3.1. FT-IR Spectroscopy
shows the FT-IR spectra of PVA, graphite, GO, HDA-GS, and OMB. The spectrum of pure graphite does not exhibit any peaks while that of PVA and GO exhibit significant and broad absorption peaks characteristic for the OH and COOH functional groups [31
]. Details of each substance are specified below.
The characteristic absorption peaks of the PVA are observed at 3330 cm−1
(O–H stretching), 2940 cm−1
and 2910 cm−1
(asymmetric stretching CH2
), 1730 cm−1
(due to water absorption), and 1256 and 1090 cm−1
(C-H bending and C-O stretching). In the case of the GO, the characteristic absorption peaks of the O-H are observed at 3220 cm−1
(stretching) and asymmetric epoxy appears in 3063 cm−1
, but does not appear to overlap with OH, 1732 and 1616 cm−1
(C=O stretching), and 1040 cm−1
(C–O stretching). The epoxide ring and C=C bond are much weaker compared to the others, whereas the O–H stretching is more intense. Figure 1
also shows the spectrum of HDA-GS: 3160 cm−1
(O–H stretching), 2921 and 2852 cm−1
(aliphatic C–H stretching), 1375 cm−1
(aromatic C–N–C symmetric stretching), and 790 cm−1
(N–H out-of-plane stretching). The OMB also exhibits a peak at 1100 cm−1
3.2. Morphology of HDA-GS
The SEM images of unmodified natural graphite, GO, and HDA-GS are shown in Figure 2
. The natural graphite exhibits a lamellar structure similar to that of graphene sheets with stacks denser than those observed for the other materials (Figure 2
a). Figure 2
b shows the translucent GO sheets that are wrinkled and folded in a manner resembling thin paper. Additionally, HDA-GS is prepared from chemically modified GO and exhibit an entangled morphology and a random distribution (Figure 2
Generally, a significant volume expansion and high porosity are observed in the FGSs, which results in low FGS bulk densities that can cause feed problems during melt compounding. Normally, master batches are required to solve the problem. However, in the study, solution blending of FGSs with polymers is successfully employed.
The AFM imaging provides more reliable information on the sheet dimensions and can also be used to probe the surface topology, defects, and bending properties. Additionally, stepped-height scans can also allow us to determine the lateral sizes and thicknesses of the particles lying on the substrates. Thus, AFM is used to indicate that the carbon sheets obtained in the present study are comprised of only a single atomic layer. Figure 3
shows the AFM image of HDA-treated GO sheets (HDA-GSs) on mica and a profile plot that reveals the average sheet thickness corresponding to 1.76 nm (Figure 3
) when the thickness of the bare graphene sheet is approximately 1 nm [32
], and the thickness of the layer of substituted HDA organic groups is approximately 0.76 nm. However, the layer of the substituted HDA organic groups exhibits a thickness of approximately 0.38 nm when the organic HDA groups on both sides of the graphene sheet in the synthesized FGS are significantly tilted. The change in thickness is associated with different orientations that are adopted by the long alkyl chains in the chemically modified HDA-GSs.
The XRD traces of the pure nanofillers and their PVA hybrid films are shown in Figure 4
. The d001
reflection of pristine SPT is present at 2θ
= 6.62°, and this corresponds to an interlayer spacing d
of 13.54 Å, as shown in Figure 4
a. Specifically, Figure 4
a also shows the XRD curves of the SPT/PVA hybrid films with clay contents in the range of 0–10 wt%. With respect to the PVA hybrids with a clay content ≤10 wt%, clay peaks did not appear in the XRD traces, which indicates that the clay particles are homogeneously dispersed in the hybrid polymer matrices. Figure 4
b also shows the XRD curves of OMB in the region 2θ
= 2°–10°. The d001
reflection of OMB is observed at 2θ
= 6.64°, and this corresponds to an interlayer distance (d
) of 13.30 Å. In the case of PVA hybrids containing 3 wt% of OMB, only a very slight peak appeared at 2θ
= 4.78°, and this corresponds to an interlayer spacing of 18.46 Å (Figure 4
b). The result indicates that a small amount of the clay is not aggregated in the PVA matrix. However, a significant increase in aggregation is observed for samples with OMB loadings reaching 10 wt%, as shown by the intensities of the XRD peaks. This suggests that perfect exfoliation of the layered structure of the clay did not occur. Further evidence of clay dispersion into PVA on a nanometer scale is obtained via TEM.
The XRD diffractograms of pure HDA-GS and HDA-GS/PVA hybrid films are shown in Figure 4
c. The d001
reflection for HDA-GS is observed at 2θ
= 2.73° and corresponds to an interlayer spacing (d
) of 32.32 Å. With respect to PVA hybrid films containing up to 7 wt% of HDA-GS, the peak observed at 2θ
= 2.73° for GS is observed to disappear from the diffraction patterns. The result indicates that the graphene layers are exfoliated and homogeneously dispersed throughout the PVA matrix and provide supporting evidence for the nanocomposite character of the HDA-GS/PVA hybrids. However, the intensities of the XRD peaks at 2θ
= 2.64 (d
= 33.42 Å) and 2θ
= 5.35 (d
= 16.50 Å) increase suddenly when the HDA-GS loading increases from 7 wt% to 10 wt%, which suggests that the dispersion is more effective at lower loadings as opposed to higher loadings of graphene.
Given the periodic arrangement of the graphite layers in the virgin GS and in intercalated hybrids, the XRD offers a convenient method to determine the interlayer spacing. However, although the XRD enables precise routine measurements of the GS layer spacings, it neither allows for the determination of the spatial distributions of GSs or the detection of any inhomogeneous sections of the hybrids. Initially, a few layered GSs do not exhibit well-defined basal reflections, and, thus, it is difficult to systematically follow peak broadening and reductions in intensity. Therefore, all conclusions on the mechanisms of hybrid formation and microstructure, based solely on XRD results, are only tentative. Therefore, further evidence of the GS dispersion in the PVA films on a nanometer scale is obtained via SEM and TEM as described in the next section.
3.4. Morphologies of PVA Hybrids
In addition to using XRD to measure the d
spacings of the nanocomposites, SEM and TEM are used to evaluate the degree of intercalation and amount of aggregation in the nanofiller clusters. The morphologies of the aggregated fillers are characterized via SEM. Large filler aggregates can be easily imaged by SEM due to the difference between the scattering densities of the filler and the PVA matrix [33
The morphologies of the hybrid films containing up to 10 wt% of SPT in the PVA matrix are examined by observing their fracture surfaces by SEM (Figure 5
a). The PVA hybrid films containing 3 wt% and 5 wt% of SPT display uniform and dispersed phases. Conversely, the films containing 7 wt% and 10 wt% of SPT exhibit large particles and some deformed regions that can result from the coarseness of the fractured surface. A similar type of behavior was observed for PVA/OMB hybrid films. For example, the fractured surface of the 5-wt% OMB hybrid film (Figure 5
b) exhibit uniform and dispersed phases. However, increased agglomeration is observed in the PVA matrix with clay content exceeding 7 wt% in the OMB system.
The fractured surfaces display increased levels of deformation for samples with higher clay contents. The trend is most likely linked to increases in the agglomeration of clay particles and indicate the lack of interfacial interactions between the clay and matrix polymers. Thus, several defects and significant agglomeration occur in interphase areas in high clay content nanocomposite PVA films. Figure 6
shows a comparative analysis of the SEM micrographs obtained for the PVA hybrids with different contents of HDA-GSs that exhibit platelet-orientation distribution morphology.
Graphene dispersions are readily observed in the SEM images due to the differences in the scattering densities of graphene and the matrix polymer. The SEM images of the fractured surfaces of PVA hybrid films containing 0–10 wt% of HDA-GS are compared in Figure 6
. The hybrid film with 5 wt% of HDA-GS display a morphology consisting of graphene domains that are well dispersed throughout the continuous PVA phase. However, deformed surfaces and voids are observed in the case of the 7-wt% FGS hybrid film. Conversely, the micrographs of the 10-wt% HDA-GS/PVA hybrid films exhibit increased levels of voids and deformed regions when compared to the 3–7 wt% HDA-GS/PVA hybrid films due to the coarseness of the fractured surface. Overall, the comparison reveals that increases in the graphene content in the hybrid films increase the level of deformation of their fractured surfaces. The finding potentially results from the agglomeration of graphene particles. A comparison of the micrographs indicate that the fractured surfaces of the hybrid films with higher graphene content are more deformed than those of hybrid films with lower graphene content, and this is possibly due to the agglomeration of graphene particles [34
]. It should be noted that Figure 6
also shows that most of the graphene remains in the form of straight and rigid platelets in the composite, which indicates that the graphene sheets are extremely stiff.
We extend the morphological analysis via TEM to evaluate the degree of intercalation and degree of aggregation of the nanofiller clusters. Additional direct evidence for the formation of a true nanocomposite is provided by the TEM analysis of ultra-microtomed sections. Figure 7
shows the micrographs of the PVA hybrid films with identical contents of the three different nanofillers. The dark lines in the photographs denote the intersections of the clays and GSs (1-nm thick) while the length between the dark lines denotes the interlayer distance. Figure 7
a shows the morphologies of the PVA hybrids with 5 wt% and 10 wt% of SPT. With respect to the hybrid containing 5 wt% of SPT, each layer in the polymer matrix consists of well-dispersed individual clay layers (dark lines), and a few of the clays aggregate to a thickness of approximately 10 nm. In a manner similar to the case for the hybrid material containing 10 wt% of SPT, these clays are mainly aggregated in the polymer matrix. However, the average particle size is observed to be below 20 nm, as calculated from the TEM images. The presence of agglomerated particles in the SEM micrographs of hybrid materials with higher SPT contents (see Figure 5
a) is attributed to the formation of aggregated layers.
Typical TEM images of PVA hybrid films containing 5 and 10 wt% OMB are shown in Figure 7
b. Evidently, the clays are well dispersed within the polymer matrix irrespective of the clay content. In contrast to the clay layers in the hybrids containing SPT, the clay layers in OMB hybrids are exfoliated within the matrix polymer.
The TEM micrographs of the 5 and 10 wt% HDA-GS hybrid films are shown in Figure 7
c. The TEM micrographs indicate that the GS in the 5 wt% HDA-GS hybrid is dispersed in the polymer matrix, which indicates the formation of a nanocomposite. The findings suggest that GS breaks down into nanoscale building blocks during the intercalative polymerization process and is homogeneously dispersed in the polymer matrix to yield a polymer/GS nanocomposite. As in the case of the 10 wt% HDA-GS (see Figure 7
c), the GSs are mostly agglomerated in the polymer matrix. In contrast to the hybrids containing 5 wt% of HDA-GS, the graphene layers of the 10 wt% hybrid exhibited agglomeration of the dispersed graphene phase and are not intercalated into the matrix polymer. The agglomeration of the dispersed graphene phase visibly increases with increases in graphene content, and the outcome is consistent with the XRD and SEM data shown in Figure 4
c and Figure 6
The XRD, SEM, and TEM results indicate that the fillers are well dispersed throughout the PVA matrix at low filler contents while aggregated structures are present at higher filler contents. Additionally, the dispersions of SPT and OMB exceed that of HDA-GS in the PVA matrix (see Figure 4
, Figure 5
, Figure 6
and Figure 7
). The unusual thermal and gas barrier properties of these hybrid films are discussed in the following sections with respect to the dispersion of the nanofillers.
3.5. Thermal Properties
A comparison of the DSC results for pure PVA and PVA hybrids with approximately 3–10 wt% of clays (SPT and OMB) and graphene (HDA-GS) are listed in Table 1
. The glass transition temperature (Tg
) of pure PVA corresponds to 69 °C. The Tg
values of PVA hybrids containing various clay contents are virtually unchanged in the DSC results when compared to that of pure PVA irrespective of the filler loading, i.e., approximately 68–71 °C for the three different fillers. Generally, Tg
increases when the filler content increases up to a critical concentration. The increase in Tg
is potentially due to the confinement of the intercalated polymer chains within the filler galleries that prevents the segmental motion of the polymer chains [36
]. However, the Tg
values of the three types of PVA hybrid films remains constant irrespective of the filler content, which suggests that the variation in the filler content does not affect the confinement of the PVA chains.
The endothermic peak of pure PVA appears at 165 °C and corresponds to its melting transition temperature (Tm
) (see Table 1
). The Tm
values of the hybrid films are observed to increase from 165 to 176 °C when the SPT loading is increased from 0 to 5 wt% and, subsequently, decreases to 166 °C at an SPT content of 10 wt%. The increase in the Tm
of the hybrid film potentially occurs as a result of the insulation effect of the clays and the interactions between the clay and PVA chains [38
]. However, the decrease in the Tm
value of the 10-wt% SPT hybrid suggests that its domains are more poorly dispersed in the PVA matrix than those in the 5-wt% SPT hybrid. Hence, increases in clay content led to the aggregation of clay particles, and this reduces the heat-insulation effect of the clay layers in the polymer matrix.
A similar trend is observed for both the OMB and HDA-GS hybrids. Specifically, the Tm
of the PVA hybrids increases from 165 to 184 °C and from 165 to 173 °C when the filler loadings increase to 5 wt% for OMB and 7 wt% for HDA-GS. As observed for the SPT hybrid material, the maximum transition peaks of the PVA hybrids increase with the addition of the nanofiller only up to a certain content level and, subsequently, decreases when the content is increased above this point. For example, when the filler content of PVA reaches 10 wt%, the Tm
decreases to 178 °C and 156 °C for materials employing OMB and HDA-GS nanofillers, respectively. The DSC thermograms of the PVA hybrids with various HDA-GS contents are shown in Figure 8
. When the nanofiller is included, the peak is stronger than that of the pure PVA, and the more the amount of the filler increased up to 7 wt%, the greater the intensity of the peak becomes. Hence, the HDA-GS appears to act as a nucleating agent [39
]. However, when the amount of HDA-GS reaches 10 wt%, the degree of dispersion decreases in conjunction with decreases in the peak intensity.
In a manner similar to the results for Tm
, the initial thermal degradation temperatures (TDi
) of the PVA hybrid films also increases linearly from 227 °C to 249 °C while increasing the SPT loading from 0 to 5 wt%, as shown in Table 1
. With respect to the hybrids containing OMB, the TDi
value varies from 227 to 252 °C when the content of the OMB organoclay increases from 0 to 5 wt% in the PVA hybrids. The highest increase of 25 °C in the TDi
relative to that of pure PVA is observed for the 5 wt% OMB hybrid (252 °C). The clay enhances the TDi
by acting as an insulator and a mass-transport barrier to the volatile products generated during decomposition [40
]. The increase in thermal stability is attributed to the high thermal stability of the clay and interactions between the clay particles and polymer matrix. In contrast to the behavior observed for clay contents ranging from 0 to 5 wt%, the TDi
values of the hybrids decreases when the clay content increases from 5 to 10 wt%. For example, the TDi
values of PVA hybrid film containing 10 wt% of clay loadings were 11 °C (238 °C) and 6 °C (246 °C) lower than those of PVA hybrids containing 5 wt% of SPT and OMB, respectively. The decrease in TDi
appeared to correspond to the result of clay aggregation that occurs when the clay content in the polymer matrix exceeds a critical value. Similar results are observed in samples containing HDA-GS, as shown in Table 1
. The TDi
values of the PVA hybrid films also increases linearly from 227 to 245 °C when the HDA-GS loading is increased from 0 to 7 wt%. At 10 wt% content of HDA-GS in the PVA, the TDi
decreases again to 238 °C.
An analysis of the weight residue at 600 °C (wtR600
) indicates that the weight increases with growth in the filler loading from 0 to 10 wt% and corresponds to a range from 3% to 15% for SPT, from 3% to 14% for OMB, and from 3% to 15% for HDA-GS (Table 1
). The enhancement in char formation with increases in filler content is ascribed to the high heat resistance of the clays and graphene.
3.6. Gas Permeation
The mobility of the polymer chain segments in the polymer nanocomposite clearly differs from that of the pure polymer due to the confined environment, and this also affects the gas permeability. The following two main factors are responsible for permeability reduction [42
]: (i) polymer chain-segment immobility and (ii) detour ratio, which is defined as the ratio of the film thickness in the nominal diffusion flow direction to the average length of the tortuous diffusion distance between nanolayers.
High aspect ratio nanolayers also lead to properties that are not possible for larger-scaled composites. The impermeable nano-sized layers mandate a tortuous pathway for a permeant to transverse the nanocomposite. Enhanced barrier characteristics, chemical resistance, reduced solvent uptake, and flame retardance of polymer hybrids benefit from the hindered diffusion pathways through the nanocomposite [43
In order to further characterize the barrier properties of the PVA hybrids fabricated by the intercalation of polymer chains in the galleries of clays and graphene, the permeability of the resulting PVA hybrid films to O2
is evaluated for various filler loadings in the range of 0–10 wt%. The results are summarized in Table 2
. The thicknesses of all the films subjected to the gas permeation measurements are in the range of 20 to 24 μm. We discuss our results in terms of relative permeability, Pc/Pp
, where Pp
denotes the permeability of the pure polymer and Pc
denotes the permeability of the composite. The results confirm that the mass-transfer process for O2
as the penetrant is highly dependent on the level of filler loading. For example, the addition of only 7 wt% of SPT results in a 95% reduction in the permeability rate of O2
/day) relative to that of the pure PVA film (5.13 cc/m2
/day). With respect to the PVA hybrid films containing 3 to 7 wt% of OMB, the relative O2
permeability rate is close to zero (see Table 2
). The outcome is attributed to the increase in the length of the tortuous paths followed by the gas molecules and interactions between the O2
and clay molecules. Furthermore, films containing higher amounts of clay appear as significantly more rigid, and this decreases their gas permeability. In a manner similar to the results observed for the clay hybrids, the addition of 5 wt% of HDA-GS results in an 81% reduction in the permeability rate of O2
/day) relative to that of the pure PVA film (5.13 cc/m2
/day). However, when the clay content increases from 7 to 10 wt%, the O2
permeability rate slightly increases from 0.24 to 0.88 cc/m2
/day for SPT and more significantly from <10−2
to 3.18 cc/m2
/day for OMB. The increase in the HDA-GS loading from 5 to 10 wt% led to a similar increase in permeability from 0.98 to 4.27 cc/m2
/day. The increases in permeability values are primarily due to the aggregation of the filler particles in materials employing loadings that exceed the critical filler content levels. The present results are further corroborated via the electron micrographs shown in Figure 5
, Figure 6
and Figure 7
Collectively, the results of the gas permeation analysis reveal that the strongest gas barrier effect of OMB is observed with respect to the O2
TR among the three types of nanofillers that are studied. The enhanced gas barrier capacity of OMB stems from its hydrophilic character, and this allows the formation of hydrogen bonds between OMB molecules and the PVA polymer matrix as well as the reinforcement of chain packing, which significantly reduces the gas permeability of the material [45