- Email: [email protected]
In vitro studies of the adhesion of diamond-like carbon thin films on CoCrMo biomedical implant alloy
Available online at www.sciencedirect.com
Acta Materialia 59 (2011) 4678–4689 www.elsevier.com/locate/actamat
In vitro studies of the adhesion of diamond-like carbon thin films on CoCrMo biomedical implant alloy C.V. Falub a,b,⇑, U. Mu¨ller a, G. Thorwarth a,c, M. Parlinska-Wojtan a, C. Voisard c, R. Hauert a b
a ¨ berlandstrasse 129, CH-8600 Du¨bendorf, Switzerland Empa, Swiss Federal Laboratories for Materials Science and Technology, U Swiss Federal Institute of Technology (ETHZ), Laboratory for Solid State Physics, Schafmattstrasse 16, CH-8093 Zu¨rich, Switzerland c Synthes GmbH, Langendorfstrasse 2, CH-4513 Langendorf, Switzerland
Received 9 October 2010; received in revised form 4 April 2011; accepted 6 April 2011 Available online 30 April 2011
Abstract Diamond-like carbon (DLC) coatings are promising candidates for improving the mechanical wear and chemical resistance of loadbearing implants. However, in view of medical data showing severe problems due to film delamination, the application of DLC layers to the medicotechnical field is not yet recognized. In an extension to an earlier publication, in which we established that coating delamination in phosphate-buffered saline (PBS) solution is driven by stress–corrosion cracking, we present for the first time a detailed structural and chemical analysis of the interface region of DLC-coated CoCrMo. High-resolution and scanning transmission electron microscopy, energy-dispersive X-ray spectroscopy and X-ray photoelectron spectroscopy shows that a 5 nm thick carbide layer with an average Me:C stoichiometry of 2:1 is formed at this interface. It is shown that the failure of DLC-coated CoCrMo in PBS is due to the instability of Co carbides at the film–implant interface. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Diamond-like carbon (DLC); Interface; Corrosion; Transmission electron microscopy (TEM); X-ray photoelectron spectroscopy (XPS)
1. Introduction Corrosion products and increased wear are the major processes responsible for the aseptic loosening of orthopedic implants after relatively short periods of time [1–3]. A feasible solution to reduce the wear and to enhance the corrosion resistance, biocompatibility and mechanical properties of currently used implant materials is by surface modification techniques, such as ion implantation or deposition of ceramic layers (e.g. diamond-like carbon (DLC), TiN, ZrO2) [4]. While ion implantation has limitations related to its shallow penetration depth (0.3 lm), the ⇑ Corresponding author. Present address: Laboratory for Solid State Physics, Swiss Federal Institute of Technology (ETHZ), Schafmattstrasse16, CH-8093 Zu¨rich, Switzerland. Tel.: +41 44 633 2261; fax: +41 44 633 1072. E-mail address: [email protected] (C.V. Falub).
use of ceramic coatings improves the wear resistance of the implants, though the adhesion of these coatings to the implant is problematic [4]. In view of their excellent mechanical properties (e.g. low wear, high hardness), as well as chemical inertness and biological compatibility, the application of DLC coatings to metallic articulating biomedical implants should provide a promising solution [5–8]. However, in spite of these outstanding properties, the application of DLC films to implants has not yet achieved widespread recognition as the physiological environment has been shown to cause unpredictable delamination of these coatings from the metal surface, which may lead to the total failure of the coated implant [9,10]. DLC layers delaminate from implant surfaces due to their high residual stresses (in the GPa range) and because their interfaces with the implant become corroded in biological fluids [11]. Hence, to minimize these problems and develop successful DLC-coated
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.04.014
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
4679
Fig. 1. Adhesion lifetime prediction of the DLC-coated CoCrMo implants: (a) 3-D profile of a Rockwell indentation in a 4 lm thick DLC-coated CoCrMo substrate; (b) (G, hvi) diagram derived from optical monitoring of the delaminating crack front and finite-element calculations [11].
implants, it is necessary not only to perform simulator tests under conditions closely resembling the in vivo situation [12], but also to design biocompatible interface layers that are wear and corrosion resistant. For this purpose, a detailed understanding of the mechanisms involved in the degradation process of the interface is required. In a previous paper, we presented a quantitative method to predict the adhesion lifetime of DLC coatings on CoCrMo biomedical alloy in phosphate-buffered saline (PBS) [11]. The method was based on the Rockwell C indentation and consisted of measuring the speed of film delamination (v) around the crater and plotting it against the strain energy release rate (G) obtained from calculations using the finite-element method (FEM) (Fig. 1). It was shown that the delamination of DLC coatings from the CoCrMo substrate follows the laws of stress–corrosion cracking (SCC), with a threshold (GTH) and a “stage 1” crack propagation (Fig. 1b) [13]. Having established a quantitative method to assess the adhesion lifetime of protective coatings on ductile implant alloys in terms of the delamination due to SCC, we present in this work a structural and chemical analysis of the interfacial region between the DLC coating and the CoCrMo substrate. The exact location of the coating delamination in a biological environment is discussed in connection with corrosion phenomena occurring in the reactive interface layer formed at the film–implant interface. 2. Materials and methods The DLC layers were deposited on a CoCrMo biomedical implant alloy using an radiofrequency plasma-activated chemical vapor deposition (PACVD) method with acetylene (C2H2) as process gas. A more detailed description of the deposition process and of the cleaning stages of the CoCrMo substrates can be found in Ref. [11]. To
investigate the delamination process and find the causes of coating failure, DLC coatings with pure metal interlayers of Co and Mo were deposited; these interlayers were grown by means of a DC magnetron sputtering process using Ar as process gas. The SCC curves were determined based on the experimental procedure described in detail in Ref. [11]. As a corrosive medium, 0.01 M PBS with pH 7.4 (Sigma–Aldrich, Switzerland) at a constant temperature of 37 °C was used. The chemical composition at the interface was determined by measuring sputter-depth X-ray photoelectron spectroscopy (XPS) profiles with a Physical Electronics (PHI) Quantum 2000. Material was etched away from the surface by a 2 keV Ar+ beam in between consecutive analysis of the elemental concentrations. Since depth resolution is reduced by the surface roughening induced by sputter-depth profiling, the interface chemistry was analyzed on a dedicated XPS sample (15 nm DLC/CoCrMo) grown in similar conditions as the thick coatings used to investigate the corrosion stability in biological fluid. The microstructure of the DLC films was determined by scanning electron microscopy (SEM) with a Philips XL30 ESEM–FEG equipped with an X-ray detector for energy dispersive X-ray (EDX) analysis. Lamellae for transmission electron microscopy (TEM) investigation were prepared with a FEI STRATA DB 235 dual beam (DB) focused-ion beam (FIB) workstation. The microstructure of the interfacial region of the coated samples was studied with a conventional JEOL JEM2200FS microscope operating at 200 keV, equipped with a field emission gun and an in-column filter. The high-resolution (HR) TEM images were zero-loss filtered to ameliorate their quality. Besides micrographs, compositional analyses were performed in bright-field scanning/transmission electron microscopy (STEM) mode with a JEOL EDX detector having a 5 cm2 window. The beam diameter during the analysis
4680
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
was 0.5 nm. Surface topography of the samples was investigated by means of a stylus profilometer (P-10, KLA Tencor) and a Reichert-Jung MeF3 metallographic microscope equipped with a Leica DC500 CCD camera and side illumination. Optical micrographs were obtained by all-in-focus image reconstruction by merging multiple differentially focused images recorded at several heights. 3. Results and discussion 3.1. Evidence and structural characterization of the reactive interface layer 3.1.1. HRTEM and STEM analyses Fig. 2a shows an optical micrograph top-view of the Rockwell indentation in a 4 lm DLC-coated CoCrMo substrate after immersion in PBS for 20 days; the inset presents the Rockwell pattern immediately after performing the indentation. It is clear that, although in the beginning the coating exhibited almost no debonding around the indentation, after immersing the sample for 20 days the delamination area has considerably grown. Such progressive, environmentally assisted delamination was quantified in a previous publication using time-dependent monitoring of the crack front position and numerical calculations based on the FEM [11]. In order to find the exact location of the coating delamination, a transverse FIB cut was performed perpendicular to the crack front (Fig. 2a) and the delamination crack was followed along the DLC–CoCrMo interface. Near the end of the crack a TEM lamella (12 lm 7 lm) was prepared and a central electron-transparent area (5 lm 7 lm 80 nm) was polished such that the crack tip lies inside this window (Fig. 2b). In Fig. 3a, which shows a TEM micrograph of an undamaged zone of the DLC–CoCrMo interface region, a very thin transition region is observed between the coating and the substrate as a medium gray area. This region corresponds to the
reactive interface layer formed in the first few seconds of the deposition process, and is mainly due to the reaction between metal and carbon atoms, as shown later in this paper. In order to perform a numerical analysis of the image contrast, a white rectangular box (length 43.1 nm, width 9.6 nm) was overlaid on the micrograph perpendicular to the DLC–CoCrMo interface, and the rows of pixels with gray levels ranging from 0 (total black) to 255 (total white) were averaged over the rectangle width. The resulting averaged row was assigned to an electron transmission ranging from 0% to 100%; the width of the interface was estimated by determining the full-width at half maximum (FWHM) of the first derivative of this transmission and was found to be approximately 7 nm. However, the numerical analysis of the TEM image contrast in Fig. 3a described above might provide an inaccurate estimate of the interface thickness, since this contrast can be influenced by the quality of lamella polishing, non-homogeneity of the CoCrMo substrate, as well as sample tilt during the TEM investigation. Therefore, HRTEM investigation of the interface region was carried out to obtain additional and more accurate information, as shown in Fig. 3b, where rows of atoms corresponding to a substrate grain are visible. Fig. 3c and d depict the selected-area electron diffraction (SAED) patterns for the substrate and DLC coating, respectively. It was observed that while the SAED pattern of the DLC film exhibited only very blurred halos, typical of an amorphous structure, the SAED pattern of the substrate showed the coexistence of the face-centered cubic (fcc) and hexagonal close-packed (hcp) phases associated with the fcc ! hcp transformation in the wrought CoCrMo alloy [14]. A close inspection of the interface region in Fig. 3a revealed the presence of an amorphous region 4.6 nm thick between the darker, polycrystalline substrate, and the lighter amorphous DLC film. However, as the TEM lamella was 80 nm thick and the direction of the electron beam could have been slightly off with respect to the normal to the surface of the TEM
Fig. 2. (a) Optical micrograph top-view of the Rockwell indentation in a 4 lm thick DLC-coated CoCrMo substrate after immersion in PBS for 20 days. The inset depicts the Rockwell pattern immediately after the indentation. (b) SEM micrograph of the TEM lamella prepared by FIB at the location indicated in (a).
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
4681
Fig. 3. TEM analysis of the interface layer between DLC and CoCrMo: (a) TEM micrograph and numerical analysis of the contrast across the interface; (b) HRTEM micrograph of the interface layer; (c and d) selected-area electron diffractions of the substrate and coating.
lamella during the measurements, the real value of the interface thickness might be lower. Moreover, careful visual inspection of the interface region could not fully exclude any local ordering, such as metal carbide clusters. The visibility of crystalline clusters in the nanometer range, however, could be obstructed by the thick (80 nm) lamella; therefore, several tens of clusters would have to be oriented in the same direction towards the electron beam, which is of course highly improbable. The reactive interface layer observed in the TEM images is inherent to coatings deposited on substrates by the PACVD process, in which interface atomic mixing takes place from the very first moment of film growth. Thus, during the in situ sputter cleaning process step using Ar+ ions with an energy of 600 eV, substrate atoms were not only removed from the substrate, but were also redeposited to a certain extent. Furthermore, at the beginning of the film deposition, when Ar gas was exchanged for C2H2, previously sputtered substrate atoms were redeposited onto the surface together with various ionic species (C+, H+ or CH+, etc.). Bombardment at an energy of 600 eV (the energy of the ion species can be lower due to the gas collisions) produced an additional type of mixing which resulted in a broadening of the interface layer thickness. Hence, the reactive interface layer was formed by redeposition and ion implantation during the first few seconds of the growth process. The depth penetration of ion species (C+, H+ or CH+, etc.) in CoCrMo can be estimated using TRIM simulations (Version SRIM 2006) [15]. It was found that the 600 eV C+ ions were implanted into CoCrMo to an average depth of 2 nm. However, some C+ ions were implanted at greater depths: e.g. 45% of the total incident ions are implanted deeper than 2 nm, 18% are implanted
deeper than 3 nm, whereas only 3.5% are implanted deeper than 4 nm. The H+ ions were implanted much deeper (e.g. 8 nm average depth), whereas heavier ion species, such as CH+, were implanted at shallower depths than C+ ions. However, since H+ ions are much lighter, they should not contribute greatly to the amorphization of the substrate and, therefore, to the formation of the reactive interface layer. In conclusion, TRIM calculations estimate that the contribution of ion implantation to the thickness of the reactive interface layer is about 2–3 nm. If one considers also the redeposition of the atoms sputtered immediately before the Ar ! C2H2 gas exchange, a thicker interface layer is expected, which is consistent with the value of 4.6 nm obtained from HRTEM. In order to assess the chemical composition of the interface, a quantitative EDX analysis with a 0.5 nm probe was performed in bright-field STEM mode over a distance of 20 nm across the DLC–CoCrMo interface (Fig. 4), as depicted in Fig. 4a by the small white dots. A decrease in the Co, Cr and Mo signals from the bulk values, e.g. 67, 27 and 6 at.%, respectively, correlating with an increase of the C signal, was observed as the beam was moved across the interface (Fig. 4b). However, this decrease was not very abrupt, i.e. there was still up to 20 at.% Co and 10 at.% Cr detected in the DLC film at 8–10 nm from the middle of the reactive interface layer. Thus, either the interface was much broader than estimated earlier from HRTEM or, more probably, the interface region was contaminated by the Ga+ ion polishing performed during TEM lamella preparation. Nevertheless, it is interesting to note that apparently there was an enrichment of Mo, and to a lesser extent of Cr, across the interface, since the Co/Mo, Cr/Mo and Co/Cr concentration
4682
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 4. STEM–EDX analysis of the interface layer between DLC and CoCrMo: (a) TEM micrograph of the region analyzed indicating with white dots the positions where the EDX scans were performed; (b) concentration vs. distance for Co, Cr, Mo and C; (c) concentration ratios vs. distance.
ratios decreased from the corresponding bulk values of 13.6, 5.2 and 2.6, respectively (Fig. 4c). This suggests that at the interface the chemical bonds between Mo and C, and Cr and C are favored. 3.1.2. XPS analysis Depth profile XPS was employed to obtain more detailed information about the chemical status (i.e. binding states) of the interfacial layer between DLC and CoCrMo that STEM–EDX was not able to reveal. However, since this technique is not accurate enough to analyze a 4 lm thick DLC coating—because the depth resolution is diminished by sputter-induced surface roughening to about 5% of the sputtered depth—the interface chemistry was analyzed on a dedicated XPS sample (15 nm DLC/CoCrMo). This sample was grown in the same conditions as the thick coatings used to investigate the corrosion stability in biological fluid. The measured depth-profile XPS spectra were separated into two, appropriately chosen chemical binding states, the states of the pure elements (e.g. C for pure DLC and Co, Cr and Mo for pure metal substrate) and their corresponding carbides, e.g. CoCrMo–Cx, Co–Cx, Cr–Cx and Mo–Cx. In the case of carbon, the combination of Co–Cx, Cr–Cx and Mo–Cx was considered for simplicity as one chemical state, namely CoCrMo–Cx. The decomposition was performed by spectral fitting using as a base the signals from the pure DLC, pure metallic CoCrMo substrate and the signals measured for CoCrMo–carbide reference films as a base. The impact of different fitting procedures on spectral decomposition has been analyzed in detail for the DLC/CoCrMo system in a separate publication [16].
Separation of the sputter spectra into two contributions, corresponding to pure elements and elemental carbides, was essential, since the sputter depth-profile XPS spectra in Figs. 5a and 6a revealed the presence of an additional chemical state at the DLC–CoCrMo interface. This chemical state was noticeable by a shift towards lower binding energies for C 1s signal (Fig. 5c and d) and higher binding energies for Mo 3d signals (Fig. 6c and d); for Co 3p and Cr 3p signals, the observed shifts were not so prominent. In the spectral decomposition, each XPS spectrum for a certain element was modeled using three Gaussian subcomponents whose intensities were the main fitting parameters as shown in Figs. 5c and d and 6c and d. During the fitting procedure, the intensity ratio (I1:I2:I3) and binding energy shifts (De21 = e2 e1, De31 = e3 e1) of these subcomponents for the pure element and metal carbide states were kept constant. Fig. 7 displays the relative concentration vs. sputtering time for all metal carbide components decomposed from the sputter depth profiles, e.g. CoCrMo–Cx, Co–Cx, Cr–Cx, Mo–Cx. On the same graph the sum (Co–Cx) + (Cr–Cx) + (Mo–Cx) of the metal carbides obtained from the sputter spectra corresponding to pure metals is also represented. The solid lines represent the Gaussian fits of the experimental data. The relative concentration of CoCrMo–Cx and (Co–Cx) + (Cr–Cx) + (Mo– Cx) reached maxima at slightly different sputtering times (6.9 min vs. 7.2 min), probably due to small but inevitable errors in the decomposition procedure. Furthermore, since the maximum of the relative concentration of the CoCrMo–Cx component is only about half that of the (Co–Cx) + (Cr–Cx) + (Mo–Cx) peak, the average
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
4683
Fig. 5. Sputter depth profile XPS analysis for carbon: (a) C 1s sputter depth profile mapping (contour interval is 5% of the peak height); (b) profile decomposition into pure DLC and CoCrMo–Cx; (c and d) modeling of the XPS spectra recorded at sputtering times of 3.1 and 7.8 min (horizontal lines 1 and 2 in (a)), corresponding to CoCrMo–Cx and pure DLC.
Fig. 6. Sputter depth profile XPS analysis for molybdenum: (a) Mo 3d sputter depth profile mapping (contour interval is 5% of the peak height); (b) profile decomposition into pure Mo and Mo–Cx; (c and d) modeling of the XPS spectra recorded at sputtering times of 6 and 9.3 min (horizontal lines 1 and 2 in (a)), corresponding to pure Mo and Mo–Cx.
metal:carbon stoichiometry seems to be close to 2:1, e.g. Me2C. However, in view of the limited depth resolution of the sputter XPS method, the stoichiometry of the CoCrMo–carbide present at the interface cannot be determined accurately. Hence, it could be that an average Me2C stoichiometry is present only in the center of the reactive interface layer, whereas towards the substrate different stoichiometries exist. Moreover, since the peak metal carbide concentration ratios (Co–Cx):(Mo–Cx) = 10.6, (Cr–Cx):(Mo–Cx) = 3.6 and (Co–Cx):(Cr–Cx) = 2.9 are similar to the data shown in Fig. 4c obtained by STEM–EDX, a decrease of Co and an increase of Mo and Cr concentrations in the reactive interface layer with respect to the relative concentration in the bulk of the CoCrMo substrate is possible.
The thickness of the carbide interface layer was estimated as the FWHM of the Gaussian fitted to all carbide components combined together, e.g. (Co–Cx) + (Cr– Cx) + (Mo–Cx), and required 1.3 min etch time. Based on the obtained average carbide stoichiometry at the interface between DLC and CoCrMo the corresponding sputter rate was estimated to be 3.9 nm min1 [16]. Consequently, the thickness of the reactive interface layer is 5 nm, which is in very good agreement with the value obtained from HRTEM. However, it should be taken into account that the depth profile XPS method can overestimate the determined thickness of the reactive interface layer by 1– 2 nm through mixing processing occurring in sputter-depth profiling, such as recoil mixing and radiation-enhanced diffusion [17].
4684
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 7. Atomic concentration vs. sputtering time for metallic carbide components decomposed from the sputter depth-profile XPS spectra.
3.2. Instability of the reactive interface layer in PBS 3.2.1. Exact location and cause of the coating failure Previously, it was shown that, if the load induced at the interface (described by the strain energy release rate, G / d r2 , where d and r are the thickness and residual stress of the coating) is higher than the threshold for crack initiation in corrosive medium (GTH), then cracks will propagate with a velocity described by the SCC laws [11]. In particular, for a 4 lm thick DLC coating (r = 3.8 GPa, compressive) the strain energy release rate is 133 J m2, which is more than twice the threshold value for crack initiation in PBS at 37 °C of 60 J m2. Hence, in the absence of external loads this coating will delaminate in PBS at 37 °C with a velocity of 47 lm day1. However, it should be mentioned that at typical DLC deposition rates of 30–50 nm min1 the reactive interface layer, evidence of which was revealed in the previous section, is formed within the first 6–10 s of the deposition. If, during this relatively short time, gases such as O2 or H2O are present in high concentration, then the formation of metal carbides is hindered in favor of metal oxides. Consequently, the threshold strain energy release rate (GTH) and adhesion lifetime of the coating can be greatly reduced for these contaminated interfaces. The objective of this section is to show where the exact location of the coating failure is in PBS at 37 °C, and to reveal what causes this failure. For this purpose, the same lamella used to investigate the reactive interface layer, described in the previous section, was analyzed by TEM, but this time in a region where coating is delaminated, i.e. near the crack tip. While the SEM image in the inset of Fig. 2b already suggested that the crack responsible for coating delamination appears at the interface between the DLC film and the CoCrMo substrate, the TEM micrograph in Fig. 8a clearly demonstrates that this crack occurs
Fig. 8. Corroded CoCrMo carbide interface in PBS at 37 °C: (a) TEM view of the interface crack; (b) STEM–EDX chemical analysis across the interface.
approximately in the middle of the reactive interface layer, leaving an 2–3 nm thick medium-contrast strip on each side of the crack. Similarly, other authors revealed by TEM analysis that the mechanical failure in air of the DLC coatings deposited by ion implantation on aluminum alloy (A5052) substrates occurs due to the breaking of the mixing layer of oxide and DLC at the interface [18]. However, in order to be confident about our results, STEM– EDX analysis was performed across the delaminated interface (Fig. 8b). The results show a decrease in the Co, Cr, Mo signals across the first medium gray strip followed by an approximately constant level of these signals corresponding to the open crack. Finally, a slight increase over the second medium gray strip was observed, indicating that the Co, Cr and Mo elements were present there in a higher concentration due to the metal carbide bonds. The reason that Co, Cr and Mo signals did not cancel out completely in the open crack was probably the contamination of the thin underlying carbon foil which supports the TEM lamella. Thus, it is evident that the metal carbide reactive interface layer was unstable in PBS at 37 °C.
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
An important cause of the coating delamination could be the instability of the Co carbides, which are present in large concentration (38 at.%) at the interface between DLC and CoCrMo as revealed by the XPS results. For example, in the case of Co2C at 300 K, the reported Gibbs energy of formation D , which is a measure of the binding energy, is 13.90 kJ mol1 (0.14 eV), showing that formation of this carbide is unfavorable [19]. On the contrary, since for Mo2C, Cr3C2 and Cr7C3 the reported values of D are 54 kJ mol1 (0.56 eV), 86.27 kJ mol1 (0.89 eV) and 166.16 kJ mol1 (1.72 eV), respectively, the formation of these carbides is favorable [20]. However, to obtain more insight into the delamination process, several DLC coatings with 90 nm Co and Mo interlayers, and thicknesses between 2 and 8 lm, were deposited on the CoCrMo substrates. These DLC coatings were grown in similar conditions of bias voltage and chamber pressure as the coating grown directly on CoCrMo, but for Co and Mo interlayers DC magnetron sputtering was used. Fig. 9 compares the (G, hvi) curves of the DLC/CoCrMo, DLC/Co/CoCrMo and DLC/Mo/CoCrMo samples. For DLC/CoCrMo, a typical SCC with a threshold and a “stage 1” crack propagation with a slope parameter m = 3.0 was obtained. Low slope parameters in the range of 2–10 are typical for SCC of ductile materials [13]. However, the threshold strain energy release rate (GTH) for the DLC/Mo/CoCrMo sample was 170 J m2, which is almost 3 times larger than for the DLC/CoCrMo sample. The DLC/Co/CoCrMo sample displayed a “stage 1” crack propagation with a slope parameter m = 3.3. However, GTH for this sample could not be estimated accurately, since the stress induced at the interface for a 2 lm DLC coating is too large and therefore thinner coatings should have also been tested. Nevertheless, one could compare the adhesion strength of these samples by looking at the
Fig. 9. (G, hvi) diagrams for DLC/CoCrMo, DLC/Co/CoCrMo and DLC/Mo/CoCrMo samples immersed in PBS at 37 °C.
4685
cross-sections G (hvi = const.) and hvi(G = const.) through the SCC curves, indicated in Fig. 9 by the horizontal and vertical dotted lines. Consequently, for G = 400 J m2 the average speed of delamination for the DLC/CoCrMo sample is 11 times larger than for the DLC/Mo/CoCrMo sample and 4.5 times smaller than for DLC/Co/CoCrMo sample. Similarly, for hvi = 100 lm day1, the strain energy release rate for the DLC/CoCrMo sample is 2.5 times larger than for DLC/Co/CoCrMo and 5 times smaller than for DLC/Mo/CoCrMo sample. Therefore, from these findings it is obvious that the adhesion of DLC on CoCrMo is greatly improved by using a Mo interlayer. Conversely, the adhesion of DLC on CoCrMo is greatly diminished with a Co interlayer. Moreover, the analysis of the XPS sputter depth profiles for the DLC/Co/CoCrMo sample after coating delamination revealed that the 90 nm thick Co interlayer was profoundly corroded. The progressively delaminating coated sample was removed first from PBS, cleaned with deionized water in an ultrasound bath, and then an area from which the coating had delaminated was sputtered through the Co layer. Consequently, oxygen was present only at the surface but, when the sputtering was performed on previously delaminated regions that had been exposed to PBS for longer times, oxygen was also found in the Co interlayer; the thickness of the oxide layer increased with increasing exposure time to PBS. Thus, pure Co was found to be unstable in PBS and gradually transformed into Co oxide. 3.2.2. Corrosion mechanisms involved during the coating delamination Having established that Co carbide was the main cause for the failure of the DLC coating on the CoCrMo alloy, it is nevertheless imperative to identify and understand the mechanisms via which delamination occurs. In order for coating delamination from CoCrMo to occur, it is necessary for the corrosive medium to reach the interface region. Since DLC coatings deposited in an unclean room environment usually exhibit many defects that traverse their entire thickness down to the substrate, such as pinholes that originate from dust particles [21] and concentrations of stress at surface imperfections (e.g. inclusions at the surface) [22], this condition is fulfilled. Fig. 10a and b illustrate the 3D surface profiles of such pinholes through a 4 lm DLC coating deposited on CoCrMo; for a better visual inspection the vertical dimensions are exaggerated. These pinholes are 4 lm deep and their lateral size varies from a few tens of nanometers up to several tens of microns. Fig. 10c shows a top view optical micrograph of the same coating exposed to PBS at 37 °C for approximately 1 day obtained by all-in-focus image reconstruction. After the breaking of the interface bonds around the pinhole by corrosion mechanisms that will be discussed below, the coating buckled due to its internal compressive stress, since it was no longer adhering to the substrate. The buckling of the coating was revealed by the circular dome with a radius of 45 lm illuminated by the side microscope light source.
4686
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 10. Pinhole initiated corrosion: (a and b) 3-D profiles of pinholes in a 4 lm thick DLC-coated CoCrMo sample; (c) optical micrograph top view showing the buckling of the coating in PBS.
Similar film buckling was also observed in the 3-D scan shown in Fig. 10b. In addition, radial cracks with lengths in the range of 10–15 lm were observed around the central pinhole; these cracks were caused by the elastic energy stored in the buckled coating which exceeded the fracture toughness of the film. The crack front under the coating will propagate outwards, as indicated by the radial arrows, until the coating delaminates, leaving behind an almost circular pit, such as the one shown in Fig. 10b with a lateral size of 200 lm.
Fig. 11 schematically displays how micrometer-sized defects, such as pinholes, gradually lead to the delamination of the DLC coating. In the beginning, the PBS liquid entering through the pinhole (Fig. 11a) starts to locally corrode the reactive interface layer, forming a small crevice under the DLC coating (Fig. 11b). Similar schematic diagrams were proposed by other authors [23,4]. In addition to the dissolution of the reactive interface layer, an etching process of the CoCrMo substrate was also observed for some pinholes, revealed by small pits in the metal substrate
Fig. 11. Simple schematics of the proposed gradual delamination mechanism of the DLC coatings from CoCrMo.
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
at the location of the pinholes. The formation of such pits is usually attributed to “passive film breakdown”, i.e. interaction of some aggressive ions (e.g. Cl) with the protective surface oxide film at locations where it is defective or weak [1,24]. Due to the gradual dissolution process of the reactive interface layer, the advancing interface crack front induces buckling of the coating (Fig. 11c) until the elastic energy stored in the coating exceeds the fracture toughness and the coating breaks leaving behind a crater (Fig. 11d). In PBS, this SCC process driven by the corrosive attack at the crack front tip and residual stress of the coating may continue until the whole DLC coating delaminates from the CoCrMo substrate. Simple schematics of a possible SCC mechanism occurring in the reactively formed metal carbide interface during the crack front expansion are proposed in Fig. 11e and f, in which the crack growth is related to oxidation reactions at the crack tip as described by the slip-oxidation and enhanced surface mobility models [25,26]. However, one has to make the distinction between the classic slip-oxidation model, in which the oxide film ruptures at the location of the bulk metal grain boundary that subsequently dissolves as the crack advances, and the present case, in which ruptures occur in the middle of a metal carbide interface layer. According to the enhanced surface mobility model, crack growth occurs due to the capture of vacancies by the stressed lattice at the crack tip. Thus, the transport of atoms takes place by surface diffusion from the highly stressed crack tip to the less stressed regions. In addition, the corrosive medium influences the crack growth rate by affecting the dissolution and surface diffusion of the atoms. After the breaking of the unstable Co carbide bonds, the oxide-free surface of the crack tip becomes exposed to the PBS solution and positively charged Co2+ ions migrate from the metal into the solution, leaving electrons behind according to the anodic reaction: Co ! Co2+ + 2e [27]. In neutral PBS solution the electrons produced can react with oxygen molecules adsorbed in the crevice from air dissolved in solution, according to the cathodic reaction O2 + 2H2O + 4e ! 4OH [1,27]. In the crevice formed at the interface, however, the pH can reach acidic values because of hydrolysis reactions such as Co2+ + 2H2O ! Co(OH)2 + 2H+ [1,27,28]. Acidification of the fluid in the crevice could also take place as a result of the anodic reactions Co + H2O ! CoO + 2H+ + 2e, Co + 2Cl ! CoCl2 + H2O ! CoO + 2H+ + 2Cl [27,28]. Consequently, the residual electrons can react with hydrogen ions, adsorbed on the metal surface from the solution, to produce hydrogen gas by the cathodic reaction 2H+ + 2e ! H2" [1]. It should be mentioned that acidification depends strongly on the exchange of the fluid inside the crack with the environment and therefore on the crack geometry. The last reaction is also relevant for in vivo conditions immediately after implantation when the conditions are most aggressive and the biological fluid becomes acidic
4687
[29]. However, it is unknown whether the H2 gas produced forms small bubbles and leaves the crevice, or if it penetrates into the reactive interface layer, inducing coating delamination through hydrogen embrittlement. 3.2.3. Effects of the PBS temperature on the coating delamination Since crack openings are usually much smaller than the pinhole dimensions as shown in Figs. 8a and 10a and b, reaction-limiting processes of the corrosion at the crack tip could play an important role in coating delamination. Hence, SCC curves for DLC/CoCrMo were determined for temperatures of 4, 23, 37, 66 and 95 °C. All curves were fitted with the reaction-controlled model [30]. The corresponding values of the GTH and crack-growth exponent m for the “stage 1” SCC are plotted in Fig. 12a. An exponential decrease in the GTH was observed, whereas the crack-growth parameter m exhibits a slight linear increase with temperature. Growth of the value of m with increasing temperature suggests that the solution is slightly more aggressive [13] or that the material properties at the crack tip are slightly modified with increasing temperature. Moreover, the decrease of GTH with increasing temperature suggests that a diffusion process of the corrosive medium to the crack tip took place. Thus, the higher the temperature, the higher the solution exchange at the crack tip, and therefore the higher the dissolution efficiency of the interface material. Indeed, by plotting the average delamination velocity hvi in PBS vs. the inverse temperature T for an interface load of 133 J m2 corresponding to a 4 lm thick DLC coating far away from the Rockwell indentation, a typical behavior described by the Arrhenius equation hvi ¼ A expðEa =RT Þ was observed up to 37 °C [31]. This suggests that at least at temperatures below 37 °C the crack growth is a reaction-limited process. In the Arrhenius equation, Ea is the activation energy, i.e. the minimum amount of energy necessary for the SCC reaction, A is the pre-exponential factor equivalent to the speed of delamination when Ea = 0, and R is the universal gas constant. It should be also noted that in the Arrhenius equation Ea and A were considered as being independent of the temperature, which is not always the case, but for a small range of temperatures it is usually a good approximation. Consequently, values of 13 ± 7 lm s1 and 26.5 ± 0.6 kJ mol1 were obtained for the pre-exponential factor and activation energy, respectively. However, from Fig. 12b it is also evident that for temperatures higher than 37 °C, the crack growth behavior deviates from the linear dependence, which suggests that the rate-limiting factor at these temperatures has different causes (special chemical reactions, variation of the pH value of the corrosive media with the temperature, mechanical stability and ductility of the interface material, atypical diffusion processes, variation of the oxide layer thickness, etc.). The gradual decrease in the slope of the Arrhenius plot with increasing temperature could also be an indication for the temperature variation of
4688
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 12. Temperature dependence of the SCC-induced coating delamination in PBS: (a) m and GTH vs. T; (b) Arrhenius plot for the temperature range of 4–95 °C and G = 133 J m2.
the pre-exponential factor (e.g. A ¼ BT n , where B and n are constants). Consequently, the modified Arrhenius equation hvi ¼ BT n expðEa =RT Þ could be used to describe the temperature variation of the SCC-induced coating delamination [31]. Using this new equation to fit the experimental data points for the temperature range from 4 to 95 °C, a value of 430 kJ mol1 was obtained for the activation energy. However, since this value is much larger than the previously reported one, of 26.5 ± 0.6 kJ mol1, which was obtained using the linear dependence for temperatures below 37 °C, one should consider it with extreme caution. First of all, in the modified Arrhenius equation the temperature dependence of the activation energy was neglected. Furthermore, in view of the complex system one deals with, of a multielement interface in a dynamic corrosive environment, where a multitude of chemicophysical processes may play a role, it could be misleading merely to treat the delamination process as a chemical reaction described by an Arrhenius behavior. Nevertheless, in general, if coating delamination is driven by SCC, then a similar analysis could be applied to different coating/substrate combinations to determine the activation energy for SCC. 4. Conclusions In the present study the exact microstructure of the interface between DLC thin films and CoCrMo biomedical implants, as well as the delamination processes occurring in PBS solution, were investigated, with the following findings:
2.
3.
4.
5.
of 5 nm, has a disordered amorphous-like structure. Moreover, sputter depth profile XPS measurements revealed its average Me:C stoichiometry to be 2:1. It was demonstrated that in the absence of any interlayer the DLC delaminates from the CoCrMo due to the instability of the reactive interface layer in PBS at 37 °C. Moreover, HRTEM revealed that the crack propagation occurs roughly in the middle of this layer. The failure of the DLC on CoCrMo was shown to be due to the metastable Co carbides (38 at.% from XPS) formed in the reactive interface layer during the film deposition, since by growing DLC layers with a 90 nm Co interlayer for the same load (G = 400 J m2) the average speed of delamination was increased by a factor of 4.5. However, by growing layers with a 90 nm Mo interlayer, the average speed of delamination was decreased by a factor 11, and the threshold for SCC was increased from 60 to 170 J m2. It was shown that delamination is initiated by the corrosive medium that penetrates the pinholes traversing the DLC layer. Subsequently, the delamination continues due to diffusion processes in the crevice formed in the carbide interface layer and chemical reactions at the crack tip. In addition, by varying the temperature of the PBS it was shown that up to temperatures of 37 °C the SCC-induced coating delamination follows an Arrhenius behavior with thermal activation energy of 26.5 kJ mol1; for higher temperatures deviations from the linear Arrhenius behavior were observed.
Acknowledgments 1. A reactive interface layer between the DLC coating and CoCrMo substrate was revealed and investigated in detail by structural and chemical analysis. HRTEM measurements showed that this layer, with a thickness
We would like to thank R. Crockett, P. Schmutz and B. Lehman for useful discussions and comments, and S. Meier for preparing the lamellae for TEM investigations. This
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
work was supported by the Swiss Commission for Technology and Innovation (CTI) and Synthes GmbH. References [1] Kruger J. Fundamental aspects of corrosion of metallic implants. In: Syrett B, Acharya A, editors. Corrosion and degradation of implant materials. ASTM STP 684; 1979. [2] Greenfield EM, Bi Y, Ragab AA, Goldberg VM, Nalepka JL, Seabold JM. J Biomed Mater Res Part B: Appl Biomater 2005;72B:179. [3] Jacobs JJ, Skipor AK, Paterson LM, Hallab NJ, Paprosky WG, Black J, et al. J Bone Joint Surg Am 1998;80:1447. [4] Lappalainen R, Santavirta SS. Clin Orthop Relat Res 2005;430:72. [5] Tiainen VM. Diamond Relat Mater 2001;10:153. [6] Hauert R. Tribol Int 2004;37:991. [7] Bruinink A, Schroeder A, Francz G, Hauert R. Biomaterials 2005;26:3487. [8] Hauert R. DLC films in biomedical applications. In: Donnet C, Erdemir A, editors. Tribology of diamond-like carbon films: fundamentals and applications. Berlin: Springer-Verlag; 2008. [9] Ta¨ger G, Podleska LE, Schmidt B, Ziegler M, Nast-Kolb D. Materialwiss Werkst 2003;34:1094. [10] Joyce TJ. Wear 2007;263:1050. [11] Falub CV, Thorwarth G, Affolter C, Mu¨ller U, Voisard C, Hauert R. Acta Biomater 2009;5:3086. [12] Thorwarth G, Falub CV, Mu¨ller U, Weisse B, Voisard C, Tobler M, et al. Acta Biomater 2010;6:2335. [13] Jones RH, Ricker RE. Mechanisms of stress–corrosion cracking. In: Jones RH, editor. Stress–corrosion cracking. Materials Park (OH): ASM International; 1996.
4689
[14] Rajan K. Metall Trans A 1982;13A:1161. [15] Biersack JP, Haggmark LG. Nucl Instrum Methods 1980;174:257. [16] Mu¨ller U, Falub CV, Thorwarth G, Voisard C, Hauert R. Acta Mater 2011;59:1150. [17] Littmark U, Hofer WO. The theory of recoil mixing in solids. In: Oechsner H, editor. Topics in current physics, thin films and depth profile analysis, 37. Berlin: Springer-Verlag; 1984. [18] Oka Y, Nishijima M, Hiraga K, Yatsuzuka M. Nucl Instrum Methods Phys Res B 2007;257:702. [19] Seng WF, Barnes PA. Mater Sci Eng B 2000;76:225. [20] Barin I. Thermochemical data of pure substances, 2nd ed., vols. 1 and 2; 1989. p. 921, 560, 561. [21] Pranevicius L. Vapor deposition coating technologies. In: Tracton AA, editor. Coatings technology handbook. Boca Raton (FL): CRC Press; 2006. [22] Outka DA, Hsu WL, Boehme DR, Yang NYC, Ottesen DK, Johnsen HA, et al. Compilation of diamond-like carbon properties for barriers and hard coatings. Sandia report, SAND94-8219, UC-404; 1994. [23] Chandra L, Allen M, Butter R, Rushton N, Lettington AH, Clyne TW. Diamond Relat Mater 1995;4:852. [24] Hoar TP, Jacob WR. Nature 1967;216:1299. [25] Andresen PL. Corrosion 1993;49:714. [26] Galvele JR. Corros Sci 1987;27:1. [27] Deltombe E, Zoubov N, Pourbaix M. Section 12.2: cobalt. In: Pourbaix M, editor. Atlas of electrochemical equilibria in aqueous solutions. Oxford: Pergamon Press; 1966. [28] Gallant D, Simard S. Corros Sci 2005;47:1810. [29] Lopez DA, Duran A, Cere´ SM. J Mater Sci Mater Med 2008;19:2137. [30] Cook RF, Liniger EG. J Am Ceram Soc 1993;76:1096. [31] McNaught AD, Wilkinson A, editors. IUPAC compendium of chemical terminology, 2nd ed. Oxford: Blackwell Science; 1997.
Acta Materialia 59 (2011) 4678–4689 www.elsevier.com/locate/actamat
In vitro studies of the adhesion of diamond-like carbon thin films on CoCrMo biomedical implant alloy C.V. Falub a,b,⇑, U. Mu¨ller a, G. Thorwarth a,c, M. Parlinska-Wojtan a, C. Voisard c, R. Hauert a b
a ¨ berlandstrasse 129, CH-8600 Du¨bendorf, Switzerland Empa, Swiss Federal Laboratories for Materials Science and Technology, U Swiss Federal Institute of Technology (ETHZ), Laboratory for Solid State Physics, Schafmattstrasse 16, CH-8093 Zu¨rich, Switzerland c Synthes GmbH, Langendorfstrasse 2, CH-4513 Langendorf, Switzerland
Received 9 October 2010; received in revised form 4 April 2011; accepted 6 April 2011 Available online 30 April 2011
Abstract Diamond-like carbon (DLC) coatings are promising candidates for improving the mechanical wear and chemical resistance of loadbearing implants. However, in view of medical data showing severe problems due to film delamination, the application of DLC layers to the medicotechnical field is not yet recognized. In an extension to an earlier publication, in which we established that coating delamination in phosphate-buffered saline (PBS) solution is driven by stress–corrosion cracking, we present for the first time a detailed structural and chemical analysis of the interface region of DLC-coated CoCrMo. High-resolution and scanning transmission electron microscopy, energy-dispersive X-ray spectroscopy and X-ray photoelectron spectroscopy shows that a 5 nm thick carbide layer with an average Me:C stoichiometry of 2:1 is formed at this interface. It is shown that the failure of DLC-coated CoCrMo in PBS is due to the instability of Co carbides at the film–implant interface. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Diamond-like carbon (DLC); Interface; Corrosion; Transmission electron microscopy (TEM); X-ray photoelectron spectroscopy (XPS)
1. Introduction Corrosion products and increased wear are the major processes responsible for the aseptic loosening of orthopedic implants after relatively short periods of time [1–3]. A feasible solution to reduce the wear and to enhance the corrosion resistance, biocompatibility and mechanical properties of currently used implant materials is by surface modification techniques, such as ion implantation or deposition of ceramic layers (e.g. diamond-like carbon (DLC), TiN, ZrO2) [4]. While ion implantation has limitations related to its shallow penetration depth (0.3 lm), the ⇑ Corresponding author. Present address: Laboratory for Solid State Physics, Swiss Federal Institute of Technology (ETHZ), Schafmattstrasse16, CH-8093 Zu¨rich, Switzerland. Tel.: +41 44 633 2261; fax: +41 44 633 1072. E-mail address: [email protected] (C.V. Falub).
use of ceramic coatings improves the wear resistance of the implants, though the adhesion of these coatings to the implant is problematic [4]. In view of their excellent mechanical properties (e.g. low wear, high hardness), as well as chemical inertness and biological compatibility, the application of DLC coatings to metallic articulating biomedical implants should provide a promising solution [5–8]. However, in spite of these outstanding properties, the application of DLC films to implants has not yet achieved widespread recognition as the physiological environment has been shown to cause unpredictable delamination of these coatings from the metal surface, which may lead to the total failure of the coated implant [9,10]. DLC layers delaminate from implant surfaces due to their high residual stresses (in the GPa range) and because their interfaces with the implant become corroded in biological fluids [11]. Hence, to minimize these problems and develop successful DLC-coated
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.04.014
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
4679
Fig. 1. Adhesion lifetime prediction of the DLC-coated CoCrMo implants: (a) 3-D profile of a Rockwell indentation in a 4 lm thick DLC-coated CoCrMo substrate; (b) (G, hvi) diagram derived from optical monitoring of the delaminating crack front and finite-element calculations [11].
implants, it is necessary not only to perform simulator tests under conditions closely resembling the in vivo situation [12], but also to design biocompatible interface layers that are wear and corrosion resistant. For this purpose, a detailed understanding of the mechanisms involved in the degradation process of the interface is required. In a previous paper, we presented a quantitative method to predict the adhesion lifetime of DLC coatings on CoCrMo biomedical alloy in phosphate-buffered saline (PBS) [11]. The method was based on the Rockwell C indentation and consisted of measuring the speed of film delamination (v) around the crater and plotting it against the strain energy release rate (G) obtained from calculations using the finite-element method (FEM) (Fig. 1). It was shown that the delamination of DLC coatings from the CoCrMo substrate follows the laws of stress–corrosion cracking (SCC), with a threshold (GTH) and a “stage 1” crack propagation (Fig. 1b) [13]. Having established a quantitative method to assess the adhesion lifetime of protective coatings on ductile implant alloys in terms of the delamination due to SCC, we present in this work a structural and chemical analysis of the interfacial region between the DLC coating and the CoCrMo substrate. The exact location of the coating delamination in a biological environment is discussed in connection with corrosion phenomena occurring in the reactive interface layer formed at the film–implant interface. 2. Materials and methods The DLC layers were deposited on a CoCrMo biomedical implant alloy using an radiofrequency plasma-activated chemical vapor deposition (PACVD) method with acetylene (C2H2) as process gas. A more detailed description of the deposition process and of the cleaning stages of the CoCrMo substrates can be found in Ref. [11]. To
investigate the delamination process and find the causes of coating failure, DLC coatings with pure metal interlayers of Co and Mo were deposited; these interlayers were grown by means of a DC magnetron sputtering process using Ar as process gas. The SCC curves were determined based on the experimental procedure described in detail in Ref. [11]. As a corrosive medium, 0.01 M PBS with pH 7.4 (Sigma–Aldrich, Switzerland) at a constant temperature of 37 °C was used. The chemical composition at the interface was determined by measuring sputter-depth X-ray photoelectron spectroscopy (XPS) profiles with a Physical Electronics (PHI) Quantum 2000. Material was etched away from the surface by a 2 keV Ar+ beam in between consecutive analysis of the elemental concentrations. Since depth resolution is reduced by the surface roughening induced by sputter-depth profiling, the interface chemistry was analyzed on a dedicated XPS sample (15 nm DLC/CoCrMo) grown in similar conditions as the thick coatings used to investigate the corrosion stability in biological fluid. The microstructure of the DLC films was determined by scanning electron microscopy (SEM) with a Philips XL30 ESEM–FEG equipped with an X-ray detector for energy dispersive X-ray (EDX) analysis. Lamellae for transmission electron microscopy (TEM) investigation were prepared with a FEI STRATA DB 235 dual beam (DB) focused-ion beam (FIB) workstation. The microstructure of the interfacial region of the coated samples was studied with a conventional JEOL JEM2200FS microscope operating at 200 keV, equipped with a field emission gun and an in-column filter. The high-resolution (HR) TEM images were zero-loss filtered to ameliorate their quality. Besides micrographs, compositional analyses were performed in bright-field scanning/transmission electron microscopy (STEM) mode with a JEOL EDX detector having a 5 cm2 window. The beam diameter during the analysis
4680
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
was 0.5 nm. Surface topography of the samples was investigated by means of a stylus profilometer (P-10, KLA Tencor) and a Reichert-Jung MeF3 metallographic microscope equipped with a Leica DC500 CCD camera and side illumination. Optical micrographs were obtained by all-in-focus image reconstruction by merging multiple differentially focused images recorded at several heights. 3. Results and discussion 3.1. Evidence and structural characterization of the reactive interface layer 3.1.1. HRTEM and STEM analyses Fig. 2a shows an optical micrograph top-view of the Rockwell indentation in a 4 lm DLC-coated CoCrMo substrate after immersion in PBS for 20 days; the inset presents the Rockwell pattern immediately after performing the indentation. It is clear that, although in the beginning the coating exhibited almost no debonding around the indentation, after immersing the sample for 20 days the delamination area has considerably grown. Such progressive, environmentally assisted delamination was quantified in a previous publication using time-dependent monitoring of the crack front position and numerical calculations based on the FEM [11]. In order to find the exact location of the coating delamination, a transverse FIB cut was performed perpendicular to the crack front (Fig. 2a) and the delamination crack was followed along the DLC–CoCrMo interface. Near the end of the crack a TEM lamella (12 lm 7 lm) was prepared and a central electron-transparent area (5 lm 7 lm 80 nm) was polished such that the crack tip lies inside this window (Fig. 2b). In Fig. 3a, which shows a TEM micrograph of an undamaged zone of the DLC–CoCrMo interface region, a very thin transition region is observed between the coating and the substrate as a medium gray area. This region corresponds to the
reactive interface layer formed in the first few seconds of the deposition process, and is mainly due to the reaction between metal and carbon atoms, as shown later in this paper. In order to perform a numerical analysis of the image contrast, a white rectangular box (length 43.1 nm, width 9.6 nm) was overlaid on the micrograph perpendicular to the DLC–CoCrMo interface, and the rows of pixels with gray levels ranging from 0 (total black) to 255 (total white) were averaged over the rectangle width. The resulting averaged row was assigned to an electron transmission ranging from 0% to 100%; the width of the interface was estimated by determining the full-width at half maximum (FWHM) of the first derivative of this transmission and was found to be approximately 7 nm. However, the numerical analysis of the TEM image contrast in Fig. 3a described above might provide an inaccurate estimate of the interface thickness, since this contrast can be influenced by the quality of lamella polishing, non-homogeneity of the CoCrMo substrate, as well as sample tilt during the TEM investigation. Therefore, HRTEM investigation of the interface region was carried out to obtain additional and more accurate information, as shown in Fig. 3b, where rows of atoms corresponding to a substrate grain are visible. Fig. 3c and d depict the selected-area electron diffraction (SAED) patterns for the substrate and DLC coating, respectively. It was observed that while the SAED pattern of the DLC film exhibited only very blurred halos, typical of an amorphous structure, the SAED pattern of the substrate showed the coexistence of the face-centered cubic (fcc) and hexagonal close-packed (hcp) phases associated with the fcc ! hcp transformation in the wrought CoCrMo alloy [14]. A close inspection of the interface region in Fig. 3a revealed the presence of an amorphous region 4.6 nm thick between the darker, polycrystalline substrate, and the lighter amorphous DLC film. However, as the TEM lamella was 80 nm thick and the direction of the electron beam could have been slightly off with respect to the normal to the surface of the TEM
Fig. 2. (a) Optical micrograph top-view of the Rockwell indentation in a 4 lm thick DLC-coated CoCrMo substrate after immersion in PBS for 20 days. The inset depicts the Rockwell pattern immediately after the indentation. (b) SEM micrograph of the TEM lamella prepared by FIB at the location indicated in (a).
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
4681
Fig. 3. TEM analysis of the interface layer between DLC and CoCrMo: (a) TEM micrograph and numerical analysis of the contrast across the interface; (b) HRTEM micrograph of the interface layer; (c and d) selected-area electron diffractions of the substrate and coating.
lamella during the measurements, the real value of the interface thickness might be lower. Moreover, careful visual inspection of the interface region could not fully exclude any local ordering, such as metal carbide clusters. The visibility of crystalline clusters in the nanometer range, however, could be obstructed by the thick (80 nm) lamella; therefore, several tens of clusters would have to be oriented in the same direction towards the electron beam, which is of course highly improbable. The reactive interface layer observed in the TEM images is inherent to coatings deposited on substrates by the PACVD process, in which interface atomic mixing takes place from the very first moment of film growth. Thus, during the in situ sputter cleaning process step using Ar+ ions with an energy of 600 eV, substrate atoms were not only removed from the substrate, but were also redeposited to a certain extent. Furthermore, at the beginning of the film deposition, when Ar gas was exchanged for C2H2, previously sputtered substrate atoms were redeposited onto the surface together with various ionic species (C+, H+ or CH+, etc.). Bombardment at an energy of 600 eV (the energy of the ion species can be lower due to the gas collisions) produced an additional type of mixing which resulted in a broadening of the interface layer thickness. Hence, the reactive interface layer was formed by redeposition and ion implantation during the first few seconds of the growth process. The depth penetration of ion species (C+, H+ or CH+, etc.) in CoCrMo can be estimated using TRIM simulations (Version SRIM 2006) [15]. It was found that the 600 eV C+ ions were implanted into CoCrMo to an average depth of 2 nm. However, some C+ ions were implanted at greater depths: e.g. 45% of the total incident ions are implanted deeper than 2 nm, 18% are implanted
deeper than 3 nm, whereas only 3.5% are implanted deeper than 4 nm. The H+ ions were implanted much deeper (e.g. 8 nm average depth), whereas heavier ion species, such as CH+, were implanted at shallower depths than C+ ions. However, since H+ ions are much lighter, they should not contribute greatly to the amorphization of the substrate and, therefore, to the formation of the reactive interface layer. In conclusion, TRIM calculations estimate that the contribution of ion implantation to the thickness of the reactive interface layer is about 2–3 nm. If one considers also the redeposition of the atoms sputtered immediately before the Ar ! C2H2 gas exchange, a thicker interface layer is expected, which is consistent with the value of 4.6 nm obtained from HRTEM. In order to assess the chemical composition of the interface, a quantitative EDX analysis with a 0.5 nm probe was performed in bright-field STEM mode over a distance of 20 nm across the DLC–CoCrMo interface (Fig. 4), as depicted in Fig. 4a by the small white dots. A decrease in the Co, Cr and Mo signals from the bulk values, e.g. 67, 27 and 6 at.%, respectively, correlating with an increase of the C signal, was observed as the beam was moved across the interface (Fig. 4b). However, this decrease was not very abrupt, i.e. there was still up to 20 at.% Co and 10 at.% Cr detected in the DLC film at 8–10 nm from the middle of the reactive interface layer. Thus, either the interface was much broader than estimated earlier from HRTEM or, more probably, the interface region was contaminated by the Ga+ ion polishing performed during TEM lamella preparation. Nevertheless, it is interesting to note that apparently there was an enrichment of Mo, and to a lesser extent of Cr, across the interface, since the Co/Mo, Cr/Mo and Co/Cr concentration
4682
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 4. STEM–EDX analysis of the interface layer between DLC and CoCrMo: (a) TEM micrograph of the region analyzed indicating with white dots the positions where the EDX scans were performed; (b) concentration vs. distance for Co, Cr, Mo and C; (c) concentration ratios vs. distance.
ratios decreased from the corresponding bulk values of 13.6, 5.2 and 2.6, respectively (Fig. 4c). This suggests that at the interface the chemical bonds between Mo and C, and Cr and C are favored. 3.1.2. XPS analysis Depth profile XPS was employed to obtain more detailed information about the chemical status (i.e. binding states) of the interfacial layer between DLC and CoCrMo that STEM–EDX was not able to reveal. However, since this technique is not accurate enough to analyze a 4 lm thick DLC coating—because the depth resolution is diminished by sputter-induced surface roughening to about 5% of the sputtered depth—the interface chemistry was analyzed on a dedicated XPS sample (15 nm DLC/CoCrMo). This sample was grown in the same conditions as the thick coatings used to investigate the corrosion stability in biological fluid. The measured depth-profile XPS spectra were separated into two, appropriately chosen chemical binding states, the states of the pure elements (e.g. C for pure DLC and Co, Cr and Mo for pure metal substrate) and their corresponding carbides, e.g. CoCrMo–Cx, Co–Cx, Cr–Cx and Mo–Cx. In the case of carbon, the combination of Co–Cx, Cr–Cx and Mo–Cx was considered for simplicity as one chemical state, namely CoCrMo–Cx. The decomposition was performed by spectral fitting using as a base the signals from the pure DLC, pure metallic CoCrMo substrate and the signals measured for CoCrMo–carbide reference films as a base. The impact of different fitting procedures on spectral decomposition has been analyzed in detail for the DLC/CoCrMo system in a separate publication [16].
Separation of the sputter spectra into two contributions, corresponding to pure elements and elemental carbides, was essential, since the sputter depth-profile XPS spectra in Figs. 5a and 6a revealed the presence of an additional chemical state at the DLC–CoCrMo interface. This chemical state was noticeable by a shift towards lower binding energies for C 1s signal (Fig. 5c and d) and higher binding energies for Mo 3d signals (Fig. 6c and d); for Co 3p and Cr 3p signals, the observed shifts were not so prominent. In the spectral decomposition, each XPS spectrum for a certain element was modeled using three Gaussian subcomponents whose intensities were the main fitting parameters as shown in Figs. 5c and d and 6c and d. During the fitting procedure, the intensity ratio (I1:I2:I3) and binding energy shifts (De21 = e2 e1, De31 = e3 e1) of these subcomponents for the pure element and metal carbide states were kept constant. Fig. 7 displays the relative concentration vs. sputtering time for all metal carbide components decomposed from the sputter depth profiles, e.g. CoCrMo–Cx, Co–Cx, Cr–Cx, Mo–Cx. On the same graph the sum (Co–Cx) + (Cr–Cx) + (Mo–Cx) of the metal carbides obtained from the sputter spectra corresponding to pure metals is also represented. The solid lines represent the Gaussian fits of the experimental data. The relative concentration of CoCrMo–Cx and (Co–Cx) + (Cr–Cx) + (Mo– Cx) reached maxima at slightly different sputtering times (6.9 min vs. 7.2 min), probably due to small but inevitable errors in the decomposition procedure. Furthermore, since the maximum of the relative concentration of the CoCrMo–Cx component is only about half that of the (Co–Cx) + (Cr–Cx) + (Mo–Cx) peak, the average
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
4683
Fig. 5. Sputter depth profile XPS analysis for carbon: (a) C 1s sputter depth profile mapping (contour interval is 5% of the peak height); (b) profile decomposition into pure DLC and CoCrMo–Cx; (c and d) modeling of the XPS spectra recorded at sputtering times of 3.1 and 7.8 min (horizontal lines 1 and 2 in (a)), corresponding to CoCrMo–Cx and pure DLC.
Fig. 6. Sputter depth profile XPS analysis for molybdenum: (a) Mo 3d sputter depth profile mapping (contour interval is 5% of the peak height); (b) profile decomposition into pure Mo and Mo–Cx; (c and d) modeling of the XPS spectra recorded at sputtering times of 6 and 9.3 min (horizontal lines 1 and 2 in (a)), corresponding to pure Mo and Mo–Cx.
metal:carbon stoichiometry seems to be close to 2:1, e.g. Me2C. However, in view of the limited depth resolution of the sputter XPS method, the stoichiometry of the CoCrMo–carbide present at the interface cannot be determined accurately. Hence, it could be that an average Me2C stoichiometry is present only in the center of the reactive interface layer, whereas towards the substrate different stoichiometries exist. Moreover, since the peak metal carbide concentration ratios (Co–Cx):(Mo–Cx) = 10.6, (Cr–Cx):(Mo–Cx) = 3.6 and (Co–Cx):(Cr–Cx) = 2.9 are similar to the data shown in Fig. 4c obtained by STEM–EDX, a decrease of Co and an increase of Mo and Cr concentrations in the reactive interface layer with respect to the relative concentration in the bulk of the CoCrMo substrate is possible.
The thickness of the carbide interface layer was estimated as the FWHM of the Gaussian fitted to all carbide components combined together, e.g. (Co–Cx) + (Cr– Cx) + (Mo–Cx), and required 1.3 min etch time. Based on the obtained average carbide stoichiometry at the interface between DLC and CoCrMo the corresponding sputter rate was estimated to be 3.9 nm min1 [16]. Consequently, the thickness of the reactive interface layer is 5 nm, which is in very good agreement with the value obtained from HRTEM. However, it should be taken into account that the depth profile XPS method can overestimate the determined thickness of the reactive interface layer by 1– 2 nm through mixing processing occurring in sputter-depth profiling, such as recoil mixing and radiation-enhanced diffusion [17].
4684
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 7. Atomic concentration vs. sputtering time for metallic carbide components decomposed from the sputter depth-profile XPS spectra.
3.2. Instability of the reactive interface layer in PBS 3.2.1. Exact location and cause of the coating failure Previously, it was shown that, if the load induced at the interface (described by the strain energy release rate, G / d r2 , where d and r are the thickness and residual stress of the coating) is higher than the threshold for crack initiation in corrosive medium (GTH), then cracks will propagate with a velocity described by the SCC laws [11]. In particular, for a 4 lm thick DLC coating (r = 3.8 GPa, compressive) the strain energy release rate is 133 J m2, which is more than twice the threshold value for crack initiation in PBS at 37 °C of 60 J m2. Hence, in the absence of external loads this coating will delaminate in PBS at 37 °C with a velocity of 47 lm day1. However, it should be mentioned that at typical DLC deposition rates of 30–50 nm min1 the reactive interface layer, evidence of which was revealed in the previous section, is formed within the first 6–10 s of the deposition. If, during this relatively short time, gases such as O2 or H2O are present in high concentration, then the formation of metal carbides is hindered in favor of metal oxides. Consequently, the threshold strain energy release rate (GTH) and adhesion lifetime of the coating can be greatly reduced for these contaminated interfaces. The objective of this section is to show where the exact location of the coating failure is in PBS at 37 °C, and to reveal what causes this failure. For this purpose, the same lamella used to investigate the reactive interface layer, described in the previous section, was analyzed by TEM, but this time in a region where coating is delaminated, i.e. near the crack tip. While the SEM image in the inset of Fig. 2b already suggested that the crack responsible for coating delamination appears at the interface between the DLC film and the CoCrMo substrate, the TEM micrograph in Fig. 8a clearly demonstrates that this crack occurs
Fig. 8. Corroded CoCrMo carbide interface in PBS at 37 °C: (a) TEM view of the interface crack; (b) STEM–EDX chemical analysis across the interface.
approximately in the middle of the reactive interface layer, leaving an 2–3 nm thick medium-contrast strip on each side of the crack. Similarly, other authors revealed by TEM analysis that the mechanical failure in air of the DLC coatings deposited by ion implantation on aluminum alloy (A5052) substrates occurs due to the breaking of the mixing layer of oxide and DLC at the interface [18]. However, in order to be confident about our results, STEM– EDX analysis was performed across the delaminated interface (Fig. 8b). The results show a decrease in the Co, Cr, Mo signals across the first medium gray strip followed by an approximately constant level of these signals corresponding to the open crack. Finally, a slight increase over the second medium gray strip was observed, indicating that the Co, Cr and Mo elements were present there in a higher concentration due to the metal carbide bonds. The reason that Co, Cr and Mo signals did not cancel out completely in the open crack was probably the contamination of the thin underlying carbon foil which supports the TEM lamella. Thus, it is evident that the metal carbide reactive interface layer was unstable in PBS at 37 °C.
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
An important cause of the coating delamination could be the instability of the Co carbides, which are present in large concentration (38 at.%) at the interface between DLC and CoCrMo as revealed by the XPS results. For example, in the case of Co2C at 300 K, the reported Gibbs energy of formation D , which is a measure of the binding energy, is 13.90 kJ mol1 (0.14 eV), showing that formation of this carbide is unfavorable [19]. On the contrary, since for Mo2C, Cr3C2 and Cr7C3 the reported values of D are 54 kJ mol1 (0.56 eV), 86.27 kJ mol1 (0.89 eV) and 166.16 kJ mol1 (1.72 eV), respectively, the formation of these carbides is favorable [20]. However, to obtain more insight into the delamination process, several DLC coatings with 90 nm Co and Mo interlayers, and thicknesses between 2 and 8 lm, were deposited on the CoCrMo substrates. These DLC coatings were grown in similar conditions of bias voltage and chamber pressure as the coating grown directly on CoCrMo, but for Co and Mo interlayers DC magnetron sputtering was used. Fig. 9 compares the (G, hvi) curves of the DLC/CoCrMo, DLC/Co/CoCrMo and DLC/Mo/CoCrMo samples. For DLC/CoCrMo, a typical SCC with a threshold and a “stage 1” crack propagation with a slope parameter m = 3.0 was obtained. Low slope parameters in the range of 2–10 are typical for SCC of ductile materials [13]. However, the threshold strain energy release rate (GTH) for the DLC/Mo/CoCrMo sample was 170 J m2, which is almost 3 times larger than for the DLC/CoCrMo sample. The DLC/Co/CoCrMo sample displayed a “stage 1” crack propagation with a slope parameter m = 3.3. However, GTH for this sample could not be estimated accurately, since the stress induced at the interface for a 2 lm DLC coating is too large and therefore thinner coatings should have also been tested. Nevertheless, one could compare the adhesion strength of these samples by looking at the
Fig. 9. (G, hvi) diagrams for DLC/CoCrMo, DLC/Co/CoCrMo and DLC/Mo/CoCrMo samples immersed in PBS at 37 °C.
4685
cross-sections G (hvi = const.) and hvi(G = const.) through the SCC curves, indicated in Fig. 9 by the horizontal and vertical dotted lines. Consequently, for G = 400 J m2 the average speed of delamination for the DLC/CoCrMo sample is 11 times larger than for the DLC/Mo/CoCrMo sample and 4.5 times smaller than for DLC/Co/CoCrMo sample. Similarly, for hvi = 100 lm day1, the strain energy release rate for the DLC/CoCrMo sample is 2.5 times larger than for DLC/Co/CoCrMo and 5 times smaller than for DLC/Mo/CoCrMo sample. Therefore, from these findings it is obvious that the adhesion of DLC on CoCrMo is greatly improved by using a Mo interlayer. Conversely, the adhesion of DLC on CoCrMo is greatly diminished with a Co interlayer. Moreover, the analysis of the XPS sputter depth profiles for the DLC/Co/CoCrMo sample after coating delamination revealed that the 90 nm thick Co interlayer was profoundly corroded. The progressively delaminating coated sample was removed first from PBS, cleaned with deionized water in an ultrasound bath, and then an area from which the coating had delaminated was sputtered through the Co layer. Consequently, oxygen was present only at the surface but, when the sputtering was performed on previously delaminated regions that had been exposed to PBS for longer times, oxygen was also found in the Co interlayer; the thickness of the oxide layer increased with increasing exposure time to PBS. Thus, pure Co was found to be unstable in PBS and gradually transformed into Co oxide. 3.2.2. Corrosion mechanisms involved during the coating delamination Having established that Co carbide was the main cause for the failure of the DLC coating on the CoCrMo alloy, it is nevertheless imperative to identify and understand the mechanisms via which delamination occurs. In order for coating delamination from CoCrMo to occur, it is necessary for the corrosive medium to reach the interface region. Since DLC coatings deposited in an unclean room environment usually exhibit many defects that traverse their entire thickness down to the substrate, such as pinholes that originate from dust particles [21] and concentrations of stress at surface imperfections (e.g. inclusions at the surface) [22], this condition is fulfilled. Fig. 10a and b illustrate the 3D surface profiles of such pinholes through a 4 lm DLC coating deposited on CoCrMo; for a better visual inspection the vertical dimensions are exaggerated. These pinholes are 4 lm deep and their lateral size varies from a few tens of nanometers up to several tens of microns. Fig. 10c shows a top view optical micrograph of the same coating exposed to PBS at 37 °C for approximately 1 day obtained by all-in-focus image reconstruction. After the breaking of the interface bonds around the pinhole by corrosion mechanisms that will be discussed below, the coating buckled due to its internal compressive stress, since it was no longer adhering to the substrate. The buckling of the coating was revealed by the circular dome with a radius of 45 lm illuminated by the side microscope light source.
4686
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 10. Pinhole initiated corrosion: (a and b) 3-D profiles of pinholes in a 4 lm thick DLC-coated CoCrMo sample; (c) optical micrograph top view showing the buckling of the coating in PBS.
Similar film buckling was also observed in the 3-D scan shown in Fig. 10b. In addition, radial cracks with lengths in the range of 10–15 lm were observed around the central pinhole; these cracks were caused by the elastic energy stored in the buckled coating which exceeded the fracture toughness of the film. The crack front under the coating will propagate outwards, as indicated by the radial arrows, until the coating delaminates, leaving behind an almost circular pit, such as the one shown in Fig. 10b with a lateral size of 200 lm.
Fig. 11 schematically displays how micrometer-sized defects, such as pinholes, gradually lead to the delamination of the DLC coating. In the beginning, the PBS liquid entering through the pinhole (Fig. 11a) starts to locally corrode the reactive interface layer, forming a small crevice under the DLC coating (Fig. 11b). Similar schematic diagrams were proposed by other authors [23,4]. In addition to the dissolution of the reactive interface layer, an etching process of the CoCrMo substrate was also observed for some pinholes, revealed by small pits in the metal substrate
Fig. 11. Simple schematics of the proposed gradual delamination mechanism of the DLC coatings from CoCrMo.
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
at the location of the pinholes. The formation of such pits is usually attributed to “passive film breakdown”, i.e. interaction of some aggressive ions (e.g. Cl) with the protective surface oxide film at locations where it is defective or weak [1,24]. Due to the gradual dissolution process of the reactive interface layer, the advancing interface crack front induces buckling of the coating (Fig. 11c) until the elastic energy stored in the coating exceeds the fracture toughness and the coating breaks leaving behind a crater (Fig. 11d). In PBS, this SCC process driven by the corrosive attack at the crack front tip and residual stress of the coating may continue until the whole DLC coating delaminates from the CoCrMo substrate. Simple schematics of a possible SCC mechanism occurring in the reactively formed metal carbide interface during the crack front expansion are proposed in Fig. 11e and f, in which the crack growth is related to oxidation reactions at the crack tip as described by the slip-oxidation and enhanced surface mobility models [25,26]. However, one has to make the distinction between the classic slip-oxidation model, in which the oxide film ruptures at the location of the bulk metal grain boundary that subsequently dissolves as the crack advances, and the present case, in which ruptures occur in the middle of a metal carbide interface layer. According to the enhanced surface mobility model, crack growth occurs due to the capture of vacancies by the stressed lattice at the crack tip. Thus, the transport of atoms takes place by surface diffusion from the highly stressed crack tip to the less stressed regions. In addition, the corrosive medium influences the crack growth rate by affecting the dissolution and surface diffusion of the atoms. After the breaking of the unstable Co carbide bonds, the oxide-free surface of the crack tip becomes exposed to the PBS solution and positively charged Co2+ ions migrate from the metal into the solution, leaving electrons behind according to the anodic reaction: Co ! Co2+ + 2e [27]. In neutral PBS solution the electrons produced can react with oxygen molecules adsorbed in the crevice from air dissolved in solution, according to the cathodic reaction O2 + 2H2O + 4e ! 4OH [1,27]. In the crevice formed at the interface, however, the pH can reach acidic values because of hydrolysis reactions such as Co2+ + 2H2O ! Co(OH)2 + 2H+ [1,27,28]. Acidification of the fluid in the crevice could also take place as a result of the anodic reactions Co + H2O ! CoO + 2H+ + 2e, Co + 2Cl ! CoCl2 + H2O ! CoO + 2H+ + 2Cl [27,28]. Consequently, the residual electrons can react with hydrogen ions, adsorbed on the metal surface from the solution, to produce hydrogen gas by the cathodic reaction 2H+ + 2e ! H2" [1]. It should be mentioned that acidification depends strongly on the exchange of the fluid inside the crack with the environment and therefore on the crack geometry. The last reaction is also relevant for in vivo conditions immediately after implantation when the conditions are most aggressive and the biological fluid becomes acidic
4687
[29]. However, it is unknown whether the H2 gas produced forms small bubbles and leaves the crevice, or if it penetrates into the reactive interface layer, inducing coating delamination through hydrogen embrittlement. 3.2.3. Effects of the PBS temperature on the coating delamination Since crack openings are usually much smaller than the pinhole dimensions as shown in Figs. 8a and 10a and b, reaction-limiting processes of the corrosion at the crack tip could play an important role in coating delamination. Hence, SCC curves for DLC/CoCrMo were determined for temperatures of 4, 23, 37, 66 and 95 °C. All curves were fitted with the reaction-controlled model [30]. The corresponding values of the GTH and crack-growth exponent m for the “stage 1” SCC are plotted in Fig. 12a. An exponential decrease in the GTH was observed, whereas the crack-growth parameter m exhibits a slight linear increase with temperature. Growth of the value of m with increasing temperature suggests that the solution is slightly more aggressive [13] or that the material properties at the crack tip are slightly modified with increasing temperature. Moreover, the decrease of GTH with increasing temperature suggests that a diffusion process of the corrosive medium to the crack tip took place. Thus, the higher the temperature, the higher the solution exchange at the crack tip, and therefore the higher the dissolution efficiency of the interface material. Indeed, by plotting the average delamination velocity hvi in PBS vs. the inverse temperature T for an interface load of 133 J m2 corresponding to a 4 lm thick DLC coating far away from the Rockwell indentation, a typical behavior described by the Arrhenius equation hvi ¼ A expðEa =RT Þ was observed up to 37 °C [31]. This suggests that at least at temperatures below 37 °C the crack growth is a reaction-limited process. In the Arrhenius equation, Ea is the activation energy, i.e. the minimum amount of energy necessary for the SCC reaction, A is the pre-exponential factor equivalent to the speed of delamination when Ea = 0, and R is the universal gas constant. It should be also noted that in the Arrhenius equation Ea and A were considered as being independent of the temperature, which is not always the case, but for a small range of temperatures it is usually a good approximation. Consequently, values of 13 ± 7 lm s1 and 26.5 ± 0.6 kJ mol1 were obtained for the pre-exponential factor and activation energy, respectively. However, from Fig. 12b it is also evident that for temperatures higher than 37 °C, the crack growth behavior deviates from the linear dependence, which suggests that the rate-limiting factor at these temperatures has different causes (special chemical reactions, variation of the pH value of the corrosive media with the temperature, mechanical stability and ductility of the interface material, atypical diffusion processes, variation of the oxide layer thickness, etc.). The gradual decrease in the slope of the Arrhenius plot with increasing temperature could also be an indication for the temperature variation of
4688
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
Fig. 12. Temperature dependence of the SCC-induced coating delamination in PBS: (a) m and GTH vs. T; (b) Arrhenius plot for the temperature range of 4–95 °C and G = 133 J m2.
the pre-exponential factor (e.g. A ¼ BT n , where B and n are constants). Consequently, the modified Arrhenius equation hvi ¼ BT n expðEa =RT Þ could be used to describe the temperature variation of the SCC-induced coating delamination [31]. Using this new equation to fit the experimental data points for the temperature range from 4 to 95 °C, a value of 430 kJ mol1 was obtained for the activation energy. However, since this value is much larger than the previously reported one, of 26.5 ± 0.6 kJ mol1, which was obtained using the linear dependence for temperatures below 37 °C, one should consider it with extreme caution. First of all, in the modified Arrhenius equation the temperature dependence of the activation energy was neglected. Furthermore, in view of the complex system one deals with, of a multielement interface in a dynamic corrosive environment, where a multitude of chemicophysical processes may play a role, it could be misleading merely to treat the delamination process as a chemical reaction described by an Arrhenius behavior. Nevertheless, in general, if coating delamination is driven by SCC, then a similar analysis could be applied to different coating/substrate combinations to determine the activation energy for SCC. 4. Conclusions In the present study the exact microstructure of the interface between DLC thin films and CoCrMo biomedical implants, as well as the delamination processes occurring in PBS solution, were investigated, with the following findings:
2.
3.
4.
5.
of 5 nm, has a disordered amorphous-like structure. Moreover, sputter depth profile XPS measurements revealed its average Me:C stoichiometry to be 2:1. It was demonstrated that in the absence of any interlayer the DLC delaminates from the CoCrMo due to the instability of the reactive interface layer in PBS at 37 °C. Moreover, HRTEM revealed that the crack propagation occurs roughly in the middle of this layer. The failure of the DLC on CoCrMo was shown to be due to the metastable Co carbides (38 at.% from XPS) formed in the reactive interface layer during the film deposition, since by growing DLC layers with a 90 nm Co interlayer for the same load (G = 400 J m2) the average speed of delamination was increased by a factor of 4.5. However, by growing layers with a 90 nm Mo interlayer, the average speed of delamination was decreased by a factor 11, and the threshold for SCC was increased from 60 to 170 J m2. It was shown that delamination is initiated by the corrosive medium that penetrates the pinholes traversing the DLC layer. Subsequently, the delamination continues due to diffusion processes in the crevice formed in the carbide interface layer and chemical reactions at the crack tip. In addition, by varying the temperature of the PBS it was shown that up to temperatures of 37 °C the SCC-induced coating delamination follows an Arrhenius behavior with thermal activation energy of 26.5 kJ mol1; for higher temperatures deviations from the linear Arrhenius behavior were observed.
Acknowledgments 1. A reactive interface layer between the DLC coating and CoCrMo substrate was revealed and investigated in detail by structural and chemical analysis. HRTEM measurements showed that this layer, with a thickness
We would like to thank R. Crockett, P. Schmutz and B. Lehman for useful discussions and comments, and S. Meier for preparing the lamellae for TEM investigations. This
C.V. Falub et al. / Acta Materialia 59 (2011) 4678–4689
work was supported by the Swiss Commission for Technology and Innovation (CTI) and Synthes GmbH. References [1] Kruger J. Fundamental aspects of corrosion of metallic implants. In: Syrett B, Acharya A, editors. Corrosion and degradation of implant materials. ASTM STP 684; 1979. [2] Greenfield EM, Bi Y, Ragab AA, Goldberg VM, Nalepka JL, Seabold JM. J Biomed Mater Res Part B: Appl Biomater 2005;72B:179. [3] Jacobs JJ, Skipor AK, Paterson LM, Hallab NJ, Paprosky WG, Black J, et al. J Bone Joint Surg Am 1998;80:1447. [4] Lappalainen R, Santavirta SS. Clin Orthop Relat Res 2005;430:72. [5] Tiainen VM. Diamond Relat Mater 2001;10:153. [6] Hauert R. Tribol Int 2004;37:991. [7] Bruinink A, Schroeder A, Francz G, Hauert R. Biomaterials 2005;26:3487. [8] Hauert R. DLC films in biomedical applications. In: Donnet C, Erdemir A, editors. Tribology of diamond-like carbon films: fundamentals and applications. Berlin: Springer-Verlag; 2008. [9] Ta¨ger G, Podleska LE, Schmidt B, Ziegler M, Nast-Kolb D. Materialwiss Werkst 2003;34:1094. [10] Joyce TJ. Wear 2007;263:1050. [11] Falub CV, Thorwarth G, Affolter C, Mu¨ller U, Voisard C, Hauert R. Acta Biomater 2009;5:3086. [12] Thorwarth G, Falub CV, Mu¨ller U, Weisse B, Voisard C, Tobler M, et al. Acta Biomater 2010;6:2335. [13] Jones RH, Ricker RE. Mechanisms of stress–corrosion cracking. In: Jones RH, editor. Stress–corrosion cracking. Materials Park (OH): ASM International; 1996.
4689
[14] Rajan K. Metall Trans A 1982;13A:1161. [15] Biersack JP, Haggmark LG. Nucl Instrum Methods 1980;174:257. [16] Mu¨ller U, Falub CV, Thorwarth G, Voisard C, Hauert R. Acta Mater 2011;59:1150. [17] Littmark U, Hofer WO. The theory of recoil mixing in solids. In: Oechsner H, editor. Topics in current physics, thin films and depth profile analysis, 37. Berlin: Springer-Verlag; 1984. [18] Oka Y, Nishijima M, Hiraga K, Yatsuzuka M. Nucl Instrum Methods Phys Res B 2007;257:702. [19] Seng WF, Barnes PA. Mater Sci Eng B 2000;76:225. [20] Barin I. Thermochemical data of pure substances, 2nd ed., vols. 1 and 2; 1989. p. 921, 560, 561. [21] Pranevicius L. Vapor deposition coating technologies. In: Tracton AA, editor. Coatings technology handbook. Boca Raton (FL): CRC Press; 2006. [22] Outka DA, Hsu WL, Boehme DR, Yang NYC, Ottesen DK, Johnsen HA, et al. Compilation of diamond-like carbon properties for barriers and hard coatings. Sandia report, SAND94-8219, UC-404; 1994. [23] Chandra L, Allen M, Butter R, Rushton N, Lettington AH, Clyne TW. Diamond Relat Mater 1995;4:852. [24] Hoar TP, Jacob WR. Nature 1967;216:1299. [25] Andresen PL. Corrosion 1993;49:714. [26] Galvele JR. Corros Sci 1987;27:1. [27] Deltombe E, Zoubov N, Pourbaix M. Section 12.2: cobalt. In: Pourbaix M, editor. Atlas of electrochemical equilibria in aqueous solutions. Oxford: Pergamon Press; 1966. [28] Gallant D, Simard S. Corros Sci 2005;47:1810. [29] Lopez DA, Duran A, Cere´ SM. J Mater Sci Mater Med 2008;19:2137. [30] Cook RF, Liniger EG. J Am Ceram Soc 1993;76:1096. [31] McNaught AD, Wilkinson A, editors. IUPAC compendium of chemical terminology, 2nd ed. Oxford: Blackwell Science; 1997.