ARTICLE IN PRESS

0042-207X/$ - s

doi:10.1016/j.va

�CorrespondE-mail addr

Vacuum 81 (2006) 479–488

www.elsevier.com/locate/vacuum

Deposition and characterization of TiAlN/Si3N4 superhardnanocomposite coatings prepared by reactive direct current

unbalanced magnetron sputtering

Harish C. Barshilia�, B. Deepthi, K.S. Rajam

Surface Engineering Division, National Aerospace Laboratories, Post Bag No. 1779, Bangalore–560 017, India

Received 19 April 2006; received in revised form 3 July 2006; accepted 9 July 2006

Abstract

Superhard nanocomposite coatings of TiAlN/Si3N4 with varying silicon contents were synthesized using reactive direct current (DC)

unbalanced magnetron sputtering. The Si and TiAl targets were sputtered using an asymmetric bipolar-pulsed DC power supply and a

DC power supply, respectively, in Ar+N2 plasma. The structural and mechanical properties of the coatings were characterized using

X-ray diffraction (XRD) and nanoindentation techniques, respectively. The elemental composition of the TiAlN/Si3N4 nanocomposite

coatings was determined using energy-dispersive X-ray analysis and the bonding structure was characterized by X-ray photoelectron

spectroscopy. The surface morphology of the coatings was studied using atomic force microscopy. The XRD data showed that the

nanocomposite coatings exhibited (1 1 1) and (2 0 0) reflections of cubic TiAlN phase. The broadening of the diffraction peaks with an

increase in the silicon content in the nanocomposite coatings, suggested a decrease in the average crystallite size. The TiAlN/Si3N4

nanocomposite coatings exhibited a maximum hardness of 43GPa and an elastic modulus of 350GPa at a silicon concentration of

approximately 11 at%. The hardness and the elastic modulus of the nanocomposite coatings decreased significantly at higher silicon

contents. Micro-Raman spectroscopy was used to characterize the structural changes as a result of heating of the nanocomposite

coatings in air (400–850 1C) and in vacuum (900 1C). The Raman data of the nanocomposite coatings annealed in air and vacuum showed

better thermal stability as compared to that of the TiAlN coatings. Similarly, the nanocomposite coatings deposited on mild steel

substrates exhibited improved corrosion resistance.

r 2006 Elsevier Ltd. All rights reserved.

Keywords: TiAlN/Si3N4 superhard nanocomposite coatings; Unbalanced magnetron sputtering; Structural and mechanical properties; Thermal stability;

Corrosion behavior

1. Introduction

Thin films of transition metal nitrides have been widelyused in many engineering applications especially due totheir high hardness, chemical inertness and excellent wearresistance. Among them, the properties and the applica-tions of TiN coatings have been studied extensively. Themain disadvantage of TiN is its limited oxidation resistance(approximately 450–500 1C). The addition of other ele-ments such as Al, Cr, Si, etc. increases the oxidationresistance of TiN [1–3]. Initially, TiAlZrN-based thin films

ee front matter r 2006 Elsevier Ltd. All rights reserved.

cuum.2006.07.003

ing author. Tel.: +91 80 2508 6494; fax: +91 80 2521 0113.

ess: harish@css.nal.res.in (H.C. Barshilia).

were developed for an electronic application in 1972 [4].Subsequently, TiAlN coatings have been developed for theengineering applications as an alternative to TiN since 1986[5]. It has been reported that, the addition of aluminum toTiN, thus forming TiAlN, improves the oxidation behaviorand the thermal stability of the coating, by forming a stableoxide layer on the surface of the film during oxidation [6].The properties of TiAlN coatings are mainly controlled bythe aluminum content. TiAlN coatings with hardnessvalues varying in the range 20–35GPa have been reportedin the literature [7,8]. Even hardness as high as 47GPa hasbeen achieved for TiAlN coatings prepared using magne-tron sputtering [9]. These superhard films were reported tobe nanocomposite films of TiAlN/AlN, having relatively

ARTICLE IN PRESSH.C. Barshilia et al. / Vacuum 81 (2006) 479–488480

large (approximately 30 nm) TiAlN grains oriented in onedirection and surrounded by an amorphous and/ornanocrystalline-AlN phase [9].

In recent years, there has been a considerable progress inthe development of newer coatings such as superhardcoatings (hardness X40GPa) [10,11]. The superhard coat-ings are the new generation futuristic materials havingexotic properties. These coatings are being developed in anattempt to replace the conventional superhard materialssuch as diamond and cubic boron nitride, which exhibitsevere limitations for high-temperature applications. Thesuperhard coatings find applications in diverse fields suchas moving mechanical components, cutting tools, dies, etc.The superhard coatings can be prepared mainly by tworoutes: (1) nanolayered multilayers (such as TiN/NbN [12],TiAlN/CrN [13], TiN/CrN [14], etc.) and (2) nanocompo-sites [11,15]. The multilayer coatings have been studiedextensively, however, the nanocomposite coatings havebeen developed very recently. These coatings exhibitsuperior mechanical properties. It is believed that theenhanced mechanical properties of the nanocompositecoatings are due to the fact that, at very small grain sizes(p10 nm), the number of atoms in a boundary region,which surrounds the grain, is comparable to that in thegrain [16]. This hinders the generation of dislocations andprevents crack propagation due to the grain boundaryenhancement and suppression of the grain boundarysliding [16].

A substantial number of nanocomposite coatings havebeen studied [17–19]. Among them, nanocomposite coat-ings of TiN/Si3N4 have been studied extensively due totheir high hardness and elastic modulus, improved wearresistance and high oxidation resistance [20–22]. Theprominent hardness of the nanocomposite TiN/Si3N4

coatings and improved thermal stability of TiAlN coatingsled to the exploration of TiAlN/Si3N4 nanocompositecoatings. Park et al. [23] have reported that hardness andYoung’s modulus of the Ti–Al–Si–N films increased withthe incorporation of silicon and exhibited a maximumhardness and a Young’s modulus of 55 and 650GPa,respectively, at a silicon content of 9 at% [23].

Nanocomposite coatings of TiAlN/Si3N4 have beenprepared using different techniques such as arc ion platingand direct current (DC)/RF sputtering [23–25]. Generally,RF sputtering has been used for the sputtering of thesilicon target because of its low electrical conductivity. Butin this case, the deposition rates are low, since sputteringtakes place only during one-half of the cycle [26]. In recentyears, pulsed plasma-reactive sputtering has been devel-oped in order to achieve oxide and nitride coatings withhigh growth rates [27]. Asymmetric bipolar-pulsed DCpower supply in which the polarity of the target voltageperiodically shifts from negative to positive is ideal forreactive sputtering of silicon target [28].

In the present investigation, the superhard nanocompo-site coatings of TiAlN/Si3N4 have been prepared usingreactive DC unbalanced magnetron sputtering of a

composite TiAl target (DC sputtering) and a silicon target(pulsed sputtering) in Ar+N2 plasma. The nanocompositecoatings have been deposited at various silicon concentra-tions. Structural and mechanical properties of the coatingshave been measured using XRD, X-ray photoelectronspectroscopy (XPS), energy-dispersive X-ray analysis(EDAX), nanoindentation, atomic force microscopy(AFM) and micro-Raman spectroscopy techniques. Thethermal stability and the corrosion behavior of the TiAlN/Si3N4 nanocomposite coatings have also been investigatedin detail.

2. Experimental details

Superhard nanocomposite coatings of TiAlN/Si3N4 weredeposited on mild steel and silicon substrates using areactive DC unbalanced magnetron-sputtering system. Thesputtering system consisted of two sputtering guns(diameter ¼ 0.075m), which were inclined at an angle tothe normal. The use of these guns allows independentsputtering of the two materials. The nanocompositecoatings were deposited using a high purity TiAl(99.95%) composite target and a high purity silicon(99.999%) target [27]. The composition of the TiAl targetwas approximately 50:50. A DC power supply was used tosputter the TiAl target and an asymmetric bipolar-pulsedDC power supply (frequency ¼ 100 kHz, pulse width ¼2976 ns, positive pulse bias ¼+37V) was used to sputterthe silicon target. The substrates were chemically cleaned inan ultrasonic agitator using acetone, isopropyl alcohol andtrichloroethylene. A base pressure of 4.0� 10�4 Pa wasmaintained in the vacuum chamber using a turbomolecularpump backed by a rotary pump. The substrates weredegassed at 300 1C for 10min and were subjected to Ar+

ion bombardment at a substrate bias of �800V and apressure of 4.0� 10�1 Pa for 30min to remove oxideimpurities. The flow rates of argon and nitrogen werecontrolled separately by mass flow controllers. Prior to thedeposition, the targets were sputter cleaned for 5min. Atitanium interlayer of approximately 0.5 mm thickness wasdeposited on the substrates in order to improve theadhesion of the coatings. A DC substrate bias of �100Vwas applied during the deposition, to improve themechanical properties of the coatings. The nanocompositecoatings were deposited at a substrate temperature of300 1C and an Ar+N2 gas pressure of 1.0� 10�1 Pa. Forall the experiments, the power density for the TiAl targetwas 4.93W/cm2. In order to obtain nanocomposite coat-ings with different silicon contents the power density forthe silicon target was varied from 0.88 to 2.19W/cm2. Thethickness of the nanocomposite coating was approximately1.5 mm. The process conditions for the deposition of theTiAlN/Si3N4 nanocomposite coatings are summarized inTable 1.X-ray diffraction (XRD) data of the films were recorded

in a Rigaku D/max 2200 Ultima X-ray powder diffract-ometer using a Cu Ka radiation (l ¼ 0:15418 nm).

ARTICLE IN PRESS

Table 1

Parameters for the deposition of TiAlN/Si3N4 nanocomposite coatings

Step Process Parameters

1. Substrate cleaning with ion

bombardment

Substrate bias ¼ �850V, Ar flow rate ¼ 17 sccm, pressure ¼ 4.0� 10�1 Pa, duration ¼ 30min

2. Target cleaning TiAl: power ¼ 30W, duration ¼ 5min

Si: power ¼ 40W, duration ¼ 5min

3. Ti interlayer Ti target: power ¼ 200W, substrate bias ¼ �100V, Ar flow rate ¼ 17 sccm, pressure ¼ 1.0� 10�1 Pa,

substrate temperature ¼ 300 1C, deposition time ¼ 30min, thickness ¼ 0.5 mm

4. TiAlN/Si3N4 nanocomposite

deposition

Nitrogen flow rate ¼ 3.5 sccm, argon flow rate ¼ 17 sccm, pressure ¼ 1.0� 10�1 Pa, substrate

bias ¼ �100V, substrate temperature ¼ 300 1C

TiAl target: power ¼ 250watts

Si target: power ¼ 40–90watts

Deposition time ¼ 150min

Thickness of the coating ¼ 1.5 mm

5. Cooling down and venting 45min

H.C. Barshilia et al. / Vacuum 81 (2006) 479–488 481

The elemental compositions of the coatings were deter-mined using EDAX is in a scanning electron microscope(Leo 440I). The bonding structure of the nanocompositecoatings was characterized by XPS using an ESCA 3000(V.G. Microtech) system with a monochromatic Al Ka

X-ray beam (energy ¼ 1486.5 eV and power ¼ 150W). Themechanical properties of the coatings (hardness and elasticmodulus) were measured in a nanoindenter (CSEMInstruments). For hardness measurements, indentationswere made on the sample using a Berkovich diamondindenter. After initial contact of the indenter on thesurface, the load was increased at a predetermined rate(10mN/min) to the desired maximum load and thendecreased at the same rate to zero. The maximum appliedload was 5mN for all the samples. From the load vs.displacement plot, the hardness was calculated using Oliverand Pharr method [29]. More details of nanoindentationmeasurements are described elsewhere [30]. Surface ima-ging of the as-deposited samples was carried out usingAFM. The maximum scan ranges for AFM in the X, Y andZ-axes were 40, 40 and 4 mm, respectively. The AFMresolution along the X, Y and Z-axes was o1 nm. TheAFM was operated in non-contact mode.

In order to test the thermal stability, TiAlN and TiAlN/Si3N4 nanocomposite coatings were simultaneously heatedin air in a resistive furnace at TA ¼ 700, 750, 800 and850 1C. Annealing involved increasing the temperature ofthe samples from room temperature to the desiredtemperature at a slow heating rate of 3 1C/min andmaintaining the desired temperature for 30min. Subse-quently, the samples were cooled down at a rate of3 1C/min. Structural changes as a result of the heating weremeasured using micro-Raman spectroscopy. The coatingswere also annealed in vacuum to separate the effect ofoxidation-induced spectral changes in the Raman data. Forthis, the samples were vacuum-sealed in a quartz tube atapproximately 3.0� 10�3 Pa pressure. The quartz tube was

then placed in the furnace, and the samples were annealedat 900 1C for 30min under identical heating and coolingrates. A Dilor–Jobin–Yvon–Spex integrated micro-Ramanspectrometer was used for the present study [31].The corrosion behavior of the coatings was studied using

potentiodynamic polarization in 3.5% NaCl solution underfree-air conditions at room temperature. The experimentalconditions for the polarization measurements are describedelsewhere [32]. The corrosion potential (Ecorr), the corro-sion current density (icorr) and the polarization resistance(Rp) were deduced from the Tafel (log i vs. E) plots.

3. Results and discussion

3.1. Structural characterization

Single-phase TiAlN coatings with a nanoindentationhardness of approximately 26GPa were obtained for anitrogen flow rate of 3.5 sccm, a substrate bias of –100V, asubstrate temperature of 300 1C and a power density of4.93W/cm2. The lower panel of Fig. 1 shows the XRDpattern of the TiAlN coating deposited under the aboveconditions. The high intensity peak centered at 2y ¼ 42:401is assigned to TiAlN (2 0 0) and the peak with a lowerintensity at 2y ¼ 36:461 is assigned to TiAlN (1 1 1) [4].Other higher angle reflections such as (2 2 0), (3 1 1), (2 2 2),(4 0 0) were very low in intensity. Subsequently, nanocom-posite coatings of TiAlN/Si3N4 were deposited at differentsilicon concentrations. The relative silicon content of thenanocomposite coatings was determined using EDAX. TheXRD patterns of the TiAlN/Si3N4 nanocomposite coatingsdeposited at various silicon contents are also shown inFig. 1. As the silicon content increased, the intensities ofthe (1 1 1) and (2 0 0) reflections decreased gradually,showing an increase in the amorphous content in thecoatings. A significant broadening of the diffraction peakswas observed at higher silicon contents. This originates

ARTICLE IN PRESSH.C. Barshilia et al. / Vacuum 81 (2006) 479–488482

from the decrease in crystallite size of the coatings.Determination of the average crystallite size using broad-ening of the XRD peaks indicated an average crystallitesize of approximately 80 A for the TiAlN coatings. Ingeneral, the crystallite size decreased with an increase in thesilicon content in the nanocomposite coatings. Forexample, the average crystallite size of the nanocomposite

2θ (deg.)

Inte

nsity

(ar

b. u

nits

)

30 35 40 45 50

TiAlN

(a)

(b)

(c)

(d)

•

♦

•

•

•

•♦

♦

♦

(111

) (200

)♦

TiAlN/Si3N4

Fig. 1. XRD patterns of TiAlN/Si3N4 nanocomposite coatings deposited

at various silicon contents: (a) 8.8 at% (b) 17.6 at% (c) 20 at%. (d) 22 at%.

Also shown is the XRD data of TiAlN coating.

Binding Energy (eV)

Inte

nsity

(ar

b. u

nits

)

390 395 400392.5 397.5 402.5

N 1s 396.5 eV

N 1s 398.1 eV

Binding En

Inte

nsity

(ar

b. u

nits

)

69 7371

(a)

(c)

Fig. 2. XPS spectra of (a) N 1s peak (b) Si 2p peak and (c) Al 2p peak for t

coatings decreased to about 36 A for a silicon content of20 at%. Similar results have been reported in the literature[33,34].It must be pointed out that the micro-strain present in

both the TiAlN and the TiAlN/Si3N4 nanocompositecoatings due to the substitution of Ti by Al or Si mighthave significantly contributed to the broadening of theXRD peaks [35]. Therefore, the crystallite size reportedabove must be smaller than the actual crystallite size of thecoatings. In addition, in general there was a shift in the 2yvalues of the diffraction peaks towards lower angles, withan increase in the silicon content. This shift can beattributed to higher residual stresses in the coatings.Residual stresses tend to shift the diffraction peaks tohigher d-values (or lower 2y values) [36]. In the presentstudy, a 0.94% increase in the d-value was observed forTiAlN/Si3N4 nanocomposite coatings when compared tothe TiAlN coating. Similar results have been reported inthe literature [37,38]. The XRD data of single-phase Si3N4

coating showed no peaks, which indicated its amorphousnature.Fig. 2 shows high-resolution XPS core level spectra of a

TiAlN/Si3N4 nanocomposite coating with a silicon contentof 11 at%. Deconvolution of the N 1s spectrum (Fig. 2(a))revealed the presence of two peaks centered at 396.5 and398.1 eV. The high-intensity peak at 396.5 eV correspondsto nitrogen in TiN or AlN, as the binding energies ofTiN and AlN have been reported to be very similar—396.7and 396.8 eV for TiN and AlN, respectively [39]. The

ergy (eV)

75 7977

Al 2p 74.05 eV

Binding Energy (eV)

Inte

nsity

(ar

b. u

nits

)

95 100 105 110

Si 2p 102.0 eV

(b)

he TiAlN/Si3N4 nanocomposite coating with a silicon content of 11 at%.

ARTICLE IN PRESS

33

36

39

42H

Har

dnes

s (G

Pa)

300

315

330

345

360

TiAlN

Ela

stic

Mod

ulus

(G

Pa)

E

H.C. Barshilia et al. / Vacuum 81 (2006) 479–488 483

low-intensity peak centered at a binding energy of 398.1 eVis a characteristic of nitrogen in Si3N4 [39]. In Fig. 2(b), thepeak centered at 102.0 eV originates from the Si 2pelectrons of Si3N4 [39]. Peaks pertaining to free silicon(99.0 eV) and other silicon compounds were not observedin the spectrum, indicating that the bonding state of siliconwas in the form of Si3N4. The Al 2p spectrum (Fig. 2(c))showed a characteristic peak at a binding energy of74.05 eV, which corresponds to AlN [40]. From the XRDand XPS data, it is clear that the nanocomposite coatingsconsisted of TiAlN nanocrystals and amorphous phaseof Si3N4.

8 10 12 14 16 18 20 22

27

30

Silicon Content (at.%)

255

270

285

TiAlN

Fig. 4. Variations of nanoindentation hardness and elastic modulus of

TiAlN/Si3N4 nanocomposite coatings with silicon content.

3.2. Mechanical characterization

Fig. 3 shows typical load vs. displacement curves forTiAlN, Si3N4 and TiAlN/Si3N4 nanocomposite coatings at5mN load. The maximum indentation depth was about88 nm for the TiAlN/Si3N4 nanocomposite coating, whichis less than 1/10th of the coating thickness. For the TiAlNand the Si3N4 coatings, the maximum indentation depthswere about 92 and 114 nm, respectively. The variations ofhardness and elastic modulus of the TiAlN/Si3N4 nano-composite coatings with silicon content are shown inFig. 4. The nanocomposite coatings exhibited a maximumhardness of 43GPa and an elastic modulus of 350GPa fora silicon content of about 11 at%. TiAlN and Si3N4

coatings prepared under identical conditions exhibitedhardness values of 27 and 26GPa, respectively. Themaximum hardness of the TiAlN/Si3N4 nanocompositecoatings is very high as compared to the rule-of-mixturesvalue (approximately 26.9GPa). The hardness and theelastic modulus of the nanocomposite coatings decreased

0 15 30 45 60 75 90 105 1200

1

2

3

4

5 (a)(b)(c)

(c) TiAlN/Si3N4

Loa

d (m

N)

Indendation Depth (nm)

(b) TiAlN

Unl

oadi

ngLoad

ing

(a) Si3N4

Fig. 3. Typical load vs. displacement curves for TiAlN, Si3N4 and TiAlN/

Si3N4 nanocomposite coatings at 5mN load.

significantly with further increase in the silicon content.For example, at a silicon concentration of approximately20 at%, the nanocomposite coating showed a hardness of32GPa and an elastic modulus of 260GPa. The decrease inthe hardness is due to the increase in the amorphous Si3N4

content in the TiAlN matrix. As mentioned earlier, thecrystallite size of the TiAlN/Si3N4 nanocomposite coatingsdecreased with an increase in the silicon content in thecoating. The small crystallite size of the nanocompositecoating hinders the generation of dislocations, preventscrack propagation and suppresses grain boundary sliding.This results in the enhanced mechanical properties of thenanocomposite coatings [16]. However, at very smallcrystallite size (i.e. when the amorphous Si3N4 content isvery high), the hardness of the nanocomposite coatingdecreases.Figs. 5(a) and (b) show the three-dimensional surface

morphologies of TiAlN coating and TiAlN/Si3N4 nano-composite coating (11 at% Si), respectively. In both thecases, the films were found to be smooth with a root meansquare roughness value of approximately 2.2 nm for theTiAlN coating and 2.23 nm for the TiAlN/Si3N4 nano-composite coating.

3.3. Thermal stability

In Fig. 6, we plot the composite Raman spectra of as-deposited TiAlN and TiAlN coatings heat-treated atdifferent temperatures. The spectrum of TiAlN showedtwo broad bands centered at 312 and 598 cm�1. Thesebands originate due to the acoustic transitions in the150–350 cm�1 region (LA and TA) and the optic modes inthe 400–650 cm�1 region (LO and TO) [31]. For sputtered

ARTICLE IN PRESS

0

5000

10000

10000

5000

40

80

nm nm

0

(a)

nm

(b)

0

5000

10000

40

80

10000

5000

0

nm

Fig. 5. Three-dimensional AFM images of 1.5-mm-thick (a) TiAlN coating and (b) TiAlN/Si3N4 nanocomposite coating. For the nanocomposite coating,

the silicon content was 11 at%.

150 300 450 600 750 900 1050

598

Raman Shift (cm-1)

312

25°C

800°C

750°C

700°C

Inte

nsity

(ar

b. u

nits

)

850°C

449

521

613

TiAlN

Fig. 6. Composite Raman spectra of as-deposited TiAlN coating and

coatings heat-treated up to 850 1C in air.

H.C. Barshilia et al. / Vacuum 81 (2006) 479–488484

TiAlN coatings, scattering in the acoustic range isprimarily determined by the vibrations of the Ti ions andthat in the optic range by vibrations of the N ions [31]. Thespectrum did not change significantly after heating thesample at 700 1C. However, at TA ¼ 750 1C additionalpeaks centered at 449 and 613 cm�1 were observed. This isdue to the formation of rutile TiO2 phase [31]. Theintensities of these peaks increased with an increase in theannealing temperature. An additional peak centered at521 cm�1 also emerged which originates from the siliconsubstrate (Table 2). The intensity of the acoustic modeband decreased with TA, whereas, the optic mode banddisappeared completely at TAX750 1C. This indicatesremoval of nitrogen atoms from TiAlN matrix andoxidation of the coating at TAX750 1C. Raman spectraof as-deposited TiAlN/Si3N4 nanocomposite coating andcoatings heat-treated up to 850 1C are shown in Fig. 7. As itis shown in the figure, the Raman spectra of thenanocomposite coatings did not change up to 750 1C andonly at TA ¼ 800 1C, peaks associated with TiO2 (rutile)started appearing (See Table 2).It has been reported that the activation energy for the

oxidation of TiN coating is 136 kJ/mol [6]. For TiAlNcoating with low aluminum content, the activation energyhas been reported to be 185 kJ/mol and at higher aluminumcontent, the activation energy has been reported to be inthe range of 260–460 kJ/mol [6]. Similarly, the activationenergy for the oxidation of chemical vapor-depositedsilicon nitride has been reported to be 461 kJ/mol [41].Therefore, the higher oxidation resistance of TiAlN/Si3N4

nanocomposite coatings can be attributed to the compara-tively high activation energies of TiAlN and Si3N4. Thesuperior oxidation behavior of the TiAlN/Si3N4 nanocom-posite coatings is also due to the strongly differing valuesof Gibbs free energy for the oxide formation (DG1 ¼

�378.2, �212.6 and �204.75 kcal/mol for Al2O3, TiO2 andSiO2, respectively) [42]. In addition, the amorphous Si3N4

phase boundary around the TiAlN crystallites suppressesdiffusion of oxygen and therefore reduces the oxidationrate of the TiAlN/Si3N4 nanocomposite coatings [15].

ARTICLE IN PRESS

Table 2

Assignment of various Raman peaks of as-deposited and heat-treated TiAlN and TiAlN/Si3N4 nanocomposite coatings deposited on silicon substrates

Sample Peak position (cm�1) Peak assignment

TiAlN

(a) As-deposited 312 (B) Longitudinal and transverse acoustic

598 (B) Transverse optic mode

(b) Heat-treated in air (850 1C) 521 (S) Si

449 (S) TiO2 (rutile)

613 (S) TiO2 (rutile)

(c) Heat treated in vacuum (900 1C) 323 (S) Longitudinal and transverse acoustic

536 (W) Not identified

607 (S) Transverse optic mode

TiAlN/Si3N4 nanocomposite (11 at% Si)

(a) As-deposited 314 (B) Longitudinal and transverse acoustic

614 (B) Transverse optic mode

(b) Heat-treated in air (850 1C) 449 (S) TiO2 (rutile)

613 (S) TiO2 (rutile)

(c) Heat treated in vacuum (900 1C) 315 (S) Longitudinal and transverse acoustic

590 (S) Transverse optic mode

B, broad; S, strong; W, weak.

H.C. Barshilia et al. / Vacuum 81 (2006) 479–488 485

In addition to the oxide formation, heating of thenanocomposite coatings in air can also lead to otherchanges like phase transformation. These changes affectthe nature of the Raman spectra. In order to study thestructural stability of the coatings at higher annealingtemperatures, they were heated in high vacuum(3� 10�3 Pa) for 30min at 900 1C. The representativeRaman spectra of the as-deposited and the vacuum-annealed coatings of TiAlN and TiAlN/Si3N4 nanocom-posite coating are given in Fig. 8. After annealing at900 1C, no significant change in the nature of the Ramanspectra was observed in the case of the TiAlN/Si3N4

nanocomposite coatings. However, for TiAlN coatingsheated under similar conditions at 900 1C, a broad andweak peak emerged at 536 cm�1 in addition to the acousticand optic modes. The origin of this peak is not clear atpresent and may be attributed to the non-stoichiometricphase of TiAlN, suggesting structural changes in TiAlNcoatings as a result of annealing at 900 1C. These resultsdemonstrate that the nanocomposite coatings studied inthe current work showed no phase transformation evenafter heat treatment in vacuum up to 900 1C.

3.4. Corrosion behavior

The corrosion behavior of TiAlN, Si3N4 and TiAlN/Si3N4 nanocomposite coatings deposited on mild steelsubstrates was investigated using potentiodynamic polar-ization in 3.5% NaCl solution under free-air conditions atroom temperature. The Tafel plots obtained for the mildsteel substrate, single-phase TiAlN, single-phase Si3N4 andTiAlN/Si3N4 nanocomposite coatings are shown in Fig. 9.The corrosion potential, the corrosion current density andthe polarization resistance obtained from the measure-

ments are given in Table 3. The corrosion potential,increased from �0.647V for the uncoated mild steelsubstrate to �0.529V for the TiAlN/Si3N4 nanocompositecoating with a silicon content of 20 at%. The corrosioncurrent density is commonly utilized as an importantparameter to evaluate the kinetics of corrosion reactions.The corrosion rate is normally proportional to thecorrosion current density. For the mild steel substrate,the corrosion current density was about 4.996 mA/cm2,which decreased to 1.973 mA/cm2 in the case of TiAlN/Si3N4 nanocomposite coating. Single-phase TiAlN andSi3N4 coatings exhibited corrosion currents of 3.603 and0.186 mA/cm2, respectively. The decrease in icorr of coatedsubstrate therefore confirms improvement in the corrosionresistance. Furthermore, the corrosion resistance was alsofound to increase with an increase in the silicon concentra-tion in the TiN/Si3N4 nanocomposite coatings. Rp in-creased from 7.862 kO cm2 for a silicon concentration of8.8 at% to 13.295 kO cm2 for a silicon concentration of20 at%. Single-phase TiAlN and Si3N4 coatings exhibiteda corrosion resistance of 5.260 and 105.78 kOcm2,respectively.Several factors are believed to be responsible for the

enhanced corrosion properties of the nanocomposite coat-ings. These include: non-columnar and dense micro-structure, reduced crystallite size of the nanocompositecoatings, amorphous micro-structure of Si3N4 and poorelectrical conductivity of Si3N4. The dense micro-structureand reduced crystallite size of the nanocomposite coatingsare attributed to the ion bombardment during depositionand incorporation of amorphous phase in the TiAlNmatrix. The densification of the PVD coatings is mainlycontrolled by ion current density. All the coatings studiedin the present work were deposited at same substrate bias

ARTICLE IN PRESS

150 300 450 600 750 900 1050

614

Raman Shift (cm-1)

314

25°C

800°C

750°C

700°C

Inte

nsity

(ar

b. u

nits

)

850°C

449 61

3

TiAlN/Si3N4

Fig. 7. Composite Raman spectra of as-deposited TiAlN/Si3N4 nano-

composite coating (11 at% Si) and coatings heat-treated up to 850 1C in

air.

150 300 450 600 750 900 105059

8

TiAlN/Si3N4

(As-deposited)

TiAlN

(Vacuum Annealed - 900ºC)

TiAlN

(As-deposited)

312

Raman Shift (cm-1)60

7

536

323

315

590

611

315

TiAlN/Si3N4

(Vacuum Annealed - 900°C)

Inte

nsity

(ar

b. u

nits

)

Fig. 8. Representative Raman spectra of as-deposited and vacuum

annealed coatings of TiAlN and TiAlN/Si3N4 nanocomposite (11 at%

Si) at 900 1C.

MS

(c) Si3N4

TiAlN

(d)

(b)(a)

E (V vs. Ag/AgCl)

i (A

/cm

2 )

1.0×10-9

1.0×10-7

1.0×10-5

1.0×10-3

1.0×10-1

-0.850 -0.750 -0.650 -0.550 -0.450 -0.350

Fig. 9. Potentiodynamic polarization curves of TiAlN/Si3N4 nanocom-

posite coatings at different silicon concentrations: (a) 8.8 at%, (b) 11 at%,

(c) 15.4 at% and (d) 20 at%. Also shown are the polarization curves of the

mild steel substrate, Si3N4 and TiAlN coatings.

H.C. Barshilia et al. / Vacuum 81 (2006) 479–488486

(�100V) and the same power density to the TiAl target(4.93W/cm2). It is well known that the substrate biascritically affects the ion current. It is also true that theplasma density affects the ion current. However, thisvariation in the ion current with the variation in the powerdensity to the silicon target was not very significant for thedeposition conditions used in the present study. It is,therefore, believed that the contribution to the ion currentis mainly due to the substrate bias. The power density tothe silicon target marginally affected the ion currentdensity. Hence, the densification caused due to ion currentis expected to be almost same for all the samples.

Addition of amorphous Si3N4 in the TiAlN matrixis known to change the surface morphology from apronounced columnar micro-structure to a dense structureat low silicon contents [34,43]. For example, it has been

ARTICLE IN PRESS

Table 3

Potentiodynamic polarization data of TiAlN, Si3N4 and TiAlN/Si3N4

nanocomposite coatings deposited on mild steel substrates in 3.5% NaCl

solution under free-air conditions at room temperature

Sample Ecorr (V) icorr (mA/cm2) Rp (kO cm2)

Mild steel �0.647 4.996 4.553

TiAlN �0.528 3.603 5.260

Si3N4 �0.514 0.186 105.780

TiAlN/Si3N4 (8.8 at% Si) �0.534 2.703 7.862

TiAlN/Si3N4 (11 at% Si) �0.536 2.546 9.022

TiAlN/Si3N4 (15.4 at% Si) �0.540 2.215 11.045

TiAlN/Si3N4 (20 at% Si) �0.529 1.973 13.295

H.C. Barshilia et al. / Vacuum 81 (2006) 479–488 487

reported that the unbalanced magnetron-sputtered TiNcoatings exhibit columns of 60–80 nm in diameter, whichalmost disappear with the addition of 9 at% Si [34].

As mentioned previously, addition of amorphous Si3N4

in the TiAlN matrix decreases the crystallite size of thenanocomposite coatings. Both the densification effect andthe reduced crystallite size observed in the nanocompositecoatings are believed to be responsible for the enhancedcorrosion behavior of the nanocomposite coatings.Furthermore, it is believed that the corrosion behavior ofthe nanocomposite coatings of TiAlN/Si3N4 wouldhave been also affected by the poor electrical conductivityof Si3N4.

A dense micro-structure in the case of TiAlN/Si3N4

nanocomposite coatings prevents the diffusion of thecorrosive medium into the substrate. Also, the additionof Al to transition metal nitrides improves the corrosionresistance as it forms a passive Al2O3 layer on the surfaceof the coating during chemical attack [44]. This preventsthe coating from further attack.

4. Conclusions

The XRD data of the TiAlN/Si3N4 nanocompositecoatings exhibited (1 1 1) and (2 0 0) reflections of cubicTiAlN phase. The diffraction peaks got broadenedsignificantly at very high silicon content, suggesting adecrease in the crystallite size of the coatings. The XPSdata indicated that nitrogen present in the nanocompositecoatings was predominantly in the form of TiN or AlNwith small amount of Si3N4. The TiAlN/Si3N4 nanocom-posite coatings exhibited a maximum hardness and elasticmodulus of 43 and 350GPa, respectively at a siliconcontent of 11 at%. The hardness and elastic modulus of thecoatings decreased with further increase in the siliconcontent. For example, at a silicon concentration ofapproximately 20 at%, the nanocomposite coating showeda hardness of 32GPa and an elastic modulus of 260GPa.The TiAlN/Si3N4 nanocomposite coatings were stable upto 800 1C in air as compared to TiAlN coatings, which gotoxidized at 750 1C. Results of vacuum annealing of thecoatings at 900 1C showed no phase transformation in thecase of nanocomposite coatings. The TiAlN, Si3N4 and

TiAlN/Si3N4 nanocomposite coatings deposited on mildsteel substrates demonstrated good corrosion resistance in3.5% NaCl solution, and the nanocomposite coatingsshowed better corrosion resistance as compared to theTiAlN coating.

Acknowledgments

The authors thank Director, NAL for giving permissionto publish these results. Dr. Anjana Jain is thanked forXRD work. Thanks are also due to Mr. V.K.W. Grips andMs. V.E. Selvi for help in potentiodynamic polarizationmeasurements. We also thank Mr. Siju and Mr. N.T.Manikandanath for their help in AFM, nanoindentationand Raman measurements. This work was supportedby the Council of Scientific and Industrial Research,New Delhi (Grant no. CMM 0022), India.

References

[1] Vaz F, Rebouta L, Andritschky M, da Silva MF, Soares JC. J Mater

Process Technol 1999;92–93:169.

[2] McIntyre D, Greene JE, Hakansson G, Sundgren JE, Munz WD.

J Appl Phys 1990;67:1542.

[3] Otani Y, Hofmann S. Thin Solid Films 1996;287:188.

[4] Wasa K, Hayakawa S. Thin Solid Films 1972;10:367.

[5] Munz WD. J Vac Sci Technol A 1986;4:2717.

[6] Ichimura H, Kawana A. J Mater Res 1993;8:1093.

[7] Shew BY, Huang JL, Lii D. Thin Solid Films 1997;293:212.

[8] Tanaka Y, Gur TM, Kelly M, Hagstrom SB, Ikeda T, Wakihira K,

et al. J Vac Sci Technol A 1992;10:1749.

[9] Musil J, Hruby H. Thin Solid Films 2000;365:104.

[10] Ehrhardt H. Surf Coat Technol 1995;74–75:29.

[11] Musil J. Surf Coat Technol 2000;125:322.

[12] Barshilia HC, Rajam KS. Surf Coat Technol 2004;183:174.

[13] Barshilia HC, Surya Prakash M, Jain A, Rajam KS. Vacuum

2005;77:169.

[14] Barshilia HC, Jain A, Rajam KS. Vacuum 2004;72:241.

[15] Veprek S, Heijman MGJV, Karvankova P, Prochazka J. Thin Solid

Films 2005;476:1.

[16] Musil J, Vlcek J. Surf Coat Technol 2001;142–144:557.

[17] Holubar P, Jilek M, Sima M. Surf Coat Technol 1999;120–

121:184.

[18] Veprek S, Niederhofer A, Moto K, Bolom T, Mannling HD,

Nesladek P, et al. Surf Coat Technol 2000;133–134:152.

[19] Niederhofer A, Nesladek P, Mannling HD, Moto K, Veprek S, Jilek

M. Surf Coat Technol 1999;120–121:173.

[20] Vaz F, Rebouta L, Ramos S, da Silva MF, Soares JC. Surf Coat

Technol 1998;108–109:236.

[21] Rebouta L, Tavares CJ, Aimo R, Wang Z, Pischow K, Alves E, et al.

Surf Coat Technol 2000;133–134:234.

[22] Ma S, Prochazka J, Karvankova P, Ma Q, Niu X, Wang X, et al. Surf

Coat Technol 2005;194:143.

[23] Park IW, Choi SR, Suh JH, Park CG, Kim KH. Thin Solid Films

2004;447–448:443.

[24] Carvalho S, Rebouta L, Cavaleiro A, Rocha LA, Gomes J, Alves E.

Thin Solid Films 2001;398–399:391.

[25] Carvalho S, Vaz F, Rebouta L, Schneider D, Cavaleiro A, Alves E.

Surf Coat Technol 2001;142–144:110.

[26] Vossen JL, Kern W, editors. Thin film processes. New York:

Academic Press; 1978.

[27] Barshilia HC, Rajam KS. Surf Coat Technol 2006; in press.

[28] Sellers J. Surf Coat Technol 1998;98:1245.

ARTICLE IN PRESSH.C. Barshilia et al. / Vacuum 81 (2006) 479–488488

[29] Oliver WC, Pharr GM. J Mater Res 1992;7:1564.

[30] Barshilia HC, Surya Prakash M, Sridhara Rao DV, Rajam KS. Surf

Coat Technol 2005;195:147.

[31] Barshilia HC, Rajam KS. J Mater Res 2004;19:3196.

[32] William Grips VK, Ezhil Selvi V, Barshilia HC, Rajam KS.

Electrochem Acta 2006;51:3461.

[33] Park IW, Kim KH, Suh JH, Park CG, Lee MH. J Korean Phys Soc

2003;42:783.

[34] Jiang N, Shen YG, Mai YW, Chan T, Tung SC. Mater Sci Eng B

2004;106:163.

[35] Cullity BD, editor. Elements of X-ray diffraction. Reading, MA:

Addison-Wesley; 1978.

[36] Perry AJ, Jagner M. Thin Solid Films 1989;171:197.

[37] Carvalho S, Ribeiro E, Rebouta L, Tavares C, Mendonca JP,

Monteiro AC, et al. Surf Coat Technol 2004;177–178:459.

[38] Carvalho S, Ribeiro E, Rebouta L, Pacaud J, Goudeau P, Renault

PO, et al. Surf Coat Technol 2003;172:109.

[39] Bertoti I. Surf Coat Technol 2002;151–152:194.

[40] Briggs D, Seah MP, editors. Practical surface analysis, vol. 1.

New York: Wiley; 1990 p. 602.

[41] Hirai T, Niihara K, Goto T. J Am Ceram Soc 1980;63:419.

[42] Weast RC, Astle MJ, editors. CRC hand book of chemistry and

physics. Boca Raton, FL: CRC Press; 1982.

[43] Diserens M, Patscheider J, Levy F. Surf Coat Technol 1998;108–109:241.

[44] Cunha L, Andritschky M, Rebouta L, Pischow K. Surf Coat Technol

1999;116–119:1152.