Accepted Manuscript

Improved space survivability of polyhedral oligomeric silsesquioxane (POSS)

polyimides fabricated via novel POSS–diamine

Xing–Feng Lei, Ming–Tao Qiao, Li–Dong Tian, Pan Yao, Yong Ma, He–Peng

Zhang and Qiu–Yu Zhang

PII: S0010-938X(14)00470-3

DOI: http://dx.doi.org/10.1016/j.corsci.2014.10.013

Reference: CS 6043

To appear in: Corrosion Science

Received Date: 11 August 2014

Accepted Date: 13 October 2014

Please cite this article as: X. Lei, M. Qiao, L. Tian, P. Yao, Y. Ma, P.Z.a. Zhang, Improved space survivability of

polyhedral oligomeric silsesquioxane (POSS) polyimides fabricated via novel POSS–diamine, Corrosion Science

(2014), doi: http://dx.doi.org/10.1016/j.corsci.2014.10.013

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Improved space survivability of polyhedral oligomeric silsesquioxane (POSS) polyimides fabricated via novel POSS–diamine†

Xing–Feng Lei, Ming–Tao Qiao, Li–Dong Tian, Pan Yao, Yong Ma, He–Peng

Zhang and Qiu–Yu Zhang�

Department of Applied Chemistry, Key Laboratory of Space Applied Physics

and Chemistry of Ministry of Education, School of Science, Northwestern

Polytechnical University, Youyi Road 127#, Xi’an, 710072, P. R. China.

Abstract: Polyhedral oligomeric silsesquioxane (POSS) surrounded by two amine

groups, namely POSS–diamine, was prepared via hydrolytic co–condensation. POSS

polyimide hybrid membranes were subsequently fabricated by co–polymerizing

POSS–diamine with imide monomers. Hybrid membranes exhibit significantly

improved space survivability. A hybrid membrane with 29.7 wt % POSS loading

shows the best atomic oxygen (AO) resistance with the lowest erosion yield of 0.9 ×

10−25 cm3·atom−1. The enhancement in AO resistance is attributed to the formation of

a SiO2 passivating layer on the membrane surface upon AO exposure. POSS

polyimides with desirable AO survivability may find wide usage in aerospace.

Keywords: A. Polymer; B. XPS; B. Weight loss; C. Interfaces; C. Oxidation

1. Introduction

Materials, not only metals but also polymer composites, tend to corrode or

degrade in almost every atmospheric condition [1–8]. To speak of, a large portion of

high–performance polymeric materials, in particular aromatic polyimides, are

extensively used in man–made satellites. However, those hydrocarbon–based

polyimides are subject to severe degradation when exposed to atomic oxygen (AO),

ultraviolet (UV) and vacuum–ultraviolet (VUV) radiation and thermal cycles, which

are present in the low earth orbit (LEO) [9–10]. Among which, AO is ubiquitous in

LEO and strong enough to induce organic bond cleavage due to its strong oxidization

and considerable translational energy of approximately 4.5~5.0 eV [11–13].

Consequently, most of the organic polymers onboard spacecraft are oxidized and

eroded and then generate free radicals and give rise to reactions that finally result in

chain scission and/or cross–linking or even releasing volatile gases, which can cause

significant reduction in physical, mechanical and optical properties of polymers

[9,10].

A variety of promising methods focusing on physically and chemically tailoring

the molecular structure of polyimides to improve oxidation resistance of polyimides,

including polymer blending [14,15], copolymerization [16–18], organic–inorganic

hybrid technology [9,19] and coating technology [20], have been attempted during

past several years. The synthesis of fluorinated polyimides has been one such

approach [21–23], for fluorinated polyimides often exhibit good resistance to AO

attack. However, remarkable increases in the erosion yield (mass loss) is typically

observed when such polyimides are exposed simultaneously to AO and UV radiation

[21,22]. In addition, poor mechanical properties and high monomer costs also restrict

its wide use.

A second approach is introducing phosphorus into polyimide molecular

backbones to fabricate phosphorus–containing polyimides by copolycondensation

[16–18,24–26]. A substantial reduction in the AO erosion yield has been obtained,

and the resulting polyimides demonstrated a combination of admirable properties. In

this approach, however, the resulting membranes were relatively brittle and exhibited

poor fracture toughness [18], which cannot meet the space requirements, either.

Strategies to making polyimides with high space survivability that access AO

erosion yield below 1% that of pristine polyimide typically introduce silicon to

improve the AO resistance [9,19,27,28]. Among a few examples of silicon–containing

polyimides, those derived from polyhedral oligomeric silsesquioxane (POSS) are the

most thoroughly investigated representatives of this new class of materials [9,19],

owing to its potential applications in various areas and the ease with several functional

groups can be attached on POSS vertex [9,29–37]. As reported by Gonzalez [38] and

Minton [9,19], POSS polyimides are likely to form a SiO2 passivating layer upon AO

exposure, and thus prevents AO from eroding the bulk matrix, thereby increasing AO

resistance. It is another noticeable problem that the previously reported work typically

involving POSS monomer looks specifically at bifunctional POSS–diamines, which

are relatively expensive and difficult to synthesize [9,19,39]. For instance, in the

approach of Minton’s group [9,19], such expensive chemical reagents as

octa–cyclopentyl POSS, (aminopropyl)–hepta–isobutyl POSS,

N–p–Lithiophenyl–N–1,1,4,4,–tetramethyldisilylazacyclopentane, n–BuLi et al. were

typically employed and the whole preparation process was relatively complicated and

difficult to operate, which is adverse to mass production and hence compromise wide

use of POSS polyimides and should be rationally solved for practical application. In

contrast, octa–functional POSS, such as octa–(aminopropyl or

aminophenyl)silsesquioxane (OAPS), is relatively easy to manufacture and not that

expensive in comparison to POSS–diamine [29,33,40], but it is difficult to get all of

the amine groups reacted and hence one could not make polymers with well–defined

structures. Additionally, insoluble polyamic acid (PAA) gels usually form [41], and

thus hybrid membranes could not be obtained or only membrane fragments were

acquired due to extreme brittleness [42].

We wonder whether, in another approach, POSS–diamine can be synthesized and

incorporated into polyimide molecular backbones to fabricate POSS polyimides

inherently withstanding the “ambitious” AO attack through a versatile method without

significantly sacrificing the other properties. Previously, we have reported that

covalently tethering amine–functionalized hyperbranched polysiloxanes (HBPSis)

into polyimide molecular skeletons results in organic–inorganic HBPSi polyimides

with good thermal stability and outstanding AO resistance [43]. Here, we adopt

phenyltriethoxysilane and γ–aminopropyltriethoxysilane to directly synthesize newly

developed POSS–diamine via moderate hydrolytic co–condensation reactions.

Finally, POSS polyimides could be easily obtained by copolymerizing POSS–diamine

with other imide monomers in a polar aprotic solvent and subsequent thermal

imidization. This copolymerization approach provides a facile method to design a

variety of polyimide nanocomposites with tunable AO resistance and mechanical

properties by varying POSS addition. Polyimide matrix maintains the advantage of

both high thermal and mechanical strength while POSS molecule imparts excellent

AO resistance to the resulting nanocomposites.

In the present study, the AO resistance of POSS polyimides was investigated by

exposing its surface to various AO fluences, and the protection/oxidation mechanism

was established. Several surface analytical techniques were herein adopted to reveal

the relationship between the erosion yields of resulting POSS polyimides and their

structures. Curve–fitting models were also used to predict the service life of POSS

polyimide in the simulated AO environment. These evaluations may contribute to

providing some effective viewpoints to the degradation behaviour and protection

mechanism of POSS polyimides in AO environment.

2. Experimental

2.1 Materials. Phenyltriethoxysilane (PTES) and γ–aminopropyltriethoxysilane

(APS) were purchased from Nanjing Chengong silicon co., Ltd (Nanjing, China) and

used as received. Tetraethylammonium hydroxide (TEAOH, 25 wt % in aqueous

solution), absolute ethanol, 4,4'–diaminodiphenyl ether (ODA) and pyromellitic

dianhydride (PMDA) were provided by Sinopharm Chemical Reagent Co., Ltd

(Beijing, China). ODA and PMDA were purified by vacuum sublimation prior to use

while the other reagents were used as received. N,N'–Dimethylacetamide (DMAc)

was purchased from Tianjin Fuyu Fine Chemicals Co. Ltd (Tianjin, China) and

freshly distilled under reduced pressure over phosphorus pentoxide and stored over 4

Å molecular sieves prior to use.

2.2 General Characterization. 1H and 13C nuclear magnetic resonance (NMR)

spectra of POSS–diamine were recorded on a BRUKER AVANCE–300 spectrometer

in DMSO–d6 at room temperature. Chemical shifts were referenced to

tetramethylsilane. 29Si–NMR spectrum of POSS–diamine was acquired on a Bruker

Avance 500 spectrometer (Bruker BioSpin, Switzerland) operating at 50.7 MHz in

DMSO–d6. Elemental analysis of POSS–diamine was run in a Heraeus VarioEL–III

CHN elemental analyzer (Germany). The molecular weight of POSS–diamine was

measured with gel permeation chromatography (GPC) by using tetrahydrofuran (THF)

as eluent solvent. Polystyrene standards were used for calibration of the GPC. Amino

content was acquired by acid–base back titration using phenolphthalein as indicator.

The purity (wt %) of POSS–diamine was estimated based on its amino content.

Fourier transform infrared spectroscopy (FT–IR) measurements were conducted on a

FT–IR spectrophotometer (BRUKER TENSOR 27). The FT–IR spectrum of

POSS–diamine was obtained by using thin KBr disk as the sample holder, while

FT–IR spectra of polyimide films were collected by using an attenuated total

reflectance (ATR) instrument. Thermal gravimetric analyses (TGA) of POSS–diamine

and polyimide membranes were carried out on a TGA/DSC 1 synchronous thermal

analyser at a heating rate of 10 °C/min under air and/or nitrogen atmosphere from

room temperature to 1000 °C. The isothermal TGA of polyimide membranes were

conducted under nitrogen atmosphere aging at two different temperatures (400 °C and

500 °C) for 10 h. The d–spacing within POSS–diamine molecule and resulting

membranes was determined by wide angle X–ray diffraction (WAXD) by using a

Shimadzu XRD–700 (Shimadzu, Japan) X–ray diffractometer (40 kV, 40 mA) with a

copper target at a scanning rate of 4 °/min and calculated from the scattering angle

(2θ) according to Bragg's equation:

2 sinn d� �� (1)

The storage modulus E’ and internal loss factors tan δ of polyimide membranes were

determined on a Mettler Toledo DMA/SDTA861e instrument. The cast membranes

were cut into 6 × 30 mm rectangular samples for dynamic mechanical thermal

analysis (DMA). All polyimide samples were subjected to the temperature scan mode

at a programmed heating rate of 10 °C /min at a frequency of 1 Hz from room

temperature to 450 °C in a tensile mode. The glass transition temperatures (Tgs) of

resulting membranes are regarded as the peak temperature in the tan δ curves.

Thermo–mechanical analysis (TMA) was performed on a TMA/SDTA 840 device at a

heating rate of 5 °C/min. The coefficients of thermal expansion (CTE) of the

membranes were measured in the temperature range of 75–150 °C. Mechanical

properties of polyimide membranes were measured with a universal testing machine

(UTM) according to GB13022–91 at a drawing rate of 10 mm/min at ambient

temperature on strips approximately 40–45 μm thick and 15 mm wide with a 60 mm

gauge length. An average of at least seven individual determinations was used.

UV–vis transmittance and absorbance of the resulting membranes was measured by

using a U–3010UV–VIS spectrophotometer and the wavelength ranged from 200 to

800 nm. Membrane thickness was measured with a thickness gauge, and densities

were obtained by liquid suspension method with a mixture of toluene and

tetrachloromethane at 25 °C. Inherent viscosities (�inh) were obtained on 0.5% (w/v)

polyamic acid (PAA) solutions in freshly distilled N,N'–dimethylacetamide (DMAc)

at 25 °C with an Ubbelohde viscometer. Presented �inh values are the average of at

least nine individual tests of each sample.

2.3 AO Exposure Testing. Ground–based AO exposure measurements were

performed with a combined space effect testing facility (CSETF) equipped with

neutral AO beam and vacuum ultraviolet ray (VUV) sources. The specific operation

procedures and details of the system were described in our previous studies [43,44].

The AO flux at the sample’s position was finally calculated to be approximately

4.89×1015 atom∙cm-2∙s-1 from mass loss of Kapton® H. The exposure area is a circular

domain with a diameter of 30 mm and the thickness of samples ranges from 40 to 45

μm. During AO exposure, all samples were handled in vacuum chamber and

irradiated to various AO fluences ranging from 0.88 to 3.87 × 1020 O atoms·cm-2. The

AO fluence was controlled by exposure period and monitored by using a reference

sample, Kapton® H, which was installed on the sample holder and exposed to AO

under identical conditions. To eliminate the measurement errors, three or more

individual tests were carried out for each sample under the same exposure conditions.

2.4 Surface Characterization. Samples of Kapton® H and POSS/polyimide

nanocomposites were exposed to a variety of AO fluences. In order to minimize the

influence of air exposure on the surface compositions, surface analyses, including

surface topography (SEM, AFM) and surface chemistry analysis (XPS) were

immediately carried out with caution once the AO–exposed samples were removed

from the vacuum chamber. Surface morphologies of polyimide membranes before and

after AO exposure were observed with Scanning Electron Microscope (SEM,

MERLIN ZEISS, Germany). TappingModeTM atomic force microscope (AFM)

images and root–mean–square surface roughness (RMS) of the polyimide membranes

before and after AO exposure were collected using an SPI3800–SPA–400 (Japan,

NSK Ltd.) scanning force microscope on silicon wafer under ambient conditions with

the scanning rate ranging from 0.5 to 1.0 Hz. An X–Ray Photoelectron Spectroscopy

(XPS) instrument (AXIS Ultra DLD, Kratos Co., UK) equipped with a

monochromatic Al Kα X–Ray source with a residual pressure of ca. 10-8 Pa was

utilized to detect the elemental components and valence variations at the membrane

surface before and after AO exposure. The pass energy during the whole measurement

was 40 eV in all cases. The shifts of binding energy of XPS curves were calibrated by

assuming that the lowest C 1s component was 284.6 eV for the unexposed polyimide

(0 wt % POSS) sample. The high resolution XPS spectra of C 1s, Si 2p and O 1s were

curve–fitted according to XPS standard spectra databases and theoretical

compositions.

2.5 Synthesis of Amine–Functionalized Polyhedral Oligomeric

Silsesquioxane (POSS–diamine). An absolute ethanol (170 mL) solution containing

tetraethylammonium hydroxide (2 mL, 25 wt % in aqueous solution) and H2O (36.11

mmol, 6.5 g) was maintained at ambient temperature in a water bath and then, the

mixture of phenyltriethoxysilane (45 mmol, 10.82 g) and

γ–aminopropyltriethoxysilane (15 mmol, 3.32 g) was added dropwise via a dropping

funnel under vigorous stirring over a period of 2 h. The mixture was allowed to react

with stirring at ambient conditions for 10 h and afterward reflux at 50 °C for an

additional 50 h. After cooling to room temperature, the crude products were obtained

by casting the reaction solution onto a large dust–free glass plate followed by

evaporation most of the solvents under air circulation at ambient temperature. Collect

the crude products and wash it with deionized water (3 × 100 mL), then treated with

cold methanol/tetrahydrofuran mixture (v/v = 1/10, 2 × 50 mL). Subsequently, the

crude products were dissolved in N,N'–dimethylacetamide (50 mL) and filtered

through a filter with a diameter of 0.22 μm and then precipitated out by adding

deionized water (200 mL). The typical dissolution–precipitation process was repeated

at least three times and finally the resultant products were collected by suction

filtration. The water was removed by a typical freeze–drying procedure to give white

powder products in ~37% yield (2.72 g, POSS–diamine with ~96.2 wt % in purity).

The synthetic procedure of POSS–diamine macromonomer is presented in scheme 1.

29Si–NMR of POSS–diamine (DMSO–d6, � ppm): –64.2, –77.8. Elemental analysis

calculated for C42H46N2Si8O12 (POSS–diamine, M = 995.72 g/mol): calculated C,

50.69%, H, 4.65%, N, 2.81%; found C, 50.39%, H, 4.66%, N, 2.68%. GPC data: Mw =

988.6 g/mol, Mn = 896.8 g/mol, polydispersity 1.10. Amino content for

POSS–diamine: theoretically 2.0086 mmol/g, acid–base titration 1.9698 mmol/g.

2.6 General Procedure for the Preparation of POSS–containing Polyimide

Membranes (POSS–PIs). POSS–PIs were produced in the form of membranes

(40–45 μm thick). The schematic production process is depicted in scheme 2.

Polyamic acids from 4,4'–diaminodiphenyl ether and pyromellitic dianhydride with

different POSS addition, varied between 4.1 wt % and 29.7 wt %, were prepared

according to table 1 as 12 wt % solid content of polyamic acid in freshly distilled

N,N'–dimethylacetamide in all cases. A representative procedure for 8.8 wt % POSS

polyamic acid precursor is as follows: Specifically, POSS–diamine (0.395 g) and

4,4'–diaminodiphenyl ether (9.53 mmol, 1.908 g) were added into a 100 mL,

absolutely dry, three–necked flask containing 35 mL of freshly distilled

N,N'–dimethylacetamide with stirring. After 4,4'–diaminodiphenyl ether and

POSS–diamine were completely dissolved, pyromellitic dianhydride (10 mmol, 2.181

g) was quickly added in one portion with stirring and reacted at room temperature

under nitrogen atmosphere for 24 h to afford a viscous polyamic acid solution.

POSS–containing polyimide membranes were produced at a bench–scale process by

casting corresponding polyamic acid solutions onto dust–free glass plates by using an

adjustable doctor blade. Subsequently, the samples were placed in a

temperature–programmable oven and heated to 360 °C in air at a heating rate of 4

°C/min and held at this temperature for a period of 60 min to yield fully imidized

hybrid membranes. To eliminate the residue stress within polyimide membranes, the

final stage was allowed to cool down to room temperature at a rate of 2 °C/min.

Polyimide membranes were finally obtained by soaking glass plates into deionized

water and then peeled off from glass plates and afterward dried in a vacuum oven at

120 °C for 4 h.

3. Results and discussion

3.1 Hydrolytic Co–condensation Reactions of Siloxanes. The synthesis of

amine–functionalized POSS is usually conducted by the method described by Wei

[30], Minton [9] and Carniato [45,46], which is a standard chemical synthesis

approach based on molecular structure design. This method generally involves several

kinds of expensive reagents as well as multiple purification procedures. The whole

operating process is complicated [9,30], although the resulting products possess

well–defined molecular structure. According to Voronkov [47], hydrolytic

co–condensation of the mixture of trifunctional silane monomers having comparable

reactivity usually gives hetero–substituted fully condensed polyhedral oligomeric

silsesquioxane (POSS) in a moderate yield. Thus, under the given experimental

conditions, phenyltriethoxysilane and γ–aminopropyltriethoxysilane mixture is

anticipated to yield POSS macromonomers functionalized with aminopropyl and

phenyl groups. Amine groups are capable of reacting with dianhydrides to form imide

rings and thus the compatibility between the POSS cage and polyimide matrix could

be significantly enhanced, as compared to adding unreactive POSS into the polymer

matrix [14]. In our study, we have reported the method presented in scheme 1. After

hydrolysis, silanols formed and reacted with the other alkoxy groups to give one

molecule of alcohol and finally produced cage–like polysilsesquioxanes via

co–condensation reactions. As is reported, the hydrolysis of alkoxy silane is a typical

stepwise hydrolysis process [48] and it usually proceeds at low speed under neutral

conditions (pH = 7), while condensation reaction is typically inhibited at pH = ~4.2

[49,50]. Therefore, we added tetraethylammonium hydroxide to adjust the pH of the

system to accelerate the reactions. Additionally, the initial hydrolysis stage of the

process should be carried out at room temperature for 10 h. Our data indicates that, a

hydrolysis time of less than 10 h usually results in a remarkable decrease of the final

product yield. In the second stage, cage structure starts to form at relatively low speed

through condensation reactions between silanols. If the second stage is carried out for

a longer time, a higher yield was obtained, but mainly in the direction of formation of

networked or branched polysiloxanes. So there is an optimum condition that is in

favour of formation of cage–structured products [51]. In this paper, the reasonable

condensation time is around 50 h.

As is reported by Feher [40] and Pavithran [31], the sol–gel process to the

direction of formation of cage structure is further proceeding on concentrating most of

the ethanol by slowly evaporation. In this paper, the solution was carefully cast on a

big dust–free glass plate so as to slowly evaporate the ethanol at ambient temperature,

finally yielding white powder crude products. However, as Feher and co–workers

pointed, formation of POSS cage typically occurs with thermodynamic control and

usually gives predominantly fully condensed POSS together with incompletely

condensed POSS [40]. Therefore, the crude products should be purified to remove the

“drop–in” networked products according to the typical procedures as described in the

experimental section. In the report of Wei and co–workers [30], POSS–diamine with

apparent conjugated structure demonstrated the rhombohedral crystal morphology. In

our formulation, however, the synthesized POSS–diamine seems amorphous, as is

evidenced by its X–ray diffraction curve depicted in figure 1a. There are only two

distinct diffraction peaks at 2� = 8.1° and 20.6° by POSS–diamine, corresponding to

d–spacings of 10.8 Å and 4.2 Å, respectively. The peak corresponding to a d–spacing

of 10.8 Å is caused by the size of bifunctional POSS molecules [30], while the other

peak is broad, implying that POSS–diamine possesses completely amorphous

structure [52]. This is well in agreement with its disordered aggregation morphology,

as is typically observed in scanning electron microscope (SEM) image shown in

figure 1b.

3.2 Structural Analysis of POSS–diamine. Figure 2 has given the fourier

transform infrared spectral features of POSS–diamine. The characteristic bands

between 3300 and 3500 cm-1 are ascribed to primary amine group (N–H stretching),

implying that POSS has been functionalized with amine groups [31]. The spectrum

also contains bands due to phenyl group 3060 and 3016 cm-1 (C–H stretching), 1440

and 1590 cm-1 (C=C stretching), and propyl group 2930 and 2860 cm-1 (C–H

stretching) [31,53]. The intense and sharp absorption at 1135 cm-1 can be assigned to

Si–O–Si stretching of cage–structured POSS, while the presence of the weak peak at

1050 cm-1 indicates that there still exists trace amounts of incompletely condensed

branched/networked or linear polysilsesquioxanes in the target products [31].

The 1H and 13C nuclear magnetic resonance (NMR) results of POSS–diamine are

illustrated in figure 3. For 1H–NMR spectrum, the distinct peaks are attributed to the

protons of phenyl and aminopropyl groups. Moreover, for 13C–NMR, the signals at

ca. 127–135 ppm are owing to the aromatic carbon atoms, and peaks at ca. 44.3, 26.8

and 10.0 ppm are assigned to carbon atoms of aminopropyl group (a, b and c),

confirming that the POSS cage is decorated with phenyl and aminopropyl groups.

Currently, the molar ratio of phenyl group to aminopropyl group is estimated

according to the 1H–NMR and 13C–NMR spectra of POSS–diamine by comparing the

integrals of peaks assignable to respective units. The calculated value is ca. 2.9 and

2.99 based on different NMR spectra, very close to the theoretical value 3, suggesting

that POSS–diamine possesses the anticipated structure.

3.3 Thermal Decomposition Behaviour of POSS–diamine. Figure 4 has

displayed the results of thermogravimetric analysis (TGA) and differential scanning

calorimetry (DSC) analysis of POSS–diamine. The TGA and DSC curves demonstrate

a typical two–step decomposition process. The first step at ca. 275–530 °C is gradual,

while the second is much more rapid at ca. 530–670 °C. For POSS–diamine, the onset

thermal decomposition temperature is approximately 275 °C and 5 wt % weight loss

decomposition temperature (Tdec) is ca. 470 °C. As is known to all, the magnitude of

Tdec for a material is mainly dominated by its chemical structure, such as the bonding

energy, defects within the molecular skeletons and bond reactivity [54]. Generally,

materials having higher bonding energy usually demonstrate better thermal stability.

POSS–diamine mainly contains four distinct chemical bonds, Si–O, C–C, C–N and

Si–C, which have different bonding energies of ca. 420, 332, 305 and 340–240

kJ/mol, respectively. Therefore, the thermal decomposition of POSS–diamine is likely

to start from the initial cleavage of C–C, C–N and Si–C bonds in the POSS corner due

to their weak bonding energies. Thus the first thermal degradation step is mainly

ascribed to the decomposition of aminopropyl, whereas the second step is most likely

owing to the decomposition of phenyl groups.

The char yield is often regarded as the weight residual percentage of samples

after TGA measurement, and this value is ca. 56.3 % for POSS–diamine in air

atmosphere, a litter higher than the theoretical value 48.3 %. This is most likely the

result of incomplete thermal degradation of thermal–stable phenyl groups in the POSS

corner and absorption of volatiles by the resulting char residuals. At high temperature,

phenyl groups begin to decompose and POSS cage gradually collapses, eventually to

give inorganic polysiloxanes (SiOx). In order to monitor the molecular changes of

POSS–diamine, the char residuals were characterized by fourier transform infrared

spectroscopy. The results are presented in figure 5. As can be clearly observed, no

evident absorption bands appear at the range of 2700–3600 cm-1 after thermal

degradation, suggesting that almost all of the organic substituents (aminopropyl and

phenyl groups) surrounding the POSS cage are completely decomposed through

thermal oxidation degradation. In addition, the strong Si–O–Si stretching absorption

band of cage–structured POSS–diamine varies from 1135 cm-1 to 1100 cm-1,

indicating that POSS cage completely collapses and is disrupted and finally converts

to networked SiOx.

3.4 XPS Analyses of POSS–diamine. Figure 6 shows the high–resolution XPS

spectra of C 1s and Si 2p for POSS–diamine. The C 1s peak is decomposed into the

C–Si, C–C, C–H and C–N sub peaks based on the assumption that each peak consists

of the Gaussian/Lorentzian sum function [55]. The energy positions of these sub

peaks has been labeled in figure 6a. The Si 2p peak is found out at approximately

102.4 eV. This is well in agreement with the electron state of silicon atoms within

POSS–diamine (C–Si–O3) [55]. Strictly speaking, the decompositions are not

sufficient to accurately acquire the fraction of the bond. Therefore, as a rough

estimation, the mole fraction of the bond was obtained from the ratio based on the

area of corresponding peak. The full width at half maximum (FWHM) of Si–C, C–C,

C–N and C–Si–O3 are 1.39, 1.86, 2.01 eV and 2.0 eV, respectively. The ratio between

Si–C and C–N bond concentration is approximately 3.6, very close to 4 in theory.

Thus the cage structure of POSS–diamine is once again evidenced by its XPS

patterns.

3.5 Polymer Characterization. The structures of the resulting polyimides were

characterized by fourier transform infrared spectroscopy (FT–IR). As shown in figure

7, the absorption peaks at ca. 1780 cm-1 (C=O asymmetrical stretching), 1715 cm-1

(C=O symmetrical stretching), 1375 cm-1 (C–N stretching) and 720 cm-1 (C=O

bending) suggest the formation of the five–membered aromatic imide rings.

Additionally, the vanishing of the polyamic acid bands at 1650 (C=O stretching of

polyamic acid) and 1535 cm-1 (C–N stretching of polyamic acid) indicates complete

conversion of the amic acid to imide groups [56]. The sharp bands around 1100 cm-1

can be attributed to the Si−O−Si asymmetric stretching vibrations of POSS cage,

suggesting that POSS–diamine is successfully incorporated into polyimide backbones

[57]. Besides, there is no existence of the typical absorption bands of the isoimide

groups around 1810 cm-1 and 980 cm-1 in the FT–IR spectra [56].

The resulting membranes are amber in colour and completely transparent. The

morphology of the resulting polyimides was evaluated by utilizing a wide–angle

X–ray diffraction (WAXD) diffractometer. The diffraction curves of all polyimides

are broad and no obvious peak features appear in all cases, implying that all

membranes possess completely amorphous structure. From the scattering angles (2θ)

in the central of the broad peaks, the d–spacing is calculated according to Eq. (1) and

the specific results are summarized in table 2. It is well known that, for crystal

polymers, d–spacing is often described as the average distances between neighbouring

repeating units, while In the case of linear polymers, such as POSS polyimides,

d–spacing is often referred to the average intersegmental distances [58,59]. As listed

in table 2, the 2θ values gradually decrease and the d–spacing increases with POSS

amount, suggesting an increasing average intersegmental distance. This is most likely

a result of inefficient packing density due to the significant steric hindrance of POSS

molecules, which is supported by the decreasing densities of POSS polyimides as

illustrated in table 2.

3.6 Optical Properties. The UV–visible absorbance and transmittance spectra of

the obtained POSS polyimides are shown in figure 8. Except 29.7 wt % POSS

polyimide, all POSS polyimides demonstrate slight absorbance and desirable optical

transparency in the visible light region. As shown in table 2, from 4.1 wt % to 21.9 wt

% POSS polyimides, the optical transparency at 600 nm of hybrid membranes

maintained at approximately 75%, and is comparable to that of the pristine one. This

indicates that good compatibility between POSS–diamine and polyimide matrix has

been achieved, owing to the reactive amine groups in the POSS corner. It is

significant to note that, once the POSS addition exceeds a certain critical value, POSS

molecules are likely to aggregate and present self–assembled characteristics [30,60].

This may be responsible for the reduced optical transparency of 29.7 wt % POSS

polyimide.

3.7 Mechanical Performances. Table 2 also shows a negligible reduction (113.0

vs 113.4) in tensile strength for the 4.1 wt % POSS polyimide membrane, as

compared to that of the pristine one. Even if the POSS loading reaches 14.4 wt %, the

tensile strength still exhibits admirable retention (105.3 vs 113.4). However, as POSS

addition increases to 21.9 wt %, the resulting hybrid membranes are quite weak and

brittle. More specifically, for 29.7 wt % POSS polyimide, it is far too brittle to get its

tensile strength, although this formulation demonstrates an acceptable intrinsic

viscosity (1.63 dL/g). A few possible causes may be responsible for the decreasing

mechanical performance of POSS polyimides. Since POSS cage is approximately 1~2

nm [42], quite large and stiff, and generates significant steric hindrance. Therefore,

POSS–diamine may demonstrate lower reactivity, as compared to

4,4'–diaminodiphenyl ether [30]. This eventually results in lower molecular weight

and inefficient chain–chain packing, which is supported by the decreasing intrinsic

viscosities and densities as summarized in table 1 and table 2. Additionally, the

resulting POSS macromonomer is likely to contain a very small amount of

mono–amine POSS, possibly inhibiting the molecular chain growth, which is also

responsible for the reduction in mechanical performance to a certain extent. Second,

POSS molecule consists of many Si–O bonds, less polar than the imide rings. This

may remarkably increase the intervals between the polyimide chains and lead to more

free volume, resulting in decreased chain–chain entanglements and weak chain–chain

interactions, which also reduce the mechanical performance of the nanocomposites

[30]. On the whole, although POSS incorporation has exerted certain effects on

mechanical properties of the resulting nanocomposites, the very least tensile strength

is still above 70 MPa, much better than that afforded by 10 mol % POSS polyimide

reported by Wei and co–workers (70.2 vs 46.4) [30].

3.8 Thermal Properties. Table 3 has summarized the dynamic mechanical

thermal analysis (DMA), Thermo–mechanical analysis (TMA) and thermal

gravimetric analysis (TGA) results of the resulting nanocomposites. The results

indicate that the storage modulus (E’, at 50 °C) of the POSS nanocomposites are

between 2610 and 2410 MPa, higher than that of the pristine polyimide (2400 MPa),

but monotonically decreased with POSS addition. The higher storage modulus of the

nanocomposites may be caused by the restricted rotational motion of the polyimide

chains afforded by POSS molecules due to its significant rigidity and steric hindrance

[61], while the decrease in storage modulus with POSS amount possibly originates

from the increase in free volume within polyimide matrix and inefficient chain–chain

entanglement tendency, which is also in accordance with the decreasing tensile

properties as summarized in table 2.

The glass transition temperatures (Tgs) of the resulting nanocomposites were

determined by DMA. The values reported in table 3 are regarded as the peak

temperature in the internal loss factor (tan δ) curves. It is clearly indicated that, the Tg

gradually increases with POSS amount and reaches its highest value of 392 °C at 21.9

wt % POSS addition, which is quite different from Wei’s [30] and Minton’s [62]

work. In their investigations, a clear trend toward decreasing Tgs of resulting

nanocomposites with POSS content is visible, possibly arising from the increase in

the free volume and the inefficient packing of the polymer chains due to the presence

POSS molecules. In fact, the Tg of a polymer is affected by many factors, such as

rigidity of polymer chains, molecular weight of polymer, interactions between

polymer chains, specific structure of polymer (hyperbranched, linear or crosslinked)

and fraction of the free volume. In the present investigation, however, Tg of the

resulting nanocomposites seems to be dominated by the rigidity of polyimide chains.

The cooperative movement of the segment may be significantly restricted due to

POSS molecules and hence lead to higher Tgs. The increasing rigidity is well in

agreement with the decreasing elongation at break as summarized in table 2.

Additionally, the POSS–diamine employed in the present study differs in molecular

structure from Minton’s and Wei’s investigations. This may also be responsible for

the contrary variations in Tgs with POSS amount.

Table 3 also shows a decreasing trend in the intensity of the tan δ at Tg. As is

known to all, the intensity of the tan δ at Tg is an evaluation of the energy–damping

characteristic of a material [61]. As indicated in table 3, POSS nanocomposites seem

not so good as the pristine polyimide at dissipating energy. This also suggests a

hindered rotation of polyimide chains arising from the rigidity of POSS molecules

[61], in accordance with the increasing Tgs.

Figure 9 has depicted the storage modulus of the resulting nanocomposites in the

rubbery plateau region (about 30~40 °C above Tg). As pointed by Minton [62] and

Menard [63], the magnitude of the modulus in this region is inversely proportional to

the molecular weight ( Mc ) between neighbouring entanglements and proportional to

the cross–linking density (ρd). Thus if a large increase in modulus is typically

observed in this region, it is indicative of the occurrence of a cross–linking reaction

during the testing. During thermal imidization, a minority of the imide rings may be

cleaved and recombination with another polyimide chain possibly occurs. This

phenomenon is known to occur at high temperatures for polyimides fabricated with

4,4'–diaminodiphenyl ether and pyromellitic dianhydride [62]. As can be clearly seen

in figure 9, the presence of POSS cage is likely to retard this cross–linking reaction at

low loading amount (4.1 and 8.8 wt %), due to the significant rigidity of POSS cage.

However, at higher loading amount (14.4 and 21.9 wt %), POSS molecule seems to

demonstrate a reinforcing mechanism, which ultimately results in an increased

modulus in the rubbery region, indicating that there is an optimal amount of POSS

addition to inhibit the cross–linking reactions [62].

In a practical application, polyimide membranes are extensively used as plastic

substrates in display devices [64] or laminated with other metals or ceramics for

further thermal processing [61]. Thus the dimensional stability should be taken into

account. The CTE value is the indicative of dimensional stability of polyimide

membranes and presently is evaluated in thermo–mechanical analysis (TMA) mode in

the temperature range of 75 °C to 150 °C. As indicated in table 3 and figure 10, the

CTE values of POSS nanocomposites are between 33.5 and 58.2 ppm�K-1, remarkably

slower than that of the pristine polyimide (68.2 ppm�K-1), but gradually increase with

POSS addition. As mentioned above, POSS incorporation possibly leads to a

restricted thermal rotational motion of the polymer segments, finally resulting in

better dimensional stability in comparison to the pristine polyimide [61,64]. In

addition, the POSS cage is likely to restrict the heat transfer through polyimide matrix

due to the exceptionally high thermal stability afforded by the high Si–O bonding

energy, which may also be responsible for the better dimensional stability of POSS

nanocomposites [65]. The gradual increase in CTE value with POSS amount is

possibly a result of the increase in the free volume within nanocomposites owing to

POSS cage [30].

The thermal gravimetric analysis (TGA) curves of POSS polyimides in air are

shown in figure 11. As indicated, the increasing residues at 1000 °C from 4.1 to 29.7

wt % POSS polyimides also suggest the successful incorporation of the POSS

molecules in the hybrid materials. The thermal decomposition behaviour of all

nanocomposites was investigated by utilizing a TGA/DSC 1 synchronous thermal

analyzer at a heating rate of 10 °C/min under air atmosphere from room temperature

to 1000 °C. The TGA and differential scanning calorimetry (DSC) curves of the select

sample, PI–21.9, are illustrated in figure 12. The DSC curves suggest two distinct

exothermal peaks in all cases, implying a typical two–step weight loss process of

resulting nanocomposites. The first step at ca. 450–550 °C is gradual and

insignificant, possibly due to the initial thermal degradation of aminopropyl groups,

while the second is much more rapid and remarkable at ca. 550–680 °C, mainly owing

to the scission of polyimide main chains and decomposition of phenyl groups. The

thermal degradation temperatures (Tds) of resulting nanocomposites in air and

nitrogen, summarized in table 3, are slightly lower than that of the pristine polyimide

ascribed to the initial decomposition of aminopropyl groups (at about 450 °C, vide

supra) and lower molecular weight due to POSS introduction. All nanocomposites

exhibit degradation temperatures at 5 wt % weight loss ranging from 544 to 563 °C in

air and 566 to 572 °C in nitrogen. In the low earth orbit (LEO) environment, the

typical thermal cycle suffers by spacecraft is approximately ± 100 °C [18]. Thus, the

admirable thermal performance of these nanocomposites promises this

newly–developed material can be used in the LEO environment.

The isothermal TGA analyses of resulting polyimides were carried out at two

different temperatures (400 °C and 500 °C) in nitrogen for 10 h. The specific weight

loss of nanocomposites aging at 500 °C for 10 h are summarized in table 3. All

nanocomposites are quite stable, exhibiting negligible weight loss ( 1%) at 400 °C

and only a few percent at 500 °C, suggesting that relatively short time usage at high

temperatures are possible. One will first notice that, almost all nanocomposites

demonstrate superior thermal stability in comparison to pristine polyimide, which is

contrary to the TGA results. This is most likely ascribed to the high fraction of the

free volume due to POSS introduction, possibly hampering the heat transfer through

polyimide matrix during isothermal process. In addition to this, the phenyl–substituted

macromonomer, namely POSS–diamine, is probably decomposed not really the

fastest in high temperature environment, which may also be responsible for the

decreased weight loss. On the other hand, molecular weight also plays an

indispensable role in thermal decomposition. As table 1 indicated, the intrinsic

viscosities of POSS polyamic acids gradually decrease with POSS addition,

suggesting that nanocomposites with high to medium molecular weight are formed,

which may be responsible for the decreasing thermal stability from 4.1 to 14.4 wt %

POSS polyimide. However, once the POSS amount exceeds a certain critical value,

the molecular weight possibly no more makes the dominant contribution to the

thermal decomposition, but the high fraction of the free volume may hinder heat

transfer through polyimide matrix and possibly dominates the thermal behaviour of

the nanocomposites. Therefore, 21.9 and 29.7 wt % POSS polyimide exhibit

decreased weight loss. However, this hypothesis is presently still under investigation.

3.9 AO Resistance Properties of the Resulting Nanocomposites. Simulated AO

exposure measurements were performed on polyimide membrane samples by utilizing

a neutral beam producing facility in our lab. The mass loss of AO–irradiated

polyimide membranes are illustrated in table S1 (Supplementary information) and

figure 13. It is clearly indicated that, the mass loss of pristine polyimide increases

linearly with AO fluence (or AO exposure duration), suggesting a constant erosion

rate, while the POSS/polyimide nanocomposites demonstrate a much lower erosion

rate under the same AO exposure conditions and, their mass loss approximately

exhibits as power function of the AO fluence throughout the whole exposure period

[28,43]. For 14.4, 21.9 and 29.7 wt % POSS polyimide, it is clearly seen that their

mass loss is significantly lower than that of pristine polyimide at a given AO fluence

and mass loss rate gradually decreases with POSS addition or AO fluence. After

exposure to an AO fluence of 0.88 × 1020 O atoms·cm-2, 29.7 wt % POSS polyimide

demonstrates a mass loss of ca. 9.7 % that of pristine polyimide. More specifically, in

an AO exposure with a total fluence of 3.87 × 1020 O atoms·cm-2, its mass loss is

merely ca. 2.9% that of pristine polyimide. These results suggest that POSS

incorporation is of significant importance in resisting AO attack in the simulated AO

environment.

To predict the service life of POSS polyimide in the simulated AO environment,

curve–fitting model is herein used to provide the mass loss as the function of AO

fluence. The best–fit curve is expressed by the following function

1.91230.0072 0.2500y x� (standard deviation = 0.15%, R2 = 0.99) for 29.7 wt % POSS

polyimide membrane. Its service life is estimated based on the mass loss of POSS

polyimide membranes in the simulated AO environment by assuming that the

thickness of 29.7 wt % POSS polyimide membranes is 40 μm and the polyimide

membranes could no more resist AO erosion when its mass loss reaches 50% that of

its original state, which corresponds to an AO fluence of 61.84 × 1020 O atoms·cm-2.

As reported, exposure to 2.0 × 1020 O atoms·cm-2 Kapton–equivalent AO fluence is

roughly equivalent to half one year exposure to AO in low earth orbit environment

[18]. Therefore, the service life of 29.7 wt % POSS polyimide membrane is estimated

to be approximately 15 years in the simulated AO environment. However, this is just

an estimate only considering the AO attack in the simulated AO environment. Since

the space factors are complex and severe, the actual service life of POSS polyimide

membranes may be much shorter in real space environment.

In order to differentiate the mass loss with POSS amount and AO fluence, the AO

erosion yields of POSS/polyimide nanocomposites were calculated according to

previous study [43] and compared with pristine polyimide. The specific results are

summarized in table S1 (Supplementary information). The erosion yields, relative to

that of pristine polyimide, when POSS polyimides are subjected to AO equivalent

fluences of 0.88, 1.76, 2.64 and 3.87 × 1020 O atoms·cm−2, are graphed in figure 14 as

the function of POSS amount. It is clearly indicated that the erosion yields of all

nanocomposites appear to rapidly decrease with increasing both AO fluence and

POSS addition, and a dramatical reduction in AO erosion yield is visible when the

POSS loading surpasses 8.8 wt %, suggesting that nanocomposites with less POSS

content are possibly insufficient to generate any significant enhancement to AO

resistance towards AO attack. Thus it seems that there is a threshold of POSS addition

to remarkably enhance the AO resistance of nanocomposites. In addition, for 29.7 wt

% POSS polyimide, the maximum reduction in AO erosion yield is about 97%,

approximately two orders of magnitude reduction in AO erosion yield by comparing

to pristine polyimide under the same conditions. From the experimental data

presented here, it is evident that the resulting nanocomposites with high POSS loading

have high survivability in the simulated AO environment.

Surface morphologies of resulting polyimides before and after AO irradiation

were observed by scanning electron microscope (SEM) and atomic force microscopy

(AFM). The SEM and AFM images of select samples, 8.8 wt % and 21.9 wt % POSS

polyimides before and after exposure to AO, are given in figures 15 and 16

respectively. It is evident that, before AO exposure, surfaces of POSS polyimides are

flat and smooth, in accordance with the small root mean square roughness (RMS)

values above the corresponding AFM images. However, the surface morphologies

have been significantly altered after exposure to AO. As indicated in figure 15 and our

previous study [43], the surface of pristine polyimide is remarkably roughened and

presents “carpet–like” morphology with the RMS values rapidly increasing from 2.66

to 327.6 nm after AO exposure from 0 to 3.87 × 1020 O atoms·cm-2. In contrast, the

surfaces of the nanocomposites are much denser and smoother albeit a certain amount

of microcracks and pinholes appeared on the surfaces. For 8.8 wt % POSS polyimide,

a clear trend toward increasing RMS values is visible after AO irradiation from 0 to

3.87 × 1020 O atoms·cm-2, but it is to a much lesser extent in comparison to pristine

polyimide. Additionally, for 21.9 wt % POSS polyimide, its root mean square surface

roughness is quite low and gradually decrease with AO fluence, suggesting that the

nanocomposites with higher POSS content becomes increasingly resistant to AO

attack with AO fluence, which corresponds to the decreasing mass loss rate illustrated

in figure 13(b). The superior AO resistance of polyimides containing POSS molecules

is once again highlighted by their surface morphologies.

As discussed above, POSS molecules play a positive role in resisting AO attack

and therefore impart high survivability to POSS/polyimide nanocomposites in AO

environment. There are a few possible causes for the enhanced AO durability of

nanocomposites. Since the surface coverage of polyimide moiety gradually decreases

with increasing POSS amount, this may result in less severe oxidation degradation of

organic portions and lead to decreased mass loss during AO irradiation and hence

enhanced AO resistance. In addition, phenyl–substituted macromonomer POSS, may

demonstrate slower degradation in the AO environment due to the admirable stability

of benzene ring and its inorganic silica–like cores. This may also be responsible for

the enhanced AO resistance of POSS polyimides to a certain extent. However, these

cannot account for the entire significant reduction in mass loss as well as that in AO

erosion yield. As reported by Minton [9,19] and Gonzalez [38], when

POSS/polyimide nanocomposites are exposed to AO environment, the organic

polyimide portions are likely to be much more easily eroded away while silicon atoms

remain and are eventually oxidized to a SiO2 passivating layer. This passivating silica

layer may protect the underlying polymer from further AO attack. Thus, one may

suspect whether the inert silica protective layer limits the erosion yield to a low value

in the current formulation. In order to reveal the protection/erosion mechanism of

POSS/polyimide nanocomposites in AO environment and develop an understanding

of the nature of the passivating layer, X–ray photoelectron spectroscopy (XPS)

measurements were adopted to detect the changes in electron states and

concentrations of surface atoms on topmost surface (approximately 0–10 nm) of the

polyimide membranes.

Currently, all sample surfaces before and after exposure to various AO fluences

have been probed by XPS and the results are presented in table S1 (Supplementary

information), figure 17 and figure 18. For pristine polyimide, its surface compositions

have changed insignificantly and only a slight decrease in carbon atomic

concentration appears. In contrast, a clear trend toward remarkably decreasing carbon

atomic concentration is visible for POSS/polyimide nanocomposites after AO

exposure. Prior to AO irradiation, the silicon atomic concentration is very low, while

a dramatical increase in silicon and oxygen atomic concentrations in all cases are

typically observed even after a small dose of AO irradiation, and these values

approximately tend to a steady–state value as the AO fluence increases, suggesting

that the main reaction path of mass loss for POSS/polyimide nanocomposites in AO

environment is likely to be the oxidation degradation of surface atoms like carbon,

hydrogen and nitrogen, which possibly generate off–gassing of volatile species

(carbon dioxide, carbon monoxide, nitrogen oxide and hydroxide) while silicon atoms

remain and are possibly oxidized to silicon oxides (SiOx) when AO reacts with a

hydrocarbon–based surface [9,19,28,62].

From the XPS analysis discussed above, it is confirmed that the AO–irradiated

POSS polyimides is covered with an inert silicon oxides layer. With the aim of

developing an understanding of the nature of the passivating layer, high–resolution

XPS survey spectra of Si 2p peaks are obtained corresponding to all polyimide

membranes before and after exposure to AO, and the Si 2p spectra of select samples,

PI–8.8 and PI–21.9 are shown in figure 19. It is clearly indicated that, the Si 2p peak

shifts from a lower binding energy of ~101.7 and ~102.8 eV, possibly corresponding

to suboxide and silsesquioxane (Si2O3) [19], to a higher binding energy of ~103.7 eV.

This indicates that the inert SiOx passivating layer is most likely SiO2 [9,28,43],

probably due to the reaction path that the AO has reacted with POSS molecules by O

atom addition to Si–O and Si–C bonds followed by O–O and Si–C bonds scission

[62], thus the POSS cage has been disrupted and finally oxidized, ultimately to give

silica [9,19]. This hypothesis is also supported by the O/Si ratio shown in figure 20.

When AO fluence increases from 0.88 to 3.87 × 1020 O atoms·cm-2, O/Si ratio of

POSS/polyimide nanocomposites gradually tends towards 2/1 and this is much more

apparent for nanocomposites with high POSS addition. It is significant to note that

O/Si ratio of 14.4, 21.9 and 29.7 wt % POSS polyimides rapidly reaches

approximately 2/1 albeit experiencing a small dose of AO irradiation of 0.88 × 1020 O

atoms·cm-2, suggesting a rapid conversion of silsesquioxane to SiO2 upon AO

exposure, which is also consistent with remarkable decrease in carbon atomic

concentration summarized in table 4. Therefore, for 14.4, 21.9 and 29.7 wt % POSS

polyimides, the organic polyimide moiety on the topmost surface may be rapidly

eroded away and a connected and dense silica passivating layer may be generated

within relatively short periods and plays the role of protective coating, consequently

leading to gentle mass loss and superior AO resistance [44]. From the SEM images

shown in figure 15, it is clearly indicated that the AO–exposed surfaces of 21.9 wt %

POSS polyimide are much denser and more connected than 8.8 wt % POSS

polyimide, indicating that higher POSS cage addition is likely to generate preferable

and more connected silica passivating layer upon AO exposure. This may be

responsible for the threshold of POSS content to remarkably improve the AO

resistance of resulting POSS polyimides.

4. Conclusions

In this paper, a novel bifunctional macromonomer, POSS–diamine, was prepared,

characterized, and subsequently copolymerized with imide monomers to prepare

POSS–based nanocomposites. The resulting POSS polyimides exhibit a combination

of admirable properties such as good thermal stability, desirable mechanical strength

and high AO survivability. Results from AO exposure experiments indicate that the

AO erosion yields of the resulting nanocomposites are approximately two orders of

magnitude reduction that of the pristine polyimide. XPS analysis demonstrates that

the surface carbon atomic concentrations of the resulting nanocomposites remarkably

decrease while the surface silicon atomic concentrations significantly increase and are

further oxidized from suboxide to SiO2, thus protecting the underlying polymer from

additional erosion. SEM and AFM images demonstrate that the surface of pristine

polyimide becomes significantly roughened and presents “carpet–like” morphology

whereas 21.9 wt % POSS polyimide exhibits less root mean square surface roughness.

Furthermore, SEM images reveals that higher POSS amount is likely to generate a

much denser and more connected silica passivating layer upon AO exposure, which

may be responsible for the threshold of POSS content to remarkably improve the AO

durability of the resulting POSS polyimides. Polyimide main chain offers the

advantage of maintaining desirable thermal stability and mechanical strength, while

POSS molecule imparts high AO survivability to resulting hybrid materials. This

indicates that POSS polyimides may find wide use as surface protective materials

onboard spacecrafts to resist AO attack in the low earth orbit environment.

Acknowledgements

The authors are grateful for the financial support provided by the National

Natural Science Foundation of China (No. 51173146), Basic Research Fund of

Northwestern Polytechnical University (JC20120248).

Author information

*Corresponding author, E–mail: qyzhang@nwpu.edu.cn; Tel: +86–029–88431675;

Fax: +86–029–88431653;

Notes

†Electronic supplementary information (ESI) is available free of charge on the

internet: http://www.sciencedirect.com.

All of the authors of this paper declare no competing financial interest.

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Figures and Tables

Scheme 1. Possible chemical structure of POSS–diamine and its synthesis route

through hydrolytic co–condensations of siloxane compounds.

Scheme 2. Fabrication of POSS/polyimide nanocomposites through

co–polycondensation and thermal imidization. (Note: The molecular structure of

POSS/polyimide nanocomposites presented here represents one of the possible

molecular structures)

Fig. 1. X–ray diffraction (XRD) pattern and surface morphology of POSS–diamine,

(a) XRD curve of POSS–diamine and (b) scanning electron microscope (SEM) image,

showing the aggregation structure of POSS–diamine. Note: the scale bar at the bottom

of SEM image indicates 1 μm.

Fig. 2. Fourier transform infrared spectrum (FT–IR) of POSS–diamine.

Fig. 3. Nuclear magnetic resonance (NMR) spectra of POSS–diamine in DMSO–d6:

(a) 1H–NMR and (b) 13C–NMR.

Fig. 4. (Top) Thermal gravimetric analysis (TGA) and (bottom) differential scanning

calorimetry (DSC) curves of POSS–diamine with a heating rate of 10 °C�min-1 in air

from room temperature to 1000 °C.

Fig. 5. Fourier transform infrared spectra (FT–IR) of POSS–diamine before (a) and

after (b) thermal decomposition under air atmosphere.

Fig. 6. Deconvoluted binding energy of POSS–diamine: (a) C 1s and (b) Si 2p peaks.

C 1s peak is fitting into C–Si, C–C, C–H and C–N sub peaks. Si 2p peak is fitting into

C–Si–O3 sub peak.

Fig. 7. Fourier transform infrared spectra (FT–IR) of the resulting polyimide

membranes obtained by utilizing an attenuated total reflectance (ATR) instrument

under ambient conditions. (Note: x wt % POSS polyimide is denoted as PI–x,

similarly hereinafter)

Fig. 8. Optical properties of POSS polyimide membranes with thicknesses ranged

from 40 to 45 μm: (a) absorbance and (b) transmittance in the ultraviolet and visible

light region.

Fig. 9. Storage modulus for the resulting polyimide membranes in the rubbery plateau

region.

Fig. 10. Coefficient of thermal expansion (CTE) of polyimide membranes determined

with a heating rate of 5 °C�min-1 over the temperature range of 75−150 °C in

thermo–mechanical analysis (TMA) mode.

Fig. 11. Thermal gravimetric analysis (TGA) curves of POSS/polyimide

nanocomposites with a heating rate of 10 °C�min-1 from room temperature to 1000 °C

in air.

Fig. 12. (Top) Thermal gravimetric analysis (TGA) and (bottom) differential scanning

calorimetry (DSC) curves of 21.9 wt % POSS polyimide with a heating rate of 10

°C�min-1 from room temperature to 1000 °C in air.

Fig. 13. Mass loss of polyimide membranes containing 0, 4.1, 8.8, 14.4, 21.9 and 29.7

wt % POSS, normalised to the exposure area (9π/4 cm2) of the polyimide membranes:

(a) normalised mass loss vs POSS addition and (b) normalised mass loss vs AO

fluence. Samples were subjected to AO equivalent fluences of 0.88, 1.76, 2.64 and

3.87 × 1020 O atoms·cm−2 at ambient conditions.

Fig. 14. AO erosion yields of polyimide membranes containing 4.1, 8.8, 14.4, 21.9

and 29.7 wt % POSS, relative to pristine polyimide (0 wt % POSS) erosion yield.

Samples were subjected to AO equivalent fluences of 0.88, 1.76, 2.64 and 3.87 × 1020

O atoms·cm−2 at ambient conditions.

0 1.76 3.87

Fig. 15. Select SEM images of polyimide membranes with increasing atomic oxygen

fluence. (A−C) pristine polyimide, (D−F) 8.8 wt % and (G−I) 21.9 wt % POSS

polyimide surfaces after exposure to AO fluences of 0, 1.76 and 3.87 × 1020 O

atoms·cm−2. Note: AO fluence is shown above corresponding images and the scale

bar at the bottom of each image indicates 1 μm.

1.55 nm 88.2 nm 112.6 nm

1.25 nm 12.1 nm 7.27 nm

Fig. 16. Select atomic force microscopy (AFM) images (5 μm × 5 μm) of polyimide

membranes with increasing AO fluence. (A−C) 8.8 wt % and (D−F) 21.9 wt % POSS

polyimide surfaces after exposure to AO fluences of 0, 1.76 and 3.87 × 1020 O

atoms·cm−2. Note: root mean square roughness (RMS) values are shown above

corresponding images.

Fig. 17. Surface atomic concentration of carbon (atom %) determined from X–ray

photoelectron spectroscopy (XPS) survey scans for POSS/polyimide nanocomposites

after exposure to AO fluences of 0, 0.88, 1.76, 2.64 and 3.87 × 1020 O atoms·cm−2, (a)

concentration vs POSS addition and (b) concentration vs AO fluence.

Fig. 18. Surface atomic concentration of silicon (atom %) determined from X–ray

photoelectron spectroscopy (XPS) survey scans for POSS/polyimide nanocomposites

after exposure to AO fluences of 0, 0.88, 1.76, 2.64 and 3.87 × 1020 O atoms·cm−2, (a)

concentration vs POSS addition and (b) concentration vs AO fluence.

Fig. 19. High–resolution X–ray photoelectron spectroscopy (XPS) spectra of Si 2p

curves corresponding to (a) 8.8 wt % and (b) 21.9 wt % POSS polyimide membranes

after exposure to AO fluences of 0, 0.88, 1.76, 2.64 and 3.87 × 1020 O atoms·cm−2.

Fig. 20. Diagram of variations in O/Si ratio for POSS/polyimide nanocomposites after

exposure to AO fluences of 0, 0.88, 1.76, 2.64 and 3.87 × 1020 O atoms·cm−2, (a) O/Si

ratio vs POSS addition and (b) O/Si ratio vs AO fluence.

Table 1. Recipe and intrinsic viscosities for the synthesis of pristine and

POSS–containing polyamic acids.

Sample ODA/g PMDA/g POSS–diamine/g DMAc/g �inha/dL·g-1

Pristine PI 2.002 2.181 / 30.68 3.32

4.1 wt % POSS PI 1.959 2.181 0.180 31.68 3.23

8.8 wt % POSS PI 1.908 2.181 0.395 32.88 2.69

14.4 wt % POSS PI 1.841 2.181 0.676 34.45 2.60

21.9 wt % POSS PI 1.740 2.181 1.100 36.82 1.99

29.7 wt % POSS PI 1.620 2.181 1.604 39.64 1.63

a Determined in N,N'–dimethylacetamide (0.5 g�dL-1) at 25 °C.

Tab

le 2

. Pro

perti

es o

f pris

tine

and

POSS

pol

yim

ide

mem

bran

es.

Sam

ple

Mem

bran

e de

nsity

(g·c

m-3

)a 2θ

(°)b

d–sp

acin

g (Å

)b Tr

ansp

aren

cy a

t

600

nm (%

)c

Tens

ile st

reng

th

(MPa

)d

Bre

ak e

long

atio

n

(%)d

Pris

tine

PI

1.41

2 19

.15

4.63

79

.2

113.

4 26

4.1

wt %

PO

SS P

I 1.

408

18.9

0 4.

70

76.0

11

3.0

20

8.8

wt %

PO

SS P

I 1.

402

18.7

5 4.

73

73.6

10

6.3

12

14.4

wt %

PO

SS P

I 1.

399

18.1

3 4.

89

74.7

10

5.3

10

21.9

wt %

PO

SS P

I 1.

395

17.9

8 4.

93

75.6

70

.2

5

29.7

wt %

PO

SS P

I 1.

378

17.5

4 5.

05

67.0

—

—

a Mea

sure

d by

liq

uid

susp

ensi

on m

etho

d at

am

bien

t co

nditi

ons.

b Det

erm

ined

by

wid

e an

gle

X–r

ay d

iffra

ctio

n (W

AX

D).

c Mem

bran

es f

or

trans

mitt

ance

mea

sure

men

t w

ith t

hick

ness

es r

ange

d fr

om 4

0 to

45

μm.

d Mea

sure

d w

ith a

uni

vers

al t

estin

g m

achi

ne (

UTM

) ac

cord

ing

to

GB

1302

2–91

at a

dra

win

g ra

te o

f 10

mm

/min

at a

mbi

ent c

ondi

tions

.

Tab

le 3

. The

rmal

pro

perti

es o

f pris

tine

and

POSS

pol

yim

ide

mem

bran

es.

Sam

ple

DM

A

TM

A

TG

A

E’a (M

Pa)

tana δ

T g

b (°C

)

CTE

c (ppm

�K-1

)

T dd in

air

(°C

) T d

d in N

2 (°C

) W

eigh

t los

se (%)

Pris

tine

PI

2400

0.

253

371

68

.2±8

.6

56

3±5

571±

6 24

.2

4.1

wt %

PO

SS P

I 26

10

0.21

8 38

0

33.5

±4.5

563±

4 57

2±5

12.6

8.8

wt %

PO

SS P

I 25

85

0.21

1 38

5

47.3

±3.6

559±

2 57

1±8

16.3

14.4

wt %

PO

SS P

I 24

56

0.19

3 38

8

49.3

±5.4

556±

4 57

0±6

24.7

21.9

wt %

PO

SS P

I 24

10

0.14

5 39

2

54.1

±3.8

549±

4 56

8±4

14.6

29.7

wt %

PO

SS P

I —

—

—

58.2

±5.2

544±

9 56

6±5

11.5

a Sto

rage

mod

ulus

(50

°C) a

nd in

tern

al lo

ss fa

ctor

(pea

k va

lue)

, det

erm

ined

in a

tens

ile d

ynam

ic m

echa

nica

l the

rmal

ana

lysi

s (D

MA

) mod

e w

ith

a he

atin

g ra

te o

f 10

°C /m

in a

t a fr

eque

ncy

of 1

Hz

from

room

tem

pera

ture

to 4

50 °

C. b T

he p

eak

tem

pera

ture

in th

e ta

n δ

curv

e is

des

igna

ted

as

T g.

c The

coe

cien

t of

the

rmal

exp

ansi

on d

eter

min

ed w

ith a

hea

ting

rate

of

5 °C�m

in-1

ove

r th

e te

mpe

ratu

re r

ange

of

75−1

50 °

C i

n

ther

mo–

mec

hani

cal a

naly

sis

(TM

A)

mod

e. d T

he th

erm

al d

ecom

posi

tion

tem

pera

ture

at 5

% w

eigh

t los

s w

ith a

hea

ting

rate

of

10 °

C�m

in-1

. e

Wei

ght l

oss a

ging

at 5

00 °C

for 1

0 h

in n

itrog

en fl

ow.

Graphical abstract

Table of Contents Entry Graph

Upon AO exposure, pristine polyimide is severely eroded and exhibits

“carpet–like” surface morphology, while POSS polyimides demonstrate high AO

survivability.

Highlights of this manuscript

Novel polyhedral oligomeric silsesquioxane (POSS)–diamine was synthesized

through a facile hydrolytic co–condensation.

POSS polyimides were fabricated by co–polymerizing POSS–diamine with imide

monomers.

The atomic oxygen (AO) resistance of POSS polyimides were studied upon AO

exposure.

POSS polyimides exhibit significantly improved space survivability.

POSS polyimides with desirable AO survivability may find wide usage in

aerospace.