Titanium carbide coatings deposited by reactive ion beam-assisted, electron beam–physical vapor...

Surface and Coatings Technology 124 (2000) 142–153www.elsevier.nl/locate/surfcoat

Titanium carbide coatings deposited by reactive ion beam-assisted,electron beam–physical vapor deposition

Douglas E. Wolfe *, Jogender SinghApplied Research Laboratory, Pennsylvania State University, State College, PA 16804, USA

Received 22 April 1999; accepted in revised form 8 November 1999

Abstract

Titanium carbide (TiC), in addition to other transition metal carbides, has excellent wear-resistant properties in machining ofstainless and high speed tool steels. TiC was successfully deposited on 304 stainless steel by reactive ion beam-assisted electronbeam–physical vapor deposition (RIBA, EB–PVD). The hardness values ranged from 1100 to 3500 VHN0.050 depending on theprocessing parameters. The surface morphology and degree of texturing were correlated with the processing conditions. Ionbombardment resulted in strong surface and growth texture in primarily the (111) orientation. A change from the (111) orientationto the (200) was observed at high substrate temperatures. Vickers hardness, scanning electron microscopy (SEM), X-ray diffraction( XRD) and electron probe microanalysis (EPMA) were used to characterize these films. © 2000 Elsevier Science S.A. Allrights reserved.

Keywords: Reactive ion beam-assisted electron beam–physical vapor deposition; EB-PVD; Titanium carbide coating

1. Introduction cathodic arc [17–19], evaporation [20–22], ion plating[23–26 ] and activated reactive evaporation (ARE) [27–30]. Depending on the deposition technique and process-In the last few decades, significant improvements

have taken place in coating processes and materials. A ing conditions, film properties vary significantly. Thispaper describes the preliminary results of TiC depositedchange from chemical vapor deposition (CVD) to physi-

cal vapor deposition (PVD) processes is underway. The by RIBA, EB–PVD.advantages and disadvantages of each process areexplained elsewhere [1,2]. PVD allows coating depos-ition at lower temperatures than CVD methods. 2. Experimental procedureHowever, with lower substrate temperature comes alower degree of adhesion between the substrate and the Titanium carbide (Ti

xC1−x) coatings were deposited

coating, which often results in delamination. To compen- by RIBA, EB–PVD in a Denton Vacuum (Modelsate for the lower degree of bonding between the coating DV-SJ/26) chamber as shown in Fig. 1. A more detailedand the substrate (as compared with high temperature description of the experimental apparatus is discussedCVD processes) bombardment processes such as bias elsewhere [31]. 304 stainless steel samplesand ion assist are employed in most PVD coating (2.5×5×0.076 cm) were mounted on a 15.24 cm diame-processes. In addition, ion bombardment processes also ter substrate holder and placed within a resistance heateraffect the morphology, composition, residual stresses shell. The source–substrate distance was 25.8 cm. Aand microstructure (i.e. density and degree of texturing) mechanical pump evacuated the chamber toof the deposited film. ~10−3 Torr, then a cryogenic pump evacuated the

TiC has been produced by CVD [3–8] as well as chamber to ~10−7 Torr. Once the chamber was at anumerous PVD processes such as sputtering [9–16 ], suitable pressure, the substrates were cleaned using

argon ion beam sputtering for 5 min at 600 V and600 mA.* Corresponding author. Tel.: +1-814-865-0316;

A 100 V negative d.c. bias was applied to the substratefax: +1-814-863-2986.E-mail address: [email protected] (D.E. Wolfe) holder rotating at 20 rpm. The samples were preheated

0257-8972/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved.PII: S0257-8972 ( 99 ) 00644-1

143D.E. Wolfe, J. Singh / Surface and Coatings Technology 124 (2000) 142–153

Fig. 1. Schematic drawing of the RIBA, EB–PVD apparatus.

and allowed to soak at the desired temperature for face morphology and coating thickness were determinedby optical microscopy and SEM.15 min (250 to 650°C). High purity (99.999%) titanium

metal was evaporated at the rate of ~0.294 g/min by abeam of electrons deflected at 270° toward the surface.Using a constant gas flow of 30 sccm argon and 90 sccm 3. Results and discussionacetylene, the pressure was maintained at2–3×10−4 Torr. During the evaporation of titanium, 3.1. Electron probe microanalysisan ion beam with current densities between 27 and162 mA/cm2 and energies between 400–500 eV was The composition of the RIBA, EB–PVD films was

determined using EPMA. The Ti/C ratio decreased asdirected towards the substrates to form TiC. The sampleswere post-heated for 15 min at the deposition temper- the temperature of the substrate increased, as shown in

Fig. 2. This strongly suggests that at constant pressureature before cooling.The films were deposited to a thickness of 15±2 mm. (~2.7×10−4 Torr) and ion beam current density

(75 mA/cm2) the composition of the films was dependentThe hardness of the coatings was determined using amicrohardness indenter (50 g load). As a result of the on the substrate temperature. This was expected as the

process consisted of both physically evaporating tita-relatively thin coatings (15±2 mm), the hardness valuesmay have been influenced by the substrate. This effect nium and relying on the dissociation of acetylene (into

hydrogen and carbon) in order for carbon to react withwas minimized by using a lower load (50 g) as well asby using identical substrate material (304 stainless steel ). titanium forming TiC (as in the CVD process). As the

temperature is increased, more energy is added to theThe composition of the films was obtained by EPMA.The crystal structure, lattice parameter, stress and pole system and the probability of dissociating acetylene is

increased. In addition, the higher temperatures decreasefigures were obtained using XRD techniques. The sur-

144 D.E. Wolfe, J. Singh / Surface and Coatings Technology 124 (2000) 142–153

carbon, TiC compound and derivatives of acetylenemolecules (C

xH1−x). Once adsorption occurs, the phy-

sisorbed species will migrate along the surface to aregion of low energy (surface mobility) before beingchemically absorbed (chemisorption) into the growingfilm. During migration, the species gain enough thermalenergy to desorb or sputter due to the constant bom-bardment of incoming ions. Surface mobility or surfacediffusion occurs as the system acts to minimize its totalfree energy. In contrast, higher energies are needed forTi and C ions to diffuse along the surface as they arebound more tightly, not only by adsorption (van derWaals forces), but also by covalent and metallic bonds(chemisorption). Therefore, sputtering of atoms wouldbe higher than for ions or compounds [1,4,32].Fig. 2. Ti/C ratio as a function of temperature at constant ion bom-

bardment (75 mA/cm2). Changing the density and energy of the bombardingspecies results in changes in composition due to preferen-tial sputtering and re-sputtering effects [32]. For exam-the activation energy barrier for absorption of theple, during ion sputtering, only 20–40 eV (depending onspecies into the growing film.the material ) is needed to sputter the surface atoms.Fig. 3 shows Ti/C ratio plotted as a function ofSince TiC is a compound, the sputtering rates of titaniumcurrent density. The Ti/C ratio increased with increasingand carbon may be different depending on the composi-current density between 27 mA/cm2 (near stoichiometriction and microstructure. Individual Ti and C atomsTiC) and 119 mA/cm2, then decreased aboveadsorbed on the surface would be ejected (sputtered) at119 mA/cm2 (Fig. 3) to produce near stoichiometric TiChigher rates than a TiC compound [15,32]. Since theagain. This was not expected. The Ti/C ratio of themass of carbon is lower than titanium, carbon wouldfilms was expected to decrease continuously with increas-preferentially sputter faster than titanium atoms, asing current density as more acetylene would decomposewould acetylene molecules. Based on the above explana-resulting in increased concentrations of carbon in thetion, Fig. 3 suggests that carbon is preferentially sput-growing film.tered more than the titanium atoms up to currentThe compositional variation in the thin films withdensities of 119 mA/cm2. At 148 and 162 mA/cm2 itvarying current densities may be explained by severalappears that the current densities are high enough tocompetitive reactions taking place simultaneously duringincrease the sputtering rates of titanium atoms resultingthe deposition process (i.e. adsorption, chemisorption,in a near stoichiometric ratio.desorption and sputtering effects). During vapor depos-

ition, the gaseous species adsorb onto the substrate3.2. Hardness: effects of temperaturesurface where they are bonded by weak van der Waals

forces. In this investigation, the physisorbed speciesThe Vickers hardness values ranged fromconsist of neutrals and ions of titanium, hydrogen, and

1301 VHN0.050 to 2274 VHN0.050 over the temperaturerange 250–750°C. For comparison, the hardness valueof 304 stainless steel is 490 VHN0.050. The general trendshows that the hardness of the TiC coating increased asthe process temperature was increased from 250 to650°C as shown in Fig. 4. The lower hardness value atTsub=750°C is most likely a combined effect of composi-tion and texturing ((200) orientation explained later).The variation in the hardness values is a direct result ofprocessing temperature (ignoring variations in thedeposition system) which influenced the compositionand microstructure (i.e. texture, density and surfacemorphology) of the TiC coatings.

The reaction between titanium and acetylene gas isan endothermic reaction, which means that energy mustbe supplied in order to overcome the activation barrierfor the chemical reaction to occur. The three mainFig. 3. Ti/C ratio as a function of current density (mA/cm2) at constant

temperature (650°C). methods of adding energy to the system are chemical,


cient). The hardness of the coating decreased withfurther increase of the substrate temperature to 750°C(Ti/C of 0.9). This confirms the findings reported byShikama et al. [34] and Jacobson et al. [35] that hardnessincreases with increasing carbon content. However, thisdoes not explain the low hardness value for the TiCcoating (hyperstoichiometric carbon) deposited at750°C. This sample should have had the highest hardnessresulting in interstitial carbon atoms [34]. One possibleexplanation could be the presence of free carbon ortitanium but XRD did not detect the presence of eitherphase. A more convincing explanation is that the hard-ness of the TiC films is a function of composition,microstructure and texturing as will be discussed later.

Fig. 4. Vickers hardness number vs. temperature (°C) for TiC depos-ited by RIBA, EB–PVD with constant ion bombardment 3.3. Hardness: effects of ion bombardment(75 mA/cm2).

Several samples were deposited under identical pro-cessing conditions (Tsub=650°C, constant pressurebombardment and temperature. No additional energy

was added through chemistry until the dissociation of (2–3×10−4 Torr), and gas flow ratio (1:3) of argon andacetylene) with the exception of the degree of ionacetylene occurred. Since the level of argon/acetylene

bombardment was constant for those coatings deposited bombardment (27–162 mA/cm2). The hardness valuesranged from 1721 to 3504 VHN0.050 over the range ofas a function of temperature, temperature was the main

process parameter used in initiating the chemical reac- 27–162 mA/cm2. The general trend shows that the hard-ness of the TiC coating increases with increasing levelstion between acetylene and titanium. The higher depos-

ition temperatures provided enough thermal energy to of ion bombardment as shown in Fig. 6. The generallyaccepted bulk hardness value for TiC is ~2800 VHNovercome the activation barrier for the formation of

TiC. These results are in good agreement with Bunshah [15]. The TiC coating deposited with a current densityof 162 mA/cm2 (3504 VHN0.050) showed over 25%[33] who reported that the hardness of TiC increased

over the temperature range 500–1000°C. improvement in the hardness value (near stoichiometry).It should be noted that, due to the limitations of theFig. 5 shows the Ti/C ratio versus Vickers hardness

number. At low deposition temperatures (250°C), the ion source, current densities above 162 mA/cm2 were notpossible under these processing conditions. This increaseaverage hardness was 1300 VHN0.050 corresponding to

a Ti/C ratio of 1.66. As the deposition temperature was in hardness can be attributed to a more dense film,texturing and intrinsic stress.increased up to 650°C, the corresponding hardness

increased and the Ti/C ratio decreased. The maximum In contrast, when the hardness of the above coatingswas plotted as a function of composition (i.e. Ti/Chardness (2274 VHN) was observed at a substrate tem-

perature of 650°C and Ti/C ratio of 1.3 (carbon defi- ratio), the highest hardness occurred at a Ti/C ratio of

Fig. 5. Vickers hardness number as a function of the Ti/C ratio for Fig. 6. Vickers hardness number as a function of current density(mA/cm2) for TiC deposited by RIBA, EB–PVD at constant temper-TiC deposited between 250 and 750°C by RIBA, EB–PVD at constant

ion bombardment (75 mA/cm2). ature (650°C).


Fig. 7. Vickers hardness number as a function of the Ti/C ratio for Fig. 8. Lattice parameter (nm) as a function of substrate temperatureTiC deposited by RIBA, EB–PVD at constant temperature (650°C). (°C ) deposited at constant ion bombardment (75 mA/cm2).

a graphitic phase, no shift in the lattice parameter would1.06 (Fig. 7). This is quite surprising, yet very informa-tive as it strongly suggests that the Ti/C ratio may not occur, except possibly due to uniform semi-coherent or

incoherent strain between the two phases. A more validbe the determining factor in the degree of hardness inthe films. These results support the finding that hardness explanation would be that the carbon was incorporated

into the film in interstitial positions, which would explainvalues are also a function of coating microstructure (i.e.degree of texturing, morphology and density), or possi- the increase in the lattice parameter. However, signifi-

cant amounts of carbon interstitial atoms are onlybly entrapped gases such as argon and hydrogen. Itshould be noted that although the coatings were depos- expected to occur when the Ti/C ratio is less than one.

In this study, only the sample deposited at 750°Cited under identical processing conditions, compositionalvariations could occur resulting from the equipment contained excess carbon. Since the level of argon bom-

bardment was kept constant (and assuming a constantitself or due to sputtering of the growing film duringdeposition. argon incorporation into the film with increasing tem-

perature) more carbon interstitial atoms occur withincreasing temperature resulting in an increase in the3.4. X-ray diffractionlattice parameter. According to the Ti–C phase diagram,no free carbon would be expected for carbon deficientThe XRD patterns of the TiC coatings were compared

to the standard Joint Committee for Powder Diffraction films in the composition range produced in thisinvestigation.files and were determined to be of the cubic phase. The

lattice parameters ranged from 0.4294 to 0.4339 nm Since no other phase was detected by XRD, it isreasonable to assume that the expansion of the latticedepending on the processing conditions. In all of the

samples, the XRD results supported the coatings being is a direct result of large amounts of lattice defects inthe TiC coating. These lattice defects may be interstitials,TiC with lattice parameters close to the accepted stoi-

chiometric TiC value (0.4327 nm). No other phase (i.e. substitutions or vacancies. TiC has a wide solubilityrange for defects. On the atomistic scale, a perfect TiCcarbon or titanium) besides TiC was detected in the

coatings. lattice would contain two titanium and two carbonatoms per unit cell. Each titanium atom is surroundedFig. 8 shows the lattice parameter plotted as a func-

tion of substrate temperature. The trend shows that the by six carbon atoms and each carbon atom is surroundedby six titanium atoms. As the lattice accommodateslattice parameter increases as the deposition temperature

is increased. Since the level of ion bombardment was defects, such as titanium and carbon vacancies, thestructure of the phase does not change but the latticekept constant at 75 mA/cm2, the only additional energy

added to the system was thermal energy via substrate parameter can. One possible explanation for thesevacancies is that the forcible injection of argon breaksheating. This additional thermal energy, combined with

the constant bombardment, resulted in increased surface the surface bonds of the growing film and producescarbon Shottky defects before becoming entrapped inmobility which contributed to surface atom rearrange-

ment in addition to the increased carbon concentration. the growing film. The impinging argon atoms force thecarbon atoms out of the equilibrium lattice positionsSince no other phase was detected by XRD, it is

believed that the increasing carbon content in the films forming a carbon interstitial.Fig. 9 shows the lattice parameter plotted as a func-with increased temperature did not precipitate into a

graphitic phase. If the additional carbon was present as tion of current density. As in Fig. 8, the trend shows


Fig. 9. Lattice parameter (nm) as a function of current density Fig. 10. Lattice parameter (nm) vs. Ti/C ratio.(mA/cm2) at constant temperature (650°C).

that the lattice parameter increases with increasing cur- as a function of the atomic ratio Ti/C. The latticeparameter increased from 0.4294 nm (Ti0.6C0.4) torent density. Sundgren et al. [16,36 ] found a similar

trend with increasing levels of bias (a form of bombard- 0.4333 nm (Ti0.52C0.48), then up to 0.4346 nm(Ti0.33C0.67). This is in good agreement with Shikamament) for sputter-deposited TiC and TiN films. At a

current density of 27 mA/cm2, the lattice parameter was et al. [11] and Sundgren et al. [16 ] who found that thelattice parameter increases with increasing carbon con-found to be 0.4333 nm, and increased to 0.4339 nm at

a current density of 162 mA/cm2. This increase in the centration even above the stoichiometric composition.Earlier studies [34,35] suggested that excess carbonlattice parameter can be explained based on argon and

hydrogen incorporation into the film. As the mixture of (above the stoichiometric composition) should have thesame lattice parameter as stoichiometric TiC as theyargon and acetylene exits the ion source, the atoms and

molecules are ionized. These energetic ionized gases are believed that the excess carbon was a separate phase inthe form of graphite, as shown by the phase diagramaccelerated towards the substrate surface both through

the accelerating voltage and by the applied negative for Ti–C [39]. However, numerous authors have recentlyproduced hyperstoichiometric carbon TiC films [40,41].substrate bias. As the energetic argon and acetylene ions

bombard the surface, sputtering of the surface occurs Fig. 11(a)–(e) show the relative intensities for thediffracting planes (h/2h) deposited as a function ofdue to momentum transfer. In addition, argon ions

break surface bonds before becoming implanted a few temperature. At low substrate temperatures (between250 and 650°C) the coatings had a (111) orientation.atomic levels into the coating [32]. The penetration

depth is highly dependent on the kinetic energy of the At a higher substrate temperature (750°C), the orienta-tion changed to (200). It is believed that at even higherbombarding species. During this ‘ion implantation’ step,

stress develops within the film as the atoms are no substrate temperatures the orientation would change to(220), as interpreted from Fig. 11. The relative intensitylonger at equilibrium lattice positions. It is assumed that

as these excess carbon, hydrogen and argon atoms get for the (200) diffracting planes increases from 59.9% atTsub=250°C to 90.4% at Tsub=650°C before changingembedded in the coating, the lattice parameter increases,

as shown in Fig. 9. from the (111) to the (200) orientation at Tsub=750°C(100% relative intensity). The same trend occurs for theThorton and Hoffman [37] suggested that it is

unlikely that argon is incorporated into metallic films (220) and (311) planes. This is explained by surfacemobility. At low temperatures, the fast growth planeswith similar atomic radii. However, this study pertained

to amorphous materials. It is believed that this assump- grow themselves out of existence (i.e. high indexed, lowenergy planes [(200) and (220)] while the low indexed,tion does not take into account the fact that TiC is a

defect structure which can accommodate large amounts high energy planes survive. As the substrate temperatureis increased, more of the high indexed, low energy planesof argon similar to the high concentration of voids in

amorphous materials trapping argon. There have been survive as shown by Fig. 11(a)–(e) (higher relative inten-sities). At a substrate temperature of 750°C, there isseveral studies discussing the possibility of argon incor-

poration occurring during film growth [14,38]. Energy enough energy for the surface atoms to move aroundallowing the low energy, high indexed planes to survive.dispersive spectroscopy of the films confirmed the pres-

ence of argon in only a few of the TiC coatings as The intensities for the diffracting planes at 750°C arevery similar to those of a randomly oriented powderpresented in the next section.

Fig. 10 shows the dependence of the lattice parameter sample (Fig. 11(f )). These results are similar to those


Fig. 11. Relative intensity of the diffraction planes (h/2h) showing the varying degree of texture for TiC coatings deposited by RIBA, EB–PVD atTsub=(a) 250°C, (b) 450°C, (c) 550°C, (d) 650°C, (e) 750°C, and (f ) a randomly oriented sample.

reported by Bunshah and Raghuram [42] who showed free energy (thermodynamic contributions) [43–45].During the early stages of film growth, a (200) orienta-TiC coatings with a (220) orientation between depos-

ition temperatures of 520 to 1080°C before changing to tion was reported. As the film thickness increased, thestrain energy associated with thicker films becomesa (200) orientation between 1120 and 1450°C. They

believed that at even higher deposition temperatures the greater than the surface free energy resulting in a (111)orientation. Others suggest that kinetic factors such asorientation of the film would become even more random

as a result of increased surface mobility. film composition and ion incident angle and energies[46,47] are more important in determining the orienta-The degree of texturing for the samples deposited

under varying degrees of ion bombardment all had (111) tion of the growing films. The present results suggestthat the preferred orientation must be a complex phe-orientation. There was no change in orientation over

the range of current densities studied. The degree of the nomenon in which both kinetic and thermodynamiceffects contribute to the preferred orientation (especially(111) orientation is discussed in the next section as

determined by the pole figures. as a function of processing conditions). Since the filmthickness in the present investigation is quite largeLittle agreement for the evolution of preferred orien-

tation exists in the literature, as it is still not fully (15 mm), the thermodynamic argument of the strainenergy being greater than the surface free energy doesunderstood. Recently, the origin of texture was explained

based on competition between strain energy and surface not hold true, unless thermal contraction stress is consid-


Table 1Crystallographic orientation as a function of temperature showing the angle (away from the normal parallel to the surface) of maximum intensityfor the (111), (200) and (220) planes within the TiC films. Samples were deposited under the following conditions: current density 75 mA/cm2mixture of 30 sccm Ar and 90 sccm acetylene; and evaporation rate 0.294 g/min

Tsub (°C) (111) angle Degree of texture (200) angle Degree of texture (220) angle Degree of texture Overall degree of texture

250 20° V. weak 35° V. weak 15° V. weak V. weak450 15° Moderate 45° Weak 15° V. weak Weak550 30° Moderate 30° Weak 15° V. weak Weak650 15° Moderate 35° Moderate 35° Weak Moderate750 35° Strong 25° V. strong 15° Moderate Strong

Table 2Crystallographic orientation as a function of current density showing the angle (away from the normal parallel to the surface) of maximum intensityfor the (111), (200) and (220) planes within the TiC films. Samples were deposited under the following conditions: Tsub=650°C; mixture of 30 sccmAr and 90 sccm acetylene; and evaporation rate 0.294 g/min

Current density (111) angle Degree of texture (200) angle Degree of texture (220) angle Degree of texture Overall degree of texture(mA/cm2)

27 0° Strong 55° Moderate 30° Moderate Moderate75 15° Moderate 35° Moderate 35° Weak Moderate

119 5° Strong 50° Moderate 35° Moderate Moderate148 5° Moderate 50° Moderate 15° Moderate Moderate162 5° Strong 60° Strong 30° Moderate Strong

ered. In this investigation, the film orientation changed list the maximum psi angles for each plane. It shouldfrom (111) to (200) with increasing temperature. The be noted that, since the samples were rotated, very littleopposite should occur according to thermodynamic texturing occurred within the textured plane (phi angle).treatment alone which predicts (200) orientation during This type of texturing is often referred to as fibre texture.the early stages of growth before changing to a (111) The films grow with a preferred orientation perpendicu-orientation. It is clear that preferred orientation is a lar to the surface parallel, but randomly oriented withincomplex phenomenon which is a function of thermal the plane parallel to the surface. As discussed previously,and kinetic energy as well as the deposition processing texturing occurred in all of the samples to some degree.parameters. The degree of texturing for each of the samples was

qualitatively described as being weak, moderate or3.5. Pole figures strong.

Fig. 12(a) and (b) show the pole figures for a TiCTo understand the degree of texturing in the deposited coating which is weakly oriented. From the wide spacing

films better, pole figures were determined. Pole figures of the pole contours in conjunction with the weakare stereographic projections which describe a three intensity values, the coating did not appear to be highlydimensional crystal orientation using a two dimensional

oriented with respect to a particular set of planes. Inexternal system of axes. The axes chosen for these figures

contrast, the pole figure for TiC which is stronglywere phi and psi. Pole figures describe the statisticaloriented is shown in Fig. 12(c) and (d). The closelyaverage distribution of the pole density (i.e. directionsspaced contour lines and high intensity values stronglydrawn at a 90° angle from the plane of interest) for asuggest that the film was strongly oriented with thecertain plane within a polycrystalline material. The polemajority of the (111) oriented parallel to the surface,density is uniform in a randomly oriented polycrystallinewith the (200) oriented 55° from the normal to thematerial. When the density of pole contours is concen-surface. The present pole figures are consistent with thetrated in the stereographic projection/circle, this iscubic phase where the angle between the (111) andreferred to as preferred orientation or texturing.(200) plane is 54.7°.Partial pole figures were obtained for the (111), (200)

The trend appears to show that the coatings becomeand (220) planes for each of the samples. The lack ofmore textured with increasing temperature as well asobserved symmetry in a few of the pole figures is thewith increasing levels of ion assist. It should also beresult of obtaining the patterns for only one quadrantpointed out that the angle of argon and acetylene(psi from 0°–90°). A true pole figure is conducted fromimpingement on the growing film was ~50°. Littlepsi=−180° to +180°. However, this is not possible for

coatings deposited on thick substrates. Tables 1 and 2 texturing within the plane parallel to the surface


Fig. 12. Measured pole figures for: (a) weak (111) orientation; (b) weak (200) orientation; (c) strong (111) orientation; and (d) strong (200)orientation for TiC deposited by RIBA, EB–PVD.

occurred due to the rotation of the samples parallel to bombardment of the coating occurs. Residual stressesrelated to thermal expansion mismatch become morethe source material.significant above 500°C.

Generally, compressive stresses are desired for wear3.6. Stressapplications [28]. However, too high a level of compres-sive stress can lead to delamination of the coating fromThe stress state of a deposited film can often have a

significant effect on the performance of the coating. the substrate. Fig. 13 shows the stress as a function ofsubstrate temperature determined by XRD analysis. TheGenerally, evaporated films have a tensile stress.

However this is very dependent on the coating process stresses are compressive in nature (most likely due tothe bias (bombardment) for all the films depositedparameters as well as substrate/coating properties. If a

high thermal expansion mismatch between the substrate between 250 and 750°C. The amount of compressivestress increases between 250 and 650°C but decreases atand the coating exists, large residual extrinsic stresses

may result on cooling. If the thermal expansion of the 750°C. This decrease in compressive stress at 750°C ismost likely the result of thermal relaxation [48].substrate is larger than that of the coating, the substrate

will contract more during cooling resulting in large Fig. 14(a) shows the stress as a function of currentdensity at constant substrate temperature (Tsub=compressive stresses in the coating. If the coating has a

larger thermal expansion coefficient than the substrate, 650°C) for TiC deposited by RIBA, EB–PVD. Asexpected, the amount of compressive stress increasestensile extrinsic stresses will develop, assuming that no


ing compositional variations and bombardment (i.e.number, type, angle and size), as well as the energy ofthe particles impinging on the surface. Due to theexperimental configuration, controlling the compositionindependently of the level of ion bombardment wasextremely difficult.

To eliminate the compositional contribution (ofcarbon and possibly hydrogen) to the amount of stressin the coatings, TiC was directly evaporated from TiCpellets with increasing levels of argon bombardment(maximum Tsub=210°C). The results are shown inFig. 14(b). The coatings were deposited at a lowersubstrate temperature (as compared to Fig. 14(a)) tominimize stresses associated with thermal expansionmismatch. The general trend shows an increase in com-Fig. 13. Stress (MPa) as a function of substrate temperature (°C) atpressive stress (Fig. 14(b)) with increasing levels ofconstant ion bombardment (75 mA/cm2) for TiC deposited by RIBA,

EB–PVD. argon bombardment. Also, EDS, Energy DispersiveSpectroscopy suggested that argon was present in eachof the films. However, quantifying the argon was diffi-cult. Therefore, it is believed that the amount of com-pressive stress is associated with the amount ofentrapped gas incorporated into the film.

3.7. Surface morphology

The surface morphologies of the samples depositedwith increasing substrate temperature are shown inFig. 15(a)–(e). It appears that the grain size increasedas the substrate temperature was increased. This isprimarily the result of coarsening. The surface morphol-ogy of the TiC films grown between 250 and 550°Cshowed no evidence of surface texturing. At highertemperatures (650°C), surface texturing was observed.The surface morphology of Fig. 15(e) is very differentfrom the textured surface of Fig. 15(f )–( j). The surfaceappears to show small amounts of surface porosity (asmarked in Fig. 15(e)) possibly due to the (200) orienta-tion, which is less closely packed than the (111) orienta-tion. This surface porosity may have also contributedto the low hardness value.

The surface morphologies of all of the samples depos-ited at 650°C with increasing levels of ion bombardmentappear to show evidence of surface texturing as shownby Fig. 15(f )–( j). The grain size was difficult to deter-mine due to the high degree of surface texturing. Theredoes not appear to be any significant change with thedegree of surface texturing by increasing the ion currentdensity. This is strongly supported by the XRD results

(a)

(b)

which show all of the films to have a (111) orientation.Fig. 14. (a) Stress (MPa) as a function of current density (mA/cm2) atconstant temperature (650°C) for TiC deposited by RIBA, EB–PVD.(b) Stress (MPa) as a function of current density (mA/cm2) for TiCdeposited by direct EB–PVD (maximum Tsub=210°C). 4. Conclusions

TiC was successfully deposited by RIBA, EB–PVDwith increasing levels of ion bombardment. The exactnature of the compressive stress is not fully understood. on 304 stainless steel. The effects of temperature and

ion bombardment were discussed and were found to beStress occurs in coatings by a number of factors includ-


Fig. 15. SEM micrographs showing the surface morphology of TiC by RIBA, EB–PVD as a function of temperature: (a) 250°C; (b) 450°C; (c)550°C; (d) 650°C; (e) 750°C; and current density: (f ) 27 mA/cm2; (g) 75 mA/cm2; (h) 119 mA/cm2; (i) 148 mA/cm2; and ( j) 162 mA/cm2.

consistent with those found in the literature. A change of 3504 VHN0.050. The lattice parameter and amount ofcompressive stress was shown to increase with increasingfrom (111) orientation to a (200) orientation was

observed at high substrate temperatures (750°C). Ion temperature and ion bombardment. The increase wasattributed to the incorporation of argon, hydrogen andbombardment resulted in strong growth and surface

texturing which contributed to the high hardness value carbon interstitials.


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