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Titlepage ignore_heading
Abstract ignore_heading
Dedication & Acknowledgements ignore_heading
COMMENT Declaration ignore_heading
TOC, Tables, Figures, Glossary ignore_heading
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COMMENT Our own publications
End of Preliminaries ignore_heading
Chapters
Introduction
TODO Paragraph: Cavitation ignore_heading
Paragraph: Introduction to Stellite Alloys for Hostile Environments ignore_heading
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What they are and where they came from.
- Identification as Cobalt-base superalloys
- Core beneficial properties high strength, corrosion resistance, high-temperature hardness
- Pioneering figure (Elwood Haynes) and seminal alloy Stellite 6 with nominal composition
Stellites are a family of cobalt-base superalloys used in aggresive service environments due to retention of strength, wear resistance, and oxidation resistance at high temperature \cite{ahmedStructurePropertyRelationships2014, shinEffectMolybdenumMicrostructure2003}. Starting with Elwood Haynes's development of alloys like Stellite 6 in the early 1900s \cite{hasanBasicsStellitesMachining2016}, stellites became critical to components used in medical implants & tools, machine tools, and nuclear components, and new variations on the original CoCrWC and CoCrMoC alloys see expanding use in sectors like oil & gas and chemical processing \cite{malayogluComparingPerformanceHIPed2003, ahmedStructurePropertyRelationships2014, raghuRecentDevelopmentsWear1997}.
Paragraph: Impact of Composition, Microstructure, and Processing on Corrosion and Cavitation Performance ignore_heading
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What they are and where they came from.
- Main alloying elements and ref to tab:stellite_composition
- Describe the microstructure briefly
- Carbon grades
The main alloying elements in Stellite alloys are cobalt (Co), chromium (25-33 wt% Cr), tungsten (0-18 wt% W), molybdenum (0-18 wt% Mo), carbon (0.1-3.3 wt% C), and trace elements iron (Fe), nickel (Ni), silicon (Si), phosphorus (P), sulphur (S), boron (B), lanthanum (La), & manganese (Mn); \tref{tab:stellite_composition} summarizes the nominal and measured composition of commonly used Stellite alloys \cite{ahmedMappingMechanicalProperties2023, alimardaniEffectLocalizedDynamic2010, ashworthMicrostructurePropertyRelationships1999, bunchCorrosionGallingResistant1989, davis2000nickel, desaiEffectCarbideSize1984, ferozhkhanMetallurgicalStudyStellite2017, pacquentinTemperatureInfluenceRepair2025, ratiaComparisonSlidingWear2019, zhangFrictionWearCharacterization2002}. Stellite alloys possess a composite-like microstructure, combining a cobalt-rich matrix strengthened by solid solutions of chromium, tungsten, & molybdenum, with embedded hard carbide phases with carbide formers Cr (of carbide type $\textrm{M}_{7}\textrm{C}_{3}$ & $\textrm{M}_{23}\textrm{C}_{6}$) and W/Mo (of carbide type $\textrm{MC}$ & $\textrm{M}_{6}\textrm{C}$), that impede wear and crack propagation \cite{ahmedSlidingWearBlended2021a, crookCobaltbaseAlloysResist1994}.
Paragraph: Co phases ignore_heading
Understanding the cobalt phase is crucial for studying structural changes in Co-based alloys widely used in industry. The fcc cobalt phase, especially its delayed transition to hcp at ambient and moderate temperatures \cite{DUBOS2020128812}, is of particular interest due to its impact on material properties in Co-based alloys \cite{Rajan19821161}. As the cobalt phase in stellite alloys is observed to consist of the fcc phase \cite{Rajan19821161}, the potential for strain-induced fcc to hcp transformation is of interest under the mechanical loading of cavitation erosion.
Cobalt exhibits a hexagonal close-packed (hcp) structure above 700 K \footnote{the theoretical transition temperature was determined to be 825 K by Lizarraga et al \cite{Lizarraga2017}} and shifts to a face-centered cubic (fcc) structure above this temperature.
At ambient conditions, the metastable FCC retained phase can be transformed into HCP phase by mechanical loading, although any HCP phase is completely transformed into a FCC phase between 673 K and 743 K \cite{DUBOS2020128812}.
The main allotropes of the Co solid solution in stellite are the hexagonal close-packed (hcp) $\epsilon−Co$ (ICDD# 01-071-4239) and the face-centered cubic (fcc) $\gamma−Co$ (ICDD 00-015-0806) \cite{wuMicrostructureEvolutionCrack2019} with the $\epsilon−Co$ phase being more thermodynamically stable below 700K \cite{lizarragaFirstPrinciplesTheory2017}. However the $\gamma$\textrightarrow$\epsilon$ transformation is difficult to achieve under normal condition, leading to cobalt alloys often retaining the metastable $\gamma$ phase \cite{davis2000nickel, frenkMicrostructuralEffectsSliding1994}. At ambient conditions, the metastable fcc retained phase in stellites can undergo a strain-induced martensitic-type transition involving partial dislocation movement \cite{HUANG2023106170}, such as during cavitation erosion as observed experimentally by Woodford \cite{woodfordCavitationerosionlnducedPhaseTransformations1972}.
The fcc to hcp transition is related to the very low stacking fault energy of the fcc structure (7 mJ/m2) of cobalt \cite{Tawancy1986337}
This promotes the FCC→HCP martensitic transformation and improves the mechanical properties of the alloy. Phase composition seriously affects the properties of alloys. The FCC phase is soft and plastic, while the HCP phase is hard and not plastic. Therefore, when an alloy undergoes strain-induced FCC→HCP martensitic transformation, the martensite HCP phase plays a “secondary hardening” effect, which greatly improves the overall hardness of the alloy.
This fcc to hcp transition is linked to the very low stacking fault energy of the fcc structure (approximately 7 mJ/m²). The addition of other elements influences these characteristics: solid-solution strengthening increases the fcc cobalt matrix strength by distorting the atomic lattice and can decrease the low stacking fault energy by adjusting the electronic structure. Solute atoms like molybdenum (Mo) and tungsten (W), due to their large atomic sizes, discourage dislocation motion in stellites. Since dislocation cross-slip is a primary deformation mode in imperfect crystals at elevated temperatures and dislocation slip is a diffusion-enhanced process at high temperatures, this contributes to high-temperature stability. Furthermore, elements like nickel (Ni), iron (Fe), and carbon (C) stabilize the fcc structure of cobalt (a = 0.35 nm), while chromium (Cr) and tungsten (W) tend to stabilize the hcp structure (a = 0.25 nm and c = 0.41 nm).
The $\epsilon-\textrm{Co}$ (ICDD# 01-071-4239)and $\gamma-\textrm{Co}$ (ICDD 00-015-0806) phases are the main phases of the Co solid solution in stellite Although the $\epsilon-\textrm{Co}$ phase is more stable at room temperature according to the phase diagram, cobalt alloys often posses a majority $\gamma-\textrm{Co}$ phase; the $\gamma$\textrightarrow$\epsilon$ transformation rarely occurs under normal cooling condiyions and the metastable $\gamma$ phase is generally retained
Cobalt and Co-Cr alloys undergo thermally induced phase transformation from the high temperature face-centered cubic (fcc) $\gamma$ phase to low temperature hexagonal close-packed (hcp) $\epsilon$ phase at 700 K and strain induced fcc-hcp transition through maretensitic-type mechanism (partial movement of dislocations) \cite{HUANG2023106170}. At ambient conditions, the metastable FCC retained phase in stellites can be transformed into HCP phase by mechanical loading, although any HCP phase is completely transformed into a FCC phase between 673 K and 743 K \cite{DUBOS2020128812}; the metastable fcc cobalt phase in stellite alloys \cite{Rajan19821161} absorbs a large part of imparted energy under the mechanical loading of cavitation erosion. The fcc to hcp transition is related to the very low stacking fault energy of the fcc structure (7 mJ/m2) \cite{Tawancy1986337}.
Solid-solution strengthening leads to increase of the fcc cobalt matrix strength (due to distortion of the atomic lattice with the addition of elements of different atomic radii), and decrease of low stacking fault energy \cite{Tawancy1986337} due to the adjusted electronic structure of the metallic lattice. Dislocation motion in stellites is discouraged by solute atoms of Mo and W, due to the large atomic sizes. Given that dislocation cross slip is the main deformation mode in imperfect crystals at elevated temperature, as dislocation slip is a diffusion process that is enhanced at high temperature, this leads to high temperature stability \cite{LIU2022294}. In addition, nickel (Ni), iron (Fe), and carbon (C) stabilize the fcc structure of cobalt (a = 0.35 nm), while chromium (Cr) and tungsten (W), stabilize the hcp structure (a = 0.25 nm and c = 0.41 nm) \cite{Vacchieri20171100, Tawancy1986337}.
\cite{woodfordCavitationerosionlnducedPhaseTransformations1972}
confirmed that $\gamma$\textrightarrow$\epsilon$ transformation occur on the surface of cobalt-base alloys during cavitation erosion, a relationship a direct relationship between . While the extent of this deformation-induced transformation was observed to increase with the severity of erosion and reach a steady state corresponding with the weight loss rate, subsequent experiments on Stellite 6B with varying initial hcp phase content and different aging treatments failed to establish a direct correlation between the transformation characteristics and erosion resistance.
Solid-solution strengthening is provided by elements not tied in secondary phases, leading to increase of the fcc cobalt matrix strength.
With the addition of elements with different atomic radiuses, the atomic lattice of the fcc cobalt matrix is distorted leading to increased strength. The already low stacking fault energy of the fcc cobalt structure (7 mJ/m2) \cite{Tawancy1986337} is further decreased, inhibiting dislocation cross slip. Given that dislocation cross slip is the main deformation mode in imperfect crystals at elevated temperature, as dislocation slip is a diffusion process that is enhanced at high temperature, this leads to high temperature stability \cite{LIU2022294}.
The addition of nickel (Ni), iron (Fe), and carbon (C) stabilize the fcc structure of cobalt, while chromium (Cr) and tungsten (W), stabilize the hcp structure. Cr guarantees hot corrosion resistance and forms M23C6 carbides, while form MC carbides \cite{Vacchieri20171100}. The fcc cobalt phase has lattice constant a = 0.35 nm while the hcp cobalt phase has lattice constant a = 0.25 nm and c = 0.41 nm \cite{Tawancy1986337}.
Paragraph: Carbides and the grades of stellite alloy ignore_heading
The proportion and type of carbides depend on carbon content and the relative amounts of carbon with carbide formers (Cr, W, Mo), which greatly influence alloy performance and intended applications as carbides provide higher strength and but may reduce corrosion resistance due to localized corrosion at carbide boundaries. High carbon alloy (>1.2 wt%) have higher hardness due to greater carbide formation and are primarily used for wear resistance, low carbon alloys (<0.5 wt%) are used for enhanced corrosion resistance, and medium carbon alloys (0.5 wt% - 1.2 wt%) are used in applications requiring a combination of wear and corrosion resistance \cite{davis2000nickel}.
The carbide and grain size can be controlled by the rate of freezing, with larger carbide sizes indicating slower freezing rates \cite{yuInfluenceManufacturingProcess2008}.
Role of HIPping vs as Cast
Paragraph: Stellite 1 ignore_heading
Paragraph: Tungsten and Molybdenum carbides
Tungsten (W) and molybdenum (Mo) are refractory elements that provide solid solution strengthening to the matrix, by virtue of their large atomic size that impedes dislocation flow when present as solute atoms \cite{boeckRelationshipsProcessingMicrostructure1985}, and also form $\textrm{M}_6\textrm{C}$ and $\textrm{M}_12\textrm{C}$ carbides along with $\textrm{MC}$ carbides and $\textrm{Co}_3\textrm{M}$ & $\textrm{Co}_7\textrm{M}_6$ intermetallics during solidification.
In carbon-rich regions, the $\textrm{MC}$ phase (of type $\textrm{WC}$ and $\textrm{MoC}$) is observed \cite{zhangThermodynamicInvestigationPhase2019}, which ca
In carbon-poor regions, ternary $\textrm{M}_6\textrm{C}$ and $\textrm{M}_{12}\textrm{C}$ carbides have been identified, where the $\textrm{M}_6\textrm{C}$ carbide (of type $\textrm{Co}_3\textrm{Mo}_3\textrm{C}$) is stable in the temperature ranges of 900C to 1300C and can vary in composition from Mo_40_Co_46C_14 to Mo56Co30C14, while the $Mo_12C$ carbide of type ($Co6Mo6C$ carbide decomposes into $Mo_6C$ and $\mu-Mo$ phases above 1100C \cite{zhangThermodynamicModelingCCoMo2016}.
When present in large quantities, W and Mo also participate in formation of W-rich or Mo-rich carbides during alloy solidification \cite{davis2000nickel, raghuRecentDevelopmentsWear1997},
leading to generation of Topologically Close-Packed (TCP) phases, such as the μ phase (of type Co_7W6 and Co7Mo6) and σ phase (pf type Co3W and Co3Mo) \cite{zhangThermodynamicInvestigationPhase2019}, which are intermetallic brittle phases that add strength to the material \cite{yuTriboMechanicalEvaluationsCobaltBased2007, ishidaIntermetallicCompoundsCobase2008} while also promoting crack initiation and propagation \cite{zhaoFirstprinciplesStudyPreferential2023}. Previous work on the Stellite 1 sample by Ahmed et al \cite{ahmedMappingMechanicalProperties2023} indicate that Co6W6C is identified as the main W-rich carbide in Stellite 1, although Co3W3C was also identified inaddition to Co3W and Co7W intermetallics.
There are two main phases in the tungsten-carbon system: the hexagonal monocarbide $\textrm{WC}$ (ICDD Card# 03-065-4539, COD:2102265), denoted as $\delta-\textrm{WC}$, and multiple variations of hexagonal-close-packed subcarbide $\textrm{W}_2\textrm{C}$ (ICDD:00-002-1134, COD:1539792) \cite{kurlovPhaseEquilibriaWC2006, tulhoffCarbides2000}
WC carbides precipitate as discrete particles distributed heterogeneously throughout the alloy intragranularly
The precipitation of the tungsten-rich phase $M_6C$ is closely related to the decomposition of the MC carbide, and the $M_6C$ only occurs in the vicinity of the MC \cite{jiangSecondaryM6CPrecipitation1999}, as $M_6C$ carbides form only when the tungsten and.or molybdenum content exceeds 4-6 a/o.
Paragraph: Chromium carbide ignore_heading
Chromium carbides have high hardness and wear resistance, as well as excellent resistance to chemical corrosion, making them often used in surface coatings \cite{liElectronicMechanicalProperties2011}
In the Cr-C binary phase diagram, there are three phases : cubic Cr23C6 (space group , melting point 1848 K), orthorhombic Cr3C2 (space group Pnma, melting point 2083 K) and Cr7C3 (space group Pnma, melting point 2038 K) \cite{medvedevaStabilityBinaryTernary2015} \cite{liElectronicMechanicalProperties2011}
The M_23C6 carbides are formed during heat treatment of carbides with a lower M/C ratio or from solid solution close to boundaries \cite{medvedevaStabilityBinaryTernary2015}. Fine M23C6 carbides act as obstacles to gliding of mobile dislocations, which result in long-term creep strength \cite{godecCoarseningBehaviourM23C62016}.
Although M23C6 can precipitate as primary carbide during solidification, it is most commonly found in secondary carbides along grain boundaries.
M7C3 is a metastable pseudo-eutectic carbide that typically forms at lower carbon-chromium ratios and effectively transforms into secondary M23C6 upon heat treatment.
In addition to being a carbide former, chromium provides solid solution strengthing and corrosion/oxidation resistance to the cobalt-based matrix.
The Cr7C3 carbide is unstable at high temperatures and transforms to M23C6 upon heat treatment. Under further temperature and time, Cr23C6 partially transforms to Cr_6C \cite{mohammadnezhadInsightMicrostructureCharacterization}.
$$ 2Cr_{7}C_{3} + 9Cr \rightarrow Cr_{23}C_{6} $$ $$ Cr_{23}C_{6} + 13Cr \rightarrow 6Cr_{6}C $$
$$ 23Cr_{7}C_{3} \rightarrow 7Cr_{23}C_6 + 27C $$ $$ 6C + 23Cr \rightarrow Cr23C6 $$
Paragraph 2: Fundamental Mechanisms of Corrosion and Cavitation Resistance ignore_heading
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Fundamental Strengthening Mechanisms
- Primary mechanism: Hard carbide precipitation (e.g., M$_{7}$C$_{3}$, M$_{23}$C$_{6}$), with dependence on carbon content and processing.
- Secondary mechanism: Solid solution strengthening by specific elements (W, Mo, Cr), also linked to carbon content.
- Additional mechanism: Stress-induced phase transformation (fcc to hcp) contributing to wear resistance via work hardening.
- How they are made and modified.
- Current research directions and future outlook.
The remarkable ability of Stellite alloys to withstand these specific challenges stems from key metallurgical features. Their corrosion resistance is primarily attributed to a high chromium content, typically 20-30 wt.%, which promotes the formation of a highly stable, tenacious, and self-healing chromium-rich passive oxide film on the material's surface; this film acts as a barrier isolating the underlying alloy from the corrosive environment. Alloying elements such as molybdenum and tungsten can further enhance this passivity, particularly improving resistance to localized corrosion phenomena like pitting and crevice corrosion in aggressive media. Concurrently, their outstanding cavitation resistance is largely derived from the unique behavior of the cobalt-rich matrix, which can undergo a stress-induced crystallographic transformation from a face-centered cubic (fcc) to a hexagonal close-packed (hcp) structure. This transformation, often facilitated by mechanical twinning, effectively absorbs the intense, localized impact energy from collapsing cavitation bubbles and leads to significant work hardening, thereby impeding material detachment and erosion.
Antony suggests that the cavitation-erosion resistance of Stellites derives from the matrix phase and is enhanced by the strain-induced fcc \textrightarrow hcp allotropic transformation \cite{antonyWearResistantCobaltBaseAlloys1983}.
Corrosion resistance of Stellites
The cavitation erosion of stellites has been investigated in experimental studies \cite{Wang2023, Szala2022741, Mitelea2022967, Liu2022, Sun2021, Szala2021, Zhang2021, Mutascu2019776, Kovalenko2019175, E201890, Ciubotariu2016154, Singh201487, Hattor2014257, Depczynski20131045, Singh2012498, Romo201216, Hattori20091954, Ding201797, Guo2016123, Ciubotariu201698}, along with investigations into cobalt-based alloys \cite{Lavigne2022, Hou2020, Liu2019, Zhang20191060, E2019246, Romero2019581, Romero2019518, Lei20119, Qin2011209, Ding200866, Feng2006558}.
Stellites achieve oxidation resistance through the formation of a passivating external Cr2O3 scale, due to the high proportion of Cr in their chemical composition \cite{pettitOxidationHotCorrosion1984}.
as seen by Zhang et al in Green Death solution \cite{zhangPittingCorrosionCharacterization2014}.
However, Cr-based carbides may be preferentialy oxidized below the external Cr2O3 scale, particularly at the boundary of carbides which are depleted of Cr \cite{zhangPittingCorrosionCharacterization2014}, where preferential attack of carbides proceed until they have been consumed \cite{pettitOxidationHotCorrosion1984}.
Lemaire et al \cite{lemaireEvidenceTribocorrosionWear2001} investigated the behavior of Stellite 6 in pressurized high temperature water and proposed an oxidative wear mechanism where wear proceeds by repeated detachment of the surface oxide spontaneously forming on the stellite surface.
Di Martino et al \cite{dimartinoCorrosionMetalsAlloys2004} also found that the protective chromium-rich film are abraded easily, leading to further corrosion.
In such lower-temperature regimes, the passive films formed are typically very thin (in the nanometer range, rather than the micrometer scale observed at high temperatures)
Molybdenum and tungsten have favorable effects on the selective oxidation of chromium until chromium has been depleted, at which point molybdenum and tungsten result in increased oxidation due to development of less protective phases \cite{pettitOxidationHotCorrosion1984}.
Analytical Investigations
Strain hardening
Cavitation bubble collapse induce a work hardening of the material surface, comparable to that obtained in conventional peening \cite{swietlickiEffectsShotPeening2022}, characterized by the thickness of the hardened layers and the shape of the strain profile below the surface.
The strain profile within the material can usually be modeled by the following power law:
\begin{equation} \epsilon\left(x\right) = \epsilon_s {\left( 1 - \frac{x}{L} \right)}^{\theta} \end{equation}where $\epsilon\left(x\right)$ is the strain at depth $x$ from the eroded surface, $\epsilon_s$ is the failure rupture strain on the eroded surface, $L$ is the thickness of the hardened layer, and $\theta$ is the shape factor of the power law.
After each cycle, the thickness of the hardened layer $L$ and the surface strain $\epsilon_s$ will increase continuously until damage is initiated at the surface ($\epsilon_s$ reaches the failure rupture strain $\epsilon_R$), at which point the strain profile is in steady-state.
\begin{equation} \epsilon_R = \epsilon_{mean} {\left( 1 - \frac{\Delta L }{L+ \Delta L} \right)}^{\theta} \end{equation}In Woodford's investigations on the $\gamma$\textrrightarrow$\epsilon$ transformation on the surface of cobalt-base alloys during cavitation erosion, the transformed layer was found to extend to a depth of 25 to 50 \um \footnote{Converted from the original 1e-3 to 2e-3 \cite{woodfordCavitationerosionlnducedPhaseTransformations1972}}.
Correlative empirical methods
Empirical methods are common for addressing complex cavitation erosion, involving lab tests to correlate cavitation erosion resistance with mechanical properties.
Karimi and Leo
The Karimi and Leo phenomenological model describes cavitation erosion rate as a function of
Karimi and Leo
Noskievic
Noskievic formulated a mathematical relaxation model for the dynamics of the cavitation erosion using a differential equation applied to forced oscillations with damping:
\begin{equation} \frac{\mathrm{d}^2 v }{\mathrm{d}t^2} + 2 \alpha \frac{\mathrm{d} v }{\mathrm{d}t} + \beta^2 v = I \end{equation}where $I$ is erosion intensity, which can vary linearly with time, $v = \frac{\mathrm{d} v }{\mathrm{d}t}$ is erosion rate, $\alpha$ is strain hardening or internal friction of material during plastic deformation, and $\beta$ is coefficient inversely proportional to material strength. The general solution of equation can be written as:
Hoff and Langbein equation
Hoff and Langbein proposed a simple exponential function for the rate of erosion, representing the normalized erosion rate requiring only the A simple exponential function for the rate of erosion was proposed by Hoff and Langbein,
$$ \frac{ \dot{e} }{ \dot{e_{max}} } = 1 - e^{\frac{-t_i}{t}} $$
$\dot{e}$ - erosion rate at any time t $\dot{e_{max}}$ - Maximum of peak erosion rate $t_i$ - incubation period (intercept on time axis extended from linear potion of erosion-time curve) $t$ - exposure time
L Sitnik model
$$ V = V_o {\left[ ln\left( \frac{t}{t_o} + 1 \right) \right]} ^ {\beta} $$
$$ \dot{V} = \frac{\beta V_o}{t + t_o} {\left[ ln \left( \frac{t}{t_o} + 1 \right) \right]}^{\beta - 1}$$
V_o > 0 t_o > 0 β >= 1
Experimental Investigations
Materials and Microstructure
The HIPed Stellite 1 alloys were manufactured by canning the gas-atomized powders, manufactured by Deloro Stellite (UK), at a temperature and pressure of 1200C and 100 MPa, for 4 hours in a HIPing vessel where the chemical composition and sieve analysis have been reported in previous work by Ahmed et al \cite{ahmedInfluenceAlloyComposition2025}.
The sieve analysis of these powders indicate that the majority of powder particles were in the size range of 45 to 180 um.
The cast alloy samples were produced via sand casting process.
| Co | Cr | W | Mo | C | Fe | Ni | Si | Mn | |
| HIPed Stellite 1 | Bal. | 31.70 | 12.70 | 0.29 | 2.47 | 2.30 | 2.38 | 1.06 | 0.26 |
| +250 | +180 | +125 | +45 | -45 | |
| HIPed Stellite 1 | 0.10 | 2.40 | 47.90 | 49.50 | 0.10 |
Materials and Microstructure
The HIPed alloy was produced via canning the gas-atomized powders at 1200C and 100 MPa pressure for 4h, while the cast alloys were produced via sand casting. % Sieve analysis and description of powders
% Refer to Table of chemical compositions of both cast and HIPed alloys.
The microstructure of the alloys were observed via SEM in BSE mode, and the chemical compositions of the identified phases developed in the alloys were determined via EDS as well as with XRD under Cu $K_{\alpha}$ radiation.
Image analysis was also conducted to ascertain the volume fractions of individual phases.
The Vickers microhardness was measured using a Wilson hardness tester under loads of BLAH. Thirty measurements under each load were conducted on each sample.
Paragraph XRD: Experimental determination of $\epsilon-Co$ ignore_heading
The microstructure phase identification was investigated out using X-ray diffraction technique with Cu-K$\alpha$ radiation ($\lambda = 1.5406 \AA{}$). The volume fraction of ε-Co can be determined using the intensity of the $(200)_{\gamma}$ and $(10\bar{1}1)_{hcp}$ peaks, using the following equation proposed by Sage and Guillaud:
\begin{equation} \textrm{hcp} (\textrm{vol}\%) = \frac{I(10\bar{1}1)_{\epsilon}}{I(10\bar{1}1)_{\epsilon} + 1.5I(200)_{\gamma}} \end{equation}Experimental determination of SFE
To experimentally determine the SFE, the XRD method proposed by Reed and Schramm was employed \cite{reedRelationshipStackingfaultEnergy1974}:
\begin{equation} SFE = \frac{K_{111} \omega_0 G_{111} a_0}{\pi \sqrt{3}} \frac{{<\epsilon_{50\AA}^2>}_{111}}{\alpha} A^{-0.37} \end{equation}where:
\begin{equation} K_{111}\omega_0 &= 6.6 \\ G_{111} = \frac{1}{3}\frac{1}{C_{44} + C_{11} - C_{12}} A &= \frac{1 C_{44}}{C_{11}-C_{12}} \end{equation}ployed \cite{reedRelationshipStackingfaultEnergy1974}:
$SFE$ = stacking fault energy $\frac{mJ}{m^2}$ $K_{111}\omega_0$ = 6.6, as obtained by $A$ is the Zener elastic anisotropy $C_{ij}$ are elastic stiffness coefficients $G_{111}$ is the shear modulus of the (111)-plane, in which stacking faults are formed $a_0$ is the lattic constant of the fcc-metal matrix ${<\epsilon^2_{111}>}_{50\AA}$ is thr root mean square microstrain in the <111> direction averaged over the distance of $50 \AA$ $\alpha$ is stacking fault probability
Elastic constant
Microhardness ignore_heading
Microhardness measurements were taken on the surfaces of the as-cast and HIPed samples. The Wilson Tukon 1102 hardness tester was used for Vickers microhardness testing with a load of 300 grams (HV0.3) for 10s, and averaged by using ten individual indentations. The specimen surface was prepared in the same fashion as for microstructural analysis.
Previous work
Indentation fracture toughness ignore_heading
The indentation fracture toughness was made with hardness equipment (AVK-A, AKASHI) at a load of 49 N for 10 s, and the value was obtained from five measurements on the cross section. The fracture toughness was evaluated based to the Evans-Wilshaw equation [21, 22].
\begin{equation} K_{IC} = 0.079 {\left( \frac{P}{a^{\frac{3}{2}}} \right)}log{\left(\frac{4.5 a }{c}\right)} \end{equation}where $P$ is indenter load $[\textrm{mN}]$, $2c$ is the crack length $\left[\mu\textrm{m}\right]$, and $2a$ is the length of indentation diagonal $\left[\mu\textrm{m}\right]$
Electrochemical instrument and experiments
A Corrtest potentiostat was used for electrochemical experiments in a conventional three electrode cell, with the sample as working electrode with exposed area 2cm2, a saturated calomel electrode (SCE) as reference electrode, and a Pt plate as counterelectrode.
All electrochemical experiments were performed at room temperature.
The open circuit potential was continuously recorded for 1 h, before the electrical impedence spectroscopy (EIS), LPR, and cyclic voltametry experiments.
Cavitation Erosion Test Apparatus ignore_heading
A magnetostrictive vibratory apparatus, operating in general accordance with ASTM Standard G32, was utilized. The system functioned at an ultrasonic frequency of 20 kHz, with a peak-to-peak displacement amplitude of the horn tip maintained at 96 µm. The horn, fabricated from a cavitation-resistant titanium alloy, featured a flat tip of 16 mm diameter. Experiments were conducted using a stationary specimen configuration, with the specimen positioned 0.5 mm below the vibrating horn tip, this distance being precisely set using a dial gauge.
Paragraph: Lit review of corrosion ignore_heading
The aqueous oxidation of Stellite 6 alloy was investigated in a 1979 study using X-ray Photoelectron Spectroscopy (XPS) \cite{mcintyreXRayPhotoelectronSpectroscopic1979}. Specimens were exposed to pH 10 water at 285°C. To understand the oxidation behavior, the study measured dissolved oxygen concentration against exposure duration.
The high-temperature corrosion resistance of stellite coatings is attributable to the formation of cobalt & chromium surface \cite{cesanekDeteriorationLocalMechanical2015}.
Heathcock et al found that carbides are selectively eroded, with the carbide-matrix interface acting as initiating erosion site \cite{heathcockCavitationErosionCobaltbased1981}.
Paragraph 4: Synergistic Challenges in Applications Prone to Corrosion and Cavitation ignore
Paragraph 5: Research and Development for Enhanced Corrosion and Cavitation Performance ignore
Paragraph 6: Influence of HIPing ignore
Compared with the case alloys, the HIPed alloys had relatively finer, rounded, and distributed carbides.
Paragraph: Cavitation Erosion Resistance
The primary result of an erosion test is the cumulative mass loss versus time, which is then converted to volumetric loss and mean depth of erosion (MDE) versus time for the purposes of comparison between materials of different densities. The calculation of the mean depth of erosion for this test method should be performed in conformity with ASTM G-32.
General Background
%cite:@Franc2004265, @Romo201216, @Kumar2024, @Kim200685, @Gao2024, @20221xix, @Usta2023, @Cheng2023, @Zheng2022
Cavitation erosion presents a significant challenge in materials degradation in various industrial sectors, including hydroelectric power, marine propulsion, and nuclear systems, stemming from a complex interaction between fluid dynamics and material response \cite{francCavitationErosion2005, romoCavitationHighvelocitySlurry2012}. Hydrodynamically, the phenomenon initiates with the formation and subsequent violent collapse of vapor bubbles within a liquid, triggered by local pressures dropping to the saturated vapor pressure. These implosions generate intense, localized shockwaves and high-speed microjets that repeatedly impact adjacent solid surfaces \cite{gevariDirectIndirectThermal2020}. From a materials perspective, these impacts induce high stresses (100-1000 MPa) and high strain rates, surpassing material thresholds and leading to damage accumulation via plastic deformation, work hardening, fatigue crack initiation and propagation, and eventual material detachment. Mitigating this requires materials capable of effectively absorbing or resisting this dynamic loading, often under demanding conditions that may also include corrosion.
% Martensitic transformation Crucially, the cobalt matrix often possesses a low stacking fault energy, facilitating a strain-induced martensitic transformation from a metastable face-centered cubic $\gamma$ phase to a hexagonal close-packed $\epsilon$ phase under the intense loading of cavitation. This transformation is a primary mechanism for dissipating impact energy and enhancing work hardening, contributing significantly to Stellite's characteristic cavitation resistance \cite{huangMicrostructureEvolutionMartensite2023, tawancyFccHcpTransformation1986}.
HIPing is a thermo-mechanical material processing technique which involves the simultaneous application of pressure (up to 200 MPa) and temperature (2000 C), which results in casting densification, porosity closure, and metallurgical bonding. \cite{yuComparisonTriboMechanicalProperties2007}
While commonly applied via casting or weld overlays, processing routes like Hot Isostatic Pressing (HIP) offer potential advantages such as microstructure refinement \cite{stoicaInfluenceHeattreatmentSliding2005} finer microstructures and enhanced fatigue resistance \cite{ahmedInfluenceReHIPingStructure2013, yuComparisonTriboMechanicalProperties2007}.
HIPing of surface coatings results in microstructure refinement, which can yield improved fatigue and fracture resistance.
HIPing leads to carbide refinement, which can yield improved impact toughness \cite{yuInfluenceManufacturingProcess2008}, and reduce carbide brittleness \cite{yuComparisonTriboMechanicalProperties2007}.
Furthermore, HIP facilitates the consolidation of novel 'blended' alloys created from mixed elemental or pre-alloyed powders, providing a pathway to potentially tailor compositions or microstructures for optimized performance. However, despite the prevalence of Stellite alloys and the known influence of processing on microstructure and properties, the specific cavitation erosion behavior of HIP-consolidated Stellites, particularly these blended formulations, remains underexplored in academic literature. Given that erosion mechanisms in Stellites often involve interactions at the carbide-matrix interface \cite{szalaEffectNitrogenIon2021}, understanding how HIP processing and compositional blending affect these interfaces and the matrix's transformative capacity under cavitation, especially when potentially coupled with corrosion, constitutes a critical knowledge gap addressed by this research.
% Need to describe Stellite 1 §ion{Stellite 1}
Stellite 1 is a high-carbon and high-tungsten alloy, making it suitable for demanding applications that require hardness & toughness to combat sliding & abrasive wear \cite{crookCobaltbaseAlloysResist1994}
§ion{Stellites} §ion{Objectives and Scope of the Research Work} §ion{Thesis Outline} §ion{Literature Survey} §ion{Cavitation Tests}
COMMENT Cavitation erosion mechanisms based on erosion particles
Discussion
§ion{Experimental Test Procedure} ⊂section{Hardness Tests} ⊂section{Cavitation}
§ion{Relationships between cavitation erosion resistance and mechanical properties} §ion{Influence of vibratory amplitude}
% Insert the whole spiel by that French dude about displacement and pressure (and then ruin it) The pressure of the solution depends on the amplitude of the vibratory tip attached to the ultrasonic device. Under simple assumptions, kinetic energy of cavitation is proportional to the square of the amplitude and maximum hammer pressure is proportional to A.
\begin{align} x &= A sin(2 \pi f t) \\ v &= \frac{dx}{dt} = 2 \pi f A sin(2 \pi f t) \\ v_{max} &= 2 \pi f A \\ v_{mean} &= \frac{1}{\pi} \int^\pi_0 A sin(2 \pi f t) = 4 f A \\ \end{align}However, several researchers have found that erosion rates are not proportional to the second power of amplitude, but instead a smaller number. Thiruvengadum \cite{thiruvengadamTheoryErosion1967} and Hobbs find that erosion rates are proportional to the 1.8 and 1.5 power of peak-to-peak amplitude. Tomlinson et al find that erosion rate is linearly proportional to peak-to-peak amplitude in copper [3]. Maximum erosion rate is approximately proportional to the 1.5 power of p-p amplitude [4]. The propagation of ultrasonic waves may result in thermal energy absorption or into chemical energy, resulting in reduced power. For the purposes of converting data from studies that do not use an amplitude of 50um, a exponent factor of 1.5 has been applied.