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Dielectric constant and refractive index of poly (siloxane–imide) block copolymer
Materials and Design 32 (2011) 3173–3182
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Dielectric constant and refractive index of poly (siloxane–imide) block copolymer Muhammad Bisyrul Hafi Othman, Mohamad Riduwan Ramli, Looi Yien Tyng, Zulkifli Ahmad ⇑, Hazizan Md. Akil School of Material and Mineral Resources Engineering, Engineering Campus, University Science of Malaysia Seri Ampangan, 14300 Nibong Tebal Pulau Pinang, Malaysia
a r t i c l e
i n f o
Article history: Received 6 December 2010 Accepted 21 February 2011 Available online 24 February 2011 Keywords: Film and sheet Electrical Optical
a b s t r a c t Recent technological advances demanded polyimides of improved versatility in terms of electronic, optical and thermal properties. In this work, a series of poly(siloxane–imide) block copolymers were synthesized in order to investigate the effect on their optical and electronic properties. The polyimide unit was derived from 3,30 ,4,40 -Biphenyltetracarboxylic dianhydride (BPDA) and 4-(4-{1-[4-(4-aminophenoxy) phenyl]-1-methylethyl} phenoxy) aniline (BAPP) while the siloxane unit was derived from 3-[3-(3-aminopropyl)-1,1,3,3-tetramethyldisiloxanyl] propylamine (DMS) and Poly(dimethylsiloxane), bis(3-aminopropyl)terminated (PDMS). The structure of the polyimide was characterized by fourier transformer infra red (FT-IR), nuclear magnetic resonance (NMR) spectroscopy, elemental analysis, solution viscosity and gas permeation chromatography (GPC). Scanning electron microscope (SEM) analysis suggested a microphase separation between the two components. The refractive index and dielectric properties showed a decreasing trend with increased silicone unit in the polyimide backbone. However ultra violet visible (UV–Vis) and optical transparency was not significantly affected. Electronic and optical properties of this copolymer were discussed in relation to the polysiloxane content. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction In recent years, polyimide has received significant attention in microelectronic applications. This has been due to its many desirable characteristics, such as excellent mechanical properties, superior chemical and weather resistance, low thermal expansion coefficient, and its formability [1–4]. Technological advances gearing towards miniaturization of electronic devices have demanded that polyimides be further optimized, in terms of their electronic and optical properties. For example, in electronic packaging, a low dielectric constant with low dissipative heat energy is desirable [5–7]. Attempts to achieve this have included the introduction of a porous structure into the polymer matrix. The fabrication of voids allows a diffusion of air into the polymer matrices, and since air has a dielectric constant of 1.0, it induces a lowering of the dielectric constant. This porosity was fabricated into the polyimide structure by preparing block copolymers, consisting of thermally stable polyimide and thermally labile polymer blocks; the latter being thermolyse to affect the formation of nanovoids [8–11]. Nanoporous polyimide was also fabricated by initially incorporating tetraethoxysilane (TEOS) via sol–gel polycondensation into the polyamic acid intermediate, followed by acid etching of the inorganic phase, to leave behind pore structures. This has been demonstrated in several previous works, where the incorporation of ⇑ Corresponding author. Tel.: +60 45995099; fax: +60 45941011. E-mail address: zulkifl[email protected] (Z. Ahmad). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.02.048
fluorine into the polyimide effectively reduces the dielectric constant [12–15]. This is due to the low polarizability of the C–F bond, in addition to inducing an increase in free volume. It is vital for numerous applications to design polyimides with a controlled refractive index, such as waveguide and passive encapsulant materials for electronic devices. In comparison to a vacuum, the refractive index rises due to the different speeds of light travelling through a media. This speed depends on the relative permittivity and permeability of the media, and is therefore, closely related to the dielectric constant. The refractive index is related to the dielectric constant at the visible region through the following Maxwell equation:
n¼
pffiffiffiffiffiffiffiffiffi er :lr
ð1Þ
where er refers to relative permeability, and lr refers to relative permittivity. One of the strategies for controlling the refractive index is the introduction of fluorine into the polyimide [16,17]. Polysiloxane is a flexible inorganic elastomer with a very low glass transition temperature (Tg), but a high thermal stability. This is due to its stronger and more flexible SiAO bonds in comparison to normal organic linkages [18–20]. A common member of this polymer, polydimethylsiloxane (PDMS), displays dielectric constant values between 2.3 and 2.8, with a refractive index of 1.4000. The refractive index can be effectively controlled by incorporation of aromatic ring, as a pendant group. It displays excellent light transmission, which is near to ultra violet (UV) and a visible (Vis) region whose absorption characteristics occur near infra red
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region (NIR) [21]. Through the monitoring of changes in the refractive index or the absorption spectrum, several optical and electronic properties of this material can be tailored. The strategy of blending polysiloxane into an organic based polymer is limited as polysiloxane is relatively immiscible, resulting in phase separation during mixing with the latter. Several works have incorporated siloxane units into polyimide, in order to fabricate a new hybrid material with synergized optical and electronic properties. Most of these works have focused on thermal and mechanical properties. These included the utilisation of amino alkyl-terminated polydimethylsiloxane. One related work studied the dielectric properties, by synthesizing the copolymers based on pyromellitic dianhydride and 4,40 -oxydianiline, with amino alkyl-terminated polydimethylsiloxane [22–24]. A proportional amount of 4,40 -diphenylmethane diisocyanate was also incorporated to induce a carbon dioxide emission for the formation of pores. This work reported a consistent reduction in dielectric constant with the increase in polysiloxane units. However, the effect of 4,40 -diphenylmethane diisocyanate in relation to porosity is not quite conclusive. Furthermore, this system employed two variables, which made exclusive monitoring for each parameter ambiguous. This study reports the effect of copolymerization of siloxane units into polyimide and its effects on the optoelectronic properties. Variation of the silicone content was achieved by utilising silicone of different disiloxane repeat units. The chemical structures and molecular weights were established using nuclear magnetic resonance (NMR), fourier transformer infra red (FT-IR), and inherent viscosity, gas permeation chromatography (GPC), and scanning electron microscope (SEM) analysis. The dielectric constant, refractive index, and UV–Vis spectroscopy, were further explored and discussed in order to establish their optoelectronic properties.
Tetrahydrofuran, THF (99.5%), was purchased from Merck Darmstadt, Germany, and was distilled over calcium hydride and a N2 flow to remove traces of water before being used. The chemical structures of the monomers used, are shown in Fig. 1. 2.2. Synthesis of polyimide 2.2.1. Synthesis of polyimide (S1) BAPP (10 mmol, 4.10 g) was diluted with a THF solution (70.0 mL) in a pre-equilibrated 100 mL 3-necks round bottom flask at 25 °C, and initially purged with nitrogen gas for 5 min. The solution was stirred until homogeneous. Then BPDA (10 mmol, 2.94 g) was added and the mixture was further stirred vigorously for 90 min. Distilled water (100 ml) was added to give a white precipitate. After several cycles of washing, the product was filtered and vacuum dried at 90 °C for 24 h, producing a polyamic acid at 98.4% yield. Curing into polyimide was performed by casting onto a Teflon mould 20% of the polyamic acid in NMP, and heated stepwise under a vacuum from 25 °C to 100 °C for 1 h, 100–150 °C for 1 h, 150–200 °C for 1 h, and 200–310 °C for 1 h. The temperature was then allowed to decrease from 300 °C to 25 °C for 2 h, to give a tough yellowish film. 2.2.1.1. Elemental analysis. Calc. for C43H28N2O6, C 77.25, H 4.19, N 4.19%, Found C 75.09, H 4.63, N 4.91%, GPC analysis: Mn = 11,636, Mw = 28,353, PD = 2.4366, FTIR spectrum (film): mmax = 1724 cm1 (weak, symmetrical C@O imide stretch), 1780 cm1 (strong, asymmetrical C@O stretch), 1374 cm1 (strong, CANAC imide stretch), 1102 cm1, 747 cm1 (medium, CAOAC imide ring deformation). 1 H NMR (400 Hz, CDCl3): d (ppm) 8.23(2Ha, singlet), 8.08(2Hb, singlet), 7.42(2Hc, doublet, J = 12.0 Hz), 7.24(4Hf, doublet, J = 10.0 Hz), 7.17(4Hd, doublet, J = 11.6 Hz), 7.11(4Hg, doublet, J = 11.6), 7.01(4He, doublet, J = 11.2 Hz), 1.71(6Hh, singlet).
2. Experimental 2.1. Material 4-(4-(1-(4-(4-aminophenoxy) phenyl)-1-methylethyl) phenoxy) aniline, BAPP (98%), 3,30 ,4,40 -Biphenyltetracarboxylic dianhydride, BPDA (98%) and Poly-(dimethylsiloxane), Bis-(3-aminopropyl) terminated, (PMDS) (Mn 2500 gmol1), was obtained from Aldrich Chemical. BPDA was recrystallized from acetic anhydride and PMDS was dehydrated with a molecular sieve (Type 4A, Aldrich). 3-[3-(3-aminopropyl)-1,1,3,3-tetramethyldisiloxanyl] propylamine terminated, DMS (95%), was purchased from Fluka Sdn Bhd.
2.2.2. Synthesis of polyimide (S2) BPDA (20 mmol, 5.88 g) was diluted with a THF solution (100 mL) in a pre-equilibrated 3-neck round bottom flask at 25 °C, and was initially purged with nitrogen gas for 5 min. After stirring until homogeneous, DMS (5 mmol, 1.4 mL) was added drop-wise and the mixture was stirred vigorously at 25 °C for 24 h. The BAPP (10 mmol, 4.10 g) was added to the homogeneous BPDA–DMS solution and the mixture was left to be stirred vigorously at 25 °C for 24 h. The crude reaction product was poured into distilled water, to form a precipitate and then washed with excess distilled water, until a colourless solution was observed. The prod-
Fig. 1. The chemical structures of monomers used in this work.
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uct was filtered followed by vacuum drying at 80 °C for 24 h, to afford a white mass of silicone-incorporated polyamic acid, of yield 98.6%. Curing of the polyimide precursor was performed as per the S1 sample, to give a transparent yellowish film. 2.2.2.1. Elemental analysis. Calc. for C97H94N6O16Si4, C 68.05, H 5.53, N 4.91%, Found C 67.80, H 4.11, N 4.39%, GPC analysis: Mn = 13,444, Mw = 22,975, PD = 1.7089, inherent viscosity 0.19 dLg1 in THF (1 mg/ml at 25 °C). FTIR spectrum (film): mmax = 1724 cm1 (weak, symmetrical C@O imide stretch), 1780 cm1 (strong, asymmetrical C@O stretch), 1374 cm1 (strong, CANAC imide stretch), 1102 cm1, 747 cm1 (medium, CAOAC imide ring deformation). 1H NMR (400 Hz, CDCl3): d (ppm) 8.23 (2Ha, singlet), 8.08 (2Hb, singlet), 7.42 (2Hc, doublet, J = 8.8 Hz), 7.24 (4Hf, doublet, J = 8.8 Hz), 7.17(4Hd, doublet, J = 8.8 Hz), 7.11(4Hg, doublet, J = 8.8 Hz), 7.01(4He, doublet, J = 8.8 Hz), 2.86 (6Hi, triplet, J = 19.8 Hz), 1.71(6Hh, singlet), 1.68 (6Hj multiplet, J = 19.6 Hz), 1.27, (6Hk, triplet, J = 19.6 Hz), 0.91 (6Hl, singlet). 2.2.3. Synthesis of polyimide precursor (S3) BPDA (20 mmol, 5.88 g) was diluted with 100 mL of the THF solution in a 3-neck round bottom flask, which was equilibrated in a water bath below 25 °C, and initially purged with nitrogen gas for 5 min and then stirred with a magnetic stirring bar, until the BPDA–THF solution was homogeneous. Then PDMS (5 mmol, 12.8 ml) was added drop-wise and the mixture was stirred vigorously below 25 °C for 24 h, under a nitrogen gas flow. After the BPDA–PDMS solution became homogeneous, the BAPP (15 mmol, 6.15 g) was added and the mixture was stirred vigorously below 25 °C for 24 h, under a nitrogen gas flow. After the reaction was
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completed, the BAPP–BPDA–PDMS mixture was poured into distilled water to form a precipitate and then washed with excess distilled water, until a colourless solution was observed. The product was obtained by filtration followed by drying in a vacuum at 90 °C for 24 h. The product was poly-(amic siloxane acid) of yield 97.4%. Stepwise curing was performed as per the S1 sample. 2.2.3.1. Elemental analysis. Calc. for C479H480N20O93Si35, C 64.01.25, H 5.38, N 3,12%, Found C 64.70, H 3.63, N 4.04%, GPC analysis: Mn = 10,055, Mw = 18,146, PD = 1.8047, inherent viscosity 0.29 dLg1 in THF (1 mg/ml at 25 °C). FTIR spectrum (film): mmax = 1724 cm1 (weak, symmetrical C@O imide stretch), 1780 cm1 (strong, asymmetrical C@O stretch), 1374 cm1 (strong, CANAC imide stretch), 1102 cm1, 747 cm1 (medium, CAOAC imide ring deformation). 1H NMR (400 Hz, CDCl3):d (ppm) 8.23 (2Ha, singlet), 8.08 (2Hb, singlet), 7.42 (2Hc, doublet, J = 8.4 Hz), 7.24(4Hf, doublet, J = 8.4 Hz), 7.17(4Hd, doublet, J = 8.4 Hz), 7.11(4Hg, doublet, J = 8.4 Hz), 7.01(4He, doublet, J = 8.4 Hz), 2.86 (6Hi, triplet, J = 19.6 Hz), 1.71(6Hh, singlet), 1.68 (6Hj multiplet, J = 20), 1.27, (6Hk triplet, J = 20), 0.094 (6Hl, singlet), 0.102 (6Hm, singlet). 2.3. Sample preparations The crude precursors were precipitated with excesses of water. Then they were exhaustively washed using soxhlet extraction, initially with methanol/water in a volume ratio of 2:1 followed by acetone/water in a volume ratio of 5:1; each for a period of 24 h. The precursor was dried at 90 °C in a vacuum for 12 h, before it was poured into a Teflon casting and heated stepwise for curing, from 25 °C to 100 °C for 1 h, 100–150 °C for 1 h, 150–200 °C for
Scheme 1. Synthetic route of PI S1.
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1 h, and 200–310 °C for 1 h. The temperature was then allowed to decrease from 300 °C to 25 °C for 2 h, to give a tough yellowish film.
2.4. Characterization Elemental Analysis (EA) was performed by using a Perkin Elmer PE 2400 CHN and CHNS elemental analyser. Molecular Weight Average (Mn) and Molecular Weight Distribution (MWD) was determined by using GPC, in accordance with the ASTM D529605 [25]. An Inherent Viscosity test was performed; by using an Ubbelohde Calibrated Viscometer Tube in accordance with ASTM D 4603-03 [26] at 25 °C and approximately 30–50 mg of sample in 10 mL of THF was used. FTIR Spectrum was obtained using a Spectrum GX Perkin Elmer Model in accordance with ASTM E 1252 [27], at a wave length 4000–400 cm1. Samples were purified by Soxhlet, prior to measurements by scanning, four times. The H NMR was performed using a Bruker 400 Ultra Shield TM Model at a frequency of 400 MHz in an ambient temperature with deuterated chloroform as a solvent and tetramethlsilane as a standard reference. The surface morphology was observed using SEM equipment; the model used was a ZESS Supra 35VP at 5 K magnifications. The refractive index was obtained by using Refractometer model ATAGO, in accordance with ASTM D542 [28], at a wavelength of 589 nm at room temperature. The dielectric constant was obtained using LCR meter 4284 A, in accordance with ASTM D150-98 [29], at room temperature. UV–VIS absorbance spectroscopy was performed using UV–VIS spectrometer Lambda 35 Perkin Elmer Model at wavelength range 300–800 nm.
3. Results and discussions 3.1. Synthesis considerations The preparation of the three polyimides is shown in the reaction, as in Schemes 1 and 2. The synthesis of S1 was performed under normal addition and polycondensation into polyamic acid, and intermediately followed by chain cyclization to afford a quantitative yield of the polyimide product (Table 1). In S2 and S3, the silicon unit was first allowed to react with the excess anhydride, to ensure the end-capping of the latter onto the former. This allows the anhydride end-capper to copolymerize further with the dianiline during the subsequent stage. The copolymerization sequence between the three units can be monitored by following this procedure. In a three monomeric system, it has been generally shown that the sequence of adding each monomer significantly affects both the product’s molecular weight and morphology. This is due to different terminal group and monomer reactivities during the copolymerization process. Ren et al., found that the sequence of adding 4,4-diamino diphenylmethane with 1,4-bis (3-amino propyldimethylsilyl) benzene and pyromellitic dianhydride (PMDA) markedly affected the dielectric constant. It would be expected that by changing the sequence, such as by reacting the three monomers in a single stage in this work, that its properties might correspondingly change. The inherent viscosity showed that the products were of a moderate molecular weight, as substantiated by the Mn and Mw values. The synthesized polyimide products displayed polydispersity between 1.7 and 2.5. Based on the average molar mass Mn, obtained from GPC and the elemental analysis, a repeat unit for S1–S3 can be proposed, as in Fig. 2.
Scheme 2. PDMS Capping reactions with the anhydride monomer and the complete reaction to form poly(imide siloxane) block copolymer S2 (n = 1) and S3 (n = 35).
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a
PI series
Yield percentage (mol/mol%)
Inherent viscosities (dLg1)a
Molecular weight distribution (gmol1)
Polydispersity PD
Mn
Mw
(Mw/Mn)
Sl S2 S3
96.78 86.07 82.46
0.32 0.19 0.29
12,736 13,444 10,055
31,799 22,975 18,146
2.4968 1.7089 1.8047
Inherent viscosity measure in NMP at 25 + 0.1 °C with concentration of 0.3 g/dL.
3.2. Structure conformation Fig. 3 shows the FTIR spectra of the three polyimides in the range of 3800 and 1200 cm1. No apparent peaks were present in
the region 3600–3200 cm1 corresponding to primary and secondary amides, which confirmed an almost complete imidization reaction, to yield polyimide from their respective polyamic acid precursors. A relatively small peak occurred at 2950–2820 cm1
Fig. 2. The propose of repeated unit of PI series based on GPC and elemental analysis.
Fig. 3. FTIR spectra: (a) representative, and (b) the pure polyimide (S1) and the silicon incorporated polyimide (S2 and S3).
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for the S1 sample, representing the methyl group of the isopropyl bridge, of the dianiline moiety. The intensity of these peaks increase correspondingly for the S2 and S3 samples, which is consistent with the increasing methyl content, derived from the incorporated silicone units. The successful incorporation of silicon units could be further substantiated by the presence of strong and broad peaks at the region 1090–1020 cm1, representing the SiAOAC and SiAOASi bonds (open chain). This characteristic peak was not detected in sample S1. Fig. 4 shows the NMR spectra of the three series of polyimides. The incorporation of the silicone unit in spectra S2 and S3 were proven by the presence of peaks at the region 1.5–0.5 ppm corresponding to the methyl hydrogen attached to the silicon atom.
These peaks were not detected in the case of polyimide S1. The incorporation of the siloxane unit was further substantiated with the presence of peaks at 2.8 ppm (triplet), 1.7 ppm (multiplet), and 1.2 ppm (triplet), which corresponds to the propyl proton of bis-(3-aminopropyl) terminated siloxane moiety. These peaks were again not detected in the S1 spectrum. Peaks corresponding to all aromatic hydrogen occurred within the region 6.8–8.3 ppm and that of methyl protons at 1.8 ppm at the polyimide moiety. Apparently, no substantial chemical shift occurred to these protons during the incorporation of siloxane units, implying that the spacer propyl unit which connect the polyimide and siloxane unit, does not pose any local electronic perturbation to either of these groups.
Fig. 4. 1H NMR (400 MHz) spectra of: (a) S1, (b) S2, and (c) S3 in CDCl3.
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3.3. Morphology of copolymers Morphology of the copolymerized polyimides was investigated using SEM. Fig. 5a of the pure polyimide sample, was essentially featureless. Fig. 5b and c revealed the existence of a very fine phase separation in the form of microspheres; presumably a siloxane segment in the polyimide matrix. These microspheres became far more significant for the S3 sample than the S2 sample. It has been established that phase separation results in a 20% silicone incorporated polyimide system [24]. In this study, the concentration of silicone was deliberately set at a 10% weight ratio to minimize this phase separation effect. The differential scanning calorimetry (DSC) scan for pure polyimide S1 revealed a thermal transition at 244 °C corresponding to its glass transition. For polysiloxane incorporated samples (S2 and S3), the thermal transitions in the 200–206 °C region corre-
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sponds to the glass transition of these copolymers. Due to the chain flexibility of the copolymerized polysiloxane unit, both S2 and S3 samples displayed a lower Tg when compared to the pure polyimide. These reduction revealed that the copolymers are less rigid compared to the pure polyimide. The glass transition for S3 is both lower (203 °C) and broader than that of S2 (Tg = 209 °C). This is mostly due to the presence of a longer polysiloxane segment in S3 which rendered its higher chain flexibility. The Tg of the polysiloxane segment, which would otherwise been detected in the DSC scan in the sub-ambient region, was not observed. This could be due to the micro phase separation morphology, whose scale of thermal transition was too small for detection. Another plausible reason would be that the siloxane unit was tightly confined between the rigid polyimide chains which limit any segmental movement to the former. As expected, no melting transition was observed for all samples within this range of temperatures (Fig. 6). 3.4. Refractive index and dielectric constant The refractive index and dielectric constant values of the three polyimides are shown in Table 2. It can be seen that S1 displayed the highest refractive index, followed by S2, then S3. These results can be attributed to increased silicone unit content. Based on the Lorenz–Lorenz equation, the refractive index n is dependent on Ru, the molar refractivity and Vu, the molar volume, as in the following equation:
Ru ¼ V u
Fig. 5. The surface morphology of polyimide series: (a) parent PI, (b) PI incorporate DMS and (c) PI incorporate PDMS.
ðn22 1Þ ðn22 þ 2Þ
ð2Þ
Ru and Vu are additive functions of the respective group’s composition in the repeating unit. Brunchi et al., verified that the increases of silicone repeating unit, would adversely affect the ratio of Ru/Vu, and subsequently the refractive index [30]. This is due to the bigger size of Si atom as compared to that of carbon which significantly affects the molar volume of the system. Hence, it was expected that with a higher content in the silicone unit, the refractive index would correspondingly decrease. A close relationship of the refractive index to the dielectric constant is shown by the relationship in Eq. (1). Both quantities involve an interaction of radiation with the local electronic environment of chemical bonds. Table 2 shows the dielectric constant for samples S1, S2, and S3 with a decreasing trend as the polysiloxane content in the polymer increased. This result was consistent with that of the refractive index. Silicon is comparatively larger than a carbon atom and the SiAO bond is more flexible than the CAC bond. This is substantiated in several DMA and mechanical property studies [24,31]. Thus, it can be affirmed that the bulky silicon units would be less mobile. Its presence affects the bulk movement of the whole polyimide network which reduces the efficiency of the dipole in reacting to polarity change during treatment with an alternating frequency. Furthermore, the molar polarization significantly decreases as the result of an increase in free volume. Several recent studies have demonstrated a similar trend of a decreasing dielectric constant, with an increasing siloxane content into polyimide structures [32,33]. The values obtained in this study of dielectric constant at 1 kHz, are within the range obtained from these other studies. These values were larger than those derived from the Maxwell equation (shown in Table 2), because the Maxwell equation is only applied at the optical frequency of approximately 1015 Hz, which only involves electronic polarization. Fig. 7 shows the variation of dielectric constant with an applied electric field frequency. The results show that with an increasing frequency, e decreases very gradually. This variation is attributed to the frequency dependence of the polarization mechanisms.
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Fig. 6. DSC thermogram (2nd scan) for three systems of PI.
Table 2 Refractive Index, dielectric constant and cutoff wavelength (/nm) from UV–Vis spectra of polyimides and copolymers.
a b c
Polymer
Refractive indexa
Dielectric constant estimationb
Dielectric constantc
Absmax (k/nm)
Sl S2 S3
1.4983 1.4936 1.4922
2.25 2.23 2.22
2.90 2.57 2.43
407 408 413
Measured at sodium wavelength 589 nm at temperature 25 °C. Estimated from Maxwell equations = n2. Measured at 1 kHz in room temperature.
The overall contribution of dielectric constant can be integrated into the following equation:
eTotal ¼ eelectronic þ eatomic þ edipole
ð3Þ
At a high frequency regime, in the range 1013–1018 Hz, the polarization mechanism involves electronic polarization, followed by the frequency range 1010–1013 Hz, which includes atomic polarizability, and finally below 109 Hz, which includes dipole polarization. Within the 103–106 Hz frequency range measured in this
work, the summation of all three modes of polarization were involved. However, as the frequency gradually increased, the atomic polarizability contribution began to diminish. Effectively, the dielectric constant gradually decreased with the increasing frequency. It can be observed that the dielectric constant slightly increases from 0 to 10 kHz for S1 and S2. This peculiarity can be attributed to a phenomenon known as electrode polarization. It originates from a high-impedance layer on the electrode surface, mainly due to an imperfect contact between electrode and the specimen during measurement. Electrolytic products further aggravate this effect. This suggests that at low frequency, there is sufficient time for any slight conduction through the specimen, to transfer the entire applied field across the very thin electrode. This anomaly is characterized by an apparent increase in relative permittivity, as observed at 0–10 kHz [34,35]. The dielectric loss of the polyimides is represented in Fig. 8. and refers to the imaginary part of a relative permittivity, as in the following equation:
er ¼ e0r ðxÞ þ ie00r ðxÞ
ð4Þ
where the first term on the right side of the equation refers to a real component and the second term refers to an imaginary component. The dielectric loss of the S1 sample increased almost proportion-
Fig. 7. Dielectric constant of PI series corresponding to frequency 1 kHz, 10 kHz, 100 kHz and 1000 kHz at room temperature.
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Fig. 8. The dielectric loss of the pure polyimides S1 and polysiloxane–polyimide copolymers S2 and S3.
ately with the frequency; at least within the measured frequency range. This was due to the inability of the polarizable units in the polyimide chains to align in phase with the increasing electric field frequency. There will be a phase lag with the oscillating electric field, which then contributes to the dielectric loss. Physically, the loss is constituted in the form of dissipative heat energy [36]. In the case of the siloxane-incorporated polyimide samples, it showed an initial decrease in dielectric loss, followed by a sinusoidal variation at a higher frequency range. However, the value obtained in the range 2.59 104–7.5 103, are reasonably low compared to some of the other organic polymers and inorganic materials. The dielectric loss for sample S3 is higher compared to that of sample S2. This was apparently due to the higher siloxane content in the S3 sample. Since the siloxane chains in S3 are longer, this induced inefficient bulk chain flexibility to the copolymers. This effect increased a phase lag with that of the applied oscillating electric field. 3.5. UV–VIS absorbance spectroscopy The UV–Vis absorbance spectroscopy of the three series of polyimides is shown in Fig. 9. All synthesized siloxane-incorporated polyimides displayed a cut-off absorption wavelength below
450 nm. Two characteristic broad peaks were detected at 360 nm and 400 nm wavelengths, which corresponds to the p–p transition due to the charge transfer complex of the aromatic rings of the polyimide chain backbone [37,38]. It is obvious that the incorporation of polysiloxane into the polyimide chain does not affect any significant change in the absorption maxima, or the transparency of the copolymer. It was reported by Kai et al., that except at NIR region, poly-(methylsiloxane) does not display any characteristic absorption in the UV–visible region. There are noticeable decreased in absorbances in S1 as compared to S2 and S3 at 400 nm region. Several factors contribute to the change in absorbance, albeit extinction coefficient during the UV–visible scan, such as complex formation, particle size distribution [39], and the extent of imidization during polyimide curing [37]. It has been shown that in addition to the red shift effect, the extinction coefficient gradually decreases with respect to the level of imidization curing in the polyimide. A decrease in absorbency for S1 at the 400 nm region could possibly reflect a complete imidization curing of the sample, while samples S2 and S3 showed a comparatively higher absorbency due to inefficient imidization curing, as the result of the presence of siloxane units which posed a steric hindrance during the curing process. Visually, all synthesized PI
Fig. 9. The UV–VIS absorbance spectroscopy of the three polyimide.
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showed a yellowish colouration, which is consistent with the tailing absorption into the visible 450 nm wavelength for all samples. 4. Conclusion Copolymerization of the silicone unit into the polyimide backbone was successfully performed. FTIR and H NMR spectra verified that polyimide and poly-(imide siloxane) were completely imidized from poly-(amic acid) and poly-(amic siloxane acid), respectively. All polyimide series were synthesized with moderate molecular weights and narrow molecular weight distributions. SEM showed a micro phase separation between the two phases. Incorporation of silicon unit into the polyimide chain significantly reduced the refractive index and dielectric constant, with a reasonably low dielectric loss. However, siloxane unit does not affect UV– visible absorption or transparency in the copolymer. Acknowledgment The authors wish to thank the USM for sponsoring this Project under USM-RU-PGRS 8032006. References [1] Ghosh M, Mittal K. Polyimides: fundamentals and applications. CRC Press; 1996. [2] Cha H, Hedrick J, DiPietro R, Blume T, Beyers R, Yoon D. Structures and dielectric properties of thin polyimide films with nano foam morphology. Appl Phys Lett 1996;68:1930–2. [3] Feger C, Khojasteh M, McGrath J. Polyimides: materials, chemistry and characterization. SPE; 1988. [4] Mittal K. Polyimides: synthesis, characterization, and applications, vols. 1 and 2. New York: Plenum Press; 1984. [5] Courtois H, Buisson O, Chaussy J, Pannetier B. Miniature low temperature high frequency filters for single electronics. Rev Sci Instrum 2009;66:3465–8. [6] Binnig G, Rohrer H, Gerber C, Weibel E. Tunneling through a controllable vacuum gap. Appl Phys Lett 2009;40:178–80. [7] Zhang Y, Ke S, Huang H, Zhao L, Yu L, Chan H. Dielectric relaxation in polyimide nanofoamed films with low dielectric constant. Appl Phys Lett 2008;92(052910):1–3. [8] Jiang L, Liu J, Wu D, Li H, Jin R. A methodology for the preparation of nanoporous polyimide films with low dielectric constants. Thin Solid Films 2006;510:241–6. [9] Mehdipour Ataei S, Saidi S. Structure–property relationships of low dielectric constant, nanoporous, thermally stable polyimides via grafting of poly (propylene glycol) oligomers. Polym Adv Technol 2008;19:889–94. [10] Zhao XY, Liu HJ. Review of polymer materials with low dielectric constant. Polym Int 2010;59:597–606. [11] Yu Y-Y, Chien W-C, Chen S-Y. Preparation and optical properties of organic/ inorganic nanocomposite materials by UV curing process. Mater Des 2010;31:2061–70. [12] Cornelius C, Marand E. Hybrid inorganic–organic materials based on a 6FDA– 6FpDA–DABA polyimide and silica: physical characterization studies. Polymer 2002;43:2385–400. [13] Dhara MG, Banerjee S. Fluorinated high-performance polymers: poly(arylene ether)s and aromatic polyimides containing trifluoromethyl groups. Progr Polym Sci 2010;35:1022–77.
[14] Ding M. Isomeric polyimides. Progr Polym Sci 2007;32:623–68. [15] Hougham G, Tesoro G, Viehbeck A. Influence of free volume change on the relative permittivity and refractive index in fluoropolyimides. Macromolecules 1996;29:3453–6. [16] Liu Y, Xing Y, Zhang Y, Guan S, Zhang H, Wang Y, et al. Novel soluble fluorinated poly (ether imide) s with different pendant groups: synthesis, thermal, dielectric, and optical properties. J Polym Sci Part A: Polym Chem 2010;48:3281–9. [17] Cho S, Allcock H. Novel highly fluorinated perfluorocyclobutane-based phosphazene polymers for photonic applications. Chem Mater 2007;19:6338–44. [18] Zou L, Anthamatten M. Synthesis and characterization of polyimide– polysiloxane segmented copolymers for fuel cell applications. J Polym Sci Part A: Polym Chem 2007;45:3747–58. [19] Kricheldorf H, Burger C. Silicon in polymer synthesis. Springer: Berlin; 1996. [20] Mark J, Allcock H, West R. Inorganic polymers. USA: Oxford University Press; 2005. [21] Su K, DeGroot JJV, Norris AW, Lo PY. Siloxane materials for optical applications. In: Lu W, Young J, editors. ICO20. Materials and Nanostructures. 1st ed. SPIE; 2006. p. 60291C-8. [22] Ghosh A, Banerjee S, Häußler L, Voit B. New fluorinated poly (imide siloxane) random and block copolymers with variation of siloxane loading. J Macromol Sci, Part A. 2010;47:671–80. [23] Allen G. Speciality polymers: prospect-retrospect. Mater Des 1986;7:179–81. [24] Rogers ME, Glass TE, Mecham SJ, Rodrigues D, Wilkes GL, McGrath JE. Perfectly alternating segmented polyimide–polydimethyl siloxane copolymers via transimidization. J Polym Sci Part A: Polym Chem 1994;32:2663–75. [25] Annual Book of ASTM Standards. Standard test method for molecular weight averages and molecular weight distribution of polystyrene by high performance size-exclusion chromatography, vol. 08; 2005. p. 03]. [26] Anual Book of ASTM Standards. Standard test method for determining inherent viscosity by glass capillary viscometer; 2003 [08.02] [27] Anual Book of ASTM Standards. Standard practice for general techniques for obtaining infrared spectra for qualitative analysis; 2007 [03.06]. [28] Anual Book of ASTM Standards. Test method for index of refraction of transparent organic plastics; 2006 [08.01]. [29] Anual Book of ASTM Standards. Standard test methods for ac loss characteristics and permittivity (dielectric constant) of solid electrical insulation; 2004 [10.01]. [30] Brunchi CE, Filimon A, Cazacu M, Ioan S. Properties of some poly(siloxane)s for optical applications. High Perform Polym 2009;21:31–47. [31] Arnold C, Summers J, Chen Y, Bott R, Chen D, McGrath J. Structure-property behaviour of soluble polyimide–polydimethylsiloxane segmented copolymers. Polymer 1989;30:986–95. [32] Ghosh A, Banerjee S. Thermal, mechanical, and dielectric properties of novel fluorinated copoly (imide siloxane) s. J Appl Polym Sci 2008;109:2329–40. [33] Tao L, Yang H, Liu J, Fan L, Yang S. Synthesis and characterization of highly optical transparent and low dielectric constant fluorinated polyimides. Polymer 2009;50:6009–18. [34] Blythe A, Bloor D. Electrical properties of polymers. Cambridge University Press; 2005. [35] Alegaonkar P, Bhoraskar V, Balaya P, Goyal P. Dielectric properties of 1 MeV electron-irradiated polyimide. Appl Phys Lett 2009;80:640–2. [36] Pethrick RA. The development and application of polymeric materials for the electronics industry. Mater Des 1986;7:173–8. [37] Pyun E, Mathisen R, Sung C. Kinetics and mechanisms of thermal imidization of a polyamic acid studied by ultraviolet-visible spectroscopy. Macromolecules 1989;22:1174–83. [38] Herminghaus S, Boese D, Yoon D, Smith B. Large anisotropy in optical properties of thin polyimide films of poly (p phenylene biphenyltetracarboximide). Appl Phys Lett 2009;59:1043–5. [39] Roy A, Sharma S, Gupta R. Frequency and size distribution dependence of visible and infrared extinction for astronomical silicate and graphite grains. J Quant Spectrosc Radiative Trans 2010;111:795–801.
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Dielectric constant and refractive index of poly (siloxane–imide) block copolymer Muhammad Bisyrul Hafi Othman, Mohamad Riduwan Ramli, Looi Yien Tyng, Zulkifli Ahmad ⇑, Hazizan Md. Akil School of Material and Mineral Resources Engineering, Engineering Campus, University Science of Malaysia Seri Ampangan, 14300 Nibong Tebal Pulau Pinang, Malaysia
a r t i c l e
i n f o
Article history: Received 6 December 2010 Accepted 21 February 2011 Available online 24 February 2011 Keywords: Film and sheet Electrical Optical
a b s t r a c t Recent technological advances demanded polyimides of improved versatility in terms of electronic, optical and thermal properties. In this work, a series of poly(siloxane–imide) block copolymers were synthesized in order to investigate the effect on their optical and electronic properties. The polyimide unit was derived from 3,30 ,4,40 -Biphenyltetracarboxylic dianhydride (BPDA) and 4-(4-{1-[4-(4-aminophenoxy) phenyl]-1-methylethyl} phenoxy) aniline (BAPP) while the siloxane unit was derived from 3-[3-(3-aminopropyl)-1,1,3,3-tetramethyldisiloxanyl] propylamine (DMS) and Poly(dimethylsiloxane), bis(3-aminopropyl)terminated (PDMS). The structure of the polyimide was characterized by fourier transformer infra red (FT-IR), nuclear magnetic resonance (NMR) spectroscopy, elemental analysis, solution viscosity and gas permeation chromatography (GPC). Scanning electron microscope (SEM) analysis suggested a microphase separation between the two components. The refractive index and dielectric properties showed a decreasing trend with increased silicone unit in the polyimide backbone. However ultra violet visible (UV–Vis) and optical transparency was not significantly affected. Electronic and optical properties of this copolymer were discussed in relation to the polysiloxane content. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction In recent years, polyimide has received significant attention in microelectronic applications. This has been due to its many desirable characteristics, such as excellent mechanical properties, superior chemical and weather resistance, low thermal expansion coefficient, and its formability [1–4]. Technological advances gearing towards miniaturization of electronic devices have demanded that polyimides be further optimized, in terms of their electronic and optical properties. For example, in electronic packaging, a low dielectric constant with low dissipative heat energy is desirable [5–7]. Attempts to achieve this have included the introduction of a porous structure into the polymer matrix. The fabrication of voids allows a diffusion of air into the polymer matrices, and since air has a dielectric constant of 1.0, it induces a lowering of the dielectric constant. This porosity was fabricated into the polyimide structure by preparing block copolymers, consisting of thermally stable polyimide and thermally labile polymer blocks; the latter being thermolyse to affect the formation of nanovoids [8–11]. Nanoporous polyimide was also fabricated by initially incorporating tetraethoxysilane (TEOS) via sol–gel polycondensation into the polyamic acid intermediate, followed by acid etching of the inorganic phase, to leave behind pore structures. This has been demonstrated in several previous works, where the incorporation of ⇑ Corresponding author. Tel.: +60 45995099; fax: +60 45941011. E-mail address: zulkifl[email protected] (Z. Ahmad). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.02.048
fluorine into the polyimide effectively reduces the dielectric constant [12–15]. This is due to the low polarizability of the C–F bond, in addition to inducing an increase in free volume. It is vital for numerous applications to design polyimides with a controlled refractive index, such as waveguide and passive encapsulant materials for electronic devices. In comparison to a vacuum, the refractive index rises due to the different speeds of light travelling through a media. This speed depends on the relative permittivity and permeability of the media, and is therefore, closely related to the dielectric constant. The refractive index is related to the dielectric constant at the visible region through the following Maxwell equation:
n¼
pffiffiffiffiffiffiffiffiffi er :lr
ð1Þ
where er refers to relative permeability, and lr refers to relative permittivity. One of the strategies for controlling the refractive index is the introduction of fluorine into the polyimide [16,17]. Polysiloxane is a flexible inorganic elastomer with a very low glass transition temperature (Tg), but a high thermal stability. This is due to its stronger and more flexible SiAO bonds in comparison to normal organic linkages [18–20]. A common member of this polymer, polydimethylsiloxane (PDMS), displays dielectric constant values between 2.3 and 2.8, with a refractive index of 1.4000. The refractive index can be effectively controlled by incorporation of aromatic ring, as a pendant group. It displays excellent light transmission, which is near to ultra violet (UV) and a visible (Vis) region whose absorption characteristics occur near infra red
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region (NIR) [21]. Through the monitoring of changes in the refractive index or the absorption spectrum, several optical and electronic properties of this material can be tailored. The strategy of blending polysiloxane into an organic based polymer is limited as polysiloxane is relatively immiscible, resulting in phase separation during mixing with the latter. Several works have incorporated siloxane units into polyimide, in order to fabricate a new hybrid material with synergized optical and electronic properties. Most of these works have focused on thermal and mechanical properties. These included the utilisation of amino alkyl-terminated polydimethylsiloxane. One related work studied the dielectric properties, by synthesizing the copolymers based on pyromellitic dianhydride and 4,40 -oxydianiline, with amino alkyl-terminated polydimethylsiloxane [22–24]. A proportional amount of 4,40 -diphenylmethane diisocyanate was also incorporated to induce a carbon dioxide emission for the formation of pores. This work reported a consistent reduction in dielectric constant with the increase in polysiloxane units. However, the effect of 4,40 -diphenylmethane diisocyanate in relation to porosity is not quite conclusive. Furthermore, this system employed two variables, which made exclusive monitoring for each parameter ambiguous. This study reports the effect of copolymerization of siloxane units into polyimide and its effects on the optoelectronic properties. Variation of the silicone content was achieved by utilising silicone of different disiloxane repeat units. The chemical structures and molecular weights were established using nuclear magnetic resonance (NMR), fourier transformer infra red (FT-IR), and inherent viscosity, gas permeation chromatography (GPC), and scanning electron microscope (SEM) analysis. The dielectric constant, refractive index, and UV–Vis spectroscopy, were further explored and discussed in order to establish their optoelectronic properties.
Tetrahydrofuran, THF (99.5%), was purchased from Merck Darmstadt, Germany, and was distilled over calcium hydride and a N2 flow to remove traces of water before being used. The chemical structures of the monomers used, are shown in Fig. 1. 2.2. Synthesis of polyimide 2.2.1. Synthesis of polyimide (S1) BAPP (10 mmol, 4.10 g) was diluted with a THF solution (70.0 mL) in a pre-equilibrated 100 mL 3-necks round bottom flask at 25 °C, and initially purged with nitrogen gas for 5 min. The solution was stirred until homogeneous. Then BPDA (10 mmol, 2.94 g) was added and the mixture was further stirred vigorously for 90 min. Distilled water (100 ml) was added to give a white precipitate. After several cycles of washing, the product was filtered and vacuum dried at 90 °C for 24 h, producing a polyamic acid at 98.4% yield. Curing into polyimide was performed by casting onto a Teflon mould 20% of the polyamic acid in NMP, and heated stepwise under a vacuum from 25 °C to 100 °C for 1 h, 100–150 °C for 1 h, 150–200 °C for 1 h, and 200–310 °C for 1 h. The temperature was then allowed to decrease from 300 °C to 25 °C for 2 h, to give a tough yellowish film. 2.2.1.1. Elemental analysis. Calc. for C43H28N2O6, C 77.25, H 4.19, N 4.19%, Found C 75.09, H 4.63, N 4.91%, GPC analysis: Mn = 11,636, Mw = 28,353, PD = 2.4366, FTIR spectrum (film): mmax = 1724 cm1 (weak, symmetrical C@O imide stretch), 1780 cm1 (strong, asymmetrical C@O stretch), 1374 cm1 (strong, CANAC imide stretch), 1102 cm1, 747 cm1 (medium, CAOAC imide ring deformation). 1 H NMR (400 Hz, CDCl3): d (ppm) 8.23(2Ha, singlet), 8.08(2Hb, singlet), 7.42(2Hc, doublet, J = 12.0 Hz), 7.24(4Hf, doublet, J = 10.0 Hz), 7.17(4Hd, doublet, J = 11.6 Hz), 7.11(4Hg, doublet, J = 11.6), 7.01(4He, doublet, J = 11.2 Hz), 1.71(6Hh, singlet).
2. Experimental 2.1. Material 4-(4-(1-(4-(4-aminophenoxy) phenyl)-1-methylethyl) phenoxy) aniline, BAPP (98%), 3,30 ,4,40 -Biphenyltetracarboxylic dianhydride, BPDA (98%) and Poly-(dimethylsiloxane), Bis-(3-aminopropyl) terminated, (PMDS) (Mn 2500 gmol1), was obtained from Aldrich Chemical. BPDA was recrystallized from acetic anhydride and PMDS was dehydrated with a molecular sieve (Type 4A, Aldrich). 3-[3-(3-aminopropyl)-1,1,3,3-tetramethyldisiloxanyl] propylamine terminated, DMS (95%), was purchased from Fluka Sdn Bhd.
2.2.2. Synthesis of polyimide (S2) BPDA (20 mmol, 5.88 g) was diluted with a THF solution (100 mL) in a pre-equilibrated 3-neck round bottom flask at 25 °C, and was initially purged with nitrogen gas for 5 min. After stirring until homogeneous, DMS (5 mmol, 1.4 mL) was added drop-wise and the mixture was stirred vigorously at 25 °C for 24 h. The BAPP (10 mmol, 4.10 g) was added to the homogeneous BPDA–DMS solution and the mixture was left to be stirred vigorously at 25 °C for 24 h. The crude reaction product was poured into distilled water, to form a precipitate and then washed with excess distilled water, until a colourless solution was observed. The prod-
Fig. 1. The chemical structures of monomers used in this work.
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uct was filtered followed by vacuum drying at 80 °C for 24 h, to afford a white mass of silicone-incorporated polyamic acid, of yield 98.6%. Curing of the polyimide precursor was performed as per the S1 sample, to give a transparent yellowish film. 2.2.2.1. Elemental analysis. Calc. for C97H94N6O16Si4, C 68.05, H 5.53, N 4.91%, Found C 67.80, H 4.11, N 4.39%, GPC analysis: Mn = 13,444, Mw = 22,975, PD = 1.7089, inherent viscosity 0.19 dLg1 in THF (1 mg/ml at 25 °C). FTIR spectrum (film): mmax = 1724 cm1 (weak, symmetrical C@O imide stretch), 1780 cm1 (strong, asymmetrical C@O stretch), 1374 cm1 (strong, CANAC imide stretch), 1102 cm1, 747 cm1 (medium, CAOAC imide ring deformation). 1H NMR (400 Hz, CDCl3): d (ppm) 8.23 (2Ha, singlet), 8.08 (2Hb, singlet), 7.42 (2Hc, doublet, J = 8.8 Hz), 7.24 (4Hf, doublet, J = 8.8 Hz), 7.17(4Hd, doublet, J = 8.8 Hz), 7.11(4Hg, doublet, J = 8.8 Hz), 7.01(4He, doublet, J = 8.8 Hz), 2.86 (6Hi, triplet, J = 19.8 Hz), 1.71(6Hh, singlet), 1.68 (6Hj multiplet, J = 19.6 Hz), 1.27, (6Hk, triplet, J = 19.6 Hz), 0.91 (6Hl, singlet). 2.2.3. Synthesis of polyimide precursor (S3) BPDA (20 mmol, 5.88 g) was diluted with 100 mL of the THF solution in a 3-neck round bottom flask, which was equilibrated in a water bath below 25 °C, and initially purged with nitrogen gas for 5 min and then stirred with a magnetic stirring bar, until the BPDA–THF solution was homogeneous. Then PDMS (5 mmol, 12.8 ml) was added drop-wise and the mixture was stirred vigorously below 25 °C for 24 h, under a nitrogen gas flow. After the BPDA–PDMS solution became homogeneous, the BAPP (15 mmol, 6.15 g) was added and the mixture was stirred vigorously below 25 °C for 24 h, under a nitrogen gas flow. After the reaction was
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completed, the BAPP–BPDA–PDMS mixture was poured into distilled water to form a precipitate and then washed with excess distilled water, until a colourless solution was observed. The product was obtained by filtration followed by drying in a vacuum at 90 °C for 24 h. The product was poly-(amic siloxane acid) of yield 97.4%. Stepwise curing was performed as per the S1 sample. 2.2.3.1. Elemental analysis. Calc. for C479H480N20O93Si35, C 64.01.25, H 5.38, N 3,12%, Found C 64.70, H 3.63, N 4.04%, GPC analysis: Mn = 10,055, Mw = 18,146, PD = 1.8047, inherent viscosity 0.29 dLg1 in THF (1 mg/ml at 25 °C). FTIR spectrum (film): mmax = 1724 cm1 (weak, symmetrical C@O imide stretch), 1780 cm1 (strong, asymmetrical C@O stretch), 1374 cm1 (strong, CANAC imide stretch), 1102 cm1, 747 cm1 (medium, CAOAC imide ring deformation). 1H NMR (400 Hz, CDCl3):d (ppm) 8.23 (2Ha, singlet), 8.08 (2Hb, singlet), 7.42 (2Hc, doublet, J = 8.4 Hz), 7.24(4Hf, doublet, J = 8.4 Hz), 7.17(4Hd, doublet, J = 8.4 Hz), 7.11(4Hg, doublet, J = 8.4 Hz), 7.01(4He, doublet, J = 8.4 Hz), 2.86 (6Hi, triplet, J = 19.6 Hz), 1.71(6Hh, singlet), 1.68 (6Hj multiplet, J = 20), 1.27, (6Hk triplet, J = 20), 0.094 (6Hl, singlet), 0.102 (6Hm, singlet). 2.3. Sample preparations The crude precursors were precipitated with excesses of water. Then they were exhaustively washed using soxhlet extraction, initially with methanol/water in a volume ratio of 2:1 followed by acetone/water in a volume ratio of 5:1; each for a period of 24 h. The precursor was dried at 90 °C in a vacuum for 12 h, before it was poured into a Teflon casting and heated stepwise for curing, from 25 °C to 100 °C for 1 h, 100–150 °C for 1 h, 150–200 °C for
Scheme 1. Synthetic route of PI S1.
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1 h, and 200–310 °C for 1 h. The temperature was then allowed to decrease from 300 °C to 25 °C for 2 h, to give a tough yellowish film.
2.4. Characterization Elemental Analysis (EA) was performed by using a Perkin Elmer PE 2400 CHN and CHNS elemental analyser. Molecular Weight Average (Mn) and Molecular Weight Distribution (MWD) was determined by using GPC, in accordance with the ASTM D529605 [25]. An Inherent Viscosity test was performed; by using an Ubbelohde Calibrated Viscometer Tube in accordance with ASTM D 4603-03 [26] at 25 °C and approximately 30–50 mg of sample in 10 mL of THF was used. FTIR Spectrum was obtained using a Spectrum GX Perkin Elmer Model in accordance with ASTM E 1252 [27], at a wave length 4000–400 cm1. Samples were purified by Soxhlet, prior to measurements by scanning, four times. The H NMR was performed using a Bruker 400 Ultra Shield TM Model at a frequency of 400 MHz in an ambient temperature with deuterated chloroform as a solvent and tetramethlsilane as a standard reference. The surface morphology was observed using SEM equipment; the model used was a ZESS Supra 35VP at 5 K magnifications. The refractive index was obtained by using Refractometer model ATAGO, in accordance with ASTM D542 [28], at a wavelength of 589 nm at room temperature. The dielectric constant was obtained using LCR meter 4284 A, in accordance with ASTM D150-98 [29], at room temperature. UV–VIS absorbance spectroscopy was performed using UV–VIS spectrometer Lambda 35 Perkin Elmer Model at wavelength range 300–800 nm.
3. Results and discussions 3.1. Synthesis considerations The preparation of the three polyimides is shown in the reaction, as in Schemes 1 and 2. The synthesis of S1 was performed under normal addition and polycondensation into polyamic acid, and intermediately followed by chain cyclization to afford a quantitative yield of the polyimide product (Table 1). In S2 and S3, the silicon unit was first allowed to react with the excess anhydride, to ensure the end-capping of the latter onto the former. This allows the anhydride end-capper to copolymerize further with the dianiline during the subsequent stage. The copolymerization sequence between the three units can be monitored by following this procedure. In a three monomeric system, it has been generally shown that the sequence of adding each monomer significantly affects both the product’s molecular weight and morphology. This is due to different terminal group and monomer reactivities during the copolymerization process. Ren et al., found that the sequence of adding 4,4-diamino diphenylmethane with 1,4-bis (3-amino propyldimethylsilyl) benzene and pyromellitic dianhydride (PMDA) markedly affected the dielectric constant. It would be expected that by changing the sequence, such as by reacting the three monomers in a single stage in this work, that its properties might correspondingly change. The inherent viscosity showed that the products were of a moderate molecular weight, as substantiated by the Mn and Mw values. The synthesized polyimide products displayed polydispersity between 1.7 and 2.5. Based on the average molar mass Mn, obtained from GPC and the elemental analysis, a repeat unit for S1–S3 can be proposed, as in Fig. 2.
Scheme 2. PDMS Capping reactions with the anhydride monomer and the complete reaction to form poly(imide siloxane) block copolymer S2 (n = 1) and S3 (n = 35).
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M.B.H. Othman et al. / Materials and Design 32 (2011) 3173–3182 Table 1 Yield percentage, Inherent viscosity and molecular weight distribution of polyimides.
a
PI series
Yield percentage (mol/mol%)
Inherent viscosities (dLg1)a
Molecular weight distribution (gmol1)
Polydispersity PD
Mn
Mw
(Mw/Mn)
Sl S2 S3
96.78 86.07 82.46
0.32 0.19 0.29
12,736 13,444 10,055
31,799 22,975 18,146
2.4968 1.7089 1.8047
Inherent viscosity measure in NMP at 25 + 0.1 °C with concentration of 0.3 g/dL.
3.2. Structure conformation Fig. 3 shows the FTIR spectra of the three polyimides in the range of 3800 and 1200 cm1. No apparent peaks were present in
the region 3600–3200 cm1 corresponding to primary and secondary amides, which confirmed an almost complete imidization reaction, to yield polyimide from their respective polyamic acid precursors. A relatively small peak occurred at 2950–2820 cm1
Fig. 2. The propose of repeated unit of PI series based on GPC and elemental analysis.
Fig. 3. FTIR spectra: (a) representative, and (b) the pure polyimide (S1) and the silicon incorporated polyimide (S2 and S3).
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for the S1 sample, representing the methyl group of the isopropyl bridge, of the dianiline moiety. The intensity of these peaks increase correspondingly for the S2 and S3 samples, which is consistent with the increasing methyl content, derived from the incorporated silicone units. The successful incorporation of silicon units could be further substantiated by the presence of strong and broad peaks at the region 1090–1020 cm1, representing the SiAOAC and SiAOASi bonds (open chain). This characteristic peak was not detected in sample S1. Fig. 4 shows the NMR spectra of the three series of polyimides. The incorporation of the silicone unit in spectra S2 and S3 were proven by the presence of peaks at the region 1.5–0.5 ppm corresponding to the methyl hydrogen attached to the silicon atom.
These peaks were not detected in the case of polyimide S1. The incorporation of the siloxane unit was further substantiated with the presence of peaks at 2.8 ppm (triplet), 1.7 ppm (multiplet), and 1.2 ppm (triplet), which corresponds to the propyl proton of bis-(3-aminopropyl) terminated siloxane moiety. These peaks were again not detected in the S1 spectrum. Peaks corresponding to all aromatic hydrogen occurred within the region 6.8–8.3 ppm and that of methyl protons at 1.8 ppm at the polyimide moiety. Apparently, no substantial chemical shift occurred to these protons during the incorporation of siloxane units, implying that the spacer propyl unit which connect the polyimide and siloxane unit, does not pose any local electronic perturbation to either of these groups.
Fig. 4. 1H NMR (400 MHz) spectra of: (a) S1, (b) S2, and (c) S3 in CDCl3.
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3.3. Morphology of copolymers Morphology of the copolymerized polyimides was investigated using SEM. Fig. 5a of the pure polyimide sample, was essentially featureless. Fig. 5b and c revealed the existence of a very fine phase separation in the form of microspheres; presumably a siloxane segment in the polyimide matrix. These microspheres became far more significant for the S3 sample than the S2 sample. It has been established that phase separation results in a 20% silicone incorporated polyimide system [24]. In this study, the concentration of silicone was deliberately set at a 10% weight ratio to minimize this phase separation effect. The differential scanning calorimetry (DSC) scan for pure polyimide S1 revealed a thermal transition at 244 °C corresponding to its glass transition. For polysiloxane incorporated samples (S2 and S3), the thermal transitions in the 200–206 °C region corre-
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sponds to the glass transition of these copolymers. Due to the chain flexibility of the copolymerized polysiloxane unit, both S2 and S3 samples displayed a lower Tg when compared to the pure polyimide. These reduction revealed that the copolymers are less rigid compared to the pure polyimide. The glass transition for S3 is both lower (203 °C) and broader than that of S2 (Tg = 209 °C). This is mostly due to the presence of a longer polysiloxane segment in S3 which rendered its higher chain flexibility. The Tg of the polysiloxane segment, which would otherwise been detected in the DSC scan in the sub-ambient region, was not observed. This could be due to the micro phase separation morphology, whose scale of thermal transition was too small for detection. Another plausible reason would be that the siloxane unit was tightly confined between the rigid polyimide chains which limit any segmental movement to the former. As expected, no melting transition was observed for all samples within this range of temperatures (Fig. 6). 3.4. Refractive index and dielectric constant The refractive index and dielectric constant values of the three polyimides are shown in Table 2. It can be seen that S1 displayed the highest refractive index, followed by S2, then S3. These results can be attributed to increased silicone unit content. Based on the Lorenz–Lorenz equation, the refractive index n is dependent on Ru, the molar refractivity and Vu, the molar volume, as in the following equation:
Ru ¼ V u
Fig. 5. The surface morphology of polyimide series: (a) parent PI, (b) PI incorporate DMS and (c) PI incorporate PDMS.
ðn22 1Þ ðn22 þ 2Þ
ð2Þ
Ru and Vu are additive functions of the respective group’s composition in the repeating unit. Brunchi et al., verified that the increases of silicone repeating unit, would adversely affect the ratio of Ru/Vu, and subsequently the refractive index [30]. This is due to the bigger size of Si atom as compared to that of carbon which significantly affects the molar volume of the system. Hence, it was expected that with a higher content in the silicone unit, the refractive index would correspondingly decrease. A close relationship of the refractive index to the dielectric constant is shown by the relationship in Eq. (1). Both quantities involve an interaction of radiation with the local electronic environment of chemical bonds. Table 2 shows the dielectric constant for samples S1, S2, and S3 with a decreasing trend as the polysiloxane content in the polymer increased. This result was consistent with that of the refractive index. Silicon is comparatively larger than a carbon atom and the SiAO bond is more flexible than the CAC bond. This is substantiated in several DMA and mechanical property studies [24,31]. Thus, it can be affirmed that the bulky silicon units would be less mobile. Its presence affects the bulk movement of the whole polyimide network which reduces the efficiency of the dipole in reacting to polarity change during treatment with an alternating frequency. Furthermore, the molar polarization significantly decreases as the result of an increase in free volume. Several recent studies have demonstrated a similar trend of a decreasing dielectric constant, with an increasing siloxane content into polyimide structures [32,33]. The values obtained in this study of dielectric constant at 1 kHz, are within the range obtained from these other studies. These values were larger than those derived from the Maxwell equation (shown in Table 2), because the Maxwell equation is only applied at the optical frequency of approximately 1015 Hz, which only involves electronic polarization. Fig. 7 shows the variation of dielectric constant with an applied electric field frequency. The results show that with an increasing frequency, e decreases very gradually. This variation is attributed to the frequency dependence of the polarization mechanisms.
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Fig. 6. DSC thermogram (2nd scan) for three systems of PI.
Table 2 Refractive Index, dielectric constant and cutoff wavelength (/nm) from UV–Vis spectra of polyimides and copolymers.
a b c
Polymer
Refractive indexa
Dielectric constant estimationb
Dielectric constantc
Absmax (k/nm)
Sl S2 S3
1.4983 1.4936 1.4922
2.25 2.23 2.22
2.90 2.57 2.43
407 408 413
Measured at sodium wavelength 589 nm at temperature 25 °C. Estimated from Maxwell equations = n2. Measured at 1 kHz in room temperature.
The overall contribution of dielectric constant can be integrated into the following equation:
eTotal ¼ eelectronic þ eatomic þ edipole
ð3Þ
At a high frequency regime, in the range 1013–1018 Hz, the polarization mechanism involves electronic polarization, followed by the frequency range 1010–1013 Hz, which includes atomic polarizability, and finally below 109 Hz, which includes dipole polarization. Within the 103–106 Hz frequency range measured in this
work, the summation of all three modes of polarization were involved. However, as the frequency gradually increased, the atomic polarizability contribution began to diminish. Effectively, the dielectric constant gradually decreased with the increasing frequency. It can be observed that the dielectric constant slightly increases from 0 to 10 kHz for S1 and S2. This peculiarity can be attributed to a phenomenon known as electrode polarization. It originates from a high-impedance layer on the electrode surface, mainly due to an imperfect contact between electrode and the specimen during measurement. Electrolytic products further aggravate this effect. This suggests that at low frequency, there is sufficient time for any slight conduction through the specimen, to transfer the entire applied field across the very thin electrode. This anomaly is characterized by an apparent increase in relative permittivity, as observed at 0–10 kHz [34,35]. The dielectric loss of the polyimides is represented in Fig. 8. and refers to the imaginary part of a relative permittivity, as in the following equation:
er ¼ e0r ðxÞ þ ie00r ðxÞ
ð4Þ
where the first term on the right side of the equation refers to a real component and the second term refers to an imaginary component. The dielectric loss of the S1 sample increased almost proportion-
Fig. 7. Dielectric constant of PI series corresponding to frequency 1 kHz, 10 kHz, 100 kHz and 1000 kHz at room temperature.
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Fig. 8. The dielectric loss of the pure polyimides S1 and polysiloxane–polyimide copolymers S2 and S3.
ately with the frequency; at least within the measured frequency range. This was due to the inability of the polarizable units in the polyimide chains to align in phase with the increasing electric field frequency. There will be a phase lag with the oscillating electric field, which then contributes to the dielectric loss. Physically, the loss is constituted in the form of dissipative heat energy [36]. In the case of the siloxane-incorporated polyimide samples, it showed an initial decrease in dielectric loss, followed by a sinusoidal variation at a higher frequency range. However, the value obtained in the range 2.59 104–7.5 103, are reasonably low compared to some of the other organic polymers and inorganic materials. The dielectric loss for sample S3 is higher compared to that of sample S2. This was apparently due to the higher siloxane content in the S3 sample. Since the siloxane chains in S3 are longer, this induced inefficient bulk chain flexibility to the copolymers. This effect increased a phase lag with that of the applied oscillating electric field. 3.5. UV–VIS absorbance spectroscopy The UV–Vis absorbance spectroscopy of the three series of polyimides is shown in Fig. 9. All synthesized siloxane-incorporated polyimides displayed a cut-off absorption wavelength below
450 nm. Two characteristic broad peaks were detected at 360 nm and 400 nm wavelengths, which corresponds to the p–p transition due to the charge transfer complex of the aromatic rings of the polyimide chain backbone [37,38]. It is obvious that the incorporation of polysiloxane into the polyimide chain does not affect any significant change in the absorption maxima, or the transparency of the copolymer. It was reported by Kai et al., that except at NIR region, poly-(methylsiloxane) does not display any characteristic absorption in the UV–visible region. There are noticeable decreased in absorbances in S1 as compared to S2 and S3 at 400 nm region. Several factors contribute to the change in absorbance, albeit extinction coefficient during the UV–visible scan, such as complex formation, particle size distribution [39], and the extent of imidization during polyimide curing [37]. It has been shown that in addition to the red shift effect, the extinction coefficient gradually decreases with respect to the level of imidization curing in the polyimide. A decrease in absorbency for S1 at the 400 nm region could possibly reflect a complete imidization curing of the sample, while samples S2 and S3 showed a comparatively higher absorbency due to inefficient imidization curing, as the result of the presence of siloxane units which posed a steric hindrance during the curing process. Visually, all synthesized PI
Fig. 9. The UV–VIS absorbance spectroscopy of the three polyimide.
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showed a yellowish colouration, which is consistent with the tailing absorption into the visible 450 nm wavelength for all samples. 4. Conclusion Copolymerization of the silicone unit into the polyimide backbone was successfully performed. FTIR and H NMR spectra verified that polyimide and poly-(imide siloxane) were completely imidized from poly-(amic acid) and poly-(amic siloxane acid), respectively. All polyimide series were synthesized with moderate molecular weights and narrow molecular weight distributions. SEM showed a micro phase separation between the two phases. Incorporation of silicon unit into the polyimide chain significantly reduced the refractive index and dielectric constant, with a reasonably low dielectric loss. However, siloxane unit does not affect UV– visible absorption or transparency in the copolymer. Acknowledgment The authors wish to thank the USM for sponsoring this Project under USM-RU-PGRS 8032006. References [1] Ghosh M, Mittal K. Polyimides: fundamentals and applications. CRC Press; 1996. [2] Cha H, Hedrick J, DiPietro R, Blume T, Beyers R, Yoon D. Structures and dielectric properties of thin polyimide films with nano foam morphology. Appl Phys Lett 1996;68:1930–2. [3] Feger C, Khojasteh M, McGrath J. Polyimides: materials, chemistry and characterization. SPE; 1988. [4] Mittal K. Polyimides: synthesis, characterization, and applications, vols. 1 and 2. New York: Plenum Press; 1984. [5] Courtois H, Buisson O, Chaussy J, Pannetier B. Miniature low temperature high frequency filters for single electronics. Rev Sci Instrum 2009;66:3465–8. [6] Binnig G, Rohrer H, Gerber C, Weibel E. Tunneling through a controllable vacuum gap. Appl Phys Lett 2009;40:178–80. [7] Zhang Y, Ke S, Huang H, Zhao L, Yu L, Chan H. Dielectric relaxation in polyimide nanofoamed films with low dielectric constant. Appl Phys Lett 2008;92(052910):1–3. [8] Jiang L, Liu J, Wu D, Li H, Jin R. A methodology for the preparation of nanoporous polyimide films with low dielectric constants. Thin Solid Films 2006;510:241–6. [9] Mehdipour Ataei S, Saidi S. Structure–property relationships of low dielectric constant, nanoporous, thermally stable polyimides via grafting of poly (propylene glycol) oligomers. Polym Adv Technol 2008;19:889–94. [10] Zhao XY, Liu HJ. Review of polymer materials with low dielectric constant. Polym Int 2010;59:597–606. [11] Yu Y-Y, Chien W-C, Chen S-Y. Preparation and optical properties of organic/ inorganic nanocomposite materials by UV curing process. Mater Des 2010;31:2061–70. [12] Cornelius C, Marand E. Hybrid inorganic–organic materials based on a 6FDA– 6FpDA–DABA polyimide and silica: physical characterization studies. Polymer 2002;43:2385–400. [13] Dhara MG, Banerjee S. Fluorinated high-performance polymers: poly(arylene ether)s and aromatic polyimides containing trifluoromethyl groups. Progr Polym Sci 2010;35:1022–77.
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