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A transmission electron microscopy study of wear of magnesia partially stabilised zirconia
A transmission electron microscopy study of wear of magnesia partially stabilised zirconia W. M. Rainforth deponent of Engineering ~ufe~~, Shefield SI 43lJ (UK)
Un~ve~~~ of She~eId, PO Box 600, Sir Robert ~ad~eid Building, copper
Street,
R. Stevens Schoof
ofpatents,
V~iversi~
of Leeds, Leeds LS2 9JT (UK)
Abstract The role which the stress assisted transformation of the tetragonal to monoclinic phase plays in the wear of zirconia ceramics remains unclear. This paper discusses the use of detailed transmission electron microscopy of the worn surface to assess the role of transformation in the shding wear of magnesia partially stabilised zirconia against a steel counterface. Three mechanisms were identified for the wear of the ceramic. Firstly, transformation occurred cooperatively in bands along the [lOtI] directions which led to extensive microcrack formation. Coalescence of the microcracks promoted preferential wear from the bands. Thus, transformation of the tetragonal to the evidence was found of a monoclinic phase had led to an increased wear rate. Secondly, clear experimentai trib~hemical wear mechanism of the zirconia. In such areas, the transformation of the tetragonal to mon~linic phase played a minor role and was only found to a depth of 200 nm. Finally, abrasive grooving was shown to cause extensive plastic deformation. Interestingly, no evidence of fracture or transformation was found at the abrasive grooves. Thus, this work has demonstrated that under these particular sliding conditions, transformation toughening can have a detrimental as well as a beneficial effect on the tribological properties of zirconia.
1. Introduction
Many workers have highlighted the importance of fracture toughness in the wear resistance of ceramics [l]. Zirconia engineering ceramics offer a high toughness (K,,) as well as high hardness and good chemical inertness [2]. Therefore, these materials should provide excellent wear resistance. Indeed, transformation toughened ceramics have found several applications as tribological materials [3-53 and have given good laboratory results L&-8].However, these materials have also shown very poor wear resistance both under mild sliding conditions and in apparently similar conditions to those where they perform well [9-111. Whilst variable wear behaviour has been attributed to low fracture toughness in many ceramics, zirconia ceramics are generally considered to be damage tolerant and therefore should be expected to give more consistent performance. Transformation toughened zirconia ceramics have received considerable attention over the last 15 years [12]. Substantial understanding exists as to the toughening mechanisms and the relationship between toughness and the transformation of the tetragonal phase to the monoclinic phase. The transformation is asso-
ciated with a volume increase of approximately 5%. Transformation occurs adjacent to an advancing crack tip, and the toughening increment results from crack tip shielding, a mechanism which exerts compressive forces on the crack wake as opposed to the crack tip. Thus, the mechanism inherently relies on the constraint and integrity of the untransformed surrounding material. Indeed, it has been shown that in the toughest MgPSZ materials where very large amounts of transformation are found, failure occurs as a result of the coalescence of the microcracks initiated by the transformation itself, in other words, transfo~at~on limits the strength in the toughest zirconia ceramics [13]. At the surface of the ceramic there is a reduced constraint because of the absence of surrounding material. Thus, spontaneous transformation occurs on cooling from sintering. During grinding or abrasion it is well established [14] that surface transformation can promote compressive stresses imparting good abrasion resistance and a damage tolerant ceramic. However, during sliding contact the surface stress state is different; in particular, adhesive forces promote a strong tensile force and a plucking action. For transformation to provide a toughening increment the surrounding ma-
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W. M. Rainforth, R. Stevens / Wear of magnesia partially stabilised zirconia
terial must remain untransformed in order to provide the constraint required to promote the compressive forces. It is not clear what effect the combination of the absence of constraint at a free surface and the stress system resulting from sliding contact will have on the transformation process. Some authors have suggested that high wear rates in zirconias may be associated with surface fracture on a very fine scale [ 111.Birkby et al. [ 1l] have demonstrated that a 2 mol.% yttria-tetragonal zirconia polycrystals (2Y-TZP) (K,, of 11 MPa m-l’*) wore at a greater rate than a 3Y-TZP (K,, of 6 MPa m-l”) drawing die. The wear mechanism was found to be different between the two ceramics, whereby material removal occurred by an intergranular fracture mechanism in the 2Y sample but a ‘plastic deformation’ mechanism in the 3Y sample. Since the higher toughness ceramic had undergone surface fracture whilst the lower toughness ceramic had not, it was suggested that the transformation of the tetragonal phase to the monoclinic phase and the associated volume increase had resulted in eruption of the surface and therefore an increase in the wear rate. In contrast, Lindberg and Richerson [7] found that the strength degradation as a result of sliding contact was less for both Y-TZP and Mg-PSZ than for Sic or Si,N,. Moreover, Fischer et al. [l] have shown that, at very slow sliding speeds, the wear rate decreased to the fourth power as the toughness increases from a 6Y-TZP to a 3Y-TZP. However, the zirconias with yttria contents greater than 3Y contain appreciable levels of cubic phase and in none of the ceramics would the same levels of transformation from tetragonal to monoclinic phase be expected as in the 2Y-TZP investigated by Birkby et al. [ll]. Clearly, there remain several unanswered questions about the role of transformation in the wear behaviour of zirconia ceramics. This paper presents a detailed investigation into the role of transformation at the worn surface and in particular reports the results of detailed transmission electron microscopy (TEM) investigation of the worn surface of an Mg-PSZ.
2. Experimental
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of the annular disc, located by a central spindle. Loads were applied by dead weight which were secured directly to the top plate containing the pins. The Mg-PSZ used was supplied by Coors UK in the maximum hardness condition with a toughness of 7.9 MPa m-l’* as measured by indentation (indentation gives a lower estimate of the toughness compared to single edge notched beam for example [16]). The microstructure of the material consisted of fine tetragonal precipitates in a coarse cubic grain structure (grain size 80 pm), Fig. 1, with some monoclinic phase along the grain boundaries. The zirconia was tested as both the disc and the pins but in each case was worn against a steel counterface. Zirconia pins were worn against a l%Cr, l%C bearing steel disc whilst 316L stainless steel pins were worn against a zirconia disc. A bearing steel disc was used in order to reduce, as far as possible, the wear rate of the disc and ensure that the major wearing component was the pins. Stainless steel pins were used in order to study the high strain deformation of a single phase steel (the results of which are reported in ref. 19). It was believed that the zirconia disc would show little or no wear. However, significant wear did occur and therefore it was considered that additional information could be derived from the study of the zirconia disc. Moreover, this provided information about the wear of zirconia against a relatively soft and a relatively hard counterface. All tests were performed at 0.24 m s-’ using loads in the range 5-55 N per pin under dry ambient conditions. All surfaces were prepared by lapping to a high standard of surface polish. The R, of the zirconia disc was measured at 0.008 pm, and the pins were believed to have a similar value but were too small to measure directly. The wear rate was measured by weight loss. A control pin was kept next to the wear rig during
procedure
Wear testing was undertaken on a tri-pin-on-disc machine, the details of which are given elsewhere [15]. The rig employs three 1 cm diameter pins with a truncated cone machined at one end, providing a 3 mm diameter contact face. The pins were held in a top plate which was prevented from rotating by two half-bridge strain gauges, whilst the disc was rotated. The design was such that only a few millimetres of the pin projected from the top plate surface in order to ensure maximum stiffness. The head was placed on top
Fig. 1. TEM image of the tetragonal precipitates in the cubic matrix. The precipitates are of a size that are metastable and the propagation of a crack would result in transformation to the monoclinic phase.
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W: M. Rainforth, R. Stevens I Wear of magnesia partially stabilised zirconia
tests and its weight measured at the same time as the pins in order to correct for errors in weight changes attributed to factors such as water absorption. The worn surfaces were examined by optical microscopy (Nomarski contrast), scanning electron microscopy (SEM) and transmission electron microscopy (TEM) (Jeol2OOCX) and X-ray diffraction. The volume fraction of monoclinic phase was determined using the Toraya et al. [17] modification to the equation proposed by Garvie and Nicholson [18]. Specimens of the worn surface for TEM were obtained by argon ion milling from below the worn surface (back thinning~, the full details of which are given elsewhere [2]. Measurements were made of the depth of the microstructural change from the worn surface. The surface was considered to be the region where the transferred layer was detected. The thinning rate of the ion beam thinner was calibrated for a given thinning angle and ion beam current. A given specimen was viewed initially when the smallest hole was found. The same specimen was then further thinned, using both ion beams (compared to the single beam used in back thinning) for a carefully monitored time and then re-examined. This permitted the change in microstructure as a function of depth to be assessed in each foil. In addition, pronounced features in the foil (e.g. large pores) and the diameter of the hole in the foil were used to further assess the amount of additional thinning which had occurred. It was found that the thinning time to perforation was very reproducible provided new cathodes were used each time and the beam current was carefully controlled. However, there were clearly errors in such a technique and of a magnitude difficult to assess but were considered to be in the region of f 150 nm.
+ +
/
/
10
20
,
I
,
30
,
1
I
1,
40
50
60
70
LOAD
iN)
Fig. 2. Wear factor (wear rate normaiised by the load and sliding distance) as a function of load for the Mg-PSZ disc worn against stainless steel pins.
t
1
IO
20
30
40
LOADINI Fig. 3. Wear factor of the Mg-PSZ pins as a function of load, worn against the hardened bearing steel disc.
3. Results 3.1. Wear results The wear rates as a function of load are given for the pins and the discs in Figs. 2 and 3. The wear behaviour of the metal pins is discussed in detail elsewhere [19]. In all tests the running-in period was quiet and was associated with a low friction coefficient (less than 0.1 irrespective of couple or load). At this stage, little damage was observed on the worn surface. After approximately 2 km of sliding a gradual increase in noise, vibration and friction coefficient was found. This increase was associated with the start of visible transfer of the metal to the ceramic. Steady state was characterised by a rapidly changing friction coefficient (0.2-0.4, irrespective of load or couple), substantial surface noise and vibration at the highest load (55 N per pin). Extensive metal transfer to the zirconia surface was observed for the zirconia
disc worn against stainless steel pins. Similar transfer had occurred for the zirconia pins against the hardened bearing steel disc, but to a much lesser extent. The transferred layer was shown to consist of metal oxide containing only very small amounts of zirconia. The role of the transfer layer will be considered in detail later. 3.2. X-Ray di~acti~n ~~~e~ts Table 1 gives the vol.% monoclinic of the starting material, a ground surface and the worn surfaces (for the case of the zirconia disc the X-ray measurements were made after the transfer layer had been removed using the method described below). The ground surface monoclinic can be taken as a measure of the maximum level of monoclinic which could be generated by stress assisted transfo~ation, and is the level expected from severe abrasive wear. The polished surface monoclinic
W. M. Rainforth, R. Stevens I Wear of magnesia partially stabilised zirconia TABLE
1. Vol.% monoclinic
Polished surface Ground surface Pin, 6.6 N Pin, 19.6 N Pin, 37.6 N Disc, 37 N
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values 10% 36.6% 18.6% 11.2% 17.7% 24.5%
is the quantity of monoclinic in the sintered material (which occurs mainly along the grain boundary regions). Thus, the sliding wear tests produced some stress assisted transformation, but the levels were fairly low; they were lower than those which could be generated by gentle rubbing on 600 grit silicon carbide, for example. 3.3. Optical and scanning electron microscopy During the initial stages of contact there was little or no evidence of metal transfer to the zirconia surface; however, transfer rapidly built up after a few km of sliding. For the zirconia disc running against the stainless steel pins, the outer regions of the wear scar (where, because of the pin profile, contact was always comparatively recent) showed the characteristics of mild wear. However, the majority of the wear scar on the disc was covered with a thick transfer layer, the majority of which had become oxidised. Frequent areas were found where the oxide layer had spalled. Thus, the wear of the metal occurred via adhesion promoting transfer, oxidation of the transfer layer and eventual spalling of the layer at some critical thickness. No evidence was found of the removal of zirconia particles during the spalling process, although small particles would have been difficult to assess because of the low volume fraction. By comparison of the wear rates the zirconia particles would have constituted less than 5% by weight of the wear debris. Moreover, as will be shown later, much of the zirconia was found in solid solution in the metal oxide debris, making detection of isolated particles even more difficult. In order to establish the exact nature of the worn surface of the disc it was necessary to remove the transferred metal layer. This was done by dissolving the metal oxide in dilute HCl (which was shown not to attack the zirconia surface in any way, by an examination of a backthinned polished surface in both the etched and unetched conditions). Once this had been undertaken it was evident that the same basic features were present on both the zirconia pins and discs. A general view of the worn surface of the disc after removal of the transfer layer is given in Fig. 4. The zirconia grains showed two different morphologies. Many grains were very smooth and the only evidence of wear was differential wear between grains as shown by the surface relief. This mechanism was dominant
Fig. 4. Optical micrograph (Nomarski contrast) of the worn MgPSZ disc after removal of the transfer layer. Note the smooth areas which show grain relief and the grains which have undergone more severe wear which exhibit parallel grooves.
on the outer region of the wear track on the discs and constituted approximately 65% of the area on the worn pin surfaces. Specifically, the mechanism was associated with regions of less metal transfer. The second wear mechanism was dominant on the worn disc, but less frequent on the pin surfaces, and constituted a much more severe wear mechanism. In these regions the surface consisted of grains which contained parallel grooves, and occasionally isolated pits, which did not cross grain boundaries and were perpendicular or within 30” to perpendicular of the sliding direction, Figs. 4-6. On the zirconia disc worn against the stainless steel the grooves tended to be broad and frequently deep, Fig. 5, whilst on the zirccnia pins worn against the bearing steel the grooves were more crack like, Fig. 6. Thus, the amount and severity of grooving corresponded to the amount and severity of the metal transfer. The grooving was clearly associated with the crystallographic orientation of the zirconia grains. In order to verity this the surface was further etched in HF. The tetragonal/monoclinic precipitates were found to be orientated along and perpendicular to the grooves,
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W. M. Rainforth,
R. Stevens 1 Wear of magnesia partially stabilised zirconia
Fig. 5. Detail of the grooves on the disc surface. Note how they do not cross the grain boundaries.
Fig. 7. Detail of the grooves after etching the precipitate orientation.
in HF which shows
Fig. 6. SEM image using back-scattered electrons showing the worn Mg-PSZ pins. The grooves were much sharper and more crack like.
Fig. 7. Those grains which did not contain any grooves failed to show any ciearly discernible orientation relationship.
Fig. 8. TJZM image of a back thinned specimen showing the interface between zirconia (upper region) and transferred metal oxide (bottom left). Much of the cubic matrix has been removed. The precipitates showed both monoclinic and tetragonal symmetry.
3.4. Transmission electron microscopy
surface of the same sample and on surfaces where the transferred layer had been removed. The interface between the transferred layer and the zirconia was shown to be diffuse rather than clear cut, which infers some tribochemical reaction had taken place between the two. ~roughout these tests no discrete zirconia wear debris particles were ever found. However, the metal oxide was shown to contain varying amounts of zirconium ions in solid solution which is further evidence of a chemical interaction behveen metal and ceramic. The quantity of zirconia in solid solution indicated that the tribochemical wear mechanism must have constituted a significant fraction of the overall wear rate. The structure of the metal oxide was found to be a mixture of three phases, namely, Fe@,, an amorphous
3.4.1.
with the transfer layer present
A view of the interface between the transferred metal oxide and the zirconia is shown in Fig. 8 taken from an area with no grooves. The oblate spheroid shape of the precipitate appears to have been unchanged by the wear mechanism. Interestingly, some of the precipitates were shown to have retained the tetragonal phase symmetry, although the majority were shown to be monoclinic. Thus, in~mplete transfo~ation had occurred at the worn surface. The significance of this point will be discussed later. Virtually all the cubic matrix phase appears to have been removed at the outer zirconia surface. This result was also found in occasional areas elsewhere on the
W. M. Rainforth, R. Stevens / Wear of magnesia partially stabilised zirconia
oxide phase and a non-equilibrium body centred cubic oxide phase. The proportions of each phase varied between each test and counterface material, but in each case, the proportion of metallic ions in the oxide was identical to the proportion of metal atoms in the steel. This result and the presence of an amorphous phase indicates that the temperature was not high enough to generate the thermodynamically predicted metal oxides. 3.4.2. Smooth areas with the oxide removed Whilst the structure was difficult to image at the worn surface because of large residual strains, it was clear that the precipitates were mainly monoclinic (as shown by diffraction analysis), Fig. 9. Interestingly, the structure of these areas retained the same general morphology as the starting material. However, some areas showed a distortion in the precipitate geometry and orientation, as illustrated in Fig. 9. In addition, small rotations in the diffraction pattern were observed across small areas of the sample (less than 2 mm). No dislocations or subgrain boundaries were observed. Thus, the distortion was considered to be mainly elastic rather than as a result of plastic deformation. The depth over which any microstructural change could be detected was shown to be only of the order of 100-200 nm from the surface, in agreement with the X-ray results. In the region more than 200 nm from the surface some residual strain tended to obscure the image but no distortion of the precipitates was found. Interestingly, the majority of precipitates at this depth were tetragonal. 3.4.3. Regions where the grooving was observed Back-thinned samples from the disc were difficult to interpret because the surface was rough and thin foils
Fig. 9. TEM image of a back thinned specimen showing the near surface structure after removal of the transferred layer. The precipitates are mainly monoclinic as shown by the twinning, but there is considerable distortion of the structure, believed to be from residual elastic strain.
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tended to break-up. This was also taken to indicate a high proportion of microcracks in this area. A general view of the sample taken from the pin worn against the bearing steel disc is shown in Fig. 10 with a detail in Fig. 11. The micrograph confirms that the cracks/grooves were orientated in a [lOO] direction, i.e. they were parallel to the long axis of one of the three precipitate orientations. The shapes of the precipitates in this region had remained unchanged, but were shown to be monoclinic, with no tetragonal phase positively identified. The depth to which transformation had occurred was greater than in those areas discussed above, although an exact measurement was difficult. Regions inbetween the major cracks showed extensive microcracking, Fig. 11, with all cracks aligned along the same [loo] direction as the larger cracks. Evidence was found of a transition between the fine microcracks and the coarse macrocracks. Thus, it would appear that the larger cracks had been formed as a result of microcrack coalescence. One additional feature of interest was found in these foils. Occasional three-body abrasive grooves were ob-
Fig. 10. TEM image of a back thinned specimen taken from an area showing grooving. The diffraction pattern shows a [lOO] type zone and demonstrates that the cracks run along a [lOO] direction.
Fig. 11. Detail from Fig. 10. Fine microcracks have initiated at the cubic/monoclinic interface which then coalesce to give the larger cracks in Fig. 10.
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W. M. Rainforth, R. Stevens I Wear of magnesia partially stabilised zirconia
Fig. 12. Detail from a region just adjacent to Fig. 10 showing the interface between the normal structure of cubic/monoclinic (upper right) and a three-body abrasive groove (lower right). The diffraction pattern from the abraded region shows considerable rotation in the structure. Further work is required to fully evaluate the microstructural changes which have taken place.
served which had totally destroyed the precipitates and led to several other complex changes, as shown in Fig. 12. Whilst the details of the microstructure require further work to clarify, it was clear that the precipitate structure had been destroyed. Considerable rotation of the diffraction pattern occurred as the selected area aperture was moved across small distances. In addition, the sharp cubic and tetragonal diffraction spots were replaced by an arc of diffraction spots extending over 5-10”. This is indicative of a very fine grain size (dark field imaging indicated of the order of 10-30 nm) within the coarse (80 pm) cubic grain. Interestingly, no evidence of cracking was found in these areas.
4. Discussion 4.1. Smooth su$ace areas and metal transfer The extent of transformation at the worn surface as shown both by X-ray results and TEM results was low, substantially lower than that expected under abrasive wear conditions. The transformation zone was shown to be less than about 200 nm at the surface, whereas transformation zone widths adjacent to a crack in the same material would be expected to be up to 8 pm in size [20]. The TEM results demonstrated that where transformation from tetragonal to monoclinic had occurred it was restricted to about the outer 200 nm of the worn surface for the smooth regions, although it extended somewhat deeper in the regions containing grooves. The smooth areas which showed differential wear between different grains (‘grain relief) is typical of ceramics which exhibit mild wear. For example, Wallbridge et al. [15] have observed a similar effect for the sliding wear of alumina on alumina under the same test geometry as the present study. Wallbridge et al. [15] used the work of Steijn [21] to explain the origin
of the grain relief. Steijn [21] found that, for single crystal alumina, resistance to wear depended strongly on crystal orientation. Very high wear resistance was found when the rubbing direction was perpendicular to the c axis on the prismatic and rhombohedral crystal faces. However, brittle type wear occurred when rubbing was parallel to the c axis. Thus, wear results correlated qualitatively with the ease of plastic deformation. On the assumption that the differences in elastic modulus with crystallographic orientation were insufficient to explain the grain relief, Wallbridge et al. [15] suggested that those grains which were orientated for slip would wear less than those grains which were not orientated for slip. In the latter case the grain was considered to undergo a microfracture mechanism. The variation in hardness with crystal orientation in ceramics is well established [22, 231, and it is also apparent that the hardness is not an inherent material property but depends on the load used to make the indent [23]. In addition, chemo-mechanical effects can have a strong effect on hardness, whereby surface absorbed species alter the ease of crystallographic slip and therefore the hardness [24]. Such arguments can also be invoked to explain grain relief on the basis that water vapour can also provide a chemo-mechanical effect that alters the surface plasticity. Clear evidence has been presented in the present study which shows that no plastic deformation type wear mechanism had occurred in this region although significant residual elastic strain was evidently present. Moreover, no evidence of microfracture type mechanisms was found, and would have been clearly visible in these experiments. Rather, the wear mechanism appears to have been predominantly one of tribochemical wear. The cubic matrix was shown to have been preferentially removed and the interface between metal oxide and zirconia was diffuse. No discrete zirconia wear debris was found, the zirconia only being found in solid solution in the transferred layer and the free wear debris. Therefore, the wear rate is dependent on the chemical reaction rates as a function of crystallographic orientation. Whilst it is possible to envisage that the preferential removal of cubic phase by the tribochemical wear mechanism would increase the rate of transformation, it is considered that such an effect would be very surface specific (of the order of the precipitate size) and would be of less significance than the contact stresses. Further work is required to investigate this mechanism in other non-transforming ceramics, particularly alumina, since grain relief is often associated with very low wear rates and the mechanism may therefore determine the ultimate wear rate that is achievable. The observation that transformation was restricted to the outer 200 nm would most probably have been
W. M. Rainforth, R. Stevens / Wear of magnesia partially stabilised zirconia
a result of the temperature rises at the interface reducing
the driving force for transformation, rather than low contact stresses. Values are available in the literature for fracture toughness (and therefore extent of stress induced transformation) as a function of temperature for Mg-PSZs with different martensite start temperatures (MS) [25]. The M, value for the material in the current work would be around 113 K, which indicates that the K,, would fall to that equivalent to the base cubic matrix at around 250 “C, i.e. that very little transformation would be expected to occur above this temperature (although the actual M, temperature was not measured for the material used in this study). Thus, the observation of transformation at the outer surface, and to a greater extent in the cracked/grooved area (discussed below) suggests that the surface temperature must have been such that some driving force for transformation still existed. Such an estimate of temperature is in line with the non-equilibrium phase composition of the metal oxide, which would have been expected to transform to equilibrium oxides at higher temperatures. Indeed, heating the wear debris resulted in rapid transformation to the equilibrium structure at temperatures as low as 300 “C. In addition, this estimate agrees well with analytical calculations of the interface temperature made elsewhere [2], of measurements made by an implanted thermocouple [2] and of the microstructure at the worn stainless steel surface [19]. 4.2. Areas with groovinglcracking The grooving/cracking had promoted a higher wear rate than that associated with the grain relief. The grooving/cracking was clearly associated with the crystallographic orientation of the zirconia grains, as shown both by the etching experiments and the TEM results, with cracking being largely restricted to the [NO] directions. Similar parallel grooves and various shaped pits have also been reported by Hannink et al. [4] for the wear of Mg-PSZ tappet facing under lubricated wear. Whilst these authors did also ascribe orientation relationships based on the appearance of the groove/ pit, they did not make any attempt to explain the mechanism of formation. Three possible mechanisms can be envisaged for the formation of the grooves/cracks. Firstly, the cracks may have formed as a result of thermal expansion stresses, so-called thermal crazing. However, this mechanism can be rejected for two reasons. Firstly, thermal cracking would cross grain boundaries whereas the grooves did not. Secondly, tests conducted at higher sliding speeds tended to produce less grooving/cracking, which is the opposite to that expected from thermal cracking. The second possible mechanism is that transformation occurred cooperatively in bands (in a similar manner to the transformation bands found adjacent to a hardness
329
indent or a bend test specimen, as reported by Marshall and James [26] for example). Such bands would lead to a local raising of the surface thereby giving preferential wear in this region. However, many bands were either too deep or too sharp to have arisen from such a mechanism. Moreover, transformation was shown to occur throughout the surface rather than in bands. The final mechanism is one where transformation of tetragonal to monoclinic led to microcracking which, with the repeated application of shear stresses, coalesced to give macro surface cracks. Coalescence of these cracks would then transform the appearance to a groove. Very fine microcracks were found at the interface between precipitate and matrix which appeared to have coalesced to give larger cracks running through the cubic matrix and along precipitate interfaces (Figs. 10, 11). The interface between the tetragonal (or monoclinic) precipitate lies mainly on the (100) planes. In addition, the twin variants of the monoclinic are highly populated along (OOl),, so that maximum strain is generated along {loo}, [27]. Consequently, microcracks occur preferentially along this direction, explaining the initiation of grooves in grains with a [NO], orientation relationship with the surface before initiation in other orientations. Thus, the third possible mechanism of microcrack coalescence was believed to be responsible for the grooving/cracking. The greater severity of the mechanism when the zirconia was worn against the stainless steel counterface compared to the harder bearing steel disc suggests that the adhesive forces played a critical role in the mechanism. The greater rate of metal transfer when worn against the softer stainless steel counter-face compared to the harder bearing steel disc was considered to reflect the levels of the adhesive forces. The adhesive forces will have increased the tractions normal to the surface, thereby increasing the crack nucleation and propagation rates. Moreover, this explains why such grooving is not observed under abrasive wear conditions. Adhesive forces are low under abrasive wear conditions which ensures that the dominant forces are compressive rather than tensile, such that transformation imparts a beneficial effect. Clearly, with the absence of the constraint of untransformed material, the state of the surface forces is critical in determining whether transformation has a beneficial or detrimental effect. In addition to the differences in adhesive forces between the two test configurations, it is also probable that the flash temperature in the zirconia surface was higher for the pins compared to the discs since the pins were in continuous contact whilst the discs were in intermittent contact. The driving force for transformation would have been reduced at higher temperature thereby reducing the extent of a transformation driven wear mechanism.
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R. Stevens / Wear of magnesia partially stabiiised zirconur
The current results are not considered to contradict the findings of either Birkby et al. [ll] or Fischer et al. [l]. Fischer et al. [l] found that the wear resistance of yttria-zirconias varied with the fourth power of the toughness, which suggests that the transformation of the tetragonal to monoclinic has a dramatic beneficial effect. However, the toughest ceramic tested was a 3YTZP which has a limited transformation zone width. The transformation zone width is a measure of the distance from a crack that stress assisted transformation occurs and would be substantially larger in the MgPSZ used in the current study, compared with the 3YTZP used by Fischer et al. [l]. Both Birkby et al. [l l] and Rainforth [2] have demonstrated that in tougher TZP (e.g. %Y-TZP) ceramics the transformation zone widths are appreciably larger and there is evidence of a transformation driven wear mechanism. Birkby et al. [ll] clearly demonstrated that such a transformation initiated mechanism does result in accelerated wear rates. Therefore, whilst transformation can impart improved wear resistance, especially in abrasive wear conditions, there is a maximum toughness in which the zirconia ceramic should be used in sliding wear. Frequently, the zirconia should not be used in the maximum toughness condition, but in the maximum hardness condition. However, the exact value of the toughness will depend on the surface stress system, in particular, the magnitude of the adhesive forces. 4.3. Abrasive grooves Occasional three-body abrasive grooves were observed on the zirconia surface which were presumed to have been formed from detached wear debris. TEM analysis of the grooves indicated that they were associated with dramatic microstructural change. The resulting structure apparently consisted of cubic phase with no positive identification of monoclinic or tetragonal. Considerable rotation of the diffraction pattern occurred over small distances, greater than could be possible by residual elastic strains. The precipitate structure had been completely destroyed and replaced by a nanocrystalline structure which was apparently cubic. This, therefore, strongly suggests that extensive dislocation flow had taken place. Temperature in itself could not be responsible since many hours at 1700 “C are required to dissolve the tetragonal precipitates and a three-body abrasive wear event is very short. However, the structure was extremely difficult to image and none of the subgrain boundaries could be imaged because of the residual strain. Interestingly, no cracking of any form could be found in the abrasive grooves even in grains that had undergone transformation initiated wear. However, the suggestion that compressive stresses from transformation are responsible for preventing crack formation is not appli-
cable in this case, rather the mechanism is apparently much more complex. Indeed, Swain and Hannink [14] have shown interesting and complex phase and morphological changes below the ground surface of a CeTZP. Further work is required to establish the actual structure of this region and the mechanism of formation.
5. Conclusions (i) A tribochemical wear mechanism between the steel (and transferred metal oxide) and the zirconia was shown to occur which promoted a smooth wear surface with grain relief. No evidence of microfracture or plastic deformation was found contradicting other theories of mild ceramic wear. (ii) In the smooth regions, minimal stress assisted transformation had occurred and therefore it played only a minor role in the wear mechanism. (iii) Greater quantities of transformation were found in regions where extensive metal transfer had occurred. Transformation of the tetragonal to monoclinic phase in these regions has been found to increase the wear rate. Microcracks were found to initiate at the matrix/ precipitate interface which coalesced to give macrocracks. (iv) The extent of the transformation initiated wear mechanism depended strongly on the adhesive forces, being more pronounced for a stainless steel compared to a bearing steel counterface. (v) The only circumstances where dislocation flow was identified was in three-body abrasive grooves.
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R. H. J. Hannink, M. J. Murray and M. Marmach, Magnesiapartially stabilised zirconias (Mg-PSZ) as wear resistant materials, in K. C. Ludema (ed.), Proc. Int. Conf. on Wear of Materials, Reston, VA April 11-14, 1983, American Society of Mechanical Engineers, New York, 1983, pp. 181-186. R. H. J. Hannink, M. J. Murray and H. G. Scott, Friction and wear of partially stabilized zirconia: basic science and practical applications, Wear, 1tW (1984) 35.5-366. H. G. Scott, Friction and wear of zirconia at very low sliding speeds, Proc. Int. Conf on Wear of Materials, Vancouver, Canada, 1985, American Society of Mechanical Engineers, New York, 1985, pp. 8-12.
W. M. Rainforth, R. Stevens I Wear of magnesia partially stabilised zirconia 7 L. J. Lindberg
8
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and D. W. Richerson, Comparison of the contact stress and friction behaviour of SIC and ZrOz materials, Proc. 2nd Int. Conf on Ceram. Mater. and Components for Engines, Lubeck-Travemunde, 1986, pp. 20-27. C. Yust and F. Car&ran, Observations on the sliding wear of ceramics, ASLE Trans., 28 (1984) 245-252. J. Breznak, E. Breval and N. H. Macmillan, Sliding friction and wear of structural ceramics, Part 1, Room temperature behaviour, J. Mater. Sci., 20 (1985) 4657-4680. E. Breval, J. Breznak and N. H. Macmillan, Sliding friction and wear of structural ceramics, Part 2, Analysis of roomtemperature wear debris, J. Mater. Sci., 21 (1986) 931-935. I. Birkby, P. Harrison and R. Stevens, The effect of surface transformation on the wear behaviour of zirconia TZP ceramics, J. Eur. Ceram. Sot., 5 (1989) 37-46. I. Nettleship and R. Stevens, Tetragonal zirconia polycrystals (TZP)-a review, Znt. J. High Tech. Ceram., 3 (1987) l-32. M. V. Swain, Inelastic deformation of Mg-PSZ and its significance for strength-toughness relationships of zirconia toughened ceramics, Acta MetaN., 33 (1985) 2083-2091. M. V. Swain and R. H. J. Hannink, Metastability of the martensitic transformation in a 12 mol% ceria-zirconia alloy: II, Grinding studies, J. Am. &ram. Sot., 72 (1989) 1358-1364. N. Wallbridge, D. Dowson and E. W. Roberts, The wear characteristics of sliding pairs of high density polycrystalline aluminium oxide under both dry and wet conditions, in K. C. Ludema (ed.), Proc. Int. Conf on Wear of Materials, Reston, VA, April 11-14, 1983, American Society of Mechanical Engineers, New York, 1983, pp. 202-211. J. Wang, M. Rainforth and R. Stevens, The grain size dependence of the mechanical properties in TZP ceramics, Btit. &ram. Trans. .I, 88 (1989) l-6.
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17 H. Toraya, M. Yoshimura and S. Somiya, Calibration curve for quantitative analysis of the monoclinic-tetragonal ZrOz system by X-ray diffraction, J. Am. Ceram. Sot., 67 (1984) c119. 18 R. C. Garvie and P. S. Nicholson, Phase analysis in zirconia systems, J. Am. Ceram. Sot., 55 (1972) 303-305. 19 W. M. Rainforth, R. Stevens and J. Nutting, Deformation structures induced by sliding contact, Philos. Mug., 66 (1992) 21. 20 M. Ruhle, B. Kruas, A. Strecker and D. Waidelich, in N. Claussen, M. Ruhle and A. H. Heuer (eds.), Science and Technology of Zirconia II, American Ceramics Society, Columbus, OH, 1984, p. 256. 21 R. P. Steijn, On the wear of sapphire, J. Appl. Phys., 32 (1961) 1951-1958. 22 J. T. Czernuszka and T. F. Page, Characterising the surface contact behaviour of ceramics, Part 1, Hardness response of glass-bonded alumina and titania, J. Mater. Sci., 22 (1987) 3907-3916. 23 J. T. Czemuszka and T. F. Page, A problem in assessing the wear behaviour of ceramics: load, temperature and environmental sensitivity of indentation hardness, Proc. Br. Ceram. Sot., 34 (1984) 145-156. 24 A. R. C. Westwood, N. H. Macmillan and R. S. Kalyoncu, Environment-sensitive hardness and machinability of A&O,, J. Am. Ceram. Sot., 56 (1973) 258-262. 25 P. F. Becher, M. V. Swain and M. K. Ferber, Relation of transformation temperature to the fracture toughness of transformation-toughened ceramics, J. Mater. Sci., 22 (1987) 76-84. 26 D. B. Marshall and M. R. James, Reversible stress-induced martensitic transformations, J. Am. Ceram. Sot., 69 (1986) 215-217. 27 G. Bansal and A. H. Heuer, On a martensitic phase transformation in zirconia-I metallographic evidence, Acta Metall., 20 (1972) 1281-1289.
Un~ve~~~ of She~eId, PO Box 600, Sir Robert ~ad~eid Building, copper
Street,
R. Stevens Schoof
ofpatents,
V~iversi~
of Leeds, Leeds LS2 9JT (UK)
Abstract The role which the stress assisted transformation of the tetragonal to monoclinic phase plays in the wear of zirconia ceramics remains unclear. This paper discusses the use of detailed transmission electron microscopy of the worn surface to assess the role of transformation in the shding wear of magnesia partially stabilised zirconia against a steel counterface. Three mechanisms were identified for the wear of the ceramic. Firstly, transformation occurred cooperatively in bands along the [lOtI] directions which led to extensive microcrack formation. Coalescence of the microcracks promoted preferential wear from the bands. Thus, transformation of the tetragonal to the evidence was found of a monoclinic phase had led to an increased wear rate. Secondly, clear experimentai trib~hemical wear mechanism of the zirconia. In such areas, the transformation of the tetragonal to mon~linic phase played a minor role and was only found to a depth of 200 nm. Finally, abrasive grooving was shown to cause extensive plastic deformation. Interestingly, no evidence of fracture or transformation was found at the abrasive grooves. Thus, this work has demonstrated that under these particular sliding conditions, transformation toughening can have a detrimental as well as a beneficial effect on the tribological properties of zirconia.
1. Introduction
Many workers have highlighted the importance of fracture toughness in the wear resistance of ceramics [l]. Zirconia engineering ceramics offer a high toughness (K,,) as well as high hardness and good chemical inertness [2]. Therefore, these materials should provide excellent wear resistance. Indeed, transformation toughened ceramics have found several applications as tribological materials [3-53 and have given good laboratory results L&-8].However, these materials have also shown very poor wear resistance both under mild sliding conditions and in apparently similar conditions to those where they perform well [9-111. Whilst variable wear behaviour has been attributed to low fracture toughness in many ceramics, zirconia ceramics are generally considered to be damage tolerant and therefore should be expected to give more consistent performance. Transformation toughened zirconia ceramics have received considerable attention over the last 15 years [12]. Substantial understanding exists as to the toughening mechanisms and the relationship between toughness and the transformation of the tetragonal phase to the monoclinic phase. The transformation is asso-
ciated with a volume increase of approximately 5%. Transformation occurs adjacent to an advancing crack tip, and the toughening increment results from crack tip shielding, a mechanism which exerts compressive forces on the crack wake as opposed to the crack tip. Thus, the mechanism inherently relies on the constraint and integrity of the untransformed surrounding material. Indeed, it has been shown that in the toughest MgPSZ materials where very large amounts of transformation are found, failure occurs as a result of the coalescence of the microcracks initiated by the transformation itself, in other words, transfo~at~on limits the strength in the toughest zirconia ceramics [13]. At the surface of the ceramic there is a reduced constraint because of the absence of surrounding material. Thus, spontaneous transformation occurs on cooling from sintering. During grinding or abrasion it is well established [14] that surface transformation can promote compressive stresses imparting good abrasion resistance and a damage tolerant ceramic. However, during sliding contact the surface stress state is different; in particular, adhesive forces promote a strong tensile force and a plucking action. For transformation to provide a toughening increment the surrounding ma-
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W. M. Rainforth, R. Stevens / Wear of magnesia partially stabilised zirconia
terial must remain untransformed in order to provide the constraint required to promote the compressive forces. It is not clear what effect the combination of the absence of constraint at a free surface and the stress system resulting from sliding contact will have on the transformation process. Some authors have suggested that high wear rates in zirconias may be associated with surface fracture on a very fine scale [ 111.Birkby et al. [ 1l] have demonstrated that a 2 mol.% yttria-tetragonal zirconia polycrystals (2Y-TZP) (K,, of 11 MPa m-l’*) wore at a greater rate than a 3Y-TZP (K,, of 6 MPa m-l”) drawing die. The wear mechanism was found to be different between the two ceramics, whereby material removal occurred by an intergranular fracture mechanism in the 2Y sample but a ‘plastic deformation’ mechanism in the 3Y sample. Since the higher toughness ceramic had undergone surface fracture whilst the lower toughness ceramic had not, it was suggested that the transformation of the tetragonal phase to the monoclinic phase and the associated volume increase had resulted in eruption of the surface and therefore an increase in the wear rate. In contrast, Lindberg and Richerson [7] found that the strength degradation as a result of sliding contact was less for both Y-TZP and Mg-PSZ than for Sic or Si,N,. Moreover, Fischer et al. [l] have shown that, at very slow sliding speeds, the wear rate decreased to the fourth power as the toughness increases from a 6Y-TZP to a 3Y-TZP. However, the zirconias with yttria contents greater than 3Y contain appreciable levels of cubic phase and in none of the ceramics would the same levels of transformation from tetragonal to monoclinic phase be expected as in the 2Y-TZP investigated by Birkby et al. [ll]. Clearly, there remain several unanswered questions about the role of transformation in the wear behaviour of zirconia ceramics. This paper presents a detailed investigation into the role of transformation at the worn surface and in particular reports the results of detailed transmission electron microscopy (TEM) investigation of the worn surface of an Mg-PSZ.
2. Experimental
323
of the annular disc, located by a central spindle. Loads were applied by dead weight which were secured directly to the top plate containing the pins. The Mg-PSZ used was supplied by Coors UK in the maximum hardness condition with a toughness of 7.9 MPa m-l’* as measured by indentation (indentation gives a lower estimate of the toughness compared to single edge notched beam for example [16]). The microstructure of the material consisted of fine tetragonal precipitates in a coarse cubic grain structure (grain size 80 pm), Fig. 1, with some monoclinic phase along the grain boundaries. The zirconia was tested as both the disc and the pins but in each case was worn against a steel counterface. Zirconia pins were worn against a l%Cr, l%C bearing steel disc whilst 316L stainless steel pins were worn against a zirconia disc. A bearing steel disc was used in order to reduce, as far as possible, the wear rate of the disc and ensure that the major wearing component was the pins. Stainless steel pins were used in order to study the high strain deformation of a single phase steel (the results of which are reported in ref. 19). It was believed that the zirconia disc would show little or no wear. However, significant wear did occur and therefore it was considered that additional information could be derived from the study of the zirconia disc. Moreover, this provided information about the wear of zirconia against a relatively soft and a relatively hard counterface. All tests were performed at 0.24 m s-’ using loads in the range 5-55 N per pin under dry ambient conditions. All surfaces were prepared by lapping to a high standard of surface polish. The R, of the zirconia disc was measured at 0.008 pm, and the pins were believed to have a similar value but were too small to measure directly. The wear rate was measured by weight loss. A control pin was kept next to the wear rig during
procedure
Wear testing was undertaken on a tri-pin-on-disc machine, the details of which are given elsewhere [15]. The rig employs three 1 cm diameter pins with a truncated cone machined at one end, providing a 3 mm diameter contact face. The pins were held in a top plate which was prevented from rotating by two half-bridge strain gauges, whilst the disc was rotated. The design was such that only a few millimetres of the pin projected from the top plate surface in order to ensure maximum stiffness. The head was placed on top
Fig. 1. TEM image of the tetragonal precipitates in the cubic matrix. The precipitates are of a size that are metastable and the propagation of a crack would result in transformation to the monoclinic phase.
324
W: M. Rainforth, R. Stevens I Wear of magnesia partially stabilised zirconia
tests and its weight measured at the same time as the pins in order to correct for errors in weight changes attributed to factors such as water absorption. The worn surfaces were examined by optical microscopy (Nomarski contrast), scanning electron microscopy (SEM) and transmission electron microscopy (TEM) (Jeol2OOCX) and X-ray diffraction. The volume fraction of monoclinic phase was determined using the Toraya et al. [17] modification to the equation proposed by Garvie and Nicholson [18]. Specimens of the worn surface for TEM were obtained by argon ion milling from below the worn surface (back thinning~, the full details of which are given elsewhere [2]. Measurements were made of the depth of the microstructural change from the worn surface. The surface was considered to be the region where the transferred layer was detected. The thinning rate of the ion beam thinner was calibrated for a given thinning angle and ion beam current. A given specimen was viewed initially when the smallest hole was found. The same specimen was then further thinned, using both ion beams (compared to the single beam used in back thinning) for a carefully monitored time and then re-examined. This permitted the change in microstructure as a function of depth to be assessed in each foil. In addition, pronounced features in the foil (e.g. large pores) and the diameter of the hole in the foil were used to further assess the amount of additional thinning which had occurred. It was found that the thinning time to perforation was very reproducible provided new cathodes were used each time and the beam current was carefully controlled. However, there were clearly errors in such a technique and of a magnitude difficult to assess but were considered to be in the region of f 150 nm.
+ +
/
/
10
20
,
I
,
30
,
1
I
1,
40
50
60
70
LOAD
iN)
Fig. 2. Wear factor (wear rate normaiised by the load and sliding distance) as a function of load for the Mg-PSZ disc worn against stainless steel pins.
t
1
IO
20
30
40
LOADINI Fig. 3. Wear factor of the Mg-PSZ pins as a function of load, worn against the hardened bearing steel disc.
3. Results 3.1. Wear results The wear rates as a function of load are given for the pins and the discs in Figs. 2 and 3. The wear behaviour of the metal pins is discussed in detail elsewhere [19]. In all tests the running-in period was quiet and was associated with a low friction coefficient (less than 0.1 irrespective of couple or load). At this stage, little damage was observed on the worn surface. After approximately 2 km of sliding a gradual increase in noise, vibration and friction coefficient was found. This increase was associated with the start of visible transfer of the metal to the ceramic. Steady state was characterised by a rapidly changing friction coefficient (0.2-0.4, irrespective of load or couple), substantial surface noise and vibration at the highest load (55 N per pin). Extensive metal transfer to the zirconia surface was observed for the zirconia
disc worn against stainless steel pins. Similar transfer had occurred for the zirconia pins against the hardened bearing steel disc, but to a much lesser extent. The transferred layer was shown to consist of metal oxide containing only very small amounts of zirconia. The role of the transfer layer will be considered in detail later. 3.2. X-Ray di~acti~n ~~~e~ts Table 1 gives the vol.% monoclinic of the starting material, a ground surface and the worn surfaces (for the case of the zirconia disc the X-ray measurements were made after the transfer layer had been removed using the method described below). The ground surface monoclinic can be taken as a measure of the maximum level of monoclinic which could be generated by stress assisted transfo~ation, and is the level expected from severe abrasive wear. The polished surface monoclinic
W. M. Rainforth, R. Stevens I Wear of magnesia partially stabilised zirconia TABLE
1. Vol.% monoclinic
Polished surface Ground surface Pin, 6.6 N Pin, 19.6 N Pin, 37.6 N Disc, 37 N
325
values 10% 36.6% 18.6% 11.2% 17.7% 24.5%
is the quantity of monoclinic in the sintered material (which occurs mainly along the grain boundary regions). Thus, the sliding wear tests produced some stress assisted transformation, but the levels were fairly low; they were lower than those which could be generated by gentle rubbing on 600 grit silicon carbide, for example. 3.3. Optical and scanning electron microscopy During the initial stages of contact there was little or no evidence of metal transfer to the zirconia surface; however, transfer rapidly built up after a few km of sliding. For the zirconia disc running against the stainless steel pins, the outer regions of the wear scar (where, because of the pin profile, contact was always comparatively recent) showed the characteristics of mild wear. However, the majority of the wear scar on the disc was covered with a thick transfer layer, the majority of which had become oxidised. Frequent areas were found where the oxide layer had spalled. Thus, the wear of the metal occurred via adhesion promoting transfer, oxidation of the transfer layer and eventual spalling of the layer at some critical thickness. No evidence was found of the removal of zirconia particles during the spalling process, although small particles would have been difficult to assess because of the low volume fraction. By comparison of the wear rates the zirconia particles would have constituted less than 5% by weight of the wear debris. Moreover, as will be shown later, much of the zirconia was found in solid solution in the metal oxide debris, making detection of isolated particles even more difficult. In order to establish the exact nature of the worn surface of the disc it was necessary to remove the transferred metal layer. This was done by dissolving the metal oxide in dilute HCl (which was shown not to attack the zirconia surface in any way, by an examination of a backthinned polished surface in both the etched and unetched conditions). Once this had been undertaken it was evident that the same basic features were present on both the zirconia pins and discs. A general view of the worn surface of the disc after removal of the transfer layer is given in Fig. 4. The zirconia grains showed two different morphologies. Many grains were very smooth and the only evidence of wear was differential wear between grains as shown by the surface relief. This mechanism was dominant
Fig. 4. Optical micrograph (Nomarski contrast) of the worn MgPSZ disc after removal of the transfer layer. Note the smooth areas which show grain relief and the grains which have undergone more severe wear which exhibit parallel grooves.
on the outer region of the wear track on the discs and constituted approximately 65% of the area on the worn pin surfaces. Specifically, the mechanism was associated with regions of less metal transfer. The second wear mechanism was dominant on the worn disc, but less frequent on the pin surfaces, and constituted a much more severe wear mechanism. In these regions the surface consisted of grains which contained parallel grooves, and occasionally isolated pits, which did not cross grain boundaries and were perpendicular or within 30” to perpendicular of the sliding direction, Figs. 4-6. On the zirconia disc worn against the stainless steel the grooves tended to be broad and frequently deep, Fig. 5, whilst on the zirccnia pins worn against the bearing steel the grooves were more crack like, Fig. 6. Thus, the amount and severity of grooving corresponded to the amount and severity of the metal transfer. The grooving was clearly associated with the crystallographic orientation of the zirconia grains. In order to verity this the surface was further etched in HF. The tetragonal/monoclinic precipitates were found to be orientated along and perpendicular to the grooves,
326
W. M. Rainforth,
R. Stevens 1 Wear of magnesia partially stabilised zirconia
Fig. 5. Detail of the grooves on the disc surface. Note how they do not cross the grain boundaries.
Fig. 7. Detail of the grooves after etching the precipitate orientation.
in HF which shows
Fig. 6. SEM image using back-scattered electrons showing the worn Mg-PSZ pins. The grooves were much sharper and more crack like.
Fig. 7. Those grains which did not contain any grooves failed to show any ciearly discernible orientation relationship.
Fig. 8. TJZM image of a back thinned specimen showing the interface between zirconia (upper region) and transferred metal oxide (bottom left). Much of the cubic matrix has been removed. The precipitates showed both monoclinic and tetragonal symmetry.
3.4. Transmission electron microscopy
surface of the same sample and on surfaces where the transferred layer had been removed. The interface between the transferred layer and the zirconia was shown to be diffuse rather than clear cut, which infers some tribochemical reaction had taken place between the two. ~roughout these tests no discrete zirconia wear debris particles were ever found. However, the metal oxide was shown to contain varying amounts of zirconium ions in solid solution which is further evidence of a chemical interaction behveen metal and ceramic. The quantity of zirconia in solid solution indicated that the tribochemical wear mechanism must have constituted a significant fraction of the overall wear rate. The structure of the metal oxide was found to be a mixture of three phases, namely, Fe@,, an amorphous
3.4.1.
with the transfer layer present
A view of the interface between the transferred metal oxide and the zirconia is shown in Fig. 8 taken from an area with no grooves. The oblate spheroid shape of the precipitate appears to have been unchanged by the wear mechanism. Interestingly, some of the precipitates were shown to have retained the tetragonal phase symmetry, although the majority were shown to be monoclinic. Thus, in~mplete transfo~ation had occurred at the worn surface. The significance of this point will be discussed later. Virtually all the cubic matrix phase appears to have been removed at the outer zirconia surface. This result was also found in occasional areas elsewhere on the
W. M. Rainforth, R. Stevens / Wear of magnesia partially stabilised zirconia
oxide phase and a non-equilibrium body centred cubic oxide phase. The proportions of each phase varied between each test and counterface material, but in each case, the proportion of metallic ions in the oxide was identical to the proportion of metal atoms in the steel. This result and the presence of an amorphous phase indicates that the temperature was not high enough to generate the thermodynamically predicted metal oxides. 3.4.2. Smooth areas with the oxide removed Whilst the structure was difficult to image at the worn surface because of large residual strains, it was clear that the precipitates were mainly monoclinic (as shown by diffraction analysis), Fig. 9. Interestingly, the structure of these areas retained the same general morphology as the starting material. However, some areas showed a distortion in the precipitate geometry and orientation, as illustrated in Fig. 9. In addition, small rotations in the diffraction pattern were observed across small areas of the sample (less than 2 mm). No dislocations or subgrain boundaries were observed. Thus, the distortion was considered to be mainly elastic rather than as a result of plastic deformation. The depth over which any microstructural change could be detected was shown to be only of the order of 100-200 nm from the surface, in agreement with the X-ray results. In the region more than 200 nm from the surface some residual strain tended to obscure the image but no distortion of the precipitates was found. Interestingly, the majority of precipitates at this depth were tetragonal. 3.4.3. Regions where the grooving was observed Back-thinned samples from the disc were difficult to interpret because the surface was rough and thin foils
Fig. 9. TEM image of a back thinned specimen showing the near surface structure after removal of the transferred layer. The precipitates are mainly monoclinic as shown by the twinning, but there is considerable distortion of the structure, believed to be from residual elastic strain.
327
tended to break-up. This was also taken to indicate a high proportion of microcracks in this area. A general view of the sample taken from the pin worn against the bearing steel disc is shown in Fig. 10 with a detail in Fig. 11. The micrograph confirms that the cracks/grooves were orientated in a [lOO] direction, i.e. they were parallel to the long axis of one of the three precipitate orientations. The shapes of the precipitates in this region had remained unchanged, but were shown to be monoclinic, with no tetragonal phase positively identified. The depth to which transformation had occurred was greater than in those areas discussed above, although an exact measurement was difficult. Regions inbetween the major cracks showed extensive microcracking, Fig. 11, with all cracks aligned along the same [loo] direction as the larger cracks. Evidence was found of a transition between the fine microcracks and the coarse macrocracks. Thus, it would appear that the larger cracks had been formed as a result of microcrack coalescence. One additional feature of interest was found in these foils. Occasional three-body abrasive grooves were ob-
Fig. 10. TEM image of a back thinned specimen taken from an area showing grooving. The diffraction pattern shows a [lOO] type zone and demonstrates that the cracks run along a [lOO] direction.
Fig. 11. Detail from Fig. 10. Fine microcracks have initiated at the cubic/monoclinic interface which then coalesce to give the larger cracks in Fig. 10.
328
W. M. Rainforth, R. Stevens I Wear of magnesia partially stabilised zirconia
Fig. 12. Detail from a region just adjacent to Fig. 10 showing the interface between the normal structure of cubic/monoclinic (upper right) and a three-body abrasive groove (lower right). The diffraction pattern from the abraded region shows considerable rotation in the structure. Further work is required to fully evaluate the microstructural changes which have taken place.
served which had totally destroyed the precipitates and led to several other complex changes, as shown in Fig. 12. Whilst the details of the microstructure require further work to clarify, it was clear that the precipitate structure had been destroyed. Considerable rotation of the diffraction pattern occurred as the selected area aperture was moved across small distances. In addition, the sharp cubic and tetragonal diffraction spots were replaced by an arc of diffraction spots extending over 5-10”. This is indicative of a very fine grain size (dark field imaging indicated of the order of 10-30 nm) within the coarse (80 pm) cubic grain. Interestingly, no evidence of cracking was found in these areas.
4. Discussion 4.1. Smooth su$ace areas and metal transfer The extent of transformation at the worn surface as shown both by X-ray results and TEM results was low, substantially lower than that expected under abrasive wear conditions. The transformation zone was shown to be less than about 200 nm at the surface, whereas transformation zone widths adjacent to a crack in the same material would be expected to be up to 8 pm in size [20]. The TEM results demonstrated that where transformation from tetragonal to monoclinic had occurred it was restricted to about the outer 200 nm of the worn surface for the smooth regions, although it extended somewhat deeper in the regions containing grooves. The smooth areas which showed differential wear between different grains (‘grain relief) is typical of ceramics which exhibit mild wear. For example, Wallbridge et al. [15] have observed a similar effect for the sliding wear of alumina on alumina under the same test geometry as the present study. Wallbridge et al. [15] used the work of Steijn [21] to explain the origin
of the grain relief. Steijn [21] found that, for single crystal alumina, resistance to wear depended strongly on crystal orientation. Very high wear resistance was found when the rubbing direction was perpendicular to the c axis on the prismatic and rhombohedral crystal faces. However, brittle type wear occurred when rubbing was parallel to the c axis. Thus, wear results correlated qualitatively with the ease of plastic deformation. On the assumption that the differences in elastic modulus with crystallographic orientation were insufficient to explain the grain relief, Wallbridge et al. [15] suggested that those grains which were orientated for slip would wear less than those grains which were not orientated for slip. In the latter case the grain was considered to undergo a microfracture mechanism. The variation in hardness with crystal orientation in ceramics is well established [22, 231, and it is also apparent that the hardness is not an inherent material property but depends on the load used to make the indent [23]. In addition, chemo-mechanical effects can have a strong effect on hardness, whereby surface absorbed species alter the ease of crystallographic slip and therefore the hardness [24]. Such arguments can also be invoked to explain grain relief on the basis that water vapour can also provide a chemo-mechanical effect that alters the surface plasticity. Clear evidence has been presented in the present study which shows that no plastic deformation type wear mechanism had occurred in this region although significant residual elastic strain was evidently present. Moreover, no evidence of microfracture type mechanisms was found, and would have been clearly visible in these experiments. Rather, the wear mechanism appears to have been predominantly one of tribochemical wear. The cubic matrix was shown to have been preferentially removed and the interface between metal oxide and zirconia was diffuse. No discrete zirconia wear debris was found, the zirconia only being found in solid solution in the transferred layer and the free wear debris. Therefore, the wear rate is dependent on the chemical reaction rates as a function of crystallographic orientation. Whilst it is possible to envisage that the preferential removal of cubic phase by the tribochemical wear mechanism would increase the rate of transformation, it is considered that such an effect would be very surface specific (of the order of the precipitate size) and would be of less significance than the contact stresses. Further work is required to investigate this mechanism in other non-transforming ceramics, particularly alumina, since grain relief is often associated with very low wear rates and the mechanism may therefore determine the ultimate wear rate that is achievable. The observation that transformation was restricted to the outer 200 nm would most probably have been
W. M. Rainforth, R. Stevens / Wear of magnesia partially stabilised zirconia
a result of the temperature rises at the interface reducing
the driving force for transformation, rather than low contact stresses. Values are available in the literature for fracture toughness (and therefore extent of stress induced transformation) as a function of temperature for Mg-PSZs with different martensite start temperatures (MS) [25]. The M, value for the material in the current work would be around 113 K, which indicates that the K,, would fall to that equivalent to the base cubic matrix at around 250 “C, i.e. that very little transformation would be expected to occur above this temperature (although the actual M, temperature was not measured for the material used in this study). Thus, the observation of transformation at the outer surface, and to a greater extent in the cracked/grooved area (discussed below) suggests that the surface temperature must have been such that some driving force for transformation still existed. Such an estimate of temperature is in line with the non-equilibrium phase composition of the metal oxide, which would have been expected to transform to equilibrium oxides at higher temperatures. Indeed, heating the wear debris resulted in rapid transformation to the equilibrium structure at temperatures as low as 300 “C. In addition, this estimate agrees well with analytical calculations of the interface temperature made elsewhere [2], of measurements made by an implanted thermocouple [2] and of the microstructure at the worn stainless steel surface [19]. 4.2. Areas with groovinglcracking The grooving/cracking had promoted a higher wear rate than that associated with the grain relief. The grooving/cracking was clearly associated with the crystallographic orientation of the zirconia grains, as shown both by the etching experiments and the TEM results, with cracking being largely restricted to the [NO] directions. Similar parallel grooves and various shaped pits have also been reported by Hannink et al. [4] for the wear of Mg-PSZ tappet facing under lubricated wear. Whilst these authors did also ascribe orientation relationships based on the appearance of the groove/ pit, they did not make any attempt to explain the mechanism of formation. Three possible mechanisms can be envisaged for the formation of the grooves/cracks. Firstly, the cracks may have formed as a result of thermal expansion stresses, so-called thermal crazing. However, this mechanism can be rejected for two reasons. Firstly, thermal cracking would cross grain boundaries whereas the grooves did not. Secondly, tests conducted at higher sliding speeds tended to produce less grooving/cracking, which is the opposite to that expected from thermal cracking. The second possible mechanism is that transformation occurred cooperatively in bands (in a similar manner to the transformation bands found adjacent to a hardness
329
indent or a bend test specimen, as reported by Marshall and James [26] for example). Such bands would lead to a local raising of the surface thereby giving preferential wear in this region. However, many bands were either too deep or too sharp to have arisen from such a mechanism. Moreover, transformation was shown to occur throughout the surface rather than in bands. The final mechanism is one where transformation of tetragonal to monoclinic led to microcracking which, with the repeated application of shear stresses, coalesced to give macro surface cracks. Coalescence of these cracks would then transform the appearance to a groove. Very fine microcracks were found at the interface between precipitate and matrix which appeared to have coalesced to give larger cracks running through the cubic matrix and along precipitate interfaces (Figs. 10, 11). The interface between the tetragonal (or monoclinic) precipitate lies mainly on the (100) planes. In addition, the twin variants of the monoclinic are highly populated along (OOl),, so that maximum strain is generated along {loo}, [27]. Consequently, microcracks occur preferentially along this direction, explaining the initiation of grooves in grains with a [NO], orientation relationship with the surface before initiation in other orientations. Thus, the third possible mechanism of microcrack coalescence was believed to be responsible for the grooving/cracking. The greater severity of the mechanism when the zirconia was worn against the stainless steel counterface compared to the harder bearing steel disc suggests that the adhesive forces played a critical role in the mechanism. The greater rate of metal transfer when worn against the softer stainless steel counter-face compared to the harder bearing steel disc was considered to reflect the levels of the adhesive forces. The adhesive forces will have increased the tractions normal to the surface, thereby increasing the crack nucleation and propagation rates. Moreover, this explains why such grooving is not observed under abrasive wear conditions. Adhesive forces are low under abrasive wear conditions which ensures that the dominant forces are compressive rather than tensile, such that transformation imparts a beneficial effect. Clearly, with the absence of the constraint of untransformed material, the state of the surface forces is critical in determining whether transformation has a beneficial or detrimental effect. In addition to the differences in adhesive forces between the two test configurations, it is also probable that the flash temperature in the zirconia surface was higher for the pins compared to the discs since the pins were in continuous contact whilst the discs were in intermittent contact. The driving force for transformation would have been reduced at higher temperature thereby reducing the extent of a transformation driven wear mechanism.
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The current results are not considered to contradict the findings of either Birkby et al. [ll] or Fischer et al. [l]. Fischer et al. [l] found that the wear resistance of yttria-zirconias varied with the fourth power of the toughness, which suggests that the transformation of the tetragonal to monoclinic has a dramatic beneficial effect. However, the toughest ceramic tested was a 3YTZP which has a limited transformation zone width. The transformation zone width is a measure of the distance from a crack that stress assisted transformation occurs and would be substantially larger in the MgPSZ used in the current study, compared with the 3YTZP used by Fischer et al. [l]. Both Birkby et al. [l l] and Rainforth [2] have demonstrated that in tougher TZP (e.g. %Y-TZP) ceramics the transformation zone widths are appreciably larger and there is evidence of a transformation driven wear mechanism. Birkby et al. [ll] clearly demonstrated that such a transformation initiated mechanism does result in accelerated wear rates. Therefore, whilst transformation can impart improved wear resistance, especially in abrasive wear conditions, there is a maximum toughness in which the zirconia ceramic should be used in sliding wear. Frequently, the zirconia should not be used in the maximum toughness condition, but in the maximum hardness condition. However, the exact value of the toughness will depend on the surface stress system, in particular, the magnitude of the adhesive forces. 4.3. Abrasive grooves Occasional three-body abrasive grooves were observed on the zirconia surface which were presumed to have been formed from detached wear debris. TEM analysis of the grooves indicated that they were associated with dramatic microstructural change. The resulting structure apparently consisted of cubic phase with no positive identification of monoclinic or tetragonal. Considerable rotation of the diffraction pattern occurred over small distances, greater than could be possible by residual elastic strains. The precipitate structure had been completely destroyed and replaced by a nanocrystalline structure which was apparently cubic. This, therefore, strongly suggests that extensive dislocation flow had taken place. Temperature in itself could not be responsible since many hours at 1700 “C are required to dissolve the tetragonal precipitates and a three-body abrasive wear event is very short. However, the structure was extremely difficult to image and none of the subgrain boundaries could be imaged because of the residual strain. Interestingly, no cracking of any form could be found in the abrasive grooves even in grains that had undergone transformation initiated wear. However, the suggestion that compressive stresses from transformation are responsible for preventing crack formation is not appli-
cable in this case, rather the mechanism is apparently much more complex. Indeed, Swain and Hannink [14] have shown interesting and complex phase and morphological changes below the ground surface of a CeTZP. Further work is required to establish the actual structure of this region and the mechanism of formation.
5. Conclusions (i) A tribochemical wear mechanism between the steel (and transferred metal oxide) and the zirconia was shown to occur which promoted a smooth wear surface with grain relief. No evidence of microfracture or plastic deformation was found contradicting other theories of mild ceramic wear. (ii) In the smooth regions, minimal stress assisted transformation had occurred and therefore it played only a minor role in the wear mechanism. (iii) Greater quantities of transformation were found in regions where extensive metal transfer had occurred. Transformation of the tetragonal to monoclinic phase in these regions has been found to increase the wear rate. Microcracks were found to initiate at the matrix/ precipitate interface which coalesced to give macrocracks. (iv) The extent of the transformation initiated wear mechanism depended strongly on the adhesive forces, being more pronounced for a stainless steel compared to a bearing steel counterface. (v) The only circumstances where dislocation flow was identified was in three-body abrasive grooves.
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