Surface & Coatings Technology 216 (2013) 46–51

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Surface & Coatings Technology

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Oxygen diffusion hardening of tantalum coatings on cp-titanium forbiomedical applications

T. Schmitz a, C. Hertl b, E. Werner b, U. Gbureck a, J. Groll a, C. Moseke a,⁎a Department for Functional Materials in Medicine and Dentistry, University of Würzburg, Pleicherwall 2, 97070 Würzburg, Germanyb Lehrstuhl für Werkstoffkunde und Werkstoffmechanik, Technische Universität München, Boltzmannstrasse 15, 85748 Garching, Germany

⁎ Corresponding author.E-mail address: claus.moseke@fmz.uni-wuerzburg.d

0257-8972/$ – see front matter © 2012 Elsevier B.V. Allhttp://dx.doi.org/10.1016/j.surfcoat.2012.11.021

a b s t r a c t

a r t i c l e i n f o

Article history:Received 28 November 2011Accepted in revised form 7 November 2012Available online 16 November 2012

Keywords:Magnetron sputteringPhysical vapor depositionSurface hardeningVickers hardness

Protective tantalum coatings on titanium substrates were produced using a two step process. At first, substrateswere coated with Ta layers of 5 μm thickness by physical vapor deposition (PVD). In a second step, the coatedsamples were hardened by oxygen diffusion for up to 3 h. During this process the samples were exposed tooxygen for 1–2 h at a pressure of 6.7·10−3 mbar at 350–450 °C, followed by 1–2 h annealing in oxygen-freeatmosphere at the same temperature. X-ray diffraction (XRD) analysis demonstrated a shift of peaks for oxygendiffusion treated samples, which was attributed to the diffusion of atomic oxygen into the Ta-layer. The herebycausedmechanical stress in the crystal lattice led to an increase in Vickers hardness of the Ta layers from 570 HVto over 900 HV. In order to compare the adhesion of untreated samples with oxygen diffusion treated samples,the coatings were investigated using Rockwell measurements. These tests demonstrated an increase of criticalforce for coating delamination from 12 N for untreated samples up to 25 N for diffusion treated samples.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

Titanium and its alloys are standard materials for implant applica-tions involving contact with both hard and soft tissues because oftheir advantageous combination of good mechanical properties andhigh biocompatibility [1]. The native oxide layer protects the metalagainst corrosion in physiological environment [2], whereas the com-paratively low elastic modulus and good fatigue strength of the bulkmaterial make them highly suitable particularly for the replacementof hard tissues, e.g. as anchoring parts in total hip and knee arthroplasty[3], as supportive devices for fracture healing [4], and as enossalimplants [5]. Adsorbed biomolecules generally undergo only few struc-tural changes on the surface of these materials [6]. Hence, the bulkmaterial is not recognized as a foreign material by the cellular environ-ment, which is the reason for the high biocompatibility of titanium.Despite the prevailing positive experiences with the use of titaniumand its alloys in biomaterial applications, under certain circumstancesan aseptic loosening of titanium prostheses can be observed aftershort implant duration. This is accompanied frequently with the forma-tion of abrasion debris from the prosthesis surface caused by relativemovements between hard tissue/bone cement and implant. The nativeoxide layer, which is only a few nanometers thick and shows lowmechanical stability, does not resist these movements [7,8]. Due totheir small size, the generated abrasive particles cause inflammatoryreactions of the surrounding tissue followed by progressive loss ofbone in the worst case [9,10].

e (C. Moseke).

rights reserved.

Over the last years numerous studies have been undertaken toimprove the tribological properties of the titanium surface by differentcoating technologies such as chemical (CVD) [11] and physical (PVD)vapor deposition [12–14] aswell as thermal and electrochemical oxida-tion techniques [15,16]. Hard material layers consisting of metal oxidesand nitrides can improve the abrasion resistance of the bulk material.However, a considerable disadvantage of these material systems is theabrupt transition from the brittle/hard mechanical properties of thesurface coating to the ductile properties of the substrate, which maylead to delamination of the coating. A gradient-like transition zonebetween the mechanical properties of the hard coating and the softsubstrate would be preferable. For the case of bulk titanium and Tialloys the enhancement of surface hardness can be achieved by oxygendiffusion hardening (ODH) [17–20].

The aim of the study at handwas to expand this technique frombulkto layer systems in order to generate self-healing gradient-like hardcoatings on titanium (Fig. 1). In the first step a thin tantalum layer ofapprox. 5 μm is deposited on the surface by radio frequency (RF) mag-netron sputtering. The refractory metal tantalum is chosen due to thefact that the reaction with oxygen not only occurs much faster thanfor titanium, but is also accompanied by a volume increase that willresult in self-induced crack healing of the surface [21]. In a secondstep, oxygen diffusion hardening is applied to these coatings to achieveboth a hardened surface with high abrasion resistance and a smoothtransition zone of mechanical properties from the surface to the bulkmaterial. The coatings then are characterized regarding their chemicalcomposition,morphology and adhesion using X-ray diffraction analysis,scanning electron microscopy (SEM), Rockwell testing and Vickershardness measurement.

Fig. 1. Concept of producing oxygen diffusion hardened Ta layers on Ti substrates.

47T. Schmitz et al. / Surface & Coatings Technology 216 (2013) 46–51

2. Materials and experimental methods

2.1. Sample preparation and coating process

Tantalum films were deposited on commercially pure titaniumdisks (grade 2, 15.5 mm diameter, 1 mm height) by RF magnetronsputtering using a Ta target (120 mm diameter, 10 mm height)with a target-to-substrate distance of 100 mm. Ti disks were mechan-ically ground and polished to mirror-like appearance. All substrateswere cleaned in an ultrasonic bath (first with acetone for 10 min,then with ethanol for 10 min) and finally dried in air. The Ti sampleswere attached to the substrate holder of a PVD-system of type PLS570 (Pfeiffer Vacuum, Asslar, Germany). The substrate holder wasespecially designed for this purpose and could be heated by a pair ofinternally installed halogen lamps. The temperature was measured bya thermocouple that was embedded between two titanium samples.The chamberwas evacuated for 15 h at a temperature of 40 °C followedby a one hour cool-down to 15 °C, with both heating and coolingachieved by means of temperature-stabilized water circulation. Thisprocedure resulted in a base pressure of 1∗10−6 mbar. 30 min beforedeposition the substrate holder was heated up to 350 °C or 450 °C.Prior to deposition the substrateswere sputter-cleaned in argon plasma(300 W, 180 sccm, 1.6∗10−2 mbar) for 10–15 min.

A wide range of process parameters for sputter deposition could bevaried, such as deposition time (180–300 min), working pressure (6.7–9.9∗10−3 mbar), negative substrate bias voltage (0–300 V) and sub-strate temperatures up to 550 °C. However, the coatings examined inthis study were deposited with a set of fixed parameters, namely adeposition time of 180 min, a working pressure of 7.0∗10−3 mbar(100 sccm Argon) and a substrate bias voltage of 0 V, while the sub-strate temperature was either 350 or 450 °C. On the basis of measuredthickness and deposition time the calculated deposition rate wasapproximately 0.5 nm/s.

2.2. Oxygen diffusion hardening

Following the deposition process the substrate temperature waskept for 3 h at either 350 or 450 °C. In order to achieve oxygen diffusionhardening directly after finishing the sputter deposition process, theworking gas argon was replaced with oxygen (100 sccm), resulting ina pressure of 6.7∗10−3 mbar. The samples were kept in this oxygenatmosphere at the elevated temperatures for 1 or 2 h. Then oxygenwas removed from the chamber and the samples were heated for addi-tional 1 to 2 h without process gas. Finally the substrates were cooleddown to room temperature before venting the chamber.

Fig. 2. X-ray diffraction patterns of tantalum coatings on titanium obtained at differentbias voltages.

2.3. Coating characterization

The surfacemorphology of the coatingswasdetermined by scanningelectron microscopy (SEM) using a microscope DSM 940 (Zeiss,Oberkochen, Germany). The film thicknesses were determined by theevaluation of SEM images of either cross sections of Ti disks or by dis-ruption of Ta coated titanium foils. The depositedmass was determinedbymeasuring themass of the samples before and after deposition usinga precision balance of type MC1 (Sartorius, Göttingen, Germany). Then

the obtained values for film thickness and mass were used to calculatethe density of the deposited coatings.

The crystal structure of the tantalum films was examined by X-raydiffraction (XRD) in Bragg–Brentano geometry with a Siemens D5005X-ray diffractometer (Bruker AXS, Karlsruhe, Germany) using Cu-Kαradiation with a voltage of 40 kV and a tube current of 40 mA. Depthprofiles of the oxygen concentration in the ODH treated samples wererecorded by sputter ablation of the surface and analysis of the releasedatoms using a glow discharge spectrometer GDS-750A (LECO, TechnikGmbH, Gilching, Germany). For comparison also depth profile analysisby means of a time-of-flight secondary ion mass spectrometer (TOFSIMS V, IONTOF, Münster, Germany) was performed.

In order to investigate the adhesion of the deposited tantalum coat-ings to the titanium substrates scratch tests were performed with ahardness tester of type 3212B (Zwick, Ulm, Germany) equipped witha Rockwell C diamond with a conical angle of 120° and a tip radius of200 μm. The scratches were induced under constant load with the tipbeing moved across the surface with a constant velocity of 1 mm/min.The scratch traces with a length of 2 mm were evaluated by opticalmicroscopy to determine the critical load atwhich the substratewas ex-posed as a result of coating delamination. Due to the strong dependencyof the critical load on film thickness and substrate roughness, onlycoatings with a thickness of 5–6 μm were tested and the Ti substrateswere polished to mirror-like appearance.

To determine the hardness of the oxygen diffusion hardened sam-ples, Vickers hardness testing was performed using a hardness tester(Micro-Duromat 4000E, Wetzlar) and a light microscope (Metaplan 2,Wetzlar). With a proof load of 490.5 mN five indents were performedon the surface of each specimen. Although this increased proof load isnot in accordance with DIN EN ISO 6507, it was chosen in order toobtain measurable indents.

3. Results and discussion

Fig. 2 shows the diffraction patterns of Ta films deposited on Tisubstrates with different bias voltages. Films deposited without nega-tive substrate bias exhibit a bcc crystal lattice, characteristics of theα-phase of tantalum, with a very strong (110) texture. The texture

Fig. 3. Calculated density of tantalum coating on titanium vs. the applied bias voltage.

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effect decreases with increasing bias voltage and the (110) peak almostcompletely disappears when the bias reaches voltages of−300 V.

Besides the influence on crystallography also an influence of thebias voltage on the density of the deposited films is observed. Ascan be seen in Fig. 3, density decreases almost linearly with increasingbias voltage. One possible reason for this is the so-called re-sputteringeffect [22]: positive Ar ions are accelerated from the plasma towardsthe negatively charged substrate surface and interfere with the filmdeposition by sputtering of already adhered Ta particles, whichresults in reduced film density as well as in reduced growth rate. Thiseffect becomes more pronounced with increasing bias voltage due tothe rising kinetic energy of the Ar ions. An additional effect that maybe responsible for the decreasing density is the incorporation of Arions into the film, as was also described by Catania et al. [22].

However, the decreasing density of the films only partially explainsthe decreasing intensity of the (110) α-Ta peak. By comparison of theSEM images presented in Fig. 3 also changes in the morphology of thedeposited films can be observed. The surface of the Ta film with strong(110) α-phase texture that was deposited without substrate bias(Fig. 4a) shows elongated grains arranged parallel to the plane of thefilm, whereas the film obtained by sputtering with −300 V substratebias voltage (Fig. 4b) shows (in accordance to XRD analysis) a differentsurface morphology with smaller pyramid-shaped grains. Fig. 4c showsa cross-section of a deposited Ta film.

Due to the negative influence of substrate biasing on the density ofthe coatings the Ta films prepared for oxygen diffusion hardeningwere deposited without an additional bias voltage. In addition, thetemperature was raised to at least 350 °C to promote the growth ofthe α-phase of tantalum, which has been found to show ductilebehavior in contrast to the metastable β-phase, which is more brittle[23,24] and therefore less suitable for biomedical applications. Tafilms deposited on titanium with these parameter settings are the

Fig. 4. Scanning electron microscopy images of the surface morphology of 5 μm thick Ta(b); (c): cross-section of the coating shown in (a).

basic material for the subsequent ODH treatment. The diffractionpattern of a Ta film prepared in this way shows the expected strong(110) α-texture (Fig. 5a).

Fig. 5b shows the influence of the ODH treatment on position andshape of the (110) peaks of α-Ta coatings. According to literaturedata, the exposition of the samples to oxygen at 350 °C or 450 °Cwith subsequent annealing in vacuum is expected to lead to the occu-pation of interstitial sites of the tantalum lattice by oxygen [25]. Thisgenerates elastic strain in the tantalum lattice and changes in latticespacing. The position of diffraction peaks is related to the lattice spacingvia Bragg's law [26]:

nλ ¼ 2dhkl sin θhkl; ð1Þ

where dhkl is the interplanar spacing of the diffracting lattice planes(here 110), θhkl is the scattering angle, λ is the wavelength and n is aninteger. The elastic strain εhkl originating from a shift Δdhkl in latticespacing can be calculated from the scattering angle θ of a diffusion-hardened sample and that of an untreated tantalum coated titaniumreference sample θ0:

εhkl ¼Δdhkld0;hkl

¼ dhkl−d0;hkld0;hkl

¼ sin θ0sin θ

−1: ð2Þ

Here d0,hkl denotes the reference lattice spacing (Bragg angle θ0) anddhkl the lattice spacing of the strained lattice (Bragg angle θ). ApplyingHooke's law the resulting stress σhkl for this direction can be calculated

σhkl ¼ E·εhkl; ð3Þ

where σhkl is the stress in direction of lattice strain εhkl and E is themod-ulus of elasticity (170 GPa) for Ta. The compressive stress caused byinterstitially dissolved oxygen could be calculated to 342 GPa for spec-imens oxidized for 1 h and 456 GPa for specimens treated for 2 h.Fig. 5b presents the corresponding peak shift for three samples. In addi-tion to diffraction patterns of the samples exposed to an oxygen atmo-sphere with a pressure of 6.7∗10−3 mbar at 350 °C for 1 and 2 h withsubsequent annealing in vacuum for two respectively 1 h, a diffractionpattern of a sample that was only annealed in vacuum for 3 h withoutany exposition to oxygen was included as a reference.

A significant peak shift was observed between samples annealedwith and without oxygen, and the peak shift slightly increased withtime of exposition to oxygen at a temperature of 350 °C. A similar,though less pronounced, behavior was observed for samples treatedat 450 °C (data not shown here). Considering Bragg's law (Eq. (1)),the shift of the peaks to smaller reflection angles can be attributedto an increase of the lattice plane distances. In addition to the shiftedpeak positions also a decrease of maximum and integrated intensityof the peaks could be observed, particularly in the samples exposedto oxygen for 2 h. This can be attributed either to a reduction of

films deposited using a zero substrate bias voltage (a) and a bias voltage of −300 V

Fig. 5. a) Typical XRD pattern of a Ta film deposited without negative substrate biasand a substrate temperature of 350 °C. b). X-ray diffraction patterns of (110) α-Tapeaks; comparison of tantalum coatings treated by ODH for 1 respectively 2 h with avacuum annealed Ta film.

Fig. 6. GDS depth profiles of oxygen concentration in Ta coatings treated with the follow-ing parameters: a: 350 °C for 1 h, b: 350 °C for 2 h, and c: 450 °C for 1 h, and 450 °C for2 h. The profiles have been shifted on the abscissa with respect to the layer thickness.

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preferential orientation or an increasing number of lattice defectscaused by the incorporation of oxygen atoms into the lattice. Howev-er, a decrease of the material density as described above is unlikely,since the tantalum atoms are not replaced by oxygen during theprocess. Similarly, the formation of tantalum oxide (with a lowerdensity) was not observed during the experiments. Hence the broad-ening of the reflection peaks was most likely caused by a gradient oflattice expansion, which decreased with increasing depth.

The distribution of oxygen in the coating system is shown in Fig. 6by means of depth profiles of oxygen concentration recorded by glowdischarge spectrometry (GDS). These measurements showed practical-ly no differences between samples treated at 350 °C for 1 or 2 h, whichsupports the above-mentioned assumption that the slight peak shiftand the reduction of peak intensity in the diffraction patterns (Fig. 5)may have only occurred due to thermal effects rather than to differ-ences in oxygen incorporation. Apparently the oxygen concentrationwas noticeably increased in the surface-near region of the coating,when the process temperature was elevated from 350 °C to 450 °C.For these samples also the prolongation of the ODH treatment from 1to 2 h resulted in further increasing oxygen concentration.

In a depth of about 5 μm,which corresponded to the average coatingthickness, all depth profiles showed a local maximum of the measuredoxygen concentration, which could not be explained by diffusion ofoxygen introduced by the ODH process. The additionally performedTOF-SIMS measurements revealed that these local oxygen accumula-tions occurred simultaneously with local concentration maxima oftitanium and tantalum oxide at the interface between coating and

substrate. Most probably the native titanium oxide layer could not becompletely removed from the substrate surface due to surface rough-ness in the range of a few hundred nanometers. The appearance oftantalum oxide at the interface can be explained by capture of residualoxygenmolecules in the process gas by Ta atoms at the beginning of thesputtering process, an effect which is characteristic for refractorymetals.

The main objective of the ODH treatment of tantalum coatings wasto develop potential protective coatings for titanium substrates inbiomedical applications where mechanical stress is involved. There-fore, besides the crystallographic properties of the deposited films,also their hardness and their adhesion to the substrate were of crucialimportance. The determination of critical loads of the coatings isstrongly influenced by the underlying substrate, i.e. by its mechanicaland topographical properties like surface hardness and roughness.The results obtained from measurements on different coating–sub-strate systems must be evaluated with caution; actually it seems ap-propriate to restrict the evaluation of film adhesion to relativeresults from measurements within one coating–substrate system. Inthis work the material system chosen for the investigation of its po-tential adhesion improvement by oxygen diffusion hardeningconsisted of Ti substrate, polished to mirror-like surface roughnessand coated with a Ta layer of approximately 5 μm thickness.

The evaluation of the adhesion between substrate and coating iscommonly done by the measurement of the critical force when coat-ing delamination occurs and the substrate becomes exposed. In thescratch tests performed the ODH treated samples partially exhibitedsignificant improvement of adhesion. The determined values for thecritical force LC were typically in the range of 9–12 N for theuntreated samples, but reached values up to 25 N for some ODHtreated samples. Fig. 7 shows the tracks induced by scratch tests onsamples that were coated with Ta in the same way, but underwentdifferent treatment afterwards. Fig. 7a–c shows the grooves obtainedfrom scratch tests performed with loads slightly below LC. Smallcracks occurred in the track, which ran perpendicular to the directionof the tip's movement. This cracking is due to sticking (initial) frictionbetween diamond and coating at the sides of the Rockwell-indenter[27]. Fig. 7e–f shows that depth and width of the cracks increasewith increasing load until reaching LC at which the titanium substratewas exposed partially.

Vickers hardness HV0.05 (Fig. 8) strongly increased with oxidationtemperature and time. While the untreated titanium substrateshowed a hardness of approx. 260 HV0.05, the coating with tantalumincreased hardness to 540–570 HV0.05 with a further improvementto more than 900 HV0.05 after 2 h ODH treatment at 450 °C. Obviouslythe influence of treatment time was significantly higher for the

Fig. 7. Tracks of scratch tests on 5 μm thick Ta films deposited on titanium. Images in the left column correspond to a samplewithout ODH, samples in themiddle and right columnswereannealed at 450 °C in oxygen for 1 or 2 h, respectively. Upper row: below critical loading, (a) 10 N, (b) 10 N and (c) 24 N. Bottom: critical loading, (d) 12 N, (e) 12 N and (f) 25 N.

50 T. Schmitz et al. / Surface & Coatings Technology 216 (2013) 46–51

samples treated at 450 °C, which was in good correlation with thedepth profiles of oxygen concentration. Apparently, the interstitiallydissolved oxygen atoms led to an increase in lattice spacing thatresulted in elastic strains. These strains are associated with stressfields around the oxygen atoms that hinder the motion of dislocationswhich are forced to interact with these fields. This interaction resultsin an increase of flow stress that can be correlated with the hardnessof the tantalum coating [20]. The values of surface hardness plotted inFig. 8 are arithmetic mean values of 5 indents per specimenwith errorbars representing the standard deviation.

4. Conclusions

The deposition of tantalum on titanium at a working pressure of7.0∗10−3 mbar (100 sccm Argon), substrate bias voltage of 0 V,with a deposition time of 180 min and substrate temperature ofeither 350 or 450 °C results in a 5 μm thick layer of alpha-tantalumwith strong (110)-texture. In relation to pure titanium, this coating

Fig. 8. Vickers hardness HV 0.05 of titanium coated with 5 μm tantalum by PVD andfurther ODH treatment at different temperatures for up to 2 h.

causes an increase of surface hardness, which can be further increasedby oxygen diffusion hardening of the tantalum layer. The results ofXRD and GDS indicate a gradient like transition zone between thehigh surface hardness and themore ductile substrate. The oxygen diffu-sion hardened tantalum coatings show higher delamination resistancein scratch tests. Therefore, the enhanced properties of oxygen diffusionhardened tantalum coatings on titanium enlarge thefield of application,i.e. for load-bearing orthopedic implants.

Acknowledgments

The authors want to thank Deutsche Forschungsgemeinschaft DFGfor financial support within projects GB 1/13-1 and WE 2351/12-1 aswell as in the financing of the field emission scanning electron micro-scope JEOL JSM 7600F. More thanks go to Dr. Martin Kamp (Lehrstuhlfür Technische Physik, Würzburg, Germany) for his experimental sup-port with the SIMS measurements.

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