Transition wear behavior of SiC-particulate- and SiC-whisker-reinforced 7091 Al metal matrix composites

Materials Science and Engineering, A147 ( 1991 ) 211-224

211

Transition wear behavior of SiC-particulate- and SiC-whiskerreinforced 7091 A1 metal matrix composites A. Wang and H. J. Rack* Department of Mechanical Engineering, Clemson University, Clemson, SC 29634-0921 (U.S.A.)

(Received December 10, 1990; in revised form May 1, 1991)

Abstract The wear behavior of unreinforced and reinforced 7091 AI, the latter containing either 20 vol.% SiC particulates (SiCp) or 20 vol.% SiC whiskers (SiCw), was studied as a function of sliding distance and sliding velocity under unlubricated conditions. At sliding velocities below 1.2 m s- 1, SiC reinforcement does not affect wear resistance. Wear debris produced from both the unreinforced and reinforced materials was predominantly metallic and was small in dimension and dark in color. The mechanism of wear under these conditions was surface-fatigue-related surface cracking. At sliding velocities greater than 1.2 m s 1, the wear rates of the reinforced materials were lower than for the unreinforced matrix. Both the unreinforced alloy and the SiC-reinforced composites exhibited elevated wear rates during the initial period of sliding, the mechanism of wear under these conditions, i.e. high velocity and short sliding distance, being controlled by subsurface-cracking-assisted adhesive transfer and by abrasion. During steady state sliding, these elevated wear rates were maintained by the unreinforced alloy, reduced wear rates being observed in the reinforced composites. The initial wear rates of the composites depend strongly upon reinforcement orientation, the highest wear rates being observed on the perpendicularly oriented SiCw composite. However, the steady state wear rates of the composites were generally independent of reinforcement geometry (particulate vs. whiskers) and orientation (perpendicular vs. parallel) with the exception of wear at 3.6 m s-~ where the parallel-oriented SiCw composite was superior.

1. Introduction Aerospace applications have generally provided the principal impetus for the development of light weight metal matrix composites (MMCs). Recently the requirement for light, strong and stiff materials has extended this interest in M M C s to many other fields [1-5]. Apart from high specific strength and stiffness, MMCs often possess excellent friction and wear properties. For exampie, ceramic-fiber- and particulate-reinforced MMCs have been employed in automotive or aircraft brakes [1-3] and in diesel piston engines [4, 5] to improve wear resistance, Several studies [6-8] suggest that MMCs under lubricated sliding against a smooth counterface exhibit superior wear resistance over unreinforced alloys. Under unlubricated dry conditions,

*Research associateandProfessor, respectively,

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however, complex and often opposite results have been reported by different researchers [9-14]. For example, Saka and Karalekes [9] found that the unlubricated wear resistance of a C u - A I : O 3 composite against a WC-Co-coated steel in a block-on-ring configuration was less than that of the unreinforced matrix alloy. You et al. [10] and Alpas and Embury [11] reported a similar detrimental effect of SiC particulate reinforcement in aluminum under dry conditions; a much improved wear resistance was achieved under lubricated sliding conditions. In contrast, Hosking et al. [12] and Milliere and Suery [13], utilizing block-on-ring and pin-on-disc configurations respectively, observed a significant enhancement in the unlubricated wear resistance of A120 3fiber-reinforced composites sliding against a steel ball bearing. Furthermore, the latter authors showed that this improvement increased with increasing reinforcement volume fraction and reinforcement size. Finally, Pan et al. [6] and © Elsevier Sequoia/Printed in The Netherlands

212

Wang and Rack [14], utilizing the pin-on-disc configuration, have recently reported that under limited loading conditions the steady state wear resistance of SiC-whisker-reinforced aluminum composites was more than an order of magnitude greater than that of the unreinforced matrix alloys, This investigation has extended our previous studies of the unlubricated wear resistance of SiC-reinforced aluminum to include the effect of sliding speed and sliding distance. The wear rates in an age-hardened high strength 7091 AI alloy, with and without SiC reinforcement, were correlated with microstructural observations; in particular, the effects of reinforcement geometry (whisker v s . particulate) and orientation (perpendicular v s . parallel) were considered to assist in establishing operative mechanisms and their relationships to specific tribological conditions.

cially aged at 125 °C for 24 h. Previous studies [16] have shown that this heat treatment results in maximum hardness, microstructurally consisting of a mixture of Guinier-Preston zones and ~/'. Mechanical characterization of this aged condition utilized Vickers microhardness measurements at 500 gf load, the values reported in Table 1 being averaged from five readings.

t~

2. Experimental methods Unreinforced 7091 A1 alloy (chemistry: 1.59 Cu; 2.28 Mg; 6.11 Zn; 0.40 Co; 0.07 Fe; balance, aluminum) and 7091-SiC composites containing either 20 vol.% grade 3 SiC particulates (SiCp) or F-9 SiC whiskers (SiCw) were fabricated by the powder metallurgy method [15]. Following vacuum hot pressing in the two-phase mushy (liquid plus solid) zone both the matrix alloy and the composites were extruded, the extrusion ratio being 11.3:1. Examination of the as-extruded microstructures (Figs. l(a) and l(b)), indicated that the SiC particulates and whiskers were uniformly distributed throughout the matrix, both tending to be closely aligned with the extrusion direction. Previous quantitative analysis of similarly processed extrusions showed that the SiC particulates had a size distribution ranging from 2 to 10 /~m, with a mean diameter of 5 # m [16], while the SiC whiskers had a uniform average diameter of 0.5 # m with a mean length-todiameter ratio of 6 [17]. Pin-on-disc wear specimens were prepared from the unreinforced and the SiCp-reinforced alloys such that the wear surface lay perpendicular (N) to the extrusion axis (Fig. 2). The SiCwreinforced samples were oriented with the whisker axis either perpendicular (N) or within and parallel (P) to the sliding direction. Prior to wear testing, all specimens were solution treated at 475 °C for 2 h, water quenched, naturally aged for 6 days at 25 °C and then artifi-

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213 TABLE 1

TABLE 2

Vickers hardness a of the materials studied

Initial wear rates ( × 10 - 3 mm 3 m - ~) of the materials studied

Material

Hv

7091 7091-20vol.%SiCp 7091-20vol.%SiC,~(P) b 7[)91-20vol.%SiCw(N) b

103 + 1 156+4 156_+4 167 _+ 2

aUnder 0.5 kgf load. bp, parallel orientation; N, normal orientation.

Velocity ( m s -~)

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(I.36 0.72 1.20 1.80 3.60

2.91 2.35 4.70 30.5 50.0

2.31 1.00 2.30 18.8 28.1

3.85 3.01 2.30 22.7 31.1

2.08 1.99 2.22 8.58 13.2

~'N, normal orientation; P, parallel orientation.

Wear tests were undertaken under constant pressure (0.43 MPa) at sliding velocities between 0.36 and 3.60 m s-1, the pin specimen being mounted in a vertical sample holder against a rotating stainless steel disc (17-4 PH stainless steel; hardness Hv0.5kgf= 345). All specimens followed a single track 115 mm in diameter, with the tangential force (friction force) during sliding being continuously monitored by a force transducer attached to the specimen holder. Finally, the bulk temperature rise within the pin specimen was monitored by spot welding a type K thermocouple on selected specimens approximately 2 mm distance from the contact surface, Prior to testing, both the 17-4 PH steel discs and the aluminum-based pins were ground in situ against 240 grit SiC paper, and then cleaned in acetone. All wear tests were carried out in air at 25 °C, relative humidities being between 50 and 70%, in an incremental manner, i.e. 1000 rotations (361 m) per increment and 10 000 rotations (3610 m) in total. After each increment, the pin specimen was removed from the tester, cleaned in acetone and weighed with a balance to an accuracy of 0.10 mg, and then remounted in the wear tester at the same location. Volume losses were calculated by dividing the mass losses by the density of the samples (2.70 and 2.82 g cm -3 for matrix and composite respectively), Both surface and subsurface observations of the pins and discs were performed by scanning electron microscopy. Subsurface observations utilized longitudinal (parallel to wear track) crosssections of the pin specimens; prior to sectioning, the surfaces of the pins were coated with a thin layer of nickel for edge retention. Wear debris was also collected after sliding for 1000 rotations and 10 000 rotations respectively, the composition and morphology of these wear particles being revealed by X-ray diffraction (Cu K a radiation at 35 kV) and scanning electron microscopy,

Finally, a Talysurf profilometer was used to characterize the roughness of wear tracks of selected disc surfaces. The direction of measurements was perpendicular to the wear tracks. 3. Experimental results

3.1. Wear rate, sliding distance and sliding velocity relationships Damage accumulation during unlubricated sliding was a function of sliding velocity and reinforcement (Fig. 3). For the unreinforced 7091 A1 alloy, a linear relationship between volume loss and sliding distance was observed at all velocities. However, for the reinforced composites an initial unstable period of sliding was observed, followed by a steady state period during which the volume loss increased linearly with increasing sliding distance. Tables 2 and 3 summarize both initial and steady state wear rates for all materials tested under different conditions. Because of the nonlinear nature of the initial unstable period of sliding observed in the reinforced composites, the initial wear rate in this instance was defined as the average volume loss per unit sliding distance during an initial sliding distance of 361 m, while the steady state wear rate was ascertained from the slopes of the linear part of the wear curves shown in Fig. 3. The initial wear rates, when plotted against sliding velocity, exhibited two different regimes (Fig. 4) with an apparent wear rate minimum at a sliding velocity between 0.72 and 1.2 m s- 1 being observed for all materials. This minimum was followed by a dramatic increase in wear rate as the sliding velocity increased to 1.8 m s -~. At sliding velocities below 1.2 m s-~ the wear resistance of the unreinforced alloy was superior to that of the perpendicularly oriented SiC w composite, the reverse effect being observed at sliding velocities

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above 1.2 m s- 1. The wear resistances of the SiCp and parallel-oriented SiCw composites were superior to those of the unreinforced alloy over the entire sliding velocity range examined (Table 2). In general, the initial wear resistance of the SiCp composite was superior to that of the SiCw composite in the perpendicular orientation, while the parallel orientation was superior to the perpendicular orientation when the same reinforcement, i.e. SiCw, was considered. The steady state wear rate, when plotted against sliding velocity (Fig. 5(a)), again exhibited two different regimes for the unreinforced matrix alloy. However, for the reinforced composites (Fig. 5(b)), increasing sliding velocity resulted in a decreasing steady state wear rate. At sliding velocities less than 1.2 m s-1, the unreinforced and reinforced alloys exhibited similar wear rates. At

sliding velocities greater than 1.2 m s- 1, the reinforced composites exhibited steady state wear rates an order of magnitude lower than those of the unreinforced matrix alloy. Finally, there appeared to be no systematic effect of reinforcement geometry and orientation on steady state wear rate at velocities below 3.6 m s- 1; at 3.6 m s-~ the parallel-oriented SiCw composite was superior (Table 3 and Fig. 5(b)).

3.2. Friction andfriction-inducedtemperaturerise Typical friction and friction-induced specimen temperature rise traces are shown in Figs. 6(a) and 6(b). These represent respectively the lower (less than 1.2 m s- 1) and upper(greater than 1.2 m s- l) velocity regimes examined in this study. Within the lower velocity regime, the coefficient of friction and the temperature rise remained at

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TABLE 3

lower constant values; in contrast, constant values of the coefficient of friction and temperature rise were maintained by the unreinforced alloy throughout the entire Period of sliding. Figure 7 shows a plot of the steady state coefficient of friction v s . sliding velocity. Similar to the steady state wear rate-sliding velocity curves (Fig. 5) the SiCp-reinforced composite exhibited a continuous decrease in the coefficient of friction with sliding velocity while the unreinforced alloy showed a minimum at 0.72 m s-1 followed by a dramatic increase at 1.2 m s 1. Finally, Fig. 8 shows a plot of specimen temperature v s . sliding velocity during steady state sliding. Both the reinforced and the unreinforced alloys showed a continuous increase in temperature, the rate of increase in temperature being higher for the unreinforced alloy than for the reinforced composite.

Steady state wear rates ( × 10- 3 mm 3 m - i ) of the materials studied Velocity

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3.64 1.99 2.22 1.92 1.11

(m s- i)

aN, normal orientation; P, parallel orientation.

constant values with increasing sliding distance, both tending to be higher for the reinforced composites than for the unreinforced alloy. Within the higher velocity regime, the reinforced composite exhibited initially greater values of the coefficient of friction and temperature rise, followed by

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3.3. Scanning electron microscopy a n d X - r a y diffraction R e g a r d l e s s o f t h e test m a t e r i a l s a n d e x p e r i mental conditions, the worn specimen surfaces showed two morphologies: one representing mild w e a r a n d t h e o t h e r s e v e r e wear. T h o s e s a m p l e s w h i c h e x h i b i t e d w e a r r a t e s o n t h e o r d e r o f 10 -3

Fig. 9. Scanning electron micrographs of worn surfaces of (a) the unreinforced 7091 alloy and (b) 20vol.%SiC,reinforced composite (normal orientation) after sliding for 3610 m at 0.36 m s- ~.

3 m - t , i.e. the u n r e i n f o r c e d a n d r e i n f o r c e d s p e c i m e n s d u r i n g b o t h initial a n d s t e a d y state sliding at v e l o c i t i e s less t h a n 1.2 m s - 1 a n d t h e r e i n f o r c e d c o m p o s i t e s d u r i n g s t e a d y s t a t e sliding mm.

217

at all velocities, showed smooth mild-wear-type surfaces (Fig. 9). Further detailed study of these mild wear tracks indicated that they exhibited, independent of reinforcement content, regions of surface cracking, flaking being occasionally observed between surface cracks (Fig. 10). Scanning electron microscopy of wear debris collected from such worn samples (Fig. 11) revealed particles that were irregular in shape and small in dimension (about 3 k~m). X-ray diffraction results (Figs. 12 and 13) indicated that these particles were predominantly metallic, containing either a-Fe or aluminum. No evidence for oxide formation was detected, the presence of SiC within the wear debris being representative of the reinforcement phase. Those specimens which exhibited wear rates on the order of 10 -2 mm 3 m -1, i.e. the unreinforced specimens during both initial and steady state sliding and the reinforced composites during the initial period of sliding, all at velocities greater than 1.2 m s-1, showed roughened severe-wear-

type surfaces (Fig. 14). Furthermore, these specimens showed surface grooving (marked g), tearing (marked t) and flaking (marked f) (Fig. 15). For the reinforced composites, fragmentation

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Fig. 15. High magnification scanning electron micrographs of worn surfaces of (a) 7091 A1 and (b) SiC.-7091 composite after sliding for 361 m at 3.60 m s-]. Fragments are present within the wear scars on the surface of the reinforced composite.

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Fig. 14. Scanning electron micrographs of worn surfaces ot (a) the unreinforced 7091 alloy and (b) 20vol.%SiCwreinforced composite (normal orientation) after sliding for 361 m at 3.60 m s-J.

or pulverization (marked p) was also observed at the bottom of the flaked areas. T h e bulk of the wear debris (darker features marked p in Fig. 16), was in the form of large flakes and chips with

linear dimensions from 10 to ,200 /~m. T h e brighter features (marked d) were debris produced from the steel disc material, the observed contrast being d u e to the large difference in atomic number between iron and aluminum. Scanning electron micrographs of longitudinal cross-sections of the unreinforced alloy worn in mild and severe wear regimes are shown in Fig. 17(a) and Fig. 17(b) respectively. W h e n sliding in the unreinforced alloy occurred in the mild wear regime, plastic deformation of the near-surface region caused locally isolated void formation around secondary particles (marked p), these having been previously identified as dispersions of C o - F e - r i c h intermetallics [17] introduced during vacuum hot pressing [15]. In contrast, when sliding occurred in the severe wear regime, these voids were linked into cracks (marked c) within a severely deformed layer extending to

219

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Fig. 16. Scanning electron micrographs of wear debris produced from (a) 7091 A1 and (b) SiCw-7091 composite after sliding for 361 m at 3.60 m s 1; p and d represent oin and disc debrisrespectively.

approximately 20 ktm below the wear surface (Fig. 17(b)). Scanning electron micrographs of longitudinal cross-sections of the perpendicularly oriented SiCw composite, worn in the mild and severe regimes, are shown in Fig. 18(a) and Fig. 18(b) respectively. When sliding in this material occurred in the mild wear regime, plastic deformation caused realignment of the SiC whiskers along the sliding direction and local fracture, as noted by arrows, of the whiskers (Fig. 18(a)). When sliding occurred in the severe wear regime, fracture was no longer limited to the whiskers, cracking again extending into the matrix (Fig. 18(b)). Subsurface shear cracking was also observed on the cross-sections of the SiCv and the parallel-oriented SiCw composites worn in the severe wear regime (Fig. 19). These shear cracks were usually hundreds of micrometers long and

Fig. 17. Scanning electron micrographs of longitudinal cross-sections of the unreinforced 7091 alloy after sliding for 3610 m at (a) 0.36 m s- ~ and (b) 3.60 m s t

propagated preferentially along reinforcementmatrix interfaces(Fig. 20). Finally, examination of worn 17-4 PH disc surfaces during and after each test revealed that, whenever wear occurred in the mild wear regime, a thin layer consisting of dark-colored and loosely compacted wear debris was formed on the 17-4 PH wear tracks, regardless of the composition and microstructure of the mating pin material. Observations also indicated that, whenever wear occurred in the severe wear regime, a shiny surface free of fine loose debris but containing larger transferred fragments was found on the 17-4 PH counterpart. Examples of those two morphologies are shown in Fig. 21(a) and Fig. 21(b) respectively. Profilometry traces of those two surfaces (Fig. 22) further revealed that the

220

Fig. 18. Scanning electron micrographs of longitudinal cross-sections of the perpendicularly oriented SiCw cornposite after sliding for 361 m at (a) 0.36 m s- ~and (b) 3.60 m s 1.

Fig. 19. Scanning electron micrographs of longitudinal cross-sections of (a) the perpendicularly oriented SiCp composite and (b) the parallel-oriented SiCw composite after slidi n g f o r 3 6 1 m a t 3 . 6 0 m s -~.

17-4 PH stainless steel surface had been smoothed under mild wear conditions, while the surface had been roughened under severe wear conditions,

suggests that surface mechanical fatigue was the predominant wear mechanism. Earlier studies [21-23] of dry sliding wear processes have suggested that wear debris is principally produced from a mechanically mixed surface layer. Further, for steady state wear processes to persist under these conditions, removal of this mixed layer must be accompanied by replenishment of fresh material from the substrate. This can be achieved only when the mechanically mixed layer is discontinuous, providing some limited degree of fresh metal-tometal contact and material transfer. Therefore the integrity and areal fraction of the mixed layer will have a strong effect on wear rates. In fact, experiments performed on an ionimplanted system [24] revealed that reduced friction and wear persisted for times longer than needed for the wear scar to exceed the depth of

4. Discussion

4.1. Mechanisms ofmiMwear The foregoing results have shown that at sliding velocities below 1.2 m s -1 the wear behavior of the reinforced composites closely resembles that of the unreinforced alloy. The morphologies of either worn surfaces or wear debris were similar to those previously described to be characteristic of mild wear [18-20]. However, in contrast to the authors of these papers, who reported that mild wear debris consisted of oxide particles, the present results indicate that the bulk of the wear debris was metallic. This observation

221

Fig. 20. An overall view of a subsurface shear crack formed in the parallel-oriented SiC~ composite after sliding for 361 m at 3.60ms-L

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50p.m Fig. 22. Profilometry traces of the 17-4 PH stainless steel disc surfaces (a) prior to sliding, (b) after sliding for 361 m at 0.36 m s i and (c) after sliding for 3 6 1 m at 3.60 m s against the 7091-SiC~ (normal orientation) composite.

Fig. 21. Scanning electron micrographs of the 17-4 PH stainless steel disc surfaces after sliding against the SiC~-7091 composite for 361 m at (a) 0.36 m s t and (b)

3.60ms-'.

the implanted layer. Others [21, 22] observed that when the hardness of the surface mixed layer is greater than that of the substrate the former can sink into the latter, yielding a smooth contact surface and low wear rate; when the mixed layer is softer than the substrate the former stands proud on the latter, forming islands and causing surface

roughening and a high wear rate. However, the present experiments suggest that the occurrence of either a smoothed or roughened surface depends upon experimental, tribological, conditions rather than on the relative hardness between the mixed layer and the substrate alone. For example, with an initial sliding distance of 361 m, a smoothed surface was observed on the steel disc sliding against a n SiCp composite at 0.36 m S-i (Fig. 21(a)), while a roughened surface was found when sliding against the same material at 3.60 m s-~ (Fig. 21(b)). Although the composition, microstructure and mechanical properties of the surface mixed layer in the present study are currently unknown and are the subject of continuing study, qualitative comparisons based on the presence of the SiC reinforcement are possible. It is obvious that the

222 surface layer containing SiC reinforcement is expected to be harder and more brittle than that without SiC. The tendency for surface cracking is therefore greater for the former than for the latter (Figs. 9(a) and 9(b)). This may explain why the reinforced composites did not improve the wear resistance under mild wear conditions, Since, under mild wear conditions, the characteristics of the surface mixed layer determine wear rates, the initial microstructure of the bulk material should have little effect on steady state wear rates. Indeed, the SiCp- and SiCw-reinforced composites generally exhibited similar wear rates under mildwearconditions, 4.2. M e c h a n i s m s o f severe wear

Severe wear occurred at sliding velocities greater than 1.2 m s- 1 under both initial (run-in) and steady state conditions for the unreinforced matrix alloy and during initial run-in for the reinforced composites. This type of wear was represented by higher wear rates and larger metallic wear debris. Severe transfer of pin material onto the disc surface was observed, indicating the inability of the surface mixed layer on the pin to protect the pin substrate. In fact, plastic deformation under severe wear conditions was so intense that cracking was observed below the contact surface, which makes it impossible for any stable mixed layer to form on the substrate. Under these conditions, adhesive transfer assisted by subsurface cracking is expected to be the rate-controlling process for wear. According to the simple adhesive theory of wear [25], wear rates should be proportional to the inverse of hardness. This prediction is not consistent with the results of this study. During initial run-in, higher initial wear rates were actually observed with the harder perpendicularly oriented SiC w composite than with the softer SiCp and parallel-oriented SiCw-reinforced composites (Fig. 4(b)). In addition, examination of wear debris revealed that both pin and disc debris were present during this initial sliding period, with the morphologies of some debris being similar to those produced by abrasion, e.g. long chips [17]. Therefore abrasive wear of both pin and disc materials must have occurred during the initial run-in period, the abrasiveness of the SiC reinforcement clearly being responsible for wear of the discs. Furthermore, those disc surfaces which had undergone severe abrasion and roughening would, in turn, be expected to cause more abra-

sion on the composite pin surfaces. This mutual abrasion process has been found on similar composite-steel systems [10, 14] and the initial wear rates were indeed very sensitive to the surface finishing of both contact bodies [14]. In summary, both adhesive transfer and abrasion were involved as wear mechanisms in severe wear. 4.3. M e c h a n i s m s o f wear transitions

The volume loss-sliding distance curves had two different forms for the unreinforced and reinforced alloys under different conditions. For sliding of the unreinforced matrix alloy, linear relations between wear and sliding distance were established almost at the start of the test at all sliding velocities used. For sliding of the reinforced composites, steady state sliding was achieved after an initial run-in period. Earlier studies on the transition wear behavior of metals were principally concerned with the effect of surface oxide film formation or breakdown on wear rates and wear mechanisms of steels. It was thought that the severe-to-mild wear transition occurs only when frictional heating is high enough for a sufficiently thick oxidized layerto form on a metal substrate. Recently it has been found that oxidation is not the only cause for mild-severe wear transition. For example, Sawa and Rigney [26] showed that some transitions in friction and wear in dual-phase steels can be directly related to changes associated with transfer and mechanical mixing, as well as to oxidation. Severe-to-mild wear transitions were also observed by Oike et al. [27] on rubbing nontransition metals when fine ceramic particles were introduced into the sliding system. They reported that ceramic particles entrapped between two metal bodies promote initially high wear rates, but reduced steady state wear rates. These authors' findings are particularly important to the present composite-steel system because loose wear debris containing SiC reinforcement may act in the way similar to the ceramic particles used by Oike et al. [27]. Indeed, the bulk of the wear debris produced by mild and severe wear was metallic in nature and fine debris was also observed on wear tracks when mild wear occurred. Therefore the transition wear phenomenon observed here is believed to be primarily due to the transition of one type of mechanical wear to another. Friction and temperature rise measurements

223

suggest that friction-induced thermal softening may also play a role in wear transitions at higher velocities. In fact, above a critical velocity of 1.2 m s- ~, plastic deformation of the subsurface was so severe that shear cracks were observed. Once this occurred, any surface layer, whether mechanically mixed and/or chemically mixed, would be destroyed. Consequently, fresh metal-to-metal contact and transfer would be significantly enhanced, leading to severe wear. The conditions for subsurface shear localization and cracking have been considered by Rosenfield [28], who proposed that under certain sliding conditions plastic deformation will be localized within a region below the contact surface where the yield strength is lower than that of the surroundings. This localized plastic flow may lead to shear instability and fracture of this region if the maximum driving force on the localized shear plane exceeds a critical value. According to Rosenfield's shear instability model, the maximum driving force for shear fracture lies at the subsurface, some distance away from the contact interface, and the critical condition for subsurface shear instability is favored by a higher coefficient of friction, heavier loading and a steeper gradient of shear strength along the depth of deformation. The present results indicate that reinforcements mitigate shear instability. Although a cornplete understanding of this phenomenon is lacking at this time, it is proposed that this phenomenon may be associated with a reduced coefficient of friction at high velocities, and with the increased high temperature stability of aluminum alloys that occurs with the incorporation of ceramic reinforcements. For example, previous studies have shown that high-temperature deformation resistance of aluminum alloys increases with increasing SiC reinforcement [29, 30]. Consequently, the composite's shear strength is less likely to be reduced by frictional heating than is that of the unreinforced alloy.

Specifically, the following conclusions can be drawn from this study. ( 1 ) Wear rates of both unreinforced and reinforced materials were of the same order of magnitude when the sliding velocity was less than 1.2 m s- J, the mechanisms of wear under these conditions being surface-fatigue-related surface cracking. Wear debris produced from both types of materials was small in dimension and dark in color, the bulk of the debris being metallic. Under these conditions, i.e. low sliding speeds, SiC reinforcement does not affectwearresistance. (2) At sliding velocities greater than 1.2 m s- ~, the wear rates of the reinforced materials were lower than for the unreinforced matrix. Both the unreinforced alloy and the SiC-reinforced composites exhibited elevated wear rates during the initial period of sliding, the mechanisms of wear under these conditions, i.e. high velocity and short sliding distance, being controlled by subsurface-cracking-assisted adhesive transfer and by abrasion. During steady state sliding, however, these elevated wear rates were maintained only by the unreinforced alloy, while much reduced wear rates were observed with the reinforced composites. (3) The initial wear rates of the composites depend strongly upon reinforcement orientation, the highest wear rates being observed on the perpendicularly oriented SiCw composite. However, the steady state wear rates of the composites were generally independent of reinforcement geometry (particulate vs. whiskers) and orientation (perpendicular vs. parallel) with the exception of wear at 3.6 m s ~ where the parallel-oriented SIC,, composite was superior. Acknowledgment The authors wish to acknowledge the financial support of the Clemson Center for Advanced Manufacturing under F. Paul, Director. References

5. Conclusions Two forms of transition wear phenomena have studied, i.e. sliding-distance-related transition and sliding-velocity-related transition. Both transitions were caused by the change from one type of mechanical wear to another, the formation

been

and breakdown of surface mechanically mixed layers being responsible for these transitions.

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