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Sputter-deposited Al–Au coatings
Intermetallics 12 (2004) 579–587 www.elsevier.com/locate/intermet
Sputter-deposited Al–Au coatings C. Mitterera,*, H. Lenharta,1, P.H. Mayrhofera, M. Kathreinb a
Department of Physical Metallurgy and Materials Testing, University of Leoben, Franz-Josef-Strasse 18, A-8700 Leoben, Austria b CERATIZIT Austria GmbH, A-6600 Reutte, Austria Received 5 January 2004; accepted 13 February 2004
Abstract Transition metal nitride-based, wear-resistant hard coatings on cutting tools and other substrates often lack self-lubricating properties at elevated temperatures and distinct colorations allowing product differentiation. In the present work, the possibility of achieving these objectives using coatings based on the purple-red Al2Au phase within the Al–Au system was investigated. Coatings within this system have been deposited onto cemented carbide and Inconel substrates using unbalanced magnetron sputtering. The coatings were characterized with respect to their topography and morphology, chemical and phase composition, hardness, optical, oxidation and tribological properties. Al2Au-containing coatings have been deposited with dense, fine-grained structures yielding a hardness of 4 GPa and pink coloration. Vacuum annealing at temperatures of 500 C results in a chance to a pronounced purple coloration. The coatings are stable up to about 850 C, where the onset of oxidation occurs. Low friction coefficients when testing against alumina counterparts were achieved in the temperature range between 500 and 700 C. The concept of applying Al2Aucontaining coatings as a colored self-lubricating layer on top of a hard coated cemented carbide tool warrants further investigations. # 2004 Elsevier Ltd. All rights reserved. Keywords: A. Aluminides, miscellaneous; B. Phase identification; B. Tribological properties; B. Thermal stability; C. Coatings, intermetallic and otherwise
1. Introduction Hard coatings, e.g. based on transition metal nitrides, are widely applied to improve the lifetime and the performance of many kinds of cutting and forming tools. Besides the functionality of the coating itself, provided by e.g. hardness, toughness, oxidation and wear resistance, there is also an increasing demand for cosmetic factors like attractive colorations of a cutting tool [1]. Fortunately, the most widely applied titanium nitride TiN is characterized by a bright golden-yellow color [2] making determination of used cutting edges of a cutting insert in poor lighting conditions of a machine shop easier. However, for other hard compounds there exists a considerable lack of attractive bright colors [3] allowing product differentiation and to distinguish used and unused cutting edges. On the other hand, during recent years tremendous effort has been laid on development * Corresponding author. Tel.: +43-3842-402-4220; fax: +43-3842402-4202. E-mail address: [email protected] (C. Mitterer). 1 Present address: Steirische Wirtschaftsfo¨rderung, Nikolaiplatz 2, A-8020 Graz, Austria. 0966-9795/$ - see front matter # 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2004.02.003
of low-friction coatings being able to reduce friction forces. Today, low-friction coatings rely on solid lubricant phases like diamond-like carbon DLC [4], molybdenum disulfide MoS2 [5] or chlorine-induced rutile formation on chlorine-containing TiN coatings [6,7]. These lubricants often begin to fail in their tribological effectiveness with increasing temperature, in humid atmosphere or due to oxidation [8,9]. Consequently, this study is focused on the development of a new type of colored low-friction coatings based on the intermetallic Al2Au phase. Representing a Zintl phase with the cubic CaF2 structure [10], Al2Au shows a distinct reflectivity minimum at 545 nm [11,12], resulting in a purple-red color. Being relatively hard and brittle at room temperature, aluminides like Al2Au show plastic deformation and possibly self-lubrication at high temperatures [13]. Within this work, we present a comprehensive overview on the sputter deposition of Al–Au coatings and on the characterization of coating composition and microstructure as well as optical, mechanical, oxidation, and tribological properties. Special emphasis will also be laid on the thermal stability of the coatings.
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2. Experimental details The coatings investigated in this work have been deposited using an unbalanced d.c. magnetron sputtering system that has been described in detail in Refs. [14,15]. Different Al/Au atomic ratios have been adjusted by placing different numbers of Au pieces (20 201 mm3) on the erosion track of an Al target (11506 mm) mounted onto an unbalanced Gencoa PP 150 magnetron. The target was positioned 9 cm from the parallel-plate substrate assembly. Coatings were deposited onto uncoated and (Ti,Al)N coated cemented carbide cutting inserts of different grades and geometries and Inconel 718 substrates (151 mm) in Ar atmosphere at 0.2 Pa. The substrates were positioned directly above the target erosion track to ensure a homogeneous Al/Au distribution within the coating. For all deposition runs, the substrate temperature was held constant at 300 C, the bias voltage was 50 V, and the target sputter power was about 380 W. Prior to deposition, all substrates used were metallographically grounded, polished and ultrasonically cleaned with ethylene and acetone. After target pre-cleaning and ion etching of the substrates within the deposition chamber, coatings in the thickness range between 0.1 and 7 mm were deposited by selecting suitable deposition times. Coatings with thicknesses of 3–7 mm were used for the basic chemical, microstructural mechanical, and tribological characterization, whereas for the investigation of the optical properties also thinner coatings were used. Coating topography and morphology were investigated using scanning electron microscopy (SEM, Cambridge Instruments Stereoscan 360). For the investigation of the coating morphology, cross-sections of coated cemented carbide inserts were prepared by brittle fracturing in liquid nitrogen. The chemical composition of the coatings was determined by energy-dispersive X-ray analysis (EDX, Link eXL) using Au and Al elemental standards. X-ray diffraction (XRD) patterns were recorded using a Siemens D 500 Bragg-Brentano diffractometer and Cu-Ka radiation. Lattice parameters were corrected over cos2y. In addition, high-temperature XRD (HT-XRD, Siemens D5005 with heating chamber Paar HTK 1200, BraggBrentano configuration, heating rate, 0.6 K.min 1) investigations in He atmosphere were performed for coatings deposited onto cemented carbide inserts to obtain information on the thermal stability of the coatings. Coating hardness was characterized using a Vickers microhardness tester (Reichert-Jung MD 4000) at a load of 25 mN. Coating coloration was measured using a Dr Lange Micro Color colorimeter (light source, D65; measurement geometry, d/8) yielding CIE-L*a*b* values, where L*, a*, and b* characterize the brightness, the red-green, and the yellow-blue component of the surface investigated. Optical properties were determined from
ellipsometric measurements using a home-made spectroscopic rotating analyzer ellipsometer (angle of incidence, 75 ; wavelength spectrum, 370–770 nm) [16]. Coating friction coefficients and their evolution at temperatures from 25 C up to 700 C in ambient air were studied during dry sliding experiments conducted with a ball-ondisc tribometer (CSM High-Temperature Tribometer). Alumina balls with 6 mm in diameter served as frictional counterparts in order to investigate the friction behavior at elevated temperatures without interference of hightemperature changes of the ball material. A load of 1 N and a sliding speed of 0.1 m.s 1 were used in all experiments. The radius of the wear track was set to 3, 5 or 7 mm, respectively, and the number of test cycles was adapted to reach a sliding distance of 50 m. Thermo-gravimetric analysis (TGA) was conducted for coatings deposited onto Inconel 718 substrates in a Netzsch-STA 409 C thermal analyzer in order to characterize the oxidation behavior of the coatings deposited. Dynamical TGA runs were performed for thermal ramping from 400 to 1000 C at 5 K.min 1 in Ar/O2 atmosphere (Ar/O2 flow ratio of 2.5).
3. Results and discussion Varying the number of Au pieces on the Al target, the Al/Au atomic ratio could be adjusted between 0.5 and 3.2. Fig. 1 shows representative SEM micrographs of the surface of Au–Al coatings with different Au/Al atomic ratios. High Au contents in the coating result in the formation of rough, open surfaces, whereas the increasing Al content yields a significant densification of the surface morphology. This could be a result of the different melting temperatures of Al and Au causing different thermal activation of the growth process for the individual condensing atoms [17], different diffusion rates of Al and Au leading to the formation of Kirkendall voids [18], the different phases formed in the coating (see below), or the difference between the atomic masses of the atoms involved where various collision behavior in the transport phase from the target to the substrate and during ion-assisted growth could be assumed [19]. The morphology found for the coating cross-sections summarized in Fig. 2 is very similar to the surface topography. For high Al contents, a relatively coarse-grained and dense structure was obtained (see Fig. 2a). High Au contents result in the formation of a porous coarsegrained structure (see Fig. 2c), whereas dense fibrous morphologies where obtained for Al/Au contents between 2.5 and 1.5 (see Fig. 2b). XRD results showed that the Al2Au phase is present in all coatings deposited within the Al/Au atomic ratio range between 0.5 and 3.2. For high Al or high Au contents, respectively, the presence of metallic Al (see Fig. 3, Al/Au=3.2) or Au (see Fig. 3, Al/Au41.0) was
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Fig. 1. SEM micrographs of the surface of Al–Au coated cemented carbide. (a) Al/Au=3.2, (b) Al/Au=1.6, (c) Al/Au=0.6.
Fig. 2. SEM micrographs of the fracture cross-sections of Al–Au coated cemented carbide. (a) Al/Au=3.2, (b) Al/Au=1.6, (c) Al/Au=0.6.
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Fig. 3. XRD spectra for Al–Au coated cemented carbide for different Al/Au atomic ratios (s. . .substrate).
also detected. For Al/Au41.0, the monoclinic AlAu phase was additionally formed. No evidence for the other phases present within the Al–Au system, i.e. AlAu2, Al2Au5 and AlAu4 was found. It should be mentioned that the formation of a single-phase coating consisting of the Al2Au phase results in the densest coating topography and morphology as shown in Figs. 1 and 2. The lattice parameter of the Al2Au phase is plotted in Fig. 4 against the Al/Au atomic ratio. For the stoichiometric composition (i.e. Al/Au=2), the measured lattice parameter is only slightly below the range of the literature values given in Refs. [10,22]. Increasing the Au content,
i.e. decreasing Al/Au, results in an increasing lattice parameter with a maximum for Al/Au=0.7 and decreasing tendency for the lower Al/Au ratio studied. The increasing lattice parameter could be interpreted by the incorporation of additional Au atoms in the Al2Au phase where these Au atoms cause an expansion of the lattice. For Al/Au40.7, metallic Au is present in the coating (see Fig. 3), enabling stress relaxation and probably reducing the excess Au content within the Al2Au phase. Analogously, for increasing Al/Au atomic ratios a decrease of the lattice parameter was found, which could be attributed to an Au deficiency within the Al2Au lattice.
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Fig. 4. Lattice parameter of the Al2Au phase and hardness of Al–Au coated cemented carbide as a function of the Al/Au atomic ratio.
The measured coating hardness shown in Fig. 4 is in good agreement with the content of the Al2Au phase within the coatings. For Al/Au=1. . .2, the hardness maximum with a value of about 4 GPa was obtained. Increasing the Al or the Au content beyond this range results in decreasing hardness, caused by the formation of metallic Al or Au phases, respectively, and by the open coating structure at low Al/Au atomic ratios. Most probably due to the formation of the intermetallic AlAu phase (see Fig. 3), the range of coatings with maximum hardness is somewhat expanded to Al/Au atomic ratios below the value of stoichiometry of Al2Au. Fig. 5a shows the spectral reflectivity of Al–Au coatings with different Al/Au contents. Coatings with Al/Au5 1.0 show a pink coloration, with the most pronounced color tone expressed by CIE-L*a*b* values for Al/Au= 1.6 (L*=53.9, a*=10.7, b*= 10.9). In good agreement, the spectral reflectivity for the coating with Al/Au=1.6 shows the most pronounced minimum, which was obtained for a wavelength of about 550 nm. While the wavelength position of this minimum is in excellent agreement with the literature value (given to 545 nm in Refs. [11,12]), the minimum is not as pronounced as reported for bulk Al2Au (given to about 10% at 545 nm [11]). This is due to the high defect density observed in films grown under intense ion irradiation, which often deteriorates optical properties [2,3]. For higher Al/Au atomic ratios, a flattening of the reflectivity minimum was observed (see Fig. 5a). The decrease of the Al/Au atomic ratio results first in a general decrease of the spectral reflectivity (Al/Au=1.0) followed by a significant increase of the reflectivity values in the short-wavelength region. The latter as well as the flattening of the reflectivity minimum can be explained by the decreasing Al2Au
content within the coating (see Fig. 3) and by the rough surfaces formed (see Figs. 1 and 2). Fig. 5b shows the influence of the coating thickness on the spectral reflectivity curves for coatings with Al/Au= 3.2 deposited onto (Ti,Al)N coated cemented carbides. The decreasing coating thickness shows a significant effect on the spectral reflectivity, with increased values for lower coating thicknesses, in particular in the low and high wavelength region, respectively. This corresponds to an increasing brightness L*. For the lowest coating thickness investigated (i.e. 0.1 mm), an additional shift of the reflectivity minimum to lower values was observed, which could be attributed to the increasing role of the (Ti,Al)N sublayer in determining the optical appearance. The response of selected Al–Au coatings against hightemperature exposure was investigated using vacuum annealing experiments (residual pressure, 10 4 Pa; annealing time, 1 h) as well as dynamical TGA in an argon/oxygen mixture and HT-XRD investigations. Fig. 6 shows the changes of the spectral reflectivity for a 0.1 mm thick Al–Au coating with Al/Au=2.4 deposited onto (Ti,Al)N coated cemented carbide with annealing temperature. Annealing at 500 C results in a significant improvement of the reflectivity minimum accompanied by increasing reflectivity at high wavelengths, most probably due to relaxation of defects within the coating [20,21]. For this coating, the coloration changed from pink to purple. Coatings annealed at 600 and 700 C are characterized by a shift of the curves to lower reflectivity values, without significant changes of the curve shape. The coating surface after annealing at 600 and 700 C was spotty. Investigation of these spots by EDX yielded a lower Al/Au atomic ratio in the center of the spots with respect to the overall coating composition, whereas
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Fig. 5. Spectral reflectivity of Al–Au coated cemented carbide in dependence of (a) Al/Au atomic ratio (no sublayer) and (b) various thicknesses of the Al–Au coating (Al/Au=3.2, with (Ti,Al)N sublayer).
Fig. 6. Spectral reflectivity of an Al–Au coated cemented carbide (Al/Au=2.4, Al–Au coating thickness 0.1 mm, (Ti,Al)N sublayer) as a function of annealing temperature.
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the surrounding zone was characterized by high Al contents with vanishing Au peaks. The appearance of these spots is attributed to the release of excess Al from the Al2Au phase; due to the higher diffusivity of Al, an Al-rich zone is formed around the Al-depleted Au-rich center [23]. Obviously, the formation of these spots causes the observed shift of the reflectivity curves to lower values (see Fig. 6). Dynamical TGA investigations were performed to obtain information about the oxidation behavior of 3–4 mm thick Al–Au coatings with different Al/Au contents deposited onto Inconel substrates. Although the coatings show the previously described tendency to decompose at elevated temperature, the observed mass gain at a given temperature is considerably lower compared to both, transition metal nitride-based hard coatings measured using the same equipment [24,25] and the Inconel substrate. With increasing Al2Au content, i.e. Al/Au atomic ratio approaching 2, the mass gain decreases. This indicates that the metallic Al present in the coating deteriorates the oxidation resistance. In addition, the open structure of the coating (see Figs. 1 and 2) is assumed to contribute to the lower oxidation resistance of high-Al containing coatings. XRD investigations after TGA up to 700 C only showed the presence of the Al2Au phase for all coatings; there was no evidence for crystalline oxide phases formed, which is in agreement with the negligible mass gain (see Fig. 7). After TGA up to 1000 C, no evidence for the Al2Au phase has been found by XRD; only peaks of a-Al2O3 and elemental Au have been detected. Fig. 8 shows the dependence of the Al2Au (200) relative peak intensity obtained during HT-XRD investigations for coatings deposited onto cemented carbide as a function of temperature. For all coatings
Fig. 7. TGA curves for thermal ramping at 5 K.min Inconel substrate.
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investigated, the Al2Au peaks vanish above 830–860 C, with a slightly higher thermal stability of the phase for coatings with higher Al/Au atomic ratios. It should be noted here that this temperature range is substantially lower than the melting temperature of Al2Au (Tm= 1060 C [10,11]), which points towards oxidation of the Al2Au phase. Although using He atmosphere, a-Al2O3 was detected after HT-XRD investigations up to 1100 C, since the residual oxygen in the heating chamber was sufficient for oxidation of the coating. In addition to the WC substrate peaks, Co and Au peaks were also found. Thus, we have to assume that above 830–860 C an a-Al2O3 layer is slowly formed [26] at the top of the Al–Au coating as indicated by the continuous slight mass gain during TGA. This is confirmed by the increasing mass gain for the higher Al content indicating pronounced Al segregation and a-Al2O3 formation (compare the different Al/Au ratios in Fig. 7). The remaining Au is assumed to be incorporated in Au-rich and Co-rich solid solutions formed in the Al-Co phase diagram. Fig. 9 shows the evolution of the friction curves for Al–Au coated cemented carbide disks (Al/Au=3.2) tested against alumina balls with increasing temperature. At room temperature, the average friction coefficient is 0.41 0.14, with the data scattering increasing with increasing sliding distance. At 500 and 700 C, the friction coefficients were determined to 0.54 0.26 and 0.46 0.32, respectively. Although for both temperatures large scattering of the friction coefficients was observed indicating stick-slip effects, the average friction coefficients are considerably lower compared to transition metal nitride-based hard coatings tested at the same temperatures [27,28] indicating easy shearing and thus self-lubricating properties of the coating.
in argon/oxygen atmosphere of Al–Au coated Inconel samples compared to the uncoated
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Fig. 8. Relative intensities of the (200) XRD peak of the Al2Au phase as a function of temperature for different Al/Au atomic ratios (heating rate, 0.6.K.min 1, He atmosphere).
Fig. 9. Friction coefficients obtained in ball-on-disc testing of Al–Au coated cemented carbide (Al/Au=3.2) sliding against alumina balls in ambient air at different temperatures using a load of 1 N and a sliding speed of 0.1 m.s 1.
4. Conclusions The aim of this work was to investigate the system Al–Au to elucidate the potential of Al2Au-containing coatings as a colored self-lubricating layer on tools. Using co-sputtering from an Al target partially covered by Au pieces, the chemical coating composition could be adjusted to obtain single-phase Al2Au coatings showing a dense fine-grained morphology. The hardness of the single-phase coating yielded a maximum value of 4 GPa. In the as-deposited state, single-phase Al2Au coatings show a pink coloration turning to purple after
annealing at 500 C. The coatings are thermally stable up to about 850 C, where the onset for oxidation and a-Al2O3 formation occurs. Friction experiments yielded relatively low friction coefficients of 0.4 to 0.5 for the temperature range up to 700 C, with increasing stick-slip effects at higher temperatures. Finally, it is concluded that Al2Au-containing coatings could be potential candidates for thin colored and self-lubricating layers deposited onto coated cemented carbide tools. In addition, these coatings may be advantageously used to protect Inconel alloys in high temperature applications.
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Acknowledgements The authors are grateful to Gerhard Hawranek for SEM characterization, to Manfred Pock for tribological testing and to Bernhard Sartory for HT-XRD investigations. References [1] North B. Surf Coat Technol 1998;106:129. [2] Mitterer C, Mayrhofer PH, Waldhauser W, Kelesoglu E, Losbichler P. Surf Coat Technol 1998;108-109:230. [3] Mitterer C, Waldhauser W, Beck U, Reiners G. Surf Coat Technol 1996;86/87:715. [4] Grill A. Diam Relat Mater 1999;8(2-5):428. [5] Gilmore R, Baker MA, Gibson PN, Gissler W, Stoiber M, Losbichler P, Mitterer C. Surf Coat Technol 1998;108-109:345. [6] Stoiber M, Badisch E, Lugmair C, Mitterer C. Surf Coat Technol 2003;163-164:451. [7] Badisch E, Mitterer C, Mayrhofer PH, Mori G, Bakker RJ, Brenner J, Sto¨ri H. Thin Solid Films, accepted. [8] Andersson J, Erck RA, Erdemir A. Wear 2003;254:1070. [9] Badisch E, Fontalvo G, Stoiber M, Mitterer C. Surf Coat Technol 2003;163-164:585. [10] Cahn RW. Nature 1998;396:523. [11] Shih HC. Z Metallkde 1980;71:577. [12] Hsu L-S, Guo G-Y, Denlinger JD, Allen JW. J Phys Chem Solids 2001;62:1047.
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[13] Eskner M, Sandstro¨m R. Surf Coat Technol 2003;165:71. [14] Losbichler P, Mitterer C. Surf Coat Technol 1997;97:568. [15] Losbichler P, Mitterer C, Gibson PN, Gissler W, Hofer F, Warbichler P. Surf Coat Technol 1997;94/95:297. [16] Jascht G, Mayrhofer PH, Mitterer C. Berg- und Hu¨ttenma¨nnische Monatshefte 1997;142(1):35. [17] Thornton JA. J Vac Sci Technol 1974;11(4):666. [18] Raymond T. Semicond Int 1989;12:152. [19] Mitterer C. J Solid State Chem 1997;133:279. [20] Mayrhofer PH, Tischler G, Mitterer C. Surf Coat Technol 2001; 142–144:78. [21] Mayrhofer PH, Kunc F, Musil J, Mitterer C. Thin Solid Films 2002;415:151. [22] Powder Diffraction File, Joint Committee on Powder Diffraction Standards, International Center for Diffraction Data, Swarthmore, PA, card 17–877. [23] de Waal HS, Pretorius R, Prozesky VM, Churms CL. Nucl Instr and Methods B 1997;130:722. [24] Mitterer C, Holler F, U¨stel F, Heim D. Surf Coat Technol 2000; 125(1–3):233. [25] Mayrhofer PH, Willmann H, Mitterer C. Surf Coat Technol 2001;146–147:222. [26] Svensson H, Angenete J, Stiller J. Surf Coat Technol 2004;177– 178:152. [27] Badisch E, Fontalvo G, Stoiber M, Mitterer C. Surf Coat Technol 2003;163–164:585. [28] Kutschej, K, Mayrhofer PH, Kathrein M., Polcik P, Tessadri R, Mitterer C. Surf. Coat. Technol., submitted.
Sputter-deposited Al–Au coatings C. Mitterera,*, H. Lenharta,1, P.H. Mayrhofera, M. Kathreinb a
Department of Physical Metallurgy and Materials Testing, University of Leoben, Franz-Josef-Strasse 18, A-8700 Leoben, Austria b CERATIZIT Austria GmbH, A-6600 Reutte, Austria Received 5 January 2004; accepted 13 February 2004
Abstract Transition metal nitride-based, wear-resistant hard coatings on cutting tools and other substrates often lack self-lubricating properties at elevated temperatures and distinct colorations allowing product differentiation. In the present work, the possibility of achieving these objectives using coatings based on the purple-red Al2Au phase within the Al–Au system was investigated. Coatings within this system have been deposited onto cemented carbide and Inconel substrates using unbalanced magnetron sputtering. The coatings were characterized with respect to their topography and morphology, chemical and phase composition, hardness, optical, oxidation and tribological properties. Al2Au-containing coatings have been deposited with dense, fine-grained structures yielding a hardness of 4 GPa and pink coloration. Vacuum annealing at temperatures of 500 C results in a chance to a pronounced purple coloration. The coatings are stable up to about 850 C, where the onset of oxidation occurs. Low friction coefficients when testing against alumina counterparts were achieved in the temperature range between 500 and 700 C. The concept of applying Al2Aucontaining coatings as a colored self-lubricating layer on top of a hard coated cemented carbide tool warrants further investigations. # 2004 Elsevier Ltd. All rights reserved. Keywords: A. Aluminides, miscellaneous; B. Phase identification; B. Tribological properties; B. Thermal stability; C. Coatings, intermetallic and otherwise
1. Introduction Hard coatings, e.g. based on transition metal nitrides, are widely applied to improve the lifetime and the performance of many kinds of cutting and forming tools. Besides the functionality of the coating itself, provided by e.g. hardness, toughness, oxidation and wear resistance, there is also an increasing demand for cosmetic factors like attractive colorations of a cutting tool [1]. Fortunately, the most widely applied titanium nitride TiN is characterized by a bright golden-yellow color [2] making determination of used cutting edges of a cutting insert in poor lighting conditions of a machine shop easier. However, for other hard compounds there exists a considerable lack of attractive bright colors [3] allowing product differentiation and to distinguish used and unused cutting edges. On the other hand, during recent years tremendous effort has been laid on development * Corresponding author. Tel.: +43-3842-402-4220; fax: +43-3842402-4202. E-mail address: [email protected] (C. Mitterer). 1 Present address: Steirische Wirtschaftsfo¨rderung, Nikolaiplatz 2, A-8020 Graz, Austria. 0966-9795/$ - see front matter # 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2004.02.003
of low-friction coatings being able to reduce friction forces. Today, low-friction coatings rely on solid lubricant phases like diamond-like carbon DLC [4], molybdenum disulfide MoS2 [5] or chlorine-induced rutile formation on chlorine-containing TiN coatings [6,7]. These lubricants often begin to fail in their tribological effectiveness with increasing temperature, in humid atmosphere or due to oxidation [8,9]. Consequently, this study is focused on the development of a new type of colored low-friction coatings based on the intermetallic Al2Au phase. Representing a Zintl phase with the cubic CaF2 structure [10], Al2Au shows a distinct reflectivity minimum at 545 nm [11,12], resulting in a purple-red color. Being relatively hard and brittle at room temperature, aluminides like Al2Au show plastic deformation and possibly self-lubrication at high temperatures [13]. Within this work, we present a comprehensive overview on the sputter deposition of Al–Au coatings and on the characterization of coating composition and microstructure as well as optical, mechanical, oxidation, and tribological properties. Special emphasis will also be laid on the thermal stability of the coatings.
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2. Experimental details The coatings investigated in this work have been deposited using an unbalanced d.c. magnetron sputtering system that has been described in detail in Refs. [14,15]. Different Al/Au atomic ratios have been adjusted by placing different numbers of Au pieces (20 201 mm3) on the erosion track of an Al target (11506 mm) mounted onto an unbalanced Gencoa PP 150 magnetron. The target was positioned 9 cm from the parallel-plate substrate assembly. Coatings were deposited onto uncoated and (Ti,Al)N coated cemented carbide cutting inserts of different grades and geometries and Inconel 718 substrates (151 mm) in Ar atmosphere at 0.2 Pa. The substrates were positioned directly above the target erosion track to ensure a homogeneous Al/Au distribution within the coating. For all deposition runs, the substrate temperature was held constant at 300 C, the bias voltage was 50 V, and the target sputter power was about 380 W. Prior to deposition, all substrates used were metallographically grounded, polished and ultrasonically cleaned with ethylene and acetone. After target pre-cleaning and ion etching of the substrates within the deposition chamber, coatings in the thickness range between 0.1 and 7 mm were deposited by selecting suitable deposition times. Coatings with thicknesses of 3–7 mm were used for the basic chemical, microstructural mechanical, and tribological characterization, whereas for the investigation of the optical properties also thinner coatings were used. Coating topography and morphology were investigated using scanning electron microscopy (SEM, Cambridge Instruments Stereoscan 360). For the investigation of the coating morphology, cross-sections of coated cemented carbide inserts were prepared by brittle fracturing in liquid nitrogen. The chemical composition of the coatings was determined by energy-dispersive X-ray analysis (EDX, Link eXL) using Au and Al elemental standards. X-ray diffraction (XRD) patterns were recorded using a Siemens D 500 Bragg-Brentano diffractometer and Cu-Ka radiation. Lattice parameters were corrected over cos2y. In addition, high-temperature XRD (HT-XRD, Siemens D5005 with heating chamber Paar HTK 1200, BraggBrentano configuration, heating rate, 0.6 K.min 1) investigations in He atmosphere were performed for coatings deposited onto cemented carbide inserts to obtain information on the thermal stability of the coatings. Coating hardness was characterized using a Vickers microhardness tester (Reichert-Jung MD 4000) at a load of 25 mN. Coating coloration was measured using a Dr Lange Micro Color colorimeter (light source, D65; measurement geometry, d/8) yielding CIE-L*a*b* values, where L*, a*, and b* characterize the brightness, the red-green, and the yellow-blue component of the surface investigated. Optical properties were determined from
ellipsometric measurements using a home-made spectroscopic rotating analyzer ellipsometer (angle of incidence, 75 ; wavelength spectrum, 370–770 nm) [16]. Coating friction coefficients and their evolution at temperatures from 25 C up to 700 C in ambient air were studied during dry sliding experiments conducted with a ball-ondisc tribometer (CSM High-Temperature Tribometer). Alumina balls with 6 mm in diameter served as frictional counterparts in order to investigate the friction behavior at elevated temperatures without interference of hightemperature changes of the ball material. A load of 1 N and a sliding speed of 0.1 m.s 1 were used in all experiments. The radius of the wear track was set to 3, 5 or 7 mm, respectively, and the number of test cycles was adapted to reach a sliding distance of 50 m. Thermo-gravimetric analysis (TGA) was conducted for coatings deposited onto Inconel 718 substrates in a Netzsch-STA 409 C thermal analyzer in order to characterize the oxidation behavior of the coatings deposited. Dynamical TGA runs were performed for thermal ramping from 400 to 1000 C at 5 K.min 1 in Ar/O2 atmosphere (Ar/O2 flow ratio of 2.5).
3. Results and discussion Varying the number of Au pieces on the Al target, the Al/Au atomic ratio could be adjusted between 0.5 and 3.2. Fig. 1 shows representative SEM micrographs of the surface of Au–Al coatings with different Au/Al atomic ratios. High Au contents in the coating result in the formation of rough, open surfaces, whereas the increasing Al content yields a significant densification of the surface morphology. This could be a result of the different melting temperatures of Al and Au causing different thermal activation of the growth process for the individual condensing atoms [17], different diffusion rates of Al and Au leading to the formation of Kirkendall voids [18], the different phases formed in the coating (see below), or the difference between the atomic masses of the atoms involved where various collision behavior in the transport phase from the target to the substrate and during ion-assisted growth could be assumed [19]. The morphology found for the coating cross-sections summarized in Fig. 2 is very similar to the surface topography. For high Al contents, a relatively coarse-grained and dense structure was obtained (see Fig. 2a). High Au contents result in the formation of a porous coarsegrained structure (see Fig. 2c), whereas dense fibrous morphologies where obtained for Al/Au contents between 2.5 and 1.5 (see Fig. 2b). XRD results showed that the Al2Au phase is present in all coatings deposited within the Al/Au atomic ratio range between 0.5 and 3.2. For high Al or high Au contents, respectively, the presence of metallic Al (see Fig. 3, Al/Au=3.2) or Au (see Fig. 3, Al/Au41.0) was
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Fig. 1. SEM micrographs of the surface of Al–Au coated cemented carbide. (a) Al/Au=3.2, (b) Al/Au=1.6, (c) Al/Au=0.6.
Fig. 2. SEM micrographs of the fracture cross-sections of Al–Au coated cemented carbide. (a) Al/Au=3.2, (b) Al/Au=1.6, (c) Al/Au=0.6.
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Fig. 3. XRD spectra for Al–Au coated cemented carbide for different Al/Au atomic ratios (s. . .substrate).
also detected. For Al/Au41.0, the monoclinic AlAu phase was additionally formed. No evidence for the other phases present within the Al–Au system, i.e. AlAu2, Al2Au5 and AlAu4 was found. It should be mentioned that the formation of a single-phase coating consisting of the Al2Au phase results in the densest coating topography and morphology as shown in Figs. 1 and 2. The lattice parameter of the Al2Au phase is plotted in Fig. 4 against the Al/Au atomic ratio. For the stoichiometric composition (i.e. Al/Au=2), the measured lattice parameter is only slightly below the range of the literature values given in Refs. [10,22]. Increasing the Au content,
i.e. decreasing Al/Au, results in an increasing lattice parameter with a maximum for Al/Au=0.7 and decreasing tendency for the lower Al/Au ratio studied. The increasing lattice parameter could be interpreted by the incorporation of additional Au atoms in the Al2Au phase where these Au atoms cause an expansion of the lattice. For Al/Au40.7, metallic Au is present in the coating (see Fig. 3), enabling stress relaxation and probably reducing the excess Au content within the Al2Au phase. Analogously, for increasing Al/Au atomic ratios a decrease of the lattice parameter was found, which could be attributed to an Au deficiency within the Al2Au lattice.
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Fig. 4. Lattice parameter of the Al2Au phase and hardness of Al–Au coated cemented carbide as a function of the Al/Au atomic ratio.
The measured coating hardness shown in Fig. 4 is in good agreement with the content of the Al2Au phase within the coatings. For Al/Au=1. . .2, the hardness maximum with a value of about 4 GPa was obtained. Increasing the Al or the Au content beyond this range results in decreasing hardness, caused by the formation of metallic Al or Au phases, respectively, and by the open coating structure at low Al/Au atomic ratios. Most probably due to the formation of the intermetallic AlAu phase (see Fig. 3), the range of coatings with maximum hardness is somewhat expanded to Al/Au atomic ratios below the value of stoichiometry of Al2Au. Fig. 5a shows the spectral reflectivity of Al–Au coatings with different Al/Au contents. Coatings with Al/Au5 1.0 show a pink coloration, with the most pronounced color tone expressed by CIE-L*a*b* values for Al/Au= 1.6 (L*=53.9, a*=10.7, b*= 10.9). In good agreement, the spectral reflectivity for the coating with Al/Au=1.6 shows the most pronounced minimum, which was obtained for a wavelength of about 550 nm. While the wavelength position of this minimum is in excellent agreement with the literature value (given to 545 nm in Refs. [11,12]), the minimum is not as pronounced as reported for bulk Al2Au (given to about 10% at 545 nm [11]). This is due to the high defect density observed in films grown under intense ion irradiation, which often deteriorates optical properties [2,3]. For higher Al/Au atomic ratios, a flattening of the reflectivity minimum was observed (see Fig. 5a). The decrease of the Al/Au atomic ratio results first in a general decrease of the spectral reflectivity (Al/Au=1.0) followed by a significant increase of the reflectivity values in the short-wavelength region. The latter as well as the flattening of the reflectivity minimum can be explained by the decreasing Al2Au
content within the coating (see Fig. 3) and by the rough surfaces formed (see Figs. 1 and 2). Fig. 5b shows the influence of the coating thickness on the spectral reflectivity curves for coatings with Al/Au= 3.2 deposited onto (Ti,Al)N coated cemented carbides. The decreasing coating thickness shows a significant effect on the spectral reflectivity, with increased values for lower coating thicknesses, in particular in the low and high wavelength region, respectively. This corresponds to an increasing brightness L*. For the lowest coating thickness investigated (i.e. 0.1 mm), an additional shift of the reflectivity minimum to lower values was observed, which could be attributed to the increasing role of the (Ti,Al)N sublayer in determining the optical appearance. The response of selected Al–Au coatings against hightemperature exposure was investigated using vacuum annealing experiments (residual pressure, 10 4 Pa; annealing time, 1 h) as well as dynamical TGA in an argon/oxygen mixture and HT-XRD investigations. Fig. 6 shows the changes of the spectral reflectivity for a 0.1 mm thick Al–Au coating with Al/Au=2.4 deposited onto (Ti,Al)N coated cemented carbide with annealing temperature. Annealing at 500 C results in a significant improvement of the reflectivity minimum accompanied by increasing reflectivity at high wavelengths, most probably due to relaxation of defects within the coating [20,21]. For this coating, the coloration changed from pink to purple. Coatings annealed at 600 and 700 C are characterized by a shift of the curves to lower reflectivity values, without significant changes of the curve shape. The coating surface after annealing at 600 and 700 C was spotty. Investigation of these spots by EDX yielded a lower Al/Au atomic ratio in the center of the spots with respect to the overall coating composition, whereas
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Fig. 5. Spectral reflectivity of Al–Au coated cemented carbide in dependence of (a) Al/Au atomic ratio (no sublayer) and (b) various thicknesses of the Al–Au coating (Al/Au=3.2, with (Ti,Al)N sublayer).
Fig. 6. Spectral reflectivity of an Al–Au coated cemented carbide (Al/Au=2.4, Al–Au coating thickness 0.1 mm, (Ti,Al)N sublayer) as a function of annealing temperature.
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the surrounding zone was characterized by high Al contents with vanishing Au peaks. The appearance of these spots is attributed to the release of excess Al from the Al2Au phase; due to the higher diffusivity of Al, an Al-rich zone is formed around the Al-depleted Au-rich center [23]. Obviously, the formation of these spots causes the observed shift of the reflectivity curves to lower values (see Fig. 6). Dynamical TGA investigations were performed to obtain information about the oxidation behavior of 3–4 mm thick Al–Au coatings with different Al/Au contents deposited onto Inconel substrates. Although the coatings show the previously described tendency to decompose at elevated temperature, the observed mass gain at a given temperature is considerably lower compared to both, transition metal nitride-based hard coatings measured using the same equipment [24,25] and the Inconel substrate. With increasing Al2Au content, i.e. Al/Au atomic ratio approaching 2, the mass gain decreases. This indicates that the metallic Al present in the coating deteriorates the oxidation resistance. In addition, the open structure of the coating (see Figs. 1 and 2) is assumed to contribute to the lower oxidation resistance of high-Al containing coatings. XRD investigations after TGA up to 700 C only showed the presence of the Al2Au phase for all coatings; there was no evidence for crystalline oxide phases formed, which is in agreement with the negligible mass gain (see Fig. 7). After TGA up to 1000 C, no evidence for the Al2Au phase has been found by XRD; only peaks of a-Al2O3 and elemental Au have been detected. Fig. 8 shows the dependence of the Al2Au (200) relative peak intensity obtained during HT-XRD investigations for coatings deposited onto cemented carbide as a function of temperature. For all coatings
Fig. 7. TGA curves for thermal ramping at 5 K.min Inconel substrate.
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investigated, the Al2Au peaks vanish above 830–860 C, with a slightly higher thermal stability of the phase for coatings with higher Al/Au atomic ratios. It should be noted here that this temperature range is substantially lower than the melting temperature of Al2Au (Tm= 1060 C [10,11]), which points towards oxidation of the Al2Au phase. Although using He atmosphere, a-Al2O3 was detected after HT-XRD investigations up to 1100 C, since the residual oxygen in the heating chamber was sufficient for oxidation of the coating. In addition to the WC substrate peaks, Co and Au peaks were also found. Thus, we have to assume that above 830–860 C an a-Al2O3 layer is slowly formed [26] at the top of the Al–Au coating as indicated by the continuous slight mass gain during TGA. This is confirmed by the increasing mass gain for the higher Al content indicating pronounced Al segregation and a-Al2O3 formation (compare the different Al/Au ratios in Fig. 7). The remaining Au is assumed to be incorporated in Au-rich and Co-rich solid solutions formed in the Al-Co phase diagram. Fig. 9 shows the evolution of the friction curves for Al–Au coated cemented carbide disks (Al/Au=3.2) tested against alumina balls with increasing temperature. At room temperature, the average friction coefficient is 0.41 0.14, with the data scattering increasing with increasing sliding distance. At 500 and 700 C, the friction coefficients were determined to 0.54 0.26 and 0.46 0.32, respectively. Although for both temperatures large scattering of the friction coefficients was observed indicating stick-slip effects, the average friction coefficients are considerably lower compared to transition metal nitride-based hard coatings tested at the same temperatures [27,28] indicating easy shearing and thus self-lubricating properties of the coating.
in argon/oxygen atmosphere of Al–Au coated Inconel samples compared to the uncoated
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Fig. 8. Relative intensities of the (200) XRD peak of the Al2Au phase as a function of temperature for different Al/Au atomic ratios (heating rate, 0.6.K.min 1, He atmosphere).
Fig. 9. Friction coefficients obtained in ball-on-disc testing of Al–Au coated cemented carbide (Al/Au=3.2) sliding against alumina balls in ambient air at different temperatures using a load of 1 N and a sliding speed of 0.1 m.s 1.
4. Conclusions The aim of this work was to investigate the system Al–Au to elucidate the potential of Al2Au-containing coatings as a colored self-lubricating layer on tools. Using co-sputtering from an Al target partially covered by Au pieces, the chemical coating composition could be adjusted to obtain single-phase Al2Au coatings showing a dense fine-grained morphology. The hardness of the single-phase coating yielded a maximum value of 4 GPa. In the as-deposited state, single-phase Al2Au coatings show a pink coloration turning to purple after
annealing at 500 C. The coatings are thermally stable up to about 850 C, where the onset for oxidation and a-Al2O3 formation occurs. Friction experiments yielded relatively low friction coefficients of 0.4 to 0.5 for the temperature range up to 700 C, with increasing stick-slip effects at higher temperatures. Finally, it is concluded that Al2Au-containing coatings could be potential candidates for thin colored and self-lubricating layers deposited onto coated cemented carbide tools. In addition, these coatings may be advantageously used to protect Inconel alloys in high temperature applications.
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