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Tailored amorphous silicon carbide barrier dielectrics by nitrogen and oxygen doping
Thin Solid Films 531 (2013) 552–558
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Tailored amorphous silicon carbide barrier dielectrics by nitrogen and oxygen doping Yusuke Matsuda a, Sean W. King b, Reinhold H. Dauskardt a,⁎ a b
Stanford University, Stanford, CA 94305, USA Intel Corporation, Hillsboro, OR 97006, USA
a r t i c l e
i n f o
Article history: Received 14 April 2012 Received in revised form 28 November 2012 Accepted 29 November 2012 Available online 21 December 2012 Keywords: Structure–property relationship Mechanical properties Fracture Elastic constant Hybrid materials Dielectrics Moisture-sensitivity Thin films
a b s t r a c t The effects of N or O doping into hydrogenated amorphous silicon carbide (a-SiC:H) films on molecular structure and resulting material properties with particular attention to elastic constant, cohesive fracture energy, and moisture-assisted cracking were investigated. Fourier transform infrared spectroscopy and x-ray photoelectron spectroscopy characterizations demonstrated that doped N primarily formed Si\N and N\H bonds, and doped O formed Si\O suboxide bonds. The elastic constant of both N-doped a-SiC:H (a-SiCN:H) and O-doped a-SiC:H (a-SiCO:H) films increased with increasing N and O atomic concentrations (at.%). The cohesive fracture energy, Gc, of the a-SiCN:H and a-SiCO:H films also increased with increasing N and O at.%. These increases in the mechanical properties of the films were attributed to film densification with increasing N and O at.%. The a-SiCN:H films exhibited a greater increase in Gc than the a-SiCO:H films, which was due to the moisture-insensitivity of the a-SiCN:H films as opposed to the a-SiCO:H films. The a-SiCN:H films exhibited no moisture-assisted fracture behavior, which was attributed to moisture-insensitivity of Si\N bonds due to their less polar nature. © 2013 Elsevier B.V. All rights reserved.
1. Introduction Barrier dielectrics are commonly used in nanoscience and energy applications, such as microelectronic devices, nano/micro electromechanical systems, and photovoltaics. In particular, barrier dielectrics play an important role in low-k/Cu microelectric interconnect technologies where they serve multiple purposes from preventing Cu diffusion into low-k interlayer dielectrics (ILD) to reducing Cu electromigration to protecting Cu and low-k ILD from process damages [1,2]. Hydrogenated amorphous silicon carbide (a-SiC:H) films are especially attractive barrier dielectrics for the low-k/Cu interconnects because of their excellent thermal and chemical stabilities [3], and thermo-mechanical properties [4]. Perhaps, their most significant advantage is a remarkable insensitivity to a phenomenon called “moisture-assisted cracking” [4,5]. This occurs well below fracture resistance and leads to a catastrophic device failure through slow crack growth over time, and therefore is a serious reliability issue in devices containing materials sensitive to the phenomenon. a-SiC:H films, however, have relatively high dielectric constant, k, and low band gap, Eg ~ 3.0 eV, compared with traditional Si3N4 (k = 8, Eg ~ 5.5 eV) and SiO2 (k = 3.9, Eg ~ 9 eV) barrier dielectrics. Toward further reduction in scaling feature sizes and improved device reliability, lowering k and increasing Eg are critical while still maintaining
⁎ Corresponding author at: 496 Lomita Mall, Durand Bldg., Rm. 121 Stanford University Stanford, CA 94305, USA. Tel./fax: +1 650 725 0679, +1 650 725 4034. E-mail address: [email protected] (R.H. Dauskardt). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2012.11.141
the desirable properties of a-SiC:H films. One possible route to achieve this is to dope either N or O atoms into a-SiC:H films. It has been demonstrated that the resulting N or O-doped a-SiC:H films, hereafter referred to as a-SiCN:H and a-SiCO:H films, respectively, had increased Eg and decreased k [6,7]. On the other hand, it is not well understood how N or O doping changes the mechanical properties of a-SiC:H films, such as elastic constant, fracture resistance, and sensitivity to moisture-assisted cracking, which are crucial for their integration into devices and the reliability. N and O are trivalent and divalent, respectively, and thus a-SiCN:H and a-SiCO:H films can have very different glass network connectivity, which affects their elastic constant and fracture resistance. From a mechanochemical point of view, doped O can form polar Si\O bonds, whereas doped N can form less polar Si\N and C\N bonds. Such difference in bond nature will lead to a marked contrast in sensitivity to moisture-assisted cracking between a-SiCN:H and a-SiCO:H films. In the present study, we investigate the effects of N or O doping on the molecular structure and resulting mechanical properties of a-SiC:H films. We first characterize changes in the molecular structure with the doping using Fourier transform infrared spectroscopy (FTIR) and x-ray photoelectron spectroscopy (XPS). The doped N primarily forms Si\N and N\H bonds, and possibly some C\N bonds, and the doped O forms only Si\O suboxide bonds. We then characterize dielectric and optical properties, and subsequently mechanical properties of the films ranging from elastic constant to cohesive fracture resistance to sensitivity to moisture-assisted cracking. We believe that fundamental understanding of the effects of N or O doping on the molecular structure
Y. Matsuda et al. / Thin Solid Films 531 (2013) 552–558
and resulting material properties of the a-SiC:H films provides design guidelines for tailoring the films as advanced barrier dielectrics for future technology nodes in microelectronics as well as emerging nanoscience and energy technologies that require barrier dielectrics. 2. Experiment Films were deposited on a 300 mm (100) silicon wafer using a manufacturing plasma enhanced chemical vapor deposition (PECVD) system at 400 °C, similar to methods described elsewhere [6–8]. Base a-SiC:H films were synthesized using methylsilane precursors diluted in He, and H2 gases. In order to dope O and N into the a-SiC:H films, additional oxidizing and nitrodizing gases were used during the deposition process, respectively. Deposition time was adjusted to produce nominal film thickness of 500 nm, and the bottom and top sides of the films were capped with 25 nm silicon carbon nitride (SiCN) layers deposited by the same PECVD system. Low frequency (100 kHz) dielectric constants of the films were measured using a Hg probe. Film density, ρfilm, was determined by X-ray reflectivity technique [9]. The Young's modulus, Esaw, of the films was determined using picosecond ultrasonic technique [10] in which the ρfilm determined from the X-ray reflectivity and Poisson's ratio of 0.25 was used. These material properties are listed in Table 1. The molecular structures of the films were characterized by FTIR (Nicolet Magna-IR 860 and Bio-Rad QS-3300 spectrometers) with a transmission mode in the wavenumber range from 500 to 4000 cm−1. Details on FTIR characterization of the films are described elsewhere [11,12]. XPS measurements of the films were performed with VersaProbe (Physical Electronics). Before the measurements, all of the films were sputtered by Ar+ ion (5 keV at 3 μA) to remove the SiCN layer. Survey scans were carried out three times to characterize chemical compositions of the films. High resolution scans of Si 2p, C 1s, N 1s, and O 1s core photoelectrons were also performed 20 times to examine the bonding state in the films. The binding energy resolution for the high resolution scans was ~0.6 eV. Shifts in the binding energy due to surface charge were corrected by referring to the C 1s line of C\Si bond (283.2–3 eV [13]). The measured peaks were fitted using the combination of Gaussian and Lorentzian functions. The cohesive fracture energy, Gc, of the films was measured by the double cantilever beam (DCB) testing with a Delaminator Test System (DTS, Menlo Park, CA). In specimen preparation, the films with additional metal layers were bonded to silicon wafers with epoxy adhesives (Fig. 1), which were then diced by a high-speed wafer saw to fabricate DCB geometry specimens, 5 mm wide, 1.56 mm in total thickness, and 50 mm in length. The DCB specimens were loaded in pure mode I, and displacement was measured to determine Gc in a laboratory air test environment at ~ 25 °C and ~ 40% relative humidity
Table 1 Material properties and chemical composition of a-SiC:H, a-SiCN:H, and a-SiCO:H films. XRR film density and dielectric constant values have maximum–minimum errors of ±0.1. Film designation
XRR film density, ρ (g cm−3)
Dielectric constant, k
Refractive index, n
XPS composition (at.%) C
Si
O
N
a-SiC:H a-SiCN:H-1 a-SiCN:H-2 a-SiCN:H-3 a-SiCN:H-4 a-SiCO:H-1 a-SiCO:H-2 a-SiCO:H-3 a-SiCO:H-4 a-SiCO:H-5 a-SiCO:H-6
1.30 1.56 1.54 1.5 1.69 1.51 1.50 1.49 1.50 1.62 1.70
3.60 3.53 3.32 3.37 – 3.69 3.65 3.58 3.63 3.69 3.75
1.77 1.81 1.78 1.76 1.77 1.75 1.73 1.71 1.70 1.70 1.69
49.9 43.3 40.7 35.8 30.0 42.1 42.6 42.3 38.9 36.9 35.6
47.6 40.1 40.5 40.3 43.6 47.6 47.3 46.6 45.5 45.1 44.4
2.6 2.8 1.5 4.1 2.8 10.3 10.1 11.1 15.6 18.1 20.0
0 13.8 17.4 19.8 23.6 0 0 0 0 0 0
553
Si (775 µm) bonding layer (1 µm) SiCN (25 nm) film of interest (500 nm) SiCN (25 nm) Si (775 µm) Fig. 1. Schematic illustration of multi-layered thin film structure (not to scale).
(RH). The method for calculating Gc is detailed elsewhere [14]. After fracture energy measurements, all fracture paths were characterized by XPS survey scan to identify the fracture paths. To characterize sensitivity to moisture-assisted fracture behavior of the films, crack growth velocity, v, was measured with the DCB technique at 25 °C at different humidities in a chamber. To measure v, DCB specimens were loaded at a constant loading rate (2 μm sec−1) to a predetermined load at which the displacement was fixed. Then, automated analysis of the load relaxation with increasing compliance of the specimens determined the v over the range from ~10−11 to ~10−4 m s−1 as a function of applied strain energy release rate, G, to produce a characteristic v-G curve. General method of crack growth measurements is described elsewhere [15]. 3. Results and discussion 3.1. Molecular structure and material characterization 3.1.1. FTIR The molecular structure of the films was characterized using FTIR in the wavenumber range from 500 to 4000 cm −1 (Fig. 2A–B). Detailed absorbance peaks between 500 and 1200 cm −1 are also shown in Fig. 3A–B. Main absorbance peaks in all of the films consist of Si\C stretching modes at ~ 790 cm −1 [11,16,17], Si\CH2\Si wagging modes and Si\O asymmetric stretching modes between 1000 and 1050 cm −1 [11,17], S\Hx stretching modes between 2000 and 2200 cm −1 [16,17], and C\Hx stretching modes between 2850 and 3100 cm −1 [16,17,18,19]. The relatively minor absorbance peaks consist of Si\CH3 symmetric bending modes at ~ 1260 cm −1 [11,17], Si\CH2\Si symmetric bending modes at ~ 1350 cm −1 [11,17], and Si\CH3 asymmetric bending modes at ~1400 cm−1 [11,17]. In a-SiCN: H films, there are additional peaks corresponding to Si\N stretching modes between 870 and 950 cm−1 [20–22], N\H bending modes at ~1150 cm−1[20–22], and N\Hx stretching modes at ~3400 cm−1 [20,22]. In the a-SiCN:H films, as N at.% increased, the peak intensity corresponding to Si\N stretching modes (~870–950 cm−1) and N\H stretching modes (~870–950, ~3400 cm−1) increased (Figs. 2A and 3A). This indicates that the incorporated N in the a-SiCN:H films primarily formed Si\N and N\H bonds. In the a-SiCO:H films, as O at.% increased, the peak intensity corresponding to Si\CH2\Si/Si\O\Si stretching modes increased (Figs. 2B and 3B), indicating that the incorporated O formed Si\O\Si bonds. In common with both of the films, N and O doping resulted in decreasing peak intensity corresponding to Si\Hx stretching modes. Additionally, as N at.% increased, the position
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A Si-N Si-CH2-Si/Si-O N-H C-H3
Absorbance (arb. units)
Si-C
N at.% Si-Hx
C-Hx
N-Hx a-SiCN:H-4 a-SiCN:H-3 a-SiCN:H-2 a-SiCN:H-1 a-SiC:H
500
1000 1500 2000 2500 3000 3500 4000
Wavenumber (cm-1)
B Si-C
Si-CH2-Si/Si-O
Absorbance (arb. units)
O at.% C-H3 C-Hx
Si-CH2-Si Si-Hx
a-SiCO:H-6 a-SiCO:H-5 a-SiCO:H-4 a-SiCO:H-3 a-SiCO:H-2 a-SiCO:H-1 a-SiC:H
500
of the peak corresponding to Si\Hx stretching modes shifted toward higher values of wavenumber. We note that this shift is not because of the formation of C`N bonds, which has been designated to ~2200 cm−1 [23], but rather because of the formation of N\Si bonds [24,25]. As shown in the following section, C`N bonds were not detected by XPS high resolution scans in the films.
1000 1500 2000 2500 3000 3500 4000
Wavenumber (cm-1) Fig. 2. FTIR spectra from 500 to 4000 cm−1 (A) a-SiCN:H; and (B) a-SiCO:H films.
A Si-N
Si-C
Absorbance (arb. units)
Si-CH2-Si/Si-O N at.% N-H a-SiCN:H-4
a-SiCN:H-3 a-SiCN:H-2 a-SiCN:H-1
3.1.2. XPS The spectrum of Si 2p, C 1s, N 1s, and O 1s core photoelectrons in the films was characterized by XPS high resolution scans (Fig. 4A–G). The XPS Si 2p core spectrum of the films with varying N and O at.% is shown in Fig. 4A–B. The Si 2p spectrum of the a-SiCN:H films are composed of three peaks corresponding to Si\H/Si\Si (~ 99.3 eV) [13], Si\C (~ 100.35 eV) [13,26], and Si\N bonds (~101.2 eV) (Fig. 4A). As N at.% increased, the peak intensity for Si\N bonds increased, which is consistent with the FTIR results (Figs. 2A and 3A). The Si 2p spectrum of the a-SiCO:H films consists of three peaks for Si\H/Si\Si (~ 99.3 eV), Si\C (~ 100.35 eV), and Si\Ox (x b 2, 101.5 eV) bonds (Fig. 4B) [27]. As O at. % increased, the peak intensity corresponding to Si\Ox bonds also increased similar to the FTIR results (Figs. 2B and 3B). The XPS C 1s core spectrum of the films with varying N and O at.% is shown in Fig. 4C–D. The C 1s spectrum of the a-SiCN:H films is decomposed into two main peaks corresponding to C\Si (~283.3 eV) [13,26] and sp3 C\C bonds (284.4–5 eV) [13], and one very weak peak possibly corresponding to C\N bonds with trivalent N configuration (~285.6 eV) [28] (Fig. 4C). The area under the weak peak at ~285.6 eV consistently increased from 2% to 6% with increasing N at.% (Table 2), suggesting that some of the doped N formed C\N bonds. The C 1s spectrum of the a-SiCO:H films is composed of two peaks corresponding to C\Si (~283.3 eV) and sp3 C\C bonds (284.4–5 eV) (Fig. 4D). The XPS N 1s core spectrum of the a-SiCN:H films with varying N at% is shown in Fig. 4e. The N 1s spectrum of the films consists of two peaks for N\Si bonds with trivalent N configuration (397.2–397.5 eV) [28] and N\H or N\C bonds or both with trivalent N configuration (398.3–398.94 eV) [28]. The XPS O 1s core spectrum of the films with varying N and O at.% is shown in Fig. 4F–G. The O 1s peak was observed at ~532.2 eV. The summary of XPS core spectrum measurements (peak position, peak area fraction, and FWHM) is listed in Table 2. These observations enable us to understand how doped N and O form bonds in the a-SiC: H films. In the a-SiCN:H films, doped N forms primarily Si\N and N\H bonds, and possibly a fraction of C\N bonds. In the SiCO:H films, doped O forms only Si\O suboxide bonds.
a-SiC:H
500
600
700
800
3.1.3. Film density The film density, ρfilm, increased with increasing N and O at.% in the films (Table 1), which clearly demonstrates that the films experienced higher degree of cross-linking in the glass network. The film densification with increasing N and O at.% is consistent with the decreasing and increasing FTIR peak intensity corresponding to Si\Hx bond stretching modes and network Si\N and Si\CH2\Si bonds/Si\O\Si bonds, respectively (Figs. 2 and 3).
900 1000 1100 1200 -1
Wavenumber (cm )
B Si-O-Si/Si-CH2-Si
Absorbance (arb.units)
Si-C
O at.%
a-SiCO:H-6 a-SiCO:H-5 a-SiCO:H-4 a-SiCO:H-3 a-SiCO:H-2 a-SiCO:H-1 a-SiC:H
600
800
1000
1200
Wavenumber (cm-1) Fig. 3. FTIR spectra from 500 to 1200 cm−1 of (A) a-SiCN:H; and (B) a-SiCO:H films.
3.1.4. Dielectric constant and refractive index Dielectric constant, k, of the films measured by the Hg probe technique is listed in Table 1. The k of the a-SiCO:H films remained almost constant with increasing O at.%, and that of the a-SiCN:H films even decreased with increasing N at.% in spite of the film densification and more polar nature of Si\O and Si\N bonds than Si\C bonds. These observations are surprising because k generally depends on film density and bond polarity [29]. This may be related to a decreasing Si\Hx bond density in both of the a-SiCN:H and a-SiCO:H films with increasing N and O at.%. It is possible that the Si\Hx bonds can react with diffused water molecules into the films to form polar
Y. Matsuda et al. / Thin Solid Films 531 (2013) 552–558
Si\OH bonds with a very high dielectric constant [29]. The refractive index of the films was also measured (Table 1). The changes in the values of the refractive index with increasing N and O at.% were consistent with those of k. 3.2. Mechanical properties The elastic constant, Esaw, of the films measured by the surface acoustic wave technique is shown as a function of ρfilm (Fig. 5a–b). The a-SiCN:H films exhibited an increase in Esaw from 8.1 to 13.1 GPa with increasing ρfilm. The a-SiCO:H films also exhibited a greater increase in Esaw from 8.1 to 23.4 GPa with increasing ρfilm. These increases in Esaw were due to increases in glass network connectivity with increasing ρfilm, similar to the cases demonstrated in
Si-C
Si 2p
C
Si-N
Si-H/Si-Si
other types of hybrid glasses [4,30,31]. However, it is unclear why the greater stiffening was observed in the a-SiCO:H films compared with the a-SiCN:H film in spite of the lower valence of O than that of N. One possible explanation may be related to differences in bond stiffness. Given the same network bond density, elastic constant is generally proportional to bond stiffness, which is usually greater in bonds having a higher bond dissociation energy. As the bond dissociation energy of Si\O suboxide bonds [32] is ~ two times higher than that of Si\N [32] and C\N bonds [33] (Table 3), which can lead to the greater Esaw of the a-SiCO:H films than that of the a-SiCN:H films. The cohesive fracture energy, Gc, of the films measured by the DCB testing is shown as a function of ρfilm (Fig. 5c–d). Both of the a-SiCN:H and a-SiCO:H films exhibited an increase in Gc with increasing ρfilm. The Gc increase in the a-SiCN:H was from 1.7 to 4.3 J m −2, which is
C 1s
C-Si
C-C
C-N a-SiCN:H-4
Intensity (counts s-1)
a-SiCN:H-4
a-SiCN:H-3
a-SiCN:H-2
a-SiCN:H-1
a-SiCN:H-3
Intensity (counts s-1)
A
555
a-SiCN:H-2
a-SiCN:H-1
a-SiC:H
a-SiC:H 98
100
102
282
104
Binding energy (eV)
B Si 2p
284
D
Si-C Si-Ox
C 1s
C-Si
a-SiCO:H-6
Si-H/Si-Si
286
C-C a-SiCO:H-6
a-SiCO:H-5
a-SiCO:H-3
a-SiCO:H-2
100
a-SiCO:H-5
Intensity (counts s-1)
Intensity (counts s-1)
a-SiCO:H-4
98
288
Binding energy (eV)
a-SiCO:H-4
a-SiCO:H-3
a-SiCO:H-2
a-SiCO:H-1
a-SiCO:H-1
a-SiC:H
a-SiC:H
102
Binding energy (eV)
104
282
284
286
288
Binding energy (cm-1)
Fig. 4. XPS spectra of a-SiCN:H and a-SiCO:H films. (A) Si 2p spectra of a-SiCN:H; (B) Si 2p spectra of a-SiCO:H; (C) C 1s spectra of a-SiCN:H; (D) C 1s spectra of a-SiCO:H; (E) N 1s spectra of a-SiCN:H; (F) O 1s spectra of a-SiCN:H; and (G) O 1s spectra of a-SiCO:H films.
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E
Table 2 Lists of peak position, peak area fraction, and FWHM of peak. Peak position shifts due to surface charge were corrected by referring to C 1s C\Si bond (283.2–3 eV). Peak area fraction was determined by integrating the area of each peak.
N-Si
N 1s
N-H/N-C
a-SiCN:H-4
Intensity (counts s-1)
Film
a-SiCN:H-3 a-SiCN:H-2 a-SiCN:H-1
394
396
398
400
402
Si 2p
a-SiC:H
a-SiCN:H-1
a-SiCN:H-2
Binding energy (eV)
F
a-SiCN:H-3
O 1s a-SiCN:H-4
a-SiCN:H-4
Intensity (counts s-1)
a-SiCO:H-1
a-SiCN:H-3
a-SiCN:H-2 a-SiCN:H-1 a-SiC:H 528
530
532
534
536
Binding energy (eV)
Peak position [eV] Peak area [%] FWHM [eV]
a-SiCO:H-2
a-SiCO:H-3
a-SiCO:H-4
a-SiCO:H-5
a-SiCO:H-6
C 1s
Si\H/Si\Si
Si\C
99.5 16.0 0.9 99.3 12.4 0.9 99.3 14.0 0.9 99.3 3.5 0.9 99.3 5.6 0.9 99.5 10.2 0.9 99.5 10.3 0.9 99.5 5.7 0.9 99.5 3.5 0.9 99.3 6.7 0.9 99.4 3.8 0.9
100.4 81.8 1.3 100.4 58.5 1.2 100.4 58.9 1.2 100.4 26.1 1.2 100.4 34.2 1.2 100.4 82.3 1.5 100.4 83.4 1.5 100.4 74.6 1.5 100.4 57.5 1.5 100.4 58.44 1.5 100.4 48.9 1.5
Si\N
Si\Ox
C\Si
C\C
–
101.5 2.2 0.8
283.3 82.8 1.4 283.3 74.7 1.5 283.2 84.6 1.6 283.3 73.3 1.68 283.2 64.7 1.5 283.3 84.4 1.4 283.2 76.8 1.4 283.3 83.1 1.5 283.3 77.4 1.5 283.2 82.3 1.5 283.3 83.5 1.5
284.5 17.2 1.2 284.5 22.9 1.2 284.4 12.7 1.1 284.5 23.6 1.2 284.4 29.2 1.3 284.5 15.7 1.2 284.4 23.2 1.3 284.4 16.9 1.3 284.4 22.6 1.4 284.4 17.8 1.2 284.4 16.5 1.3
101.2 29.1 1.3 101.2 27.1 1.3 101.2 70.4 1.6 101.2 60.2 1.5 –
–
–
–
–
–
–
–
–
– 101.6 7.5 1.0 101.5 6.3 1.1 101.6 19.8 1.4 101.6 39.0 1.6 101.5 34.9 1.6 101.5 47.3 1.7
C\N – 285.7 2.4 1.0 285.6 2.7 1.0 285.7 3.1 1.0 285.5 6.2 1.0 –
–
–
–
–
–
G O 1s
Intensity (counts s-1)
a-SiCO:H-6
a-SiCO:H-5
a-SiCO:H-4
a-SiCO:H-3 a-SiCO:H-2 a-SiCO:H-1 a-SiC:H 528
530
532
534
Binding energy (eV) Fig. 4 (continued).
536
greater than that in the a-SiCO:H films from 1.7 to 3.3 J m −2. These increases in Gc were due primarily to an increase in network bond density. In almost ideally brittle materials like the a-SiCO:H and a-SiCN:H films, Gc is proportional to the product of density of bond broken and energy required for breaking each bond. The film densification results in a higher density of bond broken, thereby increasing the Gc of the films. The greater Gc increase in the a-SiCN:H films compared with the a-SiCO:H films in spite of a lower bond dissociation energy of Si\N and C\N than Si\O bond is most likely attributed to the moistureinsensitivity of the a-SiCN:H films. As detailed in the following section, the a-SiCN:H films were found insensitive to moisture-assisted cracking as opposed to the a-SiOC:H films. It is known that materials sensitive to moisture-assisted cracking, such as silica glasses, exhibit a lower Gc in moist-environments than that in vacuum because of the chemical reaction between strained moisture-sensitive bonds and water molecules [34]. To examine how N and O doping affect the sensitivity of the films to moisture-assisted cracking, the crack growth velocity, v, was measured as a function of applied strain energy release rate, G, at 25 °C and at different humidities (20 and 70% RH for a-SiCO:H films, 40% RH for a-SiCN:H films). The a-SiCO:H films exhibited an increasing sensitivity to moisture-assisted cracking as more Si\O\Si suboxide bonds were formed with increasing O at.% [4]. Conversely, all of the a-SiCN:H films were found insensitive to moisture-assisted cracking. As shown in Fig. 6, all of the a-SiCN:H films exhibited sharp load
Y. Matsuda et al. / Thin Solid Films 531 (2013) 552–558
b
a
25 a-SiCN:H a-SiC:H
20 15 10
N at.%
5
Elastic constant, ESAW (GPa)
Elastic constant, ESAW (GPa)
25
a-SiCO:H a-SiC:H
20 15 10
O at.%
5 0
0 1.3
1.4
1.5
1.6
1.7
1.3
Film density, ρfilm (g cm-3)
1.4
1.5
1.6
1.7
Film density, ρfilm (g cm-3)
5 a-SiCN:H a-SiC:H 4 N at.%
3 2 1 0 1.3
1.4
1.5
1.6
1.7
Cohesive fracture energy, Gc (J m-2)
d
c Cohesive fracture energy, Gc (J m-2)
557
5 a-SiCO:H a-SiC:H 4 3 2
O at.%
1
0
1.3
1.4
1.5
1.6
1.7
Film density, ρfilm (g cm-3)
Film density, ρfilm (g cm-3)
Fig. 5. Mechanical properties of a-SiC:H, a-SiCN:H, and a-SiCO:H films. (a) Esaw of a-SiCN:H; (b) Esaw of a-SiCO:H; (c) Gc of a-SiCN:H; and (d) Gc of a-SiOC:H films.
4. Conclusion We have characterized the effects of O or N doping in a-SiC:H films on their molecular structure and resulting material properties with particular attention to mechanical properties. FTIR and XPS characterizations demonstrated that doped N primarily formed Si\N and N\H bonds, and doped O formed only Si\O suboxide bonds. The elastic constant of both a-SiCN:H and a-SiCO:H films increased with increasing N and O at.%, respectively. The cohesive fracture energy of the a-SiCN:H and a-SiCO:H films also increased with increasing N and O at.%, respectively. These increases in the mechanical properties of the films were due primarily to film densification. The greater increase in cohesive fracture energy of the a-SiCN:H films than that of the a-SiCO:H films was attributed to moisture-insensitivity of the a-SiCN:H films as opposed to the a-SiCON:H films. The moisture-insensitivity of the a-SiCN:H films was attributed to less polar nature of Si\N and C\N bonds. We believe that the findings provide basis for tailoring the
mechanical properties of the a-SiC:H films for nanoscience and energy applications where barrier dielectrics are needed.
Acknowledgment This work was supported by the U.S. Department of Energy, under contract no. DE-FG02-07ER46391. Y.M. was supported by a Heiwa Nakajima Foundation Fellowship and a Stanford Graduate Fellowship.
2.0
a-SiCN:H-2
1.6
Load (N)
drops at critical loads corresponding to Gc, which are clear indications of moisture-insensitivity of the films. This moisture insensitivity is mostly attributed to the insensitivity of Si\N and Si\C bonds to moisture-attack [35] and advantageous as barrier dielectrics.
1.2
0.8 Table 3 Bond dissociation energy of network bonds in a-SiC:H, a-SiCN:H, and a-SiCO:H films. Bond
Bond dissociation energy [kJ mol−1]
Si\C Si\N Si\O C\N
452 [28] 437 [28] 800 [28] 260–320 [29]
0.4 80
120
160
200
240
280
Displacement (μm) Fig. 6. Load-displacement curve of a-SiCN:H-2. Sharp load drops at critical load are indications of the insensitivity of the films to moisture-assisted cracking.
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Thin Solid Films journal homepage: www.elsevier.com/locate/tsf
Tailored amorphous silicon carbide barrier dielectrics by nitrogen and oxygen doping Yusuke Matsuda a, Sean W. King b, Reinhold H. Dauskardt a,⁎ a b
Stanford University, Stanford, CA 94305, USA Intel Corporation, Hillsboro, OR 97006, USA
a r t i c l e
i n f o
Article history: Received 14 April 2012 Received in revised form 28 November 2012 Accepted 29 November 2012 Available online 21 December 2012 Keywords: Structure–property relationship Mechanical properties Fracture Elastic constant Hybrid materials Dielectrics Moisture-sensitivity Thin films
a b s t r a c t The effects of N or O doping into hydrogenated amorphous silicon carbide (a-SiC:H) films on molecular structure and resulting material properties with particular attention to elastic constant, cohesive fracture energy, and moisture-assisted cracking were investigated. Fourier transform infrared spectroscopy and x-ray photoelectron spectroscopy characterizations demonstrated that doped N primarily formed Si\N and N\H bonds, and doped O formed Si\O suboxide bonds. The elastic constant of both N-doped a-SiC:H (a-SiCN:H) and O-doped a-SiC:H (a-SiCO:H) films increased with increasing N and O atomic concentrations (at.%). The cohesive fracture energy, Gc, of the a-SiCN:H and a-SiCO:H films also increased with increasing N and O at.%. These increases in the mechanical properties of the films were attributed to film densification with increasing N and O at.%. The a-SiCN:H films exhibited a greater increase in Gc than the a-SiCO:H films, which was due to the moisture-insensitivity of the a-SiCN:H films as opposed to the a-SiCO:H films. The a-SiCN:H films exhibited no moisture-assisted fracture behavior, which was attributed to moisture-insensitivity of Si\N bonds due to their less polar nature. © 2013 Elsevier B.V. All rights reserved.
1. Introduction Barrier dielectrics are commonly used in nanoscience and energy applications, such as microelectronic devices, nano/micro electromechanical systems, and photovoltaics. In particular, barrier dielectrics play an important role in low-k/Cu microelectric interconnect technologies where they serve multiple purposes from preventing Cu diffusion into low-k interlayer dielectrics (ILD) to reducing Cu electromigration to protecting Cu and low-k ILD from process damages [1,2]. Hydrogenated amorphous silicon carbide (a-SiC:H) films are especially attractive barrier dielectrics for the low-k/Cu interconnects because of their excellent thermal and chemical stabilities [3], and thermo-mechanical properties [4]. Perhaps, their most significant advantage is a remarkable insensitivity to a phenomenon called “moisture-assisted cracking” [4,5]. This occurs well below fracture resistance and leads to a catastrophic device failure through slow crack growth over time, and therefore is a serious reliability issue in devices containing materials sensitive to the phenomenon. a-SiC:H films, however, have relatively high dielectric constant, k, and low band gap, Eg ~ 3.0 eV, compared with traditional Si3N4 (k = 8, Eg ~ 5.5 eV) and SiO2 (k = 3.9, Eg ~ 9 eV) barrier dielectrics. Toward further reduction in scaling feature sizes and improved device reliability, lowering k and increasing Eg are critical while still maintaining
⁎ Corresponding author at: 496 Lomita Mall, Durand Bldg., Rm. 121 Stanford University Stanford, CA 94305, USA. Tel./fax: +1 650 725 0679, +1 650 725 4034. E-mail address: [email protected] (R.H. Dauskardt). 0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.tsf.2012.11.141
the desirable properties of a-SiC:H films. One possible route to achieve this is to dope either N or O atoms into a-SiC:H films. It has been demonstrated that the resulting N or O-doped a-SiC:H films, hereafter referred to as a-SiCN:H and a-SiCO:H films, respectively, had increased Eg and decreased k [6,7]. On the other hand, it is not well understood how N or O doping changes the mechanical properties of a-SiC:H films, such as elastic constant, fracture resistance, and sensitivity to moisture-assisted cracking, which are crucial for their integration into devices and the reliability. N and O are trivalent and divalent, respectively, and thus a-SiCN:H and a-SiCO:H films can have very different glass network connectivity, which affects their elastic constant and fracture resistance. From a mechanochemical point of view, doped O can form polar Si\O bonds, whereas doped N can form less polar Si\N and C\N bonds. Such difference in bond nature will lead to a marked contrast in sensitivity to moisture-assisted cracking between a-SiCN:H and a-SiCO:H films. In the present study, we investigate the effects of N or O doping on the molecular structure and resulting mechanical properties of a-SiC:H films. We first characterize changes in the molecular structure with the doping using Fourier transform infrared spectroscopy (FTIR) and x-ray photoelectron spectroscopy (XPS). The doped N primarily forms Si\N and N\H bonds, and possibly some C\N bonds, and the doped O forms only Si\O suboxide bonds. We then characterize dielectric and optical properties, and subsequently mechanical properties of the films ranging from elastic constant to cohesive fracture resistance to sensitivity to moisture-assisted cracking. We believe that fundamental understanding of the effects of N or O doping on the molecular structure
Y. Matsuda et al. / Thin Solid Films 531 (2013) 552–558
and resulting material properties of the a-SiC:H films provides design guidelines for tailoring the films as advanced barrier dielectrics for future technology nodes in microelectronics as well as emerging nanoscience and energy technologies that require barrier dielectrics. 2. Experiment Films were deposited on a 300 mm (100) silicon wafer using a manufacturing plasma enhanced chemical vapor deposition (PECVD) system at 400 °C, similar to methods described elsewhere [6–8]. Base a-SiC:H films were synthesized using methylsilane precursors diluted in He, and H2 gases. In order to dope O and N into the a-SiC:H films, additional oxidizing and nitrodizing gases were used during the deposition process, respectively. Deposition time was adjusted to produce nominal film thickness of 500 nm, and the bottom and top sides of the films were capped with 25 nm silicon carbon nitride (SiCN) layers deposited by the same PECVD system. Low frequency (100 kHz) dielectric constants of the films were measured using a Hg probe. Film density, ρfilm, was determined by X-ray reflectivity technique [9]. The Young's modulus, Esaw, of the films was determined using picosecond ultrasonic technique [10] in which the ρfilm determined from the X-ray reflectivity and Poisson's ratio of 0.25 was used. These material properties are listed in Table 1. The molecular structures of the films were characterized by FTIR (Nicolet Magna-IR 860 and Bio-Rad QS-3300 spectrometers) with a transmission mode in the wavenumber range from 500 to 4000 cm−1. Details on FTIR characterization of the films are described elsewhere [11,12]. XPS measurements of the films were performed with VersaProbe (Physical Electronics). Before the measurements, all of the films were sputtered by Ar+ ion (5 keV at 3 μA) to remove the SiCN layer. Survey scans were carried out three times to characterize chemical compositions of the films. High resolution scans of Si 2p, C 1s, N 1s, and O 1s core photoelectrons were also performed 20 times to examine the bonding state in the films. The binding energy resolution for the high resolution scans was ~0.6 eV. Shifts in the binding energy due to surface charge were corrected by referring to the C 1s line of C\Si bond (283.2–3 eV [13]). The measured peaks were fitted using the combination of Gaussian and Lorentzian functions. The cohesive fracture energy, Gc, of the films was measured by the double cantilever beam (DCB) testing with a Delaminator Test System (DTS, Menlo Park, CA). In specimen preparation, the films with additional metal layers were bonded to silicon wafers with epoxy adhesives (Fig. 1), which were then diced by a high-speed wafer saw to fabricate DCB geometry specimens, 5 mm wide, 1.56 mm in total thickness, and 50 mm in length. The DCB specimens were loaded in pure mode I, and displacement was measured to determine Gc in a laboratory air test environment at ~ 25 °C and ~ 40% relative humidity
Table 1 Material properties and chemical composition of a-SiC:H, a-SiCN:H, and a-SiCO:H films. XRR film density and dielectric constant values have maximum–minimum errors of ±0.1. Film designation
XRR film density, ρ (g cm−3)
Dielectric constant, k
Refractive index, n
XPS composition (at.%) C
Si
O
N
a-SiC:H a-SiCN:H-1 a-SiCN:H-2 a-SiCN:H-3 a-SiCN:H-4 a-SiCO:H-1 a-SiCO:H-2 a-SiCO:H-3 a-SiCO:H-4 a-SiCO:H-5 a-SiCO:H-6
1.30 1.56 1.54 1.5 1.69 1.51 1.50 1.49 1.50 1.62 1.70
3.60 3.53 3.32 3.37 – 3.69 3.65 3.58 3.63 3.69 3.75
1.77 1.81 1.78 1.76 1.77 1.75 1.73 1.71 1.70 1.70 1.69
49.9 43.3 40.7 35.8 30.0 42.1 42.6 42.3 38.9 36.9 35.6
47.6 40.1 40.5 40.3 43.6 47.6 47.3 46.6 45.5 45.1 44.4
2.6 2.8 1.5 4.1 2.8 10.3 10.1 11.1 15.6 18.1 20.0
0 13.8 17.4 19.8 23.6 0 0 0 0 0 0
553
Si (775 µm) bonding layer (1 µm) SiCN (25 nm) film of interest (500 nm) SiCN (25 nm) Si (775 µm) Fig. 1. Schematic illustration of multi-layered thin film structure (not to scale).
(RH). The method for calculating Gc is detailed elsewhere [14]. After fracture energy measurements, all fracture paths were characterized by XPS survey scan to identify the fracture paths. To characterize sensitivity to moisture-assisted fracture behavior of the films, crack growth velocity, v, was measured with the DCB technique at 25 °C at different humidities in a chamber. To measure v, DCB specimens were loaded at a constant loading rate (2 μm sec−1) to a predetermined load at which the displacement was fixed. Then, automated analysis of the load relaxation with increasing compliance of the specimens determined the v over the range from ~10−11 to ~10−4 m s−1 as a function of applied strain energy release rate, G, to produce a characteristic v-G curve. General method of crack growth measurements is described elsewhere [15]. 3. Results and discussion 3.1. Molecular structure and material characterization 3.1.1. FTIR The molecular structure of the films was characterized using FTIR in the wavenumber range from 500 to 4000 cm −1 (Fig. 2A–B). Detailed absorbance peaks between 500 and 1200 cm −1 are also shown in Fig. 3A–B. Main absorbance peaks in all of the films consist of Si\C stretching modes at ~ 790 cm −1 [11,16,17], Si\CH2\Si wagging modes and Si\O asymmetric stretching modes between 1000 and 1050 cm −1 [11,17], S\Hx stretching modes between 2000 and 2200 cm −1 [16,17], and C\Hx stretching modes between 2850 and 3100 cm −1 [16,17,18,19]. The relatively minor absorbance peaks consist of Si\CH3 symmetric bending modes at ~ 1260 cm −1 [11,17], Si\CH2\Si symmetric bending modes at ~ 1350 cm −1 [11,17], and Si\CH3 asymmetric bending modes at ~1400 cm−1 [11,17]. In a-SiCN: H films, there are additional peaks corresponding to Si\N stretching modes between 870 and 950 cm−1 [20–22], N\H bending modes at ~1150 cm−1[20–22], and N\Hx stretching modes at ~3400 cm−1 [20,22]. In the a-SiCN:H films, as N at.% increased, the peak intensity corresponding to Si\N stretching modes (~870–950 cm−1) and N\H stretching modes (~870–950, ~3400 cm−1) increased (Figs. 2A and 3A). This indicates that the incorporated N in the a-SiCN:H films primarily formed Si\N and N\H bonds. In the a-SiCO:H films, as O at.% increased, the peak intensity corresponding to Si\CH2\Si/Si\O\Si stretching modes increased (Figs. 2B and 3B), indicating that the incorporated O formed Si\O\Si bonds. In common with both of the films, N and O doping resulted in decreasing peak intensity corresponding to Si\Hx stretching modes. Additionally, as N at.% increased, the position
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A Si-N Si-CH2-Si/Si-O N-H C-H3
Absorbance (arb. units)
Si-C
N at.% Si-Hx
C-Hx
N-Hx a-SiCN:H-4 a-SiCN:H-3 a-SiCN:H-2 a-SiCN:H-1 a-SiC:H
500
1000 1500 2000 2500 3000 3500 4000
Wavenumber (cm-1)
B Si-C
Si-CH2-Si/Si-O
Absorbance (arb. units)
O at.% C-H3 C-Hx
Si-CH2-Si Si-Hx
a-SiCO:H-6 a-SiCO:H-5 a-SiCO:H-4 a-SiCO:H-3 a-SiCO:H-2 a-SiCO:H-1 a-SiC:H
500
of the peak corresponding to Si\Hx stretching modes shifted toward higher values of wavenumber. We note that this shift is not because of the formation of C`N bonds, which has been designated to ~2200 cm−1 [23], but rather because of the formation of N\Si bonds [24,25]. As shown in the following section, C`N bonds were not detected by XPS high resolution scans in the films.
1000 1500 2000 2500 3000 3500 4000
Wavenumber (cm-1) Fig. 2. FTIR spectra from 500 to 4000 cm−1 (A) a-SiCN:H; and (B) a-SiCO:H films.
A Si-N
Si-C
Absorbance (arb. units)
Si-CH2-Si/Si-O N at.% N-H a-SiCN:H-4
a-SiCN:H-3 a-SiCN:H-2 a-SiCN:H-1
3.1.2. XPS The spectrum of Si 2p, C 1s, N 1s, and O 1s core photoelectrons in the films was characterized by XPS high resolution scans (Fig. 4A–G). The XPS Si 2p core spectrum of the films with varying N and O at.% is shown in Fig. 4A–B. The Si 2p spectrum of the a-SiCN:H films are composed of three peaks corresponding to Si\H/Si\Si (~ 99.3 eV) [13], Si\C (~ 100.35 eV) [13,26], and Si\N bonds (~101.2 eV) (Fig. 4A). As N at.% increased, the peak intensity for Si\N bonds increased, which is consistent with the FTIR results (Figs. 2A and 3A). The Si 2p spectrum of the a-SiCO:H films consists of three peaks for Si\H/Si\Si (~ 99.3 eV), Si\C (~ 100.35 eV), and Si\Ox (x b 2, 101.5 eV) bonds (Fig. 4B) [27]. As O at. % increased, the peak intensity corresponding to Si\Ox bonds also increased similar to the FTIR results (Figs. 2B and 3B). The XPS C 1s core spectrum of the films with varying N and O at.% is shown in Fig. 4C–D. The C 1s spectrum of the a-SiCN:H films is decomposed into two main peaks corresponding to C\Si (~283.3 eV) [13,26] and sp3 C\C bonds (284.4–5 eV) [13], and one very weak peak possibly corresponding to C\N bonds with trivalent N configuration (~285.6 eV) [28] (Fig. 4C). The area under the weak peak at ~285.6 eV consistently increased from 2% to 6% with increasing N at.% (Table 2), suggesting that some of the doped N formed C\N bonds. The C 1s spectrum of the a-SiCO:H films is composed of two peaks corresponding to C\Si (~283.3 eV) and sp3 C\C bonds (284.4–5 eV) (Fig. 4D). The XPS N 1s core spectrum of the a-SiCN:H films with varying N at% is shown in Fig. 4e. The N 1s spectrum of the films consists of two peaks for N\Si bonds with trivalent N configuration (397.2–397.5 eV) [28] and N\H or N\C bonds or both with trivalent N configuration (398.3–398.94 eV) [28]. The XPS O 1s core spectrum of the films with varying N and O at.% is shown in Fig. 4F–G. The O 1s peak was observed at ~532.2 eV. The summary of XPS core spectrum measurements (peak position, peak area fraction, and FWHM) is listed in Table 2. These observations enable us to understand how doped N and O form bonds in the a-SiC: H films. In the a-SiCN:H films, doped N forms primarily Si\N and N\H bonds, and possibly a fraction of C\N bonds. In the SiCO:H films, doped O forms only Si\O suboxide bonds.
a-SiC:H
500
600
700
800
3.1.3. Film density The film density, ρfilm, increased with increasing N and O at.% in the films (Table 1), which clearly demonstrates that the films experienced higher degree of cross-linking in the glass network. The film densification with increasing N and O at.% is consistent with the decreasing and increasing FTIR peak intensity corresponding to Si\Hx bond stretching modes and network Si\N and Si\CH2\Si bonds/Si\O\Si bonds, respectively (Figs. 2 and 3).
900 1000 1100 1200 -1
Wavenumber (cm )
B Si-O-Si/Si-CH2-Si
Absorbance (arb.units)
Si-C
O at.%
a-SiCO:H-6 a-SiCO:H-5 a-SiCO:H-4 a-SiCO:H-3 a-SiCO:H-2 a-SiCO:H-1 a-SiC:H
600
800
1000
1200
Wavenumber (cm-1) Fig. 3. FTIR spectra from 500 to 1200 cm−1 of (A) a-SiCN:H; and (B) a-SiCO:H films.
3.1.4. Dielectric constant and refractive index Dielectric constant, k, of the films measured by the Hg probe technique is listed in Table 1. The k of the a-SiCO:H films remained almost constant with increasing O at.%, and that of the a-SiCN:H films even decreased with increasing N at.% in spite of the film densification and more polar nature of Si\O and Si\N bonds than Si\C bonds. These observations are surprising because k generally depends on film density and bond polarity [29]. This may be related to a decreasing Si\Hx bond density in both of the a-SiCN:H and a-SiCO:H films with increasing N and O at.%. It is possible that the Si\Hx bonds can react with diffused water molecules into the films to form polar
Y. Matsuda et al. / Thin Solid Films 531 (2013) 552–558
Si\OH bonds with a very high dielectric constant [29]. The refractive index of the films was also measured (Table 1). The changes in the values of the refractive index with increasing N and O at.% were consistent with those of k. 3.2. Mechanical properties The elastic constant, Esaw, of the films measured by the surface acoustic wave technique is shown as a function of ρfilm (Fig. 5a–b). The a-SiCN:H films exhibited an increase in Esaw from 8.1 to 13.1 GPa with increasing ρfilm. The a-SiCO:H films also exhibited a greater increase in Esaw from 8.1 to 23.4 GPa with increasing ρfilm. These increases in Esaw were due to increases in glass network connectivity with increasing ρfilm, similar to the cases demonstrated in
Si-C
Si 2p
C
Si-N
Si-H/Si-Si
other types of hybrid glasses [4,30,31]. However, it is unclear why the greater stiffening was observed in the a-SiCO:H films compared with the a-SiCN:H film in spite of the lower valence of O than that of N. One possible explanation may be related to differences in bond stiffness. Given the same network bond density, elastic constant is generally proportional to bond stiffness, which is usually greater in bonds having a higher bond dissociation energy. As the bond dissociation energy of Si\O suboxide bonds [32] is ~ two times higher than that of Si\N [32] and C\N bonds [33] (Table 3), which can lead to the greater Esaw of the a-SiCO:H films than that of the a-SiCN:H films. The cohesive fracture energy, Gc, of the films measured by the DCB testing is shown as a function of ρfilm (Fig. 5c–d). Both of the a-SiCN:H and a-SiCO:H films exhibited an increase in Gc with increasing ρfilm. The Gc increase in the a-SiCN:H was from 1.7 to 4.3 J m −2, which is
C 1s
C-Si
C-C
C-N a-SiCN:H-4
Intensity (counts s-1)
a-SiCN:H-4
a-SiCN:H-3
a-SiCN:H-2
a-SiCN:H-1
a-SiCN:H-3
Intensity (counts s-1)
A
555
a-SiCN:H-2
a-SiCN:H-1
a-SiC:H
a-SiC:H 98
100
102
282
104
Binding energy (eV)
B Si 2p
284
D
Si-C Si-Ox
C 1s
C-Si
a-SiCO:H-6
Si-H/Si-Si
286
C-C a-SiCO:H-6
a-SiCO:H-5
a-SiCO:H-3
a-SiCO:H-2
100
a-SiCO:H-5
Intensity (counts s-1)
Intensity (counts s-1)
a-SiCO:H-4
98
288
Binding energy (eV)
a-SiCO:H-4
a-SiCO:H-3
a-SiCO:H-2
a-SiCO:H-1
a-SiCO:H-1
a-SiC:H
a-SiC:H
102
Binding energy (eV)
104
282
284
286
288
Binding energy (cm-1)
Fig. 4. XPS spectra of a-SiCN:H and a-SiCO:H films. (A) Si 2p spectra of a-SiCN:H; (B) Si 2p spectra of a-SiCO:H; (C) C 1s spectra of a-SiCN:H; (D) C 1s spectra of a-SiCO:H; (E) N 1s spectra of a-SiCN:H; (F) O 1s spectra of a-SiCN:H; and (G) O 1s spectra of a-SiCO:H films.
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E
Table 2 Lists of peak position, peak area fraction, and FWHM of peak. Peak position shifts due to surface charge were corrected by referring to C 1s C\Si bond (283.2–3 eV). Peak area fraction was determined by integrating the area of each peak.
N-Si
N 1s
N-H/N-C
a-SiCN:H-4
Intensity (counts s-1)
Film
a-SiCN:H-3 a-SiCN:H-2 a-SiCN:H-1
394
396
398
400
402
Si 2p
a-SiC:H
a-SiCN:H-1
a-SiCN:H-2
Binding energy (eV)
F
a-SiCN:H-3
O 1s a-SiCN:H-4
a-SiCN:H-4
Intensity (counts s-1)
a-SiCO:H-1
a-SiCN:H-3
a-SiCN:H-2 a-SiCN:H-1 a-SiC:H 528
530
532
534
536
Binding energy (eV)
Peak position [eV] Peak area [%] FWHM [eV]
a-SiCO:H-2
a-SiCO:H-3
a-SiCO:H-4
a-SiCO:H-5
a-SiCO:H-6
C 1s
Si\H/Si\Si
Si\C
99.5 16.0 0.9 99.3 12.4 0.9 99.3 14.0 0.9 99.3 3.5 0.9 99.3 5.6 0.9 99.5 10.2 0.9 99.5 10.3 0.9 99.5 5.7 0.9 99.5 3.5 0.9 99.3 6.7 0.9 99.4 3.8 0.9
100.4 81.8 1.3 100.4 58.5 1.2 100.4 58.9 1.2 100.4 26.1 1.2 100.4 34.2 1.2 100.4 82.3 1.5 100.4 83.4 1.5 100.4 74.6 1.5 100.4 57.5 1.5 100.4 58.44 1.5 100.4 48.9 1.5
Si\N
Si\Ox
C\Si
C\C
–
101.5 2.2 0.8
283.3 82.8 1.4 283.3 74.7 1.5 283.2 84.6 1.6 283.3 73.3 1.68 283.2 64.7 1.5 283.3 84.4 1.4 283.2 76.8 1.4 283.3 83.1 1.5 283.3 77.4 1.5 283.2 82.3 1.5 283.3 83.5 1.5
284.5 17.2 1.2 284.5 22.9 1.2 284.4 12.7 1.1 284.5 23.6 1.2 284.4 29.2 1.3 284.5 15.7 1.2 284.4 23.2 1.3 284.4 16.9 1.3 284.4 22.6 1.4 284.4 17.8 1.2 284.4 16.5 1.3
101.2 29.1 1.3 101.2 27.1 1.3 101.2 70.4 1.6 101.2 60.2 1.5 –
–
–
–
–
–
–
–
–
– 101.6 7.5 1.0 101.5 6.3 1.1 101.6 19.8 1.4 101.6 39.0 1.6 101.5 34.9 1.6 101.5 47.3 1.7
C\N – 285.7 2.4 1.0 285.6 2.7 1.0 285.7 3.1 1.0 285.5 6.2 1.0 –
–
–
–
–
–
G O 1s
Intensity (counts s-1)
a-SiCO:H-6
a-SiCO:H-5
a-SiCO:H-4
a-SiCO:H-3 a-SiCO:H-2 a-SiCO:H-1 a-SiC:H 528
530
532
534
Binding energy (eV) Fig. 4 (continued).
536
greater than that in the a-SiCO:H films from 1.7 to 3.3 J m −2. These increases in Gc were due primarily to an increase in network bond density. In almost ideally brittle materials like the a-SiCO:H and a-SiCN:H films, Gc is proportional to the product of density of bond broken and energy required for breaking each bond. The film densification results in a higher density of bond broken, thereby increasing the Gc of the films. The greater Gc increase in the a-SiCN:H films compared with the a-SiCO:H films in spite of a lower bond dissociation energy of Si\N and C\N than Si\O bond is most likely attributed to the moistureinsensitivity of the a-SiCN:H films. As detailed in the following section, the a-SiCN:H films were found insensitive to moisture-assisted cracking as opposed to the a-SiOC:H films. It is known that materials sensitive to moisture-assisted cracking, such as silica glasses, exhibit a lower Gc in moist-environments than that in vacuum because of the chemical reaction between strained moisture-sensitive bonds and water molecules [34]. To examine how N and O doping affect the sensitivity of the films to moisture-assisted cracking, the crack growth velocity, v, was measured as a function of applied strain energy release rate, G, at 25 °C and at different humidities (20 and 70% RH for a-SiCO:H films, 40% RH for a-SiCN:H films). The a-SiCO:H films exhibited an increasing sensitivity to moisture-assisted cracking as more Si\O\Si suboxide bonds were formed with increasing O at.% [4]. Conversely, all of the a-SiCN:H films were found insensitive to moisture-assisted cracking. As shown in Fig. 6, all of the a-SiCN:H films exhibited sharp load
Y. Matsuda et al. / Thin Solid Films 531 (2013) 552–558
b
a
25 a-SiCN:H a-SiC:H
20 15 10
N at.%
5
Elastic constant, ESAW (GPa)
Elastic constant, ESAW (GPa)
25
a-SiCO:H a-SiC:H
20 15 10
O at.%
5 0
0 1.3
1.4
1.5
1.6
1.7
1.3
Film density, ρfilm (g cm-3)
1.4
1.5
1.6
1.7
Film density, ρfilm (g cm-3)
5 a-SiCN:H a-SiC:H 4 N at.%
3 2 1 0 1.3
1.4
1.5
1.6
1.7
Cohesive fracture energy, Gc (J m-2)
d
c Cohesive fracture energy, Gc (J m-2)
557
5 a-SiCO:H a-SiC:H 4 3 2
O at.%
1
0
1.3
1.4
1.5
1.6
1.7
Film density, ρfilm (g cm-3)
Film density, ρfilm (g cm-3)
Fig. 5. Mechanical properties of a-SiC:H, a-SiCN:H, and a-SiCO:H films. (a) Esaw of a-SiCN:H; (b) Esaw of a-SiCO:H; (c) Gc of a-SiCN:H; and (d) Gc of a-SiOC:H films.
4. Conclusion We have characterized the effects of O or N doping in a-SiC:H films on their molecular structure and resulting material properties with particular attention to mechanical properties. FTIR and XPS characterizations demonstrated that doped N primarily formed Si\N and N\H bonds, and doped O formed only Si\O suboxide bonds. The elastic constant of both a-SiCN:H and a-SiCO:H films increased with increasing N and O at.%, respectively. The cohesive fracture energy of the a-SiCN:H and a-SiCO:H films also increased with increasing N and O at.%, respectively. These increases in the mechanical properties of the films were due primarily to film densification. The greater increase in cohesive fracture energy of the a-SiCN:H films than that of the a-SiCO:H films was attributed to moisture-insensitivity of the a-SiCN:H films as opposed to the a-SiCON:H films. The moisture-insensitivity of the a-SiCN:H films was attributed to less polar nature of Si\N and C\N bonds. We believe that the findings provide basis for tailoring the
mechanical properties of the a-SiC:H films for nanoscience and energy applications where barrier dielectrics are needed.
Acknowledgment This work was supported by the U.S. Department of Energy, under contract no. DE-FG02-07ER46391. Y.M. was supported by a Heiwa Nakajima Foundation Fellowship and a Stanford Graduate Fellowship.
2.0
a-SiCN:H-2
1.6
Load (N)
drops at critical loads corresponding to Gc, which are clear indications of moisture-insensitivity of the films. This moisture insensitivity is mostly attributed to the insensitivity of Si\N and Si\C bonds to moisture-attack [35] and advantageous as barrier dielectrics.
1.2
0.8 Table 3 Bond dissociation energy of network bonds in a-SiC:H, a-SiCN:H, and a-SiCO:H films. Bond
Bond dissociation energy [kJ mol−1]
Si\C Si\N Si\O C\N
452 [28] 437 [28] 800 [28] 260–320 [29]
0.4 80
120
160
200
240
280
Displacement (μm) Fig. 6. Load-displacement curve of a-SiCN:H-2. Sharp load drops at critical load are indications of the insensitivity of the films to moisture-assisted cracking.
558
Y. Matsuda et al. / Thin Solid Films 531 (2013) 552–558
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