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Wear 266 (2009) 907–916

Contents lists available at ScienceDirect

Wear

journa l homepage: www.e lsev ier .com/ locate /wear

ear resistance of annealed plain carbon steels in pre-strained condition

. Toppoa, S.B. Singhb, K.K. Rayb,∗

Department of Manufacturing Engineering, National Institute of Foundry and Forge Technology, Ranchi 834003, IndiaDepartment of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, India

r t i c l e i n f o

rticle history:eceived 26 January 2008eceived in revised form 25 October 2008ccepted 2 December 2008vailable online 7 December 2008

a b s t r a c t

The primary aim of this investigation is to examine the influence of pre-strain on wear behaviour of plaincarbon steels. Experimental work comprises of detailed measurement of wear rate, microstructural char-acteristics, hardness, tensile properties of unstrained and pre-strained specimens, and SEM examinationsof the worn-out surfaces and subsurfaces of five hypo-eutectoid and one nearly eutectoid steel. The resultsassist to infer that: (1) wear rate of pre-strained specimens increases with increasing pre-strain up to a

eywords:teelsre-strainliding wearelaminationicro-cracking

critical value of the latter beyond which it drops with further increase in pre-strain, (2) wear behaviourof all the investigated hypo-eutectoid steels exhibits a critical pre-strain as 0.7 ± 0.05, normalized withrespect to uniform strain, (3) simultaneous increase of wear rate and hardness up to the critical value ofpre-strain is contrary to Archard’s law; this has been explained by the ease of nucleation of voids/micro-cracks and their subsequent propagation during sliding wear of pre-strained specimens, and (4) wearrate is found to be considerably dependent on the microstructural constituents of the steels beyond the

.

critical value of pre-strain

. Introduction

A large market for steels that are designed specifically to reducehe rate of wear exists. These steels find numerous applications inrushing machines, conveyors, screening machines, hoppers andilos, chutes and so forth. Structural components in these applica-ions are often directly put into service in cold deformed condition.old deformation is a pre-straining process, which changes dislo-ation substructure of the metallic materials, and in turn, altersheir mechanical behaviour. Extensive literature exists related tohear instabilities [1], deformation substructure [2], strength [3],ardness [4], fatigue [5,6] and fracture [2,7–9] to understand thetructure–property relations of cold deformed materials. But theres lack in organized attempts to understand wear behaviour of

aterials in pre-strained or cold deformed condition. In order toulfill this gap, the present investigation aims to examine the slidingear behaviour of a series of plain carbon steels in the pre-strained

ondition.It is well documented that sliding wear is governed by abrasion,

dhesion, and delamination [10]. One of the basic manifestations

f abrasion and adhesion is the process of fragmentation of mate-ials under sliding contact which leads to formation of wear debris.anin et al. [11] have reported that fragment formation duringliding wear is the result of increased dislocation density, which

∗ Corresponding author. Tel.: +91 3222 283278; fax: +91 3222 282280/255303.E-mail addresses: kkrmt@metal.iitkgp.ernet.in, kkrmt@iitkgp.ac.in (K.K. Ray).

043-1648/$ – see front matter © 2009 Published by Elsevier B.V.oi:10.1016/j.wear.2008.12.005

© 2009 Published by Elsevier B.V.

increases the level of lattice energy; and to reduce this lattice energyit would create missing misorientation boundaries and thus formfragments. On the other hand, the process of delamination is causedby nucleation of voids, their growth and subsequent coalescencein a severely strained material during sliding [12]. Initiation ofvoids during delamination is mainly due to dislocation/dislocationor dislocation/twin interactions, and void formation around hardinclusions [13]. Since these mechanisms are predominantly gov-erned by dislocation substructure, it is expected that pre-strainingof material will have considerable effect on their sliding wearbehaviour. An emphasis is thus laid to examine the influence oftensile pre-straining on wear behaviour of a few steels.

Increase in the degree of pre-straining or cold working on mate-rials generally increases their hardness and tensile strength. Onthe other hand, it is popularly considered that the volume wornper unit sliding distance is inversely proportional to the hardnessof the material in sliding wear [14]. Thus, it is naturally expectedthat pre-straining will enhance the wear resistance of materials.The work of Biswas and Kailas [1], however, indicates that in spiteof higher hardness, wear of titanium is greater than that of cop-per. This observation has been elucidated by Biswas and Kailas interms of strain rate response of these metals, as reflected in theirstress–strain characteristics. Further, Wang et al. [15] have shown

that microstructure of a material plays significant role than its initialhardness on its sliding wear characteristics. They have supportedtheir postulations with the evidence that wear resistance of lamel-lar pearlitic structure of 52100 and 1080 steels is higher than thatof spheroidized structure in dry sliding, even though hardness of

908 V. Toppo et al. / Wear 266 (2009) 907–916

Table 1Chemical compositions of the investigated steels (in wt%).

Material C Mn Si S P Cr Ni Cu Al Fe N (ppm)

S15 0.145 0.37 0.05 0.009 0.015 0.03 0.025 0.05 0.025 Bal 70S22 0.22 0.87 0.074 0.008 0.014 0.26 – – 0.03 Bal 70SSSS

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38 0.38 0.81 0.22 0.005 0.0147 0.47 0.77 0.22 0.016 0.01464 0.64 0.75 0.19 0.1 0.01782 0.82 0.75 0.18 – 0.015

he former is lower than that of the later. Wang et al. [15] havettributed the difference in wear resistance of various microstruc-ures to emerge from different possibilities like thermal stabilityf the microstructure, resistance to plastic deformation, and resis-ance to nucleation and propagation of micro-cracks. Thus, it isifficult to predict a-priori the nature of the influence of pre-strainn sliding wear characteristics of metallic materials.

Examinations of wear behaviour of variously cold deformed low,edium, and high carbon steels are the major content of this report.

old deformation on the materials has been applied by tensile pre-training, and the maximum magnitude of pre-strain has been keptimited within the uniform elongation of the selected steels. A seriesf pin-on-disk wear tests have been carried out in the severe wearegime of the steels; the severe wear regime has been initially eval-ated for each of the steels by determining their wear rates atifferent loads for a fixed sliding speed. The wear experiments areupplemented by studies related to tensile behaviour of the pre-trained materials, hardness measurements on specimens beforend after the wear tests, SEM observations of the worn-out sur-aces and examinations of subsurface cracks. The inherent aim forhe overall investigation is to reveal the influence of pre-strain onear resistance of materials.

. Experimental procedure

.1. Material and microstructure

Six commercial plain carbon steels with carbon contents of.15%, 0.22%, 0.38%, 0.47%, 0.64% and 0.82%, obtained in the form ofods with diameters ranging between 10 and 13 mm, are selected forhis investigation. Chemical compositions of the steels are shown inable 1. Microstructural characterizations of the steels were donen their annealed condition. The annealing treatments have beenone by soaking the as-received rods at 1173 K (0.15%C), 1165 K0.22%C), 1145 K (0.38%C), 1129 K (0.47%C), 1096 K (0.64%C), and046 K (0.82%C) for half an hour followed by furnace cooling; thennealing temperatures vary with the C-content of the steels, whichre shown in the parenthesis. The temperature of annealing forhe steels has been selected from the Fe–C diagram, consideringustenitization temperature to be approximately 50 K above Ac3

ine.

Sample blanks of approximately 10 mm height were cut fromhe annealed rods for microstructural examinations on the trans-erse sections. The specimens were ground, polished up to 0.25 �miamond finish and were then etched using 2% Nital solution for

able 2ome salient microstructural features of the selected steels.

teels Microstructure Inclusion content (%)

.15C Ferrite–pearlite 0.24 ± 0.03

.22C Ferrite–pearlite 0.21 ± 0.01

.38C Ferrite–pearlite 0.05 ± 0.02

.47C Ferrite–pearlite 0.09 ± 0.02

.64C Ferrite–pearlite 0.1 ± 0.02

.82C Full pearlitic 0.05 ± 0.01

* Not measured; because the steel exhibits mostly pearlitic structure.

0.125 0.01 0.01 0.017 Bal 640.13 0.005 – 0.003 Bal 630.02 0.02 0.04 0.002 Bal 650.03 0.002 0.03 0.002 Bal 68

metallographic examinations. Prior to metallographic examina-tions, inclusion content of the steels was determined on polishedspecimens using the Japanese standard method J0555 [16]. Vol-ume fraction of pearlite in the microstructures of the steels wasmeasured by point counting method [17] with the help of an imageanalyzer.

2.2. Specimens fabrication

A number of specimen blanks, approximately 150 mm in lengthand 8 mm in diameter, were first machined from the annealed steelrods in order to fabricate pre-strained specimens for the wear, hard-ness and tensile tests. These rods were pulled in a Shimadzu (model:AG-5000G, capacity 50 kN) testing machine at a crosshead speedof 1 mm/min at room temperature (≈300 K) to induce some pre-determined strain. The selected pre-strain levels were within theuniform elongation of the selected steels; typically for low carbonsteels these were in the range of 4–20%. Wear pins of 4 mm diameterand 30 mm length while tensile specimens of 5 mm gauge diameterwith 25 mm gauge length, were fabricated from these pre-strainedspecimen blanks. A set of wear pins and tensile specimens werealso fabricated from the unstrained rods.

2.3. Tensile test

Tensile tests were carried out on the fabricated specimens at acrosshead speed of 1 mm/min using the Shimadzu testing machinefollowing ASTM standard E8 M [18] at room temperature of ≈300 K.This crosshead speed corresponds to a nominal strain rate of6.67 × 10−4 s−1.

2.4. Wear test

Wear experiments were carried out on a pin-on-disc machine(model: Ducom TR 20) following ASTM standard G 99-03 [19].The disks were made of hardened En-31 steel and were hav-ing initial hardness and surface roughness as Rc = 61 ± 1 andRa = 0.36 ± 0.02 �m, respectively. The surfaces of the test pins wereground and were polished up to 0.25 �m diamond finish prior toeach test. Both the disk and the pin were cleaned with acetone

before each test, and each specimen was tested on a fresh trackof the disk. All wear tests were carried out using sliding speed of0.5 m/s in air at the ambient temperature of ≈300 K and at the rela-tive humidity of ≈45%. The normal load was varied between 5 and70 N on unstrained specimens to determine the transition load from

Pearlite (%) Ferrite grain size (�m) Hardness (HV10)

14.9 ± 1 14.3 ± 1.8 114.0 ± 3.327.1 ± 0.9 18.3 ± 1.9 156.2 ± 3.942.6 ± 1.4 13.8 ± 2.4 157.6 ± 3.463.0 ± 1.6 –* 202.5 ± 4.778.9 ± 1.6 –* 211.0 ± 6.0100 Not applicable 261.4 ± 7.0

V. Toppo et al. / Wear 266 (2009) 907–916 909

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steels are summarized in Table 2. Microscopic examinations indi-cate that the steels with carbon contents of 0.15%, 0.22%, 0.38%,0.47%, and 0.64% contain well-distributed mixture of ferrite andpearlite, whereas the steel with 0.82%C has almost fully pearliticstructure as shown in Fig. 1.

ig. 1. Typical microstructures of (a) 0.15%C, (b) 0.22%C, (c) 0.47%C, and (d) 0.82%C ste

ild to severe wear. The loads corresponding to severe wear regimeere considered for wear tests on the pre-strained specimens of

he selected steels. The height loss of the pins was measured usingLVDT of 1 �m accuracy and the data of height loss versus slidingistance were automatically recorded.

.5. Hardness measurement

Macro- and micro-Vickers hardness measurements have beenarried out on the investigated steels. The former tests wereone at a load of 10 kgf using a Vickers hardness-testing machineFuel Instruments & Engineers Pvt. Ltd., India, model: VM 50) onnstrained steel samples, whereas the latter tests were carriedut with indentation load of 200 gmf for loading duration of 15 sith the help of a microhardness tester (LECO, model: DM-400)

n both unstrained and pre-strained steel specimens before andfter the wear tests. In addition, Vickers microhardness measure-ents were also done on the phases of the investigated steels at a

oad of 10 gmf for the same loading duration of 15 s using a LEICAmodel: Q550MW) microhardness tester. Average hardness valuesere estimated considering 10 readings for macrohardness and0 readings for microhardness. Appropriate care was taken duringach measurement so that the distance between two successivendentations was more than thrice the average diagonal length ofhe indentations.

.6. SEM examination

Specimens were cut from the end of the worn-out surfacesor SEM examinations. Transverse sections of the worn-out sur-

aces were also ground and polished for the same examinations.he normal and transverse sections of the worn-out surfaces werexamined using a scanning electron microscope (model: JEOL, JSM-800) after appropriate ultrasonic cleaning in acetone. A series ofepresentative photographs of the surface, subsurface and wearebris of the specimens were recorded.

e bright constituent is pearlite, whereas the dark or the grayish constituent is ferrite.

3. Results and discussions

3.1. Materials and microstructure

Details of the microstructural characteristics of the investigated

Fig. 2. Influence of carbon content on the microhardness of ferrite and pearlitecolonies of the investigated steels.

910 V. Toppo et al. / Wear 26

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work hardening of the specimens during the pre-straining process.The effect of pre-strain on the stress–strain behaviour is illus-

trated using a set of representative plots for the 0.38%C steel in Fig. 4.The pre-strained specimens do not exhibit yield point phenomenon

Fig. 3. Effect of pre-strain on hardness.

Estimated hardness values of the steels are compiled in Table 2;ardness increases with increase in carbon content of the steels ingreement with natural expectation. Microhardness values of fer-ite and pearlite in the investigated steels are shown in Fig. 2. Theesults in Fig. 2 indicate that hardness of both ferrite and pearlitencreases with increase in carbon content of the steels. The increasen macrohardness of the steels with increased carbon content ishus due to the combined influence of the increase in (i) the amount

f the harder constituent (pearlite) and (ii) the hardness values ofhe constituent phases (Fig. 2). The increase in microhardness ofhe constituents, on the other hand, is attributed to the increasen the mutual hydrostatic stress on the constituents with increase

ig. 4. Typical engineering stress–strain curves for the 0.38%C steel specimens hav-ng varying magnitudes of initial pre-strains.

6 (2009) 907–916

in pearlite content as explained by Ray and Mondal [20]. In addi-tion to the effect of hydrostatic stress, microhardness of pearliteis also dependent on its interlamellar spacing [21,22], which hasnot been accounted here. The scatter associated with microhard-ness is higher than that of macrohardness because microhardnessis known [20] to vary within a grain from its centre to surface.

3.2. Effect of pre-strain on hardness and tensile properties

The variations of hardness as a function of pre-strain for all thesteels are shown in Fig. 3. The results in Fig. 3 indicate that hardnessincreases almost linearly with increased pre-strain of a specimen.Increase in hardness with increased pre-strain is attributed to the

Fig. 5. Effect of pre-strain on yield and tensile strength on (a) 0.38 and (b) 0.82%Csteels.

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s observed in the stress–strain curves for specimens without pre-train. The yield strength increases monotonically with increasingre-strain for each material. Effect of pre-strain on the ultimateensile strength of the selected steels is similar to that on yieldtrength, although the increase in ultimate tensile strength withncreased pre-strain is rather small (Fig. 5). The observed naturef variation of strength with pre-strain is in good agreement witheveral earlier reports like that on 316 stainless steel [23], AISI 4340teel [2] and plain C–Mn steel [24]. Bassim and Bayoumi [2] haveeported that the rate of change in strength of 4340 steel with pre-training has a direct correlation with its dislocation substructureeveloped during pre-straining process. These investigators haveeported that with pre-straining, mobile dislocation density slowlyncreases, and gradually results in the formation of cellular struc-ure. Once the cell structures are well developed, their size getseduced with increasing pre-stain. All these phenomena have arofound influence on the deformation path of the materials [25].he relation between dislocation substructure and pre-strain cane considered to explain the effect of pre-strain on strength of the

nvestigated steels.The nature of variation of uniform and total elongation with

mposed pre-strain is almost similar for all the steels. The uniformnd total elongation continuously decreases with increase in pre-train as illustrated in Fig. 6 by a typical plot. The results depicted inig. 6 are indicative of the fact that increased pre-strain decreasesuctility due to work hardening. Typical plots of strain hardeningxponent (n) versus imposed pre-strain levels (as shown in Fig. 7)how that n monotonically decreases with increase in pre-strain,nd hence these results support the contention that strain hard-ning is the primary reason behind the decrease in ductility of theteels due to pre-straining.

.3. Transition load selection

Wear behaviour of the steels has been examined using two setsf experiments. First, the load corresponding to the instigation of

he severe wear regime, commonly referred as transition load, wasstimated using un-deformed annealed specimens of the selectedix varieties of steels. Next, wear experiments were performed onlyn the severe wear regime (as estimated from the first set of experi-

ents) for specimens having various levels of pre-strain and having

Fig. 6. Effect of pre-strain on uniform and total elongation

6 (2009) 907–916 911

different compositions. The wear tests were done following theprocedure described in Section 2.4.

To estimate the transition load between mild and severe wear,several wear tests were done on each of the steels applying differ-ent loads ranging between 4.9 and 68.7 N. The wear rate (W) of eachspecimen was estimated from the last 75% of the obtained data ofheight loss versus sliding distance following the suggestion ren-dered by Ruff and Ives [26]. This was done in order to eliminate theeffect of run-in-wear data and to ensure that the data selected foranalyses correspond to perfect contact between pin and the disc.The specific wear rate (WS) was calculated using the equation:

WS = V

(P × SD), (1)

where V is the wear volume of the pin specimen in mm3, P is thenormal load in N, and SD is the sliding distance in mm.

The variations of W and WS versus normal load for the un-deformed specimens are plotted in Fig. 8(a) and (b), in log–log scale.Comparison of the wear regimes for the investigated steels contain-ing C > 0.15% was made with the typical nature of variation of Wversus normal load (Fig. 9), as suggested by Welsh [27]. The loadvalues at which T1 and T2 transitions occur are often described as‘mild oxidative’ and ‘severe metallic’. But the search for correlationbetween the present results and that of Welsh is apparent, becauseWelsh used different test configuration and kept the initial hardnessof the two rubbing surfaces identical. However, the results shownin Fig. 8(a) can be compared with the schematic plot portrayed inFig. 9 to estimate the magnitude of T1.

A number of investigations have provided explanations for thetransition between wear mechanisms i.e. from mild to severe wear,using wear mechanism maps [28] or using specific wear rate versusapplied load [29] data, other than the variation of wear rate withnormal load as specified by Welsh [27]. In the present investigation,the tests were carried out within the limited range of normalizedpressure (≈10−4–10−3) and normalized velocity (102), which corre-sponds to the mild oxidation regime in wear mechanism maps [28].

The obtained results are not in good agreement with the expec-tations derived from the above-mentioned background, since theresults used in wear mechanism maps were generated by severalinvestigators, and the generated results can be used only as a guide-line [28].

s of the investigated (a) 0.47 and (b) 0.64%C steels.

912 V. Toppo et al. / Wear 266 (2009) 907–916

Fig. 7. Typical effect of the magnitude of pre-strain on the strain-hardening expo-nent (n) of the investigated (a) 0.22 and (b) 0.64%C steels.

Fig. 8. Typical plot of (a) wear rate vs. normal load and (b) spec

Fig. 9. Schematic diagram of wear rate vs. normal load showing transition loads T1and T2 (following ref. [27]).

It can be concluded from the comparative assessment of theresults described in Fig. 8(a) with that in Fig. 9, that the transitionload between mild and severe wear is ≤4.9 N for 0.15%C steel, and≤29.4 N for the other selected steels. The results in Fig. 8(b) alsosupport the contention. These results are also in good agreementwith that reported by Wang [29]. On the basis of these results, weartests of the pre-strained specimens for 0.15%C steel, 0.22%C steeland steels with C > 0.22% were carried out at normal loads of 29.4,39 and 49 N respectively, so that the nature of the wear remainspredominantly in the severe wear regime.

3.4. Influence of microstructure and pre-strain on wear behaviour

Variations of wear rate (W or WS) with imposed pre-strain, forall the steels, are shown in Fig. 10. The results in Fig. 10(a) and (b)

depict that the magnitudes of W and WS increase with increas-ing pre-strain up to a critical level followed by a drop at furtherhigher pre-strains. Increase in wear rate with increasing pre-straincan be explained as follows: increase in pre-strain results in anincrease of dislocation density in metallic materials, which assists

ific wear rate vs. normal load for the investigated steels.

V. Toppo et al. / Wear 266 (2009) 907–916 913

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Fig. 10. Typical plot of (a) wear rate vs. pre

oid formation in the matrix or at the particle–matrix interface.he void formation inside the matrix occurs either at intersectionsf dislocations or at dislocation-twin boundary. Initiation of voidst the particle–matrix interface, on the other hand, occurs whenhe applied stress at a particle exceeds the cohesive strength ofhe particle–matrix interface aided by the stress field of highermount of dislocations due to pre-straining. During wear test, underhe influence of applied load on the specimens, the stress field ofhe dislocation density originated due to pre-straining, not onlyacilitates the void nucleation, but also assists their growth, andoalescence at the particle/matrix interface. The above phenomenaause an increase in the wear rate of a material with increasingre-strain, and are considered to be the primary reason for thenhancement of the wear rate of the pre-strained specimens of theelected steels. Typical photographs of the subsurface layers of theested pins provide the experimental evidence for the occurrencef the above phenomenon through the presence of elongated voidsnd subsurface cracks in Fig. 11.

An examination of Fig. 10 reveals consistent drop in wear raterom the highest pre-strains for each of the steels. The overallmount of wear in pre-strained specimens is governed by [30]:

i) rate of nucleation of voids, (ii) rate of oxidation, (iii) nature ofislocation substructure, and (iv) their interactions. In the regimef severe wear, accelerated nucleation of voids and micro-cracksggravates the process of wear by delamination. On the other hand,ear surface of the specimens having higher internal energy is

ig. 11. SEM photographs of worn surface layers of (a) 0.38%C steel subjected to 12% pre-longated voids and subsurface cracks on the wear pins.

n and (b) specific wear rate vs. pre-strain.

susceptible to higher oxidation [31], and as a consequence, dis-continuous but an adherent oxide layer forms and results in thereduction of the overall wear rate in specimens with the highestamount of pre-strain. Cracking and subsequent spalling of the oxidelayers make the wear process to continue; however, the overallwear rate is controlled by several opposing phenomena as discussedabove.

An alternative possibility which can lead to drop in wear rate atthe higher pre-strained specimens is the initial nature of disloca-tion substructure and their interaction with the forces active duringthe process of wear. With increasing pre-strain, the random array ofdislocations would tend to form stable dislocation cells [2]. Thesewould not tend to impede the nucleation of micro-voids and micro-cracks during wear but would provide resistance for their growthin the subsurface of the pre-strained specimens. In addition, for-mation of an oxide layer would also reduce the transfer of activeforces to the metal substrate and aid in retarding crack growth. Theconsequence of the above phenomena is the drop in wear rate forspecimens pre-strained beyond a critical amount.

In steels with ferrite–pearlite structure and having carbon con-tent of ≤0.4 wt%, the ferrite is almost continuous [32]. With increase

in pre-strain, the random array of dislocation tends to form stabledislocation cells [2] and would not tend to impede the nucleation ofvoids and micro-cracks during wear in presence of continuous softphase of ferrite, but may provide resistance to their growth at thehigher pre-strain levels. This could be the possible reason for small

strain and (b) 0.38%C steel subjected to 15.7% pre-strain. The photographs indicate

914 V. Toppo et al. / Wear 266 (2009) 907–916

Fig. 12. Typical plots of normalized specific wear rate vs. normalized pre-strain ofthe selected steels. The specific wear rates have been normalized with respect to thehighest value amongst the pre-strained specimens of a steel, whereas pre-strains arenTs

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mation and crack initiation behaviour during sliding wear of theselected steels, and their influence on the wear rate is manifestedin Fig. 13.

ormalized with respect to the strain corresponding to uniform elongation of a steel.he vertical lines indicate the usual range in which lie the maximum normalizedpecific wear rates of the investigated steels.

rop in wear rate for the steels with carbon content ≤0.38%. On thether hand, in steels with C content >0.38%, where the volume ofearlite, the harder constituent, is significantly more than that of

errite, the softer constituent, that impede the nucleation of voidsnd micro-cracks during sliding wear, would be the reason for theignificant drop in W or WS.

In order to understand the critical magnitudes of pre-strainsor the different steels at which drop of wear rates commences, aypical plot of normalized WS versus normalized pre-strain is con-tructed and is shown in Fig. 12; the wear rates are normalized withespect to the highest achieved value of WS for a specimen whereashe pre-strains are normalized with respect to the strain corre-ponding to uniform elongation of each of the steels. The resultsn Fig. 12 indicate that the normalized critical pre-strains up to

hich wear rate increases with increasing pre-strain for the hypo-utectoid steels are almost similar and this value can be representeds 0.7 ± 0.05; the corresponding strains fall within a narrow rangeetween 12% and 15%, closer to the uniform strain. However, theritical pre-strain of 0.82%C steel is 4%, which is significantly dif-erent from the range estimated for the hypo-eutectoid steels. Thiseviation of the critical pre-strain for the near eutectoid steel isrimarily due to the abrupt variation in the microstructural con-tituents.

The specific wear rate at critical pre-strain for all the investi-ated steels is examined as a function of the amount of pearliten Fig. 13. The plot in Fig. 13 reveals two distinct regimes of wearate versus percentage pearlite. The magnitude of WS decreasesith increasing pearlite content till its amount reaches ≈80%

corresponding to 0.64%C steel) and then it again increases withncreasing pearlite content. The nature of variation between WS

nd pearlite content can be explained with distribution of ferritend pearlite in the varied microstructures of the investigated steels,hich control the deformation behaviour/strain rate response of

he constituent phases during sliding wear test. In the low-carbon

Fig. 13. Influence of volume fraction of pearlite on the specific wear rate estimatedat critical pre-strain for the different steels.

steels (up to ≈0.4%C, which exhibit amount of ferrite to be higherthan that of pearlite) ferrite regions completely encapsulate pearlitecolonies whereas in steels with higher carbon (>0.4%C, whichexhibit amount of pearlite to be higher than that of ferrite), fer-rite regions remain embedded inside pearlite colonies. These varieddistributions of the microstructural constituents are schematicallyshown in Fig. 14; these are considered to cause different defor-

Fig. 14. Schematic diagram of the microstructural constituents for steels having car-bon content (a) up to ≈0.4%C and (b) �0.4%C. The represented hatched constituentis pearlite, whereas the white constituent is ferrite.

ar 266 (2009) 907–916 915

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.5. Correlation of wear rate with some mechanical properties

All metallic materials get work hardened when subjected toeformation. During sliding wear of metallic materials the degreef work hardening can be approximated by comparing the hardnessefore and after the wear test. The degree of work hardening of the

nvestigated steels has been estimated by this approach and theesults are shown in Fig. 15. Hardness values of the worn-out sur-aces, termed as pesudo-dynamic hardness [30], usually increasesith increasing pre-strain up to a maximum value followed by drop

t higher pre-strains as shown in Fig. 15(a). The degree of workardening for the specimens during wear tests has been indexedy the difference between pesudo-dynamic and static hardness val-

es, and is shown in Fig. 15(b). The results in Fig. 15(b) show thathe degree of work hardening of steel due to wear test does notepict any monotonic variation but can be considered to indicate aecreasing trend with increasing pre-strain of the specimens.

ig. 15. Hardness of the specimen surfaces (a) after the wear test, and (b) % increasen hardness due to wear for the selected steels.

Fig. 16. Variation of wear rate with initial hardness of the investigated steels.

The results in Fig. 3 indicate that initial static hardness increasesalmost linearly with increased pre-strain of a specimen. On theother hand, the plots of wear rate versus hardness in Fig. 16 indi-cate that the wear rate of the investigated specimens of the selectedsteels increases with increasing hardness up to critical pre-strain;these results contradict the generalized law of wear rate por-trayed by Archard [33]. Similar contradicting results have also beenreported earlier by Wang et al. for 52100 and 1080 steels with differ-ent microstructures [15], and by Biswas and Kailas for copper andtitanium [1] having different strain rate response. The results ofthe present investigation thus assist to infer that the inherent sub-structure and microstructure of a material are the dominant factors

which govern the wear resistance of metallic materials rather thantheir hardness values.

As the microstructure, substructure, and strain rate response ofmaterial are interrelated with the mechanical properties of a mate-rial, it becomes obsessive to examine also the relationship between

Fig. 17. Variation of wear rate with tensile strength.

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[31] G.E. Dieter, Mechanical Metallurgy, McGraw-Hill Book Company, London, 1988,

16 V. Toppo et al. / W

ear resistance and tensile properties of the pre-strained samples.he variations of wear rate for the steels with their tensile strengthre illustrated in Fig. 17. Fig. 17 reveals that the wear rate of pre-trained specimens of a steel increases with increase in its tensiletrength in a similar manner to the variation between wear rate andardness (Fig. 16). The nature of variation of wear rate with tensiletrength of the investigated steels can also be explained in a similaranner to the variation between wear rate and hardness.

. Conclusions

The results of the present investigation and their pertinent dis-ussion lead to the following major conclusions:

1. Wear rate of the investigated steels increases with increasingpre-strain up to some critical value and then it drops with furtherincrease in pre-strain.

. The magnitudes of the critical pre-strains of the selected steelsvary with the nature and the amount of the microstructural con-stituents. But interestingly the normalized critical pre-strain liesin a narrow range of 0.7 ± 0.05 for hypo-eutectoid steels, whennormalization is done with respect to uniform strain of a mate-rial.

. For specimens subjected to pre-strains beyond the critical value,the magnitude of drop in wear rate from its maximum value isconsiderably higher in the plain carbon steels containing carbon>0.4% than that in steels having carbon <0.4%.

. Hardness and tensile strength of all the investigated plain carbonsteels increase monotonically with increasing pre-strain, due towork hardening. But contrary to the expectation, wear rate alsoincreases with increasing hardness up to the value of critical pre-strain. These results are in direct contradiction to the popularArchard’s law.

. The phenomena of nucleation of voids or micro-cracks, andtheir subsequent rapid propagation are attributed as the gov-erning factors for the deterioration of wear resistance of thepre-strained steels up to their critical pre-strain.

. The overall process of wear in the pre-strained specimens of theinvestigated steels is governed by (i) the ease of nucleation ofvoids or micro-cracks, (ii) propagation of these cracks, whichcauses delamination and (iii) oxidative wear. The wear rate ofthe pre-strained specimens is governed by the competitive pre-dominance of these wear processes; higher pre-strain inducespredominance of oxidative wear, which results in drop in wearrate.

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