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Synthesis and characterization of quaternary Ti–Si–C–N coatings prepared by a hybrid deposition technique
Surface & Coatings Technology 188–189 (2004) 415 – 419 www.elsevier.com/locate/surfcoat
Synthesis and characterization of quaternary Ti–Si–C–N coatings prepared by a hybrid deposition technique Jun-Ha Jeon, Sung Ryong Choi, Won Sub Chung, Kwang Ho Kim* School of Materials Science and Engineering, Pusan National University, San 30, Changjeon-dong, Keumjeong Busan 609-735, South Korea Available online 11 September 2004
Abstract Quaternary Ti–Si–C–N coatings have been deposited on WC–Co substrates by a hybrid system combining the arc ion plating (AIP) and sputtering techniques using Ti and Si targets in an Ar/N2/CH4 gaseous mixture. Ti–Si–C–N coatings with a Si content of 8.9 at.% had a fine composite microstructure comprising nano-sized crystallites of Ti(C,N) surrounded by amorphous phase of Si3N4/SiC mixture. The microhardness value of the Ti–Si–C–N coatings under a load of 25 g was ~55 GPa, i.e., much larger than ~30 GPa for Ti–C–N coatings. In addition, the average friction coefficient of the Ti–Si–C–N coatings largely decreased with increasing Si content. In this work, the microstructure and mechanical properties of Ti–Si–C–N coatings were systematically investigated using instrumental analyses. D 2004 Elsevier B.V. All rights reserved. Keywords: Ti–Si–C–N coatings; Arc ion plating; Nanocomposite; Superhardness; Wear behavior
1. Introduction Ti–C–N coatings have been widespread applied to various tools and dies because of high hardness (~30 GPa), good wear property, and improved corrosion resistance compared to TiN coatings [1–5]. The higher hardness value of Ti–C–N coatings than that of TiN coatings was attributed to the solid-solution hardening by carbon atoms. Ti–C–N coatings were studied mainly by chemical vapor deposition techniques [6–8]. On the other hand, arc ion plating techniques have been adopted to improve the mechanical properties of Ti–C–N coatings [9,10]. Recently, some groups [11–13] reported ternary Ti–Si–N coatings where some amount of amorphous silicon nitride was incorporated into TiN coatings. The Ti–Si–N coatings have microstructures characterized by a nanocomposite of fine TiN crystallites encapsulated with amorphous silicon nitride [14,15]. Such a microstructural modification of TiN by the Si addition caused a superhardness (z40GPa), and other improvements in wear resistant and oxidation resistant properties [16–18]. * Corresponding author. Tel.: +82 51 510 2391; fax: +82 51 510 3660. E-mail address: [email protected] (K.H. Kim). 0257-8972/$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2004.08.042
Quaternary Ti–Si–C–N system can have superior properties tailored through the microstructural designing from both Ti–C–N and Ti–Si–N coatings. Up to now, only a few papers have been reported on Ti–Si–C–N coating systems [19–21]. However, their properties as well as the microstructure were not thoroughly investigated. In this work, quaternary Ti–Si–C–N coatings have been deposited on WC–Co substrates using a hybrid system combining the arc ion plating (AIP) and DC magnetron sputtering techniques. The microstructure and mechanical properties of Ti–Si–C–N coatings were systematically investigated.
2. Experimental 2.1. Deposition of Ti–Si–C–N coatings Ti–Si–C–N coatings have been deposited on WC–Co substrates by the hybrid coating system, where the AIP method was combined with a magnetron sputtering technique. A schematic diagram of the hybrid deposition system was depicted in previous work [22–24]. An arc cathode gun for the Ti source and a DC sputter gun for the Si source were installed on each side of the chamber wall. A
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Table 1 Typical deposition conditions for Ti–Si–C–N films by hybrid coating system Base pressure Working pressure Working gas ratio Arc material Sputter material Arc current Sputter current Substrate temperature Substrate bias voltage Rotational velocity of substrate Typical deposition rate
9.3310 3 Pa 6.6710 2 Pa CH4/N2=20:40 Ti (99.99%) Si (99.99%) 60 A 0–2.2 A 300 8C 60 V 25 rpm ~3 Am/h
rotational substrate holder was located on a straight line between two sources. The WC–Co substrates, which had been machined to disc types of 20 mm in diameter and 3 mm in thickness, were cleaned in an ultrasonic bath cleaner using acetone and alcohol, and were cleaned again by ion bombardment using a bias voltage of 600 V under Ar atmosphere of 32 Pa for 12 min just before deposition. The typical deposition conditions for Ti–Si–C–N coatings are summarized in Table 1. 2.2. Characterization of Ti–Si–C–N coatings The thickness and surface morphology of coatings were measured using scanning electron microscopy (SEM, Hitach, S-4200) and a stylus (a-STEP) instrument. Compositional analyses of the coatings to determine the Ti, Si, C, and N contents were carried out by electron probe microanalyzer (EPMA, Shimadzu, EPMA1600). In order to identify the crystallographic structure of Ti–Si–C–N coatings, X-ray diffractometer using CuKa radiation was used. The crystallite size in the coating was determined from the direct observation by a field emission-transmission electron microscope (FE-TEM, JEOL, JEM-2010F) operating at 200 kV. The other structural information on the film components was obtained from the analyses by the selected area diffraction pattern (SADP) and high-resolution transmission electron microscopy (HRTEM). X-ray photoelectron spectroscopy (XPS, VG Scientifics, ESCALAB 250) was used to observe the chemical bonding status of Ti–Si–C–N coatings. The hardness of coatings was evaluated using a microhardness tester with Knoop indentor (Matsuzawa, MMT-7) under a load of 25 g. The average friction coefficient and wear behavior were studied through sliding tests using a conventional ball-on-disc wear apparatus. Steel balls (diameter 12 mm, 700 Hv0.2) were used as a counterpart material. The sliding tests were conducted with a sliding speed of 0.2 m/s under a load of 1 N at ambient temperature (approx. 19.8 8C) and relative humidity (RH) of 25–30%. The optical microscopy and scanning electron microscopy were employed to observe the morphology of wear tracks after each sliding test. The energy dispersive spectroscopy (EDS)
was used to reveal the composition of wear debris formed during wear test.
3. Results and discussion 3.1. Synthesis and characterization of Ti–Si–C–N coatings Fig. 1 shows the changes of Ti, Si, C, and N contents in at.% for Ti–Si–C–N coatings as a function of the d.c. sputter current collected by the Si target at a fixed Ti arc current of 60 A. The Si content in Ti–Si–C–N coatings linearly increased from 0 to 15.5 at.% with raising the d.c. sputter current from 0 to 2.2 A, whereas the nitrogen and carbon contents in Ti–Si–C–N coatings were nearly constant. Fig. 2 shows the X-ray diffraction patterns of Ti–Si–C–N coatings with various Si contents. The diffraction patterns of Ti–Si–C–N coatings showed multiple orientations of (111), (200), (220), (311), and (222) of Ti(C,N) crystal planes. Ti(C,N) has same crystal structure of f.c.c. as TiN, but interplanar d spacing becomes a little higher as carbon atoms replace nitrogen ones. These multiple orientations were in good agreement with other Ti–C–N coatings [6,8,25,26]. Other diffraction peaks from crystalline phases such as Si3N4, SiC, and titanium silicide phase like TiSi2 were not detected. This result postulated that Si existed as an amorphous phase of silicon nitride, silicon carbide, or free Si. As the Si was incorporated into the Ti(C,N), the diffraction peak intensities gradually reduced and started to disappear above the Si content of 15.5 at.%. In addition, peak broadening phenomenon was observed with increasing of Si content in Ti(C,N) coatings. Similar XRD peak broadenings have been reported for other ternary Ti–Si–N and quaternary Ti–Al–Si–N coating with Si addition [22,23,27–29], and was attributed mainly to the effect of the diminution of the crystallite size [14,18,30,31]. Fig. 3 shows the cross-sectional HRTEM
Fig. 1. Compositional changes of Ti–Si–C–N coatings as a function of DC sputter current to Si target at fixed Ti arc current of 60 A.
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Fig. 2. The X-ray diffraction patterns of Ti–Si–C–N coatings with various Si contents.
image, selected area electron diffraction (SAED) pattern, and dark-field TEM image for the Ti–Si–C–N coating with Si content of 8.9 at.%. From the results included in Fig. 3, it was established that the Ti–Si–C–N coatings were composed of Ti(C,N)-based crystallites and Si-based amorphous phase, which were clearly distinguished from each other by lattice fringe contrast (Fig. 3a). The Ti–Si– C–N coating with 8.9 at.% Si was regarded as a nanocomposite having fine crystallites of approximately 8 nm in size embedded in an amorphous matrix. The dark-
Fig. 3. Cross-sectional HRTEM image, electron diffraction pattern, and dark-field TEM image for Ti–Si–C–N coating with Si content of 8.9 at.%.
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field TEM image (Fig. 3b) of the Ti–Si–C–N coatings with a Si content of 8.9 at.% also shows again that fine Ti(C,N) crystallites (white area) were almost uniformly distributed in the matrix. The maximum hardness of the coatings could be obtained in this case as shown later (Fig. 5). The present microstructural evolution of Ti(C, N) coatings with Si addition is similarly reported to other systems such as Ti–Si–N [14,24,32] and Ti–Al–Si–N coatings [23]. In order to clarify bonding status of the amorphous phase comprising Ti–Si–C–N coatings, XPS analyses were performed, and the detailed spectra near the binding energy of Si 2p are shown in Fig. 4. In Fig. 4, a and b show XPS spectra for Ti–Si–C–N coatings with two different Si contents of 8.9 and 15.9 at.%, respectively. The binding energy of Si 2p level was calibrated with one of C 1s peak at 284.5eV. The XPS peaks could be separated into those corresponding to Si3N4 at 101.8 eV and SiC at both 100.4 eV [33] and 100.7 eV [34]. The Si incorporated into Ti–C– N coatings was found from XPS analysis to bond with both nitrogen and carbon, i.e., Si3N4 and SiC. With increasing of Si content in Ti–Si–C–N coatings, the Si3N4 peak prevailed
Fig. 4. High-resolution XPS spectra near the binding energy of Si 2p for Ti– Si–C–N coatings having various Si contents. (a) 8.9 at.%, (b) 15.5 at.%.
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by far (Fig. 4b). In addition, the peaks corresponding to free Si at 99.28 eV and TiSi2 at 98.8 eV were not detected although the Si content increased up to 15.5 at.%. From the results of Fig. 3 and Fig. 4, it was concluded that Si in Ti– Si–C–N coatings existed mainly as amorphous silicon nitride with some silicon carbide. From the instrumental analyses of XRD, XPS and HRTEM, it could be summarized that our quaternary Ti– Si(8.9 at.%)–C–N coatings had composite microstructure consisting of nano-sized Ti(C,N) crystallites surrounded by amorphous phases of Si3N4/SiC. 3.2. Mechanical evaluation of Ti–Si–C–N coatings Fig. 5 shows the micro-hardness value of the Ti–Si–C–N coatings as a function of the Si content. The hardness values were obtained from samples having thickness of about 3.0 Am. The micro-hardness values of Ti–C–N and Ti–Si–N coatings in our work were about 30 and 40 GPa, respectively. As the Si content in the coatings increased, the micro-hardness steeply increased and reached the maximum value of approximately 55 GPa at the Si content of 8.9 at.%. The micro-hardness, however, reduced again with a further increase in the Si content. The hardness value of Ti–Si–C–N coatings significantly increased up to ~55 GPa with a Si content of 8.9 at.% compared with that (~30 GPa) of Ti–C–N coatings. This surprisingly enhanced hardness is believed to originate from the microstructural evolution of Ti(C,N) coatings with Si addition (Fig. 3). A nanocomposite structure with very fine crystallites resulted in the enhanced hardness of Ti–Si–C–N coatings. In addition, grain boundary hardening derived from the increased cohesive energy at interphase boundaries along with the percolation phenomenon of amorphous phase is believed to play a role in enhancing the hardness. On the other hand, the hardness reduction with a further increase in Si content after maximum hardness in Fig. 5 has been
Fig. 5. The micro-hardness values of the Ti–Si–C–N coatings with various Si contents.
Fig. 6. The friction coefficients of the Ti(C,N), Ti–Si(8.9 at.%)–C–N and Ti–Si(15.5 at.%)–C–N coatings against steel ball.
similarly explained to Ti–Si–N coating system [14,16,17, 35,36]. This reduction was explained with a thickening phenomenon of amorphous Si3N4 phase between crystallites as Si content increased. When the amorphous Si3N4 and/or SiC phases become thicker beyond a certain critical value, the interaction between nanocrystallites and amorphous phases is significantly reduced and the hardness of nanocomposite becomes strongly dependent on the property of the amorphous phase. Fig. 6 shows the friction coefficients of the Ti(C,N), Ti–Si(8.9 at.%)–C–N, and Ti–Si(15.5 at.%)– C–N coatings against steel ball. The average friction coefficient of coatings largely decreased from 0.75 of Ti(C,N) to 0.5 with increasing Si content. This result would be caused by smoother surface with increasing amorphous phase and by tribo-chemical reaction, which often takes place in many ceramics, e.g. Si3N4/SiC can react with H2O to produce SiO2 or Si(OH)2 tribo-layer. These products of SiO2 and Si(OH)2 were known to play a role as a selflubricating layer [37]. Fig. 7 shows the surface morphologies of wear track and composition analyses for the wear
Fig. 7. The surface morphology of wear tracks and composition of wear debris after the sliding wear test. (a) Ti(C,N), (b) Ti–Si(8.9 at.%)–C–N.
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debris after the sliding wear test of 30,000 cycles. The surface morphology of the wear track for the Ti–C–N coatings was somewhat rough, and the width of wear track relatively narrowed as shown in Fig. 7a. Whereas, the surface morphology for the Ti–Si(8.9 at.%)–C–N coating was relatively smooth, and the width of the wear track was widened as shown in Fig. 7b. This result would be due to the adhesive wear behavior between hard coatings and relatively soft steel (~700 Hv0.2). Thus, steel ball would be more worn and smeared (from EDS results of wear debris) for the Ti–Si(8.9 at.%)–C–N coatings because of higher hardness (~55 GPa).
[3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]
4. Conclusions Quaternary Ti–Si–C–N coatings have been deposited on WC–Co substrates by a hybrid system combining the arc ion plating (AIP) and sputtering techniques. From XRD, XPS and HRTEM analyses, it could be suggested that quaternary Ti–Si–C–N coatings had a fine composite microstructure consisting of nano-sized crystallites of Ti(C,N) surrounded by amorphous phase of Si3N4/SiC mixture. The hardness value of Ti–C–N coatings was significantly increased from ~30 GPa of Ti–C–N coatings to ~55 GPa with Si addition of 8.9 at.%. In addition, the average friction coefficient of the Ti–Si–C–N coatings largely decreased with increasing Si content. This behavior would be caused both by the smoother surface due to the increased amount of amorphous phase and by the possible formation of self-lubricating tribo-layers such as SiO2 or Si(OH)2.
[15] [16] [17]
[18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]
Acknowledgements This work was performed through National Research Laboratory (NRL) project supported by Ministry of Science and Technology of Korea (MOST). Authors also thank Dr. Won of Korea Basic Science Institute for obtaining and discussing XPS results.
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Synthesis and characterization of quaternary Ti–Si–C–N coatings prepared by a hybrid deposition technique Jun-Ha Jeon, Sung Ryong Choi, Won Sub Chung, Kwang Ho Kim* School of Materials Science and Engineering, Pusan National University, San 30, Changjeon-dong, Keumjeong Busan 609-735, South Korea Available online 11 September 2004
Abstract Quaternary Ti–Si–C–N coatings have been deposited on WC–Co substrates by a hybrid system combining the arc ion plating (AIP) and sputtering techniques using Ti and Si targets in an Ar/N2/CH4 gaseous mixture. Ti–Si–C–N coatings with a Si content of 8.9 at.% had a fine composite microstructure comprising nano-sized crystallites of Ti(C,N) surrounded by amorphous phase of Si3N4/SiC mixture. The microhardness value of the Ti–Si–C–N coatings under a load of 25 g was ~55 GPa, i.e., much larger than ~30 GPa for Ti–C–N coatings. In addition, the average friction coefficient of the Ti–Si–C–N coatings largely decreased with increasing Si content. In this work, the microstructure and mechanical properties of Ti–Si–C–N coatings were systematically investigated using instrumental analyses. D 2004 Elsevier B.V. All rights reserved. Keywords: Ti–Si–C–N coatings; Arc ion plating; Nanocomposite; Superhardness; Wear behavior
1. Introduction Ti–C–N coatings have been widespread applied to various tools and dies because of high hardness (~30 GPa), good wear property, and improved corrosion resistance compared to TiN coatings [1–5]. The higher hardness value of Ti–C–N coatings than that of TiN coatings was attributed to the solid-solution hardening by carbon atoms. Ti–C–N coatings were studied mainly by chemical vapor deposition techniques [6–8]. On the other hand, arc ion plating techniques have been adopted to improve the mechanical properties of Ti–C–N coatings [9,10]. Recently, some groups [11–13] reported ternary Ti–Si–N coatings where some amount of amorphous silicon nitride was incorporated into TiN coatings. The Ti–Si–N coatings have microstructures characterized by a nanocomposite of fine TiN crystallites encapsulated with amorphous silicon nitride [14,15]. Such a microstructural modification of TiN by the Si addition caused a superhardness (z40GPa), and other improvements in wear resistant and oxidation resistant properties [16–18]. * Corresponding author. Tel.: +82 51 510 2391; fax: +82 51 510 3660. E-mail address: [email protected] (K.H. Kim). 0257-8972/$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2004.08.042
Quaternary Ti–Si–C–N system can have superior properties tailored through the microstructural designing from both Ti–C–N and Ti–Si–N coatings. Up to now, only a few papers have been reported on Ti–Si–C–N coating systems [19–21]. However, their properties as well as the microstructure were not thoroughly investigated. In this work, quaternary Ti–Si–C–N coatings have been deposited on WC–Co substrates using a hybrid system combining the arc ion plating (AIP) and DC magnetron sputtering techniques. The microstructure and mechanical properties of Ti–Si–C–N coatings were systematically investigated.
2. Experimental 2.1. Deposition of Ti–Si–C–N coatings Ti–Si–C–N coatings have been deposited on WC–Co substrates by the hybrid coating system, where the AIP method was combined with a magnetron sputtering technique. A schematic diagram of the hybrid deposition system was depicted in previous work [22–24]. An arc cathode gun for the Ti source and a DC sputter gun for the Si source were installed on each side of the chamber wall. A
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Table 1 Typical deposition conditions for Ti–Si–C–N films by hybrid coating system Base pressure Working pressure Working gas ratio Arc material Sputter material Arc current Sputter current Substrate temperature Substrate bias voltage Rotational velocity of substrate Typical deposition rate
9.3310 3 Pa 6.6710 2 Pa CH4/N2=20:40 Ti (99.99%) Si (99.99%) 60 A 0–2.2 A 300 8C 60 V 25 rpm ~3 Am/h
rotational substrate holder was located on a straight line between two sources. The WC–Co substrates, which had been machined to disc types of 20 mm in diameter and 3 mm in thickness, were cleaned in an ultrasonic bath cleaner using acetone and alcohol, and were cleaned again by ion bombardment using a bias voltage of 600 V under Ar atmosphere of 32 Pa for 12 min just before deposition. The typical deposition conditions for Ti–Si–C–N coatings are summarized in Table 1. 2.2. Characterization of Ti–Si–C–N coatings The thickness and surface morphology of coatings were measured using scanning electron microscopy (SEM, Hitach, S-4200) and a stylus (a-STEP) instrument. Compositional analyses of the coatings to determine the Ti, Si, C, and N contents were carried out by electron probe microanalyzer (EPMA, Shimadzu, EPMA1600). In order to identify the crystallographic structure of Ti–Si–C–N coatings, X-ray diffractometer using CuKa radiation was used. The crystallite size in the coating was determined from the direct observation by a field emission-transmission electron microscope (FE-TEM, JEOL, JEM-2010F) operating at 200 kV. The other structural information on the film components was obtained from the analyses by the selected area diffraction pattern (SADP) and high-resolution transmission electron microscopy (HRTEM). X-ray photoelectron spectroscopy (XPS, VG Scientifics, ESCALAB 250) was used to observe the chemical bonding status of Ti–Si–C–N coatings. The hardness of coatings was evaluated using a microhardness tester with Knoop indentor (Matsuzawa, MMT-7) under a load of 25 g. The average friction coefficient and wear behavior were studied through sliding tests using a conventional ball-on-disc wear apparatus. Steel balls (diameter 12 mm, 700 Hv0.2) were used as a counterpart material. The sliding tests were conducted with a sliding speed of 0.2 m/s under a load of 1 N at ambient temperature (approx. 19.8 8C) and relative humidity (RH) of 25–30%. The optical microscopy and scanning electron microscopy were employed to observe the morphology of wear tracks after each sliding test. The energy dispersive spectroscopy (EDS)
was used to reveal the composition of wear debris formed during wear test.
3. Results and discussion 3.1. Synthesis and characterization of Ti–Si–C–N coatings Fig. 1 shows the changes of Ti, Si, C, and N contents in at.% for Ti–Si–C–N coatings as a function of the d.c. sputter current collected by the Si target at a fixed Ti arc current of 60 A. The Si content in Ti–Si–C–N coatings linearly increased from 0 to 15.5 at.% with raising the d.c. sputter current from 0 to 2.2 A, whereas the nitrogen and carbon contents in Ti–Si–C–N coatings were nearly constant. Fig. 2 shows the X-ray diffraction patterns of Ti–Si–C–N coatings with various Si contents. The diffraction patterns of Ti–Si–C–N coatings showed multiple orientations of (111), (200), (220), (311), and (222) of Ti(C,N) crystal planes. Ti(C,N) has same crystal structure of f.c.c. as TiN, but interplanar d spacing becomes a little higher as carbon atoms replace nitrogen ones. These multiple orientations were in good agreement with other Ti–C–N coatings [6,8,25,26]. Other diffraction peaks from crystalline phases such as Si3N4, SiC, and titanium silicide phase like TiSi2 were not detected. This result postulated that Si existed as an amorphous phase of silicon nitride, silicon carbide, or free Si. As the Si was incorporated into the Ti(C,N), the diffraction peak intensities gradually reduced and started to disappear above the Si content of 15.5 at.%. In addition, peak broadening phenomenon was observed with increasing of Si content in Ti(C,N) coatings. Similar XRD peak broadenings have been reported for other ternary Ti–Si–N and quaternary Ti–Al–Si–N coating with Si addition [22,23,27–29], and was attributed mainly to the effect of the diminution of the crystallite size [14,18,30,31]. Fig. 3 shows the cross-sectional HRTEM
Fig. 1. Compositional changes of Ti–Si–C–N coatings as a function of DC sputter current to Si target at fixed Ti arc current of 60 A.
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Fig. 2. The X-ray diffraction patterns of Ti–Si–C–N coatings with various Si contents.
image, selected area electron diffraction (SAED) pattern, and dark-field TEM image for the Ti–Si–C–N coating with Si content of 8.9 at.%. From the results included in Fig. 3, it was established that the Ti–Si–C–N coatings were composed of Ti(C,N)-based crystallites and Si-based amorphous phase, which were clearly distinguished from each other by lattice fringe contrast (Fig. 3a). The Ti–Si– C–N coating with 8.9 at.% Si was regarded as a nanocomposite having fine crystallites of approximately 8 nm in size embedded in an amorphous matrix. The dark-
Fig. 3. Cross-sectional HRTEM image, electron diffraction pattern, and dark-field TEM image for Ti–Si–C–N coating with Si content of 8.9 at.%.
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field TEM image (Fig. 3b) of the Ti–Si–C–N coatings with a Si content of 8.9 at.% also shows again that fine Ti(C,N) crystallites (white area) were almost uniformly distributed in the matrix. The maximum hardness of the coatings could be obtained in this case as shown later (Fig. 5). The present microstructural evolution of Ti(C, N) coatings with Si addition is similarly reported to other systems such as Ti–Si–N [14,24,32] and Ti–Al–Si–N coatings [23]. In order to clarify bonding status of the amorphous phase comprising Ti–Si–C–N coatings, XPS analyses were performed, and the detailed spectra near the binding energy of Si 2p are shown in Fig. 4. In Fig. 4, a and b show XPS spectra for Ti–Si–C–N coatings with two different Si contents of 8.9 and 15.9 at.%, respectively. The binding energy of Si 2p level was calibrated with one of C 1s peak at 284.5eV. The XPS peaks could be separated into those corresponding to Si3N4 at 101.8 eV and SiC at both 100.4 eV [33] and 100.7 eV [34]. The Si incorporated into Ti–C– N coatings was found from XPS analysis to bond with both nitrogen and carbon, i.e., Si3N4 and SiC. With increasing of Si content in Ti–Si–C–N coatings, the Si3N4 peak prevailed
Fig. 4. High-resolution XPS spectra near the binding energy of Si 2p for Ti– Si–C–N coatings having various Si contents. (a) 8.9 at.%, (b) 15.5 at.%.
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by far (Fig. 4b). In addition, the peaks corresponding to free Si at 99.28 eV and TiSi2 at 98.8 eV were not detected although the Si content increased up to 15.5 at.%. From the results of Fig. 3 and Fig. 4, it was concluded that Si in Ti– Si–C–N coatings existed mainly as amorphous silicon nitride with some silicon carbide. From the instrumental analyses of XRD, XPS and HRTEM, it could be summarized that our quaternary Ti– Si(8.9 at.%)–C–N coatings had composite microstructure consisting of nano-sized Ti(C,N) crystallites surrounded by amorphous phases of Si3N4/SiC. 3.2. Mechanical evaluation of Ti–Si–C–N coatings Fig. 5 shows the micro-hardness value of the Ti–Si–C–N coatings as a function of the Si content. The hardness values were obtained from samples having thickness of about 3.0 Am. The micro-hardness values of Ti–C–N and Ti–Si–N coatings in our work were about 30 and 40 GPa, respectively. As the Si content in the coatings increased, the micro-hardness steeply increased and reached the maximum value of approximately 55 GPa at the Si content of 8.9 at.%. The micro-hardness, however, reduced again with a further increase in the Si content. The hardness value of Ti–Si–C–N coatings significantly increased up to ~55 GPa with a Si content of 8.9 at.% compared with that (~30 GPa) of Ti–C–N coatings. This surprisingly enhanced hardness is believed to originate from the microstructural evolution of Ti(C,N) coatings with Si addition (Fig. 3). A nanocomposite structure with very fine crystallites resulted in the enhanced hardness of Ti–Si–C–N coatings. In addition, grain boundary hardening derived from the increased cohesive energy at interphase boundaries along with the percolation phenomenon of amorphous phase is believed to play a role in enhancing the hardness. On the other hand, the hardness reduction with a further increase in Si content after maximum hardness in Fig. 5 has been
Fig. 5. The micro-hardness values of the Ti–Si–C–N coatings with various Si contents.
Fig. 6. The friction coefficients of the Ti(C,N), Ti–Si(8.9 at.%)–C–N and Ti–Si(15.5 at.%)–C–N coatings against steel ball.
similarly explained to Ti–Si–N coating system [14,16,17, 35,36]. This reduction was explained with a thickening phenomenon of amorphous Si3N4 phase between crystallites as Si content increased. When the amorphous Si3N4 and/or SiC phases become thicker beyond a certain critical value, the interaction between nanocrystallites and amorphous phases is significantly reduced and the hardness of nanocomposite becomes strongly dependent on the property of the amorphous phase. Fig. 6 shows the friction coefficients of the Ti(C,N), Ti–Si(8.9 at.%)–C–N, and Ti–Si(15.5 at.%)– C–N coatings against steel ball. The average friction coefficient of coatings largely decreased from 0.75 of Ti(C,N) to 0.5 with increasing Si content. This result would be caused by smoother surface with increasing amorphous phase and by tribo-chemical reaction, which often takes place in many ceramics, e.g. Si3N4/SiC can react with H2O to produce SiO2 or Si(OH)2 tribo-layer. These products of SiO2 and Si(OH)2 were known to play a role as a selflubricating layer [37]. Fig. 7 shows the surface morphologies of wear track and composition analyses for the wear
Fig. 7. The surface morphology of wear tracks and composition of wear debris after the sliding wear test. (a) Ti(C,N), (b) Ti–Si(8.9 at.%)–C–N.
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debris after the sliding wear test of 30,000 cycles. The surface morphology of the wear track for the Ti–C–N coatings was somewhat rough, and the width of wear track relatively narrowed as shown in Fig. 7a. Whereas, the surface morphology for the Ti–Si(8.9 at.%)–C–N coating was relatively smooth, and the width of the wear track was widened as shown in Fig. 7b. This result would be due to the adhesive wear behavior between hard coatings and relatively soft steel (~700 Hv0.2). Thus, steel ball would be more worn and smeared (from EDS results of wear debris) for the Ti–Si(8.9 at.%)–C–N coatings because of higher hardness (~55 GPa).
[3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]
4. Conclusions Quaternary Ti–Si–C–N coatings have been deposited on WC–Co substrates by a hybrid system combining the arc ion plating (AIP) and sputtering techniques. From XRD, XPS and HRTEM analyses, it could be suggested that quaternary Ti–Si–C–N coatings had a fine composite microstructure consisting of nano-sized crystallites of Ti(C,N) surrounded by amorphous phase of Si3N4/SiC mixture. The hardness value of Ti–C–N coatings was significantly increased from ~30 GPa of Ti–C–N coatings to ~55 GPa with Si addition of 8.9 at.%. In addition, the average friction coefficient of the Ti–Si–C–N coatings largely decreased with increasing Si content. This behavior would be caused both by the smoother surface due to the increased amount of amorphous phase and by the possible formation of self-lubricating tribo-layers such as SiO2 or Si(OH)2.
[15] [16] [17]
[18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]
Acknowledgements This work was performed through National Research Laboratory (NRL) project supported by Ministry of Science and Technology of Korea (MOST). Authors also thank Dr. Won of Korea Basic Science Institute for obtaining and discussing XPS results.
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