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Materials Science and Engineering A 527 (2010) 5000–5007

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

icrostructure and tribological behavior of spark plasma sintered iron-basedmorphous coatings

shish Singha, Srinivasa R. Bakshib, Arvind Agarwalb, Sandip P. Harimkara,∗

School of Mechanical and Aerospace Engineering, Oklahoma State University, 218 Engineering North, Stillwater, OK 74078, United StatesDepartment of Mechanical and Materials Engineering, Florida International University, Miami, FL 33174, United States

r t i c l e i n f o

rticle history:eceived 19 March 2010eceived in revised form 14 April 2010ccepted 19 April 2010

a b s t r a c t

Spark plasma sintering (SPS) has been used to produce thick (∼400 �m) amorphous coatings ofFe48Cr15Mo14Y2C15B6 alloy composition on aluminum substrates. The coatings were fabricated usinguniaxial pressure of 50 MPa over a range of temperatures (550–590 ◦C) below crystallization temper-ature (∼631 ◦C) of this glassy alloy. Under the investigated SPS processing parameters, the infiltrationof the aluminum substrate material in the overlaid amorphous powder was observed resulting in com-

eywords:ulk amorphous alloysinteringowder metallurgy

posite amorphous coating. Detailed investigations on evolution of phases, microstructure, and interfacecharacteristics in the amorphous coatings are presented. The amorphous coatings exhibited high surfacehardness (∼880–1007 HV) and superior wear resistance (∼75–80% decrease in weight loss) comparedto substrate material. The wear mechanisms were dominated by ploughing of soft aluminum phase in

dislo

-ray diffractionlectron microscopyear

the interparticle regions,regions of the coatings.

. Introduction

Amorphous materials or bulk metallic glasses represent aew class of advanced materials exhibiting attractive combina-ions of properties such high strength/hardness and excellentear/corrosion resistance [1–4]. These outstanding properties arerimarily due to disordered atomic arrangement in the amorphousaterials resulting in absence of grain boundaries and defects in theicrostructure. While the non-equilibrium nature of amorphousaterials offers outstanding properties, it also presents significant

hallenges in the processing of such materials. Even though theapid solidification (casting) methods for processing amorphouslloys are well established, the need for simultaneous mold fillingnd rapid cooling rate limits the range of geometries that can beormed [5,6]. These processing difficulties in combination with lowensile ductility and toughness are likely to limit the applicationsf amorphous materials as bulk structural materials. However, themorphous materials can be good candidates for wear/corrosionesistant coatings on the crystalline substrates.

Significant efforts have been made in the past to fabricate amor-

hous coatings on crystalline substrates using various processesuch as laser surface cladding [7–10] and high velocity oxy-fueloating/thermal spraying [11,12]. Yue et al. reported laser claddingf Zr-based amorphous coating on magnesium alloy for structural

∗ Corresponding author. Tel.: +1 405 744 5900; fax: +1 405 744 7873.E-mail address: sandip.harimkar@okstate.edu (S.P. Harimkar).

921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2010.04.066

dging of amorphous particles, and microcutting (abrasion) of amorphous

© 2010 Elsevier B.V. All rights reserved.

applications [13]. The laser cladding resulted in the formation ofnanocrystalline phases in the amorphous matrix of the coating.Wong and Liang also observed the formation of crystalline phasesin Ni–Cr–B–Si coating laser clad on Al–Si alloy [14]. Basu et al.also attempted laser cladding of the Fe-based amorphous coat-ings on steel substrates and observed formation of intermetallicphases in the coating [15]. While laser surface cladding providesvery high cooling rates (up to 106 K/s), most of these approacheshave had limited success in retaining fully amorphous structurein the coatings. This can be primarily attributed to the dilution ofthe melt pool from the underlying substrate partially melted dur-ing laser processing. The melt dilution significantly influences theglass forming ability making it difficult to solidify into amorphousstructure. Furthermore, the solid substrate provides catalytic nucle-ation sites for nucleation and growth of stable crystalline phasesduring constrained solidification. However, the high cooling ratesassociated with laser surface cladding resulted in highly refinedmicrostructure in the coatings with improved hardness and wearproperties [15]. Limited efforts have been made to eliminate thedilution effects by using the substrate of glass forming compo-sition such that subsequent laser surface melting/resolidificationcauses surface amorphization [16]. However, this limits the rangeof substrates that can be used for laser surface cladding. In gen-

eral, the success in retaining fully amorphous structure in coatingsdeposited using laser processing and thermal spraying has beenlimited.

Recently, spark plasma sintering (SPS) has attracted significantinterests for the fabrication of bulk nanostructured and amorphous

d Engineering A 527 (2010) 5000–5007 5001

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radiation at 45 kV and 40 mA. The diffraction angle was variedbetween 30◦ and 70◦ 2� at a step increment of 0.02◦ 2� with acount time of 1 s. A microhardness tester (BuehlerTM) was usedfor measuring hardness by performing indentations at a load of

A. Singh et al. / Materials Science an

aterials which are otherwise difficult to process through otheronventional sintering technique like hot pressing (HP) and hotsostatic pressing (HIP) [17]. By combining the effects of uniaxialressure and pulsed direct current, the process offers enormousossibilities of sintering these materials at significantly lower tem-eratures and shorter sintering times (less than an hour) comparedo conventional hot sintering. While most of these investiga-ions dealt with spark plasma sintering of bulk nanostructured

aterials, very limited efforts have been directed towards fabri-ating nanostructured or amorphous coatings [18,19]. Recently,quab et al. reported fabrication of oxidation resistant aluminizedCrAlY (where M = Co, Ni, or Co/Ni) coatings on nickel super-

lloy substrate using spark plasma sintering [20]. The coatingsxhibited porosity-free homogeneous microstructure with gooddherence and uniform interdiffusion layer between the substratend the coating. Nishimoto and Akamatsu also reported fabrica-ion of oxidation resistant Fe3Al coating on austenitic stainlessteel using spark plasma sintering [21]. Spark plasma sinter-ng has also been used for post-spray processing of thermallyprayed coatings [22,23]. Significant improvements in the den-ities of thermally sprayed ZrO2–MgO and WC-Co coatings haveeen reported using SPS as post-spray treatment. While researchersre extending the capabilities of the SPS process for the fabrica-ion of novel compositions and microstructures, major efforts areeing undertaken by the machine manufacturers to increase theize of the sintered samples up to 350 mm in diameter [24]. Withhese developments, spark plasma sintering is expected to playn important role in the surface engineering of materials. In theresent investigation, we explored the possibility of fabricatingear resistant amorphous coatings on aluminum alloy substratesing spark plasma sintering. Detailed analysis of the developmentf microstructure and the improvement in mechanical proper-ies (hardness and wear resistance) for the amorphous coatings isresented.

. Experimental procedure

In the present investigation, we used starting amorphous pow-er of nominal composition Fe48Cr15Mo14Y2C15B6 for fabrication ofmorphous coatings using spark plasma sintering. The amorphousowder was prepared by melting mixture of high purity (greaterhan 99.9 wt.% purity) elemental powders (Fe, Cr, Mo, Y, B, and Cith nominal glass forming composition) followed by high pres-

ure gas atomization. Differential scanning calorimetry (DSC) withconstant heating rate of 20 ◦C/min was used to determine the

lass transition and crystallization temperature of the as-receivedmorphous alloy powder.

The substrates used in the present study were aluminum alloyiscs of 20 mm diameter and 5 mm thickness. The substrate sur-aces were well polished using 400 and 600 mesh abrasive papersefore coating. Each substrate disc was then placed inside theraphite die and measured quantity of amorphous powder wasniformly loaded on the substrate surface to form ∼400 �m thickoating. A commercial spark plasma sintering equipment (Thermalechnology, Inc.; Model 10-3) was used for forming amorphousoating on the aluminum alloy substrate discs. The schematic of thePS tooling arrangement used for fabricating amorphous coatingsn aluminum alloy substrate is presented in Fig. 1. The SPS exper-ments were carried out under high vacuum under a pressure of0 MPa at three different temperatures (560, 575, and 590 ◦C) well

elow the crystallization temperature of the given Fe-based amor-hous alloy. A typical processing cycle consisted of three steps:apid heating cycle with a rate of 100 ◦C/min, holding for 15 mint the processing temperature, and rapid cooling using nitrogenurging (cooling rate ∼150 ◦C/min).

Fig. 1. Schematic of the spark plasma sintering setup for fabricating amorphouscoatings on aluminum alloys substrate.

After SPS processing, the surfaces and cross-sections of theamorphous coated aluminum substrates were prepared using con-ventional metallographic techniques for further microstructuralanalysis. The X-ray diffraction (XRD) analysis of the starting amor-phous powder and coated surfaces was carried out using PhilipsNorelco X-ray diffractometer operating with Cu K� (� = 1.54178 Å)

Fig. 2. (a) DSC scan and (b) XRD pattern from the starting amorphous alloy powderof composition Fe48Cr15Mo14Y2C15B6 used for fabrication of amorphous coatingsusing SPS. The vertical arrows in Fig. 2(a) indicate the characteristic transition tem-peratures of amorphous alloy.

5002 A. Singh et al. / Materials Science and Engineering A 527 (2010) 5000–5007

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ig. 3. SEM micrographs showing: (a) typical low magnification cross-section of spnterfacial regions of the amorphous coatings prepared at 560, 575, and 590 ◦C, resp

.94 N and holding time of 10 s. The microhardness was measuredn the coating surface and also along the cross-section. Around 10icrohardness readings were taken at each location and an average

alue is reported. The wear tests were performed on the substratend amorphous coated samples using a ball-on-disc tribometerNanoveaTM, Irvin, CA) at a load of 4 N and 136.3 rpm disc rotation. Amm diameter aluminum oxide (Al2O3) ball was used as a counterody to create a wear track of 6 mm diameter on the sample sur-ace. The weight loss was recorded as a function of linear sliding

istance. The sample surfaces before and after wear were ana-

yzed using a scanning electron microscope (SEM) equipped withnergy dispersive spectroscopy (EDS) detector. Both topographicnd back scattered images were used for analysis. The roughnessf the samples before and after wear test was measured using sur-

Fig. 4. XRD patterns from spark plasma sintered amorphous coatings fabricate

sma sintered amorphous coating prepared at 590 ◦C, and (b–d) high magnificationly.

face roughness tester (Model TR200, Micro PhotonicsTM, Allentown,PA).

3. Results and discussion

3.1. Microstructure

Fig. 2 presents the DSC scan and XRD pattern of the gas atomizedFe48Cr15Mo14Y2C15B6 powder. The powder exhibited a distinct

glass transition temperature (Tg) at 575 ◦C followed by doubleexothermic crystallization (Tx1 and Tx2 ) peaks above 600 ◦C (Fig. 2a).The XRD pattern of the starting powder exhibited a characteristicbroad peak with diffused intensity (Fig. 2b). These results confirmedthat the starting alloy powder was glassy (amorphous) in nature

d with processing temperatures of (a) 560 ◦C, (b) 575 ◦C, and (c) 590 ◦C.

A. Singh et al. / Materials Science and Engineering A 527 (2010) 5000–5007 5003

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Fig. 5. EDS elemental map showing distribution of aluminu

nd hence suitable for the proposed objective of developing amor-hous coating on aluminum substrate by spark plasma sinteringSPS).

The typical microstructures of the cross-section of the amor-hous coatings prepared by spark plasma sintering are presented

Fig. 6. EDS elemental map showing distribution of aluminum and ir

iron in surface of amorphous coatings prepared at 560 ◦C.

in Fig. 3. The figure indicates that the coatings prepared at vari-ous processing temperatures (560, 575, and 590 ◦C) are fully denseand exhibit very good bonding with the underlying aluminum sub-strates, as indicated by the absence of porosities at the interface.The average thickness of the coatings is around 400 �m. In all the

on in cross-section of amorphous coatings prepared at 560 ◦C.

5004 A. Singh et al. / Materials Science and Engineering A 527 (2010) 5000–5007

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ig. 7. EDS line scans and corresponding elemental distributions across the interf90 ◦C.

ases, the substrate aluminum alloy infiltrated into overlaid amor-hous powder during SPS processing resulting in the formation ofomposite coating. The image analysis of the cross-sections of theoatings indicated that almost 25% of the particles in the coatingrocessed at 575 ◦C are bigger than 20 �m (∼20–50 �m range). Weelieve that the larger percentage of bigger particles (as large as0 �m) in the coating processed at 575 ◦C is due to handling androcessing of powder prior to sintering. The X-ray diffraction (XRD)atterns from the surfaces of the coatings are presented in Fig. 4.he XRD patterns shows broad peak corresponding to Fe-basedmorphous material and superimposed crystalline peaks corre-ponding to infiltrated aluminum. The crystalline peaks in the XRDattern of coatings prepared at 590 ◦C are strongest in intensity

ndicating that aluminum is a major phase at the surface of coat-ngs. It seems that the substrate aluminum infiltrated across the fullhickness of the amorphous powder layer and reached the surfacef the coatings indicating highest degree of aluminum infiltrationt this temperature. At lower temperatures (560 and 575 ◦C), thentensity of the crystalline peaks is relatively smaller indicating rel-

tively lesser degree of infiltration of substrate aluminum at theurface of the coatings. Furthermore, the aluminum peaks in theRD pattern for the coatings processed at 575 ◦C are not as sharps they are in the XRD pattern for the coating prepared at 560 ◦Clower temperature). The coatings prepared at 575 ◦C showed rel-

or spark plasma sintered amorphous coatings processed at (a–c) 575 ◦C and (d–f)

atively bigger particles (10–50 �m) compared to coating preparedat 560 ◦C. These bigger particles seem to retard the aluminum infil-tration effects resulting in smaller relative proportion of aluminumat the surface of the coatings. The absence of sharp peak in thecoating synthesized at 575 ◦C seems to be due to this smaller rel-ative proportion of aluminum at the surface of the coatings. Notethat the SPS processing of all coatings was carried out for sametime (15 min). Figs. 5 and 6 present the distribution of iron (majorelement in the coating) and aluminum (major element in the sub-strate) at the surface and cross-section of the coatings prepared at560 ◦C. The figures clearly indicate the preferential distribution ofaluminum between amorphous particles both at the surface andacross the cross-section of the coatings.

A careful look at the interfaces of the coatings showedinteresting features (Fig. 3). The coatings prepared at lower tem-peratures (560 and 575 ◦C) shows distinct interdiffusion layer about25 �m thick along the interface. The smaller amorphous particlesdistributed within the interdiffusion layer can be easily seen. Sur-prisingly, the coatings prepared at higher temperature (590 ◦C)

does not show distinct interdiffusion layer at the interface. How-ever, individual large particles near the interface show very thin(∼1–2 �m) reaction layer on the particle surface as indicated bylighter contrast (Fig. 7). To analyze the composition of this reactionlayer, significant efforts were made to map the elemental distri-

A. Singh et al. / Materials Science and Engineering A 527 (2010) 5000–5007 5005

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ig. 8. Microhardness profiles along the thicknesses of spark plasma sintered amorhows the SEM micrographs of the microhardness indentations in the interfacial re

ution across the cross-sections of the coatings. The elementalistribution along the line across the coating/substrate interface

s also presented in Fig. 7. The figure clearly indicates that thenterdiffusion layer at the interface in coatings prepared at low tem-erature (575 ◦C) consists of aluminum element. It seems that themorphous particles at the interface partially reacted with the alu-inum substrate forming the continuous interdiffusion layer. The

gure also indicates sharp change in the elemental aluminum con-ent across the interface indicating absence of interdiffusion layer inoatings prepared at 590 ◦C. The absence of distinct interdiffusionayer at the interface in coatings prepared at higher temperature590 ◦C) is contrary to the common intuition that reaction ratencreases with temperature. While the interdiffusion at the inter-

◦

ace is expected to be higher at higher temperature (590 C), thenhanced plastic flow of aluminum substrate (due to higher tem-erature) under applied pressure is expected to cause continuous

nfiltration of aluminum into the overlaid amorphous powder. Thisynamics of the plastic flow of aluminum at the amorphous pow-

ig. 9. Variation of coefficient of friction with sliding distance during wear testing ofluminum substrate and spark plasma sintered amorphous coatings prepared withrocessing temperatures of 560, 575, and 590 ◦C.

coatings prepared with processing temperatures of 560, 575, and 590 ◦C. The insetof coating prepared at 560 ◦C.

der surface seems to prevent the build-up of interdiffusion layerresulting in sharp coating/substrate interface at higher tempera-ture (590 ◦C). The observation of aluminum as a major phase atthe surface of the coatings prepared at 590 ◦C is indicative of thisenhanced infiltration effects (as indicated by XRD analysis).

3.2. Microhardness

Fig. 8 presents the microhardness distribution along the depthof amorphous coatings prepared at various temperatures. The aver-age hardness of the coatings prepared at 560, 575, and 590 ◦Cwere 990, 1007, and 880 HV, respectively. The higher surface hard-ness (∼1000 HV) of the coatings prepared at lower temperatures

◦

(560 and 575 C) seems to be primarily due to presence of largerfraction of amorphous phase at the surface. The relatively lowersurface hardness (880 HV) of the coatings prepared at higher tem-perature (590 ◦C) is primarily due to presence of larger fractionof soft aluminum phase in the coating due to these enhanced

Fig. 10. Variation of cumulative weight loss with sliding distance during wear test-ing of aluminum substrate and spark plasma sintered amorphous coatings preparedwith processing temperatures of 560, 575, and 590 ◦C.

5006 A. Singh et al. / Materials Science and Engineering A 527 (2010) 5000–5007

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ig. 11. Typical SEM micrograph and corresponding EDS map showing elementalnclosed by dark lines (marked A) and dotted lines (marked B) indicate amorphous

nfiltration effects. The high microhardness (800–1000 HV) wasetained along the 400 �m thickness of amorphous coatings fol-owed by steep decrease at the interface region towards substrate.

his trend is clearly indicated by the variation of indentation sizecross the interface (inset of Fig. 8). In our previous investigation,e observed that the microhardness of SPS sintered fully amor-hous/partially crystallized amorphous alloys was in the range of200–1350 HV [19]. The relatively lower hardness (880–1007 HV)

ig. 12. Typical secondary and backscattered SEM images from worn surfaces of amorphon the backscattered SEM images correspond to aluminum-rich areas on the worn surface

bution of aluminum on the worn surfaces of the amorphous coating. The regionsnd aluminum-rich areas on the worn surface, respectively.

of the amorphous coatings in the present investigation is primarilydue to infiltration of aluminum substrate material in the over-laid amorphous powder bed. The coatings thus exhibit composite

amorphous-crystalline microstructure, the crystalline phase beingthe infiltrated aluminum substrate material. However, no devit-rification of the amorphous powder was observed. The surfacehardness of the amorphous coatings is almost 10 times the hardnessof the substrate aluminum material.

us coatings showing (a) early and (b) later stages of wear process. The dark regionss.

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.3. Friction and wear

Fig. 9 presents the variation of friction coefficient as a functionf sliding distance for the substrate and the amorphous coatingsrepared at various temperatures. The aluminum alloy substratexhibited uniform friction coefficient in the range of 0.45–0.55 indi-ating stable sliding during wear test. However, amorphous coatingxhibited distinct regimes of wear process characterized by lowriction coefficient in the initial stage and high friction coefficient inhe later stages separated by a transition region. The friction coef-cient corresponding to 20 m of sliding distance is in the rangef 0.1–0.2. The lower values of friction coefficient in the initialtages of wear test can be attributed to low surface roughnessf the coatings. The initial measured roughness of all the coat-ngs was in the range of 0.065–0.075 �m. The post-wear surfaceoughness for coatings processed at 560, 575, and 590 ◦C were 13.9,1.0, and 16.3 �m, respectively. Except for the coating preparedt 575 ◦C, the friction coefficient quickly increases and stabilizesetween 0.3 and 0.5 at sliding distance longer than 50 m. Similaristinct regimes of wear processes have been observed for variousr-based amorphous alloys (as-cast, deformed, and creep-tested).iang et al. reported that these regimes correspond to running-in,ransition, and steady state [25]. The steady state friction coeffi-ient of the as-cast Zr-based amorphous alloy reported by Jiang etl. was around 0.8. Abrasive wear is often considered as a domi-ating micro-mechanism of wear in amorphous alloys. The lowerteady state friction coefficients (0.3–0.5) reported in the presentnvestigation seems to be due to presence of soft aluminum phasen the amorphous coatings (due to substrate infiltration). This softluminum phase between the amorphous particles wears out fasteresulting in softer wear debris that provides lubricating effects.

The wear response of the amorphous coatings in ball-on-disconfiguration was also studied by monitoring weight loss with sid-ng distance during test (Fig. 10). The weight loss was highest at allliding distances for the aluminum substrate compared to amor-hous coatings prepared at all temperatures (560, 575, 590 ◦C).he amorphous coatings exhibited superior wear resistance com-ared to substrate aluminum alloy as indicated by significantly

esser weight loss for amorphous coatings. The decrease in weightoss of aluminum substrate with amorphous coatings is more than5–80%. The coating prepared at 575 ◦C exhibited highest wearesistance. This coating also exhibited highest surface hardness1007 HV). Considering the significant extent of standard deviationlength of error bars) of data points in Fig. 10, there is not significantffect of SPS temperature on the wear weight loss for amorphousoatings.

The worn surfaces of the amorphous coatings were examined byEM. Fig. 11 presents a typical micrograph of the worn surface of themorphous coating. Two regions can be easily identified in the wornurfaces: amorphous particle regions (marked “A”) and aluminum-ich regions (marked “B”). It seems that the wear process initiatesy ploughing the soft aluminum phase present between the amor-hous particles resulting in formation of micropits (Fig. 12a). Theackscattered SEM image clearly indicates dark regions corre-

ponding to aluminum-rich phase due to compositional contrastetween the regions of micropits and the surrounding unwornurface. The continuous wear eventually causes dislodging or pull-ut of some amorphous particles. The underlying aluminum-richegions, after the amorphous particles are dislodged, appear as

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ineering A 527 (2010) 5000–5007 5007

dark regions (Fig. 12b). Furthermore, some amorphous particlesremain protruded out on the worn surface (no dislodging) and showscratches and microcutting marks indicative of abrasive wear. Thus,the wear mechanisms of amorphous coatings prepared by SPS aredominated by particle pull-outs and abrasive wear of amorphousparticles/regions.

4. Conclusion

Spark plasma sintering can be effectively used to fabricatethick (∼400 �m) Fe48Cr15Mo14Y2C15B6 amorphous coatings at pro-cessing temperatures below crystallization temperature of theamorphous alloy. The major effect associated with the fabricationof amorphous coatings using SPS is related to the infiltration ofsubstrate material in the overlaid amorphous powder. The amor-phous coatings exhibited high surface hardness (∼880–1007 HV)and superior wear resistance (∼75–80% decrease in weight loss)compared to substrate material. The wear mechanisms were dom-inated by ploughing of soft aluminum phase in the interparticleregions, dislodging of amorphous particles, and microcutting (abra-sion) of amorphous regions of the coatings. Major challenges forthe widespread implementation of SPS process for fabrication ofamorphous coatings are related to: (i) designing of experimentalstrategies to prevent the infiltration effects (such that fully amor-phous coatings can be formed) and (ii) developing SPS capabilitiesto scale-up the part sizes.

References

[1] A. Inoue, B. Shen, N. Nishiyama, in: M. Miller, P. Liaw (Eds.), Bulk MetallicGlasses, Springer, 2008, pp. 1–25.

[2] A. Inoue, B. Shen, A. Takeuchi, Mater. Trans. 47 (2008) 1275–1285.[3] A. Inoue, Mater. Sci. Eng. A 304–306 (2001) 1–10.[4] W.L. Johnson, MRS Bull. 24 (1999) 42–56.[5] J. Schroers, G. Kumar, T.M. Hodges, S. Chan, T.R. Kyriakides, JOM 61 (2009)

21–29.[6] J. Schroers, N. Paton, Adv. Mater. Process. 164 (2006) 61–63.[7] Q. Zhu, S. Qu, X. Wang, Z. Zou, Appl. Surf. Sci. 253 (2007) 7060–7064.[8] I. Manna, J.D. Majumdar, B.R. Chandra, S. Nayak, N.B. Dahotre, Surf. Coat. Tech-

nol. 201 (2006) 434–440.[9] X. Wu, Y. Hong, Surf. Coat. Technol. 141 (2001) 141–144.10] X. Wu, B. Xu, Y. Hong, Mater. Lett. 56 (2002) 838–841.11] J.C. Farmer, J.J. Haslam, S.D. Day, T. Lian, C.K. Saw, P.D. Hailey, J.-S. Choi, R.B.

Rebak, N. Yang, J.H. Payer, J.H. Perepezko, K. Hildal, E.J. Lavernia, L. Ajdelsztajn,D.J. Branagan, E.J. Buffa, L.F. Aprigliano, J. Mater. Res. 22 (2007) 2297–2311.

12] T. Parker, F.E. Spada, A.E. Berkowitz, K.S. Vecchio, E.J. Lavernia, R. Rodriguez,Mater. Lett. 48 (2001) 184–187.

13] T.M. Yue, Y.P. Su, H.O. Yang, Mater. Lett. 61 (2007) 209–212.14] T.T. Wong, G.Y. Liang, Mater. Charact. 38 (1997) 85–89.15] A. Basu, A.N. Samant, S.P. Harimkar, J.D. Majumdar, I. Manna, N.B. Dahotre, Surf.

Coat. Technol. 202 (2008) 2623–2631.16] D.T.A. Matthews, V. Ocelik, J.T.M. De Hosson, Mater. Sci. Eng. A 471 (2007)

155.17] Z.A. Munir, U. Anselmi-Tamburini, M. Ohyanagi, J. Mater. Sci. 41 (2006)

763–777.18] D.M. Hulbert, D. Jiang, U. Anselmi-Tamburini, C. Unuvar, A.K. Mukherjee, Mater.

Sci. Eng. A 1–2 (2008) 251–255.19] S.P. Harimkar, S.R. Paital, A. Singh, R. Aalund, N.B. Dahotre, J. Non-Cryst. Solids

355 (2009) 2179–2182.20] D. Oquab, C. Estournes, D. Monceau, Adv. Eng. Mater. 9 (2007) 413–417.21] A. Nishimoto, K. Akamatsu, Plasma Process. Polym. 6 (2009) S941–S943.22] B. Prawara, H. Yara, Y. Miyagi, T. Fukushima, Surf. Coat. Technol. 162 (2003)

234–241.23] L.G. Yu, K.A. Khor, H. Li, K.C. Pay, T.H. Yip, P. Cheang, Surf. Coat. Technol. 182

(2004) 308–317.24] R. Aalund, Thermal Technology LLC, private communication.25] F. Jiang, J. Qu, G. Fan, W. Jiang, D. Qiao, M.W. Freels, P.K. Liaw, H. Choo, Adv. Eng.

Mater. 11 (2009) 925–931.