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Surface & Coatings Technology xxx (2015) xxx–xxx

SCT-20243; No of Pages 10

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Surface & Coatings Technology

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Tribological andmechanical behavior of nanostructured Al/Ti multilayers

Sina Izadi, Hesham Mraied, Wenjun Cai ⁎Department of Mechanical Engineering, University of South Florida, 4202 E Fowler Avenue, Tampa, FL 33620, USA

⁎ Corresponding author.E-mail address: [email protected] (W. Cai).

http://dx.doi.org/10.1016/j.surfcoat.2015.04.0390257-8972/© 2015 Elsevier B.V. All rights reserved.

Please cite this article as: S. Izadi, et al., Surf.

a b s t r a c t
a r t i c l e i n f o
Article history:Received 6 February 2015Accepted in revised form 22 April 2015Available online xxxx

Keywords:Al–TiMultilayerNanowearHardnessNanoindentationTEM

Al/Ti nanostructured metallic multilayers (NMMs) with individual layer thicknesses of 2.5 and 30 nm weresputter-deposited on Si substrate. The mechanical and tribological properties of the Al/Ti NMMs were studiedby nanoidentation and nanowear tests. Microstructures of the NMMswere characterized using X-ray diffraction,scanning electronmicroscopy, electron backscattered orientation mapping, and transmission electronmicroscopybefore and after the tests. Decreasing the layer thickness from30 to 2.5 nm led to an increase of hardness from3.13to 4.94 GPa. For both NMMs, the wear rate increased with applied load and decreased with hardness. Nano-indentation tests on the worn surface revealed significant work hardening of the subsurface material. Thedeformation mechanisms of NMMs were found to strongly depend on layer thickness and interface structure.

© 2015 Elsevier B.V. All rights reserved.

1. Introduction

Nanostructured metallic multilayers (NMMs) represent a new classof engineeringmaterials that exhibit excellent physical,mechanical, andtribological properties due to the nano-scale layered structure. They arewidely used in X-ray optics, thin film magnets, wear-resistant coatings,microelectromechanical devices, and as radiation damage tolerant ma-terials. Extensive research has been carried out to study the relationshipbetween the microstructure and mechanical properties of NMMs.In particular, it was discovered that near-theoretical strength can beachieved at extremely small layer thicknesses [1]. The deformation be-havior and strengthening mechanism of NMMs are sensitive to layerthickness and interface structure. It is generally accepted that threedifferent deformation modes may operate, depending on the layerthickness h [1–3]. When h is at submicron length scales, dislocationspile up against the interfaces and the Hall–Petch hardening modelapplies, where the strength increases with decreasing layer thicknessas σ ∝ h−1/2. When h is from a few to a few tens of nanometers, thereis insufficient spacing between the interfaces for dislocation pile-up;thus, Hall–Petch hardening breaks down. The deformation mechanismnow involves the motion of single dislocation bowing between theinterfaces, where the confined layer slip (CLS) mechanism operates. Inthis regime, the yield strength is found to increase with decreasinglayer thickness as σ ∝ ln(h)/h [1]. The stress reaches a maximumvalue at approximately h ≈ 5 nm. After that, deformation mechanismsof NMMs become insensitive to layer thickness but highly sensitive tothe structure and property of interfaces, as they represent a large volumefraction with decreasing layer thicknesses. For NNMs with coherent

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interfaces, such as Cu/Ni multilayers, strength depends on the coherencystrains, and themaximumflowstress is equal to the coherency stress. ForNNMs with incoherent interfaces, the deformation mode is associatedwith the transmission of single glide dislocation across interfaces [4].

Extensive previous research highlights the fact that novel deforma-tion physics control plasticity in the NMMs. However, the vast majorityof these studies focus on multilayers consisting of face-centered cubic(fcc) and body-centered cubic (bcc) structures, while those of fcc/hcp(hexagonal close-packed) multilayers remain limited. In this study, wechose Al/Ti as a model system to study the structure-sensitive mechan-ical and tribological properties of fcc/hcp NMMs. Al/Ti NMMs have beenstudied previously for their self-propagating reactions [5], but there hasbeen very limited prior research on the mechanical and tribologicalbehavior of this system [6]. Due to the low densities of the constituentmaterials (ρAl = 2.7 g/cm3 and ρTi = 4.5 g/cm3), the Al/Ti NMMs, withnominal density of ~3.6 g/cm3, are also promising novel lightweighthigh strength materials. In the current work, Al/Ti NMMs were syn-thesized with two representative layer thicknesses, 2.5 nm and 30 nm,in order to reveal the structural evolution when different deformationmechanismoperates. The goals of the currentwork are to (1) synthesizeand characterize Al/Ti NMMs at different layer thicknesses, (2) investi-gate the mechanical and tribological properties of Al/Ti NMMs, and(3) establish the relationship between the microstructure evolutionand tribological responses as a function of layer thickness.

2. Experimental procedure

Al/Ti NMMswith equal-spaced Al and Ti layers were deposited on Si(100) substrates in a load-locked, multi-chamber SFI sputtering toolusing an Eratron 8210 DC power supply. The aluminum chamber wasinstalled with a 9″ diameter S-gun toroidal Al target (99.999%), and

dx.doi.org/10.1016/j.surfcoat.2015.04.039

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the Ti chamber was installed with a 3″ diameter S-gun toroidal Ti target(99.999%). The depositions were done sequentially under vacuumwitha robot transferring the wafer from the Ti chamber to the Al chamberand back again. Argon flowwas adjusted tomaintain a deposition pres-sure of 5 mTorr. No substrate biasing or backside heating was appliedduring deposition. Two types of NMMs were deposited, hereafter de-noted as AT2.5 and AT30 (as listed in Table 1), with individual layerthicknesses of 2.5 nm and 30 nm and total film thickness of 930 ±16 nm and 1.8 ± 0.02 μm, respectively. The arithmetic average surfaceroughness (Ra) of AT2.5 and AT30 is 5.96 nm and 8.67 nm, respectively,measured using Digital Instruments atomic force microscopy.

Grazing incidence X-ray diffraction (GIXRD) was performed usingPANalytical X'Pert PRO diffractometer (Cu Kα, 154.06 pm) at 45 kVand 40 mA with an incidence angle of 3° and step size of 0.025°. Thepenetration depth (PD, the depth at which the intensity is attenuatedby a factor of 1/e) was estimated from PD = sin α/μ [7], where α isthe grazing incidence angle and μ is the linear adsorption coefficient,~523.8 cm–1 estimated from Al50Ti50 composite [8]. The calculated PDis ~1.0 μm, comparable to the total film thickness and much largerthan the individual layer thickness of both samples. Sample surfaceswere examined using a Hitachi SU70 scanning electron microscopy(SEM) at an accelerating voltage of 30 kV. Electron backscattered dif-fraction (EBSD) analysis was performed on a JEOL ARM transmissionelectron microscope (TEM) operated at 200 kV coupled with Topspindata acquisition system. The phase and orientation mappings were ob-tained using ASTAR software with a precession angle of 0.5° and a stepsize of 5 nm. TEM analysis was performed using TECNAI G2 in bright-field (BF), dark-field (DF), and selected area diffraction (SAD) modes.TEM samples were prepared using the standard lift-out method in aFEI Quanta 200 focused ion beam (FIB) microscope. Protective carbon(~100 nm) and Pt layers (~2–3 μm)were deposited on sample surfacesprior to milling in order to minimize FIB-induced damages.

Film hardness and elastic modulus were measured using a Hysitrontriboindenter Ti900 with a Berkovich tip following the Oliver and Pharrmethod [9]. The tip area function was calibrated using a standard fusedquartz sample. For each nanoindentation test, a trapezoidal loadingprofile was used with 2 or 3 mN maximum load, 2 s holding time, and5 s loading/unloading time. This maximum load was chosen to ensurethat the total penetration depth was less than 15% of the total samplethickness to minimize the substrate effect [10]. Average spacing be-tween the indents was kept at least 8 μm apart to minimize indentationinteractions.

Nanowear testswere performed using the triboindenter with a 1 μmcono-spherical tip. Wear tests were performed under 30, 60 and 93 μNnormal loads and up to 10 cycles over areas of 2 × 2 μm2. Wear volume(V) is calculated by multiplying the wear area by the wear scar height,whichwas determined from the height difference between the unwornand worn surfaces using Hysitron Triboview software. Under a specificnormal load (F), the wear rate (w) is defined as the ratio betweenwear volume and the number of cycles (n). Post-wear nanoindentationtests were performed on selected samples with a worn area of 35 ×35 μm2 to probe the subsurface hardness, using the same loading pro-files defined previously in the nanoindentation tests.

Table 1Summary of microstructure and mechanical properties of Al/Ti multilayers. Layerthickness (h) and in-plane grain size (d) were measured from TEM analysis. Mechanicalproperties including maximum penetration depth (hmax), hardness (H), and reducedmodulus (Er) were measured from nanoindentation under 2 ̶ 3 mN load.

Sample h (nm) d (nm) hmax (nm) H (GPa) Er (GPa) Er/H

AT2.5 2.5 ± 0.3 138 ± 43 128.2 ± 5.7 4.94 ± 0.44 120.1 ± 7.6 24.3AT30 30.0 ± 1.5 140 ± 30 216.0 ± 13.6 3.13 ± 0.43 110 ± 7.2 35.1

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3. Results

3.1. Microstructure of as-deposited Al/Ti NMMs

GIXRD line scans of as-deposited AT2.5 andAT30 are shown in Fig. 1.The high intensity peak located at 2θ=56.2° is ascribed to Si (311) fromthe substrate. It can be seen that for both samples, fcc Al phase co-existswith hcp Ti phase. The peak at 2θ≈ 38.5° arises from superposition of Al(111) and Ti (00.2). The lattice constants are aAl = 4.046 Å for fcc-Al,and aTi = 2.969 Å and cTi = 4.666 Å for hcp-Ti. Significant peak-broadening was observed in AT2.5 due to its ultra-fine layer thicknessand, very likely, bi-axialmicro strain [11]. Unfortunately, a good estima-tion of themicro strain is not possible here due to insufficient Bragg dif-fractions and overlapping of Al(111) and Ti (00.2). SEM images of theas-deposited samples are shown in Fig. 2(a) and (b). The surface mor-phology of both samples consists mainly of spherical nodules around200–500 nm in diameter. Cross-sectional TEM images of AT2.5 andAT30 in Fig. 2(c) and (d) reveal the equal-spaced layered structure,where the bright and dark stripes correspond to the Al and Ti layers,respectively. The average layer thickness (h) is 2.5 ± 0.3 nm and30 ± 1.5 nm for AT2.5 and AT30, respectively, measured from cross-sectional TEM images. Low magnification TEM images (not shownhere) reveal that through thickness columnar grains are present forboth samples, similar to those reported previously [12]. The in-planegrain size (d) is ~140 nm for both samples, much larger than their re-spective layer thicknesses. Thus, in the present study, the layer thick-ness is considered as the critical structural length scale that governsthe mechanical and tribological behavior [2].

HRTEM image of AT30 in Fig. 3(a) shows an orientation relationshipof Al (111) || Ti (00.1), both parallel to the interface plane. The interfaceis semi-coherent, with out-of-plane edge dislocations separated by~2.44 ± 0.65 nm (see IFFT images in Fig. 3(b) and (d)), in agreementwith a previous report [13]. These quasi-periodically spaced edge dislo-cations also lead to a misorientation of ~3.2° between Al (111) and Ti(00.1) planes. The HRTEM image in Fig. 3(c) shows Al (200) || Ti(10.1), both inclined to the interface plane at an angle of ~55°, with ameasured misfit strain of ~8%. According to the Al/Ti multilayer bi-phase diagram [13], the interface between the hcp-Ti and fcc-Al layertransforms from incoherent to coherent when the bilayer thickness be-comes smaller than ~12.5 nm [13]. However, for AT2.5 (with a bilayerthickness of ~5 nm), a coherent interface is not observed. Instead,semi-coherent interfaces with dislocation spacing of ~2.23 ± 0.50 nmare observed. Interestingly, Fig. 4(a) shows that the hcp-Ti layer of as-deposited AT2.5 exhibited fcc structure (Fig. 4(b) and (c)) under TEManalysis. Similar polymorphic transformation was reported previously

Fig. 1. Grazing-incidence XRD line scans of as-deposited AT2.5 and AT30.


Fig. 2. SEM images of surface morphology of as-deposited (a) AT2.5 and (b) AT30. Cross-sectional TEM images of as-deposited (c) AT2.5 and (d) AT30. Arrows in (c) and (d) indicate filmgrowth direction.

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in Al/Ti [14] and Ni/Ti [15] multilayers with h Ti around 30 ̶ 500 nm,where the as-deposited hcp-Ti layers transformed to fcc structureafter ion milling during TEM sample preparation. The measured latticeconstant ratio between the fcc-Al and fcc-Ti phases is aTi/aAl = 1.08,very close to the value of 1.09 reported previously for Al/Ti multilayers[14]. In addition, the HRTEM image and SAD patterns in Fig. 4 indicatethat the fcc-Ti and fcc-Al layers were twined with noncoherent twinboundaries, as indicated by white arrows in Fig. 4(a).

3.2. Nanoindentation of as-deposited Al/Ti NMMs

Representative loading/unloading curves of AT2.5 and AT30 areshown in Fig. 5. No obvious “pop-in” was observed from the load–displacement curves of both samples, indicating the absence of macro-scopic shear band formation. A higher penetration depth is observedfor AT30 than AT2.5 under the same loading condition, indicating ahigher resistance to plastic deformation of the latter. The reducedYoung's modulus (Er) and hardness (H) of as-deposited samples arelisted in Table 1. In particular, AT2.5 exhibits a high hardness of4.94 GPa, much higher than the hardness of a Al50Ti50 composite(~2.65 GPa), estimated from nanohardness of monolithic Al (HAl =0.67 ± 0.18 GPa) and Ti (HTi = 4.62 ± 0.62 GPa). The nanohardnessof monolithic Al (2 μm) and Ti (1.2 μm) was measured using 800 μNand 1.7 mNmaximum loads respectively following the same procedureas the multilayered samples. It is also interesting to note that AT2.5exhibits a smaller Er/H value than AT30, which indicates a higher wearresistance of the former sample [16–18].

Insets in Fig. 5 show the post-indentation surface topographies ofboth samples. Whereas indents of AT30 exhibit high contrast from the


undeformed surface, those of sample AT2.5 are hardly visible due tothe small penetration depths. After careful SEM analysis, it was not pos-sible to differentiate the indents of AT2.5 from the as-deposited surfaceroughness. Thus, post-indentation microstructure characterization wasperformed only for AT30. Fig. 6(b) shows the phase map of the boxarea in Fig. 6(a), where the red and green colors represent the fcc-Aland hcp-Ti phases, respectively. It can be seen that immediately belowthe indent, Al layer thicknesses are reduced to ~10 nm (correspondsto a local compressive strain ~67%), while Ti layer thicknesses remainalmost unaffected. For multilayers consisting of two phases with differ-ent strengths, plastic deformation typically occurs in the softer phasefirst until the stress is high enough to transfer or generate dislocationsin the harder phase. Thus, in the present study, it is not surprising thatplastic deformation is mainly accommodated by the softer Al layers. Inaddition to layer compression, several shear bands (SBs) are observed,as indicated by the dashed lines in Fig. 6(b). During co-deformation ofmultilayers, shear bands are often observed due to the difference inflow stresses and hardening behaviors of the adjacent layers, suchas those seen in the Al/W [19] and Cu/Ta multilayers [20,21]. Fig. 7shows the orientation maps of AT30 after nanoindentation. FromFig. 7(b), it can be seen that (111)Al || (00.2)Ti || film growth direction(y-axis as defined in Fig. 6(a) inset). This is not surprising as preferredgrowth in b111N and b00.2N directions for fcc and hcp thin filmshave been commonly observed to minimize surface energy during de-position [11]. Just below the indent, grains with orientations differentfrom the original were observed, as indicated by the white arrows inFig. 7(c). The orientation change of these sub-grains clearly indicateslocal grain rotation and formation of cell boundaries perpendicular tothe layer interfaces.


Fig. 3. (a) and (c) HRTEM images of as-deposited AT30, (b) inverse fast Fourier transforms (IFFT) image of box area in (a), and (d) IFFT of image (c). Insets in (b) and (d) show the FFT ofimages (a) and (c), respectively. Arrow in (a) and (c) indicates film growth direction.


3.3. Nanowear tests and subsurface microstructural evolution

Fig. 8 shows typical surface topographies of the two samples after10 cycles of wear tests under 93 μN load. Surface height profiles inFig. 8(c) show that both the wear scar depth and the pile-up heightare lower for AT2.5, indicating smallermaterial loss and higher wear re-sistance than AT30. Fig. 9(a) and (b) shows that thewear rate is initiallyhigh during the running-in period, which then reaches steady-statewear after ~five cycles. Fig. 9(c) shows that the wear rates increase al-most linearly with the applied load, in agreement with Archard's law[17]. In addition, the wear rate of AT30 is at least two times that ofAT2.5, as expected due to its lower hardness and higher Er/H ratio[22]. It was noticed that increasing the wear track dimensions from2 × 2 μm2 to 35 × 35 μm2 leads to an increase of wear scar depth. Thisis not surprising given that the conical tip radius (1 μm) is close to thesmaller wear track dimensions (2 × 2 μm2). For example, the wearscar depth of AT30 after 10 cycles of wear under 93 μN load increasedfrom ~110 nm to ~650 nm when the wear track dimension waschanged from 2 × 2 μm2 to 35 × 35 μm2. Scratchwear involves two pro-cesses: indentation in thefilm normal direction,which often leads to se-vere plastic deformation of the subsurface material, and translationalong the sliding direction, which leads to material removal due to


abrasive wear [23]. For a large wear track, the translation motion dom-inates and material loss is mainly due to abrasive wear. On the otherhand, for a small wear track, the proportion of abrasive wear becomescomparable to film indentation. In this case, a significant proportion ofmaterial loss is due to layer compression, instead of abrasion. Thus asmaller amount of total material loss is expected when the wear trackdimension becomes comparable to the tip radius.

Fig. 10 shows a montage BF cross-sectional TEM image of AT30 after10 cycles of wear tests under 93 μN load. The worn and unworn areasand the pile-up can be seen clearly. Within the top ~100–150 nm, a se-verely plastically-deformed layer (SPDL) was observed. During nanocontacts, the plastic zone size c can be estimated from Johnson's spher-ical cavity model as [24]

c ¼ffiffiffiffiffiffiffiffiffiffiffiffi3P

2πσy

sð1Þ

where P is the normal load, and σy is the yield strength (approximat-ed as 1

3H from Tabor relationship). The estimated c from Eq. (1) is~207 nm for AT30, close to the size of SPDL observed by TEM. Withinthe SPDL, significant layer compression and shearing are observed,


Fig. 4. (a)HRTEM image of as-depositedAT2.5, (b) corresponding selected area diffraction of image (a), and (c) schematic of diffraction patterns from twinned fcc-Al and fcc-Ti along b110N axis.


and the well-defined interfaces between Al and Ti layers are lost. Weardebris and pile-ups are mainly generated by material removal from theSPDL. Fig. 11(a) and (b) show the BF and DF TEM images of a wear de-bris attached to the unworn surface of AT30. It can be seen that thema-jority of the debris contains severely-refined layers, with the minimumlayer thickness ~3.8 nm for both the Al and Ti layers, corresponding to alocal compressive strain of ~87%. The interfaces of the deformed layersare wavy due to shear band formation, as shown in the pile-up inFig. 11(c) and (d). A similar wavy layered structure also was reportedin dual phase Ag–Cu alloy after severe sliding wear [25]. In addition,vortex-like features are often observed in the SPDL, as shown inFig. 11(c) and (d), which resembles the Kelvin–Helmoltz instability[26]. Using molecular dynamics simulations, Kim et al. [26] have dem-onstrated that these subsurface vorticities were developed due towear-induced shear instability at the contacting surfaces.

Fig. 5. Typical nanoindentation load–displacement profiles of AT2.5 and AT30 under 3mNload. Insets showsurface topographies of the two samples after nanoindentation. Scale barin both insets is 10 μm.


The subsurface microstructure of AT2.5 was drastically differentfrom AT30. Fig. 12 shows the TEM images of AT2.5 after 10 cycles ofwear under 93 μN load. Unlike AT30, no shear band formation was ob-served and all layers remained planar. The SPDL is ~30 nm, significantlysmaller than the plastic zone size of ~170 nm estimated from Eq. (1).HRTEM image in Fig. 12(c) shows that just below the surface, the con-trast between adjacent Al and Ti layers diminished, indicating the inter-faces are more relaxed with fewer misfit dislocations. IFFT images atdepths ~10 nm (inset 1 in Fig. 12(c)) and 40 nm (inset 2 in Fig. 12(c))confirm that the lattice mismatch and the edge dislocation density arereduced near the top of the SPDL. These results suggest that the non-equilibrium interfaces of Al/Ti NMMs,which contain excess free volumeand dislocations, were relaxed during abrasive wear, similar to the re-laxation of non-equilibrium grain boundaries after cyclic mechanicalloading [11]. Finally, it is noted that while the γ-AlTi and Al3Ti interme-tallic phases were found in Al/Ti NMMs after ion irradiation [27] or highpressure torsion [28], no intermixing was observed after indentation orwear in the present work.

4. Discussion

4.1. Phase stability of Ti in Al/Ti NMMs

An interesting observation of the presentwork is that the Ti layers ofas-deposited AT2.5 transform from hcp to fcc structure after TEM sam-ple preparation. A similar behaviorwas not observed in AT30, indicatingthat the phase stability of Ti is related to the layer thickness of NMM.Even though fcc-Ti is not predicted by the equilibrium phase diagram,metastable fcc-Ti has been reported in epitaxial Ti thin films (withb50 nm film thickness) [29] and Ti powders after high energy mechan-ical milling [28], where compressive residual stresses were both pres-ent. Chakraborty et al. [11] showed that a compressive stress of morethan 2.5 GPa was developed in Ti at low film thicknesses. This highstress is likely to generate Shockley partials on basal planes that canalter the stacking sequence, similar towhat occurred duringmartensiticphase transformation. The phase stability of fcc-Ti and hcp-Ti can be fur-ther understood by considering the change of total free energy, whichincludes the change in bulk free energy, elastic strain energy, and sur-face free energy, during the hcp to fcc transformation. At large filmthicknesses, the bulk free energy term dominates and hcp-Ti is more


Fig. 6. (a) Cross-sectional BF TEM image of AT 30 after nanoindentation, and (b) orientation phase map from box area in (a). Red and green in (b) represent the Al and Ti phasesrespectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)


stable (~0.065–0.07 eV/atom lower than fcc-Ti [11]); at low film thick-nesses, the elastic strain energy and surface free energy dominate andfcc-Ti becomesmore stable. A previous study [11] showed that the strainenergy density of fcc-Ti thin film is about one order of magnitude lowerthan that of hcp-Ti at thicknesses smaller than 144 nm. Thus, there existsa criticalfilm thickness, belowwhich fcc-Ti ismore stable thanhcp-Ti. For

Fig. 7.Orientationmaps in (a) X, (b) Y, and (c) Z direction (as defined in Fig. 6(a)) of AT 30 afte(d). (For interpretation of the references to color in this figure legend, the reader is referred to


Al/Ti multilayers, this critical film thickness was found to increase withdecreasing volume fraction of Ti [13]. According to Banerjee et al. [13],for Al/Ti multilayers with 50% volume fraction of Ti, the estimated criticallayer thickness is around 1.4–1.8 nm, which is comparable to the layerthickness of AT2.5, andmuch smaller than that of AT30. Thus, in the pres-ent work, metastable fcc-Ti was seen only in the former.

r nanoindentation test. The orientations are color coded according to the triangle legend inthe web version of this article.)


Fig. 8. Surface topographies of (a) AT2.5 and (b) AT30 after 10 cycles of wear with 93 μN load. (c) Surface height profiles of samples AT2.5 and AT30 measured from images (a) and (b).

Fig. 9.Wear rates as a function of number of cycles for samples (a) AT2.5 and (b) AT30. (c) Wear rates of AT2.5 and AT30 after 10 cycles of wear test under 30, 60 and 93 μN load.


4.2. Al/Ti interface structure

Anorientation relationship of (111)Al || (00.2)Ti || interface planewasobserved in both as-deposited samples (note the (00.2)Ti plane of AT2.5transformed to (111)Ti after TEM sample preparation). The interfaceswere always semi-coherent, separated by edge dislocations with spac-ing around 2.2–2.4 nm (see Section 3.1 for details). For NMMs, the de-gree of interface coherency depends on the layer thickness, and thereexists a critical layer thickness (hc) for epitaxial growth. If h≤ hc, the in-terface is fully coherent; if h N hc, dislocationswill be formed to relax the

Fig. 10. Montage BF TEM image of AT30 after 10 cycles of wear under 93 μN load. The wear srespectively. Carbon and Pt were used as protective coatings during TEM sample preparation u


misfit strain, and semi-coherent or incoherent interfaceswill be formed.This critical layer thickness hc can be estimated as [30]:

hc ¼ b2πεm

ln

ffiffiffi2

phcb

!ð2Þ

where b is the Burgers vector and εm is the lattice misfitstrain. In the current Al/Ti system, assume baN type Burgers vector for

Ti, b!¼ 1

3b1120N = 0.17 nm, the misfit strain εm = 8%, and the

car area is 35 × 35 μm2. The bright and dark contrast corresponds to the Al and Ti layerssing FIB.


Fig. 11. (a) BF and (b) DF TEM images of a wear debris attached to the surface of AT30 after 10 cycles of wear under 93 μN load.White dashed lines represent thewear surface. (c) BF and(d) DF TEM images of pile-up, (e) and (f) BF TEM images of subsurface material from the same sample.


calculated value of hc is 0.43 nm, comparable to a few monolayers.Previous experimental work shows that Ti grows epitaxially on Al(100) up to 5.5 monolayers [15], and on Al (111) up to 2 monolayers[31]. This small hc indicates that misfit dislocations are inevitable to


accommodate for the misfit strain for Al/Ti NMMs with nanometerscale layer thicknesses. The spacing between misfit dislocations, λ,can be estimated as λ ¼ b

εm= 2.1 nm, which agrees very well with

the HRTEM observation.


Fig. 12. (a) BF, and (b) DF TEM images of AT2.5 after 10 cycles ofwear under 93 μm load. (c)Montage HRTEM image from the SPDL (as indicated in (a)). Thewear scar area is 35 × 35 μm2.Dashed lines represent the wear surface. Insets 1 and 2 in (c) correspond to the IFFT of box area 1 and 2 respectively.


4.3. Wear induced subsurface work-hardening

As presented in Section 3.3, wear leads to characteristically differentmicrostructure evolutions for AT2.5 and AT30. To evaluate how thesemicrostructure evolutions affect the subsurface mechanical properties,nanoindentation tests were performed on worn surfaces of AT2.5 andAT30 after wear tests under 60 μN normal load up to five cycles. Theresults are shown in Fig. 13. It can be seen that both AT2.5 and AT30 ex-perienced significant work-hardening, reaching a hardness of 6.80 ±0.16 GPa and 4.35 ± 0.17 GPa, respectively, after five cycles of wear.Quite remarkably, this subsurface hardening is more significant forAT2.5, where the hardness increased by ~30% despite its small micro-structural change. As discussed in the introduction, at layer thicknessof 2.5 nm, the flow stress is related to the interface barrier stress,which is highly sensitive to the interface structure. As shown inFig. 12(c),with decreasing distance from the surface, a decrease of inter-face dislocation density was observed. In other words, wear increasedthe interface coherency right below the contact surface. Thus higher co-herency stress needs to be overcome for dislocations to transmit to theother layers [29]. For AT30, the hardness increased by ~39% after five cy-cles of wear. At this layer thickness, the operative strengthening mech-anism is CLS (or modified CLS). Consider the greatest layer thickness

Fig. 13. Nanoindentation hardness measured on wear scars (35 × 35 μm2) of AT2.5 andAT30 after wear tests under 60 μN load up to five cycles.


refinement (~3.8 nm) in the SPDL of AT30, it can be estimated thatthe maximum strength is increased by ~2.1 times compared to thebulk material. However, since these ultrafine structures were only ob-served in the topmost subsurface material, the measured hardness ismuch lower and represents an average value over the entire SPDL. Finally,it is worth noting that Al/Ti NMM studied here is a promising lightweighthigh strength material. Given the low density of Al and Ti, the estimatedspecific strength of Al/Ti NMM is around 320–500 kN m/kg (assuming aTabor factor of three), well in excess of steel (130 kN m/kg), Al-6061(110 kN m/kg), and Ti–6Al–4V (240 kN m/kg) [32].

5. Conclusions

In summary, equal-spaced Al/Ti NMMs with individual layer thick-ness (h) of 2.5 and 30 nm were deposited on Si substrate by physicalvapor deposition. Semi-coherent interfaces between theAl and Ti layerswere observed in both samples, with an orientation relationship of Al(111) || Ti (00.2), parallel to the interface plane. The misfit strain was~8%, and the nominal misfit dislocation spacing was ~2.2–2.4 nm forboth samples. The hcp-Ti layer was unstable at lower layer thicknesses(2.5 nm) and transformed to metastable fcc structure after TEM samplepreparation. Both hardness and wear resistance increased with de-creasing layer thickness. The wear rate approached a steady state afterfive cycles and scaled linearly with applied load, in agreement withArchard's law. Wear induced severe plastic deformation of the subsur-face material and led to the formation of a SPDL of 100–150 nm and30 nm for AT30 and AT2.5, respectively. The deformation mechanismwas found to strongly depend on layer thickness. At h= 30 nm, severecompression and shearing of Al and Ti layers led to shear band andvorticity formation in the SPDL. At h=2.5 nm,wear-induced plastic de-formation relaxed the interface and lowered misfit dislocation density.These relaxed interfaces of AT2.5 led to an increase in surface hardnessby ~30% after five cycles of wear.

Acknowledgment

The authors thankfully acknowledge the assistance of Richard Everlyfor magnetron sputtering, Jay Bieber for SEM analysis, Robert Tufts forXRD measurement, and Yusuf Emirov for TEM analysis. The samplepreparation and material characterization were performed at theNano Research and Educational Center (NREC) at the University ofSouth Florida. The authors thankfully acknowledge the assistance ofAmith Darbal from NanoMEGAS for the orientation mapping analysis.This work was financially supported by a University of South Floridastart-up fund.



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