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High temperature sliding abrasion of a nickel-base alloy and composite
Wear 225–229 Ž1999. 154–162
High temperature sliding abrasion of a nickel-base alloy and composite Hans Berns ) , Stefan Koch Lehrstuhl Werkstofftechnik, Ruhr-UniÕersitat ¨ Bochum, D-44780 Bochum, Germany
Abstract In previous work, sliding abrasion by interfacial flint particles was studied for a nickel-base MMC under argon up to 9008C. In the present investigation flint is replaced by corundum and silicon carbide. The hot hardness of the latter is above that of the NiCrAlSi-matrix and the WCrW2 C particles, which dramatically lowers the wear resistance of the MMC compared to abrasion by flint. Corundum is softer than tungsten carbide and the drop of wear resistance is less severe. In air, WCrW2 C is readily oxidized, lowering the wear resistance to matrix level. In comparison, Cr3 C 2 shows lower hot hardness but much better resistance to oxidation. A self-protecting particle layer of material debris and abrasive particles is formed under argon as well as in air, but at 8008C the oxidation of the debris in air entails a highly dense ‘glaze’ layer. q 1999 Published by Elsevier Science S.A. All rights reserved. Keywords: High temperature; Sliding abrasion; Nickel-base alloy; Composite; Oxidation; Protective layer
1. Introduction In high temperature processing of ores the crushing or compacting tools suffer from oxidation and from sliding abrasion by metal oxides as well as by accompanying minerals such as CaO, SiO 2 and Al 2 O 3 . The refining step of solid state reduction is usually done in a protective atmosphere, which suppresses the chemical attack on the tools. This type of wear was previously studied in ring-on-disc experiments with interfacial flint particles under argon at temperatures up to 10508C w1x. A self-protecting zone of abrasive particles was found in and on the work-hardened surface of the metallic specimens, which lowered the wear rate as the temperature was raised to the beginning of recrystallisation, above which the wear rate increased rapidly. Consequently the wear rate minimum was shifted to a higher critical temperature TC by changing from a bcc ferritic steel to an fcc austenitic steel or nickel alloy to an fcc cobalt alloy of high stacking fault energy. At room temperature abrasion is impeded by hard particles such as carbides and borides embedded in a metal matrix. To evaluate their effect at elevated temperature,
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numerous matrices and hard particles were investigated by microindentation and microscratching up to 9008C w2x. A nickel alloy, solid solution strengthened with Cr and Si and precipitation hardened by gX-Ni 3 Al, was developed for service at temperatures up to 7508C. The addition of up to 50 vro of crushed eutectic carbides WCrW2 C offered a superior hot hardness and lowered the wear rate of these powder metallurgically ŽPM. produced metal matrix composites ŽMMC. against flint in an argon atmosphere w3x. For high temperature service in air the NiCrAlSi-alloy promises a high resistance to oxidation. However, it is not clear up to which temperature the tungsten carbide will remain sufficiently stable. Of course hard oxide particles such as Al 2 O 3 or Y2 O 3 would not be affected by oxygen, but their bond to the metal matrix is inferior to that of metal carbides and borides. One can expect that Cr3 C 2 or CrB 2 hard particles in the NiCrAlSi—matrix would chemically degrade less than WCrW2 C because a protective oxide film is formed. Both Cr-based hard particles are available in crushed form, which is a prerequisite for a high resistance to abrading particles. In contrast, TiC is agglomerated from sub-micron to the required size and therefore lacks inner strength. In addition, TiC is less resistant to oxidation than Cr3 C 2 w4x. Boron tends to react with Al of the MM, thus diminishing precipitation hardening by gX-Ni 3 Al w5x. Therefore, CrB 2 is excluded. The first aim of this study was to investigate the change in high temperature wear behaviour of MMC caused by
0043-1648r99r$ - see front matter q 1999 Published by Elsevier Science S.A. All rights reserved. PII: S 0 0 4 3 - 1 6 4 8 Ž 9 9 . 0 0 0 0 8 - 3
H. Berns, S. Koch r Wear 225–229 (1999) 154–162
the access of air. The second aim was to evaluate the effect of abrasive particles harder than flint. The third aim was to recommend an MMC for service up to 7508C in air. 2. Experimental procedure In preparing the PM-materials, the following powders were used: Ža. an atomized nickel alloy containing Žwro. 18.9 Cr, 4.1 Al, 3.3 Si, 0.06 C with a powder grain size of - 100 mm to form the metal matrix ŽMM., Žb. crushed hard particles ŽHP. of Cr3 C 2 and eutectic WCrW2 C with a mesh size of 35 to 75 mm to reinforce the MM. Cans were filled with pure MM powder or mixtures of MM with 30 vro HP ŽMMC., evacuated, sealed and hot isostatically pressed ŽHIP. for 3 h at 11008C ŽWCrW2 C. or 10508C ŽCr3 C 2 . under a pressure of 180 MPa to yield full density. After solution annealing at 11008C for 2 h and precipitation annealing at 7508C for 8 h, the following hardness values ŽHV 30. were measured at room temperature: MM s 265, MM q Cr3 C 2 s 580, MM q WCrW2 C s 568. The specimens were prepared by spark erosive wire cutting and subsequent grinding. An additional diamond polish was applied to the samples for static oxidation experiments. The mass gain D m of specimens 40 by 20 by 5 mm in size, hanging in a sealed vertical tube furnace, was measured by a thermal balance. After heating in an argon atmosphere the specimens were exposed to dry air with - 50 mgrm3 of water for a duration of 240 min at a constant temperature up to 11008C. The microhardness ŽHV 0.05. of microstructural constituents MM and HP as well as of the abrasive particles ŽAP. was measured up to 9008C in vacuum using a tester outlined previously w2,6x. Upon increasing the load up to 2 N, Palmqvist cracks were initiated in the HP at room temperature and the fracture toughness K Ic was evaluated according to Ref. w7x. Hardness and toughness were averaged from 5 indentations. Microscratching experiments were carried out in a testrig described in Refs. w2,8x, employing a cubic boron nitride indenter with a wedge angle of 1158 and a front angle of 908 under a normal force of 0.3 N at a speed of 2 mmrs. Measuring the groove width and the tangential force in the MM and in the HP allowed to derive the tangential stress which is equal to the specific scratch energy es . The specimens were scratched at room temperature in air, heated up in vacuum of 2.5 = 10y4 bar to 6008C, were either not oxidized, or else oxidized in air of 2.5 = 10 3 bar for 60 min, and scratched at 600, 700 and 208C, under a pressure of 2.5 = 10y4 bar. The tangential force and the groove width were taken at 9 different locations to give an average es . A higher pressure or temperature led to excessive heating and distortion of some parts in the rig which was built to serve under high vacuum. Sliding abrasion was registered in a ring-on-disc device with interfacial AP at a rotational speed of 28 mmrs,
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under a normal pressure of 0.82 MPa and at a temperature up to 9008C in argon or air. Powders of crushed flint ŽSiO 2 ., corundum ŽAl 2 O 3 . and silicon carbide ŽSiC. of 63 to 100 mm mesh size were used as AP. From the mass losses of ring and disc, D m r and D m d , made of the same material, the length of the wear path Ž L s 50 m. and the contact area of ring and disc Ž A r s 0.88 A d . a dimensionless wear rate W was derived. Ws
D mr Ar
q
D md
1
Ad
rL
Ž 1.
The reciprocal value was defined as wear resistance Wy1 which is an average of three tests w1x. The AP flowing gradually from inside the ring through the interface and two radial slots to the outside were collected for a subsequent sieve analysis. The microstructure of the bulk and changes in the subsurface zone were studied by scanning electron microscopy ŽSEM., in combination with an energy dispersive X-ray ŽEDX. analyser and by light microscopy ŽLM.. The surfaces of oxidized, scratched or worn specimens were examined as well. X-ray diffraction was employed for phase analysis using a tube with a copper anode. The recorded intensities were compared with ASTM data.
3. Results 3.1. Microstructure The microstuctures of both MMC are shown in Fig. 1. Crushed angular HP are dispersed in the MM. An effective dispersion is an important prerequisite for a good fracture toughness of an MMC because crack extension along interconnected brittle HP is subdued. The properties of the dispersion depend on the size ratio and on the volume ratio of the two powders involved. At a mean size ratio close to one and a volume ratio of 3r7, the dispersion is in accordance with Ref. w3x. At higher magnification, the MM reveals some M 23 C 6 carbides of F 2 mm in size distributed along grain boundaries. Their amount is about equal in the MM and MMC materials which is evidence of minimal carbon diffusion from the HP into the MM during HIP or heat treatment. The WCrW2 C particles exhibit a feathery eutectic structure while the Cr3 C 2 have a single phase appearance. The HPrMM interface has broadened into a zone of interdiffusion and new phases which indicates that a good bond has been formed. Around WCrW2 C particles a zone of M 6 C and M 12 C was detected w3x and around Cr3 C 2 was a zone of M 7 C 3 w9x. In Fig. 1, the zone thickness seems to change from particle to particle, but this is deceptive because the angle between the HPrMM interface and the metallographic plane changes locally. The lowest apparent zone thickness is assumed to be the real one: 2 mm for WCrW2 C and 4.2 mm for Cr3 C 2 .
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7008C. As a consequence, the hardness ratio H MM rHAP becomes larger than one in the case of flint above 5008C, but remains below one for the other abrasives. At the target temperature of 7508C, the hardness ratio H HP rHAP is above unity except for SiC. Either hardness ratio exceeding unity means that scratching of the respective material constituent is impeded or impossible. The fracture toughness derived from Palmqvist cracks at room temperature increases in the following order ŽMPa 6m.: flint s 0.7; SiC s 2.1; Al 2 O 3 s 3.2; Cr3 C 2 s 3.5; WCrW2 C s 5.0. The toughness ratio of HP to AP is above unity at least at room temperature. As the temperature is raised, cracks cease to form w2x so that no ratio is measurable at 7508C. However, it seems reasonable to start with a room temperature ratio above one to promote crack extension in the AP and not in the HP if both come into repeated abrading contact. 3.3. Oxidation
Fig. 1. Microstructure of MMC consisting of a nickel alloy and dispersed hard particles Ža. eutectic WCrW2 C, Žb. Cr3 C 2 .
3.2. Microindentation In Fig. 2 the microhardness of MMC constituents and abrasives is depicted as a function of the testing temperature. While the carbides and oxides reveal a considerable loss of hardness as the temperature increases, the precipitation hardened MM retains an almost constant level to
Fig. 2. Hot hardness of MMC constituents and abrasives used.
The mass gain during static oxidation in air is given in Fig. 3. For the MM q WCrW2 C material, a logarithmic plot of the data does not yield a straight line. Thus, the thickness s of the oxide layer does not grow along the common law found for homogeneous alloys s s c P t q with q F 1 depending on the type and denseness of the oxide layer. The mass gain per area D mrA is proportional to s via the density r of the material, i.e., D mrA r s s. As to the temperature dependence of the oxidation from 700 to 11008C, the increase of the above exponent q at the beginning is far above that at the end of the test duration. Comparing these results with the almost unmeasurably low mass gain of the MM leads to the conclusion that WCrW2 C is rapidly oxidized. This is supported by a micrograph perpendicular to the oxidized surface showing a WCrW2 C particle transforming into a voluminous oxide ŽFig. 4.. According to Ref. w4x the specific volume of the oxide is
Fig. 3. Mass gain of MMC and metal matrix MM during static oxidation in air.
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Fig. 4. Metallographic section perpendicular to the surface of MMq WCrW2 C oxidized for 5 h at 7008C in air. On the left side an oxidized WCrW2 C particle, on the right side the lightly oxidized metal matrix.
about 2.5 times that of the carbide. If all particles are oxidized the rate of oxidation approaches that of the MM. Quite in contrast to WCrW2 C the Cr3 C 2 particles in the respective MMC hardly augment the mass gain compared to the MM. 3.4. Specific scratch energy In the previous section it was demonstrated that WCrW2 C deteriorated into a porous oxide at temperatures G 7008C. This led to the question up to which temperature these HP retain their high resistance to abrasion. As the oxide layer of pre-oxidized specimens tended to spall upon heating, samples were therefore oxidized in the scratching device at 6008C and subsequently scratched at 600, 700 and 208C. In this range of temperature the specific scratch energy es of the non-oxidized MM stayed rather constant between 6 to 8 Jrmm3 and was slightly increased by the thin oxide layer. In contrast, WCrW2 C revealed a drop of es from 25 Jrmm3 at 6008C Žunoxidized. to 16.1 Jrmm3 Žoxidized. and 12 Jrmm3 at 7008C Žoxidized.. These results prove that the MM is hardly effected by oxidation up to 7008C, while WCrW2 C is markedly harmed above 6008C. Apparently, the oxidation proceeds if the oxide layer formed at 6008C is heated to 7008C. The layer thickness is in the mm range and therefore hard to quantify in metallographic sections. However, in a top view, the oxidized layer on the WCrW2 C spalls alongside the scratch, while it adheres to the MM ŽFig. 5..
Fig. 5. MMqWCrW2 C scratched after oxidation at 6008C, Ža. metal matrix, Žb. WCrW2 C.
the hardness ratio H MM rHAP is 2.86 for flint, 0.31 for corundum and 0.14 for SiC. Except for flint, the high temperature wear resistance therefore strongly depends on the HP. The hardness ratio H HP rHAP is depicted in Fig. 6.
3.5. Wear resistance 3.5.1. Effect of abrasiÕes In Fig. 6 the wear resistance Wy1 of MM q WCrW2 C in argon is plotted to 9008C using flint, corundum or silicon carbide as abrasive particles. Wy1 decreases as the hardness of the AP increases Žsee Fig. 2.. At, e.g., 8008C
Fig. 6. Effect of abrasive and testing temperature on the resistance to sliding abrasion Wy1 of MM and MMqWCrW2 C in argon, H HP and HAP are the hardness values ŽHV 0.05. of hard particles and abrasive particles.
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It is evident that a high ratio entails a high wear resistance. Against flint and corundum a maximum wear resistance is found at a critical temperature TC , while the wear resistance against SiC is almost independent of temperature. At the target temperature of 7508C, flint is softer than the MM and the HP, corundum is harder than the MM but softer than the HP and SiC is harder than both constituents of the MMC. After the wear test, the MM surface between the HP is covered by a thin layer of agglomerated debris of MM and AP. In addition, fine AP are pressed into the deformed subsurface zone similar to mechanical alloying w10x. After cleaning in an ultrasonic bath, EDX analyses were taken of the MM surface. In case of flint and SiC, the Si content was determined and in case of corundum the Al content. In spite of a large scatter, the mean values of Si or Al increase from 30 " 3 wro at room temperature to 45 " 3 wro at 8008C and drop again at 9008C almost to the initial level. Apparently the layer is adhering best at 8008C and may exert a self-protecting effect on the MM up to this temperature, in agreement with earlier findings w1x. In the case of flint, this self-protection plays a part below 6008C because flint is harder than the MM. In the case of corundum protection extends to 8008C, and the drop of Wy1 at 9008C is probably connected to a breakdown of the self-protecting layer. Despite its low H MM rHAP ratio not even SiC leads to a continuous decrease of Wy1 with temperature as one might expect. Again self-protection of the MM is involved. The AP particle size within the surface agglomerate is much finer than the initial mesh size of 100 to ) 63 mm. Therefore grinding of the AP by the wear process seems to be a prerequisite to form a protective layer. Sieve analyses gave less than 10 wro in the class of 63 to ) 36 mm and
Fig. 8. Resistance Wy1 of metal matrix MM and MMqWCrW2 C to sliding abrasion against flint and silicon carbide in argon or air as a function of testing temperature.
less of 3 wro in the class of F 36 mm after the wear test. The resistance to fracturing of AP is best characterized by the weight ratio R of the classes 100 to ) 80 mm and 80 to ) 63 mm. Fig. 7 points out that at room temperature all three AP are brittle and fracture considerably in the wear interface during sliding abrasion. As the temperature of the wear test is raised flint stays brittle while corundum becomes more resistant to fracturing at 6008C and silicon carbide at 8008C. At 9008C all AP are again equally crushed. 3.5.2. Access of air In Fig. 8, the wear resistance of MM and MM q WCrW2 C up to 9008C in air is compared to that in argon using flint and SiC as abrasive particles. At the target temperature of 7508C, in argon there is little difference in Wy1 for MM and MMC due to a well adhering self-protecting layer of flint Žsee Ref. w3x.. The access of air lowers the wear resistance against flint above 6008C, and MM and MMC are again on the same level. One reason is the rapid oxidation of WCrW2 C rendering the HP ineffective, the other is again a self-protecting layer. Its Si content is lower than in argon and amounts to 30 " 5 wro between 600 and 9008C. Instead metal oxides of type NiO, NiWO4 , WO 3 , CrWO4 and Cr2 O 3 are found by X-ray diffraction on the MMC. Thus, the layer is built up of ground flint Table 1 Abrasive wear resistance of MMC against different abrasives at elevated temperatures in air Material
Abrasive
Abrasive wear resistance at w8Cx 700
Fig. 7. Resistance of the abrasives to fracturing during sliding abrasion given by the ratio R of the sieve fractions Ž100 to )80.rŽ80 to )63. in mm.
1 MMqWCrW2 C 2 MMqCr3 C 2 3 MMqCr3 C 2 4 MMqWCrW2 C
flint corundum SiC SiC
800 6
4.11=10 3.25=10 6 0.62=10 6 0.61=10 6
7.03=10 6 2.36=10 6 1.14=10 6 1.10=10 6
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above 6008C, flint is softer than the MM and the HP. The surface layer formed is not really contributing to the wear resistance but amounts to a mass loss by oxidation. In contrast, silicon carbide is harder than both the material constituents and the products of MMC oxidation, apparently enhancing the self-protecting effect of the layer. The wear rate by SiC is an order of magnitude higher than that by flint. This entails a much larger material surface exposed to oxidation per unit time and leads to more protection or ‘lubrication’ of the wear surface. First results of wear tests with MM q Cr3 C 2 are given in Table 1. They reveal that its wear resistance against SiC is equal to that of MM q WCrW2 C. Against corundum the wear resistance of MM q Cr3 C 2 at 7008C comes close to that of the MM q WCrW2 C against the softer flint. The wear specimens were subsequently inspected microscopically as to their wear surface and subsurface zone. The apparent density of the layer increased with temperature. An example is given in Fig. 9: At 7008C flint forms a rough, compact particle layer on the MM, while at 8008C
Fig. 9. Wear surface after sliding abrasion against flint in air, Ža. metal matrix at 7008C covered by a particle layer, Žb. flint particle pressed into the metal matrix covered by a glaze layer, Žc. oxidizing WCrW2 C pressing through the glaze layer.
particles and oxidized constituents of the material. The grinding of flint by MMC is almost identical to that in argon, but it is promoted by MMC compared to MM, as is demonstrated at 6008C at which temperature the oxidation of WCrW2 C is still slow ŽFig. 7.. Compared with flint, the wear resistance against silicon carbide is much lower and the presence of air has the opposite effect. Instead of lowering Wy1 an increase occurs in air ŽFig. 8.. In the temperature range of interest
Fig. 10. Metallographic section perpendicular to the wear surface after sliding abrasion against flint at 7008C in air showing the deformed metal matrix covered by a layer of wear debris and ground abrasives, EDX scan of Ni, Cr and Si along the line drawn in SEM picture.
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the layer is smooth like a glaze. The oxidation of a WCrW2 C particle penetrates the layer. At 7008C wear grooves are seen which flatten the rough particle porous layer. At 8008C no grooves are visible. The subsurface zone is deformed in the wear direction and covered by a layer of varying thickness which is enriched by Cr, Ni and Si ŽFig. 10..
4. Discussion 4.1. Material The selection of an MMC with 30 vro of HP was based on previous work and is a compromise between good wear resistance and toughness. A higher HP content enhances the wear resistance below and above TC w3x. However, to retain a uniform dispersion of HP the mean MM powder size has to be reduced to fill the decreasing space between the HP. As fine atomized powders are hardly available, a smaller fraction is gained by sieving which increases the cost if the larger fraction cannot be used elsewhere. As shown in Fig. 1, a dispersion of HP can be achieved and an embrittling network of interconnected HP is avoided, which would have a detrimental effect on the fracture toughness of the MMC w11x. Finite element calculations have demonstrated that the angular shape of crushed HP locally raises the stress in the MM w12,13x which is detrimental to the fracture toughness of martensitic MMC at room temperature, but less harmful in the present MMC at elevated temperature. The diffusion zone around the HP consists of softer carbides with a higher metal to carbon ratio. Although the overall carbide volume increases, the harder core of the HP diminishes. It is therefore desirable to keep the zone thickness small, which can be achieved by lowering the temperature during manufacturing. In MM q WCrW2 C and MM q Cr3 C 2 the growth of the zone thickness is controlled by the diffusion of W and Cr, respectively. The smaller Cr atoms were expected to diffuse faster and therefore the HIP temperature was chosen as low as 10508C for the MMC containing Cr3 C 2 . The high pressure still guaranteed full density. But during solution annealing at 11008C the zone became thicker. This treatment, however, is appropriate for a nickel-base alloy with 4 wro Al. Decreasing the solution temperature would impair the precipitation hardening. Cr3 C 2 seems to be very stable during service in air but only sufficiently stable during manufacturing. To continue the studies, coated Cr3 C 2 will be considered.
holds true for flint above 5008C. Against corundum, the wear resistance is improved by HP of superior hot hardness and a surface layer of agglomerated debris, which protects the softer MM up to 8008C. Silicon carbide is superior in hot hardness to both material constituents, which is a situation rarely met in application ŽFig. 11.. This picture changes if air is admitted. Static oxidation of WCrW2 C becomes very rapid as the temperature is raised above 7008C. In sliding abrasion, the oxidative intervals between material removal are much shorter than those applied in the scratching experiments. Yet these experiments give a slow-motion description of what happens during sliding abrasion, i.e., a degradation of the WCrW2 C by oxidation starting already at about 6008C. As a conclusion, this type of hard particle does not match with the nickel-base MM which is quite resistant to high temperature oxidation. In the temperature range of 600 to 9008C, the addition of WCrW2 C is justified only under a protective atmosphere. In air, Cr3 C 2 is of similar oxidation resistance as the MM and therefore a better choice. Above 5008C Cr3 C 2 is harder than corundum which is among the hardest minerals in the applications. In Table 1 first results on the wear resistance of MM q Cr3 C 2 are compared with other MMCrAP combinations in the range of the target temperature. Compared to the result of the MM q WCrW2 C at 7008C against flint the performance of the MM q Cr3 C 2 against the much harder corundum looks promising. In view of all the data collected to now, an MMC consisting of 30 vro Cr3 C 2 particles dispersed in a precipitation hardened NiCrAlSi-matrix appears to be a suitable material against sliding abrasion up to 750 and even 8008C in air. 4.3. ProtectiÕe layer A self-protecting layer reducing wear is formed on the MM during sliding abrasion. It consists of wear debris and
4.2. Wear resistance In argon, the AP interact with the MM and with the HP. The highest wear resistance is reached, if the hardness ratios H MM rHAP and H HPrHAP are above unity, which
Fig. 11. Effect of the hot hardness of hard particles H HP , metal matrix H MM and abrasive particles HAP on the wear resistance Wy1 of MMq WCrW2 C and MM in argon.
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ground abrasive particles. The consistency of the layer depends on the temperature and the atmosphere. The present study has shown that not only flint w1,3x contributes to the layer but also the harder abrasives corundum and silicon carbide. The smaller their size, the more points of contact emerge which exert an attraction due to their surface energy w14x. As pointed out in Fig. 7 the resistance to grinding during sliding abrasion depends on the type of abrasive and the temperature. The radial slots in the ring specimen Žcompare insert of Fig. 6. were necessary to introduce abrasive particles to the interface but also allowed part of the particles to flow unharmed from the reservoir inside the ring to the outside. In addition, the small fraction of very fine particles found in the layer could not be evaluated. Therefore, the results of Fig. 7 are of qualitative character. They reveal, though, that the grinding process is enhanced in MMC compared to MM and indirectly affected by the surrounding atmosphere. In argon grinding was more pronounced between the wear surfaces of MM compared to MMC w15x. It is not clear if the AP fracture by shear or compression and how this is influenced by the atmosphere. The oxidation of WCrW2 C during sliding abrasion at 6008C in air is still moderate and so these HP support crushing. However, at higher temperatures they are readily oxidized and become ineffective. As the temperature of sliding abrasion is raised, the metallic wear debris in the particle layer is joined by sintering under argon which increases their coherence. In air, the material debris is interconnected by oxidation, which augments their volume at the expense of AP in the layer. This is revealed, e.g., by a lower Si-content measured by EDX on the wear surface as compared to an argon atmosphere. At 8008C in air, the agglomerated particle layer changed into a ‘glaze’ of metal oxides and incorporated fine AP covering the MM which is for example shown in Fig. 9b, c. A similar change from a particle layer in argon to a glaze layer in the oxidizing atmosphere was found in two-body sliding wear of the nickel alloy Nimonic 80A in pure oxygen already at 6008C w16x. The third body, i.e., the AP, is apparently not needed to glaze but is integrated if present such as in three-body sliding abrasion. Fine AP are worked into the deformed wear-surface zone of the MM under argon. This is the more pronounced the softer the MM. In case of a softer austenitic steel, a banded structure of AP and MM occurs in this zone, which is similar to early stages of mechanical alloying w15x. In air, no such zone was observed for the present nickel-base MM.
5. Conclusions High temperature sliding abrasion of a nickel-base alloy and composite was investigated in argon and air, which has led to the following conclusions.
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Ža. The powder metallurgical parameters chosen resulted in a beneficial dispersion of 30 vro of WCrW2 C or Cr3 C 2 particles in a NiCrAlSi-matrix subsequently precipitation hardened. Žb. Above 5008C in argon, flint was found to be softer than the metal matrix and the WCrW2 C. Thus the wear resistance is high. Silicon carbide was harder than both material constituents up to 9008C. The wear resistance is low. Corundum was softer than the WCrW2 C but harder than the matrix, leading to an intermediate wear resistance. Žc. Above about 6008C, the oxidation of WCrW2 C increased rapidly with temperature. The porous oxides offered no resistance to abrasion in air. In comparison Cr3 C 2 revealed a lower hot hardness but a much higher resistance to oxidation. Žd. The hot hardness ratio of Cr3 C 2 and corundum is above unity. Together with its resistance to oxidation this carbide is a promising candidate for wear resistant nickelbase composites in applications up to 8008C where the abrasives are rarely harder than corundum. Že. A self-protecting particle layer consisting of material debris and small abrasive particles ground in the wear interface is formed on the metal matrix during sliding abrasion. At 8008C in air, the debris is oxidized to form a ‘glaze’ layer of high density.
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ŽEd.., Hard alloys and hard composites Žin German., Springer-Verlag, Berlin, 1998. w13x D.F. Watt, X.Q. Xu, D.J. Lloyd, Effects of particle morphology and spacing on the strain fields in a plastically deforming matrix, Acta Mater. 44 Ž2. Ž1996. 789–799. w14x K.L. Johnson, K. Kendall, A.D. Roberts, Surface energy and the contact of elastic solids, Proc. R. Soc. London A 324 Ž1971. 301–313.
w15x H. Berns, S. Koch, Influence of abrasives on wear mechanism and wear resistance in a sliding abrasion test at elevated temperatures, special issue of Wear, Intern. Conference on Erosive and Abrasive Wear, Cambridge, England, Sept. 1998. w16x J. Jang, F.H. Stott, M.M. Stack, The effect of partial pressure of oxygene on the tribological behaviour of a nickel-based alloy, N80A, at elevated temperatures, Wear 203–204 Ž1997. 615–625.
High temperature sliding abrasion of a nickel-base alloy and composite Hans Berns ) , Stefan Koch Lehrstuhl Werkstofftechnik, Ruhr-UniÕersitat ¨ Bochum, D-44780 Bochum, Germany
Abstract In previous work, sliding abrasion by interfacial flint particles was studied for a nickel-base MMC under argon up to 9008C. In the present investigation flint is replaced by corundum and silicon carbide. The hot hardness of the latter is above that of the NiCrAlSi-matrix and the WCrW2 C particles, which dramatically lowers the wear resistance of the MMC compared to abrasion by flint. Corundum is softer than tungsten carbide and the drop of wear resistance is less severe. In air, WCrW2 C is readily oxidized, lowering the wear resistance to matrix level. In comparison, Cr3 C 2 shows lower hot hardness but much better resistance to oxidation. A self-protecting particle layer of material debris and abrasive particles is formed under argon as well as in air, but at 8008C the oxidation of the debris in air entails a highly dense ‘glaze’ layer. q 1999 Published by Elsevier Science S.A. All rights reserved. Keywords: High temperature; Sliding abrasion; Nickel-base alloy; Composite; Oxidation; Protective layer
1. Introduction In high temperature processing of ores the crushing or compacting tools suffer from oxidation and from sliding abrasion by metal oxides as well as by accompanying minerals such as CaO, SiO 2 and Al 2 O 3 . The refining step of solid state reduction is usually done in a protective atmosphere, which suppresses the chemical attack on the tools. This type of wear was previously studied in ring-on-disc experiments with interfacial flint particles under argon at temperatures up to 10508C w1x. A self-protecting zone of abrasive particles was found in and on the work-hardened surface of the metallic specimens, which lowered the wear rate as the temperature was raised to the beginning of recrystallisation, above which the wear rate increased rapidly. Consequently the wear rate minimum was shifted to a higher critical temperature TC by changing from a bcc ferritic steel to an fcc austenitic steel or nickel alloy to an fcc cobalt alloy of high stacking fault energy. At room temperature abrasion is impeded by hard particles such as carbides and borides embedded in a metal matrix. To evaluate their effect at elevated temperature,
) Corresponding author. Tel.: q49-234-700-5955; fax: q49-234-7094104; e-mail: [email protected]
numerous matrices and hard particles were investigated by microindentation and microscratching up to 9008C w2x. A nickel alloy, solid solution strengthened with Cr and Si and precipitation hardened by gX-Ni 3 Al, was developed for service at temperatures up to 7508C. The addition of up to 50 vro of crushed eutectic carbides WCrW2 C offered a superior hot hardness and lowered the wear rate of these powder metallurgically ŽPM. produced metal matrix composites ŽMMC. against flint in an argon atmosphere w3x. For high temperature service in air the NiCrAlSi-alloy promises a high resistance to oxidation. However, it is not clear up to which temperature the tungsten carbide will remain sufficiently stable. Of course hard oxide particles such as Al 2 O 3 or Y2 O 3 would not be affected by oxygen, but their bond to the metal matrix is inferior to that of metal carbides and borides. One can expect that Cr3 C 2 or CrB 2 hard particles in the NiCrAlSi—matrix would chemically degrade less than WCrW2 C because a protective oxide film is formed. Both Cr-based hard particles are available in crushed form, which is a prerequisite for a high resistance to abrading particles. In contrast, TiC is agglomerated from sub-micron to the required size and therefore lacks inner strength. In addition, TiC is less resistant to oxidation than Cr3 C 2 w4x. Boron tends to react with Al of the MM, thus diminishing precipitation hardening by gX-Ni 3 Al w5x. Therefore, CrB 2 is excluded. The first aim of this study was to investigate the change in high temperature wear behaviour of MMC caused by
0043-1648r99r$ - see front matter q 1999 Published by Elsevier Science S.A. All rights reserved. PII: S 0 0 4 3 - 1 6 4 8 Ž 9 9 . 0 0 0 0 8 - 3
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the access of air. The second aim was to evaluate the effect of abrasive particles harder than flint. The third aim was to recommend an MMC for service up to 7508C in air. 2. Experimental procedure In preparing the PM-materials, the following powders were used: Ža. an atomized nickel alloy containing Žwro. 18.9 Cr, 4.1 Al, 3.3 Si, 0.06 C with a powder grain size of - 100 mm to form the metal matrix ŽMM., Žb. crushed hard particles ŽHP. of Cr3 C 2 and eutectic WCrW2 C with a mesh size of 35 to 75 mm to reinforce the MM. Cans were filled with pure MM powder or mixtures of MM with 30 vro HP ŽMMC., evacuated, sealed and hot isostatically pressed ŽHIP. for 3 h at 11008C ŽWCrW2 C. or 10508C ŽCr3 C 2 . under a pressure of 180 MPa to yield full density. After solution annealing at 11008C for 2 h and precipitation annealing at 7508C for 8 h, the following hardness values ŽHV 30. were measured at room temperature: MM s 265, MM q Cr3 C 2 s 580, MM q WCrW2 C s 568. The specimens were prepared by spark erosive wire cutting and subsequent grinding. An additional diamond polish was applied to the samples for static oxidation experiments. The mass gain D m of specimens 40 by 20 by 5 mm in size, hanging in a sealed vertical tube furnace, was measured by a thermal balance. After heating in an argon atmosphere the specimens were exposed to dry air with - 50 mgrm3 of water for a duration of 240 min at a constant temperature up to 11008C. The microhardness ŽHV 0.05. of microstructural constituents MM and HP as well as of the abrasive particles ŽAP. was measured up to 9008C in vacuum using a tester outlined previously w2,6x. Upon increasing the load up to 2 N, Palmqvist cracks were initiated in the HP at room temperature and the fracture toughness K Ic was evaluated according to Ref. w7x. Hardness and toughness were averaged from 5 indentations. Microscratching experiments were carried out in a testrig described in Refs. w2,8x, employing a cubic boron nitride indenter with a wedge angle of 1158 and a front angle of 908 under a normal force of 0.3 N at a speed of 2 mmrs. Measuring the groove width and the tangential force in the MM and in the HP allowed to derive the tangential stress which is equal to the specific scratch energy es . The specimens were scratched at room temperature in air, heated up in vacuum of 2.5 = 10y4 bar to 6008C, were either not oxidized, or else oxidized in air of 2.5 = 10 3 bar for 60 min, and scratched at 600, 700 and 208C, under a pressure of 2.5 = 10y4 bar. The tangential force and the groove width were taken at 9 different locations to give an average es . A higher pressure or temperature led to excessive heating and distortion of some parts in the rig which was built to serve under high vacuum. Sliding abrasion was registered in a ring-on-disc device with interfacial AP at a rotational speed of 28 mmrs,
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under a normal pressure of 0.82 MPa and at a temperature up to 9008C in argon or air. Powders of crushed flint ŽSiO 2 ., corundum ŽAl 2 O 3 . and silicon carbide ŽSiC. of 63 to 100 mm mesh size were used as AP. From the mass losses of ring and disc, D m r and D m d , made of the same material, the length of the wear path Ž L s 50 m. and the contact area of ring and disc Ž A r s 0.88 A d . a dimensionless wear rate W was derived. Ws
D mr Ar
q
D md
1
Ad
rL
Ž 1.
The reciprocal value was defined as wear resistance Wy1 which is an average of three tests w1x. The AP flowing gradually from inside the ring through the interface and two radial slots to the outside were collected for a subsequent sieve analysis. The microstructure of the bulk and changes in the subsurface zone were studied by scanning electron microscopy ŽSEM., in combination with an energy dispersive X-ray ŽEDX. analyser and by light microscopy ŽLM.. The surfaces of oxidized, scratched or worn specimens were examined as well. X-ray diffraction was employed for phase analysis using a tube with a copper anode. The recorded intensities were compared with ASTM data.
3. Results 3.1. Microstructure The microstuctures of both MMC are shown in Fig. 1. Crushed angular HP are dispersed in the MM. An effective dispersion is an important prerequisite for a good fracture toughness of an MMC because crack extension along interconnected brittle HP is subdued. The properties of the dispersion depend on the size ratio and on the volume ratio of the two powders involved. At a mean size ratio close to one and a volume ratio of 3r7, the dispersion is in accordance with Ref. w3x. At higher magnification, the MM reveals some M 23 C 6 carbides of F 2 mm in size distributed along grain boundaries. Their amount is about equal in the MM and MMC materials which is evidence of minimal carbon diffusion from the HP into the MM during HIP or heat treatment. The WCrW2 C particles exhibit a feathery eutectic structure while the Cr3 C 2 have a single phase appearance. The HPrMM interface has broadened into a zone of interdiffusion and new phases which indicates that a good bond has been formed. Around WCrW2 C particles a zone of M 6 C and M 12 C was detected w3x and around Cr3 C 2 was a zone of M 7 C 3 w9x. In Fig. 1, the zone thickness seems to change from particle to particle, but this is deceptive because the angle between the HPrMM interface and the metallographic plane changes locally. The lowest apparent zone thickness is assumed to be the real one: 2 mm for WCrW2 C and 4.2 mm for Cr3 C 2 .
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7008C. As a consequence, the hardness ratio H MM rHAP becomes larger than one in the case of flint above 5008C, but remains below one for the other abrasives. At the target temperature of 7508C, the hardness ratio H HP rHAP is above unity except for SiC. Either hardness ratio exceeding unity means that scratching of the respective material constituent is impeded or impossible. The fracture toughness derived from Palmqvist cracks at room temperature increases in the following order ŽMPa 6m.: flint s 0.7; SiC s 2.1; Al 2 O 3 s 3.2; Cr3 C 2 s 3.5; WCrW2 C s 5.0. The toughness ratio of HP to AP is above unity at least at room temperature. As the temperature is raised, cracks cease to form w2x so that no ratio is measurable at 7508C. However, it seems reasonable to start with a room temperature ratio above one to promote crack extension in the AP and not in the HP if both come into repeated abrading contact. 3.3. Oxidation
Fig. 1. Microstructure of MMC consisting of a nickel alloy and dispersed hard particles Ža. eutectic WCrW2 C, Žb. Cr3 C 2 .
3.2. Microindentation In Fig. 2 the microhardness of MMC constituents and abrasives is depicted as a function of the testing temperature. While the carbides and oxides reveal a considerable loss of hardness as the temperature increases, the precipitation hardened MM retains an almost constant level to
Fig. 2. Hot hardness of MMC constituents and abrasives used.
The mass gain during static oxidation in air is given in Fig. 3. For the MM q WCrW2 C material, a logarithmic plot of the data does not yield a straight line. Thus, the thickness s of the oxide layer does not grow along the common law found for homogeneous alloys s s c P t q with q F 1 depending on the type and denseness of the oxide layer. The mass gain per area D mrA is proportional to s via the density r of the material, i.e., D mrA r s s. As to the temperature dependence of the oxidation from 700 to 11008C, the increase of the above exponent q at the beginning is far above that at the end of the test duration. Comparing these results with the almost unmeasurably low mass gain of the MM leads to the conclusion that WCrW2 C is rapidly oxidized. This is supported by a micrograph perpendicular to the oxidized surface showing a WCrW2 C particle transforming into a voluminous oxide ŽFig. 4.. According to Ref. w4x the specific volume of the oxide is
Fig. 3. Mass gain of MMC and metal matrix MM during static oxidation in air.
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Fig. 4. Metallographic section perpendicular to the surface of MMq WCrW2 C oxidized for 5 h at 7008C in air. On the left side an oxidized WCrW2 C particle, on the right side the lightly oxidized metal matrix.
about 2.5 times that of the carbide. If all particles are oxidized the rate of oxidation approaches that of the MM. Quite in contrast to WCrW2 C the Cr3 C 2 particles in the respective MMC hardly augment the mass gain compared to the MM. 3.4. Specific scratch energy In the previous section it was demonstrated that WCrW2 C deteriorated into a porous oxide at temperatures G 7008C. This led to the question up to which temperature these HP retain their high resistance to abrasion. As the oxide layer of pre-oxidized specimens tended to spall upon heating, samples were therefore oxidized in the scratching device at 6008C and subsequently scratched at 600, 700 and 208C. In this range of temperature the specific scratch energy es of the non-oxidized MM stayed rather constant between 6 to 8 Jrmm3 and was slightly increased by the thin oxide layer. In contrast, WCrW2 C revealed a drop of es from 25 Jrmm3 at 6008C Žunoxidized. to 16.1 Jrmm3 Žoxidized. and 12 Jrmm3 at 7008C Žoxidized.. These results prove that the MM is hardly effected by oxidation up to 7008C, while WCrW2 C is markedly harmed above 6008C. Apparently, the oxidation proceeds if the oxide layer formed at 6008C is heated to 7008C. The layer thickness is in the mm range and therefore hard to quantify in metallographic sections. However, in a top view, the oxidized layer on the WCrW2 C spalls alongside the scratch, while it adheres to the MM ŽFig. 5..
Fig. 5. MMqWCrW2 C scratched after oxidation at 6008C, Ža. metal matrix, Žb. WCrW2 C.
the hardness ratio H MM rHAP is 2.86 for flint, 0.31 for corundum and 0.14 for SiC. Except for flint, the high temperature wear resistance therefore strongly depends on the HP. The hardness ratio H HP rHAP is depicted in Fig. 6.
3.5. Wear resistance 3.5.1. Effect of abrasiÕes In Fig. 6 the wear resistance Wy1 of MM q WCrW2 C in argon is plotted to 9008C using flint, corundum or silicon carbide as abrasive particles. Wy1 decreases as the hardness of the AP increases Žsee Fig. 2.. At, e.g., 8008C
Fig. 6. Effect of abrasive and testing temperature on the resistance to sliding abrasion Wy1 of MM and MMqWCrW2 C in argon, H HP and HAP are the hardness values ŽHV 0.05. of hard particles and abrasive particles.
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It is evident that a high ratio entails a high wear resistance. Against flint and corundum a maximum wear resistance is found at a critical temperature TC , while the wear resistance against SiC is almost independent of temperature. At the target temperature of 7508C, flint is softer than the MM and the HP, corundum is harder than the MM but softer than the HP and SiC is harder than both constituents of the MMC. After the wear test, the MM surface between the HP is covered by a thin layer of agglomerated debris of MM and AP. In addition, fine AP are pressed into the deformed subsurface zone similar to mechanical alloying w10x. After cleaning in an ultrasonic bath, EDX analyses were taken of the MM surface. In case of flint and SiC, the Si content was determined and in case of corundum the Al content. In spite of a large scatter, the mean values of Si or Al increase from 30 " 3 wro at room temperature to 45 " 3 wro at 8008C and drop again at 9008C almost to the initial level. Apparently the layer is adhering best at 8008C and may exert a self-protecting effect on the MM up to this temperature, in agreement with earlier findings w1x. In the case of flint, this self-protection plays a part below 6008C because flint is harder than the MM. In the case of corundum protection extends to 8008C, and the drop of Wy1 at 9008C is probably connected to a breakdown of the self-protecting layer. Despite its low H MM rHAP ratio not even SiC leads to a continuous decrease of Wy1 with temperature as one might expect. Again self-protection of the MM is involved. The AP particle size within the surface agglomerate is much finer than the initial mesh size of 100 to ) 63 mm. Therefore grinding of the AP by the wear process seems to be a prerequisite to form a protective layer. Sieve analyses gave less than 10 wro in the class of 63 to ) 36 mm and
Fig. 8. Resistance Wy1 of metal matrix MM and MMqWCrW2 C to sliding abrasion against flint and silicon carbide in argon or air as a function of testing temperature.
less of 3 wro in the class of F 36 mm after the wear test. The resistance to fracturing of AP is best characterized by the weight ratio R of the classes 100 to ) 80 mm and 80 to ) 63 mm. Fig. 7 points out that at room temperature all three AP are brittle and fracture considerably in the wear interface during sliding abrasion. As the temperature of the wear test is raised flint stays brittle while corundum becomes more resistant to fracturing at 6008C and silicon carbide at 8008C. At 9008C all AP are again equally crushed. 3.5.2. Access of air In Fig. 8, the wear resistance of MM and MM q WCrW2 C up to 9008C in air is compared to that in argon using flint and SiC as abrasive particles. At the target temperature of 7508C, in argon there is little difference in Wy1 for MM and MMC due to a well adhering self-protecting layer of flint Žsee Ref. w3x.. The access of air lowers the wear resistance against flint above 6008C, and MM and MMC are again on the same level. One reason is the rapid oxidation of WCrW2 C rendering the HP ineffective, the other is again a self-protecting layer. Its Si content is lower than in argon and amounts to 30 " 5 wro between 600 and 9008C. Instead metal oxides of type NiO, NiWO4 , WO 3 , CrWO4 and Cr2 O 3 are found by X-ray diffraction on the MMC. Thus, the layer is built up of ground flint Table 1 Abrasive wear resistance of MMC against different abrasives at elevated temperatures in air Material
Abrasive
Abrasive wear resistance at w8Cx 700
Fig. 7. Resistance of the abrasives to fracturing during sliding abrasion given by the ratio R of the sieve fractions Ž100 to )80.rŽ80 to )63. in mm.
1 MMqWCrW2 C 2 MMqCr3 C 2 3 MMqCr3 C 2 4 MMqWCrW2 C
flint corundum SiC SiC
800 6
4.11=10 3.25=10 6 0.62=10 6 0.61=10 6
7.03=10 6 2.36=10 6 1.14=10 6 1.10=10 6
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above 6008C, flint is softer than the MM and the HP. The surface layer formed is not really contributing to the wear resistance but amounts to a mass loss by oxidation. In contrast, silicon carbide is harder than both the material constituents and the products of MMC oxidation, apparently enhancing the self-protecting effect of the layer. The wear rate by SiC is an order of magnitude higher than that by flint. This entails a much larger material surface exposed to oxidation per unit time and leads to more protection or ‘lubrication’ of the wear surface. First results of wear tests with MM q Cr3 C 2 are given in Table 1. They reveal that its wear resistance against SiC is equal to that of MM q WCrW2 C. Against corundum the wear resistance of MM q Cr3 C 2 at 7008C comes close to that of the MM q WCrW2 C against the softer flint. The wear specimens were subsequently inspected microscopically as to their wear surface and subsurface zone. The apparent density of the layer increased with temperature. An example is given in Fig. 9: At 7008C flint forms a rough, compact particle layer on the MM, while at 8008C
Fig. 9. Wear surface after sliding abrasion against flint in air, Ža. metal matrix at 7008C covered by a particle layer, Žb. flint particle pressed into the metal matrix covered by a glaze layer, Žc. oxidizing WCrW2 C pressing through the glaze layer.
particles and oxidized constituents of the material. The grinding of flint by MMC is almost identical to that in argon, but it is promoted by MMC compared to MM, as is demonstrated at 6008C at which temperature the oxidation of WCrW2 C is still slow ŽFig. 7.. Compared with flint, the wear resistance against silicon carbide is much lower and the presence of air has the opposite effect. Instead of lowering Wy1 an increase occurs in air ŽFig. 8.. In the temperature range of interest
Fig. 10. Metallographic section perpendicular to the wear surface after sliding abrasion against flint at 7008C in air showing the deformed metal matrix covered by a layer of wear debris and ground abrasives, EDX scan of Ni, Cr and Si along the line drawn in SEM picture.
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the layer is smooth like a glaze. The oxidation of a WCrW2 C particle penetrates the layer. At 7008C wear grooves are seen which flatten the rough particle porous layer. At 8008C no grooves are visible. The subsurface zone is deformed in the wear direction and covered by a layer of varying thickness which is enriched by Cr, Ni and Si ŽFig. 10..
4. Discussion 4.1. Material The selection of an MMC with 30 vro of HP was based on previous work and is a compromise between good wear resistance and toughness. A higher HP content enhances the wear resistance below and above TC w3x. However, to retain a uniform dispersion of HP the mean MM powder size has to be reduced to fill the decreasing space between the HP. As fine atomized powders are hardly available, a smaller fraction is gained by sieving which increases the cost if the larger fraction cannot be used elsewhere. As shown in Fig. 1, a dispersion of HP can be achieved and an embrittling network of interconnected HP is avoided, which would have a detrimental effect on the fracture toughness of the MMC w11x. Finite element calculations have demonstrated that the angular shape of crushed HP locally raises the stress in the MM w12,13x which is detrimental to the fracture toughness of martensitic MMC at room temperature, but less harmful in the present MMC at elevated temperature. The diffusion zone around the HP consists of softer carbides with a higher metal to carbon ratio. Although the overall carbide volume increases, the harder core of the HP diminishes. It is therefore desirable to keep the zone thickness small, which can be achieved by lowering the temperature during manufacturing. In MM q WCrW2 C and MM q Cr3 C 2 the growth of the zone thickness is controlled by the diffusion of W and Cr, respectively. The smaller Cr atoms were expected to diffuse faster and therefore the HIP temperature was chosen as low as 10508C for the MMC containing Cr3 C 2 . The high pressure still guaranteed full density. But during solution annealing at 11008C the zone became thicker. This treatment, however, is appropriate for a nickel-base alloy with 4 wro Al. Decreasing the solution temperature would impair the precipitation hardening. Cr3 C 2 seems to be very stable during service in air but only sufficiently stable during manufacturing. To continue the studies, coated Cr3 C 2 will be considered.
holds true for flint above 5008C. Against corundum, the wear resistance is improved by HP of superior hot hardness and a surface layer of agglomerated debris, which protects the softer MM up to 8008C. Silicon carbide is superior in hot hardness to both material constituents, which is a situation rarely met in application ŽFig. 11.. This picture changes if air is admitted. Static oxidation of WCrW2 C becomes very rapid as the temperature is raised above 7008C. In sliding abrasion, the oxidative intervals between material removal are much shorter than those applied in the scratching experiments. Yet these experiments give a slow-motion description of what happens during sliding abrasion, i.e., a degradation of the WCrW2 C by oxidation starting already at about 6008C. As a conclusion, this type of hard particle does not match with the nickel-base MM which is quite resistant to high temperature oxidation. In the temperature range of 600 to 9008C, the addition of WCrW2 C is justified only under a protective atmosphere. In air, Cr3 C 2 is of similar oxidation resistance as the MM and therefore a better choice. Above 5008C Cr3 C 2 is harder than corundum which is among the hardest minerals in the applications. In Table 1 first results on the wear resistance of MM q Cr3 C 2 are compared with other MMCrAP combinations in the range of the target temperature. Compared to the result of the MM q WCrW2 C at 7008C against flint the performance of the MM q Cr3 C 2 against the much harder corundum looks promising. In view of all the data collected to now, an MMC consisting of 30 vro Cr3 C 2 particles dispersed in a precipitation hardened NiCrAlSi-matrix appears to be a suitable material against sliding abrasion up to 750 and even 8008C in air. 4.3. ProtectiÕe layer A self-protecting layer reducing wear is formed on the MM during sliding abrasion. It consists of wear debris and
4.2. Wear resistance In argon, the AP interact with the MM and with the HP. The highest wear resistance is reached, if the hardness ratios H MM rHAP and H HPrHAP are above unity, which
Fig. 11. Effect of the hot hardness of hard particles H HP , metal matrix H MM and abrasive particles HAP on the wear resistance Wy1 of MMq WCrW2 C and MM in argon.
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ground abrasive particles. The consistency of the layer depends on the temperature and the atmosphere. The present study has shown that not only flint w1,3x contributes to the layer but also the harder abrasives corundum and silicon carbide. The smaller their size, the more points of contact emerge which exert an attraction due to their surface energy w14x. As pointed out in Fig. 7 the resistance to grinding during sliding abrasion depends on the type of abrasive and the temperature. The radial slots in the ring specimen Žcompare insert of Fig. 6. were necessary to introduce abrasive particles to the interface but also allowed part of the particles to flow unharmed from the reservoir inside the ring to the outside. In addition, the small fraction of very fine particles found in the layer could not be evaluated. Therefore, the results of Fig. 7 are of qualitative character. They reveal, though, that the grinding process is enhanced in MMC compared to MM and indirectly affected by the surrounding atmosphere. In argon grinding was more pronounced between the wear surfaces of MM compared to MMC w15x. It is not clear if the AP fracture by shear or compression and how this is influenced by the atmosphere. The oxidation of WCrW2 C during sliding abrasion at 6008C in air is still moderate and so these HP support crushing. However, at higher temperatures they are readily oxidized and become ineffective. As the temperature of sliding abrasion is raised, the metallic wear debris in the particle layer is joined by sintering under argon which increases their coherence. In air, the material debris is interconnected by oxidation, which augments their volume at the expense of AP in the layer. This is revealed, e.g., by a lower Si-content measured by EDX on the wear surface as compared to an argon atmosphere. At 8008C in air, the agglomerated particle layer changed into a ‘glaze’ of metal oxides and incorporated fine AP covering the MM which is for example shown in Fig. 9b, c. A similar change from a particle layer in argon to a glaze layer in the oxidizing atmosphere was found in two-body sliding wear of the nickel alloy Nimonic 80A in pure oxygen already at 6008C w16x. The third body, i.e., the AP, is apparently not needed to glaze but is integrated if present such as in three-body sliding abrasion. Fine AP are worked into the deformed wear-surface zone of the MM under argon. This is the more pronounced the softer the MM. In case of a softer austenitic steel, a banded structure of AP and MM occurs in this zone, which is similar to early stages of mechanical alloying w15x. In air, no such zone was observed for the present nickel-base MM.
5. Conclusions High temperature sliding abrasion of a nickel-base alloy and composite was investigated in argon and air, which has led to the following conclusions.
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Ža. The powder metallurgical parameters chosen resulted in a beneficial dispersion of 30 vro of WCrW2 C or Cr3 C 2 particles in a NiCrAlSi-matrix subsequently precipitation hardened. Žb. Above 5008C in argon, flint was found to be softer than the metal matrix and the WCrW2 C. Thus the wear resistance is high. Silicon carbide was harder than both material constituents up to 9008C. The wear resistance is low. Corundum was softer than the WCrW2 C but harder than the matrix, leading to an intermediate wear resistance. Žc. Above about 6008C, the oxidation of WCrW2 C increased rapidly with temperature. The porous oxides offered no resistance to abrasion in air. In comparison Cr3 C 2 revealed a lower hot hardness but a much higher resistance to oxidation. Žd. The hot hardness ratio of Cr3 C 2 and corundum is above unity. Together with its resistance to oxidation this carbide is a promising candidate for wear resistant nickelbase composites in applications up to 8008C where the abrasives are rarely harder than corundum. Že. A self-protecting particle layer consisting of material debris and small abrasive particles ground in the wear interface is formed on the metal matrix during sliding abrasion. At 8008C in air, the debris is oxidized to form a ‘glaze’ layer of high density.
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