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Microstructural characterisation of carbon doped CrAlTiN nanoscale multilayer coatings
Surface & Coatings Technology 205 (2011) 3251–3259
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s e v i e r. c o m / l o c a t e / s u r f c o a t
Microstructural characterisation of carbon doped CrAlTiN nanoscale multilayer coatings Xiaoying Li ⁎, Wenwen Wu, Hanshan Dong School of Metallurgy and Materials, The University of Birmingham, Edgbaston, Birmingham B15 2TT, UK
a r t i c l e
i n f o
Article history: Received 21 April 2010 Accepted in revised form 17 November 2010 Available online 25 November 2010 Keywords: Carbon doped CrAlTiN coating Nanoscale multilayer TEM characterisation
a b s t r a c t Carbon doped CrAlTiN coatings for reducing the coefficient of friction sliding against a WC-Co ball in air were developed using a closed field unbalanced magnetron sputtering system (CFUMS). The carbon content was controlled by increasing the current of the carbon target power from 0 A to 5 A. The surface morphology, cross-sectional microstructure, phase constituent, chemical bonding energy and mechanical properties of coatings were characterised by means of XRD, SEM, XTEM, XPS, and nanoindentation. The results showed that the hardness, Young's modulus and coefficient of friction (CoF) of the coatings sliding against a WC-Co ball in air decreased with increasing carbon content. Microstructure characterisation revealed that the carbon doped coatings consisted of four sublayers from the substrate to the surface and the main phase formed is fcc B1NaCl like (Cr, Al, Ti) N and (Cr, Al, Ti) (C, N) phase in sublayers III and IV respectively when the carbon content is below 3.96 at.%. When the carbon concentration exceeds this value, the excessive carbon will begin to form the amorphous carbon (or carbon-riched) phase which leads to a decrease in the coefficient of friction (CoF) of the coatings sliding against a WC-Co ball in air. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The application of hard coatings on cutting tools and mechanical components has become a common practice in the manufacturing industry in the last two decades [1–5]. Conventional binary hard coatings such as TiN and CrN are commonly used to enhance the performance of cutting tools [6]. Recently, ternary Cr–X–N coatings, where X is the alloying element such as Ti [7,8], Al [9], Si [10], B [11], C [12], and Ni [13], etc., have been actively investigated to improve the properties of CrN based coatings. For example, in comparison to CrN coatings, Cr–Al–N coatings have a higher hardness (25–32 GPa) and much better oxidation resistance (up to 900 °C) due to the formation of a stable oxidation barrier of an Al2O3 layer by migrating Al atoms to the surface region [14]. Most recently, a CrN-based Cr–Ti–Al–N coating system has been developed by Teer Coatings to improve the performance of twist drills and extrusion dies [15,16]. It has been revealed that quaternary Cr–Ti–Al–N coatings have been designed to achieve even higher hardness and wear resistance as compared to that of the ternary coatings [3,17,18]. However, the coefficient of friction of these so-called superhard coatings is very high (0.7～0.9) [19] at room temperature, and extra lubrication is needed in many applications, which raises environmental concerns. On the other hand, graphite-like carbon (GLC) coatings were reported to have a very low coefficient of friction against most engineering materials [20]. However, GLC
⁎ Corresponding author. Tel.: + 44 121 4145243; fax: + 44 121 4147373. E-mail address: [email protected] (X. Li). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.11.046
degraded at high-temperatures and was found to be non-functional beyond 400 °C in air [21], which eliminates their use for dry and highspeed cutting tools. A complex interactive effect often determines the point at which a new synthetic coating is expected to fulfil it manufacturing purpose. By adding solid carbon into quaternary Cr–Al–Ti–N coatings with a nanocomposite structure may possess a more desirable combination of properties in terms of high hardness and low friction. However, there are few studies on the design and synthesis of such new coatings and there is little work on the influence of the C concentrations on the structure, microstructure and mechanical properties of carbon doped CrAlTiN coatings. In the present study, new carbon doped CrAlTiN coatings were designed and deposited by magnetron sputtering ion plating with carbon content ranging from 0 to 24.34 at.%. Detailed TEM characterization has been conducted to investigate the effect of carbon content on the structure, microstructure, and mechanical properties of the CrAlTiCN coatings, thus advancing the scientific understanding of the role of carbon in these new coatings. 2. Experimental procedures 2.1. Substrate materials The substrate materials used in this study were hardened (HRC≧62) M2 high speed steel and a Si wafer. The composition of the M2 steel is given in Table 1. The M2 discs were ground using silicon carbide paper from 120 grit down to 2400 grit, polished with 6 μm and 1 μm diamond paste and finally polished with 0.25 μm colloidal silica. The samples
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Table 1 Nominal chemical composition of M2 high speed steel. Element
C
Si
Mn
Cr
Mo
V
W
Fe
Composition (wt.%)
0.78–0.88
0.20–0.40
0.20–0.40
3.75–4.50
4.50–5.50
1.60–2.20
5.50–6.75
Balance
were cleaned in soapy water and then acetone prior to and after surface polishing. 2.2. Coating equipment Cr–Al–Ti–C–N coatings with different carbon concentrations were prepared using the UDP 450 closed-field unbalanced magnetron sputtering ion plating (CFUBMSIP) system [15,16]. The schematic illustration of the system is shown in Fig. 1. It can be seen that the sputtering system contains four rectangular cathodes with Cr, Ti, graphite, and Al targets, which is operated in the unbalanced magnetron mode. The closed magnetic field coupling results in a high degree of ionization and a high bias current density. The equipment has a rotating cylindrical substrate holder with a speed of 5 rpm. After cleaning, the samples were inserted into the deposition chamber and the system was pumped down to a high vacuum condition with a background pressure less than 3.0×10− 3 Pa. Firstly argon and then a mixture of argon and nitrogen gases was introduced into the chamber during the process with a fixed Ar flow rate of 20 sccm, and the nitrogen content was controlled by a plasma optical emission monitor (OEM) with a feedback control of 60–70%. 2.3. Deposition processes A−70 V DC bias was applied on the substrate during deposition. The process started with plasma surface cleaning with argon ions. Deposition is started with a Cr adhesion layer, and then a Cr–N transition layer by gradually introducing nitrogen gas. Following
these two sublayers, a compositionally graded Cr–Al–Ti–N sublayer was deposited. For samples 1#–5#, an extra carbon doped Cr–Al–Ti–N top layer was deposited. The relative concentration of Cr, Ti, Al, and C in the coatings was adjusted by the sputtering power applied to the targets during deposition. The range of carbon content was controlled by changing carbon target current from 0 A to 5 A. Detailed deposition process conditions for the carbon doped CrAlTiN coatings are summarized in Table 2. 2.4. Chemical composition and bonding Surface morphologies and cross-sectional microstructure of the deposited samples were examined using a field emission JEOL7000 scanning electron microscope (SEM) with EDX and WDX capabilities. The composition analysis was carried out both qualitatively to determine the existence of certain elements and quantitatively to determine the amount of each element present. The chemical bonding of the carbon doped CrAlTiN coatings was studied by X-ray Photoelectron Spectroscopy (XPS), using a AXISULTRAX unit. The spectrometer was equipped with a hemispherical energy analyzer. An X-ray tube with monochromatic Al Kα radiation (hv = 1486.6 ev) was used as a signal excitation source. The investigation was carried out under vacuum of the order of 10− 8 Pa. Samples were subjected to etching by argon ions with Ar + ion bombardment of energy 4 keV and ion current of 1.0 μA to remove the top contamination layer, which can reduce the occurrence of preferential sputtering in the etching process. 2.5. Microstructure characterisation
Fig. 1. Cross-section schematic drawing of arrangement plan of magnetron targets and substrate holder in the CFUBMSIP system.
The phase constituents of the coated surfaces were analysed in normal XRD, X'Pert Philips diffractometer, and glancing angle XRD (GXRD) D8 diffractometer (Bruker AXS) with Cu-Kα radiation (λ = 0.154 nm). The diffraction patterns obtained were analysed and indexed using X'Pert High Score analytical software with a PCPDFWIN data base [22]. Crystallite size was calculated by the Rietveld method and the line broadening (the width of the XRD lines) as obtained from the line profile fitting was used to estimate the crystallite size [23]. In order to investigate the effect of carbon content on microstructure, representative CrAlTiCN-1#, 3#, and 5# samples were selected for the TEM studies. Cross-sectional TEM samples were prepared as follows [24]: (1) Two slabs were glued using G-1 epoxy with the treated surfaces facing each other; (2) The assembly was cut to a thickness of less than 1 mm with the interface of the two layers glued together in the centre followed by fine grinding and polishing to 50 μm thickness; (3) This thin assembly was then glued to a 3 mm (o.d) and 0.5 mm (i.d) brass reinforcing disc. The exposed side was further polished to a thickness of approximately 20 μm; (4) The sample was then placed in an ion beam miller (Gatan 691 Precision Ion Polisher System, PIPS™), making sure
Table 2 Deposition conditions for the synthesis of carbon doped CrAlTiN coatings. Sample code
CrAlTiCN-0#
Target current, A Cr/Al/Ti
4.5/3.5/3.5 A
C target current 0A Substrate bias Base pressure Substrate rotation speed / Temperature, °C
CrAlTiCN-1#
CrAlTiCN-2#
CrAlTiCN-3#
CrAlTiCN-4#
CrAlTiCN-5#
4.5/3.5/3.5 A
4.5/3.5/3.5 A
4.5/3.5/3.5 A
4.5/3.5/3.5 A
4.5/3.5/3.5 A
1A
2A
3A − 70 V b 3.0 × 10− 3 Pa 5 rpm / ~ 200°C
4A
5A
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that the milling was done on the interface of the glued surface; (5) Milling was terminated once a desired thin area around the interface of the glued surfaces is created. FEI Tecnai F20-FEG (EDX) was used for the TEM study on the microstructure and phase constitutions.
Table 3 Elemental concentration of carbon doped CrAlTiN coatings by EDX and WDX. Sample
Element (at%) WDX N
2.6. Mechanical properties Nanoindentation hardness (H) and Young's modulus (E) were measured using a Nano-Test 600 machine (Micromaterials, UK). A Berkovick diamond tip (3-faced pyramid) with a radius of about 200 nm at its apex was used in the tests. Based on the thickness of the coatings, E and H were measured continuously up to an indentation depth of 1200 nm for each sample, and then calculated using the Oliver and Pharr method [25]. Under non-lubrication conditions, wear tests were carried out by a pin-on-disc tribometer, using a stationary 5 mm WC-Co counter ball (Grade 10, Spheric-Trafalgar Ltd) at a speed of 0.025 m/s (420 rpm) for 1 h. The normal contact load acting on the ball was 20 N and a wear track of 11 mm in diameter was produced. A stylus profilometer, Surf Corder SE 1700, was used to measure the surface roughness of the tested samples.
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CrAlTiN-0# CrAlTiCN-1# CrAlTiCN-2# CrAlTiCN-3# CrAlTiCN-4# CrAlTiCN-5#
45.98 44.12 47.02 44.79 41.87 33.16
EDX C
Al
Ti
Cr
N
C
2.34 3.96 9.65 12.53 24.34
6.87 6.96 6.58 5.91 6.09 5.78
6.07 6.45 5.09 5.87 5.33 5.35
40.53 40.13 37.35 33.78 34.18 31.37
46.53 43.54 46.29 45.69 42.14 35.18
2.92 4.69 8.75 12.26 22.32
3. Results 3.1. General characterization SEM observation found no cracks in the interfaces between the coating layer and the substrate for all the coatings deposited by the CFUBMSIP method, which indicated a good adherence of the coatings to the metal substrates. The surface morphologies and cross-sectional microstructures for all designed coatings revealed similar features under SEM and the typical SEM images are presented in Fig. 2. The coating surfaces showed columnar boundaries and the average columnar diameter was scattered between 100 to 300 nm (Fig. 2a). All the coatings are very dense and there is no sign of cracking on the surface or within the coatings, which indicates a good compactness of the coatings. All coatings have a similar thickness of about 5 μm. However, compared with 0#, the carbon doped coatings show two sublayers with different contrasts (Fig. 2b), which were caused by the change in the chemical composition of the coatings (see Table 2). The light elements C and N and metal elements Cr, Al, Ti in the coatings were analysed by WDX and EDX respectively, and the results are shown in Table 3. It can be seen that the carbon contents in the coatings increased with the carbon target current. When increasing the carbon target current from 0 to 5 A, the carbon content of the CrAlTiCN coatings was increased to 24.34 at.%. The metallic content
Fig. 3. Nano-hardness and Young's modulus variation with CrAlTiN and carbon doped CrAlTiN coatings.
ratios, on the other hand, did not change significantly in the doped coatings, and the nitrogen content in all deposited samples is around 40～50 at.%. It can be seen from Fig. 3 that the hardness and elastic modulus for all coatings show an approximately linear trend with carbon content, both of which decreased with increasing the carbon target current. As shown in Table 4, all the carbon doped CrAlTiN coatings possess very similar roughness values. An obvious increase of Ra was observed with increasing the carbon content. However, the coefficient of friction (CoF) value of the coatings against a WC-Co ball in air was found to decrease with increasing the carbon contents and the CoF for the coating 5# is about 0.350, which is about 47% lower than that of the carbon free coating, 0# sample.
Fig. 2. Surface (a) and cross-sectional (b) SEI SEM micrographs of the CrAlTiCN-5# sample.
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Table 4 Surface roughness and friction coefficient of carbon doped CrAlTiN coatings. Sample
0#
1#
2#
3#
4#
5#
Roughness (μm) CoF
0.0303 ± 0.0013 0.655 ± 0.140
0.0315 ± 0.0005 0.420 ± 0.132
0.0378 ± 0.0018 0.410 ± 0.090
0.0372 ± 0.0028 0.375 ± 0.083
0.0405 ± 0.0035 0.400 ± 0.080
0.0388 ± 0.0022 0.350 ± 0.065
3.2. XRD phase identification Fig. 4a shows the XRD patterns obtained from the carbon-free and carbon-doped coatings. Due to the substrate contribution, it is difficult to distinguish the peaks accurately. A tentative indexing shown in Fig. 4a displayed fcc (111), (200), (220), and (311) reflections of all the coatings, and the (110), (200), (211), and (220) reflections of the bcc α-Fe substrate. The intensity of the peaks is relatively weak except the peak around 2θ = 45°, which was superimposed with α-Fe (110) from the substrate. In order to avoid substrate interference effects, glancing angle XRD was employed and the typical GAXRD charts are given in Fig. 4b. As shown in Fig. 4b, a set of fcc (111), (200), (220), (311), (222), and (400) reflections obtained from the coatings, which confirmed the indexes in Fig. 4a. For detailed phase identification, the XRD line profile analysis (Pearson VII) was applied. The results of the XRD line fitting are shown in Fig. 5a, b and c on the examples of the low-angle diffraction lines. A slight general increase of the peaks broadening with increasing the carbon contents in the coatings is observed, which might indicate an increase in stress and a decrease in grain size [26]. Table 5 summarizes the lattice parameters and average crystallite size of the fcc phase in the respective samples, which were calculated using the sin2Ψ method [27]. The calculated lattice parameters of this fcc structure cannot exactly match any known phases, such as the lattice parameter for stoichiometric TiN (0.424 nm), CrN (0.417 nm) [28,29], TiC (0.431 nm) [30], AlN (0.412 nm) [31] and TiCN (0.426 nm) [32], but is close to CrN.
layer; a CrAlTiN alternative layer; and a carbon doped CrAlTiN multilayer sublayer. A schematical illustration of the Sublayers is shown in Fig. 8, and detailed TEM and SAD analysis is detailed below.
3.3. XPS analysis Similar bonding status of Cr2p, Al2s, Ti2p, C1s, O1s and N1s were obtained for different carbon doped coatings and a typical XPS spectrum of sample 5# is shown in Fig. 6. Although cleaning was conducted at the beginning to remove the top contamination layer, the oxygen peak was still detected in all the coatings. This could be the contaminated air present in the chamber during the deposition period and also oxygen absorption from air to the coating surface cannot be ruled out. The Ti2p and Al2s XPS spectra were very weak and the strong spectrum of N1s, C1s, Cr2p and O1s of the CrAlTiCN-5# coating are plotted in Fig. 7. The spectrum of N1s at 396.9 eV can be related to Cr–N bonds [33– 35]. In the Cr2p spectrum (Fig. 7a), the peak centred at 573.9 eV can be attributed to metallic Cr and Cr–N, while the tail of the peak could be assigned to Cr–O [36]. The shoulder at a higher binding energy of 583.5 eV can be assigned to Cr–C and/or Cr–N bonds [37,38]. The C1s spectrum (Fig. 7b) shows an asymmetrical peak which can be fitted into two components located at 282.3 and 284.6 eV. The 282.3 eV peak is attributed to Cr–C and the peak at 284.6 eV is typical for amorphous free carbon [26]. The N1s (Fig. 7c) of 396.9 eV and Cr2p 575.6 eV peak positions are in agreement with the literature data for the CrN phase [39].The peaks of O1s are decomposed to two peaks of 531.4 and 576.9. The maximum peak at 531.4 ± 0.2 eV is Cr–O [40,41], which is consistent with the Cr2p peak at 576.9 eV. The secondary peak could be attributed to the presence of Al2O3 (531.8 eV) [42]. 3.4. XTEM analysis TEM observation on 1#, 3#, and 5# samples revealed that the coatings contained four sublayers from the substrate to the surface and they are: a pure chromium interface layer; a chromium nitrides
Fig. 4. XRD (a) and GAXRD (b) patterns obtained on CrAlTiN and carbon doped CrAlTiN coatings.
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Table 5 Lattice parameters and grain sizes of the coatings calculated from GAXRD. Sample
Stress-free lattice parameter
Crystallite size
CrAlTiCN-1# CrAlTiCN-3# CrAlTiCN-5#
(4.135 ± 0.004) Å (4.152 ± 0.003) Å (4.141 ± 0.004) Å
(3.0 ± 0.5) nm (3.0 ± 0.6) nm (2.5 ± 0.4) nm
3.4.1. Sublayers I, II and III Based on the deposition process (Table 2), all coatings were processed with the same deposition parameters before carbon was introduced. Thus, coatings with different carbon contents possess the same structure and thickness for Sublayers I, II, and III. As shown in Fig. 9a, Sublayer Iis a pure chromium interface Sublayer and the layer possess columnar grains of about 150 nm diameter perpendicular to the interface. When the N2 was introduced with Cr deposition, a chromium nitride Sublayer II was formed. TEM microstructure (Fig. 9a) revealed that the layer is also columnar structured like Sublayer I but the columnar diameter in Sublayer IIis much thinner (30～50 nm) than that in Sublayer I. Indexing SAD patterns from Sublayers I and II (Fig. 9b) found that the layer is composed of Cr, CrN and Cr2N. The CrAlTiN multilayer (Sublayer III) (shown in Fig. 10d) was formed when samples were rotating against the targets: chromium, aluminum, and titanium, such that a multilayer with 5–10 nm intervals was formed. The contrast between the alternated layers is related to the change in the compositions, or the higher the atomic number, the darker the layer, whereas the contrast between the columnar structure is related to the diffraction conditions of the grains. There still existed many finer columnar microstructures (～100 nm in diameter) which do not extend through the entire thickness of Sublayer III. They nucleated at the interface between Sublayers IIand III as well as in the growth direction throughout Sublayer III. The SAD patterns from this CrAlTiN layer showed a typical fcc structure, a = 0.414 nm, with a preferred orientation (PR) of [110] parallel to the coating surface and two pairs of (200) arcs parallel and normal to the coating surface with ± 15° divations, respectively. 3.4.2. Sublayer IV It can be seen from Tables 2 and 3 that the process conditions and carbon contents are different for Sublayer IV of the coating samples. TEM observation on the carbon doped CrAlTiN layer or Sublayer IV revealed some common microstructure features. Firstly, Sublayer IV is characterised by columnar and multilayer structures shown in Figs. 10–13. Secondly, Sublayer IV possess a single cubic B1 NaCl-type structure as evidenced by the SAD patterns shown
Fig. 5. GAXRD patterns obtained on carbon doped CrAlTiN: (a) 1#, (b) 3# and (c) 5# coatings.
Fig. 6. XPS survey spectra of the CrAlTiCN-5# sample.
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Fig. 8. Schematical illustration of the layer structure.
Fig. 7. Fitting spectrum by XPS for carbon doped CrAlTiN-5# coating: (a) Cr2p, (b) C1s, (c) N1s.
in Figs. 11c, 12c and 13c. However, by comparing the SAD patterns from these three samples, it can be seen that the diffraction rings changed from sharp/strong orientation, diffuse/strong orientation to diffuse/no orientation for the low, middle and high carbon content coatings respectively, which indicates a reduction of crystallization and lose of preferred orientation as clearly indicated in the high resolution TEM microstructures (Figs. 11b, 12b and 13b). It can be seen from Fig. 11c that for the low carbon content coating, arcs of (200) planes are parallel and perpendicular to the coating surface. This indicates that the grains were growing in such a way as to keep two sets of the [110] zone direction parallel to the surface and the planes of {111} intersecting 45° ± 25° to the coating surface. The
Fig. 9. TEM micrograph (a) and electron diffraction pattern (b) on Sublayers I and II from CrAlTiCN-5# coating.
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Fig. 10. TEM micrographs (a), (d) and electron diffraction patterns (b), (c) on Sublayers III and IV from CrAlTiCN-5# coating.
lattice fringes show apparent crystalline structure and the lattice ordering is also visible showing continuity across multilayer boundaries (Fig. 11b). The origin of the contrast in this case in the bright field mode of the TEM which is due to the fact that the constituent layers have different atomic numbers and the brighter regions are from the lighter element (carbon) and the dark regions are from the heavier elements Cr, Al, and Ti. For the CrAlTiCN-3# sample, although the SAD pattern showed preferred orientation of [110] parallel to the coating surface, the diffraction rings are diffused (Fig. 12c). HRTEM proved that the lattice fringes discontinue across the multi-layer boundaries, where thin layers of amorphous (bright regions) are evident. Nano-scale grains of ～10 nm were formed in the layer. When a high quantity of carbon was introduced into the CrAlTiN coating (sample 5#), the SAD pattern (Fig. 13c) revealed broad continuous rings superimposed with diffused rings, characteristic of amorphous structures. The HRTEM microstructure (Fig. 13b) shows a higher ratio of the bright region (rich in carbon) to the dark region (rich in metal elements Al, Cr, and Ti), which are composed of amorphous and nanocrystalline (Cr, Al, Ti) (CN) (～13 nm) structures, respectively. It is noted that the crystallinity and the size of the crystalline changed along the distance from the interface between Sublayer III and Sublayer IV for the CrAlTiCN-5# coating, which has the highest carbon content. It can be seen from Fig. 14a that when carbon was first introduced into the coating, the nano-scale multi-layer showed thin and relatively sharp interfaces between the crystalline (Cr, Al, Ti)(C, N) layer (darker region) and the carbon layer (brighter region). Clear lattice fringes can be observed in the darker crystalline region while the carbon rich nano-layer (bright regions) is mainly amorphous embedded with few fine (～10 nm) crystalline. In the middle areas of the (Cr, Al, Ti)(C, N) sublayer (Fig. 14b), amorphous carbon layer is thicker and the crystalline grains of (Cr, Al, Ti)(C, N) is smaller than
those shown in Fig. 14a. At the outermost area, only extreme fine crystal clusters are embedded in the amorphous structure in both darker and brighter regions (Fig. 14c). 4. Discussion 4.1. The structure of the coatings The carbon doped (CrAlTi) N coatings show an fcc structure similar to that of the CrN, i.e. B1 (NaCl) structure. As reported by Knotek and Loffer [43], the Cr–Al–Ti–N–C system has no equilibrium phase. The aluminum and titanium atoms are substituted for chromium atoms in the lattice of CrN. These substitutions are attributed to both the unbalanced magnetic field and the high pressure effect during deposition, which tend to form a metastable (CrAlTi) (CN) single phase rather than mixed phases of CrN(C), AlN and TiN(C). As shown in Fig. 6c, the XPS spectrum of the CrAlTiN(C) coatings show the binding energy peak due to Cr–N bond, but does not show the energy peaks due to Ti–N and Al–N bonds. Furthermore, the shape of the Cr– N bond from the (CrAlTi)(CN) coating is similar to that of the pure CrN film. These facts suggest that the CrAlTiCN thin films have a structure in which “B1-AlN” and TiN are automatically dissolved in the CrN lattice. In other words, Al and Ti atoms in the (CrAlTi) (CN) film seem to exist keeping their coordination numbers. Substitution of Cr atoms with Al, Ti atoms in the B1-NaCl structure of CrN should result an expansion of the lattice parameter owing to the larger atomic radii of the substituted Al, Ti atoms. However, a contracted lattice was observed for all CrAlTiN(C) coatings (Table 5). Similar trends were also reported by Suzuki et.al [44] whereas the unit cell parameter of the (Cr, Al)N coating decreased as the Al ratio increased. This decrease in the unit cell parameter with Al is considered to be due to the small ionic radius of Al. Furthermore, it
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Fig. 11. TEM micrographs and electron diffraction pattern on Sublayer IVfrom CrAlTiCN1# coating.
Fig. 12. TEM micrographs and electron diffraction pattern on Sublayer IV from CrAlTiCN-3# coating.
is connected with the metastable phase formed during the coating process as suggested by Loffler [45]. The deviations from the stoichiometric composition of the (CrAlTi) (CN) phase, i.e. with 50 at.% (N, C) and 50 at.% (Cr, Al, Ti), results in unoccupied atomic lattice positions, thus distorting the entire lattice and leading to a reduction in the lattice parameter of the coatings.
size of the carbon clusters increase and a thicker layer of amorphous carbon phase is formed on the growth surface of the MN crystals. The growth of the MN crystal is interrupted by these C nanolayers decreasing the grain size of (CrAlTi) (CN) further (Fig. 13a and b). The occurrence of coexisting nanocrystalline/amorphous phases in ternary and quaternary systems was predicted by Holleck [47] on the basis of theoretical consideration of the thermodynamics and kinetics of PVD processes. Furthermore, similar growth mechanisms and microstructures for (Ti, Al) (C, N) coatings were reported by Streber et al. [37]. The coating structure evolution along the depth observed in sample 5# (Fig. 14) has never been reported in the literature. It can be seen that the columnar crystal structure of fcc MN gradually developed into nanocrystalline from the beginning of Sublayer III to the outermost of the coating surface. This may be caused by the accumulation of the carbon content in the coating system since the tendency for structural change is similar to the influence of carbon content on the structural change discussed above. Further investigation of the carbon content along the depth is needed to confirm this assumption.
4.2. Carbon effect Although a large quantity of carbon was introduced in Sublayer IV of the 1#–5# coatings, no trace of crystalline carbides could be found in the SAD patterns taken from them (Fig. 10). As presented in Section 3.2.3, instead an fcc structured nanocrystalline with an amorphous carbon nanocomposite coating was formed. The construction of the nanocomposite is carbon content dependent, i.e., the higher the carbon content, the lower the crystallinity and the weaker the preferred orientation of the coatings. Similar observations were also reported in the Ti–Al–N–C coating system [46]. This phenomenon is most likely related to the position of the carbon atoms within the coating structures during film growth procedures. From the chemical composition of the carbon free metastable nanocrystalline fcc (CrAlTi) N coating (Table.3 0# sample), an N-deficit is evident from the (Cr + Al+ Ti)/N ratio. When a low concentration of carbon, such as sample 1#, was sputtered on to the coating surface, the carbon was fully dissolved in the fcc MN phase, resulting in the growth of a metastable solid solution single fcc (CrAlTi) (CN) phase (Fig. 11a and b). The nanoscale multilayer period favours the necessary interdiffusion of carbon to MN. However, when the coating is not in an N-deficit, such as samples 3# and 5#, carbon atoms can either substitute regular N positions and/or start to build nanoclusters or agglomerate as a thin layer (Fig. 12a and b). When increasing the carbon content further, the
5. Conclusions The effects of carbon content on the phase composition and microstructure of carbon doped CrAlTiN coatings deposited by closedfield unbalanced magnetron sputtering have been investigated by XRD, SEM, TEM, and XPS. The carbon doped CrAlTiN coatings consisted of a pure columnar chromium interface layer, a columnar CrN layer, a (CrAlTi) N alternative columnar layer and a unique multilayer with nanocrystallites embedded in amorphous carbon. It was observed that carbon content is an important process parameter in controlling the microstructure and composition of the deposited CrAlTiCN coatings.
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the amorphous phase will lead to a decrease in the coating hardness and Young's modulus. Acknowledgements The authors gratefully acknowledge to financial support received from EPSRC, UK (EP/C536061/1), NSFC, China (50735006) and ECFP7 program (NMP3-SL-2008-213600) which made this work proceeding. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]
Fig. 13. TEM micrographs and electron diffraction pattern on Sublayer IV from CrAlTiCN-5# coating.
[17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]
When the C content is low (e.g. 2.34 at.% and 3.96 at.%), C atoms exist in the crystal lattice as solid solution. If the C concentration exceeds its solid solubility in the crystal phase, the excessive C will begin to form an amorphous carbon (or carbon-riched) phase. Cross-sectional SEM and HRTEM revealed a typical nanocomposite structure comprising nanocrystals embedded in an amorphous phase. The large volume fraction of
[29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47]
Fig. 14. High-resolution TEM micrographs on Sublayer IV from CrAlTiCN-5# coating: (a) initial region, (b) middle region and (c) top region.
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s e v i e r. c o m / l o c a t e / s u r f c o a t
Microstructural characterisation of carbon doped CrAlTiN nanoscale multilayer coatings Xiaoying Li ⁎, Wenwen Wu, Hanshan Dong School of Metallurgy and Materials, The University of Birmingham, Edgbaston, Birmingham B15 2TT, UK
a r t i c l e
i n f o
Article history: Received 21 April 2010 Accepted in revised form 17 November 2010 Available online 25 November 2010 Keywords: Carbon doped CrAlTiN coating Nanoscale multilayer TEM characterisation
a b s t r a c t Carbon doped CrAlTiN coatings for reducing the coefficient of friction sliding against a WC-Co ball in air were developed using a closed field unbalanced magnetron sputtering system (CFUMS). The carbon content was controlled by increasing the current of the carbon target power from 0 A to 5 A. The surface morphology, cross-sectional microstructure, phase constituent, chemical bonding energy and mechanical properties of coatings were characterised by means of XRD, SEM, XTEM, XPS, and nanoindentation. The results showed that the hardness, Young's modulus and coefficient of friction (CoF) of the coatings sliding against a WC-Co ball in air decreased with increasing carbon content. Microstructure characterisation revealed that the carbon doped coatings consisted of four sublayers from the substrate to the surface and the main phase formed is fcc B1NaCl like (Cr, Al, Ti) N and (Cr, Al, Ti) (C, N) phase in sublayers III and IV respectively when the carbon content is below 3.96 at.%. When the carbon concentration exceeds this value, the excessive carbon will begin to form the amorphous carbon (or carbon-riched) phase which leads to a decrease in the coefficient of friction (CoF) of the coatings sliding against a WC-Co ball in air. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The application of hard coatings on cutting tools and mechanical components has become a common practice in the manufacturing industry in the last two decades [1–5]. Conventional binary hard coatings such as TiN and CrN are commonly used to enhance the performance of cutting tools [6]. Recently, ternary Cr–X–N coatings, where X is the alloying element such as Ti [7,8], Al [9], Si [10], B [11], C [12], and Ni [13], etc., have been actively investigated to improve the properties of CrN based coatings. For example, in comparison to CrN coatings, Cr–Al–N coatings have a higher hardness (25–32 GPa) and much better oxidation resistance (up to 900 °C) due to the formation of a stable oxidation barrier of an Al2O3 layer by migrating Al atoms to the surface region [14]. Most recently, a CrN-based Cr–Ti–Al–N coating system has been developed by Teer Coatings to improve the performance of twist drills and extrusion dies [15,16]. It has been revealed that quaternary Cr–Ti–Al–N coatings have been designed to achieve even higher hardness and wear resistance as compared to that of the ternary coatings [3,17,18]. However, the coefficient of friction of these so-called superhard coatings is very high (0.7～0.9) [19] at room temperature, and extra lubrication is needed in many applications, which raises environmental concerns. On the other hand, graphite-like carbon (GLC) coatings were reported to have a very low coefficient of friction against most engineering materials [20]. However, GLC
⁎ Corresponding author. Tel.: + 44 121 4145243; fax: + 44 121 4147373. E-mail address: [email protected] (X. Li). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.11.046
degraded at high-temperatures and was found to be non-functional beyond 400 °C in air [21], which eliminates their use for dry and highspeed cutting tools. A complex interactive effect often determines the point at which a new synthetic coating is expected to fulfil it manufacturing purpose. By adding solid carbon into quaternary Cr–Al–Ti–N coatings with a nanocomposite structure may possess a more desirable combination of properties in terms of high hardness and low friction. However, there are few studies on the design and synthesis of such new coatings and there is little work on the influence of the C concentrations on the structure, microstructure and mechanical properties of carbon doped CrAlTiN coatings. In the present study, new carbon doped CrAlTiN coatings were designed and deposited by magnetron sputtering ion plating with carbon content ranging from 0 to 24.34 at.%. Detailed TEM characterization has been conducted to investigate the effect of carbon content on the structure, microstructure, and mechanical properties of the CrAlTiCN coatings, thus advancing the scientific understanding of the role of carbon in these new coatings. 2. Experimental procedures 2.1. Substrate materials The substrate materials used in this study were hardened (HRC≧62) M2 high speed steel and a Si wafer. The composition of the M2 steel is given in Table 1. The M2 discs were ground using silicon carbide paper from 120 grit down to 2400 grit, polished with 6 μm and 1 μm diamond paste and finally polished with 0.25 μm colloidal silica. The samples
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Table 1 Nominal chemical composition of M2 high speed steel. Element
C
Si
Mn
Cr
Mo
V
W
Fe
Composition (wt.%)
0.78–0.88
0.20–0.40
0.20–0.40
3.75–4.50
4.50–5.50
1.60–2.20
5.50–6.75
Balance
were cleaned in soapy water and then acetone prior to and after surface polishing. 2.2. Coating equipment Cr–Al–Ti–C–N coatings with different carbon concentrations were prepared using the UDP 450 closed-field unbalanced magnetron sputtering ion plating (CFUBMSIP) system [15,16]. The schematic illustration of the system is shown in Fig. 1. It can be seen that the sputtering system contains four rectangular cathodes with Cr, Ti, graphite, and Al targets, which is operated in the unbalanced magnetron mode. The closed magnetic field coupling results in a high degree of ionization and a high bias current density. The equipment has a rotating cylindrical substrate holder with a speed of 5 rpm. After cleaning, the samples were inserted into the deposition chamber and the system was pumped down to a high vacuum condition with a background pressure less than 3.0×10− 3 Pa. Firstly argon and then a mixture of argon and nitrogen gases was introduced into the chamber during the process with a fixed Ar flow rate of 20 sccm, and the nitrogen content was controlled by a plasma optical emission monitor (OEM) with a feedback control of 60–70%. 2.3. Deposition processes A−70 V DC bias was applied on the substrate during deposition. The process started with plasma surface cleaning with argon ions. Deposition is started with a Cr adhesion layer, and then a Cr–N transition layer by gradually introducing nitrogen gas. Following
these two sublayers, a compositionally graded Cr–Al–Ti–N sublayer was deposited. For samples 1#–5#, an extra carbon doped Cr–Al–Ti–N top layer was deposited. The relative concentration of Cr, Ti, Al, and C in the coatings was adjusted by the sputtering power applied to the targets during deposition. The range of carbon content was controlled by changing carbon target current from 0 A to 5 A. Detailed deposition process conditions for the carbon doped CrAlTiN coatings are summarized in Table 2. 2.4. Chemical composition and bonding Surface morphologies and cross-sectional microstructure of the deposited samples were examined using a field emission JEOL7000 scanning electron microscope (SEM) with EDX and WDX capabilities. The composition analysis was carried out both qualitatively to determine the existence of certain elements and quantitatively to determine the amount of each element present. The chemical bonding of the carbon doped CrAlTiN coatings was studied by X-ray Photoelectron Spectroscopy (XPS), using a AXISULTRAX unit. The spectrometer was equipped with a hemispherical energy analyzer. An X-ray tube with monochromatic Al Kα radiation (hv = 1486.6 ev) was used as a signal excitation source. The investigation was carried out under vacuum of the order of 10− 8 Pa. Samples were subjected to etching by argon ions with Ar + ion bombardment of energy 4 keV and ion current of 1.0 μA to remove the top contamination layer, which can reduce the occurrence of preferential sputtering in the etching process. 2.5. Microstructure characterisation
Fig. 1. Cross-section schematic drawing of arrangement plan of magnetron targets and substrate holder in the CFUBMSIP system.
The phase constituents of the coated surfaces were analysed in normal XRD, X'Pert Philips diffractometer, and glancing angle XRD (GXRD) D8 diffractometer (Bruker AXS) with Cu-Kα radiation (λ = 0.154 nm). The diffraction patterns obtained were analysed and indexed using X'Pert High Score analytical software with a PCPDFWIN data base [22]. Crystallite size was calculated by the Rietveld method and the line broadening (the width of the XRD lines) as obtained from the line profile fitting was used to estimate the crystallite size [23]. In order to investigate the effect of carbon content on microstructure, representative CrAlTiCN-1#, 3#, and 5# samples were selected for the TEM studies. Cross-sectional TEM samples were prepared as follows [24]: (1) Two slabs were glued using G-1 epoxy with the treated surfaces facing each other; (2) The assembly was cut to a thickness of less than 1 mm with the interface of the two layers glued together in the centre followed by fine grinding and polishing to 50 μm thickness; (3) This thin assembly was then glued to a 3 mm (o.d) and 0.5 mm (i.d) brass reinforcing disc. The exposed side was further polished to a thickness of approximately 20 μm; (4) The sample was then placed in an ion beam miller (Gatan 691 Precision Ion Polisher System, PIPS™), making sure
Table 2 Deposition conditions for the synthesis of carbon doped CrAlTiN coatings. Sample code
CrAlTiCN-0#
Target current, A Cr/Al/Ti
4.5/3.5/3.5 A
C target current 0A Substrate bias Base pressure Substrate rotation speed / Temperature, °C
CrAlTiCN-1#
CrAlTiCN-2#
CrAlTiCN-3#
CrAlTiCN-4#
CrAlTiCN-5#
4.5/3.5/3.5 A
4.5/3.5/3.5 A
4.5/3.5/3.5 A
4.5/3.5/3.5 A
4.5/3.5/3.5 A
1A
2A
3A − 70 V b 3.0 × 10− 3 Pa 5 rpm / ~ 200°C
4A
5A
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that the milling was done on the interface of the glued surface; (5) Milling was terminated once a desired thin area around the interface of the glued surfaces is created. FEI Tecnai F20-FEG (EDX) was used for the TEM study on the microstructure and phase constitutions.
Table 3 Elemental concentration of carbon doped CrAlTiN coatings by EDX and WDX. Sample
Element (at%) WDX N
2.6. Mechanical properties Nanoindentation hardness (H) and Young's modulus (E) were measured using a Nano-Test 600 machine (Micromaterials, UK). A Berkovick diamond tip (3-faced pyramid) with a radius of about 200 nm at its apex was used in the tests. Based on the thickness of the coatings, E and H were measured continuously up to an indentation depth of 1200 nm for each sample, and then calculated using the Oliver and Pharr method [25]. Under non-lubrication conditions, wear tests were carried out by a pin-on-disc tribometer, using a stationary 5 mm WC-Co counter ball (Grade 10, Spheric-Trafalgar Ltd) at a speed of 0.025 m/s (420 rpm) for 1 h. The normal contact load acting on the ball was 20 N and a wear track of 11 mm in diameter was produced. A stylus profilometer, Surf Corder SE 1700, was used to measure the surface roughness of the tested samples.
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CrAlTiN-0# CrAlTiCN-1# CrAlTiCN-2# CrAlTiCN-3# CrAlTiCN-4# CrAlTiCN-5#
45.98 44.12 47.02 44.79 41.87 33.16
EDX C
Al
Ti
Cr
N
C
2.34 3.96 9.65 12.53 24.34
6.87 6.96 6.58 5.91 6.09 5.78
6.07 6.45 5.09 5.87 5.33 5.35
40.53 40.13 37.35 33.78 34.18 31.37
46.53 43.54 46.29 45.69 42.14 35.18
2.92 4.69 8.75 12.26 22.32
3. Results 3.1. General characterization SEM observation found no cracks in the interfaces between the coating layer and the substrate for all the coatings deposited by the CFUBMSIP method, which indicated a good adherence of the coatings to the metal substrates. The surface morphologies and cross-sectional microstructures for all designed coatings revealed similar features under SEM and the typical SEM images are presented in Fig. 2. The coating surfaces showed columnar boundaries and the average columnar diameter was scattered between 100 to 300 nm (Fig. 2a). All the coatings are very dense and there is no sign of cracking on the surface or within the coatings, which indicates a good compactness of the coatings. All coatings have a similar thickness of about 5 μm. However, compared with 0#, the carbon doped coatings show two sublayers with different contrasts (Fig. 2b), which were caused by the change in the chemical composition of the coatings (see Table 2). The light elements C and N and metal elements Cr, Al, Ti in the coatings were analysed by WDX and EDX respectively, and the results are shown in Table 3. It can be seen that the carbon contents in the coatings increased with the carbon target current. When increasing the carbon target current from 0 to 5 A, the carbon content of the CrAlTiCN coatings was increased to 24.34 at.%. The metallic content
Fig. 3. Nano-hardness and Young's modulus variation with CrAlTiN and carbon doped CrAlTiN coatings.
ratios, on the other hand, did not change significantly in the doped coatings, and the nitrogen content in all deposited samples is around 40～50 at.%. It can be seen from Fig. 3 that the hardness and elastic modulus for all coatings show an approximately linear trend with carbon content, both of which decreased with increasing the carbon target current. As shown in Table 4, all the carbon doped CrAlTiN coatings possess very similar roughness values. An obvious increase of Ra was observed with increasing the carbon content. However, the coefficient of friction (CoF) value of the coatings against a WC-Co ball in air was found to decrease with increasing the carbon contents and the CoF for the coating 5# is about 0.350, which is about 47% lower than that of the carbon free coating, 0# sample.
Fig. 2. Surface (a) and cross-sectional (b) SEI SEM micrographs of the CrAlTiCN-5# sample.
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Table 4 Surface roughness and friction coefficient of carbon doped CrAlTiN coatings. Sample
0#
1#
2#
3#
4#
5#
Roughness (μm) CoF
0.0303 ± 0.0013 0.655 ± 0.140
0.0315 ± 0.0005 0.420 ± 0.132
0.0378 ± 0.0018 0.410 ± 0.090
0.0372 ± 0.0028 0.375 ± 0.083
0.0405 ± 0.0035 0.400 ± 0.080
0.0388 ± 0.0022 0.350 ± 0.065
3.2. XRD phase identification Fig. 4a shows the XRD patterns obtained from the carbon-free and carbon-doped coatings. Due to the substrate contribution, it is difficult to distinguish the peaks accurately. A tentative indexing shown in Fig. 4a displayed fcc (111), (200), (220), and (311) reflections of all the coatings, and the (110), (200), (211), and (220) reflections of the bcc α-Fe substrate. The intensity of the peaks is relatively weak except the peak around 2θ = 45°, which was superimposed with α-Fe (110) from the substrate. In order to avoid substrate interference effects, glancing angle XRD was employed and the typical GAXRD charts are given in Fig. 4b. As shown in Fig. 4b, a set of fcc (111), (200), (220), (311), (222), and (400) reflections obtained from the coatings, which confirmed the indexes in Fig. 4a. For detailed phase identification, the XRD line profile analysis (Pearson VII) was applied. The results of the XRD line fitting are shown in Fig. 5a, b and c on the examples of the low-angle diffraction lines. A slight general increase of the peaks broadening with increasing the carbon contents in the coatings is observed, which might indicate an increase in stress and a decrease in grain size [26]. Table 5 summarizes the lattice parameters and average crystallite size of the fcc phase in the respective samples, which were calculated using the sin2Ψ method [27]. The calculated lattice parameters of this fcc structure cannot exactly match any known phases, such as the lattice parameter for stoichiometric TiN (0.424 nm), CrN (0.417 nm) [28,29], TiC (0.431 nm) [30], AlN (0.412 nm) [31] and TiCN (0.426 nm) [32], but is close to CrN.
layer; a CrAlTiN alternative layer; and a carbon doped CrAlTiN multilayer sublayer. A schematical illustration of the Sublayers is shown in Fig. 8, and detailed TEM and SAD analysis is detailed below.
3.3. XPS analysis Similar bonding status of Cr2p, Al2s, Ti2p, C1s, O1s and N1s were obtained for different carbon doped coatings and a typical XPS spectrum of sample 5# is shown in Fig. 6. Although cleaning was conducted at the beginning to remove the top contamination layer, the oxygen peak was still detected in all the coatings. This could be the contaminated air present in the chamber during the deposition period and also oxygen absorption from air to the coating surface cannot be ruled out. The Ti2p and Al2s XPS spectra were very weak and the strong spectrum of N1s, C1s, Cr2p and O1s of the CrAlTiCN-5# coating are plotted in Fig. 7. The spectrum of N1s at 396.9 eV can be related to Cr–N bonds [33– 35]. In the Cr2p spectrum (Fig. 7a), the peak centred at 573.9 eV can be attributed to metallic Cr and Cr–N, while the tail of the peak could be assigned to Cr–O [36]. The shoulder at a higher binding energy of 583.5 eV can be assigned to Cr–C and/or Cr–N bonds [37,38]. The C1s spectrum (Fig. 7b) shows an asymmetrical peak which can be fitted into two components located at 282.3 and 284.6 eV. The 282.3 eV peak is attributed to Cr–C and the peak at 284.6 eV is typical for amorphous free carbon [26]. The N1s (Fig. 7c) of 396.9 eV and Cr2p 575.6 eV peak positions are in agreement with the literature data for the CrN phase [39].The peaks of O1s are decomposed to two peaks of 531.4 and 576.9. The maximum peak at 531.4 ± 0.2 eV is Cr–O [40,41], which is consistent with the Cr2p peak at 576.9 eV. The secondary peak could be attributed to the presence of Al2O3 (531.8 eV) [42]. 3.4. XTEM analysis TEM observation on 1#, 3#, and 5# samples revealed that the coatings contained four sublayers from the substrate to the surface and they are: a pure chromium interface layer; a chromium nitrides
Fig. 4. XRD (a) and GAXRD (b) patterns obtained on CrAlTiN and carbon doped CrAlTiN coatings.
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Table 5 Lattice parameters and grain sizes of the coatings calculated from GAXRD. Sample
Stress-free lattice parameter
Crystallite size
CrAlTiCN-1# CrAlTiCN-3# CrAlTiCN-5#
(4.135 ± 0.004) Å (4.152 ± 0.003) Å (4.141 ± 0.004) Å
(3.0 ± 0.5) nm (3.0 ± 0.6) nm (2.5 ± 0.4) nm
3.4.1. Sublayers I, II and III Based on the deposition process (Table 2), all coatings were processed with the same deposition parameters before carbon was introduced. Thus, coatings with different carbon contents possess the same structure and thickness for Sublayers I, II, and III. As shown in Fig. 9a, Sublayer Iis a pure chromium interface Sublayer and the layer possess columnar grains of about 150 nm diameter perpendicular to the interface. When the N2 was introduced with Cr deposition, a chromium nitride Sublayer II was formed. TEM microstructure (Fig. 9a) revealed that the layer is also columnar structured like Sublayer I but the columnar diameter in Sublayer IIis much thinner (30～50 nm) than that in Sublayer I. Indexing SAD patterns from Sublayers I and II (Fig. 9b) found that the layer is composed of Cr, CrN and Cr2N. The CrAlTiN multilayer (Sublayer III) (shown in Fig. 10d) was formed when samples were rotating against the targets: chromium, aluminum, and titanium, such that a multilayer with 5–10 nm intervals was formed. The contrast between the alternated layers is related to the change in the compositions, or the higher the atomic number, the darker the layer, whereas the contrast between the columnar structure is related to the diffraction conditions of the grains. There still existed many finer columnar microstructures (～100 nm in diameter) which do not extend through the entire thickness of Sublayer III. They nucleated at the interface between Sublayers IIand III as well as in the growth direction throughout Sublayer III. The SAD patterns from this CrAlTiN layer showed a typical fcc structure, a = 0.414 nm, with a preferred orientation (PR) of [110] parallel to the coating surface and two pairs of (200) arcs parallel and normal to the coating surface with ± 15° divations, respectively. 3.4.2. Sublayer IV It can be seen from Tables 2 and 3 that the process conditions and carbon contents are different for Sublayer IV of the coating samples. TEM observation on the carbon doped CrAlTiN layer or Sublayer IV revealed some common microstructure features. Firstly, Sublayer IV is characterised by columnar and multilayer structures shown in Figs. 10–13. Secondly, Sublayer IV possess a single cubic B1 NaCl-type structure as evidenced by the SAD patterns shown
Fig. 5. GAXRD patterns obtained on carbon doped CrAlTiN: (a) 1#, (b) 3# and (c) 5# coatings.
Fig. 6. XPS survey spectra of the CrAlTiCN-5# sample.
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Fig. 8. Schematical illustration of the layer structure.
Fig. 7. Fitting spectrum by XPS for carbon doped CrAlTiN-5# coating: (a) Cr2p, (b) C1s, (c) N1s.
in Figs. 11c, 12c and 13c. However, by comparing the SAD patterns from these three samples, it can be seen that the diffraction rings changed from sharp/strong orientation, diffuse/strong orientation to diffuse/no orientation for the low, middle and high carbon content coatings respectively, which indicates a reduction of crystallization and lose of preferred orientation as clearly indicated in the high resolution TEM microstructures (Figs. 11b, 12b and 13b). It can be seen from Fig. 11c that for the low carbon content coating, arcs of (200) planes are parallel and perpendicular to the coating surface. This indicates that the grains were growing in such a way as to keep two sets of the [110] zone direction parallel to the surface and the planes of {111} intersecting 45° ± 25° to the coating surface. The
Fig. 9. TEM micrograph (a) and electron diffraction pattern (b) on Sublayers I and II from CrAlTiCN-5# coating.
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Fig. 10. TEM micrographs (a), (d) and electron diffraction patterns (b), (c) on Sublayers III and IV from CrAlTiCN-5# coating.
lattice fringes show apparent crystalline structure and the lattice ordering is also visible showing continuity across multilayer boundaries (Fig. 11b). The origin of the contrast in this case in the bright field mode of the TEM which is due to the fact that the constituent layers have different atomic numbers and the brighter regions are from the lighter element (carbon) and the dark regions are from the heavier elements Cr, Al, and Ti. For the CrAlTiCN-3# sample, although the SAD pattern showed preferred orientation of [110] parallel to the coating surface, the diffraction rings are diffused (Fig. 12c). HRTEM proved that the lattice fringes discontinue across the multi-layer boundaries, where thin layers of amorphous (bright regions) are evident. Nano-scale grains of ～10 nm were formed in the layer. When a high quantity of carbon was introduced into the CrAlTiN coating (sample 5#), the SAD pattern (Fig. 13c) revealed broad continuous rings superimposed with diffused rings, characteristic of amorphous structures. The HRTEM microstructure (Fig. 13b) shows a higher ratio of the bright region (rich in carbon) to the dark region (rich in metal elements Al, Cr, and Ti), which are composed of amorphous and nanocrystalline (Cr, Al, Ti) (CN) (～13 nm) structures, respectively. It is noted that the crystallinity and the size of the crystalline changed along the distance from the interface between Sublayer III and Sublayer IV for the CrAlTiCN-5# coating, which has the highest carbon content. It can be seen from Fig. 14a that when carbon was first introduced into the coating, the nano-scale multi-layer showed thin and relatively sharp interfaces between the crystalline (Cr, Al, Ti)(C, N) layer (darker region) and the carbon layer (brighter region). Clear lattice fringes can be observed in the darker crystalline region while the carbon rich nano-layer (bright regions) is mainly amorphous embedded with few fine (～10 nm) crystalline. In the middle areas of the (Cr, Al, Ti)(C, N) sublayer (Fig. 14b), amorphous carbon layer is thicker and the crystalline grains of (Cr, Al, Ti)(C, N) is smaller than
those shown in Fig. 14a. At the outermost area, only extreme fine crystal clusters are embedded in the amorphous structure in both darker and brighter regions (Fig. 14c). 4. Discussion 4.1. The structure of the coatings The carbon doped (CrAlTi) N coatings show an fcc structure similar to that of the CrN, i.e. B1 (NaCl) structure. As reported by Knotek and Loffer [43], the Cr–Al–Ti–N–C system has no equilibrium phase. The aluminum and titanium atoms are substituted for chromium atoms in the lattice of CrN. These substitutions are attributed to both the unbalanced magnetic field and the high pressure effect during deposition, which tend to form a metastable (CrAlTi) (CN) single phase rather than mixed phases of CrN(C), AlN and TiN(C). As shown in Fig. 6c, the XPS spectrum of the CrAlTiN(C) coatings show the binding energy peak due to Cr–N bond, but does not show the energy peaks due to Ti–N and Al–N bonds. Furthermore, the shape of the Cr– N bond from the (CrAlTi)(CN) coating is similar to that of the pure CrN film. These facts suggest that the CrAlTiCN thin films have a structure in which “B1-AlN” and TiN are automatically dissolved in the CrN lattice. In other words, Al and Ti atoms in the (CrAlTi) (CN) film seem to exist keeping their coordination numbers. Substitution of Cr atoms with Al, Ti atoms in the B1-NaCl structure of CrN should result an expansion of the lattice parameter owing to the larger atomic radii of the substituted Al, Ti atoms. However, a contracted lattice was observed for all CrAlTiN(C) coatings (Table 5). Similar trends were also reported by Suzuki et.al [44] whereas the unit cell parameter of the (Cr, Al)N coating decreased as the Al ratio increased. This decrease in the unit cell parameter with Al is considered to be due to the small ionic radius of Al. Furthermore, it
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Fig. 11. TEM micrographs and electron diffraction pattern on Sublayer IVfrom CrAlTiCN1# coating.
Fig. 12. TEM micrographs and electron diffraction pattern on Sublayer IV from CrAlTiCN-3# coating.
is connected with the metastable phase formed during the coating process as suggested by Loffler [45]. The deviations from the stoichiometric composition of the (CrAlTi) (CN) phase, i.e. with 50 at.% (N, C) and 50 at.% (Cr, Al, Ti), results in unoccupied atomic lattice positions, thus distorting the entire lattice and leading to a reduction in the lattice parameter of the coatings.
size of the carbon clusters increase and a thicker layer of amorphous carbon phase is formed on the growth surface of the MN crystals. The growth of the MN crystal is interrupted by these C nanolayers decreasing the grain size of (CrAlTi) (CN) further (Fig. 13a and b). The occurrence of coexisting nanocrystalline/amorphous phases in ternary and quaternary systems was predicted by Holleck [47] on the basis of theoretical consideration of the thermodynamics and kinetics of PVD processes. Furthermore, similar growth mechanisms and microstructures for (Ti, Al) (C, N) coatings were reported by Streber et al. [37]. The coating structure evolution along the depth observed in sample 5# (Fig. 14) has never been reported in the literature. It can be seen that the columnar crystal structure of fcc MN gradually developed into nanocrystalline from the beginning of Sublayer III to the outermost of the coating surface. This may be caused by the accumulation of the carbon content in the coating system since the tendency for structural change is similar to the influence of carbon content on the structural change discussed above. Further investigation of the carbon content along the depth is needed to confirm this assumption.
4.2. Carbon effect Although a large quantity of carbon was introduced in Sublayer IV of the 1#–5# coatings, no trace of crystalline carbides could be found in the SAD patterns taken from them (Fig. 10). As presented in Section 3.2.3, instead an fcc structured nanocrystalline with an amorphous carbon nanocomposite coating was formed. The construction of the nanocomposite is carbon content dependent, i.e., the higher the carbon content, the lower the crystallinity and the weaker the preferred orientation of the coatings. Similar observations were also reported in the Ti–Al–N–C coating system [46]. This phenomenon is most likely related to the position of the carbon atoms within the coating structures during film growth procedures. From the chemical composition of the carbon free metastable nanocrystalline fcc (CrAlTi) N coating (Table.3 0# sample), an N-deficit is evident from the (Cr + Al+ Ti)/N ratio. When a low concentration of carbon, such as sample 1#, was sputtered on to the coating surface, the carbon was fully dissolved in the fcc MN phase, resulting in the growth of a metastable solid solution single fcc (CrAlTi) (CN) phase (Fig. 11a and b). The nanoscale multilayer period favours the necessary interdiffusion of carbon to MN. However, when the coating is not in an N-deficit, such as samples 3# and 5#, carbon atoms can either substitute regular N positions and/or start to build nanoclusters or agglomerate as a thin layer (Fig. 12a and b). When increasing the carbon content further, the
5. Conclusions The effects of carbon content on the phase composition and microstructure of carbon doped CrAlTiN coatings deposited by closedfield unbalanced magnetron sputtering have been investigated by XRD, SEM, TEM, and XPS. The carbon doped CrAlTiN coatings consisted of a pure columnar chromium interface layer, a columnar CrN layer, a (CrAlTi) N alternative columnar layer and a unique multilayer with nanocrystallites embedded in amorphous carbon. It was observed that carbon content is an important process parameter in controlling the microstructure and composition of the deposited CrAlTiCN coatings.
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the amorphous phase will lead to a decrease in the coating hardness and Young's modulus. Acknowledgements The authors gratefully acknowledge to financial support received from EPSRC, UK (EP/C536061/1), NSFC, China (50735006) and ECFP7 program (NMP3-SL-2008-213600) which made this work proceeding. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]
Fig. 13. TEM micrographs and electron diffraction pattern on Sublayer IV from CrAlTiCN-5# coating.
[17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]
When the C content is low (e.g. 2.34 at.% and 3.96 at.%), C atoms exist in the crystal lattice as solid solution. If the C concentration exceeds its solid solubility in the crystal phase, the excessive C will begin to form an amorphous carbon (or carbon-riched) phase. Cross-sectional SEM and HRTEM revealed a typical nanocomposite structure comprising nanocrystals embedded in an amorphous phase. The large volume fraction of
[29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47]
Fig. 14. High-resolution TEM micrographs on Sublayer IV from CrAlTiCN-5# coating: (a) initial region, (b) middle region and (c) top region.
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