pulsed laser deposition and properties of mn+1axn phase formulated ti3sic2 thin films

Pulsed laser deposition and properties of Mnþ1AXn phase formulated

Ti3SiC2 thin films

J.J. Hua,�, J.E. Bultmanb, S. Pattonb and J.S. Zabinskic

aUES, Inc., Air Force Research Laboratory (AFRL / MLBT), Wright-Patterson Air Force Base, Ohio 45433-7750bUniversity of Dayton Research Institute, Air Force Research Laboratory (AFRL / MLBT), Wright-Patterson Air Force Base,

Ohio 45433-7750cMaterials and Manufacturing Directorate, Air Force Research Laboratory (AFRL / MLBT), Wright-Patterson Air Force Base, Dayton,

Ohio 45433-7750

Received 20 March 2003; accepted 16 June 2003

The ternary phase ceramic, Ti3SiC2; has often been synthesized through reactive hot pressing, providing bulk samples for

studying its mechanical and physical properties. Chemical vapor deposition has been the most popular route to make Ti3SiC2 films.

Recently, magnetron sputtering (MS) and pulsed laser deposition (PLD) have been used to produce good quality films. In this

paper, we present the results on the synthesis and tribological characterization of Ti3SiC2 thin films prepared by PLD. The films

had a surface roughness of 0.46 nm, a friction coefficient of 0.2 in humid air, and hardness of between 30 and 40GPa. The transfer

films were identified on the surface of counterparts using scanning electron microscopy. Anisotropic layer structure of Ti3SiC2 and

nano crystallites in the coatings observed by X-ray diffraction and transmission electron microscopy are related to the low friction

and high hardness. A specially designed sample coated with half Ti3SiC2 and half TiC was fabricated for comparing the properties

between the two materials using a lateral force microscope. Lateral force images of the coatings indicated that the lateral force

against Ti3SiC2 was lower than against TiC. PLD Ti3SiC2 coatings may be produced at near room temperature to 300 8C, which is

acceptable for many commercial applications.

KEY WORDS: Ti3SiC2, MAX phase, thin film, pulsed laser deposition, friction, hardness, tribology

1. Introduction

Titanium silicon carbide, Ti3SiC2; was first synthe-sized and characterized in 1967 by Jeitschko andNowotny using a chemical reaction of TiH2, Si, andgraphite at 2000 8C [1]. It was catalogued to the ternaryMnþ1AXnðMAXÞ phase, where M is an early transitionmetal, A is an A-group element (mostly IIIA and IVA)and X is either C and/or N. The MAX phases weregenerally determined as layered hexagonal carbides andnitrides in a sequence of n ¼ 1 or the 211 phases (e.g.,Ti2AlN;Ti2AlC) [2,3], n ¼ 2 or the 312 phases (e.g.,Ti3SiC2;Ti3AlC2) [1,4], and n ¼ 3 or the 413 phases(e.g., Ti4AlN3) [5]. In early investigations, the 312 phaseof Ti3SiC2 was expected to be useful as a soft ceramicmaterial, which was easily machinable, oxidationresistant, and lubricious [1,6–8]. After those initialreports, little work was done until the 1990s. Figure 1shows a histogram of the annual statistics of publishedpapers on Ti3SiC2 materials [1,6–26]. The interest inTi3SiC2 has surged in recent years with a good reviewfrom Barsoum’s group [9]. The renewed attentionresulted from its unusual combination of properties,such as high elastic stiffness, damage tolerance [10],fatigue toughness [11], easy machinability, good elec-

trical and thermal conductivities, excellent resistance tooxidation at high-temperatures [12–14] and resistance tothermal shock [14]. In addition, the strain-rate sensitiv-ity of Ti3SiC2 caused coatings to be brittle if loadedrapidly, but to be quite plastic if loaded slowly [15].Ti3SiC2 has a negligible thermoelectric power, whichwas defined as the open circuit electric field created by atemperature gradient [16]. The physical and chemicalproperties of MAX phase materials were related to theirlayered crystal structures that had been described asthermodynamically stable nano-laminates in Barsoum’sreview [9].

The Ti3SiC2 crystal has a layered structure of doubleTiC-like octahedrons, which are separated by planes ofsilicon where the atoms are arranged hexagonally in thebasal planes [17,18]. Compared to Ti–C bonds in theoctahedron, the Ti–Si bond between repeat layers alongthe c-axis is quite weak, and especially in shear, thebonding between the layers is weaker than along the in-plane layers. That suggests a similarity to well-knownsolid lubricants, MoS2 and graphite. The tribologicalbehavior, including friction and wear, of Ti3SiC2

samples has been reported and very low friction hasbeen observed, but challenges remain. The frictioncoefficient ð�Þ of Ti3SiC2 depends strongly on themeasurement technique, load, and surface finish. Pin-on-disc tribometer tests gave initial values of � from0.15 to 0.45 and steady-state values of 0.83 under several

�To whom all correspondence should be address. E-mail: jianjun.hu@

wpafb.af.mil

Tribology Letters, Vol. 16, Nos. 1–2, February 2004 (# 2004) 113

1023-8883/04/0200–0113/0 # 2004 Plenum Publishing Corporation

Newtons load. Sliding wear rates of between 1:3� 10�3

and 6� 10�2 mm3=Nm related to grain sizes have beenreported [19,20]. However, measurements taken by alateral force microscope (LFM) provided � values of0.003 on smooth regions, and up to 0.2 on rough regionsof the easy fracture/cleavage face of Ti3SiC2 polycrystal-line ceramics [21,22]. Therefore, the basal plane ofTi3SiC2 hexagonal cell exhibits very low friction. Thehardness of MAX phases is also known to be quiteanisotropic. Nickel et al. reported that the Ti3SiC2

hardness normal and parallel to the basal plane were12–15 and 3–4GPa; respectively [6]. Such anisotropy ofTi3SiC2 hardness may be useful in optimizing wearresistance through controlling crystal orientation [9]. Thelow friction, wear and oxidation resistance of Ti3SiC2

suggest that it has promise for tribological applications.Most of the available data were obtained from

Ti3SiC2 bulk samples that were prepared, for example,by reactively hot pressing of Ti, graphite and SiCpowders at 40MPa and 1600 8C [14]. Powder samples ofTi3SiC2 can be made by arc-melting and annealing of amixture of Ti, Si, and C powders [23]. Thin film coatingsof this material may provide surface protection and lowfriction for a number of engineering materials. Ti3SiC2

coatings may be particularly appropriate for microelec-tromechanical systems (MEMS) because the load andgeometry are quite similar to those in LFM experimentswhere the lowest � has been reported. Ti3SiC2 filmsproduced by chemical vapor depositions (CVD) [6,7],using SiCl4;TiCl4;CCl4, and H2 as precursor species,were formed at temperatures above 1300 8C. However,the high temperature employed in CVD limits theapplication of Ti3SiC2 to ceramic substrates. Using themagnetron sputtering (MS) method [24,25], epitaxialTi3SiC2 thin films have been grown at 900 8C on a TiC(111) seed layer deposited on MgO (111) substrates. Alower deposition temperature is required for steelsubstrates so that microstructure and dimensional

tolerance can be maintained. Our work focuses onphysical vapor deposition (PVD) of Ti3SiC2 thin films attemperatures below 300 8C [26]. In this paper, growthand characterization of Ti3SiC2 thin films prepared bypulsed laser deposition (PLD) is reported. Tribologicalproperties are discussed in relation to depositionconditions and microstructure.

2. Experimental

2.1. Pulsed laser deposition of Ti3SiC2 thin films

PLD is very effective for producing various coatingssuch as diamond-like carbon, oxides, and ceramics[27–29]. Combining laser deposition with magnetronsputtering (MSPLD) has enabled the preparation ofnano composites with multifunctional, self-adaptivebehavior [30–32]. Figure 2 shows the hybrid MS-PLDdeposition system, which was used for the presentTi3SiC2 film syntheses. The UHV chamber was evacu-ated to less than 10�7 Torr, and the pressure wasmeasured using a Bayert Alpert gauge. A LambdaPhysik LPX 110i excimer laser was used to provide apulsed beam of UV radiation of 248 nm wavelength,20 ns duration, 1–50Hz rate, and 200–600mJ energy.Beam steering was accomplished using a computercontrolled mirror system, which permits the laser beamto randomly strike the target over the selected ablationarea. Both the target and substrate were rotated toensure coating uniformity. A calibrated quartz crystaloscillator was used to monitor the coating thickness.Resistive heating was employed to anneal the substrate.Temperature was adjustable from room to 500 8C andwas measured using a calibrated infrared pyrometer. Acomputer was used to monitor and/or control thedeposition parameters including pressure, laser energyand pulse rate, magnetron sputtering power, gas flowand concentration, temperature, growth rate, and coat-

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Figure 1. Histogram of the statistic amount of publications on Ti3SiC2 materials per year.

J.J. Hu et al./Pulsed laser deposition and properties of Mnþ1AXn phase formulated Ti3SiC2 thin films114

ing thickness. Films prepared for this study were grownusing only the laser—the magnetron sputtering cap-ability was not employed but planned for futureexperiments.

Laser targets were fabricated by mixing Ti, C, andSiC powders at a 3 : 1 : 1 molar ratio, and were hotpressed into 1-inch disks. The target surface was cleanedin vacuum using the laser with sample shutter closed.Deposition was commenced by opening the shutter.

Ti3SiC2 films of approximate 0:25–0:50�m thicknesswere grown onto polished silicon (surface roughness of0.5 nm Ra) and steel (20 nm Ra) substrates by pulsedlaser ablation of the target. The deposition rate wascontrolled by adjusting the laser pulse frequency. Sub-strate temperatures were set at 100 and 300 8C to growTi3SiC2 films with different microstructures. The tribo-logical properties of the different coatings were evalu-ated using a ball-on-disc tribometer and a nanoindenter.

2.2. Analytical methodology

Coating chemistry and microstructure were studiedusing X-ray photoelectron spectroscopy (XPS), Ramanspectroscopy, X-ray diffraction (XRD), and transmis-sion electron microscopy (TEM). A surface ScienceInstruments M-probe XPS, equipped with a concentrichemispherical electron energy analyzer and operated ata base pressure of 10�8 Torr, was used to collectelemental and chemical information. The oxide andcontamination layers on coating surfaces were removedby sputter etching at a 4� 4mm2 area before XPS

acquisitions. Relative peak areas were used for compo-sition quantification, and corrected for a spectrometerfactor and X-ray cross-section. Raman spectra wereobtained using a Renishaw Ramascope 2000 equippedwith a 514:5 nmArþ laser and an air-cooled charge-coupled device (CCD). Raman analyses were used todetermine the bonding state of carbon atoms. Crystal-lographic data of coatings were collected using a Rigakuthin film XRD system with a monochromator in front ofthe CuK� X-ray source, which was operated in �–2�mode. XRD scans were also taken from the uncoatedside of the substrates to discriminate its peaks from thecoating. Atomic-resolution microstructures wereobserved using a Philips CM200 field emission gun(FEG) TEM operated at 200KV. The probe size ofincident electron beams was adjustable from 25 nmdown to 1 nm for chemical microanalyses using aNORAN X-ray energy dispersive spectrometer (EDS)on the TEM.

Coating topography was studied using a ParkScientific LS Auto Probe atomic force microscope(AFM). The cantilever used in this work was triangularwith a circular pyramid tip on end. The pyramid tip hasa height of 3�m; an apex angle of 158 and a circularbase of 1�m radius. The surface roughness wascalculated with 3-dimensional AFM image processing.The lateral force (resistance to sliding) was measured toevaluate friction at the so-called ‘‘single asperitycontacts’’ level that were formed between the tip andcoating surface [33,34]. Scans in contact mode (LFM)were made at a frequency of 3–4Hz with a set pointforce of 200 nN.

Pulsed laser248 nm300 mJ20 ns

Programmablemirror

Focusinglens

To downstream pressurecontrol and pump system

Substrateholder

Ar, N2

Ti-Si-CTarget

Ion-beam powersupply

Quartzoscillator

Magnetronpower supply

GroundedUHV chamber

Figure 2. Schematic illustration of the magnetron sputtering assisted pulsed laser deposition system.

J.J. Hu et al./Pulsed laser deposition and properties of Mnþ1AXn phase formulated Ti3SiC2 thin films 115

2.3. Tribology and mechanical properties

Friction coefficients were collected using a ball-on-disc tribometer run with the specimens in the horizontalposition. The measurements were made in air, at 35%relative humidity (RH), and at room temperature. A1.6mm diameter 440C steel ball was used as thecounterface. For each friction trace, a normal load of0.1N was set and the sliding speed maintained at30mms�1: The approximate rotational speed is 50RPMand the mean Hertzian normal contact stress is about500MPa. Cross-sectional wear scar area was measuredusing a KLA-Tencor profilometer, and was introducedinto wear rate calculations. After rubbing, both ball anddisc surfaces were examined using Leica 360 FEGscanning electron microscope (SEM), which wasequipped with a Link ISIS system of EDS for chemicalmicroanalyses.

The coating hardness and Young’s modulus weremeasured using a Nanoindenter IIs microprobe. Theindentation depth limit was 50 nm, and a diamondpyramid (Berkovich) indenter was loaded between1–2mN: Hardness calculations were based on Oliverand Pharr’s model using the maximum depth and loaddata [35]. The elastic modulus was calculated from theslope of the unloading segments in the load versusdisplacement curves. Fifteen indents were taken for eachspecimen and averaged.

3. Results

3.1. Microstructure characterization

Figure 3 shows the XRD 2�-scans from Ti3SiC2

coatings on silicon substrates at a deposition tempera-ture of 100 and 300 8C, respectively. Diffraction peaksfrom the coatings were indexed to the hexagonal Ti3SiC2

phase as shown in the figure. The (008) reflection wasmost prominent in the 100 8C coating, and the (102)reflection was most prominent in the 300 8C coating.Average crystallite size was estimated from the full

width value at half maximum (FWHM) of the diffractionpeaks. The calculated diameter was 5.6 nm for the 100 8Ccoating, and 12.5 nm for the 300 8C coating. High-resolution TEM images taken from the 100 and 300 8Ccoatings are shown in figure 4. The textured crystallites ascircled in the 300 8C coating are clearly larger than in the100 8C coating. Crystal lattice spacings of 2.2 A in figure4(a) and of 2.5 A in figure 4(b) correspond to the {008}and {102} planes of hexagonal Ti3SiC2 unit cells,respectively. Our XRD and TEM measurements resultedin the same conclusion of different textures formed byadjusting the deposition temperatures.

XPS measurements of the coating chemistry are givenin table 1. Both the 100 and 300 8C specimens deviatedfrom an optimum Ti3SiC2 stoichiometry. The bondingstate of carbon atoms in the coatings was further definedby Raman spectroscopy. Although the coatings were C-rich, no signal with Raman shift between 1200 to1700 cm�1 were detected, i.e., no graphite and diamond-like carbon (DLC) phases were observed in the Ramanspectra. If there was a significant amount of TiC [29], theXRD reflections (111) and (200) from TiC wouldbroaden the XRD peaks (102) and (008) from Ti3SiC2

in figure 3, respectively, on the right shoulder of theobserved peaks. In this work, XRD scans did not showthe existence of other phases in the Ti–Si–C system. Anelectron probe of 1.4 nm diameter was used to do EDSelement mappings for plan-view TEM samples of thecoatings. The mapping results showed a uniformdistribution of Ti, Si and C.

The 3-dimensional AFM images, as shown in figure5(a), were taken from the Ti3SiC2 coatings for topo-graphic information. The AFM images are compared tothe flat uniform coating shown in the cross-sectionalSEM image figure 5(b). The surface roughness of 100and 300 8C coatings was measured to be 2.32 and0.46 nm, respectively, by means of digital AFM imageanalyses. The later value is approximately equal to thesurface roughness of the Si substrate. Therefore, the300 8C coating had a much smoother surface than the100 8C coating.

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Figure 3. X-ray diffraction scans for the Ti3SiC2 coatings at deposition temperatures of 100 and 300 8C, respectively.


3.2. Friction and wear

Ball-on-disk tribometer test results are shown infigure 6. The 100 8C Ti3SiC2 coating failed in a veryshort time and the friction coefficient increased to 0.6 in50 s. A large wear scar was observed on the ball surfaceby optical microscopy as shown in figure 6(b). However,data from the 300 8C Ti3SiC2 coating showed an initialfriction coefficient as low as 0.1 (see inset in figure 6(c)).From the trace in figure 6(c), the steady-state frictioncoefficient rose to around 0.2 after 10,000 s. Examina-tion of the ball after rubbing against the 300 8C Ti3SiC2

coating using optical microscopy revealed a very smallwear scar as shown in figure 6(d). TEM observationsconfirmed that the c-axis of Ti3SiC2 nanocrystallinegrains lies on the 100 8C coating plane, as shown infigure 4(a). In that case, the sliding direction is normal tothe basal plane of the hexagonal Ti3SiC2 unit cell. The300 8C coating was grown where the h102i axis of theTi3SiC2 nanocrystalline grains lies on the coating plane,as shown in figure 4(b). Therefore, the sliding directionis closer to parallel to the basal plane than to thenormal. Referring to the crystal orientation mechanismsof super lubricity of MoS2 [36], sliding directionsparallel to the basal plane are preferred for super lowfriction, like that observed by LFM measurements ofTi3SiC2 ð� ¼ 0:003Þ [21,22]. The cross-sectional area ofthe wear scars was measured using a profilometer. Using

the values of load and sliding distance in the pin-on-disctest, an average wear rate of 1:8� 10�5 mm3=Nm wascalculated for the 300 8C coating, which is not unusualfor the wear-in period.

Transfer films were observed with high-magnificationSEM images taken from the ball surface. Figure 7(a) isan SEM image showing transferred films at a micronscale. The coating transferred from disk to ball wasfurther identified by EDS analyses. Ti, Si, and C wereidentified on the EDS spectrum as shown in figure 7(b).The weak Fe and Cr signals came from the steel ballbecause the interaction volume of electron and the X-ray transmission were both deeper than the transferredfilms. In addition, there was an oxygen peak in the EDSspectrum that could be caused by tribo-oxidation of Tiin the coatings [37]. TiO2 and TiO2�x have been reportedas lubricants, and have been applied for lubrication ofceramic materials [38]. The impact of Ti oxidization onthe tribology behavior of Ti3SiC2 required furtherinvestigation. Though oxide is in the debris, transferfilms and wear scars appear to be Ti3SiC2: Therefore, itis unlikely that TiO2�x is the active lubricant in thiswork.

Figure 8(a) shows an SEM image taken from thewear track of the 300 8C coating. A significant amountof debris was distributed along the two edges beside thewear track. At higher magnification of the wear scar, theTi3SiC2 coating exhibits plastic-like deformation, and atypical SEM micrograph showing the deformation ispresented in figure 8(b). Such a phenomenon in nanocomposite coatings has been generally attributed todiffusional creep and nanometer-sized grain boundarysliding [39]. Azimuthally disordered orientations ofTi3SiC2 nanocrystalline grains were observed in figure4 that had been suggested to minimize grain incoherencestrain and facilitate grain boundary sliding. In fact, theternary carbide Ti3SiC2 in bulk form is also plasticallyvery anisotropic [15]. Its layered crystal structure andhigh c=a ratio result in deformation occurring over-whelmingly by slip along the basal planes. The

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Figure 4. High-resolution transmission electron micrographs taken from (a) the 100 8C coating, and (b) the 300 8C coating.

Table 1

XPS measured chemical compositions of the Ti3SiC2 coatings at

deposition temperature of 100 and 300 8C.

At%

100 8C 300 8C Stoichiometry

Ti 47.20 46.49 50.00

Si 18.41 12.07 16.67

C 34.39 41.44 33.33


deformation is highly nonuniform and this nonuniformstate of stress may be relieved by decohesion at grainboundaries, microcracking, climb of dislocations, orgrain boundary sliding. Therefore, here the deformationmechanism can be theoretically modeled on the basis ofboth the Ti3SiC2 crystallographic features and the PLDcoating structure.

3.3. Hardness

Unlike the other carbides (such as TiC: 28GPa, HfC:27GPa, and WC: 23GPa), the reported hardness ofTi3SiC2 was low, while its elastic modulus was relativelyhigh. In fact, reports on polycrystalline bulk samples ofTi3SiC2 revealed a hardness of 4GPa and a Young’smodulus of 320GPa [14]. Such a high ratio of elasticmodulus to hardness had been compared to that ofductile metals [8]. Besides the fundamental materialsproperties, coating properties also depend on, forexample, coating structure, internal stress, and adhe-sion, which in turn depend on deposition conditions.Figure 9 shows the hardness and Young’s modulus datafrom the 100 and 300 8C Ti3SiC2 coatings on siliconsubstrates, respectively. The hardness and elastic mod-ulus of the 300 8C coating were both slightly higher thanthose of the 100 8C coating. The hardness values of thecoatings were between 30 and 40GPa. Therefore, theTi3SiC2 coatings grown by PLD have a 7–10 timeshigher hardness than Ti3SiC2 in bulk forms (hardnessnormal to the basal plane of Ti3SiC2 hexagonal cell isonly 3–5 times higher than that parallel to). The highhardness associated with high elastic modulus had been

theoretically desirable for wear improvement by pre-venting crack propagation. A high ratio of hardness toelastic modulus (H/E) has also been proposed as areliable indicator of good wear resistance in coatings[40]. Reasonably, the increased hardness of the coatingsresulted in a much better wear resistance than Ti3SiC2 inbulk forms [19,20]. The measured wear rate of Ti3SiC2

bulk samples was about two orders of magnitude higherthan the values we measured from the Ti3SiC2 coating.

There are several phenomena that may contribute tothe measured hardness. Different dominant orientationsand sizes of Ti3SiC2 crystallites in the 100 and 300 8CPLD coatings were observed by means of XRD andTEM, as shown in figures 3 and 4. For the 100 8Ccoating, the c-axis of crystallites lies on the coatingplane, and hence the hardness measurement is parallel tothe basal plane so that a lower hardness was observed[6]. Proper sized nanocrystalline grains can restrict cracksize, create a large volume of grain boundaries, andhence increase the coating toughness and hardness. Thecomposites consisting of nanocrystalline transitionmetal (Ti, V, W) nitride and amorphous Si3N4 canreach or even exceed the hardness of 50GPa when thecrystallite size decreases to 3:0–3:5 nm [41]. Anotherreport of the super-hardness effect for nano compositesshowed that 35 nm ZrN grains within a Cu matrix weresufficient [42]. Voevodin and Zabinski reported thathardness and toughness could be optimized by usingsomewhat larger, 10–20 nm grains [31]. By selecting thematrix and included particle chemistry correctly, coat-ings exhibited good tribological response across severalextreme environments. The size of Ti3SiC2 crystallites inthe 100 and 300 8C coatings was between 5 to 15 nm.

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Ti3SiC2 coating

Figure 5. (a) AFM topographic image, and (b) cross-sectional SEM image of the Ti3SiC2 coating on silicon substrate.


Except for the unlikely possibility of TiC nano crystal-lites, no other crystalline phase in Ti–Si–C system wasidentified. The grain boundary matrix was amorphousbecause the coating stoichiometry varied from Ti3SiC2:

The coating stress was measured between 380 and600MPa using a wafer curvature system. Therefore, theinternal stress was not too high and did not contributemuch to increased hardness. The super hardness effect

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Figure 6. Friction traces from the Ti3SiC2 coatings at deposition temperature of (a) 100 8C and (c) 300 8C, respectively, and the corresponding

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Figure 7. Transferred Ti3SiC2 film on the pin surface after rubbing: (a) SEM image, and (b) corresponding EDS spectrum.


and the amorphous matrix are more likely to beresponsible for increased hardness.

While friction was relatively good, the wear rate wasmarginal. Transition layers such as Ti and TiC havebeen considered for further investigations to relax stress,and to improve adhesion.

4. Comparison of Ti3SiC2 to TiC

There have been many publications on the synthesisand characterization of TiC based hard coatings. Inorder to directly compare the properties between PLDcoatings of Ti3SiC2 and TiC using LFM, we designedand fabricated a special sample. The two half surfaces ofa finely polished 440C steel plate were separately coatedby pulsed laser ablation of Ti3SiC2 and TiC targets,respectively, as shown on the right of figure 10. Thesubstrate temperature remained at 300 8C for thedepositions. This configuration was used to evaluatefriction so that difficulties in determining real � valuesby LFM were avoided (e.g., hard to calibrate load andlateral force). LFM images were taken at the junctionbetween the Ti3SiC2 and TiC coatings at a load of200 nN, as shown on the left of figure 10. The adhesionforce as the tip approaching to samples, which wasdetermined at a few nN, was negligible in comparison

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Figure 8. (a) SEM image of the Ti3SiC2 coating wear track and debris,

and (b) a high-magnification SEM image taken from the wear scar.

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Figure 9. Hardness and Young’s modulus of the Ti3SiC2 coatings at

deposition temperatures of 100 and 300 8C, respectively.

TiC

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Figure 10. LFM image of the PLD TiC and Ti3SiC2 coatings on a same plate with half side of TiC by half side of Ti3SiC2:Dark contrast in LFM

reveals low friction.


with the load of 200 nN. The Ti3SiC2 side exhibited amuch darker contrast than the TiC side. The lateralforce against the Ti3SiC2 coating, cited in photo-detector response (V) as shown as in figure 10, wasless than half of that against TiC. The LFM measure-ment of TiC gave a frictional coefficient of about 0.2[43]. Therefore, we may approximately estimate that thefriction coefficient of Ti3SiC2 is less than 0.1 in LFMmeasurements. A similar friction was determined using aHysitron system for the MS deposited Ti3SiC2 coatingson MgO substrates above 700 8C [44].

The hardness and Young’s modulus of Ti3SiC2 andTiC coatings were measured using a nanoindenter, andthe results are compared in figure 11. Although theTi3SiC2 is soft in bulk form, its hardness as a PLDcoating is quite close to that of TiC. The Young’smodulus of the Ti3SiC2 coating was slightly higher thanthat of TiC as shown on the right of figure 11. Ti3SiC2

coatings can have mechanical properties comparable toTiC, and offer the potential of much lower friction.

5. Conclusions

Ti3SiC2 coatings were synthesized by pulsed laserablation of Ti3SiC2 targets at 100 and 300 8C depositiontemperatures. They consisted of crystallites around 5 to15 nanometers according to the measurements of XRDspectra and TEM images. The coatings grown at 300 8Cnot only showed a smooth surface with 0.46 nm rough-ness, but also exhibited an initial friction coefficient of0.1 and had a steady-state value of 0.2. Transfer filmsfrom disk to ball suggested that the low friction was inpart due to Ti3SiC2 rubbing against itself. The hardnessof Ti3SiC2 coatings determined using a nanoindenterwas dramatically higher than that of Ti3SiC2 in bulkform. Differences in friction and hardness of thecoatings can be attributed to the anisotropy of Ti3SiC2

layered structures and to the texture of nanocrystallinegrains. In comparison to TiC, the Ti3SiC2 coatings haveadvantages of low friction plus the Ti3SiC2 propertiesdescribed in introduction, while a comparable hardness

is maintained. Ti3SiC2 may be particularly well suited toMEMS applications as under LFM conditions, frictionis less than half that of TiC. PLD synthesis provides ameans to deposit crystalline Ti3SiC2 at low depositiontemperatures ð100–300� CÞ:

Acknowledgments

The Air Force Office of Scientific Research (AFOSR)is gratefully acknowledged for financial support. Thanksto Dr. J. Nainaparampil for nanoindenter and AFMassistance, B. Philips for help with XPS, and A.Voevodin for helpful discussions.

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pulsed laser deposition and properties of mn+1axn phase formulated ti3sic2 thin films

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