Shreirs Corrosion 4 vol set

Electrochemistry and Corrosion Science
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The key issues on oxide formation are discussed below. Wear that is dependent on the tendency to oxidation is highly dependent on temperature. Jiang et al. These mechanisms and modes of formation are discussed in more detail in Section 4. At a critical thickness the oxide layers are no longer able to withstand the forces acting tangentially on them, and suffer failure and break away.

Knowing the critical thickness of the oxide [61] and the static oxidation properties of the wearing materials, wear rate prediction can be easily achieved. However, in this case, the time required to reach this critical thickness depends on the contact frequency F, which is the inverse of the elapsed time between two contacts at a given point between the contacting surfaces; whilst this can clearly be related to sliding speed, the frequency of contact can also be changed by varying the length of the wear track without any need to vary sliding speed.

It is clear that each asperity will not make contact each time the disk rotates; however, Garcia [83] comments that the probability of a contact and hence a wear particle being generated is included in the statistical meaning of the wear coefficient, Ka. For this to work, it is necessary for a wear particle to be generated each time a contact is made; this in itself is highly improbable.

The wear process was characterised by two transitions; T 1 being mild-to-severe at low load and sliding speed and T2 being severe-to-mild at high load and sliding speed. Increasing the sliding speed decreased the critical load at which these transitions occurred Fig. The variation in the upper transition from the intermediate severe wear back to higher speed mild wear, was observed to be the more sensitive to sliding speed. Low speed, low load mild wear was attributed to the presence of loose oxide debris at the sliding interface and intermediate severe wear to direct metal-to-metal wear.

The mild wear encountered at high speed and high load was attributed to hardening, accompanied by the development of an adherent oxide film, as a result of frictional heating. The hardening came about as a result of the low carbon steels undergoing phase changes, due to high localised temperatures around points of contact being sufficient to produce a transformation to austenite, followed by rapid cooling by conduction of heat into the bulk metal producing a structure at the surface not too dissimilar to martensite.

A critical hardness had to be exceeded by these phase changes for mild wear to be re-established under high speed, high load conditions; the transition back from severe wear to mild wear is in fact a two part transition, with T2 referring to the point where sufficient phase hardening occurs to suppress severe wear without the intervention of an oxide film the development of which further acts to protect the wear surface and a T3 transition approximately matching the point where permanent phase change hardening occurs.

Subramanian [85] conducted a series of sliding tests of an Al The wear rate of the Al Further increases in speed above this critical value led to progressive increases in wear. Subramanian explains the decrease in wear with increasing sliding speed due to increasing strain rates and due to increased hardness and flow strength of the wear surface. The true area of contact is thus and with a lower level of contacts between the wearing surfaces, a lower wear rate results.

This results in an increase in the true area of contact and thus an increase in the wear rate. It is not stated whether the particles produced at any particular speed are metallic or oxide. The critical speed was observed to be dependent on counterface material and a higher transition was noted for harder, more thermally conductive alloys. Decreased mutual solubility also led to a higher transition speed. Welsh [26,27] discusses the existence of lower limits of load and speed, marking the transition from mild to severe wear and also an upper limit, marking the transition back to mild wear.

So [86] on the other hand only discusses a single limit or critical value for both load and speed for the transition from mild wear at low speeds and loads, to severe wear at high speeds and loads. In one test of note, a high carbon steel sample underwent a mild-to-severe wear transition at a contact pressure of 4. A Stellite sample remained in the mild wear state at a contact pressure of 8. As So used a pin-on-disk configuration, compared to the pin-on-rotating-cylinder configuration of Welsh, this may account for the differences in results; the pin-on-disk configuration may not have generated sufficiently severe conditions for the upper transition to occur.

Most load and sliding speed work done to date has concentrated on what happens at room temperature, with little work at elevated temperature. Amongst the most comprehensive work carried out on the sliding wear of cobalt was that by Buckley [87], who compared the sliding wear in vacuum of cobalt with that of copper. The differences between the sliding behaviour of metals in hexagonal close-packed phase and face-centred cubic phase are due to the greater number of active slip systems available in face-centred cubic structures. There are twelve primary slip systems within a typical face-centred cubic metal 4 slip planes each with three slip directions , which are all crystallographically similar, compared to only three primary slip systems in cobalt, these being based on the basal plane with the highest atomic density i.

Cross slip is also more difficult, as with hexagonal close-packed structures such as cobalt, screw dislocations are required to move out of the primary basal glide plane onto planes that unlike face-centred cubic structures, are crystallographically different. Hexagonal close-packed materials are thus less deformable. The effect of these hexagonal close-packed structures on wear was further elaborated on by Persson [88,89] on studying the low friction tribological properties of Stellite 21 and Stellite 6. A thin, easily sheared layer can develop at the sliding surface due to the shear-induced alignment of the hexagonal close-packed basal plane parallel to the direction of sliding [89].

This alignment significantly reduces friction and improves galling resistance, with shear and adhesive transfer restricted to this layer. This sliding regime persists even as this layer is removed, as it is easily regenerated. Also, the removal of material in such thin sections may at least in part explain the ready generation of fine Stellite 6-sourced Co- based oxide debris observed elsewhere [, ]. These sections may be more easily commutable to a small size and oxidised, providing a ready supply of material for 'glaze' formation.

The formation of these oxides was by a similar route to that observed for iron-based and nickel-based alloys, with alloying components present in the oxides to roughly the same proportions as the original alloy. He specifies an initial period of low wear for up to an hour, followed by the production of a bright, rough metallic wear scars showing characteristics of abrasion and evidence of material transfer.

This he attributes to a probable change in phase from hexagonal close-packed to face-centred cubic and thus a loss in wear resistance. Later, the bright worn surface is lost with increasing amounts of oxide being produced, although the load-bearing areas remain metallic. In a fretting wear situation, it is difficult to see how this could occur. However, the level of alteration of temperature for any phase transitions will also depend on the effects of other alloying components in cobalt-based alloys. As already stated, chromium will raise the transition temperature quite dramatically.

Other references [68, ] suggest that tungsten and molybdenum also raise this transition, whilst nickel and iron also, magnesium and carbon [88,89] have the effect of stabilising the higher temperature face-centred cubic structure due to increases in stacking fault energy [88,89] and suppressing this transition. It is possible that the presence of nickel to Thus a much smaller increment in temperature due to frictional heating and flash temperatures may be needed to effect any phase transition, confirming the conclusion that the damage observed can be attributed to phase changes and hence a decrease in resistance to deformation.

At elevated temperatures, Stott et al. The face-centred cubic to hexagonal close-packed transformation observed in Co-based alloys is considered a martensitic transformation [91]. After sliding for up to 10, m, the Stellite 6 layer was observed to be mostly covered by an oxide layer reported to consist of W3O, CrO and Co2O3. Where this oxide layer spalled, a new oxide film was observed to replace it readily. This applied for all combinations of loads The experimental data obtained from these tests are presented in Fig.

The steels underwent increased wear compared to the Stellite 6, despite being of much greater hardness and So et al. When used as a disk material, only a thin layer of oxide material was formed on the AISI , compared to the thicker layer formed on the Stellite 6; the wear rate of the steel was seven times that of the Stellite 6 laser-clad pin. As the pin material, severe wear was observed for the AISI steel, the rate of wear being 10 times higher than that of the Stellite 6 laser-clad disk. For all but the highest load, the wear rate of the Stellite 6 pin actually decreased when the sliding speed was increased from 1 to 2 m.

At the lowest load used, the decrease in wear continued up to 4 m. The increasing wear rate for specimens under a load of The higher flash temperatures also led to changes in the oxide phases that were reported to form on the respective wear surfaces. So does not offer an explanation for the change in oxide with temperature, though as for the oxidation of iron, it appears that this can be attributed to changes in oxidation state of the chromium in Stellite 6, with preferential oxidation of tungsten and cobalt respectively at lower temperatures.

It is curious to note here that, in the work of Wood [1] and Rose [2], no such shift was observed for Stellite 6, with Cr2O3, Co3O4 or a combined oxide of the two being consistently observed from XRD results. So comments that, under the most severe conditions This implies that So has underestimated the temperature at the sliding interface in this case. The softening may again be attributable to phase changes from hexagonal close-packed to face-centred cubic. Where no cobalt was present within the alloy, wear rates were observed to be highest.

Increases in wear were observed for all combinations with increased contact pressure, though at high load, increases became less severe for cobalt-chromium and cobalt-iron-chromium alloys. Of particular note is the response to increasing the sliding speed by an order of ten from the 7. For the high cobalt-chromium alloys including Stellite 1 and Stellite 6, there was a slight decrease in the observed wear rate.

Where cobalt levels were low or non-existent, the converse was true and increases in wear were observed. Crook and Li [69] attributed the superior wear resistance of the cobalt-chromium alloys, firstly to the superior galling resistance and secondly to the tendency of alloys when in the face-centred cubic form to undergo phase changes and become hexagonal close-packed, which as discussed earlier is less prone to deformation, due to a smaller number of available slip planes. Conversely, they point out that high nickel alloys have a poor galling resistance, yet specifically quote the work of Stott et al.

However, Stott et al. In both cases, it is not possible to say that in an extreme high wear environment e. Even in a like-on-like situation, wear properties of Nimonic 80A are inferior to those of Stellite 6 [2,93]. In the cast form, carbon combines with chromium to form a chromium carbide phase at the grain boundaries; in Stellite alloys, these are of the form M7C3 and M23C6 [29].

The presence of these hard, difficult-to-deform carbides may have had a number of effects on both sets of experimental work. Firstly, they may have further inhibited deformation of the mainly cobalt matrix during sliding wear, over and above the effect expected from the hexagonal close-packed structure, blocking the operation of the fewer slip planes present. Secondly, the removal of material from the Stellite alloys may have released some of these carbides into the sliding interface, increasing the levels of wear observed due to increased abrasion effects.

The possible effects of carbides are discussed in more detail in Sections 2. In many, experimentation has concentrated only on single-phase alloys.

However, second phases are used in many alloys for various reasons, including enhancement of strength and creep resistance, especially in high temperature systems where the properties of the metallic matrix can become less robust. Vardavoulias [94] studied a number of steels into which hard ceramic phases of various sizes were introduced; these included titanium carbide modified to a much finer carbon nitride phase by nitrogen annealing , copper phosphide and alumina.

They may not protrude above the surface of the oxide layer and thus cannot directly protect the matrix or impinge on the counterface material. The particles are lost as the oxide layer breaks up at the critical thickness to form debris. The only contribution may be to enhance the load-carrying capacity of the metallic matrix in supporting the oxide film. Whilst a small quantity of the second phase may be removed with the oxide, as it breaks up, most will remain embedded in the substrate or matrix.

These particles protrude above the nominal surface of the interface and the counterface slides over them. Enhanced wear of the counterface material by abrasion may occur in this stage.

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A further possibility is detachment of second phase particles as the oxide breaks up, with these particles acting as third body abrasives; this occurs where cohesion between matrix and second phase is poor. If the mean particle size is much greater than the critical oxide thickness, the particles show increased efficiency in providing oxidational wear protection to the material subject to wear.

Again, the main interaction is between the particles and the counterface and this process controls the wear mechanism; the matrix plays no direct part. The majority of the particles are surrounded by the matrix, thus break up is more difficult and detachment is almost impossible. The inference here is that the first stage cannot resume until these larger particles wear to near to the level of the rest of the sample surface; as other particles will continue to be exposed elsewhere on the surface, first stage wear with protective oxide layer formation cannot readily happen and severe wear will continue.

Tribaloy alloys, consisting of a hard Laves intermetallic second phase in a Ni-based or more notably a Co-based matrix [95] are possible examples not discussed in the current work. The section starts with a brief introduction to ODS alloys to provide a context and facilitate interpretation of the experimental observations on their HT wear characteristics. They derive their strength from the insoluble, deformation resistant and thermally stable dispersoids such as Y2O3 as used in Incoloy MA [96] introduced during fabrication by mechanical alloying. This process offers great opportunities in the selection of dispersoids-matrix combination.

Strengthening by insoluble, deformation resistant and inert dispersoids is far more effective than other methods such as solid solution and precipitation hardening as thermodynamics imposes limitation on their continued effectiveness. During high temperature deformation, most ODS alloys exhibit a threshold stress o below which creep becomes negligible; the threshold stress o is less than Orowan stress or. Several models such as dislocation climb, dislocatio n detachment local climb have been advanced to explain the existence of the threshold stress.

None of these models have been universally accepted. To increase the temperature capability of ODS alloys, the powder produced by mechanical alloying is subjected to hot extrusion and the fine grained extruded structure is then subjected to a high temperature secondary recrystallisation anneal.

The overall enhancement of HT capability may be associated with: Large grain size produced by secondary recrystallisation; High GAR Grain Aspect Ratio produced by recrystallisation under a high temperature gradient and minimising the detrimental effect of transverse grain boundary effect; Formation of serrated grain boundaries allowing grain inter-locking; and Minimum grain boundary hardening. The apparent improved wear resistance observed at room temperature is attributable to transfer [] and work hardening of a layer of Incoloy HT or back-transferred Incoloy MA [3], protecting the ODS alloy surface from sustained wear Fig.

The loss of strength and hardness suffered by the ODS alloys at higher temperatures [] Fig. Severe wear continues Fig. Some transfer of Incoloy HT material still occurs at high temperature; however, continued sliding removes this layer [1,3]. Such transfer tends to be greater at higher sliding speeds and, hence, the weight loss is lower at 0. The relative contributions of sample- and counterface-sourced debris at least for Incoloy MA versus Stellite 6 have been observed to depend on sliding speed [3,22]. A slow sliding speed -1 0. Faster sliding speeds 0. Below these temperatures, the oxide separates the surfaces even in the form of the aforementioned loose debris.

No severe wear is observed at 0. No significant increases in weight loss were observed with increasing load up to 20N load Fig. From such bonding comes crystal structure, ordering, high strength at low and high temperature, low ductility and low K1c, particularly at low temperature. The low ductility, low K1c and high strength of intermetallics stem from such critical factors as complex crystal structures, the large Burgers Vector, high lattice stress, the in- adequate slip systems and the inability to cross slip. The expected improvement in wear resistance from high strength, ordering and an adequate slip system may be off set by low K1c and low fracture strain.

The advantage in providing high wear resistance due to ordering may eventually be lost by the destruction of this ordering due to elemental diffusion from the counterface into the intermetallic lattice. When worn against Stellite 6 7N, 0. The high TiAl wear has been attributed to the elements of the counterface material interfering with the ordered structure of the intermetallic, as small atomic radii Si and N enter the lattice and increase its susceptibility to wear.

A mixture of abrasive and adhesive wear with some evidence of stick-slip, possibly due to high mutual chemical compatibility between the Si3N4 and the TiAl, has been reported [98] Si3N4 has been previously observed to have high mutual chemical compatibility with both Ti and Al [99] , leading to moderate wear of the TiAl and high wear of the Si3N4. The formation of this layer, sourced from both the TiAl and Si3N4, prevents substantial wear.

It is suggested that any high mutual compatibility enhances the very early stages of wear, providing the necessary material for the wear-resistant layer. TiAl undergoes higher wear when worn against Al2O3 at room temperature [97] with enhanced material removal by abrasion, the debris from which does not form a protective layer.

The results of room temperature tests seemed to be greatly influenced by the degree of material transfer to the Nimonic surface, this transfer itself being influenced by counterface hardness; the softer was the counterface the higher was the amount of transfer.

Metallic transfer readily occurred from the relatively soft Incoloy HT counterface onto the surfaces of each of the Nimonic materials at room temperature. The transferred material formed a wear resistant layer which protected the Nimonic material surfaces and led to very small weight changes. Inman [3] later reported the development of metallic transfer layers at 0. When slid against Incoloy HT, mainly Incoloy HT-sourced layers that protected the Nimonic alloys against wear, together with low weight changes, were observed. Rose 0.

Limited compacted oxide layers formed at 7N Fig 23a and 10N load, protecting the Nimonic 80A surface. At loads between 15N to 25N Fig. This was reflected by a rapid increase in weight loss above 15N Fig. Thin mixed oxide layers were formed on the Nimonic surfaces; these contained material from both the Nimonic and Stellite 6 wear surfaces and provided only limited protection. Testing of Nimonic 80A against a Stellite 6 counterface at 0. This debris separated the wear surfaces, preventing metallic contact; a low temperature mild wear regime resulted.

Behaviour at intermediate and high temperature depended significantly on sliding speed. Such severe wear was observed at 0. A similar pattern was observed at 0. However, high wear rates continued; although technically a mild wear regime, the high wear rates indicated that the oxide was not protective. Two slightly different forms of oxidational behaviour were observed, depending on researcher. These latter observations were confirmed by Inman [3,21,23] at both 0. Rose [2] stated that layers were unable to form due to insufficient debris adhesion to the Nimonic 80A sample surfaces and lack of debris cohesion, caused by the ploughing of sample surfaces by hard carbide particles in the Stellite 6 counterface.

This is discussed further in Section 5. At all loads 7N to 25N , the wear mechanism remained predominantly oxidational, with elements of abrasive wear, regardless of applied load Fig. There was no evidence that changing load had any significant effect on wear regime in this case; however, the grooves on the damaged surface became more pronounced at high load.

This was reflected by increased weight loss at high load Fig. Thin mixed oxide layers were formed on the Nimonic surfaces; these contained material from both wear surfaces that provided only limited protection. Under vacuum, increasing pressure from high vacuum conditions to Pa was enough to result in a decrease in friction in sliding of an iron-chromium alloy. Buckley [] noted during the like-on-like sliding of clean iron that a pressure of Pa or 3 Torr was sufficient to prevent seizure.

Barnes et al. On raising the partial pressure to Pa, however, significant amounts of oxide were observed and areas of compacted debris had developed. The compacted debris was either completely oxidised or oxide-covered metallic debris; its formation was accompanied by decreases in friction Fig. However, despite the presence of this oxide debris, the wear rate remained high until the oxygen partial pressure reached 1 Pa or above.

Changes in partial pressure were also made during sliding tests [], with oxygen in some cases being removed from the wear system the pressure was from Pa to Pa. When this occurred, the oxide debris and the compacted oxide layers remained at the wear interface, showing continued stability and wear resistance even without a continued supply of oxygen.

For mild steel a decrease in wear was observed with increasing relative humidity under fretting conditions [,] and similarly with carbon steel under sliding wear conditions [15]. It has been suggested that adsorbed moisture might have a dual effect [], in that on the debris surface, it might act as a lubricant, promoting speedier debris dispersal and, thus, less abrasive wear.

From this, it was proposed that the hydrated form of the iron oxide that develops in the presence of the moisture might be a less abrasive medium. Such debris does, however, have the potential to enhance interface contact and bring about adhesive wear. Experimental work by Bill [] demonstrated that this could be the case, with the relationship between relative humidity and wear rates becoming quite complex. Oh et al. At low relative humidity, severe wear was encountered, with total losses amounting to between 0.

The amounts of carbon in the steel were observed to affect this transition, which occurred at higher values with increasing carbon content. Friction was also observed to fall rapidly, from between 0. In most practical situations, corrosion product will not form in a reducing non-oxidising atmosphere and, thus, the formation of wear protective layers is not possible.

Only the presence of adsorbed gases or other volatiles will act to separate the wear surfaces, adhesion and, therefore, levels of wear and friction []. In other oxidising atmospheres, only carbon dioxide has been examined to any significant extent. Research in other environments is extremely limited.

The possibility of titanium nitride is extremely unlikely, if not impossible, due to the tests being carried out at room temperature. However, no attempt was made to analyse the debris or explain the result. Elimination of metal-to-metal contact was observed immediately on commencement of sliding in the case of Jethete M and also the immediate establishment of compacted oxide in the case of the stainless steel. Iwabuchi et al. Times of 5 minutes and 1 hour did result in progressive decreases in wear Fig. The observed decreases in wear occurred because of break-down of the oxide layer to form debris, the presence of which prevented metal contact and adhesion.

Although there was some scatter in the data after m of sliding, there was no evidence of any effect of pre-oxidation on overall wear, regardless of the time of pre-oxidation and it was concluded that pre-oxidation had no effect under the prescribed test conditions. Thus, although pre-oxidation can decrease or eliminate early metal-metal contact, this cannot be guaranteed. Subsequent tests indicated the complete elimination of the severe wear regime Fig. A decrease in severe wear was also observed when stainless steel underwent pre-sliding for m at room temperature, due again to the presence of an accumulated loose oxide layer.

In both cases, the availability of pre-existing oxide debris acted to prevent contact between the metallic interfaces. The presence of accumulated oxide from pre-sliding did not, however, lead to a decrease in the rate of wear during mild wear. A sliding speed of 28 x m. This was due to the oxygen in the surface layers assisting the formation of oxidised debris and, thus, decreasing the initial severe wear period.

Certain aspects of heat treatment of the alloys and the sliding conditions were observed to affect this. For example, in the case of the chromium and carbon steels, the improvement in wear resulting from oxygen ion implantation was noticeably less for AISI and AISI B in the annealed form compared to the martensitic form; this can be seen from the data in Tables 3 and 4.

Relative humidity has a marked effect, as in the case of AISI steel, where the oxygen ion implanted material actually undergoes a higher level of wear than the untreated material. These observations were attributed to the higher plasticity of the annealed samples. A change in the form of the debris was also observed, from a smooth oxide layer for the martensitic samples to loose debris for the annealed samples.

The one exception was for AISI B steel, where the decrease was greater in the annealed state and sliding was accompanied by a change in the state of the oxide debris from the loose form to the oxide layer form. In comparison to standard pre-oxidation treatment Table 4 , wear tended to be less, with the exception of annealed AISI , where the pre-oxidised samples produced superior results, regardless of the levels of relative humidity.

Langguth et al. The generation of HT wear resistant surfaces in situ overcomes the serious limitations on materials and coatings imposed by HT wear conditions. These events, under certain conditions of temperature, pressure and speed [,12,,23,24,] lead to the formation of surfaces with self functionalised HT wear resistance. It was thus concluded that the observed low wear and friction arise from the physical properties and condition of the glaze, rather than their chemical compositions.

Further research [4,49,] allowed the identification of the following modes of compacted layer formation. The first mode is characterised by the formation of transient oxides, followed by the oxide thickening by continued oxidation by oxygen diffusion to the substrate-oxide interface and through physical defects. The second mode of formation is characterised by two stages. Stage one involves the formation of an insufficiently thick layer due to unfavourable temperature and low alloy strength, possibly involving an extended pre-glaze or severe run-in period. Stott et al.

These mechanisms [49] were seen as limiting cases for oxide debris generation, after which the build-up of oxide to form compacted layers continued: Oxidation — scrape — reoxidation: This involves a two-stage process. In the first step, oxide generation takes place in the areas of contact between the two sliding surfaces, with general oxidation over the apparent sliding area of contact and, also, at asperity contacts where temperatures exceed the general temperature in the region of the sliding area of contact.

In the second stage, this oxide is removed by subsequent traversals of the sliding interfaces, exposing fresh metal for further oxidation. The debris formed may then be either completely removed from the interface, act as a third body abrasive, thereby contributing to the wear process or be compacted to form a wear-protective oxide layer. Total oxidation: Under certain conditions, particularly high ambient temperatures, oxide generated during sliding or even present prior to the commencement of sliding, is not completely removed by subsequent traversals of the sliding interfaces, allowing the oxide to thicken with time.

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Provided this layer is coherent and adherent to the metal substrate and can withstand the stresses of sliding, a plastically deformed wear-protective oxide layer can develop. Metal debris: Debris particles generated during the early stages of wear are broken up by the sliding action, with any fresh areas of exposed metal being subject to further oxidation. There may be a high level of oxidation of the debris surfaces, due to the relatively large exposed surface area of metal.

Enhanced oxidation is promoted by heat of deformation and increased energy of the particles due to increased defect density and surface energy the exposed surface area of debris material will increase as particle size decreases. The resulting oxide can later develop into a wear-protective layer. The formation of compacted layers has also been observed by Wood et al. The one major difficulty with these mechanisms is that they were developed from work on low speed reciprocating sliding wear, where frictional heating is not such an important factor [60]. At sliding speeds of greater that 1 m.

While this adhesion force is weak, friction levels for mechanism 4 are lower than those for mechanism 2 ; however, increasing this adhesion force above a critical level results in a situation where the reverse is the case. Skidding then becomes the dominant mechanism, with no relative movement between neighbouring particles — an increase in adhesion force locks them in place. There is a transition from metal-metal wear to contact between these primarily oxide particle layers, at which point increases in contact resistance and decreases in levels of wear are coincident.

Removal of this more loosely compacted material by ultrasonic cleaning in acetone left behind the more compacted debris, load-bearing areas. These observations clearly indicate that temperature is a major driving force for adhesion between particles and formation of load-bearing compacted debris layers. The compacted layers formed during the sliding phase of the test became solidly sintered together as a result of the subsequent heating of the samples. The effect of a very small particle size would be to increase the available surface energy, due to the resultant increase in relative surface area.

This would act to drive the adhesion and sintering processes and allow for observable sintering at temperatures where sintering of the larger particles used in powder technology applications would not be noticeable. As adhesion itself is temperature-dependent, increases in temperature due to ambient or frictional heating would accelerate the adhesion and, therefore, the sintering process.

From experimental observations, Jiang et al. Generation of wear particles due to the relative movement of the metal surfaces; 2. Removal of some particles from the wear tracks to form loose wear particles; 3. Retention of other particles within the wear track; 4. Comminution of the retained particles by repeated plastic deformation and fracture, with particles freely moving between the rubbing surfaces and undergoing partial or even complete oxidation, due to continued exposure of fresh metallic surfaces during comminutation; 5. Continued fragmentation and agglomeration at various sites on the wear surfaces, due to adhesion forces between solid surfaces originating from surface energy and the formation of relatively stable compact layers.

This has two effects, viz: i.

Special order items

Firstly, material loss is reduced by a material recycling effect of the wear debris particles. Material breaking away from the compacted debris may rejoin it. Secondly, due to heavy deformation and oxidation of the wear debris particles, the layers formed are hard and wear-protective.

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Two competitive processes then occur during subsequent sliding, i. The compacted layers are continually broken down, the debris generated promoting wear though, again, reincorporation may occur. Continuing sintering and cold welding between particles within the layers, leading to further consolidation.

For the latter case to predominate, the temperature must be high enough in excess of a critical temperature to encourage the sintering processes required to ensure the formation of a solid wear-protective layer on top of the compacted particle layers before the layers are broken-down. The effects of this can be seen in the experimental work of Jiang et al.

However, Jiang does account for debris removal from the system leading to wear as would occur in, for example, high speed unidirectional sliding [,,] or debris removed by introduced interfacial airflow [45]. Incoloy MA versus Stellite 6 , despite the more adverse sliding conditions. Weight changes after 4, m of sliding were extremely low for all temperatures, with the largest mean change being 0.

The following discussion focuses on the situation at room temperature and oC. The data in Fig. Coefficient of friction levels Fig. The results indicate elemental transfer from the counterface to the specimen and mixing of the transferred and host element oxides. The associated selected area diffraction SAD pattern Fig. Dark-field images Fig. This is consistent with the structural variations observed by TEM. The latter exhibits a fine-grained structure 5 - 20 nm with irregular shaped grain boundaries. The particles are larger close to the interface up to about 50 nm.

The quantification is based on theoretical k-factors and uses a thickness correction for an estimated nm sample thickness EELS thickness measurements indicate a thickness of nm mean free path. Quantitative analysis of the Nimonic 80A layer Area 1 gives the characteristic composition of the bulk alloy Table 5 , apart from a slightly higher silicon concentration and a small amount of cobalt.

The interface layer consists of a mixture of Nimonic 80A and Stellite 6, with a higher than average titanium concentration []. There are some variations in the chromium, cobalt and oxygen concentrations as well as a few distinct particles with a high nickel concentration.

Furthermore, this line trace also confirms aluminium enrichment at the interface; moreover, this Al2 O3 layer is between the TiO2 layer and the Nimonic 80A substrate. The atomic numbers of chromium 24 , cobalt 27 and nickel 28 are similar and compositional variations of these main elements do not explain the strong contrast observed in the HAADF-STEM image Fig.

Furthermore, a low oxygen concentration coincides with a low chromium concentration; local EDX analysis shows a low chromium and oxygen concentration for the particles that appear bright in the image. This implies that some of the nickel and cobalt particles are not completely oxidised. EDX line traces in an area up to about one micron below the interface reveal the preferential segregation of light elements, especially aluminium and in some cases titanium Fig. These results clearly indicate the formation of a wear resistant nanostructured surface during sliding wear of Nimonic 80A against Stellite 6 at 20oC using a speed of 0.

The analyses reveal the complex structure of this surface, which consists of multiple layers: 1 A loose, uncompacted, highly oxidised layer. The nickel concentration from the Nimonic 80A is uniform throughout the whole layer, except for the presence of several larger non- oxidised particles that were randomly dispersed in the glaze.

The compositions indicate the presence of elements from both Stellite 6 and Nimonic 80A, but chromium- enriched. Slight chromium depletion was observed. The larger dark areas perpendicular to the interface were also due to aluminium enrichment. Additionally, it is apparent that the high Cr activity in both counterface and sample materials leads to the formation of Cr2O3 preferentially in the early stages of the process.

Quantification of the results on average gives The interfacial layer consists of grains of nm and has a higher dislocation density. The layer just beneath the interfacial layer shows subsurface deformation and elongated grains. The poorly- defined irregular boundaries indicate non-equilibrium high-energy configuration. Sub-surface deformation is illustrated in Fig. Dislocations, present as networks inside the deformed elongated grains, have been observed in the deformed substrate. Shearing deformation took place in the substrate as a response to the sliding process.

The creation of nano-structures is confirmed by the STM topography, indicating grains of between 5 and 10 nm Fig. It has been indicated by various authors [3,19,20,] that, in many systems, surfaces with ultra-fine structure are generated during high temperature sliding wear. Mechanical mixing involving repeated welding, fracture and re-welding of the debris generated from both contacting surfaces is responsible for the generation of the ultra- fine structured surfaces.

Moreover, the detailed TEM studies presented here has enabled understanding of the formation mechanisms of wear resistant nano-structured surfaces. These processes are aided by high temperature oxidation and diffusion. The positron annihilation studies confirmed the presence of vacancy clusters consisting of five vacancies [].

The next stage in the process involves deformation of oxides and generation of dislocations, leading to the formation of sub-grains. High internal stress is created inside the grains; the dislocation density and arrangement depend on the grain size, with smaller grains containing fewer dislocations.

The process leads to the formation of high energy grain boundaries with a high defect density []. Several authors have also constructed wear maps in an attempt to present wear data in a more easily understood format, allowing prediction of likely wear mode under specified sliding conditions. Lim [,] Fig. Adachi et al. The following section discusses recent developments in high temperature wear maps made by the present authors, following a brief review of earlier work by Lim.

This showed that the selection of sliding conditions and configuration can greatly affect the wear behaviour and transitions observed, with load, sliding speed and temperature potentially having a large influence on the boundaries between the different modes of wear. For example, it can be seen that, if a relatively high fixed load is used, with increasing sliding speed, a transition from severe-to-mild wear is observed Welsh [26].

More complex forms result at lower loads when sliding speeds are at much higher values Archard and Hirst [31], and Welsh [26]. To summarise, when two surfaces are worn against each other, the outcome can be a variety of apparently contradictory results. Thus, mapping is necessary to understand the relative behaviours in different tests and the potential outcome under a given set of sliding conditions.

The weight change data are presented in Figs. For both systems [22], mild wear with low weight loss dominates at 0. At all temperatures, there was virtually no initial severe wear period, with sufficient Co-Cr oxide debris forming extremely rapidly. The Nimonic 80A versus Stellite 6 system [21] is characterised by three distinct wear regimes at 0. The sliding behaviour at 0. The behaviour for the Incoloy MA versus Stellite 6 system follows the same general pattern at 0. However, in this case, the high temperature mild wear regime confers protection, with the mixed Fe-Cr and Co-Cr oxides at 0.

Such sliding studies indicate the potential for complex behaviour during sliding of dissimilar materials. In such cases, the necessity for mapping wear behaviour to assist prediction of mild or severe wear is important, if potentially catastrophic material failure is to be avoided. At the outset a review of some of the well-known and relevant wear theories and models, supported by experimental findings on conventional and advanced materials, has been presented. This background information has provided a framework to discuss new areas of high temperature wear. In this context the high temperature wear behaviour of those materials which have provided new information has been considered.

Particular attention has been focussed on high temperature wear behaviour of Oxide Dispersion Strengthened and Nimonic alloys, and intermetallic materials involving like-on-like and unlike-on-unlike combinations. The most significant part of this chapter includes the exposition of the phenomena of glaze formation at fundamental levels. The second solid may be in the form of a second body the opposing sliding surface or third body wear debris.

Shreir's Corrosion, Four-Volume Set

The hard particles or surface must be 1. Such forces may include chemical bonding chemical adhesion , inter-solubility diffusive adhesion , Van-der-Waals forces dispersive adhesion and electrostatic forces electrostatic adhesion. Surface interlocking may also occur mechanical adhesion via material filling surface voids or pores. Adhesion is of greater influence during contact of clean metallic surfaces and thus during severe wear , as there are no contaminants to prevent this contact. Adhesion is also more effective in a vacuum, where there is no surrounding atmosphere to affect it.

Adhesion is normally used to describe attraction between dissimilar molecules and cohesion between like molecules. The higher the value, the greater the ability to resist movement. Beyond this, the oxide becomes unstable, generating a wear particle. This assists the propagation of sub-surface cracks, which link together and allow surface material to break away. Such a process leads to the generation of the large, flat, angular debris seen during severe wear. HIPped hot A method of alloy or material production by powder metallurgy isostatically methods, using a combination of temperature and isostatic pressed pressure to produce the final item.

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Shreir's Corrosion: Four-Volume Set

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