Effects of composition and thermal annealing on the mechanical properties of silicon oxycarbide films

Sensors and Actuators A 176 (2012) 90–98

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Effects of composition and thermal annealing on the mechanical properties of silicon oxycarbide films Ping Du, Xiaoning Wang, I.-Kuan Lin, Xin Zhang ∗ Department of Mechanical Engineering, Boston University, Boston, MA 02215, USA

a r t i c l e

i n f o

Article history: Received 21 May 2011 Received in revised form 15 December 2011 Accepted 2 January 2012 Available online 21 January 2012 Keywords: Silicon carbide Silicon dioxide Sputtering Modulus Hardness Residual stress Thermal annealing Nanoindentation Microstructure

a b s t r a c t There is an increasing trend to incorporate silicon carbide (SiC) into silicon oxides to improve the mechanical properties, thermal stability, and chemical resistance. In this work the silicon oxycarbide (SiOC) films were deposited by RF magnetron co-sputtering from silicon dioxide and silicon carbide targets. Subsequently rapid thermal annealing was applied to the as-deposited films to tune the mechanical properties. Energy dispersive spectroscopy, scanning electron microscopy, Fourier transform infrared spectroscopy and ellipsometry were employed to characterize the compositions and microstructure of the films. The residual stress of the films was calculated from the film–substrate curvature measurement using Stoney’s equation. The film stress changed from compressive to tensile after annealing, and it generally increased with carbon contents. The Young’s modulus and hardness were investigated by the depth-sensing nanoindentation, which were found to increase with the carbon content and annealing temperature. A thorough microstructural analysis was conducted to investigate the effect of carbon content and annealing temperature on the mechanical properties of SiOC films. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Silicon-based ceramics (silicon oxides, silicon nitrides, silicon carbides and their combinations) are important structural and electronic materials for various applications in semiconductor and microelectromechanical systems (MEMS) [1]. For example, silicon carbide (SiC) is considered as a promising structural material with high strength at elevated temperatures [2,3], excellent chemical stability and is inert to corrosive atmosphere [4]. Recently, there has been an increasing trend to incorporate nitrogen or carbon into silicon oxides [5,6]. These higher valence atoms can be expected to increase the bond density and thus improve the mechanical properties, thermal stability, and chemical resistance of silicon oxide; as has been experimentally shown in multiple systems. In this work, we focused on silicon oxycarbide (SiOC) films. The fabrication technique and mechanical properties characterization were presented. Common fabrication techniques for these ceramics include melting, chemical vapor deposition (CVD), and sol–gel pyrolysis [5,7]. However, these approaches are typically conducted under high temperature. Therefore they often lead to phase separation, high residual stress and hydrogen incorporation which may have

∗ Corresponding author. Tel.: +1 617 358 2702; fax: +1 617 353 5548. E-mail address: [email protected] (X. Zhang). 0924-4247/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.sna.2012.01.002

detrimental effects to the final film properties. Sputtering is a more promising technique because of several intrinsic advantages such as: low temperature processing, low hydrogen content, elimination of toxic gases and a simpler deposition system [8,9]. Ech-chamikh et al. had synthesized SiOC films by reactive radio-frequency (RF) sputtering from a composite target (a silicon disc with graphite chips being placed) with oxygen gas. However, the films exhibited strong phase separation into silicon and/or carbon clusters [10]. Ryan and Pantano had synthesized SiOC films by the similar technique, using a silicon carbide target with oxygen gas [11]. The variation of composition was obtained by controlling the amount of oxygen present in the plasma. In this work, the SiOC films were deposited by RF magnetron co-sputtering from silicon dioxide and silicon carbide targets. The ratio of oxygen to carbon content, ranging from silicon oxide (SiOx ) to silicon carbide (SiCy ), was achieved by varying the power applied to each target. Using this co-sputtering technique, more homogeneous and broader range of composition could be expected in the films. In semiconductor and MEMS applications, mechanical properties such as modulus, hardness and residual stress play key roles to determine the device’s performance and structural reliability [12–14]. Nanoindentation has recently attracted intensive attention to measure modulus and hardness of materials at small scales. The mechanical properties can be determined directly from indentation load and displacement data using appropriate models [15].

P. Du et al. / Sensors and Actuators A 176 (2012) 90–98

Nanoindentation has proven particularly useful in probing the properties of thin films since indentations as shallow as several nanometers can be used to make measurements. Residual stresses, usually induced during the film formation, may have deleterious effects in thin-film processing: excessive compressive stress can result in film buckling and delamination, whereas excessive tensile stress may lead to film cracking [16]. In either case, residual stresses may damage the structures or reduce their service lifetime. Therefore, it is of vital importance to identify a reliable technique to minimize the residual stress in a controlled manner during the fabrication of MEMS. Rapid thermal annealing (RTA) can be employed to effectively alter the mechanical properties and residual stress of SiOC with small thermal budget [17]. The distinct response of residual stresses in silicon oxides and carbides will make it possible to tune the stress state in the film by properly controlling the composition of the films. In this work, we fabricate silicon oxycarbide films, and study the effect of composition on the mechanical properties of the films before and after rapid thermal annealing. The SiOC films with different compositions of oxygen and carbon contents, were deposited by RF magnetron co-sputtering from silicon dioxide and silicon carbide targets. Subsequently RTA was applied to these films in order to alter the mechanical properties and stress states. Energy dispersive spectroscopy (EDS), scanning electron microscopy (SEM), Fourier transform infrared spectroscopy (FTIR) and ellipsometry were employed to characterize the compositions and microstructure of the films. The residual stress of the films was calculated from the film–substrate curvature measurement using Stoney’s equation, and the Young’s modulus and hardness were investigated by the depth-sensing nanoindentation. Moreover, a thorough microstructural study was conducted to investigate the effect of carbon content and annealing temperature on the measured mechanical properties of SiOC films.

2. Experimental 2.1. Specimen preparation SiOC films were deposited on n-type (100) silicon wafer substrates at room temperature (23 ◦ C) by a Discovery 18 RF magnetron sputter (Denton Vacuum Inc.). All silicon substrates were precleaned with piranha solution (hydrogen peroxide and sulfuric acid with a volume ratio of 1:3) for 10 min, then rinsed with deionized water and dried with nitrogen. SiOx or SiCy films were directly sputtered from pure (>99.9% purity) silicon dioxide or silicon carbide targets (Kurt J. Lesker Inc.) at a fixed RF power of 300 W, respectively. The deposition rate of SiOx was 12.4 nm/min, and that of SiCy was 4.6 nm/min. The rate of SiOx was about three times of the SiCy . In other word, using a RF power of 100 W on the silicon dioxide target should result in comparable amount of SiOx versus SiCy obtained from 300 W RF power on the silicon carbide target. Subsequently, three SiOC films with different carbon/oxygen ratios were deposited by co-sputtering the SiO2 and SiC targets. The composition variation was achieved by varying the RF power applied to SiO2 target at 50 W, 100 W and 200 W, while maintaining the RF power of SiC target at 300 W (listed in Table 1). The flow rate of argon was 25 sccm (standard cubic centimeters per minute) to maintain a sputtering pressure at 4 mTorr during the deposition. The thickness of each film was controlled in ∼1 ␮m. After the deposition, rapid thermal annealing was applied to the specimens using a RTP-600S rapid thermal processing system (Modular Process Technology Corp.). Each specimen was subjected to 400 ◦ C, 600 ◦ C and 800 ◦ C annealing for 10 min in a nitrogen ambient environment, respectively.

91

Table 1 Sputtering conditions and atomic concentration of SiOC films. Specimen

SiOx SiOx Cy 1 SiOx Cy 2 SiOx Cy 3 SiCy a b

RF power (W)

Atomic concentration (at.%)

SiO2

SiC

Si

O

C

C/(C + O)

300 50 100 200 0

0 300 300 300 300

24.02 33.83 37.41 40.45 46.41

73.76 49.17 37.72 25.41 1.74b

2.22a 17.00 24.88 34.15 51.86

2.92 25.69 39.74 57.33 96.75

Could be due to the contamination of carbon. Could be due to the absorption of free oxygen in the film surface.

2.2. Energy dispersive spectroscopy A JSM-6100 scanning electron microscope (JEOL Ltd.) equipped with an energy dispersive spectroscopy (EDS, from Oxford Instruments) was used to quantitatively analyze the composition of the sputtered SiOC films. An accelerating voltage of 5 kV was selected for the electron beam. This was the minimum voltage required to obtain successful element analysis for Si, O and C. In addition, the low voltage ensured the interaction between beam and samples was confined within the 1 ␮m top layer of the films [6]. A Phi-Rho-Z correction procedure was used for all specimens in the quantitative analysis. 2.3. Scanning electron microscope and Fourier transform infrared spectroscope A SUPRA 55VP scanning electron microscope (SEM, from Carl Zeiss) was employed to obtain high-resolution surface micrographs of the SiOC films. The accelerating voltage was set to 2 kV, the aperture was 30 ␮m, and the working distance was 4.9 mm. The signals from both the Inlens and Second-electron (SE2) detectors were mixed to obtain optimal image quality. The chemical bonding structures were identified by Fourier transform infrared spectroscopy (FTIR) with an IFS 66 FTIR spectrometer (Bruker Inc.). A MCT-B detector and a K–Br beam splitter were employed in the analysis to obtain the mid-infrared (MIR) absorption spectra (4000–400 cm−1 ). A background spectrum of bare silicon wafer was first obtained to normalize the spectra of SiOC films, eliminating the contributions from the substrate, instrumentation and atmosphere. A total of 256 scans were performed with a resolution of 4 cm−1 for the background and each SiOC specimen. The obtained spectra were analyzed using the OPUS 5.5 software package (also from Bruker Inc.). 2.4. Thickness measurement The change of thickness before and after the thermal annealing was used to represent the microstructure revolution in the films. The thickness was measured by a VASE ellipsometer (J.A. Woollam Co., Inc.). The ellipsometer featured variable wavelength and angle of incidence allowing flexible measurement capabilities of thickness and optical constants of thin films. The change in polarization as the light reflects or transmits from the material. The polarization is expressed by amplitude ratio and phase difference between the incident and reflective lights. The wavelength used for the SiOC films was in the range of 600–1100 nm, in which the majority interference pattern occurred. The sampling rate was 10 nm/point, and the measurements were conducted in three incident angles, i.e. 65◦ , 70◦ and 75◦ . The model used to fit the experimental data consists of three layers: a Si wafer substrate, a Cauchy layer for the film and a surface roughness layer. The thickness non-uniformity option was applied in the Cauchy layer to take into account the uneven thickness in the

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P. Du et al. / Sensors and Actuators A 176 (2012) 90–98

2

Si

Loading

1.5

SiOx

1

Unloading S=dP/dh

SiOxCy1 Intensity (a.u.)

Load (mN)

C O

Pmax

0.5

hmax

hf 0 0

50

100

Disp (nm)

SiOxCy2

Fig. 1. Schematic diagram of load–displacement curve for nanoindentation tests.

SiOxCy3 sputtered films. Both the thickness and optical constants (refractive index and absorption coefficient) were determined by data regression, in which the object was to find the global minimum of mean square error (MSE).

SiCy 0

2.5. Residual stress measurement

(1)

where  is the residual stress, t is the thickness,  is the curvature after deposition, and 0 is the curvature of the substrate before deposition. E is defined as the biaxial modulus and equals E/(1 − ), where E is Young’s modulus and  is the Poisson’s ratio. The subscript s stands for substrate and f for film. Because the thickness ratio to the wafer is only 0.2% in this work, the Stony’s equation should be sufficient to calculate the residual stress without causing serious errors [19]. In addition, the modulus of film is not involved in Eq. (1), therefore the measured residual stress is independent of mechanical properties of the film.

The Young’s moduli and hardness of the SiOC films were characterized by a TI 900 Triboindenter (Hysitron Inc.). A standard Berkovich tip (also from Hysitron Inc.) with a tip radius of approximately 150 nm and a half-angle of 65.35◦ was chosen for the experiments. The tip area function was calibrated with a standard fused silica sample. The tip was installed overnight before the day of testing in order to reach the thermal equilibrium and minimize the drift. Regular constant-rate loading nanoindentation tests (loading/unloading rate: 400 ␮N/s) were performed on SiOC films without the maximum load holding period. Each sample was indented at 6 different locations and the results were averaged. The Young’s modulus was measured by fitting the unloading segment of load–displacement (P–h) curve to an empirical power-law relation (Fig. 1) [15,20]: m

1.5

2

2.5

3

Fig. 2. EDX spectra of sputtered SiOC films (a.u.: arbitrary unit).

contact area and the measured unloading stiffness through the relation √ 1  S Er = (3) √ ˇ 2 A where Er is the reduced modulus, ˇ is the geometric constant for the tip, S = dP/dh is the experimentally measured stiffness of the upper portion of the unloading data, and A is contact area. For a Berkovich tip, ˇ equals 1.05 [20]. The reduced modulus Er is related to the Young’s modulus and Poisson’s ratio of the indented material (E, ) and the tip (Ei , i ) by 1 − i2 1 1 − 2 + = Er E Ei

(4)

Meanwhile, the hardness of the films can be derived from the peak load Pmax divided by the projected contact area at peak load A:

2.6. Young’s modulus and hardness measurements

P = a(h − hc )

1

Energy (keV)

The curvature of the film–substrate system was measured by an Alpha-Step 500 Surface Profiler (KLA Tencor). Subsequently the curvature was used to calculate the residual stress in the film by the Stony’s equation [18] E  s ts2 f = ( − 0 ) 6tf

0.5

(2)

where hc is the contact depth, a and m are fitting constants. Measurement of the elastic modulus follows from its relationship to

H=

Pmax A

(5)

3. Results and discussion 3.1. Film composition analysis The stoichiometric composition of the deposited films was investigated by EDS technique. As the accelerating voltage used in this work 5 kV is different from the default in the quantitative analysis, it is crucial to use the correct accelerating voltage in the spectra condition during analysis. The EDS spectra of all five films are shown in Fig. 2. The elements C, O and Si were easily identified from the absorbance peaks at 0.25 keV, 0.5 keV and 1.7 keV, respectively. The corresponding atomic concentrations of Si, O and C in atomic concentration (at.%) were also summarized in Table 1. The Si content is in the range of 24.2–46.41 at.%, which agrees well with the 25–40 at.% reported by Ryan and Pantano [11]. The analysis indicates that with increasing carbon content in the SiOC films, the

P. Du et al. / Sensors and Actuators A 176 (2012) 90–98

93

Fig. 3. SEM micrographs of five as-deposited films (scale bar = 100 nm).

oxygen peaks continuously decrease while the carbon peaks continuously increase in the EDS spectra. The small amount of carbon (2.22 at.%) in the SiOx specimen could be due to the contamination which always contributes to the carbon peak. The small amount of oxygen (1.74 at.%) in the SiCy spectrum could be due to the absorption of free oxygen in the SiCy film surface. 3.2. Microstructure analysis The SEM micrographs of the SiOC films were obtained with the magnification of 100k. As shown in Fig. 3, the SiOx and SiCy have distinct surface characteristics. The surface of SiOx consists of relatively large clusters, while SiCy is closely packed with much finer clusters. The three SiOC films show a noticeable transition from silica-like loose structure to silicon carbide-like dense structure. The results agree with the density measurement from Ryan and Pantano [11]. In their report, the density of the SiOC films reduced

from 2.8 g/cm3 to 1.9 g/cm3 as the oxygen content increased. Therefore carbon-rich structures should be more rigid and denser than oxygen-rich structures, which will in turn influence the mechanical properties of the films. The detailed mechanical characterization will be discussed in Section 3.3. Since sputtering is a relatively low temperature process, the deposited films on unheated substrates are generally amorphous [11,21]. For the non-stoichiometric amorphous silicon oxynitride films, the random bonding model (RBM) has been proposed to depict the chemical structures [22]. A similar concept can be adopted to SiOx Cy films, according to the high-resolution X-ray photoelectron spectroscopy (XPS) [11]. In this model, silicon atoms randomly bond with oxygen and carbon atoms to form a homogeneous O–Si–C network with five different tetrahedral coordinates, in which y = 4 − x and x may vary from 0 to 4. Fig. 4 presents the FTIR absorption spectra of SiOC films with various carbon contents. The spectra exhibit characteristic peaks

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P. Du et al. / Sensors and Actuators A 176 (2012) 90–98

(a) Residual stress (GPa)

Absorption (a.u.)

Si-O-Si stretching

Si-O-Si bending SiOx

SiOxCy1 SiOxCy2

0.5

0

-0.5

SiOxCy3

0

200

1000

800

Wavenumber (cm-1)

600

SiCy

Fig. 4. FTIR spectra of the as-deposited SiOC films (the spectra were vertically shifted for the sake of clarity).

centered at 780, 810, and 1060 cm−1 , corresponding to Si–C–Si stretching mode [23], Si–O–Si bending mode and Si–O–Si stretching mode [24], respectively. The peak associated with the stretching mode of Si–C–Si bonds is clearly visible in the SiCy film; this spectrum shows no detectable Si–O–Si peaks. While in the SiOx film, the spectrum is dominated by the stretching and bending modes of Si–O–Si bonds. As the carbon content decreases, the spectra show a continuous shift of the peak from silicon-carbide-like to silica-like structures. However, for the three SiOC films, the broad absorption peaks indicate that they are actually the combinations of overlapping peaks with different Si–O and Si–C variations. Therefore, the FTIR spectra confirm that the co-sputtered SiOC films could be represented by the RBM structure. 3.3. Residual stress The residual stress in the films was calculated from the curvature of the film–substrate system by using the Stoney’s equation. Before film deposition, the averaged initial curvature of the silicon wafer substrate was measured from the profiler as −2.05 × 10−2 m−1 . After deposition, the curvatures of SiOC films subjected to three RTA temperatures (400 ◦ C, 600 ◦ C and 800 ◦ C) were measured. The residual stress was calculated based on the properties of the wafer (E s = 180 GPa [25] and ts = 500 ␮m). The resulting residual stresses for different carbon contents under different annealing temperatures were plotted in Fig. 5. For the as-deposited films, the residual stresses are generally compressive for all films. The development of compressive stress is associated with “atomic shot peening” mechanism [26]. During sputtering the sputtered atoms are ejected with energies on the order of 10 eV. Moreover, the surface of the growing film is also subjected to bombardment by argon atoms, namely ions that are neutralized at the cathode and backscattered. Both of these types of bombardment are prominent at low argon pressure. With sufficient energy atoms may be forced into spaces too small to accommodate them under thermal equilibrium conditions. In such a film growing under bombardment, the number of bonds under compression is higher than that under tension. Consequently such a process puts the whole film under a state of compression.

(b) Residual stress (GPa)

1200

400

600

800

Temperature (oC)

Si-C-Si stretching

SiCy 1400

2.92 % 25.69 % 39.74 % 57.33 % 96.75 %

1

400 oC

1

600 oC

0.8

800 oC

0.6 0.4 0.2 0 -0.2 0

20

40

60

80

100

C/(C+O) in SiOC (%) Fig. 5. The influence of (a) annealing temperature and (b) carbon content [C/(C + O)] on the residual stresses of SiOC films (the films at minimum stress states are marked by black solid ellipses).

For the effect of annealing temperature, all five films experience increases in residual stress from compressive to tensile (Fig. 5(a)). Similar behavior has been reported by Chaker et al. for PECVD SiC films [27] and Cao et al. for PECVD SiOx films [13]. This is accompanied by a change in the material microstructure and correspondingly, the thermomechanical properties. To be more specific, we know that the sputtered films are generally amorphous due to the relative low deposition temperature. The atoms deposited on the substrate lack sufficient kinetic energy to reach a thermodynamic equilibrium state and therefore result in excess vacancies. Once they experience high temperature annealing, the increased mobility tends to drive the loosely ordered atoms to re-arrange to a relatively more ordered structure, and the concentration of the vacancies will be reduced [28]. Macroscopically the total volume decreases and this represents a net increase of tensile stress. This was confirmed from the thickness measurements. The three-layer model successfully captured the measured data, and the MSE was assured to be within the limit. The results show that the surface roughness of the films is less than 8 nm, and the thickness non-uniformity is less than 5%. The averaged decreases in thickness relative to the as-deposited films are plotted in Fig. 6. It implies that the films under different carbon contents all experience the reduction of thickness after the thermal annealing, and the amount of reduction increases with higher annealing temperatures. Moreover, SiCy is more sensitive to the post-deposition thermal treatment. The residual stress changes from compressive

P. Du et al. / Sensors and Actuators A 176 (2012) 90–98 Table 2 Modulus and hardness of as-deposited SiCy films.

0

Thickness decrease(nm)

95

SiCy

-10

Sputtered

Bulk single crystal [35]

-20 2.92 % 25.69 % 39.74 % 57.33 % 96.75 %

-30

-40

This work Ref. [23] Ref. [35]

0

200

E (GPa)

H (GPa)

272.08 ± 5.64 231 ± 25 280 ± 5.0 440–499

23.37 ± 1.02 28.1 ± 1.1 22.2 ± 0.2 32–36

3.4. Young’s modulus and hardness

400

600

800

Temperature (oC) Fig. 6. The decrease in thickness versus the annealing temperatures from the ellipsometry measurements. The thickness was relative to the thickness of as-deposited film at each carbon content.

(−290.6 MPa) at the as-deposited state to high tensile (770.9 MPa) after the 400 ◦ C RTA treatment. It indicates that using only thermal treatment is not sufficient to control the residual stress of silicon carbide films. For the effect of carbon content, the residual stresses of all RTA treated films show monotonic increases from compressive to tensile with increasing carbon contents (Fig. 5(b)). The thermal annealing consists of both heating and cooling steps, in other words, the net change of temperature is zero for a complete cycle. Therefore the thermal stress is zero for the annealed films, and any change in stress reflects the change in the intrinsic stress. In the RBM model, silicon atoms randomly bond with oxygen and carbon atoms to form a homogeneous O–Si–C network with five different tetrahedral coordinates, in which y = 4 − x and x may vary from 0 to 4. The O–Si–O covalent bonds inside the Si–O tetrahedral structure have a fixed arrangement. In comparison, the Si–O–Si bridging bonds connecting the tetrahedral units are flexible. Without inducing change in potential energy, the Si–O–Si bond angle (normally about 144◦ ) can be varied in a wide range from 110◦ to 160◦ , and the bond can also rotate freely around the Si–O axis [6]. Although the silicon carbide shares the common tetrahedral structure with silicon oxide, the Si–C–Si bridging bonds connecting the silicon–carbon tetrahedral units are very rigid due to the necessity of the carbon atom forming four equivalent bonds to silicon atoms rather than two bonds in Si–O–Si. The support of such statement is the FTIR peak width, which is a measure of the tetrahedral bond strain [29]. The broader peak reflects the broader distribution of vibration energies, and consequently more bond strain within the amorphous network. The peak width was represented by full width at half maximum (FWHM) obtained from the peak analyzer in Origin software package. The resulting FWHM of the stretching mode for Si–O–Si bond in SiOx and Si–C–Si bond in SiCy were 87.46 cm−1 and 261.93 cm−1 , respectively. Therefore, as the carbon content increases, more Si–C–Si bonds are formed. The strengthening of the interaction among Si and C atoms in the film leads to film contracting [30]. Accordingly, the film enrichment in Si–C–Si bonds results in a tensile stress. It is worth noting that the minimum stresses occurred for the 53.8% carbon content film at the as-deposited state (−36.93 MPa), 31.7% and 41.7% carbon content films after 400 ◦ C RTA treatments (−21.16 MPa and 20.43 MPa, respectively). In other words, we realized a “stress engineering”, through which the stress can be manipulated by controlling compositions and annealing temperatures.

Constant-rate of loading nanoindentation tests without the holding period were performed on all SiOC films. The representative load–displacement curves of the nanoindentation on SiOC films are shown in Fig. 7(a) and (b). The maximum indentation depth at the peak load of 2 mN in the SiOC films was less than 130 nm, which counts for approximately 13% of the film thickness (∼1 ␮m). Therefore, the effect of silicon substrate on the Young’s modulus and indentation-hardness could be negligible [6]. The discrete displacement bursts, or “pop-ins”, generally associated with structural changes in materials were not observed in the indentation loading segments. The Young’s modulus and hardness were extracted from load–displacement data of nanoindentation tests using Eqs. (3)–(5). The elastic constants of diamond Berkovich tip were Ei = 1141 GPa and vi = 0.07 [31]. The Poisson’s ratio of silicon oxide was 0.17 [32], and that of silicon carbide was 0.23 [23]. There are no reported Poisson’s ratios for silicon oxycarbides, therefore in this work we assume that they change linearly according to the carbon content of the samples. The effect of residual stress on the modulus and hardness varies for different types of materials. For high modulus-to-yield strength (E/ y ) materials (such as metals and alloys), the residual stress has large effect on the “pile-up” of contact area [33]. When pile-up occurs, the real contact area is greater than that predicted by the depth-dependent area function [20]. This can be rationalized using principles of mechanics applied to the initial stage of yielding during [31]. According to Eqs. (3) and (5), the change of contact area will consequently affect the modulus and hardness. However, the “pile-up” phenomenon rarely happens in low E/ y materials (such as glasses and ceramics). Bolshakov and Pharr proposed a convenient measurable parameter to identify the amount of pile-up [34]: the ratio of the final indentation depth hf to the maximum indentation depth hmax (hf /hmax ). Pile-up is large only when hf /hmax is close to 1, while very little pile-up is observed when hf /hmax < 0.7. In this work, hf /hmax of all indentation tests is less than 0.5, indicating that the “pile-up” is negligible and hence the effect of residual stress can be omitted. The calculation results were presented in Fig. 8(a) and (b). Both the Young’s modulus and hardness of as-deposited SiOC films show monotonic increases with increasing carbon contents. For instance, the Young’s modulus of SiOC films increased from 68.14 GPa to 272.08 GPa, and the hardness increased from 6.89 GPa to 23.37 GPa. The measured Young’s modulus of SiOx 68.14 ± 1.79 GPa, is in good agreement with that of silicon dioxide as obtained by the microindentation test (E = 67 ± 1.5 GPa) [32]. The measured Young’s modulus of SiCy 272.08 ± 5.64 GPa, is comparable to that of sputtered SiC reported by Khakani et al. (231 ± 25 GPa) [23] and Kulikovsky et al. (280 ± 5.0 GPa) [35]. The hardness falls into the similar range of the literatures (Table 2). However, both the modulus and hardness are lower than those of bulk SiC single crystal. Since the sputtered deposited films are generally amorphous, it is therefore more likely to form numerous point defects. In addition, the reduction of the short-range order and strength of inter-atomic bonding due to the ion bombardment could also contribute to the lower modulus and hardness.

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(a)

2

2 ASD o

400 C

1.5

Load (mN)

1.5

Load (mN)

(b)

0 % % 2.92 31.7 25.69%% 41.7 39.74%% 57.33% % 53.8 96.75% % 93.6

1

o

600 C o

800 C

1

0.5

0.5

0

25

50

75

100

0

125

0

10

20

30

40

50

60

70

Displacement (nm)

Displacement (nm)

Fig. 7. Representative load–displacement curves of (a) as-deposited SiOC films at different carbon contents and (b) SiCy film at different annealing temperatures (ASD denotes the as-deposited film).

(b)

300

SiCy

Hardness (GPa)

Young's modulus (GPa)

(a) 350

250

SiOxCy3 200

SiOxCy2 150

SiOxCy1 100

40

30

SiCy SiOxCy3

20

SiOxCy2 SiOxCy1

10

SiOx

50

0

SiOx 200

400

600

0

800

Temperature (oC)

0

200

400

600

800

Temperature (oC)

Fig. 8. (a) The Young’s modulus and (b) hardness of SiOC films as functions of annealing temperature.

(a)

(b)

ASD o

600 C o

800 C

1200

1000

800

400 C

Absorption (a.u.)

Absorption (a.u.) 1400

ASD o

o

400 C

600

o

600 C o

800 C

1400

1200

830

Frequency shift (cm )

820

-1

(c)

1000

800

600

Wavenumber (cm-1)

Wavenumber (cm-1)

SiOxCy1

SiOx

810 800

SiOxCy2

790

SiOxCy3 780

SiCy

770 760

0

200

400

600

800

Temperature (oC) Fig. 9. Representative FTIR spectra for as-deposited and annealed (a) SiOx and (b) SiCy films. (c) The shift of the absorption peaks for all SiOC films.

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For the effect of thermal annealing, both the modulus and hardness generally increase with increasing annealing temperatures. SiCy is more sensitive to annealing than SiOx in both modulus and hardness. For instance, the relative increase of modulus (from the as-deposited state to 800 ◦ C annealing) for SiOx is 4.83%, and that for SiCy is 18.81%; the increase of hardness for SiOx is 13.35%, and that for SiCy is 54.81%. The increase of annealing temperature results in the growth of modulus and hardness could be due to the microstructure evolution. Therefore, FTIR absorption spectra of both the as-deposited and annealed SiOC films were measured and compared. The representative spectra of SiOx and SiCy were plotted in Fig. 9(a) and (b). It clearly shows that the absorption peaks shift to higher wavenumber (blue shift) in both spectra. The peak shifting for all SiOC films are summarized in Fig. 9(c). SiOx shifts the least with a slope of 4.8 × 10−3 cm−1 /◦ C, while SiCy has much larger shifting with a slope of 2.41 × 10−2 cm−1 /◦ C. Awad et al. has reported a similar behavior for CVD deposited SiC films [36]. The blue shift is correlated to the increase of Si–C bond density, which can be determined using the inverse cross section and the integrated absorbance of the Si–C stretching band. Moreover, the Si–C bond density can be related to the modulus and hardness by a constant-plus-linear relation. It indicates that thermal annealing at increasing temperatures leads to a more dense microstructure in the amorphous matrix. Further support of such hypothesis is the reduction of film thickness observed as a result of the increasing annealing temperature, since the decreased point defects and increased density could contribute to the enhancement of mechanical properties. It is worth noting that the hardness is more sensitive to annealing than the modulus, for both SiOx and SiCy films. Hardness depends not only on the bonding strength between atoms, but also on the microstructure that prevents plastic flow. Due to the much smaller cluster size and more densely packed structure, the strengthening cluster boundaries of SiCy can be responsible for the enhanced hardness. Indeed, we find that the hardness of the SiCy film at 600 ◦ C and 800 ◦ C annealing (34.53 GPa and 36.18 GPa, respectively), are comparable to the bulk SiC single crystal [35]. 4. Conclusion Amorphous SiOC films with varied composition of oxygen and carbon content (ranging from SiOx to SiCy ) were deposited by RF magnetron sputtering. The residual stress, Young’s modulus and hardness were measured and found to be closely correlated to the film composition and post-thermal annealing temperatures. The change from compressive to tensile stress after thermal annealing was attributed to the structural relaxation. The increasing stress with higher carbon content was due to the bonding configuration of Si–O–Si and Si–C–Si bonds. The modulus and hardness generally increased with higher carbon content and annealing temperatures, which can be related to the bond density change and film densification. This work presents a versatile approach to control the mechanical properties of SiOC films, which will provide more opportunities for SiOC to be integrated into the MEMS field. We believe this microstructure-based theory will not only contribute to a better understanding of thermomechanical issues concerning the sputtered SiOC films, but also provide insights into the analysis of similar responses of other amorphous physical vapor deposited (PVD) or CVD materials. Acknowledgements This work was supported by the National Science Foundation in part under Grants CMMI-0700688 and ECCS-0901702.

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