Material-specific wear mechanisms: relevance to wear modelling

169

Wear, 141 (1990) 169-183

Material-specific wear mechanisms: relevance to wear modeUing* J. K. Lancaster School of Engineering and Informatio Reading RG6 2AY (VK.]

Scimces, University of Reading, Whiteknights,

(Received February 26, 1990; accepted June 19, 1990)

Abstract Wear can be classified in a variety of ways but none has yet proved to be completely satisfactory. It is now becoming evident that a clear distinction must be made between the different processes responsible for detaching particles from a surface and those involved in the production of debris. Particle detachment processes are reviewed for four groups of materials-metals, polymers, ceramics/glasses and carbon/graphites-with particular emphasis on the role of the properties of the various materials. Particle production via crack propagation appears to be a common feature with many materials, but appropriate models still await development. These phenowmm am all interrelated with causes and co-n&ions, are brought abou by them, and disappear as the causes and ctm.ditions change (The Teach&g of Bra 175th edn., Togo, 1981).

1. Introduction There are two separate routes to the modelling of wear, depending upon the requirement. If the primary need is for design information, relevant experimental data can be extracted from the literature or derived from experiments sprang the intended application as closely as possible. These data are then fitted to empirical equations and extrapolated or interpolated as necessary via conventional statisticdl procedures. Examples of this type of approach have been used, for example, in work on brake materials [ 1 ] and on self-lubricated rolling-element bearings [ 21. The alternative and more fundamental approach is to seek information on wear from experiments in which the variables involved have been identified and controlled. The resulting wear models are then usually rather simple and somewhat restrictive in scope. Subsequent expe~en~tion almost invariably reveals discrepancies and when the reasons for these have identified, the model can then be refined. This approach is essentially a long-term one and tends to be open ended. It does, however, have the great advantage that the individual contributions are all *Modified version of a paper presented at the International Workshop on Wear ModeRing, Argonne National laboratory, Chicago, U.S.A., June 16-17, 1968.

160

additive and lead to a gradually increasing level of understanding. Typical examples are the early studies of mild and severe wear by Archard [3] and by Archard and Hirst [4], and of abrasive wear by Kruschov and Babichev [51. This review of wear processes follows the philosophy of the second approach described above: that is, it is concerned primarily with the types of wear which occur in well-controlled conditions of sliding and with the simple models to which they lead. It is not intended to provide details of all the various types of wear which have been identified for different groups of materials, but merely to illustrate some of the different trends in behaviour which more sophisticated models will ultimately have to take into account. To begin with, however, it is appropriate to consider the various ways in which wear can be categorized.

2. Wear classifications Attempts to agree on a formal classification of wear processes have been continuing for over 30 years. The problem is complicated because wear is a property of the system [6]; that is, it depends as much on the imposed sliding conditions as on the properties of the materials concerned. The various classification schemes which have been proposed in the past appear to fall into four main groups, as shown in Table 1. The first two of these arose during early research into wear phenomena (e.g. refs. 4, 5, 7 and 8) whilst the last two have become more prominent following the advent of sophisticated techniques for the examination and characterization of worn surfaces (see e.g. refs. 9 and 10). Tabor’s suggestion [ 111 of dividing wear processes into “adhesive”, “non-adhesive” and “mixed” effectively spans the second and third of the groups in Table 1. All of these previous schemes, however, are unsatisfactory in one fundamental way: they fail to make a clear distinction between the primary processes involved in the detachment of particles from a surface and the subsequent ones which ultimately lead to the production of debris [ 121. It is now becoming widely recognized that the particles TABLE

1

Types of wear classification conditions Rough/smooth surfaces,

Imposed

corrosive

environments,

abrasive

particles

PhenomeruL Mild/severe,

oxidation,

Debris production Delamination, Interfacial/bulk

seizure

route

cutting,

Mode of asperity

transfer, tearing,

fatigue,

cavitation,

spalling,

contact

deformation,

elastic/plastic

displacement

pitting,

adhesion

161

detached from a surface seldom emerge directly as debris. They are usually trapped within the contact zone for varying periods of time during which they are subjected to a variety of processes, such as comminution, chemical reactions and aggregation, and ultimately form a so-called “third body” within which the relative motion of the surfaces must be accommodated [ 131. Debris emerges from within the third body via processes which may or may not be the same as those originally involved in particle detachment. An example from early work illustrating this behaviour will be given later. The wear classification still most widely used at the present time is the one based on the ideas of Burwell and Strang [ 141 with four principal wear mechanisms: abrasion, adhesion, corrosion and fatigue. Each of these terms has now been given [ 151 a formal definition but it is important to remember that they do not uniquely define the various ways in which particles can be produced. In any one set of imposed sliding conditions, particle detachment and debris frequently result from a mixture of two or more of the above processes operating either simultaneously or successively for limited periods of time. An attempt to show some of the interrelationships involved is given in Fig. 1 which starts with an identification of the mode of asperity deformation on the left-hand side and ends, on the right-hand side, with more descriptive terminology relating to different wear phenomena. Even this, however, is still not wholly satisfactory because it fails to take into account the properties of the individual materials concerned. For example, abrasion is defined [ 151 as “wear by displacement of material caused by hard particles or protuberances”. The consequences of abrasion, expressed in this way, can vary very widely with different groups of materials, as shown in Table 2. The mechanisms of particle detachment can range from tearing and fatigue for rubbers, through cutting and chip formation for ductile metals and plastics, DEFORMATION

PARTICLE DETACHMENT

INTERFACE MODIFICATION

DEBRIS PRCDUCTN3N

PHENOMENA

PLASTIC MATERIALS

ADHESCN

/ \

/ ELASTIC \

FATIGUE

DELAMINATION

Fig.

1. Interrelationships between wear processes.

162 TABLE 2 Abrasive wear mechanisms for different groups of materials Material

Deformation mode

Particle detachment processes

Material properties

Metals

Plastic-elastic

Plastic grooving Prow formation Cutting (chip formation)

Hardness Ductility

Polymers

Elastic-plastic

Plastic grooving cutting Fatigue

Toughness (se) Fatigue properties

Elastomers

Elastic

Tearing Roll formation Fatigue

Tensile strength Elastic modulus Fatigue properties

Ceramics (carbons)

Elastic Brittle fracture

Crack propagation Flaking Fatigue

H/E ratio

Fracture toughness

to crack formation and brittle fracture for ceramics and glasses. A similar situation is also likely to apply to wear involving adhesion and corrosion. Terms such as “adhesive wear” etc. are thus very imprecise when used in isolation. The long-term goal must therefore be to devise a wear classification scheme based primarily on the mechanisms by which particles are detached, coupled with details on how specific geometrical arrangements then respond in producing debris. The review which follows is concerned mainly with the first part of this, i.e. particle detachment mechanisms. Information on the second part is now emerging from recent work by Godet and coworkers [13, 16, 171. 3. Metals In the 1950s and 1960s considerable attention was paid to the types of wear occurring in relatively soft metals sliding against a smooth harder counterface. Two broad types of wear phenomena were identified [4] and described as “severe” and “mild”. The characteristic features of severe wear were high wear rates, extensive plastic deformation, transfer to the harder counterface and flake-like metallic wear debris. Those of mild wear, in contrast, were low wear rates, minimal plastic deformation, surface film formation protecting against metal-to-metal contact and oxide wear debris. Details of the mechanisms of both types of wear were elucidated and are now briefly described. 3.1. Severe wear This type is typically found at high loads and low speeds (and low temperatures). The wear rate is generally directly proportional to the applied

163

load and inversely propo~ion~ to the hardness when this is changed by varying the temperature [ 181. These relationships are in full agreement with the original Archard [3] model: v -=s

KL4 H

where V is the wear volume, s is the sliding distance, L is the load and W is the hardness. A wide range of merent combinations of materials can exhibit severe wear under suitable sliding conditions, as shown in Fig. 2 where the wear rates are plotted against hardness. Most of these combinations satisfy the various criteria which have been proposed to relate so-&led “adhesive wear” to metallurgical compatibility [ 19-211 but there are exceptions: e.g. neither cadmium nor bismuth exhibit much solid solubility in iron. This suggests that the transfer of fragments to a counterface may not be attributable solely to adhesion, and some form of mechanical asperity 1 O1 \

)
Wear Rate (mm3JNm)

l-

ll5

lli

.l _

2,

1

in/Pb

14

60-40

2

PbJCu

15

PtJNi

brassITS

3

CdJTS

16

ZrJTS

4

AglCu

17

FeJTS

5

AZITS

16

NiJTS

6

BiJTS

1g

At AlloyJTS PUTS

7

ZnJTS

20

8

CulNi

21

10% A&CuiNi

8

MQJCU

22

MoJCr

10

MgJA( Alloy

23

TilTS

11

AuJNi

24

MoJTS

12

MgJTS

25

WJStellite

13

SnJTS

26

Stellite

12

(TS = 18%W tool steel, 700 VPN 0.07pm Ra)

11 pj \ 13

V=O.lSm/s,

L=i-250N

14 15 \ IO

,g::,S

320

2’\

22 ?3 24

4,

\

1;

2526\

Indentation IO

hardness

Slope -1.70

(VPN)

5 1

10

10”

100

104

F’ig. 2. “Severe” wear rates of various combinations VS. hardness.

20

164

interlocking [22] is also likely to play a part. For all these metal combinations, Fig. 2 shows that the wear rate decreases more rapidly than proportionately with hardness. Rabinowicz [23] has also reported a similar, more rapid, decrease for unlubricated similar metal combinations. These deviations from the predictions of the simple Archard model imply that changing from one metal to another brings in other parameters which influence wear apart from the hardness, e.g. crystal structure, toughness, ductility and strain hardening. The model therefore needs extending to incorporate these aspects. A detail@ examination of the severe wear process for one particular combination-60/40 brass with 2% lead sliding on “Stellite” -provided one of the earliest examples of the distinction between the processes of particle detachment and debris production [ 241. Initially, relatively small fragments are transferred to the harder counterface and these transferred lumps then increase in size by successive preferential transfer until a critical size is reached beyond which they are detached as flake-like debris. Histograms illustrating the relative sizes of the original asperity contacts (static), the transferred fragments and the wear debris are given in Fig. 3. The primary process leading to transfer has not been unambiguously established but adhesion coupled with junction growth must almost certainly play a major part. It may be recalled that the original interpretation of the proportionality constant K in the Archard wear model was that it represented the probability of detaching a fragment during a single asperity encounter. In the light of later evidence it is now more reasonable to envisage adhesion being supplemented by crack propagation and fatigue during repeated stress cycles. The large size of the transferred fragments compared with the initial contact sizes (Fig. 3) is consistent with the idea of a contribution from fatigue. Halling [25] has developed a wear model incorporating low cycle fatigue concepts, together with metallic properties such as ductility and strain hardening, which in its simplest form reduces to the same predictions as the original Archard model. The mechanism of removal of the transferred aggregates from the counter-face as debris is also somewhat ambiguous. There is evidence I261

Fig. 3. Sizes of (a) asperity contacts, (b) initially transferred fragments and (c) wear debris; brass on stellite; 2.5 kgf load; from ref. 24.

165

for crack propagation at or near the original interface with the counterface and crack propagation may also be enhanced by oxidation (stress-corrosion) [ 111. Alternatively, Rabinowicz [ 271 suggests that the gradual accumulation of strain energy during repeated contacts could ultimately provide sufficient interfacial energy to enable the transferred aggregate to be detached.

3.2. Mild wear This regime is less well defined than severe wear and a range of different mechanisms has been identified [28]. In the simplest, e.g. brass on tool steel [26], the formation of an oxide film limits the extent of inter-metallic contact and the wear rate is lessened through a reduction in the scale of the severe wear process. This mechanism is likely to be most prevalent with reactive (to oxygen) metals at relatively low speeds where frictional heating is slight. Under these conditions, the variation of mild wear rate with hardness (Fig. 4) is not greatly different from that of severe wear. Such a trend would be most improbable if contacts were occurring only between completely oxidized surfaces. At higher speeds, where frictional heating becomes significant, the role played by inter-metallic contact diminishes, except during the early stages of sliding. For soft tool steel sliding on hard tool steel radio tracer experiments (291 have again been able to distinguish between particle detachment and debris production. The wear process involves three stages: transfer of metal fragments to the counterface, oxidation of these fragments and removal of this oxide by attrition. The rate-determining stage is the rate of oxidation. For hardened tool steel sliding on itself [30], however, intermetallic contact and transfer are restricted mainly to the very early stages of sliding. Following this, the surfaces gradually acquire a high degree of conformity and, after reaching steady state conditions, the rate-determining process then becomes one of three-body abrasion by oxide debris. A variety of models have been proposed for the growth and removal of the oxide film during oxidational wear and reviews have been given by Quinn [31] and Sullivan [ 321. With steels, the consensus of opinion seems to be that the oxide film grows over a number of discrete plateaux to some critical thickness and is then detached more or less in its entirety. The reason for the sudden detachment of oxide flakes has not been fully established but is most probably associated with a combination of mechanical and thermal fatigue [33]. In so far as oxide growth is concerned there are two main problems. First, the rate constants governing the growth of oxide during sliding will be very different from those appropriate to static conditions and can only be estimated indirectly [34]. Secondly, there is conflicting evidence as to whether the most important temperature governing oxidation is the localized contact ‘flash’ temperature or the mean surface temperature at which a longer time is available for oxidation. The former view is supported by experimental work on the oxidational wear of steels [35] whilst the latter appears more appropriate to the wear of ahiminium alloys [ 361 and the copper-iron system [37]. In general terms, out-of-contact oxidation is likely

166

WEAR

RATE

(mm3/Nm)

10-l 2

\ 1

3

4 5

d

‘.\..‘6

InlTS Pbl NI Pb/ TS

4 5 6 1 8 9 10 11 12 13 14 15 16 17 18 19

Sn/ Cu Sn/ TS Sn/ Ni A! / TS Agl A! alloy Agl TS MglCd Pt ITS Cu I Ni Cu /TS NI/TS Fe / TS Zrl TS lO%AP-Cu/Ni va ITS MO/ TS

( TS: 18%W TOOL STEEL 700 VPN, 0.07m Ra) V. O.l8m/s, Lx l-250N

7 10

8

1 2 3

1o-3 9

\

15 I4 16

11

17 18 \

12 13

19

d \

SLOPE - 1.42

I

10 INDENTATION Fig. 4. “Mild” wear rates of various

lo2 HARDNESS

combinations

IO3 (VPN)

vs. hardness.

to be most important when the ambient temperatures are high and the

imposed loads and speeds are relatively low. In view of all the above complexity, it is hardly surprising that many of the relationships between mild (oxidational) wear and the parameters controlling sliding also tend to be complex. One such example is ahuninium sliding on hardened tool steel. At low speeds of sliding, the wear volume-time relationships are irregular and typically of the form shown in the insert (a) in Fig. 5. When the mean slope of such lines (the wear rate) obtained at different loads is plotted against load, the result is the non-linear relationship given by the solid line in Fig. 5. The wear debris is a mixture of metal and

167

WEAR 1 mm3/m

RATE I SEVERE /

1o-2

10-j

r,

lo-'t

5

c

lo-!

(a)

t

(b)

t

L

1

10 LOAD

100 (N)

Fig. 5. Wear-time and wear rate-load relationships for alurninium on tool steel.

oxide particles and the wear process seems to involve mild and severe wear regimes occurring successively over short intervals of time. At higher speeds of sliding it becomes possible to resolve these separate regimes, as shown by insert (b) in Fig. 5. During the severe wear regime, the friction is high (CL=0.8), the surface temperature rises, the rate of oxidation increases and a protective f&n is ultimately formed largely preventing metal-to-metal contact. Mild wear then ensues but because the coefficient of friction is now lower (cc= 0.4) the temperature and rate of oxidation are no longer sticient to maintain the oxide film and penetration eventually occurs leading again to severe wear. When the wear rates in the two regimes are separately plotted

168

against load the resulting relationships are linear (hatched lines in F’ig. 5) and evenly dispersed above and below the non-linear curve characteristic of low speeds. It is tempting to speculate that many other non-linear relationships between wear and time or wear rate and load might be interpreted in a similar way but experimental proof will depend on being able to achieve an adequate resolution of the Merent wear modes as they occur.

4. Polymers The dominant property of polymers, as far as wear is concerned, is their relatively low moduli of elasticity compared with metals. During sliding against all but extremely rough surfaces, the localized deformation of the asperity contacts for polymer/metal or polymer/polymer combinations is thus likely to fall within the elastic regime [38]. The most general attempt at classi@ing polymer wear is that by Briscoe [39], shown in Fig. 6, and each of these two broad regimes-bulk deformation wear and interfacial wear-will now be discussed. 4.1. Bulk dqfmtion wear Abrasion falls within this category and the simplest model of abrasive wear in polymers is the one by Ratner et at, [do]. They suggest that three consecutive stages are involved in the detachment of a particle: defo~ation ENERGY

DISSIPATION

*d

INTERFACE

BULK

g

COUNTERFACE

MOTlON

-L----/I

/ pLA7c I

LOSSES,,

SLIDING I

v’SCoELIAST’C

TRPE I

\ lNlERFACE,ZO~E 1

1

&

ik

THERMAL

F3g.

j

OEGRADAI8ON, MELTING

6. Modes of energy dissipation and wear for polymers; from ref. 39.

169

by the penetration of an indenting asperity which is opposed by the hardness rr; relative motion opposed by the frictional force F= &; and tial de~c~ent involving the work to fracture approximated by the product se (s is the rupture stress and e the strain at rupture). The volume removed per unit sliding distance is thus

v 1 -cxpd S

se

The ways in which these parameters each vary with temperat~e are shown in Fig. 7(a) and Fig. (7b) for amorphous and crystalline poiymers respectively, and the predicted wear rate variations are given by the full lines. There has been some, albeit limited, verification of these trends [42]. The dominant material properties lie in the product se and the wear rates of a range of polymers sliding in single traversals over rough steel correlated generally with l/se. There are, however, complications in some instances. For a series of y-damaged pol~e~~uoroethylenes (PTFEs), Briscoe et al. 1431 found that the wear rates against abrasive papers increased less rapidly than proportionately with l/se, or even with l/s’e where the extra s was introduced as an approximation for the hardness. To explain this trend, they modified the original Ratner argument and introduced an extra term- the damage efficiency- to account for the fact that only some of the strain at rupture a

AMORPHOUS

POLYMERS

Fig. 7. Predicted wear rate-temperature semicrystalline polymers; from ref. 41.

variations for (a) amorphous polymers and (b)

170

is involved in particle detachment and a significant part can be recovered elastically. The damage efficiency term also takes account of the fact that debris within the contact -even during single-pass abrasion-can reduce the effective load and, in turn, the wear rate. The wider applicability of this damage efficiency term to other polymer systems still remains to be explored. There is considerable circumstantial evidence suggesting that during sliding against counterfaces which are insufhciently rough for the elastic limit to be exceeded during asperity deformation, particle detachment arises from fatigue. Direct evidence, however, is scarce; developing cracks in polymers on an asperity scale are difficult to detect because of elastic recovery, although they have been observed during fretting [ 441. Two examples of the circumstantial evidence are the following. When polymers slide in single traversals over metal surfaces of different roughnesses, the wear rate increases very rapidly with increasing roughness (Fig. 8(a)) and this trend is wholly explicable in terms of a fatigue wear model [45]. The simplest, based on elastic contact between a single rigid hemisphere of radius r indenting a polymer flat of

I

NYLON, 4

2

3

POLYACETAL,

POLYETHYLENE,

7

WEAR

POLY

(METHYL

s POLYPROPYLENE,

METHACRYLATE), 6

POLYSTYRENE,

PTFE

RATE

mm31 Nm

IO-;

16:

IO-

IO-

IO-

8. Variation of wear rate with counterface roughness R. and average radius of counterface asperities r,, during single traversals over mild steel; from ref. 46.

Fig.

171

modulus E under a load L, leads to _V c[ r-

2(t - I)/3

~‘t

+ 2)/3

~2(t

- 1)/3

S

where t is the exponent in the Wohler relations~p

n being the number of cycles to failure, ~7the applied stress and u. the faihrre stress during a single cycle. From computer analysis of the topography of metal counterfaces, the average values of r can be determined, and from the slopes of the Iines in Fig. 8(b) values of t can be derived. These are generally broadly similar in magnitude to those found in co~vention~ fatigue tests. More sophisticated wear models have been developed by Kragelski and Nepomnyaschi [47J based on topographies characterized by the bearing area curve and, more recently, by Jain and Bahadur [48] who assume a gaussian height distribution of the contacting surfaces. It may be noted that the latter model, in contrast with the earlier ones, predicts a Iinear relationship between wear rate and load, as is usually observed experimentally [38]. The second model of fatigue wear in polymers approaches the problem from the standpoint of fracture mechanics [49j. The wear rate is assumed to be inversely proportional to the number of cycles to failure and this can be derived by integrating the Paris equation for fatigue crack growth. The final result for the wear rate shows that

where a,, is the- initial crack length, Ag is the imposed stress range, and A and n are the constants in the Paris equation da dnr =A(AK)" and

AK= Aa(rra)‘”

This model has been verified experimentally by determining A and n from cyclic loading .of polyethersulphone samples immersed in various fluids. Values of A range over several orders of magnitude but those of n are more restricted. After inserting these values into the above equation and e~at~g the constant of propo~io~~ by fitting the results to one particular fluid, it is possible to compare the predicted wear rates in the other fluids with those determined experimentally. As shown in Fig. 9, the agreement is very reasonable. By interpreting wear in this way, it now becomes possible to explain why polymer wear rates are affected by material properties such as the degree of crystallinity, molecular weight and, for thermosets, cross-linking density. All these influence fatigue crack propagation rates by modifying the ener~-~sipat~ mechanisms which affect dustily at the crack tip. For very low modulus polymers - elastomers-deformation wear against rough surfaces is generally associated with elongation of the material at the

172

WEAR 0

RATE

(~o~‘md/Nm)

I

2

3

4

n-HEXANE

n- HEXADECANE

0

THEORETICAL

tZ.l

EXPERIMENTAL

CYCLOHEXANE

TETRACHLOROMETHANE

XYLENE

TOLUENE

ACETONE

n -PROPANOL

ETHANOL

FORMAMIDE

t INCREASING SOLUBILITY PARAMETER

Fig. 9. Comparison of experimental wear rates with predictions from crack-growth polyethersulphone against stainless steel in various liquids; from ref. 49.

data:

rear of individual asperity contacts followed by tearing and recovery of the remaining material into a lip [50]. Repeated events of this type eventually lead to the formation of rows of ridges, transverse to the direction of motion- the so-called “abrasion patterns” [ 511. The tearing process during successive contacts can be modelled as a propagating crack and it thus becomes possible to apply fatigue and fracture mechanics concepts to rubber wear [52]. Various attempts have been made to relate the wear of elastomers to their physical and mechanical properties, such as Shore hardness, tensile strength, resilience etc. [53-551 but the empirical relationships proposed only have limited applicability. The main reason for this is that rubber surfaces become extensively degraded during sliding, by processes such as chain scission or free-radical reactions, leading to surface layers whose properties are very different from those of the bulk materials. Such surface changes also tend to be a major feature of the wear of many materials in the inter-facial wear regime, as will now be described. 4.2. Interfacial wear This regime is typical in general of polymers sliding against smooth hard counterfaces where energy is dissipated primarily by overcoming adhesion at the asperity contacts. The primary mechanism involved in detaching fragments from the polymer is probably fatigue, as discussed earlier, but with the localized stress distributions modified by traction. The overall wear

173

process, however, is seriously complicated by the changes which occur in the surfaces, and primarily in that of the counterface, during sliding. Some of these are shown in Fig. 10 and their main effect is to modify the topography, the localized stresses and, in turn, the rate of crack propagation and wear. The two main types of counterface modification are transfer and abrasion/ polishing. Amorphous polymers below their glass transition temperatures, e.g. polymethylmethacrylate (PMMA), usually transfer only slightly, if at all, and these materials are amongst the few which may exhibit true interfacial sliding. For semicrystaline polymers, transfer is of two types. With nylon 6.6, polypropylene and similar materials, fragments adhere to the counterface and lead to irregular films of the order of 0.1-l pm thick. When deposited on metal surfaces which are initially very smooth, such films may lead to increased stresses and wear [ 421. It has not yet been unequivocally established whether these transfer films are in dynamic equilibrium, i.e. they are being replenished at the same rate as material is being lost, nor whether the debris originates only from within the film as in the severe wear of metals. The latter does, however, appear very probable as the debris particles are frequently relatively large in size and flake-like in appearance [ 561. The second type of transfer occurs with PTFF and high density, or ultrahigh molecular weight, polyethylene (HDPE). Here the films tend to be much thinner (10-50 run) and very much more uniform. They probably

COUNTERFACE WEAR

RATE

ASPERITIES

PENETRATE

CHARACTERISTIC

(MICRO-CUTTI

POLYMER;

OF INITIAL

NC, Low

CYCLE

ROUGHNESS

FATIGUE)

I COUNTERFACE

I ROUGHNESS INCREASES

ABRASION/ CORROSION -

FRICTION POLYMER/ REACTION PRODUCT _ It’ ROUGi-iNESS DECREASES

I WEAR INCREASES nlCnOcuTTIH~ (

MODIFIED

TRANSFER I ROUGHNESS INCREASES

I

WEAR DECREASES (FATICUL)

BY

DECREASES c

HYohOoYnAHIC LUBRICAlIOW

Fig. 10. Changes occurring in a counterface during sliding and their effects on the wear of polymers.

174

originate via the extrusion of sheets or fibrils from these polymers, which is possible because of their smooth molecular profiles [ 57). Again, it is probable, although not conclusively proven, that the origin of the debris is within these films. In addition to reducing the roughness of the counterface, the uniform transfer films from HDPE and PTFE generally reduce the coefficients of friction via preferential orientation of the molecular chains. This will also lead to lower wear by reducing the contact stresses. Damage to a metal counterface during sliding against polymers normally arises either from the ingress of adventitious abrasive particles or from the presence of fillers or reinforcing fibres in the polymer. In view of the extreme sensitivity of fatigue wear to the level of the counterface roughness (Fig. 8) it is clear that even very small improvements in the counter-face finish will lead to large reductions in wear rate [46]. In developing any general model for inter-facial wear, a topic rather neglected so far, it will thus be necessary to take into account those parameters which can lead to counter-face abrasion or polishing, e.g. the shape and size of filler particles and their hardness relative to that of the metal.

6. Ceramics and glasses

Following a considerable amount of research over the past few years, (e.g. refs. 58-62) understanding of the wear processes in ceramics is now approaching a level similar to that achieved for other groups of materials. Detachment mechanisms appear to be of four main types. (a) When the contacting surfaces are relatively smooth and deformation is localized within a very small volume, the superimposed hydrostatic stress field [63] permits plastic flow and grooving. Whether or not this makes a significant contribution to the total wear is arguable. (b) For contact stresses still below the elastic limit crack propagation occurs during repeated contacts leading to particle detachment via fatigue or delayed fracture. (c) With rough surfaces where the depth of deformation exceeds some critical value [64], cracking and brittle fracture may occur on a macro-scale. Radial cracks develop during the loading part of the cycle followed by lateral cracks during the unloading part. It has been suggested [65] that material is removed when these two types of cracks intersect. Crack propagation may also be thermally activated and empirical relationships have been found [66] between the wear rates of a range of ceramics and a thermal stress resistance parameter. (d) Finally, with certain materials such as Sic and Si3N4,there is a type of corrosive wear in which products of reactions with the environment are continuously removed during sliding [ 671. The first two of these processes may be illutrated by some recent work on the wear of glass [68]. When soda-lime glass slides against a smooth

175

tungsten carbide/cobalt counterface, the worn surfaces frequently exhibit extremely smooth areas crossed by fine parallel grooves. These grooves generally originate from small pits and it is thought that they result from plastic grooving by the debris particles coming from the pits. Of much more signilicance for the overall wear rate, however, is the damage resulting from crack propagation and brittle fracture. A simple model has been derived by considering crack growth beneath an elastically deformed hemispherical asperity during successive contacts at a rate controlled by the stress level and the frequency of contact. It should be noted that glass is not subject to dynamic fatigue in the conventional sense; cyclic stressing does not enhance the crack growth rate and the failure process is one of “delayed fracture” or so-called “static fatigue” [69]. Under constant sliding conditions, the analysis predicts direct proportionality between the wear rate and the rate of slow crack growth. Slow crack growth in glass is known [ 701 to be very sensitive to environmental humidity and temperature, and Figs. 1 l(a) and 1 l(b) show that the wear rates of soda-lime glass vary with these parameters in almost the same way as the crack growth rates. The way in which the crack growth rate varies with stress intensity further leads to the prediction that the wear rate should increase rapidly with increasing counterface roughness, as in the fatigue wear of polymers (Fig. 8(a)). Figure 11(c) confirms this trend experimentally, although only up to some limiting roughness. Beyond the limit it is suggested that the stress intensity has reached the point at which the mode of crack propagation changes from being reactionrate controlled to diffusion controlled. The latter is largely independent of the stress intensity [ 701. Wear between rough surfaces, supposedly resulting from the intersection of radial and lateral cracks, has been modelled on several occasions in terms of fracture toughness Kc. One model, for example [65], relates the rate of wear to the inverse of Kc314Hln. However, as Ajayi and Ludema [ 711 have pointed out, these models are not in general well supported by the experimental evidence for two main reasons. Firstly, the scale of damage occurring during wear is generally appreciably smaller than that likely to result from the linking of lateral and radial cracks. Secondly, the mechanisms of material removal from ceramics are strongly dependent on their composition and microstructure and thus tend to be specific to individual materials. Grain boundary phases, for example, control the deformation and microfracture of Si3N4whereas with A1203 there is plastic deformation and slip within the grains themselves. Chemical reactivity with the environment and its influence on wear also tend to be material specific, the effects of water being particularly confused. With alumina, for example, water is known to accelerate subcritical crack growth [72] but rates of wear in the presence of water can be either higher [73] or lower [ 741 than the values in air. Similar ambiguity exists for zirconia; water can either increase [75] or decrease [76] wear. Only with the silicon-based ceramics-Si,N,, Sic and Sialon-is the effect of water on wear reasonably consistent. Wear rates are reduced in water because tribochemical reactions produce a relatively soft and deformable surface layer

176

(b)

10-6 00% t”l!ldi

Ccl

.urtaw

01 :eughncr*

01 ctdllde

10 Ccumer’arc

I!.mRal

Fig. 11. Variation of the wear rate of glass on tungsten carbide with (a) relative humidity; (h) temperature fin water); (c) counterface roughness; from ref. 68.

which Iowers the coefficient of friction and inhibits microfracture [67]. At the limit, wear can result entirely from the removal of reaction prociucts-corrosive wear in conventional termindogy.

177

6. Carbons

and graphites

These are an interesting subgroup of materials which, because of their relatively low elastic moduli compared with metals, present similarities to both polymers and ceramics in their wear behaviour. Taking the similarities with polymers first of all, graphitic materials transfer to a counter-face and the transfer film plays an important role in determining the magnitude of the friction and wear rate 1771. There is some direct evidence that the films are in dynamic eq~b~~, as shown by the radio tracer experiments illustrated in F’ig. 12. The transfer films develop from particles which are broken down almost to the size of the crystalline unit cell and which are subsequently aggregated and oriented to give a degree of graphitic order which is often greater than that of the original material itself [78]. Large-scale pieces of debris apparently come from within the films, as a consequence of the gradual build-up of compressive stresses leading to “blistering” and fracture. The trends in wear rate with counter-face roughness are also broadly the same as those found with polymers 1791 and can be interpreted in terms of fatigue in a similar way. Direct evidence of fatigue wear in carbons and graphites has been obtained from experiments using reciprocating crossed cylinders [781. The failure modes observed during sliding are generally similar to those

for ceramics and involve brittle fracture; there is little evidence of plasticity. Various typical features of worn surfaces (for a low graphite carbon sliding against itself) are illustrated in F’ig. 13 which shows, in order, (a) flaking of surface films; (b) preferential removal of intergranular material; (c) crack propagation within grains; (d) chipping at the edges of cracks or grams; (e) subsurface fracture of whole grams; Counts/s above background

xcXA

0.4 -

0.3 -

x

x

1

0.2 -

C

B

0

0.1 - I

F

’

8

L

2

4

6

I

t

8 IO Time (h)

t

,

12

14

FTg. 12. Formation and replacement of the transfer film for a carbon (C8) pin on a copper ring (3 N, 1.9 m s-l); radioactivity of Na” impurity monitored. A, Active pin on fresh copper; B, active pin on inactive transfer filrq and C, inactive pin on active transfer film.

178

X600

b.

X280

c. x1000

d.

X5700

e. XI00

f.

x12000

a.

F’ig. 13. Worn surfaces of a low graphite carbon after sliding against itself; from ref. 80.

179

(f) removal of material on an extremely small scale of size to give very smooth areas. It might be supposed that process (e) - the removal of whole grains - will make the greatest contribution to the total wear and that all the other processes, particularly process (f), can therefore be ignored. This is a dangerous oversimplification, however. Process (f) makes an important indirect contribution to wear because, as a result of the development of such smooth areas, the real area of contact increases and the ac~ornp~~g increase in friction greatly increases the likelihood of grain fracture. Through recognizing this fact, it has proved possible to reduce wear appreciably by preventing the development of these smooth areas [8 1 J. The unique feature of the wear of carbons and graphites lies in their sensitivity to adsorbed vapours from the environment. Below a critical partial pressure of such vapours (water or hydrocarbons), the wear rate increases dramatically by many orders of magnitude - “dusting” [82]. The sequence of events involved is given in Fig. 14. Above the critical partial pressure, which varies with the type of vapour involved, dusting is prevented by SLIDING

ENVIRONMENT

increased severity of sliding conditions

Reduction in partial pressure of water vapor, p

p/pa falls below critical value

c Increased adhesion t Increased friction # Disruption of surface fifms Increased roughness Fewer con&t

areas

I Higher contact temperatures

z Higher contact stresses & Brittle fracture of asperities 4

Carbon properties fncreased surface roughness

3 : STEADY-STATE

I

DUSTING

Fig. 14. Sequence of events involved in the transition to dusting of graphitic carbons; from ref. 83.

180

adsorption on to active sites generated at the crystallite edges by wear. Water is adsorbed directly on to these sites but hydrocarbons, which are adsorbed more readily on the basal planes, can migrate to the edge sites as required [ 841. The main parameter governing the onset of dusting is the asperity contact temperature which controls the amount of vapour which can be adsorbed [85]. The mechanism of wear within the dusting regime is also believed to involve crack propagation which is enhanced by the presence of vapours [86]. The dusting wear rate increases with increasing partial pressure of vapour (provided that this is always less than the critical value at which dusting ceases) and the debris size and the scale of surface damage increase similarly. The effectiveness of a particular vapour in promoting crack growth and increasing the rate of dusting wear is directly related to its absorption energy.

7. Concluding comments Despite the very wide disparity in the processes by which particles can be detached and debris produced from different materials, it is tempting to see if any common features can nevertheless be discerned. The first such feature appears to be that at one stage or another crack propagation comes into virtually all the wear processes which have been described, as well as into some, such as delamination wear [lo] which, for lack of space, it has not been possible to cover. In some instances wear rates correlate with fatigue data, but in many they do not because the composition and structure of worn surface layers can be very different from those of the bulk material and will therefore be associated with different fatigue properties. Although a more f~d~en~ approach is to relate crack propagation to wear by the fracture mechanics route, here again there are problems [87]. In the first place, the critical crack length, below which it is no longer justifiable to apply linear elastic fracture mechanics, is often much greater than the size of the asperity contacts, the transferred fragments and even the wear debris. Secondly, there is the difficulty of the crack initiation phase. With polymers this appears to be negligible [491 but with steels it can become very significant ]88]. Thirdly and finally, is the role played by the microstructure in crack propagation. This must become increasingly spout when, as in sliding, the crack lengths associated with particle detachment are commensurate in size with the asperity contacts and with the individual grains. Microstructural aspects still await inclusion in more sophisticated fracture mechanics models of wear processes involving crack propagation and fatigue. The second, more or less common feature in most of the wear processes in non-metallic materials, is the marked sensitivity of the wear rate to the level of surface roughness. Whilst an initial roughness value can easily be ~co~o~ted into any wear model, it is much more dBcult to include the roughness obtained during steady state conditions. To achieve this implies being able to predict how surfaces will run in, or what sort of third-body

181

films will develop, from a knowledge of the original composition and microstructure of the material and the geometrical conilguration of the contact. The generalrole of thirdbodiesin wearprocesseshas been extensively studied in recent years by Godet and coworkers (e.g. refs. 89 and 90) to elucidate the mechanics of their formation and destruction and how they influenceaspectssuch as load support and velocity adaptation. Understanding of the materials science relating to third-body films, however, is less well advanced. Crucial to further progress in this area is some means for characterizing the appropriate chemical, physical and mechanical properties of very thin surface films. The final point worth making is on the ubiquitous and largely indirect influence of adhesion on wear. Indirect, because as pointed out on many occasions, overcoming adhesive forces during sliding leads only to transfer; the formation of loose particles requires some additional mechanism. The way in which friction arising from adhesion intluences the stresses around asperity contacts and in turn the direction of shearing or crack growth, is comparatively straightforward to model [91]. Much more difllcult, however, is how to include the role of adhesive forces between particles in the formation and eventual rheology of third-body films. Relevant information on this aspect is now emerging from recent attempts to provide cross-fertilization between tribology and particle technology [92].

Acknowledgments This paper was prepared in its present form during a period spent at the Institut National des Sciences Appliquees de Lyon. Thanks go to Professor R. Hamelin, Directeur de I’INSA for provision of facilities, to Professor M. Godet for helpful comments and to the Region Rhone-Alpes and the Fondation Scientiflque de Lyon et Sud-Est for financial support.

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