Polyimide Nanofoams Based on Ordered Polyimides Derived from

Chem. Mater. 1997, 9, 105-118

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Polyimide Nanofoams Based on Ordered Polyimides Derived from Poly(amic alkyl esters): PMDA/4-BDAF Kenneth R. Carter,* Richard A. DiPietro, Martha I. Sanchez, and Thomas P. Russell† IBM Research Division, Almaden Research Center, 650 Harry Rd., San Jose, California 95120-6099

Priya Lakshmanan‡ and James E. McGrath Virginia Polytechnic Institute and State University, Department of Chemistry and NSF Science and Technology Center: High Performance Polymeric Adhesives and Composites, Blacksburg, Virginia 24061-0344 Received April 19, 1996. Revised Manuscript Received September 17, 1996X

A means of generating foams of high-temperature polymers, polyimides, has been developed for use in dielectric layers in microelectronics. In these systems, the pore sizes generated are in the tens of nanometers range, thus the term “nanofoams”. The foams are generated by preparing phase-separated block copolymers with the majority phase comprised of polyimide and the minor phase consisting of a thermally labile block. Films are cast, solvent is removed, and the copolymers are cured, causing phase separation of the two blocks. The labile blocks are subsequently removed via thermal treatments leaving pores having a size and shape commensurate with the size and shape of the original copolymer morphology. The polyimide derived from pyromellitic dianhydride (PMDA) and 2,2-bis[4-(4-aminophenoxy)phenyl]hexafluoropropane (4-BDAF) was used as the matrix materials for the generation of nanofoams, and poly(propylene oxide) oligomers were used as the thermally labile constituent. The synthesis and characterization of the copolymers were performed, and the process for obtaining nanofoams was optimized. The foams were characterized by a variety of techniques including TEM, SAXS, WAXD, DMTA, density, and refractive index measurements. Thin-film, high-modulus foams with good mechanical properties can be synthesized using the copolymer/nanofoam approach.

Introduction High-performance polymers, such as polyimides, polyesters, and aramids, have found a number of uses in a variety of electronic devices and components.1 The use of polyimides, in particular, can be attributed to their possession of many favorable properties including high thermal stability, ease of processing, low stress/coefficient of thermal expansion (CTE), and very good electrical properties. The favorable electrical properties include low dielectric constant, high resistivity, and high breakdown voltage. The importance of polyimide possessing a low dielectric constant cannot be understated. In a microelectronic device, the propagation velocity, V, of a pulsed signal in a line, is inversely proportional to the square root of the dielectric constant, , of the propagation medium:

V ) c/x The lower the dielectric constant of the medium, the faster the signal velocity and the shorter the device cycle time. In addition, lowering of the dielectric constant of the medium decreases signal “crosstalk”, thereby allowing denser circuit patterns. †

Current address: Department of Polymer Science & Engineering, University of Massachusetts, Amherst, MA 01003. ‡ Current address: Bell Labs, 2000 NE Expressway, Norcross, GA 30071. X Abstract published in Advance ACS Abstracts, December 1, 1996. (1) Lai, J. H., Ed. Polymers for Electronic Applications; CRC Press: Boca Raton, FL, 1989.

S0897-4756(96)00242-6 CCC: $14.00

Generally speaking, many common polyimides, such as the polyimide derived from pyromellitic dianhydride (PMDA) and oxydianiline (ODA), PMDA/ODA, have measured dielectric constants around 3.1 ( 0.4 (typically measured at 1 MHz). PMDA/ODA has found widespread use due to its high thermal stability, ease of processing, good mechanical properties, and solvent resistance. These properties are, in part, due to the ordering of the polymer and the alignment of polymer chains parallel to the surface.2-4 This orientation leads to an anisotropy in the dielectric constant observed for this and other polyimides,5 which can limit the use of these materials when an isotropic, low dielectric constant is essential. A variety of approaches have been studied to synthesize polyimides with lower dielectric constants. One approach has been to develop materials high in organofluorine content, often by the incorporation of pendant perfluoroalkyl groups. For example, Hoechst6 developed Sixef polyimide, which contains hexafluoroisopropylidene linkages, with a dielectric constant of 2.6, while DuPont7 has reported the synthesis of polyimides based (2) Russell, T. P.; Grugger, H.; Swalen, J. D. J. Polym. Sci., Polym. Phys. 1983, 21, 1745. (3) Takahashi, N.; Yoon, D. Y.; Parrish, W. Macromolecules 1984, 17, 2585. (4) Russell, T. P. J. Polym. Sci., Polym. Phys. Ed. 1986, 22, 1105. (5) Boese, D.; Herminghaus, S.; Yoon, D. Y.; Swalen, J. D.; Rabolt, J. F. Mater. Res. Soc. Symp. Ser. 1991, 227, 379. (6) Haider, M.; Chenevey, E.; Vora, R. H.; Cooper, W.; Glick, M.; Jaffe, M. Mater. Res. Soc. Symp. Proc. 1991, 227, 379.

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on 9,9-disubstituted xanthene dianhydrides, 6FXDA/ 6FDAm (9,9-bis(trifluoromethyl)xanthenetetracarboxylic dianhydride/2,2-bis(4-aminophenyl)-1,1,1,3,3,3-hexafluoropropane) which has a dielectric constant of ∼2.5. The methodology to develop highly fluorinated polyimides is limited, to a certain extent, by synthetic difficulties associated with the incorporation of greater amounts of pendant perfluoroalkyl groups. Incorporating bulky, nonaromatic groups into the polymer chain can also lead to a lowering of the dielectric constant. For example, Volksen et al. reported the synthesis of a cycloaliphatic polyimide, PMDA/DACH (pyromellitic dianhydride/1,4diaminocyclohexane), based upon a cyclohexyl repeating group, which has a dielectric constant of 2.6.8 A low dielectric constant material can also be made by forming a composite where the added products (which possess a low dielectric constant) are submicron in size and distributed uniformly throughout the matrix. Alternatively a block copolymer can be made where one of the blocks has a low dielectric constant. For example PMDA/ODA-perfluoroalkylene aryl ether block copolymers have been prepared where the dielectric constant of PMDA/ODA was lowered from 3.1 to 2.8 without significant reduction of thermal or mechanical properties.9 Of course, this approach is limited by the dielectric constant of the added coblock. A potentially better route is to incorporate pores within the polymer matrix, i.e., generating a foamed polymer. For microelectronic applications it is mandatory that the size of the pores be much smaller than the film thickness and any features, i.e., much less than 1 µm. The foam must have good mechanical properties for stability under fabrication conditions and intended uses and preferably have a closed cell structure to minimize solvent penetration and retention during processing. A variety of approaches have been reported for the foaming of polyimides, such as the use of foaming agents,10-12 partial degradation generating a foaming agent,13,14 the inclusion of hollow microspheres,15 and microwave processing.16 Most methods result in materials with large void sizes and open-pore structures, making them unsuitable for microelectronic applications. High-temperature thermoplastic foams can also be made by a novel block copolymer approach.17-19 Here the pore sizes are on the nanometer size scale, hence (7) Auman, B. C.; Trofimenko, S. Polym. Prepr. (Am. Chem. Soc., Div. Polym. Chem.) 1992, 34, 244. (8) Volksen, W.; Sanchez, M. I.; Chah, H. J.; Yoon, D. Y. Polym. Prepr. (Am. Chem. Soc., Div. Polym. Chem.) 1994, 36, 709. (9) Labadie, J. W.; Sanchez, M. I.; Cheng, Y. Y.; Hedrick, J. L. Mater. Res. Soc. Symp. Proc. 1991, 227, 43. (10) Smearing, R. W.; Floryan, D. C. U.S. Patent 4,543,368 to General Electric, 1985; CA:104:34973. (11) Krutchen, C. M.; Wu, W. P. U.S. Patent 4,535,100 to Mobil Oil, 1985; CA:103:161477. (12) Hoki, T.; Matsuki, Y. European Patent 186308 to Asahi Chem., 1986; CA:105:192247. (13) Meyers, R. A. J. Polym. Sci. A-1 1969, 7, 2757. (14) Carleton, P. S.; Farrisey, W. J.; Rose, J. S. J. Appl. Polym. Sci. 1972, 16, 2983. (15) Narkis, M.; Paterman, M.; Boneh, H.; Kenig, S. Polym. Eng. Sci. 1982, 22, 417. (16) Gagliani, J.; Supkis, D. E. Adv. Astronaut. Sci. 1979, 38, 193. (17) Labadie, J. W.; Hedrick, J. L.; Wakharkar, V.; Hofer, D. C.; Russell, T. P. IEEE Trans Compon., Hybrids, Manuf. Technol. 1992, 15, 925. (18) Hedrick, J. L.; Labadie, J. W.; Russell, T. P.; Hofer, D. C.; Wakharkar, V. Polymer 1993, 34, 4717. (19) Hedrick, J. L.; Labadie, J. W.; Russell, T. P.; Wakharkar, V.; Hofer, D. C. In Adv. Polyimide Sci. Technol., Proc. Int. Conf. Polyimides; Feger, C. L., Khojasteh, M. M., Htoo, M. S., Eds.; Technomic: Lancaster, PA, 1993.

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Figure 1. Schematic representation of foam formation via phase separated imide/labile block copolymer precursors.

the term “nanofoam”. The foams are prepared from block copolymers consisting of thermally stable and thermally labile blocks, the latter being the dispersed phase (Figure 1). Foam formation is effected by thermolysis of the thermally labile block, leaving pores with a size and shape corresponding to that of the initial copolymer morphology. The utility of this approach was first demonstrated using triblock copolymers consisting of the high-Tg block, poly(phenylquinoxaline) (PPQ), and either poly(propylene oxide) (PO) or poly(methyl methacrylate) (PMMA) as the thermally labile block.17 A series of PPQ/PO or PMMA copolymers were made where the weight percent of the thermally labile block was 14-24% of the copolymer. After appropriate thermal treatment and decomposition of the labile blocks, an 8-12% reduction in density was observed along with a reduction in dielectric constant from 2.8 for PPQ to 2.4 for the foam. Scattering and electron microscopy measurements clearly showed the presence of pores tens of nanometers in size. The foamed polymers showed lower residual stress values, but the foams were found to collapse upon heating above the Tg. In addition, PPQ foams would have limited use in microelectronic applications since the copolymers are soluble only in selected chlorinated solvents and m-cresol, making processing difficult. The block copolymer approach has been extended to foamed polyimide structures. Amorphous high-temperature polymer nanofoams comprised of oxydiphthalic anhydride (ODPA) and 9,9′-bis(4-aminophenyl)fluorene (FDA) and PMDA/FDA20 as well as PMDA/1,1-bis(4aminophenyl)-1-phenyl-2,2,2-trifluoroethane (3FDA)21 have been successfully prepared in a manner similar to that described for PPQ. To meet solvent resistance criteria, ordered polyimides such as PMDA/ODA, PMDA/ PDA, and BPDA/PDA were studied as matrix materials for nanofoams.22 Both PO and PMMA block copolymers were produced via the corresponding poly(amic alkyl ester) routes, thereby yielding a solvent-processable polyimide precursor. While it was possible to foam these copolymers at temperatures near 300 °C, the porous structures started to collapse at higher temperatures and were fully collapsed before reaching the Tg near 400 °C. One possible origin for the lack of stability is that these polymers undergo ordering at elevated (20) Charlier, Y.; Hedrick, J. L.; Russell, T. P.; Volksen, W. Polymer 1995, 36, 987. (21) Hedrick, J. L.; Hawker, C. J.; DiPietro, R. A.; Je´roˆme, R.; Charlier, Y. Polymer 1995, 36, 4855. (22) Hedrick, J. L.; Russell, T. P.; Labadie, J. W.; Lucas, M.; Swanson, S. A. Polymer 1995, 36, 2685.

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temperatures. The progressively increasing ordering at temperatures in excess of 300 °C forces a volume contraction which, in turn, leads to foam instability even below the Tg of the polymer. However, if ordered structures can be made to behave as effective crosslinks, the foam stability and solvent resistance should be enhanced. Therefore, it is important to determine if other ordered polyimide systems could be used to generate stable nanofoams. Polyimides derived from PMDA and 2,2-bis[4-(4aminophenoxy)phenyl]hexafluoropropane (4-BDAF), PMDA/4-BDAF, a tough, moderately high Tg (305 °C), polymer with low water absorption and a dielectric constant of 2.8 (1 MHz)23 were used in this study. The good solvent resistance and high thermal stability of PMDA/4-BDAF has been ascribed, in part, to its semicrystallinity.24 The polymer has a very high melting point, ca. 470 °C with calorimetric evidence that ordering begins at temperatures above 300 °C. The intrinsically low dielectric constant of PMDA/4-BDAF could be lowered significantly if a foamed structure could be produced. Consequently, copolymers of PMDA/4-BDAF with PO were studied to examine the effect of ordering on the stability of nanofoam structures. Experimental Section Materials. All materials were commercially available and used as received unless otherwise noted. Pyromellitic dianhydride (PMDA, Chriskev Co.) and 2,2-bis[4-(4-aminophenoxy)phenyl]hexafluoropropane (4-BDAF, Central Glass Co., Ltd.) were sublimed twice prior to use. Hydroxy-terminated poly(propylene oxide) of nominal molecular weights 5000 and 10 000 g/mol were kindly supplied by Dow Chemical Co. mand p-diethyldichloropyromellitate, aminobenzoate-terminated poly(propylene oxide), aminophenyl carbonate terminated poly(propylene oxide), and 3,5-diaminobenzoate terminated poly(propylene oxide) were prepared as described by Hedrick et al.20 Characterization. NMR spectra (in DMSO-d6 or CDCl3) were recorded on an IBM WP 250 spectrometer operating at 250.1 MHz (1H) and 62.9 MHz (13C) with chemical shifts reported in ppm downfield from tetramethylsilane. Polymer films for mechanical and thermal analysis were cast from N-methyl-2-pyrrolidinone (NMP) and heated to remove solvent and to cycloimidize the polymer. Glass transition temperatures, taken as the midpoint of the change in slope of the baseline, were measured on a DuPont DSC 1090 instrument at a heating rate of 10 °C/min (N2). Thermal gravimetric analyses (TGA, N2 and air) of the polymer films were conducted on a Perkin-Elmer Model TGA-7 at a heating rate of 10 °C/min or isothermal weight loss was measured at different temperatures over 10 h. Dynamic mechanical measurements were made with a Polymer Laboratories dynamic mechanical thermal analyzer (DMTA) in the tension mode with a heating rate of 10 °C/min (10 Hz, N2). Density measurements were obtained with a density gradient column composed of water and calcium nitrate. The column was calibrated against a set of beads of known densities (Scientific Glass & Instruments, Inc.) and maintained at 25 °C (d ) 1.051.45). At least two specimens were used for each density measurement. Wide-angle X-ray diffraction, WAXD, measurements were performed in both reflection and transmission geometries using a sealed tube X-ray source with 1.542 Å incident radiation. The diffracted X-rays were detected with a scintil(23) Jones, R. J.; Chang, G. E.; Powell, S. H.; Green, H. E. In Polyimides: Synthesis, Characterization, and Applications; Mittal, K. L., Ed.; Plenum: New York, 1984; Vol. 2. (24) Rogers, M. E.; Brink, M. H.; McGrath, J. E.; Brennan, A. Polymer 1993, 34, 849.

Chem. Mater., Vol. 9, No. 1, 1997 107 lation counter mounted behind an analyzer crystal that was step scanned over the angular range of interest. Small-angle X-ray scattering studies were performed on Beamline I-4 at the Stanford Synchrotron Radiation Laboratory. Details of the beamline optics are described elsewhere.25 In these experiments, X-rays of wavelength 1.53 Å were passed through a Mettler FP85 hot stage where the sample was mounted. X-ray scattering profiles were obtained either as the sample temperature was ramped at 10 °C/min or at a specific temperature as a function of time. Experiments were performed under a steady flow of N2 or in air. Samples were prepared by placing stacks of thin films in a sample cell with Kapton windows. Standard procedures were used to correct the scattering profiles for parasitic scattering, electronic noise, and sample absorption. Scattering profiles are presented as a function of scattering vector, Q, which is given by (4π/λ)sin θ, where λ is the wavelength and 2θ is the scattering angle. Transmission electron microscopy (TEM) was performed on a Philips 420T scanning transmission electron microscope operating at 100 kV. The polymer samples were embedded in an epoxy resin and microtomed in cross section. The resulting sections were between 500 and 1000 Å thick. The residual stress in the thin films was obtained on a Tencor Flexus F2300 stress analyzer where the radius of curvature of a silicon wafer is measured, and the stress is calculated from the change in the wafer curvature before and after polymer deposition. The wafers are mounted on a hot stage for in situ high-temperature stress measurements. Refractive index values were obtained on 2-3 µm thick samples spin coated onto quartz or gold-coated quartz wafers which were then cured and foamed. Measurements were made on a Metricon Prism Coupler (PC-2000) which is an optical waveguide technique providing both in-plane and out-of-plane refractive indexes. Preparation of 4,6-Dicarbethoxyisophthalic Acid (Meta Isomer), m-DADE. A 250 mL flask equipped with a stirrer, reflux condenser, and heating mantle was charged with 150 mL of absolute ethanol. PMDA (52.3 g, 0.24 mol) was then added as a solid. The reaction mixture was stirred under argon for 4 h with gentle heating, during which time the temperature rose to approximately 60 °C and the mixture became homogeneous. (Occasionally, complete solubility does not occur; nonetheless, the reaction goes to completion.) When the reaction temperature started to decline, 50 mL of ethyl acetate was added, and the mixture was allowed to cool with stirring overnight. The resulting slurry was vacuum filtered to give fraction 1 as a white crystalline solid which was washed once with 15 mL of ethyl acetate. The combined filtrate and wash were slowly evaporated at reduced pressure until turbidity ensued. The suspension was heated to dissolution and allowed to cool slowly to room temperature with stirring. The resulting white crystalline solid was filtered (fraction 2), washed once with 15 mL of ethyl acetate, and the combined filtrate and washings were treated as before. The white, crystalline precipitate was filtered (fraction 3) and washed with 15 mL of 1:1 ethyl acetate/ hexane and 15 mL of hexane. The combined filtrate and washings were again treated as before. The resulting slurry was filtered (fraction 4) and washed with 15 mL of hexane. The final filtrate was discarded. All fractions were dried in a vacuum oven overnight at 50 °C (0.2 Torr). HPLC analysis revealed that fractions 1 and 2 were primarily the para diacid diester and fractions 3 and 4 were primarily the meta isomer. The weights and m/p ratio (normalized to a combined 100%) of a typical set of four fractions is illustrated in Table 1. Pure meta isomer can be prepared by recrystallization of fractions 3 and 4 from hexane/ethyl acetate. Diethyldichloropyromellitate, DPDC. The various DPDC monomers used in this study, p-DPDC,m-DPDC, and mixed isomer DPDC, were prepared from the appropriate diester diacids of PMDA. The preparation of m-DPDC is outlined here. (25) Russell, T. P. In Handbook of Synchrotron Radiation; Brown, G. S., Moncton, D. E., Eds.; North-Holland: Amsterdam, 1991; Vol. 3.

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Table 1. Synthesis and Fractional Precipitation of PMDA Dialkyl Esters fraction

wta (g)

m/p ratio

percentage yield

1 2 3 4

25.4 6.8 17.5 22.1

6.4/93.6 3.0/97.0 96.1/3.9 84.1/15.9

34.1 9.1 23.5 29.6

a

Total ) 71.8 g (96.3%).

A 250 mL flask equipped with a stirrer, chilled water reflux condenser, addition funnel, and heating mantle was charged with the m-diester diacid of PMDA (DADE, 23.3 g, 0.075 mol). Ethyl acetate (100 mL) was added, and the reaction mixture was heated to 55-58 °C. Oxalyl chloride (27.3 g, 18.7 mL, 0.215 mol) was slowly added to the reaction mixture over 3-4 h.26a The reaction temperature was maintained at 55-58 °C for an additional 12 h. When the reaction was judged complete (HPLC), the batch was cooled, a vacuum was applied to the reactor, and the excess oxalyl chloride was removed (along with ethyl acetate) by vacuum distillation (32-55 °C at 660 Torr). The batch was distilled to a thick syrup, fresh dry ethyl acetate (100 mL) was added via vacuum transfer, and the resulting solution was again distilled to a thick syrup.26b At this point, all of the residual oxalyl chloride had been removed (as determined by HPLC) and the batch was evaporated to a yellowish oil which was crystallized twice from hexane to yield the m-DPCD in 70-80% yield, mp 53-55 °C. 1H NMR (CDCl3) δ 8.30 and 7.92 (d, 1H, m-Ar-H), 8.10 (s, 1H, p-Ar-H), 4.504.38 (q, 4H), 1.39-1.36 (t, 6H). 3,5-Diaminobenzoate-Terminated Poly(propylene oxide). A 50 mL flask was charged with 20.0 g (8.65 mmol OH) of hydroxy-terminated poly(propylene oxide) (10 000 g/mol nominal molecular weight) which had been dried by azeotropic distillation of toluene. The toluene was subsequently removed under reduced pressure. To this flask, 0.91 g (9 mmol) of triethylamine was added, followed by 2.01 g (8.7 mmol) of distilled 3,5-dinitrobenzoyl chloride. The mixture was allowed to react overnight with stirring and diluted with 25 mL of diethyl ether. The suspension was filtered to remove the precipitated triethylamine hydrochloride salt, and the solvent was evaporated. The amorphous residue was then dissolved in 50 mL of ethyl acetate and hydrogenated at 3 atm in the presence of 150 mg of Pearlman’s catalyst (Pd(OH)2 on activated carbon, Aldrich) for 48 h. The mixture was filtered, solvent was removed in vacuo, and the resulting diamino derivative was dried azeotropically with toluene. The toluene was removed under high vacuum yielding the product as a viscous, clear brown liquid with a molecular weight of 7900 g/mol as determined by potentiometric titration with perchloric acid. 4-Aminophenyl Carbonate Terminated Poly(propylene oxide). The 4-aminophenyl carbonate terminated poly(propylene oxide) (PO) oligomers were prepared by reacting a 5-fold excess of nitrophenyl chloroformate with hydroxyterminated propylene oxide oligomers (5000 and 10 000 g/mol nominal molecular weight) in THF using pyridine as an acid acceptor. After several hours, the insoluble amine hydrochloride salt was removed by filtration and the 4-nitrophenyl chloroformate complex was reduced with Pearlman’s catalyst. After the reduction, the oligomer was washed with water. This procedure yielded oligomers of 4000, 5600, and 6500 g/mol as determined by 1H NMR and potentiometric titration with perchloric acid. PMDA/4-BDAF Polyimide: Two-Stage Poly(amic acid) Route. A 50 mL three-neck flask fitted with an overhead stirrer was charged with 4.495 g (8.66 mmol) of 4-BDAF and (26) (a) There was a significant amount of off-gases throughout the addition and until the reaction was complete. These gases contain HCl, CO, and oxalyl chloride. They should be scrubbed with a caustic scrubber or dissolved in running water by using a nonmetallic aspirator. (b) The distillate contains oxalyl chloride which should be carefully quenched with an appropriate alcohol and neutralized prior to disposal.

25 mL of NMP. The flask was heated with stirring under an argon blanket to dissolve the diamine. After a homogeneous solution was obtained, the flask was cooled to 5 °C and 1.891 g (8.66 mmol) of PMDA was added and washed into the flask with 13 mL of NMP. The solution was allowed to warm to room temperature and stir for 24 h. The resulting viscous poly(amic acid) solution was filtered (1 µm) and cast or spun onto the appropriate substrate and thermally cured to give clear, tough thin films. The cure cycle used for all of the copolymers consisted of several ramp and hold cycles as follows: heating to 80 °C and hold for 1 h, heat to 130 °Cshold 1 h, heat to 250 °Cshold 1 h, and final heating to 305 °Cshold 305 °C. The heating rate was maintained at 5 °C min-1, and processing was carried out under an argon atmosphere. PMDA/4-BDAF Polyimide: Two-Stage Poly(amic ethyl ester) Route, with Stoichiometric Imbalance. A 100 mL three-neck flask fitted with an overhead stirrer and addition funnel was charged with 3.613 g (6.97 mmol) of 4-BDAF, 60 mL of NMP, and 1.18 g (14.95 mmol) of pyridine. The flask was heated with stirring under an argon blanket to dissolve the diamine. After a homogeneous solution was obtained, the flask was cooled to 5 °C and a solution of 2.759 g (6.79 mmol) of p-DPDC in 15 mL of THF was added dropwise. The solution was allowed to warm to room temperature with stirring for 24 h. The resulting viscous poly(amic ethyl ester) solution was precipitated into methanol/water (1:1), filtered, and washed three times with water and twice with methanol. The polymer was vacuum dried to constant weight (50 °C at 26 mmHg). PMDA/4-BDAF-co-Poly(propylene oxide) Triblock Copolymer. The triblock copolymers were prepared in a twostage process in which the corresponding poly(amic ethyl ester) coblock could be isolated, analyzed, and subsequently processed into imide triblock copolymers. A representative example follows: A 100 mL three-neck flask fitted with an overhead stirrer and addition funnel was charged with 3.547 g (6.84 mmol) of 4-BDAF, 1.05 g (0.16 mmol) of a 6.5K MW 4-aminobenzoate-terminated poly(propylene oxide) oligomer, 60 mL of NMP and 1.1 g (13.84 mmol) of pyridine. The flask was heated with stirring under an argon blanket to achieve a homogeneous solution. The flask was cooled to 5 °C, and a solution of 2.403 g (6.92 mmol) of p-diethyl pyromellitate diacyl chloride in 15 mL of THF was added dropwise. The solution was allowed to warm to room temperature and stirred for 24 h. The resulting viscous poly(amic ethyl ester) solution was precipitated into methanol/water (1:1), filtered, and washed three times with water and twice with methanol/water (1:1). The polymer was vacuum dried to constant weight (50 °C at 26 mmHg). Foam Formation. The PMDA/4-BDAF-co-poly(propylene oxide) triblock copolymers were dissolved in NMP at concentrations of 10-20% solids. Thin films could be formed by spin coating (1-10 µm) or doctor blading (10-40 µm). The removal of solvent and imidization was accomplished by heating the polymer to 305 °C, as described above, under an argon atmosphere. The poly(propylene oxide) (PO) blocks were decomposed by heating the films to 250 °C for 10 h in air, resulting in PMDA/4-BDAF polyimide foams.

Results and Discussion The block copolymer approach used to generate polyimide nanofoams requires well-defined block segments, both the thermally stable matrix block (imide) and the thermally labile block, which in this study was PO. A critical requirement of the PO blocks is that they must undergo a controllable and quantitative degradation into low molecular weight products which must in turn diffuse easily through the polyimide matrix. Because the decomposition of the PO blocks is thermally initiated, it is equally important that the PO decomposition temperature, Td, be below the Tg of the PMDA/4-BDAF matrix and also high enough to permit film imidization, annealing, and solvent removal. The temperature range between PO degradation, Td, and imide Tg is referred

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Chem. Mater., Vol. 9, No. 1, 1997 109

Scheme 1. Synthetic Route to Aminophenyl-Terminated Poly(propylene oxide) Oligomers

to as the “processing window”. PO possesses sufficient thermal stability at 305 °C, under an inert atmosphere, to permit film preparation and annealing. The PO can then be thermally degraded at temperatures greater than 250 °C in the presence of atmospheric oxygen. Therefore, in this case, the processing window is between 250 and 305 °C. The successful incorporation of poly(propylene oxide) (PO) into these copolymers required the synthesis of the proper amino-terminated PO oligomers. The availability of suitable hydroxy-terminated PO allowed us to explore methods of functionalizing the hydroxy end groups to give amino-terminated oligomers (Scheme 1). Two amino-terminated PO derivatives were prepared for this study, an aminophenyl carbonate end group, which was used to generate imide/PO triblock copolymers, and a 3,5-diaminobenzoate-terminated PO oligomer, which was used to make imide/PO graft copolymers. The aminophenyl carbonate terminated PO oligomers were synthesized by reacting hydroxy-terminated PO with 4-nitrophenyl chloroformate in THF in the presence of pyridine. This reaction mixture was filtered to remove the pyridine hydrochloride salts and then reduced with Pearlman’s catalyst under H2, yielding 4-aminophenyl carbonate-terminated PO oligomers. The 3,5-diaminobenzoate-terminated PO oligomer was prepared by reacting hydroxy-terminated PO with 3,5dinitrobenzoyl chloride in the presence of triethylamine. Triethylamine hydrochloride salt was filtered from the reaction mixture, and the oligomer was reduced using Pearlman’s catalyst, yielding the diamino oligomer. The degree of functionalization and molecular weight was determined by both 1H NMR and potentiometric titration of amine end groups with perchloric acid. Previous studies have shown that PMDA/4-BDAF is an insoluble polyimide that can be prepared via a twostage synthesis in which the poly(amic acid) formation is followed by thermal imidization.23,24 To prepare the proper polyimide/propylene oxide copolymers for use in this study, a poly(amic alkyl ester) precursor route was selected. The synthesis and isolation of the proper block

copolymers via the corresponding poly(amic alkyl ester)s permitted easy analysis and storage of these polymer intermediates. Unlike poly(amic acid) polyimide precursors, poly(amic alkyl ester)s have been shown to be hydrolytically stable and soluble in a variety of organic solvents,27 permitting samples to be isolated, analyzed, and stored for long periods of time. Additionally, there is evidence that the higher imidization temperatures of poly(amic alkyl ester)s give rise to linear chain extension and may influence the development of well-defined microphase-separated morphologies in block copolymers.28 Synthesis of poly(amic alkyl ester)s required the preparation of the proper diethyl pyromellitate diacyl chloride (DPDC) which can be isolated as either the meta or para isomers (Scheme 2). PMDA was allowed to react with an excess of ethanol to give an isomeric mixture of the diester diacids of PMDA (DADE) in high yield. This mixture can be separated into the two isomers by repeated recrystallization from ethyl acetate, with the para isomer recrystallizing first. The diethyl pyromellitate diacids were then reacted with oxalyl chloride to give the desired diacyl chlorides. The diacyl chlorides were sensitive to moisture and were stored under an inert atmosphere. The synthesis of the copolymers, outlined in Scheme 3, involved reacting pyromellitate diacyl chloride, 4-BDAF, and the amine-terminated PO oligomer, in NMP. Pyridine was also added to the mixture as a scavenger of the HCl evolved during the formation of the poly(amic alkyl ester). These reactions were carried out at a solids content of 10-15% (w/w) and the reaction mixture became a viscous, yellow, homogeneous solution after stirring for 24 h. The poly(amic alkyl ester) copolymers were precipitated into methanol/water 1:1 and washed repeatedly with methanol and water to (27) Volksen, W.; Yoon, D. Y.; Hedrick, J. L.; Hofer, D. Mater. Res. Soc. Symp. Proc. 1991, 227, 23. (28) Volksen, W.; Yoon, D. Y.; Hedrick, J. L. Proc. Electron. Compon. Technol. Conf. 1991, 41, 572.

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Scheme 2. Synthesis of Meta and Para Diethyl Pyromellitate Diacyl Chlorides (DPDC)

remove any unreacted PO oligomer. The copolymers were then filtered and vacuum-dried to constant weight. The synthesis of two different types of PO copolymers was attempted. As shown in Scheme 3, if a monofunctional amine-terminated PO block is reacted at the proper stoichiometry with DPDC and 4-BDAF an A-B-A type copolymer is obtained, with PO comprising the -A- blocks. The molecular weight of the PO oligomer remains unchanged during the formation of the copolymers. The precisely controlled stoichiometric imbalance of these condensation reactions provides for good control of the molecular weight in these block copolymers. Since monofunctional PO oligomer acts as a polyimide chain terminator, we focused our attention on commercially available PO of higher molecular weight (4000-8000). The use of higher molecular weight oligomers provided incorporation of a large amount of PO into the copolymers while concurrently giving a longer polyimide, -B-, segment. A graft copolymer approach was attempted to synthesize PO-containing copolymers similar to a previous study.20 In these experiments, DPDC, 4-BDAF, and 3,5diaminobenzoate-terminated PO oligomer were used. Unfortunately, only low molecular weight oligomeric products could be isolated from these reactions. Similar low molecular weight products enriched in PO content were also observed in the previous study. The A-B-A block copolymers, isolated as the poly(amic alkyl esters), could be stored, characterized (NMR, DSC, TGA, GPC), and redissolved and processed into thin films. The thin films of the poly(amic alkyl esters) were made by casting or spin coating them from NMP solutions (10-20% solids) onto appropriate substrates. Solvent can be removed by heating to 150 °C, but the resulting poly(amic alkyl esters) formed very brittle thin films. The films were cured to PMDA/4-BDAF-PO polyimide copolymers by heating, in steps, to 305 °C at 5 °C/min and holding at 305 °C for 1 h. Unlike the corresponding poly(amic alkyl esters), the PMDA/4BDAF-PO polyimides formed flexible, tough films. A variety of PMDA/4-BDAF homopolymers (1, 1a, 1b), as well as a series of PMDA/4-BDAF-PO polyimides (212) were synthesized. Both triblock and graft copoly-

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mers were synthesized and both isomeric prepolyimide poly(amic alkyl esters) (meta and para) were investigated. Table 2 summarizes the polymers prepared. The homopolymers 1, 1a, and 1b were prepared in order to examine any differences that might be present due to the use of different synthetic approaches to make PMDA/4-BDAF polyimide, namely, the poly(amic acid) route (1), and the poly(amic alkyl ester) routes, para (1a) and meta (1b). It was desirable to study copolymers containing 10-30% PO, which would most likely yield isolated phase-separated domains of PO in the polyimide matrix. Because the poly(amic alkyl esters) were soluble, 1H NMR was used to determine the amount of PO incorporated into each copolymer (Figure 2). By comparison of the area under the resonances due to the aromatic protons (8.1-7.0 ppm) to the area of the methyl resonance of the poly(propylene oxide) segment (1.0 ppm), the compositions of the copolymers could be determined. Cured films were also studied by TGA and the amount of PO was determined by measuring the weight loss upon thermolysis of the copolymers. There was very good agreement on the fraction of PO incorporated into the copolymer as measured by these two techniques. In most reactions, the amount of PO found was observed to be slightly less than the amount of PO charged in the reaction. This was to be expected since low molecular weight fractions, rich in PO, are lost due to fractionation during the polymer isolation procedure. The solubility of the poly(amic alkyl esters) enabled the use of gel permeation chromatography (GPC) to determine relative molecular weights. The weightaverage molecular weights of several of the polymer samples were measured and found to range between 17 000 and 54 000 g/mol relative to polystyrene standards. The minimum molecular weight needed for good film formation was found to be 20 000 g/mol. Samples with lower molecular weights gave films of poor mechanical properties when cast and cured. The two graft copolymer samples made with the diamino-terminated PO, 11 and 12, are seen to have GPC molecular weights of 19 000 and 17 000, respectively. These polymers did not form good free-standing polymer films. Thin films of PMDA/4-BDAF homopolymer and copolymers were made by either casting or spin coating solutions of the poly(amic acid) or poly(amic alkyl esters). Films ranging in thickness from 1 to 40 µm could be made in this manner. After the initial coating of the substrate, the samples were subjected to a ramped heating cycle to effect solvent removal, imidization, and thermal curing. Typically, the sample was heated from room temperature to 80 °C at 5 °C/min, then 130 °C, then 250 °C, and finally 305 °C. The sample was held at each of the temperatures indicated for 1 h. After curing and full imidization, the PMDA/4-BDAF homopolymer and copolymers formed tough, clear yellow films that were solvent resistant and insoluble in common organic solvents. This development of solvent resistance is thought to be a ramification of the ordering and crystallization of the polymer above 300 °C. A crystallization exotherm can be clearly seen in DSC thermograms of the polymer, with the exotherm peak at 320 °C. The Tg of the homopolymer is observed at 305 °C. Neither the crystallization exotherm nor the

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Chem. Mater., Vol. 9, No. 1, 1997 111 Scheme 3. Synthesis of PMDA/4-BDAF-co-PO Triblock Copolymers

Table 2 no.

DPDC isomer

1 1a 1b 2 3 4 5 6 7 8 9 10 11 12

PMDA para meta meta meta meta para para para para para para para para

block type

triblock triblock triblock triblock triblock triblock triblock triblock triblock graft graft

PO block Mw

% PO in reaction

4000 4000 4000 5600 6500 6500 13300 5600 5600 7900 7900

0 0 0 10 18 22 15 17.9 17.9 22 24.4 28.9 18 25.5

Tg are observed in annealed samples, which is consistent with other semicrystalline polymers.23,24 X-ray diffraction studies showed that crystallization occurred in all samples, including the homopolymers and copolymers derived from poly(amic acids) and poly(amic alkyl esters). The cured copolymers were foamed by heating the films to 250 °C in air for 6-10 h. The PO underwent oxidative thermolysis to yield low molecular weight byproducts, which quantitatively diffused from the polymer leaving voids in the polymer matrix. Thermogravimetric analysis (TGA) was performed to monitor the thermolysis of PO from copolymer samples. No detectable residual PO or PO byproducts were observed as examined by TGA or IR. A typical PMDA/4-BDAFco-PO TGA thermogram is shown in Figure 3. The

% PO (NMR)

% PO (TGA)

poly(amic alkyl ester) Mw 54,000

9.6 17.9 21 9.7 7 5.3 18.9 18.7 28 2.9 18.47

7.3 17.2 21 6 7.4 19.25 25.3

20 000 46 000 33 000 29 000 44 000 33 000 19 000 17 000

onset of PO degradation is about 300 °C when thermolysis is attempted in an inert atmosphere. The thermolysis is monomodal with complete PO removal from the sample occurring by 450 °C. The polyimide matrix begins to degrade at 500 °C. When thermolysis was performed in air, the onset of PO degradation was observed at 220 °C and was bimodal, suggesting two competing thermolysis mechanisms, indicating that there is significant oxidative degradation at this lower temperature followed by thermally induced degradation at 300 °C. Isothermal TGA (Figure 4) was performed to establish the best conditions for complete thermolysis and removal of PO. The TGA experiments helped establish that heating the samples at 250 °C for >6 h was sufficient for quantitative removal of PO from the samples.

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Figure 2.

1H

Carter et al.

NMR of imide/PO copolymer 9.

Figure 3. Dynamic thermogravimetric analyses of imide/PO copolymer 9 in air and high-purity N2.

Porosity and void formation were determined from density, refractive index, and small-angle X-ray scattering measurements. The characteristics of several PMDA/4-BDAF polyimide foams are shown in Table 3. Foams were prepared from a variety of PO oligomers and at levels of porosity up to 24%. In most of the samples, the PO volume fraction in the copolymers before foaming produced a similar volume fraction of pores after foaming. One notable exception was sample 5. Sample 5 had a PO content of 36%, as measured by NMR prior to film preparation. After foaming, sample 5 had a measured porosity of only 14%. The reasons for this very low porosity/loading ratio was most probably due to the copolymer morphology. The morphology

of the foam is expected to retain the initial morphology present in the copolymer. When the initial copolymer morphology consists of isolated domains of PO in a polyimide matrix, one expects voids of similar size and shape in the foamed polyimide. Sample 5 had a PO content of 36%, which would force the morphology to be essentially bicontinuous. Upon foam formation the matrix network is too weak to withstand the stress generated during foam formation, resulting in a collapse of the foam structure in this sample. The storage modulus and tan δ of the polymers were determined as a function of temperature by DMTA. In Figure 5, one can see that the homopolymer has a modulus of about 500 MPa with Tg values that depend

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Chem. Mater., Vol. 9, No. 1, 1997 113

Figure 4. Isothermal thermogravimetric analysis of imide/PO copolymer 9 in air at 250 °C. Table 3. PMDA/4-BDAF Foam Density Measurements no.

polymer

1 1a 2f 5f 6f 7f 8f 9f 10f 11f 5f

PMDA/4-BDAF from amic acid PMDA/4-BDAF from p-amic ester PMDA/4-BDAF 4K PO, meta PMDA/4-BDAF 5.6K PO, para PMDA/4-BDAF 6.5K PO, para PMDA/4-BDAF 6.5K PO, para PMDA/4-BDAF 13.3K PO, para PMDA/4-BDAF 5.6K PO, para PMDA/4-BDAF 5.6K PO, para PMDA/4-BDAF 7.9K PO, graft PMDA/4-BDAF 7.9K PO, graft

PO vol %

cure temp (°C)

13 13 10 7 25 25 36 4 4

300 300 300 300 300 300 300 300 300 300 300

upon the polyimide precursor. PMDA/4-BDAF polyimide derived from the poly(amic acid) precursor has the highest Tg at 325 °C followed by the polyimide derived from the p-poly(amic alkyl ester) at 305 °C and finally the polyimide derived from the m-poly(amic alkyl ester) at 275 °C. While the chemical structure of the cured polyimide is independent of the isomeric ratio of the alkyl-ester precursor (meta or para), there are significant differences in the Tg values observed for polymers depending upon this precursor. These differences may arise from differences in the chain packing of the parent poly(amic ester)s, which are not affected by the intramolecular imidization process. This argument also holds for the different Tg values found for the polyimide prepared from the poly(alkyl ester) and poly(amic acid). Intermolecular association of the amic acid would be higher and, hence, give rise to a more densely packed material. The DMTA traces of a PMDA/4-BDAF copolymer sample before and after foaming are shown in Figure 6. Indication of imide/PO microphase separation is evident by the relaxation at -80 °C, the Tg of PO. The Tg of the matrix is, however, approximately 30 °C lower than the Tg observed for the polyimide homopolymer. The difference in the Tg values may arise from the finite solubility of the PO block in the matrix or residual amounts of NMP present after curing. The

foaming temp (°C)

density (g/cm3)

porosity (%)

250 250 250 250 250 250 250 250 250

1.47 1.47 1.37 1.3 1.31 1.32 1.21 1.13 1.26 1.36 1.3

7 12 11 10 18 24 14 8 12

overall modulus of the copolymer is lower than that of the corresponding polyimide, again most likely due to residual solvent and partial plasticization by the PO phase. After foaming, the PO relaxation is no longer present and a modest modulus increase is observed. The Tg of the system is also observed to increase to 300 °C as measured by the increase in tan δ. In a comparison of the DMTA of the foamed polymer sample and the homopolymer sample (Figure 7), one can see that the initial modulus of the two samples are nearly identical up to 250 °C, at which point the foamed sample experiences a drop in modulus to about 150 MPa, clearly becoming an elastomeric-like material retaining mechanical integrity. While the homopolymer experiences a dramatic loss of modulus above the Tg (>300 °C), the loss in modulus of the foamed polyimide is not as dramatic above 300 °C. The increased modulus in the foamed sample suggests that the porous polyimide structure may have enough integrity to prevent complete collapse even though the sample has been heated above its Tg. Previous studies of other polyimide systems have shown that collapse begins as the polymer is heated to its Tg and is complete at temperatures above the Tg.17-22 Foam thickness as a function of temperature was examined (Figure 8). A 23 µm thick film of a 25% PO

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Figure 5. Dynamic thermal mechanical analyses of PMDA/4-BDAF homopolymers.

Figure 6. Dynamic thermal mechanical analyses of imide/PO copolymer 9 before and after foaming.

copolymer was cast onto a microscope slide and heated to 305 °C under argon. Samples of the film were then heated to various temperatures between 250 and 400 °C for 2 h at each temperature, and then film thickness was measured. The first observation was that a 5% thickness change was observed after heating the sample to 250 °C because of loss of residual solvent and ambient water. The foamed structure formed at 250 °C is fairly stable up to the Tg of the polyimide. By 300 °C a 6% reduction in film thickness and, consequently, porosity

occurs. By 400 °C the film thickness has been reduced by 25% corresponding to a complete collapse of the foam. Typical small-angle X-ray scattering, SAXS, data for a PMDA/4-BDAF copolymer containing 20% PO that has been annealed at 300 °C under helium and then heated in air at a rate of 10 °C/min from 40 to 400 °C are shown in Figure 9. The SAXS profile for the imidized, but not foamed, copolymer decreases monotonically as a function of the scattering vector. This result is typical for a disordered, microphase separated

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Chem. Mater., Vol. 9, No. 1, 1997 115

Figure 7. Dynamic thermal mechanical analyses of PMDA/4-BDAF homopolymer and foam derived from imide/PO copolymer 9.

Figure 8. Foam thickness as a function of thermal aging.

material where there are no long-range correlations in the spatial distributions of the PO domains on the polyimide matrix. For such a system, the scattered intensity, I(Q), can be described in a Debye-Bueche manner29 as

I(Q) ) Ca/(1 + a2Q2)2 where C is a constant and a is a correlation length describing the average size of heterogeneities in the system. Consequently, a plot of [I(Q)]-1/2 as a function of Q2 will yield the correlation length when the square root of the slope to intercept ratio is taken. If the volume fractions of the phases are known, then the average size of the domains of each component, li, is given as a/φj, where φj is the volume fraction of the other phase. From such an analysis, a correlation length of 80 Å is obtained that translates into an average PO microdomain size of ∼100 Å. Increasing the temperature results in a slight decrease in the overall scattering up to ∼190 °C. As the (29) Debye, P.; Bueche, A. M. J. Appl. Phys. 1940, 20, 518.

Figure 9. Small-angle X-ray scattering SAXS results for a PMDA/4-BDAF foam.

temperature is increased further, a dramatic increase in the scattering is seen up to ∼280 °C, whereupon the scattering remains relatively constant up to ∼300 °C and then a dramatic drop in the intensity is seen. It is important to note that the angular dependence of the scattering up to ∼280 °C does not change significantly with temperature. This result indicates that the size and spatial distribution of the microdomains, be these voids or PO microdomains, does not change. Rather it is only the contrast between the microdomains that changes, i.e., the electron density difference between the matrix and dispersed microdomains, ∆F, changes. This difference increases rapidly with increasing temperature up to ∼280 °C. At higher temperatures, both ∆F and the correlation length change. As the total amount of scattering decreases, the dependence of the scattering on Q becomes weaker. Consequently, the correlation length is decreasing along with a decrease in the contrast or the number of scatterers. This behavior is precisely as expected for a nanofoam undergoing collapse. As the size of the dispersed phase decreases, the

116


Figure 10. Small-angle X-ray scattering intensity as a function of temperature.

correlation length decreases. If one has a true nanofoam, then the contrast between the phases will be independent of the phase size, i.e., small voids have the same electron density as large voids, i.e., zero. If a foam collapses, then the total number of voids dispersed in the matrix will decrease and give rise to a reduction in the scattering. If the collapse of the foam proceeded by eliminating the smaller pores first, then the falloff in the scattering would become much sharper as the intensity dropped. This is not the case. In fact, the opposite occurs. The angular dependence of the scattering is weaker, indicating that the average size distribution of the pores becomes smaller. This means that the larger sized pores must get smaller as the collapse occurs. Thus, the collapse of the foam occurs by a continual shift of the average pore size to smaller and smaller sizes with the smallest pores vanishing due to the high internal pressure required to maintain them. One can also examine the total integrated scattering or invariant, Qmv:

∫0∞ I(Q)Q2 dQ

Qmv ) φ1φ2(∆F)2 ) k

The constant, k, allows one to place the data on the absolute level. As indicated, the invariant is related to the mean-square electron density difference, (∆F)2 ) (F1 - F2)2, and the volume fraction of the phases. The integration of the intensity extends to infinity, which requires the use of Porod’s law where, at large angles, I(Q) ∝ 1/Q4, for a system with infinitely sharp phase boundaries. If the phase boundaries are diffuse, then the interfacial width must be determined and the data must be corrected to perform the extrapolation. Over the entire temperature range, the interfacial width between either the polyimide matrix and the PO or the polyimide and pore was small, ∼10 Å. The interfacial width did not vary in any systematic manner with temperature. The sharp interface also indicates that the PO and the polyimide are incompatible, otherwise a broader interface would be observed. After correction of the data for the phase boundary, the invariant was determined and is shown in Figure 10. Here one sees a slight decrease in the invariant up to

Carter et al.

∼180 °C, whereupon it increases quite sharply, levels off, and then begins to decrease rapidly with increasing temperature. The rapid increase in the invariant is associated with the formation of the foam, i.e., a marked increase in (∆F)2, while the decrease arises from a decrease in the number of pores, causing the product of φ1φ2 to decrease. These results quite nicely parallel the film thickness studies discussed previously. Demonstration of successful void formation can also be studied by transmission electron microscopy (TEM). Figure 11 is a TEM micrograph of polyimide foam derived from a PMDA/4-BDAF-co-PO copolymer, 4. The pores are the small light regions in the micrograph, with the darker areas corresponding to the polyimide matrix. It can be difficult to see the foamed regions in the micrographs because there is not a great deal of electron contrast between the small voids and the polyimide matrix which was microtomed to a thickness of ∼500 Å. The voids are randomly distributed and though they are not strictly spherical, there is no evidence of any significant pore interconnectivity. Interconnectivity is difficult to assess because the micrograph represents a projection of the morphology in the thin section onto one plane. The average pore diameter is 100 Å, which corresponds well to the results from the X-ray scattering experiments. Stress measurements, using a wafer curvature method, were first made on a 6 µm thick film of the homopolymer which was cured to 250 °C. In situ heating of the sample to 400 °C and holding for an hour resulted in a final film stress of 49 MPa. A copolymer with 20% PO which was cured to 300 °C in argon and foamed at 250 °C in air had a stress value of 28 MPa. It is believed that the pores are acting as stress concentrators, fortuitously lowering the overall stress of the system. Figure 12 shows stress as a function of temperature for the foamed polymer as it is heated above the cure temperature to 350 °C, held for 1 h, and cooled. Once above the cure temperature, mobility occurs reducing the stress of the film and resulting in collapse of the foam. By the time the film has cooled to room temperature, the stress has now increased to the value of the homopolymer, indicating complete collapse of the foam. To obtain a measure of the optical anisotropy of the thin films, their in-plane and out-of-plane refractive indexes were measured using a commercially available Metricon prism coupler. In addition to obtaining an accurate measurement of the difference between the inplane and out-of-plane refractive indexes, an estimate of the respective in-plane and out-of-plane dielectric constants can be obtained by applying the Maxwell equation ≈ (n)2 and correcting for frequency dispersions (∆ = +0.3).30 This type of analysis is very useful for screening series of polyimide compositions, obviating the need for complex capacitance dielectric measurements on samples especially when making comparative measurements on samples of similar chemical composition. We observed a drop in refractive index of samples upon foaming. PMDA/4-BDAF homopolymer has an inplane refractive index, nTE, of 1.63 and an out-of-plane refractive index, nTM, of 1.56, showing that this poly(30) Boese, D.; Lee, H.; Yoon, D. Y.; Swalen, J. D.; Rabolt, J. F. J. Polym. Sci., Part B: Polym. Phys. 1992, 30, 1321.

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Figure 11. TEM micrograph of a PMDA/4-BDAF foam.

Figure 12. Thin-film stress of a PMDA/4-BDAF foam as a function of temperature.

imide is anisotropic (∆n ) 0.07). A foamed sample of PMDA/4-BDAF polyimide (derived from a 23% PO copolymer) had a measured in-plane refractive index nTE of 1.51 and an out-of-plane refractive index nTM of 1.46 (∆n ) 0.05), a small, but noticeable drop in the films anisotropy. Thus, utilizing the modified Maxwell equation, we would expect the foamed polyimide to have an in-plane dielectric constant of approximately 2.4 relative to the homopolyimide dielectric constant of 2.7. For comparison, the reported measured in-plane dielectric constant of PMDA/4-BDAF polyimide was 2.8.23 If these foamed materials were to be considered as candidates

for integration into actual devices as thin-film insulators, an in-depth study of their dielectric properties would have to be performed. The porosity of the foamed samples can also be studied indirectly by examining the refractive index of thin-film samples.31 The experimentally measured drop in refractive index for the foamed polymer corresponds to what one would expect for a sample with 18-19% porosity. Therefore the foaming efficiency, as measured (31) Sanchez, M. I.; Hedrick, J. L.; Russell, T. P. J. Polym. Sci., Part B: Polym. Phys. 1995, 33, 253.

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by this refractive index technique, is similar to those measured by TGA and density measurements. Conclusions The ability to prepare thin-film foams of PMDA/4BDAF, with pores in the nanometer size range has been demonstrated. A drop in the refractive index of the foamed polymers was observed which is expected to correspond to a proportional lowering of the dielectric constant. The premise for utilizing PMDA/4-BDAF was to examine whether polyimides with a significant degree of structural order could serve as matrixes for nanofoams and whether the semicrystalline order in these materials would act to stabilize the foamed structure. In previous attempts to use highly ordered polyimides such as PMDA/ODA as foam matrixes, foamed structures were foamed, but they were observed to collapse as the polymers were heated to temperatures well below the softening point of the polyimide matrix (softening point > 450 °C). This collapse was found to be caused by stresses induced in the polymer film due to ordering of the polyimide at high temperatures. In the case of PMDA/4-BDAF, we were able to cure the copolymers at the Tg of the polyimide, 305 °C under argon, and then

Carter et al.

the copolymers were foamed at 250 °C in air. The initial curing at 305 °C allowed for solvent removal, imidization, and the development of order in the polyimide matrix as evidenced by DSC and WAXD measurements. These nanofoam films could be subjected to further thermal treatments approaching the polyimide Tg, with a negligible amount of foam collapse. When heated above the matrix Tg, foam collapse became more prominent, though complete foam collapse was not achieved until heating to >350 °C, nearly 50 °C above the matrix Tg. Therefore, when compared to foams derived from other ordered polyimides (i.e., PMDA/ODA), the PMDA/ 4-BDAF system afforded a small degree of foam stabilization above the polymer’s Tg. The stabilization of the porous matrix was not sufficient to prevent complete collapse of the foamed structure when samples were heated to greater than 50 °C above the polyimide Tg for extended times. Acknowledgment. The authors would like to acknowledge that this work was funded, in part, by NIST (ATP Contract 70NAAB-3H1-365). CM960242R

Polyimide Nanofoams Based on Ordered Polyimides Derived from

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