Materials Science and Engineering A 422 (2006) 298–308

Mechanical properties and interface toughness of FeCothin films on Ti–6Al–4V

K.S. Chan a,∗, H. Ji b, X. Wang b, S.J. Hudak Jr. a, B.R. Lanning a

a Southwest Research Institute, San Antonio, TX 78238, United Statesb University of Texas at San Antonio, Mechanical Engineering and Biomechanics, San Antonio, TX 78249, United States

Received 27 January 2006; accepted 14 February 2006

Abstract

Magnetostrictive materials such as FeCo are attractive for applications as embedded thin-film sensors for detecting strains or cracks in structuralcomponents. For these applications, FeCo thin films must possess good adhesion and adequate mechanical strength such that the embedded sensorswould have adequate durability, and not degrade the structural integrity of the underlying components. In this paper, we investigate the effects ofthin-film thickness and the surface condition of the substrate on the mechanical properties and interface toughness of FeCo thin films deposited toa Ti–6Al–4V substrate. The Young’s modulus, hardness, and yield strength of FeCo thin films were measured by nanoindentation. Furthermore,ttcao©

K

1

ts[tsiia

t[ibn

0d

he interface toughness of FeCo/Ti–6Al–4V was determined by cross-section nanoindentation on the edge and by brale indentation on the face ofhe thin films. The results indicate that the interface toughness exhibit spatial variations as well as dependence on layer thickness and the surfaceondition of the substrate. The experimental results are combined with a theoretical analysis to predict the debonding strength of the thin filmsnd to assess the propensity of thin-film cracking versus interface debonding. The theoretical predictions are evaluated against experimental dataf the debonding strength of FeCo thin films and notch fatigue strength of Ti–6Al–4V.

2006 Elsevier B.V. All rights reserved.

eywords: Interface toughness; Thin film; Brale indentation; Cross-sectional nanoindentation; FeCo; Ti alloys; Durability

. Introduction

Materials damage prognosis is a revolutionary concept inhe management of the health of aircraft structures, propul-ion systems, as well as the entire system of military vehicles1]. This revolutionary concept requires the integration of real-ime in situ interrogation of the material damage states duringervice, physics-based models of the life-limiting processes inndividual components and subsystems, and automated reason-ng to make a robust prediction of the future capability of anircraft.

One of the key requirements of this revolutionary approacho vehicle health management is the use of on-board sensors2,3]. One advantage of real-time continuous health monitor-ng of critical engine components for fatigue damage or defectsy on-board sensors is that significant improvement in compo-ent reliability can be achieved with relatively low interrogation

∗ Corresponding author. Tel.: +1 210 522 2053; fax: +1 210 522 6965.E-mail address: kchan@swri.edu (K.S. Chan).

sensitivities [3]. Recently, magnetostrictive thin films have beenshown to exhibit the required performance as embedded sen-sors for direct detection and monitoring of fatigue cracks atambient temperatures [3]. Enhanced magnetostrictive sensorperformance appears to be achievable through optimization ofthe thin-film architecture and composition [3]. On the otherhand, very little is known about the durability of magnetostric-tive thin films under monotonic and cyclic loading conditions atambient or elevated temperatures.

The pertinent material properties for achieving durabilityin thin films include adhesion, interface toughness and fatigueresistance. Studies of metallic thin films have revealed that thefatigue life of Cu and Ag thin films depend on film thickness[4,5], with thin films being more fatigue-resistant than thickerfilms [5]. The stress amplitude at a constant fatigue life increaseswith decreasing film thickness in Cu and Ag thins film, as wellas in TiCuTi tri-layer thin films [4,5]. A theoretical model pre-dicted that the presence of a surface oxide layer on a metalsubstrate can delay fatigue crack initiation in the substrate [6,7].Experimental data on multilayered Cu–Ni coatings showed thata nanoscale layered structure retarded fatigue crack initiation

921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.02.035

K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308 299

and improved fatigue life [8,9]. The magnetostrictive materialof interest to this study is FeCo thin films, whose magnetostric-tive properties are well known [10]. The fatigue life responses ofbulk Fe–50Co [11] and Fe–49Co–2V [12–14] were investigatedby several investigators. There are, however, no studies on themechanical properties, interface adhesion and fatigue resistanceof FeCo thin films on a ductile substrate.

The objective of this article is to report the results of an inves-tigation on the mechanical properties and interface adhesion ofFeCo thin films on Ti–6Al–4V. This investigation focused onthe effects of film thickness and surface condition on mechani-cal properties such as Young’s modulus, hardness, yield strength,and interface toughness of FeCo thin films vapor deposited ontoTi–6Al–4V using magnetron sputtering. First, theoretical anal-yses are presented to provide a background on determining theinterface toughness of thin films using (1) a brale indentationtechnique in a Rockwell C harness tester, and (2) a cross-sectional nanoindentation (CSN) technique. The experimentalprocedure describes deposition processing of FeCo thin filmson Ti–6Al–4V, nanoindentation and CSN tests, and the braleindentation tests. The interface toughness data obtained by thetwo indentation techniques are then presented and compared toassess the effects of layer thickness and surface condition onthe adhesion of FeCo thin films on Ti–6Al–4V. In addition, theexperimental results are combined with a theoretical analysis topredict the debonding strength of the thin films and to assesstimc

2

c[omia[amStdnucataptsfp

the substrate. The interface toughness was then computed viaa theoretical analysis [19] and the measured load–displacementcurve.

2.1. Mechanics of brale indentation

The mechanics of interface debonding of a thin film ofthickness t bonded on a ductile substract by brale indentationhas been analyzed by Drory and Hutchinson [17]. For filmsthat are very thin compared to the indentation depth into thesubstrate, the film has little effect on the deformation of thesubstrate by the indenter. The essential information needed tocarry out a thin-film debonding analysis is the radial surfacedisplacement of an elastic–plastic half space loaded by a coni-cal indenter. In a recent publication [17], Drory and Hutchinsoncomputed the surface displacements of a half-space due to abrale indenter, which is a conical indenter with a 60◦ semi-angle. The radial surface displacements were used to gener-ate the strains induced on the thin film due to indention. Theresults were then utilized to compute the energy release rateavailable to drive an axisymmetric interface crack expandingoutward from the indentation. In this paper, we utilized theresults of Drory and Hutchinson [17] for the material character-istics that are relevant to brale indentation of FeCo thin film onTi–6Al–4V.

di

l

wifi((σ

tTwoo

σ

war

G

w

G

w

he propensity of thin-film cracking versus interface debond-ng. The durability and functionality of FeCo as thin-film sensor

aterial is examined under monotonic and fatigue loading in aompanion paper [16].

. Theoretical background

Various experimental techniques have been developed forharactering the adhesion of thin films on ductile substrates15]. Among the various methods, indentation appears to bene of the simplest and more popular techniques for deter-ining the adhesion of thin films on a ductile substrate. Both

ndentation techniques via a conventional hardness tester andnanoindenter have been developed. Drory and Hutchinson

17] used a diamond brale indenter (120◦ diamond cone withslightly rounded point) in a Rockwell hardness tester to deter-ine the interface toughness of a diamond-coated titanium alloy.ubsequently, Vlassak et al. [18], proposed a wedge indenta-

ion technique for measuring the adhesion of brittle films touctile substrate. In both brale and wedge indentation tech-iques, an indenter is driven through the brittle coating into thenderlying ductile substrate. Plastic deformation of the substrateauses the thin film to debond from the substrate. Theoreticalnalyses [17,18] were then utilized to compute the interfaceoughness on the basis of the dimensions of the indentationnd the debonded areas. More recently, Sanchez et al. [19]roposed a cross-sectional nanoindentation (CSN) techniquehat involved making a nanoindentation normal to the cross-ection of a thin-film substrate at a distance close to the inter-ace of interest. The nanoindentation produced a wedge thatushed on the interface, causing the interface to debond from

The analysis of Drory and Hutchinson [17] showed that theisplacement field (uI) for a brale indentation can be describedn terms of a third order polynomial given by [17]:

n

(uI

d

)= b0 + b1

( r

d

)+ b2

( r

d

)2 + b3

( r

d

)3(1)

here r is the radial distance measured from the center of thendent, d the diameter of the indent, and bi the polynomial coef-cients whose values, which depend on the ratio of yield stressσy) to elastic modulus (E) and the strain hardening exponentN), have been obtained by Drory and Hutchinson for a range ofy/E and N values [17]. Once the displacement field is known,

he radial (εr) and circumferential (εθ) strains can be obtained.he radial component, σr(rd), of the stress in the film at r = rd,here rd is the distance from the center of the indent to the tipf the debonded interface crack, can be expressed in the termsf the strains by [17]:

r(rd) = E

1 − ν2 [εr(rd) + νεθ(rd)] (2)

here εr(rd) and εθ(rd) are the radial and circumferential strainst r = rd, respectively, and ν is Poisson’s ratio. The elastic energyelease rate G is given by [17]:

= (1 − ν2)tσ2r

2E(3)

hich can be combined with Eq. (2) to give [17]:

= Et

2(1 − ν2)[εr(rd) + νεθ(rd)]2 (4)

here t is the thickness of the thin film.

300 K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308

2.2. Cross-sectional nanoindentation

Sanchez et al. [19] analyzed the mechanics of cross-sectionindentation by applying a wedge load to an axisymmetric circu-lar plate of radius a with its edge clamped and an inner ring ofradius b and a fixed vertical displacement u0. The vertical dis-placement, u0, is related to the maximum thin film deflection,δd, at debonding by

u0 = δd tan θ (5)

where θ = 65.3◦ based on the geometry of the Berkovich inden-ter. Sanchez et al. [19] obtained the solution for the criticalenergy release rate for an interface crack formed by CSN. Forsmall values of b/a ratios (b/a < 0.2), the wedge load can betreated as a point load and the critical energy release rate isgiven by [19]:

Gic = 8Et3(δd tan θ)2

12(1 − ν2)a4 (6)

where E is the Young’s modulus, ν the Poisson’s ratio and t is thethickness of the thin film. To compute the critical energy releaserate, the load–displacement curve must be recorded during CSNto determine the maximum thin-film deflection, δd, at interfacedebonding. In addition, the half length, a, of the interface crackmust also be measured, together with the elastic properties andt

3

3

Cmpwawc9

wdiwStstttt(aoc

19 mm × 6.4 mm per specimen. Base pressure of the chamberprior to deposition was between 1 × 10−5 and 7 × 10−5 Pa andFeCo was sputtered at an argon pressure between 4 × 10−2 and0.2 Pa. Residual stresses of the films were determined by mea-suring the radius of curvature of a thin film deposited onto a lowmodulus (Kapton) substrate to determine the processing temper-ature. Once the deposition temperature was known, the residualstress in FeCo thin film was computed on the basis of the coeffi-cients of thermal expansions. The film thickness values were 0.9and 4.1 �m. Film thickness was first measured by profilometryand subsequently verified by scanning electron microscopy.

The FeCo thin films deposited on Ti–6Al–4V substrate tookon the topography of the Ti–6–4 substrate after deposition.On rough, mechanically polished surfaces, the thin films con-formed to the underlying features and showed polishing marks.On electro-polished and chemically etched surfaces, the thinfilms conformed to and depicted the � + � microstructure ofTi–6–4. At high magnification, chemically etched FeCo thinfilms showed a fine-grained microstructure, as shown in Fig. 1(a)and (b) for 0.9 and 4.1 �m films, respectively. To confirmthe film microstructure, FeCo thin films were deposited ina Si substrate and subsequently etched. For 2 �m FeCo thinfilm, the microstructure was consisted mostly of 0.1–0.3 �mequiaxed grains, as shown in Fig. 1(c). For 7 �m thick films,the microstructure contained 0.4–0.62 �m equiaxed grains, asshown in Fig. 1(d).

3

ad(yaieopayltla

3

NtasT22T

hickness of the thin film.

. Experimental procedure

.1. Thin-film processing

A Ti–6Al–4V pancake forging from the U.S. Air Force Highycle Fatigue Program [20] was utilized in the program. Thisaterial was chosen because its microstructure and fatigue

roperties were fully characterized [20]. As described else-here [20,21], the test material was solution-treated at 932 ◦C

nd aged at 704 ◦C for 2 h, resulting a bimodal microstructureith ≈60 vol.% primary � grains and ≈40 vol.% lamellar � + �

olonies. The yield strength in the longitudinal direction was30 MPa and the ultimate tensile strength was 978 MPa [21].

Ti–6Al–4V coupon specimens 19 mm in length, 19 mm inidth, and 6.4 mm in thickness were prepared by electro-ischarge machining (EDM), followed by mechanically polish-ng via 600 grit (17 �m) SiC abrasive papers. Selected specimensere electropolished at −40 ◦C in a solution containing 15%O4, 5% butyl alcohol, and 80% methanol by volume, while

he remaining specimens retained the mechanically polishedurfaces. The FeCo films were produced using magnetron sput-ering deposition from a pure cobalt target. The composition ofhe FeCo film was established by placing a pure iron foil over thearget with a predefined surface area such that the final composi-ion of the film, as determined by energy dispersive spectroscopyEDX), was ∼50% Fe and 50% Co (by weight) or Fe–49Co (intomic percent). Films were deposited through shadow masksnto rough or smooth electro-polished surfaces of Ti–6Al–4Voupons that had been ion sputter cleaned. The coated area was

.2. Nanoindentation

Nanoindentation tests of FeCo thin films were carried out onNanoIndenter® XP Nanoindentation system with a Berkovichiamond tip using the MTS® Continuous Stiffness MeasurementCSM) technique for measuring elastic modulus, hardness, andield strength. The surface approach velocity was set to be 5 nm/snd the indentation depth limit was set to be 2000 nm. The load-ng direction was perpendicular to the FeCo film surface. Forach specimen, an array of five indents was made at a distancef 3–4 mm apart. Indentation load was recorded as a function ofenetration depth. The results were analyzed using MTS® datanalysis software to obtain the elastic modulus, hardness, andield strength of FeCo thin films. The data analysis software uti-ized the method developed by Giannakopoulos and Suresh [22]o estimate the yield strength, σy, from the measured indentationoads and penetration depth data. Details of this analysis methodre summarized in the Appendix.

.3. Cross-sectional nanoindentation

The CSN tests were performed using a NanoIndenter® XPanoindentation system with a Berkovich indenter. To identify

he debonding load, several indents were made at a distance ofbout 3 �m from the interface using various indentation loadstarting at 50 mN and ending at 350 mN at 50 mN increments.he load at interface debonding was determined to be around00–300 mN. Subsequent indentation was performed at either00 or 300 mN using a surface approach velocity of 5 nm/s.he loading rate was sufficiently high that creep would not be

K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308 301

Fig. 1. Microstructures of FeCo thin films: (a) 0.9 �m film on Ti–6Al–4V, (b) 4.1 �m film on Ti–6Al–4V, (c) 2 �m film on Si, and (d) 7 �m film on Si.

expected to occur in either the Ti–6Al–4V substrate or FeCothin film at ambient temperature. The indents were placed alongthe edge of the thin film at a distance of 3 �m from the interface.The spacing between indents was about 3 mm. Indentation loadwas measured as a function of displacement into the surfaceduring both loading and unloading. Optical images were takenbefore and after each indentation and the position of each indentwas recorded. After indentation, the specimens were examinedin a scanning electron microscope to determine whether or notinterface debonding had occurred and to measure the size andshape of the debonded regions.

3.4. Brale indentation

Indentation interface toughness tests were performed at ambi-ent temperature using a brale indenter (120◦ diamond cone witha slightly rounded point) in a Rockwell C hardness tester. Thebrale (Rockwell C) indenter was pushed into the FeCo thin filmon the Ti–6Al–4V substrate under a constant applied load of15 kg. For an individual FeCo/Ti–6Al–4V specimen, at least 10indentations were made along the length of the film. After inden-tation, the diameter of the indent, the size of the debonded regionand the cracked region on the thin film were measured by scan-ning electron microscopy (SEM). The sizes of the indented andthe debonded/cracked region, as well as the residual stresses inthe thin film, which were determined separately by curvaturemud

strain energy release rate (G) of FeCo thin films deposited onTi–6Al–4V. The Young’s modulus of FeCo thin film was deter-mined to be 1.7 × 105 MPa by the nanoindentation techniqueand the Poisson’s ratio was taken to be 0.33. The critical elasticenergy release rate, Gi, at the tip of the interface crack was takento be the interface toughness of the thin films.

4. Results

4.1. Nanoindentation results

Typical results of Young’s modulus and hardness data ofFeCo thin films obtained by the nanoindentation technique arepresent in Fig. 2(a) and (b), respectively. As shown in Fig. 2, fiveindents were made for each specimen to determine the Young’smodulus and hardness. The hardness data were then used toobtain the yield strength data via MTS data analysis software.A summary of the mean values and standard deviations of thesemeasurements are presented in Table 1 for different values offilm thickness and various surface conditions.

4.2. Cross-sectional nanoindentation

Fig. 3(a) shows a typical load–displacement curve showingboth the loading and unloading cycle for CSN at 200 mN wherethe interface adjacent to the indent did not debond. For compari-saw

easurements of the thin film deposited on Kapton, were thentilized as input to the interface fracture model, Eqs. (1)–(4),eveloped by Drory and Hutchinson [17] to compute the elastic

on, Fig. 3(b) shows the load–displacement curve for the loadingnd unloading cycle of CSN at a maximum load of 200 mNhere the interface debonded. The load–displacement curves

302 K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308

Table 1Summary of nanoindentation measurements of Young’s modulus, hardness, and yield strength of FeCo thin film on Si and Ti–6Al–4V substrates with various filmthickness and surface conditions

System Film thickness (�m) Substrate surface Young’s modulusof film (GPa)

Hardness offilm (GPa)

Yield strengthof film (MPa)

FeCo/Si 2 – 165.5 (15.9)a 5.5 (1.0)a 1830 (330)a

FeCo/Si 7 – 170.0 (6.2) 5.4 (0.2) 1800 (70)FeCo/Ti–6Al–4V 1 Mechanically polished 187.3 (50.5) 8.8 (2.4) 2930 (800)FeCo/Ti–6Al–4V 4 Mechanically polished 176.0 (53.8) 8.2 (4.0) 2730 (1330)FeCo/Ti–6Al–4V 0.9 EDM (rough) 164.8 (22.2) 10.4 (2.3) 3470 (770)FeCo/Ti–6Al–4V 0.9 Electropolished 167.7 (11.7) 10.7 (1.3) 3570 (430)FeCo/Ti–6Al–4V 4.1 EDM (rough) 206.1 (33.1) 8.7 (21) 2830 (700)FeCo/Ti–6Al–4V 4.1 Electropolished 201.4 (14.7) 8.5 (1.4) 2900 (700)

a Standard deviations are shown in parentheses.

Fig. 2. Mechanical properties of 4.1 �m FeCo thin film on Ti–6Al–4V deter-mined by nanoindentation: (a) Young’s modulus as a function of indenter dis-placement into substrate, and (b) hardness as a function of indenter displacementinto substrate.

Fig. 3. Load–displacement curves for cross-sectional nanoindentation onFeCo/Ti–6Al–4V: (a) without interface debonding, and (b) with interfacedebonding.

K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308 303

shown in Fig. 3(a) and (b) exhibited significant differences inboth the shape and the penetration depth, but no evidence ofcreep during loading or unloading. For an intact interface, boththe loading and unloading portions of the load–displacementcurve are rather smooth, as shown in Fig. 3(a). The penetrationdisplacement into the substrate at 200 mN is about 1490 nm butit decreases to 1160 nm after unloading to zero load. In contrast,the load–displacement curve associated with a debonded inter-face, Fig. 3(b), exhibits several steps or discontinuities duringloading to a maximum load of 200 mN. These steps or disconti-nuities indicate increases in the penetration distance at constantloads, which have been previously identified by Sanchez et al.[19], as the point at which the onset of interface debondingoccurs during CSN. These authors also showed that the totalpenetration depth could be distinguished into two components;one arising from plastic flow while another arising from inter-face debonding. In Fig. 3(b), the maximum penetration distanceat the maximum load is about 2500 nm. During unloading, thepenetration depth decreases drastically when the load is reducedbelow 20 mN, resulting in a penetration distance of 1200 nm atzero load. The increase in the penetration distance,∆, in Fig. 3(b)appears to be the result of interface debonding and its value istaken to be the maximum deflection, δd, at the onset of interfacedebonding. Note that the debond distance ∆ is measured from anew reference point where the unloading curve goes to zero loadin Fig. 3(b). For an intact interface, ∆ is zero and the interfacetE

Fs

Fig. 4. SEM micrograph shows the cross-sectional view of a FeCo thin film onTi–6Al–4V with a CSN indent that causes interface debonding and detachmentfrom the substrate.

from the appropriate load–displacement curve and used to com-pute the interface toughness via Eq. (6) using ∆ = δd, providingthat the elastic properties of the thin film and the length (a) ofthe interface debonds are known.

SEM observation of the debonded areas indicated that CSNcaused not only interface debonding but also fracture and detach-ment of the FeCo thin film from the Ti–6Al–4V substrate. Fig. 4shows the cross-sectional view of an indent that caused debond-

oughness of the thin film cannot be computed on the basis ofq. (6). For debonded interface, the value of ∆ can be measured

ig. 5. Top views of interface debonds film detachment produced by CSN testing: (urface.

a) and (b) 4.1 �m film on a rough surface, (c) and (d) 0.9 �m film on a rough

304 K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308

Fig. 6. CSN interface toughness of 0.9 �m FeCo film on either a rough or pol-ished surface as a function of distance along the edge of the coupon specimen.RS denotes residual stress.

ing and detachment of a patch of FeCo thin film from thesubstrate. The length and depth of the interface debond wasdetermined from SEM images taken normal to the surface of theFeCo thin film. Fig. 5 depicts several debonded regions in theFeCo thin film, which are generally not semi-circular. The lengthof the debond on the edge was taken to be 2a, while the depth ofthe debond was taken to be c. Both a and c were measured fromSEM micrographs of the interface debonds. The half-length anddepth of the interface debond were averaged and used in con-

Fig. 7. CSN interface toughness of 4.1 �m FeCo film on a rough surface as afunction of distance along the edge of the couple specimen.

junction with elastic constants from nanoindentation, maximumdeflection, δd (=∆), from the load–displacement curve, and Eq.(6) to compute the average interface toughness.

The interface toughness values determined by the CSN tech-nique for 0.9 �m FeCo thin film on polished and rough surfacesare compared in Fig. 6 as a function of location along the edge ofthe specimen. The interface toughness values show substantialvariation with location. The mean value of the interface tough-ness is in the range of 1–10 J/m2 for 0.9 �m FeCo films on a

Fro

ig. 8. Brale indentation of FeCo thin film bonded to Ti–6Al–4V: (a) 0.9 �m FeCo fiough surface film with fair amounts of interface debonding, (c) 4.1 �m FeCo on a pn a rough surface with extensive interface debonding.

lm on a polished surface with little interface debonding, (b) 0.9 �m FeCo on aolished surface with extensive interface debonding, and (d) 4.1 �m FeCo film

K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308 305

polished or rough surface. The CSN interface toughness for4.1 mm films on rough surfaces are presented in Fig. 7, whichshow that the interface toughness varies with location and hasa value ranging from 1 to 10 J/m2. The CSN tests on polishedsurfaced did not result in any interface debonding because thecoatings covered part of the edge of the specimen where mea-surements were made.

4.3. Brale indentation

Fig. 8(a) and (b) shows brale indents produced on 0.9 �mFeCo films deposited on polished and rough surfaces, whileFig. 8 (c) and (d) presents those on 4.1 �m thin films. Braleindentation did not always produce interface debonding on0.9 �m FeCo thin film. This is illustrated in Fig. 8(a), whichdepicts a brale indent with little or no visible interface debondingfor the 0.9 �m film on a polished surface. Interface debondingappeared to occur more readily for the 0.9 �m film depositedon a rough surface, as shown in Fig. 8(b). In this case, thedebonded region is relatively small. In contrast, brale indentationon 4.1 mm FeCo films produced substantial interface debond-ing for both polished and rough surfaces, which are depicted inFig. 8(c) and (d), respectively.

Ff4

The maximum length of the interface debond and the diam-eter of the indent were measured on individual indents usingSEM. The results were used in conjunction with Eq. (4) to com-pute the interface toughness of FeCo/Ti–6Al–4V. For limitedcases, such as that shown in Fig. 8(a), where interface debond-ing was not obvious, the maximum length of the cracked regionswas used in conjunction of Eq. (4) to get an estimate of the lowerbound of the interface toughness. The computed interface tough-ness values are presented in Fig. 9(a) and (b) for 0.9 and 4.1 �mFeCo films, respectively. Fig. 9(a) shows that for the 0.9 �mfilm, the interface toughness ranges from 1.4 to 15.8 J/m2 forfilm on a rough surface, compared to an interface toughness inthe range of 6.8–22 J/m2 for film on a polished surface. In bothcases, the interface toughness shows spatial variations but it ismore pronounced for the rough surface. Similarly, the interfacetoughness for 4.1 �m FeCo film shows spatial variation, rangingfrom 1.8 to 6.3 J/m2 for film on a rough surface and from 2.5 to6.7 J/m2 for film on a polished surface. Thus, the surface con-dition appears to affect the interface toughness of 0.9 �m film,but not that of the 4.1 �m thick FeCo film on Ti–6Al–4V.

5. Discussion

5.1. CSN versus brale indentation

Both the CSN and brale indentation techniques are capableoTlrIsttuoiaCnacie

5

F

ig. 9. Interface toughness of FeCo film on a rough or polished surface as aunction of distance along the length of the specimen: (a) 0.9 �m film, and (b).1 �m film. RS denotes residual stress.

Rrot(sai

f determining the interface toughness of FeCo thin films oni–6Al–4V. It is worthwhile to note that the CSN technique is a

ocal measurement technique that interrogates a small interfacialegion on the order of 10–50 �m in length for FeCo thin films.n contrast, the brale indentation technique is a more macro-copic measurement technique since the size of the indent andhe debonded region are on the order of 200–500 �m in diame-er, respectively. The brale indentation technique is also easier tose since special precision equipment is not needed (it requiresnly a Rockwell C hardness tester), a load–displacement curves not required, and the indents are larger and thus easier to locatend analyze compared to those produced by the CSN technique.ompared to the brale indentation technique, the CSN tech-ique tends to give a lower interface toughness value because oflarger interface debond length along the edge of the specimenompared to those computed on the basis of the debond depth. Its unknown whether or not the lower interface toughness at thedge is a true material characteristic or merely an edge effect.

.2. Effects of film thickness and surface condition

In a companion study [16], the effect of layer thickness ofeCo thin films deposited on low-stress ground surfaces (0.8 �mMS) of Ti–6Al–4V was examined by brale indentation. The

esults, shown in Fig. 10, indicate that the interface toughnessf FeCo thin films is essentially independent of the thin-filmhickness. The previous results on low-stress ground surfaces0.8 �m RMS) are compared against those from the presenttudy for rough mechanically polished (600 grit, 17 �m SiC)nd electropolished surfaces in Fig. 10. For 4.1 �m films, thenterface toughness for films on electropolished and rough sur-

306 K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308

Fig. 10. Interface toughness of FeCo films as a function of film thickness forthree surface conditions including rough mechanically polished (600 grit), elec-tropolished, and low-stress ground (0.8 �m RMS) surfaces.

faces (600 grit) are in agreement with those for low-stress groundsurfaces. For 0.9 �m films, the interface toughness appears todepend on the surface condition. The highest interface tough-ness is observed on the thin film deposited on the electropolishedsurface, followed by that on the mechanically polished surface,and the lowest for the low-stress ground surface. For all threesurfaces, the film was deposited after ion cleaning of the deposi-tion surfaces under identical conditions to remove any titaniumoxides that may have formed on the surfaces. The difference inthe topography of the deposition surface may lead to differentdegrees of oxide removal from the surface and lead to varia-tions in the observed interface toughness. Different degrees ofsurface oxide removal may also be responsible for the spatialvariations observed in the interface toughness. Since thin filmstend to adopt the topography of the substrate more closely, theeffect of the surface condition on interface toughness becomemore pronounced as the film thickness decreases.

5.3. Film cracking versus interface debonding

The fracture mechanics of crack deflection versus crack pen-etration has been analyzed by He and Hutchinson [23,24], whopresented the results in terms of a plot of Gi/Gp as a function ofthe Dundurs parameter [25], α = (E1 − E2)/(E1 + E2) [23], whereGi is the interface toughness, or the critical energy release rateortFTtCaabtFw

Fig. 11. Measured values of Gi/Gp for FeCo/Ti–6Al–4V compared against thetheoretical boundary of He et al. [24], for transition of interface debondingto crack penetration into the substrate. The Gi/Gp ratio for FeCo/Ti–6Al–4Vexceeds the transition boundary at locations where the interface toughnessexceeds ≈6 J/m2.

toughness values from Fig. 10 to obtain the Gi/Gp ratios forFeCo/Ti–6Al–4V. Fig. 11 shows that the interface toughness ofFeCo films is sufficiently low so that the ratio of Gi/Gp is gener-ally below the boundary (solid line) for the transition of interfacedebonding to crack penetration into the substrate. Thus, interfacedebonding is expected to be the dominant fracture mechanismin FeCo thin films bonded to Ti–6Al–4V. On the other hand,for 0.9 �m thick FeCo thin film deposited on electropolishedsurfaces, the interface toughness is raised sufficiently to causethe Gi/Gp ratio to exceed the transition boundary that an inter-face crack may deflect from the interface and penetrate into thesubstrate. Since crack penetration into the substrate can lead tofailure in the substrate, the transition boundary dictates that thereis a maximum interface toughness above which further increasein the interface toughness of the thin film is undesirable andshould be avoided.

5.4. Interface debonding strength

Under steady-state interface crack growth condition, thedebonding stress, σd, of an interface crack depends on the inter-face toughness, Gi, according to the relation given by [17]:

σd =[

2EGi

(1 − ν2)t

]1/2

(7)

wte(iwcCsTiF

f the interface crack, and Gp is the critical energy releaseate for crack penetration into the substrate; E1 and E2 arehe Young’s moduli of the thin film (E1 = 1.7 × 105 MPa foreCo) and the Ti–6Al–4V substrate (E2 = 1.7 × 105 MPa fori–6Al–4V), respectively. The theoretical crack morphology

ransition boundary [24] is shown in Fig. 11 as the solid curve.rack penetration into the substrate occurs for Gi/Gp ratiosbove the solid curve, while interface cracks remain to propagatelong an interface for Gi/Gp values that are less than the solidoundary. A possible measure of Gp is the fatigue crack growthhreshold, �Kth, of small cracks or large cracks at a high R ratio.or Ti–6–4, �Kth = 2 MPa

√m [26] and it leads to Gp = 21 J/m2

hich can be used in conjunction with the measured interface

here t is the film thickness, E the Young’s modulus, and ν ishe Poisson’s ratio. The interface toughness Gi is the criticalnergy release rate at the onset of interface crack growth. Eq.7) indicates the debonding stress increases with t−1/2 when thenterface toughness, Gi, is independent of the film thickness,hich was observed in FeCo thin films deposited on mechani-

ally polished surfaces of Ti–6Al–4V substrate [16]. Taken fromhan et al. [16], Fig. 12 shows a summary of the debonding

tresses for FeCo thin films deposited on mechanically polishedi–6Al–4V substrates as a function of film thickness. The results

ndicate that the debonding stress increases with decreasingeCo thickness for thickness less than 4 �m. For film thickness

K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308 307

Fig. 12. Interface debonding as a function of film thickness for FeCo/Ti–6Al–4V. From Chan et al. [16].

less than 3 �m, the thin films did not debond at 748 MPa. Forthin films deposited on surfaces without ion cleaning, the inter-face toughness was about 2 J/m2, while Gi was about 4 J/m2 forthin films deposited on ion-cleaned surfaces. The correspond-ing debonding stresses computed based on Eq. (7) and theseinterface toughness values are in reasonable agreement withthe experimental data, as shown in Fig. 12. The lower interfacetoughness (2 J/m2) gives debonding stresses that are comparableto the lower bound data of the thicker thin films (>4.2 �m), whilethe higher interface toughness (4 J/m2) gives debonding stressesthat are comparable to the upper bound data of the thin filmsover the entire range of film thickness investigated. The resultsof this study indicate that the interface toughness can increasewith decreasing layer thickness when 0.9 mm films are depositedon electropolished surfaces. According to Eq. (7), the interfacedebonding stresses are expected to further increase beyond thoseobserved in FeCo films deposited on mechanically polished sur-faces.

Based on the measured interface toughness, the debondingstrength of FeCo thin films are expected to be in the range of440–750 MPa for thin films less than 4 �m without and withion-cleaned surfaces, respectively. These debonding strengths ofFeCo thin films are compared against the notch fatigue strengthsof Ti–6Al–4V at the stress ratio, R, of 0.1, 0.5, and −1.0 inFig. 13. The comparison indicates that the debonding strength

FT

of FeCo with an interface toughness of 2 J/m2 is high and exceedthe notch strengths of Ti–6Al–4V at all three R ratios. Becauseof the relatively low interface toughness, interface cracks inFeCo/Ti–6Al–4V are expected to remain on the interface andwould not penetrate into the substrate. Thus, FeCo thin films arenot expected to reduce the fatigue strength of the Ti–6Al–4Vsubstrate. In contrast, an increase of the interface toughnessfrom 2 to 10 J/m2 would increase the debonding stress by afactor of

√5 to 1230–1680 MPa. Unfortunately, the Gi/Gp ratio

is also increased to 0.48, which exceeds the crack penetrationboundary and may lead to a reduction in fatigue life. Therefore,electropolishing of the substrate prior to deposition of FeCothin films are unnecessary and can be undesirable for thin filmsless than 1 �m. In summary, both the surface condition of thesubstrate and the film thickness need to be considered to opti-mize the interface toughness and debonding strength in orderto ensure that any interface cracks formed in FeCo/Ti–6Al–4Vwould not penetrate into the substrate and cause a debit on fatiguelife.

6. Conclusions

The conclusions reached in this study are as follows:

1

2

3

4

5

A

RtTitSi

ig. 13. Interface strength of FeCo/Ti–6Al–4V exceeds the notch strengths ofi–6Al–4V. RS denotes residual stress.

. Cross-section nanoindentation measures the interface tough-ness of FeCo thin film at a local scale with a larger scatterin the measured interface toughness values as the techniqueinterrogates only an interfacial area on the order of 10–50 �min radius.

. Brale indentation in a Rockwell hardness tester measuresinterface toughness of FeCo thin film over a larger interfacialarea (200–300 �m radius) with a smaller experimental scatterin the interface toughness measurements.

. The roughness of the deposition surfaces affects the interfacetoughness of 0.9 �m thick FeCo thin films, but not 4.1 �mfilms.

. A high interface toughness for FeCo/Ti–6Al–4V may beundesirable since it can lead to deflection of an interfacecrack to penetrate into the substrate.

. Optimum interface debonding strength can be achieved inFeCo/Ti–6Al–4V by controlling the interface toughness andthe layer thickness of the FeCo thin films so that the thinfilm does not degrade the fatigue strength of the Ti–6Al–4Vsubstrate.

cknowledgements

This work was supported by the Air Force Office of Scientificesearch through Grant No. FA9550-05-1-0154 and a subcon-

ract with The University of Texas at San Antonio (UTSA).he technical assistance by Mr. J. Spencer and Mr. B. Chapa

n performing the indentation tests and microstructural charac-erization of test specimens, and the clerical assistance by Ms. L.alas and Ms. A. Matthews, all of Southwest Research Institute,

s acknowledged.

308 K.S. Chan et al. / Materials Science and Engineering A 422 (2006) 298–308

Appendix A

The yield strength of FeCo thin films was calculated fromthe nanoindentation indentation hardness data using the methoddeveloped by Giannakopoulos and Suresh [22]. In this proce-dure, the effective elastic modulus of the indenter-specimensystem, E*, was first calculated according to the following equa-tion:

hr

hmax= 1 − d∗ pav

E∗ (A1)

where hr is the residual depth of penetration determined from themeasured load–penetration depth (P–h curve), hmax the maxi-mum penetration depth prior to unloading, Pav the hardness ofthe thin films measured by the nanoindentation hardness data,and d* is a constant with d* = 4.678 for a Berkovich indenter[22]. Then, the load–penetration depth (P–h) curve, and themaximum indentation load (Pmax ≈ Ch2

max) were used to esti-mate the value C (indentation curvature). The yield strength, σy,and the flow stress, σ0.29 at a plastic strain of 0.29 were thenestimated by simultaneously solving the following two equa-tions derived by Giannakopoulos and Suresh [22] for describ-ing the plastic stress–strain response of the coating under theindenter:

σ0.29 − σy = 1 − 0.142hr − 0.957

(hr

)2

(A2)

C

0

waafi

R

Russ (Eds.), Materials Damage Prognosis, TMS, Warrendale, 2004, pp.3–10.

[2] J. Littles Jr., R.G. Pettit, J.J. Schirra, B.A. Cowles, R.A. Holmes, S.M.Russ, A.H. Rosenberger, J.M. Larsen, in: J.M. Larsen, L. Christodoulou,J.R. Calcaterra, M.L. Dent, M.M. Derriso, W.J. Hardman, J.W. Jones,S.M. Russ (Eds.), Materials Damage Prognosis, TMS, Warrendale, 2004,pp. 191–201.

[3] S.J. Hudak Jr., B.R. Lanning, G.M. Light, J.M. Major, J.A. Moryle,E.Q. Enright, R.C. McClung, in: J.M. Larsen, L. Christodoulou, J.R.Calcaterra, M.L. Dent, M.M. Derriso, W.J. Hardman, J.W. Jones, S.M.Russ (Eds.), Materials Damage Prognosis, TMS, Warrendale, 2004, pp.157–166.

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491–496.[9] M.R. Stoudt, R.C. Cammarata, R.E. Ricker, Int. J. Fatigue 23 (2001)

5215–5223.[10] R.S. Sundar, S.C. Deevi, Int. Mater. Rev. 50 (3) (2005) 157–192.[11] L. Zhao, I. Baker, Acta Metall. Mater. 42 (1994) 1953–1958.[12] K.R. Jordan, N.S. Stoloff, Trans. AIME 245 (1969) 2027–2034.[13] N.S. Stoloff, S.J. Choe, K. Rajan, Scripta Metall. Mater. 26 (1992)

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Light, Metall. Mater. Trans. A (2006), in press.[

[

[

[

[[[

[

[[

0.29E∗ hmax hmax

= P

h2 = M1σ0.29

{1 + σy

σ0.29

} {M2 + ln

(E∗

σy

)}for

.5 ≤ pav

σy≤ 3.0 (A3)

here M1 and M2 are constants (M1 = 6.02 and M2 = −0.875 forBerkovich indenter [22]). Based on measured hardness result

nd the estimated yield strength, the Pav/σy ratio for FeCo thinlms was estimated to be about 3.

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