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Influence of silicon content on the microstructure, mechanical and tribological properties of magnetron sputtered Ti–Mo–Si–N films
Accepted Manuscript Influence of silicon content on the microstructure, mechanical and tribological properties of magnetron sputtered Ti-Mo-Si-N films Junhua Xu, Hongbo Ju, Lihua Yu PII:
S0042-207X(14)00278-4
DOI:
10.1016/j.vacuum.2014.08.010
Reference:
VAC 6406
To appear in:
Vacuum
Received Date: 8 July 2014 Revised Date:
12 August 2014
Accepted Date: 13 August 2014
Please cite this article as: Xu J, Ju H, Yu L, Influence of silicon content on the microstructure, mechanical and tribological properties of magnetron sputtered Ti-Mo-Si-N films, Vacuum (2014), doi: 10.1016/j.vacuum.2014.08.010. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Influence of silicon content on the microstructure, mechanical and tribological properties of magnetron sputtered Ti-Mo-Si-N films
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Junhua Xu*, Hongbo Ju, Lihua Yu School of Materials Science and Engineering, Jiangsu University of Science and
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Technology, Mengxi Road 2, Zhenjiang, Jiangsu Province, 212003, China
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Abstract: Ti-Mo-Si-N films with various Si content (0-17.2 at.%) were deposited by reactive magnetron sputtering and the effects of Si content on the microstructure, mechanical and tribological properties of Ti-Mo-Si-N films were investigated. The results showed that the face-centered cubic (fcc) interstitial solid solution of
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Ti-Mo-Si-N was formed by dissolution of Si into Ti-Mo-N lattice with the Si content in the range of 3.1-5.0 at.%. With a further increase in Si content, the films consisted of fcc-Ti-Mo-Si-N and amorphous Si3N4 phases. The hardness and fracture toughness
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of Ti-Mo-Si-N films first increased and then decreased with the increase of Si content and the highest values were 34.5 GPa and 2.6 MP.m1/2, respectively, at 5.0 at.% Si.
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The average friction coefficient and wear rate of Ti-Mo-Si-N films first decreased and then increased with the increase of Si content and the lowest values were 0.35 and 7.8×10-8 mm3/Nmm, respectively, at 5.0 at.% Si. Key words: reactive magnetron sputtering; Ti-Mo-Si-N films; mechanical and tribological properties *
Corresponding author, email: [email protected]
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ACCEPTED MANUSCRIPT 1. Introduction Since 1980s, transition metal nitride (TMN) films deposited by physical vapor deposition have aroused considerable interest because of their excellent properties [1].
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TiN film, which has been extensively researched, is widely used in industry of cutting tools and human replacement organs. Although the application of TiN film is very successful, its further demands like decreased average friction coefficient in cutting
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processes are driving force for further development of TiN film [2].
TiN-based films containing the element which is able to form certain self-lubricity
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tribo-oxidation were widely studied in recent years [2, 3]. For instance, Ti-Mo-N films were reported to show a low average friction coefficient because Mo can react with the oxygen/moisture in the air and form Magnéli phase (MoO3) during wear test [2, 4]. However, the improvement of the friction property is at the expense of wear
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property because MoO3 on the wear track can be worn away easily by the counterpart [3]. High wear rate means shorter life cycle, so it is of crucial importance to improve the wear resistance of Ti-Mo-N films. At present, attention introducing one more
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element into ternary films was paid to enhance the hardness and tribological properties of the films. According to the Ref. [5, 6], the incorporation of Si into the
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TMN films generally leads to the formation of nanocomposite TM-Si-N films, in which TMN nanocrystallites are surrounded by an amorphous Si3N4. In recent years, the study of Ti-Si-N films, in which Si atoms dissolve into interstitial sites of TiN lattice or formed TiN/Si3N4 nanocomposite films, shows excellent mechanical properties, such as high hardness [7, 8, 9]. It is believable that Ti-Mo-Si-N films can exhibit good mechanical properties and excellent wear resistance as is the case for the similar films, such as Ti-Al-Si-N [7] and Mo-Al-Si-N [10]. However, there is little 2
ACCEPTED MANUSCRIPT information available in the literature about the microstructure, mechanical and tribological properties of Ti-Mo-Si-N films. In addition, the fracture toughness of nitride films is one of the important mechanical properties and has been aroused considerable researchers’ interest. The radial cracking indentation method is usually
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used to evaluate the toughness of nitride films quantitatively [11]. As for this method, the indentation depth should be less than 10 % of the film thickness. Even for very brittle nitride films, the indentation depth is still a few hundred nanometers in order to
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induce a radial crack [12]. Therefore, it is difficult to introduce well-developed radial cracks under the condition that the depth limitation of nanoindentation excludes the
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substrate effect. We calculated the fracture toughness using a scratch tester quantitatively.
In our paper, by analogy with Ti-Mo-N and Ti-Si-N films, Mo and Si atoms were incorporated into TiN film simultaneously using reactive magnetron sputtering, and
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the effects of Si content on the microstructure, compressive residual stress, hardness, fracture toughness and tribological properties of Ti-Mo-Si-N films were investigated.
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2. Experiment details
Ti-Mo-Si-N films with a thickness of about 2 µm were deposited on both AISI 304
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stainless steel and Si (100) wafer substrates using a multi-target magnetron sputtering system. Ti (99.9 %), Mo (99.95 %) and Si (99.9 %) targets with a diameter of 75 mm were sputtered by three radio frequency powers. Mirror polished stainless steels and Si (100) wafer substrates were ultrasonically cleaned in acetone and alcohol, and then mounted on the substrate holder in the vacuum chamber. Substrate holder was electrically grounded with no applied bias voltage. After base pressure reached 6×10-4 Pa, Ar (99.999 %) and N2 (99.999 %) were introduced into the chamber by means of 3
ACCEPTED MANUSCRIPT two separate gas manifolds. The Ti-Mo-Si-N films with different Si content were achieved by fixing the powers of Ti target at 250 W and Mo target at 150 W and adjusting the power of Si target from 0 W to 130 W while constantly keeping the working pressure at 0.3 Pa and the nitrogen to argon ratio (flow ratio) at 10:3. Besides
achieved using temperature transducer after deposition.
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this, the substrate was not heated during deposition and its temperature value was
The infrared analysis was carried out in transmission mode at normal incidence by
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GIGILAB FT5200 Fourier transform infrared spectrometer (FTIR). The elemental
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composition and chemical bonds of the film were characterized by X-ray photoelectron spectroscopy (XPS) with Al Kα irradiation at a pass energy of 160 eV after removing the surface contaminants on the films by sputtering with Ar+ ion beam at a primary energy of 3 keV for 3 min. The spectra were calibrated by the C 1s line with a binding energy of 285.0 eV, and then corrected for the linear emission
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background and decomposed into peaks with Gaussian-Lorentzian line shapes by a non-linear least-square fitting method. The elemental compositions of the film were
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calculated from the areas under their XPS peaks by considering the relative sensitivity factors based on Wagner data analysis system. The microstructure of the films was
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characterized by glancing incident X-ray diffraction (XRD) at 2-degree incidence angle with a Siemens X-ray diffractometer using Cu Kα radiation. The compressive residual stress (σ) of the films was calculated by Stoney’s equation [3]:
σ=
E ts 2 1 1 − 1 − υ 6t f R Rs
(1)
where E is Elastic modulus of the substrate (E=170 GPa), υ is Poisson’s ratio of the substrate ( υ =0.3), ts is the thickness of the substrate, tf is the thickness of the film, R is the substrate curvature radii of the Si wafer and Rs is the curvature radii of the film 4
ACCEPTED MANUSCRIPT on the Si wafer. R and Rs were measured by Bruker DEKTAK-XT profilometer. Hardness and elastic modulus of the films were determined with nanoindenter CPX+NHT2+MST, which was equipped with a diamond Berkovich indenter tip (3-side pyramid). An automatic indentation mode was programmed to place
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indentations in a 3×3 array. The maximum load of 3 mN was used to meet the criterion of d/h<0.1, which guarantees a minimal influence from the substrate on the hardness measure, where d and h are the indenter penetration depth into the film and
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the film thickness, respectively. Before making indents on the films, the indenter was calibrated with respect to a reference sample of fused silica. The fracture toughness
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(KIC) of the films was measured by a scratch tester, which was equipped with a diamond Rockwell tip. The fracture toughness was calculated by the following equation [13]: 2 pf g a 1/2 −1 R ( ) sin R2 cotθ π a
(2)
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K IC =
where p is the pressure required to open the crack, R is the radius of the indenter cone into the groove, 2a is the total crack length, fg is the coefficient of grooving friction
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which is obtained from scratch test, θ (θ=65.03 °) is a angle between the axis of the
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diamond pyramid and the three faces. The 30 min wear test was carried out along a circular track of 8mm diameter against a 9 mm diameter Al2O3 counterpart at 50 rpm under a constant normal load of 3 N in the atmosphere (the relative humidity of about 25-30 %) at room temperature (about 25 °C) using a UTM-2 CETR tribometer. After the wear tests, the wear tracks were examined using a profilometer (Bruker DEKTAK-XT) to measure the wear loss of the films (V). The wear rate of the films (W) was calculated by Archard’s classical wear equation [14]:
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ACCEPTED MANUSCRIPT W=
V S×L
(3)
where S is the total sliding distance and L is the applied load. For FTIR, XPS, XRD and nanoindentation measurements, the Ti-Mo-Si-N samples
3. Results and discussion 3.1 Elemental compositions and microstructure
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were deposited on AISI 304 stainless steel substrates.
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were deposited on Si (100) wafer substrates; for scratch and wear tests, the samples
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Fig. 1 shows the elemental compositions of the Ti-Mo-Si-N films as a function of the power of Si target. As the power of Si target increases from 0 W to 130 W, the Si content in the Ti-Mo-Si-N films increases almost linearly from 0 at.% to 17.2 at.%. Although FTIR is not usually employed to analyze the structure of nitride films, it
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can give precious information about the chemical bonds involved in the amorphous phase as for example in the Zr-Si-N [15] and Ti-B-N [16] films. Fig. 2 illustrates the
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FTIR spectra of Ti-Mo-Si-N films with various Si content. In order to clearly identify the chemical bonding involved in Ti-Mo-Si-N films, Si3N4 film deposited under the
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same experimental conditions was also characterized by FTIR. As shown in Fig. 2, for Si3N4 film, four absorption peaks are detected at approximately 480cm-1,700cm-1, 1500cm-1 and 2400cm-1. The absorption peaks at about 480cm-1 and 700cm-1 are attributed to the symmetric stretching of Si-N bonds in amorphous Si3N4 [17], and other absorption peaks at about 1500cm-1 and 2400cm-1 are similar to the results of amorphous Si3N4 characterized by FTIR in Ref. [15]. No absorption peak is detected for Ti-Mo-Si-N films with a Si content in the range between 3.1 at.% and 5.0 at.%. However, as the Si content increases from 8.8 at.% to 17.2 at.%, four absorption 6
ACCEPTED MANUSCRIPT peaks which is similar with the peaks of Si3N4 film are detected by FTIR. Besides, the intensity of those peaks increases with increasing Si content in the films, indicating that the amount of Si3N4 increases with increasing Si content in the films. In order to clarify further bonding status of the amorphous phase in the Ti-Mo-Si-N
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films, Si 2p spectra with various Si content was carried out by XPS and the results are shown in Fig. 3. The Si 2p spectra reveal that no obvious peak is detected for Ti-Mo-N and Ti-Mo-Si-N (5.0 at.% Si) films. The peak corresponding to 101.8 eV,
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which is in good agreement with the silicon fourfold coordinated to nitrogen as Si3N4
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[18, 19], is detected for Ti-Mo-Si-N film at 8.8 at.% Si. This is consistent with the results from Fig. 2.
Fig. 4 shows the XRD pattern (a) and lattice constant (b) of the Ti-Mo-Si-N films with various Si content. As shown in Fig. 4(a), the diffraction pattern of Ti-Mo-N film
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presents multiple orientations of (111), (200), (220), (311) and (222) crystal planes of face-centered cubic (fcc) TiN phase (PDF card 65-0715) where some Ti atoms in TiN lattice are replaced by Mo atoms [2, 4]. For each XRD pattern of Ti-Mo-Si-N films, it
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shows a single of fcc-TiN structure. Since no X-ray diffraction peaks corresponding to crystalline phases such as Si3N4, MoSi2 or TiSi2 are observed, it can be concluded
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that the Si3N4 phase in Ti-Mo-Si-N films with a Si content in the range between 8.8 at.% and 17.2 at.% detected by FTIR and XPS is amorphous. As shown in Fig. 4(b), the lattice constant of the films increases with increasing Si content, this suggests that the Si atoms dissolve into the interstitial sites of Ti-Mo-N lattice. It is well known that the phase formation is affected by the thermodynamic and kinetic issues. Compared to the amorphous, the crystalline phase is preferred to form because of its higher potential energy [20]. The formation of TiN is easier than that of 7
ACCEPTED MANUSCRIPT amorphous Si3N4 because the difference of electronegativity (Ti=1.5, Si=2.5 and N=3.0) between Ti and N is larger than that between Si and N. Many Ti-Si-N (nc-TiN/a-Si3N4) films with a low Si content were deposited at the substrate temperature of >500 °C [21]. However, solid solution of Ti-Si-N is easy to form at
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low temperature of <100 °C because TiN crystals are able to kinetically hinder the formation of the nc-TiN/a-Si3N4 at this temperature [20]. The Ti-Mo-Si-N films were deposited at uncontrolled temperature in our experiment. So the TiN is preferentially
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formed compared to the Si3N4 in our case. For low Si content in the Ti-Mo-Si-N films (<5.0 at.%), Si atoms dissolved into interstitial sites in Ti-Mo-N lattice and formed
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interstitial solid solutions because no amorphous Si3N4 is detected in Fig. 2 and 3. With a further increase in Si content, the dissolution of Si is too weak to consume all Si atoms deposited on the films. So the amorphous Si3N4 is detected in Fig. 2 and 3 and the films are consisted of interstitial solid solutions and amorphous Si3N4 phases.
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To sum up, interstitial solid solution of Ti-Mo-Si-N is formed by dissolution of Si atoms into fcc-Ti-Mo-N lattice when the Si content is below 5.0 at.%. With a further
phases.
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increase in Si content, the films consist of fcc-Ti-Mo-Si-N and amorphous Si3N4
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3.2 Mechanical properties
The residual stress of Ti-Mo-Si-N films with various Si content is presented in Fig.
5. As shown in Fig. 5, all films regardless of Si content are in compressive stress state. The compressive residual stress of Ti-Mo-N film is about -1.5 GPa. The compressive residual stress of Ti-Mo-Si-N films first increases gradually and then decreases with increasing Si content, after reaching the maximum value of -2.4 GPa, at 8.8 at.% Si.
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ACCEPTED MANUSCRIPT The measured compressive residual stress here is a combination of both thermal stress and intrinsic stress [3]. The thermal stress is determined by the temperature of substrate and the difference of thermal expansion coefficient between the film and substrate. The thermal stress (σT) can be calculated by the following equation [22]: EF (α F − α S )(TD − TM ) 1 − nF ………………….(3)
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σT =
where nF is the Poisson ratio of the film (nF=0.3); αF and αS are the thermal expansion
°
C). In our experiment, the
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the temperature during stress measurement (25
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coefficient of the film and the substrate; TD is the deposition temperature; and TM is
temperature of substrate increases from 89 °C at 0 at.% Si to 162 °C at 17.2 at.% Si because the number of ions bombardment on the substrate increases with increasing the power of Si target. The thermal expansion coefficient of Si wafer, TiN and amorphous Si3N4 is 2.7×10-6/°C [3], 6.0×10-6/°C [23] and 3.2×10-6/°C [24],
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respectively. Thus, the thermal stress of all films is in tensile stress state and its value is in the range of 0.1-0.2 GPa. The compressive residual stress of Ti-Mo-Si-N films is
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mainly attributed to the intrinsic stress. The lattice expansion induced by incorporation of Si atoms into the interstitial sites in Ti-Mo-N crystal lattice (as
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shown in Fig. 2b) leads to the increase of intrinsic stress in the films. Besides this, the compressive residual stress can be released by the amorphous Si3N4 due to its low density and disordered network [10]. On the basis of the above analysis, the lattice expansion induced by incorporation
of Si atoms into the interstitial sites in Ti-Mo-N crystal lattice leads to the increase of compressive residual stress of the films when the Si content is below 5.0 at.%; although the amorphous Si3N4 appears in the film at 8.8 at. % Si, the compressive residual stress of the film further increases because the amount of amorphous Si3N4 is 9
ACCEPTED MANUSCRIPT too few to weaken the effect of Si dissolution; with a further increase in Si content, the decrease of compressive residual stress is mainly attributed to the increase of the amount of amorphous Si3N4. Fig. 6 shows the load-displacement curves of Ti-Mo-Si-N films with various Si
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content obtained from the nanoindentation measurement. The maximum load of 3 mN was used to meet the criterion of d/h<0.1, which guarantees a minimal influence from the substrate on the hardness measurement, where d and h are the indenter penetration
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depth into the film and the film thickness, respectively. The hardness (H), elastic
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modulus (E), H/E and H3/E*2 ratios are shown in Fig. 7. As shown in Fig. 7(a), the hardness and elastic modulus of Ti-Mo-N film are 27.4 GPa and 401 GPa, respectively. The hardness and elastic modulus of Ti-Mo-Si-N films first increase gradually and then decrease with increasing Si content, after reaching the maximum
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values of 34.5 GPa and 412 GPa, respectively, at 5.0 at. % Si.
The initial increase in hardness is mainly attributed to the effect of solid solution strengthening and the increase of compressive residual stress. As the Si content
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increases from 8.8 at.% to 17.2 at.%, amount of amorphous Si3N4 in the films could be result in the decrease of the hardness because the hardness of Si3N4 film deposited
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under the same experimental conditions is about 21.3 GPa. Besides this, decreasing compressive residual stress also leads to the decrease of the hardness of the films when the Si content is above 8.8 at.%. The elastic modulus is influenced by the compressive residual stress [25]. Higher compressive residual stress usually leads to higher elastic modulus [3, 25, 26]. The elastic modulus of the films increases from 401 GPa at 0 at.% Si to 412 GPa at 5.0 at.% Si because of the increase of the compressive residual stress. The elastic 10
ACCEPTED MANUSCRIPT modulus of Si3N4 film that we deposited under the same experimental conditions is 190 GPa. Thus, the decrease of elastic modulus is mainly attributed to the appearance of the amorphous Si3N4 with a further increase in Si content in Ti-Mo-Si-N films. An interesting attention is being given to the relationship between the hardness and
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elastic modulus during the development of the hard films for mechanical applications recently. The ratio of hardness to elastic modulus (H/E) is regarded as an important factor to describe the resistance of material against elastic strain to failure [3]. The
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H3/E*2 ratio (E* is the effective elastic modulus and E*=E/(1-µ2) where µ is Poisson’s ratio) relates to the resistance to plastic deformation and the tribological properties of
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the films [15]. As shown in Fig. 7(b), the ratios of H/E and H3/E*2 of Ti-Mo-N film are 0.069 and 0.137 GPa, respectively. The H/E and H3/E*2 ratios of Ti-Mo-Si-N films, all of which are higher than that of Ti-Mo-N film, firstly increase and then decrease with increasing Si content in the films, after reaching the maximum values of
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0.085 and 0.204 GPa, respectively, at 5.0 at.% Si. The elastic recovery is the ratio of the recovered displacement after unloading to
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the total indentation displacement and can be calculated from load-displacement curves (Fig. 6). The elastic recovery of Ti-Mo-N film is 57 %. The elastic recovery of
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Ti-Mo-Si-N films, which is in the range between 46 % and 43 %, is lower than that of Ti-Mo-N film indicating higher fracture toughness in a qualitative sense [10]. The fracture toughness (KIC) of Ti-Mo-Si-N films with various Si content was calculated using Eq. (2) and the results are shown in Fig. 8. As shown in Fig. 8, the fracture toughness of Ti-Mo-N film is about 1.3 MPa.m1/2. The fracture toughness of Ti-Mo-Si-N films initially increases and then decreases with the increase of Si content, after reaching the maximum value of about 2.7 MPa.m1/2, at 5.0 at.% Si.
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ACCEPTED MANUSCRIPT Fig. 9 illustrates the fracture toughness of Ti-Mo-Si-N films as a function of compressive residual stress (a) and H3/E*2ratio (b). It indicates that the fracture toughness increases with the increase of compressive residual stress because the compressive residual stress allows films to sustain more tensile strain before fracture
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[12]. Since H3/E*2 ratio is connected with the resistance to plastic deformation, the fracture toughness of Ti-Mo-Si-N films increases with increasing H3/E*2 ratio.
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3.3 Tribological properties
Fig. 10(a) illustrates a few examples of friction coefficient curves and wear track
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profiles against Al2O3 counterpart at RT after ball-on-disc wear test. Each curve records the friction coefficient variation from running-in to the long-term steady state friction. The wear track of Ti-Mo-Si-N films becomes shallower and narrower as the Si content increases from 3.1 at.% to 5.0 at.%. With a further increase in Si content,
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the wear track becomes deeper and wider. In all cases, the depth of the wear tracks is lower than the thickness of the films, indicating all the films in our paper were not worn out.
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The average friction coefficient and wear rate of Ti-Mo-Si-N films with various Si content are presented in Fig. 10(b). As shown in Fig. 10(b), the average friction
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coefficient and wear rate of Ti-Mo-N film are 0.38, 7.1×10-7 mm3/N.mm, respectively. As the Si content increases, the average friction coefficient and wear rate firstly decrease and then increase, after reaching the minimum values of 0.35, 7.8×10-8 mm3/N.mm, respectively, at 5.0 at.% Si. Ti-Mo-Si-N films with Si content in the range between 3.1 at. % and 5.0 at.% present lower average friction coefficients and wear rates compared with Ti-Mo-N film.
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ACCEPTED MANUSCRIPT In order to study the effects of Si content on the tribological properties of Ti-Mo-Si-N films, the 3D and optical microscope images of wear track on Ti-Mo-Si-N films are presented in Fig. 11. As shown in Fig. 11(a), the wear track of Ti-Mo-Si-N film at 3.1 at.% Si shows a relatively narrow and smooth. There are many
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fine scratches along the counterpart sliding direction. It reveals that the main wear mechanism is abrasive wear. By comparison, as the Si content increases to 5.0 at.%, improved wear resistance is observed and the wear track shows much narrow and
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smooth. It is attributed to the dissolution of Si in the Ti-Mo-N lattice which could effectively inhibit the cracks to propagate and reduce the contact stress between
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sliding surface and counterpart [6]. With a further increase in Si content, the film at 17.2 at.% Si exhibits a relatively poor wear resistance and a lot of obvious scratches and wear debris are detected. The increase of average friction coefficient and wear rate is related to the formation of the amorphous Si3N4 in the films.
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In dry ball-on-disc wear test, the tribo-film plays an important role in determining the tribological properties [2, 3]. In addition, tribo-oxidation was reported to be a
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general phenomenon in the wear test at room temperature. It was reported that MoO3, SiO2 and Si(OH)4 could be formed easily by a reaction between the wear track and
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oxygen/moisture in the air for transition metal nitride films, e.g. Ti-Mo-N [2, 4] and Mo-Al-N films [3]. Although MoO3 can reduce the average friction coefficient, it can be worn away easily by the counterpart during the wear test because MoO3 has low shear strength [3, 4]. SiO2 and Si(OH)4 were known to play a role as a self-lubricating and protection layer on the wear track of the films and result in lower wear rate compared with MoO3. When the Si content is in the range between 3.1 at.% and 5.0 at.%, the addition of Si in the films results in the decrease of MoO3, so the wear tracks become shallower with increasing Si content. 13
ACCEPTED MANUSCRIPT The tribological properties of transition metal nitride films are also influenced by their mechanical properties [3, 4]. Hardness is one of the mechanical properties relating to the tribological properties. Higher hardness leads to an improved load
According to the Archard’s equation [27]: V W =K L H
(4)
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carrying capacity and a smaller real contact area with the counterpart [3, 4].
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where V is the wear volume, L is the sliding distance (38 m), K is Archards wear
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coefficient, W is the normal load (3 N), H is the hardness of the film. So the wear rate of Ti-Mo-Si-N films is reverse ratio to the hardness. So the wear rate shows the nearly opposite variation trend to that of hardness with the increase of Si content. Besides this, it is considered that the parameters of H/E and H3/E*2 are more important than hardness alone in valuing the tribological properties of the films [3].The high values
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of H/E and H3/E*2 confirm an improved tribological properties. To sum up, as the Si content increases from 3.1 at.% to 5.0 at.%, the decrease of
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average friction coefficient and wear rate could be attributed to the increase of Si content in the films, the hardness and the ratios of H/E and H3/E*2. With a further
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increase in Si content, the increase of average friction coefficient and wear rate could be mainly attributed to the amorphous Si3N4. The decrease of the hardness, the ratios of H/E and H3/E*2 also attribute the increase of the average friction coefficient and wear rate.
4. Conclusions Ti-Mo-Si-N films with various Si content were deposited by reactive magnetron sputtering. The main results could be concluded as follows: 14
ACCEPTED MANUSCRIPT (1)
When the Si content was below 5.0 at.%, Si dissolved into interstitial sites in
fcc-Ti-Mo-N lattice and formed interstitial solid solutions. With a further increase in Si content, the films were consisted of fcc-Ti-Mo-Si-N and amorphous Si3N4. (2)
The hardness and fracture toughness of the films first increased and then
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decreased with increasing Si content, after reaching the maximum values of 34.5 GPa and 2.6 MPa.m1/2, respectively, at 5.0 at.% Si. The hardness enhancement was mainly attributed to the solid solution strengthening and the increase of compressive residual
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stress. The formation of amorphous Si3N4 could be result in the decrease of the hardness. The fracture toughness was influenced by the compressive residual stress
(3)
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and the H3/E*2 ratio.
The average friction coefficient and wear rate first decreased and then
increased with increasing Si content, after reaching the minimum values of 0.35 and
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7.8×10-8 mm3/Nmm, respectively, at 5.0 at.% Si. The average friction coefficient and wear rate were influenced by the hardness, amorphous Si3N4 and the ratios of H/E and H3/E*2.
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Acknowledgment
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This project was supported by National Natural Science Foundation of China (51074080, 51374115) and Research Innovation Program for College Graduates of Jiangsu Province (CXZZ12-0717).
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[10] J. F. Yang, B. Prakash, Y. Jiang, X. P. Wang, Q. F. Fang: ‘Effect of Si content on the microstructure and mechanical properties of Mo–Al–Si–N coatings’, Vacuum, 2012, 86, 2010-2013.
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[11] S. Zhang, X. Zhang: ‘Toughness evaluation of hard coatings and thin films’, Thin
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[12] J. Musil: ‘Hard nanocomposite coatings: thermal stability, oxidation resistance and toughness’, Surface and Coatings Technology, 2012, 207, 50-65. [13] J. W. Hoehn, S. K. Venkataraman, H. Huang, W. W. Gerberich:
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‘Micromechanical toughness test applied to NiAl’, Materials Science and Engineering: A, 1995, 192-193, 301-308.
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[14] J. Shi, C. M. Muders, A. Kumar, X. Jiang, Z. L. Pei, J. Gong, C. Sun: ’Study on nanocomposite Ti-Al-Si-Cu-N films with various Si contents deposited by cathodic
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vacuum arc plating’, Applied Surface Science, 2012, 258, 9642-9649. [15] D. Pilloud, J. F. Pierson, A. P. Marques, A. Cavaeiro: ‘ Structural changes in Zr-Si-N films vs. their silicon content’, Surface and Coatings Technology, 2004, 180-180, 352-356. [16] J. F. Pierson, E. Tomasella, Ph. Bauer: ‘Reactively sputtered Ti-B-N nanocomposite films: correlation between structure and optical properties’, Thin Solid Films, 2001, 384, 46. 17
ACCEPTED MANUSCRIPT [17] S. Peter, S. Bernutz, S. Berg, F. Richiter: ‘FTIR analysis of a-SiCN: H films deposited by PECVD’, Vacuum, 2013, 98, 81-87. [18] J. L. Bischoff, F. Lutz, D. Bolmont, L. Kubler: ‘Use of multilayer techniques for XPS identification of various nitrogen environments in the Si/NH3 system’, Surface
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[19] Q. Liu, Q. F. Fang, F. J. Liang, J. X. Wang, J. F. Yang, C. Li: ‘Synthesis and properties of nanocomposite MoSiN hard films’, Surface and Coatings Technology,
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[20] C. K. Chung, H. C. Chang, S. C. Chang, M. W. Liao, C. C. Lai: ‘Evolution of enhanced crystallinity and mechanical property of nanocomposite Ti–Si–N thin films using magnetron reactive co-sputtering’, Journal of Alloys and Compounds, 2012, 537, 318-322.
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[21] R.F. Zhang, S. Veprek, On the spinodal nature of the phase segregation and formation of stable nanostructure in the Ti-Si-N system, Materials Science and
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Engineering A, 424, 2006, 128-137.
[22] M. Bielawski: ‘Residual stress control in TiN/Si coatings deposited by
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unbalanced magnetron sputtering’, Surface and Coatings Technology, 2006, 200, 3987-3995.
[23] F. A. Almeida, M. Belmonte, A. J. S. Fernandes, F. J. Oliveira, R. F. Silva: ‘MPCVD diamond coating of Si3N4-TiN electroconductive composite substrates’, Diamond and Related Materials, 2007, 16, 987-982. [24] C. Tien, T. Lin: ‘Thermal expansion coefficient and thermomechanical properties of SiNx thin films prepared by plasma-enhanced chemical vapor deposition’, Applied Optics, 2012, 51, 7229-7235. 18
ACCEPTED MANUSCRIPT [25] J. Malzbender: ‘The use of theories to determine mechanical and thermal stresses in monolithic, coated and multilayered materials with stress-dependent elastic modulus or gradient in elastic modulus exemplified for thermal barrier coatings’, Surface and Coatings Technology, 2004, 186, 416-422.
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[26] S. S. Manson, National advisory committee for aeronautics—behavior of materials under conditions of thermal stress, NACA, Washington,1953.
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[27] J. Yang, Y. Jiang, Z. Yuan, X. Wang, Q. Fang: ‘Effect of carbon content on the microstructure and properties of W-Si-C-N coatings fabricated by magnetron
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sputtering’, Materials Science and Engineering B, 2012, 177, 1120-1125.
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ACCEPTED MANUSCRIPT Figure captions Fig. 1 Elemental compositions of Ti-Mo-Si-N films as a function of the power of Si target
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Fig. 2 The FTIR spectra of Ti-Mo-Si-N films with various Si content Fig. 3 XPS spectra of Si 2p of Ti-Mo-Si-N films with various Si content
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Fig. 4 XRD patterns (a) and lattice constant (b) of the Ti-Mo-Si-N films
Fig. 5 Compressive residual stress of Ti-Mo-Si-N films with various Si content
at.%, (b) 5.0 at.%, (c) 17.2 at.%
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Fig. 6 Loading-displacement curve of Ti-Mo-Si-N films with various Si content: (a) 0
Fig. 7 Hardness (H), elastic modulus (E), H/E and H3/E*2 ratios of Ti-Mo-Si-N films
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with various Si content
Fig. 8 The toughness of Ti-Mo-Si-N films with various Si content Fig. 9 The fracture toughness of Ti-Mo-Si-N films as a function of compressive
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residual stress (a) and H3/E*2 (b)
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Fig. 10 The friction coefficient curves, wear track profiles (a), the average friction coefficients and wear rates (b) of Ti-Mo-Si-N films with various Si content Fig. 11 3D and optical microscope images of wear tracks on Ti-Mo-Si-N films with various Si content: (a) 3.1 at.%, (b) 5.0 at.%, (c) 17.2 at.%
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ACCEPTED MANUSCRIPT Highlights 1. Ti-Mo-Si-N films were deposited by reactive magnetron sputtering. 2. Hardness of Ti-Mo-Si-N film (5.0 at.% Si) reached a maximum value of 34.5 GPa.
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3. Addition of Si (<5.0 at.% Si) led to the increase of fracture toughness.
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4. The films (3.1-5.0 at.% Si) were found to be optimized for wear resistance tools.
S0042-207X(14)00278-4
DOI:
10.1016/j.vacuum.2014.08.010
Reference:
VAC 6406
To appear in:
Vacuum
Received Date: 8 July 2014 Revised Date:
12 August 2014
Accepted Date: 13 August 2014
Please cite this article as: Xu J, Ju H, Yu L, Influence of silicon content on the microstructure, mechanical and tribological properties of magnetron sputtered Ti-Mo-Si-N films, Vacuum (2014), doi: 10.1016/j.vacuum.2014.08.010. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Influence of silicon content on the microstructure, mechanical and tribological properties of magnetron sputtered Ti-Mo-Si-N films
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Junhua Xu*, Hongbo Ju, Lihua Yu School of Materials Science and Engineering, Jiangsu University of Science and
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Technology, Mengxi Road 2, Zhenjiang, Jiangsu Province, 212003, China
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Abstract: Ti-Mo-Si-N films with various Si content (0-17.2 at.%) were deposited by reactive magnetron sputtering and the effects of Si content on the microstructure, mechanical and tribological properties of Ti-Mo-Si-N films were investigated. The results showed that the face-centered cubic (fcc) interstitial solid solution of
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Ti-Mo-Si-N was formed by dissolution of Si into Ti-Mo-N lattice with the Si content in the range of 3.1-5.0 at.%. With a further increase in Si content, the films consisted of fcc-Ti-Mo-Si-N and amorphous Si3N4 phases. The hardness and fracture toughness
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of Ti-Mo-Si-N films first increased and then decreased with the increase of Si content and the highest values were 34.5 GPa and 2.6 MP.m1/2, respectively, at 5.0 at.% Si.
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The average friction coefficient and wear rate of Ti-Mo-Si-N films first decreased and then increased with the increase of Si content and the lowest values were 0.35 and 7.8×10-8 mm3/Nmm, respectively, at 5.0 at.% Si. Key words: reactive magnetron sputtering; Ti-Mo-Si-N films; mechanical and tribological properties *
Corresponding author, email: [email protected]
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ACCEPTED MANUSCRIPT 1. Introduction Since 1980s, transition metal nitride (TMN) films deposited by physical vapor deposition have aroused considerable interest because of their excellent properties [1].
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TiN film, which has been extensively researched, is widely used in industry of cutting tools and human replacement organs. Although the application of TiN film is very successful, its further demands like decreased average friction coefficient in cutting
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processes are driving force for further development of TiN film [2].
TiN-based films containing the element which is able to form certain self-lubricity
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tribo-oxidation were widely studied in recent years [2, 3]. For instance, Ti-Mo-N films were reported to show a low average friction coefficient because Mo can react with the oxygen/moisture in the air and form Magnéli phase (MoO3) during wear test [2, 4]. However, the improvement of the friction property is at the expense of wear
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property because MoO3 on the wear track can be worn away easily by the counterpart [3]. High wear rate means shorter life cycle, so it is of crucial importance to improve the wear resistance of Ti-Mo-N films. At present, attention introducing one more
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element into ternary films was paid to enhance the hardness and tribological properties of the films. According to the Ref. [5, 6], the incorporation of Si into the
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TMN films generally leads to the formation of nanocomposite TM-Si-N films, in which TMN nanocrystallites are surrounded by an amorphous Si3N4. In recent years, the study of Ti-Si-N films, in which Si atoms dissolve into interstitial sites of TiN lattice or formed TiN/Si3N4 nanocomposite films, shows excellent mechanical properties, such as high hardness [7, 8, 9]. It is believable that Ti-Mo-Si-N films can exhibit good mechanical properties and excellent wear resistance as is the case for the similar films, such as Ti-Al-Si-N [7] and Mo-Al-Si-N [10]. However, there is little 2
ACCEPTED MANUSCRIPT information available in the literature about the microstructure, mechanical and tribological properties of Ti-Mo-Si-N films. In addition, the fracture toughness of nitride films is one of the important mechanical properties and has been aroused considerable researchers’ interest. The radial cracking indentation method is usually
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used to evaluate the toughness of nitride films quantitatively [11]. As for this method, the indentation depth should be less than 10 % of the film thickness. Even for very brittle nitride films, the indentation depth is still a few hundred nanometers in order to
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induce a radial crack [12]. Therefore, it is difficult to introduce well-developed radial cracks under the condition that the depth limitation of nanoindentation excludes the
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substrate effect. We calculated the fracture toughness using a scratch tester quantitatively.
In our paper, by analogy with Ti-Mo-N and Ti-Si-N films, Mo and Si atoms were incorporated into TiN film simultaneously using reactive magnetron sputtering, and
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the effects of Si content on the microstructure, compressive residual stress, hardness, fracture toughness and tribological properties of Ti-Mo-Si-N films were investigated.
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2. Experiment details
Ti-Mo-Si-N films with a thickness of about 2 µm were deposited on both AISI 304
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stainless steel and Si (100) wafer substrates using a multi-target magnetron sputtering system. Ti (99.9 %), Mo (99.95 %) and Si (99.9 %) targets with a diameter of 75 mm were sputtered by three radio frequency powers. Mirror polished stainless steels and Si (100) wafer substrates were ultrasonically cleaned in acetone and alcohol, and then mounted on the substrate holder in the vacuum chamber. Substrate holder was electrically grounded with no applied bias voltage. After base pressure reached 6×10-4 Pa, Ar (99.999 %) and N2 (99.999 %) were introduced into the chamber by means of 3
ACCEPTED MANUSCRIPT two separate gas manifolds. The Ti-Mo-Si-N films with different Si content were achieved by fixing the powers of Ti target at 250 W and Mo target at 150 W and adjusting the power of Si target from 0 W to 130 W while constantly keeping the working pressure at 0.3 Pa and the nitrogen to argon ratio (flow ratio) at 10:3. Besides
achieved using temperature transducer after deposition.
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this, the substrate was not heated during deposition and its temperature value was
The infrared analysis was carried out in transmission mode at normal incidence by
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GIGILAB FT5200 Fourier transform infrared spectrometer (FTIR). The elemental
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composition and chemical bonds of the film were characterized by X-ray photoelectron spectroscopy (XPS) with Al Kα irradiation at a pass energy of 160 eV after removing the surface contaminants on the films by sputtering with Ar+ ion beam at a primary energy of 3 keV for 3 min. The spectra were calibrated by the C 1s line with a binding energy of 285.0 eV, and then corrected for the linear emission
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background and decomposed into peaks with Gaussian-Lorentzian line shapes by a non-linear least-square fitting method. The elemental compositions of the film were
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calculated from the areas under their XPS peaks by considering the relative sensitivity factors based on Wagner data analysis system. The microstructure of the films was
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characterized by glancing incident X-ray diffraction (XRD) at 2-degree incidence angle with a Siemens X-ray diffractometer using Cu Kα radiation. The compressive residual stress (σ) of the films was calculated by Stoney’s equation [3]:
σ=
E ts 2 1 1 − 1 − υ 6t f R Rs
(1)
where E is Elastic modulus of the substrate (E=170 GPa), υ is Poisson’s ratio of the substrate ( υ =0.3), ts is the thickness of the substrate, tf is the thickness of the film, R is the substrate curvature radii of the Si wafer and Rs is the curvature radii of the film 4
ACCEPTED MANUSCRIPT on the Si wafer. R and Rs were measured by Bruker DEKTAK-XT profilometer. Hardness and elastic modulus of the films were determined with nanoindenter CPX+NHT2+MST, which was equipped with a diamond Berkovich indenter tip (3-side pyramid). An automatic indentation mode was programmed to place
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indentations in a 3×3 array. The maximum load of 3 mN was used to meet the criterion of d/h<0.1, which guarantees a minimal influence from the substrate on the hardness measure, where d and h are the indenter penetration depth into the film and
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the film thickness, respectively. Before making indents on the films, the indenter was calibrated with respect to a reference sample of fused silica. The fracture toughness
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(KIC) of the films was measured by a scratch tester, which was equipped with a diamond Rockwell tip. The fracture toughness was calculated by the following equation [13]: 2 pf g a 1/2 −1 R ( ) sin R2 cotθ π a
(2)
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K IC =
where p is the pressure required to open the crack, R is the radius of the indenter cone into the groove, 2a is the total crack length, fg is the coefficient of grooving friction
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which is obtained from scratch test, θ (θ=65.03 °) is a angle between the axis of the
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diamond pyramid and the three faces. The 30 min wear test was carried out along a circular track of 8mm diameter against a 9 mm diameter Al2O3 counterpart at 50 rpm under a constant normal load of 3 N in the atmosphere (the relative humidity of about 25-30 %) at room temperature (about 25 °C) using a UTM-2 CETR tribometer. After the wear tests, the wear tracks were examined using a profilometer (Bruker DEKTAK-XT) to measure the wear loss of the films (V). The wear rate of the films (W) was calculated by Archard’s classical wear equation [14]:
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V S×L
(3)
where S is the total sliding distance and L is the applied load. For FTIR, XPS, XRD and nanoindentation measurements, the Ti-Mo-Si-N samples
3. Results and discussion 3.1 Elemental compositions and microstructure
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were deposited on AISI 304 stainless steel substrates.
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were deposited on Si (100) wafer substrates; for scratch and wear tests, the samples
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Fig. 1 shows the elemental compositions of the Ti-Mo-Si-N films as a function of the power of Si target. As the power of Si target increases from 0 W to 130 W, the Si content in the Ti-Mo-Si-N films increases almost linearly from 0 at.% to 17.2 at.%. Although FTIR is not usually employed to analyze the structure of nitride films, it
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can give precious information about the chemical bonds involved in the amorphous phase as for example in the Zr-Si-N [15] and Ti-B-N [16] films. Fig. 2 illustrates the
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FTIR spectra of Ti-Mo-Si-N films with various Si content. In order to clearly identify the chemical bonding involved in Ti-Mo-Si-N films, Si3N4 film deposited under the
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same experimental conditions was also characterized by FTIR. As shown in Fig. 2, for Si3N4 film, four absorption peaks are detected at approximately 480cm-1,700cm-1, 1500cm-1 and 2400cm-1. The absorption peaks at about 480cm-1 and 700cm-1 are attributed to the symmetric stretching of Si-N bonds in amorphous Si3N4 [17], and other absorption peaks at about 1500cm-1 and 2400cm-1 are similar to the results of amorphous Si3N4 characterized by FTIR in Ref. [15]. No absorption peak is detected for Ti-Mo-Si-N films with a Si content in the range between 3.1 at.% and 5.0 at.%. However, as the Si content increases from 8.8 at.% to 17.2 at.%, four absorption 6
ACCEPTED MANUSCRIPT peaks which is similar with the peaks of Si3N4 film are detected by FTIR. Besides, the intensity of those peaks increases with increasing Si content in the films, indicating that the amount of Si3N4 increases with increasing Si content in the films. In order to clarify further bonding status of the amorphous phase in the Ti-Mo-Si-N
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films, Si 2p spectra with various Si content was carried out by XPS and the results are shown in Fig. 3. The Si 2p spectra reveal that no obvious peak is detected for Ti-Mo-N and Ti-Mo-Si-N (5.0 at.% Si) films. The peak corresponding to 101.8 eV,
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which is in good agreement with the silicon fourfold coordinated to nitrogen as Si3N4
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[18, 19], is detected for Ti-Mo-Si-N film at 8.8 at.% Si. This is consistent with the results from Fig. 2.
Fig. 4 shows the XRD pattern (a) and lattice constant (b) of the Ti-Mo-Si-N films with various Si content. As shown in Fig. 4(a), the diffraction pattern of Ti-Mo-N film
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presents multiple orientations of (111), (200), (220), (311) and (222) crystal planes of face-centered cubic (fcc) TiN phase (PDF card 65-0715) where some Ti atoms in TiN lattice are replaced by Mo atoms [2, 4]. For each XRD pattern of Ti-Mo-Si-N films, it
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shows a single of fcc-TiN structure. Since no X-ray diffraction peaks corresponding to crystalline phases such as Si3N4, MoSi2 or TiSi2 are observed, it can be concluded
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that the Si3N4 phase in Ti-Mo-Si-N films with a Si content in the range between 8.8 at.% and 17.2 at.% detected by FTIR and XPS is amorphous. As shown in Fig. 4(b), the lattice constant of the films increases with increasing Si content, this suggests that the Si atoms dissolve into the interstitial sites of Ti-Mo-N lattice. It is well known that the phase formation is affected by the thermodynamic and kinetic issues. Compared to the amorphous, the crystalline phase is preferred to form because of its higher potential energy [20]. The formation of TiN is easier than that of 7
ACCEPTED MANUSCRIPT amorphous Si3N4 because the difference of electronegativity (Ti=1.5, Si=2.5 and N=3.0) between Ti and N is larger than that between Si and N. Many Ti-Si-N (nc-TiN/a-Si3N4) films with a low Si content were deposited at the substrate temperature of >500 °C [21]. However, solid solution of Ti-Si-N is easy to form at
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low temperature of <100 °C because TiN crystals are able to kinetically hinder the formation of the nc-TiN/a-Si3N4 at this temperature [20]. The Ti-Mo-Si-N films were deposited at uncontrolled temperature in our experiment. So the TiN is preferentially
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formed compared to the Si3N4 in our case. For low Si content in the Ti-Mo-Si-N films (<5.0 at.%), Si atoms dissolved into interstitial sites in Ti-Mo-N lattice and formed
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interstitial solid solutions because no amorphous Si3N4 is detected in Fig. 2 and 3. With a further increase in Si content, the dissolution of Si is too weak to consume all Si atoms deposited on the films. So the amorphous Si3N4 is detected in Fig. 2 and 3 and the films are consisted of interstitial solid solutions and amorphous Si3N4 phases.
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To sum up, interstitial solid solution of Ti-Mo-Si-N is formed by dissolution of Si atoms into fcc-Ti-Mo-N lattice when the Si content is below 5.0 at.%. With a further
phases.
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increase in Si content, the films consist of fcc-Ti-Mo-Si-N and amorphous Si3N4
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3.2 Mechanical properties
The residual stress of Ti-Mo-Si-N films with various Si content is presented in Fig.
5. As shown in Fig. 5, all films regardless of Si content are in compressive stress state. The compressive residual stress of Ti-Mo-N film is about -1.5 GPa. The compressive residual stress of Ti-Mo-Si-N films first increases gradually and then decreases with increasing Si content, after reaching the maximum value of -2.4 GPa, at 8.8 at.% Si.
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σT =
where nF is the Poisson ratio of the film (nF=0.3); αF and αS are the thermal expansion
°
C). In our experiment, the
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the temperature during stress measurement (25
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coefficient of the film and the substrate; TD is the deposition temperature; and TM is
temperature of substrate increases from 89 °C at 0 at.% Si to 162 °C at 17.2 at.% Si because the number of ions bombardment on the substrate increases with increasing the power of Si target. The thermal expansion coefficient of Si wafer, TiN and amorphous Si3N4 is 2.7×10-6/°C [3], 6.0×10-6/°C [23] and 3.2×10-6/°C [24],
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respectively. Thus, the thermal stress of all films is in tensile stress state and its value is in the range of 0.1-0.2 GPa. The compressive residual stress of Ti-Mo-Si-N films is
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mainly attributed to the intrinsic stress. The lattice expansion induced by incorporation of Si atoms into the interstitial sites in Ti-Mo-N crystal lattice (as
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shown in Fig. 2b) leads to the increase of intrinsic stress in the films. Besides this, the compressive residual stress can be released by the amorphous Si3N4 due to its low density and disordered network [10]. On the basis of the above analysis, the lattice expansion induced by incorporation
of Si atoms into the interstitial sites in Ti-Mo-N crystal lattice leads to the increase of compressive residual stress of the films when the Si content is below 5.0 at.%; although the amorphous Si3N4 appears in the film at 8.8 at. % Si, the compressive residual stress of the film further increases because the amount of amorphous Si3N4 is 9
ACCEPTED MANUSCRIPT too few to weaken the effect of Si dissolution; with a further increase in Si content, the decrease of compressive residual stress is mainly attributed to the increase of the amount of amorphous Si3N4. Fig. 6 shows the load-displacement curves of Ti-Mo-Si-N films with various Si
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content obtained from the nanoindentation measurement. The maximum load of 3 mN was used to meet the criterion of d/h<0.1, which guarantees a minimal influence from the substrate on the hardness measurement, where d and h are the indenter penetration
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depth into the film and the film thickness, respectively. The hardness (H), elastic
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modulus (E), H/E and H3/E*2 ratios are shown in Fig. 7. As shown in Fig. 7(a), the hardness and elastic modulus of Ti-Mo-N film are 27.4 GPa and 401 GPa, respectively. The hardness and elastic modulus of Ti-Mo-Si-N films first increase gradually and then decrease with increasing Si content, after reaching the maximum
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values of 34.5 GPa and 412 GPa, respectively, at 5.0 at. % Si.
The initial increase in hardness is mainly attributed to the effect of solid solution strengthening and the increase of compressive residual stress. As the Si content
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increases from 8.8 at.% to 17.2 at.%, amount of amorphous Si3N4 in the films could be result in the decrease of the hardness because the hardness of Si3N4 film deposited
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under the same experimental conditions is about 21.3 GPa. Besides this, decreasing compressive residual stress also leads to the decrease of the hardness of the films when the Si content is above 8.8 at.%. The elastic modulus is influenced by the compressive residual stress [25]. Higher compressive residual stress usually leads to higher elastic modulus [3, 25, 26]. The elastic modulus of the films increases from 401 GPa at 0 at.% Si to 412 GPa at 5.0 at.% Si because of the increase of the compressive residual stress. The elastic 10
ACCEPTED MANUSCRIPT modulus of Si3N4 film that we deposited under the same experimental conditions is 190 GPa. Thus, the decrease of elastic modulus is mainly attributed to the appearance of the amorphous Si3N4 with a further increase in Si content in Ti-Mo-Si-N films. An interesting attention is being given to the relationship between the hardness and
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elastic modulus during the development of the hard films for mechanical applications recently. The ratio of hardness to elastic modulus (H/E) is regarded as an important factor to describe the resistance of material against elastic strain to failure [3]. The
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H3/E*2 ratio (E* is the effective elastic modulus and E*=E/(1-µ2) where µ is Poisson’s ratio) relates to the resistance to plastic deformation and the tribological properties of
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the films [15]. As shown in Fig. 7(b), the ratios of H/E and H3/E*2 of Ti-Mo-N film are 0.069 and 0.137 GPa, respectively. The H/E and H3/E*2 ratios of Ti-Mo-Si-N films, all of which are higher than that of Ti-Mo-N film, firstly increase and then decrease with increasing Si content in the films, after reaching the maximum values of
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0.085 and 0.204 GPa, respectively, at 5.0 at.% Si. The elastic recovery is the ratio of the recovered displacement after unloading to
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the total indentation displacement and can be calculated from load-displacement curves (Fig. 6). The elastic recovery of Ti-Mo-N film is 57 %. The elastic recovery of
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Ti-Mo-Si-N films, which is in the range between 46 % and 43 %, is lower than that of Ti-Mo-N film indicating higher fracture toughness in a qualitative sense [10]. The fracture toughness (KIC) of Ti-Mo-Si-N films with various Si content was calculated using Eq. (2) and the results are shown in Fig. 8. As shown in Fig. 8, the fracture toughness of Ti-Mo-N film is about 1.3 MPa.m1/2. The fracture toughness of Ti-Mo-Si-N films initially increases and then decreases with the increase of Si content, after reaching the maximum value of about 2.7 MPa.m1/2, at 5.0 at.% Si.
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ACCEPTED MANUSCRIPT Fig. 9 illustrates the fracture toughness of Ti-Mo-Si-N films as a function of compressive residual stress (a) and H3/E*2ratio (b). It indicates that the fracture toughness increases with the increase of compressive residual stress because the compressive residual stress allows films to sustain more tensile strain before fracture
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[12]. Since H3/E*2 ratio is connected with the resistance to plastic deformation, the fracture toughness of Ti-Mo-Si-N films increases with increasing H3/E*2 ratio.
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3.3 Tribological properties
Fig. 10(a) illustrates a few examples of friction coefficient curves and wear track
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profiles against Al2O3 counterpart at RT after ball-on-disc wear test. Each curve records the friction coefficient variation from running-in to the long-term steady state friction. The wear track of Ti-Mo-Si-N films becomes shallower and narrower as the Si content increases from 3.1 at.% to 5.0 at.%. With a further increase in Si content,
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the wear track becomes deeper and wider. In all cases, the depth of the wear tracks is lower than the thickness of the films, indicating all the films in our paper were not worn out.
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The average friction coefficient and wear rate of Ti-Mo-Si-N films with various Si content are presented in Fig. 10(b). As shown in Fig. 10(b), the average friction
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coefficient and wear rate of Ti-Mo-N film are 0.38, 7.1×10-7 mm3/N.mm, respectively. As the Si content increases, the average friction coefficient and wear rate firstly decrease and then increase, after reaching the minimum values of 0.35, 7.8×10-8 mm3/N.mm, respectively, at 5.0 at.% Si. Ti-Mo-Si-N films with Si content in the range between 3.1 at. % and 5.0 at.% present lower average friction coefficients and wear rates compared with Ti-Mo-N film.
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ACCEPTED MANUSCRIPT In order to study the effects of Si content on the tribological properties of Ti-Mo-Si-N films, the 3D and optical microscope images of wear track on Ti-Mo-Si-N films are presented in Fig. 11. As shown in Fig. 11(a), the wear track of Ti-Mo-Si-N film at 3.1 at.% Si shows a relatively narrow and smooth. There are many
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fine scratches along the counterpart sliding direction. It reveals that the main wear mechanism is abrasive wear. By comparison, as the Si content increases to 5.0 at.%, improved wear resistance is observed and the wear track shows much narrow and
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smooth. It is attributed to the dissolution of Si in the Ti-Mo-N lattice which could effectively inhibit the cracks to propagate and reduce the contact stress between
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sliding surface and counterpart [6]. With a further increase in Si content, the film at 17.2 at.% Si exhibits a relatively poor wear resistance and a lot of obvious scratches and wear debris are detected. The increase of average friction coefficient and wear rate is related to the formation of the amorphous Si3N4 in the films.
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In dry ball-on-disc wear test, the tribo-film plays an important role in determining the tribological properties [2, 3]. In addition, tribo-oxidation was reported to be a
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general phenomenon in the wear test at room temperature. It was reported that MoO3, SiO2 and Si(OH)4 could be formed easily by a reaction between the wear track and
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oxygen/moisture in the air for transition metal nitride films, e.g. Ti-Mo-N [2, 4] and Mo-Al-N films [3]. Although MoO3 can reduce the average friction coefficient, it can be worn away easily by the counterpart during the wear test because MoO3 has low shear strength [3, 4]. SiO2 and Si(OH)4 were known to play a role as a self-lubricating and protection layer on the wear track of the films and result in lower wear rate compared with MoO3. When the Si content is in the range between 3.1 at.% and 5.0 at.%, the addition of Si in the films results in the decrease of MoO3, so the wear tracks become shallower with increasing Si content. 13
ACCEPTED MANUSCRIPT The tribological properties of transition metal nitride films are also influenced by their mechanical properties [3, 4]. Hardness is one of the mechanical properties relating to the tribological properties. Higher hardness leads to an improved load
According to the Archard’s equation [27]: V W =K L H
(4)
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carrying capacity and a smaller real contact area with the counterpart [3, 4].
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where V is the wear volume, L is the sliding distance (38 m), K is Archards wear
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coefficient, W is the normal load (3 N), H is the hardness of the film. So the wear rate of Ti-Mo-Si-N films is reverse ratio to the hardness. So the wear rate shows the nearly opposite variation trend to that of hardness with the increase of Si content. Besides this, it is considered that the parameters of H/E and H3/E*2 are more important than hardness alone in valuing the tribological properties of the films [3].The high values
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of H/E and H3/E*2 confirm an improved tribological properties. To sum up, as the Si content increases from 3.1 at.% to 5.0 at.%, the decrease of
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average friction coefficient and wear rate could be attributed to the increase of Si content in the films, the hardness and the ratios of H/E and H3/E*2. With a further
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increase in Si content, the increase of average friction coefficient and wear rate could be mainly attributed to the amorphous Si3N4. The decrease of the hardness, the ratios of H/E and H3/E*2 also attribute the increase of the average friction coefficient and wear rate.
4. Conclusions Ti-Mo-Si-N films with various Si content were deposited by reactive magnetron sputtering. The main results could be concluded as follows: 14
ACCEPTED MANUSCRIPT (1)
When the Si content was below 5.0 at.%, Si dissolved into interstitial sites in
fcc-Ti-Mo-N lattice and formed interstitial solid solutions. With a further increase in Si content, the films were consisted of fcc-Ti-Mo-Si-N and amorphous Si3N4. (2)
The hardness and fracture toughness of the films first increased and then
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decreased with increasing Si content, after reaching the maximum values of 34.5 GPa and 2.6 MPa.m1/2, respectively, at 5.0 at.% Si. The hardness enhancement was mainly attributed to the solid solution strengthening and the increase of compressive residual
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stress. The formation of amorphous Si3N4 could be result in the decrease of the hardness. The fracture toughness was influenced by the compressive residual stress
(3)
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and the H3/E*2 ratio.
The average friction coefficient and wear rate first decreased and then
increased with increasing Si content, after reaching the minimum values of 0.35 and
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7.8×10-8 mm3/Nmm, respectively, at 5.0 at.% Si. The average friction coefficient and wear rate were influenced by the hardness, amorphous Si3N4 and the ratios of H/E and H3/E*2.
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Acknowledgment
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This project was supported by National Natural Science Foundation of China (51074080, 51374115) and Research Innovation Program for College Graduates of Jiangsu Province (CXZZ12-0717).
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sputtering’, Materials Science and Engineering B, 2012, 177, 1120-1125.
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ACCEPTED MANUSCRIPT Figure captions Fig. 1 Elemental compositions of Ti-Mo-Si-N films as a function of the power of Si target
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Fig. 2 The FTIR spectra of Ti-Mo-Si-N films with various Si content Fig. 3 XPS spectra of Si 2p of Ti-Mo-Si-N films with various Si content
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Fig. 4 XRD patterns (a) and lattice constant (b) of the Ti-Mo-Si-N films
Fig. 5 Compressive residual stress of Ti-Mo-Si-N films with various Si content
at.%, (b) 5.0 at.%, (c) 17.2 at.%
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Fig. 6 Loading-displacement curve of Ti-Mo-Si-N films with various Si content: (a) 0
Fig. 7 Hardness (H), elastic modulus (E), H/E and H3/E*2 ratios of Ti-Mo-Si-N films
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with various Si content
Fig. 8 The toughness of Ti-Mo-Si-N films with various Si content Fig. 9 The fracture toughness of Ti-Mo-Si-N films as a function of compressive
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residual stress (a) and H3/E*2 (b)
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Fig. 10 The friction coefficient curves, wear track profiles (a), the average friction coefficients and wear rates (b) of Ti-Mo-Si-N films with various Si content Fig. 11 3D and optical microscope images of wear tracks on Ti-Mo-Si-N films with various Si content: (a) 3.1 at.%, (b) 5.0 at.%, (c) 17.2 at.%
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ACCEPTED MANUSCRIPT Highlights 1. Ti-Mo-Si-N films were deposited by reactive magnetron sputtering. 2. Hardness of Ti-Mo-Si-N film (5.0 at.% Si) reached a maximum value of 34.5 GPa.
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3. Addition of Si (<5.0 at.% Si) led to the increase of fracture toughness.
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4. The films (3.1-5.0 at.% Si) were found to be optimized for wear resistance tools.