Acta Materialia 50 (2002) 4913–4924www.actamat-journals.com

Ti–6Al–4V strengthened by laser melt injection of WCp

particles

J.A. Vreeling, V. Ocelı´k, J.T.M. De Hosson∗

Department of Applied Physics, Materials Science Center and the Netherlands Institute for Metals Research, University ofGroningen, Nijenborgh 4, 9747 AG Groningen, The Netherlands

Received 11 August 2002; received in revised form 19 August 2002; accepted 19 August 2002

Abstract

The laser melt injection (LMI) process has been explored to create a metal–matrix composite consisting of 80µmsized WC particles embedded in a Ti–6Al–4V alloy. In particular the influences of the processing parameters, e.g.power density, scanning speed and powder flow rate, on the dimensions and microstructure of the laser track havebeen examined. The microstructure was investigated by advanced transmission electron microscopy including energyfiltering techniques and scanning electron microscopy with an integrated electron back-scatter diffraction/orientationimaging microscopy. Typical dimensions of a single laser track are a width of 1.8 mm and a depth of 0.7 mm. Thevolume fraction of the WC particles is about 0.25–0.30. An important finding is that the particle distribution is homo-geneous and that the particles are injected over the whole depth and whole width of the melt pool. Only occasionallya crystal orientation relation between WC, W2C and TiC is observed. A substantial increase in wear resistance wasobserved, i.e. 0.5× 10�6mm3/Nm for the WCp laser embedded and 269× 10�6mm3/Nm for the untreated Ti–6Al–4Valloy at the same contact stress (20 MPa). 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved.

Keywords: Laser treatment; Scanning electron microscopy; Transmission electron microscopy; Electron back-scatter diffraction;Titanium; Carbides; Microstructure; Wear

1. Introduction

Although titanium alloys exhibit excellentmaterial properties like a high strength to weightratio, good corrosion resistance, good creep,fatigue and toughness properties [1], the high fric-tion and poor wear resistance hamper many poten-

∗ Corresponding author. Tel.:+31-503634898; fax:+31-503634881.

E-mail address: hossonj@phys.rug.nl (J.T.M. De Hosson).

1359-6454/02/$22.00. 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved.PII: S1359 -6454(02 )00366-X

tial applications. Protecting Ti with a coating is anadequate solution to improve these properties,while keeping the advantageous bulk propertiesunaffected. In this work a high power laser is usedto protect Ti–6Al–4V by means of the so-calledlaser melt injection (LMI) process [2–5]. In thisprocess the laser beam is used to locally melt thetop layer of a metal (Ti–6Al–4V) substrate, whilesimultaneously a ceramic powder (80µm sizedWC particles) is injected into the melt pool. Duringrapid solidification of the substrate the particles aretrapped and a metal–matrix composite (MMC)layer is formed, while the bulk is unimpaired.

4914 J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

The properties of the produced MMC layerdepend, among other things, on the amount of theinjected powder, the microstructure of the meltpool matrix and the bonding between particle andmatrix. Since the wetting properties between cer-amic and metal materials are usually poor [6], thebonding between particle and matrix may cause asevere problem. To circumvent that problem, wedecided to choose the combination of a Ti-sub-strate with WC particles. Because of the highaffinity of Ti for C, a chemical reaction betweenWC particle and Ti matrix occurs during the laserprocess, where TiC is formed [7]. The consequenceis that reaction layers between the particles andmatrix may be formed [8,9], which may improvethe wetting and may therefore improve the bond-ing [6,10].

The purpose of the present work is to describethe influence of the LMI process parameters on theappearance of the laser track, to characterize themicrostructure of the WCp/Ti–6Al–4V MMC usingX-ray diffraction (XRD), scanning electronmicroscopy (SEM), (energy-filtered) transmissionelectron microscopy (TEM) and electron back-scatter diffraction (EBSD) [11] and finally to studywear properties of prepared layer in comparisonwith the substrate alloy.

2. Experimental

Spherically shaped, granular WC particles, witha typical diameter of 80 µm were injected in Ti–6Al–4V slices (40 mm in diameter and 5 mmthick) by using a 2 kW continuous wave RofinSinar Laser Nd:YAG laser. The laser beam istransported by means of a fiber, resulting in ahomogeneous intensity distribution. After leavingthe fiber, that has a diameter of 0.8 mm, the beamis collimated before being focused to a spot sizeof 3.6 mm by setting the focal point of the laser 9mm out of focus. The lens system is water-cooledand has a focal length of 120 mm. A nozzle islocated between the lens and the specimen to sup-ply a shielding gas (5 l/min Argon) that protectsthe lens and prevents oxidation of the sample.Moving the specimen with respect to the laserbeam results, a laser track is formed. A computer

numerically controlled (CNC) X-Y table controlsthe movement of the specimen.

A powder feeding apparatus (Metco 9MP) pro-vides a constant powder supply, which is necessaryto produce a constant coating over the entire lasertrack. A carrier gas (1.5 l/min Argon) transportsthe powder into a cyclone, where the major gasflow escapes through an upper outlet. The powderis fed through a nozzle into the melt pool.

After injection, cross-sections were made andcarefully polished. A coarse view of the track isobtained by standard optical microscopy (OlympusVanox-AHMT). The microstructure is analyzed inmore detail in SEM (Philips XL30 FEG, equippedwith an energy dispersive X-ray detector (EDX)).Furthermore, crystal orientation information isacquired by orientation imaging microscopy (OIM)attachment to a Philips XL30s field emission gun(FEG) SEM. EBSD patterns are collected duringrectangular area scans and automatically indexedusing the software of TexSEM Laboratories Inc.[12]. Specimens for TEM observations were pre-pared by grinding, dimpling and ion milling round3 mm discs to electron transparency. The surfacesof the starting discs correspond to the surface afterlaser embedding. A JEOL 4000-EX/II high-resol-ution TEM (operating at 400 kV), JEOL 200 CXTEM (operating at 200 kV) and a JEOL 2010FTEM (0.5 nm probe, FEG and operating at 200kV) are also used to characterize the microstructureto greater details. Electron diffraction is used forphase identification and crystal orientation deter-mination. The JEOL2010F microscope is equippedwith an EDX system and a post-column energy fil-ter (Gatan image filter, GIF) with a 1 k by 1 kMultiScan Camera CCD camera, to performchemical analysis. The overall phase compositionis obtained by means of XRD (Philips PW-1830)using Cu-Kα radiation.

Samples for wear tests were cut out from 5 mmthick substrates by spark erosion with dimensionsof 5 × 10 × 1mm3. Tests were performed at a Plinttribometer TE 67 using a pin-on-disk setup.Stationary specimen was pressed under a knownnormal force against the surface of rotating diskmade of 100Cr6 hardened (60 HRc) steel with aninitial surface roughness of Ra � 0.1µm. Specimenand ring were completely submerged in a oil bath

4915J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

(BP Transcal N) at room temperature. The frictionforce and the change of relative position betweenpin and ring were measured continuously during20 h of the wear test. To exclude hydrodynamic lifteffects and to perform wear tests under lubricatingboundary conditions, wear experiments were per-formed at a low constant sliding speed (0.01 m/s)and with a contact pressure of 20 and 100 MPa.Wear surfaces were inspected after the test by SEMand confocal optical microscope (µ Surf Nano-focus Messtechnik).

3. Results

Cross-sections of single laser tracks, producedwith different laser and powder parameters areexamined so as to obtain a suitable set of laser pro-cessing parameters. A successful combination ofparameters is found to be: power density P �79W/mm2, scanning speed v � 8.3mm/s and

powder feeding rate m � 125mg/ s. The width ofa single laser track is about 1.8 mm, the maximumdepth is about 0.7 mm and the volume fraction ofWC inside the track lies in the range of 0.25–0.30.It is important to note that the particle distributionis homogeneous and that the particles are injectedover the whole depth and whole width of themelt pool.

An interesting and advantageous point is thatthis particular material system allows a certainvariation of the processing parameters in order toobtain different track characteristics. In general thefollowing parameter ranges result in suitabletracks: P ranging between 39 and 84 W/mm2, v inthe range from 5.0 to 11.7 mm/s and the powderfeeding rate m varying from 83 to 167 mg/s. Inaddition, to cover larger areas with an MMC layer,more adjacent laser tracks can be produced. Anexample is shown in Fig. 1. The parameters of asingle track are P � 59W/mm2, v � 8.3mm/sand m � 83mg/s. The five tracks, with a distanceof 1.6 mm between the centers of two adjacenttracks apart, have an overlap of about 10%. Thisoverlap region is treated twice, but shows, never-theless, no distinct difference with the regions thatare treated only once. As a result this makes itpossible to produce larger areas than the widths of

the laser beam of a single track, which is animportant aspect for applications.

Before analyzing the microstructure in moredetail it is worthwhile to know which phases arepresent in the coatings. Therefore, XRD measure-ments are performed on single laser tracks. Fig. 2shows a typical 2q scan of a track. Besides theexpected α-Ti and WC, W2C, TiC, W and β-Tipeaks are identified. An overview of the crystalstructures of these phases is given in Table 1.

The melt pool matrix, i.e. the regions betweenthe injected particles in the resolidified melt pool,can roughly be grouped in two regions: the bottomand center/top part. In the bottom part Ti cells,containing also Al and V are present, and furtherTiC dendrites are observed. In the center and toppart of the laser track many TiC dendrites as wellas W cells and dendrites are present (Figs. 3 and4a). A back-scatter electron (BSE) detector is usedin Fig. 3 to obtain phase contrast due to the differ-ence in atomic weight. In addition, Ti grains arepresent, having a martensitic structure (Fig. 4b).

The reaction zone consists of different reactionlayers. Fig. 5 shows a typical particle surroundedby the reaction layers. Both EDX measurements inSEM (chemical composition) and electron diffrac-tion in TEM (crystal structure and latticeparameters) were used to identify the phases of thedifferent regions. A W2C layer surrounds the WCparticle and next to the W2C a TiC reaction layeris present, which on its turn is itself, surroundedby W and TiC grains. The phase identificationagrees with the gray level in the SEM BSE micro-graph of Fig. 5, where the order of intensity isW � W2C � WC � TiC, according to theiratomic numbers. In some cases a small layer of Wis found to be present between the W2C and TiCreaction layers of which a transmission electronmicrograph is shown in Fig. 6.

The thickness of the reaction layers depends onthe laser processing parameters and on the positionin the melt pool. An increasing power and decreas-ing scanning velocity result in a thicker reactionlayer. In the top region the reaction layer becomesthicker than in the areas at the bottom. Forexample, for P � 79W/mm2, v � 8.3mm/s andm � 125mg/s the thickness of the W2C reactionlayer decreases approximately linearly with the

4916 J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

Fig. 1. SEM micrograph obtained in the secondary electron (SE) mode of a cross-section of five adjacent laser tracks. The distancebetween the centers of two neighbor tracks is 1.6 mm, which results in an overlap of about 10%. In this overlap region the LMIprocess is still successful and not changing much with the central areas of the individual tracks. The process parameters of a singletracks were P � 59W/mm2, v � 8.3mm/s and m � 83mg/s. The WC particles are the brighter features.

Fig. 2. XRD spectrum obtained by Cu-Kα radiation of a lasertrack. α-Ti, WC, W2C, W, TiC, β-Ti peaks are labeled.

Table 1Crystal structures of the phases those are present in the WC/Ti–6Al–4V layers

Phase Structure Lattice parameters (A)

WC Hex.(P6m2) a0 � 2.9,c0 � 2.8W2C Trig.(P3m1) a0 � 3.0,c0 � 4.7TiC Cub.(Fm3m) a0 � 4.3W Cub.(Im3m) a0 � 3.2α-Ti Hex.(P63/mmc) a0 � 3.3,c0 � 4.7β-Ti Cub.(Im3m) a0 � 3.2

depth of the particle (thickness is 3.2 µm at the topand 0.6 µm at a depth of 0.7 mm).

In most of the cases, no crystal orientationrelationship between the reaction layers is foundwith electron diffraction in TEM. Occasionally,however, an orientation relationship is observedbetween the TiC and the W2C layer, and the W2Clayer and the WC particle. An example is displayedin Fig. 7. The closed-packed planes and directionsare parallel, according to:

[110]TiC / / [1210]W2C , (111)TiC / / (0001)W2C

[1210]W2C / / [1210]WC, (0001)W2C / / (0001)WC

Automated EBSD measurements [12] were per-formed to analyze more grains in a larger area.With this technique electron back-scatter patterns(EBSP), that are essentially Kikuchi patterns, aregenerated by the interaction between a primaryelectron beam and a steeply tilted (70°) sample[11]. From the Kikuchi lines in these patterns theorientation of a grain can be obtained and, becausethe phases WC, W2C, TiC and W have differentcrystal structures, the phases can be discriminated.The method and preparation techniques involvedfor an MMC can be found in more detail else-where [13].

Fig. 8 shows a phase map of an area scan ofa corner of a particle with accompanying reactionlayers. In this map the grain boundaries are alsodepicted. Two WC grains are inside the scannedarea. The W2C and TiC reaction layers consist ofmany grains. The areas outside the TiC reactionlayer are indexed as W (dark grains in Fig. 8), butsince W has a similar crystal structure and latticeconstant as β-Ti, also β-Ti grains should be inde-xed as W. Nevertheless, because we are interested

4917J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

Fig. 3. SEM micrograph obtained in the back-scattering elec-tron (BSE) mode of the central/top region of a cross-section.The melt pool matrix shows the high amounts of TiC dendrites(dark, pointed by arrow (a) and W cells and dendrites (light,pointed by arrow (b) present in the melt pool matrix. There arealso a few WC particles in this micrograph visible.

Fig. 4. TEM micrographs of the melt pool matrix showing:(a) an area where TiC dendrites and W grains are present; (b)the martensitic structure of rapidly solidified Ti–Al–V alloy.

in the reaction zone it is not necessary to dis-tinguish between this two phases using OIM. Atsome places W is indexed between the W2C andTiC layer. Indeed, this is W and not β-Ti on thebasis of TEM observations (Fig. 6).

The 3D hexagonal unit cells are schematicallydepicted in the measured orientation, i.e. by 2Daxonometric projections, in order to visualize theorientations of the WC and the W2C grains (Fig.8). The upper left WC/W2C interface shows theorientation relationship as stated before with paral-lel basal planes. In addition, the TiC (not visualizedin Fig. 8) is oriented according the orientation

Fig. 5. SEM micrograph obtained in BSE mode of a particleand surrounding reaction zone. The WC particle is surroundedby a W2C layer, which on its turn is itself surrounded by a TiClayer. Next to this TiC layer W grains are present, alternatedby TiC grains.

Fig. 6. TEM micrograph of a W2C/TiC interface in the reac-tion zone where a small W grain is present.

relation observed in TEM. It is striking to see thatthese W2C and TiC are much larger and more regu-larly shaped, i.e. have constant thickness, than theother W2C and TiC grains. For all the other smallerW2C grains this orientation relation does not hold.However, it can be noticed that most W2C grainstend to have the [0001] direction perpendicular tothe WC/W2C interface.

The C distribution in the reaction zone is ana-lyzed with energy filtered TEM (EFTEM). An

4918 J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

Fig. 7. TEM micrograph, with corresponding diffraction pat-terns, of a reaction zone in the [1210] zone axis. In this casethere is an orientation relationship between the WC particle andthe W2C layer, and the W2C layer and TiC layer.

energy window with a width of 20 eV was usedto select an energy range of core-loss electrons,associated with C to form an elemental image. TheC K-edge at 283 eV was used to obtain a C imageof a WC–W2C–TiC reaction zone (Fig. 9a). Thezero-loss image is also given (Fig. 9b) to show thecorresponding microstructure. The C intensity pro-file divided by the zero-loss intensity and inte-grated over the marked width across the reaction

Fig. 8. Phase map and grain boundaries of the reaction zone (left). The orientation of the WC and some W2C grains is visualizedby plotting the axonometric projections of the hexagonal cells. The WC/W2C interface shows an orientation relationship in the leftupper corner. An SEM micrograph of the tilted sample is also shown (right).

Fig. 9. (a) Energy filtered carbon map of a WC–W2C–TiCreaction zone. In the brighter regions more C is present. (b) Thezero-loss image of the same area. (c) Integrated C intensity pro-file divided by zero-loss intensity across the reaction zone, cor-responding to the marked areas in a and c.

zone is shown in Fig. 9c. Fig. 9c shows that the Ccontent in W2C gradually decreases with decreas-ing distance from the WC particle. At the W2C/TiC interface the C content abruptly increases.

Wear test were performed on samples, the sur-face of which had been treated with a single laser

4919J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

track using optimum LMI conditions. The wearbehavior of these surfaces was compared with thatof the untreated substrate and with the surfacetreated by a simple laser remelting, without injec-tion of WC particles. Fig. 10 shows the wear lossduring wear tests for all three types of samples ata pressure of 20 MPa (curves 1,2 and 3). Moreover,the wear loss for WC/Ti–Al–V layer at contactstress of 100 MPa (curve 4) is also shown. Becauseof the very intense wear the tests for untreated andlaser beam remelted substrates were aborted after3 and 4 h, respectively. The common character-istics of wear behavior of all samples is that aftersome transition period a constant wear rate hasbeen established. Therefore two specific wear ratesk [10�6 mm3/Nm] shown in Table 2 were calcu-lated by two different methods: k1 was estimatedfrom the mass loss of pin determined by weightingit before and after wear experiment, using knowndensities of substrate alloy and WC particles andassuming about 25% volume fraction of WC in thecomposite; k2 was estimated from the slope of wearcurves in Fig. 10 in a regime of constant wear rate.Although values k1 and k2 are rather different forthe same samples, they both clearly demonstrate asubstantial increase in wear resistance of com-posite layer in comparison with the substrate alloy.Values of friction coefficient in Table 2 show, thataccording to the aim lubricating boundary con-

Fig. 10. Wear versus time for Ti–Al–V alloy (1); Ti–Al–Valloy remelted by laser beam (2) at 20 MPa contact stress; forWC/Ti–Al–V MMC at 20 MPa; (3) and 100 MPa (4) contactstress.

Table 2Specific wear rate calculated from the sample mass loss (k1)and from the slope of wear curve in Fig. 10 (k2); the averagevalue of friction coefficient fa

Sample Contact k1 (10�6 k2 (10�6 fastress mm3/Nm) mm3/Nm)(MPa)

Ti–Al–V 20 189 269 0.21Ti–Al–V 20 40 240 0.18remeltedWC/Ti–Al– 20 0.13 0.5 0.11V layerWC/Ti–Al– 100 0.08 0.05 0.12V layer

ditions prevailed [14]. In all the experiments thewear loss of disks in all experiments were underdetectable value, e.g. kdisk � 0.01 × 10�6mm3/Nm.

With SEM (Fig. 11) the morphology and quali-tative composition of wear surfaces of WC/Ti–Al–V layers tested at 20 and 100 MPa contact stresswere examined. The wear surface of the layertested at 100 MPa contact stress (Fig. 11a) showsa relatively flat morphology near the front edge ofwear surface (the right side). In the second part ofthe wear surface the rougher area is present withdeeply eroded regions between embedded WC par-ticles. BSE micrograph on the right side clearlydemonstrates that fragments of heavy particles(WC) continuously cover the flat part of the wearsurface and that inside the rough zone individualWC particles are upraised from eroded Ti alloysubstrate. The number of particles presented on thewear surface is still quite high. On the other hand,the wear surface of the layer tested at 20 MPa con-tact stress shows only ‘eroded’ wear surface mor-phology, with just a few towering WC particles.Fig. 11b shows a couple of such particles. Theshaded arrow indicates the highest place of theuphill area, which was in the direct contact withthe disk just before the wear test was stopped. Thetotal size of such areas forms only a small fractionof the whole wear surface and therefore, the actualcontact stress was much higher than the value of20 MPa calculated from the applied normal forceand sample dimension. Fig. 11b also demonstrates,that wear of WC particles proceeds by inter-granu-

4920 J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

Fig. 11. SEM micrographs of wear surfaces of WC/Ti–Al–V layers tested at contact stress of (a) 100 MPa; and (b) 20 MPa. Theleft side shows SE image and the right side BSE image from the same place. White arrows indicate the wear direction.

lar brittle fracture and that ‘ the most wear resistant’area is the hard surroundings of the WC particle.

4. Discussion

This paper reports on the exploration of the lasercladding technique, in which a progressive changein both microstructure and related properties areachieved over the molten boundary of laser poolin such a way that the usual interfacial problemsbetween a clad layer and a substrate are beingsolved. The materials we investigated fall withinthe class of ‘ functionally graded materials’(FGMs), a terminology that is now widely used bythe materials community for a class of materialsexhibiting spatially inhomogeneous microstruc-tures and properties. Graded materials in them-selves are not something new, but what is excitingabout them is the realization that gradients can bedesigned at a microstructural level to tailor specificmaterials for their functional performance in parti-cular applications [15]. Recent theoretical andexperimental work has established that controlledgradients in mechanical properties offer attractivechallenges for the design of surfaces with resist-ance to contact deformation and damage. In parti-cular the damage and failure resistance of surfaces

to normal and sliding contact or impact can bechanged substantially through such gradients. Fora recent review reference is made to Ref. [15].Therefore, a possible approach for eliminating theusually sharp interfaces between a protecting cladlayer and substrate material is to introduce the con-cept of FGMs into the LMI methodology for a bet-ter design of coatings [16–18].

Although the melting temperature of Ti–6Al–4Vis quite high (1650 °C [19]), a relatively modestlaser power density is needed to generate a meltpool. The reason is the relatively low reflectivityof Ti for a wavelength of 1.06 µm, compared toAl for instance [20]. The ease to create a melt poolin Ti is the principal reason there exists a relativelylarge process parameter window in which LMI issuccessful. This makes it possible to adjust thelaser track dimensions and particle volume fractionby changing the processing parameters. This is incontrast to SiCp injection in an aluminum substrate,where the parameter window is too small for largevariations of the parameters [5].

The fact that the injection process is successfulwithin a larger parameter window makes it alsopossible to produce adjacent tracks with an overlapso as to cover a larger area. When a second trackis produced partially over an already producedtrack the substrate surface of the later track consists

4921J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

of different areas. The surface of the to be treatedarea consists of an injected track but also partiallya heat affected zone and partially the unimpairedsubstrate. Therefore, the laser interaction with thesubstrate surface is changing within a laser track.However, the margins in the process parameterwindow make the LMI process still successful andrelatively constant within the whole treated area.This is the key to cover areas with multiple tracksthat are larger than laser and powder beam dimen-sion of a single track.

The microstructure is mainly determined by thechemical reactions in the melt pool. Due to thermo-dynamic reasons TiC will be formed when WCparticles will get in contact with liquid Ti [7],according to:

WC � Ti→TiC � W (1)

The extraction of C from WC due to the TiCformation may also degrade WC to W2C [21–23]:

2WC � Ti→W2C � TiC (2)

At the places where the temperature is higherand the liquid state of the melt pool is longerretained these reactions will develop more inten-sively. Therefore more TiC and W may be foundand consequently reaction layers are thicker in thetop–center of the melt pool. According to the thick-ness of the W2C layers and the fact that the grainboundaries of WC particles do not go over in W2Cpoint out that W2C is formed from the liquid state.Most likely, the surface layer of the WC particles(melting temperature of WC is 2870 °C [24]) ismelted when the laser beam interacts with the par-ticles during the melt injection process. Probablythe surface temperature of the WC particles is±3000 °C, which is still below the boiling point ofTi (3287 °C [24]).

The top part of this liquid WC layer will havea higher temperature than the part, which is closerto the liquid/solid interface, and therefore the vis-cosity will be higher at the surface of the WC par-ticle. When the particle penetrates in the melt poolthe top part of the melted WC may drop of theparticle and be left in the upper parts of the meltpool, while the particle penetrates further in themelt pool. Therefore we can distinct two differentcases of reaction during the laser process: between

small liquid droplets WC and Ti melt (mostly intop part of melt pool matrix) and between liquidWC, having a lower temperature than the droplets,at the solid WC particle and Ti melt.

The liquid WC that is left behind will react withliquid Ti (temperature probably around 2000 °C)according to reaction (1) to form TiC and W. Thetemperature of the reaction products will be ±3000°C, which is lower than the melting temperaturesof W (3410 °C [24]) and stoichiometric TiC (3140°C [24]), therefore W and TiC are rapidly solid-ifying. This is the reason why in the top part ofthe melt pool matrix many TiC and W dendritesare present (Fig. 3).

The liquid part of WC, which stays at the solidWC particle during the penetration towards thebottom of the melt pool, will react with Ti accord-ing to reaction (2) to form the reaction layers. TheTi melt will attract C to form a TiC layer andbecause of the lower C content, WC will transformto a W2C layer. In some cases the C content is solow that a small W layer is formed. Because thereaction and subsequent solidification takes placeat and near the WC particles, less reaction productsare found in the bottom of the melt pool matrix.There are TiC dendrites present in the matrix, incontrast to W2C that is only present at the WC par-ticles. The TiC dendrites in the bottom are formeddue to convection in the melt pool.

The particles at the top of the melt pool matrix,which are injected later (or with a lower injectionvelocity) and meet the solidification front at ahigher position in the melt pool, will experiencethe same reactions. Their reaction layers are thickerbecause less liquid WC has been lost during theirshort passage in the melt pool.

Ti that has not reacted with C rapidly solidifies,which explains the martensitic structures observed[25]. The metastable β-Ti is formed at higher tem-peratures (above 955 °C). This metastable phasecan be retained after rapid cooling of Ti containinga higher amount (�14.2 at%) of the β-stabilizer V[26]. Therefore, β-Ti can be found close to thereaction layers, where the relative amount of V inthe remaining Ti is comparatively higher becauseof the high amount of TiC that is formed there.

Whenever an orientation relationship is observedit is the [0001] direction in W2C perpendicular to

4922 J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

the WC/W2C interface (and (0001)WC

//(0001)W2C//WC/W2C interface). For the W2Cgrains that do not satisfy the orientation relationthe [0001] direction is in general tending to be per-pendicular to the WC/W2C interface as well.Therefore we may conclude that W2C prefers togrow in the [0001] direction. The crystal orien-tation of the W2C layer may be imposed by differ-ent reasons, for example interfacial strain (latticemisfits), heat flux (thermal conduction) and interfa-cial free energy. These various possibilities will bediscussed in the following paragraphs.

The interfacial strain between WC and W2C maybe a reason for a certain orientation. The latticeparameter mismatch between WC and W2C is ~3%for a0 and ~68% for c0 and therefore it is favorablefor the W2C to have the (0110) planes fitted onthe (0110) of WC. This is the case for the observedorientation relationship, as can be seen in Figs. 7and 8. However, because this orientation relation-ship is not frequently observed this is not the mainmechanism that determines the W2C orientation. Ifthe crystallographic configuration of the WC sur-face allows the orientation relation without chang-ing the growing direction from [0001] too much,i.e. the WC/W2C interface should be close to(0001), W2C will minimize the lattice misfits. Ifthis match is not possible the orientation of theW2C grain is solely determined by preferred grow-ing [0001]. In the latter case the orientation is notvery restricted, which results in many small W2Cgrains compared to the large W2C grains where theorientation relationship is fixed by both growingdirection and minimal lattice misfit.

Because the heat flux is, just as the [0001] direc-tion in W2C, perpendicular to the particle interface,it is reasonable to ascribe the preferred growingdirection to this heat flux. However, accordingthermal conductance in the W2C crystal, growingperpendicular to [0001],i.e. [1100] or [1120], isexpected. Therefore we conclude that the heat fluxis not the principal reason for the growth in the[0001] direction.

A principal reason for the preferred growingdirection might be the interfacial free energy. W2Cis growing in the [0001] direction because the for-mation of (0001) planes is, due to lower interfaceenergy, favorable. These planes consist of mono-

atomic C and W layers (one layer of C followedby two layers of W) and have a minimal numberof broken bonds.

Under lubricating boundary conditions MMCmaterials could compete easily with qualified tri-bomaterials [2,14,18]. The rather small differencein wear behavior of substrate and laser remeltedalloy can be explained by the presence of a thinand more wear resistant layer which is probablyformed during laser remelting on the top of lasertrack due to non-ideal gas shielding. After removalof this layer, the wear resistance of substrate andlaser remelted materials are almost identical, asvalues of specific wear rate k2 in Table 2 clearlyindicate. The value of k1 in Table 2 reflects alsothe initial stages of wear processes and this is themain reason why these two samples show a bigdifference in this quantity k1. The inspection ofworn surfaces of WC/Ti–Al–V alloy layer by SEMreveals that not embedded WC particles them-selves are responsible for the extreme increase ofthe wear resistance (530–1400 times). Actually thehard phases formed on the particle/substrate inter-face and also inside the substrate itself are respon-sible for this substantial enhancement. Once afused WC particle appears at the surface, it startsto crumble through the intergranular brittle fractureof individual grains inside one WC particle. There-fore fused WC particle becomes to be the weakestlink in WC/Ti–Al–V layer microstructure. This issimilar to the tensile properties of this material[27], when the crack nucleation process is con-trolled mainly by intergranular brittle fracture inWC particles. Completely different mechanicalcharacteristics can be expected for the same MMClayer prepared by the injection of single grain WCparticles [28]. Because of a good anchorage of WCparticles into matrix a further improvement can beexpected. In the wear experiment with high contactstress (100 MPa) small WC grains became disinte-grated from WC particles at the first stages of thewear process and subsequently these were incor-porated again into the matrix due to the high con-tact stress. This phenomenon contributes to theprolongation of initial stage of wear process andpartially suppresses the formation of wear-erodedareas, as Fig. 11a indicates. However, the thicknessof the layer removed by wear is in both cases

4923J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

smaller than a mean diameter of injected WC par-ticle. It means that only the most upper part ofWC/Ti–Al–V layer is playing the role during thewear test. This part of the layer is full of the newphases formed during the LMI process as micro-structural analysis reveals.

5. Conclusions

Laser melt injection is a suitable technique toform a WCp/Ti MMC in the top layer of Ti–6Al–4V. The laser processing window allows a certainvariation of the laser track dimensions and volumefraction of WCp. In addition it is possible to coata larger surface area with multiple adjacent trackswith an overlap of about 10%. An important find-ing is that the particle distribution is homogeneousand that the particles are injected over the wholedepth and whole width of the melt pool.

During the laser process new phases are formed:TiC, W and W2C. In the top of the resolidifiedmatrix large amounts of TiC and W dendrites arepresent because of the reaction between liquid WC,which is dropped of the WC particle during pen-etration in the melt pool, and the Ti melt. Theliquid WC that stays at the particle during pen-etration reacts with the Ti melt to form a W2C andTiC reaction layer around the WC particles.

W2C tends to grow in the [0001] direction.When the surface of the WC particle is close tothe (0001) planes, W2C minimizes the misfits atthe WC/W2C interface by growing according to anorientation relation, where the [0001] growth direc-tion is preserved and the (0110) planes of WC andW2C realize a good fit.

Substantial improvement of wear resistance isobserved for WC/Ti–Al–V surface layer preparedby LMI technique in comparison with substratealloy. Furthermore the wear of surface layer waseven lower, when higher load was applied. Wearbehavior demonstrates the successful adhesion ofWC particles in Ti alloy matrix and therefore thesignificance potential interests of LMI as techniquefor local wear resistance improvement of Ti andits alloys.

Acknowledgements

Financial support from the foundation for Fun-damental Research on Matter (FOM-Utrecht) andthe Netherlands Institute for Metals Research aregratefully acknowledged.

References

[1] American Society for Metals. In: 9th ed. Metals handbook,vol. 3. Metals Park, OH: American Society for Metals;1980. p. 35-3.

[2] Abboud JH, West DRF. J Mater Sci Lett 1991;10:1149.[3] Baker TN, Xin H, Hu C, Mridha S. Mater Sci Technol

1994;10:536.[4] Kloosterman AB, Kooi BJ, De Hosson JTM. Acta mater

1998;46:6205.[5] Vreeling JA, Ocelık V, van Agterveld DTL, Pei YT, De

Hosson JTM. Acta mater 2000;48:4225.[6] Zhou XB, De Hosson JTM. Acta mater 1996;44:421.[7] Barin I, Knacke O. Thermochemical properties of inor-

ganic substances. Berlin: Springer-Verlag, 1973.[8] Kooi BJ, Kabel M, Kloosterman AB, De Hosson JTM.

Acta mater 1999;47:3105.[9] Cooper KP, Slebodnick PL, Lucas KE, Hogan EA. J Mater

Sci 1998;33:3805.[10] Delannay F, Froyen L, Deruyterre A. J Mater Sci

1987;22:1.[11] Dingley DJ, Randle V. J Mater Sci 1992;27:4545.[12] Wright SI. In: Zwartz AJ, Kumar M, Adams BL, editors.

Electron backscatter diffraction in materials science. NewYork: Kluwer Academic/Plenum Publishers; 2000. p.51–64.

[13] Ocelık V, Vreeling JA, De Hosson JTM. J Mater Sci2001;36:4845.

[14] Mens JWM, Huis in’ t Veld AJ, de Gee AWJ, Sixth Inter-national Conference on Tribology, Boedapest; 1993. p.76–81.

[15] Suresh S. Science 2001;292:2447.[16] De Hosson JTM, Kooi BJ. In: Nalwa H, editor. Handbook

of surfaces and interfaces. New York: Academic Press;2001. p. 1–114 [chapter 1].

[17] Pei YT, De Hosson JTM. Acta mater 2000;48:2617; PeiYT, De Hosson JTM. Acta mater 2001;49:561.

[18] Tjong SC, Ma ZY. Mater Sci Eng 2000;29:49.[19] American Society for Metals. In: 9th ed. Metals handbook,

vol. 3. Metals Park, OH: American Society for Metals;1980. p. 38-9.

[20] American Society for Metals. In: 9th ed. Metals handbook,vol. 3. Metals Park, OH: American Society for Metals;1980. p. 66-1.

[21] Verdon C, Karami A, Martin JL. Mater Sci Eng1998;A246:11.

[22] Yongqiang Yang. Appl Surf Sci 1999;140:19.

4924 J.A. Vreeling et al. / Acta Materialia 50 (2002) 4913–4924

[23] Guilemany JM, De Paco JM, Nutting J, Miguel JR. MetallMater Trans 1999;30A:1913.

[24] Weast RC, editor. Handbook of chemistry and physics.36rd ed. Bocan Raton, FL: CRC Press; 1982–1983.

[25] Ahmed T, Rack HJ. Mater Sci Eng 1998;A243:206.[26] Folkes JA, Shibata K. J Laser Appl 1994;6:88.[27] Ocelık V, Vreeling JA, De Hosson JTM. Failure of

WCp/Ti–Al–V layer prepared by laser melt injection. In:

Brebbia CA, editor. Surface treatment V; computermethods and experimental measurements for surfacetreatment effects. Southampton: WIT Press; 2001. p.151–60.

[28] Ocelık V, De Hosson JTM. In: Dahotre N, editor. Proceed-ings 13th IFHTSE/ASM Surface Engineering Congress,October 2002, Columbus, OH, American Society for Met-als, 2002.