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Fracture and Residual Stress Characterization of Tungsten Carbide 17% Cobalt Thermal Spray Coatings Applied to High Stre...


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FRACTURE AND RESIDUAL-STRESS CHARACTERIZATION OF TUNGSTENCARBIDE 17%-COBALT THERMAL-SPR AY COATINGS APPLIED TO HIGHSTRENGTH STEEL FATIGUE SPECIMENS BY DONALD SCOTT PARKER A THESIS PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLOR IDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF MASTER OF SCIENCE UNIVERSITY OF FLORIDA 2003

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Copyright 2003 by Donald S. Parker

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To my peers, team members, friends, and co-w orkers in the aircraft and aerospace field who helped to advance my know ledge in the field of materi als science and thermal spray coatings; and to develop my research skills to a level I never though t possible. Without their support, encouragement, and critique I neve r would have been able to complete this project.

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iv ACKNOWLEDGMENTS I wish to thank my advisor (Dr. Darryl Butt) for his guidance and patience on this project while I tried to work full time, complete the research, and write the final thesis. I also thank the other members of my comm ittee, (Dr. Gerhard Fuchs and Dr. Jack Mecholsky) for their knowledge and guidance. I also tha nk the following individuals for their contributions: Mr. Bruce Sartwell fo r providing funding, test specimens, and a forum for peer review of the work; Dr. Philip E. Bretz, Mr. Jerry Schell, and Dr. JeamGabriel Legoux for their expert advice in evaluating the coating materials and fundamental mechanical properties; Dr. Thom as Watkins for his assistance and guidance on residual stress measurement and data inte rpretation; Mr. Peter Marciniak for his photographic talents; and Ms. Virginia Cummin gs for her assistance in scanning electron microscopy, metallography, and proof-reading. I also thank the Kennedy Space Center, La bs Division management (Mr. Timothy Bollo, Mr. Scott Murray, and Mr. Steve McDa nels) for their ongoing support throughout my graduate studies. I especially thank my wife, Ms. Pennie Parker; and daughter, Ms. Alli Brown for their encouragement and tolerance of the time needed to complete my studies.

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v TABLE OF CONTENTS Page ACKNOWLEDGMENTS.................................................................................................iv LIST OF TABLES............................................................................................................vii LIST OF FIGURES.........................................................................................................viii ABSTRACT......................................................................................................................x ii CHAPTER 1 INTRODUCTION AND BACKGROUND.................................................................1 Review of the Chrome-Plating Replacement Effort.....................................................1 High-Velocity Oxygen Fuel (HVOF), Thermal-spray Process....................................2 Selection of Tungsten Carbide Coatings......................................................................4 Coating Characterization..............................................................................................7 2 LITERATURE SURVEY.............................................................................................9 Introduction................................................................................................................... 9 Overview of Fatigue.....................................................................................................9 Overview of Residual Stress.......................................................................................11 Overview of Shot Penning..........................................................................................14 Electroplated Chromium and HVOF Applied Coatings.............................................15 Microstructure of the Tungsten Carbide Cobalt Coatings..........................................17 3 EXPERIMENTAL PROCEDURE.............................................................................21 Specimen Manufacture...............................................................................................21 Coating Quality Control.............................................................................................26 Residual Stress Measurement.....................................................................................28 Fatigue Testing...........................................................................................................32 Scanning Electron Microscopy...................................................................................33 4 RESULTS AND DISCUSSION.................................................................................34 Residual Stress Measurements...................................................................................34 Optical and Scanning Electron Microscopy...............................................................39 Low-Applied-Stress Specimens..................................................................................40

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vi Medium Applied Stress Specimens............................................................................50 High Applied Stress Specimens..................................................................................56 5 CONCLUSIONS........................................................................................................65 APPENDIX A EXPERIMENTAl DATA FOR DE VELOPING HVOF SPRAY PARAMETERS FOR TUNGSTEN CARBIDE 17%-COBALT COATINGS.....................................68 Coating-Process Response Measurements.................................................................70 Coating-Deposition Rate............................................................................................70 Microhardness.............................................................................................................70 Rockwell 15N Superficial Surface Hardness.............................................................71 Tensile Bond Strength................................................................................................71 Substrate Temperature................................................................................................71 Almen Strip Deflection...............................................................................................72 B FATIGUE DATA FOR 4340, 300M, AND AERMET 100.......................................84 C RESIDUAL STRESS SCAN DATA........................................................................102 LIST OF REFERENCES.................................................................................................139 BIOGRAPHICAL SKETCH...........................................................................................142

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vii LIST OF TABLES Table page 3-1 Composition for the th ree alloys evaluated................................................................23 3-2 Particle size distribution fo r the Sulzer Metco, Diamalloy 2005 ...........................25 3-3 Experimental conditions for the x-ray measurements................................................32 4-1 Shot-peened bare 4340 steel substrate material..........................................................35 4-2 Shot-peened bare 4340 steel substrate........................................................................35 4-3 Residual stress data for the WC 17% Co HVOF coating in the as-sprayed condition...................................................................................................................36 4-4 Residual stress data for the WC 17%Co HVOF coating in the finish ground, and polished condition....................................................................................................37 4-5 Low –applied stress specimen....................................................................................38 4-6 Medium-applied stress specimen................................................................................38 4-7 High-applied stress specimen.....................................................................................39 A-1 Optimization parameters for the Tungsten-carbide 17% Cobalt coatings.................69

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viii LIST OF FIGURES Figure page 1-1 Exploded view of HVOF gun and spray process.........................................................3 1-2 Morphology of a thermal-spray deposited coating.......................................................3 1-3 Fatigue results for in itial coating evaluation................................................................6 2-1 Model of shot peened surface showi ng the stress profile created by the impact dimples on the surface .............................................................................................15 2-2 Cross-section showing the coating micr ostructure of an as-sprayed WC-17%Co coating applied to a high-s trength 4340 substrate....................................................18 2-3 Polished WC 17%Co coating cross-se ction showing the la mellar structure and uniform distribution of WC particles.......................................................................18 2-4 Higher magnification micrograph of the same sample in Figure 2-3 showing the white spherical WC particles suspended in the lighter gray cobalt matrix..............19 3-1 Fracture toughness data fo r 4340, 300M, and Aermet 100........................................22 3-2 Ductility properties for tensile el ongation and reduction in area data for Aermet 100 versus 4340, and 300M........................................................................23 3-3 Hourglass fatigue specimen and coated section that was used for the study.............24 3-4 Screw type holder used for re sidual stress measurem ent during spray operations.................................................................................................................27 3-5 Scintag PTS Goniometer at Oak Ridge National Laboratory.....................................28 3-6 Theta/2-theta scan results for the coated specimen....................................................29 3-7 Specimen installation and alignment in the Goniometer............................................30 3-8 Orientation of the sample holder and the lead foil masking in the goniometer..........31 4-1 Specimen before fatigue testing.................................................................................40

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ix 4-2 Surface morphology of a finish ground coating that was polished to a 2-4 Ra roughness..................................................................................................................40 4-3 Fatigue specimen tested at 110 ks i cyclic stress for greater than 107 cycles..............41 4-4 Circumferential coating cracks and the branches that are propagating along machine lines on the surface....................................................................................42 4-5 Micrograph of the crack tip fr om the delaminated coating area.................................42 4-6 Coating and substrate at the fract ure surface for an Aermet 100 specimen tested at 135 ksi maximum applied stress................................................................43 4-7 Subsurface defect and small slow crack growth region for the 4340 steel specimen tested at 125 ksi maximum stress.............................................................44 4-8 Subsurface defect at the origin, and the small slow crack growth region for a 300M steel specimen tested at 130 ksi maximum applied stress.............................45 4-9 Aermet 100 specimen showing the much larger slow crack growth region...............45 4-10 Subsurface inclusion and the radial lin es leading to the crack origin for a 4340 steel specimen tested at 125 ksi maximum applied stress...............................46 4-11 4340 specimen showing the secondary cracking and tear ridges along parallel fronts progressing aw ay from the crack origin............................................47 4-12 Faint striations detected near the outer edge of the slow crack growth region in a 300M specimen......................................................................................47 4-13 Microvoid coalescence in the region of fast fracture for a 300M specimen. .........48 4-14 Low magnification micrograph image of the subsurface origin of the Aermet 100 sample tested at 135 ksi........................................................................48 4-15 Tear ridges and worked surface w ith concentric parallel lines on the Aermet 100 sample...................................................................................................49 4-16 Fatigue striations shown on the fracture surface approximately halfway between the origin and the transition zone between slow and rapid crack growth..49 4-17 Very fine circumferential coati ng cracks formed on the 150 ksi maximum applied stress run-out specimens..............................................................................51 4-18 Higher magnification mi crograph showing the very fine circumferential cracks propagating betw een surface defects............................................................51 4-19 Aermet 100 sample tested at 145 ksi maximum applied stress showing the origin to at the surface..............................................................................................52

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x 4-20 300M specimen tested at 145 ksi maximum applied stress......................................52 4-21 Origin area on Aermet 100 fatigue specimen tested at 145 ksi max stress..............53 4-22 Fatigue crack origin at a subsurface inclusion for a 300M specimen tested at 145 ksi maximum applied stress..........................................................................54 4-23 Transition region along the edge of the slow crack growth region for the Aermet 100 specimen...............................................................................................54 4-24 Coating/substrate interface for the Aermet 100 specimen........................................55 4-25 Cross-section through the coating at the fracture surface just below the substrate origin.........................................................................................................56 4-26 Macroscopic image showing the gro ss subsurface inclusion in the 4340 steel specimen tested at 190 ksi max stress at 59 hz and R = 0.1.....................................57 4-27 Subsurface inclusion in the 4340 steel specimen approximately 0.005” in diameter....................................................................................................................58 4-28 4340 steel, 190 ksi maximum applied stress specimen............................................58 4-29 Much smaller slow crack growth regions and the multiple origins for the 4340 steel, 220 ksi maximum applied stress specimen............................................59 4-30 Imbedded aluminum oxide particle at the fracture surface origin for the 220 ksi 4340 steel specimen.....................................................................................59 4-31 Aermet 100 specimen tested at 180 ksi maximum applied stress at 5 hz with R = -1................................................................................................................60 4-32 300M steel specimen tested at 180 ksi maximum applied stress at 5 hz with R = -1................................................................................................................61 4-33 Smeared surface features and no visi ble fatigue indications for the Aermet 100 sample tested at 180 ksi maximum applie d stress at 5 hz, and with R = -1.............61 4-34 Interspersed regions of MVC within the slow crack growth region.of a 300M specimen tested at 180 ksi, 5 hz, and R = -1............................................................62 4-35 Intact coating of the 300M specimen tested at 180 ksi.............................................62 4-36 Delaminated coating around the fr acture surface of the 4340 specimen tested at 220 ksi, 59 hz and R=0.1............................................................................63 4-37 Imbedded particle origin for the 300M sample tested at 180 ksi maximum applied stress............................................................................................................64

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xi 4-38 Origin area for the Aermet 100 specimen tested at 180 ksi......................................64 A-1 Almen strip deflection vs. substrate part temperature...............................................77 A-2 Substrate part temperature vs. nu mber of mils of coating deposited per pass of the torch.................................................................................................78 A-3 Final coating hardness vs. % porosity.......................................................................79 A-4 Powder size trends.....................................................................................................82

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xii Abstract of Thesis Presen ted to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Master of Science FRACTURE AND RESIDUAL-STRESS CHARACTERIZATION OF TUNGSTENCARBIDE 17%-COBALT THERMAL-SPR AY COATINGS APPLIED TO HIGHSTRENGTH STEEL FATIGUE SPECIMENS By Donald Scott Parker May 2003 Chair: Darryl Butt Department: Materials Science and Engineering Under an internationally funded research program directed by the United States Naval Research Laboratory, thermally spraye d coatings of tungste n-carbide 17%-cobalt are being qualified as a replacement for hexa valent chrome plating in commercial and military aircraft applications. A complete understanding of the performance characteristics and applied properties of thes e coatings (including wear, corrosion, fatigue and residual stress) is criti cal for both component repair and new design configuration. Parallel programs are underway to evaluate these coatings. Our study addresses the residual stress state of the appl ied coatings, the effect of fatigue on the initial coating condition, and crack-initiation and crack-propaga tion behavior at vari ous stress levels. These coatings are essentially anisotropic composite structures with aggregates of tungsten-carbide particles bonded to both amor phous and crystalline cobalt phases (with some free tungsten and cobalt suspended within the matrix). Because of the amorphous structure and the complex nature of the meta stable cobalt phases with in the coating, x-ray

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xiii diffraction techniques were used to char acterize only the resi dual stress conditions surrounding the tungsten-carbide pa rticles in the coating. Di ffraction was also used to establish a baseline stress state of the highstrength steel fatigue specimen substrates to determine interfacial effects of the coating. Stress states were evaluated for specimens that were as-sprayed. Stress states were th en compared to those of specimens that had been coated; finish machined; and then subjec ted to low, medium, and high stress fatigue conditions. Triaxial stress calculation results showed significant reduction in compressive residual stress (even a transition to tensile stress) in the radial direction within the coatings because of the applied axial fatigue stresses. Scanning electron microscopy was used to de termine that coating cracks initiated at surface defects present after finish grind; a nd propagated radially toward the substrate along interfacial boundarie s within the cobalt matrix. In this region, cracks propagate along splat and phase boundaries around WC particles that have high residual compressive stress. High-magnification inspec tion also confirmed that substrate fatigue cracks initiate at defects al ong the coating substrate inte rface (where aluminum oxide particles from grit-blasting the substrate ar e imbedded). Cracks that formed in the coating due to the applied axial fatigue st ress propagated from the surface to the interfacial bond line between th e coating and substrate; but did not provide preferential sites for substrate fatigue initiation.

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1 CHAPTER 1 INTRODUCTION AND BACKGROUND Review of the Chrome-Plating Replacement Effort Replacing hexavalent hard chromium plati ng operations has become a high priority for commercial and military aircraft. This is because of the rising environmental cost of handling this material (a carcinogen); and the poor long-term performance of this material in critical mechanical applicati ons. Hard-chromium electroplating is mostly used on aircraft for landing gear compone nts (including axle journals, hydraulic cylinders, locking and support pins, races, lugs and nose gear steering collars, flight controls and access-door actuators). It is also used on slidi ng wear surfaces of bearing journals, dove-tail and mid-span bl ade supports in turbine engines. The Environmental Protection Agency ha s issued reduced allowable discharge concentrations for hexavalent chromium from 1.71 mg/L to 0.55mg/L for existing permitted industrial waste streams (including Department of Defense aerospace facilities) and to 0.07 for new-permit industrial waste st reams. This has created a significant compliance issue for military aircraft overhaul depots and commercial aircraft maintenance facilities where increased contai nment and treatment of the chromium waste stream will add significant cost to long-term maintenance operations. A technical report by the Oklahoma Air L ogistics Facility [SAIC 1994] and an Advanced Technology Report by Northwestern University, both funded under a Defense Advance Research Project Agency (DARPA) [Northwestern 1996] c ontract established that thermal-sprayed coatings are the leadi ng candidates to replace hexavalent chromium

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2 plating for line-of-sight applications [S artwell 2002]. In 1996, the Hard Chrome Alternatives Team was established by the Environmental Security Technology Certification Program under the funding and direction of the Deputy Undersecretary of Defense for Installations a nd Environment to conduct adva nced development work for qualification of high-velocity oxygen-fuel ( HVOF), thermal-spray coatings for direct replacement of chromium plating on military aircraft. After 2 years of coating-property evaluations, tungsten-carbide 17%-cobalt was se lected as the material with the best chance of success. It was also determined at that time that the scale of the program had implications far beyond just the US military. Th us, an international team of commercial airlines, component manufact urers, subcontractors, a nd government engineers was integrated to form the technical base for full qualification of the selected coatings. Boeing, Lockheed, Naval Air Command, a nd the Air Force cognizant engineering organizations all set forth specific evaluation criteria, fatigue-test parameters, wear-test requirements, and flight-test conditions fo r each component being evaluated [Sartwell 2002]. High-Velocity Oxygen Fuel (H VOF), Thermal-spray Process Thermal-spray coatings are applied by f eeding a uniform-sized powder or wire (metal or ceramic) through a combustion plume or electric arc field that projects molten and semi-molten particulates through a superson ic jet stream onto a substrate as shown in Figure 1-1. The resulting coating has a lame llar grain structure of interlocking splats resulting from rapid solidification of small globules. Each thermal-spray process forms distinctly different coatings with unique m echanical and physical properties. For the combustion process, both thermal and kineti c (particle-velocity) energy is transferred

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3 from the particle to the substrate during solidificati on, creating a mechanical and interlocking diffusion (metal lurgical) bond (Figure 1-2). Figure 1-1. Exploded view of HVOF gun and sp ray process. Courtesy of Sulzer Metco Inc. 1101 Prospect Ave. Westbury, NY. Chemical reactions between the particle s in the gas stream can also create exothermic reactions that dramatically increa se both the tensile bond strength at the base metal interface, and the interlemellar streng th of the coating. High velocity oxygen fuel thermal-spray, as its name implies, mixe s oxygen and a fuel gas (usually hydrogen, propylene, kerosene or even natural gas) th at are mixed and ignites them to create Figure 1-2. Morphology of a thermal-sp ray deposited coa ting. [England 1997]

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4 a supersonic combustion plume through which powders are fed axially by a high-pressure inert gas. The molten and semi-molten par ticles are accelerated (to approximately 1850 ft/sec) and propelled toward the substrat e surface [Kim et al.2001]. The high-kinetic low-thermal energy particles create a dense, uniform coating structure with low porosity, high hardness, and high-strength. The unifo rm nature of as-sprayed HVOF coatings allows deposition of complex carbides, cerme ts, and oxide-dispersion materials with high hardness and consistent thickne ss; that can be ground or supe rfinished to provide very smooth surface roughness and low bearing ratio co atings for sliding-wear applications [Sartwell 2002]. Selection of Tungsten Carbide Coatings Tungsten-carbide materials have been widely used historically to protect surfaces from adhesive and abrasive wear in many di fferent applications (for aerospace, paper, and oil and gas industries). Functionally, th e carbides have been used as sprayed coatings, composite coatings, and sintered cermets; all with different fundamental mechanical properties and behavior. For th e aerospace application of replacing hardchrome electroplating, the benefici al wear properties of the ca rbides were very attractive as compared to other materials. However, durability of the brittl e coatings in fatiguesensitive areas was a concern. As discovere d under the initial DAR PA project, thermalsprayed coatings were ideal because of the ease and repeatability of application; and because of the limited heat treatment of th e underlying substrate required for sintered coatings. Once established, the Hard Ch rome Alternative Te am (HCAT) began evaluating several metallurgical coating co mbinations selected for their likely performance regarding sliding wear, seal compatibility, atmospheric corrosion, axial fatigue, and application-process repeatability. Process repeatability was of particular

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5 importance because the structure and chemistry of the coating will have dramatic effects on the as-sprayed residual stress state, distri bution carbide particles in the matrix, oxide formation, carbide dissolution, macro and micro hardness, and interfaci al bond strength. Accurate characterizatio n cannot be performed if the chem ical and mechanical properties are not repeated for each component sprayed. The plasma spray process was eliminated early on because of poor corrosion performance from the porous coatings; and because of the inability to successfully de velop a repeatable coating pr ocess for the carbides that resulted in uniform coating chemistry and di stribution of carbide particles. Thus, the high velocity oxygen fuel (HVOF) process was se lected as the most likely candidate to succeed. From there, generic fatigue and w ear testing was used to further reduce the candidates. Figure 1-3 shows fatigue-test resu lts generated for the tungsten-carbide 17% cobalt coating as compared to the baseline chromium plating cu rrently in use. In this example all of the specimens were shot-p eened 4340 steel hourgla ss fatigue specimens heat-treated to 280-300 ksi and coated with either the HVOF WC-17%Co or electroplated hard chromium. The fatigue behavior of the tungsten-carbide-cobalt coating is at least as good as the chrome-plated specimen (in the axial fatigue conditions set forth in this test, for all stress levels evaluated). Similar, da ta were generated for all candidate coatings. The fatigue behavior of electr oplated chromium was used as a baseline gauge of fatigue performance. Test criteria re quired that to be considered successful, the average number of cycles to failure at each stress level must meet or exceed that of the chromium electroplating.

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6 4340, SMALL HOURGLASS SPECIMEN (0.003" COATING) R = -1, AIR100.0 110.0 120.0 130.0 140.0 150.0 160.0 170.0 180.0 190.0 200.0 1.E+031.E+041.E+051.E+061.E+071.E+08CYCLES TO FAILURE, Nf EHC/Peened EHC/Peened/FIT WCCo/Peened WCCo/Peened/FIT Figure 1-3. Fatigue results for initial coating evaluation. EHC is defined as electroplated chromium and WC-Co refers to the tungsten-carbide 17%-cobalt HVOF sprayed coating [Sartwell 2002]. The coatings that performed well in the initial screening were the tungsten-carbide, 17% cobalt, and the Stellite material known as Tribaloy T-40 0, an agglomerated blend of cobalt, chrome, and molybdenum. Various co mbinations of the tungsten-carbide with lower percentages of binder material were e liminated due to the brittle nature of the coatings in wear tests and unacceptable fatigue debit as compared to the electroplated chromium. Spalling and delamination of the co ating occurred during th e tests that further indicated that the materials were poor candida tes. Both of the remaining HVOF sprayed coating materials, WC-17%Co and T400, were applied to high-strength steel landing gear material hourglass coupons of 4340, 300M, and Aermet 100 and then subjected to axial

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7 fatigue testing at R= -1, at constant amplitude with stre ss values from 110 to 175 ksi maximum applied stress. Block-on-shoe w ear tests with both lubricated and unlubricated conditions; corrosi on tests that included atmo spheric exposure, ASTM B117 salt fog, GM alternate wet/dry with cons tant UV exposure; and electrochemical impedance spectroscopy. The WC-17%-Co co ating had higher fa tigue strength and better corrosion resistance than the T-400 coatings in every te st and performed at least as well or better than the chromi um plating baseline coupons. Th is provided sufficient data to show that the WC-17%-Co HVOF sprayed co ating was the best ca ndidate to replace hexavalent chromium for the landing gear and hydraulic actuator components in aerospace applications [Sartwell 2002]. Coating Characterization Once the coating was selected, final anal ysis was to include a comprehensive evaluation of the sprayed coa ting conditions that would pr ovide a repeatable coating process, minimal fatigue debit, and mechan ical properties to optimize sliding wear performance. Critical to this was underst anding the stress conditions for which coating cracks form, method and direction of propagati on, and what effect these coating cracks have on initiation of substrat e fatigue cracks. The appro ach was to characterize the material condition of both the substrate a nd coating before during and after applied fatigue conditions. Standard hourglass fatigue specimens were coated with WC-17%-Co coatings using the spray parameters devel oped by an L12 design of experiments that optimized the coating for maximu m fatigue life [S artwell 2002]. The specimens were again tested at va rious stress levels from 110 to 175 ksi maximum applied stress at R=-1, and then al so evaluated at stress levels from 180 to 220 ksi maximum applied stress at R=0.1. Then th ey were evaluated using X-ray diffraction

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8 techniques to characterize the residual stress state of specimens subject to increasing axial fatigue loads. The values were compared to initial conditions obtained from virgin samples. Additionally, scanning electron micr oscopy was performed to evaluate both the coating cracks, and the fatigue fractures in th e substrate materials. A detailed literature search was also conducted to compare and contra st the data generated with past work in this field.

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9 CHAPTER 2 LITERATURE SURVEY Introduction The review of technical lit erature germane to this research encompassed fatigue, residual stress, chromium electroplating and tungsten-carbide coba lt HVOF coatings, and microstructural evaluatio n of tungsten-carbide cobalt coatin gs. The intent was to provide a technical basis for the characterizati on study with an emphasis on understanding the underlying aspects of tungsten-carbide-cobalt coating-substrate behavior under axial fatigue conditions. Overview of Fatigue Fatigue is described as th e process by which a compone nt fails due to repeated cyclic loading below the static tensile strength of the base a lloy [Bannantine et al. 1990]. It has been well established, in the literatu re and in text that slip is the primary mechanism by which the deformation process occu rs in metals and as is seen in static testing, the fatigue strength of a metal varies as a function of heat treatment, availability of slip systems in the alloy, and the type of slip. Crack initiation is followed by slip band crack growth and then growth along planes of highest tensile stress. Initiation is also controlled in part by the surface morphology and the availability of su rface defects like porosity, machining lines, and mechanical or environmental damage that acts as a stress concentration. There are also two distinct types of fa tigue: low cycle, referred to as load controlled, usually described as less than 104 cycles of elastic plus plastic strain where

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10 failure results from cyclic strain at high stre ss levels, even approach ing the yield strength of the material. High cycle fatigue refers to a load controlled condition usually at lower stress levels with higher number of cycles to failure (usually > 105). To accurately classify a material’s fatigue st rength, a statistical value is de rived based on the probability of a specimen attaining a specific number of cy cles to failure for a given stress. Testing is performed on multiple standardized specimens for each stress state and load condition where a sinusoidal varying stress is applied a nd recorded over a specif ic number of cycles [Gilliam 1969]. When generating fatigue data for specific materials, the standardized test specimens are manufactured from the same a lloy material, in the same orientation and heat-treated condition as the component for which they are intended to simulate (axial, rotational, bending stress etc.). Testing is performed in a servo-hydraulic machine with loads applied along the same axis as the repr esented component. A computer interface records real time test data that is plotted on a chart displaying stress amplitude vs. number of cycles. Software modeling programs perfor m a mathematical analysis to create what is known as an S-N curve from the best f it models for the data generated. For highstrength steel, as with most ferrous all oy based materials, as the number of cycles increases, the curve will theoretic ally approach an infinite value at a specific stress level, known as the fatigue limit, where a stress am plitudes below this limit will not result in fatigue crack initiation for a given set of conditions [Bannantine et al. 1990]. Environmental effects, as well as any surf ace modification, like grinding, polishing, or coating application can significantly alter this theoretical value. The literature shows however, that when there are circumstances whereby the fatigue behavior is altered, for design purposes, a fatigue debit, or enhancement is

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11 recorded for a specific applie d stress value and test condi tions must be quantified to accurately manufacture a component for its intended service conditions [McGrann et al. 1998]. For example, the aircraft industry ha s used hexavalent chromium electroplating for sliding wear applications and the compone nt designers use a theoretical fatigue debit approximately equal to a 40% reduction in predicted fatigue life where hexavalent chromium electroplating is applied to highstrength steel components heat-treated to greater than 220 ksi tensile stre ngth. The fatigue process is hi ghly sensitive to the surface condition and the chromium electroplating process significantly alters the surface morphology of the substrate due to the electr ochemical bonding process. Shot peening of the surface prior to a plating operation works to mitigate the effects of the tensile stresses; however as the tensile strength of the substr ate increases, the reduction in fatigue life will again increase in spite of the effects of the peening operation [Wang 2002]. This affect is also more pronounced at higher applied stresse s, and at longer lives (higher cycles); therefore a range of fatigue stre ngth values are assigned to ea ch part during full spectrum component fatigue testing at new manufact ure. Consequently, when trying to successfully replace chromium electropla ting with HVOF applied WC-17%Co thermalspray coatings, an entirely new set of fatigue limit levels must be established, along with a fundamental understanding of crack init iation and propagation in the coating, the residual stress state of the coating, and the e ffect it has at the substrate interface with regards to the fatigue crack initiation in the pa rent metal [Nascimento et al. 2001]. Overview of Residual Stress Residual stress is defined as the stress that remains in a material that is at equilibrium with its surroundings without any su stained applied loads. Residual stresses can be introduced into components by a vari ety of forging, welding, heat-treatment and

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12 surface treatment processes, and during th e deposition of multi-phase and composite coating materials. These stresses can significan tly affect the fatigue strength and fracture characteristics of engineeri ng components; therefore, it is essential that accurate information is available of the residual stre ss distributions present for reliable estimates of their useful lifetimes. Residual stress st ates are characterized by the scale over which they self-equilibrate and can be broken down in to three categories: type I are defined as macrostresses that arise from thermal or el astic mismatch and va ry over long distances relative to the size and shape of the component. Type II re fers to intergranular stresses that vary from grain to grain like coherency strains resulting from alloying elements, or multiple phase materials like precipitation harden ing steels or the distribution of carbide particles in a cobalt coating matrix. Type III residual stresses are those that affect areas smaller than grain size to the atomic le vel and encompass par ticle boundaries, phase interfaces, and intersplat boundari es. Type II and type III stresses must balance over their characteristic small distances and are referred to as microstresses. Surface microstresses are often induced to counteract solidification or thermal tensile stresses, or to enhance the localized susceptibility for fatigue crack initiation [Whithers and Bhadeshia 2001]. There are several methods of evaluating these types of residual stresses including x-ray diffraction, hole drilling, boring, deflection (modified layer removal) sectioning, and neutron diffraction and each has positive and negative attributes th at affect resolution and accuracy of the calculated stress. X-ray di ffraction is used because it can detect type I, II, and III stresses non-destructi vely by calculating the elastic strains through the change in the Bragg scattering angle where =2dsin Solving for strain = d/d, and with accurate knowledge of the initial value of d, the stress free interplanar spacing, the

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13 strains can be converted to stress using the st iffness equations. For metals and coatings, this can provide a very accurate determination of the stress values associated with type II and III stresses [Whithers and Bhadeshia 2001] One of the drawbacks of this technique is the limited penetration depth of the x-ray beam, which rest ricts evaluation to the near surface structure, however understanding the n ear surface stress behavior is important to understand fatigue crack formati on and propagation. For x-ray diffraction stress analysis, roughly half of the diffracted radiation or iginates from less th an 0.0004” beneath the surface. However, the x-ray beam is attenu ated exponentially as a function of depth and this rate of attenuation is c ontrolled by the linear absorpti on coefficient, the composition and density of the specimen, and the type of radiation used. Therefore, surface microstress profiles tend to be exponentially weighted averages of the stress at the surface and in the coating layers immediately beneath it. For coating materials, the net macrostress may appear to be continuous over the entire thickness of the coating, but the microstresses form a stress gradient from the surface, along individual splats, phase boundaries, and particle interfaces to the substrate bond line. To compensate for potential errors from different grain orie ntations, and increase the accuracy of the calculated stress, a series of scans are perfor med to locate at least six independent strain measurements for rotations of both and angles. There are othe r inherent difficulties in this technique for coatings, which have multiple elements, oxides, carbides, amorphous phases, and potentially multiple crystallographic structures as outlined by Prevey (1991). However, measurement of stress values asso ciated with known constituents within a coating and along the substrat e surface can provide a rela tive understanding of the internal stress state of the i ndividual phases and particles a nd their relationship to crack

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14 initiation and propagation. Fo r coating materials, internal gradients of tensile and compressive stress will tend to direct the crack path and the propagation rate of a coating fracture. [Nascimento et al. 2001]. Another non-destructive t echnique using neutron diffr action employs the same mathematical model for determining the strain s and has increased penetration depths for a more accurate stress profile as a function of de pth. However, certain materials like the cobalt matrix in tungsten-carbide coati ngs become radioactive upon bombardment by neutrons thus rendering the samples unusable after exposure. Destructive methods of determining residual stresses re quire measurement of the elastic change in a sample as layers of material are removed from the surf ace. The two dimensiona l elastic strain is recorded by strain gauges placed on the opposite surface and a stress gradient profile is established as layers are removed. Meas urement errors can be introduced by the mechanical or chemical methods of rem oving the surface layers and thus accurate readings are difficult to quantify [Prevey 1991]. Overview of Shot Penning Shot peening of the base metal is a cold working process that works to mitigate the effects of surface tensile stre sses created by the machining, forming, or coating process by creating a compressive residual stress zone at and slightly below the metal surface that retards fatigue crack initiation. The impact of the shot media acts as a tiny hammer, imparting to the surface small indentations or dimples. In order for the dimples to be created, the surface fibers of the material must yield in tension whereas below the surface, the fibers try to restore the origin al shape, thereby producing a hemisphere of cold-worked material below the dimple in highl y compressive stress, as much as half the yield strength of the material. Overlapping di mples develop an even layer of metal in

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15 residual compressive stress that tends to inhi bit the formation and growth of any surface cracks or defects as shown in Figure 2-1. Figure 2-1. Model of shot peened surface showing the stress profile created by the impact dimples on the surface [Wang and Platts 2002]. The higher strength materials (higher har dness) resist the plastic deformation of peening and therefore have a much shallo wer compressive layer than lower strength materials. Consequently, as the tensile strength of the substrate increases, the benefits of shot peening will diminish under applied load s because elastic deformation will facilitate relaxation of the induced residual stresses at the surface. This affect is also more pronounced at much higher applied stresses, and at longer lives (higher number of fatigue cycles) [Bannantine et al. 1990]. Electroplated Chromium and HVOF Applied Coatings Hard Chrome plating is an electrolytic process uti lizing a chromic acid-based electrolyte. The part is made the cathode and, with the passage of a DC current via lead

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16 anodes, chromium metal builds on the compone nt surface. The chromium electroplating process generates significant tensile stresses both at the substrate su rface, and within the plated structure. Upon solid ification, the shrinking of chrome deposits due to hydrogen gas diffusing away and as the decompositi on of the intermediate chromium hydride structure creates volume change s that result in tremendous tensile stresses that are in excess of the ultimate tensile strength of the chromium. As a result, web-like cracks form throughout the plated structure to relieve the inte rnal stresses in the coatings, however, this also generates net tensile stresses at the substrate interface that will act to exaggerate surface defects and become pref erential sites for fatigue crack initiation [Nascimento et al. 2001]. Unlike to chromium electroplating, hi gh velocity oxygen fuel (HVOF) applied thermal-spray coatings can generate net residual macrostresses, or type I stresses, that can be compressive, neutral or tensile, depending on the spray parameters used to apply the coatings, and will be relatively uniform across the bulk of the coating. The solidification stresses from unmelted particles, porosit y, carbide decomposition, or incomplete splat formation can constrain the coating matrix in tension; whereas, an optimized coating with low carbon dissolution, and uniform melting of the cobalt binder will create net compressive macrostresses within the coati ng and along the substrate interface. Fatigue testing confirms the relationship between low levels of net compressive macrostresses in applied tungsten-carbide coba lt coatings and improved fatigue performance in axial stress conditions. During fatigue crack growth, the near threshold a nd high growth rate regimes are strongly affected by mean compressive type I stresses which may delay the onset of plastic deformation and crack formation on the surface [Whithers and Bhadeshia 2001].

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17 Therefore, when combined with shot peen ing of the substrate prior to application, tungsten-carbide cobalt coatings on high-streng th steel with a net compressive residual macrostress can approach fatigue strengths ne ar uncoated base metal levels because both the substrate interface, and the free coating su rface are in compression [Tajiri et al. 1998]. It has also been shown in th e literature that when HVOF coatings are applied with a net tensile stress, the effect on the fatigue behavi or of a high-strength steel substrate is as detrimental as that of electroplated chromi um. Since any residual stresses can raise or lower the mean stress experienced over a few fatigue cycles, the tensile stresses at a free surface, or at the coating substrate interface, will accelerate the onset of fatigue crack formation and is therefore unde sirable [Reiners et al. 1998]. Microstructure of the Tungsten Carbide Cobalt Coatings The structure of the initial powder c onsists of angular particles of WC agglomerated and sintered to a cobalt binde r with nearly spherical net powder grains. When the powder is fed through an HVOF combustion plume, the particles experience a very short dwell time in the flame, wh ich allows for maximu m retention of WC, however, the kinetic energy imparted causes th e particles to deform and become almost spherical in the final coating. The highly oxidizing environment of the plume also reacts with the cobalt to form multiple metastable oxides including and amorphous phase that are retained due to the extremely rapid solid ification [Verdon et al. 1998]. Analysis of the microstructure of an HVOF sprayed WC17%Co coating using optical and scanning electron microscopy shows the lamellar morphology of the multiphase cobalt binder and the uniform distribution of tungs ten-carbide particles within the coating matrix (Figures 2-2, 2-3, and 2-4). From reviewing the availa ble literature, it was determined that there

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18 are several techniques that in corporate x-ray diffraction a nd mathematical deconvolution that are used to predict wh at phases and structures ar e present in the coating Figure 2-2. Cross-section showing the coati ng microstructure of an as-sprayed WC17%Co coating applied to a high-stre ngth 4340 substrate. 500x magnification. Figure 2-3. Polished WC 17%Co coating cro ss-section showing the lamellar structure and uniform distribution of WC particles (High magnification SEM micrograph).

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19 Figure 2-4. Higher magnification micrograph of the same sample in Figure 2-3 showing the white spherical WC particles suspended in the lighter gray cobalt matrix. The dark gray wavy structures are bands that form in between splat layers and around WC particles during so lidification that contain complex cobalt oxides, free cobalt and minute amounts of free tungsten. microstructure. A Rietveld refinement me thod of least squares deconvolution can be performed in an attempt to resolve and iden tify overlapped peaks in masked regions or where peak broadening is observe d. This analysis incorporates a statistical probability in an attempt to speculate what phases and stru ctures should be present based on the physics of the coating process and the chemistry of the powder materials sprayed. Precise x-ray mapping of the coating structure chemistry is su bject to interpretation of statistical results with a probability model pred iction and identification of most likely compounds [Prevey 1991]. The results are listed here based on th e data from various x-ray mapping attempts by the referenced authors as well as in the wo rk performed for this thesis. The spherical and semi-spherical white particles are the di stributed tungsten-carbi des (WC), some of which have very small, dark (almost black) patches of W2C sub-carbide that results from dissolution of the WC duri ng the coating deposition pr ocess. The undesirable W2C

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20 particles should be minimized due to their brittle nature and the adverse affect on the overall fracture toughness of the coating. The constituents of varying shades of gray are the different metastable structures of c obalt and cobalt oxides, including an amorphous cobalt phase that form and are trapped dur ing rapid solidification. X-ray diffraction failed to adequately identify the exact chem istry of the oxides, partially due to the amorphous structure of some and partially due to overlapping peaks of the metastable cobalt compounds listed in the JCPDS database The very small white particles are cobalt carbides that take up the remaining ca rbon when dissolution of WC occurs and formation of W2C is either retarded or kinetically unfavorable [S emant 1998]. Porosity is seen as very dark black angular holes and is located along triple points where multiple phases intersect carbide partic le boundaries. Very small quantities of free tungsten and cobalt are indicated by XRD but are difficult to identify optically [Verdon et al. 1997].

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21 CHAPTER 3 EXPERIMENTAL PROCEDURE This chapter discusses the experimental procedures performed in the characterization of the WC-17%-Co HVOF sprayed coatings on high-strength steel specimens. The study was divided into two pa rts. The first involved x-ray diffraction residual stress measurements of coated specim ens before and after fatigue testing. The second part involved optical and scanning elec tron microscopy of the coupons after being subjected to fatigue conditions at lo w, medium and high stress levels. Specimen Manufacture The hourglass fatigue bars were manufactured by Metcut Resear ch of Cincinnati, OH from three common landing gear alloy mate rials, AISI 4340, 300M, and AerMet 100; the compositions are outlined in Table 3-1. Both 4340 and 300M are considered to be high-strength low alloy steels with excelle nt hardenability due to their appreciable amounts of carbon, nickel, chromi um, and molybdenum. Both materials have a similar combination of strength and toughness (50-55 ksi root inch at 280 ksi) over a wide range of section sizes and display uniform micros tructures throughout th e hardenability range [ARP 1631]. However, 300M is essentially a modified 4340 with higher contents of silicon to retard any cementite formation and reduce temper embrittlement; molybdenum to reduce grain boundary segregation; and additions of vanadium to improve the resistance to softening during tempering ope rations and to form carbides which reduce austenite grain growth. Aermet 100 is a high cobalt alloy designed by Carpenter Technology with improved fracture toughness (100 ks i root inch) and higher resistance to

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22 stress corrosion cracking at strength levels greater than 280 ksi, Figure 3-1 [AMS 6532B 2001]. When heat-treated to this high tensile strength condition, both 4340 and 300M have a uniform martensitic microstructure whereas Aermet 100 is a highly alloyed martensitic age hardened steel that is vacuum melted and refined for a contaminant free, ultrafine grain size microstructure that is significantly more ductile than the other two materials. Standard published tensile strength data shows that Aermet 100 has nearly 15% elongation and 55% reduction in area; whereas both 4340 and 300M will experience less than 10% elongation and less than 30% reduction in area, Figure 3-2 [AIR5052 1997]. Figure 3-1. Fracture tou ghness data for 4340, 300M, and Aermet 100 showing the significant improvement for Aermet 100 ove r the other alloys. Data is for 250 ksi tensile strength sp ecimens [AIR5052 1997]. The hourglass fatigue specimens were cut to length from bar stoc k then turned on a lathe into standard configur ation (Figure 3-3). Next th ey were heat-treated to 280–300 ksi verified tensile strength with the 4340 and 300M specimens heat-treated in accordance with Military Specification MI L-H-6875 while the Aermet 100 was heattreated in accordance with Carpenter Technology Process Specification 15169.

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23 Figure 3-2. Ductility properties for tensil e elongation and reducti on in area data for Aermet 100 versus 4340, and 300M (250 ksi) [AIR5052 1997]. Table 3-1. Composition of the three alloys. Element Alloy 4340 Alloy 300M Alloy Aermet 100 Carbon 0.38-0.43 0.40-0.45 0.21-0.25 Manganese 0.65-0.85 0.60-0.90 0.10 max Silicon 0.15-0.35 1.45-1.80 0.10 max Chromium 0.70-0.90 0.70-0.95 2.90-3.30 Nickel 1.65-2.00 1.65-2.00 11.0-12.0 Molybdenum 0.20-0.30 0.30-0.50 1.10-1.30 Copper 0.35 max 0.35 max Phosphorus 0.015 max 0.010 max 0.0080 max – P+S<0.010 Sulfur 0.008 max 0.010 max 0.0050 max – P+S,0.010 Vanadium 0.05-0.10 Cobalt 13.0-14.0 Aluminum 0.015 Titanium 0.015 Oxygen 0.0020 (20 ppm) Nitrogen 0.0015 (15 ppm)

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24 Grinding to final dimension was performed in accordance with the low stress procedures outlined in Mili tary Standard MIL-STD-866 w ith all specimens undergoing Nital Etching as specified in Military Standard MIL-STD 867 to examine for any grinding burns. Upon completion of the non-de structive inspection, the specimens were then baked at 350o F for 8 hours to remove any poten tial residual hydrogen from the etching process. The center gage section was then shot peened to an intensity of 8-12 Almen “A” using S230 wrought steel shot in accordance with AMS 2432 using computer control for accurate c overage [Sartwell 2002]. Figure 3-3. Hourglass fatigue specimen and coated section that was used for the study. [Sartwell 2002] The coated center gage section was prep ared for spraying by abrasively blasting with 54-grit aluminum oxide pa rticulate propelled at 60-80 psi through a vacuum blaster to coarsen the surface to 120-150 Ra for coati ng adhesion. The specimens were then set aligned on centers in a lathe type fixture for rotation during the spray operation. Hard

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25 shadow masking was offset on each end of th e gage section to produce a tapered coating at the termination point. A Sulzer-Metco diamond jet HVOF spray system was used with and x-y gun traverse unit and hydrogen gas for the combustion fuel. The powder material was also from Sulzer-Met co and was designated Diamalloy 2005 83% WC and 17%-Co and was manufactured by the a gglomeration and sintering process [AMS 7881 2003]. The powder particle si ze distribution is shown in Table 3-2 and the process parameters are outlined in Appendix A. The material was fed axially through the center of the combustion gun at 80 psi positive pressu re with nitrogen gas where it mixed with the high pressure combustion gases. The combustion plume of burning gases creates a high velocity stream of accelerated molten cobalt and semi-molten tungsten-carbide particulate towards the surface. Each part icle creates a splat upon impact with the substrate building a lamellar coating structur e as the HVOF gun is traversed. The cobalt dissociates from the agglomerated powder particles and oxidizes in the flame creating multiple cobalt rich metastable phases that encompass the binder of the coating. The spherical tungsten-carbide is uniformly dist ributed throughout the coating matrix upon solidification Table 3-2. Particle si ze distribution for the Su lzer Metco, Diamalloy 2005 agglomerated and sintered WC 17%Co powder by sieve size. Sieve Size Minimum % Maximum % +270 (2.0876 mils) 6% +325 (1.7716 mils) 25% The coating was deposited at a rate of 0.0001” thickness per for pass with 10second intervals between each deposition pass to keep the substrate temperature below 275o F (135o C) during spraying. The final as-s prayed coating thickness was 0.006” and

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26 was uniform over the gage length to within 0.000 5” concentric relative to the cylindrical substrate as measured with a Fisher Scientif ic eddy current thickness gauge. The coating was finish ground between centers to 0.0035” and a surface roughness of 6 Ra using a conforming, 320 grit, diamond impregnated gr inding wheel on a Cincinnati Milicron bench grinder. The final coating th ickness of 0.003” +/0.0005 was achieved by polishing between centers with diamond abrasi ve paper on a conforming hub to a surface finish of 2-4 Ra. Coating Quality Control Quality control coupons were produced with the coated fatigue specimens to verify the coating properties. The results are as follows: hardness of the coating averaged HV300g 1050 across the coating thickness as meas ured on a polished cr oss-section using a diamond indenter and Wilson microhardness te ster in accordance with ASTM E384. Tensile bond strength was measured in accordance with ASTM C633 with the coating sprayed onto the face of a 1” diameter by ” thick cylindrical blank to 0.005” thickness. The coupons are then bonded to a set of inte rnally threaded 1” diameter by 2” long cylindrical blanks using 1” di ameter wafers of Cytec Industries FM 1000 film adhesive. The bonded cylinders were then aligned vertica lly in the threaded fixtures of a servo hydraulic tension/compression machine and then pulled at 0.05 inch/min in tension until failure or the coating or adhesive. The re sults averaged greater than 13,000 psi, with ultimate tensile failure occurring in the film adhesive. A coating is considered acceptable if the cohesive tensile strength (inter-lamella r, inter-particle and inter-splat structural bond) and adhesive strength (substrate bond inte grity) exceeds 13,000 psi. The coatings were evaluated to gauge the degree of re sidual stress in the as-sprayed condition by coating standard Almen “N” strips in identi cal fashion as the fatigue specimens. The

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27 strips are grit blasted on bot h sides prior to spraying to both prepare the surface for coating, and minimize the pre-stressed deflecti on in the strip. The pre-sprayed deflection should be within 0.002” of being flat to mini mize error in reading final arc deflection. The coupons are mounted in the transverse direction on a screw type fixture in accordance with AMS S13165, as shown in Figure 3-4, and then sprayed to a thickness of 0.005”. An unrestrained convex deflection indicates compressi ve residual stress and for these specimens a 0.010” positive net arc heig ht deflection was achieved. It was shown through an L12 design of experiments correlat ing the spraying parameters to almen arc height and fatigue performance that an acceptable deflection is between 0.003” and 0.012”. The L12 design of experiments is cont ained in appendix-a along with the final spray parameter set. Figure 3-4. Screw type holde r used for residual stress measurement during spray operations [AMS S13165 1997]. Th e convex deflection confirms compressive stress and the unrestrai ned arc height provides a degree of residual stress obtained.

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28 Residual Stress Measurement The residual stress experiment was c onducted jointly with Oak Ridge National Laboratory under a facility User Agreemen t with the UT-Battelle High Temperature Materials Laboratory Group, HTML Proposal No. 2002-032. The equipment utilized is a Scintag PTS Goniometer, with a liquid nitrogen cooled Ge de tector. This experimental setup also contained a MAC Science, 18 kW rotating anode generator with a Osmic CMF multiplayer mirror (Figure 3-5). Figure 3-5. Scintag PTS Goniometer at Oa k Ridge National Laboratory used for the residual stress measurements. Baseline /2 scans were run on the bare metal uncoated sample, on both a coated specimen as sprayed, and sprayed and finish ground, virgin WC-17%-Co powder, and virgin WC powder These scans showed that the best peaks to use when evaluating the internal coating residual stress were th e WC 117 (211 diffracti on plane), and 121 (103 diffraction plane) o2 peaks because they were clearly id entified in the database as WC peaks, they did not have any overlapping w ith other peaks, and they had sufficient intensity for an accurate analysis. The cobalt peaks were not well defined in relation to

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29 the database for known CoO compounds, and th e clearly identified peaks for cobalt and its oxides were overlapped by other peaks like the amorphous region between 30 and 50 degrees 2 as shown in Figure 3-6. The amor phous region is a metastable phase consisting of ternary Co, W, and C that is suspended due to the rapid undercooling. The molten material could experience a glass tr ansition since it has been shown that upon annealing above 1000o F (540o C) the binder will recrys tallize [Verdon et al. 1997]. Figure 3-6. Theta/2-theta scan results for the coated specimen. The 117 and 121 2-theta peaks were resolved for both the virgin WC powder and the sprayed coatings so stress scans were performed at these 2-theta angles. Therefore, it was decided that the experime nt would determine the three-axis stress tensors for the WC particles in the matrix for the as-sprayed, finish ground (both untested), and three fatigue tested specimen s subjected to low (110 ksi), medium (150 ksi), and high (220 ksi) axial fatigue stress levels. Lattice spacing values were obtained

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30 from the initial scans on the virgin WC pow der and reduced potential errors in the mathematical calculations. The stress scan samples were mounted such that when the calculations were performed, the three-axis stre ss tensor would indicate axia l, hoop, and radial stresses present in the coating (Figure 3-7). Figure 3-7. Specimen installation and ali gnment in the Goniometer at Oak Ridge National Laboratory. The height alignment was accomplished w ith a dial gauge probe and telescope, which was accurate to +/5 microns. The alignment was confirmed by rotation of the diffracting surface of the specimen 180o and observing that the surface was coincident with the horizontal axis at +/90o of angle Goniometer alignment was verified by examining a LaB6 powder on a zero background plate, and measuring the peak shift of the (510) reflection. The maximum peak shift was 0.01o 2 for tilting and the final xray beam size at the aperture measured 0.6mm x 4mm. Lead foil tape was used at the

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31 ends of the specimen gage section to bloc k returns outside of the desired diffraction region (Figure 3-8). Figure 3-8. Orientation of the sample holder and the lead foil masking in the goniometer. Table 3-3 lists all of the experimental c onditions for the x-ray measurements. Once the specimen and goniometer were aligned, the /2 stress scans were performed from 115o 2 123o and increments of at 0o, 45o, and 90o. The goniometer is also rotated through for the angle +/55, 42, 28.2,and 0, for each increment of angle resulting in 21 scans per sample at 0.02 2 /step; 10 s/step. The total time required was approximately 23.5 hours per scan (6 samples ev aluated) plus sample alignment time. Calculations were performed using Dolle-H auk method assuming a uniaxial stress state [Krause and Hasse 1986]. The mathematical equations are as follows: = (1+ )/E • 11 • sin2 11/E assuming 12= 22= 13= 23= 33= =0. The variables d, E and are the strain, interplana r spacing, Poisson’s ratio, Young’s modulus and stress, respectivel y. The variables and subscripts and 0 refer

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32 to the azimuthal angle, tilt angle and strain-free, respectively and d was taken as the strain free inte rplanar spacing, d0; however, the unstrained value was determined experimentally by x-ray diffraction for the WC virgin powder. Values for Poisson’s ratio and Young’s modulus were obtained from the literature [Noyan and Cohen 1987]. Table 3-3. Experimental conditi ons for the x-ray measurements Parameter Condition Equipment MAC Science 18 kW rotating anode generator Osmic CMF multilayer mirror Scintag PTS goniometer Scintag liquid N2-cooled Ge detector Power 8 kW; 40 kV, 200 mA Radiation Cu, = 1.54059 Incidence slit width 0.5 mm Receiving slit acceptance 0.25; radial divergence limiting (RDL) Soller slit Source to specimen distance 360 mm Specimen to back slit distance 280 mm Tilt axis and angles ; 0, 28.2, 42, 55 (equal steps of sin2 ) Scans 0.02 2 /step; 10 s/step Fatigue Testing Fatigue testing of the specimens was pe rformed by Metcut Research at their Cincinnati Testing Laboratory as part of a parallel research project for the Naval Research Laboratory. For this thesis all samp les provided were tested in accordance with ASTM E466-96 to generate standard S-N curves for high cycle axial fatigue (load controlled) with constant amplitude. This evaluation classifies low, medium a nd high stress levels as 110 – 135 ksi low; 140 – 175 ksi medium; and 180 – 220 ksi high. Each material was subjected to fatigue loads corresponding with the test matrix designed by Naval Air Command Stress Engineering Group. The assigned values are based on actual aircraft components manufactured from specific alloys that were being evaluated for re placement of chrome

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33 plating with tungsten-carbide cobalt coatings. After testing, the data were plotted in the standard manner with stress on the vertical axis and cycles-to-failure on the horizontal axis. A least squares regression was used to produce each S-N curve. The regression involved a linear fit to the data in ln ( N ) vs. ln ( ) space, then calculating a best fit curve in the traditional S-N space. All samples we re manufactured and prepared as indicated above and were tested until failure, or run out at 107 cycles. The raw data set from which the fatigue results were extracted for this ev aluation is contained in Appendix-b [Sartwell 2002]. Scanning Electron Microscopy Scanning electron microscopy was performed on a JEOL JSM 6400 at the University of Florida, Major Analytical In strumentation Center, along with a JEOL JSM 5900LV at NASA, Kennedy Space Center, Florid a. The samples were evaluated at various magnifications to evaluate coa ting surface morphology, coating crack origin analysis, coating fracture surf ace feature identification, and su bstrate fatigue crack origin analysis.

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34 CHAPTER 4 RESULTS AND DISCUSSION Residual Stress Measurements When materials are subjected to deforma tion, slip occurs and the microstructure deforms in the direction of the applied stress Each grain in the structure therefore has regions of the lattice structure that are elas tically strained in a state of tension or compression. In x-ray diffraction, this causes a shift, or broadening, or both to the diffraction lines associated with the diffracti on planes on a plot of intensity vs. 2-theta angle. The elastic strain (strain tensors) su rrounding this deformation can then be calculated from the change in interplanar spacing ( d/d) obtained from experimental xray diffraction experiments [Prevey 1991]. For the d spacing, plotted vs. sin2 at all angle tilts that are linear, the mathematical matrix of stress tensors can be solved by inputting data from +/3 distinct angle tilts (0, 45, and 90 degrees), plotted for all values. The values for the stra in tensors determined from th e x-ray scans, Poisson’s ratio, and Young’s modulus are then i nput into the mathematical eq uations (Excel spreadsheet) to calculate the tri-axial stresses from Hooke’s Law [Noyan and Cohen 1987]. For this experiment, alignment of the samp le was defined such that the axial, hoop, and radial stresses could be extracted for th e base metal gage area that was shot-peened, and then blasted prior to sprayi ng; and for the WC particles in the coating matrix (2-theta diffraction peaks at 117, and 121 degrees) both before and after the samples were subjected to fatigue testing. Th e data for all of the x-ray s cans is compiled in Appendix C

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35 along with the spreadsheet ca lculations for the directi onal stresses. The 4340 highstrength steel substrate materi al was evaluated first to ascer tain the baseline condition of the substrate prior to spraying and to ascertain the baseline stress c ondition on the surface of the shot-peened material. The results are shown in table 4-1. Multiple alloy combinations were not evaluated due to the amount of time required for each scan. Table 4-1. Shot-peened bare 4340 steel substrate material Condition Axial Hoop Radial Shot-peened, 4340 steel substrate 133 ksi Compressive 137 ksi Compressive 22 ksi Compressive For the bare metal shot-peened specimens, the calculated values showed compressive residual stress in all three axes, with the radial direction being the lowest value. This measure for the radial stress is expected, because the shallow penetration of the x-ray beam is measuring a value for a radial stress that is defined as being zero at the free surface and most of the effects of the shot peening on the radial stress will be seen as a gradient beneath the free surface [Prevey 1991 ]. Once the substrates were blasted with aluminum oxide, the three di mensional surface stress state wa s altered such that the net residual stress was slightly more compressive in the axial direction, slightly less in the hoop direction, and near neutra l in the radial directi on as shown in Table 4-2. Table 4-2. Shot-peened bare 4340 steel substr ate material that was blasted with 54 grit aluminum oxide prior to HVOF coating. Condition Axial Hoop Radial Shot-peened, 4340 steel substrate blasted with Aluminum oxide grit 154 ksi Compressive 89 ksi Compressive 3.1 ksi Compressive The effect of the grit-bla sting altered the stress state at the surface due to the change in surface morphology from the smooth dimpled appearance to the coarse surface

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36 formed by the blast media. As a result of the impact, the blast media creates small regions of tensile stresses between the peaks and valleys of the t ooth-like profile on the surface. In the case of a shot -peened surface, the impact velo city of the grit particles will act to increase the peening e ffect on the stress in certain di rections while the cutting of the free surface will tend to relieve some of th e near surface stresses in other directions. The near surface hoop and radial stresses, which are critical to crack initiation resistance during axial fatigue, were significantly effect ed and will tend to reduce the effects of the shot-peening on the substrate. Next, the surface of the as-sprayed and finish ground coated specimen was evaluated prior to fatigue testing in an effort to determine the effects of the grinding and polishing process on the residual stress st ate of the particles within the coating matrix. The as-sprayed specimens showed a surf ace compressive stress for the coating in all three directions as shown in Table 4-3, however, the radial stress value is higher than expected and could be attributed some errors due to the coarse su rface profile of the assprayed coating (150Ra surface roughness). Table 4-3. Residual stress da ta for the WC 17%Co HVOF coating in the as-sprayed condition. Condition Axial Hoop Radial As-sprayed WC-17%Co coating. 117o 2 211 peak 85 ksi Compressive 116 ksi Compressive 99 ksi Compressive As-sprayed WC-17%Co coating. 121o 2 103 peak 50 ksi Compressive 73 ksi Compressive 76 ksi Compressive The finish ground coated specimens showed a significant jump in two directions increasing in both the axial and hoop directi ons; however the radial compressive stress decreased, as would be expected since the as-s prayed radial stress value seemed slightly

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37 high. The surface roughness of the finish ground coating is 2-4 Ra as measured with Fisher stylus profilometer and the smoother finish should provide a more accurate measure of the near surface three-dimensiona l stress field. This indicates that the grinding and polishing process used to ach ieve the desired surface finish have a significant effect on the residual coati ng stresses imparting significantly more compressive stress in the axial and hoop directions. Table 4-4. Residual stress data for the WC 17%Co HVOF coating in the finish ground, and polished condition. Condition Axial Hoop Radial Finish ground WC-17%Co coating. 117o 2 211 peak 264 ksi Compressive 224 ksi Compressive 33 ksi Compressive Finish ground WC-17%Co coating. 121o 2 103 peak 193 ksi Compressive 245 ksi Compressive 22 ksi Compressive The coated 4340 steel specimen subjected to high cycle fatigue testing at low maximum applied stresses loads (110 ksi) showed decreased compressive residual coating stresses in all three axes for the 117o 2 211 peak as shown in Table 4-5. This result was expected, most likely due to re laxation of the residual stresses during the elastic deformation experienced during fatigue testing. The 4340 specimens tested in the range did not fracture and were run-out specimens that reached 107 cycles. The coatings only showed very fine cracks distributed ac ross the surface that were not interconnected and likely caused the reduction in residual comp ressive stress at the surface. The results for the 121o 2 103 peak were skewed due to the shallow peak derived from the scan. The 121 peak was very shallow with low intensity and therefore the data shown indicating a tensile stress is likely to be in accurate. The specimen subjected to fatigue

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38 testing at medium applied stresses loads (150 ksi) showed increased compressive residual coating stresses relative to the 110 ksi Table 4-5. Low-applie d stress specimen. Condition Axial Hoop Radial Fatigue Tested at 110 ksi WC-17%Co coating. 117o 2 211 peak 158 ksi Compressive 146 ksi Compressive 12 ksi Compressive Fatigue Tested at 110 ksi WC-17%Co coating. 121o 2 103 peak 65 ksi Tensile 44 ksi Tensile 56 ksi Tensile specimen in all three axes for the 117o 2 211 peak as shown in Table 4-6. The 121o 2 103 peak showed compressive stress in the axial and hoop direct ions but slightly tensile in the radial direction. The 4340 speci mens tested in this range did not fracture and were run-out specimens that reached 107 cycles. The coatings only showed very fine cracks distributed across the surface that were not interc onnected and likely caused the reduction in residual compressive stress at the surface. The coated 4340 steel specimen fatigue tested the highest st ress level of 220 ksi maximum applied stress were evaluated at a section of coating adjacent to a delamina ted region on the same side as the fracture surface origin. The stresses were slightly lower compressive values than the 150 ksi specimen for the axial and hoop di rections with slightly highe r compressive stress values for the radial direction as shown in table 4-7. Table 4-6. Medium-applied stress specimen. Condition Axial Hoop Radial Fatigue Tested at 150 ksi WC-17%Co coating. 117o 2 211 peak 280 ksi Compressive 215 ksi Compressive 21 ksi Compressive Fatigue Tested at 150 ksi WC-17%Co coating. 121o 2 103 peak 140 ksi Compressive 130 ksi Compressive 27 ksi Tensile

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39 Calculations for the 117o 2 211 peak again showed slightly higher values for compressive stress than the121o 2 103 peak in the axial and hoop directions, but comparable values for the radial direction. Table 4-7. High-applied stress specimen. Condition Axial Hoop Radial Fatigue Tested at 220 ksi WC-17%Co coating. 117o 2 211 peak 219 ksi Compressive 125 ksi Compressive 32 ksi Compressive Fatigue Tested at 220 ksi WC-17%Co coating. 121o 2 103 peak 100 ksi Compressive 33 ksi Compressive 25 ksi Compressive Optical and Scanning Electron Microscopy High magnification inspection began with evaluation of the surface morphology of the sprayed and ground coatings for comparis on with the fatigue tested specimens. Under applied stress, fatigue crack initia tion will occur at a site of high stress concentration and may well occur at stress va lues below theoretically expected values [DePalo et al. 2000]. The image in Figure 41 shows the ground and polished coating in the gage section of the specimen prior to fatigue testing. The process uses diamond abrasives to achieve a 2-4 Ra microfinish that implies a very smooth, mirror-like surface. However, concentric machining lines, surface po rosity and particle pullout are visible on the coating surface, as shown in Figure 4-2, before any stress loads are applied. These inherent defects can provide sufficient stre ss concentration and pr eferential sites for coating cracks to initiate [A hmed and Hatfield 1999]. The objective was to determine the origin of the fatigue cracks for the coated samples subjected to low, medium and high ranges of applied axial fatigue stress and determine what e ffect the coating application had on the fatigue life of the three substrate materials.

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40 Figure 4-1. Specimen before fatigue testing Figure 4-2. Surface morphology of a finish groun d coating that was polished to a 2-4 Ra roughness. Defects and machine lines ar e still evident (Scanning electron microscope micrograph). Low-Applied-Stress Specimens Two results were obtained for this range of applied loads. The first included specimens that did not fracture, but endur ed with fatigue lives greater than 107 cycles. These were classified as run-outs and were removed from testing and evaluated only for coating performance and to see what effect the applied loads had on crack formation and propagation within the coating matrix. Thes e included all of the 4340 samples subjected

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41 to 110 ksi maximum applied stress at 29 hz c onstant amplitude, as well as two samples subjected to 125 ksi maximum applied stress and 29 hz constant amplitude. There were no samples from the lot of 300M or Aermet 100 that achieved the run out condition. The second set included the specimens that fractured at some point during the testing and the fracture surfaces of both the coating substrate were evaluated. For the specimens that had reached the run out condition, the coating was mostly intact except for three 4340 specimens tested at 110 ksi maximum applie d stress and 29 hz that had very small regions of coating delamination at the very edges where the surface was tapered during the spray operation, as shown in Figure 4-3. Figure 4-3. Fatigue specimen tested at 110 ksi cyclic stress for greater than 107 cycles. The coating shows one small delamination near the edge of the coating. High magnification optical and SEM examin ation of one of th e specimens showed that the coating had additional cracks propa gating circumferentia lly adjacent to the delaminated areas, perpendicular to the axis of loading. These cracks appeared to migrate from initial surface defects and propagate towards other cracks and along machining lines from other defects as shown in Figure 4-4 and Figure 4-5. The behavior was consistent for the three run-out specimens that had coating failure.

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42 Figure 4-4. Circumferential co ating cracks and the branches that are propagating along machine lines on the surface shown in the scanning electron microscope micrograph. Figure 4-5. Micrograph of the crack tip from the delaminate d coating area showing the crack relation to the defect s on the coating surface. The coated specimens that fractured during the testing failed in the range from 1.7 million cycles to just below the run out lim it of 10 million cycles. All of the Aermet 100

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43 samples in this range were tested at stre ss levels of 135 ksi maxi mum applied stress and 29hz constant amplitude had cracks that initiat ed at subsurface defects with an average fatigue life of 5.43 million cycles. For 300M, th e samples in this range were tested at both 125 ksi and 59 hz constant amplitude and 130 ksi maximum applied stress at 29 hz constant amplitude and all of these also fr actured due to crack initiation at subsurface defects. The fatigue life for the 125 ksi set of 300M specimens averaged 2,748 million cycles and the 130 ksi set averaged 4.39 million cycles. For 4340, the specimens that fractured were tested at 125 ksi maximum app lied stress at 29 hz constant amplitude and the failure mode was again due to a subsurf ace defect with an aver age fatigue life of 4.78 million cycles. Figure 4-6. Coating and subs trate at the fracture surface for an Aermet 100 specimen tested at 135 ksi maximum applied stress. The coatings had circumferential surface cr acks, similar to the run-out specimens, as well as some degree of interfacial crack s that propagated along splat boundaries and around carbide particle interfaces perpendicula r to the axis of loading; however, the

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44 substrate bond line was relatively intact with no visible spalling or delamination (Figure 4-6). .Optical examination of the substrat e fracture surfaces for the three alloys, as shown in Figures 4-7, 4-8, and 4-9, revealed almost identical features for the 4340, and 300M, with a slow crack growth region of approximately 0.0625” across the surface before unstable crack growth predominates ; whereas, the Aermet 100 has a slow crack growth region of approximately 0.126”, over half the diameter of the specimen. Figure 4-7. Subsurface defect and small sl ow crack growth region for the 4340 steel specimen tested at 125 ksi maximum stress. For 4340 and 300M materials SEM inspection revealed that the mode of crack initiation was from a sub-surface inclusi on as shown in Figure 4-10. The higher magnification images shown in Figure 4-11 display the worked surface morphology of secondary cracks and tear ridges, which were common to both alloys. Fatigue striations were very difficult to locate due to the R = -1 loading condition being fully reversed

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45 which tends to obliterate the fine features in the slow crack growth area during the compression phase. Figure 4-8. Subsurface defect at the origin, and the small slow crack growth region for a 300M steel specimen tested at 130 ksi maximum applied stress. Figure 4-9. Aermet 100 specimen showing the much larger slow crack growth region and the high shear lip. Specimen also failed at a subsurface defect, which is not clearly visible in this image.

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46 However faint striations were detected as s hown in Figure 4-12, near the transition zone on the outer third of the slow growth se mi-circular region for both 4340 and 300M. The fast fracture region for 4340 and 300M wa s dominated by microvoid coalescence, as shown in Figure 4-13 indicating ductile fr acture once the crack became unstable. Figure 4-10. Subsurface inclusion and the radi al lines leading to the crack origin for a 4340 steel specimen tested at 125 ksi ma ximum applied stress. Magnification 150x. The Aermet 100 sample (Figures 4-14 and 4-15) shows similar microscopic fracture features of radial lines emanating away from a subsurface origin and a highly worked surface morphology with secondary cr acks and tear ridges. The difference between the Aermet 100 and the 4340 or 300M was that fatigue striations were much more evident and easier to re solve in the Aermet 100 specime ns and they are shown in Figure 4-16.

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47 Figure 4-11. 4340 specimen showing the s econdary cracking and tear ridges along parallel fronts progressing away fr om the crack orig in. 800x Magnification. Figure 4-12. Faint striations de tected near the outer edge of the slow crack growth region in a 300M specimen. Magnification is 2500x.

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48 Figure 4-13. Microvoid coalescence in the regi on of fast fracture for a 300M specimen. Magnification 2500x. Figure 4-14. Low magnification micrograph image of the subsurface origin of the Aermet 100 sample tested at 135 ksi. Ra dial lines are visible emanating away from the origin area. Magnification 50x.

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49 Figure 4-15. Tear ridges and worked surface wi th concentric parallel lines on the Aermet 100 sample showing the fractur e progression. Magnification 850x. Figure 4-16. Fatigue striations shown on th e fracture surface approximately halfway between the origin and the transition zone between slow and rapid crack growth.

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50 Medium Applied Stress Specimens In this range, all specimens were again test ed at an R = -1, or fully reversed load condition. The 4340 specimens were subjected to 150 ksi maximum applied stress with a constant amplitude of 59hz, 300M specimens were tested at 140 ksi maximum applied stress and 5hz constant amplitude, and the Aermet 100 specimens were tested at 145 ksi at 29hz constant amplitude. All of the 4340 samp les tested in this range of applied loads reached the run-out condition of 107 cycles without fracturing. The Aermet 100 samples averaged 3.3 million cycles to failure, and the 300M specimens averaged 506,000 cycles to failure showing the strong dependence of fatigue life on the fre quency of the applied loads. The coating surface of the run-out specimens appeared intact macroscopically with no apparent spalled regions of de laminations. However, high magnification inspection revealed that cracks had initiated on the surface of the coating perpendicular to the axis of loading. The cracks can be s een propagating between defects on the surface as shown in Figure 4-17 and 4-18. The cracks are located near the center of the coated gage section at the minimum specimen diamet er and location of highest concentrated stress. Optical examination of the Aermet 100 sp ecimens that fractured showed a flat, slow crack growth region with radial lines emanating away from th e origin area that penetrated over half the diameter of the specimen as shown in Figure 4-19. The fast fracture overload region was fibrous and dull gray in a appearance and propagated at distinct 45o angle from the initial fracture surface plane in the direction of axial loading and encompassed the remaining area of the surface. The 300M specimen had a

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51 Figure 4-17. Very fine circumferential coating cracks formed on the 150 ksi maximum applied stress run-out sp ecimens. Magnification 37x. Figure 4-18. Higher magnificat ion micrograph showing the ve ry fine circumferential cracks propagating betw een surface defects. smaller slow crack growth region by comparison (Figure 4-20) and was similar in size to the lower applied stress specimens. The fast fr acture region was divide d into two distinct regions, the first being a coarse dull gray appearance with visi ble radial lines continuing

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52 from the slow crack growth region becoming more pronounced as they reached the outer diameter of the fracture surface. The second region was a final overload shear lip that propagated at a shallow angle away from the fracture surface plane and encompassed most of the circumference of the specimen. Figure 4-19. Aermet 100 sample tested at 145 ksi maximum applied stress showing the origin to at the surface. Figure 4-20. 300M specimen tested at 145 ksi maximum applied stress.

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53 The noticeable difference at this stress level was that for the one Aermet 100 specimen out of the five total at this stress le vel (Figure 4-19) the fracture initiated at a surface defect instead of a subsurface defect. The crack origin was at an imbedded aluminum oxide particle tr apped during the surface prep aration prior to coating deposition as shown in Figure 4-21. There wa s virtually no difference in the number of cycles to failure of this specimen as compared to the others as this one lasted 3.96 million cycles and the average was 3.29 million. The fr acture surface for the 300M specimens at this stress level all initiated a subsurface inclusion as shown in Figure 4-22. Figure 4-21. Origin area on Aermet 100 fati gue specimen tested at 145 ksi max stress. The microscopic fracture surface features for both the Aermet 100, and the 300M materials were very similar to what was seen on the lower applied stress level samples with the slow crack growth areas having a c oncentric semicircle pattern of secondary cracks and tear ridges. The fast fracture, or overload regions including the shear lip were dominated by microvoid coalescence features w ith small interspersed cleavage facets, probably along grain boundary edges, indicati ng a ductile fracture once the crack became unstable (Figure 4-23).

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54 Figure 4-22. Fatigue crack orig in at a subsurface inclusion for a 300M specimen tested at 145 ksi maximum applied stress. Figure 4-23. Transition region along the edge of the slow crack growth region for the Aermet 100 specimen showing the tear ri dges with interspersed regions of microvoid coalescence.

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55 Figure 4-24. Coating/substrate interface for the Aermet 100 specimen showing the coating fracture surface and the crack along the interface. The coatings in this applied stress range showed some surface delamination and spalling near the fracture surface of the 300M specimen (Figure 4-22). However the Aermet 100 sample showed good bond line integrity even after final fr acture but coating cracks were visible along the substrate inte rface and in the first several splat layers deposited. The coating cracks propagated al ong intersplat boundari es and around carbide particles through the thickness of the coating before reaching the substrate interface as shown in Figure 4-24. The through thickness coating cracks do not propagate into the base metal or align near the substrate fatigue origin indicating that they do not provide a preferential site for substrate crack initiation Higher magnification inspection of a cro ss-section through the coating fracture surface showed that coating cracks propagate primarily around carbide particles through the cobalt binder matrix. Regions of carbide pa rticle pullout and intersplat cracking were evident as shown in Figure 4-25.

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56 Figure 4-25. Cross-section through the coat ing at the fracture surface just below the substrate origin showing the intersplat crack propagation through the binder phase and around the carbide pa rticles along the crack front. High Applied Stress Specimens Because of the aggressive nature of the R = -1 test condition, specimens tested at very high stress levels with the load applie d in fully reversed tension and compression fail in a much shorter time period and would be considered low cycle fatigue if not for the load control condition. Th erefore, in this stress range two different R values were evaluated with both Aermet 100 and 300M specimens tested at 180 ksi maximum applied stress at 5hz with R = -1, while the 4340 st eel specimens were ev aluated at 190 ksi and 220 ksi at 59hz, and R = 0.10. The Aermet 100 specimens averaged approximately 43,000 cycles to failure and the 300M specime ns averaged approximately 14,000 cycles to failure showing the increased fract ure toughness of the Aermet 100. The 4340 specimens subjected to the 190 ksi maximum a pplied stress showed three samples with fatigue lives averaging 4.46 million cycles, agai n showing the effects of the R ratio, while two of the specimens failed at much lower lives, 59,000 and 118,000 cycles.

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57 Macroscopic examination showed gross subsurface inclusions approximately 0.005” in diameter in both specimens with a star-like radial pattern em anating from this origin as shown in Figure 4-26, and 4-27. These specimens were deemed exceptions and characterization of their fracture properties wa s unrelated to the coating process. The remaining 4340 specimens showed fracture behavi or similar to what has been shown for the lower applied stress range with a small slow crack growth re gion and radial lines emanating from the origin area. The 220 ksi and 190 ksi specimens are shown in Figures 4-28 and 4-29. There are multiple origins and slow crack growth regions shown for the 220 ksi specimen. This did not reduce the numbe r of cycles to fracture as this specimen recorded the highest fatigue life for the 220 ksi maximum applied stress specimens at 50,779 cycles. Both the 190 ksi and 220 ksi spec imens had surface defects at the origins as shown in Figure 4-30, which were identifi ed as imbedded aluminum oxide particles by energy dispersive spectroscopy (EDS). Figure 4-26. Macroscopic image showing th e gross subsurface inclusion in the 4340 steel specimen tested at 190 ksi max stress at 59 hz and R = 0.1.

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58 Figure 4-27. Subsurface inclusion in the 4340 steel specimen approximately 0.005” in diameter. Figure 4-28. 4340 steel, 190 ksi ma ximum applied stress specimen.

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59 Figure 4-29. Much smaller slow crack growth regions and the multiple origins for the 4340 steel, 220 ksi maximum applied stress specimen. Figure 4-30. Imbedded aluminum oxide particle at the fracture surface origin for the 220 ksi 4340 steel specimen.

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60 The Aermet 100 and 300M specimens are shown in Figure 4-31 and 4-32 with similar features as shown for the lower applied stress specimens. The Aermet 100 showed a much smaller slow crack growth re gion in this applied stress range covering less than half the diameter of the specime n, also evidenced by the reduced cycles to failure. The 300M also showed a smaller sl ow crack growth regi on but also showed a much larger shear lip. Microscopic features were much more obliterated on the Aermet 100 specimen with no visible fatigue striations and smearing of the worked surface as shown in Figure 4-33. The 300M sample had no visible fatigue striations, however very small regions of microvoid coal escence were seen we ll into the slow crack growth region as shown in Figure 4-34 indicating progressive in stability of the crack. Figure 4-31. Aermet 100 specimen tested at 180 ksi maximum applied stress at 5 hz with R = -1.

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61 Figure 4-32. 300M steel specimen tested at 180 ksi maximum applied stress at 5 hz with R = -1. Figure 4-33. Smeared surface features and no visible fatigue indica tions for the Aermet 100 sample tested at 180 ksi maximum app lied stress at 5 hz, and with R = -1.

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62 Figure 4-34. Interspersed regions of MVC within the slow crack growth region.of a 300M specimen tested at 180 ksi, 5 hz, and R = -1. Figure 4-35. Intact coating of th e 300M specimen tested at 180 ksi.

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63 The coatings for the Aermet 100 specimens and the 300M were relatively intact, even near the fracture surface as shown in Figure 4-35. Small circumferential microcracks were visible on th e gage section surface, and interlamellar cracks were evident on the fracture surface plane. The 4340 specimen also showed good adhesion at 190 ksi, but the 220 ksi specimen had circumfe rential delamination of an entire band of coating as shown in Figure 4-36. The bond line interface showed good coating adhesion near the origin, but cracks and delaminations were visible along this bond line in several areas adjacent to the substrate shear lip. Both of these materials had surface defects at the origin, as shown in Figure 4-37, and 4-38 that were confirmed with EDS to be imbedded aluminum oxide particles. Figure 4-36. Delaminated coating around th e fracture surface of the 4340 specimen tested at 220 ksi, 59 hz and R=0.1

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64 Figure 4-37. Imbedded particle origin for the 300M sample tested at 180 ksi maximum applied stress. Figure 4-38. Origin area for the Aermet 100 specimen tested at 180 ksi showing the imbedded aluminum oxide particle.

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65 CHAPTER 5 CONCLUSIONS Cracks in WC 17%Co high velocity oxygen fuel (HVOF) thermal-spray coatings initiate at defects in the coating structure at the surface, and propagate along intersplat regions of high tensile stress within the co ating structure. Almen strip deflections confirm that the bulk macrostress is a net residual compressive stress and x-ray diffraction confirms that the tungsten-carbi de particles are also in compression. Therefore, the cobalt binder must be comprise d of tensile stresses within the matrix of individual phases. Optical and scanning electron microscopy showed that crack propagation will proceed along these inte rfacial boundaries and around the high compressive stressed WC particles to regions of high tensile stress along splat interfaces in the radial direction, perpe ndicular to the axis of loadi ng. Fractographic analysis of specimens that were subjected to low and me dium applied stress fati gue stresses, but did not fracture showed very fine coating cracks that initiated at defects on the surface and propagated circumferentially towards other de fects. As the stresses were increased a continuous, through thickness network of coa ting cracks was formed to relieve the internal stress in the coating. If continuous cracks reached the edge of the coated gage section, delamination or spalling could occu r however there was no evidence that these coating cracks provided a prefer ential site for substrate fatig ue crack initiation. Optical and SEM examination of cross-sections and of the fracture surfaces showed that when the coating cracks reached the substrate, they would either propaga te along the bond-line

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66 interface, or turn back into the coating and propagate along splat boundaries in the first few layers of deposited coating. Fractographic analysis of the fatigue speci mens that fractured at higher applied stresses clearly showed that the substrate origins were located at either a subsurface inclusion, or at a coating-s ubstrate interfacial bond line de fect. The subsurface defects were the predominant mode of initiation at applied fatigue stresses below 180 ksi for R = -1, fully reversed conditions, as well as at lower applied stresses with higher fatigue lives. The substrate-coating interface defects predom inated for the specimens tested at 180 ksi maximum applied stresses where R = -1, a nd at higher stresses of 190 ksi and 220 ksi where R = 0.1. Energy dispersive spectroscopy c onfirmed that the defects at the coatingsubstrate interface were imbedded aluminum oxide particles from the grit-blasting operation prior to coating deposition. The s ubsurface defects could not be identified but could be M2C carbide particles at grain boundaries or sulfur inclusi ons in the matrix. The reduction in fatigue life at low to moderate applied stresses associated with the HVOF coating application can be attributed to the change in the compressive residual stress state of the substrate from the grit-bla sting process prior to coating deposition. At higher stresses, coating substrate surface defect s will control the reduction in fatigue life. Coating cracks do not propagate in to the s ubstrate or provide pr eferential sites for substrate fatigue cracks to initiate. The co mpressive residual stress in the substrate surface acts to divert the coating crack path along the coating-substrate bond-line. The tungsten-carbide cobalt coatings had si milar affects on the three high-strength steel substrates and further investigation of the HVOF coating deposition process is warranted to determine if a change in surf ace preparation for coating deposition including

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67 a reduction in the quantity of imbedded particul ate can significantly decrease the amount of fatigue life reductiont. Also, surface finish ing procedures that further reduce carbide particle pullout as well as machining lines could reduce the tendency for crack formation at low and moderate applied stress levels.

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APPENDIX A EXPERIMENTAl DATA FOR DEVELOPING HVOF SPRAY PARAMETERS FOR TUNGSTEN CARBIDE 17%-COBALT COATINGS

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69Table A-1. Optimization parameters for the Tungsten-carbide 17%-Cobalt coatings Design 1: Use L8 design plus Center Points, 11 runs total FIXED: Levels 54 grit alumina grit blast at 40 psi, 6 inches FACTORS: -1 +1 C Pt Substrate is 4340 steel, 260-280 ksi A Surf Speed,Feed Rate 1335, 5.1 1835, 3.5 1585 ipm, 4.3 Powder size/type is WC-17Co, Diamalloy 2005, Lot 54480 B Combustion Gas 1525 scfh 1825 scfh 1675 scfh Powder Feed Rate** 8.5 lbs/hr C Stoic Ratio 0.405 0.485 0.445 Spray angle is 90 degrees D Spray Distance 10 inch 13 inch 11.5 inch 100 psi cooling air, 4 AJs @ 6 inch spaced over coupon area Carrier gas N2 at 148 psi, 55 flow, air vib @ 20 psi Turntable Robot Spd Robot % @ Spray pattern length Approximately 13 inch A Factor: RPM ipm mm/sec 750 mm/sec Spots/Rev Fixture diameter 2 inch (-1) 212 25 10.6 1.41% 5.1 C Pt 252 35 14.8 1.98% 4.3 (+1) 292 50 21.2 2.82% 3.5 RESPONSES: RELATED CTG FUNCTION: (B,C) Factor Combinations: 1) Part temperature Fatigue Comb Gas Stoic Ratio Hyd SCFH Oxy SCFH Air SCFH Point (CG,SR) 2) Almen strip Fatigue, ctg residual stress 1675 0.445 1159 332 920 ( 0, 0) 3) Hardness, HV300 Wear 1525 0.405 1085 258 920 (-1,-1) 4) Coating dep/pass Cost 1525 0.485 1027 314 920 (-1,+1) 5) Porosity Ctg quality, corrosion 1825 0.405 1299 342 920 (+1,-1) 6) Oxides Ctg quality 1825 0.485 1229 412 920 (+1,+1) 7) Carbides Ctg quality, wear 8) Tensile bond Adhesion/cohesion

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70 Coating-Process Response Measurements A number of process responses were tr acked for the deposited coatings. These included deposition rates, DPH mi crohardness, R15N superficia l surface hardness, tensile bond strength by a modified ASTM C633 method, Almen strip arc height as an indication of coating residual stresses, and substrate temperature during spraying. The specimens for these measurements were rotated on a turntable while mounted on a cylinder as the gun traversed parallel to the cylinder axis of rotation in an oscillating stroke. Coating-Deposition Rate The deposition rate was determined as the measured coating thickness divided by the recorded number of gun passes. The fi nal coating thickness was measured by hand held flat anvil micrometers on the coating microstructure coupons, typically a 0.25 in cube or one square inch piece of 0.060 in or thicker sheet metal. The number of gun passes was recorded as determined by an automatic cycle counter during the spray operation. One cycle consisted of an initial st roke and its return st roke, thus two passes for any given position along the cylinder’s length. Microhardness The microhardness measurements were ta ken on a polished cross-section of the microstructure coupon. A sta ndard commercially availabl e diamond pyramid indentor and test machine were used with a 300 gm load. The specimen was sectioned parallel to the rotation direction and perp endicular to the gun traverse direction, then mounted and polished by standard procedures. Ten readings were taken in a standard pattern which diagonally traversed the coating thickness from the substrate to the surface and back to the substrate again taking care to maintain adequate distances from coating edges and between indentations.

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71 Rockwell 15N Superficial Surface Hardness This data was taken by making indentati ons directly on the coated coupon surface after lightly hand polishing the surface on 400 grit SiC metallographic polishing paper to remove any loosely adherent particles or sm all asperities. Tes ting was done with a standard commercial tester with a dial indi cator or automatic printout of the hardness reading which used a 120 degree diamond cone indentor and a 15 kg load. It is possible that at the coating thickness used, there is some influence of the substrate hardness on the absolute values obtained for the coating hardne ss, but since the substr ates were all IN718 typically of Rockwell C 40+/-2 and the coati ngs all .010+/-1 in thic kness, it was assumed that changes in hardness were reflections of process effect s. In general, the Rockwell 15N hardness showed minimal variability in most of the HVOF process studies so statistical analyses were not performed for these data. Tensile Bond Strength Tensile bond strengths were determined by a modified ASTM C633 method. The coatings were deposited on one inch diamet er buttons which were 0.25 in thick. This allows buttons to be conveniently mounted al ong with the other spec imens, or in the case of coating actual parts, often lends itself to locating buttons right on or adjacent to the part. The button s were then bonded between the normal ASTM C633 mandrels with FM 1000 film adhesive (Cytek) and tested per the specification. Th e HVOF WC-Co always resulted in ultimate failure of the film adhesive at greater than 12,000 psi.. Substrate Temperature The temperatures measured for the spra y process trials were taken by an IR pyrometers used on rotating specimens which were XY traversed. This has shown

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72 general agreement w ithin about +/25 oF on the absolute values with contact thermocouples and agreement in trends between individual process trials. Almen Strip Deflection The Almen strip was a sta ndard type N SAE 1070 st eel strip 3”x0.75”x0.030” held by 4 round head screws (see Mil-S-13165C for mo re details). The Almen strip was grit blasted on one side only and the arc height due to the i nduced bending stresses recorded before and after spraying of the HVOF coati ng. The difference in the two measurements was reported as the Almen strip deflecti on. A negative sign indicated a change representing a compressive stress in the co ating and a positive sign indicated a tensile stress in the coating. (Note: Shot peening work at GE ha s demonstrated the validity of taking the difference in curvature without the n eed to return the Almen strip to a perfectly flat condition; i.e. the deflec tion is cumulative in linear fash ion in the elastic region. The single side grit blast procedure typically result ed in a starting arc he ight of -0.004 in; i.e. compressive.) The Almen strip is restrained in the flat position duri ng coating. No effort was made to calculate actual coating stresses since the exact coating modulus was unknown.

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73 A factor (B,C) Combined Factors D factor Std.Ord Turn Table Robot Trav Hydrogen Oxygen Air Sp Dist Run No. RPM Sp mm/s psi/FMR psi/FMR psi, FMR inches 1 9 252 14.8 135 psi, 50.4 148 psi, 23.1 105 psi, 50.5 11.5 2 1 212 10.6 135 psi, 47.2 148 psi, 17.8 105 psi, 50.5 10 3 2 292 21.2 135 psi, 47.2 148 psi, 17.8 105 psi, 50.5 13 4 3 212 10.6 135 psi, 56.5 148 psi, 23.8 105 psi, 50.5 13 5 4 292 21.2 135 psi, 56.5 148 psi, 23.8 105 psi, 50.5 10 6 10 252 14.8 135 psi, 50.4 148 psi, 23.1 105 psi, 50.5 11.5 7 5 212 10.6 135 psi, 44.6 148 psi, 21.8 105 psi, 50.5 13 8 6 292 21.2 135 psi, 44.6 148 psi, 21.8 105 psi, 50.5 10 9 7 212 10.6 135 psi, 53.4 148 psi, 28.7 105 psi, 50.5 10 10 8 292 21.2 135 psi, 53.4 148 psi, 28.7 105 psi, 50.5 13 11 11 252 14.8 135 psi, 50.4 148 psi, 23.1 105 psi, 50.5 11.5

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74 StdOrder Mils/pass T 6 cy Norm Alm Porosity DPH 300 1 0.790 349 4.6 0.67 1166 2 0.408 236 4.1 0.50 984 3 0.830 347 6.5 0.50 1121 4 0.386 350 11.1 0.25 1163 5 0.870 264 2.9 0.75 925 6 0.407 281 6.1 0.25 1106 7 0.670 359 15.3 0.10 1177 8 0.420 261 10.7 0.37 1011 9 0.583 330 6.4 0.37 1127 CPs 10 0.558 295 7.5 0.37 1110 11 0.600 290 6.6 0.25 1131

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75 Design of Experiments Matrix Actual Run Orde r StdOrder Run No. A B C D 99.069 6 0 0 0 0 99.021 2 -1 -1 -1 -1 99.112 3 1 -1 -1 1 99.053 4 -1 1 -1 1 34 5 1 1 -1 -1 710 6R 0 0 0 0 85 7 -1 -1 1 1 46 8 1 -1 1 -1 99.127 9 -1 1 1 -1 99.098 10 1 1 1 1 99.1311 11 0 0 0 0 99.01not used 1 Differs from 6,7,13 99.10not used 10R Equals 10 Repeat 99.14not used 7R Equals 7 Repeat 99.15 1 1 1 -1 99.16 1 1 1 -1

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76 Design of Experiments Results Analyses (All broke % used as %in epoxy) DPH 300 Cycles Thk, mils Mils/pass Tmax T6 Almen gb Almen sp De lta Alm Norm Alm Porosity Porosity Tensile,psihardness 6 7.0 0.583 320 330 1.75 10.75 9. 0 6.4 .25-.5 0.37 125601127 5 7.9 0.790 344 349 0.25 7.5 7. 25 4.6 .5-.75 0.67 125131166 6 4.9 0.408 236 236 1.0 5.0 4.0 4.1 0.5 0.50 12475984 5 8.3 0.830 341 347 1.0 11.8 10.8 6.5 0.5 0.50 133711121 7 5.4 0.386 344 350 3.0 15.0 12.0 11.1 0.25 0.25 119991163 6 6.7 0.558 295 295 1.0 11.0 10. 0 7.5 .25-.5 0.37 119861110 5 8.7 0.870 264 264 1.5 6.5 5.0 2.9 0.75 0.75 11143925 7 5.7 0.407 273 281 1.0 8.0 7.0 6.1 0.25 0.25 110991106 5 6.7 0.670 353 359 1.5 22.0 20.5 15.3 0.1 0.10 122691177 5 4.2 0.420 253 261 0.0 9.0 9.0 10.7 .25-.5 0.37 111121011 6 7.2 0.600 290 290 1.5 11.0 9.5 6.6 0.25 0.25 117161131 6 7.2 0.600 278 278 0 4.8 4.8 3.3 0.5 123161073 6 4.4 0.367 248 248 1.1 10.6 9.5 10.8 .1-.25 129641083 5 8.8 0.880 258 264 1.0 6.5 5.5 3.1 .5-.75 11849974 8 5.2 0.325 322 322 1.5 1614.5 13.9 6 4.2 0.350 333 333 0 1212 14.3

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77 Almen vs Temp (CY 6)0.0 2.0 4.0 6.0 8.0 10.0 12.0 14.0 16.0 18.0 200220240260280300320340360380 Temp, Deg FNorm. N Almen # 8 # 1 # 2 # 5 # 4 # 9 # 10 # 11 # 6 # 3 # 7 Figure A-1. Almen strip deflection vs. substrate part temperature.

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78 Temp vs mils/pass200 220 240 260 280 300 320 340 360 380 0.3000.4000.5000.6000.7000.8000.900 mils/pasTemp cy6 # 4 # 5 # 9 # 10 # 11 # 7 # 1 # 3 # 6 # 8 # 2 Figure A-2. Substrate part temper ature vs. number of mils of coa ting deposited per pass of the torch.

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79 Hardness DPH 300 vs % Porosity800 850 900 950 1000 1050 1100 1150 1200 0.000.100.200.300.400.500.600.700.80 % PorosityDPH 300, kg/mm2 # 1 Figure A-3. Final coatin g hardness vs. % porosity.

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80 ANALYSES: !2 RUNS Run-1 mils/pass T cy6 Almen Porosity DPH 300 Term EFFE1 COEF1 EFFE2 COEF2 EFFE3 COEF3 EFFE4 COEF4 EFFE6 COEF6 Constant 0.597625 305.875 7.6625 0.42375 1081.63 A -0.38475 -0.19238 -47.75 -23.875 0.675 0.3375 -0.1625 -0.08125 -31.25 -15.62 B -0.04225 -0.02113 46.75 23.375 6.475 3.2375 -0.2375 -0.11875 72.75 36.37 C -0.01175 -0.00588 -29.25 -14.625 2.175 1.0875 -0.1125 -0.05625 -53.75 -26.87 D 0.06875 0.034375 -57.75 -28.875 -3.225 -1.6125 0.2125 0.10625 -142.75 -71.37 AB 0.03775 0.018875 0.25 0.125 -0.675 -0.3375 0.1725 0.08625 -30.75 -15.38 AC 0.02825 0.014125 7.25 3.625 -1.375 -0.6875 0.0475 0.02375 38.75 19.37 AD -0.05125 -0.02563 -9.25 -4.625 2.025 1.0125 -0.0275 -0.01375 5.75 2.88 Ct Pt -0.01729 -0.875 -0.82917 -0.09375 41.04

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81 Final Data Set Almen 11.5 1081.63 1081.63 1 -15.62 -15.62 1 36.37 36.37 1 -26.87 -26.87 0 -71.37 0 1 -15.38 -15.38 1 19.37 19.37 0 2.88 0 1079.5 0.42375 0.42375 1 -0.08125 -0.08125 1 -0.11875 -0.11875 1 -0.05625 -0.05625 0 0.10625 0 1 0.08625 0.08625 1 0.02375 0.02375 0 -0.01375 0 0.2775 7.6625 7.6625 1 0.3375 0.3375 1 3.2375 3.2375 1 1.0875 1.0875 -0.33 -1.6125 0.532125 1 -0.3375 -0.3375 1 -0.6875 -0.6875 -0.33 1.0125 -0.33413 11.498 305.875 305.875 1 -23.875 -23.875 1 23.375 23.375 1 -14.625 -14.625 0 -28.875 0 1 0.125 0.125 1 3.625 3.625 0 -4.625 0 294.5 0.597625 0.597625 1 -0.19238 -0.19238 1 -0.02113 -0.02113 1 -0.00588 -0.00588 0 0.034375 0 1 0.018875 0.018875 1 0.014125 0.014125 0 -0.02563 0 0.41125 Almen = 11.5; Hardness = 1079 DPH; Poro sity =0.277%; Max Substrate Temp=294oF; 0.4 mils per pass deposit rate.

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82 Hitemco Powder Lots vs L12 Lots0 20 40 60 80 100 120 0102030405060708090100 Powder Size, micronsCumulative % Diamalloy 2005 Lot 53792 Diamalloy 2005 Lot 54327 Diamalloy 2005 Lot 53791 Diamalloy 2005 Lot 54627 Diamalloy 2005 Lot 54480 Diamalloy 2005 Lot 53558 Stark 526.062 Lot 10362 Figure A-4. Powder size trends Powder size was also tracked based on a matrix of lots to determine if was a critical variable. Analysis displayed above shows that the deviation in particle size distribution for each powder lot was insignificant and not a factor in any of the DOE runs for parameter optimization. The combustion gas ratios, flow rates, a nd spray distance are th e major factors in the spray process. The response data for ha rdness, almen deflection (residual stress) and substrate temperature identified these as cr itical parameters for controlling final mechanical properties. The stoichiometry of the combustion fuels affects the melting of

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83 the cobalt binder and dissolution of the carbi des. Non-optimum gas ratios result in secondary carbide formation and reductions in hardness and toughness. Spray distance has a substantial affect on the residual stress state and substrate part temperature due to the amount of heat transfer be tween the molten particulate and the base metal. This also affects solidification rates and thus th e net porosity and volumetric changes. Equipment: Sulzer Metco Inc. Model 2600 hybrid gun Powder: Sulzer Metco Inc. Diamalloy 2005, 83%-WC 17%Co Agglomerated/Sintered Powder Feeder: Sulzer Metco Inc. Model Single 10, Positive Pressure, Screw Feed with Vibration. Powder Feed Rate: 8.5 lb/hr (325 rpm, 6 pitch feeder screw) Vibrator Setting: 30 Powder Carrier Gas: Nitrogen at 148 psi, and 28 scfh flow. Combustion Gas: Oxygen at 148 psi; 44+/-2 Flow Meter Reading (FMR) – Not in scfh Combustion Fuel: Hydrogen at 135 psi console pressure and 1229 scfh flow. Combustion Chamber Pressure: 100-102 psi Gun Cooling Water: 9.3 – 8.7 gph Water Temperature to the gun: In, 6472o F; Out, 117-125o F; Delta, 5154o F Specimen Rotation: 2,636 rpm for round bars (0.25 inch dia.) – 16,560 in/min surface speed. Gun Traverse Speed: 400 linear in/min for round bars Spray Distance: 11.5” Cooling Air: 2 gun mounted Air Jets at 14 inches, 90-110 psi 1 stationary Air Jet at 4-6 inches, 90-110 psi Part Temperature: Max temperature as read with a infrared pyrometer is 275 deg F trailing the plume spot. Injector Number 9 Insert Number 9 Shell Number 9 Siphon Plug Number 9 Air Cap 2701

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APPENDIX B FATIGUE DATA FOR 4340, 300M, AND AERMET 100

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85 4340, SMALL HOURGLASS SPECIMEN (0.003" COATING) R = -1, AIR100.0 110.0 120.0 130.0 140.0 150.0 160.0 170.0 180.0 190.0 200.0 1.E+031.E+041.E+051.E+061.E+071.E+08CYCLES TO FAILURE, NfENGINEERING STRESS MAX, KSI EHC/Peened EHC/Peened/FIT WCCo/Peened WCCo/Peened/FIT

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86 300M, SMALL HOURGLASS SPECIMEN (0.003" COATING) R = -1, AIR100 110 120 130 140 150 160 170 180 190 200 1.E+031.E+041.E+051.E+061.E+071.E+08CYCLES TO FAILURE, NfENGINEERING STRESS MAX, KSI Bare/UNPeened EHC/Peened EHC/Peened/FIT WCCo/Peened WCCO/Peened/FIT

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87 A 100, SMALL HOURG LASS SPECIMEN (0.003" COATING) R = -1, AIR90 100 110 120 130 140 150 160 170 180 1901.E+031.E+041.E+051.E+061.E+071.E+08CYCLES TO FAILURE, NfENGINEERING STRESS MAX, KSI Bare/UNPeened Bare/UNPeened/FIT EHC/Peened EHC/Peened/FIT WCCo/Peened WCCo/Peened/FIT

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882360-70249-27 ( Small Hour Glass ) Room Temperature Sine Frequency Vary Test Specimen Assigned Nf Results Fr equency Test Test Failure Failure Number Number Condition Cycles Hours Frame origin Location ( ksi max ) ( ) end of specimen R = -1.0 4340 coated with WCCo and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 128 01--43 125.0 4,144,354Failure Gage 29 hz 39.7 60046 1.72" sub-surface 138 01--47 125.0 10,005,992Removal 29 hz 95.8 60046 Removal 145 01--51 125.0 2,109,898Failure Gage 29 hz 20.2 60046 1.80 sub-surface 151 01--55 125.0 8,111,814Failure Gage 29 hz 27.7 60046 1.75" sub-surface 163 01--59 125.0 10,009,109Removal 29 hz 95.9 60046 Removal 221 01--44 110.0 10,005,254Removal 29 hz 95.8 60046 Removal 228 01--48 110.0 10,000,057Removal 29 hz 95.8 60046 Removal 233 01--52 110.0 10,163,865Removal 29 hz 97.4 60046 Removal 251 01--56 110.0 10,001,713Removal 29 hz 95.8 60055 Removal 254 01--60 110.0 10,031,800Removal 29 hz 96.1 60055 Removal R = .10 4340 coated with WCCo and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 194 01--61 220.0 50,779Failure Gage 59 HZ 0.2 60046 1.80" multi-surface 202 01--65 220.0 49,343Failure Gage 59 HZ 0.2 60055 1.70" surface 203 01--69 220.0 45,586Failure Gage 59 HZ 0.2 60055 1.70" surface 204 01--73 220.0 41,472Failure Gage 59 HZ 0.2 60055 1.76" surface 205 01--77 220.0 32,302Failure Gage 59 HZ 0.2 60055 1.78" surface

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89144 01--62 190.0 5,502,426Failure Ga ge 59 HZ 25.9 60052 1.80" surface 150 01--66 190.0 5,877,207Failure Ga ge 59 HZ 27.7 60052 1.80" surface 152 01--70 190.0 2,091,209Failure Gage 59 HZ 9.8 60052 1.68" surface 165 01--74 190.0 59,966Failure Gage 59 HZ 0.3 60055 1.80 sub-surface 166 01--78 190.0 118,312Failure Gage 59 HZ 0.6 60055 1.75" sub-surface 177 01--63 165.0 10,000,180Removal 59 HZ 47.1 60055 Removal 186 01--67 165.0 10,140,923Removal 59 HZ 47.7 60052 Removal 187 01--71 165.0 13,763,440Removal 59 HZ 64.8 60055 Removal 188 01--75 165.0 179,017Failure Gage 59 HZ 0.8 60046 1.78" surface 189 01--79 165.0 10,006,148Removal 59 HZ 47.1 60046 Removal 235 01--64 150.0 10,003,094Removal 59 HZ 47.1 60052 Removal 249 01--68 150.0 10,042,559Removal 59 HZ 47.3 60052 Removal 250 01--72 150.0 10,000,138Removal 59 HZ 47.1 60052 Removal 252 01--76 150.0 10,128,246Removal 59 HZ 47.7 60052 Removal 253 01--80 150.0 10,000,149Removal 59 HZ 47.1 60052 Removal R = -1.0 4340 coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Un-peened 53 01--101 175 5,260Failure Gage 5 hz 0.3 60055 1.80 multi-surface 54 01--105 175 5,852Failure Gage 5 hz 0.3 60055 1.80 multi-surface 55 01--109 175 6,466Failure Gage 5 hz 0.4 60052 1.80 multi-surface 56 01--113 175 6,435Failure Gage 5 hz 0.4 60055 1.80 multi-surface 57 01--117 175 5,197Failure Gage 5 hz 0.3 60052 1.80 multi-surface 75 01--102 150 15,090Failure Gage 5 hz 0.8 60052 1.80" surface 76 01--106 150 16,834Failure Gage 5 hz 0.9 60055 1.80" surface 77 01--110 150 18,371Failure Gage 5 hz 1.0 60052 1.80" surface 78 01--114 150 16,426Failure Gage 5 hz 0.9 60055 1.77" multi-surface 79 01--118 150 12,711Failure Gage 5 hz 0.7 60052 1.80 multi-surface

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90 127 01--103 125 55,781Failure Gage 29 hz 0.5 60052 1.80" surface 129 01--107 125 53,413Failure Gage 29 hz 0.5 60052 1.80" surface 130 01--111 125 51,649Failure Gage 29 hz 0.5 60052 1.75" surface 131 01--115 125 39,252Failure Gage 29 hz 0.4 60055 1.72" surface 132 01--119 125 43,374Failure Gage 29 hz 0.4 60052 1.75" surface 133 01--104 110 121,017Failure Gage 29 hz 1.2 60055 1.75" surface 134 01--108 110 183,132Failure Gage 29 hz 1.8 60052 1.70" surface 135 01--112 110 10,024,747Removal 29 hz 96.0 60055 na Removal 136 01--116 110 10,027,264Removal 29 hz 96.0 60052 na Removal 137 01--120 110 139,047Failure Gage 29 hz 1.3 60046 1.70" surface R = -1.0 4340 coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Un-peened 190 01--121 220.0 8,683Failure Gage 59 HZ 0.1 60052 1.80" surface 206 01--125 220.0 9,010Failure Gage 59 HZ 0.1 60055 1.75" surface 208 01--129 220.0 9,957Failure Gage 59 HZ 0.1 60055 1.80" mutli-surface 210 01--133 220.0 8,754Failure Gage 59 HZ 0.1 60055 1.80" surface 211 01--137 220.0 8,910Failure Gage 59 HZ 0.1 60055 1.80" mutli-surface 146 01--122 190.0 20,136Failure Gage 59 HZ 0.1 60052 1.78" mutli-surface 147 01--126 190.0 20,769Failure Gage 59 HZ 0.1 60052 1.80" surface 148 01--130 190.0 15,117Failure Gage 59 HZ 0.1 60052 1.80" surface 149 01--134 190.0 18,968Failure Gage 59 HZ 0.1 60052 1.80" mutli-surface 153 01--138 190.0 16,567Failure Gage 59 HZ 0.1 60052 1.75" mutli-surface 154 01--123 165.0 26,710Failure Gage 59 HZ 0.1 60052 1.75" surface 155 01--127 165.0 36,345Failure Gage 59 HZ 0.2 60052 1.80" surface 156 01--131 165.0 32,139Failure Gage 59 HZ 0.1 60052 1.80" mutli-surface 157 01--135 165.0 30,147Failure Gage 59 HZ 0.1 60052 1.80" surface

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91158 01--139 165.0 27,470Failure Gage 59 HZ 0.1 60052 1.78" surface 238 01--124 150.0 59,246Failure Gage 59 HZ 0.3 60052 1.80" surface 239 01--128 150.0 44,020Failure Gage 59 HZ 0.2 60052 1.75" surface 240 01--132 150.0 47,251Failure Gage 59 HZ 0.2 60052 1.80" surface 241 01--136 150.0 49,330Failure Gage 59 HZ 0.2 60052 1.80" surface 242 01--140 150.0 50,633Failure Gage 59 HZ 0.2 60052 1.80" surface R = -1.0 4340 coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 178 01--141 175.0 9,459Failure Gage 5 hz 0.5 60052 1.80" surface 73 01--145 175.0 10,747Failure Gage 5 hz 0.6 60052 1.80" mulit-surface 179 01--149 175.0 9,321Failure Gage 5 hz 0.5 60052 1.80" surface 74 01--153 175.0 9,019Failure Gage 5 hz 0.5 60055 1.80" mulit-surface 180 01--157 175.0 11,172Failure Gage 5 hz 0.6 60052 1.80" mulit-surface 184 01--142 150.0 68,476Failure Gage 5 hz 3.8 60052 1.80" surface 110 01--146 150.0 58,800Failure Gage 5 hz 3.3 60052 1.80" surface 185 01--150 150.0 74,437Failure Gage 5 hz 4.1 60052 1.78" surface 201 01--154 150.0 56,109Failure Gage 5 hz 3.1 60055 1.80" mulit-surface 112 01--158 150.0 55,078Failure Gage 5 hz 3.1 60055 1.77" surface 232 01--143 125.0 8,035,894Failure Gage 29 hz 77.0 60055 1.75" sub-surface 236 01--147 125.0 3,831,691Failure Gage 29 hz 36.7 60055 1.75" sub-surface 237 01--151 125.0 10,000,041Removal 29 hz 95.8 60055 Removal 248 01--155 125.0 10,234,077Removal 29 hz 98 60055 Removal 255 01--159 125.0 4,925,684Failure Gage 29 hz 47.2 60052 1.75" sub-surface 256 01--144 110.0 11,241,424Removal 29 hz 107.8 60052 Removal 257 01--148 110.0 10,015,843Removal 29 hz 95.9 60055 Removal 258 01--152 110.0 10,001,486Removal 29 hz 95.8 60052 Removal

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92259 01--156 110.0 1,748,245Failure Gage 29 hz 16.7 60055 1.80" sub-surface 260 01--160 110.0 7,688,528Failure Gage 29 hz 73.6 60055 1.80" sub-surface R = -1.0 4340 coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 213 01--161 220.0 16,527Failure Gage 59 hz 0.1 60055 1.80" multi-surface 214 01--165 220.0 15,895Failure Gage 59 hz 0.1 60055 1.80" multi-surface 215 01--169 220.0 12,879Failure Gage 59 hz 0.1 60055 1.80" surface 216 01--173 220.0 16,311Failure Gage 59 hz 0.1 60055 1.75" surface 217 01--177 220.0 16,418Failure Gage 59 hz 0.1 60055 1.80" surface 169 01--162 190.0 36,184Failure Gage 59 hz 0.2 60055 1.80" multi-surface 167 01--166 190.0 38,291Failure Gage 59 hz 0.2 60055 1.75" surface 168 01--170 190.0 41,430Failure Gage 59 hz 0.2 60055 1.80" surface 170 01--174 190.0 32,225Failure Gage 59 hz 0.2 60055 1.80" surface 171 01--178 190.0 27,904Failure Gage 59 hz 0.1 60055 1.80" surface 172 01--163 165.0 82,487Failure Gage 59 hz 0.2 60055 1.80" surface 173 01--167 165.0 73,961Failure Gage 59 hz 0.3 60055 1.80" surface 174 01--171 165.0 66,317Failure Gage 59 hz 0.3 60055 1.85" surface 175 01--175 165.0 74,226Failure Gage 59 hz 0.3 60055 1.80" surface 176 01--179 165.0 97,457Failure Gage 59 hz 0.5 60055 1.75" surface 243 01--164 150.0 10,033,497Removal 59 hz 47.2 60052 Removal 244 01--168 150.0 214,289Failure Gage 59 hz 1.0 60052 1.80" surface 245 01--172 150.0 110,845Failure Gage 59 hz 0.5 60052 1.80" surface 246 01--176 150.0 163,243Failure Gage 59 hz 0.8 60052 1.80" surface 247 01--180 150.0 12,965,904Removal 59 hz 61.0 60052 Removal R = -1.0 300 M non-coated Tested as 2360-70249-21 ( qty 20 ) not -27 code 3/4" Hydralic Hourglass Specimen Note: Smooth Specimen 4340 WCCo tested as -21 ( TN 19-21 SN 02-18 TN 20-21 SN 02-19 & TN 21-21 SN 02-20 )

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93 Un-peened 17 M1-17 180 9,686Failure Gage 5 hz 0.8 60052 1.75" surface 18 M1-18 180 12,876Failure Gage 5 hz 0.7 60052 1.80" multi-surface 22 M1-19 180 8,430Failure Gage 5 hz 0.5 60052 1.78" surface 16 M1-16 170 14,347Failure Gage 5 hz 0.8 60052 1.80" surface 15 M1-15 160 22,648Failure Gage 5hz 1.3 60052 1.75" surface 14 M1-14 150 21,734Failure Gage 5hz 1.2 60052 1.80" surface 11 M1-11 145 33,873Failure Gage 5hz 1.9 60052 1.75" surface 12 M1-12 140 61,866Failure Gage 5hz 3.4 60052 1.80" surface 13 M1-13 140 87,378Failure Gage 5hz 4.9 60052 1.75" surface 23 M1-20 140 74,231Failure Gage 5 hz 4.1 60052 1.80" surface 3 M1-03 130 822,677 Shank 29 hz 7.9 60052 Na surface 4 M1-04 130 263,063Failure Gage 59/29 hz 2.5 60052 1.70" surface 7 M1-07 130 267,886Failure Gage 29 hz 2.6 60052 1.75" surface 9 M1-10 130 2,855,038Failure Gage 29 hz 6.5 60052 1.70" surface 5 M1-05 125 2,249,183Failure Gage 59/29 hz 10.6 60052 1.80" surface 6 M1-06 125 356,096Failure Gage 29 hz 3.4 60052 1.80" surface 8 M1-08 125 4,913,210Failure Gage 29 hz 47.1 60052 1.70" surface 10 M1-09 125 6,334,285Failure Gage 29 hz 60.7 60052 1.80" sub-surface 2 M1-02 120 8,590,400Failure Gage 29 hz 82.3 60052 1.80" sub-surface 1 M1-01 100 1,532,833 Shank 59 hz 7.2 60052 Na surface Note 1 : ( MI-19 & M1-20 ) Decision was made in late February 00 to run these two tests at 180.0 ksi and 140.0 ksi ( PEB / TAC / DCW ) R = -1.0 300 M coated with WCCo and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Un-peened 265 M1-21 180.0 8,515Failure 5 0.5 60055 1.80" surface 268 M1-25 180.0 14,610Failure 5 0.8 60052 1.80" surface

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94269 M1-29 180.0 8,758Failure 5 0.5 60055 1.80" surface 270 M1-33 180.0 10,592Failure 5 0.6 60052 1.80" mutli-surface 271 M1-37 180.0 11,460Failure 5 0.6 60055 1.75" surface 285 M1-22 140.0 660,790Failure 5 36.7 60055 1.75" sub-surface 292 M1-26 140.0 377,499Failure 5 21.0 60055 1.75" surface 294 M1-30 140.0 432,594Failure 5 24.0 60055 1.80" surface 297 M1-34 140.0 396,417Failure 5 22.0 60055 1.80" surface 299 M1-38 140.0 408,708Failure 5 22.7 60055 1.80" sub-surface 261 M1-23 130.0 1,359,296Failure 29 13.0 60055 1.80" sub-surface 262 M1-27 130.0 1,371,314Failure 29 13.1 60052 1.75" surface 291 M1-31 130.0 1,227,526Failure 29 11.8 60052 1.80" sub-surface 293 M1-35 130.0 2,093,314Failure 29 20.1 60052 1.75" sub-surface 295 M1-39 130.0 3,109,224Failure 29 29.8 60052 1.80" sub-surface 309 M1-24 125.0 3,372,762Failure 59 15.9 60055 1.75" sub-surface 313 M1-28 125.0 1,872,153Failure 59 8.8 60055 1.75" sub-surface 315 M1-32 125.0 1,659,482Failure 59 7.8 60072 1.75" sub-surface 316 M1-36 125.0 3,354,132Failure 59 15.8 60052 1.80" sub-surface 317 M1-40 125.0 2,098,302Failure 59 9.9 60072 1.75" sub-surface R = -1.0 300 M coated with WCCo and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 3 M1-41 180 16,675Failure Gage 5 hz 0.9 60050 1.75" surface 9 M1-45 180 13,305Failure Gage 5 hz 0.7 60047 1.75" multi-surface 10 M1-49 180 16,019Failure Gage 5 hz 0.9 60047 1.75" surface 12 M1-53 180 11,928Failure Gage 5 hz 0.7 60047 1.75" surface 13 M1-57 180 12,039Failure Gage 5 hz 0.7 60047 1.80" multi-surface

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955 M1-42 140 694,702Failure Gage 5 hz 38.6 60050 1.75" sub-surface 15 M1-46 140 550,425Failure Gage 5 hz 30.6 60047 1.70" sub-surface 19 M1-50 140 532,624Failure Gage 5 hz 29.6 60047 1.80" sub-surface 20 M1-54 140 351,970Failure Gage 5 hz 19.6 60050 1.75" sub-surface 28 M1-58 140 405,551Failure Gage 5 hz 22.5 60050 1.75" sub-surface 18 M1-43 130 4,075,390Failure Gage 5 hz 39.0 60046 1.80" sub-surface 21 M1-47 130 2,618,078Failure Gage 29 hz 25.1 60050 1.75" sub-surface 29 M1-51 130 3,400,427Failure Gage 29 hz 32.6 60050 1.75" sub-surface 30 M1-55 130 1,721,840Failure Gage 29 hz 16.5 60050 1.75" sub-surface 31 M1-59 130 1,926,256Failure Gage 29 hz 18.3 60050 1.80" sub-surface 16 M1-44 125 5,981,633Failure Gage 59 hz 29.3 60046 1.80" sub-surface 37 M1-48 125 1,898,601Failure Gage 59 hz 8.9 60050 1.75" sub-surface 38 M1-52 125 3,050,916Failure Gage 59 hz 14.4 60050 1.75" sub-surface 39 M1-56 125 3,455,400Failure Gage 59 hz 16.3 60050 1.80" sub-surface 40 M1-60 125 7,571,921Failure Gage 59 hz 35.6 60050 1.80" sub-surface R = -1.0 300 M coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Un-peened 64 M1-81 180 8,009Failure Gage 5 hz 0.4 60055 1.80" surface 63 M1-85 180 6,654Failure Gage 5 hz 0.4 60052 1.80" multi-surface 65 M1-89 180 9,303Failure Gage 5 hz 0.5 60052 1.80" multi-surface 66 M1-93 180 7,035Failure Gage 5 hz 0.4 60055 1.80" surface 67 M1-97 180 23,464Failure Gage 5 hz 1.3 60052 1.78" surface 80 M1-82 140 22,862Failure Gage 5 hz 1.3 60055 1.80" multi-surface 81 M1-86 140 59,143Failure Gage 5 hz 3.3 60052 1.80" surface 82 M1-90 140 24,989Failure Gage 5 hz 1.4 60055 1.80" surface 97 M1-94 140 22,115Failure Gage 5 hz 1.2 60055 1.77" surface 98 M1-98 140 37,476Failure Gage 5 hz 2.1 60055 1.78" surface

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96 94 M1-83 130 58,560Failure Gage 29 hz 0.6 60046 1.73" surface 96 M1-87 130 1,202,961Failure Gage 29 hz 11.5 60046 1.80" sub-surface 99 M1-91 130 61,028Failure Gage 29 hz 0.6 60046 1.80" surface 101 M1-95 130 49,290Failure Radius 29 hz 0.5 60046 1.65" surface 104 M1-99 130 80,701Failure Gage 29 hz 0.8 60046 1.75" surface 60 M1-84 125 8,375,010Failure Gage 59 hz 39.4 60046 1.80" surface 86 M1-88 125 144,764Failure Gage 59 hz 0.7 60046 1.75" surface 88 M1-92 125 80,529Failure Gage 59 hz 0.4 60046 1.80" surface 90 M1-96 125 154,827Failure Radius 59 hz 0.7 60046 1.60" surface 93 M1-100 125 122,823Failure Gage 59 hz 0.6 60046 1.77" surface Test Note 1: MI-96 The failure was not in the minimum diameter. Test Note 2: MI-95 The failure was not in the minimum diameter. R = -1.0 300 M coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 48 M1-101 180 18,357F, G S 5 hz 1.0 60052 1.80" surface 182 M1-105 180 17,022F, G S 5 hz 0.9 60052 1.80" surface 49 M1-109 180 18,510F, G S 5 hz 1.0 60052 1.80" surface 183 M1-113 180 16,408F, G S 5 hz 0.9 60052 1.80" surface 52 M1-117 180 17,087F, G S 5 hz 0.9 60052 1.80" surface 89 M1-102 140 324,709F, G S 5 hz 18.0 60052 1.75" sub-surface M1-106 140 1.80" sub-surface 92 M1-110 140 265,489F, G S 5 hz 14.7 60055 1.80" sub-surface 91 M1-114 140 705,255F, G S 5 hz 39.2 60052 1.75" sub-surface 218 M1-118 140 759,976F, G S 5 hz 42.2 60055 1.75" dob. sub-surface 212 M1-103 130 1,693,639F, G S 29 hz 16.2 60052 1.80" sub-surface

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97106 M1-107 130 1,134,818F, G S 29 hz 10.8 60046 1.80" sub-surface 113 M1-111 130 1,986,436F, G S 29 hz 19.0 60046 1.75" sub-surface 124 M1-115 130 1,162,639F, G S 29 hz 11.1 60046 1.80" sub-surface 126 M1-119 130 1,560,383F, G S 29 hz 14.9 60055 1.75" sub-surface 223 M1-104 125.0 3,609,869Failure Gage 59 hz 17.0 60055 1.80" sub-surface 224 M1-108 125.0 4,666,345Failure Gage 59 hz 22.0 60052 1.80" sub-surface 225 M1-112 125.0 5,443,061Failure Gage 59 hz 25.6 60055 1.70" sub-surface 42 M1-116 125 2,398,806F, G S 59 hz 11.3 60046 1.80" sub-surface 45 M1-120 125 3,101,098F, G S 59 hz 14.6 60046 1.80" sub-surface R = -1.0 A100 ( uncoated ) Tested as 2360-70249-22 ( qty 20 ) not -27 code 3/4" Hydralic Hourglass Specimen Un-peened 5 A1-05 160 32,610step 5 hz 1.8 60050 5--a A1-05 180 4,568Failure Gage 5 hz 0.3 60050 1.80" multi-surface 7 A1-07 180 13,948Failure Gage 5 hz 0.8 60050 1.80" surface 16 A1-16 180 19,710Failure Gage 5 hz 1.1 60050 1.80" multi-surface 20 A1-20 180 11,915Failure Gage 5 hz 0.7 60052 1.80" multi-surface 6 A1-06 175 18,889Failure Gage 5 hz 1 60050 1.80" surface 3 A1-03 150 52,437Failure Gage .5/5hz 20 60050 1.75" surface 4 A1-04 150 48,887Failure Gage 5 hz 2.7 60050 1.80" surface 17 A1-17 150 65,753Failure Gage 5 hz 3.7 60050 1.75" surface 12 A1-12 145 154,621Failure Gage 29 hz 1.5 60050 1.80" surface 13 A1-13 145 158,756Failure Gage 29 hz 1.5 60050 1.80" surface 18 A1-18 145 961,327Failure Gage 29 hz 9.2 60050 1.80" surface 10 A1-10 140 613,137Failure Gage 29 hz 6.4 60050 1.80" surface 11 A1-11 140 872,507Failure Gage 29 hz 8.4 60050 1.75" surface 14 A1-14 135 697,229Failure Gage 29 hz 6.7 60050 1.80" surface 15 A1-15 135 253,315Failure Gage 29 hz 2.4 60050 1.80" surface 19 A1-19 135 6,799,993Failure Gage 29 hz 65.1 60052 1.80" surface

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982 A1-02 130 1,500,347 Shoulder 29 hz 14.4 60050 Shoulder 9 A1-09 130 9,081,697Failure Gage 59 hz 42.8 60050 1.80" sub-surface 8 A1-08 120 10,107,505Removal 59 hz 47.6 60050 Removal 1 A1-01 100 10,030,040Removal 59 hz 47.2 60050 Removal Note : 1 : A1-20 decision to test at 180.0 ksi ( PEB / DCW ) 03-02-00. Test Note : A1-09 assigned at + / 130.0 ksi Gradual stress increase occurred around 1,500,000 cycles to 133 ksi min Testing continued at + 130 ksi max / 133 ksi min R = -1.0 A100 coated with WCCo and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Un-peened 263 A1-21 180.0 26,600Failure 5 1.5 60052 1.77" surface 264 A1-25 180.0 14,944Failure 5 0.8 60055 1.70" surface 276 A1-29 180.0 18,442Failure 5 1.0 60055 1.77" surface 277 A1-33 180.0 21,840Failure 5 1.2 60052 1.75" multi-surface 279 A1-37 180.0 23,815Failure 5 1.3 60055 1.77" surface 286 A1-22 150.0 239,823Failure 5 13.3 60052 1.80" surface 287 A1-26 150.0 55,111Failure 5 3.1 60052 1.80" surface 288 A1-30 150.0 205,010Failure 5 11.4 60055 1.75" surface 289 A1-34 150.0 57,927Failure 5 3.0 60052 1.80" surface 290 A1-38 150.0 272,468Failure 5 15.1 60055 1.75" surface 301 A1-23 145.0 399,159Failure 29 3.8 60052 1.77" surface 302 A1-27 145.0 99,792Failure 29 1.0 60055 1.80" surface 303 A1-31 145.0 81,780Failure 29 0.8 60055 1.77" surface 304 A1-35 145.0 69,042Failure 29 0.7 60055 1.75" surface 305 A1-39 145.0 1,817,100 Shank 29 17.4 60055 Na Na 306 A1-24 135.0 873,490Failure 29 8.4 60052 1.80" surface 310 A1-28 135.0 102,901Failure 29 1.0 60052 1.80" surface 311 A1-32 135.0 917,391Failure 29 8.8 60052 1.77" surface 312 A1-36 135.0 6,174,376Failure 29 59.1 60052 1.80" sub-surface

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99314 A1-40 135.0 3,011,527Failure 29 28.8 60052 1.80" sub-surface R = -1.0 A100 coated with WCCo and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 2 A1-41 180 22,662Failure Gage 5 hz 1.3 60052 1.80" surface 1 A1-45 180 43,139Failure Gage 5 hz 2.4 60047 1.80" surface 6 A1-49 180 50,507Failure Gage 5 hz 2.8 60047 1.80" surface 11 A1-53 180 38,557Failure Gage 5 hz 2.1 60052 1.80" surface 14 A1-57 180 59,733Failure Gage 5 hz 3.3 60052 1.80" surface 4 A1-42 150 351,135Failure Gage 5 hz 19.5 60052 1.80" surface 8 A1-46 150 104,695Failure Gage 5 hz 5.8 60052 1.80" surface 7 A1-50 150 342,458Failure Gage 5 hz 19.0 60047 1.80" sub-surface 17 A1-54 150 1,130,661Failure Gage 5 hz 62.8 60052 1.80" sub-surface 24 A1-58 150 1,544,045Failure Gage 5 hz 85.8 60052 1.80" sub-surface 22 A1-43 145 4,003,977Failure Gage 29 hz 38.4 60050 1.80" sub-surface 23 A1-47 145 2,665,181Failure Gage 29 hz 25.5 60050 1.75" sub-surface 25 A1-51 145 2,300,418Failure Gage 29 hz 22.0 60050 1.80" sub-surface 26 A1-55 145 3,956,070Failure Gage 29 hz 37.9 60050 1.80" surface 27 A1-59 145 3,567,031Failure Gage 29 hz 34.2 60050 1.80" sub-surface 32 A1-44 135 6,551,857Failure Gage 29 hz 62.8 60050 1.80" sub-surface 33 A1-48 135 4,874,457Failure Gage 29 hz 46.7 60050 1.80" sub-surface 34 A1-52 135 5,159,358Failure Gage 29 hz 49.4 60050 1.80" sub-surface 35 A1-56 135 4,990,165Failure Gage 29 hz 47.8 60050 1.80" sub-surface 36 A1-60 135 5,603,160Failure Gage 29 hz 53.7 60050 1.80" sub-surface R = -1.0 A100 coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Un-peened 68 A1-81 180 9,439Failure Gage 5 hz 0.5 60055 1.80" surface 69 A1-85 180 7,372Failure Gage 5 hz 0.4 60052 1.80" surface

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10070 A1-89 180 9,117Failure Gage 5 hz 0.5 60055 1.80" surface 71 A1-93 180 8,193Failure Gage 5 hz 0.5 60052 1.80" surface 72 A1-97 180 8,164Failure Gage 5 hz 0.5 60055 1.80" surface 103 A1-82 150 21,401Failure Gage 5 hz 1.2 60055 1.80" multi-surface 105 A1-86 150 17,872Failure Gage 5 hz 1.0 60055 1.85" multi-surface 107 A1-90 150 14,819Failure Gage 5 hz 0.8 60055 1.80" multi-surface 108 A1-94 150 18,147Failure Gage 5 hz 1.0 60055 1.80" surface 109 A1-98 150 16,262Failure Gage 5 hz 0.9 60055 1.80" surface 114 A1-83 145 26,872 Failure Gage 29 hz 0.3 60055 1.80" surface 116 A1-87 145 28,001Failure Gage 29 hz 0.3 60055 1.75" surface 117 A1-91 145 27,614Failure Gage 29 hz 0.3 60055 1.75" surface 118 A1-95 145 20,516Failure Gage 29 hz 0.2 60055 1.80" surface 119 A1-99 145 18,708Failure Gage 29 hz 0.2 60055 1.80" surface 120 A1-84 135 48,865Failure Gage 29 hz 0.5 60055 1.80" multi-surface 121 A1-88 135 21,045Failure Gage 29 hz 0.2 60055 1.80" surface 122 A1-92 135 43,801Failure Gage 29 hz 0.4 60055 1.80" surface 123 A1-96 135 39,622Failure Gage 29 hz 0.4 60055 1.80" surface 125 A1-100 135 19,279Failure Gage 29 hz 0.2 60055 1.80" surface Test Note : A1-83 assigned at + / 145.0 ksi. Overload occurred upon start-up Testing continued at + / 150.0 ksi R = -1.0 A100 coated with EHC and ground to 0.003" thick Environment Air 3/4" Hydralic Hourglass Specimen Peened 58 A1-101 180.0 20,683Failure Gage 5 hz 1.1 60052 1.80" muti-surface 59 A1-105 180.0 16,490Failure Gage 5 hz 0.9 60055 1.80" surface 181 A1-109 180.0 22,441Failure Gage 5 hz 1.2 60052 1.80" muti-surface 61 A1-113 180.0 19,956Failure Gage 5 hz 1.1 60052 1.80" muti-surface 62 A1-117 180.0 15,475Failure Gage 5 hz 0.9 60055 1.80" surface

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10183 A1-102 150.0 76,983Failure Gage 5 hz 4.3 60055 1.80" surface 84 A1-106 150.0 270,555Failure Gage 5 hz 15.0 60052 1.80" surface 193 A1-110 150.0 76,340Failure Gage 5 hz 4.2 60055 1.75" sub-surface 85 A1-114 150.0 85,908Failure Gage 5 hz 4.8 60055 1.80" surface 87 A1-118 150.0 72,406Failure Gage 5 hz 4.0 60055 1.80" surface 111 A1-103 145.0 1,353,375Failure Gage 29 hz 13.0 60052 1.80" sub-surface 192 A1-107 145.0 71,563Failure Gage 29 hz 0.7 60055 1.63" surface 197 A1-111 145.0 1,447,664Failure Gage 29 hz 13.9 60046 1.80" surface 115 A1-115 145.0 2,426,207Failure Gage 29 hz 23.2 60052 1.75" sub-surface 207 A1-119 145.0 64,695Failure Gage 29 hz 0.6 60052 1.80" surface 191 A1-104 135.0 4,280,726Failure Gage 29 hz 41.0 60052 1.78" sub-surface 209 A1-108 135.0 4,686,027Failure Gage 29 hz 44.9 60046 1.80" sub-surface 219 A1-112 135.0 2,450,433Failure Gage 29 hz 23.5 60046 1.75" surface 227 A1-116 135.0 5,073,914Failure Gage 29 hz 48.6 60055 1.80" sub-surface 230 A1-120 135.0 2,319,412Failure Gage 29 hz 22.2 60055 1.72" surface

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APPENDIX C RESIDUAL STRESS SCAN DATA

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103 Data for the shot-peened 4340 steel substrate. Filename Observed Std dev Peak FWHM C11962 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 116.36670.011710.021.5555 0.671 OO2 0 -42 116.16730.010511.831.6213 0.448 OO3 0 -28.2 115.96330.010212.791.6252 0.223 OO4 0 0 115.77050.0093141.6726 0.000 OO5 0 28.2 115.97770.009513.611.6745 0.223 OO6 0 42 116.15920.011811.111.5783 0.448 OO7 0 55 116.35520.01787.661.7883 0.671 Filename Observed Std dev Peak FWHM C11962 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 116.440.01299.761.7671 0.671 OO9 45 -42 116.20130.016511.471.7119 0.448 O10 45 -28.2 116.03880.010612.831.5885 0.223 O11 45 0 115.82530.013311.361.665 0.000 O12 45 28.2 116.02310.014810.191.6168 0.223 O13 45 42 116.18740.0158.261.6296 0.448 O14 45 55 116.33640.01715.961.6284 0.671 Filename Observed Std dev Peak FWHM C11962 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 116.52050.06676.251.5397 0.671 O16 90 -42 116.33260.01738.711.6498 0.448 O17 90 -28.2 116.11580.01679.471.7796 0.223 O18 90 0 115.84540.013710.491.7054 0.000 O19 90 28.2 116.01490.013710.971.6664 0.223 O20 90 42 116.23510.01679.231.8909 0.448 O21 90 55 116.38840.0167.591.8203 0.671

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104 Stress calculations for the s hot-peened 4340 steel substrate. Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.9075 116.166 0.004 0.0000167.79E-10 PHI a1 slope a1 slope Y-I=Eps33CORR. COEFF. 0 Eps11-Eps33 -0.0047830.0021511.000 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0045050.0018220.999 90 Eps22-Eps33 -0.0049710.0016920.998 ave 0.001888 stdev 0.000237 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.000014-0.0000060.201 45 0.707*(Eps13+Eps23) 0.000133-0.0000020.485 90 Eps23 0.0003180.0000040.957 Eps13 -0.000129 Eps23 0.000175 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.457 psi* 42.552 strain* -3.68E-05 d0* 0.9075

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105 Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 2.07E+11 Poisson's ratio 0.291 S2/2 (1+nu)/E zzzzzz2 6.24E-12 S1 -nu/E zzzzzz1 -1.41E-12 -nu/(1-2nu)QUE -0.696172 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-axial Sig12 Sig13 Eps22 Eps23 Sig22-hoopSig23 Eps33 Sig33-radial -0.002895 0.000372 -0.000129 -921 60 -21 -0.003082 0.000318 -951 51 0.001888 -154 Std Dev. Std Dev. 0.000265 0.000416 0.000131 98 67 21 0.000586 0.000350 163 56 0.000102 77 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -921 60 -21-0.0028950.000372-0.000129 -951 51 -0.0030820.000318 -154 0.001888

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106 Data for the shot-peened and gr it blasted 4340 steel substrate. Filename Observed Std dev Peak FWHM C11962 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 116.47050.01598.241.7028 0.671 OO2 0 -42 116.27110.012411.251.7476 0.448 OO3 0 -28.2 116.03440.010213.811.6241 0.223 OO4 0 0 115.82610.015112.831.658 0.000 OO5 0 28.2 116.06330.012111.791.7161 0.223 OO6 0 42 116.5065016.231.6831 0.448 OO7 0 55 116.7351013.521.8078 0.671 Filename Observed Std dev Peak FWHM C11962 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 116.31020.01529.731.6512 0.671 OO9 45 -42 116.11250.015411.081.628 0.448 O10 45 -28.2 115.91220.014610.551.5785 0.223 O11 45 0 115.75280.02438.981.6378 0.000 O12 45 28.2 116.08570.01757.341.5662 0.223 O13 45 42 116.25460.01716.441.7207 0.448 O14 45 55 116.38110.01825.291.6777 0.671 Filename Observed Std dev Peak FWHM C11962 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 116.25770.03245.561.6795 0.671 O16 90 -42 116.1470.02467.061.7003 0.448 O17 90 -28.2 115.93940.02397.841.7098 0.223 O18 90 0 115.82510.02347.981.5878 0.000 O19 90 28.2 116.11020.02057.341.7652 0.223 O20 90 42 116.23110.01717.891.3575 0.448 O21 90 55 116.29020.01417.71.0964 0.671

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107 Stress calculations for the shot-peened and grit blasted 4340 steel substrate. C11962 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.9072 116.227 0.004 0.0000167.76E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 -0.0064790.002237 0.996 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0047840.002461 0.995 90 Eps22-Eps33 -0.0036820.002047 0.985 ave 0.002248 stdev 0.000208 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 -0.0005830.000047 0.737 45 0.707*(Eps13+Eps23) -0.000355-0.000017 0.784 90 Eps23 -0.000241-0.000029 0.552 Eps13 -0.000262 Eps23 0.000081

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108 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 2.07E+11 Poisson's ratio 0.291 S2/2 (1+nu)/E zzzzzz2 6.24E-12 S1 -nu/E zzzzzz1 -1.41E-12 -nu/(1-2nu)QUE -0.696172 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-axial Sig12 Sig13 Eps22 Eps23 Sig22-hoopSig23 Eps33 Sig33-radial -0.004231 0.000297 -0.000262 -1060 48 -42 -0.001434 -0.000241 -611 -39 0.002248 -21 Std Dev. Std Dev. 0.000263 0.000416 0.000129 88 67 21 0.000449 0.000260 126 42 0.000101 64 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -1060 48 -42-0.0042310.000297-0.000262 -611 -39 -0.001434-0.000241 -21 0.002248

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109 Data for the as-sprayed WC17%Co coated specimen (117o 2-theta. 211 peak). Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 117.36810.005826.950.52 0.671 OO2 0 -42 117.38980.005928.440.607 0.448 OO3 0 -28.2 117.38750.006329.640.614 0.223 OO4 0 0 117.38120.006728.560.5819 0.000 OO5 0 28.2 117.35490.007727.420.6733 0.223 OO6 0 42 117.3440.005532.960.5529 0.448 OO7 0 55 117.34360.006126.360.56 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 117.35040.006624.80.6143 0.671 OO9 45 -42 117.38350.00724.360.6767 0.448 O10 45 -28.2 117.39580.009324.170.7521 0.223 O11 45 0 117.41770.008823.750.7087 0.000 O12 45 28.2 117.37720.009220.740.7022 0.223 O13 45 42 117.38440.010915.140.6628 0.448 O14 45 55 117.3880.010114.150.6353 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 117.38030.008417.910.5378 0.671 O16 90 -42 117.37920.007621.930.5977 0.448 O17 90 -28.2 117.3710.00725.740.5794 0.223 O18 90 0 117.40830.008722.190.6324 0.000 O19 90 28.2 117.38710.00822.510.6491 0.223 O20 90 42 117.36610.011217.430.6646 0.448 O21 90 55 117.36820.006918.380.5202 0.671

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110 Stress calculations for the as-sprayed WC-17%Co coated specimen (117o 2-theta. 211 peak). C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.9021 117.276 0.004 0.0000167.37E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 0.000191-0.000557 0.988 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)0.000351-0.000719 0.937 90 Eps22-Eps33 0.000257-0.000657 0.838 ave -0.000644 stdev 0.000082 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.000100-0.000001 0.911 45 0.707*(Eps13+Eps23) -0.0000280.000006 0.212 90 Eps23 0.000019-0.000007 0.246 Eps13 -0.000059 Eps23 -0.000140

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111 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 6.39E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.94E-12 S1 -nu/E zzzzzz1 -3.76E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-axial Sig12 Sig13 Eps22 Eps23 Sig22-hoopSig23 Eps33 Sig33-radial -0.000454 0.000127 -0.000059 -587 65 -31 -0.000387 0.000019 -553 10 -0.000644 -685 Std Dev. Std Dev. 0.000172 0.000221 0.000064 138 114 33 0.000189 0.000081 149 42 0.000066 79 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -587 65 -31-0.0004540.000127-0.000059 -553 10 -0.0003870.000019 -685 -0.000644

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112 Data for the as-sprayed WC17%Co coated specimen (121o 2-theta. 103 peak). Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 120.94220.011614.550.5547 0.671 OO2 0 -42 120.94260.007919.840.473 0.448 OO3 0 -28.2 120.96410.011716.090.6022 0.223 OO4 0 0 120.97780.011117.770.6323 0.000 OO5 0 28.2 120.95080.010219.540.553 0.223 OO6 0 42 120.93590.010917.70.5261 0.448 OO7 0 55 120.9460.010915.620.5628 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 120.94740.016111.720.8506 0.671 OO9 45 -42 120.98550.010516.510.6688 0.448 O10 45 -28.2 121.01310.013416.340.7159 0.223 O11 45 0 120.96610.014614.050.6027 0.000 O12 45 28.2 121.01250.015712.390.7136 0.223 O13 45 42 120.95860.016910.420.7092 0.448 O14 45 55 120.94590.011511.130.5124 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 120.93760.02218.170.8168 0.671 O16 90 -42 120.95270.010215.750.5065 0.448 O17 90 -28.2 120.94020.013914.090.687 0.223 O18 90 0 120.97120.009419.540.4885 0.000 O19 90 28.2 120.92020.012813.160.578 0.223 O20 90 42 120.98970.018610.70.6039 0.448 O21 90 55 120.9680.01528.760.5464 0.671

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113 Stress calculations for the as-sprayed WC-17%Co coated specimen (121o 2-theta.103 peak). C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.8855 120.895 0.004 0.0000166.14E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 0.000263-0.000383 0.893 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)0.000219-0.000466 0.461 90 Eps22-Eps33 0.000031-0.000314 0.093 ave -0.000388 stdev 0.000076 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.0000110.000002 0.281 45 0.707*(Eps13+Eps23) 0.000033-0.000005 0.469 90 Eps23 -0.0000600.000012 0.427 Eps13 0.000106 Eps23 0.000035 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.166 psi* 24.032 strain* -3.30E-04 d0* 0.8852

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114 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 8.25E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.5E-12 S1 -nu/E zzzzzz1 -2.91E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-Axial Sig12 Sig13 Eps22 Eps23 Sig22-HoopSig23 Eps33 Sig33-Radial -0.000124 0.000072 0.000106 -349 48 71 -0.000357 -0.000060 -504 -40 -0.000388 -525 Std Dev. Std Dev. 0.000218 0.000288 0.000104 230 191 69 0.000271 0.000144 273 96 0.000084 135 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -349 48 71-0.0001240.0000720.000106 -504 -40 -0.000357-0.000060 -525 -0.000388

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115 Data for the ground and polished WC -17%Co coated specimen (117o 2-theta. 211 peak). Filename Observed Std dev Peak FWHM C11920 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 117.4849 0.024 10.32 1.4456 0.671 OO2 0 -42 117.3281 0.014 17.08 0.9847 0.448 OO3 0 -28.2 117.2027 0.015 18.16 1.1452 0.223 OO4 0 0 117.0772 0.015 18.07 1.2171 0.000 OO5 0 28.2 117.1971 0.015 18.16 1.1492 0.223 OO6 0 42 117.3696 0.016 17.08 1.2584 0.448 OO7 0 55 117.4376 0.024 10.87 1.3303 0.671 Filename Observed Std dev Peak FWHM C11920 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 117.5241 0.022 12.54 1.578 0.671 OO9 45 -42 117.3079 0.018 15.5 1.333 0.448 O10 45 -28.2 117.1997 0.015 18.48 1.078 0.223 O11 45 0 117.102 0.015 16.76 1.244 0.000 O12 45 28.2 117.247 0.017 13.51 1.082 0.223 O13 45 42 117.3539 0.020 10.16 1.263 0.448 O14 45 55 117.493 0.037 6.88 1.662 0.671 Filename Observed Std dev Peak FWHM C11920 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 117.456 0.034 7.06 1.548 0.671 O16 90 -42 117.346 0.023 10.71 1.376 0.448 O17 90 -28.2 117.217 0.021 12.96 1.293 0.223 O18 90 0 117.132 0.017 13.79 1.297 0.000 O19 90 28.2 117.315 0.021 12.69 1.427 0.223 O20 90 42 117.453 0.026 10.22 1.441 0.448 O21 90 55 117.433 0.058 6.56 2.316 0.671

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116 Stress calculations for the ground and po lished WC-17%Co coated specimen (117o 2theta.211 peak). C11920 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.9021 117.276 0.004 0.0000167.367E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 -0.0030940.001062 -0.999 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0031530.000979 -0.994 90 Eps22-Eps33 -0.0025480.000674 -0.980 ave 0.000905 stdev 0.000205 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.0000000.000008 -0.002 45 0.707*(Eps13+Eps23) -0.000053-0.000005 -0.243 90 Eps23 -0.000166-0.000007 -0.439 Alternate Eps13 0.000000 Eps23 -0.000074 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.347 psi* 36.068 strain* -9.94E-06 d0* 0.9021 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 # of psi tilts/phi 7 Young's Mod 6.39E+11

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117 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.94E-12 S1 -nu/E zzzzzz1 -3.76E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11 Sig12 Sig13 Eps22 Eps23 Sig22 Sig23 Eps33 Sig33 -0.002189 -0.000332 0.000091 -1824 -171 47 -0.001643 -0.000166 -1543 -86 0.000905 -230 Std Dev. Std Dev. 0.000212 0.000180 0.000082 167 93 42 0.000192 0.000126 154 65 0.000071 87 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -1824 -171 47-0.002189-0.0003320.000091 -1543 -86 -0.001643-0.000166 -230 0.000905

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118 Data for the ground and polished WC -17%Co coated specimen (121o 2-theta. 103 peak). Filename Observed Std dev Peak FWHM C11920 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 120.7905 0.056 6.84 2.3559 0.671 OO2 0 -42 120.8433 0.026 9 1.0096 0.448 OO3 0 -28.2 120.7089 0.030 9.68 1.1737 0.223 OO4 0 0 120.5888 0.026 10.12 1.0994 0.000 OO5 0 28.2 120.7092 0.028 10.07 1.1809 0.223 OO6 0 42 120.759 0.036 8.48 1.6716 0.448 OO7 0 55 120.8559 0.050 6.22 1.5719 0.671 Filename Observed Std dev Peak FWHM C11920 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 120.8767 0.041 6.44 1.702 0.671 OO9 45 -42 120.7712 0.034 8.34 1.575 0.448 O10 45 -28.2 120.731 0.030 9.79 1.319 0.223 O11 45 0 120.6563 0.030 9.07 1.358 0.000 O12 45 28.2 120.726 0.031 8.35 1.104 0.223 O13 45 42 120.824 0.040 5.36 1.289 0.448 O14 45 55 120.9091 0.070 3.8 2.080 0.671 Filename Observed Std dev Peak FWHM C11920 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 120.945 0.068 3.87 1.772 0.671 O16 90 -42 120.831 0.041 6.12 1.278 0.448 O17 90 -28.2 120.744 0.027 8.69 0.965 0.223 O18 90 0 120.681 0.035 7.57 1.533 0.000 O19 90 28.2 120.772 0.036 7.64 1.404 0.223 O20 90 42 120.827 0.045 5.63 1.560 0.448 O21 90 55 121.066 0.080 4.42 2.074 0.671

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119 Stress calculations for the ground and po lished WC-17%Co coated specimen (121o 2theta.103 peak). C11920 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.8860 120.781 0.004 0.0000166.178E-10 PHI a1 slope a1 slope Y-I=Eps33CORR. COEFF. 0 Eps11-Eps33 -0.0017660.000843-0.963 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0017270.000640-0.997 90 Eps22-Eps33 -0.0023100.000591-0.973 ave 0.000691 stdev 0.000134 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.000035-0.0000130.108 45 0.707*(Eps13+Eps23) -0.0000890.000012-0.608 90 Eps23 -0.000123-0.000004-0.402 Alternate Eps13 0.000035 Eps23 -0.000161 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.476 psi* 43.624 strain* -2.38E-06 d0* 0.8860

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120 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 # of psi tilts/phi 7 Young's Mod 8.25E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.5E-12 S1 -nu/E zzzzzz1 -2.91E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11 Sig12 Sig13 Eps22 Eps23 Sig22 Sig23 Eps33 Sig33 -0.001074 0.000312 -0.000002 -1330 207 -1 -0.001619 -0.000123 -1692 -82 0.000691 -155 Std Dev. Std Dev. 0.000404 0.000312 0.000146 405 207 97 0.000296 0.000204 317 136 0.000135 202 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -1330 207 -1-0.0010740.000312-0.000002 -1692 -82 -0.001619-0.000123 -155 0.000691

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121 Data for the ground and polished WC -17%Co coated specimen (117o 2-theta. 211 peak) fatigue tested at 110ksi maximum applied stress. Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 117.35840.02973.332.49 0.671 OO2 0 -42 117.32310.03493.511.8434 0.448 OO3 0 -28.2 117.14640.03173.881.0628 0.223 OO4 0 0 117.130.03543.661.1096 0.000 OO5 0 28.2 117.19480.03463.641.1625 0.223 OO6 0 42 117.26040.02963.870.9714 0.448 OO7 0 55 117.36640.04252.491.413 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 117.33080.0354.011.5378 0.671 OO9 45 -42 117.20030.0225.151.8222 0.448 O10 45 -28.2 117.16630.03344.021.7181 0.223 O11 45 0 117.11620.0264.811.7648 0.000 O12 45 28.2 117.20550.04133.131.7458 0.223 O13 45 42 117.3420.02443.451.3885 0.448 O14 45 55 117.36630.06061.760.8184 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 117.25580.03833.611.166 0.671 O16 90 -42 117.20970.03623.751.244 0.448 O17 90 -28.2 117.16840.02265.181.1599 0.223 O18 90 0 117.07850.02444.440.9821 0.000 O19 90 28.2 117.26790.03463.121.3136 0.223 O20 90 42 117.32970.04982.190.6625 0.448 O21 90 55 117.36960.05691.570.5653 0.671

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122 Stress calculations for the ground and po lished WC-17%Co coated specimen (117o 2theta. 211 peak) fatigue tested at 110ksi maximum applied stress. C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.9023 117.234 0.004 0.0000167.38E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 -0.0019470.000631 0.984 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0018610.000646 0.999 90 Eps22-Eps33 -0.0017950.000681 0.956 ave 0.000653 stdev 0.000026 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.000041-0.000025 0.158 45 0.707*(Eps13+Eps23) -0.0002340.000018 0.671 90 Eps23 -0.0003210.000000 1.000 Eps13 -0.000010 Eps23 -0.000372 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.350 psi* 36.296 strain* -3.80E-05 d0* 0.9023

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123 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 6.39E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.94E-12 S1 -nu/E zzzzzz1 -3.76E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-axial Sig12 Sig13 Eps22 Eps23 Sig22-hoopSig23 Eps33 Sig33-radial -0.001294 0.000010 -0.000010 -1091 5 -5 -0.001143 -0.000321 -1013 -165 0.000653 -88 Std Dev. Std Dev. 0.000653 0.000784 0.000366 522 404 188 0.000693 0.000393 548 203 0.000252 295 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -1091 5 -5-0.0012940.000010-0.000010 -1013 -165 -0.001143-0.000321 -88 0.000653

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124 Data for the ground and polished WC -17%Co coated specimen (121o 2-theta. 103 peak) fatigue tested at 110ksi maximum applied stress. Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 120.83780.15891.311.8434 0.671 OO2 0 -42 120.80180.09351.721.0628 0.448 OO3 0 -28.2 120.67010.04352.981.1096 0.223 OO4 0 0 120.73960.05392.071.1272 0.000 OO5 0 28.2 120.73960.05392.071.1272 0.223 OO6 0 42 120.70410.04992.470.9714 0.448 OO7 0 55 120.58140.06082.311.413 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 120.82090.06132.781.5378 0.671 OO9 45 -42 120.71950.08781.811.8222 0.448 O10 45 -28.2 120.67030.07122.021.7181 0.223 O11 45 0 120.59170.07232.431.7648 0.000 O12 45 28.2 120.66710.080821.7458 0.223 O13 45 42 120.81350.10231.011.3885 0.448 O14 45 55 120.77660.05871.510.8184 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 120.68010.05612.321.166 0.671 O16 90 -42 120.7870.05732.191.244 0.448 O17 90 -28.2 120.75650.05332.211.1599 0.223 O18 90 0 120.68960.042.940.9821 0.000 O19 90 28.2 120.79270.06591.811.3136 0.223 O20 90 42 120.69970.04781.90.6625 0.448 O21 90 55 120.75650.05431.180.5653 0.671

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125 Stress calculations for the ground and po lished WC-17%Co coated specimen (121o 2theta. 103 peak) fatigue tested at 110ksi maximum applied stress. C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.8862 120.736 0.004 0.0000166.19E-10 PHI a1 slope a1 slope Y-I=Eps33CORR. COEFF. 0 Eps11-Eps33 0.0000940.0000120.235 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0015970.0006810.984 90 Eps22-Eps33 -0.0001210.0000610.195 ave 0.000251 stdev 0.000373 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.000317-0.0000430.422 45 0.707*(Eps13+Eps23) -0.0000650.0000160.208 90 Eps23 0.000009-0.0000220.024 Eps13 -0.000101 Eps23 -0.000409 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.356 psi* 36.640 strain* 5.14E-05 d0* 0.8862

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126 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 8.25E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.5E-12 S1 -nu/E zzzzzz1 -2.91E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-Axial Sig12 Sig13 Eps22 Eps23 Sig22-HoopSig23 Eps33 Sig33-Radial 0.000345 -0.001583 -0.000101 453 -1053 -67 0.000130 0.000009 310 6 0.000251 390 Std Dev. Std Dev. 0.001570 0.001397 0.000895 1570 930 595 0.001030 0.000511 1127 340 0.000606 824 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 453 -1053 -670.000345-0.001583-0.000101 310 6 0.0001300.000009 390 0.000251

PAGE 140

127 Data for the ground and polished WC -17%Co coated specimen (117o 2-theta. 211 peak) fatigue tested at 150ksi maximum applied stress. Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 117.4750.05022.591.2393 0.671 OO2 0 -42 117.22430.03843.021.5291 0.448 OO3 0 -28.2 117.14530.0373.131.2982 0.223 OO4 0 0 117.01050.02414.360.8569 0.000 OO5 0 28.2 117.16610.03123.861.1004 0.223 OO6 0 42 117.28830.03223.671.2174 0.448 OO7 0 55 117.4560.04652.481.5283 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 117.4730.03953.531.2603 0.671 OO9 45 -42 117.25160.03283.991.2041 0.448 O10 45 -28.2 117.11170.02114.720.7754 0.223 O11 45 0 116.99950.02924.031.1021 0.000 O12 45 28.2 117.16690.03573.021.1149 0.223 O13 45 42 117.33560.03362.941.0259 0.448 O14 45 55 117.45340.03982.061.0117 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 117.27830.04173.081.1424 0.671 O16 90 -42 117.22870.03044.021.0457 0.448 O17 90 -28.2 117.11340.03283.61.0952 0.223 O18 90 0 117.0480.03563.51.1639 0.000 O19 90 28.2 117.18810.03873.231.2437 0.223 O20 90 42 117.34550.03952.831.1306 0.448 O21 90 55 117.45670.05971.721.3801 0.671

PAGE 141

128 Stress calculations for the ground and po lished WC-17%Co coated specimen (117o 2theta. 211 peak) fatigue tested at 150ksi maximum applied stress. C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.9024 117.214 0.004 0.0000167.39E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 -0.0034860.001127 0.990 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0036770.001181 0.999 90 Eps22-Eps33 -0.0026060.000877 0.996 ave 0.001061 stdev 0.000162 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 -0.0000720.000006 0.354 45 0.707*(Eps13+Eps23) -0.0001160.000001 0.423 90 Eps23 -0.0003670.000008 0.862 Eps13 0.000204 Eps23 -0.000092 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.336 psi* 35.429 strain* -3.66E-05 d0* 0.9024

PAGE 142

129 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 6.39E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.94E-12 S1 -nu/E zzzzzz1 -3.76E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-axial Sig12 Sig13 Eps22 Eps23 Sig22-hoopSig23 Eps33 Sig33-radial -0.002425 -0.000631 0.000204 -1941 -325 105 -0.001544 -0.000367 -1487 -189 0.001061 -145 Std Dev. Std Dev. 0.000744 0.000753 0.000419 593 388 216 0.000762 0.000431 604 222 0.000287 333 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -1941 -325 105-0.002425-0.0006310.000204 -1487 -189 -0.001544-0.000367 -145 0.001061

PAGE 143

130 Data for the ground and polished WC -17%Co coated specimen (121o 2-theta. 103 peak) fatigue tested at 150ksi maximum applied stress. Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 120.74170.13531.172.2242 0.671 OO2 0 -42 120.65280.08331.681.9676 0.448 OO3 0 -28.2 120.62310.0811.521.4107 0.223 OO4 0 0 120.52970.05881.891.133 0.000 OO5 0 28.2 120.59950.09231.531.5127 0.223 OO6 0 42 120.79260.06461.91.3066 0.448 OO7 0 55 120.76670.09171.751.8434 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 120.88620.0851.71.4888 0.671 OO9 45 -42 120.78850.1051.491.5949 0.448 O10 45 -28.2 120.64830.05032.281.2652 0.223 O11 45 0 120.47670.05342.281.0421 0.000 O12 45 28.2 120.66710.0452.270.7046 0.223 O13 45 42 120.85630.08321.431.4805 0.448 O14 45 55 120.7790.12020.931.7189 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 120.74740.07751.671.2012 0.671 O16 90 -42 120.66690.0632.091.1099 0.448 O17 90 -28.2 120.51050.07281.811.2575 0.223 O18 90 0 120.45830.0781.811.52 0.000 O19 90 28.2 120.6650.05632.171.08 0.223 O20 90 42 120.6480.13111.192.2491 0.448 O21 90 55 120.61360.17180.671.8954 0.671

PAGE 144

131 Stress calculations for the ground and po lished WC-17%Co coated specimen (121o 2theta. 103 peak) fatigue tested at 150ksi maximum applied stress. C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.8865 120.667 0.004 0.0000166.21E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 -0.0017450.000650 0.981 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0027380.000771 0.950 90 Eps22-Eps33 -0.0016390.000906 0.951 ave 0.000776 stdev 0.000128 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 -0.0001720.000031 0.446 45 0.707*(Eps13+Eps23) 0.0000130.000004 0.033 90 Eps23 0.000065-0.000046 0.103 Eps13 -0.000047 Eps23 0.000190 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.363 psi* 37.071 strain* 1.04E-05 d0* 0.8865

PAGE 145

132 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 8.25E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.5E-12 S1 -nu/E zzzzzz1 -2.91E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-Axial Sig12 Sig13 Eps22 Eps23 Sig22-HoopSig23 Eps33 Sig33-Radial -0.000969 -0.001046 -0.000047 -969 -696 -31 -0.000864 0.000065 -899 43 0.000776 192 Std Dev. Std Dev. 0.001595 0.001748 0.000909 1648 1163 605 0.001706 0.000984 1740 654 0.000616 934 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -969 -696 -31-0.000969-0.001046-0.000047 -899 43 -0.0008640.000065 192 0.000776

PAGE 146

133 Data for the ground and polished WC -17%Co coated specimen (117o 2-theta. 211 peak) fatigue tested at 220ksi maximum applied stress. Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 117.43310.02772.531.0058 0.671 OO2 0 -42 117.32540.0263.331.2887 0.448 OO3 0 -28.2 117.21520.02263.971.1464 0.223 OO4 0 0 117.090.02024.070.9751 0.000 OO5 0 28.2 117.14430.02243.781.0333 0.223 OO6 0 42 117.26570.02153.831.0003 0.448 OO7 0 55 117.36850.03082.481.3148 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 117.41970.02712.991.0092 0.671 OO9 45 -42 117.2860.02194.011.2331 0.448 O10 45 -28.2 117.17270.02184.141.0705 0.223 O11 45 0 117.1070.01894.451.0808 0.000 O12 45 28.2 117.23260.02253.611.0371 0.223 O13 45 42 117.3110.02732.61.0897 0.448 O14 45 55 117.3360.04191.691.3608 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 117.24890.02133.641.0772 0.671 O16 90 -42 117.16650.02194.291.1654 0.448 O17 90 -28.2 117.13730.0244.081.2478 0.223 O18 90 0 117.13650.01684.70.8546 0.000 O19 90 28.2 117.290.02383.581.1112 0.223 O20 90 42 117.29490.02692.711.0164 0.448 O21 90 55 117.36110.03612.031.2927 0.671

PAGE 147

134 Stress calculations for the ground and po lished WC-17%Co coated specimen (117o 2theta. 211 peak) fatigue tested at 220ksi maximum applied stress. C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.9024 117.214 0.004 0.0000167.39E-10 PHI a1 slope a1 slope Y-I=Eps33CORR. COEFF. 0 Eps11-Eps33 -0.0024930.0006890.999 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0021600.0005500.999 90 Eps22-Eps33 -0.0012430.0003760.975 ave 0.000538 stdev 0.000157 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 0.0001800.0000050.959 45 0.707*(Eps13+Eps23) 0.000019-0.0000140.055 90 Eps23 -0.000360-0.0000120.933 Eps13 0.000387 Eps23 -0.000153 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.286 psi* 32.336 strain* -1.74E-05 d0* 0.9024

PAGE 148

135 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 6.39E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.94E-12 S1 -nu/E zzzzzz1 -3.76E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-axial Sig12 Sig13 Eps22 Eps23 Sig22-hoopSig23 Eps33 Sig33-radial -0.001955 -0.000292 0.000387 -1512 -151 200 -0.000705 -0.000360 -868 -186 0.000538 -227 Std Dev. Std Dev. 0.000486 0.000571 0.000268 386 294 138 0.000474 0.000259 378 134 0.000188 215 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -1512 -151 200-0.001955-0.0002920.000387 -868 -186 -0.000705-0.000360 -227 0.000538

PAGE 149

136 Data for the ground and polished WC -17%Co coated specimen (121o 2-theta. 103 peak) fatigue tested at 220ksi maximum applied stress. Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO1 0 -55 120.7230.09431.212.0685 0.671 OO2 0 -42 120.70420.05831.921.5858 0.448 OO3 0 -28.2 120.61530.05242.041.5587 0.223 OO4 0 0 120.63170.05661.581.3871 0.000 OO5 0 28.2 120.64690.03831.990.9824 0.223 OO6 0 42 120.6980.02692.490.5917 0.448 OO7 0 55 120.72570.06331.441.3666 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi OO8 45 -55 120.96620.02492.750.5511 0.671 OO9 45 -42 120.67620.04881.951.4757 0.448 O10 45 -28.2 120.68210.04731.881.1548 0.223 O11 45 0 120.64130.02393.010.7835 0.000 O12 45 28.2 120.67340.05691.821.3529 0.223 O13 45 42 120.77720.03971.760.9301 0.448 O14 45 55 120.7570.09891.012.0924 0.671 Filename Observed Std dev Peak FWHM C11941 Phi Chi 2theta 2theta In tensity (cps) Sin^2Psi O15 90 -55 120.70030.04281.971.1982 0.671 O16 90 -42 120.65610.04632.071.3279 0.448 O17 90 -28.2 120.65610.04451.991.1663 0.223 O18 90 0 120.64950.04431.911.084 0.000 O19 90 28.2 120.69180.06981.511.6975 0.223 O20 90 42 120.70670.06991.231.5344 0.448 O21 90 55 120.61720.11990.752.0537 0.671

PAGE 150

137 Stress calculations for the ground and po lished WC-17%Co coated specimen (121o 2theta. 103 peak) fatigue tested at 220ksi maximum applied stress. C11941 Calculate 3-D Stress Tensor a1 & a2 Analysis (see p119 of Noyan & Cohen) Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Lambda 1.54060 Std dev Variance Variance dzero 2theta 2theta 2theta d (^2) 0.8866 120.645 0.004 0.0000166.22E-10 PHI a1 slope a1 slope Y-I=Eps33 CORR. COEFF. 0 Eps11-Eps33 -0.0007730.000123 0.938 45 0.5*(Eps11+2Eps12+Eps22-2Eps33)-0.0015720.000120 0.951 90 Eps22-Eps33 -0.000078-0.000079 0.314 ave 0.000055 stdev 0.000116 PHI a2 slope a2 slope Y-I=ZIP CORR. COEFF. 0 Eps13 -0.000012-0.000009 0.139 45 0.707*(Eps13+Eps23) 0.0000850.000013 0.123 90 Eps23 0.000001-0.000003 0.003 Eps13 0.000119 Eps23 0.000133 if you wish calculate d0 by Hauk method (plane stress) & revise assumed d0 if so you need data from phi=0&90 sin sq psi* 0.198 psi* 26.454 strain* -1.03E-05 d0* 0.8866

PAGE 151

138 Calculate 3-D Stress Tensor Calculate Sig11, Sig22, Sig33, Sig12, Sig23 & Sig13 Young's Mod 8.25E+11 Poisson's ratio 0.24 S2/2 (1+nu)/E zzzzzz2 1.5E-12 S1 -nu/E zzzzzz1 -2.91E-13 -nu/(1-2nu)QUE -0.461538 D-H Strain Analysis D-H Stress Analysis Strain Results Stress Results(MPa) Eps11 Eps12 Eps13 Sig11-Axial Sig12 Sig13 Eps22 Eps23 Sig22-HoopSig23 Eps33 Sig33-Radial -0.000718 -0.001147 0.000119 -689 -763 79 -0.000023 0.000001 -226 1 0.000055 -174 Std Dev. Std Dev. 0.001142 0.001155 0.000649 1172 769 432 0.001144 0.000651 1174 433 0.000441 656 Stresses calc from D-H strains Strains calc from D-H Stresses Stress Results (MPa) Strain Results Sig11 Sig12 Sig13 Eps11 Eps12 Eps13 Sig22 Sig23 Eps22 Eps23 Sig33 Eps33 -689 -763 79-0.000718-0.0011470.000119 -226 1 -0.0000230.000001 -174 0.000055

PAGE 152

139 LIST OF REFERENCES Ahmed R., Hadfield M. Failure Modes of Plasma Sprayed WC-15% Co Coated Rolling Elements. WEAR 1999; 30: 39-55. AIR5052. Crack Initiation and Growth Consid erations for Landing Gear Steel with Emphasis on Aermet 100. Aerospace Information Report. Society of Automotive Engineers; 1997. AMS 6532B. Steel, Bars and Forgings, 3.1C r 11.5Ni 13.5Co 1.2Mo (0.21 0.25C), Vacuum Melted, Annealed Heat Treata ble to 280, 000 psi Tensile Strength. Aerospace Materials Specification. Soci ety of Automotive Engineers; 2001. AMS 7881. Tungsten Carbide Cobalt Powder, A gglomerated and Sintered. Aerospace Materials Specification. Society of Automotive Engineers; 2003. AMS S13165. Shot Peening of Metal Parts. Aerospace Materials Specification. Society of Automotive Engineers; 1997. ARP 1631. Manufacturing Sequence for Fabricat ion of High-Strength Steel Parts 270,000 psi (1860 MPa) Tensile St rength and Higher, 300M or 4340 Modified Low-Alloy Steels. Aerospace Recommended Practice. Society of Automotive Engineers; 1984. ASTM C633. Adhesive or C ohesive Strength of Flame Sprayed Coatings. Industry Standard. American Society fo r Testing and Materials; 2001. ASTM E384. Standard Test Method for Microindentation Har dness of Materials. Industry Standard. 9, American Societ y for Testing and Materials; 1999. Bannantine J.A., Comer J.J., and Handrock J.L. Fundamentals of Metal Fatigue Analysis. New Jersey: Prentice Hall Inc.; 1990. p. 1-28. De Palo S., Mohanty M., Marc-Charles H ., Dorfman M. Fracture Toughness of HVOF Sprayed WC-Co Coatings. J. Surface Engi neering via Applied Research 2000; May 1; 245-250. England G. 1997 The Nature of Thermal Spray Coatings. Available from URL: http://www.gordonengland.co.uk/xtsc.htm. Site last visited April 2003. Gilliam E. Materials Under Stress. Clevel and: CRC Press Interna tional Series; 1969. p. 67-68.

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140 Kim B.H., Suhr D.S. Charac teristics of HVOF-Sprayed WC-Co Cermet Coatings Affected by WC Particle Size and Fue l/Oxygen Ratio. Materials Transactions 2001; 42 (5): 833-837. Krause A.H., Haase A. X-ray Diffraction Sy stem PTS for Powder Texture and Stress Analysis. In: H.J. Bunge, Editor. Experi mental Techniques of Texture Analysis. DGM Informationsgesellschaft Verlag 1986, p. 405-408. McGrann R., Shadley J.R., Rybicki E., Bodge r B., Emory W., Somerville D., Greving D. Evaluation of Residual Stresses and Fati gue Life of Tungsten Carbide Thermal Spray Coated Aircraft Landing Gear Mate rials. In Coddet C. editor: ITSC 98. Proceedings of the International Th ermal Spray Conference; 1998 May 25-29; Nice, France; p.557-562. Nascimento M., Souza R.C., Miguel I.M., Pigatin W.L., Voorwald H. Effects of Tungsten Carbide Thermal Spray Coating by HP/HVOF and Hard Chromium Electroplating on AISI 4340 High Streng th Steel. Surface and Coatings Technology 2001; 138: 113-124. Northwestern University. Hard Chrome Coatings: Advanced Technology for Waste Elimination. Final Report.Evanston, IL, under Defense Advance Research Project Agency (DARPA) 1996 C ontract MDA972-93-1-0006. Noyan I.C., Cohen J.B. Residual Stress Meas urement by Diffraction and Interpretation. New York: Springer-Verlag; 1987. p.101-102; p.125-126. Prevey P.S. Problems with Non-destructive Surface X-ray Diffraction Residual Stress Measurement. In: Ruud C. editor. Practi cal Applications of Residual Stress Technology. Cleveland: ASM Int’l; 1991. p. 47-54. Reiners G., Kreye H., Schwetzke R. Propert ies and Characterization of Thermal Spray Coatings. In: Coddet C. editor ITSC 98. Pr oceedings of the In ternational Thermal Spray Conference; 1998 May 2529; Nice, France; p.629-634. Sartwell B. Joint Test Report, Part I, Materials Testing. Washington (D.C.): Environmental Security Technology Cer tification Program (ESTCP); 2002 Nov. Science Applications International Corporat ion (SAIC). High Velocity Oxy Fuel Results. Final Report. Oklahoma City: Air Logist ics Center, Tinker Air Force Base 1994 May Report No. F09603-90-D2215. Semant J. Materials Science Of Thermal Sp ray Deposits. J. of Thermal Spray Coatings 1998; 454 (1): 15-35 Tajiri T., Watanabe S., Amano J., Sakoda N. Fatigue Strength and Fracture Mechanism of Thermally Sprayed WC-Co Coating on a Medium Carbon Steel by HVOF. In: Coddet C. editor ITSC 98. Proceedings of the International Thermal Spray Conference; 1998 May 25-29; Nice, France; p. 611-616.

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141 Verdon C., Karimi A., Martin J.L. A St udy of High Velocity Oxy-Fuel Thermally Sprayed Tungsten Carbide Based Coatings. Part I: Microstruc tures in Materials Science and Engineering 1998; A246: 11-24. Verdon C., Karimi A., Martin J.L. Microstr uctural and Analytical Study of Thermally Sprayed WC-Co Coatings in Connection wi th their Wear Resistance. Materials Science and Engineering 1997; A234-236: 731-734. Wang T., Platts M.J. A Computer-aided Bl ank Design Method for the Peen Forming Process in Journal of Materials Pr ocessing Technology 2002; 122: 374-380. Watanabe S., Tajiri T., Sakoda N., Ama no J. Fatigue Cracks in HVOF Thermally Sprayed WC-Co Coatings. Journal of Th ermal Spray Technology 1998; March 1; 7 (1): 93-96. Whithers P., Bhadeshia H. Residual Stress. Journal of Materials Science and Technology 2001; April 1; 17: 355-375.

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142 BIOGRAPHICAL SKETCH Donald Scott Parker was born in Indianapolis, Indiana, on June 4, 1965. He attended Pike High School and played severa l sports while taking a scientific college preparatory curriculum. He graduated in Ma y 1983. He went on to the University of Florida the next semester, joining the Materials Science program in 1986. His undergraduate degree was awar ded in August 1988 with a specialization in ceramic materials. Donald has worked in several fields since receiving his undergraduate degree (construction materials with Professional Service Industries; computer board and integrated-circuit manufacturing with Am erican Semiconductor; failure analysis engineering with Naval Aviati on Depot in both Pensacola and Jacksonville, Florida). In 1999 he accepted a position with NASA at Kennedy Space Center as an aerospace Structural Materials Engineer working in the Spaceport Engineering and Technology Directorate. Don performs failure anal ysis on space-shuttle hardware, launch and ground-support equipment, space-station payloa ds and hardware, and expendable launch vehicles. He also performs applied research on new material technol ogies for integration into launch platforms and environments. D on has published five peer-reviewed technical papers related to his research and will be publishing excerpts from his thesis work for several industrial projects. He has co mpleted work on three aerospace-materials specifications related to cermet coatin gs (being published by SAE/AMS). Upon graduation Donald will continue working fo r NASA at Kennedy Space Center, Florida.


Permanent Link: http://ufdc.ufl.edu/UFE0000763/00001

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Title: Fracture and Residual Stress Characterization of Tungsten Carbide 17% Cobalt Thermal Spray Coatings Applied to High Strength Steel Fatigue Specimens
Physical Description: Mixed Material
Copyright Date: 2008

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Holding Location: University of Florida
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Permanent Link: http://ufdc.ufl.edu/UFE0000763/00001

Material Information

Title: Fracture and Residual Stress Characterization of Tungsten Carbide 17% Cobalt Thermal Spray Coatings Applied to High Strength Steel Fatigue Specimens
Physical Description: Mixed Material
Copyright Date: 2008

Record Information

Source Institution: University of Florida
Holding Location: University of Florida
Rights Management: All rights reserved by the source institution and holding location.
System ID: UFE0000763:00001


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FRACTURE AND RESIDUAL-STRESS CHARACTERIZATION OF TUNGSTEN-
CARBIDE 17%-COBALT THERMAL-SPRAY COATINGS APPLIED TO HIGH-
STRENGTH STEEL FATIGUE SPECIMENS
















BY

DONALD SCOTT PARKER


A THESIS PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
MASTER OF SCIENCE

UNIVERSITY OF FLORIDA


2003


































Copyright 2003

by

Donald S. Parker
































To my peers, team members, friends, and co-workers in the aircraft and aerospace field
who helped to advance my knowledge in the field of materials science and thermal spray
coatings; and to develop my research skills to a level I never thought possible. Without
their support, encouragement, and critique I never would have been able to complete this
project.














ACKNOWLEDGMENTS

I wish to thank my advisor (Dr. Darryl Butt) for his guidance and patience on this

project while I tried to work full time, complete the research, and write the final thesis. I

also thank the other members of my committee, (Dr. Gerhard Fuchs and Dr. Jack

Mecholsky) for their knowledge and guidance. I also thank the following individuals for

their contributions: Mr. Bruce Sartwell for providing funding, test specimens, and a

forum for peer review of the work; Dr. Philip E. Bretz, Mr. Jerry Schell, and Dr. Jeam-

Gabriel Legoux for their expert advice in evaluating the coating materials and

fundamental mechanical properties; Dr. Thomas Watkins for his assistance and guidance

on residual stress measurement and data interpretation; Mr. Peter Marciniak for his

photographic talents; and Ms. Virginia Cummings for her assistance in scanning electron

microscopy, metallography, and proof-reading.

I also thank the Kennedy Space Center, Labs Division management (Mr. Timothy

Bollo, Mr. Scott Murray, and Mr. Steve McDanels) for their ongoing support throughout

my graduate studies.

I especially thank my wife, Ms. Pennie Parker; and daughter, Ms. Alli Brown for

their encouragement and tolerance of the time needed to complete my studies.
















TABLE OF CONTENTS
Page

A C K N O W L E D G M E N T S ................................................................................................. iv

LIST O F TA B LE S ......................................... ......... ............ .............. .. vii

L IST O F FIG U R E S .................. ............. ........................... ................ viii

ABSTRACT .............. ..................... .......... .............. xii

CHAPTER

1 INTRODUCTION AND BACKGROUND ..................................... ...............

Review of the Chrome-Plating Replacement Effort................... ................................1
High-Velocity Oxygen Fuel (HVOF), Thermal-spray Process..................................2
Selection of Tungsten C arbide C oatings ........................................... .....................4
C eating C characterization ........................................... ................................. 7

2 LITER A TU RE SU RVEY .................................................. ............................... 9

In tro d u ctio n .................................................................................. 9
O v erview of F atigu e ................... .......................................... ........ .......... .. ....
Overview of R esidual Stress...................................................... ...................... 11
O verview of Shot Penning ................................................................................... 14
Electroplated Chromium and HVOF Applied Coatings............... .. ............... 15
Microstructure of the Tungsten Carbide Cobalt Coatings........................................17

3 EXPERIM ENTAL PROCEDURE ................................... ............................. ....... 21

Specim en M manufacture .......................................................................... ............... 21
C eating Q u ality C control ..................................................................... ..................26
Residual Stress M easurem ent ............................................................................... 28
F atig u e T estin g ...................................................... ................ 3 2
Scanning Electron M icroscopy .............................................................. ............... 33

4 RESULTS AND DISCU SSION ........................................... .......................... 34

Residual Stress M easurem ents ............................................................................34
Optical and Scanning Electron M icroscopy .................................... ............... 39
Low -A pplied-Stress Specim ens............................................ ........... ............... 40


v









M medium Applied Stress Specim ens...................................................................... 50
H igh A applied Stress Specim ens........................................... ........................... 56

5 CON CLU SION S ....................................... ... ......... ........ ..... ...... 65

APPENDIX

A EXPERIMENTAl DATA FOR DEVELOPING HVOF SPRAY PARAMETERS
FOR TUNGSTEN CARBIDE 17%-COBALT COATINGS .................................68

Coating-Process Response M easurements ...................................... ............... 70
C oating-D position R ate ................................................. ............................... 70
M icroh ardn ess........................... ..................................70
Rockwell 15N Superficial Surface Hardness .................................. ............... 71
Tensile B ond Strength ........................................ ................... ..... .... 71
Substrate Tem perature .................. ..................................... .. ........ .... 71
A lm en Strip D election ...................................................................... ...................72

B FATIGUE DATA FOR 4340, 300M, AND AERMET 100....................................84

C RESIDUAL STRESS SCAN DATA ................................................ ............... 102

LIST OF REFEREN CES ........................................................... .. ............... 139

BIOGRAPHICAL SKETCH ............................................................. ............... 142
















LIST OF TABLES


Table page

3-1 Composition for the three alloys evaluated. ............. .........................................23

3-2 Particle size distribution for the Sulzer Metco, Diamalloy 2005 TM...........................25

3-3 Experimental conditions for the x-ray measurements .......................................32

4-1 Shot-peened bare 4340 steel substrate material............................... ............... 35

4-2 Shot-peened bare 4340 steel substrate................................... ......................... 35

4-3 Residual stress data for the WC 17%Co HVOF coating in the as-sprayed
condition ............................................................................................. .36

4-4 Residual stress data for the WC 17%Co HVOF coating in the finish ground, and
polished condition. ....................... ...................... ................... .. ..... 37

4-5 Low -applied stress specim en. ...... ......................................................................38

4-6 M edium -applied stress specim en......................................... ........................... 38

4-7 H igh-applied stress specim en. ............................................. ........................... 39

A-1 Optimization parameters for the Tungsten-carbide 17% Cobalt coatings ...............69
















LIST OF FIGURES


Figure p

1-1 Exploded view of HVOF gun and spray process.. ............... ......... ......................3

1-2 M orphology of a thermal-spray deposited coating..................................................3

1-3 Fatigue results for initial coating evaluation.. ..................................... ............... 6

2-1 Model of shot peened surface showing the stress profile created by the impact
dim ples on the surface .................................................... .. ........ .... 15

2-2 Cross-section showing the coating microstructure of an as-sprayed WC-17%Co
coating applied to a high-strength 4340 substrate...................... ....................18

2-3 Polished WC 17%Co coating cross-section showing the lamellar structure and
uniform distribution of W C particles. ................... ..... ............................. 18

2-4 Higher magnification micrograph of the same sample in Figure 2-3 showing the
white spherical WC particles suspended in the lighter gray cobalt matrix ............19

3-1 Fracture toughness data for 4340, 300M, and Aermet 100. .....................................22

3-2 Ductility properties for tensile elongation and reduction in area data for
Aerm et 100 versus 4340, and 300M ............................................. ............... 23

3-3 Hourglass fatigue specimen and coated section that was used for the study. ............24

3-4 Screw type holder used for residual stress measurement during spray
o p e ratio n s ......................................................................... 2 7

3-5 Scintag PTS Goniometer at Oak Ridge National Laboratory...............................28

3-6 Theta/2-theta scan results for the coated specimen. .................................................29

3-7 Specimen installation and alignment in the Goniometer ..........................................30

3-8 Orientation of the sample holder and the lead foil masking in the goniometer ..........31

4-1 Specim en before fatigue testing ........................................... .......................... 40









4-2 Surface morphology of a finish ground coating that was polished to a 2-4 Ra
ro u g h n e ss.............................. ......................................................... ............... 4 0

4-3 Fatigue specimen tested at 110 ksi cyclic stress for greater than 107 cycles..............41

4-4 Circumferential coating cracks and the branches that are propagating along
m machine lines on the surface. ............................................................................. 42

4-5 Micrograph of the crack tip from the delaminated coating area.............................42

4-6 Coating and substrate at the fracture surface for an Aermet 100 specimen
tested at 135 ksi m aximum applied stress. .................................... .................43

4-7 Subsurface defect and small slow crack growth region for the 4340 steel
specimen tested at 125 ksi maximum stress ........................................................44

4-8 Subsurface defect at the origin, and the small slow crack growth region for a
300M steel specimen tested at 130 ksi maximum applied stress. ............................ 45

4-9 Aermet 100 specimen showing the much larger slow crack growth region..............45

4-10 Subsurface inclusion and the radial lines leading to the crack origin for a
4340 steel specimen tested at 125 ksi maximum applied stress.............................46

4-11 4340 specimen showing the secondary cracking and tear ridges along
parallel fronts progressing away from the crack origin.........................................47

4-12 Faint striations detected near the outer edge of the slow crack growth
region in a 300M specim en.. ............................. ............................................... 47

4-13 Microvoid coalescence in the region of fast fracture for a 300M specimen. .........48

4-14 Low magnification micrograph image of the subsurface origin of the
A erm et 100 sam ple tested at 135 ksi..................................... ....................... 48

4-15 Tear ridges and worked surface with concentric parallel lines on the
A erm et 100 sam ple............ ... ........................................................ .......... ...... 49

4-16 Fatigue striations shown on the fracture surface approximately halfway
between the origin and the transition zone between slow and rapid crack growth..49

4-17 Very fine circumferential coating cracks formed on the 150 ksi maximum
applied stress run-out specim ens .......... .......................................... ........ ..... 51

4-18 Higher magnification micrograph showing the very fine circumferential
cracks propagating between surface defects. ................................... ..................... 51

4-19 Aermet 100 sample tested at 145 ksi maximum applied stress showing the
origin to at th e su rface ...................................................................... .................. 52









4-20 300M specimen tested at 145 ksi maximum applied stress................. ......... 52

4-21 Origin area on Aermet 100 fatigue specimen tested at 145 ksi max stress ............53

4-22 Fatigue crack origin at a subsurface inclusion for a 300M specimen tested
at 145 ksi m aximum applied stress. .............................................. ............... 54

4-23 Transition region along the edge of the slow crack growth region for the
A erm et 100 specim en ........................................................ ... ...... .... 54

4-24 Coating/substrate interface for the Aermet 100 specimen............... .......... 55

4-25 Cross-section through the coating at the fracture surface just below the
substrate origin. .......................................................................56

4-26 Macroscopic image showing the gross subsurface inclusion in the 4340 steel
specimen tested at 190 ksi max stress at 59 hz and R = 0.1............................... 57

4-27 Subsurface inclusion in the 4340 steel specimen approximately 0.005" in
d ia m ete r .......................................................................... 5 8

4-28 4340 steel, 190 ksi maximum applied stress specimen. ........................................58

4-29 Much smaller slow crack growth regions and the multiple origins for the
4340 steel, 220 ksi maximum applied stress specimen.........................................59

4-30 Imbedded aluminum oxide particle at the fracture surface origin for the
220 ksi 4340 steel specim en .................................. ............... ............... 59

4-31 Aermet 100 specimen tested at 180 ksi maximum applied stress at 5 hz
w ith R = 1 ....................................................................... 6 0

4-32 300M steel specimen tested at 180 ksi maximum applied stress at 5 hz
w ith R = 1 ....................................................................... 6 1

4-33 Smeared surface features and no visible fatigue indications for the Aermet 100
sample tested at 180 ksi maximum applied stress at 5 hz, and with R = -1.............61

4-34 Interspersed regions of MVC within the slow crack growth region.of a 300M
specimen tested at 180 ksi, 5 hz, and R = -1. ................................. ............... 62

4-35 Intact coating of the 300M specimen tested at 180 ksi............... .. ............... 62

4-36 Delaminated coating around the fracture surface of the 4340 specimen
tested at 220 ksi, 59 hz and R=0.1..... ..... ................. ......... .............. ............... 63

4-37 Imbedded particle origin for the 300M sample tested at 180 ksi maximum
applied stress. ...................................................... ................. 64









4-38 Origin area for the Aermet 100 specimen tested at 180 ksi...................................64

A-1 Almen strip deflection vs. substrate part temperature. ............................................77

A-2 Substrate part temperature vs. number of mils of coating deposited
per pass of the torch. .................. ................................... ... .......... 78

A-3 Final coating hardness vs. % porosity. ........................................... ............... 79

A -4 Pow der size trends .................................................. ........ .. ...... .... 82














Abstract of Thesis Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Master of Science

FRACTURE AND RESIDUAL-STRESS CHARACTERIZATION OF TUNGSTEN-
CARBIDE 17%-COBALT THERMAL-SPRAY COATINGS APPLIED TO HIGH-
STRENGTH STEEL FATIGUE SPECIMENS

By

Donald Scott Parker

May 2003


Chair: Darryl Butt
Department: Materials Science and Engineering

Under an internationally funded research program directed by the United States

Naval Research Laboratory, thermally sprayed coatings of tungsten-carbide 17%-cobalt

are being qualified as a replacement for hexavalent chrome plating in commercial and

military aircraft applications. A complete understanding of the performance

characteristics and applied properties of these coatings (including wear, corrosion, fatigue

and residual stress) is critical for both component repair and new design configuration.

Parallel programs are underway to evaluate these coatings. Our study addresses the

residual stress state of the applied coatings, the effect of fatigue on the initial coating

condition, and crack-initiation and crack-propagation behavior at various stress levels.

These coatings are essentially anisotropic composite structures with aggregates of

tungsten-carbide particles bonded to both amorphous and crystalline cobalt phases (with

some free tungsten and cobalt suspended within the matrix). Because of the amorphous

structure and the complex nature of the metastable cobalt phases within the coating, x-ray









diffraction techniques were used to characterize only the residual stress conditions

surrounding the tungsten-carbide particles in the coating. Diffraction was also used to

establish a baseline stress state of the high-strength steel fatigue specimen substrates to

determine interfacial effects of the coating. Stress states were evaluated for specimens

that were as-sprayed. Stress states were then compared to those of specimens that had

been coated; finish machined; and then subjected to low, medium, and high stress fatigue

conditions. Triaxial stress calculation results showed significant reduction in

compressive residual stress (even a transition to tensile stress) in the radial direction

within the coatings because of the applied axial fatigue stresses.

Scanning electron microscopy was used to determine that coating cracks initiated at

surface defects present after finish grind; and propagated radially toward the substrate

along interfacial boundaries within the cobalt matrix. In this region, cracks propagate

along splat and phase boundaries around WC particles that have high residual

compressive stress. High-magnification inspection also confirmed that substrate fatigue

cracks initiate at defects along the coating substrate interface (where aluminum oxide

particles from grit-blasting the substrate are imbedded). Cracks that formed in the

coating due to the applied axial fatigue stress propagated from the surface to the

interfacial bond line between the coating and substrate; but did not provide preferential

sites for substrate fatigue initiation.














CHAPTER 1
INTRODUCTION AND BACKGROUND

Review of the Chrome-Plating Replacement Effort

Replacing hexavalent hard chromium plating operations has become a high priority

for commercial and military aircraft. This is because of the rising environmental cost of

handling this material (a carcinogen); and the poor long-term performance of this

material in critical mechanical applications. Hard-chromium electroplating is mostly

used on aircraft for landing gear components (including axle journals, hydraulic

cylinders, locking and support pins, races, lugs, and nose gear steering collars, flight

controls and access-door actuators). It is also used on sliding wear surfaces of bearing

journals, dove-tail and mid-span blade supports in turbine engines.

The Environmental Protection Agency has issued reduced allowable discharge

concentrations for hexavalent chromium from 1.71 mg/L to 0.55mg/L for existing

permitted industrial waste streams (including Department of Defense aerospace facilities)

and to 0.07 for new-permit industrial waste streams. This has created a significant

compliance issue for military aircraft overhaul depots and commercial aircraft

maintenance facilities where increased containment and treatment of the chromium waste

stream will add significant cost to long-term maintenance operations.

A technical report by the Oklahoma Air Logistics Facility [SAIC 1994] and an

Advanced Technology Report by Northwestern University, both funded under a Defense

Advance Research Project Agency (DARPA) [Northwestern 1996] contract established

that thermal-sprayed coatings are the leading candidates to replace hexavalent chromium









plating for line-of-sight applications [Sartwell 2002]. In 1996, the Hard Chrome

Alternatives Team was established by the Environmental Security Technology

Certification Program under the funding and direction of the Deputy Undersecretary of

Defense for Installations and Environment to conduct advanced development work for

qualification of high-velocity oxygen-fuel (HVOF), thermal-spray coatings for direct

replacement of chromium plating on military aircraft. After 2 years of coating-property

evaluations, tungsten-carbide 17%-cobalt was selected as the material with the best

chance of success. It was also determined at that time that the scale of the program had

implications far beyond just the US military. Thus, an international team of commercial

airlines, component manufacturers, subcontractors, and government engineers was

integrated to form the technical base for full qualification of the selected coatings.

Boeing, Lockheed, Naval Air Command, and the Air Force cognizant engineering

organizations all set forth specific evaluation criteria, fatigue-test parameters, wear-test

requirements, and flight-test conditions for each component being evaluated [Sartwell

2002].

High-Velocity Oxygen Fuel (HVOF), Thermal-spray Process

Thermal-spray coatings are applied by feeding a uniform-sized powder or wire

(metal or ceramic) through a combustion plume or electric arc field that projects molten

and semi-molten particulates through a supersonic jet stream onto a substrate as shown in

Figure 1-1. The resulting coating has a lamellar grain structure of interlocking splats

resulting from rapid solidification of small globules. Each thermal-spray process forms

distinctly different coatings with unique mechanical and physical properties. For the

combustion process, both thermal and kinetic (particle-velocity) energy is transferred










from the particle to the substrate during solidification, creating a mechanical and

interlocking diffusion (metallurgical) bond (Figure 1-2).


Compressed air Air enrelpe

Oxy-fuel mixture
Compressed air
Powder with--
nitrogen
carriergas
Shookdarcr:hds
Air cap" -
Meltir- pcirt

Figure 1-1. Exploded view of HVOF gun and spray process. Courtesy of Sulzer Metco
Inc. 1101 Prospect Ave. Westbury, NY.

Chemical reactions between the particles in the gas stream can also create

exothermic reactions that dramatically increase both the tensile bond strength at the base

metal interface, and the interlemellar strength of the coating. High velocity oxygen fuel

thermal-spray, as its name implies, mixes oxygen and a fuel gas (usually hydrogen,

propylene, kerosene or even natural gas) that are mixed and ignites them to create



2 jIr a A


ok. liDP-U S-ON


Figure 1-2. Morphology of a thermal-spray deposited coating. [England 1997]









a supersonic combustion plume through which powders are fed axially by a high-pressure

inert gas. The molten and semi-molten particles are accelerated (to approximately 1850

ft/sec) and propelled toward the substrate surface [Kim et al.2001]. The high-kinetic

low-thermal energy particles create a dense, uniform coating structure with low porosity,

high hardness, and high-strength. The uniform nature of as-sprayed HVOF coatings

allows deposition of complex carbides, cermets, and oxide-dispersion materials with high

hardness and consistent thickness; that can be ground or superfinished to provide very

smooth surface roughness and low bearing ratio coatings for sliding-wear applications

[Sartwell 2002].

Selection of Tungsten Carbide Coatings

Tungsten-carbide materials have been widely used historically to protect surfaces

from adhesive and abrasive wear in many different applications (for aerospace, paper,

and oil and gas industries). Functionally, the carbides have been used as sprayed

coatings, composite coatings, and sintered cermets; all with different fundamental

mechanical properties and behavior. For the aerospace application of replacing hard-

chrome electroplating, the beneficial wear properties of the carbides were very attractive

as compared to other materials. However, durability of the brittle coatings in fatigue-

sensitive areas was a concern. As discovered under the initial DARPA project, thermal-

sprayed coatings were ideal because of the ease and repeatability of application; and

because of the limited heat treatment of the underlying substrate required for sintered

coatings. Once established, the Hard Chrome Alternative Team (HCAT) began

evaluating several metallurgical coating combinations selected for their likely

performance regarding sliding wear, seal compatibility, atmospheric corrosion, axial

fatigue, and application-process repeatability. Process repeatability was of particular









importance because the structure and chemistry of the coating will have dramatic effects

on the as-sprayed residual stress state, distribution carbide particles in the matrix, oxide

formation, carbide dissolution, macro and micro hardness, and interfacial bond strength.

Accurate characterization cannot be performed if the chemical and mechanical properties

are not repeated for each component sprayed. The plasma spray process was eliminated

early on because of poor corrosion performance from the porous coatings; and because of

the inability to successfully develop a repeatable coating process for the carbides that

resulted in uniform coating chemistry and distribution of carbide particles. Thus, the

high velocity oxygen fuel (HVOF) process was selected as the most likely candidate to

succeed. From there, generic fatigue and wear testing was used to further reduce the

candidates. Figure 1-3 shows fatigue-test results generated for the tungsten-carbide 17%

cobalt coating as compared to the baseline chromium plating currently in use. In this

example all of the specimens were shot-peened 4340 steel hourglass fatigue specimens

heat-treated to 280-300 ksi and coated with either the HVOF WC-17%Co or electroplated

hard chromium. The fatigue behavior of the tungsten-carbide-cobalt coating is at least as

good as the chrome-plated specimen (in the axial fatigue conditions set forth in this test,

for all stress levels evaluated). Similar, data were generated for all candidate coatings.

The fatigue behavior of electroplated chromium was used as a baseline gauge of fatigue

performance. Test criteria required that to be considered successful, the average number

of cycles to failure at each stress level must meet or exceed that of the chromium

electroplating.











4340, SMALL HOURGLASS SPECIMEN
(0.003" COATING) R= -1, AIR
200.0

190.0. + EHC/Peened
EHC/Peened/FIT
80.0.
WCCo/Peened
70.0 WCCo/Peened/FIT

60.0

50.0

40.0

30.0

20.0

110.0o-

100.0 I I I I
1.E+03 1.E+04 1.E+05 1.E-+06 1.E+07 1.E+08
CYCLES TO FAILURE, N,




Figure 1-3. Fatigue results for initial coating evaluation. EHC is defined as electroplated
chromium and WC-Co refers to the tungsten-carbide 17%-cobalt HVOF
sprayed coating [Sartwell 2002].



The coatings that performed well in the initial screening were the tungsten-carbide,

17% cobalt, and the Stellite material known as Tribaloy T-400, an agglomerated blend of

cobalt, chrome, and molybdenum. Various combinations of the tungsten-carbide with

lower percentages of binder material were eliminated due to the brittle nature of the

coatings in wear tests and unacceptable fatigue debit as compared to the electroplated

chromium. Spalling and delamination of the coating occurred during the tests that further

indicated that the materials were poor candidates. Both of the remaining HVOF sprayed

coating materials, WC-17%Co and T400, were applied to high-strength steel landing gear

material hourglass coupons of 4340, 300M, and Aermet 100 and then subjected to axial









fatigue testing at R= -1, at constant amplitude with stress values from 110 to 175 ksi

maximum applied stress. Block-on-shoe wear tests with both lubricated and un-

lubricated conditions; corrosion tests that included atmospheric exposure, ASTM B 117

salt fog, GM alternate wet/dry with constant UV exposure; and electrochemical

impedance spectroscopy. The WC-17%-Co coating had higher fatigue strength and

better corrosion resistance than the T-400 coatings in every test and performed at least as

well or better than the chromium plating baseline coupons. This provided sufficient data

to show that the WC-17%-Co HVOF sprayed coating was the best candidate to replace

hexavalent chromium for the landing gear and hydraulic actuator components in

aerospace applications [Sartwell 2002].

Coating Characterization

Once the coating was selected, final analysis was to include a comprehensive

evaluation of the sprayed coating conditions that would provide a repeatable coating

process, minimal fatigue debit, and mechanical properties to optimize sliding wear

performance. Critical to this was understanding the stress conditions for which coating

cracks form, method and direction of propagation, and what effect these coating cracks

have on initiation of substrate fatigue cracks. The approach was to characterize the

material condition of both the substrate and coating before during and after applied

fatigue conditions. Standard hourglass fatigue specimens were coated with WC-17%-Co

coatings using the spray parameters developed by an L12 design of experiments that

optimized the coating for maximum fatigue life [Sartwell 2002].

The specimens were again tested at various stress levels from 110 to 175 ksi

maximum applied stress at R=-1, and then also evaluated at stress levels from 180 to 220

ksi maximum applied stress at R=0.1. Then they were evaluated using X-ray diffraction






8


techniques to characterize the residual stress state of specimens subject to increasing axial

fatigue loads. The values were compared to initial conditions obtained from virgin

samples. Additionally, scanning electron microscopy was performed to evaluate both the

coating cracks, and the fatigue fractures in the substrate materials. A detailed literature

search was also conducted to compare and contrast the data generated with past work in

this field.














CHAPTER 2
LITERATURE SURVEY

Introduction

The review of technical literature germane to this research encompassed fatigue,

residual stress, chromium electroplating and tungsten-carbide cobalt HVOF coatings, and

microstructural evaluation of tungsten-carbide cobalt coatings. The intent was to provide

a technical basis for the characterization study with an emphasis on understanding the

underlying aspects of tungsten-carbide-cobalt coating-substrate behavior under axial

fatigue conditions.

Overview of Fatigue

Fatigue is described as the process by which a component fails due to repeated

cyclic loading below the static tensile strength of the base alloy [Bannantine et al. 1990].

It has been well established, in the literature and in text that slip is the primary

mechanism by which the deformation process occurs in metals and as is seen in static

testing, the fatigue strength of a metal varies as a function of heat treatment, availability

of slip systems in the alloy, and the type of slip. Crack initiation is followed by slip band

crack growth and then growth along planes of highest tensile stress. Initiation is also

controlled in part by the surface morphology and the availability of surface defects like

porosity, machining lines, and mechanical or environmental damage that acts as a stress

concentration.

There are also two distinct types of fatigue: low cycle, referred to as load

controlled, usually described as less than 104 cycles of elastic plus plastic strain where









failure results from cyclic strain at high stress levels, even approaching the yield strength

of the material. High cycle fatigue refers to a load controlled condition usually at lower

stress levels with higher number of cycles to failure (usually > 105). To accurately

classify a material's fatigue strength, a statistical value is derived based on the probability

of a specimen attaining a specific number of cycles to failure for a given stress. Testing

is performed on multiple standardized specimens for each stress state and load condition

where a sinusoidal varying stress is applied and recorded over a specific number of cycles

[Gilliam 1969]. When generating fatigue data for specific materials, the standardized test

specimens are manufactured from the same alloy material, in the same orientation and

heat-treated condition as the component for which they are intended to simulate (axial,

rotational, bending stress etc.). Testing is performed in a servo-hydraulic machine with

loads applied along the same axis as the represented component. A computer interface

records real time test data that is plotted on a chart displaying stress amplitude vs. number

of cycles. Software modeling programs perform a mathematical analysis to create what

is known as an S-N curve from the best fit models for the data generated. For high-

strength steel, as with most ferrous alloy based materials, as the number of cycles

increases, the curve will theoretically approach an infinite value at a specific stress level,

known as the fatigue limit, where a stress amplitudes below this limit will not result in

fatigue crack initiation for a given set of conditions [Bannantine et al. 1990].

Environmental effects, as well as any surface modification, like grinding, polishing, or

coating application can significantly alter this theoretical value.

The literature shows however, that when there are circumstances whereby the

fatigue behavior is altered, for design purposes, a fatigue debit, or enhancement is









recorded for a specific applied stress value and test conditions must be quantified to

accurately manufacture a component for its intended service conditions [McGrann et al.

1998]. For example, the aircraft industry has used hexavalent chromium electroplating

for sliding wear applications and the component designers use a theoretical fatigue debit

approximately equal to a 40% reduction in predicted fatigue life where hexavalent

chromium electroplating is applied to high-strength steel components heat-treated to

greater than 220 ksi tensile strength. The fatigue process is highly sensitive to the surface

condition and the chromium electroplating process significantly alters the surface

morphology of the substrate due to the electrochemical bonding process. Shot peening of

the surface prior to a plating operation works to mitigate the effects of the tensile stresses;

however as the tensile strength of the substrate increases, the reduction in fatigue life will

again increase in spite of the effects of the peening operation [Wang 2002]. This affect is

also more pronounced at higher applied stresses, and at longer lives (higher cycles);

therefore a range of fatigue strength values are assigned to each part during full spectrum

component fatigue testing at new manufacture. Consequently, when trying to

successfully replace chromium electroplating with HVOF applied WC-17%Co thermal-

spray coatings, an entirely new set of fatigue limit levels must be established, along with

a fundamental understanding of crack initiation and propagation in the coating, the

residual stress state of the coating, and the effect it has at the substrate interface with

regards to the fatigue crack initiation in the parent metal [Nascimento et al. 2001].

Overview of Residual Stress

Residual stress is defined as the stress that remains in a material that is at

equilibrium with its surroundings without any sustained applied loads. Residual stresses

can be introduced into components by a variety of forging, welding, heat-treatment and









surface treatment processes, and during the deposition of multi-phase and composite

coating materials. These stresses can significantly affect the fatigue strength and fracture

characteristics of engineering components; therefore, it is essential that accurate

information is available of the residual stress distributions present for reliable estimates

of their useful lifetimes. Residual stress states are characterized by the scale over which

they self-equilibrate and can be broken down into three categories: type I are defined as

macrostresses that arise from thermal or elastic mismatch and vary over long distances

relative to the size and shape of the component. Type II refers to intergranular stresses

that vary from grain to grain like coherency strains resulting from alloying elements, or

multiple phase materials like precipitation hardening steels or the distribution of carbide

particles in a cobalt coating matrix. Type III residual stresses are those that affect areas

smaller than grain size to the atomic level and encompass particle boundaries, phase

interfaces, and intersplat boundaries. Type II and type III stresses must balance over their

characteristic small distances and are referred to as microstresses. Surface microstresses

are often induced to counteract solidification or thermal tensile stresses, or to enhance the

localized susceptibility for fatigue crack initiation [Whithers and Bhadeshia 2001].

There are several methods of evaluating these types of residual stresses including

x-ray diffraction, hole drilling, boring, deflection (modified layer removal) sectioning,

and neutron diffraction and each has positive and negative attributes that affect resolution

and accuracy of the calculated stress. X-ray diffraction is used because it can detect type

I, II, and III stresses non-destructively by calculating the elastic strains E through the

change in the Bragg scattering angle AO, where X=2dsinO. Solving for strain E=Ad/d, and

with accurate knowledge of the initial value of d, the stress free interplanar spacing, the









strains can be converted to stress using the stiffness equations. For metals and coatings,

this can provide a very accurate determination of the stress values associated with type II

and III stresses [Whithers and Bhadeshia 2001] One of the drawbacks of this technique

is the limited penetration depth of the x-ray beam, which restricts evaluation to the near

surface structure, however understanding the near surface stress behavior is important to

understand fatigue crack formation and propagation. For x-ray diffraction stress analysis,

roughly half of the diffracted radiation originates from less than 0.0004" beneath the

surface. However, the x-ray beam is attenuated exponentially as a function of depth and

this rate of attenuation is controlled by the linear absorption coefficient, the composition

and density of the specimen, and the type of radiation used. Therefore, surface

microstress profiles tend to be exponentially weighted averages of the stress at the

surface and in the coating layers immediately beneath it. For coating materials, the net

macrostress may appear to be continuous over the entire thickness of the coating, but the

microstresses form a stress gradient from the surface, along individual splats, phase

boundaries, and particle interfaces to the substrate bond line. To compensate for

potential errors from different grain orientations, and increase the accuracy of the

calculated stress, a series of scans are performed to locate at least six independent strain

measurements for rotations of both D and angles. There are other inherent difficulties

in this technique for coatings, which have multiple elements, oxides, carbides, amorphous

phases, and potentially multiple crystallographic structures as outlined by Prevey (1991).

However, measurement of stress values associated with known constituents within a

coating and along the substrate surface can provide a relative understanding of the

internal stress state of the individual phases and particles and their relationship to crack









initiation and propagation. For coating materials, internal gradients of tensile and

compressive stress will tend to direct the crack path and the propagation rate of a coating

fracture. [Nascimento et al. 2001].

Another non-destructive technique using neutron diffraction employs the same

mathematical model for determining the strains and has increased penetration depths for a

more accurate stress profile as a function of depth. However, certain materials like the

cobalt matrix in tungsten-carbide coatings become radioactive upon bombardment by

neutrons thus rendering the samples unusable after exposure. Destructive methods of

determining residual stresses require measurement of the elastic change in a sample as

layers of material are removed from the surface. The two dimensional elastic strain is

recorded by strain gauges placed on the opposite surface and a stress gradient profile is

established as layers are removed. Measurement errors can be introduced by the

mechanical or chemical methods of removing the surface layers and thus accurate

readings are difficult to quantify [Prevey 1991].

Overview of Shot Penning

Shot peening of the base metal is a cold working process that works to mitigate the

effects of surface tensile stresses created by the machining, forming, or coating process

by creating a compressive residual stress zone at and slightly below the metal surface that

retards fatigue crack initiation. The impact of the shot media acts as a tiny hammer,

imparting to the surface small indentations or dimples. In order for the dimples to be

created, the surface fibers of the material must yield in tension whereas below the

surface, the fibers try to restore the original shape, thereby producing a hemisphere of

cold-worked material below the dimple in highly compressive stress, as much as half the

yield strength of the material. Overlapping dimples develop an even layer of metal in










residual compressive stress that tends to inhibit the formation and growth of any surface

cracks or defects as shown in Figure 2-1.





I

I A I -el







/ Elst- Defonned
Elastic -.- Configuration
region .-
c eLiginal
Strnu, 3tleS
Hardness Ue. M
Macroscopic Deformation
(Forming Effects)


Figure 2-1. Model of shot peened surface showing the stress profile created by the
impact dimples on the surface [Wang and Platts 2002].

The higher strength materials (higher hardness) resist the plastic deformation of

peening and therefore have a much shallower compressive layer than lower strength

materials. Consequently, as the tensile strength of the substrate increases, the benefits of

shot peening will diminish under applied loads because elastic deformation will facilitate

relaxation of the induced residual stresses at the surface. This affect is also more

pronounced at much higher applied stresses, and at longer lives (higher number of fatigue

cycles) [Bannantine et al. 1990].

Electroplated Chromium and HVOF Applied Coatings

Hard Chrome plating is an electrolytic process utilizing a chromic acid-based

electrolyte. The part is made the cathode and, with the passage of a DC current via lead









anodes, chromium metal builds on the component surface. The chromium electroplating

process generates significant tensile stresses both at the substrate surface, and within the

plated structure. Upon solidification, the shrinking of chrome deposits due to hydrogen

gas diffusing away and as the decomposition of the intermediate chromium hydride

structure creates volume changes that result in tremendous tensile stresses that are in

excess of the ultimate tensile strength of the chromium. As a result, web-like cracks form

throughout the plated structure to relieve the internal stresses in the coatings, however,

this also generates net tensile stresses at the substrate interface that will act to exaggerate

surface defects and become preferential sites for fatigue crack initiation [Nascimento et

al. 2001].

Unlike to chromium electroplating, high velocity oxygen fuel (HVOF) applied

thermal-spray coatings can generate net residual macrostresses, or type I stresses, that can

be compressive, neutral or tensile, depending on the spray parameters used to apply the

coatings, and will be relatively uniform across the bulk of the coating. The solidification

stresses from unmelted particles, porosity, carbide decomposition, or incomplete splat

formation can constrain the coating matrix in tension; whereas, an optimized coating with

low carbon dissolution, and uniform melting of the cobalt binder will create net

compressive macrostresses within the coating and along the substrate interface. Fatigue

testing confirms the relationship between low levels of net compressive macrostresses in

applied tungsten-carbide cobalt coatings and improved fatigue performance in axial stress

conditions. During fatigue crack growth, the near threshold and high growth rate regimes

are strongly affected by mean compressive type I stresses which may delay the onset of

plastic deformation and crack formation on the surface [Whithers and Bhadeshia 2001].









Therefore, when combined with shot peening of the substrate prior to application,

tungsten-carbide cobalt coatings on high-strength steel with a net compressive residual

macrostress can approach fatigue strengths near uncoated base metal levels because both

the substrate interface, and the free coating surface are in compression [Tajiri et al. 1998].

It has also been shown in the literature that when HVOF coatings are applied with a net

tensile stress, the effect on the fatigue behavior of a high-strength steel substrate is as

detrimental as that of electroplated chromium. Since any residual stresses can raise or

lower the mean stress experienced over a few fatigue cycles, the tensile stresses at a free

surface, or at the coating substrate interface, will accelerate the onset of fatigue crack

formation and is therefore undesirable [Reiners et al. 1998].

Microstructure of the Tungsten Carbide Cobalt Coatings

The structure of the initial powder consists of angular particles of WC

agglomerated and sintered to a cobalt binder with nearly spherical net powder grains.

When the powder is fed through an HVOF combustion plume, the particles experience a

very short dwell time in the flame, which allows for maximum retention of WC,

however, the kinetic energy imparted causes the particles to deform and become almost

spherical in the final coating. The highly oxidizing environment of the plume also reacts

with the cobalt to form multiple metastable oxides including and amorphous phase that

are retained due to the extremely rapid solidification [Verdon et al. 1998]. Analysis of the

microstructure of an HVOF sprayed WC-17%Co coating using optical and scanning

electron microscopy shows the lamellar morphology of the multiphase cobalt binder and

the uniform distribution of tungsten-carbide particles within the coating matrix (Figures

2-2, 2-3, and 2-4). From reviewing the available literature, it was determined that there









are several techniques that incorporate x-ray diffraction and mathematical deconvolution

that are used to predict what phases and structures are present in the coating


Figure 2-2. Cross-section showing the coating microstructure of an as-sprayed WC-
17%Co coating applied to a high-strength 4340 substrate. 500x magnification.


Figure 2-3. Polished WC 17%Co coating cross-section showing the lamellar structure
and uniform distribution of WC particles (High magnification SEM
micrograph).




























Figure 2-4. Higher magnification micrograph of the same sample in Figure 2-3 showing
the white spherical WC particles suspended in the lighter gray cobalt matrix.
The dark gray wavy structures are bands that form in between splat layers and
around WC particles during solidification that contain complex cobalt oxides,
free cobalt and minute amounts of free tungsten.

microstructure. A Rietveld refinement method of least squares deconvolution can be

performed in an attempt to resolve and identify overlapped peaks in masked regions or

where peak broadening is observed. This analysis incorporates a statistical probability in

an attempt to speculate what phases and structures should be present based on the physics

of the coating process and the chemistry of the powder materials sprayed. Precise x-ray

mapping of the coating structure chemistry is subject to interpretation of statistical results

with a probability model prediction and identification of most likely compounds [Prevey

1991]. The results are listed here based on the data from various x-ray mapping attempts

by the referenced authors as well as in the work performed for this thesis. The spherical

and semi-spherical white particles are the distributed tungsten-carbides (WC), some of

which have very small, dark (almost black) patches of W2C sub-carbide that results from

dissolution of the WC during the coating deposition process. The undesirable W2C


20kU X3,700 5mm 0000 26/FEB/03









particles should be minimized due to their brittle nature and the adverse affect on the

overall fracture toughness of the coating. The constituents of varying shades of gray are

the different metastable structures of cobalt and cobalt oxides, including an amorphous

cobalt phase that form and are trapped during rapid solidification. X-ray diffraction

failed to adequately identify the exact chemistry of the oxides, partially due to the

amorphous structure of some and partially due to overlapping peaks of the metastable

cobalt compounds listed in the JCPDS database. The very small white particles are

cobalt carbides that take up the remaining carbon when dissolution of WC occurs and

formation ofW2C is either retarded or kinetically unfavorable [Semant 1998]. Porosity is

seen as very dark black angular holes and is located along triple points where multiple

phases intersect carbide particle boundaries. Very small quantities of free tungsten and

cobalt are indicated by XRD but are difficult to identify optically [Verdon et al. 1997].














CHAPTER 3
EXPERIMENTAL PROCEDURE

This chapter discusses the experimental procedures performed in the

characterization of the WC-17%-Co HVOF sprayed coatings on high-strength steel

specimens. The study was divided into two parts. The first involved x-ray diffraction

residual stress measurements of coated specimens before and after fatigue testing. The

second part involved optical and scanning electron microscopy of the coupons after being

subjected to fatigue conditions at low, medium and high stress levels.

Specimen Manufacture

The hourglass fatigue bars were manufactured by Metcut Research of Cincinnati,

OH from three common landing gear alloy materials, AISI 4340, 300M, and AerMet 100;

the compositions are outlined in Table 3-1. Both 4340 and 300M are considered to be

high-strength low alloy steels with excellent hardenability due to their appreciable

amounts of carbon, nickel, chromium, and molybdenum. Both materials have a similar

combination of strength and toughness (50-55 ksi root inch at 280 ksi) over a wide range

of section sizes and display uniform microstructures throughout the hardenability range

[ARP 1631]. However, 300M is essentially a modified 4340 with higher contents of

silicon to retard any cementite formation and reduce temper embrittlement; molybdenum

to reduce grain boundary segregation; and additions of vanadium to improve the

resistance to softening during tempering operations and to form carbides which reduce

austenite grain growth. Aermet 100 is a high cobalt alloy designed by Carpenter

Technology with improved fracture toughness (100 ksi root inch) and higher resistance to










stress corrosion cracking at strength levels greater than 280 ksi, Figure 3-1 [AMS 6532B

2001]. When heat-treated to this high tensile strength condition, both 4340 and 300M

have a uniform martensitic microstructure, whereas Aermet 100 is a highly alloyed

martensitic age hardened steel that is vacuum melted and refined for a contaminant free,

ultrafine grain size microstructure that is significantly more ductile than the other two

materials. Standard published tensile strength data shows that Aermet 100 has nearly

15% elongation and 55% reduction in area; whereas both 4340 and 300M will experience

less than 10% elongation and less than 30% reduction in area, Figure 3-2 [AIR5052

1997].

Fracture Toughness


1410
120





0
AF1410 AwrMl 100 HyTOt Mar al 260 3)OM 4A40 Htl
Alloy



Figure 3-1. Fracture toughness data for 4340, 300M, and Aermet 100 showing the
significant improvement for Aermet 100 over the other alloys. Data is for 250
ksi tensile strength specimens [AIR5052 1997].

The hourglass fatigue specimens were cut to length from bar stock then turned on a

lathe into standard configuration (Figure 3-3). Next they were heat-treated to 280-300

ksi verified tensile strength with the 4340 and 300M specimens heat-treated in

accordance with Military Specification MIL-H-6875 while the Aermet 100 was heat-

treated in accordance with Carpenter Technology Process Specification 15169.










Elongation
E lor nand % R.A.


and Reduction in Area


AFU11) I r1fio1 Ehrt16darn 21r 5 [Wt icfri In 11 aam

Ekongall-Ron = %W Reditntion In Aria


Figure 3-2. Ductility properties for tensile elongation and reduction in area data for
Aermet 100 versus 4340, and 300M (250 ksi) [AIR5052 1997].


Table 3-1. Composition of the three alloys.
Element Alloy 4340 Alloy 300M Alloy Aermet 100


Carbon
Manganese
Silicon
Chromium
Nickel
Molybdenum
Copper
Phosphorus


Sulfur


Vanadium
Cobalt
Aluminum
Titanium
Oxygen
Nitrogen


0.38-0.43
0.65-0.85
0.15-0.35
0.70-0.90
1.65-2.00
0.20-0.30
0.35 max
0.015 max


0.008 max


0.40-0.45
0.60-0.90
1.45-1.80
0.70-0.95
1.65-2.00
0.30-0.50
0.35 max
0.010 max


0.010 max


0.05-0.10


0.21-0.25
0.10 max
0.10 max
2.90-3.30
11.0-12.0
1.10-1.30


0.0080 max
P+S<0.010
0.0050 max
P+S,0.010


13.0-14.0
0.015
0.015
0.0020 (20 ppm)
0.0015 (15 ppm)










Grinding to final dimension was performed in accordance with the low stress

procedures outlined in Military Standard MIL-STD-866 with all specimens undergoing

Nital Etching as specified in Military Standard MIL-STD 867 to examine for any

grinding burns. Upon completion of the non-destructive inspection, the specimens were

then baked at 3500 F for 8 hours to remove any potential residual hydrogen from the

etching process. The center gage section was then shot peened to an intensity of 8-12

Almen "A" using S230 wrought steel shot in accordance with AMS 2432 using computer

control for accurate coverage [Sartwell 2002].





0.5" coa tng
patch

















Figure 3-3. Hourglass fatigue specimen and coated section that was used for the study.
[Sartwell 2002]











aligned on centers in a lathe type fixture for rotation during the spray operation. Hard
N_ W-\tc ut Recem-ch Inc



Figure 3-3. Hourglass fatigue specimen and coated section that was used for the study.
[Sartwell 2002]

The coated center gage section was prepared for spraying by abrasively blasting

with 54-grit aluminum oxide particulate propelled at 60-80 psi through a vacuum blaster

to coarsen the surface to 120-150 Ra for coating adhesion. The specimens were then set

aligned on centers in a lathe type fixture for rotation during the spray operation. Hard









shadow masking was offset on each end of the gage section to produce a tapered coating

at the termination point. A Sulzer-Metco diamond jet TM HVOF spray system was used

with and x-y gun traverse unit and hydrogen gas for the combustion fuel. The powder

material was also from Sulzer-Metco and was designated Diamalloy 2005 TM, 83% WC

and 17%-Co and was manufactured by the agglomeration and sintering process [AMS

7881 2003]. The powder particle size distribution is shown in Table 3-2 and the process

parameters are outlined in Appendix A. The material was fed axially through the center

of the combustion gun at 80 psi positive pressure with nitrogen gas where it mixed with

the high pressure combustion gases. The combustion plume of burning gases creates a

high velocity stream of accelerated molten cobalt and semi-molten tungsten-carbide

particulate towards the surface. Each particle creates a splat upon impact with the

substrate building a lamellar coating structure as the HVOF gun is traversed. The cobalt

dissociates from the agglomerated powder particles and oxidizes in the flame creating

multiple cobalt rich metastable phases that encompass the binder of the coating. The

spherical tungsten-carbide is uniformly distributed throughout the coating matrix upon

solidification

Table 3-2. Particle size distribution for the Sulzer Metco, Diamalloy 2005 TM,
agglomerated and sintered WC 17%Co powder by sieve size.
Sieve Size Minimum % Maximum %

+270 (2.0876 mils) -6%
+325 (1.7716 mils) 25%


The coating was deposited at a rate of 0.0001" thickness per for pass with 10-

second intervals between each deposition pass to keep the substrate temperature below

2750 F (135 C) during spraying. The final as-sprayed coating thickness was 0.006" and









was uniform over the gage length to within 0.0005" concentric relative to the cylindrical

substrate as measured with a Fisher Scientific eddy current thickness gauge. The coating

was finish ground between centers to 0.0035" and a surface roughness of 6 Ra using a

conforming, 320 grit, diamond impregnated grinding wheel on a Cincinnati Milicron

bench grinder. The final coating thickness of 0.003" +/- 0.0005 was achieved by

polishing between centers with diamond abrasive paper on a conforming hub to a surface

finish of 2-4 Ra.

Coating Quality Control

Quality control coupons were produced with the coated fatigue specimens to verify

the coating properties. The results are as follows: hardness of the coating averaged

HV300g 1050 across the coating thickness as measured on a polished cross-section using a

diamond indenter and Wilson microhardness tester in accordance with ASTM E384.

Tensile bond strength was measured in accordance with ASTM C633 with the coating

sprayed onto the face of a 1" diameter by 14" thick cylindrical blank to 0.005" thickness.

The coupons are then bonded to a set of internally threaded 1" diameter by 2" long

cylindrical blanks using 1" diameter wafers of Cytec Industries FM 1000 film adhesive.

The bonded cylinders were then aligned vertically in the threaded fixtures of a servo

hydraulic tension/compression machine and then pulled at 0.05 inch/min in tension until

failure or the coating or adhesive. The results averaged greater than 13,000 psi, with

ultimate tensile failure occurring in the film adhesive. A coating is considered acceptable

if the cohesive tensile strength (inter-lamellar, inter-particle and inter-splat structural

bond) and adhesive strength (substrate bond integrity) exceeds 13,000 psi. The coatings

were evaluated to gauge the degree of residual stress in the as-sprayed condition by

coating standard Almen "N" strips in identical fashion as the fatigue specimens. The










strips are grit blasted on both sides prior to spraying to both prepare the surface for

coating, and minimize the pre-stressed deflection in the strip. The pre-sprayed deflection

should be within 0.002" of being flat to minimize error in reading final arc deflection.

The coupons are mounted in the transverse direction on a screw type fixture in

accordance with AMS S13165, as shown in Figure 3-4, and then sprayed to a thickness of

0.005". An unrestrained convex deflection indicates compressive residual stress and for

these specimens a 0.010" positive net arc height deflection was achieved. It was shown

through an L12 design of experiments correlating the spraying parameters to almen arc

height and fatigue performance that an acceptable deflection is between 0.003" and

0.012". The L12 design of experiments is contained in appendix-a along with the final

spray parameter set.





Almen test
strip
Clamping
bolts

Mount

Initial substrate strip



Post spray Coating


Resulting I
Deflection Substrate

Figure 3-4. Screw type holder used for residual stress measurement during spray
operations [AMS S13165 1997]. The convex deflection confirms
compressive stress and the unrestrained arc height provides a degree of
residual stress obtained.









Residual Stress Measurement

The residual stress experiment was conducted jointly with Oak Ridge National

Laboratory under a facility User Agreement with the UT-Battelle High Temperature

Materials Laboratory Group, HTML Proposal No. 2002-032. The equipment utilized is a

Scintag PTS Goniometer, with a liquid nitrogen cooled Ge detector. This experimental

setup also contained a MAC Science, 18 kW rotating anode generator with a Osmic CMF

multiplayer mirror (Figure 3-5).



















Figure 3-5. Scintag PTS Goniometer at Oak Ridge National Laboratory used for the
residual stress measurements.

Baseline 0/20 scans were run on the bare metal uncoated sample, on both a coated

specimen as sprayed, and sprayed and finish ground, virgin WC-17%-Co powder, and

virgin WC powder. These scans showed that the best peaks to use when evaluating the

internal coating residual stress were the WC 117 (211 diffraction plane), and 121 (103

diffraction plane) 20 peaks because they were clearly identified in the database as WC

peaks, they did not have any overlapping with other peaks, and they had sufficient

intensity for an accurate analysis. The cobalt peaks were not well defined in relation to










the database for known CoO compounds, and the clearly identified peaks for cobalt and

its oxides were overlapped by other peaks like the amorphous region between 30 and 50

degrees 20 as shown in Figure 3-6. The amorphous region is a metastable phase

consisting of ternary Co, W, and C that is suspended due to the rapid undercooling. The

molten material could experience a glass transition since it has been shown that upon

annealing above 10000 F (5400 C) the binder will recrystallize [Verdon et al. 1997].

[C 11860 rawj Coating on fatigue sample # A1-42

:' ?' 51-0939> Unnamed mineral, syn [NRJ WC
2000 o 35-0776> W2C Tungsten Carbide
S05-0727> Co Cobalt
15-0806> Co Cobalt


1500
0

0
0
Z1000- C

C
1 0
C I






10 20 30 40 50 60 70 80 90 100 110 120 130 140 151
2-Theta(wM



Figure 3-6. Theta/2-theta scan results for the coated specimen. The 117 and 121 2-theta
peaks were resolved for both the virgin WC powder and the sprayed coatings
so stress scans were performed at these 2-theta angles.

Therefore, it was decided that the experiment would determine the three-axis stress

tensors for the WC particles in the matrix for the as-sprayed, finish ground (both

untested), and three fatigue tested specimens subjected to low (110 ksi), medium (150

ksi), and high (220 ksi) axial fatigue stress levels. Lattice spacing values were obtained









from the initial scans on the virgin WC powder and reduced potential errors in the

mathematical calculations.

The stress scan samples were mounted such that when the calculations were

performed, the three-axis stress tensor would indicate axial, hoop, and radial stresses

present in the coating (Figure 3-7).























Figure 3-7. Specimen installation and alignment in the Goniometer at Oak Ridge
National Laboratory.

The height alignment was accomplished with a dial gauge probe and telescope,

which was accurate to +/- 5 microns. The alignment was confirmed by rotation of the

diffracting surface of the specimen 180 and observing that the surface was coincident

with the horizontal axis at +/- 900 of angle X. Goniometer alignment was verified by

examining a LaB6 powder on a zero background plate, and measuring the peak shift of

the (510) reflection. The maximum peak shift was 0.010 20, for co tilting and the final x-

ray beam size at the aperture measured 0.6mm x 4mm. Lead foil tape was used at the









ends of the specimen gage section to block returns outside of the desired diffraction

region (Figure 3-8).





















Figure 3-8. Orientation of the sample holder and the lead foil masking in the goniometer.

Table 3-3 lists all of the experimental conditions for the x-ray measurements. Once

the specimen and goniometer were aligned, the 0/20 stress scans were performed from

1150 < 20 < 1230 and increments of 4 at 00, 450, and 900. The goniometer is also rotated

through x for the angle +/- 55, 42, 28.2,and 0, for each increment of angle 4, resulting in

21 scans per sample at 0.02 020/step; 10 s/step. The total time required was

approximately 23.5 hours per scan (6 samples evaluated) plus sample alignment time.

Calculations were performed using Dolle-Hauk method assuming a uniaxial stress state

[Krause and Hasse 1986]. The mathematical equations are as follows:

l = (1+v)/E y11 sin2 vol1/E assuming o12 =22= 13= 23= 33=4=0.

The variables s, d, v, E and c are the strain, interplanar spacing, Poisson's ratio,

Young's modulus and stress, respectively. The variables and subscripts 4, W and 0 refer









to the azimuthal angle, tilt angle and strain-free, respectively and d=0o,ro was taken as the

strain free interplanar spacing, do; however, the unstrained value was determined

experimentally by x-ray diffraction for the WC virgin powder. Values for Poisson's ratio

and Young's modulus were obtained from the literature [Noyan and Cohen 1987].

Table 3-3. Experimental conditions for the x-ray measurements
Parameter Condition
Equipment MAC Science 18 kW rotating anode generator
Osmic CMF multilayer mirror
Scintag PTS goniometer
Scintag liquid N2-cooled Ge detector
Power 8 kW; 40 kV, 200 mA
Radiation Cu, = 1.54059 A
Incidence slit width 0.5 mm
Receiving slit acceptance 0.250; radial divergence limiting (RDL) Soller slit
Source to specimen distance 360 mm
Specimen to back slit distance 280 mm
Tilt axis and angles co; 0, 28.2, 42, 55 (equal steps of sin2y)
Scans 0.02 020/step; 10 s/step


Fatigue Testing

Fatigue testing of the specimens was performed by Metcut Research at their

Cincinnati Testing Laboratory as part of a parallel research project for the Naval

Research Laboratory. For this thesis all samples provided were tested in accordance with

ASTM E466-96 to generate standard S-N curves for high cycle axial fatigue (load

controlled) with constant amplitude.

This evaluation classifies low, medium and high stress levels as 110 135 ksi low;

140 175 ksi medium; and 180 220 ksi high. Each material was subjected to fatigue

loads corresponding with the test matrix designed by Naval Air Command Stress

Engineering Group. The assigned values are based on actual aircraft components

manufactured from specific alloys that were being evaluated for replacement of chrome









plating with tungsten-carbide cobalt coatings. After testing, the data were plotted in the

standard manner with stress on the vertical axis and cycles-to-failure on the horizontal

axis. A least squares regression was used to produce each S-N curve. The regression

involved a linear fit to the data in In (N) vs. In (c) space, then calculating a best fit curve

in the traditional S-N space. All samples were manufactured and prepared as indicated

above and were tested until failure, or run out at 107 cycles. The raw data set from which

the fatigue results were extracted for this evaluation is contained in Appendix-b [Sartwell

2002].

Scanning Electron Microscopy

Scanning electron microscopy was performed on a JEOL JSM 6400 at the

University of Florida, Major Analytical Instrumentation Center, along with a JEOL JSM

5900LV at NASA, Kennedy Space Center, Florida. The samples were evaluated at

various magnifications to evaluate coating surface morphology, coating crack origin

analysis, coating fracture surface feature identification, and substrate fatigue crack origin

analysis.














CHAPTER 4

RESULTS AND DISCUSSION

Residual Stress Measurements

When materials are subjected to deformation, slip occurs and the microstructure

deforms in the direction of the applied stress. Each grain in the structure therefore has

regions of the lattice structure that are elastically strained in a state of tension or

compression. In x-ray diffraction, this causes a shift, or broadening, or both to the

diffraction lines associated with the diffraction planes on a plot of intensity vs. 2-theta

angle. The elastic strain (strain tensors) surrounding this deformation can then be

calculated from the change in interplanar spacing (Ad/d) obtained from experimental x-

ray diffraction experiments [Prevey 1991]. For the d spacing, plotted vs. sin2W at all y

angle tilts that are linear, the mathematical matrix of stress tensors can be solved by

inputting data from +/- 3 distinct D angle tilts (0, 45, and 90 degrees), plotted for all y

values. The values for the strain tensors determined from the x-ray scans, Poisson's ratio,

and Young's modulus are then input into the mathematical equations (Excel spreadsheet)

to calculate the tri-axial stresses from Hooke's Law [Noyan and Cohen 1987].

For this experiment, alignment of the sample was defined such that the axial, hoop,

and radial stresses could be extracted for the base metal gage area that was shot-peened,

and then blasted prior to spraying; and for the WC particles in the coating matrix (2-theta

diffraction peaks at 117, and 121 degrees) both before and after the samples were

subjected to fatigue testing. The data for all of the x-ray scans is compiled in Appendix C









along with the spreadsheet calculations for the directional stresses. The 4340 high-

strength steel substrate material was evaluated first to ascertain the baseline condition of

the substrate prior to spraying and to ascertain the baseline stress condition on the surface

of the shot-peened material. The results are shown in table 4-1. Multiple alloy

combinations were not evaluated due to the amount of time required for each scan.

Table 4-1. Shot-peened bare 4340 steel substrate material
Condition Axial Hoop Radial
Shot-peened, 4340 steel 133 ksi 137 ksi 22 ksi
substrate Compressive Compressive Compressive


For the bare metal shot-peened specimens, the calculated values showed

compressive residual stress in all three axes, with the radial direction being the lowest

value. This measure for the radial stress is expected, because the shallow penetration of

the x-ray beam is measuring a value for a radial stress that is defined as being zero at the

free surface and most of the effects of the shot peening on the radial stress will be seen as

a gradient beneath the free surface [Prevey 1991]. Once the substrates were blasted with

aluminum oxide, the three dimensional surface stress state was altered such that the net

residual stress was slightly more compressive in the axial direction, slightly less in the

hoop direction, and near neutral in the radial direction as shown in Table 4-2.

Table 4-2. Shot-peened bare 4340 steel substrate material that was blasted with 54 grit
aluminum oxide prior to HVOF coating.
Condition Axial Hoop Radial
Shot-peened, 4340 steel 154 ksi 89 ksi 3.1 ksi
substrate blasted with
substrate blasted with impressive Compressive Compressive
Aluminum oxide grit



The effect of the grit-blasting altered the stress state at the surface due to the

change in surface morphology from the smooth dimpled appearance to the coarse surface









formed by the blast media. As a result of the impact, the blast media creates small

regions of tensile stresses between the peaks and valleys of the tooth-like profile on the

surface. In the case of a shot-peened surface, the impact velocity of the grit particles will

act to increase the peening effect on the stress in certain directions while the cutting of

the free surface will tend to relieve some of the near surface stresses in other directions.

The near surface hoop and radial stresses, which are critical to crack initiation resistance

during axial fatigue, were significantly effected and will tend to reduce the effects of the

shot-peening on the substrate. Next, the surface of the as-sprayed and finish ground

coated specimen was evaluated prior to fatigue testing in an effort to determine the

effects of the grinding and polishing process on the residual stress state of the particles

within the coating matrix.

The as-sprayed specimens showed a surface compressive stress for the coating in

all three directions as shown in Table 4-3, however, the radial stress value is higher than

expected and could be attributed some errors due to the coarse surface profile of the as-

sprayed coating (150Ra surface roughness).

Table 4-3. Residual stress data for the WC 17%Co HVOF coating in the as-sprayed
condition.
Condition Axial Hoop Radial
As-sprayed WC-17%Co 85 ksi 116 ksi 99 ksi
coating. 117 29, 211 peak
coating. 117 20, 211 peak Compressive Compressive Compressive

As-sprayed WC-17%Co 50 ksi 73 ksi 76 ksi
coating. 121 20, 103 peak Compressive Compressive Compressive


The finish ground coated specimens showed a significant jump in two directions

increasing in both the axial and hoop directions; however the radial compressive stress

decreased, as would be expected since the as-sprayed radial stress value seemed slightly









high. The surface roughness of the finish ground coating is 2-4 Ra as measured with

Fisher stylus profilometer and the smoother finish should provide a more accurate

measure of the near surface three-dimensional stress field. This indicates that the

grinding and polishing process used to achieve the desired surface finish have a

significant effect on the residual coating stresses imparting significantly more

compressive stress in the axial and hoop directions.

Table 4-4. Residual stress data for the WC 17%Co HVOF coating in the finish ground,
and polished condition.
Condition Axial Hoop Radial
Finish ground WC-17%Co 264 ksi 224 ksi 33 ksi
coating. 117 29, 211 peak
coating. 117 20, 211 peak Compressive Compressive Compressive

Finish ground WC-17%Co 193 ksi 245 ksi 22 ksi
coating. 121 20, 103 peak Compressive Compressive Compressive


The coated 4340 steel specimen subjected to high cycle fatigue testing at low

maximum applied stresses loads (110 ksi) showed decreased compressive residual

coating stresses in all three axes for the 117 20, 211 peak as shown in Table 4-5. This

result was expected, most likely due to relaxation of the residual stresses during the

elastic deformation experienced during fatigue testing. The 4340 specimens tested in the

range did not fracture and were run-out specimens that reached 107 cycles. The coatings

only showed very fine cracks distributed across the surface that were not interconnected

and likely caused the reduction in residual compressive stress at the surface. The results

for the 121 20, 103 peak were skewed due to the shallow peak derived from the scan.

The 121 peak was very shallow with low intensity and therefore the data shown

indicating a tensile stress is likely to be inaccurate. The specimen subjected to fatigue









testing at medium applied stresses loads (150 ksi) showed increased compressive residual

coating stresses relative to the 110 ksi

Table 4-5. Low-applied stress specimen.
Condition Axial Hoop Radial
Fatigue Tested at 110 ksi 158 ksi 146 ksi 12 ksi
WC-17%Co coating. 117
SCompressive Compressive Compressive
20, 211 peak
Fatigue Tested at 110 ksi 65 ksi 44 ksi 56 ksi
WC-170Co coating. 1210 Tensile Tensile Tensile
20, 103 peak


specimen in all three axes for the 117 20, 211 peak as shown in Table 4-6. The

121 20, 103 peak showed compressive stress in the axial and hoop directions but slightly

tensile in the radial direction. The 4340 specimens tested in this range did not fracture

and were run-out specimens that reached 107 cycles. The coatings only showed very fine

cracks distributed across the surface that were not interconnected and likely caused the

reduction in residual compressive stress at the surface. The coated 4340 steel specimen

fatigue tested the highest stress level of 220 ksi maximum applied stress were evaluated

at a section of coating adjacent to a delaminated region on the same side as the fracture

surface origin. The stresses were slightly lower compressive values than the 150 ksi

specimen for the axial and hoop directions with slightly higher compressive stress values

for the radial direction as shown in table 4-7.

Table 4-6. Medium-applied stress specimen.
Condition Axial Hoop Radial
Fatigue Tested at 150 ksi 280 ksi 215 ksi 21 ksi
WC-17%Co coating. 117
SCompressive Compressive Compressive
20, 211 peak
Fatigue Tested at 150 ksi 140 ksi 130 ksi 27 ksi
WC-17%Co coating. 121
WC-17Co coating. 121 Compressive Compressive Tensile
20, 103 peak









Calculations for the 117 20, 211 peak again showed slightly higher values for

compressive stress than thel21 20, 103 peak in the axial and hoop directions, but

comparable values for the radial direction.

Table 4-7. High-applied stress specimen.
Condition Axial Hoop Radial
Fatigue Tested at 220 ksi 219 ksi 125 ksi 32 ksi
WC-17%Co coating. 117
SCompressive Compressive Compressive
20, 211 peak
Fatigue Tested at 220 ksi 100 ksi 33 ksi 25 ksi
WC-17%Co coating. 121
WC-7Co coating. 121 Compressive Compressive Compressive
20, 103 peak


Optical and Scanning Electron Microscopy

High magnification inspection began with evaluation of the surface morphology of

the sprayed and ground coatings for comparison with the fatigue tested specimens.

Under applied stress, fatigue crack initiation will occur at a site of high stress

concentration and may well occur at stress values below theoretically expected values

[DePalo et al. 2000]. The image in Figure 4-1 shows the ground and polished coating in

the gage section of the specimen prior to fatigue testing. The process uses diamond

abrasives to achieve a 2-4 Ra microfinish that implies a very smooth, mirror-like surface.

However, concentric machining lines, surface porosity and particle pullout are visible on

the coating surface, as shown in Figure 4-2, before any stress loads are applied. These

inherent defects can provide sufficient stress concentration and preferential sites for

coating cracks to initiate [Ahmed and Hatfield 1999]. The objective was to determine the

origin of the fatigue cracks for the coated samples subjected to low, medium and high

ranges of applied axial fatigue stress and determine what effect the coating application

had on the fatigue life of the three substrate materials.






















Figure 4-1. Specimen before fatigue testing



















Figure 4-2. Surface morphology of a finish ground coating that was polished to a 2-4 Ra
roughness. Defects and machine lines are still evident (Scanning electron
microscope micrograph).

Low-Applied-Stress Specimens

Two results were obtained for this range of applied loads. The first included

specimens that did not fracture, but endured with fatigue lives greater than 107 cycles.

These were classified as run-outs and were removed from testing and evaluated only for

coating performance and to see what effect the applied loads had on crack formation and

propagation within the coating matrix. These included all of the 4340 samples subjected









to 110 ksi maximum applied stress at 29 hz constant amplitude, as well as two samples

subjected to 125 ksi maximum applied stress and 29 hz constant amplitude. There were

no samples from the lot of 300M or Aermet 100 that achieved the run out condition. The

second set included the specimens that fractured at some point during the testing and the

fracture surfaces of both the coating substrate were evaluated. For the specimens that had

reached the run out condition, the coating was mostly intact except for three 4340

specimens tested at 110 ksi maximum applied stress and 29 hz that had very small

regions of coating delamination at the very edges where the surface was tapered during

the spray operation, as shown in Figure 4-3.














(a) 110 ksi


Figure 4-3. Fatigue specimen tested at 110 ksi cyclic stress for greater than 107 cycles.
The coating shows one small delamination near the edge of the coating.

High magnification optical and SEM examination of one of the specimens showed

that the coating had additional cracks propagating circumferentially adjacent to the

delaminated areas, perpendicular to the axis of loading. These cracks appeared to

migrate from initial surface defects and propagate towards other cracks and along

machining lines from other defects as shown in Figure 4-4 and Figure 4-5. The behavior

was consistent for the three run-out specimens that had coating failure.










LQ~dmo


- -L Cfrcumferinmial
ciucb


Figure 4-4. Circumferential coating cracks and the branches that are propagating along
machine lines on the surface shown in the scanning electron microscope
micrograph.


Figure 4-5. Micrograph of the crack tip from the delaminated coating area showing the
crack relation to the defects on the coating surface.

The coated specimens that fractured during the testing failed in the range from 1.7

million cycles to just below the run out limit of 10 million cycles. All of the Aermet 100


W1


(









samples in this range were tested at stress levels of 135 ksi maximum applied stress and

29hz constant amplitude had cracks that initiated at subsurface defects with an average

fatigue life of 5.43 million cycles. For 300M, the samples in this range were tested at

both 125 ksi and 59 hz constant amplitude and 130 ksi maximum applied stress at 29 hz

constant amplitude and all of these also fractured due to crack initiation at subsurface

defects. The fatigue life for the 125 ksi set of 300M specimens averaged 2,748 million

cycles and the 130 ksi set averaged 4.39 million cycles. For 4340, the specimens that

fractured were tested at 125 ksi maximum applied stress at 29 hz constant amplitude and

the failure mode was again due to a subsurface defect with an average fatigue life of 4.78

million cycles.

















Cz=f ar ina co tm F cc
r _- -' ..- __
U.-- .----" .,.. .* ...


Figure 4-6. Coating and substrate at the fracture surface for an Aermet 100 specimen
tested at 135 ksi maximum applied stress.

The coatings had circumferential surface cracks, similar to the run-out specimens,

as well as some degree of interfacial cracks that propagated along splat boundaries and

around carbide particle interfaces perpendicular to the axis of loading; however, the









substrate bond line was relatively intact with no visible spelling or delamination (Figure

4-6). .Optical examination of the substrate fracture surfaces for the three alloys, as

shown in Figures 4-7, 4-8, and 4-9, revealed almost identical features for the 4340, and

300M, with a slow crack growth region of approximately 0.0625" across the surface

before unstable crack growth predominates; whereas, the Aermet 100 has a slow crack

growth region of approximately 0.126", over half the diameter of the specimen.





1/16"









I ] _







Figure 4-7. Subsurface defect and small slow crack growth region for the 4340 steel
specimen tested at 125 ksi maximum stress.

For 4340 and 300M materials SEM inspection revealed that the mode of crack

initiation was from a sub-surface inclusion as shown in Figure 4-10. The higher

magnification images shown in Figure 4-11 display the worked surface morphology of

secondary cracks and tear ridges, which were common to both alloys. Fatigue striations

were very difficult to locate due to the R = -1 loading condition being fully reversed









which tends to obliterate the fine features in the slow crack growth area during the

compression phase.


ame


1/16"
MMM"


Figure 4-8. Subsurface defect at the origin, and the small slow crack growth region for a
300M steel specimen tested at 130 ksi maximum applied stress.


-.r-


S T~


Figure 4-9. Aermet 100 specimen showing the much larger slow crack growth region
and the high shear lip. Specimen also failed at a subsurface defect, which is
not clearly visible in this image.









However faint striations were detected as shown in Figure 4-12, near the transition zone

on the outer third of the slow growth semi-circular region for both 4340 and 300M. The

fast fracture region for 4340 and 300M was dominated by microvoid coalescence, as

shown in Figure 4-13 indicating ductile fracture once the crack became unstable.



.. -. 7..




~.-. ""_ -. c .
-'. .. ... -' "
















Figure 4-10. Subsurface inclusion and the radial lines leading to the crack origin for a
4340 steel specimen tested at 125 ksi maximum applied stress. Magnification
150x.

The Aermet 100 sample (Figures 4-14 and 4-15) shows similar microscopic

fracture features of radial lines emanating away from a subsurface origin and a highly

worked surface morphology with secondary cracks and tear ridges. The difference

between the Aermet 100 and the 4340 or 300M was that fatigue striations were much

more evident and easier to resolve in the Aermet 100 specimens and they are shown in

Figure 4-16.





























Figure 4-11. 4340 specimen showing the secondary cracking and tear ridges along
parallel fronts progressing away from the crack origin. 800x Magnification.


Figure 4-12. Faint striations detected near the outer edge of the slow crack growth region
in a 300M specimen. Magnification is 2500x.





























Figure 4-13. Microvoid coalescence in the region of fast fracture for a 300M specimen.
Magnification 2500x.


Figure 4-14. Low magnification micrograph image of the subsurface origin of the
Aermet 100 sample tested at 135 ksi. Radial lines are visible emanating away
from the origin area. Magnification 50x.





























Figure 4-15. Tear ridges and worked surface with concentric parallel lines on the Aermet
100 sample showing the fracture progression. Magnification 850x.


Figure 4-16. Fatigue striations shown on the fracture surface approximately halfway
between the origin and the transition zone between slow and rapid crack
growth.









Medium Applied Stress Specimens

In this range, all specimens were again tested at an R = -1, or fully reversed load

condition. The 4340 specimens were subjected to 150 ksi maximum applied stress with a

constant amplitude of 59hz, 300M specimens were tested at 140 ksi maximum applied

stress and 5hz constant amplitude, and the Aermet 100 specimens were tested at 145 ksi

at 29hz constant amplitude. All of the 4340 samples tested in this range of applied loads

reached the run-out condition of 107 cycles without fracturing. The Aermet 100 samples

averaged 3.3 million cycles to failure, and the 300M specimens averaged 506,000 cycles

to failure showing the strong dependence of fatigue life on the frequency of the applied

loads. The coating surface of the run-out specimens appeared intact macroscopically

with no apparent spalled regions of delaminations. However, high magnification

inspection revealed that cracks had initiated on the surface of the coating perpendicular to

the axis of loading. The cracks can be seen propagating between defects on the surface

as shown in Figure 4-17 and 4-18. The cracks are located near the center of the coated

gage section at the minimum specimen diameter and location of highest concentrated

stress.

Optical examination of the Aermet 100 specimens that fractured showed a flat,

slow crack growth region with radial lines emanating away from the origin area that

penetrated over half the diameter of the specimen as shown in Figure 4-19. The fast

fracture overload region was fibrous and dull gray in a appearance and propagated at

distinct 450 angle from the initial fracture surface plane in the direction of axial loading

and encompassed the remaining area of the surface. The 300M specimen had a





























Figure 4-17. Very fine circumferential coating cracks formed on the 150 ksi maximum
applied stress run-out specimens. Magnification 37x.


Figure 4-18. Higher magnification micrograph showing the very fine circumferential
cracks propagating between surface defects.

smaller slow crack growth region by comparison (Figure 4-20) and was similar in size to

the lower applied stress specimens. The fast fracture region was divided into two distinct

regions, the first being a coarse dull gray appearance with visible radial lines continuing


15kU X37 500mm 0000 OB/MAR/03


15kU X550 20mm 0000 OS/MAR/03






from the slow crack growth region becoming more pronounced as they reached the outer
diameter of the fracture surface. The second region was a final overload shear lip that
propagated at a shallow angle away from the fracture surface plane and encompassed
most of the circumference of the specimen.









Figure 4-19. Aermet 100 sample tested at 145 ksi maximum applied stress showing the
origin to at the surface.


0 25"Duiam et


ft


K


Figure 4-20. 300M specimen tested at 145 ksi maximum applied stress.









The noticeable difference at this stress level was that for the one Aermet 100

specimen out of the five total at this stress level (Figure 4-19) the fracture initiated at a

surface defect instead of a subsurface defect. The crack origin was at an imbedded

aluminum oxide particle trapped during the surface preparation prior to coating

deposition as shown in Figure 4-21. There was virtually no difference in the number of

cycles to failure of this specimen as compared to the others as this one lasted 3.96 million

cycles and the average was 3.29 million. The fracture surface for the 300M specimens at

this stress level all initiated a subsurface inclusion as shown in Figure 4-22.


















Figure 4-21. Origin area on Aermet 100 fatigue specimen tested at 145 ksi max stress.

The microscopic fracture surface features for both the Aermet 100, and the 300M

materials were very similar to what was seen on the lower applied stress level samples

with the slow crack growth areas having a concentric semicircle pattern of secondary

cracks and tear ridges. The fast fracture, or overload regions including the shear lip were

dominated by microvoid coalescence features with small interspersed cleavage facets,

probably along grain boundary edges, indicating a ductile fracture once the crack became

unstable (Figure 4-23).








MOMPlt Coating















Figure 4-22. Fatigue crack origin at a subsurface inclusion for a 300M specimen tested at
145 ksi maximum applied stress.


Figure 4-23. Transition region along the edge of the slow crack growth region for the
Aermet 100 specimen showing the tear ridges with interspersed regions of
microvoid coalescence.





























Figure 4-24. Coating/substrate interface for the Aermet 100 specimen showing the
coating fracture surface and the crack along the interface.

The coatings in this applied stress range showed some surface delamination and

spelling near the fracture surface of the 300M specimen (Figure 4-22). However the

Aermet 100 sample showed good bond line integrity even after final fracture but coating

cracks were visible along the substrate interface and in the first several splat layers

deposited. The coating cracks propagated along intersplat boundaries and around carbide

particles through the thickness of the coating before reaching the substrate interface as

shown in Figure 4-24. The through thickness coating cracks do not propagate into the

base metal or align near the substrate fatigue origin indicating that they do not provide a

preferential site for substrate crack initiation

Higher magnification inspection of a cross-section through the coating fracture

surface showed that coating cracks propagate primarily around carbide particles through

the cobalt binder matrix. Regions of carbide particle pullout and intersplat cracking were

evident as shown in Figure 4-25.





























Figure 4-25. Cross-section through the coating at the fracture surface just below the
substrate origin showing the intersplat crack propagation through the binder
phase and around the carbide particles along the crack front.

High Applied Stress Specimens

Because of the aggressive nature of the R = -1 test condition, specimens tested at

very high stress levels with the load applied in fully reversed tension and compression

fail in a much shorter time period and would be considered low cycle fatigue if not for

the load control condition. Therefore, in this stress range two different R values were

evaluated with both Aermet 100 and 300M specimens tested at 180 ksi maximum applied

stress at 5hz with R = -1, while the 4340 steel specimens were evaluated at 190 ksi and

220 ksi at 59hz, and R = 0.10. The Aermet 100 specimens averaged approximately

43,000 cycles to failure and the 300M specimens averaged approximately 14,000 cycles

to failure showing the increased fracture toughness of the Aermet 100. The 4340

specimens subjected to the 190 ksi maximum applied stress showed three samples with

fatigue lives averaging 4.46 million cycles, again showing the effects of the R ratio, while

two of the specimens failed at much lower lives, 59,000 and 118,000 cycles.


1 Tv F t

r~ii)Jk










Macroscopic examination showed gross subsurface inclusions approximately 0.005" in

diameter in both specimens with a star-like radial pattern emanating from this origin as

shown in Figure 4-26, and 4-27. These specimens were deemed exceptions and

characterization of their fracture properties was unrelated to the coating process. The

remaining 4340 specimens showed fracture behavior similar to what has been shown for

the lower applied stress range with a small slow crack growth region and radial lines

emanating from the origin area. The 220 ksi and 190 ksi specimens are shown in Figures

4-28 and 4-29. There are multiple origins and slow crack growth regions shown for the

220 ksi specimen. This did not reduce the number of cycles to fracture as this specimen

recorded the highest fatigue life for the 220 ksi maximum applied stress specimens at

50,779 cycles. Both the 190 ksi and 220 ksi specimens had surface defects at the origins

as shown in Figure 4-30, which were identified as imbedded aluminum oxide particles by

energy dispersive spectroscopy (EDS).


4*''


'i





I,


Figure 4-26. Macroscopic image showing the gross subsurface inclusion in the 4340
steel specimen tested at 190 ksi max stress at 59 hz and R = 0.1.
































Figure 4-27. Subsurface inclusion in the 4340 steel specimen approximately 0.005" in
diameter.


Figure 4-28. 4340 steel, 190 ksi maximum applied stress specimen.




























hMulllh Surfac




Figure 4-29. Much smaller slow crack growth regions and the multiple origins for the
4340 steel, 220 ksi maximum applied stress specimen.


Figure 4-30. Imbedded aluminum oxide particle at the fracture surface origin for the 220
ksi 4340 steel specimen.









The Aermet 100 and 300M specimens are shown in Figure 4-31 and 4-32 with

similar features as shown for the lower applied stress specimens. The Aermet 100

showed a much smaller slow crack growth region in this applied stress range covering

less than half the diameter of the specimen, also evidenced by the reduced cycles to

failure. The 300M also showed a smaller slow crack growth region but also showed a

much larger shear lip. Microscopic features were much more obliterated on the Aermet

100 specimen with no visible fatigue striations and smearing of the worked surface as

shown in Figure 4-33. The 300M sample had no visible fatigue striations, however very

small regions of microvoid coalescence were seen well into the slow crack growth region

as shown in Figure 4-34 indicating progressive instability of the crack.


Figure 4-31.
R


Aermet 100 specimen tested at 180 ksi maximum applied stress at 5 hz with
= -1.

































Figure 4-32. 300M steel specimen tested at 180 ksi maximum applied stress at 5 hz with
R=-1.


Figure 4-33. Smeared surface features and no visible fatigue indications for the Aermet
100 sample tested at 180 ksi maximum applied stress at 5 hz, and with R = -1.


L~:~Fj~l~Ci~





























Figure 4-34. Interspersed regions of MVC within the slow crack growth region.of a
300M specimen tested at 180 ksi, 5 hz, and R = -1.


Figure 4-35. Intact coating of the 300M specimen tested at 180 ksi.









The coatings for the Aermet 100 specimens and the 300M were relatively intact,

even near the fracture surface as shown in Figure 4-35. Small circumferential

microcracks were visible on the gage section surface, and interlamellar cracks were

evident on the fracture surface plane. The 4340 specimen also showed good adhesion at

190 ksi, but the 220 ksi specimen had circumferential delamination of an entire band of

coating as shown in Figure 4-36. The bond line interface showed good coating adhesion

near the origin, but cracks and delaminations were visible along this bond line in several

areas adjacent to the substrate shear lip. Both of these materials had surface defects at the

origin, as shown in Figure 4-37, and 4-38 that were confirmed with EDS to be imbedded

aluminum oxide particles.


Delaminated Coating


Figure 4-36. Delaminated coating around the fracture surface of the 4340 specimen
tested at 220 ksi, 59 hz and R=0.1















































Figure 4-37. Imbedded particle origin for the 300M sample tested at 180 ksi maximum

applied stress.


Figure 4-38. Origin area for the Aermet 100 specimen tested at 180 ksi showing the

imbedded aluminum oxide particle.


........... ........ ..... ................
.. ................
.. ...............
............................
..................
...... .............
.........
.......... ................

50mm 0000 09/MAR/03














CHAPTER 5
CONCLUSIONS

Cracks in WC 17%Co high velocity oxygen fuel (HVOF) thermal-spray coatings

initiate at defects in the coating structure at the surface, and propagate along intersplat

regions of high tensile stress within the coating structure. Almen strip deflections

confirm that the bulk macrostress is a net residual compressive stress and x-ray

diffraction confirms that the tungsten-carbide particles are also in compression.

Therefore, the cobalt binder must be comprised of tensile stresses within the matrix of

individual phases. Optical and scanning electron microscopy showed that crack

propagation will proceed along these interfacial boundaries and around the high

compressive stressed WC particles to regions of high tensile stress along splat interfaces

in the radial direction, perpendicular to the axis of loading. Fractographic analysis of

specimens that were subjected to low and medium applied stress fatigue stresses, but did

not fracture showed very fine coating cracks that initiated at defects on the surface and

propagated circumferentially towards other defects. As the stresses were increased a

continuous, through thickness network of coating cracks was formed to relieve the

internal stress in the coating. If continuous cracks reached the edge of the coated gage

section, delamination or spelling could occur however there was no evidence that these

coating cracks provided a preferential site for substrate fatigue crack initiation. Optical

and SEM examination of cross-sections and of the fracture surfaces showed that when the

coating cracks reached the substrate, they would either propagate along the bond-line









interface, or turn back into the coating and propagate along splat boundaries in the first

few layers of deposited coating.

Fractographic analysis of the fatigue specimens that fractured at higher applied

stresses clearly showed that the substrate origins were located at either a subsurface

inclusion, or at a coating-substrate interfacial bond line defect. The subsurface defects

were the predominant mode of initiation at applied fatigue stresses below 180 ksi for R =

-1, fully reversed conditions, as well as at lower applied stresses with higher fatigue lives.

The substrate-coating interface defects predominated for the specimens tested at 180 ksi

maximum applied stresses where R = -1, and at higher stresses of 190 ksi and 220 ksi

where R = 0.1. Energy dispersive spectroscopy confirmed that the defects at the coating-

substrate interface were imbedded aluminum oxide particles from the grit-blasting

operation prior to coating deposition. The subsurface defects could not be identified but

could be M2C carbide particles at grain boundaries or sulfur inclusions in the matrix.

The reduction in fatigue life at low to moderate applied stresses associated with the

HVOF coating application can be attributed to the change in the compressive residual

stress state of the substrate from the grit-blasting process prior to coating deposition. At

higher stresses, coating substrate surface defects will control the reduction in fatigue life.

Coating cracks do not propagate in to the substrate or provide preferential sites for

substrate fatigue cracks to initiate. The compressive residual stress in the substrate

surface acts to divert the coating crack path along the coating-substrate bond-line.

The tungsten-carbide cobalt coatings had similar affects on the three high-strength

steel substrates and further investigation of the HVOF coating deposition process is

warranted to determine if a change in surface preparation for coating deposition including






67


a reduction in the quantity of imbedded particulate can significantly decrease the amount

of fatigue life reduction. Also, surface finishing procedures that further reduce carbide

particle pullout as well as machining lines could reduce the tendency for crack formation

at low and moderate applied stress levels.















APPENDIX A
EXPERIMENTAl DATA FOR DEVELOPING HVOF SPRAY PARAMETERS FOR
TUNGSTEN CARBIDE 17%-COBALT COATINGS














Table A-1. Optimization parameters for the Tungsten-carbide 17%-Cobalt coatings


Design 1: Use L8 design plus Center Points, 11 runs
total


FACTORS:
A Surf Speed,Feed
B Combustion Gas
C Stoic Ratio
D Spray Distance


A
m Factor:
(-1)
C Pt
(+1)


Levels
-1 +1
Rate 1335, 5.1 1835. 3.5
1525 scfh 1825 scfh
0.405 0.485
10 inch 13 inch


Turntable Robol Spd


RPM
212
252
292


ipll mmisec
25 10.6
35 14.8
50 21.2


Robol Oo @

750 mmisec
1.410o
1.980o
2.8200


C Pt
1585 ipm, 4.3
1675 scfh
0.445
11.5 inch


Spols/Rev
5.1
4.3
3.5


FIXED:
54 grit alumina grill blast al 40 psi, 6 inches
Substrate is 4340 sleel, 260-280 ksi
Powder size/type is WC-17Co, Diamalloy 2005, Lot 54480
Powder Feed Rate"" 8.5 Ibs/hr
Spray angle is 90 degrees
100 psi cooling air, 4 AJs @ 6 inch spaced over coupon area
Carrier gas N: al 148 psi, 55 flow, air vib @ 20 psi
Spray pattern length Approximately 13 inch


Fixture diarneler


2 inch


(B,C) Factor Combinations:
Comb
Gas Stoic Ratio Hyd SCFH Oxy SCFH


1675
1525
1525
1825
1825


0.445
0.405
0.485
0.405
0.485


1159
1085
1027
1299
1229


RESPONSES: RELATED CTG FUNCTION:
1) Palr lemperalure Fatigue


Air SCFH Point (CG.SR)
920 ( 0, 0)
920 (-1,-1)
920 (-1,+1)
920 (+1,-1)
920 (+1,+1)


2) Almen strip


Fatigue, ctg residual stress


3) Hardness. HV.,,,,:, Wear
4) Coating dep/pass Cost
5) Porosity Ctg quality, corrosion
6) Oxides Ctg quality
7) Carbides Ctg quality, wear
8) Tensile bond Adhesion/cohesion









Coating-Process Response Measurements

A number of process responses were tracked for the deposited coatings. These

included deposition rates, DPH microhardness, R15N superficial surface hardness, tensile

bond strength by a modified ASTM C633 method, Almen strip arc height as an indication

of coating residual stresses, and substrate temperature during spraying. The specimens

for these measurements were rotated on a turntable while mounted on a cylinder as the

gun traversed parallel to the cylinder axis of rotation in an oscillating stroke.

Coating-Deposition Rate

The deposition rate was determined as the measured coating thickness divided by

the recorded number of gun passes. The final coating thickness was measured by hand

held flat anvil micrometers on the coating microstructure coupons, typically a 0.25 in

cube or one square inch piece of 0.060 in or thicker sheet metal. The number of gun

passes was recorded as determined by an automatic cycle counter during the spray

operation. One cycle consisted of an initial stroke and its return stroke, thus two passes

for any given position along the cylinder's length.

Microhardness

The microhardness measurements were taken on a polished cross-section of the

microstructure coupon. A standard commercially available diamond pyramid indentor

and test machine were used with a 300 gm load. The specimen was sectioned parallel to

the rotation direction and perpendicular to the gun traverse direction, then mounted and

polished by standard procedures. Ten readings were taken in a standard pattern which

diagonally traversed the coating thickness from the substrate to the surface and back to

the substrate again taking care to maintain adequate distances from coating edges and

between indentations.









Rockwell 15N Superficial Surface Hardness

This data was taken by making indentations directly on the coated coupon surface

after lightly hand polishing the surface on 400 grit SiC metallographic polishing paper to

remove any loosely adherent particles or small asperities. Testing was done with a

standard commercial tester with a dial indicator or automatic printout of the hardness

reading which used a 120 degree diamond cone indentor and a 15 kg load. It is possible

that at the coating thickness used, there is some influence of the substrate hardness on the

absolute values obtained for the coating hardness, but since the substrates were all IN718

typically of Rockwell C 40+/-2 and the coatings all .010+/-1 in thickness, it was assumed

that changes in hardness were reflections of process effects. In general, the Rockwell

15N hardness showed minimal variability in most of the HVOF process studies so

statistical analyses were not performed for these data.

Tensile Bond Strength

Tensile bond strengths were determined by a modified ASTM C633 method. The

coatings were deposited on one inch diameter buttons which were 0.25 in thick. This

allows buttons to be conveniently mounted along with the other specimens, or in the case

of coating actual parts, often lends itself to locating buttons right on or adjacent to the

part. The button s were then bonded between the normal ASTM C633 mandrels with FM

1000 film adhesive (Cytek) and tested per the specification. The HVOF WC-Co always

resulted in ultimate failure of the film adhesive at greater than 12,000 psi..

Substrate Temperature

The temperatures measured for the spray process trials were taken by an IR

pyrometers used on rotating specimens which were XY traversed. This has shown









general agreement within about +/- 25 F on the absolute values with contact

thermocouples and agreement in trends between individual process trials.

Almen Strip Deflection

The Almen strip was a standard type N SAE 1070 steel strip 3"x0.75"x0.030" held

by 4 round head screws (see Mil-S-13165C for more details). The Almen strip was grit

blasted on one side only and the arc height due to the induced bending stresses recorded

before and after spraying of the HVOF coating. The difference in the two measurements

was reported as the Almen strip deflection. A negative sign indicated a change

representing a compressive stress in the coating and a positive sign indicated a tensile

stress in the coating. (Note: Shot peening work at GE has demonstrated the validity of

taking the difference in curvature without the need to return the Almen strip to a perfectly

flat condition; i.e. the deflection is cumulative in linear fashion in the elastic region. The

single side grit blast procedure typically resulted in a starting arc height of -0.004 in; i.e.

compressive.) The Almen strip is restrained in the flat position during coating. No effort

was made to calculate actual coating stresses since the exact coating modulus was

unknown.
















A factor (B,C) Combined Factors D factor

Std.Ord Turn Table Robot Tray Hydrogen Oxygen Air Sp Dist


psi/FMR

135 psi, 50.4

135 psi, 47.2

135 psi, 47.2

135 psi, 56.5

135 psi, 56.5

135 psi, 50.4

135 psi, 44.6

135 psi, 44.6

135 psi, 53.4

135 psi, 53.4

135 psi, 50.4


psi/FMR

148 psi, 23.1

148 psi, 17.8

148 psi, 17.8

148 psi, 23.8

148 psi, 23.8

148 psi, 23.1

148 psi, 21.8

148 psi, 21.8

148 psi, 28.7

148 psi, 28.7

148 psi, 23.1


psi, FMR

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5

105 psi, 50.5


Run No.

1

2

3

4

5

6

7

8

9

10

11


RPM

252

212

292

212

292

252

212

292

212

292

252


Sp mm/s

14.8

10.6

21.2

10.6

21.2

14.8

10.6

21.2

10.6

21.2

14.8


inches

11.5

10

13

13

10

11.5

13

10

10

13

11.5
















StdOrder

1

2

3

4

5

6

7


CPs


Mils/pass

0.790

0.408

0.830

0.386

0.870

0.407

0.670


0.420

0.583

0.558

0.600


T 6 cy

349

236

347

350

264

281

359


261

330

295

290


Norm Aim

4.6

4.1

6.5

11.1

2.9

6.1

15.3


10.7

6.4

7.5

6.6


Porosity


DPH 300

1166

984

1121

1163

925

1106

1177


0.67

0.50

0.50

0.25

0.75

0.25

0.10


0.37

0.37

0.37

0.25


1011

1127

1110

1131
















Design of Experiments Matrix


99.01 not used

99.10not used

99.14not used

99.15

99.16


Differs from 6,7,13

Equals 10 Repeat

Equals 7 Repeat

-1

-1


Actual

Run Order

99.06

99.02

99.11

99.05

3

7

8

4

99.12

99.09

99.13


StdOrder

9

1

2

3

4

10

5

6

7

8

11


Run No.

6

2

3

4

5

6R

7

8

9

10

11













Design of
Experiments
Results Analyses (All broke

% used as % in epoxy) DPH 300

Cycles Thk, mils Mils/pass Tmax T6 Almen gb Almen sp Delta Alm Norm Alm Porosity Porosity Tensile,psi hardness


6.4

4.6

4.1

6.5

11.1

7.5

2.9

6.1

15.3

10.7

6.6


3.3

10.8

3.1

13.9
14.3


.25-.5

.5-.75

0.5

0.5

0.25

.25-.5

0.75

0.25

0.1

.25-.5

0.25


0.5

.1-.25

.5-.75


0.37

0.67

0.50
0.50

0.25

0.37

0.75

0.25
0.10

0.37

0.25


12560 1127

12513 1166

12475 984
13371 1121

11999 1163

11986 1110

11143 925

11099 1106c1
12269 1177

11112 1011

11716 1131


0.583

0.790

0.408

0.830

0.386

0.558

0.870

0.407

0.670

0.420

0.600


0.600

0.367

0.880

0.325
0.350


10.75

7.5

5.0

11.8

15.0

11.0

6.5

8.0

22.0

9.0

11.0


4.8

10.6

6.5


1.75

0.25
1.0

1.0

3.0

1.0

1.5

1.0

1.5

0.0

1.5


0

1.1

1.0

1.5
0


9.0

7.25

4.0

10.8

12.0

10.0

5.0

7.0

20.5

9.0

9.5


4.8

9.5

5.5

14.5
12


12316

12964

11849


1073

1083

974

















Almen vs Temp (CY 6)


18.0 -


16.0 -


14.0 -


12.0 -


10.0


8.0


6.0


4.0


2.0


0.0 -
200


Figure A-1. Almen strip deflection vs. substrate part temperature.


220 240 260 280 300 320 340 360 380

Temp, Deg F

















Temp vs mils/pass


380


360


340


320


300


280


260


240


220


200
0.300


Figure A-2. Substrate part temperature vs. number of mils of coating deposited per pass of the torch.


0.400 0.500 0.600 0.700 0.800
mils/pas


0.900

















Hardness DPH 300 vs % Porosity


1200 -


1150 -


1100


1050


1000


950


900


850


800 -
0.00


Figure A-3. Final coating hardness vs. % porosity.


0.10 0.20 0.30 0.40 0.50 0.60 0.70
% Porosity
















ANALYSES: !2
RUNS



Run-1 mils/pass



Term EFFE1 COEF1


Constant

A

B

C

D

AB

AC

AD


Ct Pt


Almen Porosity



EFFE3 COEF3 EFFE4 COEF4


-0.38475

-0.04225

-0.01175

0.06875

0.03775

0.02825

-0.05125


0.675

6.475

2.175

-3.225

-0.675

-1.375

2.025


-0.01729 -0.875


7.662

O.337

3.237

1.087

-1.612

-0.337

-0.687

1.O12


-0.829 17


-0.1625

-0.2375

-0.1125

0.2125

0.1725

0.0475

-0.0275


0.42375

-0.08125

-0.11875

-0.05625

0.10625

0.08625

0.02375

-0.01375


-0.09375


41.04












Final Data Set


0.42375
1 -0.08125
1 -0.11875
1 -0.05625
0 0.10625
1 0.08625
1 0.02375
0 -0.01375


7.6625
1 0.3375
1 3.2375
1 1.0875
-0.33 -1.6125
1 -0.3375
1 -0.6875
-0.33 1.0125


0.597625
1 -0.19238
1 -0.02113
1 -0.00588
0 0.034375
1 0.018875
1 0.014125
0 -0.02563


Almen = 11.5; Hardness =
0.4 mils per pass deposit r


0.597625
-0.19238
-0.02113
-0.00588
0
0.018875
0.014125
0
0.41125
1079 DPH; Porosity =0.277%; Max Substrate Temp=2940F;
ite.


0.42375
-0.08125
-0.11875
-0.05625
0
0.08625
0.02375
0
0.2775
7.6625
0.3375
3.2375
1.0875
0.532125
-0.3375
-0.6875
-0.33413


Almen 11.5







82




Hitemco Powder Lots vs L12 Lots


100 -


0 10 20 30 40 50 60
Powder Size, microns


70 80 90 100


Figure A-4. Powder size trends

Powder size was also tracked based on a matrix of lots to determine if was a critical

variable. Analysis displayed above shows that the deviation in particle size distribution

for each powder lot was insignificant and not a factor in any of the DOE runs for

parameter optimization.

The combustion gas ratios, flow rates, and spray distance are the major factors in

the spray process. The response data for hardness, almen deflection (residual stress) and

substrate temperature identified these as critical parameters for controlling final

mechanical properties. The stoichiometry of the combustion fuels affects the melting of


S- -o*- Diamalloy 2005 Lot 53792
--- Diamalloy 2005 Lot 54327
S-A- Diamalloy 2005 Lot 53791
S- Diamalloy 2005 Lot 54627
Diamalloy 2005 Lot 54480
-0'- Diamalloy 2005 Lot 53558
-A- Stark 526.062 Lot 10362


-M










the cobalt binder and dissolution of the carbides. Non-optimum gas ratios result in

secondary carbide formation and reductions in hardness and toughness. Spray distance

has a substantial affect on the residual stress state and substrate part temperature due to

the amount of heat transfer between the molten particulate and the base metal. This also

affects solidification rates and thus the net porosity and volumetric changes.


Equipment: Sulzer Metco Inc. Model 2600 hybrid gun
Powder: Sulzer Metco Inc. Diamalloy 2005, 83%-WC 17%-
Co Agglomerated/Sintered
Powder Feeder: Sulzer Metco Inc. Model Single 10, Positive
Pressure, Screw Feed with Vibration.
Powder Feed Rate: 8.5 lb/hr (325 rpm, 6 pitch feeder screw)

Vibrator Setting: 30
Powder Carrier Gas: Nitrogen at 148 psi, and 28 scfh flow.

Combustion Gas: Oxygen at 148 psi; 44+/-2 Flow Meter Reading
(FMR) Not in scfh
Combustion Fuel: Hydrogen at 135 psi console pressure and 1229 scfh
flow.
Combustion Chamber Pressure: 100-102 psi
Gun Cooling Water: 9.3 8.7 gph
Water Temperature to the gun: In, 64- 720 F; Out, 117-125 F; Delta, 51- 540 F

Specimen Rotation: 2,636 rpm for round bars (0.25 inch dia.) 16,560
in/min surface speed.
Gun Traverse Speed: 400 linear in/min for round bars
Spray Distance: 11.5"
Cooling Air: 2 gun mounted Air Jets at 14 inches, 90-110 psi
1 stationary Air Jet at 4-6 inches, 90-110 psi
Part Temperature: Max temperature as read with a infrared pyrometer
is 275 deg F trailing the plume spot.
Injector Number 9
Insert Number 9
Shell Number 9
Siphon Plug Number 9
Air Cap 2701















APPENDIX B
FATIGUE DATA FOR 4340, 300M, AND AERMET 100





















4340, SMALL HOURGLASS SPECIMEN
(0.003" COATING) R = -1, AIR
200.0


190.0
+ EHC/Peened
EHC/Peened/FIT
180.0 WCCo/Peened
S-- WCCo/Peened/FIT
170.0 -


c1l 60.0 -
LCJ

9150.0 -


2140.0 -

z
(130.0 -
z

120.0


110.0 +


100.0
1.E+03 1.E+04 1.E+05 1.E+06 1.E+07 1.E+08
CYCLES TO FAILURE, Nf














300M, SMALL HOURGLASS SPECIMEN
(0.003" COATING) R = -1, AIR


* Bare/UNPeened

+ EHC/Peened

-EHC/Peened/FIT

* WCCo/Peened

--WCCO/Peened/FIT


. . . .. .... ......... .I ..


1.E+05 1.E+06
CYCLES TO FAILURE, Nf


1.E+07


200


190-


180-


"70 -


c(160
w
150 -
(D
z
E40-

z
3130-
z

120


110.


Inn


* U


1.E+03


1.E+04


1.E+08















A100, SMALL HOURGLASS SPECIMEN
(0.003" COATING) R = -1, AIR


Bare/UNPeened
- Bare/UNPeened/FIT
+ EHC/Peened
EHC/Peened/FIT
WCCo/Peened
WCCo/Peened/FIT


t


190


180


170


160,
X

S150'
LU

S140

z
R 130
LU
LU
z
(9 120
Z
LU

110


100


%


1.E+05 1.E+06
CYCLES TO FAILURE, Nf


1.E+07


1.E+08


q


on =


S1.E+
1.E+03


1.E+04


' ''' '' ''" '"" ''U"