Hand and Flatt Finger R e s to r a V i t a l l i u m F in g e rs and Thumb t i o n o f P r o s th e s is F in g e r J o in t F u n c tio n Finger Bone F i x a t i o n Type 316 Plates and o f S m a ll SS Screws Bone F r a c tu r e s 11 H ip H ip P ro s R e p la ce - V ita l1ium, thesis Several m ent o f T ita n iu m Types F e m o ra l Head Hip Nails, H ip Type 316 Pins, Plates, F r a c tu r e SS Staples, and F ix a t io n Screws 12 Femur Femoral Blade F i x a t io n Type 316 Plates and o f S in g le SS Screw Combina and M u l t i tions Bone p le F r a c P la te and tu r e s o f Screw or Wire th e Femur D e v ic e s Intram edullary Nail-Several Types 251. Anatomy Location Implant Name Implant Type Metal N o. Name Function 13 Knee Knee R e p la c e V i t a l l i u m Prosthesis - m ent o f Several Types D is e a s e d ( e . g . , arthritic) Knees 14 T i b ia Tibia Pros R e p la c e thesis-Several m ent o f Types T i b ia V i t a l l i u m S h e lf Townley Tibia R e p la c e Type 316 Plateau Plate m ent SS and Screws a r t h r o p l a s t y o f th e T i b ia S h e lf Tibia Bolt F r a c tu r e Type 316 and T i b ia F i x a t i o n SS (Intramedul- o f th e l a r y ) N a il T i b ia S h a ft 15 F ib u la K u n ts c h e r F r a c tu r e Type 316 O le c ra n u r F i x a t i o n SS Fibula Flap o f th e N a il F ib u la 16 T a r s a l Small Bone F r a c tu r e Type 316 and Plates and F i x a t i o n SS M e t a t a r Screw s o f S m a ll s a l Bones Table VII A.S.T.M. CHEMICAL SPECIFICATION FOR IMPLANT METALS TYPE METAL MAJOR CHEMICAL CONSTITUENT (WT.%) C* ______Cr______Ni Mo ^e co . — W Ti i Stainless Type A 0.08 17.0-20.0- 10.0-14.0 2.0-4.0 Balance Steel Type B! 0.03 17.0-20.0: 10.0-14.0 2.0-4.0 Balance ; Cast Co-Cr Alloy 0.35 27.0-30.0: 5.0-7.0 2.5 Max. Balance - - Wrought Co-Cr Allcv .05-.15 19.0-21.0i 9.0-11.0 - I 3.0 Max. Balance 14.0-16.0 i Titanium .10 ; 0.30 Max. - - Balance 1 i *>faximm oontent unless otherwise specified. MECHANICAL PROPERTIES OF IMPLANT METALS AND BONE Type Material Condition Ultimate Str. Yield Str. Elon Endur. Lim* Modulus of (PSI) (PSI) (%) (PSI) Elasticity (16b PSI) Type A or B Stainless Steel Annealed 75.000 ! 30,000 38 35,000 28 Cold Worked 125.000 • 100,000 12 Cast Co-Cr Alloy - 95,000 65,000 8 38,000 36 wrought Co-Cr Alloy Annealed 130,000 55,000 - 60,000 35 Titaniixn Grade 3 60,000 50.000 18 16 Grade 4 80,000 70.000 15 35,000 16 252. Bone Ctxpact 12 to 20,000 • - - 1 to 3 ♦Typical values determined in air fatigue tests. TABLE VIII Standard E&F Series of Ketals* Ketal-metal ion Electrode ootential equilibrium vs. normal hydrogen (unit activity) electrode at 25°C, volts T Au-Au+3 +1.498 1 Pt-Pt+2 +1.2 Noble or Fd-Pd+2 +0.987 cat'aodic Ag-Ag+ +0.799 Hg-Hg2+2 +0.788 Cu-Cu+2 +0.337 h 2-h + 0*000 Pb-Pb+2 -0.126 Sn-Sn+2 -0.136 Ni=Ni+2 -0.250 Co-Co+2 -0.277 Cd-Cd+2 -0.403 Fe-Fe+2 -0.440 Cr-Cr+3 -0.744 Zn-Zn+2 -0.763 Active or A1-A1+3 -1.662 anodic Kg-i:g+2 -2.363 1 Ra-Na+ -2.714 1 K-K+ -2.925 * itenroduced from, CCtLcCSLN E.NGIIiE-.HIIJG, page 31. by K.G .Fontana & H.D. Greene TABLE X X Galvanic Series of Some Commercial Metals and Alloys in Seawater Platinum Gold Noble or Graphi te cathodic Titanium Silver | Chlorimet 3 (62 Ni, 18 Cr, 18 Mo) Hastelloy C (62 Ri, 17 Cr, 15 Mo) 13-e M q st inless steel (passive) lb-8 stainless steel(passive) . Chromium stainless steel 11-30# Cr (passive) T Inconel (passive) (60 Ni, 13 Cr, 7 Fe) L Nickel (passive) Silver solder Konel (70 Ni, 30 Cu) Cupronickels (60-D0 Cu, 40-10 Ni) Bronzes (Cu-Sn) Copner Brasses (Cu-in) " Chlorimet 2 (66Ni, 32 ’’O, IPe) Hastelloy B (60 Ri, 30 No, 6 Pe, 1 N.n) ' Inconel (active) Nickel (active) Tin Lead Lead-tin solders p 18-6 Mo stainless steel (active) L 16— 8 stainless (active) Ni-Resist (high Ni cast iron) Chromium stainless ste-1, 13# Cr (active) I* Cast iron I. Steel or iron 2024 aluminum (4.5 Cu, 1.5t.g» O.fc Mr.) Active or Cadmium anodic Commerically pure aluminum (1100) Zinc Magnesium and magnesium alloys 254 rp « t : i? V Anodic Back SI.? (AES) o f P e ta ls in Equina Serum I.etal Identification P o te n tia l • (m illivolts) T ita n iu m 3,500 Tantalum 1,650 P la tinu m 1,450 P alladium 1,350 Rhodium 1,150 Ir id iu m 1,150 Gold 1,000 Cr-Si-io Al.loy Vinertia "C" 8 dO C r-N i-K o A llo y L a n a llo y "5R" 875 Cr-Co-Ko Alloy Virillium 750 Cr-Co-LIo A lloy V ita lliu m 650 lb-8S 8o( 31fc) 480 Ni-Cp-Fe Alloy In c o n e l 350 Zirconium 320 Tyoe 304 300 S ilv e r 110 Vanadium -70 C ob alt -350 255. Aluminium -600 256. TABLE XI Predicted Electrolytio Crovioe Corrosion Heniutunce* (A Guideline Only) (H2 outurated, N »2304 , 25 C) Representative i Eol'eronoe Metal t / l i /cm2 ) 0 (for i„ value) T1-6A1-4V 2 . 122 Hhstelloy C 9 123 304-L Stainless Steel 35 123 31fc Stainless Steel 45 123 304 Stainless Steel 75 40 20Cr-80Ni 300 121 Chromium 4,000 121 Nickel 12,000 121 Fe-lbCr 25,000 124 Iron 120,000 127 Fe-15Co 200,000 12b ’'Ranked in the order of decreasing resist ince. *#ic= critic al i * from France's derived eq(l). L p=^> E W r (ietive-pu38ive crevices) (paaoive crevicoe) where Lp= the pa g rive length, A E p= passive potential range, W = crevice width,/* = solution resistivity, ic= critical anodio current density, ipap-e lve current density, and^Ea= active potential range. TAP Lei XII Relative Crevice Corrosion Resistance of Ketals and Alloys in Sea i/ater* Aetal or Alloy Resist- nee Hastelloy C Titanium Inert 90 Cu-lORi-1.5 Re 70 Cu-30Ri-0.5 Fe Bronze Best Brass Austenitic Nickel Cast Iron Cast Iron Neutral Carbon Steel Incoloy 825 Carpenter 20 Ni-Cu Alloy Less Conner 31c Stainless Steel Ni-Cr Alloys Bit Initiaion 304 Stainless Steel At Crevices Series 400 Stainless Steels 7 * .’T * r *B;, .7. D. France, Jr., CRBVTJB JCii.XSlvI. u. i. i>. L - i. Lt'J 1 257. an invitee review pa oer for nresert •. ticn -t toe 3yn.corium on Localized Corrosion, nJl'I. Annual 7eetinr» Atl nooic City. June 2 9, 191. T.iELJi XIII Comoarison of Strength Properties of Bone to Intrnnedullary Fixation Devices. Type Katerial Bending Modulus Torsional Modulus (/ - ■ * -jf> (Kq/LIv 2 X 104) (K(j/r.T2 X io3 ) B o n e ...... 210,150,170 Sagittal Plane (Ely). . .' . 24,22,17 Coronl Plane (EIx) .... 31,16,26 Intramedullary Kail CAS Cloverleaf...... 2.o4, 2.64, 2.11 . . . 1.6 Diamond ...... 1.91* ••••» 1.16 . . . 11.6 f c ? & S c h n e i d e r ...... 1.93, ...., 1.93 . . . 14.6 G refers to the shear modulus while J is the polar moment of inertia. xA. I.E XIV Results of Visual Examination :na .'.eight r-.eusurements - or 3tress-Corrosion Scecimens. Visual .’.'eight Gain or Lo. Soecimen Tre tmer.t 8 ILT. 11 HJT: O TTiT. 11 cur. Snecinens 3seCi~ ens or.ecir.er.s See ciner.s i Cold rolled no crocking slight cracair.g nil nil 700P stress relief no cr eking no cracking nil nil 75-? stress relief no creeping no cr eking nil nil SQOF stress relief no cracking no cracking r.il r.il 100? stress relief no cr eking no crackin; nil nil 259 * TA F'Lii XV Exam ale of Tent Konultc in Measuring Material Resist- -nee To 3.C.C. Alloy Nor; inal Compost i on Range of Time to Description of Designation (Fe base) Breaking (MIN) o.C.C.Resist nee T1976 lUCr lONi O.IBe 1645-104,505 ' Good Tl 980 lOCr lONi l.ODe 7315-13 3,095 Good Tl 9 JO 12.5Cr'l5Ni 1.0A1 5 360-27,075 Poor Tl 392 12.5Cr 15Hi O.IBe 6250-58,275 Good Tl J 31 12.5Cr 150i 1.0A1 0 . 1 C 27,075 Poor Tl 398 12.5Cr lONi O.IBe 3065-23,275 Poor T 2002 15Cr 15Ni 1.25Be 1.0C 86500-7560,000 Best Tent coruiscte on wire specimens, constant lo-.d= JO '/* o f y i e l o str-;,-rt!i ir hoi Ling h p C l at 1 t>4G°. * • Froir; Otnehle etui. "Effect of Alio;/ Coir, nos it i on or> otress Corrosion ro Crue king of’ Fe-Cr-Ni P.n.-.c Alloys". Corrosion, o Vol. 26, No. 11. November 1 370, p 451. TABLE XVI Effect of Alloy Composition on the Resistance of Metals to Stress Corrosion Cracking as Measured by Fracture Toughness Kjx, Kj3qc Alloy Composition Environment Approximate Tyoe Variation KIx KISCC & Condition or Heat Treatment (KsiTcvv) Mara;ing Steel 195ksi (y.s.) 0.3C Salt water (3.5# NAC1) 115 20 Kn-3teel 199ksi (y.s.) 3/6 Mn " 40 10 31> C Steel 199ksi (y.s) 31* Cr 30 20 yfo C Steel 195ksi (y.s.) Ni 115 40 Martensitic S3 Q 14 Cr 141 Co 5l.o 700 F Temper. 160 50 Kartensitic SS 14 Cr 141 Co 5Ks 900°F Temcer. 70 35 4340(172ksi,y.s.) 0.15C 111 73 4340(195ksi,y.s.) 0.28C 111 58 4340(225ksi»y.s.) 0.39C " 06 .8 For correct alloy composition see reference.sited. T allis x v i i Results of Fatigue and Stress-Corrosion Fatigue Tests Alloy ■ I Stainless 10Jo Platinum- bC> FI a tin UE- Test Faliney-7 Elgiloy Steel-304 10;,*- Iridium 40^ Iridium Stress (psi)* 100,000 100,000 100,000 50,000 1 7 0 , 0 0 0 Fatigue test (cycles)f In air 704.000 543.000 27,532,000 976,000 4 7 4 . 0 0 0 In 0.3,* KaCl 265.000 72,000 56,000 918,000 2 1 7 . 0 0 0 In distilled water 311,000 1 7 7 . 0 0 0 1 0 6 , 0 0 0 663,000 290.000 Stress tests for platinum-iridium alloys could not ie run at 10u,0dG psi» For JC$ nlntinum- 10;° iridium, this v»as above toe yield strength •■na c u:ed t nglim- when the c hue Sc was rotated. For t>0/° platinum-4Q>& iridium, this was fcelow Lhe endurance limit, and specimens la tea almost in definitely. tAll values are the median of five determination of cycles to failure. TA2LI XVIII ' Strength. Seduction Factors and Fratting Fatigue Licit. Strength Fr e tting Katerial Hardness reduction Katerial Hardness fatimue* factor limit VHK 1 0 0 0 cycles 0 . 1 C steel 137 1.41 Aluminium 31 12 5 0 0 . 3 3 C steel 165 1.48 Cooper 66 325 0 . 4 C steel 420 2.00 Silver 76 115 0 .7C steal (normalised) 270 2.08 A1 4Cu 142 1 0 0 0 .7C steel (cold dra:.vn) 365 3.1- A 1 6Zn 3K g 182 55 0.6Cr 2.5Ni 0.5*o steel 330 4.38 Cu 302n 140 40 1.4Cr 4.GXi 0 .3- 0 steel 510 3.5o C u 1CA1 283 80 A1 4Cu 117 1.59 A1 4.4Cu 0. 5Kn 1.5Lg 140 1.92 A1 4.4Cu 0. 6:.n O.TKg 160 2.72 Cu 302n 175 1.50 • t o O' * OJ Fretting fatigue lir.it of 0.10C steel in cont ct with the indicated Materials. - TAPI T. A 1 A # ST\TlST|f\l. IV.MISfUW OK *1TAI, Y* ’^O Kl/.VK IN P«TIA*IT A rm KVTJ CMS c l> k. 1 u R I t — A 1 Tvpe .« d PcsrrlpKuo of R e fe re n c e ■ n tfh g r h ir p o w I n p le a t Type • (?) Affected f i l e s r ^ d o a ln a n t f V t . i l (I) fr.iclure F r a c tu r e o f o l •rsctlptltn IW t s l C i use ( n o . ) 6 Type related D—age o f h .ia .ir r f n i t l J t t i m l ^ l M t U w stljiitin) . r * P .im ^ e __ 7) lcn *llr Brut In aut?ery C r c e o e .H 1 1 1 rtv*I M ere gone K-'ue L r v t i u l , t t B l Z 2 11 C l i n i c a l Rip P nalhrtlt II (15) Vne Ra. to J y lm N y l on * 7pendagc r o * u i t j r i’ndef io rJ C o h e n . 4-3 I S 1 1301 A l l Types A l l (1 0 0 ) C o s t .6 frsae. MoVal t r e v fe e 4) tsideffned Types S u rg . Te h - P i t t I n y MIC. ( . t t v i n l c Cerroslc-. in it C a h o c * . iI 4 4 17 H eullurilcil (61 Kip VjlI-PU tC 316L ( 10) i9rr«i free. D e s ig n C r v v lr e 30) Corr. Fall trc law itlg. ef (9) Undefined C a st Possible S.C.C. M r.:. T i t l i n g (.) Possible S.C.v iated by poor VII. p ro c e s s P o t o M t r M » t. o r None liste d a aex c a ie w .4 0 5 91 !a p )jo C Single icre« H i I d S. (41) Corrosion H u ta l F r e t t in g performance loo* plate 6 s c re w ST. 3 1 6 . d e s ig n Fe— r a l h a l l p U t e 3 3 4 .3 2 1 .2 C a n t V l t ta b s c e .1 1 6 4 2 1 2 1 X-ray study of (1580) hip mat la Type 316 (72) A ll rvsstns Marne g iv e n H/A Tlsian reaction 583 hip prosthese 317 (15) Tissue lafcei to implants 334 p in s T lt a n lw I f f l - 39 platetscrevs V lt S l I l U l 48 intraard. rods 7 1 la a lfili of (1) Singe knee e a s t 4 ( 100 ) Corrosioa •fetal r r e t c l s g Sane 0 /A G l r f ^ n s , b y l n t p r o s t h e s is w ro u g h t 4 115 re m o va l v t t a l l t u : d e s ig n • SOI Corrosion o f tome plate 6 H tc k s 4 orthopedic screws Tibia 6 C a t e r , 117 1* 1*U F o re a rm (3 D 7 )- V l t . (3) tissue fafeci P i t t i n g ’ H o o t s t a t e d MIX ( 9 ) - T ype 304 ( 0 ) C r e v ic e ( 5 2 ) - 30 4 .3 0 3 • (54) tissue iofei t . F r e t t i n g ( 1 1 5 ) - Ty 316 (20) Corrcsloi ( 22 )- 3 1 6 + V U . (18) Corrosion t 4 6 7 K a t c titlr> A l l Types Fatfgup-cor*f ob teal foray of (4 7 8 3 - 314 SS -(62.6) Corns too M e ta l F r e t t i n g (o.r)-3l546- Iori|M -kectinc S c a le s ,54 ■plants in m m (4 3 3 - V i t a l i i — -(21.7) Corrosion 6 C r e v ic e ( 6 .5) Vic. (9. of sellt XC 3 7 Jm slity check (S )flnw plate H fg . S u rfa c e C c h o s a .1 1 6 m ■—sfactttres IJewett ta ll p r o c a t d e f e c t s “ 4 T l t d l l m (4 0 ) P o o r e l f a n a l / Washer Q u a lit y y s i s h ig h Vfclanghlta I n c lu s io n c o u n t , 1s t * (32)(tame pUtes g r a in a I r e f l i p m a il 3141 SS V/tager plates 11 3 oepeetlao of ( 2 ) tome plate* 'X B f * . P i t t i n g C o rr o s io n Stress concent raff i lu e s e s w r r l c e (1) Spinal 4lata ( 314 S t (1 0 0 ) C o r r . 4 p ro e e s littdtgran- (80) fa tlt— in desten-sp(mat J e r M i n llo x e o [I) Rip malt 4 plate/ T r a c tn r a d e s ig n u l a r C o rr S .C .C . p i. Poor U rt 119 ( 1 ) l!p~pr*ethese* - T lt a s in m H r t a l C orr..Fat. la p ac t treat—tin a n F re e . C le s seniitieod :« (. * * ■ C orr.-Facli— 264 U 3 7 •nrfocmancn (421 Bene plate* 4 314 SS - < 10O) C errotia D e s ig n C r e v lc e - C o r r o e i— Corrosion acceier- fe^lastelu, ‘ ■v|lu(tao of ■ Ip n a i l s H s ta l 314 ( 5 ) b | I g « ated at itm i :k ' ■ g lo a te d [4 ) S crew s 314 SS -(211 f r a c t s r * Fretting - c o n c e n . O rttofrtle V *l. [ij. Screw* | V l t a l l o o *(29) Fracture VII. * bUu(M ikula F. R. Thompson Type EICHER HIP PROSTHESIS IN Trapezoidal Stem CFRTIFIED STAINLESS STEEL MICHELE HIP PROSTHESIS Austin T. Moore Long Straight Stem -5 -- til ' PohrihvU-nr < uj»' ft mnr « ompnnrnl ( rm* nl Figure 61 - An illustration of vitallium hip prosthesis, in situ and the various designs available for this p u r p o s e . 12 Figure 62 - Illustration of the mandible (jaw) prostheses-made of stainless steel- acrylic composite. This implant is also available in vitallium. SHOULDER PROSTHESIS Figure 63 - Illustrated application of Type 316 stainless steel in a shoulder prosthetic 19 device. y£- FLATT FINGER AND THUMB PROSTHESES CERTIFIED STAINLESS STEEL j" '-r\.i Metacarpal Phalangeal Prottheala-Smatl InterphAfangeal Proatheaia—Large am*** Thumb Pfoathaai* Figure 64 - Illustration of type 316 stainless steel in the Flatt finger prosthesis.12 total elbow replacement Figure 65 - (a) Cast vitallium elbow prostheses (b) Same as (a) but in situ 269. Figure 66 - (a) Titanium knee protheses in situ (b) Illustrated design of the knee prostheses 270. Figure 67 - Illustration of stainless steel wire fixation, used for laminectomy and fusion of cerical spine. (a) Wire located at appropriate point on pinal col. (b) Bone splint implaced (c) Wire tightened in position SHERMAN TYPE BONE PIATES THE EOOERS PLATES A 9 POKK Illustration of the use of type 316 Figure 68 stainless steel in various bone plate and s c r e w designs. ^ SMo McLaughlin Type Platec | I (a), Kuntscher Cloverleaf Intramedullary Nails V III ilu HI (b ) Hansen-Street Intramedullary Nails Figure 6 9 - (a) Illustration of the use of type 316 stainless steel in two hip nail designs and a x-ray of the device in situ (b) Illustration of the use of the same metal in two different designs of intramedullary n a i l . 12 SACRAL ROO, iue 0 Ilsrto o tp Jb stainless Jib type or Illustration - 70 Figure ARNTN PN INSTRUMENTATION SPINE HARRINGTON te i te ein fHrigo Rods Harrington of design the in steel HOOKS " ahr o rtht end ratchet for washer C" ARL EYELET SACRAL 273. Figure 71 - An x-ray showing the Harrington Rod in the spinal area for treatment of scolios 275 Figure 72 - Illustration of the use of tantalum and methyl methacrylate in cranioplasty. (a) Exposure of defect and separation of pericranium from dura. (b) Tantalum plate being screwed into place, (c-d) Different methods of attaching peri cranial layer to tantalum plate. (e) Wedge cut out of plate to secure better fit. (f-g) Cross section of repair, showing initial dead space between plate and dura (f) and absence of dead space with methyl methacrylate repair (g). 275. 276. Figure 73 - X-ray of stainless steel intramedullary nail for internal fixation of a multiple fracture of the femur. V I \ -J / / \ < o- * Figure 74 - X-ray of stainless steel (Kirschner) wire fixation for a fracture of the neck of the radius in a child, (from Campbell's Operative Orthopedics, p. 539). 278. i c ;t ^iv.o; ■\KCDIC REAU'l'1 uN CATHUjIC HE.CilCNS let'll Oxid tion: H'.-*Ii'.+n +ne 1. Hydrogen Evolution 2lf++2e»H2 Examples! Ag*Ag++e 2. Oxygen Reduction (acid oolutiona) +2 tn+Zn +2e 02+4H++4e*2H20 Al*Al+^+3© 3. Oxygon Reduction (neutral or bruiic solutions) Q2+2H20+4e*4cH“ 4. I.'etul ion reductions H = 3+e-I..+2 5. A.etal Deposition: It++e*fc Pi Figure 75 - The corrosion process showing the anodic reaction and five possible cathodic reactions. Figure 76 - Short circuited cell showing the election flow from the zinc electrode (anodic reaction) to the inert platinum electrode (where the cathodic reaction occurs by hydrogen evolution) 279. Figure 77 - A stainless steel (type 301) clip and Olivercrona silver clip after being immersed together 1 week in 3 per cent saline in agar. The dark color surrounding the devices shows an intense ioniza tion of iron due to galvanic action between silver and steel. ^ 6 Figure 78 - A type 301 stainless steel clip coupled with a type 420 stainless clip and immersed in 3% saline and agar for one week. The dark color surrounding the couple indicate galvanic action between the martensitic and austenic stainless steels. 280. CJ CM CJ E o CM a O t; Q) CM CM ■P m; o Ih rs o o CM C\' CM CM o o o o o o I I I I i + tn 01 cm\ > X Cxi Figure 79 - Typical example of electrode kinetic behavior (Polarization study) of pure zinc in an acid solution.85 Tr iiumnivo Inncd vo (a) 02 ir.crc ;in p - - ci- E (b) (c) be. o o Figure 80 - Polarization studies for active-passive systems? ’ (a) Typical anodic dissolution behavior of active-passive systems. (b) Typical anodic dissolution with increasing concentration of chloride-shows increases in Cl” allows dissolution (pitting corrosion) to occur at more active potential. (c) Corrosive condition for active-passive systems. 282. HIncorn .'3oln. ilxi o. oii 'I'o Air *31tL 30 (col ci worked) A c t i v e — — 31 CL 3j (■ u.no■ 1 cu) » « . . . Incor.ol Cut) 0.0 C .in t V .i t. 1 1 I Mi. !•< ^1' VO • •• ‘S 1 .0 - Koble __ V' 9 10' 10 J 1.0 10 10 10' Currert (Jcnn.it;- F. A /CI. ‘ Figure 81 - Polarization behavior of Inconel 600, annealed and cold worked type 3161, Stainless Steel, Titanium and Vitallium. -^l — WVWiVvVN' tt.eCTftOMC.TER AUX WE POTENTIOMETER L EC.TRICAL LEADS IS A L T BRIDGE TO ^ • “'JREFERENCE ELEC TRODE INSULATED SOLID PROBE [ „ J 1 I _%PLID AUXILIARY PROBE * ' l \ —HOLLOA’ REFERENCE PROBE _s,_ \ V X- \ ^MUSCLE X t is s u e x - C - ' FIXATION ------pSDEYDEVICE -4 7 , __ *' '* "***^ •'* * i'* , * "L** '• *' * ** * ’*■''*£j *.* ****•• ’ * BONE r = : * » • (b) Figure 82 - Circuit diagrams for conducting linear polarization measurements.102 (a) In vitro (b) In vivo 1C- C Corr. Kate lor L n- or;. / ? ^ 0.1 .01 .001 .01 0.1 1.0 10 1(X; Tine (Lor t, hrc) Figure 83 - Comparison of in-vitro corrosion rates to those in-vivo (thigh muscles of anesthetized dogs) for type 304 stainless steel and titanium plates. In-vitro corrosion rates were determined on specimens maintained at constant potentials (-.lv and .5v versus SCE for type 304 and titanium, respectively) near the midpoints of their passive regions in nitrogen-saturated 0.9% saline.1^3 285. time,t Figure 84 - Schematic potential time measurements showing three types of behavior. (a) General corrosion (b) Pitting corrosion (c) Film remains intact-desirable 0.4. * — - m ■> X8Cr-8Ni 25Cr-20Ni ' Ti T1-16I. o ' 0.2 o.4 i • Inconel-600 Ui— 0-16Pd Ti-5Nb 31 fc SS - 0.2 - 0.2 0.4 (V) • l8Cr-12Ni Klmonic 75 Ti-4Al-41iio-4Si Ti-‘3Ta - 0.2 - 0.2 0.4 J——- ( 16Cr-20Ni Vital Hu m Ti-2Ko -0.21 i-- l,— i- i—• “ 16 16 16 time (d!'.vt>) Figure 85 - Potential time measurements for various alloy systems in simulated body solution; (1) Hanks' solution and 0.17M NaCl solution (2). Note the desired behavior (type C, Figure 22) for titanium and its alloys. 286. Figure 86 The occurrance of crevice corrosion in a type 316 stainless steel Thornton nail and plate after a four year exposure period. 85 % Figure 87 - Schematic representation of the mechanism of crevice corrosion in two stages. (a) Stage I (b) Stage II 400 350 300 250 200 150 100 Type 31 CL Stalnleoo 3toel / (12cir,2) Miyoiolorlc'il Teat Solution bO°C 150 mV/hr 10" 10' 10 10' 10 CUH.U'.riT Dfc'.f.IITY ( f i A /c m 2 ) © 200 DPH (annealed) High pCO^ (l02irn)He), Low p02 (5mmHe), pH 7.33 • 400 DPH (cold-worked) High nCOp (lOOnrniHg), Low d02 (5iur.Hc,) t PH 7.33 a 200DPH (annealed) Low nCC2 ( &irjr.H/-) , Low pCp (llnunHg), pH 7.33 ■ 400EPH (cold-worked) Low pCC2 ( fcmnillg) , Low p02 (lOrr.ir.Hg), pf? 7.38 Figure 88 - The effect of pCC^ and pO£ on the potentiostatic anodic polarization of 200 and 400 DPH type 316L stainless steel s'pecimen. 289. 400 350 300 2 50 P 200 0 T 150 fi N 100 - T I A L -50 ; passive (400' mV -100 vs SCE Typo 310L St'iinloss Steal/ ( 1 2 c j ii2 ) Physiological Te3t Sol'n. 60°G 150ir.V/hr -350- .-2 1 ,2 10' 10- 10' CUHitKNT DENSITY (j* A/cm2) O 200. DPH High pC02 (l02mmHf), Low p02(5mmHg), pH 7.33 • 400 DPH High pC02 (lOOmmHg), Low p02 (5mmHg), pHl 7.33 Figure 89 Schematic representation of crevice corrosion mechanism showing polarizing effect, of passive current for cold-worked metal on the annealed stainless steel. 0 . 4 Type 316 SS Hanks's solution 8 1 6 time (hr) 0 . 4 eH (V) 0- Vitullium Hunks’s solution 0.21 16 time (hr) 0.4 . li Hanks * olution 8 16 time (hr) Figure 90 - Potential/time curves of scratch tests of implant alloys in simulated body fluids (Hanks' solution). Arrow indicates time of scratch initiation. Note the rapid repassivation of titanium in comparison to type 316 stainless steel or vitallium. 291. Co Eh (V) -1 T / i c F 5 “ o Anode Current (A/CIV ). Figure 91 - Schematic polarization curves showing the effect of a scratch or abraded surface inside a crevice will cause a shift in the oxygen-reduction curve (Co) to the load line curve (Co,j[£ ). This results in shift from passive dissolution (pt a) to active corrosion (pt b) for metal Ax (Ax could be a metal like 316 stainless steel). How ever, in metals of more negative passivation potentials (A^x) like titanium, the damaging effect is avoided. 104 Figure 92 - Microstructure of cast vitallium showing the directional morphology of spherical carbides. ^ 3 0 C - ?oc IOC ,l , (nv) c/r —1 Ov. 0 Copper in Winder's Sol'n -?oc v./c cystine. • Copper in Hinter's ool'n with cystine. 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 lor ■'■anodic Figure 93 - The effect of cystine on the polarization behavior of copper in Ringer's solution. Note that the presence of cystine in-vivo concentrations enhances the passivation of copper. without oyetine ool ’n with cystine. ?00 100 I.Mi e -IOC . 0.5 Figure 94 - The effect of cystine on the polarization behavior of nickel in Ringer's solution. Note the presence of cystine in-vivo concentrations inhibits passivation.^ 295. 1 ■ Figure 95 - Harrington spine instrumentation. An internal stainless steel (type 316) device which obtains spinal correction by means of distract ion applied to the concave side of the curve and compression to the convex side, (from p.96 Treatment of Scoliosis, by Harrington). Figure 96 - The use of strain gage instrumentation in the surgical application of a Harrington Rod device. Figure 97 - Illustration of the combined effect of pitting corrosion and fatigue. Surface view shows a lateral extension of the pitting attack in the direction of stress: (above) Transverse view shows that a root like extension of the pits occurs under dynamic loading. (below) Both views show that the damage is both stress and time dependent.153,154 > 297. A 3 y ko i 120 o 30 kni 100 CUii;iK*:T DKN3JTY, pi irn/cn;^ 60 20 6/10 10 10 CYCL1SJ T(. P-\ ILUitiO Figure 98 - The effect of corrosion rate on the endurance of 0.187, C Steel in 3% NaCl solution (deaerated) Note a minimum corrosion rate of 2 amp/cm^ is established below which fatigue life is maximum.156 KiKoV Steel looq 0.2'/ Yield 8tren^th= Si.lplcoi T e s t Temp.-- 79°F T e s t F re q .= lOoOcpm- Iv'ax. C y c lic Load • - 40o0 lb. o - 5000 lb. • - 6000 lb. A - 7000 lb. o - 8000 lb . D- P 00 l b . A- jJOO lb. 100 CKACH' G;iO, 10 Sj| = i.b X io"19a k 3. •For ^ in in./Cycle A K in pciy/TrT 10 20 40 100 700 Stress Intensity Factor RanpefAKji ks.i J i n. Figure 99 - Crack growth rate as a function of stress- intensity range for Ni-Mo-V steel. ^ 2 299. TYPE A TYiG D TYl-ii C AGGH.lJ 'IVit /•■Gil :t ".3GIVE AGGu.',3J3VE log da dN Kc see Kc Iscc or or or Figure 100- Types of fatigue crack growth.^5 10- 2 IN 2 6 0 C-'.VATER HooT AdTf. A533B A -v, f^K I=140k<:» nun • ^KI=130l:g«mm CHACK G iiO ..'T4 R\T£, IN AI. BIi'il'IT A I R ^a/dN, ^KI=140 kg.nuiT^2 mm/cycle k ------j-j— TftKI=l 30 kr.n-jr. J/<; -4 10 10w 10 10 FREQUENCY, cycles/ir.in Figure 101 - Effect of loading frequency on the rate of crack propagation of A533B (HSST) steel under load controlled tension-tension loading in 260°C-water. ^-64 301. Heol Too Uool Ntri ; 0.8 NL'en.'iur.vd Hip 0.6 re-ictior. ir.-Vivo NOR:. ALT7,N1) PCRCii, Vertic-n hip re ^t i o n bodV v. eifrht 0.2 - 0.2 0 0.5 1.0 T] F. N, ONCCNDS Figure 102 - Load-time profile for vertical hip reaction during level walking.166 ©An uocolvod 3.0 ! oor.od 2.0 !> G r 0 r m n t i o n 1.0 A COH.;C:JJ(.I. F .TIGI •. t ■' n n HAHH1 flGTui: At !;;5 I '. * ; !■,(. O’; G.JTAGT i.VljX.iAATKi; .! I HTNGGli ’vi 3GL*I. AT 3 7 ° C . 2$ (ii n ) 0 1.0 10" 10- 10l 10 M Figure 103 - In-vitro testing of Harrington Rods in Ringer's solution and the effect of shot peening. 303. i o - V iol n Ten t (] iri iUnrer 'a 3oIn C e.? >^l O " 21 1 n - 3.3«? \ 9i 10" F’r ti r i;c-Cr ck Orov. th Knte Ve d n / a t ! t 1KCH/CYCL -t 10 10" " i 10 Figure 104 - The effect of simulated body fluid (Ringer's solution) on the fatigue crack growth rate of type 316 stainless steel.166 iue 0 - ceai rpeetto o te influence the of representation Schematic - 105 Figure log CCRKCSiCN FATIGUE C^AC.K GROV’TH KATE A O G K l ' - S 1 V Cm v l X O N M F N T tes nest n h got rt of rate growth the corrosion-fatigue on intensity stress C Y C L I CS f K E SIN S T E N S I T YR A N G E ,A K I K I . Q U r . N U Y MAI TRIAL ! REGION I! , s k c a r c 83 REGION II 305. 600 . 500 Freti.in^+i'Viticue-rtenulta ir Air 400 Alterrv' t irtp Stress 200 150 100 10 ioC Cycles to F'-il uro Figure 106 - The effect of fretting on the fatigue limit of 1090 s t e e l . 172 APPENDIX B The Corrosion Fatigue Properties of Biomdedical Alloys D.F. Bowers* and R.W. Staehle* Department of Metallurgical Engineering The Ohio State University Columbus, Ohio, USA 43210 Submitted as part of the annual report, "Development of Advanced Metallic Materials for Implants and Prosthetic Devices," sponsored by the National Science Foundation, 1971-72. * Research assistant and Professor, respectively. 306 The Corrosion Fatigue Properties of Biomedical Alloys D.F. Bowers and R.W. Staehle Department of Metallurgical Engineering The Ohio State University Summary The high strain, low frequency corrosion fatigue properties of current implant metals have been determined. The metals were as follows: Type 316LVM (low carbon, vacuum melted) stainless steel, wrough vitallium (Stellite 25), titanium, Ti-6A1-4V, and multiphase alloy MP-35N. Tests were conducted in aerated Ringer's solution (lactated) at 37°C and at pH7 and 2. Cantilever bending of plate-type specimens at a variety of stress levies were carried out at 50 cpm. The results of this study are presented by standard S-N (stress vs log number of cycles) type plots. The progressive surface deterioration which occurs during the test life of each fatigue specimen was continuously monitored by open circuit potential measurements. The relative speed of this degradation is shown in potential-time plots. In addition, 308. the results of a study to determine crack initiation by corrosion potential measurement are presented. A series of controlled potential tests (both anodic and cathodic applied potentials) were conducted on titanium and Ti-6A1-4V. Endurance results and current density measurements will demonstrate the effect of fatigue stress on the surface film stability of these metals at the various applied potentials. Finally, microstructural and fractographic analyses will show how the pattern of fatigue crack initiation and propagation varies for each of these metals. It is concluded from these results that titanium and its alloy, Ti-6A1-4V, are superior in fatigue performance in this simulated body environment. INTRODUCTION The suitability of metals as repair or replacement parts in human or animal bodies depends on a number of important medical and metallurgical considerations. One of the most important from the standpoint of metal properties is corrosion fatigue strength. This property realtes the ability of a given metal to withstand dynamic stresses in a particular environment. In material selection for the design of biomedical implants it is concerned with metal performance under cyclic stress in human or animal body en vironments . A review of the literature indicates very little infor mation concerning the corrosion fatigue endurance of current biomedical alloys. In fact, most authors rely on the endurance values for alloys of similar compositions and results from tests conducted in air. For example, in a report concerning metal fatigue of orthopedic alloys, Grover cites the air endurance limits of aircraft metals similar 1 ) Q or equivalent to biomedical compositions. y These endurance limits are reported to be only 35 to 40 per cent 309 310. of the ultimate tensile strength. Grover, cautions that this reduction in strength may be even greater when notches are present on the metal surface, e.g., from corrosion damage. The only other reference to fatigue endurance of actual implant metals was given by Morral in a report on cobalt alloys for implants. -^5 Comparing the properties of sintered and cast vitallium, he cites the endurance limit of cast vitallium at 27 ksi. Since no reference is made to environment or cylcic frequency, it is assumed that these results were also obtained from high frequency tests in air. Perhaps the first study concerning the effect of cyclic loading on implant metals in simulated body environ ments was done by Cohen and co-workers. These investi gators were more concerned with the effect of cyclic stresses on corrosion rate rather than corrosion fatigue strength per se. Bone plate and screw combinations of Type 316 stainless steel and vitallium immersed in Ringer's solution and subjected to a 10 lb-50 cpm load showed a much higher corrosion rate than those held statically in solution. Furthermore, the use of vitallium in combination with Type 316 gave greater weight loss than when like metals were combined in a given plate-screw assemblage. 311. Interest in the corrosion fatigue strength of Type 316 stainless steel was again stimulated when fatigue failures of Harrington rods were reported by Hirsch and co- 1 fiR workers. ° When the critical failure region of the rod device was subjected to bending fatigue (at constant deflec tion) in Ringer's solution, Hirsch, et al., found fatigue failures to occur with 100 pound loads in 100,000 cycles. This result is significant in view of the fact that vivo measurements of applied loads on these devices are reported to vary from 44 to 100 pounds. In addition, to corrosion induced notch effects in lowering the fatigue endurance of metals, corrosion during the fatigue process can accelerate the speed of crack growth. There have been two studies reported on the effect of environment on the fatigue crack propagation in bio medical alloys. Both involve tests on Type 316 stainless steel in the annealed condition. This type of testing employs a specimen containing a predetermined and well defined notch size and the velocity of crack propagation is measured from the crack emanating from that notch. The results of such fatigue tests can be interpreted from an 2 empirical relationship by Paris: where da/dN = crack growth rate, in/cycle A K = range of stress intensity, ksi/in C,m = material-environment constants A log - log plot yields a linear relationship between these parameters. Conlangelo reported an increase in fatigue crack growth rate for specimens exposed to 4 point bending in 0.9% NaCl solution; the value of m increased from 3.23 to 4.16 in 1 84 going from air to the saline solution. Although these room temperature tests were at high frequency (1800 cpm), the author notes that test conducted at more consistent vivo conditions would probably show even a greater change in crack velocity. Wheeler and James conducted such a test on metal of 1 fifi similar composition and metallurgical condition. In addition, they used a load form more consistent to that reported for hip action in normal walking. The results also show an increase in fatigue crack rate in going from air to a simulated body fluid. However the lower frequency (51 cpm) and higher temperature (98°F) did not produce the 313. expected increase in crack velocity predicted by Colangelo. Wheeler and James stated that the discrepancy could be due to the difference in specimen design. In any case, both studies show an increase in the rate of crack propagation in simulated body fluid environments. Since there are two damaging processes continually occurring in corrosion fatigue, it is impossible to define a definite limit in cyclic stress below which no failure is predicted. Rather it becomes necessary to compare the behavior of various metals at a given stress or at a given exposure time. For implant applications, this requires knowledge of imposed stress, cyclic rate of applied stress, and possible changes in environment during the time of implantation. The present state of medical knowledge in this regard denotes.minimal information on the magnitude of stress imposed on an implant device but it is generally accepted to be in the high strain range. The cyclic rate is low and implantation time varies because both long time and short time metal appliances are in use. Even with this limited information, it is possible to at least compare the behavior of various implant alloys under high strain - low cycle fatigue in bodylike environments. 314. It is the purpose of this paper to present the results of such an investigation. This study will show the comparative corrosion fatigue behavior of five biomedical alloys presently used in implant devices. These include Type 316LVM stainless steel, titanium, vitallium, Ti-6A1- 4V, and multiphase alloy MP-35N. The in-vitro testing was conducted by subjecting these metals to high strain-low frequency bending fatigue in a simulated body environment. In addition, electrochemical measurements were taken to monitor the response of the metal to this type of testing. Materials and Methods Although the type and form of loading is not well defined in all implant applications, the reversed bending of plate-type specimens, used in these tests, is similar to the applied loading of certain implant devices, e.g., hip nails, intramedullary nails, bone plates, etc. The test machines are constant deflection devices. This means that work hardening of the specimen results in an increased load applied in each cycle until saturation is achieved. In other words, while implant applications undergo the same load on each cycle, these devices give constant strain. A plexiglas test cell was used to contain the simulated 315. body fluid environment (Figure 107). The design and construction of this cell allows for the placement of necessary equipment to measure the electrochemical response of the metal either at open circuit or applied potential. The open circuit potential was measured relative to a saturated calomel electrode while a specially designed logarithmic converter was used to monitor the current in controlled potential tests. Lactated Ringer's solution was used in these tests and its composition is shown in Table 20. In addition, the solution was modified to pH 2 in order to simulate suspected environmental conditions that operate within the crevices created by implantation. This modification was carried out by hydrochloric acid additions to insure the higher chloride concentrations prevalent in these crevices. An immersion heater was used to maintain solution tempera ture at 37°C and the solution was aerated using an immersed gas diffuser. Two types of geometry were used in the design of test specimens (Figure 108). Type I design was used in initial baseline tests on titanium and 316LVM stainless steel. How ever, the apex and base fillet of this triangular design Specimen Adapter Connection Rod Air + COa Input - Combination Heater and Stirrer Oxygen and C02 Measurement — Reference Electrode Loctated Temperature Ringer Measurement Solution c ° SPECIMEN - Teflon Insulation (bolh ends of specimen) O-Ring Seal Vise Assembly Figure 107-Schematic diagram of corrosion fatigue testing cell Table XX 317. Composition of Lactated Ringer's Solution Amount Amount Constituent (g/100 ml) Element (meg./I iter) Sodium Chloride 0.6 Na. 130 Cl. 109 Potassium Chloride 0.03 K. 4 Calcium Chloride 0.02 C a . 3 Sodium Lactate, Anh. 0.31 Lactate 28 318. Figure 108 Plan view photographs of Types I (top) and Type I I specimens used in corrosion fatigue studies. 319. served as points of stress concentration and virtually all failures occurred at one of these sites. The type I specimens were purchased from Zimmer, USA and it was arranged that the material should be tak n from the same stock as commercial implants and the surfaces prepared in the manner specified for commercial implants. The Type II specimens were designed to obtain greater certainty of stress dependance in corrosion fatigue. The specimens were prepared in our laboratory from material purchased or donated by suppliers of implant materials. The surface of Type II specimens were prepared according to ASTMF86-68 method 2 which is similar to that used by Zimmer (Type I) but avoids the electropolishing step (see Table XXI). The chemical composition and mechanical properties for all specimens is shown in Tables XXII and XXIII. The heat treatment and processing of all materials is shown in Table XXIV, The metallurgical structure of each metal is shown by the metallographic cross sections in Figure 109. The photomicrographs of Type 316LVM show that the cold worked condition was essentially the same from both sources of supply and the stress relieve heat treatment does not alter the microstructure. The work strengthened MP-35N Table XXI 320„ Fatigue Specimen Processing Procedure Step # Design 1-Zimmer Process Step § Design 2-ASTM F-86 Method 2 1 Mechanical Grinding & 1 Mechanical grinding (metallo- Buffing graphic papers 120-0000). 2 Vapor degrease 2 Polish metallographically (5 and 0.3 alumina) wet process. 3 Electro-clean (hot 3 Cold and hot water rinse alkaline cleaner 71-82 C) 4 Acid dip (10% sulfuric 4 Dry-blown hot a ir acid) 5* Electropolish (38-76°C 5 Vapor degrease with time and current density variable dependent on part size). 6 Cold water rinse 6 Wash in hot (85-90°C alconox cleaner by stirring 5 minutes) 7 Passivate (in 50% n itric 7 Cold water rinse acid at room temperature) 8 Cold water rinse 8 Dip in 10% phosphoric acid 9 Hot water rinse 9 Cold water rinse 10 S till air or blown air dry 10 Passivate 30 to 40 minutes in 20% nitric acid at 49 to 60 C. 11 Cold water rinse 12 Hot water rinse 13 Still air dry *This step is omitted on titanium parts. Table XXII Chemical Composition of Materials Used for Fatigue Specimens and Comparison with ASTM Implant Standards Type Metal Source of Material Chemical ComDosition (*) Analysis C Mn Si SP Cr WV Ni Mo Co Fe N 0 H Ti A Rr Type 316LVM ASTM FI 39-71 , .03 2.0 0.75 .01 .025 17.0-20.0 - - 10.0-14.0 2.0- - Bal. - - - - - Stainless Grade 8 max. max. max. max. max. 4.0 Steel Zimmer, U.S.A. .02 1.67 0.62 .021 .021 17.62 - - 13.21 2.28 _ Bal. - .. * Titanium ASTM F67-66, 0.10 ------0.50 0.07 0.45 150ppm Bal. -- Comm. ,Grade 4 max. max. max. max. max. Pure Zimmer, U.S.A. .02 - - -- _ _ .35 ,009 .35 72 ppm Bal. _ Wrought ASTM F90-68 .05-.15 2.0 1.0 - - 19.0-21.0 14.0- - 9.0-11 .0 - Bal . 3.0 ------V i tall i urn max. m a x . 16.0 m a x . H. Stellite Advanced .092 1.42 0.30 .006 .010 19.75 15.02 - 9.57 - 8a 1. 1.90 “ #25 A 1 1 oys ' Ti-6AT-4V ASTM FI35-70 .08 .03 -- -- - 3.5- -- - 0.25 .05 0.13 125ppm Bal. 5.5- - max. max. 4.5 • ma x . m a x . m a x . 6.5 . Reactive Metals, Inc. .01 . . 4.0 _ _ 0.24 .012 .119 90 Dpm Ba l. 6.3 - ! Reactive ■ Titanium Metals. Inc. 1-02 . _ • _ _ 0.37 .215 41 ppm - - i Type 316LVM Carpenter .017 1.66 0.50 .006 .019 17.53 - - 13.47 2.28 - Bal . ------~ i Steel Co. . i MP-35f! Latrcbe j - - - -- 20* - 35* 10* 35* ------Steel Co. i ■J---- t ^Nominal Composition only - not analysed. Table X X III Mechanical Properties of Materials Used for Fatigue Specimens and Comparisons with ASTM Implant Specifications Ultimate Yield Str. Type Metal Condition Tensile (0.2% o ffse t) El on. (%) Reduction Hardness Bend Strenqth (psi) psi. of Area (%' . 100,000 min. 12.0 min. “ ASTM FI39, Gr. B. 125,000 95Rb Cold Worked Type 316 LVM Stainless Cold Worked 122,300 105,300 20.5 54.3 24RC — Zimmer Titanium ASTM F67, Gr.4 80,COO min. 70,000 min. 15.0 min. - - 5.0 TO Zirmier Annealed 102,000 87,000 24.0 - - 6.0 TD V itallium (H.S. #25) ASTM F90, 130,000 min. 55,000 min. - -- - S trip Annealed Annealed 150,000 75,600 54.0 - 24RC - Ti-6A1-4V ASTM FI 36, 130,000 min. 120,000 min. 10.0 min. - - 10 TD _ Annealed Annealed 149,200 136,300 12.0 - - 4.5 TR Titanium RMI - 101 to 104.000 843 to 88.500 22/23.0 3.0 TR Type 316LVM Carpenter C.W. ND* ND ND ND |33/34Rp MP-35N c.w. 225,000 170,000* 42RC - 1 j oj *ND = not determined, properties cited are nominal. t o Table XXIV 323. Heat Treatment and Processing of Fatigue Specimens Source and Specimen Processing and Heat ITEM Type Material Condition Treat Cycle 1 Type 316LVM - Cold Worked Cold ro lle d sheet, annealed and Zimmer and Stress rerolled to 2B* fin is h , Machined Relieved then stress relieved @ 755 F-3hr.— a ir cool. 2 Titanium-Zirnner Annealed 1300°F--2.--air cool. 3 Type 316LVM- Cold Worked Same as Item 1 above with and Design 2 without stress relief 4 Titanium-Design 2 Annealed Same as Item 2 5 V ita llium Annealed Design 2 6 T1-6A1-4V- Annealed 1300-1525°F—2 h r.- air cool Design 2 MP-35N C.W. 42% reduction 7 *2B finish is a bright, cold rolled finish as specified in ASTM F-56. J'T* ) * . ’ /''i i BHHBB ■ ’t 100 pm 200fim ' ■ *ft &\ ** ■ r . 'S’ > (a) Type 316LUM" (Zimmer) Type 1 Spec. (b) Titanium (Zimmer) Type I Spec. 24 Rc ^ /?* •/ - *•* /■■ «* r ■',• ? ■ ■ K■? ' ' ’ ( •S* / / y r> ' * £ . . v . -•Jf « • . % - i r y - \ \ • • • -V*-- \ , > $ r , U-! >. / m -'vv N ’ . . • * ■i.-V r . * > ~ : :C. k *** i* j , ,*■ ► w , ^ * ' k '>’> , •! ■r r L.’ ' . — — — — —. .'• • ; '. !■■ .:»**'■. ii. (c) Type 3T61LVM Carpenter 33.5R .(d ) Titanium RMI Type I I Spec. v Figure 10-9-Metallographic cross sections of alloys used in corrosion fatigue studies. 325. '•_A l ^ . r I l v ~ . /' V . .• >* . JJl • ■ fc '•-i.k^t-w *'^<;i ’ r'A ipr.V vS 200|im \7 200/im (a) H. Stellite V ital! ium 25Rc (b) ADV A lloys Vi t a ll ium HR t ♦«» ' v. 7 , *“ * -V. \ . .. , ^ »V- O .V . * .; i v: • • '. '\ ■'"* v * *. \*' •' «V '•' ■* '■ \ •V ■‘- 7 v Y - V ^ V f . . . -. : : ',.• : , i . '. v ‘ * -V . • ■„>.• : \ , V7.V' •*• ■ , ••'•• ' * f b •' ■, ( «■* , • •/ - ^ *• • f .vJ ■ • i~.< -J* 1 J i* * A . \ 9 : - * ■ ;T- ' - ‘'•iiiibs’b • X’ *' ••,- V '<-‘ • ’.. ■'.'■• 1> . ‘I «■ * •. « ^ » •* ,( ' • .v ' /■_ . - ' •i'.s'sftv■■.,■ .■*•■■.•>■ ■ 1 1 V-. » ^ v >. -r- • V b * .» ' ‘ .. **•.?.,, • ..'■ •• ;v>- »:• b b - 'v ,■■••* jfc •*••;»* • >«.•'• • • I , • . - t 1 ! . • •. i-V - ‘ ‘ 7 f . 'X ■ . : yj v • •• i:.A ■• * ■' rf-v v ‘ v ;- ■ • :. 3 1 « i S w b | »-V • .. ; ; *;•>,.-y:•••'^!4:£h ______V:- .Vi 2 0 0 jjm 2 0 0 jjm b .V' ■' (c) Latrobe MP-35N 42Rc (d) RMI T i-6A-4V Figure 109-Metallographic cross-sections of alloys used in corrosion (e o n t.; fatigue studies. 326. microstrueture contains a fee matrix with strain induced hep platelets. The large grain size in vitallium is typical of this single phase wrought alloy in the annealed condition. In contrast, an extremely fine grain size is depicted in the microstructure of titanium and its alloy Ti-6A1-4V. RESULTS S-N Behavior Results from the testing of Type I stainless steel and titanium specimens in air and lactated Ringer's solution at high and low cyclic frequencies are summarized in Figure 4. This figure shows a substantial difference between the behavior of the stainless steel and titanium, with the latter being less resistant than the former. There was a substantial decrease in fatigue endurance for either material in going from high to low frequency in either test environment. Furthermore, the Ringer's solution relative to air produced a substantial debilitation which is greater in the case of titanium than for stainless steel. The results of testing type 316LVM stainless steel and titanium in Type I and II specimen geometries are shown in Figure 110. In either geometry, small changes in applied stress affect titanium endurance much more than 316LVM. Also titanium is inferior to 316LVM at the highest stress levels in either design geometry but its life exceeds that of stainless steel at intermediate or low 327 Alternating Stress, Ksi 40 50 60 70 80 0 9 20 30 10 \ 329. applied stress in the Type II design. The behavior in the higher stress range is probably dominated more by the relative strength of the two materials rather than corrosion resistance. Figure 110 also shows the effect of the stress concentration produced by the change in cross-section (top and bottom fillets) in Type I design. A substantial notch sensitivity is indicated for titanium while 316LVM shows similar behavior in either design. This has important inferences to implant design but valid proof of notch sensitivity must be established by other tests planned for this purpose. The intent of this study was to compare material endurance under implant conditions rather than material sensitivity to design. Therefore, the remainder of this report summarizes results obtained from Type II specimens. The S-N behavior of all materials is shown in Figure 111. This plot depicts the endurance results obtained under low frequency strain fatigue conditions in Ringer’s solution at pH levels of 2 and 7. On the basis of the S-N behavior depicted, the best choice of implant material in corrosion fatigue at any stress is Ti-6A1-4V. At intermediate or low Maximum Applied Stress ( k s i) 0 8 0 6 0 7 0 4 0 5 ye I Type Titanium iue11Efc f pcmn ein n h orso fatigue corrosion the on design specimen of 111-Effect Figure ye II Type ye D Type nuac of ye 11M tils sel n ttnum m titaniu and steel 3161VM stainless Type f o endurance n attd igrs uton, H, 7C a 5 cpm. 50 at 37°C, pH7, , n tio lu o s Ringer's lactated in 1 LVM 316 a a Edrne f ye 1 LM tils Sel n Titanium and Steel Stainless LVM 316 Type of Endurance _ o Nme o Cycles of Number Log (e n attd igr Slto, H7 3 C, t cpm. 0 5 at , °C 37 7, pH Solution, Ringers Lactated in V )(Pe fet f pcmn ein n h Croin Fatigue Corrosion the on Design Specimen of Effect 7 170 Maximum Applied Stress (ksi ) 150 130 110 0 9 0 7 50 0 3 Figure 113-C crrosicr fa tig u e endurance o f biomedical a llo y s in lactated R inger's solu tio n 27'C, 27'C, n tio solu inger's R lactated in s y llo a biomedical f o endurance e u tig fa crrosicr 113-C Figure C p. Sld ybl -H; pn ybl -pK7. symbols open -pH2; symbols Solid EC cpm.. o Nme f Cycles of Number Log n attd igrs ouin C, 0p . 50cpm , °C 7 3 Solution Ringer's Lactated in orso Ftge nuac o Boeia Alloys Biomedical of Endurance Fatigue Corrosion oi Smbl-H. pn Symbols-pH7. Open bols-pH2. Sym Solid MP-SSN — ^ ^ . 1 LVM 316 Vitallium - -4V 4 - I A .-6 T •Titanium 331 332. applied stress commercially pure titanium appears better than vitallium or stainless steel. The effect of lowering the solution pH on corrosion fatigue endurance is shown in greater detail in Figures 112 and 113. This effect is more detimental to 316LVM and vitallium than the other more corrosion resistant materials. In this regard, a lower pH has its most damaging effect at the lowest applied stress range on any material. Figure 114 shows the results of a stress-relief heat treatment on 316LVM. A beneficial effect on corrosion fatigue endurance is indicated but the improvement is quite small. Electrochemical Behavior The potential-time plot for test specimens of each metal at rest in Ringer's solution (pH 7 and 2) is shown in Figure 115. The behavior at pH 7 indicates that each metal exhibits an initial 2 to 4 hour period of passivation (potential rise) before a near steady state condition is reached. Based on polarization tests, conducted in a separate study by this laboratory, the near steady state corrosion potential indicates that each metal is in the passive state prior to fatigue testing. Ficure Ficure Maximum Applied Stress (ksi ) 80 0 9 0 7 0 6 50 0 4 113-Effect 113-Effect 10 37 ” C. p o, orso f i nuac o viali n 36V, 50 316LVM, and m lliu ita v of endurance e u tig fa corrosion or, pH f o J ______1 LM pH2 LVM 316 L l L I l l o Nme o Cycles of Number Log 10 E ffe c t of pH on Corrosion Corrosion on pH of t c ffe E ialu ad 1 LM, 0p i Rne' Slto, 37°C. Solution, Ringer's in . 50cpm , LVM 316 and Vitallium 1 LM pH7 H p LVM 316 iolu p 7-^ pH Vitollium PH2 Fatigue J ______cpm nuac of o Endurance I L I J L n ne' sol i t , n tio lu o s inger's R in IOb 3 3 3 f pi' or. cf t c e f f E - 4 1 1 MAX. APPLIED STRESS (KSI) 140 120 100 40 20 80 60 eoe Lcae Ringer's Lactated Aeroted ouin 3° 5 cpm. 50 , 37°C Solution, corrGsicr. fatigue endurance- endurance- fatigue corrGsicr. to YLS —> CYCLES o f Ti Ti f o i-6I V H -*“ — “0 * - pH7 4V - 6AI - Ti and and iaim pH Titanium V 4 - 1 A 6 - i T pH 2 H2 * — pH2 at at 7 ------0 n , n o i t u l o s s ' r e e r f R in r p c 50 - • ----- Maximum Applied Stre 6 6 (k s i) 0 9 0 8 0 4 0 5 0 6 0 7 ,4 iue15Efc c pir tes eif n orso ftge nuac o 316LV''. on endurance fatigue corrosion cn relief stress prior cf Ficure 115-Effect o tes Relief Stress No tes Reliefed Stress o Nme o Cycles of Number Log Endurance of 316 LVM, 50cpm. inSolution,37°C,pH7. Ringers 50cpm. Fatigue LVM, Corrosion 316 on of Relief Endurance Stress Prior of Effect I0 6 335 OC Potential ( mVH) 300 0 -3 -200 -too 0 0 6 700 0 0 4 0 0 5 0 0 3 200 100 Fir ;^e 116-Crer - M V L 6 1 3 2 H p , N 5 3 - P M n 2). and circj'i i Titanium pci 35N,pH7 H , N p 5 -3 P M t H I I I t f o ts spt-ciners e-acr. test for ct iX.a J 32 8 2 Time(hours) Vitallium, pH 7 pH Vitallium, Titanium, pH i pH Titanium, Rest Potential m l a i t n e t o P t s e R n e m i c e p S t s e T n o i t u l o S s , ' r d e e g t n i a r R e d A e t a t c a L 732 3 7 H p , C ° 7 3 ;n - Mr, 6 3 3 337. The behavior at pH 2 is similar except the magnitude of potential is always more noble than at pH 7. This is indicative of the more noble potentials associated with ox,idation-reduction reactions in aqueous solutions at low pH. The corrosion potential is less stable at the low pH for all materials. Furthermore, it appears to be in the transpassive region for 316LVM from related polarization tests on this material at pH 2.^-* Once the fatigue process is initiated the potential changes and eventually becomes more negative until it reaches its most active potential at fracture. This pattern is demonstrated for each metal at different applied stress in Figures 116 through 120. The time in test prior to assuming a more negative potential is a function of both the applied stress and the corrosion resistance of the respective metal. When the stress is high, the fall in potential is instantan eous. At lower stresses, the corrosion fatigue potential behavior maybe one of steady state, gradual rise or gradual fall prior to rapid descent to the most active potential at fracture. Aging four hours at 1000°F is reported to increase the yield strength from 170 to 250 ksi.^3 338. Consider the results for 316LVM at intermediate and low stresses. Figure 116 shows a gradual decline in potential prior to the rapid descent at failure. Figures 117 and 118 depicts a gradual rise in potential for MP-35N and Ti-6A1-4V. The potential-time behavior for titanium and vitallium at these stress levels appears to be one of near steady state as shown in Figures 119 and 120. The above results can be related to the strength of the metal's passive film and its repassivation characteristics. For example, the gradual decline in corrosion fatigue potential for 316LVM is interpreted as an early passive film breakdown and continued surface degradation without repassivation. The steady state results for titanium in dicates a stable passive film throughout the fatigue campaign while the near steady state results for vitallium indicate film breakdown and repair. The potential rise with time demonstrated by MP-35N and Ti-6A1-4V implies a ( strong passivating film (film build-up). These results confirm previous work concerning film stability of implant metals-*-®^ as well as the endurance results established in this study. 700 316 L VM ( C.W. 3 Stress Relieved ♦ * a © J 6 0 0 - Fatigue Tested in Ringer's Solution 5 0 cpm, 3 7 °C, Aerated, pH 7 3 2 5 0 0 pH 7(x 3 Time Scale) 7 0 k si \ Scale / 14 16 18 Tim e( hours) Figure 1 1 7 -Open circuit potential (KV^ ) vs. Time (HR.) for Type 316LVM in corrosion fatigue. 300 0 -3 0 C Potential ( mV -200 -100 0 0 6 0 0 4 0 0 7 0 0 5 0 0 3 200 100 7 ,pH7 H p i, s k 170 !20ksi,pH2 - 120ks/, H 7 p V 95 ks. 95 ~ ~ /< X. X / \ Ti hours)me ( p,37° Aerated °C 7 3 cpm, 0 5 SolutionRinger's in MP-35N 9C k 9C si, 2 pH 20 2624 36 Ti~6~f’ 4 / m Corrosion Fatigue Lactated Rmgers Sotufion pH 2 8 7 , 37°C,50cpm v c :8 5c •e c 8 4 3 8 Fic.’-e 119-Gper. circuit potential ( w j vs. tiire (hr.) for Ti-cAl-4V during corrosion faticje. 3 4 2 96 IOC32 36 40 96 IOC32 84 88 80 64 68 76 37°C, 5 0 cpm Lactated Ringer's Solution (pH2 & 7 ) Titanium m Corrosion Fatigue 48 52 56 60 6 0 -Si, pH 7 pH 6 0 -Si, Tim e ( hours) 4 4 20 24 28 circuit potential (MV^,) vs. time (HR.) for titanium in corrosion fatigue. 120-Cc en Fig.'--;- Fig.'--;- A m ) | ()i i iM j m , ) ) o Open Circuit Potential (M VH) vs. Time (Hour) for Vitallium in Corrosion Fatigue Lactated Ringer's Solution (p H 7 S pH2) 37°C , 50cpm . 5 0ksi,p H 2 ^-55ksi, pH7 65 ksi, pHTf" 75ksi,pH7 ______I______I______I______I______I______I______I______I______1______I______I______L_ 0 4 8 12 -16 20 24 28 32 36 40 44 48 Time (Hour) F i g u r e 1-21-Open circuit potential (MV H ) vs. time (HR.) f o r vitallium in c o r r o s i o n rati cue. 344. M etallography and Fractography Optical metallography disclosed a common transgranular mode of crack propagation for each type of implant metal regardless of environmental pH. Typical examples of this behvaior are shown in Figure 121. In test specimen design II the maximum stress is applied to the center region of the specimen. At the same applied stress, the degree of secondary cracking in this region varied among the different metal types. For example, Type 316LVM stainless steel specimens displayed a multi tude of fine cracks within this area as well as isolated instances outside this region. Vitallium and MP-35N specimens contained a small number of coarse secondary cracks which were confined to the area just adjacent to the final fracture plane. The degree of secondary cracking in the titanium metals was minimal, in fact, in many cases, no cracks were detected (Figure 122). In all cases no incidence of corrosion induced cracking could be detected, in fact, with the exception of some random pitting in all the metals at pH 2, corrosion appeared to be only a small part of the total damage (Figure 122). Thus, crack initiation appears to be more of 345. 200(im (a) 200/ym 200/;m (b) (c) Figure 1 2 2 - A common transgranual mode of crack propagation was noted in all implant metal corrosion fatigue failures. Typical examples are shown above: (a) Vitallium (b) Ti-6A1-4V (c) Type 316LVM stainless steel 346. / I « • 200pm 20 0 i/m 3 1 6 L V M % . r 200p m 'S*Sa‘ p H 7 pH 2 V i t a l l i u m V — 200pm 200pm Ti pH 2 T i - 6 A l - p H 2 F i g u r e 123 -Plane surface characterization after exposure to corrosion fatigue conditions. Note the difference in the tendency toward surface rupture or secondary cracking. 31 347. a function of the mechanical strength of the surface film. Furthermore, since the degree of secondary cracking without corrosion damage in each metal type was similar at both pH levels while failure times were substantially reduced at pH 2, the role of corrosion appears to be one that affects crack propagation. The contribution of corrosion damage to crack propagation is further exemplified by fractographic examina tions of specimens tested in both the pH 7 and 2 environ ments. For 316LVM and MP-35N, the corrosion effect is one of excessive dissolution so as to obliterate fatigue striations and induce a high propensity of surface rupture in the form of many crack-like fissures. This effect is intensified in changing from pH 7 to the more aggressive pH 2 environment (Figure 123). Less metal dissolution appears in fracture surfaces of vitallium and titanium where well defined fatigue striations are observed but a definite tendency toward pitting is seen in the striated fracture surfaces exposed to the pH 2 solution (Figure 124). With the exception of isolated areas where small surface fissures appear, corrosion damage was absent on the fracture surface of Ti-6A1-4V in either test environment (Figure 125). 348, pH ? pH 2 (316LVM) pH 7 MP-35N pH 2 F igu r e 1*24 The effect of changing the environmental pH on the corrosion fatigue fractures of Type 316LVM stainless steel and MP-35N. Top 316LVM, left-pH 7, 60 ksi, 72,665 cycles; right-pll 2, 75 ksi, 36,401 cycles. Bottom--MP-35N, left-pH 7, 90 ksi, 98,083 cycles; right-pll 2, 80 ksi, 71,225 cycles. pH 7 VitaIlium pH 2 Figure 125 Top Vitallium, left-pH 7, 55 ksi, 147,378 cycles; right-pH 2, 80 ksi, 35,867 cycles. Bottom— Titanium, left-pH 7, 80 ksi, 73,115 cycles; right-pH 2, 80 ksi, 27,006 cycles. 350. (b) F i g u r e 126 Corrosion fatigue fracture surface of Ti-6A1-4V. (a) 110 ksi, pH 7, 92,683 cycles (b) 80 ksi, pH 2, 190,401 cycles 351. R.esults for Corrosion Fatigue Crack Initiation Study The potential time behavior for each metal in tests which were stopped at various points in the fatigue life is shown in Figures 127 through 13 0. Also shown is the progressive surface damage occurring at these points by photomicrographs and/or scanning electron micrographs which are superimposed on each respective potential decay curve. The characteristic pattern of surface film integrity, demonstrated previously, is also shown by the E vs t behavior. For example, each test on 316LVM ( F ig u r e 126) shows a gradual decline in potential w ith time while that for titanium (Figure 127) was more steady state, etc. Microscopic evidence of fatigue cracks was observed in all 316LVM stainless steel specimens and the potential decay in each case was a continuous rate of decline. This early development of a rapid rate of decline in potential is probably the result of a weak surface film in the presence of a rather high applied load (90% YS). Thus, points of surface breakdown occurred very early in the test perhaps even during the static preload to establish applied stress prior to actual fatigue testing. The severity of the damage by this surface disruption and subsequent pitting in the 352. early stages of fatigue is then indicated by a continual decline in potential. Crack development easily proceeded from highly deformed regions where dense slip bands, intrusion and pits were noted as early as 1500 cycles and by 2750 cycles, cracks were definitely visible by optical microscopy (Figure 126). The results are more promising for metals w ith stronger surface film s as in the case of titanium (Figure 127). The crack observed in the micrograph fo r the specimen stopped at 4427 cycles is definitely indicated by its more active potential-tim e behavior. The fact that its behavior was a departure from that observed for the other tests is probably related to some in itia l flow in this specimen that was not present in the other test specimens. More im portant, however, is that the more active decline in potential is indicative of the fatigue crack developed. The steady state response in potential w ith time observed in a ll the other tests are indicative of the higher surface film strength of this metal and its excellent resistance to breakdown in this chloride environment. It also demonstrates a long in itia tio n stage for fatigue crack development as no cracks were observed for these specimens even at SEM m agnification. OC Pote/v.'ial ( mVH) •400 -300 -200 ■100 600 0 - 400 0 - 500 200 300 100 Figure - - OPTICAL) (O 127 ,0 2O 300 .0 500 ,0 700 ,0 900 000 100 200 300 400 17,000 14,000 13,000 12,000 11,000 10,000 9,000 8,000 7,000 6,000 5,000 4.000 3,000 2POO 1,000 (OPTICAL) 1,500 eye. Stop 1,500 pn ici ptnil s tn pos o ye 1LKosre uig rc iiito study. initiation crack during observed 316LVK Type for plots tine vs. potential circuit Open 750eye 0 5 ,7 2 (SEM) . Stop (OPTICAL) SM) (SEM W W *\ i * ,0 eye. Stop 6,200 ubro CyclesNumberof . r* ry :' t , 8.650 cyeStop 8.650 °,50cpm m at90ksi(90%Y.S.) p c 7 0 °C, 3 5 i igr ouinpH 7 H M in Solution,p V Ringers L 6 1 3 OPTICAL (OPTICAL) Z6024 eye. Z6024 Fracture to 353 O.C. Potential (m VH) 128 ircuit ptnil s tm po fr berved drn cack iiito study. initiation k c cra during d e v r obse m u i n a t i T for plot time vs. potential t i u c r ci n e p 0 ~ 8 2 1 e r u g i F a s(90% % 0 9 ksi( 0 at 8 m p c 0 5 , C ° 7 3 Titanium in Ringer's Solution,pH 7 Solution,pH in Ringer's Titanium Y.S.) ubro Cycles Number of 02 0 8 4 0 0 4 OPTI ) L A IC T P (O 0 0 0 6 eye. 4 0 Stop 6 7 I ) L A C I T P O ( 400 9500 0 5 9 0 0 84 The surface damage sustained by vita lliu m specimens was sim ilar in appearance to that of 316LVM stainless steel in that heavily deformed regions w ith dense slip bands and intrusions developed quite early in the fatigue process. However, the rate of corrosion assisted crack development w ithin these regions occurred at a much later stage in the fatigue life and this was demonstrated by a less drastic decline in potential w ith time (Figure 128). For example, although intrusions were observed as early as 3000 a n d 5500 cycles, deep microcracking at intrusion sites was not observed u n til 12,500 cycles where the potential was in a steeper rate of decline. A more active value of potential r e q u i r e d 6500 additional cycles of fatigue, where optically visib le crack formation was evident (19,000 cycle tests, F ig u r e 128). This behavior, improved to that of stainless steel, is no doubt related to the better corrosion resistance of vitallium as little or no corrosion damage, e°g*> by pits, could be detected in these te s ts . The damage observed on Ti-6A1-4V specimens appeared to be one of surface roughening or tearing (Figure 129). Deep roughening manifested by high areas in combination w ith deep depressions in the surface p ro file could be detected 0 C. Potential (mV 400 0 -4 300 0 -3 -200 600 - 0 0 4 0 0 3 0 0 5 200 iue 29 Oe crut oeta v. ie lt o vtlim bevd uig rc iiito study initiation crack during observed vitallium for plot time vs. potential circuit -Open 9 12 Figure (OPTICAL) 37°C,50cpm at 75ksi (90% YSJ (90% 75ksi at 37°C,50cpm Vitallium in Ringers Solution Ringers in Vitallium 3 ,0 0 0 eye. Stop 5 ,5 0 0 eye. Stop 0 0 ,5 5 eye. Stop 0 0 ,0 3 , pH 7 pH (OPTICAL) ubro Cce x 10 Cycles of Number : .„ •:V jU-'-vXV/ ' jU-'-vXV/ 12,500 eye. Stop 12,500 (OPTICAL) (OPTICAL) 0,0 eye. 100,000 o Fracture to (OPTICAL) 19,000 eye. 19,000 l V r r A r - 168 Stop 0 0 2 0 0 0 , 1 . 6 5 3 350 (SEM) 9 0 0 eye. Stop 300 250 tp 5 0 eye. S/op : 200 (OPTICAL) 150 2,450 eye.Stop 6,667cyc. to Fracture ----'-s's: r£ * f A (OPTICAL) 100 50 0 (SEM 77- 6 Ah 4Vin Ringers Solution, pH 7 (SEM) -5 0 37 ° C , 50 cp m at f 15 ksi (9 0 % Y.S.) -100 ( 1,800 2,400 3,000 3,600 4,200 4,800 5400 Number of Cycles gure 130-Open circuit potential vs. time plot for Ti-6A1-4V observed during crack initiation study. 358. on all of these test specimens. In tests stopped at a time when the corrosion potential was in a definite state of decline, actual surface rupture appeared as tears (1350 and 2450 cycle tests, Figure 129). When little or no change in corrosion potential occurred (the 900 and 5796 cycle tests) there Were no such sites of surface rupture. In the interpretation of all these results on a rather limited number of tests, it is important to note that open circuit potential measurements are so sensitive to prior specimen preparation and imposed experimental conditions that reproducibility in magnitude or decay rate of potential values from one test to the next is impossible. Therefore, any conclusion regarding a "critcal potential", or "critical potential decay", for corrosion fatigue crack initiation is left for future testing. Fractography Once the major crack front had developed in 316LVM, there were three different regions in its propagation path to failure. The first region (extending from the front or back plane surface to a point 1/4 of the way through the specimen thickness) is characterized by excessive surface 359. deterioration and corrosion damage which completely obliterates any evidence of fatigue striations (Figure 130). This would be expected since this region has the longest exposure to the corrosive environment. The next region (from 1/4 t to center) is marked by fatigue striations so that damage in this region appears to be more mechanical fatigue rather than corrosion fatigue. In this region there appears to be insufficient time for corrosion to dominate the damage. The last region at the center of the specimen thickness is the last to fail and this appears to occur instantaneously in a ductile manner as evidenced by the characteristic dimple structure. The regions may be defined in terms of crack growth rate, the first region is that of slow crack growth while the second and third are regions of intermediate and fast crack growth respectively. Such regions of crack growth could not be distinguished as readily on the other implant metal fracture surfaces. For example, vitallium and MP-35N fracture surfaces are dominated by the slow crack growth region where corrosion damage and secondary cracking in the main fracture plane predominates up to the small region at center thickness which is the last to fail (Figure 131). The titanium metals show Fatigue Striations 2 0 0 /im F i g u r e 131-S.E.M. Montage of the fracture surface of 316LVM stainless steel showing three distinct regions of crack propagation. 361. Vitallium Figure 1 $ . 2 s.E.M. Montage of the fracture surface of Vitallium and MP-35N depicting extensive region of slow crack growth (from plane surface to center) where both stress and corrosion participate in total damage to failure. 362. little or no evidence of corrosion damage and the fracture structure is the same from edge to center (Figure 132). The major difference between titanium and its alloy Ti-6A1-4V appears to be in the morphology of the striated networks comprising the fracture surface. These networks are more coarse in pure Ti and contain greater number of fine microcracks. Thus, the failure mode of the single phase pure titanium is more brittle in fracture characteristics than the two phase Ti-6A1-4V material (Figure 132). Results of Applied Potential Tests on Titanium and Ti-6A1-4V The initial results of a study evaluating the effect of applied potential on the corrosion fatigue response of Ti and Ti-6A1-4V is shown in Figures 13 -135. The corrosion fatigue endurance as a function of applied potential is shown in Figure 133, No tests were conducted in the potential range -800 to +800 mV hydrogen but both metals (considered passive in this range) would be expected to show improved endurance as indicated by the dash lines in this figure. It is interesting to note that the improvement in corrosion fatigue endurance for titanium at high cathodic potentials was not demonstrated by Ti-6A1-4V. Since the number of tests on Ti-6A1-4V in this potential range were sufficient to Titanium \V !W T M d t . Ti-6A1-4V Figure 1 3 3 - Typical areas on the fracture surface of titanium and Ti-6A1-4V. The fracture morphology in these metals did not change across the fracture plane. -16 Applied Potential (MV^) 0 0 4 2 1600 -300 0 0 8 00 0 134 Efc o applied of on potentialEffect the - corrosion endurance fatigue Titanium of anc Ti-cAl-4V n itrs ouin 37cC, solution, Applied Rioter's cpm.in aerated, stress at-50 90.. = Y.S. - 4V — V -4 A .-6 T 10 20 ubro Cce xO ) (xIO3 Cycles of Number fet f ple Ptnilo te orso Fatigue Corrosion the on Potential Applied of Effect nuac o Ttnu ' - n igr Solution, Ringers in V -4 I A j'6 T 8 Titanium of Endurance 7C, eae,a 50 p. ple Stess=90% Y. S. % 0 9 = s s tre S Applied cpm. 0 5 at Aerated, , 37°C Titanium - 0 3 0 4 0 5 364 365. insure that the results are genuine, more tests are planned to validate the titanium results. Both metals show reduced endurance when tested at high anodic applied potent ials . A plot of the average current density measured as a function of endurance life of each specimen is shown in Figures 134 and 135. The magnitude of positive (+) current measured at anodic potentials was always much less than negative (-) current magnitudes recorded at cathodic potent ials. This behavior is related to the substantial amount of water reduction occurring under cathodic applied potentials in comparison to rather low metal dissolution rates at applied anodic potentials. In either case, a sharp increase in current occurred just prior to failure corresponding to additional metal dissolution created by new crack surfaces. A negative to positive current reversal occurred near the value of +800 mV hydrogen (Figure 136) though accurate determination of this value requires further testing. Since the increase in current was very small throughout the major portion of the fatigue life no matter the magnitude of potential applied, the current values in this plot are an average of those recorded prior to the point where failure Current Density, j,+ -10 -ioH ijr 135-urn dniy - v. yls n corrosion in cycles f.ii.ino, of vs. -Current (-) density 5 ricjurc 3 1 10 — Fatigue of Titanium and T (-6 A I -4 V Under Applied Under V -4 I A (-6 T and Titanium of Fatigue — O Titanium ( '8 0 0 MV^j) V Titanium (-1 8 4 0 MVH) 0 4 8 (-1 Titanium V MV^j) 0 0 '8 ( Titanium O - - - V A Tf6AI 380 V) - - -480MVH) V M 0 8 (-I4 V I-4 A ,-6 T 0 MVH) 0 8 I3 - ( V 4 I- A 6 f T A ) MVh 5 6 (-9 V I-4 A ,-6 T O Cathodic Potentials, 50cpm ., 37°C Ringers, Stress=90% Y. S Stress=90% Ringers, 37°C ., 50cpm Potentials, Cathodic urn Dniy()v.Cce nCroin Fatigue Corrosion in vs. Cycles (-) Density Current nnrl ijrs srs - 90"' stress - Rimjer's, M-fAl-W under «ippl icd catliudic po tent in 1 s : >. , c p:i., I0a ubr f Cycles of Number Y.S. I0H - - -880MVH) V M 0 8 (-|8 V I-4 A ,-6 T A 1 1- i,i i 1 I J Li'e; 366, ier I03 Current D e n s ity ,j,+ cm2 iue136-Cretdniy -) s nme o yls orso i.uiip.c ■ corrosion i: cycles (-t) of number vs.density Current - 6 3 1 Figure Fatigue of T ,-6A I -4 V Under Applied Anodic Potentials Anodic Applied Under V -4 I ,-6A T of Fatigue urn Dniy()v. ubro Cce i Corrosion in Cycles of Number vs. (+) Density Current 0p , C Rne' Slto ple Sfress=90%Y.S. Applied Solution Ringer's °C 7 3 ., 50cpm - - +1200 Mh) MVh 0 0 2 1 (+ H) V V M I-4 0 A 5 0 ,-6 2 T + ( V -4 O I T,-6A □ yX ouin ple srs - 0 Y.S. 90. - stress potentials applied anodic solution applied under T1-6A1-4V Fracture ubr f Cycles—*■ of Number ‘V / j | c Pii.ger’s 367. -1600 Applied Potentiol (MVH) 0 0 4 2 Fi-.re 1600 0 0 8 0 0 8 13.7 Applies : potential (KVy average vs. ) current density (pp) o-ir 370C 5C Appliedtol-tier, aeratec, cpn. pH stress 9C. 7. - .6I- V -4 T.-6AI 010 10 bevd urn Dniy Aeae cm (Average) Density Current Observed 02 0 . I03 2 I0 Aerated, pH7. 50cpm. Applied Stress=90% Y. S. Stress=90% Applied 50cpm. pH7. Aerated, ple Ptnil MH v.Aeae urn Density Current vs.Average (MVH) Potential Applied M*. cm2 for for Titanium T,a T,a T,-6A I -4 V in Ringers Solution, 37°C Solution, Ringers in V -4 I T,-6A A ± j for Y.S. i ar.dTi Ti-.--'. ,4 if.— -.irger's 10 .5 ' 368. 369. was imminent. Metal dissolution in the form of pitting attack was prevalent on Ti-6A1-4V specimens which were subjected to high anodic potentials during the fatigue process. No such attack could be found on specimens exposed to cathodic potentials (Figure 137). The same effect was found for titanium except that the attack at anodic potentials was in the form of deep microcracks (Figure 138). These observa tions account for the observed reduction in fatigues life at anodic applied potentials. Discussion The comparative corrosion fatigue behavior of five biomedical alloys has been established by low frequency high strain bending fatigue tests in a simulated body environ ment. Precaution was taken to achieve implant conditions both respect to the metallurgical structure and surface condition of the metal and the proper conditions of the environment. In the latter case, two solutions were studied; one to simulate normal body fluid at pH 7 and a second at pH 2 to assume the condition found in crevices which may be created by the metal device, bone and tissue as a consequence of implantation. 370. Cathodic , Anodic -2240 mV S.H.E. +2050 mV S.H.E. Ti-6A1-4V Figure 13 8 The effect of applied potential on the corrosion fatigue fracture of Ti-6A1-4V. Note the increased corrosion damage (pitting) that occurred when anodic potentials were used in these tests. The applied stress was 90% of the yield strength. 371. Cathodic Test (-1840 MVh ) Anodic Test (+1500 MVh ) Titanium Figure 1 3 9 - Fractography of titanium specimens used in potentiostatic tests. Metal dissolution in the form of deep micro-cracking occurred at high anodic applied potentials (right). 372. The Ti-6A1-4V alloy was superior to all other metals in corrosion fatigue endurance in either environment. The mill annealed condition of this alloy produced an extremely fine grain two phase structure which appears very adequate for long implant life at intermediate or high stress levels in fatigue. Preliminary results also indicate that multi phase alloy MP-35N in the cold worked condition provides good fatigue strength under these implant conditions, but additional testing is necessary both in the cold worked condition and the cold worked and aged condition. The order of suitability in corrosion fatigue endurance for the other metals was as follows: titanium (commercially pure grade) followed by vitallium (rought version, i.e., Hayes Stellite No. 25) and cold worked 316LVM stainless steel. While the fatigue endurance results overlap somewhat at the highest applied stress levels, the performance of these metals indicate adequate life only as short time implants in fatigue. Changing the environment pH from 7 to 2 caused greater corrosion damage to all metals tested; however, its effect on corrosion fatigue endurance was not considered significant. The applied stresses which were predominantly used in these tests were greater than 50 per cent of the metal yield 373. strength. In this regard, it appears that a lower pH has a much greater effect on fatigue endurance when the applied stress is lower, e.g., in the range of 30 to 50 per cent yield strength. By continually monitoring the open circuit (corrosion) potential during the fatigue tests, it was possible to study the progressive surface deterioration of each respective metal type. Prior to the onset of failure, three patterns of potential behavior were determined and associated to metal surface film strength. One was that of steady potential decline representing continual surface film breakdown (316LVM). The second was that of near steady state potential with time representing either no breakdown or surface rupture (titanium) or breakdown followed by momentary repair (vitallium). The third pattern was one of potential rise and near steady state representing a strong passive film with build-up rather than break down (Ti-6A1-4V and MP-35N). These patterns in electrochemical response and their associated surface film integrity were confirmed by the endurance results obtained and the metallographic and fractographic examinations presented. The observed behavior of open circuit potential during 374. fatigue led to a stong possibility of using this technique to evaluate the crack initiation stage in corrosion fatigue. In a separate study on 4 of the 5 biometals investigated a good correspondance between surface rupture and crack formation could be related to the observed behavior of open circuit potential. The potential decay pattern characteris tic to each different metal seen earlier was also demonstrat ed in these tests and the progressive surface damage was recorded by metallographic or SEM examination. The establishment of a critical magnitude of potential or rate of potential decay for fatigue crack initiation could not be determined in these tests. The information from the crack initiation study together with SEM fractographic evaluations of crack propa gation provided some interpretation of the corrosion fatigue failure mode among the various implant metals. Type 316LVM stainless steel shows a rather fast crack initiation stage followed by a period of slow crack growth wherein both corrosion and stress contribute to the total damage. This is followed by intermediate to fast crack growth to failure with little or no corrosion damage. The multiphase alloy MP-35N and vitallium showed a somewhat longer initiation 375. time than 316LVM but a slow crack growth to failure. Titan ium and Ti- 6 A1-4V demonstrated very long times before crack initiation flowed by an intermediate rate of crack growth with little or no evidence of corrosion damage. The corrosion fatigue results for the titanium metals in potentiostatic tests were presented as a final part of this investigation. It was shown that severe corrosion damage which substantially reduces fatigue endurance could only be attained under the application of very high anodic potentials. These potentials were in the range of 2.0V hydrogen which is well above those conditions believed to be operating under practical implant conditions. Thus, material selection from a corrosion fatigue standpoint points toward the use of titanium alloys for dependable life as long term implants. CONCLUSIONS The following conclusions seem appropriate to this investigation: 1, The results of this study show that reliable information regarding the expected implant performance under fatigue conditions should be based on tests conducted at low frequency, under the proper environmental conditions. Precaution is also advised in determining the surface finish imparted to the material used in these tests. 2. Under high strain low frequency bending fatigue in a simulated body environment, the order of endurance performance from best to worse that was determined from this study was as follows: a) Ti-6A1-4V (mill annealed condition) b) MP-35N (cold worked condition)* c) Titanium (mill annealed-commercially pure grade) d) Vitallium (annealed, wrought version--HS No.25) e) 316LVM stainless steel (cold worked and stress relieved) *Requires additional testing. 377. 3. The commercially pure titanium used in these tests appears to be a notch-sensitive material under implant conditions. Subject to tests which validate this observa tion, precaution is advised in the design and fabrication of titanium implants to avoid such notches. The same condition may exist for Ti-6A1-4V though no information could be gained from this investigation. 4. The effect of a lower environmental pH as may be found in crevice sites of implants did not lower fatigue strength significantly; however, a large reduction in strength was indicated at lower applied stress, e.g., in the range of 30 to 50 per cent of the yield strength. 5. Open circuit potential measurements are good in dications of the surface deterioration that occurs during the corrosion fatigue process. When the applied stress does not exceed the static mechanical film strength of the metal, this technique could become a useful tool in studying the crack initiation stage in corrosion fatigue. 6. 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