- Email: [email protected]
Hard and superhard TiAlBN coatings deposited by twin electron-beam evaporation
Surface & Coatings Technology 201 (2007) 6078 – 6083 www.elsevier.com/locate/surfcoat
Hard and superhard TiAlBN coatings deposited by twin electron-beam evaporation C. Rebholz a,⁎, M.A. Monclus b , M.A. Baker b , P.H. Mayrhofer c , P.N. Gibson d , A. Leyland e , A. Matthews e a
Department of Mechanical and Manufacturing Engineering, University of Cyprus, 1678 Nicosia, Cyprus The Surface Analysis Laboratory, School of Engineering, University of Surrey, Guildford GU2 7XH, UK Department of Physical Metallurgy and Materials Testing, University of Leoben, Franz-Josef-Strasse 18, A-8700 Leoben, Austria d Institute for Health and Consumer Protection, Joint Research Centre, I-21020 Ispra (VA), Italy e Department of Engineering Materials, University of Sheffield, Sheffield S1 3JD, UK b
c
Available online 2 October 2006
Abstract Superhard nanostructured coatings, prepared by plasma-assisted chemical vapour deposition (PACVD) and physical vapour deposition (PAPVD) techniques, such as vacuum arc evaporation and magnetron sputtering, are receiving increasing attention due to their potential applications for wear protection. In this study nanocomposite (TiAl)BxNy (0.09 ≤ x ≤ 1.35; 1.07 ≤ y ≤ 2.30) coatings, consisting of nanocrystalline (Ti,Al)N and amorphous BN, were deposited onto Si (100), AISI 316 stainless steel and AISI M2 tool steel substrates by co-evaporation of Ti and hot isostatically pressed (HIPped) Ti–Al–B–N material from a thermionically enhanced twin crucible electron-beam (EB) evaporation source in an Ar plasma at 450 °C. The coating stoichiometry, relative phase composition, nanostructure and mechanical properties were determined using X-ray photoelectron spectroscopy (XPS) and X-ray diffraction (XRD), in combination with nanoindentation measurements. Aluminium (∼10 at.% in coatings) was found to substitute for titanium in the cubic TiN based structure. (Ti,Al)B0.14N1.12 and (Ti,Al)B0.45N1.37 coatings with average (Ti,Al)N grain sizes of 5–6 nm and either ∼70, or ∼90, mol% (Ti,Al)N showed hardness and elastic modulus values of ∼40 and ∼340 GPa, respectively. (Ti,Al)B0.14N1.12 coatings retained their ‘as-deposited’ mechanical properties for more than 90 months at room temperature in air, comparing results gathered from eight different nanoindentation systems. During vacuum annealing, all coatings examined exhibited structural stability to temperatures in excess of 900 °C, and revealed a moderate, but significant, increase in hardness. For (Ti,Al)B0.14N1.12 coatings the hardness increased from ∼40 to ∼45 GPa. © 2006 Elsevier B.V. All rights reserved. Keywords: Superhard nanocomposites; Nanoindentation; Thermal stability; Electron-beam evaporation; Physical vapour deposition (PVD)
1. Introduction In the last few years, advances in coating deposition technologies have led to the development of nanostructured coating materials with unique properties, such as superhardness (defined as hardness H N 40 GPa) in combination with high toughness [1,2]. Coatings prepared by plasma-assisted chemical vapour deposition (PACVD) and physical vapour deposition (PAPVD) are often known to exhibit increased hardness compared to their corresponding bulk counterparts, due to factors such as grain size refinement, high growth-defect density and
⁎ Corresponding author. Tel.: +352 22 892282; fax: +357 22 892254. E-mail address: [email protected] (C. Rebholz). 0257-8972/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2006.08.121
residual compressive stress, induced by the deposition process [3]. In binary coatings with large single-phase fields these hardness enhancement effects can diminish rapidly during exposure to temperatures above the deposition temperature (e.g. in severe machining conditions), as recovery and recrystallisation processes occur [3]. The growing number of nanocomposite coating systems that exhibit superhardness together with high thermal stability include coatings within the ternary Ti–Si–N [4–6] and Ti–B–N [7–9] material systems, deposited by PACVD, and by PAPVD techniques such as magnetron sputtering or vacuum arc evaporation. A detailed recent review of the different approaches to the design and synthesis of superhard nanostructured coatings can be found in ref. [10]. A continuing issue, that superhard coating deposition particularly has brought to prominence, is how to measure
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
correctly the mechanical properties of these and of ultrahard (H N 70 GPa) coatings using sub-micron indentation testing, and Fischer–Cripps has highlighted the limitations and common sources of error in two recent reviews [11,12]. In contrast to superhard coatings in the Ti–B–N ternary system (i.e. TiN/TiB2; TiN/BN), only a few studies have been reported on quaternary TiAlBN coatings — especially those deposited by electron-beam PVD (EBPVD) [13–16]. A maximum hardness value of 37 GPa and superior wet-cutting performance (compared to standard TiN and TiAlN) were observed for a dual-phase coating effectively consisting of nanocrystalline (nc-) (Ti,Al)N and amorphous (a-) BN, in approximate proportions 90 and 10 mol%, respectively [15]. No data on thermal and long-term mechanical/structural stability for such quaternary coatings are available in the literature; however, it is known that coatings within the Ti–B– N ternary system possess high-temperature thermal stability to above 900 °C [8–10] and long-term structural stability at room temperature in air [10]. The main objective of the present work was therefore to determine the dependence of the structure, density, mechanical properties and thermal stability on composition for twincrucible EBPVD nc-(Ti,Al)N/a-BN coatings with Al and N contents of approximately 10 and 50 at.%, respectively, and varying B and Ti contents. Furthermore, the present work was also directed towards a desire to compare the mechanical property data of superhard nc-(Ti,Al)N/a-BN coatings (obtained from different nanoindentation systems) during a long-term stability study, thereby gaining a better understanding of the limitations of the various measuring systems used. 2. Experimental details TiAlBN coatings, 2.3 ± 0.1 μm thick, were deposited onto Si (100), AISI 316 stainless steel and AISI M2 tool steel substrates by co-evaporating Ti and HIPped Ti–Al–B–N material from a thermionically enhanced twin-crucible EB evaporation source system (TECVAC IP35L) with a base pressure of b 5 × 10− 4 Pa. A combination of optical emission spectroscopy (OES) and partial pressure control was utilised to control the evaporation rates and hence the composition in the coating. Fig. 1 shows a schematic of the deposition system, illustrating crucible location and evaporant materials employed. Commercially available hot-pressed Ti–Al–B–N evaporation boats [17] with dimensions of 10 mm × 15 mm × 120 mm (Fig. 2a), consisting of a mixture of 50 wt.% TiB2, 30 wt.% BN and 20 wt.% AlN, were used to create the Ti–Al–B–N evaporant material; a scanning electron microscopy (SEM) image of the as-received morphology can be seen in Fig. 2b. These boats were hot isostatically pressed (HIPped) to densify the material (Fig. 2c) and reduce spitting during EB evaporation [13], and then broken up into smaller granules (of approximately 7 mm mean diameter) to place in the crucible. The crucible is shown in operation in Fig. 2d, including the filaments used for plasma enhancement; the source material evaporates by sublimation in the middle of the crucible, where the EB gun scans over the
6079
Fig. 1. Schematic drawing of the TECVAC IP35L twin EB evaporation system used for nc-(Ti,Al)N/a-BN coating deposition.
material. Prior to deposition, the AISI 316 stainless steel substrates were ultrasonically cleaned in acetone and isopropanol and then fixed in a single-rotation holder mounted 300 mm above the vapour source. The deposition process consisted of six steps: (i) radiation heating to 450 °C (ii) diode sputter cleaning and plasma heating in an Ar atmosphere at −1000 V (iii) deposition of a thin 0.2 μm Ti interlayer (iv) deposition of a 0.5 μm TiN layer (N2 as reactive gas) and monitoring of the OES signal to establish a stable (and reproducible) Ti vapour–flux proportion; (v) reduction of the Ti evaporation rate to a selected percentage of the initial Ti OES signal and removal of N2 flow, followed by (vi) twin-EB evaporation of Ti and Ti–Al–B–N. Stable evaporation of the Ti (at a selected percentage of the initial rate) and of the Ti– Al–B–N material were achieved using combined OES and partial pressure control. The power on both crucibles was adjusted during deposition to keep the Ti evaporation rate constant and the total pressure constant at 0.35 Pa. Substrate temperature and bias voltage were kept constant at 450 °C and −100 V, respectively. Both chemical and phase composition of the various coatings were determined by X-ray photoelectron spectroscopy (XPS), using a VG-Scientific Sigma Probe spectrometer employing a monochromated Al–Kα source and a hemispherical analyser [15,16]. The crystallographic structure and texture of the films were analysed by glancing-angle X-ray diffraction (GAXRD), using CuKα radiation at incident angles ranging from 0.5 to 4.0°. The X-ray generator settings were 35 kV and 30 mA, the step angle being 0.2°. The laserinduced surface acoustic wave (L-SAW) technique [18] was used to obtain coating density values from coated M2 and Si (100) samples. Fracture cross-sections of coated samples were prepared for SEM morphology and topography studies. Film thickness was measured using a Veeco DEKTAK 3ST
6080
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
Fig. 2. (a) Photograph of hot-pressed Ti–Al–B–N evaporation boats, SEM images of (b) hot-pressed and (c) HIPped Ti–Al–B–N boats, and (d) photograph of a crucible with broken up HIPped Ti–Al–B–N evaporation material.
profilometer with a vertical resolution of 50 nm and verified by SEM cross-sectional measurements. As-deposited coatings on AISI 316 substrates were annealed in vacuum for 30 min at temperatures Ta of 600, 700, 800 and 900 °C and a pressure of ≤10− 3 Pa. The hardness H and elastic modulus E values were obtained from nanoindentation measurements at room temperature using a Hysitron Triboscope instrument attached to an atomic force microscope. The area function of the triangular Berkovich diamond tip was calibrated following the procedure described in Ref. [19]. The maximum load was varied from 3 to 6 mN with a sequence of 30 indents at each load; the maximum indentation depth was always less than 10% of the top TiAlBN layer thickness. For long-term stability studies at room temperature in air, a superhard nc-(Ti,Al)N/a-BN coated M2 substrate was analysed over a period of 92 months using different nanoindentation systems at various European research institutes. The systems used, in chronological order, were: (i) Nanoindenter XP, Department of Physics, Aristotle University of Thessaloniki, Greece; (ii) Nanoindenter II, Department of Physics, University of Linköping, Sweden; (iii) Hysitron Triboscope, Department of Engineering Materials, University of Sheffield, UK; (iv) Fischerscope H100, Robert Bosch GmbH, Stuttgart, Germany; (v) Fischerscope H100, Helmut Fischer GmbH & Co.KG, SindelfingenMaichingen, Germany; (vi) CSM Nanohardnesstester (NHT), CSM Instruments SA, Peseux, Switzerland; (vii) Nanoindenter XP, Robert Bosch GmbH, Stuttgart, Germany and (viii) Hysitron Triboscope, Materials Chemistry, RWTH Aachen, Germany. With the exception of the Fischerscope H100 instruments, all systems were equipped with a Berkovich diamond indenter; the maximum indentation depth was 150 nm, i.e. ≤10% thickness of top TiAlBN layer, and the load frame compliance and tip correction function were estimated using sapphire and/or fused silica control samples [19]. Vickers diamond indenters were used for the Fischerscope H100 instruments, with linear extrapolation
of the unloading curve to determine the contact stiffness and the corrected contact depth; a load of 30 mN was used in each case. 3. Results and discussion Coating compositions, determined from XPS Ti 2p, Al 2p, B 1s and N 1s peak areas [15,16], are given in Table 1, and are plotted on the Ti–B–N phase diagram of Nowotny et al. [20] modified for the inclusion of Al into the Ti–B–N system (Fig. 3). All coatings possess approximately constant Al (∼ 10 at.%) and N (∼ 50 at.%) content, but varying B and Ti concentration, and are located close to the (Ti,Al)N–BN tie line. In accordance with their position in the phase diagram, all coatings exhibit a three-phase composition of (Ti,Al)N and BN with very small amounts of TiB2 (≤ 3 mol%), i.e. the microstructure is effectively that of a (Ti,Al)N and BN dualphase coating [15,16]. With increasing B concentration, a colour change from dark gold to metallic black was observed. Fig. 4a shows a typical SEM fracture cross-section representative of twin EB nc-(Ti,Al)N/a-BN coatings produced at 450 °C. The individual Ti, TiN and TiAlBN layers are clearly visible. In contrast to the columnar morphology of the Ti layer, the TiN and TiAlBN layers exhibit a dense, featureless ‘glassy’ Table 1 Elemental concentrations and stoichiometry for the TiAlBN coatings determined by XPS [15,16] Elemental concentration [at.%]
Stoichiometry
Ti
Al
B
N
39.1 36.2 25.3 16.8 10.1
7.2 8.0 10.1 9.0 11.4
4.2 6.2 16.0 26.4 29.1
49.5 49.5 48.6 47.8 49.4
Sum of O, C and Ar b 3 at.% in all coatings.
(Ti,Al)B0.09N1.07 (Ti,Al)B0.14N1.12 (Ti,Al)B0.45N1.37 (Ti,Al)B1.02N1.85 (Ti,Al)B1.35N2.30
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
Fig. 3. Chemical composition (as determined by XPS) of deposited nc-(Ti,Al)N/ a-BN coatings within a modified Ti–B–N equilibrium phase diagram [20].
structure. Differences in surface topography for coatings with low B concentrations deposited from hot-pressed (Fig. 2b) and HIPped (Fig. 2c) Ti–Al–B–N source material can be seen in Fig. 4b and c, respectively. The hot-pressed material yielded a significantly higher surface roughness, while using the denser HIPped material resulted in virtually ‘spit-free’ coatings and improved coating integrity. GAXRD patterns at 0.5° incident angle for the nc-(Ti,Al)N/aBN coatings are plotted in Fig. 5. At this angle of incidence only the surface structure (in the TiAIBN top layer) is detected. The only peaks obtained over the 2Θ range 10–90° were for (111), (200), (220) and (311) reflections of a face-centred cubic TiNlike phase, in which Al is assumed to substitute for Ti [14,15]. No evidence of other crystalline phases (TiB2 and/or BN) was found, in agreement with previously reported transmission electron microscopy (TEM) studies [15]. Shtansky et al. [21] also only observed the TiN phase in TEM studies for a TiAl0.3B0.5N1.9 (equivalent to (Ti,Al)B0.38N1.46) coating located close to the (Ti,Al)N–BN tie line and deposited by magnetron sputtering from a Ti–Al–B–N target (manufactured by selfpropagating high-temperature synthesis from a mixture of Ti, Al
Fig. 4. (a) SEM micrograph of fracture cross-section representative for nc-(Ti, Al)N/a-BN coatings, and surface topographies for coatings with low B contents deposited from (b) hot-pressed and (c) HIPped Ti–Al–B–N source material.
6081
Fig. 5. GAXRD patterns at 0.5° for nc-(Ti,Al)N/a-BN coatings.
and B, with 66.7, 20 and 13.3 wt.%, respectively) in Ar/N2 mixtures at a substrate temperature of 250 °C. Grain size estimations were made by adopting a single line method described by de Keijser [22], based on the least-squares fitting of broadened peaks to a pseudo-Voigt function. By increasing the incident angle to 4.0°, the intermediate TiN layer could be detected; its grain size was estimated as approximately 9 nm, with a strong (111) preferred texture. Average (Ti,Al)N grain sizes as a function of boron concentration are plotted in Fig. 6a. A crystallite size of approximately 4 nm was observed at the lowest B concentration. The mean grain size increased to
Fig. 6. (a) Average (Ti,Al)N grain size (as determined by GAXRD) and coating density (as determined by L-SAW), and (b) hardness H and elastic modulus E as a function of B content for nc-(Ti,Al)N/a-BN coatings.
6082
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
Fig. 7. Long-term stability of a nc-(Ti,Al)N/a-BN coating (6.2 at.% B), comparing hardness H and elastic modulus E values gathered from eight different nanoindentation systems.
above 5 nm at intermediate B concentrations, while at the higher concentrations (≥ 26.4 at.%) it decreased to ≤ 3 nm. The crystallite size observed for coatings with 16.0 at.% B (5–6 nm) is in good agreement to values (5–8 nm) reported for a coating of similar composition, containing 13.2 at.% B [21]. The grain sizes reported here are in the range expected for superhard nanocomposites [2,10], but significantly smaller than those measured previously for such coatings, when evaluating the average size of nc-(Ti,Al)N grains in dark-field TEM images [15]. As mentioned in Ref. [15], subgrain boundaries were identified and it is clear now from GAXRD data that the original TEM measurement values correspond to nano-agglomerates of smaller grains within individual coating columns. Coating density versus boron concentration is also displayed in Fig. 6a. The density decreased from 4.2 to 2.6 g/cm3 with increasing B concentration (i.e. higher amounts of a-BN), due to the lower density of BN compared to (Ti,Al)N. Similar density values (between 3.1 to 3.8 g/cm3, dependent on phases observed in XPS studies) have been recently reported for (mainly X-ray amorphous) TiAlBN coatings deposited by magnetron sputtering in Ar/N2 mixtures at 150 °C [23]. H and E as a function of boron content for the deposited nc(Ti,Al)N/a-BN coatings are plotted in Fig. 6b. Similar trends for H and E were observed, and maximum values of ∼ 40 and ∼ 340 GPa were found, respectively, for coatings containing 6.2 and 16 at.% B (and consisting, respectively, of ∼ 90 and ∼ 70 mol% (Ti,Al)N, [15,16]), with average (Ti,Al)N grain sizes of 5–6 nm (Fig. 6a). In good agreement with these results, highest hardness values for nc-TiN/a-Si3N4 [10] and nc-TiN/aC:H [5] coatings have been reported at 8–10 at.% silicon and ∼ 20 mol% a-C:H contents, respectively. The incorporation of 4.2 at.% B resulted in H values similar to TiN coatings, but with a much lower E of ∼ 300 GPa (compared to ∼ 500 GPa for TiN). At higher B concentrations (≥ 26.4 at.%) H and E decreased to ∼ 22 and ∼250 GPa, respectively, due to the higher amounts of the soft a-BN phase present in the coatings. High H / E values (– a good indicator of coating ‘resilience’; a long elastic strain-to-failure [23]) of 0.09–0.12 were calculated from the H and E values. H / E has been suggested as a reliable predictor of good wear-resistance [24] and of coating fracture toughness [25]. This appears to be confirmed in practice, where
twin EB TiAlBN coatings with high H / E, and with a phase fraction of ∼ 90 mol% (Ti,Al)N and ∼ 10 mol% BN, exhibit excellent wear resistance in wet-cutting drill tests, demonstrating a 150% increase in lifetime compared to standard TiAlN [14,15]. The long-term stability of a nc-(Ti,Al)N/a-BN coating with 6.2 at.% B on an M2 tool steel substrate was analysed at room temperature in air using eight nanoindentation systems at different European research institutes; the results are summarised in Fig. 7. The coating retained its mechanical properties for more than 90 months, comparing results gathered from different nanoindentation systems. Higher values were observed using Fischerscope H100 instruments, due mainly to the different evaluation method (i.e. a linear extrapolation of the unloading curve, rather than the power law fitting procedure [19] used in other systems). Excluding these data, averaged H and E values of 40 ± 1 and 360 ± 30 GPa, respectively, were measured — with the expected slight variations seen between the different systems. Using a Fischerscope instrument for mechanical property evaluation does not necessarily mean that the displayed values are always too high. Veprek et al. used a Fischerscope 100 system for H evaluation of their nc-TiN/Si3N4 [4] and nc-TiN/a-BN [7,8] coatings, where they took the upper 30% of the unloading curve for the linear fitting, since H values were found for this percentage which were in agreement with those calculated on the basis of measurement of the projected area of the remaining plastic deformation, by means of a calibrated SEM [26]. In a critical examination comparing nanoindentation results from nc-TiN/a-Si3N4 and nc-TiN/a-BN coatings with previously reported data (using a Fischerscope 100 and the above described method) Fischer-Cripps [27] concluded that H values in the 35–45 GPa range for these coatings are in general reliable, despite some concerns regarding the test methodology itself. The mechanical properties for two nc-(Ti,Al)N/a-BN coatings with 6.2 and 26.4 at.% B (and approximately 90 and 50 mol% (Ti,Al)N, respectively) are displayed as a function of annealing temperature, Ta in Fig. 8. A maximum Ta of 900 °C was used, to avoid significant interdiffusion between the AISI 316 substrate and the coatings. The hardness of the low-B content coating increases from 39.5 ± 3.4 GPa in the as-deposited state to a maximum of 45.1 ± 3.3 GPa at Ta = 700 °C, decreasing slightly
Fig. 8. Hardness H and elastic modulus E for two nc-(Ti,Al)N/a-BN coatings as a function of annealing temperature Ta.
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
with higher annealing temperature to 42.5 ± 1.9 GPa, at Ta = 900 °C. For the coating with high-B content, H also increased slightly, from 23.0 ± 1.0 GPa in the as-deposited state to a maximum of 24.6 ± 0.9 GPa at Ta = 800 °C, decreasing again to 22.1 ± 0.9 GPa at Ta = 900 °C. So-called ‘self-hardening’, i.e. an increase in H observed during annealing at the abovementioned temperatures, tends to occur when phase segregation is incomplete during deposition at the chosen conditions but is allowed to proceed to completion during annealing at a somewhat higher temperature. This behaviour is in contrast to stoichiometric binary nitride and carbide layers, in which the hardness typically decreases during annealing, due to relief of compressive stress [3]. The nanoindentation modulus E increases for both coatings from ∼305 and ∼ 253 GPa in the as-deposited state to ∼ 340 and ∼ 260 GPa for coatings 6.2 and 26.4 at.% B, respectively. A similar behaviour during annealing experiments was observed for nanocrystalline Ti–B–N films, where the increase in E was attributed to a reduction in volume fraction of the disordered phase [9,28]. 4. Conclusions In the present study it was demonstrated that hard and superhard coatings within the quaternary Ti–Al–B–N system, consisting of an essentially dual-phase (pseudo-binary) nc-(Ti, Al)N/a-BN crystalline-cubic/amorphous nanocomposite structure, can be deposited by co-evaporation of Ti and HIPped Ti– Al–B–N material in thermionically-enhanced twin-EB PAPVD system, employing a combination of OES and partial pressure control to adjust the coating composition. Measured hardness (H) and elastic modulus (E) values depend on the average (Ti, Al)N grain size and maximum values of ∼ 40 and ∼ 340 GPa, respectively, were observed for coatings with 5–6 nm grain size and nc-(Ti,Al)N phase fractions of both ∼ 70 and ∼ 90 mol%. It was also established that nc-(Ti,Al)N/a-BN coatings exhibit long-term structural stability at room temperature in air and thermal stability to temperatures in excess of 900 °C, demonstrating a moderate self-hardening effect and an increase in H from b40 to N45 GPa for some coatings. Comparing mechanical property data of superhard nc-(Ti,Al) N/a-BN coatings, obtained from eight different nanoindentation systems, has shown that H and E values obtained are very similar for systems where a power law fit of the unloading curve is used for data evaluation, while ∼50% higher H values are obtained for measurements made with Fischerscope H100 systems; modifications to the standard evaluation method (i.e. linear extrapolation of the unloading curve) are necessary in this case to avoid ‘artificial’ superhard values. Acknowledgements The Incoming Short Visit Program Project (Evaluation of the properties of novel titanium-based PVD nanostructured films)
6083
of the Royal Society is gratefully acknowledged by C. Rebholz and A. Leyland for funding support. The authors at Sheffield University express their gratitude to the UK Engineering and Physical Sciences Research Council (EPSRC), and to the European Commission, for funding support under EPSRC Platform Grant GR/S94469/01 and FP6 Coordination Action Project CA-505549 (DESHNAF), respectively. P.H. Mayrhofer acknowledges the financial support by the Erwin Schrödinger Program (project J2469-N02) of the Austrian Science Fund (FWF). References [1] S. Veprek, J. Vac. Sci. Technol., A 17 (5) (1999) 2401. [2] A.A. Voevodin, J.S. Zabinski, C. Muratore, Tsinghua Sci. Technol. 10 (6) (2005) 665. [3] L. Hultman, Vacuum 57 (2000) 1. [4] H.-D. Männling, D.S. Patil, K. Moto, M. Jilek, S. Veprek, Surf. Coat. Technol. 146–147 (2001) 263. [5] J. Patschneider, T. Zehnder, M. Diserens, Surf. Coat. Technol. 146–147 (2001) 201. [6] A. Fink, T. Larson, J. Sjölén, L. Karlsson, L. Hultman, Surf. Coat. Technol. 200 (2005) 1535. [7] P. Karvankova, M.G.J. Veprek-Heijman, O. Zindulka, A. Bergmaier, S. Veprek, Surf. Coat. Technol. 163–164 (2003) 149. [8] P. Karvankova, M.G.J. Veprek-Heijman, D. Azinovic, S. Veprek, Surf. Coat. Technol. 200 (2006) 2978. [9] P.H. Mayrhofer, C. Mitterer, J.G. Wen, I. Petrov, J.E. Greene, J. Appl. Phys. 100 (2006) 044301. [10] S. Veprek, M.G.J. Veprek-Heijman, P. Karvankova, J. Prochazka, Thin Solid Films 476 (2005) 1. [11] A.C. Fischer-Cripps, Vacuum 58 (2000) 569. [12] A.C. Fischer-Cripps, Surf. Coat. Technol. 200 (2006) 4153. [13] C. Rebholz, H. Ziegele, A. Leyland, A. Matthews, J. Vac. Sci. Technol., A 16 (5) (1998) 2851. [14] C. Rebholz, A. Leyland, A. Matthews, Thin Solid Films 343–344 (1999) 242. [15] M.A. Baker, S. Klose, C. Rebholz, A. Leyland, A. Matthews, Surf. Coat. Technol. 151–152 (2002) 338. [16] M.A. Baker, C. Rebholz, A. Leyland, A. Matthews, Vacuum 67 (2002) 471. [17] U. Goetz, Innovation in Ceramics, Sintec-Group, Halblech, Germany, 2005 www.sintec-keramik.com/sintec-en/evaporator-boats_properties.html. [18] D. Schneider, T. Schwarz, Surf. Coat. Technol. 91 (1997) 136. [19] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [20] H. Nowotny, F. Benesovsky, C. Brukl, O. Schob, Monatsch. Chem. 92 (1961) 403. [21] D.V. Shtansky, K. Kaneko, Y. Ikuhara, E.A. Levashov, Surf. Coat. Technol. 148 (2001) 206. [22] T.H. de Keijser, E.J. Mittemeijer, H.C. Rozendaal, J. Appl. Crystallogr. 16 (1983) 309. [23] C. Rebholz, A. Leyland, A. Matthews, C. Charitidis, S. Logothetidis, D. Schneider, Thin Solid Films 514 (2006) 81. [24] A. Leyland, A. Matthews, Wear 246 (2000) 1. [25] G.M. Pharr, Mater. Sci. Eng., A 253 (1998) 151. [26] S. Veprek, S. Mukherjee, P. Karvankova, H.-D. Männling, J.L. He, J. Xu, J. Prochazka, A.S. Argon, A.S. Li, Q.F. Fang, S.Z. Li, H. Manghnani, S. Tkachev, P. Zinin, Mater. Res. Soc. Symp. Proc. 750 (2003) 7. [27] A.C. Fischer-Cripps, P. Karvakova, S. Veprek, Surf. Coat. Technol. 200 (2006) 5645. [28] P.H. Mayrhofer, M. Stoiber, C. Mitterer, Scr. Mater. 53 (2005) 241.
Hard and superhard TiAlBN coatings deposited by twin electron-beam evaporation C. Rebholz a,⁎, M.A. Monclus b , M.A. Baker b , P.H. Mayrhofer c , P.N. Gibson d , A. Leyland e , A. Matthews e a
Department of Mechanical and Manufacturing Engineering, University of Cyprus, 1678 Nicosia, Cyprus The Surface Analysis Laboratory, School of Engineering, University of Surrey, Guildford GU2 7XH, UK Department of Physical Metallurgy and Materials Testing, University of Leoben, Franz-Josef-Strasse 18, A-8700 Leoben, Austria d Institute for Health and Consumer Protection, Joint Research Centre, I-21020 Ispra (VA), Italy e Department of Engineering Materials, University of Sheffield, Sheffield S1 3JD, UK b
c
Available online 2 October 2006
Abstract Superhard nanostructured coatings, prepared by plasma-assisted chemical vapour deposition (PACVD) and physical vapour deposition (PAPVD) techniques, such as vacuum arc evaporation and magnetron sputtering, are receiving increasing attention due to their potential applications for wear protection. In this study nanocomposite (TiAl)BxNy (0.09 ≤ x ≤ 1.35; 1.07 ≤ y ≤ 2.30) coatings, consisting of nanocrystalline (Ti,Al)N and amorphous BN, were deposited onto Si (100), AISI 316 stainless steel and AISI M2 tool steel substrates by co-evaporation of Ti and hot isostatically pressed (HIPped) Ti–Al–B–N material from a thermionically enhanced twin crucible electron-beam (EB) evaporation source in an Ar plasma at 450 °C. The coating stoichiometry, relative phase composition, nanostructure and mechanical properties were determined using X-ray photoelectron spectroscopy (XPS) and X-ray diffraction (XRD), in combination with nanoindentation measurements. Aluminium (∼10 at.% in coatings) was found to substitute for titanium in the cubic TiN based structure. (Ti,Al)B0.14N1.12 and (Ti,Al)B0.45N1.37 coatings with average (Ti,Al)N grain sizes of 5–6 nm and either ∼70, or ∼90, mol% (Ti,Al)N showed hardness and elastic modulus values of ∼40 and ∼340 GPa, respectively. (Ti,Al)B0.14N1.12 coatings retained their ‘as-deposited’ mechanical properties for more than 90 months at room temperature in air, comparing results gathered from eight different nanoindentation systems. During vacuum annealing, all coatings examined exhibited structural stability to temperatures in excess of 900 °C, and revealed a moderate, but significant, increase in hardness. For (Ti,Al)B0.14N1.12 coatings the hardness increased from ∼40 to ∼45 GPa. © 2006 Elsevier B.V. All rights reserved. Keywords: Superhard nanocomposites; Nanoindentation; Thermal stability; Electron-beam evaporation; Physical vapour deposition (PVD)
1. Introduction In the last few years, advances in coating deposition technologies have led to the development of nanostructured coating materials with unique properties, such as superhardness (defined as hardness H N 40 GPa) in combination with high toughness [1,2]. Coatings prepared by plasma-assisted chemical vapour deposition (PACVD) and physical vapour deposition (PAPVD) are often known to exhibit increased hardness compared to their corresponding bulk counterparts, due to factors such as grain size refinement, high growth-defect density and
⁎ Corresponding author. Tel.: +352 22 892282; fax: +357 22 892254. E-mail address: [email protected] (C. Rebholz). 0257-8972/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2006.08.121
residual compressive stress, induced by the deposition process [3]. In binary coatings with large single-phase fields these hardness enhancement effects can diminish rapidly during exposure to temperatures above the deposition temperature (e.g. in severe machining conditions), as recovery and recrystallisation processes occur [3]. The growing number of nanocomposite coating systems that exhibit superhardness together with high thermal stability include coatings within the ternary Ti–Si–N [4–6] and Ti–B–N [7–9] material systems, deposited by PACVD, and by PAPVD techniques such as magnetron sputtering or vacuum arc evaporation. A detailed recent review of the different approaches to the design and synthesis of superhard nanostructured coatings can be found in ref. [10]. A continuing issue, that superhard coating deposition particularly has brought to prominence, is how to measure
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
correctly the mechanical properties of these and of ultrahard (H N 70 GPa) coatings using sub-micron indentation testing, and Fischer–Cripps has highlighted the limitations and common sources of error in two recent reviews [11,12]. In contrast to superhard coatings in the Ti–B–N ternary system (i.e. TiN/TiB2; TiN/BN), only a few studies have been reported on quaternary TiAlBN coatings — especially those deposited by electron-beam PVD (EBPVD) [13–16]. A maximum hardness value of 37 GPa and superior wet-cutting performance (compared to standard TiN and TiAlN) were observed for a dual-phase coating effectively consisting of nanocrystalline (nc-) (Ti,Al)N and amorphous (a-) BN, in approximate proportions 90 and 10 mol%, respectively [15]. No data on thermal and long-term mechanical/structural stability for such quaternary coatings are available in the literature; however, it is known that coatings within the Ti–B– N ternary system possess high-temperature thermal stability to above 900 °C [8–10] and long-term structural stability at room temperature in air [10]. The main objective of the present work was therefore to determine the dependence of the structure, density, mechanical properties and thermal stability on composition for twincrucible EBPVD nc-(Ti,Al)N/a-BN coatings with Al and N contents of approximately 10 and 50 at.%, respectively, and varying B and Ti contents. Furthermore, the present work was also directed towards a desire to compare the mechanical property data of superhard nc-(Ti,Al)N/a-BN coatings (obtained from different nanoindentation systems) during a long-term stability study, thereby gaining a better understanding of the limitations of the various measuring systems used. 2. Experimental details TiAlBN coatings, 2.3 ± 0.1 μm thick, were deposited onto Si (100), AISI 316 stainless steel and AISI M2 tool steel substrates by co-evaporating Ti and HIPped Ti–Al–B–N material from a thermionically enhanced twin-crucible EB evaporation source system (TECVAC IP35L) with a base pressure of b 5 × 10− 4 Pa. A combination of optical emission spectroscopy (OES) and partial pressure control was utilised to control the evaporation rates and hence the composition in the coating. Fig. 1 shows a schematic of the deposition system, illustrating crucible location and evaporant materials employed. Commercially available hot-pressed Ti–Al–B–N evaporation boats [17] with dimensions of 10 mm × 15 mm × 120 mm (Fig. 2a), consisting of a mixture of 50 wt.% TiB2, 30 wt.% BN and 20 wt.% AlN, were used to create the Ti–Al–B–N evaporant material; a scanning electron microscopy (SEM) image of the as-received morphology can be seen in Fig. 2b. These boats were hot isostatically pressed (HIPped) to densify the material (Fig. 2c) and reduce spitting during EB evaporation [13], and then broken up into smaller granules (of approximately 7 mm mean diameter) to place in the crucible. The crucible is shown in operation in Fig. 2d, including the filaments used for plasma enhancement; the source material evaporates by sublimation in the middle of the crucible, where the EB gun scans over the
6079
Fig. 1. Schematic drawing of the TECVAC IP35L twin EB evaporation system used for nc-(Ti,Al)N/a-BN coating deposition.
material. Prior to deposition, the AISI 316 stainless steel substrates were ultrasonically cleaned in acetone and isopropanol and then fixed in a single-rotation holder mounted 300 mm above the vapour source. The deposition process consisted of six steps: (i) radiation heating to 450 °C (ii) diode sputter cleaning and plasma heating in an Ar atmosphere at −1000 V (iii) deposition of a thin 0.2 μm Ti interlayer (iv) deposition of a 0.5 μm TiN layer (N2 as reactive gas) and monitoring of the OES signal to establish a stable (and reproducible) Ti vapour–flux proportion; (v) reduction of the Ti evaporation rate to a selected percentage of the initial Ti OES signal and removal of N2 flow, followed by (vi) twin-EB evaporation of Ti and Ti–Al–B–N. Stable evaporation of the Ti (at a selected percentage of the initial rate) and of the Ti– Al–B–N material were achieved using combined OES and partial pressure control. The power on both crucibles was adjusted during deposition to keep the Ti evaporation rate constant and the total pressure constant at 0.35 Pa. Substrate temperature and bias voltage were kept constant at 450 °C and −100 V, respectively. Both chemical and phase composition of the various coatings were determined by X-ray photoelectron spectroscopy (XPS), using a VG-Scientific Sigma Probe spectrometer employing a monochromated Al–Kα source and a hemispherical analyser [15,16]. The crystallographic structure and texture of the films were analysed by glancing-angle X-ray diffraction (GAXRD), using CuKα radiation at incident angles ranging from 0.5 to 4.0°. The X-ray generator settings were 35 kV and 30 mA, the step angle being 0.2°. The laserinduced surface acoustic wave (L-SAW) technique [18] was used to obtain coating density values from coated M2 and Si (100) samples. Fracture cross-sections of coated samples were prepared for SEM morphology and topography studies. Film thickness was measured using a Veeco DEKTAK 3ST
6080
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
Fig. 2. (a) Photograph of hot-pressed Ti–Al–B–N evaporation boats, SEM images of (b) hot-pressed and (c) HIPped Ti–Al–B–N boats, and (d) photograph of a crucible with broken up HIPped Ti–Al–B–N evaporation material.
profilometer with a vertical resolution of 50 nm and verified by SEM cross-sectional measurements. As-deposited coatings on AISI 316 substrates were annealed in vacuum for 30 min at temperatures Ta of 600, 700, 800 and 900 °C and a pressure of ≤10− 3 Pa. The hardness H and elastic modulus E values were obtained from nanoindentation measurements at room temperature using a Hysitron Triboscope instrument attached to an atomic force microscope. The area function of the triangular Berkovich diamond tip was calibrated following the procedure described in Ref. [19]. The maximum load was varied from 3 to 6 mN with a sequence of 30 indents at each load; the maximum indentation depth was always less than 10% of the top TiAlBN layer thickness. For long-term stability studies at room temperature in air, a superhard nc-(Ti,Al)N/a-BN coated M2 substrate was analysed over a period of 92 months using different nanoindentation systems at various European research institutes. The systems used, in chronological order, were: (i) Nanoindenter XP, Department of Physics, Aristotle University of Thessaloniki, Greece; (ii) Nanoindenter II, Department of Physics, University of Linköping, Sweden; (iii) Hysitron Triboscope, Department of Engineering Materials, University of Sheffield, UK; (iv) Fischerscope H100, Robert Bosch GmbH, Stuttgart, Germany; (v) Fischerscope H100, Helmut Fischer GmbH & Co.KG, SindelfingenMaichingen, Germany; (vi) CSM Nanohardnesstester (NHT), CSM Instruments SA, Peseux, Switzerland; (vii) Nanoindenter XP, Robert Bosch GmbH, Stuttgart, Germany and (viii) Hysitron Triboscope, Materials Chemistry, RWTH Aachen, Germany. With the exception of the Fischerscope H100 instruments, all systems were equipped with a Berkovich diamond indenter; the maximum indentation depth was 150 nm, i.e. ≤10% thickness of top TiAlBN layer, and the load frame compliance and tip correction function were estimated using sapphire and/or fused silica control samples [19]. Vickers diamond indenters were used for the Fischerscope H100 instruments, with linear extrapolation
of the unloading curve to determine the contact stiffness and the corrected contact depth; a load of 30 mN was used in each case. 3. Results and discussion Coating compositions, determined from XPS Ti 2p, Al 2p, B 1s and N 1s peak areas [15,16], are given in Table 1, and are plotted on the Ti–B–N phase diagram of Nowotny et al. [20] modified for the inclusion of Al into the Ti–B–N system (Fig. 3). All coatings possess approximately constant Al (∼ 10 at.%) and N (∼ 50 at.%) content, but varying B and Ti concentration, and are located close to the (Ti,Al)N–BN tie line. In accordance with their position in the phase diagram, all coatings exhibit a three-phase composition of (Ti,Al)N and BN with very small amounts of TiB2 (≤ 3 mol%), i.e. the microstructure is effectively that of a (Ti,Al)N and BN dualphase coating [15,16]. With increasing B concentration, a colour change from dark gold to metallic black was observed. Fig. 4a shows a typical SEM fracture cross-section representative of twin EB nc-(Ti,Al)N/a-BN coatings produced at 450 °C. The individual Ti, TiN and TiAlBN layers are clearly visible. In contrast to the columnar morphology of the Ti layer, the TiN and TiAlBN layers exhibit a dense, featureless ‘glassy’ Table 1 Elemental concentrations and stoichiometry for the TiAlBN coatings determined by XPS [15,16] Elemental concentration [at.%]
Stoichiometry
Ti
Al
B
N
39.1 36.2 25.3 16.8 10.1
7.2 8.0 10.1 9.0 11.4
4.2 6.2 16.0 26.4 29.1
49.5 49.5 48.6 47.8 49.4
Sum of O, C and Ar b 3 at.% in all coatings.
(Ti,Al)B0.09N1.07 (Ti,Al)B0.14N1.12 (Ti,Al)B0.45N1.37 (Ti,Al)B1.02N1.85 (Ti,Al)B1.35N2.30
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
Fig. 3. Chemical composition (as determined by XPS) of deposited nc-(Ti,Al)N/ a-BN coatings within a modified Ti–B–N equilibrium phase diagram [20].
structure. Differences in surface topography for coatings with low B concentrations deposited from hot-pressed (Fig. 2b) and HIPped (Fig. 2c) Ti–Al–B–N source material can be seen in Fig. 4b and c, respectively. The hot-pressed material yielded a significantly higher surface roughness, while using the denser HIPped material resulted in virtually ‘spit-free’ coatings and improved coating integrity. GAXRD patterns at 0.5° incident angle for the nc-(Ti,Al)N/aBN coatings are plotted in Fig. 5. At this angle of incidence only the surface structure (in the TiAIBN top layer) is detected. The only peaks obtained over the 2Θ range 10–90° were for (111), (200), (220) and (311) reflections of a face-centred cubic TiNlike phase, in which Al is assumed to substitute for Ti [14,15]. No evidence of other crystalline phases (TiB2 and/or BN) was found, in agreement with previously reported transmission electron microscopy (TEM) studies [15]. Shtansky et al. [21] also only observed the TiN phase in TEM studies for a TiAl0.3B0.5N1.9 (equivalent to (Ti,Al)B0.38N1.46) coating located close to the (Ti,Al)N–BN tie line and deposited by magnetron sputtering from a Ti–Al–B–N target (manufactured by selfpropagating high-temperature synthesis from a mixture of Ti, Al
Fig. 4. (a) SEM micrograph of fracture cross-section representative for nc-(Ti, Al)N/a-BN coatings, and surface topographies for coatings with low B contents deposited from (b) hot-pressed and (c) HIPped Ti–Al–B–N source material.
6081
Fig. 5. GAXRD patterns at 0.5° for nc-(Ti,Al)N/a-BN coatings.
and B, with 66.7, 20 and 13.3 wt.%, respectively) in Ar/N2 mixtures at a substrate temperature of 250 °C. Grain size estimations were made by adopting a single line method described by de Keijser [22], based on the least-squares fitting of broadened peaks to a pseudo-Voigt function. By increasing the incident angle to 4.0°, the intermediate TiN layer could be detected; its grain size was estimated as approximately 9 nm, with a strong (111) preferred texture. Average (Ti,Al)N grain sizes as a function of boron concentration are plotted in Fig. 6a. A crystallite size of approximately 4 nm was observed at the lowest B concentration. The mean grain size increased to
Fig. 6. (a) Average (Ti,Al)N grain size (as determined by GAXRD) and coating density (as determined by L-SAW), and (b) hardness H and elastic modulus E as a function of B content for nc-(Ti,Al)N/a-BN coatings.
6082
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
Fig. 7. Long-term stability of a nc-(Ti,Al)N/a-BN coating (6.2 at.% B), comparing hardness H and elastic modulus E values gathered from eight different nanoindentation systems.
above 5 nm at intermediate B concentrations, while at the higher concentrations (≥ 26.4 at.%) it decreased to ≤ 3 nm. The crystallite size observed for coatings with 16.0 at.% B (5–6 nm) is in good agreement to values (5–8 nm) reported for a coating of similar composition, containing 13.2 at.% B [21]. The grain sizes reported here are in the range expected for superhard nanocomposites [2,10], but significantly smaller than those measured previously for such coatings, when evaluating the average size of nc-(Ti,Al)N grains in dark-field TEM images [15]. As mentioned in Ref. [15], subgrain boundaries were identified and it is clear now from GAXRD data that the original TEM measurement values correspond to nano-agglomerates of smaller grains within individual coating columns. Coating density versus boron concentration is also displayed in Fig. 6a. The density decreased from 4.2 to 2.6 g/cm3 with increasing B concentration (i.e. higher amounts of a-BN), due to the lower density of BN compared to (Ti,Al)N. Similar density values (between 3.1 to 3.8 g/cm3, dependent on phases observed in XPS studies) have been recently reported for (mainly X-ray amorphous) TiAlBN coatings deposited by magnetron sputtering in Ar/N2 mixtures at 150 °C [23]. H and E as a function of boron content for the deposited nc(Ti,Al)N/a-BN coatings are plotted in Fig. 6b. Similar trends for H and E were observed, and maximum values of ∼ 40 and ∼ 340 GPa were found, respectively, for coatings containing 6.2 and 16 at.% B (and consisting, respectively, of ∼ 90 and ∼ 70 mol% (Ti,Al)N, [15,16]), with average (Ti,Al)N grain sizes of 5–6 nm (Fig. 6a). In good agreement with these results, highest hardness values for nc-TiN/a-Si3N4 [10] and nc-TiN/aC:H [5] coatings have been reported at 8–10 at.% silicon and ∼ 20 mol% a-C:H contents, respectively. The incorporation of 4.2 at.% B resulted in H values similar to TiN coatings, but with a much lower E of ∼ 300 GPa (compared to ∼ 500 GPa for TiN). At higher B concentrations (≥ 26.4 at.%) H and E decreased to ∼ 22 and ∼250 GPa, respectively, due to the higher amounts of the soft a-BN phase present in the coatings. High H / E values (– a good indicator of coating ‘resilience’; a long elastic strain-to-failure [23]) of 0.09–0.12 were calculated from the H and E values. H / E has been suggested as a reliable predictor of good wear-resistance [24] and of coating fracture toughness [25]. This appears to be confirmed in practice, where
twin EB TiAlBN coatings with high H / E, and with a phase fraction of ∼ 90 mol% (Ti,Al)N and ∼ 10 mol% BN, exhibit excellent wear resistance in wet-cutting drill tests, demonstrating a 150% increase in lifetime compared to standard TiAlN [14,15]. The long-term stability of a nc-(Ti,Al)N/a-BN coating with 6.2 at.% B on an M2 tool steel substrate was analysed at room temperature in air using eight nanoindentation systems at different European research institutes; the results are summarised in Fig. 7. The coating retained its mechanical properties for more than 90 months, comparing results gathered from different nanoindentation systems. Higher values were observed using Fischerscope H100 instruments, due mainly to the different evaluation method (i.e. a linear extrapolation of the unloading curve, rather than the power law fitting procedure [19] used in other systems). Excluding these data, averaged H and E values of 40 ± 1 and 360 ± 30 GPa, respectively, were measured — with the expected slight variations seen between the different systems. Using a Fischerscope instrument for mechanical property evaluation does not necessarily mean that the displayed values are always too high. Veprek et al. used a Fischerscope 100 system for H evaluation of their nc-TiN/Si3N4 [4] and nc-TiN/a-BN [7,8] coatings, where they took the upper 30% of the unloading curve for the linear fitting, since H values were found for this percentage which were in agreement with those calculated on the basis of measurement of the projected area of the remaining plastic deformation, by means of a calibrated SEM [26]. In a critical examination comparing nanoindentation results from nc-TiN/a-Si3N4 and nc-TiN/a-BN coatings with previously reported data (using a Fischerscope 100 and the above described method) Fischer-Cripps [27] concluded that H values in the 35–45 GPa range for these coatings are in general reliable, despite some concerns regarding the test methodology itself. The mechanical properties for two nc-(Ti,Al)N/a-BN coatings with 6.2 and 26.4 at.% B (and approximately 90 and 50 mol% (Ti,Al)N, respectively) are displayed as a function of annealing temperature, Ta in Fig. 8. A maximum Ta of 900 °C was used, to avoid significant interdiffusion between the AISI 316 substrate and the coatings. The hardness of the low-B content coating increases from 39.5 ± 3.4 GPa in the as-deposited state to a maximum of 45.1 ± 3.3 GPa at Ta = 700 °C, decreasing slightly
Fig. 8. Hardness H and elastic modulus E for two nc-(Ti,Al)N/a-BN coatings as a function of annealing temperature Ta.
C. Rebholz et al. / Surface & Coatings Technology 201 (2007) 6078–6083
with higher annealing temperature to 42.5 ± 1.9 GPa, at Ta = 900 °C. For the coating with high-B content, H also increased slightly, from 23.0 ± 1.0 GPa in the as-deposited state to a maximum of 24.6 ± 0.9 GPa at Ta = 800 °C, decreasing again to 22.1 ± 0.9 GPa at Ta = 900 °C. So-called ‘self-hardening’, i.e. an increase in H observed during annealing at the abovementioned temperatures, tends to occur when phase segregation is incomplete during deposition at the chosen conditions but is allowed to proceed to completion during annealing at a somewhat higher temperature. This behaviour is in contrast to stoichiometric binary nitride and carbide layers, in which the hardness typically decreases during annealing, due to relief of compressive stress [3]. The nanoindentation modulus E increases for both coatings from ∼305 and ∼ 253 GPa in the as-deposited state to ∼ 340 and ∼ 260 GPa for coatings 6.2 and 26.4 at.% B, respectively. A similar behaviour during annealing experiments was observed for nanocrystalline Ti–B–N films, where the increase in E was attributed to a reduction in volume fraction of the disordered phase [9,28]. 4. Conclusions In the present study it was demonstrated that hard and superhard coatings within the quaternary Ti–Al–B–N system, consisting of an essentially dual-phase (pseudo-binary) nc-(Ti, Al)N/a-BN crystalline-cubic/amorphous nanocomposite structure, can be deposited by co-evaporation of Ti and HIPped Ti– Al–B–N material in thermionically-enhanced twin-EB PAPVD system, employing a combination of OES and partial pressure control to adjust the coating composition. Measured hardness (H) and elastic modulus (E) values depend on the average (Ti, Al)N grain size and maximum values of ∼ 40 and ∼ 340 GPa, respectively, were observed for coatings with 5–6 nm grain size and nc-(Ti,Al)N phase fractions of both ∼ 70 and ∼ 90 mol%. It was also established that nc-(Ti,Al)N/a-BN coatings exhibit long-term structural stability at room temperature in air and thermal stability to temperatures in excess of 900 °C, demonstrating a moderate self-hardening effect and an increase in H from b40 to N45 GPa for some coatings. Comparing mechanical property data of superhard nc-(Ti,Al) N/a-BN coatings, obtained from eight different nanoindentation systems, has shown that H and E values obtained are very similar for systems where a power law fit of the unloading curve is used for data evaluation, while ∼50% higher H values are obtained for measurements made with Fischerscope H100 systems; modifications to the standard evaluation method (i.e. linear extrapolation of the unloading curve) are necessary in this case to avoid ‘artificial’ superhard values. Acknowledgements The Incoming Short Visit Program Project (Evaluation of the properties of novel titanium-based PVD nanostructured films)
6083
of the Royal Society is gratefully acknowledged by C. Rebholz and A. Leyland for funding support. The authors at Sheffield University express their gratitude to the UK Engineering and Physical Sciences Research Council (EPSRC), and to the European Commission, for funding support under EPSRC Platform Grant GR/S94469/01 and FP6 Coordination Action Project CA-505549 (DESHNAF), respectively. P.H. Mayrhofer acknowledges the financial support by the Erwin Schrödinger Program (project J2469-N02) of the Austrian Science Fund (FWF). References [1] S. Veprek, J. Vac. Sci. Technol., A 17 (5) (1999) 2401. [2] A.A. Voevodin, J.S. Zabinski, C. Muratore, Tsinghua Sci. Technol. 10 (6) (2005) 665. [3] L. Hultman, Vacuum 57 (2000) 1. [4] H.-D. Männling, D.S. Patil, K. Moto, M. Jilek, S. Veprek, Surf. Coat. Technol. 146–147 (2001) 263. [5] J. Patschneider, T. Zehnder, M. Diserens, Surf. Coat. Technol. 146–147 (2001) 201. [6] A. Fink, T. Larson, J. Sjölén, L. Karlsson, L. Hultman, Surf. Coat. Technol. 200 (2005) 1535. [7] P. Karvankova, M.G.J. Veprek-Heijman, O. Zindulka, A. Bergmaier, S. Veprek, Surf. Coat. Technol. 163–164 (2003) 149. [8] P. Karvankova, M.G.J. Veprek-Heijman, D. Azinovic, S. Veprek, Surf. Coat. Technol. 200 (2006) 2978. [9] P.H. Mayrhofer, C. Mitterer, J.G. Wen, I. Petrov, J.E. Greene, J. Appl. Phys. 100 (2006) 044301. [10] S. Veprek, M.G.J. Veprek-Heijman, P. Karvankova, J. Prochazka, Thin Solid Films 476 (2005) 1. [11] A.C. Fischer-Cripps, Vacuum 58 (2000) 569. [12] A.C. Fischer-Cripps, Surf. Coat. Technol. 200 (2006) 4153. [13] C. Rebholz, H. Ziegele, A. Leyland, A. Matthews, J. Vac. Sci. Technol., A 16 (5) (1998) 2851. [14] C. Rebholz, A. Leyland, A. Matthews, Thin Solid Films 343–344 (1999) 242. [15] M.A. Baker, S. Klose, C. Rebholz, A. Leyland, A. Matthews, Surf. Coat. Technol. 151–152 (2002) 338. [16] M.A. Baker, C. Rebholz, A. Leyland, A. Matthews, Vacuum 67 (2002) 471. [17] U. Goetz, Innovation in Ceramics, Sintec-Group, Halblech, Germany, 2005 www.sintec-keramik.com/sintec-en/evaporator-boats_properties.html. [18] D. Schneider, T. Schwarz, Surf. Coat. Technol. 91 (1997) 136. [19] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [20] H. Nowotny, F. Benesovsky, C. Brukl, O. Schob, Monatsch. Chem. 92 (1961) 403. [21] D.V. Shtansky, K. Kaneko, Y. Ikuhara, E.A. Levashov, Surf. Coat. Technol. 148 (2001) 206. [22] T.H. de Keijser, E.J. Mittemeijer, H.C. Rozendaal, J. Appl. Crystallogr. 16 (1983) 309. [23] C. Rebholz, A. Leyland, A. Matthews, C. Charitidis, S. Logothetidis, D. Schneider, Thin Solid Films 514 (2006) 81. [24] A. Leyland, A. Matthews, Wear 246 (2000) 1. [25] G.M. Pharr, Mater. Sci. Eng., A 253 (1998) 151. [26] S. Veprek, S. Mukherjee, P. Karvankova, H.-D. Männling, J.L. He, J. Xu, J. Prochazka, A.S. Argon, A.S. Li, Q.F. Fang, S.Z. Li, H. Manghnani, S. Tkachev, P. Zinin, Mater. Res. Soc. Symp. Proc. 750 (2003) 7. [27] A.C. Fischer-Cripps, P. Karvakova, S. Veprek, Surf. Coat. Technol. 200 (2006) 5645. [28] P.H. Mayrhofer, M. Stoiber, C. Mitterer, Scr. Mater. 53 (2005) 241.