laser surface engineered tic coating on 6061 al alloy...

Ž .Applied Surface Science 153 2000 65–78www.elsevier.nlrlocaterapsusc

Laser surface engineered TiC coating on 6061 Al alloy:microstructure and wear

Lalitha R. Katipelli, Arvind Agarwal, Narendra B. Dahotre )

Department of Materials Science and Engineering, Center for Laser Applications, UniÕersity of Tennessee Space Institute,Tullahoma TN 37388, USA

Received 23 July 1999; accepted 2 September 1999

Abstract

Ž .Hard and refractory TiC has been deposited on 6061 Al alloy by Laser Surface Engineering LSE . A ‘‘composite’’coating is obtained with TiC particles of various shapes and sizes embedded in Al alloy–Ti matrix. The coating is uniform,continuous and free of cracks. The various reactions occurring during laser processing were thermodynamically analyzed andrelated to the experimental observations. Microhardness measurements suggested high hardness values in the coating regionand a strong bonding at the coatingrsubstrate interface. Dry sliding wear tests were performed to measure the wearresistance and the coefficient of friction of the coating. Wear resistance of the coated surface was found to be high whencompared to the substrate side. The coefficient of friction was found to be 0.64. q 2000 Elsevier Science B.V. All rightsreserved.

Keywords: Laser surface engineering; TiC particles; Wear; Coating; Bonding; Interface

1. Introduction

Aluminum alloys are used in many applicationsdue to their excellent properties such as highstrength-to-weight ratio, good ductility, lightweight,

w xavailability and low cost 1 . However, the surfaceproperties of Al alloys the hardness and wear resis-tances in particular are insufficient to fulfil manyindustrial requirements. Surface coating with ceram-

w xics can be a promising approach to this problem 2 .Hard TiC ceramics are well known for combining a

) Corresponding author. Tel.: q1-931-393-7495; fax: q1-931-454-2271.

Ž .E-mail address: [email protected] N.B. Dahotre

number of special properties that have made them ofparticular interest for a wide variety of applications.They are used as a wear-resistant coating for cuttingtools and inserts and as diffusion barriers in semi-

w xconductor technology 3 . TiC exhibits a very highmelting point and thermal stability, high hardnessand excellent wear resistance, low coefficient offriction, and high electrical and thermal conductivi-ties. Because of its high melting point, TiC is apromising material to be used as first-wall material

w xin fusion reactors 3,4 .There are several processes of depositing ceram-

ics on Al alloys such as flame spraying, plasmaspraying, screen printing, electroplated coating, PVD,

w xCVD, etc. 5 . These methods are not widely used, asthey do not offer metallurgical bonding to the base

0169-4332r00r$ - see front matter q 2000 Elsevier Science B.V. All rights reserved.Ž .PII: S0169-4332 99 00368-2

( )L.R. Katipelli et al.rApplied Surface Science 153 2000 65–7866

material. In such case, Laser Surface EngineeringŽ .LSE offers to improve the surface properties of Alalloys while keeping the bulk properties more or lessintact. By means of a high power laser, very hightemperatures can be reached to melt both the metalsubstrate and, sometimes, the ceramic particles. Thismay promote chemical reaction and wetting betweenceramic and metal, and, as a result, may lead to astrongly bonded ceramic–metal interface after reso-

w xlidification 2 . Moreover, LSE involves high coolingŽ 3 8 .rates 10 –10 Krs which produce meta-stable

phases, leading to the development of a wide varietyof microstructures with novel properties that cannotbe produced by any conventional processing tech-

w xnique 6 .Many researchers have studied the effect of ce-

w xramic coating on Al alloys 7–11 . However, studieson the influence of laser treatment of Al alloys for

w xceramic coating are very limited 2,12,13 . In thepresent study, hard TiC is deposited on Al 6061alloy by LSE technique.

2. Experimental procedure

2.1. Materials

ŽCommercially available TiC powder 99.5% pu-.rity and average powder size -15 mm and 6061 Al

alloy were used in the present study. The TiC pow-der was supplied by CERAC Milwaukee, WI. Thechemical composition of 6061 Al alloy is given inTable 1.

2.2. Coating process

Plates of 6061 Al alloy with dimensions 12=12=1 in3 were cleaned using sand blasting. The pow-der precursor made of 90 wt.% TiCq10 wt.% Sisuspended in a 10-wt.% hydroxyl methyl cellulose

Žwater-based organic binder proprietary formulation

Table 1Chemical composition of 6061 Al alloy in wt.%

Mg Si Cu Cr Mn Fe Al

0.8–1.2 0.4–0.8 0.05–0.4 0.04–0.35 -0.15 -0.7 Bal.

made from commercially available resins used in the.paint industry was spray deposited on the 6061 Al

substrate. The addition of 10 wt.% Si to the TiCpowder was intended to increase the wettability and

w xfluidity of Al 6 . The average sprayed precursorthickness was 150 mm. Sprayed coupons were driedat 708C for 1 h prior to laser processing.

A 2-kW Rofin Sinar continuous wave Nd:YAGlaser equipped with a fiber optic beam deliverysystem was employed for laser treatment of thesprayed substrates. The optical fiber was 17 m longand 600 mm in diameter. The laser beam was fo-cused at 0.5 mm above the surface of the substrate.The lenses within the output-coupling module offiber optic were configured to provide a beam of 3.5mm wide line in spatial distribution onto the samplesurface. Such configuration provides a rapid process-ing speed and reduces the overlap between the laserpasses. There was a 20% overlap between two con-secutive laser tracks. The laser beam power andtraverse speed were maintained constant at 1.8 kWand 120 cmrmin, respectively. The coated sampleswere additionally cooled by mounting them on awater-cooled copper plate during laser processing.

2.3. Characterization

Phase identification was carried out on a Philips˚ŽNorelco X-ray diffractometer with CuK 1.54 Aa

.wavelength radiation, operating at 40 kV and 20Ž .mA. X-ray diffraction XRD analysis was per-

formed on both the worn and unworn surfaces of thecoating to determine the possible changes in phasesas a result of wear tests. Microstructural characteriza-tion was performed on an ISI super III-A scanning

Ž .electron microscope SEM . The samples for metal-lography were prepared by sectioning the coatedplate perpendicular to the laser tracks. The sectionedfaces were polished on a Buehler Isomet 2000 clothto a diamond finish. Keller’s reagent was used as the

Ž .etchant. Energy dispersive spectroscopy EDS andX-ray mapping equipped with SEM were utilized todetermine the elemental distribution in a semi-quantitative manner. Macroscopic observations ofthe cross-section of the coating were also conductedunder BX60M Olympus microscope. This high-reso-lution optical microscope was equipped with differ-

( )L.R. Katipelli et al.rApplied Surface Science 153 2000 65–78 67

ent modes of light to observe the cross-section indesired light conditions. Microhardness tests wereperformed on a Buehler Isomet microhardness testerwith a Knoop indentor under a normal load of 200 gapplied for 15 s. Dry sliding wear tests were carriedout on a Block-on-Disc tribometer to determine boththe weight loss and coefficient of friction with re-spect to time. The speed of the rotating disc wasmaintained at 4.4 mrs and a normal load of 4 lb wasapplied to the sample while sliding. These tests werecarried out on coated coupons of size 20=25 mm2

for a total 20-min duration. Weight loss measure-ments were made after successive 2 min. Coefficientof friction was also computed simultaneously by aninterfaced computer that acquired data in the form ofvoltage and current of the motor as a function oftime.

3. Results and discussion

3.1. Thermodynamic predictions

It is essential to study the thermodynamics of thesystem to understand the nature of chemical reac-tions taking place during the coating process. There-fore, in the present study, attempts are made firstlyto thermodynamically predict the possible reactionsand, secondly, to verify the existence of these reac-tion products using analytical techniques such asEDS and XRD. Such combined thermodynamic andanalytical approach provides an insight into the me-chanical, chemical and thermophysical properties ofthe coated sample. However, it is pointed out thatoften the thermodynamic conditions that prevail dur-ing laser processing deviate from equilibrium condi-tions, thereby making equilibrium phase diagramconsideration less applicable and more difficult topredict the process products. Due to the fact thatlaser processing is a very complex non-equilibriumprocess and also because of the lack of data in theopen literature on various reaction products duringsuch non-equilibrium processes, a first approxima-tion that generation of products based on equilibriumphase diagrams can be considered. These equilibriumproducts can be considered as sources for further

generation of stoichiometric andror non-stoichio-metric derivations. Primarily, in the system underinvestigation, the following reactions are possible.The free energy of formation and their temperature

w xranges are obtained from various sources 14–18 ,

TiqC™TiC 1Ž .

Ž .DGsy182.9q0.01T kJrmole 298-T-115 K ,ŽDGsy186.4q0.013T kJrmole 1155-T-2000

. w xK 14 . From the Ti–C phase diagram, it has beenŽ .corroborated that TiC xF1 exists as a singlex

homogeneous carbide phase with a wide range ofw xstoichiometry 15 . The reaction is reversible and,

hence, it is possible for Ti and C dissociated fromTiC to combine with other reacting elements. Ac-cordingly, the following reactions may be consid-ered:

SiqC™SiC 2Ž .

ŽDGsy71.258q0.0078T kJrmole 0-T-1700. ŽK , DGsy120.3q0.0368T kJrmole 1700-T-

. w x3200 K 16 ;

4r3 AlqC™1r3 Al C 3Ž .4 3

ŽDGsy71.315 q 0.013T kJrmole 0-T-900. ŽK , DGsy91.055q0.033T kJrmole 900-T-

. w x3200 K 16 ;

Tiq3Al™TiAl 4Ž .3

ŽDGsy52.503q0.021T kJrmole 0-T-3200. w xK 17,18 ;

Tiq2Si™TiSi 5Ž .2

ŽDGsy134.19q0.0067T kJrmole 298.15-T-. Ž1700 K , DGsy1.5738q1.0063T kJrmole 1700

. w x-T-1813 K 19 . With the obtained values of freeenergy of formation, a graph of DG vs. T is plottedin Fig. 2. From the figure, it is clearly observed thatTiC is the most stable compound followed by TiSi .2

This fact is also evident from the XRD peaks of TiCand TiSi shown in Fig. 1.2

As mentioned earlier, titanium carbide exists as asingle homogeneous phase, TiC with a wide homo-x

geneity range. The crystal nature of such phase iscubic of NaCl type and for TiC of stoichiometric


Fig. 1. XRD spectrum of the TiC-coated 6061 Al sample.

composition, the lattice parameter varies between˚4.315 and 4.324 A. Similarly, silicon carbide exists

in the form of cubic b-SiC and hexagonal a-SiC.b-SiC with cubic crystal structure of ZnS type has a

˚ w xlattice parameter of 4.358 A 20 . The values oflattice parameters for both TiC and SiC being close,the X-ray reflection peaks of these phases overlap asobserved in Fig. 1.

Even though thermodynamically Al C is ex-4 3Ž .pected to exist up to 29278C Fig. 2 in vacuum, it is

stable up to 12008C and, at 22008C, it sublimesw xwithout melting 21 . In addition, the solubility of

carbon in aluminum is only 0.02–0.04 wt.% at1300–15008C and, at 1000–11008C, it is practically

w xzero 21 . Also, the possibility of formation of Al C4 3

through carbothermic reduction of SiC exists. Thus,w xAl C may form in the reaction 224 3

4Alq3SiC™Al C q3Si 6Ž .4 3

However, the content of SiC and the duration ofcontact between SiC and the molten Al controls the

w xformation of Al C 23 and, for complete stability4 3

of Al C , Si levels above 8 wt.% are required. In the4 3

present experiments, though Si in the precursor was10 wt.%, it was not sure how much was available asfree Si in the melt during processing. In view of allthese scenarios, it is less likely that Al C will form.4 3

ŽThis was further confirmed by XRD analysis Fig..1 .

Ž .In addition to reaction 4 , TiAl can form via the3w xfollowing reaction 7 :

TiC s q3Al s ™TiAl s qC 7Ž . Ž . Ž . Ž .3

In the temperature range of 1150–1800 K, the GibbsŽ .energy of formation DG for this reaction is posi-

tive, which indicates that TiAl is unstable for an3

ideally dilute carbon concentration of 1 wt.%. How-ever, TiAl is stable in the presence of TiC when3

DG becomes negative at concentration of dissolvedy5 w xC lower than the 10 wt.% 17 . Therefore, TiAl3

formation can only be expected during processing ifTi concentration is very high. Such extremely lowconcentration of C may be possible during the pre-

Ž .sent process due to formation of SiC via reaction 2 .


Fig. 2. Gibbs free energy of formation of different compounds thatmay possibly form during laser processing.

3.2. Microstructural characterization

Fig. 3 shows an overview of the cross-section ofthe TiC-coated 6061 Al sample. The coating is denseand adherent to the substrate. The thickness of thecoating is uniform and was found to be approxi-mately 125 mm. The reduction of about 20% from150 mm thickness of the precursor deposit to thefinal coating thickness of 125 mm could be attributedprimarily to evaporation of the entire binder materialalong with possible evaporation of a small amount ofthe precursor powder mixture material. The evapora-tion of these materials is also expected to reduce theporosity between the powder particles, thereby pro-viding a dense and less thick coating. The coating iscomposite in nature with TiC particles embedded inthe Al alloy matrix. It has been experimentally proventhat refractory carbides are practically not wetted byGroup III elements and are, in general, wetted by

w xtransition elements 24 . Such difference in the wet-ting nature is explained by the electronic structure of

the metals. The better wetting of refractory carbidesby transition elements is attributed to the possiblecapture of non-localized valence electrons of moltenmetal atoms by the disturbed configuration of carbon

w xatoms in the carbides 25 . Table 2 lists the wettingangles of certain refractory carbides by molten met-als. The wetting between TiC and Al seems to haveimproved due to the addition of Si in the precursor

Ž .and Mg present in 6061 Al substrate Table 1 ,which is evident from the strong and adherent coat-ing of the sample. Table 2 shows the low contactangle values suggesting high wetting of TiC bymolten Si and Mg. It has been earlier observed thatthe addition of Mg also leads to the formation ofspinel at the metalrceramic interface, which en-

w xhances the wetting 26 . In the present study, forma-tion of spinel formation was not observed by XRD.However, a detailed investigation by transmissionelectron microscope may reveal localized spinel for-mation that promotes wetting.

Ž .A thick laser melt zone LMZ comprising of thefine dendrites of Al alloy is formed in the substrateunderneath the coating. In LSE, the rapid solidifica-tion process can result in a considerable supercoolingresulting in the refinement of microstructure from

w xcoarse to fine dendritic structure in the melt zone 6 .

Fig. 3. SEM micrograph of the overview of the cross-section ofthe TiC-coated 6061 Al sample.


Table 2Wetting of some refractory carbides by molten metals

Ž .Carbide Wetting melt Temperature 8C Atmosphere Contact angle in degrees Reference

w xTiC Fe 1550 Hydrogen 39 24w xCo 1500 Hydrogen 36 24w xNi 1450 Hydrogen 17 24w xSi 1500 Vacuum 32 14w xMg 1300 Vacuum 50 14w xAl 700 Vacuum 118 14w xWC Fe 1500 Vacuum 0 24w xCo 1500 Hydrogen 0 24w xNi 1500 Vacuum 0 24w xSi 1500 Hydrogen 0 14

Mg Not available Not available Not available Not availablew xAl 1000 Vacuum 135 14

The LMZ has an approximate depth of 300 mm. Fig.4 shows the high magnification micrograph of theinterface of LMZ with the 6061 Al substrate. Theinterfacial region shows the growth of dendrites fromthe LMZ towards the coating.

Ž .Fig. 5 a shows the SEM micrograph of thecross-section of the coating. TiC particles of varioussize and shape within the matrix of 6061 Al areobserved. The volume fraction of TiC particles in thecoating is found to be approximately 65%. Theelemental X-ray maps of the coating corresponding

Ž .to Ti, Si and Al distribution are shown in Fig. 5 b ,

Fig. 4. SEM micrograph of LMZ-substrate interface.

Ž . Ž .c , and d , respectively. Ti-rich zone shows thepresence of TiC particles whereas the coating matrixis distinguished by the Al rich zone. Si is distributeduniformly over the entire surface. Table 3 providesthe results of the EDS quantitative analysis con-ducted at various locations in the coating. Theseresults suggest that the matrix within the coating

Ž . Ž .contains Al 53.24 at.% and Ti 35.86 at.% alongŽ .with some Si 10.88 at.% . Such high content of Ti

in the matrix could possibly be attributed to fragmen-tation andror dissolution of TiC particles. Fragmen-tation of such TiC particles is observed in both Figs.

Ž .3 and 5 a . Partial dissolution of TiC and, thereby,availability of free Ti in the matrix is desirable as Ti,being a highly reactive element, tends to modify thesurface properties of the carbide particles for en-hanced wettability with the substrate molten mate-rial. Similar is the case with the interface betweenTiC particles and the coating matrix, which contains

Ž . Ž .Ti 35.19 at.% along with Al 59.053 at.% and aŽ .trace amount of Si 5.763 at.% . Such nature of the

reaction product is expected to provide chemicalbonding in addition to mechanical bonding betweenthe TiC particles and the matrix.

Ž .Fig. 6 a shows the SEM micrograph of the coat-ing-substrate interface. The corresponding X-ray ele-mental distribution of Al, Ti and Si is shown in Fig.Ž . Ž . Ž .6 b , c and d , respectively. Due to a large differ-

ence in the coefficient of the thermal expansions ofmetals and ceramics, high residual stresses may oc-cur at the interface during the rapid solidificationafter laser processing. These residual stresses canresult in the delaminationrbuckling of the coating


Ž . Ž . Ž .Fig. 5. a High magnification SEM micrograph of the cross-section of the coating and corresponding X-ray elemental map of b Al, c Si,Ž .and d Ti.

w xfrom the substrate 2 . However, in the present case,no such delamination was observed. This indicatesthat the coating is sound and adherent to the sub-strate suggesting a metallurgical bond. A sharp inter-face in all these figures suggests no transferrdiffu-sion of Ti into the region of the substrate materialsurrounding the interface. On the contrary, Al has

Ž .flown into the matrix of the coating Fig. 6b .

Table 3EDS analysis in the coating

Location At.%

Ti Al Si

Matrix within the coating 35.86 53.24 10.88TiC particles 96.589 0.484 2.927Interface of TiC particles 35.19 59.053 5.763

and coating matrix

Also, interfacial adhesion strength between ce-ramic particle and metallic melt is often representedby the work of adhesion, W , which is defined asa

work per unit area of interface, necessary to separatereversibly a solid–liquid interface to create a solid–

w xvapor interface 27 . Work of adhesion, W , foraw xseveral combination of 28 metalrceramic system is

tabulated in Table 4.The work of adhesion, W , reflects directly thea

importance of energetical interactions between thesolid and liquid phases. A higher value W suggestsa

w xstronger interactions 27 . W for TiCrSi andaŽTiCrMg is significantly higher than TiCrAl Table

.4 . Thus, in the present study, the presence of Si andMg in the precursor and 6061 Al is expected toincrease the wettability and, hence, interfacialstrength between TiC and Al melt, thereby resultingin an adherent coating.


Ž . Ž . Ž . Ž .Fig. 6. a SEM micrograph of the coating–substrate interface and corresponding X-ray elemental map of b Al, c Si and d Ti.

3.3. Mechanical characterization

Mechanical features of the TiC coated 6061 Alalloy were studied using a microhardness tester. A

SEM micrograph of the cross-section of thecoating–substrate interface with microhardness in-

Ž .dentations is shown in Fig. 7 a . X-ray elementaldistribution of Al, Si, and Ti corresponding to

Table 4Work of adhesion for metalrceramic systems

Ceramic Metallic Surface energy, g of Contact angle Work of adhesion,1yÕ2 2Ž . Ž .melt liquid–vapor interface mJrm in degrees W mJrma

TiC Al 914 118 485Si 860 32 1720Mg 583 50 960Fe 1900 39 3578

TiB Al 915 160 552

Si 860 34 1575Mg Not available Not available Not availableFe 1785 49 2955


Ž .Fig. 7. a SEM micrograph of the coating, coating–substrate interface, and substrate regions with indentations and corresponding X-rayŽ . Ž . Ž .elemental map of b Al, c Si, and d Ti.

Ž . Ž . Ž . Ž .Fig. 7 a are represented in Fig. 7 b , c and d ,respectively. The relative difference in size of micro-hardness indentations explains the variation in hard-ness. Considerably high microhardness values wereobserved in the coating. The two indentations takenin the coating showed different values of microhard-

Ž . Žness Knoop , with the top part of the coating 572. Ž"60 relatively harder than the bottom part 454"

.30 . This is due to the presence of SiC, whichapparently segregated to the top of the coating dur-

Žing laser processing. SiC, being less denser 3.1 to3. Ž 3.3.2 grcm in comparison to TiC 4.92 grcm ,

floats to the top of molten coating during laserprocessing and remains at the top of the coating afterresolidification. This fact is also evident from thecorresponding X-ray elemental map of Si shown in

Ž .Fig. 7 c . The hardness of the LMZ is 72"5 andthat of unreacted substrate is 84"10. The relatively

low hardness value of LMZ is attributed to thedissolutionrredistribution of precipitates formed dur-ing the precipitation hardening of 6061 Al alloy.

In the present study, attempts were made to semi-quantitatively determine the interfacial strength. Inthe process, microhardness indentations were made

Table 5SEM analysis of the indentations at the interface

Applied Length of the longer SEM predictionsŽ . Ž .load g diagonal mm

50 32 No interfacial cracks100 61 No interfacial cracks200 72 No interfacial cracks300 107 No interfacial cracks500 116 No interfacial cracks

1000 150 No interfacial cracks


Žalong the interface at various loads 50, 100, 200,.300, 500, and 1000 g to check for initiation of any

cracks or delaminations at the site of indentations.The loads at which the microhardness tests are con-ducted and the corresponding SEM observations areillustrated in Table 5. Fig. 8 shows the corresponding

SEM micrographs of the indentations taken at all thetest loads within the limitations of the indentor beingused in the present study. SEM observations showedthat there were no visible cracks or delaminations atthese indentations. However, the possibility of finestationary cracks within the coating is not ruled out

Fig. 8. SEM micrographs of the indentations taken at the coating–substrate interface at various loads.


since this can only be observed at very high magnifi-cation. The dynamic behavior of these stationarycracks is being further investigated using a four-pointbending test. Such tests will characterize the strengthof ceramic–metal interface in a quantitative manner.

3.4. Tribological characterization

3.4.1. Wear testsThe tribological performance of a ceramic coating

depends on the characteristics such as hardness,thickness, internal stress level, and load-bearing ca-pacity. When a coating cools from the high tempera-ture of the melt zone in the deposition process, theinternal compressive stress developed within thecoating reduces the extent of surface fracture and,thus, increases the sliding wear resistance of the

w xcoated surface 29 . In addition, wear behavior ofany material system is expected to be influenced by

Ž .environment, such as dry air or inert atmosphereŽ . Žand wet lubricated . The thermo–physical force

.and temperature conditions at the point of contactare extreme during wear testing. These extreme con-ditions are favorable to cause degradationrdissocia-

Ž .tion physical and chemical of the material and,Ž .hence, occurrence of localized reaction s between

the material and surrounding environment, therebyproviding different wear responses.

Wear results in the loss of material, elevation ofsurface temperature, surface activation, and the re-moval of protective film, and, thus, may accelerate

w xoxidation or corrosion 30 . Hence, it is important to

Fig. 9. Weight loss vs. time for the wear tests.

study the wear behavior of the coated samples todetermine the nature of the coating. Fig. 9 shows aplot of weight loss vs. time for the wear tests con-

Ž .ducted on TiC coated and uncoated substrate sur-faces of the laser-treated samples. It is clearly ob-served that the weight loss for the coated side waslower than the substrate side. The wear rate is alsosignificantly improved. The total duration of thewear test was 20 min. However, while conductingthe tests on the uncoated substrate side, the substan-tial weight losses were calculated for a time durationof only 30 s due to high seizure of the substratematerial during sliding. Comparison of XRD analy-ses of the worn and unworn-coated surfaces as illus-trated in Fig. 1 indicated that there are no newphases formed during the wear process. This sug-gested that the coating was chemically and physi-cally stable under the dry sliding conditions used inthe present study.

Ž .Surface roughness profiles shown in Fig. 10 a ,Ž . Ž .b and c indicated that the average surface rough-ness, R , of the test sample increased from 0.43 toa

13.53 mm after laser processing. The average surfaceroughness of the worn surface is 6.51 mm.

Topographic features of the worn surface wereŽ .analyzed using SEM. Fig. 11 a shows a low magni-

fication micrograph of the TiC coated 6061 Al sam-ple. X-ray elemental distributions of Al, Ti, and Fe

Ž . Ž .corresponding to Fig. 11 a are shown in Fig. 11 b ,Ž . Ž .c and d , respectively. The worn surface of thecoating is identified by continuous wear scars paral-lel to the laser tracks in some isolated regions ofcontact between the sample and wear-disc. Duringsliding, there is a material-to-material contact, fol-lowed by welding or fusing of the contacting asperi-ties. When sliding continues, the asperities of theweaker material may shear off and transfer to the

w xopposite surface 31 . In the present work, the hard-Ž .ened steel sliding disc 62 Rc , whose main con-

stituent is Fe, shears off the Al present as the matrixof the coating, and Al is collected as the loose debrisat the end of the scars. On the contrary, TiC presentin the coating, being harder than Fe, shears off thematerial from the sliding disc, and, as a result, Fe isdeposited in those scars. Such composite nature ofwear and segregation of Al and Fe in particularregions of the worn surface are clearly seen from

Ž . Ž . Ž .Fig. 11 a , b and d , respectively.


Ž . Ž . Ž .Fig. 10. Optical profilometer figures showing the surface roughness of a 6061 Al substrate, b coated side and c worn coated surface.

3.4.2. Coefficient of friction measurementThe measurement of the coefficient of friction

provides direct information about the work done todeform the surface of the material. However, it is notnecessarily a direct indication of material loss or

w xseparation of loose debris from the surface 15 . Inthe present study, the tribometer interfaced with acomputer recorded the wear test parameters such as

Ž . Ž .Voltage V and Current I as a function of time.Based on the principle of energy conservation, the


Ž . Ž . Ž . Ž .Fig. 11. a SEM micrograph of the worn surface and corresponding X-ray elemental map of b Al, c Ti and d Fe.

frictional force is equal to the electrical work doneby the motor given by

W sVoltage DV =Current D I 8Ž . Ž . Ž .f

Fig. 12. Coefficient of friction vs. time for the wear tests.

The frictional theory states that

W smNÕ 9Ž .f

where N is the normal load, Õ is the linear speed ofthe sliding disc and m is the coefficient of friction.By equating the above equations, the coefficient offriction can be computed. The computed coefficientof friction for TiC coating has been plotted for theentire test time of 20 min as shown in Fig. 12. Thebest-fit line shows the coefficient of friction of theTiC-coated surface to be approximately 0.64.

4. Conclusions

1. TiC has been deposited on 6061 Al alloy usingthe LSE technique. The coating is uniform, adher-ent, and free of cracks and porosities.


2. The coating is composite in nature with the mi-crostructure consisting primarily of hard TiC par-ticles reinforced in a matrix of Al–Ti mixture.

3. The average surface roughness of the coating wasfound to be 13.5 mm.

4. Considerably high hardness values are obtained inthe coating.

5. The coating is wear resistant and the wear resis-tance is highly influenced by the composite natureof the coating.

6. The coefficient of friction was computed to beapproximately 0.64.

Acknowledgements

Authors acknowledge the partial support from theSurface Technology Division of ALCOA.

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