Keywords: Test methodology; Martens hardness; Ratner–Lancaster correlation; Plastics
1. Introduction
Polymeric materials are being increasingly used in a widenumber of applications where resistance to wear is impor-tant. These range from its use as a bearing material (inapplications such as machinery parts and biomedical jointreplacements) to its use as a glazing material where damageresults in loss of optical properties. Polymers are ideal mate-rials for bearing applications due to their general resistanceto corrosion, galling and seizure, their tolerance to smallmisalignments and shock loading and their low coefficientsof friction; as glazing materials, their low density and hightoughness (compared to traditional glass) along with hightransparency are desirable properties. In many applications,polymers may be subjected to abrasive wear, often due tocontaminants within a system, and such abrasion may resultin loss of function.
The abrasive wear of polymers and polymer-based com-posites is the subject of a large body of literature. A numberof test methods have been employed; both two-body andthree-body abrasion have been examined, the former withboth abrasive papers and rough metal counterfaces. Mosttest programmes have employed two-body abrasion tests
whereas in practical applications, three-body abrasion is farmore prevalent[1]. It is widely recognised that the processesof wear in polymers are not well understood[2]. In a reviewof some of the early literature concerning abrasive wear ofpolymers, Evans and Lancaster[3] show that in tests cov-ering eighteen polymer types, low density polyethylene ex-hibited the lowest wear rate in abrasion against a rough mildsteel but the highest wear rate in abrasion with coarse car-borundum paper. Thus, it can be seen that abrasive wear be-haviour of polymeric material is complex. Indeed, Budinski[4] notes that most of the studies on the abrasion resistanceof plastics are inconclusive and tend to recommend furtherstudy.
A number of models which attempt to relate the abrasivewear resistance of polymers to other mechanical propertieshave been proposed. One of the earliest of these is com-monly known as the Ratner–Lancaster correlation[5]. Itpredicts that the wear rate of a polymer,W is given by:
W = kµ
Hσbεb(1)
whereµ is the coefficient of friction,H the hardness andσb and εb are the stress and strain at tensile break. More-over, Lancaster[6] has demonstrated experimentally thatthe wear rate of a range of polymers (with rates coveringnearly two-orders of magnitude) is inversely proportional
to the product ofσb andεb. Other models have been pro-posed and Budinski[4] examined five such models. Heindicated that the correlations proposed by all the modelsbetween the abrasive wear behaviour and other relevantproperties of twenty-one polymeric materials was poor.Other workers have also demonstrated poor correlation be-tween experimental data and one or more of the models[7,8]. Larsen-Basse[7] argued that the mechanisms of weardiffered depending upon the polymer type. For example,in PMMA, two-body grooving wear resulted from abrasionwith SiC paper with tearing and cracking in the grooves;however, in ductile polypropylene, grooving still occurred,but unlike PMMA, no tearing and cracking was observed.In his review paper, Briscoe[2] concludes that the modelssuppose a certain mechanism of material removal to pre-vail, and that changes in mechanism will tend to make themodel predictions invalid.
Most of the published literature in this area which has con-cerned abrasion against particulates has examined behaviourwith particles normally greater than 100�m in size. How-ever, Roberts and Chang[9] indicated that the mechanism ofwear of the polymers examined in their study changed as thesize of the abrasive particles dropped below approximately10�m. More recently, a number of workers[1,10–12]haveemployed a micro-scale abrasion test to examine the be-haviour of polymeric materials with abrasives of the order of4�m in size. In this test method, abrasion is produced by therotation of a ball (with no translation) against a flat surfaceof the test material with an abrasive slurry being entrainedinto the contact zone. This results in wear of the test mate-rial, with the material removed being in the form of a spher-ical cap of the same geometry as the ball. In a polymericmaterial, the wear scar would normally be up to approxi-mately 3 mm in diameter and 60�m deep. In cases wherethe properties of polymer may vary with depth from a sur-face (due to, for example, environmental degradation), therelatively shallow wear scar allows the abrasion resistanceto be easily examined as a function of depth. Also, such atest allows the investigation of the behaviour of thin poly-mer coatings (such as paints) whereas most other tests re-quire substantially thicker samples. The micro-abrasion testhas so far only been applied to a limited range of polymericmaterials, namely, ultra-high molecular weight polyethylene(UHMWPE), polymethylmethacrylate (PMMA) and paint
systems. The current paper examines the micro-scale abra-sion behaviour of a range of polymeric materials in orderto further understand the mechanisms of wear dominant inthese materials when examined in this way.
2. Materials and experimental methods
The microabrasion behaviour of seven commerciallyavailable polymers was examined. The polymer types alongwith tensile mechanical property data supplied by the man-ufacturers are listed inTable 1. The experiments from whichthe tensile test data have been derived have been conductedaccording to the standard DIN 53455. PC, PETG, PMMAand PS were all transparent polymers employed in variousapplications for their optical properties.
Microabrasion testing of the polymers was performedwith a commercially available apparatus, the TE66Micro-Scale Abrasion Tester (Phoenix Tribology Ltd., New-bury, UK). A schematic diagram of the apparatus is shownin Fig. 1. Samples were placed in a holder block which was
Fig. 1. Schematic diagram of the microscale abrasion test apparatus.
rotated around its pivot until the sample came into contactwith the ball. A load was applied via a dead-weight system;the loads applied varied from 0.5 to 4.0 N. The ball wasrotated about a horizontal axis parallel to the plane of thespecimen surface while abrasive slurry was dripped ontothe ball and entrained into the gap between the ball andspecimen resulting in wear of the specimen. Specimen wearresults in an impression, which takes the form of a spheri-cal cap with a geometry similar to that of the ball. In lightof previous research on UHMWPE[12], 25.4 mm diameternylon balls (Du Pont Zytel PA66 balls supplied by DejayDistribution Ltd., Wokingham, UK) were employed in thiswork in order to ensure good entrainment of particles intothe wear zone. The microabrasion tests were conductedwith a slurry of SiC (grade C5, F1200, approximately 4�mparticle size, Washington Mills Ltd., Manchester) sus-pended in distilled water and delivered to the ball at a rateof approximately 38�l s−1. The slurry employed was at aconcentration of approximately 17.2 vol.% SiC. An SEMmicrograph of the abrasive particles (Fig. 2) clearly showstheir highly angular morphology. The ball-specimen slidingspeed was maintained at 0.112 m s−1 throughout the tests(this was at the top end of the range examined by Ruther-ford and Hutchings[13] where they demonstrated that wearrate was independent of sliding speed) and was chosen tominimise the time required for testing. A new ball was usedfor each test, but the ball was run in against a mild steel testcoupon for 200 revolutions before being employed in thetest to ensure that its surface was reproducible[14]. Testswere run for a range of sliding distances ranging from 1 to100 revolutions of the ball. The wear rates quoted are theaverage of three separate tests.
Examination of the wear scars following testing was madeby optical microscopy (Nikon Optiphot) to allow measure-ment of the scar diameter and to check the roundness ofthe scar. To enable the volume of the wear scar to be calcu-lated, spherical cap geometry of the wear scar was assumed(as has generally been demonstrated to be the case[10,13]).The volume,V, is related to the depth,d, of the wear scar
by the following equation:
V = πd2(
r − d
3
)(2)
wherer is the radius of the spherical cap (assumed to bethe radius of the ball). The depth,d, of the cap can becalculated from the optical measurements of the scar radius,a as follows:
d = r −√
r2 − a2 (3)
Further examination of the wear scars produced in thepolymers by micro-scale abrasion was performed with a Sur-fcom stylus-type profilometer (Advanced Metrology Sys-tems, Leicester, UK) and a scanning electron microscope(JEOL 6400).
The Martens hardness (HM) of the polymer samples (loaddivided by area of the indenter penetrating below the orig-inal sample surfacewhilst under load) were measured us-ing a TE76 depth sensing microhardness indenter (PhoenixTribology Ltd., Newbury, UK). All indentations were madeto a depth of 20�m, and the values quoted inTable 1arethe average of 25 separate indentations. The standard errorin the mean of the Martens hardness values was less than0.8 MPa in all cases.
3. Results
3.1. Test characteristics
Initially, all the polymers were subjected to wear at arange of loads with a fixed sliding distance of 100 ball revo-lutions. The wear volume was calculated from optical mea-surements of the scar. The resulting data are shown inFig. 3.The polymers can be divided into two distinct groups basedupon their wear behaviour: HDPE, PP and PVC all showed amonotonic increase in wear rate with increasing load; how-ever, the other polymers exhibited an increase in wear rateas the load was initially increased from its lowest level, butthen a sharp decrease in wear rate as the load was furtherincreased. A decrease in wear rate with increasing load hasbeen observed previously in a number of materials[1,14,15]and is associated with the formation of ridges within thewear scar.Fig. 4 shows an SEM micrograph and a profileacross a scar formed in PMMA by wear under a 2 N loadwhich exhibits a ridge.Fig. 4ashows the ridge towards thecentre of the wear scar. SEM combined with EDX analysisindicated that there is a ridge of PMMA itself which has notworn down with an additional layer of compacted SiC ontop of the ridge. The profile across the scar (Fig. 4b) indi-cated that the ridge was higher than the original height ofthe unworn polymer due to the deposition of compacted SiCabrasive. The profile of the ball following such tests exhib-ited only minimal wear and showed no grooves to match theridges on the test coupon.
Such ridged wear scars indicate that the test is invalid interms of a measurement of the wear behaviour of the test-
Fig. 3. Microscale abrasion rates of various polymers as a function of load.
Fig. 4. Detail of scar in PMMA following wear under a 2 N load for 100ball revolutions: (a) SEM micrograph of ridge within wear scar (slidingdirection from top to bottom); (b) profile across scar showing geometryof ridge.
piece itself, with wear rate being instead governed by abra-sive entrainment. For the polymers which did exhibit the for-mation of ridges in the wear scars at high loads, two regimesmay be defined, namely, one where ridging does not occurand one where ridging does occur. A number of tests acrossthe range of polymers were performed to examine the devel-opment of wear volume with sliding distance; behaviour wasonly examined where ridging did not occur. It was found thata linear increase in wear volume with sliding distance wasgenerally observed;Fig. 5 shows two examples of such be-haviour for PMMA. It can be seen that in these cases, wear isproportional to sliding distance even at small distances (i.e.there is little evidence of any running-in behaviour).Fig. 6shows SEM images of the surface of PMMA following wearunder a load of 1.0 N as a function of sliding distance (these
Fig. 5. Wear volume of PMMA vs. sliding distance under two appliedloads.
Fig. 6. Detailed morphology of wear surface of PMMA following microabrasion under 1.0 N load as a function of sliding distance: (a) 1 rev (0.080 m);(b) 5 rev (0.399 m); (c) 20 rev (1.596 m); (d) 100 rev (7.98 m). Sliding direction from top to bottom in all cases.
images correspond to one set of data inFig. 5). As the num-ber of revolutions increased the surface became more worn,as expected. After just one revolution, there was little evi-dence of directionality in the wear scar, with pitting as thepredominant wear mechanism (Fig. 6a). At greater slidingdistances, there was evidence for wear due to both particlegrooving (parallel lines) and particle rolling (pitting)[16].It may be argued that three-body wear with rolling parti-cle motion predominates at long sliding distances (Fig. 6d).Fig. 7 shows profiles through the centres of the wear scars
Fig. 7. Measured and predicted profiles across wear scars in PMMA fol-lowing microscale abrasion under a load of 1 N for two sliding distances.
perpendicular to the direction of sliding (PMMA, 1.0 N ap-plied load) generated by both one and five ball revolutions.Neither of these scars exhibited ridging. For each scar, theprofile predicted on the assumption that the wear scar is aspherical cap with the same radius of curvature as the nylonball is also presented. It is notable that even after just onerevolution of the ball, a wear scar almost 5�m deep has beendeveloped. In this case, there is a deviation of the geometryof the scar from its assumed spherical cap. However, follow-ing only five revolutions of the ball, a wear scar more than16�m deep has been developed with a geometry very closeto that assumed, as also observed by other workers[10,13].The pressure distribution in the initially non-conformingsurface contact between the ball and testpiece will tend topromote wear to form a conformal contact; it is clear thatsuch a development of a conformal contact can be veryrapid.
3.2. Microabrasive wear behaviour of different polymers
Since the formation of ridges within a wear scar inval-idated the measure of wear resistance, it was not possibleto make direct comparisons between materials under theseconditions. Instead, comparisons of the microscale abrasivewear behaviour of the seven polymer types was made fol-lowing wear under a 0.5 N load where ridging did not occurin any case. Wear consisted of 100 revolutions of the ball,following which the resultant scars were measured an ex-amined. The wear rates of the polymers are listed inTable 1.Fig. 8shows SEM micrographs of the wear scar morphology
Fig. 8. Detailed morphology of wear scars following microabrasion under 0.5 N load and 100 ball revolutions: (a) HDPE; (b) PC; (c) PETG; (d) PMMA;(e) PP; (f) PS; (g) PVC. Sliding direction from top to bottom in all cases.
of the seven polymers. A range of different surface mor-phologies can be observed. The wear scar in HDPE (Fig. 8a)is distinctly different from all the other scars. No evidenceof rolling is present, with instead long, deep parallel groovesin the direction of sliding being observed. The scar in PP(Fig. 8e) is also distinct from the others; here, very littledirectionality at all is observed and also little evidencefor particle indentation is seen. All the other samples ex-
hibit similar morphologies with a combination of shallowgrooving, parallel to the direction of sliding (two-bodydamage) and also pitting due to particle indentation inrolling (three-body damage). The degree of damage differsthroughout the samples as does the proportion of two- andthree-body abrasion. The PMMA and PS (Fig. 8d and f,respectively) show the greatest level of damage, dominatedby three-body abrasion. As the level of damage decreases,
Fig. 9. Polymer wear rate (under a 0.5 N load) vs. the reciprocal of theproduct of tensile breaking stress and tensile breaking strain (the Ratner–Lancaster correlation).
so does the proportion of three-body abrasion until in PETG(Fig. 8c) there is little damage at all. It is noticeable thatthe wear scars in PMMA under a 1 N load (Fig. 6d) anda 0.5 N load (Fig. 8d) are very similar indicating that thesame mechanism of wear operated under the two loads.
Figs. 9 and 10show the wear rates of the various poly-mers (0.5 N, 100 ball revolutions) plotted against 1/σbεb (theRatner–Lancaster correlation) and againstHM, respectively.In both cases, there is a reasonable correlation. It can be seenthat the polymers which exhibited the highest wear rates alsoexhibited the lowest values of (σbεb) and the highest hard-nesses. Conversely, those with the lowest wear rates showedhigh values of (σbεb) and low hardness.
Fig. 10. Polymer wear rate (under a 0.5 N load) vs. Martens hardness.
4. Discussion
4.1. Ridge formation in polymers
Fig. 3 shows the wear rate of the various polymers as afunction of load. It is well established[1,15] that the for-mation of ridges within a wear scar during the microscaleabrasion test is due to slurry starvation, promoted either bylow slurry viscosity or high applied loads. The formation ofa ridge in a scar is associated with a significant reduction inthe wear rate as the dominating wear mechanism changesfrom abrasion to sliding of the ball against the counterface.Trezona and Hutchings[1] performed microscale abrasiontests on PMMA, aluminium, hardened steel and aluminawith a smooth hard steel ball and found that only the softermaterials, namely, PMMA and aluminium exhibited theformation of ridges. Of the four materials examined, thesetwo materials were significantly softer than the ball. Theysuggested that when the ball is harder than the test coupon,the abrasive particles embed into the test coupon and arethus not carried through the contact zone to induce wear.Fig. 3shows that in the current test programme, three of thepolymers (HDPE, PP, PVC) did not exhibit ridge formationat any of the loads examined; examination of the hardnessesof the polymers (Table 1) shows that two of these (HDPEand PP) have verylow hardness compared to the othersand here it is in fact theharder materials that exhibit ridgeformation. Thus, the explanation of Trezona and Hutchings[1] is not applicable in this case. Other factors must alsogovern the tendency for ridge formation in these materials.Wettability of the testpiece by the slurry has been previouslysuggested[15] where slurry entrainment is encouraged bygood wetting of the contact area. It is also notable that all thepolymers which exhibited ridge formation were glassy,
transparent materials whereas the non-ridging polymerswhere all partially crystalline; rationalisation of such anobservation requires further work.
A micrograph of a ridge formed in PMMA is shown inFig. 4 along with a profile across it. The ridge is composedof unworn polymer; however, in this case, SiC particles haveformed an agglomerate (identified by EDX analysis) uponthe ridge. Normally, ridges that are observed in such testsare of the same order of magnitude in height as the sizeof the abrasive particles so that, despite the ball exhibitingvery little wear, general wear of the testpiece can proceedby abrasion even though there is ball-testpiece contact onthe face of the ridge itself[14,15]. However, in this case, theprofile (Fig. 4b) shows that the height of the ridge (includ-ing the SiC agglomerate) is now above that of the originalsurface of the PMMA; the ball will run against this and littlefurther wear of the polymer will occur as the gap betweenthe ball and sample surface is now so large.
Despite the uncertainties concerning the formation ofridges in this test system, it is clear fromFig. 3 along withevidence from the examination of the wear scars that ridgesdid not form at the lower loads examined and care wastaken in further testing (where it was hoped to ascertain thewear rate of the material itself) to be operating in a regimewhere ridge formation did not occur.
4.2. Development of wear with sliding distance
Fig. 5 shows examples of the development of wear withsliding distance for a polymer that was not exhibiting theformation of ridges, namely, PMMA under 0.5 and 1 N load.The linear behaviour indicates that the test is well-behaved,and that wear rates of the polymers may be derived fromplots such as these. It is of interest to note that there isno significant running-in period (i.e. the wear develops lin-early from the origin).Figs. 6 and 7show the develop-ment of wear in PMMA under a 1 N load (non-ridgingregime) with sliding distance.Fig. 6ashows the scar follow-ing only one ball revolution (∼80 mm); little obvious direc-tionality was observed in the scar, indicating that three-bodyabrasion was operating. As wear proceeded, the degree ofdamage increased significantly. Both pitting damage and di-rectional grooving can be observed in the later wear scars;however,Fig. 6d indicates that the wear damage is domi-nated by pitting indicating that the three-body mechanism isthe primary mode of wear. It may be argued that three-bodywear becomes more dominant as the test proceeds; suchthree-body rolling of abrasive particles is promoted by analready roughened surface of the testpiece. However, sucha shift in the dominant mechanism has not resulted in anysignificant changes in wear rate (Fig. 5). Whilst a changein mechanism with sliding distance has been observed, themechanism was generally the same at any point in a givenwear scar and the spatial separation of two-body abrasionand three-body abrasion within a wear scar, as observedby Trezona et al.[16], was not observed. After one revo-
lution (Fig. 6a), the degree of damage appears slight; how-ever, Fig. 7 shows that even after such small sliding dis-tances, material has been removed to a depth of∼5�m.The profile in this case shows significant deviation from thatof the assumed spherical cap (Fig. 7); however, after fiverevolutions of the ball, the spherical cap geometry is fullydeveloped with small deviations close to the surface associ-ated with sample scuffing[1]. Again, despite the fact that thedepth of wear was up to∼16�m after five ball revolutions(Fig. 7), the damage observed is relatively slight (Fig. 6b).
4.3. Microabrasive wear behaviour of different polymers
Fig. 3shows the wear rates of the different polymers undera load of 0.5 N which was the test load employed to ensureno ridging of any of the polymers. It can be observed that thetwo softest polymers (HDPE and PP, seeTable 1) also exhibitthe lowest wear rates and the two hardest polymers (PSand PMMA) also exhibit the highest wear rates. Classicalwear theory does not account for this. The Ratner–Lancastercorrelation between the polymer wear rates and (1/σbεb) isshown inFig. 9. Whilst the correlation is reasonable, someanomalies are noticeable. PC and HDPE have similar valuesof (1/σbεb), as do PMMA and PS. However, in each pair,there are very significant differences in wear rate, with thelowest wear rate in each pair being exhibited by the materialwith the largest tensile breaking strain.
Budinski [4] examined the correlation between hardnessand abrasion rate of a range of polymers, using both Shorehardness and Rebound hardness. He found no correlationbetween hardness and wear rate over the range of materi-als examined. However, in this work, a plot of the wearrate versus Martens hardness (Fig. 10) exhibits a reason-able correlation (better than the Ratner–Lancaster correla-tion) with the wear rate increasing with increasing polymerhardness. Early work by Lancaster[6] demonstrated a gen-eral decrease in abrasive wear rate of a range of polymers(two body abrasion against coarse carborundum paper) withincreasing hardness. It is difficult to rationalise these twoconflicting observations. In Lancaster’s work[6], the abra-sion was very aggressive and the wear rate of the polymersmay have been dominated by the depth of indentation of thecoarse abrasive particles. However, in the current work, it isobserved that a high hardness is linked with a low elongationto fracture. It appears to be this latter attribute which moststrongly governs the wear behaviour in microscale abrasiontesting.
Fig. 8 illustrates the morphology of the wear scars in theseven polymers. The three samples which show the highestwear rates (PMMA, PS and PVC) all exhibit a similar scarmorphology in which indentation type damage has occurredwith little directionality. These three materials also show thelowest values of tensile strain to failure. Such indentationtype damage may be associated with low ductility and atendency to crack under repeated indentation during testing.PC and PETG exhibit decreasing damage with an increasing
level of two body grooving and a reduction in the presenceof indentation type cracking. The wear rates of these twomaterials are very similar despite the fact that the strain tofailure for PC is significantly larger than that for PETG. Thewear scar morphology for the two softest polymers (HDPEand PP) are significantly different from the others and fromeach other. Both show high ductility with evidence of sub-stantial plastic flow. The HDPE in particular exhibits onlytwo-body grooving. Such grooving is promoted by the lowhardness and high tensile strain to failure.
The good correlation between hardness and resistance tomicroscale abrasion indicates that for resistance to damageby small-scale contaminants, softer polymers with largerstrains to failure are required. This work has demonstratedthat these can be selected on the basis of indentation hard-ness with some level of confidence. It is notable that amongstthe transparent polymers used for glazing applications, PCand PETG are significantly more resistant to abrasion thanPMMA or PS, and as such offer better long-term life with-out loss of optical function.
5. Conclusions
It has been demonstrated that, with care in choosing thetest conditions, the microscale abrasion test can be employedto investigate the abrasion rate of polymeric materials withabrasives of very small size. In spite of concerns that theabrasive wear behaviour of polymers may be very differentwith small abrasives than with large, traditional correlations(such as the Ratner–Lancaster) are reasonably successfulwith these test conditions.
The wear behaviour and rates of polymers depended crit-ically on the polymer type. High wear was associated withindentation-type morphology in the wear scar and low val-ues of tensile strain to failure. Polymers with high strains tofailure exhibited less indentation dominated wear. In the caseof HDPE, a solely grooving type mechanism of wear wasidentified. It was demonstrated that there was a good corre-lation between polymer hardness and polymer wear rate.
Under certain conditions, ridges formed in the wear scarsinvalidating the test. The tendency for ridge formation dif-fered for the different polymers. Whilst no explanation wasoffered for these differences, it was noted that ridging wasnot observed in the partially crystalline materials, whereas
ridging always occurred at intermediate loads for the trans-parent polymers.
Acknowledgements
Support from the Engineering and Physical Sciences Re-search Council (UK) (Grant Number GR/M92836), Plintand Partners Ltd. and Phoenix Tribology Ltd. is gratefullyacknowledged.
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