fracture and residual-stress characterization of tungsten- carbide ... - Fcla

FRACTURE AND RESIDUAL-STRESS CHARACTERIZATION OF TUNGSTENCARBIDE 17%-COBALT THERMAL-SPRAY COATINGS APPLIED TO HIGHSTRENGTH 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. JeamGabriel 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.

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TABLE OF CONTENTS Page ACKNOWLEDGMENTS ................................................................................................. iv LIST OF TABLES............................................................................................................ vii LIST OF FIGURES ......................................................................................................... viii ABSTRACT...................................................................................................................... xii 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

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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

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EXPERIMENTAL PROCEDURE.............................................................................21 Specimen Manufacture ...............................................................................................21 Coating Quality Control .............................................................................................26 Residual Stress Measurement .....................................................................................28 Fatigue Testing ...........................................................................................................32 Scanning Electron Microscopy...................................................................................33

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RESULTS AND DISCUSSION.................................................................................34 Residual Stress Measurements ...................................................................................34 Optical and Scanning Electron Microscopy ...............................................................39 Low-Applied-Stress Specimens..................................................................................40 v

Medium Applied Stress Specimens ............................................................................50 High Applied Stress Specimens..................................................................................56 5

CONCLUSIONS ........................................................................................................65

APPENDIX A

EXPERIMENTAl DATA FOR DEVELOPING 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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LIST OF TABLES page

Table

3-1 Composition for the three alloys evaluated. ...............................................................23 3-2 Particle size distribution for 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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LIST OF FIGURES page

Figure

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 initial coating evaluation.. ..............................................................6 2-1 Model of shot peened surface showing the stress profile created by the impact dimples 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 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 for 4340, 300M, and Aermet 100. .......................................22 3-2 Ductility properties for tensile elongation 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 residual stress measurement 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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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 ksi 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 from the delaminated coating area.................................42 4-6 Coating and substrate at the fracture 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 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 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 with 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 coating cracks formed on the 150 ksi maximum applied stress run-out specimens..............................................................................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 the surface. .............................................................................................52 ix

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 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 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 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

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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 Powder size trends .....................................................................................................82

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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 TUNGSTENCARBIDE 17%-COBALT THERMAL-SPRAY 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 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 xii

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.

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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 1

2 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

3 from the particle to the substrate during solidification, creating a mechanical and interlocking diffusion (metallurgical) bond (Figure 1-2).

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

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

4 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 hardchrome electroplating, the beneficial wear properties of the carbides were very attractive as compared to other materials. However, durability of the brittle coatings in fatiguesensitive areas was a concern. As discovered under the initial DARPA project, thermalsprayed 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

5 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.

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

190.0

EHC/Peened/FIT

180.0

WCCo/Peened WCCo/Peened/FIT

170.0 160.0 150.0 140.0 130.0 120.0 110.0 100.0 1.E+03

1.E+04

1.E+05

1.E+06

1.E+07

1.E+08

CYCLES TO FAILURE, N f

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

7 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 unlubricated conditions; corrosion tests that included atmospheric exposure, ASTM B117 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

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10 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 highstrength 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

11 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 thermalspray 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

12 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 ε through the change in the Bragg scattering angle ∆θ, where λ=2dsinθ. Solving for strain ε=∆d/d, and with accurate knowledge of the initial value of d, the stress free interplanar spacing, the

13 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 Φ and Ψ 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

14 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

15 residual compressive stress that tends to inhibit 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 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

16 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].

17 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

18 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 WC17%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).

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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

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 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 of W2C 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 21

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 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 heattreated in accordance with Carpenter Technology Process Specification 15169.

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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

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