- Email: [email protected]
Heat treatment of nanocrystalline TiN films deposited by unbalanced magnetron sputtering
Surface & Coatings Technology 200 (2006) 4291 – 4299 www.elsevier.com/locate/surfcoat
Heat treatment of nanocrystalline TiN films deposited by unbalanced magnetron sputtering Jia-Hong Huanga,T, Kae-Jy Yua, P. Sitb, Ge-Ping Yua a
Department of Engineering and System Science, National Tsing Hua University, Hsinchu 300, Taiwan Department of Physics and Materials Science, City University of Hong Kong, Kowloon, Hong Kong
b
Received 12 October 2004; accepted in revised form 5 February 2005 Available online 11 April 2005
Abstract TiN thin films were deposited on Si wafers using an unbalanced magnetron (UBM) sputtering technique. Heat treatment was implemented after deposition to improve the properties of the TiN thin films. Combinations of mixing gas atmospheres and Ti getter were evaluated by thermodynamic and kinetic estimations to find a heat treatment environment with minimum oxidation. Results showed that at 700 8C the oxidation of TiN thin film could be minimized under an atmosphere with a partial pressure ratio of Ar/H2 = 9 and attached with a Ti getter. The specimens were heat-treated at 500 and 700 8C for 1 h under Ar/H2 atmosphere. Heat treatment showed significant effect on the microstructure, crystallinity, texture, grain size, roughness, and residual stress of the nanocrystalline TiN films, while less distinct effect was found on N/Ti ratio, packing factor and electrical resistivity. In addition, the changes in structure and mechanical properties were more effective at 700 8C than those at 500 8C. The changes in texture, grain size, and crystallinity were more obvious in the specimens with a thickness less than 214 nm. The residual stress and hardness were distinctly decreased by the heat treatment especially at 700 8C. This could be attributed to the reduction of lattice defects by the heat treatment. D 2005 Elsevier B.V. All rights reserved. Keywords: Heat treatment; TiN; Unbalanced magnetron sputtering; Residual stress
1. Introduction Owing to its high hardness, wear resistance, and thermal stability, titanium nitride (TiN) thin film is commercially used in the tool industry as a protective film to prolong the service life of substrate materials. The progress of microelectromechanical systems (MEMS) is driving the demand for ultra-thin (less than 0.5 Am) protective and functional nitride coatings. Our previous studies found that as the film thickness of TiN is less than 400 nm, the film properties may become inferior due to less compact structure [1,2]. To meet the requirements of the ultra-thin coatings in MEMS, especially less than 0.5 Am, the film properties must be enhanced not from increasing thickness but through the
T Corresponding author. Tel.: +886 35715131x4274; fax: +886 35720724. E-mail address: [email protected] (J.-H. Huang). 0257-8972/$ - see front matter D 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2005.02.147
basic improvement of the film structure, which has become a new challenge to coating and thin film technology. Besides the optimization of the coating processes, heat treatment after deposition has been widely used in industry to improve the properties of hard coatings by reducing the residual stress and defects in the deposited film [3–6]. However, earlier studies have reported that one of the factors that reduce the hardness of the TiN film is the surface oxidation during heat-treating process [7,8]. Recently, to overcome the oxidation issue, heat treatments proceeding in the controlled atmosphere have been reported [9–11]. Chou et al. [10] found that for the TiN films deposited by ion-plating, both residual stress and hardness are decreased after heat treatment at 700 8C in Ar/H2 atmosphere, while the microstructure and packing factor inside the TiN films are not significantly changed. Mayrhofer et al. [11] annealed TiN films deposited by DC unbalanced magnetron sputtering at 700 8C in vacuum. They also observed the decrease in both residual stress and
4292
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
hardness of the films after heat treatment. In contrast, they reported that grain size increases at 700 8C in the nanocrystalline TiN films. Since titanium oxides are thermodynamically more stable than TiN at high temperature, which implies that annealing is usually accompanied with the oxidation of TiN, it is very difficult to totally suppress the oxidation reaction by reducing oxygen partial pressure even down to 10 20 Pa at 700 8C. Therefore, to understand the effect of heat treatment on the structure and properties of TiN film, it is necessary to search a proper environment with least oxidation for avoiding the complication of oxidation in the heat treating process. Recent interest in nanotechnology stimulates research in the nanocrystalline nitride thin films. One of the attracting characteristics is that the nanocrystalline (grain size b 100 nm) materials may exhibit very different properties from micro-crystalline (grain size N 100 nm) materials. Huang et al. [12,13] found that unlike normal micro-crystalline TiN films, the hardness of nanocrystalline TiN or ZrN thin films does not significantly vary with (111) texture coefficient. The texture-insensitive behavior was attributed to the fact that the deformation mechanisms of the nanocrystalline nitride films are possibly from grain rotation or grain boundary sliding, instead of a dislocation slip mechanism. For the nanocrystalline thin film, due to very large ratio of atoms on grain boundary to atoms in grain, the thermally activated processes may be greatly enhanced by short circuit diffusion through grain boundaries. This effect is expected to be pronounced in the heat treatment process. In fact, Mayrhofer et al. [11] found that grain growth by subgrain growth and grain boundary migration occurs at 700 8C in the TiN films with nanograin size (20–30 nm). Although many studies [9–11,14] have been performed on the heat treatment of nitride coatings in controlled atmospheres, there is little information available on the heat treatment of nanocrystalline (grain size b 20 nm) thin films (thickness b 400 nm). Using a DC unbalanced magnetron (UBM) sputtering system, we successfully deposited nanocrystalline TiN thin films on (100) silicon wafers in an earlier research [12]. To continue the study, the UBM sputtering method was chosen to deposit TiN films on silicon substrates, and the deposition conditions were slightly changed to produce the specimens suitable for the study of heat treatment. The purposes of this study were to find a heat treatment environment with minor oxidation, and to investigate the effect of heat treatment in the controlled atmosphere on the microstructure and mechanical properties of TiN films.
2. Experimental procedures The deposition of TiN films was carried out using a UBM sputtering system. The substrate material was Si (100) wafer with dimensions of 45 mm 45 mm 0.7 mm. Prior to the coating process, the Si specimens were
undergone ultrasonic cleaning progressively in acetone and ethanol and then dried in a vacuum dryer. Before deposition, the substrates were gradually preheated to a temperature of 300 8C. Meanwhile, the chamber was evacuated to 8 10 4 Pa (6 10 6 Torr) to avoid contamination during deposition process. Prior to deposition, the substrate was pre-sputtered by argon discharge at a bias of 1000 V for 10 minutes to remove the surface oxide layer. The argon pressure was fixed at 0.8 Pa (6 10 3 Torr) and the current density at the substrate was 0.15 mA/cm2 during pre-sputtering. The target-to-substrate distance was 20 cm. After pre-sputtering, high purity working gas and reactive gas were introduced, using mass flow controllers to regulate both gas flows. Argon (99.9995% in purity) gas and nitrogen gas (99.9995% in purity) flow rates were fixed at 15 sccm and 1 sccm, respectively. The total gas pressure was controlled at 0.133 Pa (1 10 3 Torr). The DC power supply of the Ti target was operated at 0.9 A using constant-current mode; in other words, the target current density was 2.8 mA/cm2. During deposition, a negative substrate bias voltage of 50 V was applied and the deposition temperature was maintained at 350 8C, monitored using a thermocouple near the substrate. Film thickness was controlled by varying deposition duration. Four sets of specimens with different film thickness, denoted as A0(86 nm), B0(214 nm), C0(364 nm), and D0(430 nm), were deposited for further heat treatment. A LINDBERG high-temperature tube furnace equipped with a mechanical pumping system was used to perform the heat treatment. After inserting the specimens, the quartz tube was sealed and pumped down to 1.33 Pa. Then, the tube was purged with working gas for several times and the final pressure was controlled at 10 Pa by flowing high purity mixing gas, including gases with partial pressure ratios of Ar/H2 = 9 and N2/H2 = 9. The oxygen partial pressure was monitored using a zirconia oxygen gauge (15% CaO-doped ZrO2) in the gas mixture. In addition to the controlled atmosphere, Si wafers with pure titanium coatings were used as the oxygen getter. The getters were placed closely to the samples for preventing the uptake of oxygen by the TiN films. The specimens were heat treated at 500 and 700 8C for 1 h. In the duration of heat treatment, the oxygen partial pressures were maintained at 5.05 10 21 Pa and 1.01 10 17 Pa at 500 and 700 8C, respectively. The crystal structure of the TiN films was characterized by X-ray diffraction (XRD). The Cu Ka line at 0.15405 nm was used as the source for diffraction. The extent of (111) preferred orientation is quantified by a texture coefficient (TC) defined as I(111) / [I(111) + I(200) + I(220)], where I is the integrated intensity of the corresponding Bragg peak. The (111) texture coefficient of the TiN powder sample, from JCPDS 38-1420, is 0.33. The grain size of the TiN thin films was estimated from the results of h/2h scans. The position of the (111)
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
diffraction peak and the full width of the peak at half maximum (FWHM) were used to estimate the grain size by Scherrer’s equation [15] t¼
0:9k B cosh
where k is the wavelength of Cu Ka, B is the FWHM of the diffraction peak, h is the peak position, and t is the grain size of the thin film. The residual stress of the TiN films was determined by modified XRD sin2w method [16,17] using a 4-circle diffractometer with psi-goniometer geometry. X-ray was incident at an angle of 28 to increase the diffraction volume of the thin film specimen. The columnar structure and the film thickness of the deposited TiN film were observed and measured from the crosssectional specimens using a field-emission gun scanning electron microscope (FE-SEM) operated at 15 keV. The N/Ti ratio and packing factor were obtained from the results of Rutherford backscattering spectroscopy (RBS). The calculation of the packing factor from the RBS data was described elsewhere [1]. To access the oxygen influence, the compositional depth profiles were determined using Auger electron spectroscopy (AES). Since large substrate effect may occur for the TiN film with a thickness less than 500 nm when using normal microhardness tester, the film hardness was measured using a Hysitron nanoindenter attached on a Digital atomic force microscope (AFM). The applying loads were dependent on the depth displacement that should be less than one tenth of the film thickness. For the specimens with a thickness larger than 300 nm (series C and D), the indentation depth was maintained at 30 nm; for those smaller than 214 nm (series A and B), one tenth of the film thickness was adopted as the indentation depth. The procedures of applying load and calibration of the nanoindenter were detailed in our previous paper [1]. The hardness was calculated using Oliver–Pharr technique [18], which could be executed from the software provided by Hysitron Inc. Ten indentations were made and the average value was reported. The precision of each value was within 10%. The surface roughness of the TiN films were determined by AFM and the root mean square roughness value (R rms) was reported. The electrical resistivity of the TiN films was measured using a four-point probe.
3. Results and discussion 3.1. Thermodynamic and kinetic estimations of the annealing environment A typical oxygen partial pressure along with furnace temperature with respect to time under the atmosphere containing hydrogen is illustrated in Fig. 1. It can be seen that the oxygen content in the system could be lowered down to 1.01 10 17 Pa at 700 8C using the mixing gas
4293
Temperature Celsius 700.00
Oxygen atm 1E-20 1E-21 1E-22
600.00
1E-23 1E-24 1E-25
500.00 1E-26 1E-27 1E-28
400.00
1E-29 1E-30 300.00
1E-31 1E-32 1E-33
200.00 1E-34 1E-35 1E-36
100.00
1E-37 1E-38 0.00
1E-39 0.00
4000.00
8000.00
12000.00
16000.00
Time (sec)
Fig. 1. The typical oxygen partial pressure along with furnace temperature with respect to time under the atmosphere containing hydrogen.
environment. In addition, since the reaction of hydrogen and oxygen is exothermic, according to the Le Chartelier’s principle, the oxygen partial pressure is increased with increasing temperature. Samples with the thinnest TiN coatings (A-series) were used to investigate the composition variation under different atmospheres at 700 8C. Fig. 2 shows the AES results for the specimens treated at 700 8C under different environments. Although the N/Ti ratio cannot be determined by the AES analysis, the depth profiles can reveal the distribution of O, Si, C, and (Ti + N). As shown in Fig. 2(a), the compositions are uniformly distributed in the as deposited sample A0. Fig. 2(b) shows the profiles of the sample A700 heat-treated at 700 8C in Ar/H2 atmosphere attached with Ti getter. Apparently, there is no significant change in compositions before and after annealing. In contrast, for the specimen directly exposing to hydrogen mixing gas without Ti getter, more oxide (~ 24 nm) is formed on the surface of the TiN coating, as shown in Fig. 2(c). For the specimen heat treated in air (Fig. 2(d)), the TiN film has been thoroughly penetrated by oxygen. The results suggest that Ti getter plays an important role in preventing oxidation of the TiN thin film. Without Ti getter, oxidation occurs even at a controlled atmosphere with extremely low oxygen content.
4294
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
(a)
Ti N C O Si
N 6
2.0x10
Counts
Counts
1.6x10
Ti
6
1.2x10
2.0x10
6
1.6x10
6
1.2x10
6
8.0x10
5
4.0x10
5
N Si
Si
6
5
8.0x10
Ti
O
O 5
4.0x10
C
C
0.0
0.0 0
200
Ti N C O Si
(b)
400
600
800
0
1000
200
Sputtering Time (s)
400
600
800
(c)
Ti N C O Si
N 6
2.0x10
(d) Ti
Ti N C O Si
6
2.0x10
Si
O Si
6
1.6x10
1.6x106
6
Ti
1.2x10
Counts
Counts
1000
Sputtering Time (s)
5
8.0x10
1.2x106
N Ti
8.0x105
O 5
4.0x105
4.0x10
C
C 0.0
0.0 0
200
400
600
800
1000
1200
Sputtering Time (s)
0
200
400
600
800
1000
1200
1400
Sputtering Time (s)
Fig. 2. AES results of TiN thin films under different annealing environments. (a) As deposited (specimen A0); (b) annealed at 700 8C in Ar/H2 = 9 atmosphere attached with Ti getter (specimen A700); (c) annealed at 700 8C in N2/H2 = 9 atmosphere without Ti getter (specimen A2); and (d) annealed at 700 8C in air (Specimen A1).
The oxidation reaction of TiN involved in the heat treatment process can be understood by the thermodynamic and kinetic estimations. The main reaction of the oxidation of TiN is TiN þ O2 YTiO2 þ 1=2N2 DG ¼ RT lnK ¼ 604672 þ 84:308T ðJ=molÞ [19] where G is the Gibbs free energy, T is the absolute temperature, R is the gas constant (= 8.314 J/K mol), and K is constant. The reaction constant K= the reaction 1=2 PN2 =PO2 , where P is the gas partial pressure. Let P N2 ~ 1 Torr = 133 Pa, and then at 700 8C, P O2 ~3.21 10 25 Pa. The calculation shows that TiO2 is more stable than TiN as the oxygen partial pressure is higher than 3.21 10 25 Pa at 700 8C. Consequently the oxidation of TiN cannot be completely avoided at 700 8C under the experimental atmosphere. On the other hand, the rate of reaction can be estimated by the kinetic theory, P U ¼ 3:513 1022 pffiffiffiffiffiffiffiffi molecules=cm2 s [20] MT
where U is the flux of gas molecules impinging on the specimen surface, P is the oxygen partial pressure (= 7.6 10 20 Torr), M is the molecular weight of O2 (= 32), and T is the absolute temperature (= 973 K). By plugging in the values of P, M, T, and multiply the duration of heat treatment 3600 s, the collisions per unit area is 5.5 104 molecules/cm2. Assuming the sticking coefficient = 1, and a monolayer i 1015 molecules/cm2, the adsorbed oxygen molecule is 5.5 10 11 layer/cm2. Therefore, the kinetic estimation indicates that the oxide layer formed during heat treatment should be much less than 1 monolayer, which is not fully consistent with the AES results of the specimen A2. Since the specimens were not sealed in vacuum, in addition to the oxide formed from heat treatment, the oxide layer may be formed before and after heat treatment. In summary, mixing gas with a partial pressure ratio Ar/ H2 = 9 is found to be an effective atmosphere with extremely low oxygen partial pressure and the attaching Ti getter could further block the oxygen penetration path to the film surface. Both gas environment and Ti getter are important to prevent the oxidation of the TiN specimens during heat treatment.
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
(a)
4295
treated specimen ranges from 20 to 60 nm. The specimen annealed at 500 8C shows the similar microstructure as the as deposited specimen; in other words, there is no distinct microstructure change for the specimen heat treated at 500 8C for 1 h. 3.3. Crystallinity and preferred orientation Fig. 4 depicts the typical XRD patterns for the B-series specimens including three conditions: as deposited, heattreated at 500 and at 700 8C for 1 h, respectively. No oxidation product of TiN is found in the XRD patterns, which is consistent with the thermodynamic and kinetic estimations. From the XRD patterns, the (111) texture coefficient and the FWHM of (111) diffraction are determined. The results are listed in Table 1. As listed in Table 1, the FWHMs of the four groups of samples are generally decreased after heat treatment. Combined with the fact shown in Fig. 4 that certain XRD peaks become distinct, indicating the reorientation of grains, for the specimens after heat treating, it obviously reveals that thermal energy provided by the heat treatment facilitates the atom rearrangement in the thin films and thereby improving the crystallinity of all specimens. Similar observations that the decrease of FWHMs and enhancement of preferred orientation due to heat treatment have been reported in previous studies on TiN films [10,11]. The XRD results also show that the as deposited TiN films with a thickness of 214 nm are not well-crystallized, which has been observed in our previous study [12], and nearly amorphous for the film with a thickness of 86 nm. It is found that the effect of heat treatment on the crystallinity is more distinct than that in the thicker specimens. The grain sizes of the TiN films calculated from the XRD data are listed in Table 1. The TiN grain size generally increases after heat treatment at 500 and 700 8C.
100 nm
(b)
100 nm
Fig. 3. SEM cross-sectional images (a) specimen D0 (as deposited) and (b) specimen D700 (700 8C heat treatment).
3.2. Microstructure Figs. 3(a) and (b) show the typical cross-sectional SEM images for the specimens before and after heat-treating at 700 8C for 1 h, respectively. Compared these two images, it can be seen that the columnar structure in the as deposited film has been changed to granular structure after heat treatment. The columnar width of the as deposited thin film is approximately 50 nm, while the grain size of the heat-
(111)
(220) (200)
Counts
B700
B500
B0
25
30
35
40
45
50
55
60
65
70
2θ Fig. 4. XRD patterns of the B group specimens at as deposited, heat-treated at 500 and 700 8C conditions.
4296
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
Table 1 Summary of the experimental results Specimen
Thickness (nm)
N/Ti ratio
Texture coefficient (111)
(111) FWHM (deg.)
Grain size (nm)
Packing factor
Resistivity (AV cm)
Roughness R rms (nm)
Hardness (GPa)
Residual stress (GPa)
A0 A500 A700 B0 B500 B700 C0 C500 C700 D0 D500 D700
86 86 86 214 214 214 364 364 364 430 430 430
0.95 0.90 0.90 0.85 0.85 0.90 0.90 0.85 0.85 0.95 0.95 0.90
N/A 0.74 0.50 N/A 0.85 0.62 1.00 0.98 0.99 1.00 1.00 1.00
N/A 1.36 0.56 N/A 0.58 0.41 0.76 0.55 0.42 0.65 0.43 0.38
N/A 7 17 N/A 17 24 13 18 23 15 23 26
0.8 0.8 0.8 0.7 0.7 0.7 0.7 0.7 0.7 0.7 0.8 0.8
111 102 106 94 91 91 97 100 100 106 107 114
0.31 0.45 0.53 0.30 0.36 0.50 0.23 0.38 0.53 0.33 0.42 0.60
9.4 9.6 9.4 23.6 17.1 13.1 30.3 26.7 23.9 33.9 30.1 25.4
8.0 4.3 1.1 8.0 4.6 1.0 8.9 4.5 1.5 10.7 – 0.8
Again, the increase of grain size is more obvious in the thinner films. It can be seen that for the film thickness smaller than 214 nm, the grain size increases evidently from nearly amorphous to 17~24 nm after 700 8C heat treatment. In contrast, for the films thicker than 364 nm the grain size increases from 13 to 26 nm after 700 8C-treatment. Although the grain size measured from SEM images is larger than that measured by XRD, the results do not conflict each other. The grain size determined from XRD is an average value while that from SEM is the result of a random section. Nevertheless, the SEM images shown in Figs. 3(a) and (b) also display the increase of grain size, which is consistent with the XRD results, and both results confirm the TiN films are nanocrystalline. Fig. 5 shows the (111) and (200) texture coefficients with respect to the film thickness for the as deposited and the two different heat-treating conditions. It can be seen that the TC111 generally increases and the corresponding TC200 decreases with increasing film thickness. The variation of texture coefficient with heat treatment temperature is depicted in Fig. 6. From Figs. 5 and 6, it is found that
there is a critical thickness 214 nm, below this thickness (200) preferred orientation increases with heat treatment temperature while (111) texture coefficient decreases as the annealing temperature increases from 500 to 700 8C; for the thickness above 346 nm, (111) is prevailed and heat treatment shows subtle effect on (111) preferred orientation. Again, the effect of heat treatment on the variation of texture coefficient is more pronounced at a film thickness smaller than 214 nm. Previous studies [21,22] proposed that the preferred orientation of TiN film is associated with the lowest overall energy conditions. The crystallographic plane with the lowest surface energy in TiN is (200) [23,24], which is expected to be the preferred orientation at small film thickness. On the other hand, there are some hypotheses on the growth of (111) texture at larger thickness. Earlier studies by Pelleg et al. [21] and Je et al. [22] suggested that due to the smaller elastic constant at b111N than that at b200N, (111) becomes dominant at large film thickness; however, Greene et al. [25] and Li et al. [26] proposed that the governing process for the development of (111)
as deposited
1.0
1.1
500°C annealed
1.0
700°C annealed
0.9
Texture Coefficient
Texture Coefficient
0.8
TC111 0.6
0.4
TC200 0.2
0.8
TC111
0.7 0.6
A(86nm) B(214nm) C(364nm) D(430nm)
0.5 0.4 0.3 0.2
TC 200
0.1
0.0
0.0
50
100
150
200
250
300
350
400
450
Film Thickness (nm) Fig. 5. (111) and (200) texture coefficients vs. film thickness for the TiN films at as deposited, heat-treated at 500 and 700 8C conditions.
0
100
200
300
400
500
600
700
800
Temperature (°C)
Fig. 6. (111) and (200) texture coefficient vs. heat treatment temperature for the TiN films with four different film thicknesses.
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
3.4. Residual stress The residual stresses for all specimens are listed in Table 1. Fig. 7 shows the effect of heat treatment on the residual stress of the TiN films. For the as deposited specimens, the compressive residual stress slightly increases as the film thickness increases from 214 to 364 nm and then increases more obviously as the film thickness increases to 430 nm. In
Total residual stress
-12
Intrinsic residual stress -10
Residual Stress (GPa)
preferred orientation is due to competitive growth instead of lowering strain energy. Recent research on the thickness dependence of TiN texture by Abadias and Tse [27] supported the competitive growth mechanism. Their results also showed that the preferred orientation of the TiN film changes from (200) to (111) as the film thickness increases, with a crossover occurring between 150 and 200 nm. This is consistent with the results obtained in the present study, which indicated that below a critical thickness 214 nm, heat treatment showed significant effect on the increase of (200) preferred orientation, while (111) is prevailed as the film thickness is above 364 nm. Due to the higher surface area to volume ratio, surface diffusion may play a significant role for the film with a thickness less than 214 nm; in addition, for the same heat treatment duration of 1 h, the thermal energy per unit volume delivering into the thin film is higher for the thinner films. Consequently, the extent of atom rearrangement is expected to be higher in the thinner films than that in the thicker films, which leads to the fact that the effect of heat treatment on the variation of crystallinity, grain size and texture coefficient is more pronounced at a film thickness smaller than 214 nm. In addition, it is kinetically unfavorable to grow (200) planes in a highly 111 textured TiN film, which may also explain the subtle effect of heat treatment on the preferred orientation of the thicker specimens (series C and D). It is noted that TC111 first increases at 500 8C and then decreases at 700 8C for the two thinner specimens A and B; on the other hand, the corresponding TC200 increases with increasing temperature for the two series of specimens, as shown in Fig. 6. Compared with the XRD patterns shown in Fig. 4, the decrease of TC111 from 500 to 700 8C is due to the growth of (200) peak but not from the decrease of (111) peak. This can be explained by the competitive growth mechanism [25,26]. Owing to lower activation energy, planes with 111 direction grow faster than those at 200 direction. At lower temperature, 500 8C, the input thermal energy is sufficient for the growth of (111) plane and hence (111) planes grow much faster than (200) planes. As the temperature increases to 700 8C, all major low-index planes, (111), (200), and (220) planes, acquire sufficient thermal energy to grow, which leads to the decrease of TC111 while the increase of TC200. In other words, this is a result of competitive growth between planes at different orientations but not an orientation switch from the existing (111) to (200).
4297
as deposited
-8 -6
500°C -4 -2 0
700°C
2 50
100
150
200
250
300
350
400
450
Thickness (nm) Fig. 7. The variation of total residual stress and the intrinsic residual stress with the film thickness for the specimens before and after heat treatment.
general, the stress maintains at a high level ranging from 8 to 10.7 GPa. After heat treatment, the residual stress can be relieved more than 40% at 500 8C, and almost 100% at 700 8C for 1 h. The residual stress is originated from two sources, the intrinsic stress (grown-in stress) and the extrinsic stress (thermal stress). The thermal stress is due to the difference in thermal expansion coefficient between film and substrate material, and can be formulated as E rthermal stress ¼ 1v ðaf as ÞDT [28], where E is the Young’s modulus, m is the Poisson’s ratio, a f and a s are the thermal expansion coefficients of the film and the substrate, respectively, and DT is the temperature difference between the deposition or heat treatment temperature and the temperature at which the stress is measured. Since the residual stress was measured at 25 8C, the thermal stresses calculated from this equation are 0.93 GPa, 1.47 GPa, and 2.09 GPa for the as deposited, 500 8C-treated, and 700 8Ctreated specimens, respectively. This tensile thermal stress has been reported by Mayrhofer et al. [11] on TiN films deposited on Si after 700 8C heat treatment. The dash lines plotted in Fig. 7 are the intrinsic stresses calculated from total residual stresses subtracting the thermal stresses at the corresponding temperatures. The high compressive intrinsic stress is associated with the incorporation of metal atoms on nitrogen sites, the nitrogen interstitials, and the trapping Ar atoms [3,29]. From Fig. 7, it can be seen that the heat treatment effectively relieves the compressive intrinsic stress; with increasing temperature, more compressive stress is relieved. The heat treatment mainly provides thermal energy to facilitate the atom rearrangement by diffusion in the thin film. The results indicate that at 700 8C most of the defects, considered responsible for the intrinsic stress, have been annihilated and about 90% of the intrinsic stress is relieved. The tensile residual stresses in the specimens A, B, and C after 700 8C heat treatment
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299 1.0 as deposited 500°C annealing 700°C annealing
0.9
Roughness (nm)
0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 50
100
150
200
250
300
350
400
450
500
Thickness (nm)
Fig. 8. The variation of roughness with film thickness for as deposited, heat-treated at 500 and 700 8C conditions.
are attributed to the large thermal stress 2.09 GPa added to the relatively small compressive intrinsic stress ~ 1 GPa. 3.5. Packing factor, electrical resistivity, and roughness Table 1 lists the values of packing factor, electrical resistivity, and roughness of the four groups of specimens. Previous studies reported that since the electrical conductivity of TiN is from metallic origin [30], the electrical resistivity is related to the packing factor of thin film [1]. The results show that in general the packing factors were not appreciably affected by the heat treatment and therefore the electrical resistivity is not supposed to be considerably varied. Roughness values (R rms) measured by AFM are also listed in Table 1. Fig. 8 depicts the effect of heat treatment on the roughness of TiN thin films. The roughness slightly increases from 0.23 to 0.6 nm with increasing heat treatment temperature. Apparently, the increase of heat treatment temperature leads to an increase of the average protrusion size and a broadening of their size distribution and consequently the surface roughness increases. This finding is opposite to our earlier results [10]. In our previous research, TiN films produced by hollow cathode discharge ion plating (HCD-IP) were heat treated at the same atmosphere. It was found that the roughness (R rms) of the TiN films decreases with increasing heat treatment temperature. The disparity may be due to the difference in grain size. The TiN films in our earlier study had a columnar structure, with a width ranging from 60 to 100 nm, which did not change with heat treatment; in contrast, the microstructure of the films in this study experienced a change from columnar to nano-granular structure ranging from 20 to 60 nm after heat treatment. The nanocrystalline TiN films produced in the present study have much
smaller roughness, about one tenth of those previously made. Therefore, the roughness is related to the grain structure in thin film, and the roughness decreases with decreasing grain size. Comparing Fig. 3(a) and (b), one can find that after heat treating at 700 8C which is definitely lower than the recrystallization temperature of TiN (~ 1500 8C), the TiN film shows the occurrence of recrystallization. Previous study reported that grain growth by subgrain growth and grain boundary migration occurs at 700 8C in the TiN films with nanograin size of 20–30 nm [11]. This indicates that in nanocrystalline materials, the recrystallization temperature becomes uncertain, because the temperature may depend on grain size. During heat treatment, grain growth in the nanocrystalline thin film is more pronounced and thereby increasing the roughness; on the other hand, for the film with larger columnar structure, the reduction of surface energy may play an important role in smoothing the film surface. 3.6. Hardness Fig. 9 shows the variation of hardness of the TiN thin film with respect to film thickness at the three heat treatment conditions. It can be seen that the hardness generally decreases with increasing heat-treated temperature, except for the group A specimens. As mentioned in the experimental procedures, the depth displacement of the nanoindentation should be less than one tenth of the film thickness. K.W. Lee et al. [31] suggested that for coatings in the 15–35 GPa hardness range deposited on silicon, substrate effects can be ignored at indentation depths b 10–15% of the film thickness, if the film thickness is not significantly less than 130–200 nm. Therefore, except for sample series A, there will be no substrate effect in the hardness measurements under the experimental conditions. For the A-group specimens with a thickness of 86 nm, the depth displacement should be about 8~10 nm which is too 40 as-deposited
35
500°C annealing 700°C annealing
30
Hardness (GPa)
4298
25 20 15 10 5 0 50
100
150
200
250
300
350
400
450
500
Thickness (nm) Fig. 9. Hardness vs. film thickness for TiN samples for as deposited, heattreated at 500 and 700 8C conditions.
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
small to reveal the film hardness without substrate influence. Since the TiN thin films are nanocrystalline, the dislocation slip mechanism is no longer active in plastic deformation due to the limited number of dislocations in the grain; instead, grain rotation and grain boundary sliding [32,33] may account for the deformation of the material. The results in our recent studies [12,13] suggested that the hardness of nanocrystalline TiN or ZrN thin films is likely contributed from the grain boundary strength and defect density. As discussed, the relief of residual stress by heat treatment is due to the reduction of defect density, and the improvement of crystallinity and texture may from the same reason. Therefore, the decrease of hardness can also be attributed to the decrease of defect densities.
4. Conclusions (1) The oxidation of TiN thin film could be minimized as the heat treatment carried out at 700 8C, under the atmosphere with a ratio of Ar/H2 = 9 and attached with a Ti getter. (2) Heat treatment showed great effect on the microstructure, crystallinity, texture, grain size, roughness, and residual stress of the nanocrystalline TiN films, while less effect on N/Ti ratio, packing factor, and electrical resistivity. The heat treatment is more effective at 700 8C than at 500 8C. (3) The changes in texture, grain size, and crystallinity were more obvious in the specimens with a thickness less than 214 nm. (4) Heat treatment at 700 8C changed the microstructure of the nanocrystalline TiN films from columnar to nano-granular structure. (5) The residual stress and hardness of TiN thin films decreased after heat treatment, which was likely due to the reduction of the defect densities after heat treatment.
Acknowledgements This research was funded by the National Science Council of the Republic of China under the contract NSC 92-2216-E-007-016. Residual stress measurements were carried out in the Department of Physics and Materials Science, City University of Hong Kong. The authors also acknowledged Dr. Wen-Jun Chou for assisting the operation of UBM sputtering machine.
4299
References [1] W.-J. Chou, G.-P. Yu, J.-H. Huang, Surf. Coat. Technol. 140 (2001) 206. [2] W.-J. Chou, G.-P. Yu, J.-H. Huang, Surf. Coat. Technol. 149 (2002) 7. [3] H. Oettel, R. Wiedemann, S. Preigler, Surf. Coat. Technol. 74–75 (1995) 273. [4] Z. Cheng, H. Peng, G. Xie, Y. Shi, Surf. Coat. Technol. 138 (2001) 237. [5] K. Xu, J. Chen, R. Gao, J. He, Surf. Coat. Technol 58 (1993) 37. [6] W. Mader, H.F. Fischmeister, E. Bergmann, Thin Solid Films 182 (1989) 141. [7] M. Tamura, H. Kubo, Surf. Coat. Technol. 49 (1991) 194. [8] I. Milosev, H.-H. Strehblow, B. Navinsek, Thin Solid Films 303 (1997) 246. [9] F.-H. Lu, H.-Y. Chen, Thin Solid Films 388–389 (2001) 368. [10] W.-J. Chou, G.-P. Yu, J.-H. Huang, Surf. Coat. Technol. 168 (2003) 43. [11] P.H. Mayrhofer, F. Kunc, J. Musil, C. Mitterer, Thin Solid Films 415 (2002) 151. [12] J.-H. Huang, K.-W. Lau, G.-P. Yu, Surf. Coat. Technol. 191 (2005) 17. [13] J.-H. Huang, H.-C. Yang, X.-J. Guo, G.-P. Yu, Surf. Coat. Technol. 195 (2005) 204. [14] P.H. Mayrhofer, A. Hfrling, L. Karlsson, J. Sjfle´n, C. Mitterer, L. Hultman, Appl. Phys. Lett. 83 (10) (2003) 2049. [15] P. Scherrer, Gftt. Nachr. 2 (1918) 98. [16] C.H. Ma, J.H. Huang, Haydn Chen, Thin Solid Films 418 (2002) 73. [17] I.C. Noyan, J.B. Cohen, Residual Stress Measurement by Diffraction and Interpolation, Spring-Verlag, New York, 1987. [18] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [19] O. Kubaschewski, L.L. Evans, C.B. Alcock, Metallurgical Thermochemistry, 4th ed., Pergamon, New York, 1967, p. 428. [20] M. Ohring, Materials Science of Thin Films, Deposition and Structure, 2nd edition, Academic Press, San Diego, 2002, p. 62. [21] J. Pelleg, L.Z. Zevin, S.L. Croitoru, Thin Solid Films 197 (1991) 117. [22] J.H. Je, D.Y. Noh, H.K. Kim, K.S. Liang, J. Appl. Phys. 81 (9) (1997) 6126. [23] L. Hultmann, J.-E. Sundgren, J. Appl. Phys. 66 (1989) 536. [24] T.C. Chou, J. Appl. Phys. 67 (1990) 2670. [25] J.E. Greene, J.E. Sundgren, L. Hultman, I. Petrov, D.G. Bergstrom, Appl. Phys. Lett. 67 (1995) 2928. [26] T.Q. Li, S. Noda, Y. Tsuji, T. Ohsawa, H. Komiyama, J. Vac. Sci. Technol., A 20 (2002) 1270. [27] G. Abadias, Y.Y. Tse, J. Appl. Phys. 95 (2004) 2414. [28] D. Burgreen, Elements of Thermal Stress Analysis, C.P. Press, New York, 1971, p. 462. [29] H. Ljungcrantz, L. Hultman, J.-E. Sundgren, J. Appl. Phys. 78 (2) (1995) 832. [30] J.E. Sundgren, B.O. Johansson, A. Rockett, S.A. Barnett, J.E. Greene, in: W.D. Sproul, J.E. Greene, J.A. Thornton (Eds.), Physics and Chemistry of Protective Coatings, American Institute of Physics, MA, 1986, p. 95. [31] K.W. Lee, Y.W. Chung, C.Y. Chan, Surf. Coat. Technol. 168 (2003) 57. [32] I.A. Ovid’ko, Science 295 (2002) 2386. [33] M. Murayama, J.M. Howe, H. Hidaka, S. Takaki, Science 295 (2002) 2433.
Heat treatment of nanocrystalline TiN films deposited by unbalanced magnetron sputtering Jia-Hong Huanga,T, Kae-Jy Yua, P. Sitb, Ge-Ping Yua a
Department of Engineering and System Science, National Tsing Hua University, Hsinchu 300, Taiwan Department of Physics and Materials Science, City University of Hong Kong, Kowloon, Hong Kong
b
Received 12 October 2004; accepted in revised form 5 February 2005 Available online 11 April 2005
Abstract TiN thin films were deposited on Si wafers using an unbalanced magnetron (UBM) sputtering technique. Heat treatment was implemented after deposition to improve the properties of the TiN thin films. Combinations of mixing gas atmospheres and Ti getter were evaluated by thermodynamic and kinetic estimations to find a heat treatment environment with minimum oxidation. Results showed that at 700 8C the oxidation of TiN thin film could be minimized under an atmosphere with a partial pressure ratio of Ar/H2 = 9 and attached with a Ti getter. The specimens were heat-treated at 500 and 700 8C for 1 h under Ar/H2 atmosphere. Heat treatment showed significant effect on the microstructure, crystallinity, texture, grain size, roughness, and residual stress of the nanocrystalline TiN films, while less distinct effect was found on N/Ti ratio, packing factor and electrical resistivity. In addition, the changes in structure and mechanical properties were more effective at 700 8C than those at 500 8C. The changes in texture, grain size, and crystallinity were more obvious in the specimens with a thickness less than 214 nm. The residual stress and hardness were distinctly decreased by the heat treatment especially at 700 8C. This could be attributed to the reduction of lattice defects by the heat treatment. D 2005 Elsevier B.V. All rights reserved. Keywords: Heat treatment; TiN; Unbalanced magnetron sputtering; Residual stress
1. Introduction Owing to its high hardness, wear resistance, and thermal stability, titanium nitride (TiN) thin film is commercially used in the tool industry as a protective film to prolong the service life of substrate materials. The progress of microelectromechanical systems (MEMS) is driving the demand for ultra-thin (less than 0.5 Am) protective and functional nitride coatings. Our previous studies found that as the film thickness of TiN is less than 400 nm, the film properties may become inferior due to less compact structure [1,2]. To meet the requirements of the ultra-thin coatings in MEMS, especially less than 0.5 Am, the film properties must be enhanced not from increasing thickness but through the
T Corresponding author. Tel.: +886 35715131x4274; fax: +886 35720724. E-mail address: [email protected] (J.-H. Huang). 0257-8972/$ - see front matter D 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2005.02.147
basic improvement of the film structure, which has become a new challenge to coating and thin film technology. Besides the optimization of the coating processes, heat treatment after deposition has been widely used in industry to improve the properties of hard coatings by reducing the residual stress and defects in the deposited film [3–6]. However, earlier studies have reported that one of the factors that reduce the hardness of the TiN film is the surface oxidation during heat-treating process [7,8]. Recently, to overcome the oxidation issue, heat treatments proceeding in the controlled atmosphere have been reported [9–11]. Chou et al. [10] found that for the TiN films deposited by ion-plating, both residual stress and hardness are decreased after heat treatment at 700 8C in Ar/H2 atmosphere, while the microstructure and packing factor inside the TiN films are not significantly changed. Mayrhofer et al. [11] annealed TiN films deposited by DC unbalanced magnetron sputtering at 700 8C in vacuum. They also observed the decrease in both residual stress and
4292
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
hardness of the films after heat treatment. In contrast, they reported that grain size increases at 700 8C in the nanocrystalline TiN films. Since titanium oxides are thermodynamically more stable than TiN at high temperature, which implies that annealing is usually accompanied with the oxidation of TiN, it is very difficult to totally suppress the oxidation reaction by reducing oxygen partial pressure even down to 10 20 Pa at 700 8C. Therefore, to understand the effect of heat treatment on the structure and properties of TiN film, it is necessary to search a proper environment with least oxidation for avoiding the complication of oxidation in the heat treating process. Recent interest in nanotechnology stimulates research in the nanocrystalline nitride thin films. One of the attracting characteristics is that the nanocrystalline (grain size b 100 nm) materials may exhibit very different properties from micro-crystalline (grain size N 100 nm) materials. Huang et al. [12,13] found that unlike normal micro-crystalline TiN films, the hardness of nanocrystalline TiN or ZrN thin films does not significantly vary with (111) texture coefficient. The texture-insensitive behavior was attributed to the fact that the deformation mechanisms of the nanocrystalline nitride films are possibly from grain rotation or grain boundary sliding, instead of a dislocation slip mechanism. For the nanocrystalline thin film, due to very large ratio of atoms on grain boundary to atoms in grain, the thermally activated processes may be greatly enhanced by short circuit diffusion through grain boundaries. This effect is expected to be pronounced in the heat treatment process. In fact, Mayrhofer et al. [11] found that grain growth by subgrain growth and grain boundary migration occurs at 700 8C in the TiN films with nanograin size (20–30 nm). Although many studies [9–11,14] have been performed on the heat treatment of nitride coatings in controlled atmospheres, there is little information available on the heat treatment of nanocrystalline (grain size b 20 nm) thin films (thickness b 400 nm). Using a DC unbalanced magnetron (UBM) sputtering system, we successfully deposited nanocrystalline TiN thin films on (100) silicon wafers in an earlier research [12]. To continue the study, the UBM sputtering method was chosen to deposit TiN films on silicon substrates, and the deposition conditions were slightly changed to produce the specimens suitable for the study of heat treatment. The purposes of this study were to find a heat treatment environment with minor oxidation, and to investigate the effect of heat treatment in the controlled atmosphere on the microstructure and mechanical properties of TiN films.
2. Experimental procedures The deposition of TiN films was carried out using a UBM sputtering system. The substrate material was Si (100) wafer with dimensions of 45 mm 45 mm 0.7 mm. Prior to the coating process, the Si specimens were
undergone ultrasonic cleaning progressively in acetone and ethanol and then dried in a vacuum dryer. Before deposition, the substrates were gradually preheated to a temperature of 300 8C. Meanwhile, the chamber was evacuated to 8 10 4 Pa (6 10 6 Torr) to avoid contamination during deposition process. Prior to deposition, the substrate was pre-sputtered by argon discharge at a bias of 1000 V for 10 minutes to remove the surface oxide layer. The argon pressure was fixed at 0.8 Pa (6 10 3 Torr) and the current density at the substrate was 0.15 mA/cm2 during pre-sputtering. The target-to-substrate distance was 20 cm. After pre-sputtering, high purity working gas and reactive gas were introduced, using mass flow controllers to regulate both gas flows. Argon (99.9995% in purity) gas and nitrogen gas (99.9995% in purity) flow rates were fixed at 15 sccm and 1 sccm, respectively. The total gas pressure was controlled at 0.133 Pa (1 10 3 Torr). The DC power supply of the Ti target was operated at 0.9 A using constant-current mode; in other words, the target current density was 2.8 mA/cm2. During deposition, a negative substrate bias voltage of 50 V was applied and the deposition temperature was maintained at 350 8C, monitored using a thermocouple near the substrate. Film thickness was controlled by varying deposition duration. Four sets of specimens with different film thickness, denoted as A0(86 nm), B0(214 nm), C0(364 nm), and D0(430 nm), were deposited for further heat treatment. A LINDBERG high-temperature tube furnace equipped with a mechanical pumping system was used to perform the heat treatment. After inserting the specimens, the quartz tube was sealed and pumped down to 1.33 Pa. Then, the tube was purged with working gas for several times and the final pressure was controlled at 10 Pa by flowing high purity mixing gas, including gases with partial pressure ratios of Ar/H2 = 9 and N2/H2 = 9. The oxygen partial pressure was monitored using a zirconia oxygen gauge (15% CaO-doped ZrO2) in the gas mixture. In addition to the controlled atmosphere, Si wafers with pure titanium coatings were used as the oxygen getter. The getters were placed closely to the samples for preventing the uptake of oxygen by the TiN films. The specimens were heat treated at 500 and 700 8C for 1 h. In the duration of heat treatment, the oxygen partial pressures were maintained at 5.05 10 21 Pa and 1.01 10 17 Pa at 500 and 700 8C, respectively. The crystal structure of the TiN films was characterized by X-ray diffraction (XRD). The Cu Ka line at 0.15405 nm was used as the source for diffraction. The extent of (111) preferred orientation is quantified by a texture coefficient (TC) defined as I(111) / [I(111) + I(200) + I(220)], where I is the integrated intensity of the corresponding Bragg peak. The (111) texture coefficient of the TiN powder sample, from JCPDS 38-1420, is 0.33. The grain size of the TiN thin films was estimated from the results of h/2h scans. The position of the (111)
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
diffraction peak and the full width of the peak at half maximum (FWHM) were used to estimate the grain size by Scherrer’s equation [15] t¼
0:9k B cosh
where k is the wavelength of Cu Ka, B is the FWHM of the diffraction peak, h is the peak position, and t is the grain size of the thin film. The residual stress of the TiN films was determined by modified XRD sin2w method [16,17] using a 4-circle diffractometer with psi-goniometer geometry. X-ray was incident at an angle of 28 to increase the diffraction volume of the thin film specimen. The columnar structure and the film thickness of the deposited TiN film were observed and measured from the crosssectional specimens using a field-emission gun scanning electron microscope (FE-SEM) operated at 15 keV. The N/Ti ratio and packing factor were obtained from the results of Rutherford backscattering spectroscopy (RBS). The calculation of the packing factor from the RBS data was described elsewhere [1]. To access the oxygen influence, the compositional depth profiles were determined using Auger electron spectroscopy (AES). Since large substrate effect may occur for the TiN film with a thickness less than 500 nm when using normal microhardness tester, the film hardness was measured using a Hysitron nanoindenter attached on a Digital atomic force microscope (AFM). The applying loads were dependent on the depth displacement that should be less than one tenth of the film thickness. For the specimens with a thickness larger than 300 nm (series C and D), the indentation depth was maintained at 30 nm; for those smaller than 214 nm (series A and B), one tenth of the film thickness was adopted as the indentation depth. The procedures of applying load and calibration of the nanoindenter were detailed in our previous paper [1]. The hardness was calculated using Oliver–Pharr technique [18], which could be executed from the software provided by Hysitron Inc. Ten indentations were made and the average value was reported. The precision of each value was within 10%. The surface roughness of the TiN films were determined by AFM and the root mean square roughness value (R rms) was reported. The electrical resistivity of the TiN films was measured using a four-point probe.
3. Results and discussion 3.1. Thermodynamic and kinetic estimations of the annealing environment A typical oxygen partial pressure along with furnace temperature with respect to time under the atmosphere containing hydrogen is illustrated in Fig. 1. It can be seen that the oxygen content in the system could be lowered down to 1.01 10 17 Pa at 700 8C using the mixing gas
4293
Temperature Celsius 700.00
Oxygen atm 1E-20 1E-21 1E-22
600.00
1E-23 1E-24 1E-25
500.00 1E-26 1E-27 1E-28
400.00
1E-29 1E-30 300.00
1E-31 1E-32 1E-33
200.00 1E-34 1E-35 1E-36
100.00
1E-37 1E-38 0.00
1E-39 0.00
4000.00
8000.00
12000.00
16000.00
Time (sec)
Fig. 1. The typical oxygen partial pressure along with furnace temperature with respect to time under the atmosphere containing hydrogen.
environment. In addition, since the reaction of hydrogen and oxygen is exothermic, according to the Le Chartelier’s principle, the oxygen partial pressure is increased with increasing temperature. Samples with the thinnest TiN coatings (A-series) were used to investigate the composition variation under different atmospheres at 700 8C. Fig. 2 shows the AES results for the specimens treated at 700 8C under different environments. Although the N/Ti ratio cannot be determined by the AES analysis, the depth profiles can reveal the distribution of O, Si, C, and (Ti + N). As shown in Fig. 2(a), the compositions are uniformly distributed in the as deposited sample A0. Fig. 2(b) shows the profiles of the sample A700 heat-treated at 700 8C in Ar/H2 atmosphere attached with Ti getter. Apparently, there is no significant change in compositions before and after annealing. In contrast, for the specimen directly exposing to hydrogen mixing gas without Ti getter, more oxide (~ 24 nm) is formed on the surface of the TiN coating, as shown in Fig. 2(c). For the specimen heat treated in air (Fig. 2(d)), the TiN film has been thoroughly penetrated by oxygen. The results suggest that Ti getter plays an important role in preventing oxidation of the TiN thin film. Without Ti getter, oxidation occurs even at a controlled atmosphere with extremely low oxygen content.
4294
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
(a)
Ti N C O Si
N 6
2.0x10
Counts
Counts
1.6x10
Ti
6
1.2x10
2.0x10
6
1.6x10
6
1.2x10
6
8.0x10
5
4.0x10
5
N Si
Si
6
5
8.0x10
Ti
O
O 5
4.0x10
C
C
0.0
0.0 0
200
Ti N C O Si
(b)
400
600
800
0
1000
200
Sputtering Time (s)
400
600
800
(c)
Ti N C O Si
N 6
2.0x10
(d) Ti
Ti N C O Si
6
2.0x10
Si
O Si
6
1.6x10
1.6x106
6
Ti
1.2x10
Counts
Counts
1000
Sputtering Time (s)
5
8.0x10
1.2x106
N Ti
8.0x105
O 5
4.0x105
4.0x10
C
C 0.0
0.0 0
200
400
600
800
1000
1200
Sputtering Time (s)
0
200
400
600
800
1000
1200
1400
Sputtering Time (s)
Fig. 2. AES results of TiN thin films under different annealing environments. (a) As deposited (specimen A0); (b) annealed at 700 8C in Ar/H2 = 9 atmosphere attached with Ti getter (specimen A700); (c) annealed at 700 8C in N2/H2 = 9 atmosphere without Ti getter (specimen A2); and (d) annealed at 700 8C in air (Specimen A1).
The oxidation reaction of TiN involved in the heat treatment process can be understood by the thermodynamic and kinetic estimations. The main reaction of the oxidation of TiN is TiN þ O2 YTiO2 þ 1=2N2 DG ¼ RT lnK ¼ 604672 þ 84:308T ðJ=molÞ [19] where G is the Gibbs free energy, T is the absolute temperature, R is the gas constant (= 8.314 J/K mol), and K is constant. The reaction constant K= the reaction 1=2 PN2 =PO2 , where P is the gas partial pressure. Let P N2 ~ 1 Torr = 133 Pa, and then at 700 8C, P O2 ~3.21 10 25 Pa. The calculation shows that TiO2 is more stable than TiN as the oxygen partial pressure is higher than 3.21 10 25 Pa at 700 8C. Consequently the oxidation of TiN cannot be completely avoided at 700 8C under the experimental atmosphere. On the other hand, the rate of reaction can be estimated by the kinetic theory, P U ¼ 3:513 1022 pffiffiffiffiffiffiffiffi molecules=cm2 s [20] MT
where U is the flux of gas molecules impinging on the specimen surface, P is the oxygen partial pressure (= 7.6 10 20 Torr), M is the molecular weight of O2 (= 32), and T is the absolute temperature (= 973 K). By plugging in the values of P, M, T, and multiply the duration of heat treatment 3600 s, the collisions per unit area is 5.5 104 molecules/cm2. Assuming the sticking coefficient = 1, and a monolayer i 1015 molecules/cm2, the adsorbed oxygen molecule is 5.5 10 11 layer/cm2. Therefore, the kinetic estimation indicates that the oxide layer formed during heat treatment should be much less than 1 monolayer, which is not fully consistent with the AES results of the specimen A2. Since the specimens were not sealed in vacuum, in addition to the oxide formed from heat treatment, the oxide layer may be formed before and after heat treatment. In summary, mixing gas with a partial pressure ratio Ar/ H2 = 9 is found to be an effective atmosphere with extremely low oxygen partial pressure and the attaching Ti getter could further block the oxygen penetration path to the film surface. Both gas environment and Ti getter are important to prevent the oxidation of the TiN specimens during heat treatment.
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
(a)
4295
treated specimen ranges from 20 to 60 nm. The specimen annealed at 500 8C shows the similar microstructure as the as deposited specimen; in other words, there is no distinct microstructure change for the specimen heat treated at 500 8C for 1 h. 3.3. Crystallinity and preferred orientation Fig. 4 depicts the typical XRD patterns for the B-series specimens including three conditions: as deposited, heattreated at 500 and at 700 8C for 1 h, respectively. No oxidation product of TiN is found in the XRD patterns, which is consistent with the thermodynamic and kinetic estimations. From the XRD patterns, the (111) texture coefficient and the FWHM of (111) diffraction are determined. The results are listed in Table 1. As listed in Table 1, the FWHMs of the four groups of samples are generally decreased after heat treatment. Combined with the fact shown in Fig. 4 that certain XRD peaks become distinct, indicating the reorientation of grains, for the specimens after heat treating, it obviously reveals that thermal energy provided by the heat treatment facilitates the atom rearrangement in the thin films and thereby improving the crystallinity of all specimens. Similar observations that the decrease of FWHMs and enhancement of preferred orientation due to heat treatment have been reported in previous studies on TiN films [10,11]. The XRD results also show that the as deposited TiN films with a thickness of 214 nm are not well-crystallized, which has been observed in our previous study [12], and nearly amorphous for the film with a thickness of 86 nm. It is found that the effect of heat treatment on the crystallinity is more distinct than that in the thicker specimens. The grain sizes of the TiN films calculated from the XRD data are listed in Table 1. The TiN grain size generally increases after heat treatment at 500 and 700 8C.
100 nm
(b)
100 nm
Fig. 3. SEM cross-sectional images (a) specimen D0 (as deposited) and (b) specimen D700 (700 8C heat treatment).
3.2. Microstructure Figs. 3(a) and (b) show the typical cross-sectional SEM images for the specimens before and after heat-treating at 700 8C for 1 h, respectively. Compared these two images, it can be seen that the columnar structure in the as deposited film has been changed to granular structure after heat treatment. The columnar width of the as deposited thin film is approximately 50 nm, while the grain size of the heat-
(111)
(220) (200)
Counts
B700
B500
B0
25
30
35
40
45
50
55
60
65
70
2θ Fig. 4. XRD patterns of the B group specimens at as deposited, heat-treated at 500 and 700 8C conditions.
4296
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
Table 1 Summary of the experimental results Specimen
Thickness (nm)
N/Ti ratio
Texture coefficient (111)
(111) FWHM (deg.)
Grain size (nm)
Packing factor
Resistivity (AV cm)
Roughness R rms (nm)
Hardness (GPa)
Residual stress (GPa)
A0 A500 A700 B0 B500 B700 C0 C500 C700 D0 D500 D700
86 86 86 214 214 214 364 364 364 430 430 430
0.95 0.90 0.90 0.85 0.85 0.90 0.90 0.85 0.85 0.95 0.95 0.90
N/A 0.74 0.50 N/A 0.85 0.62 1.00 0.98 0.99 1.00 1.00 1.00
N/A 1.36 0.56 N/A 0.58 0.41 0.76 0.55 0.42 0.65 0.43 0.38
N/A 7 17 N/A 17 24 13 18 23 15 23 26
0.8 0.8 0.8 0.7 0.7 0.7 0.7 0.7 0.7 0.7 0.8 0.8
111 102 106 94 91 91 97 100 100 106 107 114
0.31 0.45 0.53 0.30 0.36 0.50 0.23 0.38 0.53 0.33 0.42 0.60
9.4 9.6 9.4 23.6 17.1 13.1 30.3 26.7 23.9 33.9 30.1 25.4
8.0 4.3 1.1 8.0 4.6 1.0 8.9 4.5 1.5 10.7 – 0.8
Again, the increase of grain size is more obvious in the thinner films. It can be seen that for the film thickness smaller than 214 nm, the grain size increases evidently from nearly amorphous to 17~24 nm after 700 8C heat treatment. In contrast, for the films thicker than 364 nm the grain size increases from 13 to 26 nm after 700 8C-treatment. Although the grain size measured from SEM images is larger than that measured by XRD, the results do not conflict each other. The grain size determined from XRD is an average value while that from SEM is the result of a random section. Nevertheless, the SEM images shown in Figs. 3(a) and (b) also display the increase of grain size, which is consistent with the XRD results, and both results confirm the TiN films are nanocrystalline. Fig. 5 shows the (111) and (200) texture coefficients with respect to the film thickness for the as deposited and the two different heat-treating conditions. It can be seen that the TC111 generally increases and the corresponding TC200 decreases with increasing film thickness. The variation of texture coefficient with heat treatment temperature is depicted in Fig. 6. From Figs. 5 and 6, it is found that
there is a critical thickness 214 nm, below this thickness (200) preferred orientation increases with heat treatment temperature while (111) texture coefficient decreases as the annealing temperature increases from 500 to 700 8C; for the thickness above 346 nm, (111) is prevailed and heat treatment shows subtle effect on (111) preferred orientation. Again, the effect of heat treatment on the variation of texture coefficient is more pronounced at a film thickness smaller than 214 nm. Previous studies [21,22] proposed that the preferred orientation of TiN film is associated with the lowest overall energy conditions. The crystallographic plane with the lowest surface energy in TiN is (200) [23,24], which is expected to be the preferred orientation at small film thickness. On the other hand, there are some hypotheses on the growth of (111) texture at larger thickness. Earlier studies by Pelleg et al. [21] and Je et al. [22] suggested that due to the smaller elastic constant at b111N than that at b200N, (111) becomes dominant at large film thickness; however, Greene et al. [25] and Li et al. [26] proposed that the governing process for the development of (111)
as deposited
1.0
1.1
500°C annealed
1.0
700°C annealed
0.9
Texture Coefficient
Texture Coefficient
0.8
TC111 0.6
0.4
TC200 0.2
0.8
TC111
0.7 0.6
A(86nm) B(214nm) C(364nm) D(430nm)
0.5 0.4 0.3 0.2
TC 200
0.1
0.0
0.0
50
100
150
200
250
300
350
400
450
Film Thickness (nm) Fig. 5. (111) and (200) texture coefficients vs. film thickness for the TiN films at as deposited, heat-treated at 500 and 700 8C conditions.
0
100
200
300
400
500
600
700
800
Temperature (°C)
Fig. 6. (111) and (200) texture coefficient vs. heat treatment temperature for the TiN films with four different film thicknesses.
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
3.4. Residual stress The residual stresses for all specimens are listed in Table 1. Fig. 7 shows the effect of heat treatment on the residual stress of the TiN films. For the as deposited specimens, the compressive residual stress slightly increases as the film thickness increases from 214 to 364 nm and then increases more obviously as the film thickness increases to 430 nm. In
Total residual stress
-12
Intrinsic residual stress -10
Residual Stress (GPa)
preferred orientation is due to competitive growth instead of lowering strain energy. Recent research on the thickness dependence of TiN texture by Abadias and Tse [27] supported the competitive growth mechanism. Their results also showed that the preferred orientation of the TiN film changes from (200) to (111) as the film thickness increases, with a crossover occurring between 150 and 200 nm. This is consistent with the results obtained in the present study, which indicated that below a critical thickness 214 nm, heat treatment showed significant effect on the increase of (200) preferred orientation, while (111) is prevailed as the film thickness is above 364 nm. Due to the higher surface area to volume ratio, surface diffusion may play a significant role for the film with a thickness less than 214 nm; in addition, for the same heat treatment duration of 1 h, the thermal energy per unit volume delivering into the thin film is higher for the thinner films. Consequently, the extent of atom rearrangement is expected to be higher in the thinner films than that in the thicker films, which leads to the fact that the effect of heat treatment on the variation of crystallinity, grain size and texture coefficient is more pronounced at a film thickness smaller than 214 nm. In addition, it is kinetically unfavorable to grow (200) planes in a highly 111 textured TiN film, which may also explain the subtle effect of heat treatment on the preferred orientation of the thicker specimens (series C and D). It is noted that TC111 first increases at 500 8C and then decreases at 700 8C for the two thinner specimens A and B; on the other hand, the corresponding TC200 increases with increasing temperature for the two series of specimens, as shown in Fig. 6. Compared with the XRD patterns shown in Fig. 4, the decrease of TC111 from 500 to 700 8C is due to the growth of (200) peak but not from the decrease of (111) peak. This can be explained by the competitive growth mechanism [25,26]. Owing to lower activation energy, planes with 111 direction grow faster than those at 200 direction. At lower temperature, 500 8C, the input thermal energy is sufficient for the growth of (111) plane and hence (111) planes grow much faster than (200) planes. As the temperature increases to 700 8C, all major low-index planes, (111), (200), and (220) planes, acquire sufficient thermal energy to grow, which leads to the decrease of TC111 while the increase of TC200. In other words, this is a result of competitive growth between planes at different orientations but not an orientation switch from the existing (111) to (200).
4297
as deposited
-8 -6
500°C -4 -2 0
700°C
2 50
100
150
200
250
300
350
400
450
Thickness (nm) Fig. 7. The variation of total residual stress and the intrinsic residual stress with the film thickness for the specimens before and after heat treatment.
general, the stress maintains at a high level ranging from 8 to 10.7 GPa. After heat treatment, the residual stress can be relieved more than 40% at 500 8C, and almost 100% at 700 8C for 1 h. The residual stress is originated from two sources, the intrinsic stress (grown-in stress) and the extrinsic stress (thermal stress). The thermal stress is due to the difference in thermal expansion coefficient between film and substrate material, and can be formulated as E rthermal stress ¼ 1v ðaf as ÞDT [28], where E is the Young’s modulus, m is the Poisson’s ratio, a f and a s are the thermal expansion coefficients of the film and the substrate, respectively, and DT is the temperature difference between the deposition or heat treatment temperature and the temperature at which the stress is measured. Since the residual stress was measured at 25 8C, the thermal stresses calculated from this equation are 0.93 GPa, 1.47 GPa, and 2.09 GPa for the as deposited, 500 8C-treated, and 700 8Ctreated specimens, respectively. This tensile thermal stress has been reported by Mayrhofer et al. [11] on TiN films deposited on Si after 700 8C heat treatment. The dash lines plotted in Fig. 7 are the intrinsic stresses calculated from total residual stresses subtracting the thermal stresses at the corresponding temperatures. The high compressive intrinsic stress is associated with the incorporation of metal atoms on nitrogen sites, the nitrogen interstitials, and the trapping Ar atoms [3,29]. From Fig. 7, it can be seen that the heat treatment effectively relieves the compressive intrinsic stress; with increasing temperature, more compressive stress is relieved. The heat treatment mainly provides thermal energy to facilitate the atom rearrangement by diffusion in the thin film. The results indicate that at 700 8C most of the defects, considered responsible for the intrinsic stress, have been annihilated and about 90% of the intrinsic stress is relieved. The tensile residual stresses in the specimens A, B, and C after 700 8C heat treatment
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299 1.0 as deposited 500°C annealing 700°C annealing
0.9
Roughness (nm)
0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 50
100
150
200
250
300
350
400
450
500
Thickness (nm)
Fig. 8. The variation of roughness with film thickness for as deposited, heat-treated at 500 and 700 8C conditions.
are attributed to the large thermal stress 2.09 GPa added to the relatively small compressive intrinsic stress ~ 1 GPa. 3.5. Packing factor, electrical resistivity, and roughness Table 1 lists the values of packing factor, electrical resistivity, and roughness of the four groups of specimens. Previous studies reported that since the electrical conductivity of TiN is from metallic origin [30], the electrical resistivity is related to the packing factor of thin film [1]. The results show that in general the packing factors were not appreciably affected by the heat treatment and therefore the electrical resistivity is not supposed to be considerably varied. Roughness values (R rms) measured by AFM are also listed in Table 1. Fig. 8 depicts the effect of heat treatment on the roughness of TiN thin films. The roughness slightly increases from 0.23 to 0.6 nm with increasing heat treatment temperature. Apparently, the increase of heat treatment temperature leads to an increase of the average protrusion size and a broadening of their size distribution and consequently the surface roughness increases. This finding is opposite to our earlier results [10]. In our previous research, TiN films produced by hollow cathode discharge ion plating (HCD-IP) were heat treated at the same atmosphere. It was found that the roughness (R rms) of the TiN films decreases with increasing heat treatment temperature. The disparity may be due to the difference in grain size. The TiN films in our earlier study had a columnar structure, with a width ranging from 60 to 100 nm, which did not change with heat treatment; in contrast, the microstructure of the films in this study experienced a change from columnar to nano-granular structure ranging from 20 to 60 nm after heat treatment. The nanocrystalline TiN films produced in the present study have much
smaller roughness, about one tenth of those previously made. Therefore, the roughness is related to the grain structure in thin film, and the roughness decreases with decreasing grain size. Comparing Fig. 3(a) and (b), one can find that after heat treating at 700 8C which is definitely lower than the recrystallization temperature of TiN (~ 1500 8C), the TiN film shows the occurrence of recrystallization. Previous study reported that grain growth by subgrain growth and grain boundary migration occurs at 700 8C in the TiN films with nanograin size of 20–30 nm [11]. This indicates that in nanocrystalline materials, the recrystallization temperature becomes uncertain, because the temperature may depend on grain size. During heat treatment, grain growth in the nanocrystalline thin film is more pronounced and thereby increasing the roughness; on the other hand, for the film with larger columnar structure, the reduction of surface energy may play an important role in smoothing the film surface. 3.6. Hardness Fig. 9 shows the variation of hardness of the TiN thin film with respect to film thickness at the three heat treatment conditions. It can be seen that the hardness generally decreases with increasing heat-treated temperature, except for the group A specimens. As mentioned in the experimental procedures, the depth displacement of the nanoindentation should be less than one tenth of the film thickness. K.W. Lee et al. [31] suggested that for coatings in the 15–35 GPa hardness range deposited on silicon, substrate effects can be ignored at indentation depths b 10–15% of the film thickness, if the film thickness is not significantly less than 130–200 nm. Therefore, except for sample series A, there will be no substrate effect in the hardness measurements under the experimental conditions. For the A-group specimens with a thickness of 86 nm, the depth displacement should be about 8~10 nm which is too 40 as-deposited
35
500°C annealing 700°C annealing
30
Hardness (GPa)
4298
25 20 15 10 5 0 50
100
150
200
250
300
350
400
450
500
Thickness (nm) Fig. 9. Hardness vs. film thickness for TiN samples for as deposited, heattreated at 500 and 700 8C conditions.
J.-H. Huang et al. / Surface & Coatings Technology 200 (2006) 4291–4299
small to reveal the film hardness without substrate influence. Since the TiN thin films are nanocrystalline, the dislocation slip mechanism is no longer active in plastic deformation due to the limited number of dislocations in the grain; instead, grain rotation and grain boundary sliding [32,33] may account for the deformation of the material. The results in our recent studies [12,13] suggested that the hardness of nanocrystalline TiN or ZrN thin films is likely contributed from the grain boundary strength and defect density. As discussed, the relief of residual stress by heat treatment is due to the reduction of defect density, and the improvement of crystallinity and texture may from the same reason. Therefore, the decrease of hardness can also be attributed to the decrease of defect densities.
4. Conclusions (1) The oxidation of TiN thin film could be minimized as the heat treatment carried out at 700 8C, under the atmosphere with a ratio of Ar/H2 = 9 and attached with a Ti getter. (2) Heat treatment showed great effect on the microstructure, crystallinity, texture, grain size, roughness, and residual stress of the nanocrystalline TiN films, while less effect on N/Ti ratio, packing factor, and electrical resistivity. The heat treatment is more effective at 700 8C than at 500 8C. (3) The changes in texture, grain size, and crystallinity were more obvious in the specimens with a thickness less than 214 nm. (4) Heat treatment at 700 8C changed the microstructure of the nanocrystalline TiN films from columnar to nano-granular structure. (5) The residual stress and hardness of TiN thin films decreased after heat treatment, which was likely due to the reduction of the defect densities after heat treatment.
Acknowledgements This research was funded by the National Science Council of the Republic of China under the contract NSC 92-2216-E-007-016. Residual stress measurements were carried out in the Department of Physics and Materials Science, City University of Hong Kong. The authors also acknowledged Dr. Wen-Jun Chou for assisting the operation of UBM sputtering machine.
4299
References [1] W.-J. Chou, G.-P. Yu, J.-H. Huang, Surf. Coat. Technol. 140 (2001) 206. [2] W.-J. Chou, G.-P. Yu, J.-H. Huang, Surf. Coat. Technol. 149 (2002) 7. [3] H. Oettel, R. Wiedemann, S. Preigler, Surf. Coat. Technol. 74–75 (1995) 273. [4] Z. Cheng, H. Peng, G. Xie, Y. Shi, Surf. Coat. Technol. 138 (2001) 237. [5] K. Xu, J. Chen, R. Gao, J. He, Surf. Coat. Technol 58 (1993) 37. [6] W. Mader, H.F. Fischmeister, E. Bergmann, Thin Solid Films 182 (1989) 141. [7] M. Tamura, H. Kubo, Surf. Coat. Technol. 49 (1991) 194. [8] I. Milosev, H.-H. Strehblow, B. Navinsek, Thin Solid Films 303 (1997) 246. [9] F.-H. Lu, H.-Y. Chen, Thin Solid Films 388–389 (2001) 368. [10] W.-J. Chou, G.-P. Yu, J.-H. Huang, Surf. Coat. Technol. 168 (2003) 43. [11] P.H. Mayrhofer, F. Kunc, J. Musil, C. Mitterer, Thin Solid Films 415 (2002) 151. [12] J.-H. Huang, K.-W. Lau, G.-P. Yu, Surf. Coat. Technol. 191 (2005) 17. [13] J.-H. Huang, H.-C. Yang, X.-J. Guo, G.-P. Yu, Surf. Coat. Technol. 195 (2005) 204. [14] P.H. Mayrhofer, A. Hfrling, L. Karlsson, J. Sjfle´n, C. Mitterer, L. Hultman, Appl. Phys. Lett. 83 (10) (2003) 2049. [15] P. Scherrer, Gftt. Nachr. 2 (1918) 98. [16] C.H. Ma, J.H. Huang, Haydn Chen, Thin Solid Films 418 (2002) 73. [17] I.C. Noyan, J.B. Cohen, Residual Stress Measurement by Diffraction and Interpolation, Spring-Verlag, New York, 1987. [18] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [19] O. Kubaschewski, L.L. Evans, C.B. Alcock, Metallurgical Thermochemistry, 4th ed., Pergamon, New York, 1967, p. 428. [20] M. Ohring, Materials Science of Thin Films, Deposition and Structure, 2nd edition, Academic Press, San Diego, 2002, p. 62. [21] J. Pelleg, L.Z. Zevin, S.L. Croitoru, Thin Solid Films 197 (1991) 117. [22] J.H. Je, D.Y. Noh, H.K. Kim, K.S. Liang, J. Appl. Phys. 81 (9) (1997) 6126. [23] L. Hultmann, J.-E. Sundgren, J. Appl. Phys. 66 (1989) 536. [24] T.C. Chou, J. Appl. Phys. 67 (1990) 2670. [25] J.E. Greene, J.E. Sundgren, L. Hultman, I. Petrov, D.G. Bergstrom, Appl. Phys. Lett. 67 (1995) 2928. [26] T.Q. Li, S. Noda, Y. Tsuji, T. Ohsawa, H. Komiyama, J. Vac. Sci. Technol., A 20 (2002) 1270. [27] G. Abadias, Y.Y. Tse, J. Appl. Phys. 95 (2004) 2414. [28] D. Burgreen, Elements of Thermal Stress Analysis, C.P. Press, New York, 1971, p. 462. [29] H. Ljungcrantz, L. Hultman, J.-E. Sundgren, J. Appl. Phys. 78 (2) (1995) 832. [30] J.E. Sundgren, B.O. Johansson, A. Rockett, S.A. Barnett, J.E. Greene, in: W.D. Sproul, J.E. Greene, J.A. Thornton (Eds.), Physics and Chemistry of Protective Coatings, American Institute of Physics, MA, 1986, p. 95. [31] K.W. Lee, Y.W. Chung, C.Y. Chan, Surf. Coat. Technol. 168 (2003) 57. [32] I.A. Ovid’ko, Science 295 (2002) 2386. [33] M. Murayama, J.M. Howe, H. Hidaka, S. Takaki, Science 295 (2002) 2433.