tantalum coatings

Surface & Coatings Technology 202 (2008) 5001–5007

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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t

Interface interactions between porous titanium/tantalum coatings, produced by Selective Laser Melting (SLM), on a cobalt–chromium alloy P. Fox a,⁎, S. Pogson a, C.J. Sutcliffe a, E. Jones b a b

MSERC, Department of Engineering, The University of Liverpool, Liverpool, L69 3BH, UK Stryker Orthopaedics, Co Limerick, Ireland

A R T I C L E

I N F O

Article history: Received 27 November 2007 Accepted 6 May 2008 Available online 13 May 2008 Keywords: Selective Laser Melting Porous coatings Materials interactions

A B S T R A C T Porous titanium and tantalum coatings were produced on cast cobalt–chromium alloy substrate plates (Co28 Cr-6 Mo ASTM designation F75)) using the additive manufacturing process Selective Laser Melting (SLM). Both tantalum and titanium coatings where successfully produced, however, a poor interface bond was observed with the titanium coatings on the cobalt–chrome alloy. This was due to a eutectic reaction leading to the formation of a low melting point phase βTi(CoCr) which cracks during cooling, rather than the formation of titanium carbide, as previously reported. This cracking makes titanium an unsuitable material to coat cobalt–chromium alloys using SLM. Tantalum coatings, however, showed considerably improved performance in terms of interface compatibility when compared to titanium and therefore a Co–Cr/Ta system would seem feasible for the manufacture of porous structured devices when a bi-material approach is required. This would allow the advantages of a highly biocompatible structured coating to be combined with the mechanical performance of a less biocompatible substrate. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Selective Laser Melting (SLM) is an additive Rapid Prototyping (RP) technology based on Selective Laser Sintering (SLS). However, whereas SLS utilises a low power laser (CO2, λ ≈ 10 μm and power b50 W) to sinter the powder particles together [1], SLM uses a higher power laser (Ytterbium fibre/Nd:YAG, λ ≈ 1 μm and power b200 W) to fully melt the powders [2]. Although much of the early work on SLM concentrated on reducing porosity and producing fully dense material [3–5], recent works have shown that this technology could be ideal for producing coatings or full components where a controlled or designed-in interconnected porosity is required [6–8]. This allows the production of bespoke porous coatings or structures where exterior shape, pore size and relative density are controlled. Many methods exist for the application and construction of porous biomaterials, these are discussed elsewhere [9], it is also common practice in the orthopaedic industry to produce coatings on surgical implants where the coating and substrate are metallurgically different. Typically plasma spraying of Ti on Ti alloy is used to produce a rough rather than porous biocompatible surface for tissue ongrowth [10]. Cold spray processes have also been considered for the production of fully porous coatings but this technology has not yet been fully realised [11].

⁎ Corresponding author. Tel.: +44 1517945371; fax: +44 1517944675. E-mail address: [email protected] (P. Fox). 0257-8972/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2008.05.003

SLM [12] however, is able to produce components with fully interconnecting porosity (required to simulate the nature and mechanical modulus of cancellous bone). The process also has the capacity for manufacture with dissimilar metals, particularly if the substrate on which the component is built is one metal and the powder material is the other. Thus, by combining materials in the SLM process it is possible to choose the most suitable properties for specific purposes. For example, the strength, wear resistance and notch insensitivity of cobalt–chrome alloys could be used for load bearing and articulation surfaces combined with the improved biocompatibility of titanium as the porous coating for bone ingrowth. For this system of manufacture to be useful, the potential problems with reactions between the dissimilar materials must be overcome. Such interface problems have been reported in the literature when diffusion bonding wear-resistant cobalt alloys to titanium [13]. In this case titanium carbide was shown to form at the coating/substrate interface, greatly reducing the ductility of the interface leading to premature failure. The problem was solved by the introduction of intermediate bond layers of nickel and tantalum [14,15]. The production of such multi-material components has therefore been shown in the literature for the diffusion bonding process, however; the behaviour at this interface is less well understood for SLM, as the bonding of the coating to the substrate requires the melting of the substrate leading to the mixing of the alloys within the melt pool. This paper addresses these interface properties from a metallographic perspective for a Co–Cr alloy substrate with Ti and Ta coatings.

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P. Fox et al. / Surface & Coatings Technology 202 (2008) 5001–5007

2. Experimental procedure To test the suitability of the SLM technique for producing porous coatings on a dissimilar substrate; titanium and tantalum coatings were deposited on a cast cobalt–chromium alloy (Co-28 Cr-6 Mo ASTM designation F75). The porous coatings were built in a layerwise manner by applying powder in 100 μm thick layers to the Co– Cr substrate. This powder layer was then selectively melted over many layers by a scanning laser beam to form the porous constructs. The laser used was a Rofin-Sinar Krypton flash lamp pumped Nd: YAG laser (100 μm spot size, maximum power 90 W continuous wave). The system is fully described elsewhere [16], but here argon (as opposed to nitrogen) was used as the protective gas as both titanium and tantalum react with nitrogen forming metallic nitrides. Environmental control is important as oxygen affects the mechanical properties of solid titanium and the flow and wettability of the molten metal. Simple linear single direction laser scans were used to form the coating, the scan line spacing controlling the pore size (a 100 μm laser spot size and a negative 200% overlap producing a 300 μm spacing). As the line spacing was used to control pore size it could not be used to control the specific energy density this being controlled by the laser scanning speed alone. Although the scan spacing will affect the specific energy density if the pore size is altered this effect is small, as the laser scans do not overlap. Typical processing parameters used a laser power of 82 W and a scanning speed range of 100–260 mm s− 1. The powders used were obtained from a number of sources, the tantalum being a spherical gas atomised powder (particle size distribution 80% b 75 μm, purity 99.85%), the titanium powder being spherical gas atomised commercially pure titanium (particle size distribution of 80% b 45 μm and purity of 99.5%). The Co–Cr substrates were turned from 25.4 mm diameter Co–Cr bar stock the ends being surface ground to a finish of 15 μm Ra. Standard metallographic and optical microscopy techniques were used with the samples being etched with suitable etchants for the materials. A 5 g NH4FHF, 100 ml distilled H20H2O, and 2 ml HCl etch for 30 s was used for the titanium and tantalum materials, a 50 vol.% HCl, 50% H20 H2O heated to 80 °C for the cobalt–chrome alloy and a 5%

Fig. 1. Single layer of Ti on Co–Cr (power 82 W CW, − 40% beam overlap, scan speed 160 mm s− 1).

Fig. 2. Ti-multilayers on Co–Cr (power 82 W CW, −40% beam overlap, scan speed 160 mm s− 1).

Nital solution to reveal the interface microstructure without affecting the other regions. Scanning electron microscopy (SEM) in backscattered electron (BSE) imaging mode and chemical microanalysis was performed using a Hitachi-S2460N at 25 kV. X-ray diffraction (XRD) measurements were recorded at room temperature using a Rigaku Miniflex diffractometer (Cu Kα radiation). XRD analysis was carried out at different depths though the coating by grinding away material using 1200 grade silicon carbide paper. Micro-hardness testing was undertaken using the standard Vickers micro-hardness indenter test.

Fig. 3. Single layer of Ta on Co–Cr (power 82 W CW, −40% beam overlap, scan speed 160 mm s− 1).


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3. Results Optical and SEM analysis of single layer coatings on the substrate showed that titanium and tantalum behaved very differently. The first layer of titanium (Fig. 1) strongly wetted the substrate, widening the tracks and fusing them together producing a thin coating with little or no sign of balling. However, subsequent layers did not show this effect (Fig. 2) and looked similar to the tantalum coatings, which showed significant balling (Fig. 3) forming a non-continuous porous coating. Examination of the coatings in section revealed further dissimilarity between the two materials, as fine cracks were seen to occur at the coating substrate interface (Fig. 4) for the titanium coating, whilst no cracking was seen in the tantalum coating. The cracks within the titanium coated parts were confined to a thin layer at the interface between the Co–Cr substrate and the titanium coating, and did not spread into the remaining layers of the titanium coating or into the substrate. If the outer layers of the titanium coating were removed the highly fractured nature of the first layer could be observed (Fig. 5). Micro-hardness testing of the sections through the Ti coating showed that the Co–Cr substrate hardness (410Hv) was somewhat softer than the first Ti layer (520Hv). The hardness was shown to drop rapidly within the first few layers reaching a minimum of 200Hv at the outer layer. In comparison the tantalum coating showed no cracking near the interface (Fig. 6) but again the coating close to the interface was significantly harder (600Hv) than the coating further away (160Hv). To determine the phases present within the coatings and at the coating substrate interface, a combination of SEM/EDS and XRD techniques were used, with any phase variation through the coatings being determined by grinding away some of the coating between XRD analyses. The substrate material was examined first so its XRD fingerprint could be determined. The Co–Cr (Co, 28Cr, 6Mo) substrate was a standard, cast medical grade alloy conforming to ASTM F75 and as such consists of the two major phases, fcc α-Co and M23C6 carbides. However, XRD (Fig. 7) analysis of the base material ground with 1200 silicon carbide grit detected two major cobalt-rich phases fcc α-Co and hexagonal ε-Co, as well as the carbide. The appearance of this second cobalt phase is probably caused by a strain-induced martensite reaction within the ground surface [17]. As the XRD analysis was to finger-print the underlying substrate so that peaks from it could be ignored from the analysis, grinding rather than polishing was used for the micro-section preparation. XRD analysis (Fig. 8) of the titanium coatings showed the outer regions to be the low temperature αTi phase. Further into the coatings, approaching the coating/substrate interface, a bcc phase

Fig. 5. Highly fractured interface layer formed when processing Ti on Co–Cr (power 82 W CW, − 40% beam overlap, scan speed 160 mm s− 1).

was detected with lattice parameters similar to the intermetallic phase TiCo. This TiCo phase is an ordered version of the high temperature β phase of titanium [18] and the only observable difference between the XRD plots of the ordered and disordered structure is the presence of small peaks which are normally forbidden in a disordered structure. These small peaks are not generally visible within the XRD data, as there are peak overlaps and significant noise in the data. SEM/EDS analysis (Fig. 9) did not reveal any regions with compositions similar to TiCo, and therefore it is reasonable to assume that the phase is β-Ti but with a lattice parameter similar to the intermetallic and without any ordering. The only other phase detected appears to be that of TiCr2 present in small quantities. XRD analysis (Fig. 10) of the tantalum coating ground to different depths revealed the coating to consist mainly of the bcc β-Ta phase. Also observed were the phases from the substrate, and one peak that could not be identified. This peak was not observed within the

Fig. 4. Ti on Co–Cr in cross section showing cracks in the interface layer, outer layers of coating mechanically removed (power 82 W CW, −40% beam overlap, scan speed 160 mm s− 1).

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Fig. 6. Ta on Co–Cr alloy in cross section showing that although wetting is limited there is no observable cracking (power 82 W CW, − 40% beam overlap, scan speed 160 mm s− 1).

substrate or in the outer layers of the coating. It is considered to be attributable to either a small amount of an intermetallic or a carbide. 4. Discussion As can be seen, the behaviour of titanium and tantalum during SLM processing onto this Co–Cr alloy is very different, in that titanium wets the substrate much more readily, spreading across the surface, while although tantalum wets the surface, bonding to the substrate, it does not spread appreciably. The other major difference is the development of cracks within the interface region with titanium. The most likely cause of these cracks is the presence of a hard, brittle intermetallic phase at the interface. However, the composition of this region is not that of the intermetallic TiCo, and the hardness of the phases near the interface are lower for titanium than they are for tantalum. Therefore, another theory must be found to explain this behaviour. Although some limited phase diagrams are available for Co–Cr–Ti [19] they do not consider the titanium-rich section of these diagrams. Regarding the structure Co–Cr–Ta, no ternary diagrams have been found in the literature. Therefore, it is necessary to use the binary Ti–Co, Ti–Cr, Ta–

Co and Ta–Cr phase diagrams [20] to develop a clearer insight regarding the behaviour of these materials. The most important feature in the Ti–Co system is the very deep eutectic valley at the titanium-rich side of the diagram where the eutectic melting point is (1020 °C) (significantly below the melting points of both titanium (1668 °C) and cobalt (1495 °C)). Also present are a number of intermetallics, but the only high temperature intermetallic identified is TiCo (approx. 1450 °C). The titanium–chromium binary diagram shows that at high temperatures a solid solution (βTi) exists across the complete composition range, with the melting point dropping to a minimum of approximately (1400 °C). In comparison the Ta–Co and Ta–Cr diagrams do not contain any deep eutectic valleys although the melting points of Ta (3017 °C), Co (1495 °C) and Cr (1907 °C) are very different. It also has to be kept in mind that the melting point of tantalum is higher than the boiling points of cobalt (2927 °C) and chromium (2671 °C). In the sample preparation process, titanium powder was spread as a thin layer on the surface of the cobalt–chromium alloy, and powder, was then melted by scanning the laser beam across the surface of the powder bed. At the start of the scan, titanium powder will be heated to above its melting point and form a melt pool. The underlying substrate will then dissolve into the melt pool rapidly suppressing the freezing point. As this occurs the laser moves forward across the surface adding unmelted titanium to the front of the melt pool, while at the melt pool tail material is frozen out. As the melt pool is moving rapidly, it is probable that the front of the melt pool will contain material that melts at a temperature close to the melting point of titanium (1668 °C) while the tail will be highly alloyed with elements from the substrate that lower the freezing point to about (1020 °C). This laser scanning process produces an extended melt pool and extreme superheating of the molten material. A consequence of this is that the melt pool is larger than the laser beam diameter, as wetting/fluidity are increased by superheating of the liquid, and the melt pool is present for longer. The presence of the alloying additions not only suppress the melting point, but also stabilise the βTi phase and hardens the material by solid solution strengthening. This produces the microstructures observed in the first layer but does not explain the tendency to crack. After the first layer has been deposited and selectively melted, a second layer of powder is then deposited, and the powder melted by

Fig. 7. XRD of substrate Co–Cr alloy.


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Fig. 8. XRD of Ti on Co–Cr multi scans (power 82 W CW, −40% beam overlap, scan speed 160 mm s− 1).

the scanning laser beam. The titanium powder on the surface now has a significantly higher melting point than the underlying layer. This will allow the melt pool to melt a significant part of the previous layer but probably not through to the underlying substrate. As the layers mix, then the concentration of alloying additions will reduce and there will be an increase in the freezing point. A similar effect will occur with each subsequent layer but with each layer the concentration of alloying additions from the substrate will reduce until the layer composition is the same as the original powder. Associated with this change in composition will be a reduction in hardness and a change in microstructure as the βTi phase is no longer stabilised to room temperature. The cracks observed near the coating/substrate interface are probably due to the titanium-rich coating layer close to the interface being much weaker at high temperature, due to its melting point being lower than either the substrate below or coating material above. Thus the thermal stresses that develop because of the differences in thermal expansion coefficient will lead to hot tearing. It is also worth remembering that the titanium coating further away

from the interface will undergo a martensite reaction on cooling from the β region, and the imposed strains will favour certain martensite orientations allowing the thermal strains to be accommodated without cracking, thus explaining why the cracks observed are confined to the interfacial layer of the titanium. In contrast to titanium the first layer in the tantalum coating will probably form in a different manner as the substrate has a significantly lower melting point than the coating material. The powder layer was only 100 μm and therefore on heating the tantalum, the laser will also heat the substrate melting the surface of the substrate before the tantalum. The alloying of the tantalum by the substrate material will reduce the melting point producing a melt pool containing significant levels of elements from the substrate along with tantalum. However, in this case the freezing point of the liquid will not vary significantly from one end of the melt pool to the other. These results suggest that the application of a porous coating on to the Co–Cr substrate is better achieved with the use of tantalum metal as the coating material. From a metallographic perspective, the

Fig. 9. EDS analysis of the cracked region shown in Fig. 6.

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Fig. 10. XRD of Ta on Co–Cr (power 82 W CW, −40% beam overlap, scan speed 160 mm s− 1).

interface between the Co–Cr is devoid of any cracking that would weaken this interface with subsequent tantalum layers building well. A more comprehensive analysis of this bi-metal combination is warranted. In contrast, the cracking at the interface of the Co–Cr/ titanium combination is of concern. Applications of the technology discussed here are wide ranging with several medical devices being developed. Fig. 11 shows a porous knee implant in which the porous coating and solid/porous keel structure has been deposited on a dissimilar metal substrate plate. 5. Conclusions Previous published work has shown that the processability of materials by Selective Laser Melting is affected by many process parameters, but one which is often not considered is the alloying that occurs in the melt pool when one metal is deposited onto another. For alloy systems, like the Co–Cr/titanium system examined here, that show deep eutectic valleys, alloying in the melt pool will under certain circumstances significantly suppress the freezing point of the melt pool, affecting the morphology and mechanical behaviour of the deposited materials. Acknowledgements This work was carried out within the framework of a project sponsored by the UK Engineering and Physical Sciences Research Council (GR/S99020/01). References

Fig. 11. Porous and solid structured knee implant deposited on a dissimilar material substrate plate.

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