Functionalized TiO2 Nanoparticles Tune the Aggregation Structure

Article pubs.acs.org/JPCC

Functionalized TiO2 Nanoparticles Tune the Aggregation Structure and Trapping Property of Polyethylene Nanocomposites Simin Peng, Bin Dang, Yao Zhou, Jun Hu,* and Jinliang He* State Key Laboratory of Power Systems, Department of Electrical Engineering, Tsinghua University, Beijing 100084, China S Supporting Information *

ABSTRACT: The interfacial region between a nanoparticle and the polymer matrix is considered to have an important effect on the properties of nanocomposites. In this experimental study, the tuning effects of surface-modified TiO2 nanoparticles on the aggregation structure and trapping property of the nanocomposites are investigated. The addition of TiO2 nanoparticles increases the number of spherulites and decreases their sizes. It is also found that TiO2 nanoparticles can suppress the mobility of chain segments in the interfacial region, suppress crystallization, and reduce the crystallinity, depending on the surface modification and loading levels of nanoparticles. On the basis of the conclusions, the model of a spherulite and the interface is established and the trap distribution of the interface is analyzed according to the results of TSC measurements. It is assumed that there is a strong correlation between traps and the charge transport of nanocomposites, and the mechanism of charge transport is discussed with respect to the results of the volume conductivity measurement. It is believed that this study would provide an important hint to the research of the interface between nanoparticles and the polymer matrix in future research.

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INTRODUCTION

which makes the study of polymers far more complicated than that of metals and other crystalline materials. Many studies have tried to characterize the properties of the interface by detecting the macroscopic properties of nanocomposites;12−14 however, the influence of the interface microstructure is often neglected. We assume that the unique properties of the interface are due to its special microstructure, for example, many electric properties of nanodielectrics are closely related to the charge transport mechanism,15,16 which is deeply affected by the physical and chemical defects in the materials. In this experimental study, we focus on the microstructure of the interface. A series of nanocomposites were prepared by adding different weight fractions of different surface-modified TiO2 nanoparticles into LDPE via a melt blending method. The microstructure, differential scanning calorimetry experiment, 1H solid-state nuclear magnetic resonance, thermally stimulated currents, and dc volume conductivity of the nanocomposites were measured. Some new insights into the aggregation structure and trap effect were achieved through the understanding of the bound effect in the interface of nanocomposites.

Polymers filled with inorganic nanoparticles have drawn a great amount of attention in recent years because they can dramatically improve the properties of the original polymers,1−4 and it seems that polymer nanocomposites are able to solve current problems in many fields. The application of nanocomposites in electrical insulation is also called nanodielectrics.5,6 Many researchers believed that the interface, which is the nanoscale transition region between nanoparticles and the polymer matrix, plays an important role in modifying the properties of nanodielectrics.7,8 Some models of the interface have been proposed. For example, Tanaka9 proposed the multicore model of the interface by dividing the interface into three layers, named the bonded layer, the bound layer, and the loose layer, respectively. These models successfully explain some experimental results; however, the multicore structure of the interface described in the model has not been directly validated by experiments, and the mechanism of the interface has not been understood. On the one hand, the nanoscale exceeds the limit of spatial resolution of most conventional analytical measurements, and the direct detection of the interface is still considered to be a difficult thing. On the other hand, solid polymers usually have multilevel and multiscale structures.10,11 It is well known that the properties of polymer materials are directly influenced by the physical combination and stacking structure of their molecular chains or the aggregation structure of polymer,10,11 © XXXX American Chemical Society

Received: July 24, 2016 Revised: September 21, 2016 Published: October 18, 2016 A

DOI: 10.1021/acs.jpcc.6b07408 J. Phys. Chem. C XXXX, XXX, XXX−XXX

Article

The Journal of Physical Chemistry C

Figure 1. (a) FTIR spectra of three kinds of TiO2 nanoparticles. (b) XPS spectra of TiO2 with and without modification. (c, d) TEM images of LDPE/KHT and LDPE/PFT nanocomposites, respectively.

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EXPERIMENTAL SECTION Materials Preparation. The polyethylene (LL4004EL) pellets were obtained from ExxonMobil, which were additivefree, melt flow 3.5 g/10 min, density 0.924 g/cm3. The TiO2 nanoparticles were supplied by Aladdin Industrial Corporation with an average diameter of about 50 nm. The specific surface of the nanoparticles is 86.96 m2/g. The coupling agent, including (3-aminopropyl)triethoxysilane (H2N(CH2)2 Si(OCH2CH3)3, code name KH550) and 1H,1H,2H,2H-perfluorodecyltriethoxysilane (CF3(CF2)7(CH2)2Si(OCH2CH3)3), denoted as KH and PF, were chosen to modify the nanoparticles (Figure S1 in the Supporting Information). Generally, the surface modification is a widely recognized step in the preparation process of the nanocomposite just to improve the dispersion of nanoparticles. In fact, the coupling agent molecules on the surface of the nanoparticle are an important component of the interface. Among two coupling agents, KH is a kind of common coupling agent and is widely used in the surface modification of the nanoparticles. PF is thought to possess good chemical inertia and is able to reduce the surface energy. For the sake of comparison, TiO2 modified with KH and PF are denoted as KHT and PFT, respectively, whereas unmodified TiO2 is simply denoted as TiO2. The modified TiO2 nanoparticles were blended with LDPE in a HAPRO torque rheometer with a 60 mL roller-cone mixer. The rotor speed was 60 rpm, and the mixing time was 10 min. Film samples (Table S1 in the Supporting Information) with different thicknesses were obtained by using compression melding at a temperature of 140 °C, lasting 6 min under a pressure of about 15 MPa. After that, the films were cooled to room temperature under the same pressure. All films were annealed in the vacuum oven at 90 °C for 12 h to eliminate the thermal history and internal stress. Materials Characterization. The FTIR spectra were obtained with a Nicolet 6700 FTIR spectrometer in the range of 600 to 4000 cm−1. TiO2 nanoparticle samples with and

without surface modification were tested by using reflection mode. All samples were analyzed at 1 cm−1 resolution. The background of the atmosphere was measured and subtracted from each spectrum. The changes in the absorption peaks of FTIR spectra are shown in Figure 1a, where some new peaks can be observed in the FTIR spectra of TiO2 nanoparticles with surface modification. The peaks at 1012, 1248, and 2922 cm−1 represent the typical stretching vibrations of the Si−O bond, C−F bond, and C−H bond, respectively.17 The X-ray photoelectron spectra (XPS) of the samples were recorded using a Kratos analytical spectrometer with a monochromatic Al Kα source to investigate the detailed chemical structures of TiO2 after the functionalization with KH and PF. As shown in Figure 1b, compared to the XPS spectrum of TiO2, the appearance of the Si 2p peak at about 100 eV, the Si 2s peak at about 150 eV, the C 1s peak at about 285 eV, and the N 1s peak at about 400 eV in KHT confirms the successful modification of KH. In addition, the appearance of a strong F 1s peak at about 687 eV reveals that the PF is successfully grafted onto TiO2 nanoparticles.18 TEM images, as shown in Figure 1c,d, show that TiO2 nanoparticles with an average diameter of about 100 nm have good dispersion in nanocomposites. The spherulite microstructure of LDPE and nanocomposites were observed. The samples for SEM observation were broken in liquid nitrogen in order to obtain the cross sections. The thickness of the samples was about 1 mm. The cross sections were etched at room temperature for 4 h in a 1% w/v solution of potassium permanganate in 5 parts concentrated sulfuric acid to 2 parts orthophosphoric acid to 1 part water.19 Then the cross sections were sputter-coated with gold in order to avoid charge accumulation. A differential scanning calorimetry (DSC) experiment was conducted to study the crystallinity of LDPE and LDPE/TiO2 samples under a nitrogen atmosphere at a heating/cooling rate of 10 K/min between 300 and 430 K. B


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Figure 2. (a, b) SEM images of spherulites in LDPE. (c, d) SEM images of spherulites in LDPE/2%KHT and LDPE/2%PFT. (e, f) DSC melting curves of LDPE/KHT and LDPE/PFT nanocomposites. (g) 1H wide-line SSNMR spectra and fitting results of LDPE. 1 H solid-state nuclear magnetic resonance (SSNMR) spectra were obtained using a Bruker AVANCE III NMR spectrometer at a proton frequency of 400.25 MHz. Results were collected for nonspinning LDPE and nanocomposites. A 1.27 μs 90° pulse with a recycle delay of 5 s was used for experiments. The trap-level properties were tested via the thermally stimulated current (TSC) measurement. Film samples with an average thickness of 350 μm were poled with a 2.5 kV/mm dc voltage at 50 °C for 30 min in a dry nitrogen atmosphere and then cooled rapidly with the application of the poling voltage. When the sample had been cooled to −100 °C, the poling voltage was removed and the external circuit was shorted for 10 min. The TSC curves were measured by increasing the temperature under the short circuit condition from −100 to +100 °C at an increasing rate of 3 °C/min. The dc volume conductivity measurement was performed with an electrometer with a standard three-electrode system (Keithley 6517A). The thin-film samples with a thickness of about 120 μm for measurements were electrically short circuited for 5 min to eliminate internal charges, and then a constant dc electric field of 40 kV/mm was applied to the samples for 5 min. For each sample, five specimens were used

for repeat measurements, and the charging current was recorded.

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RESULTS AND DISCUSSION The morphology of the etched cross sections of pure LDPE and nanocomposites is shown in Figure 2a−d, and the black holes represent the original location of TiO2 nanoparticles, which were removed during the etching process. Figure 2a,b shows that LDPE has well-defined spherulites with an average diameter of 8 μm. With the addition of KHT and PFT, which act as heterogeneous nucleating agents, the diameter of spherulites decreases to 3 to 4 μm and their number increased. It is easy to draw a conclusion that nanoparticles can affect the crystallization of the polyethylene matrix, reduce the proportion of the crystalline region, and increase the proportion of the interphase region between the crystalline region and the noncrystalline region.20,21 The melting curves of DSC are shown in Figure 2e,f. Table 1 summarizes the melting peak temperatures Tm, melting enthalpies ΔHm, crystallinity Xc, and crystal thickness Lc of two samples, among which Xc is obtained from the ratio of the C


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The 1H wide-line spectrum of LDPE is shown in Figure 2g. The hydrogen atoms in polyethylene are in the same chemical environment, so there is only one peak in the 1H wide-line spectrum and the analysis is simplified. Although polyethylene is widely used in the research of nanocomposites, we have drawn interesting conclusions from the experiments. The analysis of the spectrum is based on the extent of constrained molecular motion in different polymer domains.26,27 The spectrum can be decomposed into three components by fitting the spectrum to the sum of a Gaussian function, a Lorentzian function, and a combined Gaussian and Lorentzian function via free DMFIT software,26 representing the rigid phase, the amorphous phase, and the intermediate phase, respectively, according to the chain segment mobility. The phase proportion can be obtained via calculating the composition of each spectrum integration, as shown in Table 2. The NMR rigid phase proportion is significantly higher than the DSC crystallinity. We know that the 1H SSNMR test is based on the mobility of molecular chain segments whereas DSC is based on the melting enthalpies in the crystalline region. Some chain segments near nanoparticles, that is, the interface, and structured chain segments in the crystalline region do not form a lattice, so they cannot contribute to the melting enthalpies in the crystalline region; however, their motion is hindered with the effect of nanoparticles and structured chain segments. As a result, they behave like chain segments in the crystalline phase and contribute to the proportion of the rigid phase. The result of ΔX = Xrigid − Xc reveals the existence and change in bound chain segments in LDPE and its nanocomposites. In pure LDPE, ΔX is completely contributed by bound chain segments near the crystalline phase, which are divided by the noncrystalline phase in the DSC experiment. After the addition of nanoparticles, ΔX is mainly contributed by bound chain segments around nanoparticles. The increase in ΔX means that the quantity of bound chain segments increases with the increment of nanoparticles. The NMR rigid phase proportion decreases after the addition of both KHT and PFT nanoparticles, compared to pure LDPE. This is mainly due to the decrease in DSC crystallinity, which contributes to the NMR rigid phase proportion. In contrast, the amorphous phase proportion increases as nanoparticles affect the crystallization process, increase the number of spherulites, and reduce their dimension. However, with the increment of KHT and PFT loading, the bound effect of nanoparticles causes more chain segments of polyethylene to transfer from the amorphous phase to the rigid phase, so the proportion of the rigid phase increases and the proportion of the amorphous phase decreases.

Table 1. Summary of DSC Results of LDPE and LDPE/TiO2 Samples sample

Tm (K)

ΔHm (J/g)

Xc (%)

Lc (nm)

LDPE LDPE/1%KHT LDPE/2%KHT LDPE/3%KHT LDPE/1%PFT LDPE/2%PFT LDPE/3%PFT

393.74 394.53 395.03 395.89 394.21 394.36 394.59

114.89 111.13 109.21 106.98 114.45 113.48 112.01

39.21 37.93 37.27 36.51 39.06 38.73 38.23

16.01 16.82 17.39 18.45 16.48 16.64 16.89

enthalpy and Lc is calculated by using the Thomson−Gibbs equation22,23 Xc = Lc =

ΔHm Δhf0

(1)

2σTm0 (Tm0 − Tm)Δhf0ρc

(2)

−2

Δh0f

where σ (93 mJ/m ) is the fold surface free energy, (293 J/g−1) is the melting enthalpy of 100% crystal-core polyethylene and T0m (410 K) is its equilibrium melting point, and ρc (1 g/cm3) is the density of crystal polyethylene. It is obvious that the melting peak temperature increases with the increase in the number of nanoparticles, indicating that KHT and PFT nanoparticles act as heterogeneous nucleating agents, increase the crystallization temperature, and accelerate the formation of the crystal nucleus.24,25 However, DSC crystallinity decreases with increases in the number of both KHT and PFT nanoparticles, which means that both of them can suppress the crystallization of polyethylene. The reason for the decrease in crystallinity is likely to be the absorption of nanoparticles on polyethylene chain segments and the space steric effect of nanoparticles; meanwhile, the bound chain segments make it difficult to crystallize, and as a result, large spherulites are almost impossible to formed, just as shown in SEM images. In addition, the crystal density may decrease as the crystal thickness increases and the crystallinity decreases. By comparing the results of two kinds of nanocomposites, we find that the effect of KHT on crystallinity and melting peak temperature is more significant than PFT. It is assumed that coupling agent PF reduces the effect of TiO2 nanoparticles on polyethylene chain segments. The C−F bond possesses good chemical inertia and can significantly reduce the surface energy and bound effect of TiO2 nanoparticles. As a result, the increase in the melting peak temperature and the decline of DSC crystallization in LDPE/PFT are not so obvious.

Table 2. Phase Composition and Line Width Decomposed from 1H SSNMR Spectra (R for Rigid, I for Intermediate, and A for Amorphous) phase composition (%)

half-peak width (kHz)

sample

R

I

A

ΔX (%)

R

I

A

LDPE LDPE/1%KHT LDPE/2%KHT LDPE/3%KHT LDPE/1%PFT LDPE/2%PFT LDPE/3%PFT

44.74 42.98 43.72 44.54 43.94 44.38 44.72

39.01 39.05 39.07 39.17 38.91 38.99 39.09

16.25 17.97 17.21 16.29 17.15 16.63 16.19

5.53 5.05 5.79 8.03 4.88 5.65 6.49

62.04 63.28 63.33 63.46 62.23 62.46 62.94

13.69 14.05 14.14 14.22 14.00 14.16 14.35

2.85 2.98 3.08 3.18 2.95 3.03 3.09

D


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Figure 3. Spherulite and interface model of the two nanocomposites. (a) LDPE/KHT and (b) LDPE/PFT.

Figure 4. (a, b) Thermally stimulated current of LDPE and nanocomposites. (c, d) Trap energy level and trap density of LDPE and nanocomposites. (e, f) Time dependence of dc leakage currents of LDPE and nanocomposites.

these bound chain segments. However, the bound effect on the chain segments decreases with distance, so it is not difficult to speculate that the strength of the bound force of nanoparticles will affect the number of bound chain segments, which means that the interface has a certain range. Although we are not yet sure of the exact size of this range, we know that PFT with a reduced bound effect will produce a smaller interface range. On the basis of the conclusions drawn above, the spherulite and interface model of the two nanocomposites is established, as shown in Figure 3. The nanoparticle, acting as the heterogeneous nucleating agent, becomes the center of the spherulite. The interface can be divided into two layers according to the chemical composition. The inner layer, denoted as the coupling agent layer, is composed of coupling agent molecules, which are connected to the nanoparticle by chemical bonds. The outer layer, denoted as the bound segments layer, is composed of bound polyethylene chain

We can also easily draw some conclusions via the half-peak width ΔυH of each phase. ΔυH is inversely proportional to the apparent spin−spin relaxation time of a proton T2*, which is an important parameters related to the molecular motion. The increase in the half-peak width in rigid, intermediate, and amorphous phases is found with the KHT and PFT load, suggesting that T2* decreases and the mobility of the chain segments is hindered, which is also consistent with previous analysis. By comparing the results of two kinds of nanocomposites, we find that the effect of KHT on the NMR phase composition and half-peak width is more significant than that of PFT. As in the previous analysis, PFT can significantly reduce the surface energy and bound effect of TiO2 nanoparticles, compared to KHT. From the DSC and NMR results, we have understood that nanoparticles can bind polyethylene chain segments and suppress their mobility,28−31 and the interface is formed by E


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decay very quickly and reach their steady state within 5 min. It is obvious that the leakage current increases with the increment of nanoparticle loading. For two kinds of nanocomposites, the volume conductivity increases with the nanoparticle content, and the volume conductivity of LDPE/KHT is lower than that of LDPE/PFT with the same nanoparticle loading. We know that the volume conductivity is mainly affected by charge transport in polymer materials and that charge transport is determined by the charge carrier mobility. The charge transport in polymer includes band transport in the molecular chain and hopping transport between molecular chains. The atoms in polyethylene molecules are bounded by strong covalent bonds, and the electronic wave functions overlap substantially, resulting in a high charge carrier mobility. However, polyethylene molecular chains are combined with van der Waals forces and the electronic wave functions overlap slightly, so the charge carrier mobility is low between polyethylene molecular chains. As a result, the charge transport of polyethylene is mainly dependent on the hopping transport between molecular chains,41 and the corresponding conductivity is called the hopping conductivity. Because there is a close relationship between the hopping conductivity and the thermal excitation of electrons, traps with different energy levels play different roles in charge transport. We know that the probability of thermal excitation in deep traps is much lower than that in shallow traps whereas the lifetime of charge carriers in deep traps is much longer than that in shallow traps, so if the charge carriers are captured by shallow traps, then it is easy for them to get enough energy via thermal excitation to get out. In contrast, it is hard for charge carriers to escape if they are captured by deep traps. TSC results reveal that the addition of nanoparticles will introduce large quantities of shallow traps into the bound segments layer of the interface. The energy level of these traps is lower than that of traps from the interphase between the crystalline region and the noncrystalline region, so they are helpful for charge transport, compared to pure LDPE, whereas a small number of charge carriers may be trapped in deep traps in the coupling agent layer of the interface. That is the reason that the volume conductivity increases with the increment of nanoparticle loading. The schematic illustration of charge transport in LDPE and nanocomposites is shown in Figure 5. In pure LDPE, charge carriers mainly depend on shallow traps in the interphase for transport, but in nanocomposites, shallow traps with lower energy levels in the bound segments layer make charge transport much easier. Besides the energy level,

segments. The microstructure and mobility of the chain segments in the bound segments layer are different from those of chain segments in crystalline and noncrystalline regions in the matrix. In addition, the range of the bound segments layer is determined by the bound effect of the nanoparticle. In this article, we assume that KHT has a strong bound effect whereas the bound effect of PFT is reduced. As a result, the range of the bound segments layer in LDPE/KHT is larger than that in LDPE/PFT. The model is supposed to help us understand the microstructure of the interface and the interaction between nanoparticles and polymer molecules. The trap-level distribution is an important property of materials. Figure 4c,d reveals the trap energy level and trap density of LDPE and its nanocomposites, which are calculated from thermally stimulated current results32 described in Figure 4a,b. One can see from Figure 4c,d that pure LDPE has one peak corresponding to a characteristic trap energy level at about 0.99 eV, which is consistent with the results reported in refs 32−34. These traps are reported to be the cavity traps from the interphase region between the crystalline region and the noncrystalline region of LDPE.35,36 In these cavities, the distance between polyethylene molecules increases and the short-range repulsive force decreases rapidly, resulting in the electrons bounded by polarized polyethylene molecules, that is, the mechanism of cavity traps.37−39 After the addition of KHT and PFT, the trap density of cavity traps from the interphase increases with the increment of nanoparticle loading, which may correspond to the increased proportion of the interphase region observed in SEM images. However, new peaks corresponding to lower characteristic trap energy levels are observed in two kinds of nanocomposites. LDPE/KHT has a characteristic trap energy level at about 0.79 eV, and LDPE/PFT has a characteristic trap energy level at about 0.72 eV, which are higher than the trap energy level of the noncrystalline region reported.40 We know that the main sources of traps in polymer nanocomposites include the traps from the interphase between the crystalline region and the noncrystalline region, the traps from the interface between the nanoparticles and the matrix, and the traps from chemical bonds on the surfaces of nanoparticles and other chemical defects. Research37−39 reports that traps produced by chemical bonds and chemical defects are discrete deep traps with an energy level of about 2.0−5.0 eV, whereas traps from the interphase and interface are usually cavity traps corresponding to an energy level lower than 1.0 eV. Therefore, we can confirm that the new peak of traps is from the cavity traps in the interface, accurately said, the bound segments layer. The coupling agent layer of the interface may produce deep traps because of its chemical bonds and chemical defects; however, it is difficult for charges trapped in these deep traps to escape via a thermal stimulation method. In addition, the trap density increases with the increment of nanoparticle loading, and the trap density is approximately proportional to the nanoparticle loading, which also reveals that the new peak of traps is produced by the interface. On the other hand, we notice that with the same nanoparticle loading the trap density of LDPE/ KHT is higher than that of LDPE/PFT. One possible reason is that PFT produces a smaller interface range because of its reduced bound effect, compared to that of KHT, so the quantity of traps in the bound segments layer of LDPE/PFT is smaller than that in LDPE/KHT. Figure 4e,f describes the time dependence of dc leakage currents of LDPE and nanocomposites. The leakage currents all

Figure 5. Schematic illustration of charge transport in LDPE and a nanocomposite. (a) Pure LDPE and (b) LDPE nanocomposite. F


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the quantity of shallow traps is also important in charge transport. We know that the bound effect of PFT is weaker, compared to KHT, resulting in a small interface range and a relatively higher energy level of shallow traps, so the influence of traps on charge transport is reduced and the leakage currents are small.

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CONCLUSIONS In this article, we reported the experimental results of the influence of functionalized TiO 2 nanoparticles on the aggregation structure of polyethylene nanocomposites. We constructed a microstructural model of the interface according to the conclusions. DSC and solid-state NMR results show that TiO2 nanoparticles can hinder crystallization, resulting in a decrease in crystallinity. The polyethylene molecular chain segments around nanoparticles are bound and their mobility is suppressed, forming the interfacial region that is different from the matrix with respect to the microstructure and molecular chain mobility. However, the influence on the aggregation structure is determined by the bound effect of nanoparticles, which is affected by the surface modification. Silane coupling agent PF possess good chemical inertia and can significantly reduce the surface energy of TiO2 nanoparticles, so the influence of the interface in LDPE/PFT is weaker than in LDPE/KHT. On the basis of the microstructural model of the interface, further research revealed the relationship between the trap energy level distribution and the special microstructure of the interface. The chain segments in the interface can produce cavity traps with an energy level of about 0.7−0.8 eV, which is lower than that of cavity traps in the matrix. The presence of interfacial shallow traps increases the volume conductivity of the nanocomposites, compared to that of pure polyethylene. However, the interface effect of PFT is weak, so the influence of the interface in LDPE/PFT on the charge carrier mobility and charge transport is reduced. The experimental results would provide an important hint to the research of the interface between nanoparticles and polymer matrixes in the future.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.6b07408. Detailed description of the experimental setup, methods for the preparation and characterization of the probe and samples, methods for data acquisition and analysis, and supplementary data (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. Tel: (86)-10-62795423. *E-mail: [email protected]. Tel: (86)-10-62775585. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors are grateful for the financial support of the National Basic Research Program of China (973 Project, grant no. 2014CB239505).

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REFERENCES

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