Double glass transitions in exfoliated poly(methyl methacrylate)/organically modified MgAl layered double hydroxide nanocomposites

Abstract

A series of nanocomposites based on poly(methyl methacrylate) (PMMA) and organically modified MgAl layered double hydroxides (O-LDH) were synthesized via in situ polymerization. The modification of LDH by sodium dodecylbenzenesulfonate (SDBS) resulted in an enlarged interlayer distance and a nearly complete substitution of the original NO₃⁻ in LDH layers. The polymerization of methyl methacrylate led to a disorderly exfoliated LDH in the PMMA matrix. The results of dynamic mechanical analysis (DMA) showed that there were two tan δ peaks in the PMMA/O-LDH nanocomposites at low frequencies. And the second peak at about 80 °C higher than the main peak (glass transition of PMMA matrix) gradually disappeared with increasing frequency, corresponding to the glass transition of PMMA chains confined on the O-LDH surface. Both of the two peaks shifted to higher temperature with increasing heating rate due to the strain lag of the sample.

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1. Introduction

Polymer/inorganic nanocomposites have attracted extensive attention for their remarkable or often unique properties due to nanometer effects of inorganic fillers and their homogeneous dispersion in the polymeric matrix.^1,38,40 The glass transition temperature (T_g), a very important technological parameter for polymer properties, is the temperature at which the polymers transform from the frozen glassy state into a liquid. However, the introduction of inorganic fillers into a polymer matrix would result in complex T_g dynamics. Because of various interfacial interactions or confinement effects of inorganic fillers, the glass transition temperature of polymer nanocomposites has been reported to increase,^2,3 decrease,^4,5 or remain the same.^6,7 Double glass transitions were identified by Tsagaropoulos and Eisenberg⁸ by using Dynamic Mechanical Analyses (DMA) to investigate several polymers filled with silica particles. Similarly, Li et al.⁹ have reported two distinct T_gs after slow cooling PMMA melt in the nanopores and fast cooling would only result in a single glass transition for PMMA. However, the second glass transition observed by Tsagaropoulos and Eisenberg has caused great controversy. Salaniwal et al.¹⁰ suggested that the higher temperature loss tangent peak may be due to polymer-filler “gel” with long structural lifetime. For Robertson and Rackaitis, it was assigned to the terminal relaxation process of un-cross-linked polymer restricted by particles.⁶ Unavoidablely, there is an ongoing debate on whether and how inorganic fillers influence the glass transition temperature of a polymer till now, the glass transition dynamics mechanism of a polymer under the confinement of different nanocomposites still deserves more attention. As generally accepted, both the size and geometrical structure of inorganic fillers are considered of great significance for the glass transition behavior.⁹ Previous researchers have well studied the influence of various shapes of inorganic fillers on the glass transition behavior of the filled polymers, such as spherical silica,^8,11,12,39 cuboid nanochannal,¹³ nanopores,^9,14,15 and nanoglobules.³⁷ The purpose of our work is to use a new type of nanofiller with layered structure like layered silicate clays, layered double hydroxides (LDHs), to investigate the T_g behaviour of PMMA/O-LDH nanocomposites and its impact on the thermal properties.

LDHs are well studied anion exchanging materials with excellent physical and performance properties.^16–21 They can be modified by the exchange of interlayer anions through introducing organic anion. This transforms hydrophilic LDH nanolayers into more hydrophobic and enlarges the interlayer distance, so that monomer or polymer chains can easily intercalate into the LDH galleries and homogeneously dispersed PMMA/O-LDH nanocomposites can be prepared.

In our work, the structures of LDHs and PMMA/O-LDH nanocomposites were well investigated by infrared spectrum (FI-TR), X-ray diffraction (XRD) and transmission electron microscopy (TEM). DMA was used to investigate the mobility of polymer chains restricted by LDH and the influence of these restrictions on the glass transition behavior of LDH particles filled polymer. We observed two glass transitions in the PMMA nanocomposites. DSC were applied to characterize thermal glass transition temperature of PMMA influenced by the addition of LDH particles.

2. Experimental section

2.1. Materials

Methyl methacrylate (MMA), benzoyl peroxide (BPO) were purchased from Shanghai Titan Scientific Co. Ltd (Shanghai, China). MMA was washed with NaOH aqueous solution (5 wt%) for several times to remove inhibitors, followed by deionized water until neutralization, and then dried by anhydrous sodium sulfate to remove the residual water before use. The metal nitrate salts Mg(NO₃)₂·6H₂O, Al(NO₃)₃·9H₂O and SDBS purchased from Sinopharm were used for the synthesis of organic MgAl-LDH.

2.2. Preparation of organo-modified MgAl-LDH

The synthesis of O-LDH was carried out by co-precipitation. Experimentally, 0.04 mol Mg(NO₃)₂·6H₂O and 0.02 mol Al(NO₃)₃·9H₂O were dissolved in 50 mL deionized water respectively. The two solutions were slowly added to a SDBS (0.02 mol) solution in a flask (500 mL) with a reflux condenser under continuous stirring at 40 °C. Then, 6% (volume fraction) ammonia was dropwise added to the mixture until white precipitate appeared. After that, added 30 mL 25% (volume fraction) ammonia in 2–3 second and maintained PH value at 9–10. The resulting slurry was continuously stirred at 40 °C for 30 min and aged at room temperature for 1 hour. Filtered and washed the products several times with deionized water until the solution pH was about 7. The final products were redissolved in 1000 mL deionized water and aged at 80 °C in a water bath for 8 hours, then dried at 80 °C until a constant weight. For comparison, the unmodified MgAl-LDH (U-LDH) was also synthesized without the addition of SDBS.

2.3. Preparation of PMMA/O-LDH nanocomposites

A wide-mouth reactor with a condenser was charged with 35 mL of MMA, 0.045 g of BPO and different amounts of O-LDH. The mixture was stirred and heated up to 85 °C under a N₂ atmosphere until the viscous PMMA/O-LDH solution was obtained and then heated up to 100 °C for 3 hours to complete the reaction. The concentrations of O-LDH in the prepared PMMA/O-LDH nanocomposites are 0, 1.9, 4.1, 7.6 wt% and 12.5 wt% and they are labeled as PMMA0, PMMA1.9, PMMA4.1, PMMA7.6 and PMMA12.5 respectively.

2.4. Characterization techniques

The morphology of samples was observed under a JEM-2010 transmission electron microscopy (JEOL, Japan) at room temperature, 120 kV acceleration voltage and bright field illumination.

The particle size distribution and average particle size of U-LDH and O-LDH were analyzed using Mastersizer (Malvern Instrument, UK) fitted with Scirocco 2000 unit.

DSC data were obtained from a Pyris Diamond DSC (PE, USA), and all measurements were made at a scan rate of 20 °C min⁻¹ under continuous nitrogen gas flow.

The FT-IR spectrum was recorded using Perkin Elmer Fourier transform infrared spectroscope. The measurements were taken at the room temperature in the wave region ranging from 3200 to 400 cm⁻¹.

DMA was used to measure the mechanical properties of the samples as a function of temperature. The experiments were performed using a Perkin-Elmer Diamond DMA instrument in tension mode at the heating rate of 1 °C min⁻¹ under a dry nitrogen atmosphere under a dry nitrogen atmosphere. The preparation process of the specimen for DMA was as follow: the polymer samples were added into a prepared mould (the size of mould is 40 mm × 5 mm × 1.5 mm) and then cured through the compression molding in a platen press under 5 MPa pressure at the temperature of 190 °C. After that, the samples were cooled under 5 MPa pressure at room temperature.

3. Results and discussion

3.1. Characterization of LDHs and its nanocomposites

Fig. 1 shows TEM images of U-LDH and O-LDH particles. The nanosized LDH nanoparticle are obtained before and after modification, displaying the typical plate-like morphology. Both of the LDH sheets are homogeneously distributed and exhibit irregular hexagon. After modification, the lamellar structure size slightly increases, which is further confirmed by the mastersizer analysis, as shown in Fig. 2.

Fig. 1 TEM images of (a) U-LDH and (b) O-LDH.

Fig. 2 Particle size distribution of U-LDH and O-LDH. The U-LDH and O-LDH powders are dispersed in deionized water before measuring.

Fig. 3 shows the FT-IR spectra of SDBS, U-LDH and O-LDH. The new peak appears at 2960 cm⁻¹ in O-LDH as compared to U-LDH attributing to asymmetric stretching vibration of –CH₃ of the long alkyl chain of SDBS anions. The new peaks at 2921 cm⁻¹ and 2857 cm⁻¹ correspond to asymmetric and symmetric stretching vibration of –CH₂ of the long alkyl chain of SDBS respectively. The weak peak at 1462 cm⁻¹ is associated with in-plane rocking vibration peak of –CH₃ and –CH₂ of SDBS. The absorption peaks at 1188, 1133, 1044, 1011 cm⁻¹ and 690 cm⁻¹ also originate from the modifying surfactant. The broad peak at 660 cm⁻¹ and sharp peak at 447 cm⁻¹ are associated with stretching vibration ν(M–O) and deformation vibration δ(O–M–O) originating from in the LDH sheets.²² Moreover, the absorption peak at 1383 cm⁻¹ correspond to NO₃⁻ in the U-LDH,²³ however, in O-LDH, this peak is very weak. These FT-IR results suggest that SDBS anions enter the LDH gallery and completely substitute the original anions NO₃⁻. This would result in an enlarged interlayer distance as discussed in the following part.

Fig. 3 IR spectra of SDBS, U-LDH and O-LDH.

Fig. 4 presents the XRD patterns of U-LDH, O-LDH and the polymer samples. The XRD pattern of U-LDH shows the (003) reflection of U-LDH at 2θ = 10.24° which is characteristic for a layered compound corresponding to a lamellar repeat distance of d₍₀₀₃₎ = 0.86 nm and indicative of long-range ordering in the stacking dimension. Subtraction of the thickness of the brucite-like LDH sheet (0.48 nm)²⁴ gives the effective interlayer distance to be 0.40 nm. Previous researches^25,26 have revealed that the interlayer distance of LDH depends on the length of interlayer anion and the charge density of the layered LDH sheets. The two factors determine the configuration and packing of the interlayer anions in the intergallery and, therefore, the interlayer distance. For SDBS, the whole molecular length of SDBS is 2.18 nm,²⁷ obviously, much bigger than that of NO₃⁻. Therefore, the modification of the LDH results in a shift of the lamellar reflection toward lower 2θ as expected and the position of (003) plane reflection in O-LDH is located at 2θ = 3.40° corresponding to d₍₀₀₃₎ = 2.60 nm. The calculated interlayer distance is 2.12 nm. However, in PMMA/O-LDH nanocomposites, the LDH characteristic reflections disappear thoroughly and the broad peaks centered on 2θ = 14.6°, 30.0° and 41.7° indicate amorphous nature of the PMMA matrix.^28,41 This implies that the regular periodicity of O-LDH in nanocomposites has been broken by the kinetic driving force of the polymerization of MMA.²⁹ On the other hand, the molecular chains of PMMA could have been effectively intercalated and well dispersed into the interlayer of the O-LDH. The TEM equipment provides the straightforward morphologies of nanocomposites, as shown in Fig. 5a, O-LDH particles are homogeneously distributed in the PMMA matrix as stacks. Fig. 5b shows the TEM image of PMMA12.5 at high magnification, dark lines in the micrograph represent the LDH layers and the LDH nanolayers comprised mainly of highly exfoliated structures lose their ordered stacking structure, which is consistent with the X-ray diffraction analysis. Generally speaking, stacking Mg/Al nanolayers are not easily exfoliated but just intercalated in nanocomposites at high filler content and the LDH characteristic reflections would be reflected in the XRD curves,^3,30 Therefore, the disappearance of the LDH characteristic reflections in all the nanocomposites indicates that the preparation of the exfoliated PMMA/O-LDH nanocomposites is performed successfully via in situ free radical polymerization.

Fig. 4 XRD patterns of U-LDH, O-LDH, PMMA0, PMMA1.9, PMMA4.1, PMMA7.6, PMMA12.5.

Fig. 5 TEM images of PMMA12.5 at low (a) and high magnification (b).

3.2. Dynamic mechanical behavior of pure PMMA and its nanocomposites

Fig. 6 shows the tan δ vs. temperature curves for pure PMMA and PMMA/O-LDH nanocomposites with different O-LDH contents at the frequency of 5 HZ. tan δ is the ratio of loss modulus (E′′) to storage modulus (E′) and its peak maximum marks the location of T_g in a system. The storage modulus (E′) correlates with the elastic modulus of the materials, and the loss modulus (E′′) is related with the energy lost due to the friction of polymer chain movement. The relaxation peak with a maximum at about 124 °C corresponds to the glass transition of pure PMMA and the T_g shifts to higher temperature with the growth of O-LDH content due to more severe restriction of nanoparticles. However, a slight reduction in T_gonset could be observed for PMMA7.6 within the experiment error. What's more, the obvious changes of the tan δ curves of the nanocomposites are the reduction in tan δ peak area as well as the broadening of the peaks, indicating that the fraction of chains participating in the glass transition decreases as the composite becomes richer in O-LDH particles due to the nanofiller-induced confinement effects. These confined polymer chains around O-LDH particles surface don't take part in the glass transition, which forms the interfacial layer between nanoparticles and polymer matrix. Combined with the TEM results (Fig. 5), a schematic of the confined polymer in nanocomposites is illustrated in Fig. 7. The red part represents the interfacial layer which are restricted by LDH particles and the white part, unconfined polymer matrix which can experience glass transition freely.

Fig. 6 Temperature-dependent mechanical relaxation spectra (tan δ) of pure PMMA and its nanocomposites.

Fig. 7 Schematic of PMMA chains confined by O-LDH layers in nanocomposites.

Fig. 8 shows the tan δ vs. temperature curves of all the samples at different frequency. The main peak due to glass transition of PMMA matrix without restrictions shifts to higher temperature with the increase of frequency because of the strain lag of the sample.¹² No peaks could be observed in the five tan δ curves of pure PMMA (Fig. 8a) at higher temperature side of the main peak. But, it is interesting that for nanocomposites, another peak with an increased intensity appears at the temperature about 80 °C higher than the main peak and the peak gradually disappears with the increase of the frequency. This suggests that the second peak is probably associated with motions of the PMMA chains confined by the O-LDH particles surface. As discussed above in Fig. 6, a fraction of PMMA chains of nanocomposites in the interfacial layer didn't experience the relaxation process at the frequency of 5 Hz. For polymer/inorganic nanocomposites, segmental relaxation times for the chains close to the filler surface are significantly longer than those of the unconfined polymer.^11,12 Therefore, as the frequency decreased, the confined polymer chains would have enough time to respond to the probe of DMA, showing a second glass transition peak in all the composites.

Fig. 8 tan δ vs. temperature curves of (a) pure PMMA; (b) PMMA1.9; (c) PMMA4.1; (d) PMMA7.6; (e) PMMA12.5 at different frequency. The heating rate is 1 °C min⁻¹.

The heating rate associated with the relaxation time of polymer chains also influences glass transition behaviour of polymers. Fig. 9 shows the tan δ vs. temperature curves of PMMA12.5 at different heating rate. With the increase of the heating rate, the two glass transition temperatures shift to higher temperature. According to the free volume theory, at the beginning of the glass transition, it needs to reach a certain temperature or higher and provide heat energy for polymers to expand the free volume until the free volume is large enough for polymer molecules to rotate and move, then adjust their conformation. As the heating rate increases, the free volume expansion lags behind the increasing heat flow and is not sufficient for polymer chains to move at the glass transition temperature of lower heating rate, therefore, the glass transition must be realized at a higher temperature. However, the second tan δ peak, firstly reported in inorganic/polymer nanocomposites by Tsagaropoulos and Eisenberg,⁸ attributed to the glass transition of polymer chains has caused a great controversy, which we have discussed in the introduction section. Representative findings from the investigations of Robertson and Rackaitis⁶ are shown in Fig. 10. In their work, the unfilled PB (U-PB) displayed terminal flow, with tan δ diverging toward infinity as the temperature increased. The introduction of carbon black (CB) led to a second tan δ peak at higher temperature and a similar tan δ peak can be produced in partial cross-linked U-PB. It is interesting that this peak disappeared after complete cross-linking of the polybutadiene (Fig. 10a and b). Therefore, the second tan δ peak was suggested to be the suppression of flow relaxation of polymer chains for the interaction with the particles rather than glass transition of the polymer confined to the filler surfaces. However, Robertson and Rackaitis' work⁶ didn't provide the direct evidence that the confined polymer chains by fillers surface can't cause a distinct, higher temperature glass transition. And recently, Li⁹ et al. have observed a second glass transition for polymers under nanoconfinement by DSC. Fortunately, in our work, chain diffusion and flow at high temperature don't occur as shown in Fig. 8a, what's more, the addition of O-LDH induces another tan δ peak with an increased intensity. Therefore, we can reasonably identify the higher temperature tan δ peak in PMMA/O-LDH system as the glass transition response of the immobilized polymer layer.

Fig. 9 tan δ vs. temperature curves of PMMA12.5 at different heating rate.

Fig. 10 (a) tan δ vs. temperature curves of pure polybutadiene (PB), carbon black (CB) filled PB and completely cross-linked carbon black (CB) filled PB by using 2 wt% peroxide. (b) tan δ vs. temperature curves of pure PB, partially cross-linked PB and completely cross-linked PB. The figure originated from paper by Robertson and Rackaitis.⁶

3.3. Thermal glass transition of pure PMMA and its nanocomposites

Fig. 11 shows the DSC curves of pure PMMA and its nanocomposites in the glass transition region. The curves show a step-like increase with the increase of temperature, indicating the thermal glass transition. Since the sensitivity of DSC is lower than that of DMA, the double glass transitions are not found in the DSC curves.¹² However, due to the nanoparticle-confinement effects, polymer chains would have more difficulty in experiencing glass transition under the heating, this would be reflected by the increase of the glass transition temperature. To evaluate the thermal T_g, the standard methods is using tangents at the DSC curve in the glassy and liquid state and the T_g is defined as the temperature at which the measured heat flow curve equals the half distance between the tangents. The interaction of the tangents of the DSC curve at T_g and the glass state can be defined as T_gonset. Similarly, T_gend was evaluated, as is shown in Fig. 11. The results are listed in Table 1. It can be seen that the incorporation of O-LDH slightly increases the T_g of PMMA, indicating a restriction of the mobility of polymer chains. This restriction here demonstrated by DMA and DSC may be because hydroxyl groups of O-LDH could intensely interact with the ester functional group (COOCH₃) of the PMMA polymeric matrix in its side chain, and therefore hinder the motion of polymeric chains.^31–33,42 Many researchers have identified the confinement effects of inorganic fillers on polymer chains,^11,34–36 however, few revealed a double glass transitions behavior, nor did Robertson and Rackaitis.⁶ The DSC curves in Robertson and Rackaitis's work showed that the filler had nearly no impact on the glass transition of polymer even after solvent extraction of weakly interacted polymer chains. This would be ascribed to the strength of interfacial interaction between inorganic and organic phase. In addition, more confined polymer chains in polymer/inorganic filler nanocomposites can be obtained due to the intense interfacial interaction, which would lead to a significant change in the thermal glass transition and contribute to observe the second glass transition. However, the glass transition temperature values of the pure PMMA and its nanocomposites obtained via DMA and DSC are different, this discrepancy can be understood considering the different probes measured by both methods.

Fig. 11 Glass transitions (DSC) of the pure PMMA and its nanocomposites.

Table 1 DSC data for pure PMMA and its nanocomposites

Sample	DSC data
	Tgonset (°C)	Tg (°C)	Tgend (°C)
PMMA0	102.2	112.8	119.4
PMMA1.9	108.6	116.6	122.7
PMMA4.1	109.7	118.3	124.7
PMMA7.6	106.7	116.5	122.8
PMMA12.5	103.5	113.4	119.6

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4. Conclusions

Nanocomposites based on PMMA and nanosized O-LDH were prepared via in situ free radical polymerization with MMA as monomer and Benzoyl peroxide (BPO) as initiator. LDH was successfully modified by SDBS, leading to a complete substitution of the original interlayer anion and enlarged LDH interlayer distance. The polymerization of methyl methacrylate resulted in a disorderly exfoliated LDH which was homogeneously dispersed in PMMA matrix with stack structure. Double tan δ peaks were found in PMMA/O-LDH nanocomposites, corresponding to the glass transitions of PMMA matrix and the interfacial immobilized layer on the O-LDH surface for the effect of confinement respectively. The two glass transition peaks shifted to higher temperature as the heating rate increased due to the strain lag of the polymer samples. The confinement effects imposed by nanoparticles were further demonstrated by DSC which showed an increased glass transition temperature after the introduction of inorganic fillers into PMMA matrix.

Acknowledgements

The authors gratefully acknowledge the financial support from the National Natural Science Foundation of China (No. 21204090 and No. 51303182).

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