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The role of subsurface deformation and strain localization on the sliding wear behaviour of laminated composites
Wear, 146 (1991)
285
285-300
The role of subsurface deformation and strain localization on the sliding wear behaviour of laminated composites A. T. Alpas
J. D. Embury
(Received
June 29, 1990; accepted
November
9, 1990)
Abstract The sliding wear behaviour of laminated composites of copper and amorphous Ni7&3i10Bi2 fabricated by diffusion bonding was investigated. Unlubricated sliding tests were performed using a block-on-ring type wear machine in order to provide a direct comparison of the wear mechanisms in the metallic glass (Ni78Si10B12) layers and polycrystalhne matrix (copper) and to determine the influence of the metallic glass layers on the overall wear resistance of the composite. Wear in the copper layer was characterized by extensive plastic deformation and by the presence of subsurface cracks. These cracks originated at the shear bands within the plastic zones under the contact surfaces and caused delamination of the material adjacent to these surfaces. The wear of metallic glass layers also involved the locahzation of plastic deformation into shear bands. However, due to the high hardness and fracture strength of Ni&3i10B12,its wear resistance was higher than the copper matrix. Localized strain and temperature gradients generated during sliding contact led to the crystallization of amorphous layers. Metallic glass layers were effective in increasing the wear resistance of the composite by supporting the load with less defo~ation and thereby obstructing the damage process initiated in the copper layers.
1. ~troduction
The production of wear debris is an extremely complex process involving mechanical, thermal and chemical interactions between the surfaces of materials in sliding and rolling contact. Under the service conditions the range of local pressures and temperatures generated during the wear may be very broad depending on the operating parameters such as applied load and sliding velocity, and on the detailed geometry of contacting surfaces [l-3]. The development of laminated composite materials provides a method of controlling the geometry of contact and also creates a possibility of designing microstructures with enhanced wear resistance baaed on the control of intrinsic mechanical properties and spatial distribution for the constituent phases. Some principal features of abrasive and adhesive wear resistance of
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laminated composites were discussed by Hornbogen [4] and Zum-Gahr [5]. In these studies, consideration was given to the description of the wear behaviour of lamellar solids in terms of rules of mixtures. The role of plastic deformation in copper-silver and copper-tin type laminated structures were investigated by Moore and Douthwaite [6] and Kennedy et aZ. [ 71. In the present work laminates of amorphous Ni,$i,,B,, and polycrystalline copper were fabricated by diffusion bonding [8]. It is known that metallic glasses exhibit a useful combination of mechanical properties such as high hardness (5-10 GPa) and fracture strengths that approach theoretical cohesive strength [9, 101. However, these materials do not possess a mechanism for distributing plastic deformation and at ambient temperatures plastic flow proceeds by nucleation and propagation of inhomogeneous shear bands [ 11, 121. Although this behaviour limits the tensile ductilities of metallic glasses to low values they show ductilities comparable with those of crystalline alloys when deformed by uniaxial compression and cold rolling [ 131. These properties make metallic glasses of interest for tribological applications and their friction and wear behaviour have been the subject of several studies that were reviewed by Lee and Evetts [ 141 and Kishore et al. [15]. It has been shown that amorphous alloys generally show lower values of coefficients of friction and wear rates [ 16,. 171 than their crystalline counterparts and this has been attributed to surface oxidation, dynamic recrystallization events taking place during sliding and abrasive contacts [ 14, 15, 18, 191. This study aims to clarify the role of metallic glass layers on the wear resistance of laminated composite materials by investigating the rate controlling wear mechanisms in Ni78Si10B12layers, and in the polycrystalline copper matrix, in relation to the deformation processes occurring during sliding wear. 2. Experimental
procedures
Laminated composite
structures were fabricated using thin ribbons
(t =57 + 2 pm) of amorphous Ni78Si‘10B r2 alloy and annealed commercial
purity copper sheets of various thicknesses. Metallic glass ribbons were produced using a melt spinning process (Vacuumschmelze, F.R.G.); their tensile strength and Vickers hardness were 2100 MPa and 9100 MPa respectively. The copper sheets had a tensile strength of 220 MPa and hardness of 670 MPa. To produce laminated composites, surfaces of metallic glass ribbons were electroplated by copper (approximately 100 pm thick). The coated ribbons were then sandwiched between annealed copper sheets and diffusion bonded in air under a pressure of 200 MPa at 560 K. Details of the diffusion bonding process are described in ref. 8. The cross-section of the composite is shown in Fig. 1. X-ray diffraction and scanning diffraction calorimetric (DSC) analyses indicated that the metallic glass layers completely maintained their non-crystalline structure during diffusion bonding. Dry sliding wear tests were performed using a block-on-ring type wear rig. Specimens in the form of 5 x 10 x 10 mm3 rectangular blocks were worn
287
Fig. 1. Microstructure of amorphous Ni78Sl‘,0B 12-Cu laminated composite. The featureless layers are metallic glass, the electrodeposited copper coatings have a columnar structure, and the annealed copper sheets show equiaxed grains. Fig. 2. Testing configurations of laminated composites. The slider moves (a) parallel and (b) perpendicular to the laminates.
against an SAE 52100 bearing steel slider ring (diameter, 38 mm). These tests were conducted such that the slider moved in a direction either parallel or transverse to the laminates as shown in the schematic diagram in Fig. 2. A normal load of 80 N and a sliding velocity of 0.1 m s-’ were used. All the experiments were performed at room temperature (21 “C) and in laboratory air (approximately 55% humidity). The volume losses during wear were determined following the procedures described in ASTM standard G7783. The morphologies of the wear tracks, loose debris and the microstructures of the material under the worn surfaces were examined by a Philips 551 scanning electron microscope.
3. Results Results of wear tests performed on laminates containing various volume fractions of Ni&i10B12 layers are shown in Fig. 3. The figure indicates that the lowest wear losses occurred in 100% metallic glass specimens tested in monolithic form and the incorporation of metallic glass layers into the copper matrix resulted in a reduction in the amount of material lost during wear. No significant difference could be detected between the wear losses of laminates lying parallel and transverse to the sliding direction so that the curves in this figure represent combined results of tests carried out using both configurations shown in Fig. 2. The volumetric wear rates of composites (i.e. the volume of material removed per unit sliding distance) are calculated from the slopes of the curves in Fig. 3 and these are shown as a function of vohune fraction of Ni&i10B12 in Table 1. The wear rates of the composites decrease rapidly with the addition of small volume fractions of metallic glass and the wear resistance of the composite containing 20% Ni,,Si,,,B,z becomes as good as that of the 100% metallic glass samples. The beneficial effect
289
Figures 4(a) and (b) show differences in the topographies of the worn surfaces of metallic glass and copper layers. It can be seen that the rate of material removal was higher in copper so that the top portions of Ni78Si10B12 layers protruded from the contact surfaces. The copper layers were characterized by severe plastic deformation and surface damage compared with metallic glasses whose surfaces appear to be smoother and relatively featureless at low magnifications. The cavities and plateau regions formed within the wear tracks of copper layers were wider than the total thickness of metallic glass layers. Observations made on cross-sections parallel to the sliding direction (Fig. 4(c)) indicated that large plastic strain gradients were developed in the copper grains adjacent to the contact surfaces as revealed by the shear displacements at the diffusion bonded Cu-Cu interfaces. Plastic flow in the copper matrix was impeded by Ni78Si10B12 layers and a large number of voids were created in front of metallic glass-Cu interfaces. These observations underline two salient features of the role of Ni78Si10B12on the sliding wear behaviour of the composite. These layers serve as load carrying constituents and they also act as barriers to the propagation of the damage process. The differences in surface topographies reflect the difliculty of producing loose debris in metallic glass layers. To clarify the micromechanism of debris
Fig. 4. Worn surface topographies of composites: (a) sliding direction is parallel to laminates; (b) sliding direction is perpendicular to laminates; (c) cross-section of composite psrallel to the slidiig direction.
290
formation in the composite, detailed metallographic examinations were performed on both matrix and Ni78Si10B12 layers. It is of value to examine these results in two separate sections. 3.2. Metallography of copper layers The microstructure of copper close to the wearing surfaces was composed of highly deformed and distorted grains that assumed a “tear drop” shape with their narrow edges lying almost parallel to these surfaces (Fig. 5). The deformed zones had a fairly uniform thickness of 60-70 ,um and included 2-5 pm grams which were initially equiaxed (Fig. 1). The plastic strain distribution due to this deformation can be estimated using a method proposed by Dautzenberg and Zaat [ 2 1 I. Shear angles of grain boundaries (and diffusion bonded interfaces) were measured at various cross-sections and these are related to equivalent plastic strain 1221. The variation of plastic strain E as a function of depth below the worn surfaces 2 is plotted in Fig. 6(a). At the contact surfaces, strains as high as 30 were reached; these decreased
Fig. 5. Cross-section of copper layers parallel to the sliding direction: (a) microstructure of the deformation zones under the worn surfaces and the subsurface cracks propagating parallel to the worn surfaces; (b) fracture and delamination of material adjacent to the contact surface.
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40
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( pm)
450
400
2 I350 z ii F v) 300
250 0
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10
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STRAIN Fig. 6. (a) Microhardness and equivalent strain vs. depth below the worn surface of copper layers; (b) stress-strain curves for deformed copper layers.
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to levels of about 2-3 at 10 pm below the surfaces. Figure 6(a) also shows the microhardness gradients generated during wear. Microhardness measurements were made (at a load of 50 gf) on tapered sections sectioned at an angle of 10” to the worn surfaces. Flow stresses were calculated from the microhardness data using a simple expression derived by Marsh [ 231 and a stress-strain curve for worn copper layers is presented in Fig. 6(b). The curve exhibits an initial stage of work hardening, but at large strains (er3), the curve flattens out, indicating either the exhaustion of the work hardening capacity or local recovery mechanisms in the grains closest to the contact surfaces. Figure 7(a) shows the microstructure of material under the wear track on a plane perpendicular to the sliding direction (and to the sections in Fig. 5). Within the plastically deformed zone, two orthogonal sets of shear bands, each inclined at about f35” to the wear surface, were developed. These bands crossed the grain boundaries without deviation except those which had undergone rotation and re-orientation (Fig. 7(b)). With increasing strain (i.e. on approaching the surface) the shear bands increased in number and the spacing between them became smaller. The intersection of these bands
(al
@I
cc>
Fig. 7. Cross-sections of copper layers transverse to the sliding direction: (a) structure of shear bands aligned 35” to the contact surface; (b) heavily deformed grains acijacent to the contact surface; (c) crack nucleation at the intersection of shear bands.
293
provided nucleation sites for crack formation as shown in Fig. 7(c). Cracks were also nucleated at some grain boundaries (Fig. 7(b)), possibly owing to the incompatibility effects between adjacent grams which were subject to different degrees of rotation during wear. Scanning electron microscopy observations also revealed that severe subsurface damage was often associated with deIamination parallel to the contact surface. Delamination events led to the formation of loose plate-like debris whose thicknesses (5-20 pm> were in close agreement with the loci of subsurface cracks (F’ig. 5@)). Thus, wear in copper layers proceeds by a damage accumulation process which involves the generation of large strain gradients below the contact surfaces. Under the influence of these subsurface strains, copper grams adjacent to the wearing surfaces lose their ability to sustain work hardening and become subject to an inhomogeneous deformation in the form of shear band formation. Cracks nucleated at these bands and at the grain boundaries propagate parallel to the contact surfaces and result in the creation of debris particles from the contact surfaces.
Wear tracks of Ni7$i10B1a were characterized by a lower level of damage and high magnification microscopic examinations were required to observe the features of the worn surfaces. These involved fine parallel grooves of approximately 0.5 pm depth and 1-2 pm width which were elongated in the direction of sliding (Fig. 8(a)). These grooves were indicative of an abrasive type of wear process. In contrast, investigations on the cross-sections parallel and perpendicular to the sliding direction (F’igs. 8(b) and (c)) revealed the development of a plastically deformed zone under the sliding contact surfaces. The deformation zone was almost continuous ~oughout the layers and penetrated to a depth of 10 pm which is about 6-7 times smaller than the depth of plastically deformed zones established in copper layers. Deformation was confined into narrow shear bands that acted as crack nucleators. Longitudinal cracks in Fig. 8(c) provide examples of the crack propagation process. However, these cracks did not reach the contact surfaces and did not cause delamination. In fact, a typical wear debris (Fig. 8(d)) was an aggregate composed of small, roughly equiaxed particles of about 0.1-l pm diameter. An X-ray examination of the debris particles revealed that the material was partially crystallized during wear. The peaks in the Xray diffraction spectrum of debris particles (Fig. 9) indicated that the crystalline particles consisted of an f.c.c. a-Ni phase. These peaks were superimposed on a broad and diffuse peak of the parent amorphous phase. It was also observed that some of the debris particles were oxidized and produced an NiO phase. X-ray analysis on the worn Ni78Si10B12 surfaces showed the presence of some a-Ni and oxide phases but to a lesser extent. In order to determine crystallization temperature of metallic glass DSC analysis was performed on undeformed Ni,sSilOB,z samples. The dynamic crystallization temperature (at a heating rate of 20 K s-‘) was found to be 770 K.
Fig. 8. (a) Morphology of worn surface of Ni78Si,0B,2; (b) structure of shear bands on a cross-section parallel to the sliding direction; (c) structure of shear bands on a section transverse to the sliding direction; (d) morphology of an aggregate of loose debris.
295
5160 320 z w 60 2 z 40 ti = 0
20
40
60
80
28 (DEGREES) Fig. 9. X-ray diffraction
pattern
of the Ni7$i10B12 debris particles.
4. Discussion The results presented in the previous section indicate that the sliding wear resistance of the laminated composites is a complex, structure-sensitive property which depends not only on the volume fraction, spatial distribution of its components, but also on the intrinsic mechanisms of debris formation in each of its constituents. In both the copper matrix and “reinforcing” Ni78Si10B12layers, distinct deformation ~crost~ct~es were developed below the contact surfaces. The strains imposed on these regions were very large and these were distributed inhomogeneously in the form of localized shear bands. Metallographic examination revealed that there is a close relationship between the properties of heterogeneously deformed subsurface regions and the ~~romecha~sms of debris formation. It is therefore of value to discuss the wear mechanism in each of the components of the laminates prior to analysing the role of factors such as the volume fraction, distribution of metallic glass layers on the overall wear resistance of the laminated composites. Let us consider the wear behaviour of the copper layer first. Wear in the copper matrix proceeds by a damage accumulation process. The increasing grain aspect ratio on approaching the contact surfaces (Fig. 5) reveals the presence of a plastic strain gradient of a form in accord with those observed by Moore and Douthwaite 161. As indicated by the stress-strain curve for the plastically deformed zone (Fig. 6), the local flow stress of the material within this region increases to a level (400 MPa) more than twice that of the strength of the undeformed bulk material. The increase in the flow stress is due to the work hardening process accompan~g the defo~ation of copper grains. However, at larger strains the material loses its ability to sustain further work hardening and consequently plastic deformation becomes heterogeneous and instabilities in the form of shear bands start to develop. The relationship between the exhaustion of work hardening capacity and the strain localization phenomenon in polycrystalhne materials has been documented by several authors 124-27 J. The details of microst~ct~~ processes leading to shear localization are not fully understood, but there is accumulating evidence that the shear bands represent volumes in which dynamic recovery events can occur concomitantly with continued deformation so that the strain
296
hardening rate inside the bands becomes close to zero or negative depending on the rate and temperature of straining [%I. It is however clear that in copper layers, the shear localization phenomenon marks the onset of the crack nucleation and propagation processes that lead to the delamination of subsurface layers during the sliding wear. Examples of shear bands acting as crack nucleators are shown in Fig. 7(c). The existence of several subsurface cracks ZOO-800 pm long at various cross-sections investigated suggests that the crack propagation is an important and possibly rate controlling stage of debris formation [29]. These cracks finally reach the contact surface and cause detachment of thin plate-like debris particles in a manner similar to that described by the delamination theory proposed by Suh [30]. Two other aspects of the heterogeneous deformation, namely, grain rotation and dynamic recrystallization should also be noted. Due to the orientation incompatibility effects between the grains subject to different degrees of rotation (Fig. 7(b)) grain boundaries can provide additional sites for crack nucleation. As seen in Fig. 5, the structure of the material between the propagating cracks is merent than the rest of the deformed zone. Detailed transmission electron microscopy analyses 131, 321 indicated that these regions consist of recrystallized grains with diameters less than 1 ,um. The origin of recrystallization can be attributed to the localization of strain energy within the bands. The rate of formation of loose debris is about six times lower in Ni7aSi10B1a layers in comparison with copper (Table 1) and the volume of the material deformed below the worn surfaces is confined to a depth of 10 pm which is also 6-7 times smaller than that observed in copper layers. These results refIect the ~c~ty of producing plastic flow and wear damage in Ni7&i10B1a due to its high hardness and yield strength. The plastically deformed zones are characterized by the formation of a shear band structure but the properties of these zones are different from those in copper because metallic glasses are deprived of a metallurgical microstructure and they do not possess a mechanism of spreading plastic deformation [g-13]. Thus the nucleation of shear bands coincides with the onset of yielding of the subsurface regions unlike polyc~sta~~e copper layers which are subject to a high rate of work hardening prior to strain localization. ~icrohar~ess measurements indicated that the average hardness (and the flow strength) of the material inside the deformed zones in Ni,aSilOB1a was similar to that of the bulk. Although shear bands act as crack initiators (Fig. 8(c)), there is no evidence that this process leads to the formation of delaminated wear fragments. In fact, the loose debris was not in the form of thin flakes but exhibited roughly spherical particles of approximately 1 pm diameter that were partially crystallized during sliding (Fig. 9). The presence of crystalline products in the wear debris implies that at some points of the contact the temperature increase is substantial. The complete crystallization sequence of Ni,sSilOB1a involves nucleation and growth of f.c.c. a-Ni particles followed by the formation of tetragonal NiaB phase. However, only primary cr-Ni particles were found in the wear debris even after 22 h sliding. Therefore, crystallization was not complete and, based on the calorimetric and X-ray measurements, it is
estimated that the temperature rise at the contact surface of the metallic glass layers should be of the order of 650 K which is the minimum temperature below which no NiaB phase is nucleated during isothermal annealing in DSC (up to 40 h). It is likely that the crystallization process is assisted by the inhomogeneous plastic deformation and the shear bands may act as favourable nucleation sites for these crystals. Partially crystalline alloy has a lower ductility and fracture toughness than amorphous phase owing to the decohesion of particle/amorphous matrix interfaces (331. The morphologies of both loose debris and the worn surfaces (Figs. S(a) and (c)) suggest that the fracture of the crystallized regions from the amorphous surfaces is the main mechanism leading to debris formation in Ni78Si10B,2 layers. As sliding proceeds, the abrasive action of the debris particles entrapped between the contact surfaces leads to the formation of plowing grooves observed on the worn surfaces of the metallic glass. It should be noted that some copper and metallic glass debris was adhered on the surface of the steel slider. A small amount of iron, and other elements, from the slider was transferred to the wearing surfaces, particularly to the surface of metallic glass layers. However, this mutual transfer phenomenon plays a secondary role as a rate controlling wear process. From the design point of view it is beneficial to be able to predict the wear resistance of the composite in terms of the wear resistance of its individual components and their volume fractions using analytical models such as rule of mixtures. Figure 10 shows the sliding wear rate of the laminated composite according to the two commonly used rules of mixtures (i.e. linear and inverse rules of mixtures [4, 51) as a function of the volume fraction of Ni&ilOB1a layers. It is seen that neither of the models is in accord with the experimental results that can be described by the following expression W= (WC, - WJ
exp( - PVJ + WA
where IV, Wc,, W, are the wear rates of the composite, copper and Ni78Si10B12 respectively. The value of the numerical constant p is 24.6 for the present experimental conditions. It should also be noted that the wear rates of the composites are independent of the orientation of lamellae on the sliding plane. Thus this expression is applicable to the composites tested in both configurations shown in Fig. 2. The important conclusion that can be drawn from Fig. 10 is that the metallic glass layers contribute to the wear resistance of the composite more efficiently than could be expected from their volume fraction. Their contribution to the wear resistance is twofold: they decrease the wear rate of the composite by supporting the applied load with less deformation than the copper matrix, and they retard the wear damage initiated in copper (Fig. 4(c)). In effect, for volume fractions larger than 0.2, almost all the load is carried by the Ni78Si10B12. Thus these layers present an effective way of controlling the sliding contact geometry and improving the wear properties of the composite.
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VOLUME
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FRACTION
0.6
OF
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Fig. 10. Wear rates W of the laminated composite as a function of volume fraction VA of Ni78Si10B12 (subscripts A and 0.1 refer to metallic gIass and copper respcctivcly): curve 1, experimental; curve 2, linear rule of mixtures, W= V,W, C (I- VJW,,; cwve 3, invcrsc rule of mixtures,
w=
6. Conclusions
(1) The homogeneous nature of strain dist~butjon within the plastically deformed zones below the sliding contact surfaces of both copper and metallic glass layers plays an important role in determining the mechanisms of wear debris formation in polycrystalline and amorphous alloys. (2) In copper layers, shear bands adjacent to the contact surfaces act as crack nucleation sites and propagation of these cracks leads to material removal by delamination of subsurface layers. (3) Localized strain gradients accelerate c~s~~zation of metallic glass layers during sliding. Wear proceeds by decohesion or fracture of crystallized particles at the contact surfaces and by microgrooving of these surfaces due to the abrasive action of detached crystallization products. (4) Metallic glass layers carry the applied loads during sliding with less deformation and wear damage than copper layers. They also impede the propagation of wear damage started in copper layers. Thus Ni,,Si,013,2 layers serve to control the contact geometry and the wear resistance of the laminates more effectively than the predictions of the rules of mixtures.
299
Acknowledgments The authors wish to acknowledge Natural Sciences and Engineering Research Council of Canada for support during the period of this work. The authors would like to thank Mr. J. Zhang and Mr. H. Hu of the University of Windsor for their help in completing some experiments.
References 1 J. F. Archard and W. Hirst, The wear of metals under unlubricated conditions, Proc. R. Sot. London, Ser. A, 236 (1956) 397-410. 2 M. A. Moore, R. C. D. Richardson and D. G. Attwood, The limiting strength of worn metal surfaces, Metall. Trans. A, 3 (1972) 2485-2491. 3 D. A. Rigney, L. H. Chen, M. G. Naylor and A. R. Rosenfield, Wear processes in sliding systems, Wear, 100 (1984) 195-219. 4 E. Hombogen, Description ofwear of materials with isotropic and anisotropic microstructures. In K. C. Ludema (ed.), Wear of Materials, ASME, New York, 1985, pp. 477-484. 5 K. H. Zum-Gal-u,Microstructure and Wear of Materials, Elsevier, Amsterdam, 1987, pp. 313328. 6 M. A. Moore and R. M. Douthwaite, Plastic deformation below worn surfaces, Metall. Trans. A, 7 (1976) 1833-1839. 7 F. E. Kennedy, L. A. Hatman, K. E. Hanck and V. A. Surprenant, The role of near surface plastic deformation in the wear of lame&r solids. In K. C. Ludema (ed.), Wear of Materials, ASME, New York, 1985, pp. 273-279. 8 A. T. Alpas and J. D. Embury, Flow localization in thin layers of amorphous alloys in laminated composites, Scripta Metall., 22 (1988) 265-270. 9 C. A. PampilIo, Flow and fracture in amorphous alloys, J. Mater. Sci., 10 (1975) 1194-1227. 10 J. C. M. Li, Mechanical properties of amorphous metals and alloys. In H. Herman (ed.), Treatise on Materials Science and Technology, Vol. 20, Academic Press, New York, 1981, pp. 326-343. 11 T. Masumato and R. Maddin, The mechanical properties of palladium silicon ahoy quenched from liquid state, Acta Metall., 19 (1971) 725-741. 12 A. S. Argon, Plastic deformation in metallic glasses, Acta Metall., 27 (1979) 47-58. 13 S. Takayama and R. Maddin, Rolling and bending deformation of Ni-Pd-P metal glasses, Acta Metall., 23 (1975) 943-952. 14 D. H. Lee and J. E. Evetts, Sliding friction and structural relaxation of metallic glasses, Acta Metall., 32 (1984) 1035-1043. 15 Kishore, U. Sudarsan, N. Chandran and K. Chattopadhyay, On the wear mechanisms of iron and nickel-based metallic glasses, Acta MetalL, 35 (1987) 1463-1473. 16 R. Klinger and H. G. Feller, Sliding friction and wear resistance of the metallic glass Fe.,aNilOBzO, Wear, 86 (1983) 287-297. 17 R. Moreton and J. K. Lancaster, The friction and wear behaviour of various metallic glasses, J. Muter. Sci. L&t., 4 (1985) 133-137. 18 N. Dolezal and G. Hausch, Adhesive wear of some Fe-%3 amorphous alioys, Rapidly Quenched Metals, (1985) pp. 1767-1770. 19 K. Miyoshi and D. H. Buckley, Microstructure and surface chemistry of amorphous alloys important to their friction and wear behaviour, Wear, I10 (1986) 295-313. 20 J. F. Archard, Contact and rubbing of flat surfaces, J. Appl. Phus., 24 (1953) 981-988. 21 J. H. Dautzenberg and J. H. Zaat, Quantitative determination of deformation by sliding wear, Wear, 23 (1973) 9-19. 22 R. Hill, Th.e Mathematical Theory of Plasticity, Clarendon, Oxford, 1950. 23 D. M. Marsh, Plastic flow in glass, Proc. R. Sot. London, Ser. A, 279 (1964) 420-435.
300 24 G. R. Canova, U. F. Kocks and M. G. Stont, On the origin of shear bands in textured polycrystals, Scripta Metall., 18 (1984) 437-442. 25 A. S. Malin and M. Hatherly, Microstructure of cold-rolled copper, Met. Sci., (1979) 463-472. 26 W. Y. Yeung and B. J. Duggan, FIow Iocalization in cold rolled brass, Mater. Sci. Technol., 2 (1986) 552-558. 27 J. F. Knott, Micromech~~ms of fibrous crack extension in engineering alloys, Met. Sci., (1980) 327-336. 28 A. Korbel, J. D. Embury, M. Hatherly, D. L. Martin and H. W. Erbsloh, Microstructural aspects of strain localization in Al-Mg alloys, Acta MetaEl., 34 (1986) 1999-2009. 29 J. R. Fleming and N. P. Suh, The relationship between crack propagation rates and wear rates, Wear, 44 (1977) 57-64. 30 N. P. Suh, Pibophysics, Prentice-Hall, Englewood Clifts, NJ, 1986. 31 J. H. Dautzenberg, The role of dynamic recrystallization in dry sliding wear, Wear, 60 (19803 401-411. 32 J. J. Wert and W. M. Cook, The influence of stacking fault energy and the wear of copper and duminum bronze, Wear, I.23 (1988) 171-192. 33 A. T. Alpas, L. Edwards and C. N. Reid, Fatigue crack propagation in an Ni&+,B,a metallic glass with amorphous and partially crystalline structures, Mater. 5%. Eng., 98 (1988) 501-504.
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285-300
The role of subsurface deformation and strain localization on the sliding wear behaviour of laminated composites A. T. Alpas
J. D. Embury
(Received
June 29, 1990; accepted
November
9, 1990)
Abstract The sliding wear behaviour of laminated composites of copper and amorphous Ni7&3i10Bi2 fabricated by diffusion bonding was investigated. Unlubricated sliding tests were performed using a block-on-ring type wear machine in order to provide a direct comparison of the wear mechanisms in the metallic glass (Ni78Si10B12) layers and polycrystalhne matrix (copper) and to determine the influence of the metallic glass layers on the overall wear resistance of the composite. Wear in the copper layer was characterized by extensive plastic deformation and by the presence of subsurface cracks. These cracks originated at the shear bands within the plastic zones under the contact surfaces and caused delamination of the material adjacent to these surfaces. The wear of metallic glass layers also involved the locahzation of plastic deformation into shear bands. However, due to the high hardness and fracture strength of Ni&3i10B12,its wear resistance was higher than the copper matrix. Localized strain and temperature gradients generated during sliding contact led to the crystallization of amorphous layers. Metallic glass layers were effective in increasing the wear resistance of the composite by supporting the load with less defo~ation and thereby obstructing the damage process initiated in the copper layers.
1. ~troduction
The production of wear debris is an extremely complex process involving mechanical, thermal and chemical interactions between the surfaces of materials in sliding and rolling contact. Under the service conditions the range of local pressures and temperatures generated during the wear may be very broad depending on the operating parameters such as applied load and sliding velocity, and on the detailed geometry of contacting surfaces [l-3]. The development of laminated composite materials provides a method of controlling the geometry of contact and also creates a possibility of designing microstructures with enhanced wear resistance baaed on the control of intrinsic mechanical properties and spatial distribution for the constituent phases. Some principal features of abrasive and adhesive wear resistance of
0043-1648/91/$3.50
Q Elsevier Sequois/Printedin The Netherlands
286
laminated composites were discussed by Hornbogen [4] and Zum-Gahr [5]. In these studies, consideration was given to the description of the wear behaviour of lamellar solids in terms of rules of mixtures. The role of plastic deformation in copper-silver and copper-tin type laminated structures were investigated by Moore and Douthwaite [6] and Kennedy et aZ. [ 71. In the present work laminates of amorphous Ni,$i,,B,, and polycrystalline copper were fabricated by diffusion bonding [8]. It is known that metallic glasses exhibit a useful combination of mechanical properties such as high hardness (5-10 GPa) and fracture strengths that approach theoretical cohesive strength [9, 101. However, these materials do not possess a mechanism for distributing plastic deformation and at ambient temperatures plastic flow proceeds by nucleation and propagation of inhomogeneous shear bands [ 11, 121. Although this behaviour limits the tensile ductilities of metallic glasses to low values they show ductilities comparable with those of crystalline alloys when deformed by uniaxial compression and cold rolling [ 131. These properties make metallic glasses of interest for tribological applications and their friction and wear behaviour have been the subject of several studies that were reviewed by Lee and Evetts [ 141 and Kishore et al. [15]. It has been shown that amorphous alloys generally show lower values of coefficients of friction and wear rates [ 16,. 171 than their crystalline counterparts and this has been attributed to surface oxidation, dynamic recrystallization events taking place during sliding and abrasive contacts [ 14, 15, 18, 191. This study aims to clarify the role of metallic glass layers on the wear resistance of laminated composite materials by investigating the rate controlling wear mechanisms in Ni78Si10B12layers, and in the polycrystalline copper matrix, in relation to the deformation processes occurring during sliding wear. 2. Experimental
procedures
Laminated composite
structures were fabricated using thin ribbons
(t =57 + 2 pm) of amorphous Ni78Si‘10B r2 alloy and annealed commercial
purity copper sheets of various thicknesses. Metallic glass ribbons were produced using a melt spinning process (Vacuumschmelze, F.R.G.); their tensile strength and Vickers hardness were 2100 MPa and 9100 MPa respectively. The copper sheets had a tensile strength of 220 MPa and hardness of 670 MPa. To produce laminated composites, surfaces of metallic glass ribbons were electroplated by copper (approximately 100 pm thick). The coated ribbons were then sandwiched between annealed copper sheets and diffusion bonded in air under a pressure of 200 MPa at 560 K. Details of the diffusion bonding process are described in ref. 8. The cross-section of the composite is shown in Fig. 1. X-ray diffraction and scanning diffraction calorimetric (DSC) analyses indicated that the metallic glass layers completely maintained their non-crystalline structure during diffusion bonding. Dry sliding wear tests were performed using a block-on-ring type wear rig. Specimens in the form of 5 x 10 x 10 mm3 rectangular blocks were worn
287
Fig. 1. Microstructure of amorphous Ni78Sl‘,0B 12-Cu laminated composite. The featureless layers are metallic glass, the electrodeposited copper coatings have a columnar structure, and the annealed copper sheets show equiaxed grains. Fig. 2. Testing configurations of laminated composites. The slider moves (a) parallel and (b) perpendicular to the laminates.
against an SAE 52100 bearing steel slider ring (diameter, 38 mm). These tests were conducted such that the slider moved in a direction either parallel or transverse to the laminates as shown in the schematic diagram in Fig. 2. A normal load of 80 N and a sliding velocity of 0.1 m s-’ were used. All the experiments were performed at room temperature (21 “C) and in laboratory air (approximately 55% humidity). The volume losses during wear were determined following the procedures described in ASTM standard G7783. The morphologies of the wear tracks, loose debris and the microstructures of the material under the worn surfaces were examined by a Philips 551 scanning electron microscope.
3. Results Results of wear tests performed on laminates containing various volume fractions of Ni&i10B12 layers are shown in Fig. 3. The figure indicates that the lowest wear losses occurred in 100% metallic glass specimens tested in monolithic form and the incorporation of metallic glass layers into the copper matrix resulted in a reduction in the amount of material lost during wear. No significant difference could be detected between the wear losses of laminates lying parallel and transverse to the sliding direction so that the curves in this figure represent combined results of tests carried out using both configurations shown in Fig. 2. The volumetric wear rates of composites (i.e. the volume of material removed per unit sliding distance) are calculated from the slopes of the curves in Fig. 3 and these are shown as a function of vohune fraction of Ni&i10B12 in Table 1. The wear rates of the composites decrease rapidly with the addition of small volume fractions of metallic glass and the wear resistance of the composite containing 20% Ni,,Si,,,B,z becomes as good as that of the 100% metallic glass samples. The beneficial effect
289
Figures 4(a) and (b) show differences in the topographies of the worn surfaces of metallic glass and copper layers. It can be seen that the rate of material removal was higher in copper so that the top portions of Ni78Si10B12 layers protruded from the contact surfaces. The copper layers were characterized by severe plastic deformation and surface damage compared with metallic glasses whose surfaces appear to be smoother and relatively featureless at low magnifications. The cavities and plateau regions formed within the wear tracks of copper layers were wider than the total thickness of metallic glass layers. Observations made on cross-sections parallel to the sliding direction (Fig. 4(c)) indicated that large plastic strain gradients were developed in the copper grains adjacent to the contact surfaces as revealed by the shear displacements at the diffusion bonded Cu-Cu interfaces. Plastic flow in the copper matrix was impeded by Ni78Si10B12 layers and a large number of voids were created in front of metallic glass-Cu interfaces. These observations underline two salient features of the role of Ni78Si10B12on the sliding wear behaviour of the composite. These layers serve as load carrying constituents and they also act as barriers to the propagation of the damage process. The differences in surface topographies reflect the difliculty of producing loose debris in metallic glass layers. To clarify the micromechanism of debris
Fig. 4. Worn surface topographies of composites: (a) sliding direction is parallel to laminates; (b) sliding direction is perpendicular to laminates; (c) cross-section of composite psrallel to the slidiig direction.
290
formation in the composite, detailed metallographic examinations were performed on both matrix and Ni78Si10B12 layers. It is of value to examine these results in two separate sections. 3.2. Metallography of copper layers The microstructure of copper close to the wearing surfaces was composed of highly deformed and distorted grains that assumed a “tear drop” shape with their narrow edges lying almost parallel to these surfaces (Fig. 5). The deformed zones had a fairly uniform thickness of 60-70 ,um and included 2-5 pm grams which were initially equiaxed (Fig. 1). The plastic strain distribution due to this deformation can be estimated using a method proposed by Dautzenberg and Zaat [ 2 1 I. Shear angles of grain boundaries (and diffusion bonded interfaces) were measured at various cross-sections and these are related to equivalent plastic strain 1221. The variation of plastic strain E as a function of depth below the worn surfaces 2 is plotted in Fig. 6(a). At the contact surfaces, strains as high as 30 were reached; these decreased
Fig. 5. Cross-section of copper layers parallel to the sliding direction: (a) microstructure of the deformation zones under the worn surfaces and the subsurface cracks propagating parallel to the worn surfaces; (b) fracture and delamination of material adjacent to the contact surface.
291
40
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20 u, z z E v)
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DEPTH
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( pm)
450
400
2 I350 z ii F v) 300
250 0
2
4
6
6
10
12
14
16
16
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22
STRAIN Fig. 6. (a) Microhardness and equivalent strain vs. depth below the worn surface of copper layers; (b) stress-strain curves for deformed copper layers.
292
to levels of about 2-3 at 10 pm below the surfaces. Figure 6(a) also shows the microhardness gradients generated during wear. Microhardness measurements were made (at a load of 50 gf) on tapered sections sectioned at an angle of 10” to the worn surfaces. Flow stresses were calculated from the microhardness data using a simple expression derived by Marsh [ 231 and a stress-strain curve for worn copper layers is presented in Fig. 6(b). The curve exhibits an initial stage of work hardening, but at large strains (er3), the curve flattens out, indicating either the exhaustion of the work hardening capacity or local recovery mechanisms in the grains closest to the contact surfaces. Figure 7(a) shows the microstructure of material under the wear track on a plane perpendicular to the sliding direction (and to the sections in Fig. 5). Within the plastically deformed zone, two orthogonal sets of shear bands, each inclined at about f35” to the wear surface, were developed. These bands crossed the grain boundaries without deviation except those which had undergone rotation and re-orientation (Fig. 7(b)). With increasing strain (i.e. on approaching the surface) the shear bands increased in number and the spacing between them became smaller. The intersection of these bands
(al
@I
cc>
Fig. 7. Cross-sections of copper layers transverse to the sliding direction: (a) structure of shear bands aligned 35” to the contact surface; (b) heavily deformed grains acijacent to the contact surface; (c) crack nucleation at the intersection of shear bands.
293
provided nucleation sites for crack formation as shown in Fig. 7(c). Cracks were also nucleated at some grain boundaries (Fig. 7(b)), possibly owing to the incompatibility effects between adjacent grams which were subject to different degrees of rotation during wear. Scanning electron microscopy observations also revealed that severe subsurface damage was often associated with deIamination parallel to the contact surface. Delamination events led to the formation of loose plate-like debris whose thicknesses (5-20 pm> were in close agreement with the loci of subsurface cracks (F’ig. 5@)). Thus, wear in copper layers proceeds by a damage accumulation process which involves the generation of large strain gradients below the contact surfaces. Under the influence of these subsurface strains, copper grams adjacent to the wearing surfaces lose their ability to sustain work hardening and become subject to an inhomogeneous deformation in the form of shear band formation. Cracks nucleated at these bands and at the grain boundaries propagate parallel to the contact surfaces and result in the creation of debris particles from the contact surfaces.
Wear tracks of Ni7$i10B1a were characterized by a lower level of damage and high magnification microscopic examinations were required to observe the features of the worn surfaces. These involved fine parallel grooves of approximately 0.5 pm depth and 1-2 pm width which were elongated in the direction of sliding (Fig. 8(a)). These grooves were indicative of an abrasive type of wear process. In contrast, investigations on the cross-sections parallel and perpendicular to the sliding direction (F’igs. 8(b) and (c)) revealed the development of a plastically deformed zone under the sliding contact surfaces. The deformation zone was almost continuous ~oughout the layers and penetrated to a depth of 10 pm which is about 6-7 times smaller than the depth of plastically deformed zones established in copper layers. Deformation was confined into narrow shear bands that acted as crack nucleators. Longitudinal cracks in Fig. 8(c) provide examples of the crack propagation process. However, these cracks did not reach the contact surfaces and did not cause delamination. In fact, a typical wear debris (Fig. 8(d)) was an aggregate composed of small, roughly equiaxed particles of about 0.1-l pm diameter. An X-ray examination of the debris particles revealed that the material was partially crystallized during wear. The peaks in the Xray diffraction spectrum of debris particles (Fig. 9) indicated that the crystalline particles consisted of an f.c.c. a-Ni phase. These peaks were superimposed on a broad and diffuse peak of the parent amorphous phase. It was also observed that some of the debris particles were oxidized and produced an NiO phase. X-ray analysis on the worn Ni78Si10B12 surfaces showed the presence of some a-Ni and oxide phases but to a lesser extent. In order to determine crystallization temperature of metallic glass DSC analysis was performed on undeformed Ni,sSilOB,z samples. The dynamic crystallization temperature (at a heating rate of 20 K s-‘) was found to be 770 K.
Fig. 8. (a) Morphology of worn surface of Ni78Si,0B,2; (b) structure of shear bands on a cross-section parallel to the sliding direction; (c) structure of shear bands on a section transverse to the sliding direction; (d) morphology of an aggregate of loose debris.
295
5160 320 z w 60 2 z 40 ti = 0
20
40
60
80
28 (DEGREES) Fig. 9. X-ray diffraction
pattern
of the Ni7$i10B12 debris particles.
4. Discussion The results presented in the previous section indicate that the sliding wear resistance of the laminated composites is a complex, structure-sensitive property which depends not only on the volume fraction, spatial distribution of its components, but also on the intrinsic mechanisms of debris formation in each of its constituents. In both the copper matrix and “reinforcing” Ni78Si10B12layers, distinct deformation ~crost~ct~es were developed below the contact surfaces. The strains imposed on these regions were very large and these were distributed inhomogeneously in the form of localized shear bands. Metallographic examination revealed that there is a close relationship between the properties of heterogeneously deformed subsurface regions and the ~~romecha~sms of debris formation. It is therefore of value to discuss the wear mechanism in each of the components of the laminates prior to analysing the role of factors such as the volume fraction, distribution of metallic glass layers on the overall wear resistance of the laminated composites. Let us consider the wear behaviour of the copper layer first. Wear in the copper matrix proceeds by a damage accumulation process. The increasing grain aspect ratio on approaching the contact surfaces (Fig. 5) reveals the presence of a plastic strain gradient of a form in accord with those observed by Moore and Douthwaite 161. As indicated by the stress-strain curve for the plastically deformed zone (Fig. 6), the local flow stress of the material within this region increases to a level (400 MPa) more than twice that of the strength of the undeformed bulk material. The increase in the flow stress is due to the work hardening process accompan~g the defo~ation of copper grains. However, at larger strains the material loses its ability to sustain further work hardening and consequently plastic deformation becomes heterogeneous and instabilities in the form of shear bands start to develop. The relationship between the exhaustion of work hardening capacity and the strain localization phenomenon in polycrystalhne materials has been documented by several authors 124-27 J. The details of microst~ct~~ processes leading to shear localization are not fully understood, but there is accumulating evidence that the shear bands represent volumes in which dynamic recovery events can occur concomitantly with continued deformation so that the strain
296
hardening rate inside the bands becomes close to zero or negative depending on the rate and temperature of straining [%I. It is however clear that in copper layers, the shear localization phenomenon marks the onset of the crack nucleation and propagation processes that lead to the delamination of subsurface layers during the sliding wear. Examples of shear bands acting as crack nucleators are shown in Fig. 7(c). The existence of several subsurface cracks ZOO-800 pm long at various cross-sections investigated suggests that the crack propagation is an important and possibly rate controlling stage of debris formation [29]. These cracks finally reach the contact surface and cause detachment of thin plate-like debris particles in a manner similar to that described by the delamination theory proposed by Suh [30]. Two other aspects of the heterogeneous deformation, namely, grain rotation and dynamic recrystallization should also be noted. Due to the orientation incompatibility effects between the grains subject to different degrees of rotation (Fig. 7(b)) grain boundaries can provide additional sites for crack nucleation. As seen in Fig. 5, the structure of the material between the propagating cracks is merent than the rest of the deformed zone. Detailed transmission electron microscopy analyses 131, 321 indicated that these regions consist of recrystallized grains with diameters less than 1 ,um. The origin of recrystallization can be attributed to the localization of strain energy within the bands. The rate of formation of loose debris is about six times lower in Ni7aSi10B1a layers in comparison with copper (Table 1) and the volume of the material deformed below the worn surfaces is confined to a depth of 10 pm which is also 6-7 times smaller than that observed in copper layers. These results refIect the ~c~ty of producing plastic flow and wear damage in Ni7&i10B1a due to its high hardness and yield strength. The plastically deformed zones are characterized by the formation of a shear band structure but the properties of these zones are different from those in copper because metallic glasses are deprived of a metallurgical microstructure and they do not possess a mechanism of spreading plastic deformation [g-13]. Thus the nucleation of shear bands coincides with the onset of yielding of the subsurface regions unlike polyc~sta~~e copper layers which are subject to a high rate of work hardening prior to strain localization. ~icrohar~ess measurements indicated that the average hardness (and the flow strength) of the material inside the deformed zones in Ni,aSilOB1a was similar to that of the bulk. Although shear bands act as crack initiators (Fig. 8(c)), there is no evidence that this process leads to the formation of delaminated wear fragments. In fact, the loose debris was not in the form of thin flakes but exhibited roughly spherical particles of approximately 1 pm diameter that were partially crystallized during sliding (Fig. 9). The presence of crystalline products in the wear debris implies that at some points of the contact the temperature increase is substantial. The complete crystallization sequence of Ni,sSilOB1a involves nucleation and growth of f.c.c. a-Ni particles followed by the formation of tetragonal NiaB phase. However, only primary cr-Ni particles were found in the wear debris even after 22 h sliding. Therefore, crystallization was not complete and, based on the calorimetric and X-ray measurements, it is
estimated that the temperature rise at the contact surface of the metallic glass layers should be of the order of 650 K which is the minimum temperature below which no NiaB phase is nucleated during isothermal annealing in DSC (up to 40 h). It is likely that the crystallization process is assisted by the inhomogeneous plastic deformation and the shear bands may act as favourable nucleation sites for these crystals. Partially crystalline alloy has a lower ductility and fracture toughness than amorphous phase owing to the decohesion of particle/amorphous matrix interfaces (331. The morphologies of both loose debris and the worn surfaces (Figs. S(a) and (c)) suggest that the fracture of the crystallized regions from the amorphous surfaces is the main mechanism leading to debris formation in Ni78Si10B,2 layers. As sliding proceeds, the abrasive action of the debris particles entrapped between the contact surfaces leads to the formation of plowing grooves observed on the worn surfaces of the metallic glass. It should be noted that some copper and metallic glass debris was adhered on the surface of the steel slider. A small amount of iron, and other elements, from the slider was transferred to the wearing surfaces, particularly to the surface of metallic glass layers. However, this mutual transfer phenomenon plays a secondary role as a rate controlling wear process. From the design point of view it is beneficial to be able to predict the wear resistance of the composite in terms of the wear resistance of its individual components and their volume fractions using analytical models such as rule of mixtures. Figure 10 shows the sliding wear rate of the laminated composite according to the two commonly used rules of mixtures (i.e. linear and inverse rules of mixtures [4, 51) as a function of the volume fraction of Ni&ilOB1a layers. It is seen that neither of the models is in accord with the experimental results that can be described by the following expression W= (WC, - WJ
exp( - PVJ + WA
where IV, Wc,, W, are the wear rates of the composite, copper and Ni78Si10B12 respectively. The value of the numerical constant p is 24.6 for the present experimental conditions. It should also be noted that the wear rates of the composites are independent of the orientation of lamellae on the sliding plane. Thus this expression is applicable to the composites tested in both configurations shown in Fig. 2. The important conclusion that can be drawn from Fig. 10 is that the metallic glass layers contribute to the wear resistance of the composite more efficiently than could be expected from their volume fraction. Their contribution to the wear resistance is twofold: they decrease the wear rate of the composite by supporting the applied load with less deformation than the copper matrix, and they retard the wear damage initiated in copper (Fig. 4(c)). In effect, for volume fractions larger than 0.2, almost all the load is carried by the Ni78Si10B12. Thus these layers present an effective way of controlling the sliding contact geometry and improving the wear properties of the composite.
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/
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I
0
0.2
VOLUME
0.4
FRACTION
0.6
OF
0.8
A
I
Ni78~ilO~12
Fig. 10. Wear rates W of the laminated composite as a function of volume fraction VA of Ni78Si10B12 (subscripts A and 0.1 refer to metallic gIass and copper respcctivcly): curve 1, experimental; curve 2, linear rule of mixtures, W= V,W, C (I- VJW,,; cwve 3, invcrsc rule of mixtures,
w=
6. Conclusions
(1) The homogeneous nature of strain dist~butjon within the plastically deformed zones below the sliding contact surfaces of both copper and metallic glass layers plays an important role in determining the mechanisms of wear debris formation in polycrystalline and amorphous alloys. (2) In copper layers, shear bands adjacent to the contact surfaces act as crack nucleation sites and propagation of these cracks leads to material removal by delamination of subsurface layers. (3) Localized strain gradients accelerate c~s~~zation of metallic glass layers during sliding. Wear proceeds by decohesion or fracture of crystallized particles at the contact surfaces and by microgrooving of these surfaces due to the abrasive action of detached crystallization products. (4) Metallic glass layers carry the applied loads during sliding with less deformation and wear damage than copper layers. They also impede the propagation of wear damage started in copper layers. Thus Ni,,Si,013,2 layers serve to control the contact geometry and the wear resistance of the laminates more effectively than the predictions of the rules of mixtures.
299
Acknowledgments The authors wish to acknowledge Natural Sciences and Engineering Research Council of Canada for support during the period of this work. The authors would like to thank Mr. J. Zhang and Mr. H. Hu of the University of Windsor for their help in completing some experiments.
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300 24 G. R. Canova, U. F. Kocks and M. G. Stont, On the origin of shear bands in textured polycrystals, Scripta Metall., 18 (1984) 437-442. 25 A. S. Malin and M. Hatherly, Microstructure of cold-rolled copper, Met. Sci., (1979) 463-472. 26 W. Y. Yeung and B. J. Duggan, FIow Iocalization in cold rolled brass, Mater. Sci. Technol., 2 (1986) 552-558. 27 J. F. Knott, Micromech~~ms of fibrous crack extension in engineering alloys, Met. Sci., (1980) 327-336. 28 A. Korbel, J. D. Embury, M. Hatherly, D. L. Martin and H. W. Erbsloh, Microstructural aspects of strain localization in Al-Mg alloys, Acta MetaEl., 34 (1986) 1999-2009. 29 J. R. Fleming and N. P. Suh, The relationship between crack propagation rates and wear rates, Wear, 44 (1977) 57-64. 30 N. P. Suh, Pibophysics, Prentice-Hall, Englewood Clifts, NJ, 1986. 31 J. H. Dautzenberg, The role of dynamic recrystallization in dry sliding wear, Wear, 60 (19803 401-411. 32 J. J. Wert and W. M. Cook, The influence of stacking fault energy and the wear of copper and duminum bronze, Wear, I.23 (1988) 171-192. 33 A. T. Alpas, L. Edwards and C. N. Reid, Fatigue crack propagation in an Ni&+,B,a metallic glass with amorphous and partially crystalline structures, Mater. 5%. Eng., 98 (1988) 501-504.