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The structures of low dielectric constant SiOC thin films prepared by direct and remote plasma enhanced chemical vapor deposition
Thin Solid Films 515 (2007) 5035 – 5039 www.elsevier.com/locate/tsf
The structures of low dielectric constant SiOC thin films prepared by direct and remote plasma enhanced chemical vapor deposition Jaeyeong Heo a , Hyeong Joon Kim a,⁎, JeongHoon Han b , Jong-Won Shon b a
School of Materials Science and Engineering, Seoul National University, San #56-1, Sillim 9-dong, Gwanak-gu, Seoul, 151-744, South Korea b Jusung Engineering, # 49, Neungpyeong-ri, Opo-eup, Gwangju-si, Kyunggi-do, 464-890, South Korea Available online 6 December 2006
Abstract Two structures of low dielectric constant (low-k) SiOC films were elucidated in this work. Low-k thin film by remote plasma mode was mainly composed of inorganic Si–O–Si backbone bonds and some oxygen atoms are partially substituted by CH3, which lowers k value. The host matrix of low-k thin films deposited by direct plasma mode, however, was mainly composed of organic C–C bonds and “M” and “D” moieties of organosilicate building blocks, and thus the low dipole and ionic polarizabilities were the important factors on lowering k value. © 2006 Elsevier B.V. All rights reserved. Keywords: Low-k; Structure; XPS; FT-IR
1. Introduction The development of reliable low-k thin films has received widespread attention from the semiconductor industry over the past few years for their application in ultra large scale integration (ULSI) inter-metal dielectric (IMD) [1]. As the circuit feature size shrinks below 0.18 μm, the increase of signal propagation delay due to interconnects becomes important. The smaller line dimensions increase the resistivity (R) of the metal lines and the narrower interline spacing increases the parasitic capacitance (C) between the lines. To reduce the RC delays, low-k dielectrics (k b 3.0) are being considered to replace the currently used SiO2 (k ∼ 4.2) for interlayer dielectric insulation. Considerable efforts have been recently spent on developing low-k materials using either plasma-enhanced chemical vapor deposition (PECVD) or spin-on deposition (SOD) techniques [2]. Among the CVD candidates, organosilicate glass (OSG) or SiOC materials appear to be well suited so that various organosilicon precursors, such as trimethylsilane (3MS) [3], tetramethylsilane (4MS) [4], bis-trimethylsilylmethane (BTMSM) [5], and tetramethylcyclotetrasiloxane (TMCTS) [6], have been examined to be adoptable to the back-end process ⁎ Corresponding author. E-mail address: [email protected] (H.J. Kim). 0040-6090/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2006.10.095
integration. The reported k values of these films are as low as 2.1, but the films are very weak because of their inherently porous structure; there is currently no dense low-k dielectric material with a dielectric constant below 2.5 [7]. In this respect, the understanding of low-k film's structure is one of the most important issues to solve the integration problems. This paper is focused on two different mechanisms of lowering dielectric constant of the films which were deposited in the same system but with different process conditions. 2. Experimental Low-k SiOC films were prepared on p-type (100) silicon substrates by radio frequency (13.56 MHz) Inductively Coupled Plasma (ICP) PECVD using a CH3-containing organosilicate precursor, bis-trimethylsilylmethane (BTMSM, C7H20Si2). The schematic drawing of the PECVD set-up we used throughout this study is shown in Fig. 1. Plasma was generated around the 3-turn coil connected to the r.f. generator. The distance between the coil and the substrate was 34 cm and that between the injector and the substrate was 5 cm. The chamber was pumped by a booster and a dry pump to a base pressure of 10− 3 Torr. The precursor was vaporized and carried with inert Ar gas from a thermostatic bubbler maintained at 30 °C to the reaction chamber.
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Fig. 1. The schematic drawing of ICP PECVD set-up.
Two different plasma excitation modes can be applied in the system. In low r.f. power range (60 W–300 W), the plasma is confined within the plasma chamber so that only the excitation gas is plasma excited. We call it as “remote plasma mode” hereafter. O2 as an oxidant gas and Ar as a dilute gas were fed into the chamber. We expect that low or moderate dissociation of the precursor will occur in this mode. In high plasma power region (700 W), the plasma extends into the reaction chamber and direct plasma dissociation of the monomer molecules will play an important role [8]. We call it as “direct plasma mode” hereafter. Ar gas was fed into the chamber through the quartz tube. Detailed experimental process condition is summarized in the Table 1. All films were post-deposition-annealed (PDA) at 450 °C in N2 ambient for 30 min, and p-Si/ low-k dielectric /Pt MetalInsulator-Semiconductor (MIS) capacitor structure was fabricated by depositing Pt-top electrodes using an electron beam evaporation method through a shadow mask. After Al was deposited as a back contact metal, annealing was performed at 400 °C for 30 min in a forming gas (5% H2 + 95% N2) ambient. A HP 4194A impedance meter and a HP 4140B picoammeter were used for capacitance–voltage (C–V) and current-density– voltage (J–V) measurements, respectively. The C–V measurement frequency was fixed at 1 MHz. We elucidated the structure of each film by XPS (Sigma Probe, ThermoVG) and FT-IR (JASCO FT/IR-660 plus) measurements. 3. Results and discussion
Fig. 2. FT-IR absorbance spectrum of BTMSM precursor.
assigned to Si–CH2–Si wagging vibration. This wagging peak is masked by Si–O–Si asymmetric stretching peak when it comes to oxygen-containing polymers, but the stretching peak is not as sharp as CH2 wagging peak [8,9]. In addition, no oxygen atom exists in BTMSM precursor. We can consequently conclude that the sharp peak at 1051 cm− 1 originated from Si–CH2–Si bonds. The inset of Fig. 2 is an enlargement ranging from 550 to 1000 cm− 1. It is identified that rocking vibration mode of CH3 and stretching vibration mode of Si–C also exists (686, 761, 782, 833, and 871 cm− 1), which confirms the structure of the precursor [6,8,9]. Fig. 3 shows the deposition rate behaviors of two different plasma modes. An Arrhenius plot for the deposition rate of both plasma mode films as a function of deposition temperature is shown in Fig. 3(a). The activation energy is calculated to be − 0.077 eV and − 0.070 eV for direct and remote plasma modes, respectively; relatively smaller energy for remote plasma mode film is ascribed to the fact that Si–O bond has larger tendency to be adsorbed to the growing film than Si–C bond at high temperature. The relatively high activation energy is attributed to larger desorption tendency of Si–C and C–C bonds at high temperature. The negative activation energy indicates that the reaction is absorption–desorption controlled [10,11]. Fig. 3(b) shows deposition rate change with plasma power for remote plasma mode films. The slope of deposition rate with plasma
FT-IR absorbance spectrum of the precursor in liquid phase at room temperature is shown in Fig. 2. Peaks at 833, 1249, 2896, and 2952 cm− 1 are CH3-related and a sharp peak at 1051 cm− 1 is Table 1 Deposition parameters for low-k films Power [W] Direct 700 mode Remote 60–300 Mode
Pressure [Torr]
Dilute Ar [sccm]
O2 [sccm]
Sub. temp. [°C]
0.2
50
b1
25–280
0.2
50
50
25
Fig. 3. The deposition rate behavior of (a) direct plasma mode films and (b) remote plasma mode films.
J. Heo et al. / Thin Solid Films 515 (2007) 5035–5039
Fig. 4. FT-IR spectra of (a) direct plasma mode films and (b) remote plasma mode films.
power is 6.9 up to 100 W region, but it decreased down to 2.0 from 100 W to 300 W. The change of deposition behavior seems to occur around 100 W. Fig. 4 shows FT-IR spectra of both plasma modes with different deposition conditions. The main building blocks for Organosilicate Glass (OSG) materials are “M ((CH3)3SiO1/2)”, “D ((CH3)2SiO2/2)”, “T (CH3SiO3/2)v, and “Q (SiO4/2)” groups corresponding to mono-, di-, tri-, and quad-oxygen-substituted silicon atoms [12]. The peak positions of the first three building blocks appear at 1250, 1260, and 1270 cm− 1, respectively. As shown in Fig. 4(a), Si(CH3)X symmetric stretching peak position for direct plasma mode films increases from 1255 cm− 1 to 1263 cm− 1 and the peak area decreases from 0.07 (a.u.) to 0.04 (a.u.) with deposition temperature. Strong absorptions at 790– 850 cm− 1 are assigned to Si(CH3)X (x = 1, 2, or 3) vibrations [6]. The absorptions at 804 cm− 1 and 844 cm− 1 are attributed to “D” and “M” vibrations, respectively. “M” to “D” peak intensity ratio also decreases from 90.7% to 58.4%. Abrupt decrease of FT-IR spectra at 1100 cm− 1 seems to be due to an artifact. FT-IR spectra for remote plasma mode films as a function of plasma power are shown in Fig. 4(b). Si(CH3)X absorption peaks at 804 and 844 cm− 1 increase with decreasing plasma power while relative intensities of Si–O–Si backbone structure peak decrease. Slight decrease in wavenumber of backbone peak (red shift) corresponds well with the result of Y. H. Kim et al. [13]; it means that carbon content incorporated in the film increases with decreasing plasma power. As an arrow depicted in Fig. 4(b), Si (CH3)X symmetric stretching peak at 1276 cm− 1 also follows the same increasing trend of Si(CH3)X peaks at 804 and 844 cm− 1. A peak at 1133 cm− 1 is attributed to larger angle Si–O–Si bonds in a cage structure with a bond angle of approximately 150°, which
Fig. 5. Changes of refractive indices and thickness loss before and after N2 PDA of (a) direct plasma mode films and (b) remote plasma mode films.
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Fig. 6. Changes of dielectric constant of (a) direct plasma mode films and (b) remote plasma mode films.
can lead to micropores and consequently a lower film density and dielectric constant [6,14]. From the FT-IR spectra comparison, it is concluded that two films for different plasma states have different main building blocks. Fig. 5 shows the changes of refractive indices and thickness loss before and after annealing at 450 °C in N2. Fig. 5(a) is for the direct plasma mode films as a function of plasma power. As deposition temperature increases, refractive index slightly increases up to 170 °C, but it abruptly soars above 170 °C. After PDA, refractive indices are unchanged in low temperature region, but it increases slightly above 150 °C. Thickness loss also shows drastic change; thickness loss of the film at room temperature is as high as 13.5% and it decreases down to 7–8% and eventually to about 3%. The generally reported thickness loss of low-k film is about 5–6% [5]. But A. Grill and Neumayer reported that thermally unstable CHx phase could be removed
Fig. 7. Si 2p XPS spectra of (a) direct plasma mode film (25 °C, k = 2.54), (b) remote plasma mode film (80 W, k = 2.9), and (c) thermal oxide.
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Fig. 8. The relative percentage change of Si building blocks.
and over 20% thickness loss could be obtained [6]. It seems that thermally unstable CHx radicals exist in the film at relatively low deposition temperature and well removed during N2 PDA so that thickness loss is higher. The changes of refractive indices and thickness loss for remote plasma mode films are shown in Fig. 5(b). The refractive indices of the films prepared by remote plasma mode are smaller than those of the films by direct mode. It has the lowest value of 1.424 at 80 W condition and gradually increases up to 1.438. All films show similar decrease in refractive index by about 0.3 after PDA and the thickness loss also follows the same V-shape change with plasma power. The highest thickness loss, however, is about 6.5% and it is smaller than those of the direct plasma mode films with deposition temperature below 100 °C. Meanwhile, the refractive indices of N2 annealed low-k films prepared by direct plasma mode are higher (n = 1.5–1.6) than those (n = 1.39–1.42) of films by remote plasma mode. It is explained from the replacement of skeleton SiO2 (RI = 1.46 at 633 nm) by SiC (RI ∼ 2.0 at 633 nm) [15]. The dielectric constant measured by C–V method after PDA is illustrated in Fig. 6. As shown in Fig. 6(a), dielectric constant of direct plasma mode films increases as increasing the deposition temperature. The film prepared at room temperature shows the lowest k value of 2.54. The dielectric constant change as a function of plasma power for remote plasma mode films is shown in Fig. 6(b). It also follows the V-shape trend as the refractive index and thickness loss do. The lowest k value obtained is 2.9 at plasma power 80 W and it increases drastically as plasma power increases from 100 to 300 W. It is explained as follows; as the plasma power increases to 100 W, the number of radicals generated by plasma increases so that deposition rate rapidly increases as well; the number eventually saturates at 100 W and it decreases above 100 W because cracking of BTMSM precursor and chemical etching by plasma start to dominate the whole process. To further confirm the structure of the films, the deposited films were characterized by XPS measurement very carefully. Fig. 7(a) is curve fitting of the Si 2p core level of each film. Fits were obtained by constraining the full width at halfmaximum (FWHM) equal to 1.45 [16]. To calibrate the
photoelectron binding energy, the C 1 s peak attributed to C–C bonds was set to 284.9 eV. The Si 2p spectrum was deconvoluted into five different moieties which are assigned as follows: “M” (101.2 eV), “D” (102.1 eV), “T” (102.8 eV), “Q” (103.4 eV) and OH related Si (104.5 eV). Binding energies that several authors [16–19] have reported are a little different from each other and fitting procedure is somewhat arbitrary, but it is meaningful enough to distinguish silicon in ‘organic’ or ‘inorganic’ states. XPS Si 2p curve fitting of direct plasma mode film (25 °C) is shown in Fig. 7(a) and that of remote plasma mode film (80 W) is shown in Fig. 7(b). Peak fitting of thermal oxide is inserted as a reference in Fig. 7(c). In both cases of Fig. 7(a) and (b), the binding energy is shifted toward lower binding energy compared to reference thermal oxide and the film of direct plasma mode is more shifted to lower energy. The chemical shift of Si 2p core level is highly dependent on the presence of surrounding oxygen atoms [20]. The presence of oxygen atom makes the higher binding energy shift compared to the presence of carbon. Fig. 8 summaries the result more clearly. “M” and “D” moieties are the main building blocks (85.7%) of direct plasma mode film, and “T” and “Q” moieties are the main building blocks (58.5 %) of remote plasma mode film. Regarding FT-IR spectra and XPS fitting results, the structures of two plasma mode films are understood as follows; films by direct plasma mode are mainly composed of organic C–C bonds and “M” and “D” moieties. Although electronic k value is higher than that of remote plasma mode films, ionic and dipolar proportions of direct plasma mode films are much smaller than those of remote plasma mode films; it is speculated from the smaller electronegativity difference of Si–C (0.7) than that of Si–O (1.7). X. Li et al. also insisted that the large number of low atomic mass C–C bonds leads to small atomic polarization, therefore low dielectric constants [21]. On the other hand, films by remote plasma mode are mainly composed of inorganic Si–O–Si bonds with some oxygen atoms substituted by CH3, resulting in micropores. The decrease of ionic dipolar contribution is the predominant factor on decrease in the dielectric constant [22]. 4. Conclusion We have demonstrated that the films with structure different from generally considered organosilicate glass can also have low dielectric constant as low as 2.54 by optimizing plasma mode and process condition. The host matrix of the film was mainly composed of organic C–C bonds and “M” and “D” moieties, and thus the low dipole and ionic polarizabilities were the important factors on lowering k value. The low-k film of remote plasma mode, however, is mainly composed of Si–O–Si backbone and some of oxygen atoms are substituted by CH3 groups, which is the general mechanism of lowering dielectric constant. Acknowledgement The first author wishes to thank the Jusung Engineering for a financial support.
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References [1] K. Maex, M.R. Baklanov, D. Shamiryan, F. Iacopi, S.H. Brongersma, Z.S. Yanovitskaya, J. Appl. Phys. 93 (2003) 8794. [2] L. Peters, Semicond. Int. (Sept. 1998) 56. [3] M.J. Loboda, Microelectron. Eng. 50 (2000) 15. [4] A. Grill, V. Patel, J. Appl. Phys. 85 (1999) 3314. [5] Y.H. Kim, S.K. Lee, H.J. Kim, J. Vac. Sci. Technol., A 18 (2000) 1216. [6] A. Grill, D.A. Neumayer, J. Appl. Phys. 94 (2003) 6697. [7] K. Mosig, T. Jacobs, K. Brennan, M. Rasco, J. Wolf, R. Augur, Microelectron. Eng. 64 (2002) 11. [8] C. Rau, W. Kulisch, Thin Solid Films 249 (1994) 28. [9] A. Lee Smith (Ed.), Analysis of Silicones, Wiley–Interscience, New York, 1974, p. 265. [10] B.K. Hwang, M.J. Loboda, G.A. Cerny, R.F. Schneider, J.A. Seifferly, T. Washer, Proc. IEEE, IITC, 2002, p. 52. [11] J.M. Shieh, K.C. Tsai, B.T. Dai, Appl. Phys. Lett. 81 (2002) 1294. [12] A.D. Ross, K.K. Gleason, J. Appl. Phys. 97 (2005) 113707-1.
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[13] Y.H. Kim, M.S. Hwang, H.J. Kim, J. Appl. Phys. 90 (2001) 3367. [14] M.R. Wang, Rusli, J.L. Xie, N. Babu, C.Y Li, K. Rakesh, J. Appl. Phys. 96 (2004) 829. [15] D. Shamiryan, K. Weidner, W.D. Gray, M.R. Baklanov, S. Vanhaelemeersch, K. Maex, Microelectron. Eng. 64 (2002) 361. [16] L.M. Han, J. Pan, S. Chen, N. Balasubramanian, J. Shi, L.S. Wong, P.D. Foo, J. Electrochem. Soc. 148 (2001) F148. [17] L. O'Hare, B. Parbhoo, S.R. Leadley, Surf. Interface Anal. 36 (2004) 1427. [18] M.R. Alexander, R.D. Short, F.R. Jones, W. Michaeli, C.J. Blomfield, Appl. Surf. Sci. 137 (1999) 179. [19] S. Roualdes, R. Berjoan, J. Durand, Sep. Purif. Technol. 25 (2001) 391. [20] L.C. Feldman, J.W. Mayer, Fundamentals of Surface and Thin Film Analysis, Elsevier Sicence, New York, 1986, p. 224. [21] X. Li, T.K.S. Wong, Rusli, D. Yang, Diamond Relat. Mater. 12 (2003) 963. [22] J.Y. Kim, M.S. Hwang, Y.H. Kim, H.J. Kim, Y. Lee, J. Appl. Phys. 90 (2001) 2469.
The structures of low dielectric constant SiOC thin films prepared by direct and remote plasma enhanced chemical vapor deposition Jaeyeong Heo a , Hyeong Joon Kim a,⁎, JeongHoon Han b , Jong-Won Shon b a
School of Materials Science and Engineering, Seoul National University, San #56-1, Sillim 9-dong, Gwanak-gu, Seoul, 151-744, South Korea b Jusung Engineering, # 49, Neungpyeong-ri, Opo-eup, Gwangju-si, Kyunggi-do, 464-890, South Korea Available online 6 December 2006
Abstract Two structures of low dielectric constant (low-k) SiOC films were elucidated in this work. Low-k thin film by remote plasma mode was mainly composed of inorganic Si–O–Si backbone bonds and some oxygen atoms are partially substituted by CH3, which lowers k value. The host matrix of low-k thin films deposited by direct plasma mode, however, was mainly composed of organic C–C bonds and “M” and “D” moieties of organosilicate building blocks, and thus the low dipole and ionic polarizabilities were the important factors on lowering k value. © 2006 Elsevier B.V. All rights reserved. Keywords: Low-k; Structure; XPS; FT-IR
1. Introduction The development of reliable low-k thin films has received widespread attention from the semiconductor industry over the past few years for their application in ultra large scale integration (ULSI) inter-metal dielectric (IMD) [1]. As the circuit feature size shrinks below 0.18 μm, the increase of signal propagation delay due to interconnects becomes important. The smaller line dimensions increase the resistivity (R) of the metal lines and the narrower interline spacing increases the parasitic capacitance (C) between the lines. To reduce the RC delays, low-k dielectrics (k b 3.0) are being considered to replace the currently used SiO2 (k ∼ 4.2) for interlayer dielectric insulation. Considerable efforts have been recently spent on developing low-k materials using either plasma-enhanced chemical vapor deposition (PECVD) or spin-on deposition (SOD) techniques [2]. Among the CVD candidates, organosilicate glass (OSG) or SiOC materials appear to be well suited so that various organosilicon precursors, such as trimethylsilane (3MS) [3], tetramethylsilane (4MS) [4], bis-trimethylsilylmethane (BTMSM) [5], and tetramethylcyclotetrasiloxane (TMCTS) [6], have been examined to be adoptable to the back-end process ⁎ Corresponding author. E-mail address: [email protected] (H.J. Kim). 0040-6090/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2006.10.095
integration. The reported k values of these films are as low as 2.1, but the films are very weak because of their inherently porous structure; there is currently no dense low-k dielectric material with a dielectric constant below 2.5 [7]. In this respect, the understanding of low-k film's structure is one of the most important issues to solve the integration problems. This paper is focused on two different mechanisms of lowering dielectric constant of the films which were deposited in the same system but with different process conditions. 2. Experimental Low-k SiOC films were prepared on p-type (100) silicon substrates by radio frequency (13.56 MHz) Inductively Coupled Plasma (ICP) PECVD using a CH3-containing organosilicate precursor, bis-trimethylsilylmethane (BTMSM, C7H20Si2). The schematic drawing of the PECVD set-up we used throughout this study is shown in Fig. 1. Plasma was generated around the 3-turn coil connected to the r.f. generator. The distance between the coil and the substrate was 34 cm and that between the injector and the substrate was 5 cm. The chamber was pumped by a booster and a dry pump to a base pressure of 10− 3 Torr. The precursor was vaporized and carried with inert Ar gas from a thermostatic bubbler maintained at 30 °C to the reaction chamber.
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Fig. 1. The schematic drawing of ICP PECVD set-up.
Two different plasma excitation modes can be applied in the system. In low r.f. power range (60 W–300 W), the plasma is confined within the plasma chamber so that only the excitation gas is plasma excited. We call it as “remote plasma mode” hereafter. O2 as an oxidant gas and Ar as a dilute gas were fed into the chamber. We expect that low or moderate dissociation of the precursor will occur in this mode. In high plasma power region (700 W), the plasma extends into the reaction chamber and direct plasma dissociation of the monomer molecules will play an important role [8]. We call it as “direct plasma mode” hereafter. Ar gas was fed into the chamber through the quartz tube. Detailed experimental process condition is summarized in the Table 1. All films were post-deposition-annealed (PDA) at 450 °C in N2 ambient for 30 min, and p-Si/ low-k dielectric /Pt MetalInsulator-Semiconductor (MIS) capacitor structure was fabricated by depositing Pt-top electrodes using an electron beam evaporation method through a shadow mask. After Al was deposited as a back contact metal, annealing was performed at 400 °C for 30 min in a forming gas (5% H2 + 95% N2) ambient. A HP 4194A impedance meter and a HP 4140B picoammeter were used for capacitance–voltage (C–V) and current-density– voltage (J–V) measurements, respectively. The C–V measurement frequency was fixed at 1 MHz. We elucidated the structure of each film by XPS (Sigma Probe, ThermoVG) and FT-IR (JASCO FT/IR-660 plus) measurements. 3. Results and discussion
Fig. 2. FT-IR absorbance spectrum of BTMSM precursor.
assigned to Si–CH2–Si wagging vibration. This wagging peak is masked by Si–O–Si asymmetric stretching peak when it comes to oxygen-containing polymers, but the stretching peak is not as sharp as CH2 wagging peak [8,9]. In addition, no oxygen atom exists in BTMSM precursor. We can consequently conclude that the sharp peak at 1051 cm− 1 originated from Si–CH2–Si bonds. The inset of Fig. 2 is an enlargement ranging from 550 to 1000 cm− 1. It is identified that rocking vibration mode of CH3 and stretching vibration mode of Si–C also exists (686, 761, 782, 833, and 871 cm− 1), which confirms the structure of the precursor [6,8,9]. Fig. 3 shows the deposition rate behaviors of two different plasma modes. An Arrhenius plot for the deposition rate of both plasma mode films as a function of deposition temperature is shown in Fig. 3(a). The activation energy is calculated to be − 0.077 eV and − 0.070 eV for direct and remote plasma modes, respectively; relatively smaller energy for remote plasma mode film is ascribed to the fact that Si–O bond has larger tendency to be adsorbed to the growing film than Si–C bond at high temperature. The relatively high activation energy is attributed to larger desorption tendency of Si–C and C–C bonds at high temperature. The negative activation energy indicates that the reaction is absorption–desorption controlled [10,11]. Fig. 3(b) shows deposition rate change with plasma power for remote plasma mode films. The slope of deposition rate with plasma
FT-IR absorbance spectrum of the precursor in liquid phase at room temperature is shown in Fig. 2. Peaks at 833, 1249, 2896, and 2952 cm− 1 are CH3-related and a sharp peak at 1051 cm− 1 is Table 1 Deposition parameters for low-k films Power [W] Direct 700 mode Remote 60–300 Mode
Pressure [Torr]
Dilute Ar [sccm]
O2 [sccm]
Sub. temp. [°C]
0.2
50
b1
25–280
0.2
50
50
25
Fig. 3. The deposition rate behavior of (a) direct plasma mode films and (b) remote plasma mode films.
J. Heo et al. / Thin Solid Films 515 (2007) 5035–5039
Fig. 4. FT-IR spectra of (a) direct plasma mode films and (b) remote plasma mode films.
power is 6.9 up to 100 W region, but it decreased down to 2.0 from 100 W to 300 W. The change of deposition behavior seems to occur around 100 W. Fig. 4 shows FT-IR spectra of both plasma modes with different deposition conditions. The main building blocks for Organosilicate Glass (OSG) materials are “M ((CH3)3SiO1/2)”, “D ((CH3)2SiO2/2)”, “T (CH3SiO3/2)v, and “Q (SiO4/2)” groups corresponding to mono-, di-, tri-, and quad-oxygen-substituted silicon atoms [12]. The peak positions of the first three building blocks appear at 1250, 1260, and 1270 cm− 1, respectively. As shown in Fig. 4(a), Si(CH3)X symmetric stretching peak position for direct plasma mode films increases from 1255 cm− 1 to 1263 cm− 1 and the peak area decreases from 0.07 (a.u.) to 0.04 (a.u.) with deposition temperature. Strong absorptions at 790– 850 cm− 1 are assigned to Si(CH3)X (x = 1, 2, or 3) vibrations [6]. The absorptions at 804 cm− 1 and 844 cm− 1 are attributed to “D” and “M” vibrations, respectively. “M” to “D” peak intensity ratio also decreases from 90.7% to 58.4%. Abrupt decrease of FT-IR spectra at 1100 cm− 1 seems to be due to an artifact. FT-IR spectra for remote plasma mode films as a function of plasma power are shown in Fig. 4(b). Si(CH3)X absorption peaks at 804 and 844 cm− 1 increase with decreasing plasma power while relative intensities of Si–O–Si backbone structure peak decrease. Slight decrease in wavenumber of backbone peak (red shift) corresponds well with the result of Y. H. Kim et al. [13]; it means that carbon content incorporated in the film increases with decreasing plasma power. As an arrow depicted in Fig. 4(b), Si (CH3)X symmetric stretching peak at 1276 cm− 1 also follows the same increasing trend of Si(CH3)X peaks at 804 and 844 cm− 1. A peak at 1133 cm− 1 is attributed to larger angle Si–O–Si bonds in a cage structure with a bond angle of approximately 150°, which
Fig. 5. Changes of refractive indices and thickness loss before and after N2 PDA of (a) direct plasma mode films and (b) remote plasma mode films.
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Fig. 6. Changes of dielectric constant of (a) direct plasma mode films and (b) remote plasma mode films.
can lead to micropores and consequently a lower film density and dielectric constant [6,14]. From the FT-IR spectra comparison, it is concluded that two films for different plasma states have different main building blocks. Fig. 5 shows the changes of refractive indices and thickness loss before and after annealing at 450 °C in N2. Fig. 5(a) is for the direct plasma mode films as a function of plasma power. As deposition temperature increases, refractive index slightly increases up to 170 °C, but it abruptly soars above 170 °C. After PDA, refractive indices are unchanged in low temperature region, but it increases slightly above 150 °C. Thickness loss also shows drastic change; thickness loss of the film at room temperature is as high as 13.5% and it decreases down to 7–8% and eventually to about 3%. The generally reported thickness loss of low-k film is about 5–6% [5]. But A. Grill and Neumayer reported that thermally unstable CHx phase could be removed
Fig. 7. Si 2p XPS spectra of (a) direct plasma mode film (25 °C, k = 2.54), (b) remote plasma mode film (80 W, k = 2.9), and (c) thermal oxide.
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Fig. 8. The relative percentage change of Si building blocks.
and over 20% thickness loss could be obtained [6]. It seems that thermally unstable CHx radicals exist in the film at relatively low deposition temperature and well removed during N2 PDA so that thickness loss is higher. The changes of refractive indices and thickness loss for remote plasma mode films are shown in Fig. 5(b). The refractive indices of the films prepared by remote plasma mode are smaller than those of the films by direct mode. It has the lowest value of 1.424 at 80 W condition and gradually increases up to 1.438. All films show similar decrease in refractive index by about 0.3 after PDA and the thickness loss also follows the same V-shape change with plasma power. The highest thickness loss, however, is about 6.5% and it is smaller than those of the direct plasma mode films with deposition temperature below 100 °C. Meanwhile, the refractive indices of N2 annealed low-k films prepared by direct plasma mode are higher (n = 1.5–1.6) than those (n = 1.39–1.42) of films by remote plasma mode. It is explained from the replacement of skeleton SiO2 (RI = 1.46 at 633 nm) by SiC (RI ∼ 2.0 at 633 nm) [15]. The dielectric constant measured by C–V method after PDA is illustrated in Fig. 6. As shown in Fig. 6(a), dielectric constant of direct plasma mode films increases as increasing the deposition temperature. The film prepared at room temperature shows the lowest k value of 2.54. The dielectric constant change as a function of plasma power for remote plasma mode films is shown in Fig. 6(b). It also follows the V-shape trend as the refractive index and thickness loss do. The lowest k value obtained is 2.9 at plasma power 80 W and it increases drastically as plasma power increases from 100 to 300 W. It is explained as follows; as the plasma power increases to 100 W, the number of radicals generated by plasma increases so that deposition rate rapidly increases as well; the number eventually saturates at 100 W and it decreases above 100 W because cracking of BTMSM precursor and chemical etching by plasma start to dominate the whole process. To further confirm the structure of the films, the deposited films were characterized by XPS measurement very carefully. Fig. 7(a) is curve fitting of the Si 2p core level of each film. Fits were obtained by constraining the full width at halfmaximum (FWHM) equal to 1.45 [16]. To calibrate the
photoelectron binding energy, the C 1 s peak attributed to C–C bonds was set to 284.9 eV. The Si 2p spectrum was deconvoluted into five different moieties which are assigned as follows: “M” (101.2 eV), “D” (102.1 eV), “T” (102.8 eV), “Q” (103.4 eV) and OH related Si (104.5 eV). Binding energies that several authors [16–19] have reported are a little different from each other and fitting procedure is somewhat arbitrary, but it is meaningful enough to distinguish silicon in ‘organic’ or ‘inorganic’ states. XPS Si 2p curve fitting of direct plasma mode film (25 °C) is shown in Fig. 7(a) and that of remote plasma mode film (80 W) is shown in Fig. 7(b). Peak fitting of thermal oxide is inserted as a reference in Fig. 7(c). In both cases of Fig. 7(a) and (b), the binding energy is shifted toward lower binding energy compared to reference thermal oxide and the film of direct plasma mode is more shifted to lower energy. The chemical shift of Si 2p core level is highly dependent on the presence of surrounding oxygen atoms [20]. The presence of oxygen atom makes the higher binding energy shift compared to the presence of carbon. Fig. 8 summaries the result more clearly. “M” and “D” moieties are the main building blocks (85.7%) of direct plasma mode film, and “T” and “Q” moieties are the main building blocks (58.5 %) of remote plasma mode film. Regarding FT-IR spectra and XPS fitting results, the structures of two plasma mode films are understood as follows; films by direct plasma mode are mainly composed of organic C–C bonds and “M” and “D” moieties. Although electronic k value is higher than that of remote plasma mode films, ionic and dipolar proportions of direct plasma mode films are much smaller than those of remote plasma mode films; it is speculated from the smaller electronegativity difference of Si–C (0.7) than that of Si–O (1.7). X. Li et al. also insisted that the large number of low atomic mass C–C bonds leads to small atomic polarization, therefore low dielectric constants [21]. On the other hand, films by remote plasma mode are mainly composed of inorganic Si–O–Si bonds with some oxygen atoms substituted by CH3, resulting in micropores. The decrease of ionic dipolar contribution is the predominant factor on decrease in the dielectric constant [22]. 4. Conclusion We have demonstrated that the films with structure different from generally considered organosilicate glass can also have low dielectric constant as low as 2.54 by optimizing plasma mode and process condition. The host matrix of the film was mainly composed of organic C–C bonds and “M” and “D” moieties, and thus the low dipole and ionic polarizabilities were the important factors on lowering k value. The low-k film of remote plasma mode, however, is mainly composed of Si–O–Si backbone and some of oxygen atoms are substituted by CH3 groups, which is the general mechanism of lowering dielectric constant. Acknowledgement The first author wishes to thank the Jusung Engineering for a financial support.
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