The influence of nanofillers migration on the mechanical property of PA6/chitosan nanocomposites

Abstract

Blends of chitosan (CS) with polyamide 6 (PA6) were prepared via the solution casting technique using formic acid as the common solvent. Polyhedral oligomeric silsesquioxane (POSS), montmorillonite (PGN), and modified montmorillonite (TCN and TL) were incorporated into the polymer matrix. The dispersion state of the nanofillers in the polymer matrix was characterized by polarized optical microscopy (POM), attenuated total reflection Fourier transform infrared spectroscopy (FTIR-ATR), X-ray diffraction measurements (XRD) and transmission electron microscopy (TEM). It was shown that POSS could achieve the best dispersion state in the polymer matrix with only several tens of nanometers. The thermal and mechanical properties of the synthesized polymer nanocomposites were studied by differential scanning calorimetry (DSC) and an Instron single column material testing system. The morphology and element analysis of the cross-section of the sample film were characterized by scanning electron microscopy with energy dispersive X-ray (SEM/EDX) to further explain the mechanical property results. It was shown that CS tended to sediment during the solvent evaporation step whereas PA6 rich domains filled the upper zones of the samples. Interestingly, POSS preferred to migrate to the upper zone of the PA6 rich domain, while PGN, TCN and TL preferred to migrate to the CS rich domain. The migration of the nanofillers was found to be very important to improve the tensile strength of PA/CS blends.

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1. Introduction

Phase inversion is a process by which a polymer is transformed from a liquid or soluble state to a solid state.¹ The phase inversion procedure is the most versatile technique, which can be used to prepare polymeric membranes. A variety of morphologies can be obtained, which are suitable for different applications, from microfiltration membranes with very porous structures, to more dense reverse osmosis membranes, to gas separation and pervaporation membranes, with complete defect-free structures.^2,3

Chitosan (CS) is one of the most abundant polysaccharides found in nature.^4,5 This material has good solubility in formic acid and can be synthesized to form a membrane according to the phase inversion process. The modification of CS by means of blending nylon is a convenient and effective method of improving its physical properties for practical utilization. Films of chitosan-nylon-4 blends showed good mechanical properties and retained the excellent chelating ability of chitosan.⁶ Blends of chitosan with strongly crystalline polyamides (nylon-4 and nylon-6) were also investigated. Their results suggested the partial miscibility of chitosan with nylon-4 and its lack of miscibility in other polyamides.⁷ Dureesne A⁸ synthesized blends of chitosan with PA6 via the phase inversion technique using formic acid as the common solvent. Due to the different densities of PA6 and CS, the CS phase tended to sediment and to form a continuous phase on the lower face of the film. The PA6 rich domains filled the upper face of the sample. A similar observation was reported for thermoplastic nanocomposites filled with wheat straw cellulose whiskers,⁹ and the behavior of these materials was modeled by subdividing the sample into layers with different whisker contents, lying parallel to the film surface. This sedimentation phenomenon was of great importance to the mechanical behavior of the blend.

The incorporation of a small amount of inorganic nanofillers during the formation of a polymer membrane can endow the polymer with new and much more improved mechanical, thermal and electrical properties.^10,11 Recently, the migration behavior of nanofillers to the surface of polymer nanocomposites (PNCs) has attracted interest. This phenomenon has been reported to be greatly correlated to the preparation method of PNCs, especially during the phase inversion procedure.^12,13 The migration of POSS¹⁴ and MMT¹⁵ was reported in PA6,¹⁴ polypropylene¹⁶ and polystyrene¹⁷ synthesized by melt blending. The decreased surface free energy was considered to be the critical issue for migration.^16,17 Our group studied the migration behaviors of POSS in POSS/PA6 nanocomposites during the phase inversion process.¹⁸ It was found that octaammonium POSS (OA-POSS) had the greatest migration behavior compared with other functional POSS. We suggested that the strong interaction between the nanofiller and the polymer matrix was another important issue for the migration behavior of the nanofiller, including the surface free energy. Increasing the interaction between the nanofiller and the polymer matrix could hinder the aggregation of the filler, which would improve the migration behavior of the filler. However, it should be mentioned that there are no publications focused on the migration of nanofillers in polymer blends with a sedimentation structure, neither on their performance thereof.

In this work, blends of CS with PA6 were prepared via the phase inversion process. Four kinds of nanofillers were introduced into the polymer matrix, including POSS, polymer grade montmorillonite (PGN), bis(2-hydroxyethyl)methyl tallow ammonium modified MMT (TCN) and 12-aminododecyl acid modified MMT (TL). The phase structure and morphology of the chitosan/PA6 nanocomposites were studied. In particular, the migration behaviors of the nanofillers in the PA6/CS blends are characterized. Our aim is to know the influence of the migration of the nanofiller on the mechanical properties of polymer blends.

2. Experimental

2.1 Materials

PA6 pellets (M_w = 30 000 g mol⁻¹) and CS ((C₆H₁₁NO₄)N) were purchased from Sigma-Aldrich. The samples were dried in a vacuum oven for more than 24 h at 80 °C and then stored in a desiccator before use. Octaammonium POSS (POSS), PGN: polymer grade montmorillonite, TCN: bis(2-hydroxyethy)methyl tallow ammonium modified MMT, TL: 12-aminododecyl acid modified MMT were purchased from Hybrid Plastics Company. The density of PA6, chitosan, POSS and PGN is 0.89 g cm⁻³, 1.75 g cm⁻³, 1.40 g cm⁻³ and 2.5 g cm⁻³, respectively. Fig. 1 shows the structure of POSS and PGN. All the nanofillers were dried in a vacuum oven for more than 24 h at 60 °C separately, and then stored in a desiccator. Formic acid was purchased from Sigma-Aldrich with a purity of 99.5% and was used as received.

Fig. 1 Microstructure of OA-POSS (a) and PGN (b).

2.2 Preparation of PA6/chitosan nanocomposites

The nanocomposites were synthesized by a solution-casting method. First, PA6 and chitosan were dissolved in formic acid (1 wt%) for at least 24 h (1.0 g of PA6 in 20 ml of formic acid). The nanofillers were then added to the solution, which was stirred for at least 48 h. The appropriate amount of POSS was measured in order to form polymer nanocomposites with the final compositions of 2 wt%. After mixing, the mixture was transferred into an ultra-plate watch-glass and maintained at 40 °C. The glass substrates with the sample films were left in the fume hood overnight to allow the formic acid solvent to evaporate. All the films were then dried in a vacuum oven at 80 °C for more than 12 h and subsequently kept in a desiccator before use.

2.3 Characterization

2.3.1 X-ray diffraction measurements (XRD). XRD measurements were carried out on a Bruker GADDS diffractometer with an area detector operating under 40 kV and 40 mA, using Cu Kα radiation (λ = 0.154 nm).

2.3.2 Polarized optical microscopy (POM). Surface morphology was analyzed by polarized optical microscopy (POM, BX51, OLYMPUS) at room temperature. The lens magnification was 10×.

2.3.3 Scanning electron microscopy (SEM) and energy dispersive X-ray (EDX). The morphology of the surfaces and cross sections of the films were analyzed by a HITACHI S-4800 scanning electron microscope using an acceleration voltage of 10 keV. Samples were put into liquid nitrogen for several minutes, and were then snapped into half to obtain the cross section. EDX analysis was also conducted on the same instrument using the EDX apparatus.

2.3.4 Transmission election microscopy (TEM). TEM (Tecnai F20, USA) was used to measure the morphology of the samples. Samples were cut by a freezing microtome before analysis by TEM.

2.3.5 Differential scanning calorimetry (DSC). The heat enthalpy, melting point and glass transition temperature of PA6 and its nanocomposites were measured by a differential scanning calorimeter (Paris 1, PerkinElmer Co., Norwalk, CT) in a nitrogen atmosphere. The samples were analyzed by the following temperature program: heating from 50 °C to 260 °C at 10 °C min⁻¹; maintaining for 3 min at 260 °C; cooling to 50 °C at 10 °C min⁻¹; maintaining at 50 °C for another 3 min before heating up again to 260 °C at 10 °C min⁻¹.

2.3.6 Tensile test. The tensile properties of the membranes were measured by an Instron model 5543 single column material testing system. Films with a gauge length of 40 mm, and width of 10 mm were stretched at a crosshead speed of 2 mm min⁻¹ at 20 °C. The resulting data were taken from the average of at least five samples.

2.3.7 Attenuated total reflectance infrared resonance (ATR). FTIR-ATR spectroscopy was performed on a Nicolet Nexus spectrometer. The detector is a liquid nitrogen cooled mercury cadmium telluride (MCT) with a range of 650–4000 cm⁻¹. The resolution was 4 cm⁻¹.

3. Results and discussions

3.1 Dispersion state of nanofillers in PA6/CS matrix

The miscibility of PA6 and CS was greatly changed due to their different crystalline behaviors. Fig. 2 shows that the PA6/CS blends with 50 wt% of PA6 can achieve a totally transparent morphology, indicating that a uniform dispersion of PA6/CS blends can be achieved. This can be attributed to the improved co-crystallization behaviors between PA6 and CS. Hence, the PA6/CS nanofiller nanocomposites were synthesized based on this weight ratio.

Fig. 2 POM images for PA6/CS blends. Weight ratio of PA6 to CS: (a) 50 : 50; (b) 75 : 35; (c) 95 : 5. The magnification is 10 times.

POM is an effective way to observe the dispersion state of the nanofiller in a polymer matrix.¹⁹ It is observed that the PA6/CS–POSS nanocomposites have a homogeneous surface morphology, as shown in Fig. 3b. This indicates that the POSS particles are uniformly dispersed in the polymer matrix. Furthermore, the presence of large black domains, which are indicative of nanofiller aggregates, can be clearly observed on the surface of the PA6/CS–PGN, PA6/CS–TCN and PA6/CS–TL nanocomposites (Fig. 3c–e). It can be attributed to the poor miscibility between the clay and the polymer matrix.

Fig. 3 POM images for (a) PA6/CS blends and PA6/CS with 2 wt% nanofiller: (b) PA6/CS–POSS; (c) PA6/CS–PGN; (d) PA6/CS–TCN; (e) PA6/CS–TL. The magnification is 10 times.

Fig. 4 shows the XRD spectra of PA6/CS and its nanocomposites with different nanofillers. It is shown that two main peaks are observed at 20.5° and 24°, indicating that the morphology of the films is mainly α-crystals in all of the nanocomposites.²⁰ Moreover, the peak ranging from 4° to 8° is generated by the aggregation of the nanofillers. In the PA6/CS–POSS nanocomposites, there is no obvious peak in this region, indicating the POSS is exfoliated. It may be due to the similar solubility parameters between POSS and the polymer matrix.¹⁸ However, the peak ranging from 4° to 8° can be found in the PA6/CS–PGN nanocomposites, suggesting that the aggregation of PGN is notable. Moreover, the peak ranging from 4° to 8° is difficult to be seen in the PA6/CS–TCN and PA6/CS–TL nanocomposites. It indicates that TCN and TL can achieve a good dispersion state in the polymer matrix.^15,18 This may be due to the fact that the modification of PGN by ammonium and 12-aminododecyl acid increases the interaction between the clay and the polymer matrix, through enhanced interface interactions.

Fig. 4 XRD spectra of PA6/CS and its nanocomposites with different nanofillers.

TEM images were further used to investigate the dispersion state of the nanofiller in the polymer matrix (Fig. 5). It is shown that POSS achieves the best dispersion state in the PA6/CS matrix with only several tens of nanometers. This shows great consistency with that of the XRD results. However, the PGN aggregates in the polymer matrix are larger in size. This is the reason for the new diffraction peak generated from 4° to 8° in the XRD spectra. After chemical modification of PGN, the introduction of amino, hydroxyl and acid groups increases the interface interaction between the nanoparticles (TCN and TL) and polymer matrix. This will increase the miscibility between the nanofiller and the polymer matrix. Thus, the nanofiller aggregates will decrease in size (Table 1).

Fig. 5 TEM images for polymer nanocomposites: (a) PA6/CS–POSS; (b) PA6/CS–PGN; (c) PA6/CS–TCN; (d) PA6/CS–TL.

Table 1 DSC results of PA6/CS and PA6/CS nanocomposites

	Tm^1 °C	Tc^1 °C	ΔHm^1 J g^−1	ΔHc^1 J g^−1
PA6/CS	220.4	183.7	28.7	25.8
PA6/CS–POSS	221.3	187.6	30.4	28.3
PA6/CS–PGN	221.0	187.4	24.1	22.4
PA6/CS–TCN	221.1	186.0	26.4	26.3
PA6/CS–TL	221.2	185.9	27.1	18.7

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3.2. DSC characterization

Fig. 6 shows the effects of the nanofiller additives on the crystallization behavior of the PA6/CS membranes, measured by DSC. A broad endotherm is observed around 90 °C, which is the typical melting point of CS.^21,22 In addition, all the samples have obvious melting points around 220 °C, which are mainly generated by the α form crystals of PA6.²¹ The melting peak of the CS phase changes indistinctly with the incorporation of PGN compared with the PA6/CS blends. It can be attributed to the poor dispersion state of PGN in the polymer matrix. However, a gradual increase of the melting peak of CS is seen with the incorporation of POSS, TCN and TL, indicating that the intermolecular hydrogen bond interaction between the nanoparticles of CS increased in the order of POSS, TCN and TL.^22,23 Table 2 shows the DSC data of PA6 phase. The variance of the melting enthalpy upon the incorporation of nanoparticles is not obvious. Thus, it is difficult to discuss the influence of the crystalline behavior of the nanocomposites generated by the incorporation of nanoparticles. However, the melting peaks and crystallization peaks (Fig. 6a and b) of the nanocomposites slightly increased in the PA6 phases when the nanoparticles were incorporated. It may be due to the mildly enhanced interface interactions between the PA6 phase and the nanoparticles.

Fig. 6 (a) DSC thermograms of PA6/CS nanocomposites during first heating scan, (b) DSC thermograms of PA6/CS nanocomposites during cooling scan.

Table 2 Silicon atom contents of PA6/CS and PA6/CS nanocomposites

Samples	Cross-section upper zone	Cross-section middle zone	Cross-section lower zone	Top side	Bottom side
PA6/CS–POSS	0.68	0.37	0.04	0.26	0.12
PA6/CS–PGN	0.08	0.10	0.22	0.08	0.23
PA6/CS–TCN	0.15	0.11	0.33	0.26	0.38
PA6/CS–TL	0.10	0.15	0.49	0.19	0.45

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3.3 Tensile properties

Fig. 7 shows the tensile strength (Fig. 7a) and breaking strains (Fig. 7b) of the polymer nanocomposites. It is clearly observed that CS has higher tensile strength and breaking strains than that of PA6. The tensile strength of the PA6/CS nanocomposites is in between that of PA6 and CS. The inter-hydrogen bonds between chitosan and the polyamide make the molecules bind tightly, so that its tensile strength increases compared with that of PA6.²⁴ Interestingly, the breaking strains of PA6/CS obviously decrease. Further incorporation of nanoparticles (POSS, PGN, TGN and TL) leads to an increment of tensile strength and breaking strains compared with that of pure PA6/CS. The incorporation of nanoparticles into the PA6/CS matrix increases the interface interactions between the nanoparticles and the polymer matrix, resulting in an enhancement of its stress transfer property. In addition, the exfoliated nanoparticles can also transfer stress during the tensile process. Thus, the tensile property of the nanocomposites improved compared with that of the PA6/CS blends.

Fig. 7 Tensile strength (a) and breaking strains (b) of CS, PA6, PA6/CS and PA6/CS nanocomposites. Data are expressed as the mean ± SD, n = 5.

However, the tensile strength and breaking strains decrease in the order of PA6/CS–POSS, PA6/CS–PGN, PA6/CS–TGN and PA6/CS–TL. It was shown by experiments and model studies that the high modulus and high aspect ratio of nanofillers could achieve a good mechanical performance in pure PA6/nanofiller nanocomposites. In this work, PGN, TCN and TL are typical nanofillers with high moduli and high aspect ratios. The reinforcing efficiency should increase if an exfoliated dispersion state can be achieved.²⁵ However, our results seem to be contradictory with the previous observation of a dispersion state and DSC results. It should be mentioned that the above study on the dispersion state of the nanofillers (TEM, XRD and POM) is mainly characterized from the bulk of the polymer matrix. The morphology of the polymer films, especially along the thick direction, will be further studied.

3.4 Morphology studies

3.4.1 Dispersion of CS in the polymer matrix. Fig. 8 shows the SEM images of the PA6/CS and PA6/CS nanocomposites on the top and bottom side. The top side (Fig. 8a) of PA6/CS appears as an undulating surface because of the evaporation process.⁸ In contrast, the bottom face is scored (Fig. 8a′), probably because of the surface effect of the box in which the blend solution is cast. The top face was still undulated with the incorporation of PGN. However, the top face gradually turns smooth in the order of PA6/CS–TCN, PA6/CS–TL and PA6/CS–POSS. This may indicate that the increased interaction between solvent, nanofiller and polymer matrix can influence the phase separation procedure.

Fig. 8 SEM images of PA6/CS and PA6/CS nanocomposites. (x) denotes the SEM images of the top side; (x′) assigns the SEM images of the bottom side. (a) PA6/CS; (b) PA6/CS–PGN; (c) PA6/CS–TCN; (d) PA6/CS–TL; (e) PA6/CS–POSS.

Fig. 9 shows the cross-section images of the PA6/CS and PA6/CS nanocomposites. It can be found that the polymer matrix is rather condensed and nonporous compared to the images of the top and bottom side. Interestingly, the morphologies are different across the red line as shown in the image. The upper zone looks transversal, while the lower zone seems longitudinal. This is an indication that CS and PA6 in the blend form mainly a two-phase system. Dufresne A⁸ synthesized blends of chitosan with PA6 via the phase inversion process. They found that the CS phase tended to sediment and to form a continuous phase on the lower face of the film, whereas PA6 rich domains filled the upper face of the sample. Thus, the different morphologies across the red line may be due to the different polymer compositions.

Fig. 9 SEM images of the cross-section of PA6/CS and PA6/CS nanocomposites. (a) PA6/CS; (b) PA6/CS–PGN; (c) PA6/CS–TCN; (d) PA6/CS–TL; (e) PA6/CS–POSS.

One of the main characteristics of PA or CS is the strong inter-chain interaction that arises from the hydrogen bonding between the amide groups on adjacent chains. Generally, some nanofillers can form hydrogen bonds with amide groups at interfaces, which can be clearly characterized by FTIR-ATR.²² Fig. 10 shows that the FTIR-ATR spectra of the samples on the top side and bottom side are quite different. In Fig. 10a, the FTIR-ATR spectra of the polymer on top side have more PA6 signals,²⁶ where the characteristic assignments of PA6 are 3300 cm⁻¹ (sharp peak, H-bonded N–H stretch vibration), 3080 cm⁻¹ (N–H in-plane bending), 1640 cm⁻¹ (amide I, C=O stretch), 1540 cm⁻¹ (amide II, C–N stretch and CO–N–H bend). In Fig. 10b, the ATR spectra of the polymer on the bottom side are similar to CS.^27,28 The absorption that occurs around 898 and 1150 cm⁻¹ are assigned to the saccharine structure in CS. The bands at 1060 and 1029 cm⁻¹, which involve C=O stretching belongs to CS as well. These indicate that the sedimentation of CS exist in the synthesized polymers. The lower face of the film seems to present CS rich domains. This gradient of CS concentration is probably induced by the processing technique itself. The CS domains tend to sediment during the solvent evaporation step. This phenomenon is due to the difference between the density of CS (ρ ≈ 1.75) and that of PA6 (ρ ≈ 0.89).⁸

Fig. 10 FTIR-ATR spectra of PA6/CS and PA6/CS nanocomposites. (a) Top side; (b) bottom side.

In addition, the FTIR-ATR spectra of the PA6 rich phase and nanocomposites in the hydrogen-bonded N–H stretch region (3100–3400 cm⁻¹) show little variance in the peak width. It indicates that the interface interaction between the nanoparticles and the PA6 rich phase is not strong.²² However, the peak width of the hydrogen-bonded N–H stretch region in the CS rich phase obviously increased, indicating a strong interface interaction between the nanoparticles and CS rich phase. This shows great consistency with the DSC results.

3.4.2 Migration of nanofillers in the polymer matrix. Table 2 summaries the silicon atom content of the polymer and the polymer nanocomposites with different nanofillers, qualitatively measured by EDX.¹⁸ Typically selected zones for EDX are shown in Fig. 11. One can image that the measured silicon atom content should be close to each other in all parts, if the dispersion of the nanofillers is uniform in the three-dimensions of the polymer matrix. However, these contents, measured by EDX, are obviously different as shown in Table 2. It can be found that the content of the silicon atom is much higher on the cross-section upper zone and top side than the lower zone for PA6/CS–POSS. It indicates that POSS prefers to migrate to the top surface of the membrane during the phase inversion procedure.¹⁸ This could be due to the lower surface free energy and strong interactions between POSS and the PA6 matrix.^15,18 However, PGN, TCN and TL show the opposite migrating direction, where a much higher silicon concentration can be measured on the lower zone. In other words, the PGN, TCN and TL tend to migrate to the bottom side. This could be due to the large density of PGN, TCN and TL (≈2.5 g cm⁻³). The densities of PGN, TCN and TL are larger than that of PA6 and CS. As a result, PGN, TCN and TL are prone to sedimentation.

Fig. 11 Typical selected zone for EDX (The sample is PA6/CS–POSS nanocomposites). (a) Cross-section upper zone; (b) cross-section middle zone; (c) cross-section lower zone; (d) top side; (e) bottom side.

3.5 The influence of tensile strength generated by the migration of nanofiller

Fig. 12 shows the proposed mechanism of the influence of migration of the nanofiller on the tensile property of the composite. From the morphology studies, CS domains tend to sediment during the solvent evaporation step. PA6 rich domains fill the upper zone of the samples. This heterogenous distribution state makes the sample easier to break when a symmetric force is loaded. As a result, the decrease in the variance of tensile strength between the PA6 rich domains and the CS rich domains is very important to improve the tensile strength of the PA/CS blends. The POSS is prone to migrate to the PA6 rich phase with a good dispersion state. Thus, more nanoparticles can be used to transfer the stress in the PA6 rich phase during the tensile measurement. These will increase the tensile strength and breaking strains of the PA6 rich domains. The variance of the mechanical properties between the PA6 rich domains and the CS rich domains decreases thereof. Thus, the tensile strength and breaking strains of the PA6/CS–POSS nanocomposites are higher than that of the PA6/CS blends. For the PA6/CS–clay systems, the mechanical properties of the nanocomposites can be improved due to the good dispersion state of the clay in the PA6/CS matrix. However, PGN, TCN and TL can achieve more loading in the CS rich domains. In addition, chemical modification of PGN increases the interface interaction between the clay and the CS rich phase, resulting in an increment of the tensile strength of the CS rich domains. This will enlarge the variance of the mechanical properties between the PA6 rich domains and the CS rich domains. Thus, the tensile strength and breaking strains gradually decreases in the order of PA6/CS–PGN, PA6/CS–TGN and PA6/CS–TL.

Fig. 12 Proposed mechanism of the influence of migration of the nanofiller to the tensile strength.

4. Conclusions

PA6/CS blended with POSS, PGN, TCN and TL nanocomposites were synthesized according to the phase inversion procedure. TEM, POM and XRD were used to characterize the dispersion state of nanofillers in the polymer matrix. The results show that POSS can achieve the best dispersion state in the polymer matrix with only several tens of nanometers. Modification of PGN (TCN and TL) can improve the dispersion state of the nanofillers. However, the dispersion state results cannot well explain the results of tensile strength. Further study on the morphologies of the polymer along the thick direction shows that CS domains tend to sediment during the solvent evaporation step. PA6 rich domains fill the upper zone of the samples. POSS is prone to migrate to the upper zone. The strong interaction between POSS and PA6 increases the tensile strength of the PA6 rich domains. This makes the variance of mechanical properties between the PA6 rich domains and the CS rich domains decrease, resulting in the improved tensile strength of the polymer matrix entirety. PGN, TCN and TL sediment to the CS rich phase, which results in the increased tensile strength of the CS rich domains. This enlarges the variance of the mechanical properties between the PA6 rich domains and the CS rich domains. The tensile strength of the polymer matrix entirety decreased when stronger interactions between the nanofillers and the CS rich domains occurred.

Acknowledgements

Funding from the project of Natural Science Foundation of China (no. 51203081, 21206078), the Natural Science Foundation of Ningbo (2012A610087) and Ningbo Science and Technology Innovation Team (2011B2002) are gratefully acknowledged.

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