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Microstructure, mechanical and tribological properties of TiN-Ag films deposited by reactive magnetron sputtering
Vacuum 141 (2017) 82e88
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Vacuum journal homepage: www.elsevier.com/locate/vacuum
Microstructure, mechanical and trobological properties of TiN-Ag films deposited by reactive magnetron sputtering Hongbo Ju, Lihua Yu, Dian Yu, Isaac Asempah, Junhua Xu* School of Materials Science and Engineering, Jiangsu University of Science and Technology, Mengxi Road 2, Zhenjiang, Jiangsu Province, 212003, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 January 2017 Received in revised form 21 March 2017 Accepted 21 March 2017 Available online 23 March 2017
Composite TiN-Ag films with various Ag content (Ag/(Ti þ Ag)) were deposited using reactive magnetron sputtering and the influence of Ag content on the crystal structure, mechanical and tribological properties were investigated. The result showed that face-centered cubic (fcc) TiN and fcc-Ag co-existed in the films and TiN had a columnar-growth and the Ag nanoparticles embedded in the boundary of columnar crystal. The hardness of TiN-Ag films initially increased gradually and reached an optimum value, and then decreased with an increase in Ag content in the films. The maximum hardness value was 29 GPa at 0.8 at.% Ag. The addition of Ag into TiN film could enhance the fracture toughness (KIC) and critical load (LC1) of the films because the nanoparticle Ag provides a large volume fraction of grain boundaries. As Ag was added into TiN film, the average friction coefficient decreased from 0.78 at 0 at.% Ag to 0.20 at 41.1 at.% Ag; however, the wear rate of TiN-Ag films initially decreased and then increased, after reaching the minimum value of about 1.3 107 mm3/(mm.N), at 0.8 at.% Ag. © 2017 Elsevier Ltd. All rights reserved.
Keywords: Reactive magnetron sputtering TiN-Ag films Microstructure Mechanical and tribological properties
1. Introduction Advances in machining technology in recent times have led to an increase in demand for materials for severe applications associated with high temperature and the stresses that are generated at the tool/work piece interface [1]. Nano-structured films, which usually exhibit a unique blend of hardness, thermal stability, wear resistance and low average friction coefficient, have gained increasing interest in recent years for a number of different applications. Transition metal nitride (TMN) films deposited by physical vapor deposition techniques, which are widely used for wear reduction on the cutting and die tools, have aroused considerable interest in the past few decades because of some of their excellent properties [1e8]. As modern manufacturing industry develop at an unprecedented rate, its further demands such as relative lower average friction coefficient are driving force for further development of TMN matrix films [3,8]. Seeking effective lubricants for transition metal nitride films is a challenging task for the tribology community [6]. Silver considered as a traditional lubricating material was combined into TMN matrix recently and a large number of related
studies has been carried out [6,9e13]. TiN-Ag film is one of the most representative self-lubriacting films. For example, H. Kostenbauer et al. [14] studied a series of TiN-Ag films' tribological properties at elevated testing temperature systematically and found that TiN films doped Ag could adjust their surface chemistry according to the changing of the service condition. D. Du et al. [15] reported that the fretting fatigue resistance of a Ti-6Al-4V alloy could be improved significantly by deposition of TiN-Ag film on the surface of the alloy. It is widely accepted that metallic Ag phase and TiN phase co-exist in the TiN-Ag films with higher Ag content. However, as the Ag content in the films is very low (eg. <2 at.% Ag), it is difficult to study what state the Ag exists in the TiN films, since only TiN peaks appear in the XRD pattern of the films. In our paper, the most representative TMN filmdTiN film was opted as the object and Ag element was introduced into TiN film to create the TiN-Ag composite films by reactive magnetron sputtering. High resolution transmission electron microscopy (HRTEM) and back scattered electron (BSE) were performed to investigate the crystal structure of TiN-Ag films. Besides this, the effects of Ag content on the mechanical (hardness and fracture toughness) and tribological properties of TiN-Ag films were also discussed in this paper. 2. Experimental details
* Corresponding author. E-mail address: [email protected] (J. Xu). http://dx.doi.org/10.1016/j.vacuum.2017.03.026 0042-207X/© 2017 Elsevier Ltd. All rights reserved.
TiN-Ag films with a thickness of about 2 mm were deposited on
H. Ju et al. / Vacuum 141 (2017) 82e88
mirror polished stainless steel (AISI 304 SS, 192 Hv, 15 mm 15mm 2.5 mm) with surface roughness of ~70 nm and silicon (100) wafer substrates using a multi-target magnetron sputtering system. Ti (99.9%) and Ag (99.9%) targets with a diameter of 75 mm were sputtered by two 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 6 104 Pa, Ar (99.999%) and N2 (99.999%) were introduced into the chamber by means of two separate gas manifolds. The TiN-Ag films with different Ag content were achieved by fixing the powers of Ti target at 280 W and adjusting the power of Ag target from 0 W to 150 W while constantly keeping the working pressure at 0.3 Pa, the deposition temperature at 200 C and the nitrogen to argon ratio (flow ratio) at 10:5. The substrates were not biased. After the deposition, the deposition rate was calculated and its value was ~17 nm/min. The TiN-Ag films deposited on the AISI 304 SS substrates were used to test for adhesion and wear test of the film, while film's elemental composition, crystal structure, hardness, elastic modulus and fracture toughness were evaluated using the film on the silicon (100) wafer substrates. The elemental compositions of the films were determined by energy dispersive spectroscopy (EDS) on an EDAX DX-4 energy dispersive analyzer. The microstructure of the films was evaluated using a Siemens X-ray diffractometer using Cu Ka radiation, operated at 40 kV, 35 mA. The grain size of the films was calculated by using Debye-Scherrer's formula for most intense peak having orientation (111) in the face-centered cubic (fcc) structure. The back scattered electron (BSE) image was obtained by the scanning electron microscopy (SEM, JEM-6480). High resolution transmission electron microscopy (HRTEM) was performed using a JEOL JEM-2010F microscope operated at an accelerating voltage of 200 kV. The residual stress (s) of the films was calculated by Stoney's equation [16].
KIC
! 0:5 E P ¼a 1:5 H Cm
83
(2)
where E is the elastic modulus of the film, H is the hardness of the film, P is 200 mN in our paper, Cm is the crack length, a is the empirical constant and its value is 0.016 in our paper. The adhesion of the films was measured by a scratch tester, which was equipped with a diamond Rockwell tip. The load was increased linearly from 0 N up to 40 N. The critical loads (Lc1) and complete delamination load (LC2) were measured by analyzing the failures events in the scratch track by optical microscopy. A 30 min wear test was carried out along a circular track of 8 mm 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 25e30%) at room temperature (about 25 C) using a UTM-2 CETR tribometer. Raman spectroscopy using the 514.5 nm Arþ laser with a backscattering optical configuration was used to study the tribo-film on the surface of wear track. 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:
W¼
V SL
(3)
where S is the total sliding distance; L is the applied load.
3. Results and discussion
E ts2 s¼ 1 y 6tf
1 1 R Rs
3.1. Crystal structure
(1)
where E–elastic modulus of the substrate (E ¼ 170 GPa).
y–Poisson's ratio of the substrate (y¼0.3)
ts–the thickness of the substrate tf–the thickness of the film R–the substrate curvature radii of the Si wafer Rs–the curvature radii of the film
The curvature radii were done using data obtained from Bruker DEKTAK-XT profilometer. Five (5) areas of each sample were made for the curvature radii and the mean value taken. In determining the hardness and elastic modulus of the films, an indenter calibrated using fused silica as reference was used in a nanoindentation tester system (nano-indenter CPX þ NHT þ MST), equipped with a diamond Berkovich indenter tip. The maximum load of 3 mN was used and a minimum of nine (9) indentations were made for each sample and the mean value taken. Before making indentation on the films, the indenter was calibrated with respect to a reference sample of fused silica. The fracture toughness (KIC) of the films was also measured by the nano-indentation tester at 200 mN and the fracture toughness was calculated by the following equation [17]:
Fig. 1 shows the XRD patterns of TiN-Ag films with various Ag content (Ag/(Ti þ Ag), at.%). As shown in Fig. 1, the diffraction of binary TiN film shows multiple orientations of (111), (200) and (222) crystal plane of face-centered cubic (fcc) TiN phase (JCPDF card 65e0715). For each XRD pattern of ternary TiN-Ag films at<1.4 at.% Ag, it shows three peaks corresponding to fcc-TiN phase. Increasing Ag content furtherly, besides the fcc-TiN diffraction peaks, two peaks with low-intensity at ~39 and ~45 are detected. The two peaks correspond to fcc-Ag. The solid solubility of Ag in the transition metal films is very limited based on the result in Refs. [6,18], therefore, for the films at<1.4 at.% Ag, no peak corresponding to fcc-Ag is attributed to the limited precision of the XRD. The grain size of the films was calculated using Debye-Scherrer's formula for most intense peak having orientation (111) in the fccTiN. The addition of Ag drops the grain size form ~35 nm at 0 at.% Ag to ~7 nm at 41.1 at.% Ag. The nanoparticle Ag embedded in the TiN films could prevent grain growth and similar results were also reported in Refs. [14,19]. The resolution transmission electron microscopy (TEM) and its selected area electron diffraction (SAED) pattern of TiN-Ag film at 0.7 at.% Ag were carried out to study the crystal structure of TiN-Ag films at<1.4 at.% Ag. The result is shown in Fig. 2. The cross-section TEM image of TiN-Ag film at 0.7 at.% Ag (Fig. 2(a)) reveals that the film has a columnar-growth and dense structure. Its corresponding
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3.2. Mechanical and tribological properties
Fig. 1. XRD patterns of TiN-Ag films with various Ag content.
SAED pattern is illustrated in Fig. 2(b), from inside to outside, the diffraction rings can be assigned to the lattice planes of fcc-TiN (111), fcc-Ag (111), fcc-TiN (200), fcc-TiN (220) and fcc-TiN (311). It suggests that fcc-TiN and fcc-Ag coexist in the film. Fig. 2(c) shows the HRTEM image of inner region of columnar crystal, a lattice fringe with a lattice spacing of about 0.2445 nm is detected. This lattice fringe belongs to the fcc-TiN (111), since the value of lattice spacing of fcc-TiN (111) (JCPDF card 65e0715) is 0.2448 nm. The HRTEM image of boundary of columnar crystal is shown in Fig. 2(d), Ag nanoparticle with a radius of about 5 nm is detected in the field view besides TiN lattice fringe. From Fig. 2(d), it can also be observed that the Ag nanoparticle seems to be randomly oriented with respect to the TiN lattice fringe. This suggests that Ag nanoparticle is not chemically bonded on the boundary of columnar crystal TiN, such as metallic Ag-Ti. Similar results were also reported in Refs. [6,19,20]. Fig. 3 illustrates the back scattered electron (BSE) image of TiNAg film at 23.5 at.% Ag. The grey cylinders in the image are TiN and the small white particles are Ag. TiN exhibits a columnar-growth and Ag particles evenly dispersed along the boundary of TiN columnar grain. Based on above results, TiN has a columnar-growth, dense structure and Ag nanoparticles embedded in boundary of the crystal.
Table 1 illustrates the residual stress of TiN-Ag films. All films regardless of Ag content are in compressive state and the compressive residual stress of the films drops gradually with the rise of Ag content in the films. The drop of the compressive residual stress is attributed to the inhibition of titanium nitride grain by Ag addition into TiN matrix. The hardness (H) and elastic modulus (E) of TiN-Ag films with various Ag are shown in Fig. 4. As shown in Fig. 4, the hardness and elastic modulus of binary TiN film are 21 GPa and 320 GPa, respectively. The addition of Ag into TiN film rises the hardness from 23 GPa at 0.7 at.% Ag to 29 GPa at 0.8 at.% Ag, and then drops gradually with a further rise of Ag content. The elastic modulus of TiN-Ag films drops gradually with an increase in Ag content. As the Ag content increases from 0 at.% to 0.8 at.%, the hardness enhancement could be attributed to the effect of fine-grain strengthening. With a further increase in Ag content, the decrease of hardness of the films is attributed to the increasing in the very soft and highly mobile silver in the film [6,21,22]. The decrease of the compressive residual stress also drops the hardness of the films at>0.8 at.% Ag. In addition, the elastic modulus of Ag film that we deposited under the same experimental conditions is only about 70 GPa, so the decrease in elastic modulus of TiN-Ag is most probably associated with the incorporation of Ag in the TiN matrix. 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 [7]. It is considered that the hard surfaces can resist abrasive wear and low elastic modulus allows elastic deformation to easy occur when the contact stress is applied. So the H/E ratio is proposed as a key parameter controlling the wear resistance. In addition, the plastic deformation resistance factor (H3/E2) also can be related with the wear rate [7]. The H/E and H3/E2 ratios of TiN-Ag films are shown in Fig. 5. As shown in Fig. 5, the ratios of H/E and H3/E2 of the films rise gradually from 0.066 to 0.090 GPa at 0 at.% Ag to 0.097 and 0.27 GPa at 0.8 at.% Ag, and then drop gradually with a further rise of Ag content. Fracture toughness (KIC) is another vital mechanical property of materials. The radial cracking indentation method is widely accepted to calculate KIC of thin films [23]. A maximum load of 200 mN was chosen to investigate KIC in our paper, since the visible cracks is hard to induced by small load. Although the displacement caused by the 200 mN is greater than ten percent (10%) of the film thickness, the value of KIC obtained from the 200 mN load could reflect the exact value of the film in some sense because all TiN-Ag films have same interlayer and substrate. The indentation morphology of TiN-Ag films with various Ag content is shown in Fig. 6. As shown in Fig. 6(a), the film at 0.7 at.% Ag exhibits the obvious radial cracks paralleled to the indentation diagonal. Rising the Ag content to 23.5 at.% enhances the fracture toughness considerably, so no radial cracks is detected after the 200 mN indentation, as shown in Fig. 6(b). KIC of TiN-Ag films are shown in Table 1 and its value rises gradually from 0.4 MPa m1/2 at 0 at.% Ag to 1.6 MPa m1/2 at 1.4 at.% Ag. Rising Ag content furtherly, KIC could not be calculated due to the disappearance of radial cracks. The nanoparticle Ag provides a large volume fraction of grain boundaries, this limits initial crack sizes and helps to deflect and terminates the growth of cracks. Therefore, the addition of Ag into TiN film enhances KIc. Fig. 7 illustrates the scratch track optical images of TiN-Ag films with various Ag content after the scratch test. As shown in Fig. 7, adding Ag into TiN film induces an obvious change of the deformation behavior. LC1 is considered as the critical load of the film and its value is obtained from the crack initiation. LC2 is obtained from
H. Ju et al. / Vacuum 141 (2017) 82e88
85
Fig. 2. TEM images of TiN-Ag film at 0.7 at.% Ag: (a) The cross-section TEM image of TiN-Ag film, (b) SAED pattern of TiN-Ag film, (c) HRTEM image of inner region of columnar crystal, (d) HRTEM image of boundary of columnar crystal.
Fig. 4. Hardness and elastic modulus of TiN-Ag films with various Ag content.
Fig. 3. The back scattered electron image of TiN-Ag film at 23.5 at.% Ag.
the occurrence of complete delamination. Table 1 shows the value of LC1 and LC2. The value of LC1 rises gradually from ~4 N at 0 at.% Ag
to ~12 N at 1.4 at.% Ag, the failure of the film is detected in the brittle behavior, and then as the Ag content rises furtherly, the deformation behavior is changed and scratch of the film at 23.5 at.% Ag exhibits no obvious cracks because of the rise of high plasticity and ductility metal Ag in the film. LC1 is attributed to the fracture
Table 1 Compressive residual stress, fracture toughness (KIC), critical load (LC1), complete delamination load (LC2) and surface energy of interfacial crack of TiN-Ag films with various Ag content. Ag content(at.%)
Compressive residual stress (GPa)
KIC (MPa.m1/2)
LC1 (N)
LC2 (N)
Gc (J/m2)
0 0.7 0.8 1.4 23.5 41.1
þ0.24 þ0.20 þ0.20 þ0.19 þ0.13 þ0.08
0.4 0.5 0.9 1.6 No cracks No cracks
4 5 9 12 No cracks No cracks
12 18 21 26 23 17
170 220 290 525 370 267
± ± ± ± ± ±
0.01 0.01 0.01 0.01 0.01 0.01
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H. Ju et al. / Vacuum 141 (2017) 82e88
Gc ¼
d2c t 2Ef
(4)
Where t is the thickness of the film, Ef is the elastic modulus of the film, dc is the film critical stress of complete delamination and it can be calculated by the following equation:
dc ¼ Fig. 5. The H/E and H3/E2 ratios of TiN-Ag films with various Ag content.
3 !2 2Lc2 4 4 þ vf 3pm 1 2vf 5 8 pd2c
(5)
Where dc is the width of the scratch track obtained from the occurrence of complete delamination,
m is the friction coefficient from the scratch tester, vf is the Pisson rate of the film.
Fig. 6. Indentation morphology of TiN-Ag films at 200 mN: (a) 0.7 at.% Ag and (b) 23.5 at.% Ag.
toughness of the films. The value of LC2 rises from ~12 N at 0 at.% Ag to 26 N at 1.4 at.% Ag, and then it drops gradually with a further increase in Ag content. Dropping of H/E ratio might be attributes to the reduction of LC2. According to Ref. [24,25], the relationship between LC2 and surface energy of the interfacial crack (Gc) could be established by Griffith energy balance method and the equation is as follows:
Table 1 also illustrates the Gc of the TiN-Ag films with various Ag content. As shown in Table 1, The value of Gc rises from 170 J/m2 N at 0 at.% Ag to 525 J/m2 at 1.4 at.% Ag, and then it drops gradually with a further increase in Ag content. Therefore, the drop of LC2 might be attributed to the decrease of Gc of the films. Similar results were also in Refs. [26e29]. Fig. 8 illustrates a few examples of friction coefficient curves at room temperature (RT). As shown in Fig. 8, the friction curve of binary TiN film fluctuates markedly when the testing time is below 380 s, and then stays constant at about 0.79 for the remaining test times. For the TiN-Ag film at 0.8 at.% Ag, the friction coefficient had a steady value of about 0.62 following a dramatic move, according to the friction curve. With a further increase in Ag content, the friction coefficient of the film does not fluctuate and a lowest steady value of about 0.20 is achieved for the TiN-Ag film at 41.1 at.% Ag. The average friction coefficient (m) of TiN-Ag films, which is calculated from the friction coefficient data, is illustrated in Fig. 9. The m of binary TiN film is about 0.78. The m of TiN-Ag, all of which are lower than that of binary TiN film, is dependent on Ag content in the film. It decreases from about 0.75 at 0.7 at.% Ag to about 0.20 at 41.1 at.% Ag. This result clearly demonstrates the lubrication effect of Ag in TiN. A similar result was also reported in other Ref. [6,14]. Fig. 10 illustrates the wear rate of TiN-Ag films with various Ag content. As illustrated in Fig. 8, the wear rate of binary TiN film is about 1.73 106 mm3/(mm.N). The wear rate of TiN-Ag films initially decreases and then increases, after reaching the minimum value of about 1.3 107 mm3/(mm.N), at 0.8 at.% Ag. As Ag content increases from 0 at.% to 0.8 at.%, Ag has good lubricity property and could play a role lubrication. Therefore, the m and wear rate of the films decrease with the increasing Ag content. Besides this, the increasing ratios of H/E and H3/E2 also attributes to the decrease in the wear rate. With a further increase in Ag content, the m further decreases due to the increase in lubricant Ag content in the films. However, the soft lubricant Ag can be easily wiped off during the wear test. In addition, the ratios of H/E and H3/E2 decrease with the increase of Ag content. So the wear rate increases with the increase in Ag content. 4. Conclusion TiN-Ag films with various Ag content were deposited by reactive magnetron sputtering and the effects of Ag content on the crystal structure, mechanical and tribological properties of TiN-Ag films were investigated. The main results could be concluded as follows:
H. Ju et al. / Vacuum 141 (2017) 82e88
87
Fig. 7. Optical images of scratch track of TiN-Ag films with various Ag content: (a) 0.7 at.% Ag (b) 1.4 at.% Ag and (c) 23.5 at.% Ag.
Fig. 9. Average friction coefficient of TiN-Ag films with various Ag content at room temperature.
Fig. 8. Friction coefficient curves of TiN-Ag films with various Ag content at room temperature.
(1) A two-phase structure, consisting of fcc-TiN and fcc-Ag, coexisted in the TiN-Ag films. TiN film had a columnar-growth and dense structure and Ag nanoparticles were embedded in boundary of columnar crystal. (2) The hardness of TiN-Ag films initially increased gradually and reached an optimum value, and then decreased with an increase in Ag content in the films and the maximum value was 29 GPa at 0.8 at.% Ag. The hardness enhancement could be attributed to the effect of fine-grain strengthening. (3) The addition of Ag into TiN film enhanced the fracture toughness (KIC) and critical load (LC1) of the films because the nanoparticle Ag in the film provided a large volume fraction of grain boundaries.
Fig. 10. Wear rate of TiN-Ag films with various Ag content at room temperature.
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H. Ju et al. / Vacuum 141 (2017) 82e88
(4) As Ag was added into TiN film, the average friction coefficient decreased from 0.78 at 0 at.% Ag to 0.20 at 41.1 at.% Ag; however, the wear rate of TiN-Ag films initially decreased and then increased, after reaching the minimum value of about 1.3 107 mm3/(mm.N), at 0.8 at.% Ag. The tribological properties of the films were influenced by the lubricating Ag phase and the ratios of H/E and H3/E2. Acknowledgements Supported by National Natural Science Foundation of China (51074080, 51374115, 51574131), Research Fund of Jiangsu University of Science and Technology (1062931609). The authors also would like to thank Jing Luan for her valuable help in improving our manuscript. References [1] L. Hultman, Vacuum 57 (2000) 1e30. [2] Z.B. Qi, F.P. Zhu, Z.T. Wu, B. Liu, Z.C. Wang, D.L. Peng, C.H. Wu, Surf. Coat. Technol. 231 (2013) 102. [3] J. Xu, H. Ju, L. Yu, Acta Metall. Sin. 48 (2012) 1132e1138. [4] J. Xu, H. Ju, L. Yu, Vacuum 103 (2014) 21e27. [5] J. Xu, H. Ju, L. Yu, Vacuum 110 (2014) 47e53. [6] H. Ju, J. Xu, Appl. Surf. Sci. 355 (2015) 878. [7] H. Ju, J. Xu, Mater. Charact. 107 (2015) 411e418. [8] H. Ju, J. Xu, Surf. Coat. Technol. 283 (2015) 311e317.
[9] S.M. Aouadi, D.P. Singh, D.S. Stone, K. Polychronopoulou, F. Nahif, C. Rebholz, C. Muratore, A.A. Voevodin, Acta Mater. 58 (2010) 5326. [10] D.S. Stone, J. Migas, A. Martini, T. Smith, C. Muratore, A.A. Voevodin, S.M. Aouadi, Surf. Coat. Technol. 206 (2012) 4316. [11] D.S. stone, S. Harbin, H. Mohseni, J.E. Mogonye, T.W. Scharf, C. Muratore, A.A. Voevodin, A. Martini, S.M. Aouadi, Surf. Coat. Technol. 217 (2013) 140e146. [12] W. Gulbinski, T. Suszko, Surf. Coat. Technol. 201 (2006) 1469. [13] P.J. Kelly, H. Li, P.S. Benson, K.A. Whitehead, J. Verran, R.D. Arnell, I. Iordanova, Surf. Coat. Technol. 205 (2010) 1606e1610. [14] H. Kostenbauer, G.A. Fontalvo, G. Mitterer, J. Keckes, Tribol. Lett. 30 (2008) 53e60. [15] D. Du, D. Liu, X. Zhang, J. Tang, D. Xiang, Appl. Surf. Sci. 365 (2016) 47e56. [16] G. Janssen, M. Abdalla, F. Keulen, B. Pujada, B. Venrooy, Thin Solid Films 517 (2009) 1858e1867. [17] S. Zhang, X. Zhang, Thin Solid Films 520 (2012) 2375e2389. [18] S. Yao, Y. Su, W. Kao, K. Chen, Surf. Coat. Technol. 201 (2006) 2520. [19] C.C. Tseng, J.H. Hsieh, S.C. Jang, Y.Y. Chang, W. Wu, Thin Solid Films 517 (2009) 4970e4974. [20] K. Kitawaki, K. Kaneko, K. Inoke, J.C. Hernandez-Garrido, P.A. Midgley, Micro 40 (2009) 308e312. [21] C.P. Mulligan, D. Gall, Surf. Coat. Technol. 200 (2005) 1495. [22] S.M. Aouadi, P. Basnyat, Y. Zhang, Q. Ge, P. Filip, Appl. Phys. Lett. 88 (2006) 219. [23] B.R. La, A.G. Evans, J. Am. Ceram. Soc. 59 (7e8) (1976) 371. [24] M.T. Laugier, J. Mater. Sci. 21 (1986) 2269. [25] S.J. Bull, D.S. Rickerby, Surf. Coat. Technol. 42 (1990) 149. [26] S. Chang, Y. Huang, Microelectron. Eng. 84 (2007) 319. [27] S. Chang, Y. Lee, C. Lu, J. Electrochem. Soc. 154 (2007) 241. [28] Y. Huang, S. Chang, C. Chang, Thin Solid Films 517 (2009) 4857. [29] Q. Wang, F. Zhou, J. Yan, Surf. Coat. Technol. 285 (2016) 203.
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Vacuum journal homepage: www.elsevier.com/locate/vacuum
Microstructure, mechanical and trobological properties of TiN-Ag films deposited by reactive magnetron sputtering Hongbo Ju, Lihua Yu, Dian Yu, Isaac Asempah, Junhua Xu* School of Materials Science and Engineering, Jiangsu University of Science and Technology, Mengxi Road 2, Zhenjiang, Jiangsu Province, 212003, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 January 2017 Received in revised form 21 March 2017 Accepted 21 March 2017 Available online 23 March 2017
Composite TiN-Ag films with various Ag content (Ag/(Ti þ Ag)) were deposited using reactive magnetron sputtering and the influence of Ag content on the crystal structure, mechanical and tribological properties were investigated. The result showed that face-centered cubic (fcc) TiN and fcc-Ag co-existed in the films and TiN had a columnar-growth and the Ag nanoparticles embedded in the boundary of columnar crystal. The hardness of TiN-Ag films initially increased gradually and reached an optimum value, and then decreased with an increase in Ag content in the films. The maximum hardness value was 29 GPa at 0.8 at.% Ag. The addition of Ag into TiN film could enhance the fracture toughness (KIC) and critical load (LC1) of the films because the nanoparticle Ag provides a large volume fraction of grain boundaries. As Ag was added into TiN film, the average friction coefficient decreased from 0.78 at 0 at.% Ag to 0.20 at 41.1 at.% Ag; however, the wear rate of TiN-Ag films initially decreased and then increased, after reaching the minimum value of about 1.3 107 mm3/(mm.N), at 0.8 at.% Ag. © 2017 Elsevier Ltd. All rights reserved.
Keywords: Reactive magnetron sputtering TiN-Ag films Microstructure Mechanical and tribological properties
1. Introduction Advances in machining technology in recent times have led to an increase in demand for materials for severe applications associated with high temperature and the stresses that are generated at the tool/work piece interface [1]. Nano-structured films, which usually exhibit a unique blend of hardness, thermal stability, wear resistance and low average friction coefficient, have gained increasing interest in recent years for a number of different applications. Transition metal nitride (TMN) films deposited by physical vapor deposition techniques, which are widely used for wear reduction on the cutting and die tools, have aroused considerable interest in the past few decades because of some of their excellent properties [1e8]. As modern manufacturing industry develop at an unprecedented rate, its further demands such as relative lower average friction coefficient are driving force for further development of TMN matrix films [3,8]. Seeking effective lubricants for transition metal nitride films is a challenging task for the tribology community [6]. Silver considered as a traditional lubricating material was combined into TMN matrix recently and a large number of related
studies has been carried out [6,9e13]. TiN-Ag film is one of the most representative self-lubriacting films. For example, H. Kostenbauer et al. [14] studied a series of TiN-Ag films' tribological properties at elevated testing temperature systematically and found that TiN films doped Ag could adjust their surface chemistry according to the changing of the service condition. D. Du et al. [15] reported that the fretting fatigue resistance of a Ti-6Al-4V alloy could be improved significantly by deposition of TiN-Ag film on the surface of the alloy. It is widely accepted that metallic Ag phase and TiN phase co-exist in the TiN-Ag films with higher Ag content. However, as the Ag content in the films is very low (eg. <2 at.% Ag), it is difficult to study what state the Ag exists in the TiN films, since only TiN peaks appear in the XRD pattern of the films. In our paper, the most representative TMN filmdTiN film was opted as the object and Ag element was introduced into TiN film to create the TiN-Ag composite films by reactive magnetron sputtering. High resolution transmission electron microscopy (HRTEM) and back scattered electron (BSE) were performed to investigate the crystal structure of TiN-Ag films. Besides this, the effects of Ag content on the mechanical (hardness and fracture toughness) and tribological properties of TiN-Ag films were also discussed in this paper. 2. Experimental details
* Corresponding author. E-mail address: [email protected] (J. Xu). http://dx.doi.org/10.1016/j.vacuum.2017.03.026 0042-207X/© 2017 Elsevier Ltd. All rights reserved.
TiN-Ag films with a thickness of about 2 mm were deposited on
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mirror polished stainless steel (AISI 304 SS, 192 Hv, 15 mm 15mm 2.5 mm) with surface roughness of ~70 nm and silicon (100) wafer substrates using a multi-target magnetron sputtering system. Ti (99.9%) and Ag (99.9%) targets with a diameter of 75 mm were sputtered by two 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 6 104 Pa, Ar (99.999%) and N2 (99.999%) were introduced into the chamber by means of two separate gas manifolds. The TiN-Ag films with different Ag content were achieved by fixing the powers of Ti target at 280 W and adjusting the power of Ag target from 0 W to 150 W while constantly keeping the working pressure at 0.3 Pa, the deposition temperature at 200 C and the nitrogen to argon ratio (flow ratio) at 10:5. The substrates were not biased. After the deposition, the deposition rate was calculated and its value was ~17 nm/min. The TiN-Ag films deposited on the AISI 304 SS substrates were used to test for adhesion and wear test of the film, while film's elemental composition, crystal structure, hardness, elastic modulus and fracture toughness were evaluated using the film on the silicon (100) wafer substrates. The elemental compositions of the films were determined by energy dispersive spectroscopy (EDS) on an EDAX DX-4 energy dispersive analyzer. The microstructure of the films was evaluated using a Siemens X-ray diffractometer using Cu Ka radiation, operated at 40 kV, 35 mA. The grain size of the films was calculated by using Debye-Scherrer's formula for most intense peak having orientation (111) in the face-centered cubic (fcc) structure. The back scattered electron (BSE) image was obtained by the scanning electron microscopy (SEM, JEM-6480). High resolution transmission electron microscopy (HRTEM) was performed using a JEOL JEM-2010F microscope operated at an accelerating voltage of 200 kV. The residual stress (s) of the films was calculated by Stoney's equation [16].
KIC
! 0:5 E P ¼a 1:5 H Cm
83
(2)
where E is the elastic modulus of the film, H is the hardness of the film, P is 200 mN in our paper, Cm is the crack length, a is the empirical constant and its value is 0.016 in our paper. The adhesion of the films was measured by a scratch tester, which was equipped with a diamond Rockwell tip. The load was increased linearly from 0 N up to 40 N. The critical loads (Lc1) and complete delamination load (LC2) were measured by analyzing the failures events in the scratch track by optical microscopy. A 30 min wear test was carried out along a circular track of 8 mm 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 25e30%) at room temperature (about 25 C) using a UTM-2 CETR tribometer. Raman spectroscopy using the 514.5 nm Arþ laser with a backscattering optical configuration was used to study the tribo-film on the surface of wear track. 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:
W¼
V SL
(3)
where S is the total sliding distance; L is the applied load.
3. Results and discussion
E ts2 s¼ 1 y 6tf
1 1 R Rs
3.1. Crystal structure
(1)
where E–elastic modulus of the substrate (E ¼ 170 GPa).
y–Poisson's ratio of the substrate (y¼0.3)
ts–the thickness of the substrate tf–the thickness of the film R–the substrate curvature radii of the Si wafer Rs–the curvature radii of the film
The curvature radii were done using data obtained from Bruker DEKTAK-XT profilometer. Five (5) areas of each sample were made for the curvature radii and the mean value taken. In determining the hardness and elastic modulus of the films, an indenter calibrated using fused silica as reference was used in a nanoindentation tester system (nano-indenter CPX þ NHT þ MST), equipped with a diamond Berkovich indenter tip. The maximum load of 3 mN was used and a minimum of nine (9) indentations were made for each sample and the mean value taken. Before making indentation on the films, the indenter was calibrated with respect to a reference sample of fused silica. The fracture toughness (KIC) of the films was also measured by the nano-indentation tester at 200 mN and the fracture toughness was calculated by the following equation [17]:
Fig. 1 shows the XRD patterns of TiN-Ag films with various Ag content (Ag/(Ti þ Ag), at.%). As shown in Fig. 1, the diffraction of binary TiN film shows multiple orientations of (111), (200) and (222) crystal plane of face-centered cubic (fcc) TiN phase (JCPDF card 65e0715). For each XRD pattern of ternary TiN-Ag films at<1.4 at.% Ag, it shows three peaks corresponding to fcc-TiN phase. Increasing Ag content furtherly, besides the fcc-TiN diffraction peaks, two peaks with low-intensity at ~39 and ~45 are detected. The two peaks correspond to fcc-Ag. The solid solubility of Ag in the transition metal films is very limited based on the result in Refs. [6,18], therefore, for the films at<1.4 at.% Ag, no peak corresponding to fcc-Ag is attributed to the limited precision of the XRD. The grain size of the films was calculated using Debye-Scherrer's formula for most intense peak having orientation (111) in the fccTiN. The addition of Ag drops the grain size form ~35 nm at 0 at.% Ag to ~7 nm at 41.1 at.% Ag. The nanoparticle Ag embedded in the TiN films could prevent grain growth and similar results were also reported in Refs. [14,19]. The resolution transmission electron microscopy (TEM) and its selected area electron diffraction (SAED) pattern of TiN-Ag film at 0.7 at.% Ag were carried out to study the crystal structure of TiN-Ag films at<1.4 at.% Ag. The result is shown in Fig. 2. The cross-section TEM image of TiN-Ag film at 0.7 at.% Ag (Fig. 2(a)) reveals that the film has a columnar-growth and dense structure. Its corresponding
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3.2. Mechanical and tribological properties
Fig. 1. XRD patterns of TiN-Ag films with various Ag content.
SAED pattern is illustrated in Fig. 2(b), from inside to outside, the diffraction rings can be assigned to the lattice planes of fcc-TiN (111), fcc-Ag (111), fcc-TiN (200), fcc-TiN (220) and fcc-TiN (311). It suggests that fcc-TiN and fcc-Ag coexist in the film. Fig. 2(c) shows the HRTEM image of inner region of columnar crystal, a lattice fringe with a lattice spacing of about 0.2445 nm is detected. This lattice fringe belongs to the fcc-TiN (111), since the value of lattice spacing of fcc-TiN (111) (JCPDF card 65e0715) is 0.2448 nm. The HRTEM image of boundary of columnar crystal is shown in Fig. 2(d), Ag nanoparticle with a radius of about 5 nm is detected in the field view besides TiN lattice fringe. From Fig. 2(d), it can also be observed that the Ag nanoparticle seems to be randomly oriented with respect to the TiN lattice fringe. This suggests that Ag nanoparticle is not chemically bonded on the boundary of columnar crystal TiN, such as metallic Ag-Ti. Similar results were also reported in Refs. [6,19,20]. Fig. 3 illustrates the back scattered electron (BSE) image of TiNAg film at 23.5 at.% Ag. The grey cylinders in the image are TiN and the small white particles are Ag. TiN exhibits a columnar-growth and Ag particles evenly dispersed along the boundary of TiN columnar grain. Based on above results, TiN has a columnar-growth, dense structure and Ag nanoparticles embedded in boundary of the crystal.
Table 1 illustrates the residual stress of TiN-Ag films. All films regardless of Ag content are in compressive state and the compressive residual stress of the films drops gradually with the rise of Ag content in the films. The drop of the compressive residual stress is attributed to the inhibition of titanium nitride grain by Ag addition into TiN matrix. The hardness (H) and elastic modulus (E) of TiN-Ag films with various Ag are shown in Fig. 4. As shown in Fig. 4, the hardness and elastic modulus of binary TiN film are 21 GPa and 320 GPa, respectively. The addition of Ag into TiN film rises the hardness from 23 GPa at 0.7 at.% Ag to 29 GPa at 0.8 at.% Ag, and then drops gradually with a further rise of Ag content. The elastic modulus of TiN-Ag films drops gradually with an increase in Ag content. As the Ag content increases from 0 at.% to 0.8 at.%, the hardness enhancement could be attributed to the effect of fine-grain strengthening. With a further increase in Ag content, the decrease of hardness of the films is attributed to the increasing in the very soft and highly mobile silver in the film [6,21,22]. The decrease of the compressive residual stress also drops the hardness of the films at>0.8 at.% Ag. In addition, the elastic modulus of Ag film that we deposited under the same experimental conditions is only about 70 GPa, so the decrease in elastic modulus of TiN-Ag is most probably associated with the incorporation of Ag in the TiN matrix. 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 [7]. It is considered that the hard surfaces can resist abrasive wear and low elastic modulus allows elastic deformation to easy occur when the contact stress is applied. So the H/E ratio is proposed as a key parameter controlling the wear resistance. In addition, the plastic deformation resistance factor (H3/E2) also can be related with the wear rate [7]. The H/E and H3/E2 ratios of TiN-Ag films are shown in Fig. 5. As shown in Fig. 5, the ratios of H/E and H3/E2 of the films rise gradually from 0.066 to 0.090 GPa at 0 at.% Ag to 0.097 and 0.27 GPa at 0.8 at.% Ag, and then drop gradually with a further rise of Ag content. Fracture toughness (KIC) is another vital mechanical property of materials. The radial cracking indentation method is widely accepted to calculate KIC of thin films [23]. A maximum load of 200 mN was chosen to investigate KIC in our paper, since the visible cracks is hard to induced by small load. Although the displacement caused by the 200 mN is greater than ten percent (10%) of the film thickness, the value of KIC obtained from the 200 mN load could reflect the exact value of the film in some sense because all TiN-Ag films have same interlayer and substrate. The indentation morphology of TiN-Ag films with various Ag content is shown in Fig. 6. As shown in Fig. 6(a), the film at 0.7 at.% Ag exhibits the obvious radial cracks paralleled to the indentation diagonal. Rising the Ag content to 23.5 at.% enhances the fracture toughness considerably, so no radial cracks is detected after the 200 mN indentation, as shown in Fig. 6(b). KIC of TiN-Ag films are shown in Table 1 and its value rises gradually from 0.4 MPa m1/2 at 0 at.% Ag to 1.6 MPa m1/2 at 1.4 at.% Ag. Rising Ag content furtherly, KIC could not be calculated due to the disappearance of radial cracks. The nanoparticle Ag provides a large volume fraction of grain boundaries, this limits initial crack sizes and helps to deflect and terminates the growth of cracks. Therefore, the addition of Ag into TiN film enhances KIc. Fig. 7 illustrates the scratch track optical images of TiN-Ag films with various Ag content after the scratch test. As shown in Fig. 7, adding Ag into TiN film induces an obvious change of the deformation behavior. LC1 is considered as the critical load of the film and its value is obtained from the crack initiation. LC2 is obtained from
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85
Fig. 2. TEM images of TiN-Ag film at 0.7 at.% Ag: (a) The cross-section TEM image of TiN-Ag film, (b) SAED pattern of TiN-Ag film, (c) HRTEM image of inner region of columnar crystal, (d) HRTEM image of boundary of columnar crystal.
Fig. 4. Hardness and elastic modulus of TiN-Ag films with various Ag content.
Fig. 3. The back scattered electron image of TiN-Ag film at 23.5 at.% Ag.
the occurrence of complete delamination. Table 1 shows the value of LC1 and LC2. The value of LC1 rises gradually from ~4 N at 0 at.% Ag
to ~12 N at 1.4 at.% Ag, the failure of the film is detected in the brittle behavior, and then as the Ag content rises furtherly, the deformation behavior is changed and scratch of the film at 23.5 at.% Ag exhibits no obvious cracks because of the rise of high plasticity and ductility metal Ag in the film. LC1 is attributed to the fracture
Table 1 Compressive residual stress, fracture toughness (KIC), critical load (LC1), complete delamination load (LC2) and surface energy of interfacial crack of TiN-Ag films with various Ag content. Ag content(at.%)
Compressive residual stress (GPa)
KIC (MPa.m1/2)
LC1 (N)
LC2 (N)
Gc (J/m2)
0 0.7 0.8 1.4 23.5 41.1
þ0.24 þ0.20 þ0.20 þ0.19 þ0.13 þ0.08
0.4 0.5 0.9 1.6 No cracks No cracks
4 5 9 12 No cracks No cracks
12 18 21 26 23 17
170 220 290 525 370 267
± ± ± ± ± ±
0.01 0.01 0.01 0.01 0.01 0.01
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Gc ¼
d2c t 2Ef
(4)
Where t is the thickness of the film, Ef is the elastic modulus of the film, dc is the film critical stress of complete delamination and it can be calculated by the following equation:
dc ¼ Fig. 5. The H/E and H3/E2 ratios of TiN-Ag films with various Ag content.
3 !2 2Lc2 4 4 þ vf 3pm 1 2vf 5 8 pd2c
(5)
Where dc is the width of the scratch track obtained from the occurrence of complete delamination,
m is the friction coefficient from the scratch tester, vf is the Pisson rate of the film.
Fig. 6. Indentation morphology of TiN-Ag films at 200 mN: (a) 0.7 at.% Ag and (b) 23.5 at.% Ag.
toughness of the films. The value of LC2 rises from ~12 N at 0 at.% Ag to 26 N at 1.4 at.% Ag, and then it drops gradually with a further increase in Ag content. Dropping of H/E ratio might be attributes to the reduction of LC2. According to Ref. [24,25], the relationship between LC2 and surface energy of the interfacial crack (Gc) could be established by Griffith energy balance method and the equation is as follows:
Table 1 also illustrates the Gc of the TiN-Ag films with various Ag content. As shown in Table 1, The value of Gc rises from 170 J/m2 N at 0 at.% Ag to 525 J/m2 at 1.4 at.% Ag, and then it drops gradually with a further increase in Ag content. Therefore, the drop of LC2 might be attributed to the decrease of Gc of the films. Similar results were also in Refs. [26e29]. Fig. 8 illustrates a few examples of friction coefficient curves at room temperature (RT). As shown in Fig. 8, the friction curve of binary TiN film fluctuates markedly when the testing time is below 380 s, and then stays constant at about 0.79 for the remaining test times. For the TiN-Ag film at 0.8 at.% Ag, the friction coefficient had a steady value of about 0.62 following a dramatic move, according to the friction curve. With a further increase in Ag content, the friction coefficient of the film does not fluctuate and a lowest steady value of about 0.20 is achieved for the TiN-Ag film at 41.1 at.% Ag. The average friction coefficient (m) of TiN-Ag films, which is calculated from the friction coefficient data, is illustrated in Fig. 9. The m of binary TiN film is about 0.78. The m of TiN-Ag, all of which are lower than that of binary TiN film, is dependent on Ag content in the film. It decreases from about 0.75 at 0.7 at.% Ag to about 0.20 at 41.1 at.% Ag. This result clearly demonstrates the lubrication effect of Ag in TiN. A similar result was also reported in other Ref. [6,14]. Fig. 10 illustrates the wear rate of TiN-Ag films with various Ag content. As illustrated in Fig. 8, the wear rate of binary TiN film is about 1.73 106 mm3/(mm.N). The wear rate of TiN-Ag films initially decreases and then increases, after reaching the minimum value of about 1.3 107 mm3/(mm.N), at 0.8 at.% Ag. As Ag content increases from 0 at.% to 0.8 at.%, Ag has good lubricity property and could play a role lubrication. Therefore, the m and wear rate of the films decrease with the increasing Ag content. Besides this, the increasing ratios of H/E and H3/E2 also attributes to the decrease in the wear rate. With a further increase in Ag content, the m further decreases due to the increase in lubricant Ag content in the films. However, the soft lubricant Ag can be easily wiped off during the wear test. In addition, the ratios of H/E and H3/E2 decrease with the increase of Ag content. So the wear rate increases with the increase in Ag content. 4. Conclusion TiN-Ag films with various Ag content were deposited by reactive magnetron sputtering and the effects of Ag content on the crystal structure, mechanical and tribological properties of TiN-Ag films were investigated. The main results could be concluded as follows:
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Fig. 7. Optical images of scratch track of TiN-Ag films with various Ag content: (a) 0.7 at.% Ag (b) 1.4 at.% Ag and (c) 23.5 at.% Ag.
Fig. 9. Average friction coefficient of TiN-Ag films with various Ag content at room temperature.
Fig. 8. Friction coefficient curves of TiN-Ag films with various Ag content at room temperature.
(1) A two-phase structure, consisting of fcc-TiN and fcc-Ag, coexisted in the TiN-Ag films. TiN film had a columnar-growth and dense structure and Ag nanoparticles were embedded in boundary of columnar crystal. (2) The hardness of TiN-Ag films initially increased gradually and reached an optimum value, and then decreased with an increase in Ag content in the films and the maximum value was 29 GPa at 0.8 at.% Ag. The hardness enhancement could be attributed to the effect of fine-grain strengthening. (3) The addition of Ag into TiN film enhanced the fracture toughness (KIC) and critical load (LC1) of the films because the nanoparticle Ag in the film provided a large volume fraction of grain boundaries.
Fig. 10. Wear rate of TiN-Ag films with various Ag content at room temperature.
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(4) As Ag was added into TiN film, the average friction coefficient decreased from 0.78 at 0 at.% Ag to 0.20 at 41.1 at.% Ag; however, the wear rate of TiN-Ag films initially decreased and then increased, after reaching the minimum value of about 1.3 107 mm3/(mm.N), at 0.8 at.% Ag. The tribological properties of the films were influenced by the lubricating Ag phase and the ratios of H/E and H3/E2. Acknowledgements Supported by National Natural Science Foundation of China (51074080, 51374115, 51574131), Research Fund of Jiangsu University of Science and Technology (1062931609). The authors also would like to thank Jing Luan for her valuable help in improving our manuscript. References [1] L. Hultman, Vacuum 57 (2000) 1e30. [2] Z.B. Qi, F.P. Zhu, Z.T. Wu, B. Liu, Z.C. Wang, D.L. Peng, C.H. Wu, Surf. Coat. Technol. 231 (2013) 102. [3] J. Xu, H. Ju, L. Yu, Acta Metall. Sin. 48 (2012) 1132e1138. [4] J. Xu, H. Ju, L. Yu, Vacuum 103 (2014) 21e27. [5] J. Xu, H. Ju, L. Yu, Vacuum 110 (2014) 47e53. [6] H. Ju, J. Xu, Appl. Surf. Sci. 355 (2015) 878. [7] H. Ju, J. Xu, Mater. Charact. 107 (2015) 411e418. [8] H. Ju, J. Xu, Surf. Coat. Technol. 283 (2015) 311e317.
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