Fire and thermal resistance properties of chemically treated ligno-cellulosic coconut fabric–reinforced polymer eco-nanocomposites

Materials used

The unsaturated isophthalic PE resin obtained from M/s Vasivibala resins (P) Ltd, Chennai, India, was used as the polymer matrix. Nanomer 1.31 PS, the OMMT nanoclay surface modified with 15–35 wt% of octadecylamine and 0.5–5 wt% of aminopropytriethoxysilane was procured from M/s Sigma Aldrich (P) Ltd, Bangalore, India. The naturally woven CS in the form of woven mat obtained from the outer basks of the coconut tree was used as a primary reinforcement. The CS mats used throughout this work were extracted from the same agricultural land near to Watrap in Virdhunagar (District), Tamil Nadu, India. The sheath mat contained thick and thin fibers arranged in random architecture. Thick fibers were oriented like fish bone shape with main fiber at the centre. Thin fibers were present in between the thick fibers and the resulting system appeared in the form of a randomly oriented woven like mat.

Process of chemical modification

The CS fiber mats were treated with 4% aq. NaOH solution in a water bath at the room temperature for 60 min. The treated mats were washed several times with fresh water until all the NaOH deposition over the fiber surface was removed. The fiber mats were then left to dry at the room temperature for 24 hours followed by the drying in an oven at 80 ℃ for next 1 h. The 0.4 wt% of trichlorovinyl silane [C₂H₃Cl₃Si] solution was made in acidified water (pH 3.5). The CSs were immersed in that solution for 1 h. The treated mats were conditioned at room temperature for about 24 h prior to use. The chemical analysis of the untreated, alkali- and silane-treated CS fiber was already reported in our previous work [23].

Fabrication of hybrid nanocomposites

In this work, the fabrication of CS fiber/MMT nanoclay reinforced hybrid polymer composites was carried out in two steps. In the first step, the proportionate wt% of nanoclay was taken according to the wt% of unsaturated isophthalic PE resin and the above mixture was stirred by the high-speed mechanical shear mixer at 500 r/min for about 2 h with radial turbine blade [24]. During the fabrication of pristine PE, the 10-mm thick rubber sheet was cut by having rectangular cavity with one side open. In order to avoid the leakage of resin, the 10-mm plywood mold was placed between the glass plates by keeping the runners on the top side and tightened by pin clips. The synthetic PE sheets were placed adjacent to the glass plate and were extended beyond the height of the mold. The resin-catalyst mixture was poured through the runner on the top of the mold. The polishing wax was applied at the sides of the cavities of the plywood mold for the easy removal of formed product.

In the second step, the CS/MMT clay/PE hybrid composites were fabricated using compression molding machine by the following procedure. The final product collected from the high-shear mixer was allowed for the degassing process. The mats were impregnated with the clay–PE mixture to which 1.5 vol% of cobalt naphthenate (as an accelerator) and MEKP (as a catalyst) were added. Then the clay–resin mixture with the curing agent was spread over the CS fiber mat with a steel brush in each layer, which were placed inside the mold cavity of size 300 × 125 × 10 mm³ by keeping one over another up to 13 layers (48 ± 2 wt%). The CS fibers were rolled with the help of a roller in a way to remove the air bubbles. The top plate was placed over the middle plate in a way to cover the CS fibers and positioned by the circular pins at the corners. The mold was closed completely by applying 150 kgf/cm² pressure in order to obtain the 10-mm-thick plate using a compression molding machine. The compressed mold was allowed at room temperature curing for another 24 h. Finally, the split mold was separated by removing the pins at the corners. The rectangular composite plate of size 300 × 125 mm² and thickness of 10 mm was carefully removed through the cavity of middle plate. Then the test specimens of the required size were cut from the fabricated composite plate.

Characterization studies

Thin slices with thickness in the range of 50–100 nm of the nanocomposite specimens were made using an ultra-microtome and their morphology was studied using the TEM Technai Sprit, FEI, (Netherlands) operating at 20–120 kV. The chemically treated fiber surface of the CSs was scanned using an FEI model XL30 ESEM SEM.

Fire testing details

The test for the fire resistance properties of the CS-reinforced PE composites was carried out on a cone calorimeter (FTT, East Grinstead, UK) according to the standard method prescribed in ISO 5660-1 [25]. The composite samples were cut into 100 × 100 mm size with a thickness of 10 mm [26]. A 50 kW/m² heat flux was applied to the sample. This is a widely used heat flux as it matches the test to developing fires [27]. All the measurements were made in a horizontal position with the conical radiant electric heater located 25 mm above the specimen. The samples were wrapped with aluminum foil around the back and edges before placing the sample onto the holder and then into the cone calorimeter. Ignitability was determined by using the standard 5s criteria for observing the time for sustained ignition of the specimen. The average values of three samples are reported in all the output measurements of fire-testing results.

The following test measurements such as the time to flameout (s), total heat release (MJ/m²), total smoke release (m²/m²), mass loss (g), specific extinction area (m²/kg), total oxygen consumed (g), average specific MLR (g/s.m²) and the effective heat of combustion (MJ/kg) were calculated to study the effectiveness of these composites against fire. All the data collected from the cone calorimeter test are expressed in engineering units in this paper.

TGA

The thermal stability of the PE nanocomposites was evaluated using SII EXSTAR TG 6000 module thermogravimetric analyzer. About 3–6 mg of sample was subjected to dynamic TGA scans at an increased heating rate of 5℃/min in the range of ambient temperature to 800℃ in N₂ atmosphere. The TG curves were analyzed as the percentage of weight loss as a function of temperature. The average of three powder samples is reported in each case (obtained from the three different locations of same specimen) using a diamond cutter.

Morphological analysis

To understand the distribution of nanoclay in PE matrix, TEM images were recorded and are shown in Figure 1. In the TEM photographs, the white region represents the matrix phase whereas the dark region corresponds to the clay distributed phase. In Figure 1, the thick dark band zones can be seen which correspond to the agglomeration and clustering of nanoclay in nanocomposite containing 5 wt% of clay. The viscosity of the resin–clay mixture increased due to the higher clay addition in the PE matrix. This was attributed to an increase in the shear force that led to improper mixing which in turn resulted in agglomeration of particles. TEM with 5 wt% of nanoclay did not show any layer dispersions. However, the presence of clay itself is expected to give enhanced fire-retardant properties in the case of hybrid composites as reported by others [28,29]. OPEN IN VIEWER

The morphology of the untreated, alkali- and silane-treated fibers of CS is presented in Figure 2(a–c). Figure 2(a) reveals the appearance of cuticle waxy layer over the CS, which is identified from the white color ripen-like structure and it leads to poor compatibility between the fiber and matrix (Figure 2a). Inherently deposited waxy layers and extended globular protrusions over the fiber surface disappeared after the alkali treatment, leading to increased fiber roughness with regularly arranged circular pores that are partially present as white dots in Figure 2(b). The magnified view of fiber surface indicates the visual observations of circular pores as well as partially existing globular white dots, which was attributed to the increase in silica content as reported by other authors [30]. The rough surface permits sufficient wetting to take place within the PE matrix which further increases the adhesive force between the fiber and the matrix by mechanical interlocking phenomenon. A non-homogenous thin film deposition of coating over the surface of silane-treated CS could be seen in Figure 2(c). The layer of trichlorovinyl silane forms the chemical bonding with the matrix which promotes a good adhesion between fiber and the PE matrix. OPEN IN VIEWER

HRR

Initially the PE/nanoclay composites were tested for flammability and later the PE/CS/nanoclay hybrid composites were tested. The key property of HRR which is the single-most parameter for fire performance [31,32] was determined in each case. The flammability characteristics of the natural reinforced composites involve a number of factors which include the orientation of fibers, chemical composition of plant fiber, degree of polymerization, thermal conductivity and the interfacial adhesion between the fiber and the matrix [8]. Figure 3 shows the HRR curves for the PE, untreated (UTC), silane-treated (STC) and the alkali-treated (ATC) CS-reinforced composites [without clay (WOC)] that were obtained from cone calorimeter, as a function of time. From Figure 3, it can be observed that both untreated and treated CS fiber-reinforced composites showed a significant (50%) decrease in the HRR compared to the pristine PE sample with the same thickness. It could be probably due to the incorporation of complex architecture of fiber orientation in the naturally woven CS fiber mat. This fiber orientation is expected to control the amount of oxygen penetration into the fiber. In general, the higher the orientation, lower will be the permeability of the fiber to oxygen [33]. OPEN IN VIEWER

Only one peak in HRR curve of PE matrix was observed in Figure 3 and it can be due to the greater rate of combustion of matrix by quicker thermal decomposition. In all CS-reinforced composites, the UTC is one which shows a higher HRR value compared to the other two composites. It may be due to the poor interfacial adhesion between the fiber and the matrix by the presence of the waxy layer and the inorganic contaminants over the fiber surface. In general, a weak interface can separate the two phases (reinforcement and polymer) apart in case of fire and as a result the PE matrix burns vigorously and the fiber which can no longer act as insulator sometimes acts rather as heat conductor, thus increasing flammability [34]. A high content of cellulose in UTC (68%) also could tend to increase the flammability of the fiber [6]. But the alkali treated CS which had low cellulose (52.8%) and high lignin (33.29%) content [23] is expected to have lower flammability. High silica content evidenced by high ash content (5.35 wt%) also tends to increase the flame resistance in the case of the alkali-treated composites. This increasing silica content was also confirmed from the presence of ordered white dots found on the alkali-treated CS fiber surface as seen in SEM image (Figure 2d). The white dots were identified as corresponding to silica-rich material using energy dispersive X-ray spectroscopy (EDS) of the alkali-treated coir fiber that was already reported by Calado and Barreto [35]. Further, it indicates that the removal of cuticle waxy layer could expose a rough fiber surface with more orderly placed pits and white globular protrusions. The observed white globular protrusions found on the surface of the treated fiber were identified as a silica-rich material [35]. The mean (average) HRR was calculated from the integral of the HRR over the entire time period that heat was liberated by each of the specimen. However, as far as HRR is concerned, the mean (average) HRR value for ATC sample was 147.19 kW/m², whereas it rose to 182.24 kW/m² for UTC. Moreover, the ATC sample had burnt more slowly than UTC that was further confirmed from the value of time to flameout. The time to flameout value of ATC was 1178 s and for UTC 765 s. This result can be interpreted with thermal conductivity results which were presented in our previous work [36]. In general, a material with a lower thermal conductivity allows heat to be dissipated slowly through the bulk and the material will burn very slowly. Accordingly, the thermal conductivity of the ATC composite (0.159 W/m-K) was found to be lower than the UTC composite (0.234 W/m-K).

The improved surface roughness of the fiber was observed due to the removal of amorphous hemicelluloses, fatty acids and their condensation products that form the waxy cuticle layer after the alkali treatment. This in turn can increase the surface-to-contact ratio between the fiber and the matrix and also enhances the mechanical interlocking capability. Thus it can provide better adhesion between fiber/matrix interfaces. In alkali-treated condition, the removal of organic impurities can form the openings between the thin fibers which form spider-web-like porous structure. The liquid PE resin holds the alternative fiber mats through these openings and hence attributes to the physical coupling of matrix between the fiber mats after cross-linking. It was also noticed during the experiments that the CS fiber samples were delaminated and buckled outwards. It could be possibly due to the tendency of fiber separation from the interlocked matrix. The results of the same are visible in the photograph of CS/PE composite in its initial shape and the delaminated burned samples as shown in Figure 4(a and b). The samples also experienced shrinkage of approximately 20% in the horizontal plane. The CS contained small pores at the surface during the alkali treatment which facilitated the impregnation of PE in them. The easy propagation of fire in the matrix that was present in-between the pores of the fiber tend to increase the fire rate. Hence, the decomposition of interlinked fiber resulted, which was reflected in a slight increase of HRR value compared to the silane-treated reinforced composites. But that was overcome by the deposition of the silane-coupling agent over the fiber surface. OPEN IN VIEWER

The silane-coupling agent treatment on CS facilitates a better bonding than the untreated CS with PE when used in the composites. Further, the PE/coupling agent-treated CS composites experienced higher flame retardancy when compared to the composites with NaOH-treated CS. Deposition of silane formed small patches over the fiber and reduced the effective surface-to-surface contact area between the sheaths. The deposition of silane formed a chemical bond between the fiber surface and the matrix through a siloxane bridge. In general, during the fire, the polymer layer present at the surface tends to burn first and the siloxane layer acts as a barrier improving the flameout time of the composite. After the burning of siloxane layer it results in the formation of char which acts as an insulating barrier at the interface of the composite thus resulting in the decreasing of HRR. The time to ignition and the time to flameout also played an important role in the HRR since the PE WOC took 51 s to start ignition and for UTC WOC, STC WOC and ATC WOC these values were 24 s, 38 s and 37 s, respectively. So, from these results, it can be inferred that when biodegradable fiber was incorporated into the polymer matrix, the time to ignition decreased. This might be due to exothermic nature of cellulose constituents in the natural fibers. Further, the time to flameout obtained from the cone calorimetry test was low for the PE matrix (435 s) and maximum for the silane-treated CS fiber-reinforced composites (1597 s).

Based on the earlier literature [37], the addition of higher weight percentage of nanoclay is often required for the enhancement of flame resistance as it offers significant advantages over the conventional composites. Figure 5(a–d) depicts the variation of HRR without and with the addition of 5 wt% of nanoclay to pure PE-, UTC-, ATC- and STC-reinforced nanocomposites, respectively. Nanoclay addition was found to have significant reduction in HRR on the conventional composite in all the conditions of the CS up to the time period of 600 s; beyond that a slight increase in pattern was noticed. This slight raise in HRR resulted due to the exposure of free fiber surface. However, the inclusion of clay in the nanocomposite greatly decreased the HRR and MLR due to its insulating behavior that acted as a barrier for mass transportation of decomposed volatile products [38–40]. It promoted the formation of the charred residue like a porous structure, that covered the fiber surface and acted as a protection barrier against the heat conduction and the mass transport, leading to an enhanced flame retardancy [41,42]. Further, inclusion of clay, in general, can also change the melt flow index of the polymer and the increase in the viscosity result is in a reduced flame spread. Even the clay agglomeration at higher wt% of nanoclay in the matrix could create the torturous path of fire advancement and it could increase the flameout time of composites. OPEN IN VIEWER

CO yield and MLR

One of the most important fire hazards is the evolution of the toxic gases such as CO and CO₂. Most of the deaths in a fire accident are caused by the inhalation of the combustion gases and the particulates and hence the monitoring and possibility of their production need to be considered. However, the gas products released by decomposing polymer composites depend on the chemical nature of the organic constituents, oxygen availability and the temperature of fire [43]. Figure 6 shows the peak values of (a) CO yield, (b) CO₂ yield for PE (C1), PE + 5 wt% (C2) , UTC WOC (C3), UTC + 5 wt% (C4), STC WOC (C5), STC + 5 wt% (C6), ATC WOC (C7) and ATC + 5 wt% (C8). From Figure 6(a), it can be inferred that the pristine PE emitted more CO than the composites. Generally, an increase in the HRR is attributed to the increase in the yield of CO gas [8]. When clay was incorporated into the pristine PE, the CO yield was reduced significantly. This shows clearly that the addition of clay to the polymer reduced the toxic CO emission. This behavior is as expected because the clay is inherently a fire-retardant and the clay particles present in the polymer improved the flameout time of the PE composite leading to reduced CO emission. However, with the addition of combustible CS fiber, the amount of CO got reduced when compared to PE. The lowest value of CO yield was observed for silane-treated CS composites, which in turn led to their complete combustion. So from Figure 6(a), it is clear that the addition of clay with the biodegradable CS fiber led to decrease in the emission of CO for all the cases. It could be due to the natural by-product obtained from the decomposition of the chemical constituents present in the fiber and the clay. OPEN IN VIEWER

The higher amount of CO₂ was evolved in the PE matrix as shown in Figure 6(b) and it may be due to the breakdown of the ester bonds within the resin [44]. In the case of the clay unfilled CS fiber-reinforced composites, the lower value of CO₂ was observed for UTC composites. However, unlike CO yield, the controlling mechanism of CO₂ yield was complex and depends on many factors other than just the HRR [32]. With the addition of clay, the CO₂ yield was found to decrease for UTC, STC and ATC fiber-reinforced composites. The PE showed the largest mass loss and the STC WOC had the least mass loss in all the clay unfilled composites as given in Table 1. The STC composites could provide a strong structural integrity by the enhanced adhesion between the fiber and the matrix which in turn restricted the rate of mass loss. In all the fiber-reinforced composites except the UTC, the addition of nanoclay decreased the MLR due to the protective barrier arising from the dispersion of the clay. In the UTC composites, the addition of clay could create a weak interface and thus led to more mass loss by separating the fiber from the matrix.

OPEN IN VIEWER

Table 1.

Cone calorimetry data for CS and CS/MMT nanocomposites at 50 kW/m².

Properties	Sample type
	PE WOC	PE + 5 wt%	UTC WOC	UTC + 5 wt%	STC WOC	STC + 5 wt%	ATC WOC	ATC + 5 wt%
Time to ignition (s)	51	68	24	49	38	68	37	50
Time to flameout (s)	435	584	765	1059	1597	1680	1678	1762
Peak HRR (kW/m2)	837.7	571.8	272.7	205.3	186.8	201.5	252.3	154.6
Mean HRR (kW/m2)	390.2	349.8	182.2	155.5	103.2	92.1	147.1	81.9
Total heat release (MJ/m2)	153.1	182.1	108.9	160.6	62.1	55.6	88.5	50.1
Total smoke release (m2/m2)	6718.7	7467.7	3391.9	4248.3	2398.2	1951.2	2927.9	1800.6
Mass lost (g)	99.2	97.2	85.4	104.8	60.1	68.3	81.1	56.6
Specific extinction area (m2/kg)	677.7	768.5	397.2	426.1	416.6	297.2	378.3	334.1
Total oxygen consumed (g)	122.3	142.0	87.2	134.9	53.7	49.5	75.1	44.4
Average specific MLR (g/(s m2))	27.8	22.2	14.2	9.9	9.5	10.4	12.1	8.5
Effective heat of comb. (MJ/kg)	15.3	18.5	12.6	15.9	10.6	8.3	11.3	9.1

TGA

The TGA measures the amount and the rate of change in the weight of a material as a function of temperature or time in a controlled atmosphere. The TGA results are usually displayed as curves of weight loss variation with temperature and referred to as thermograms. In general, the thermal analysis for the material is carried out in order to evaluate the chemical, physical and structural changes occurring in a material under an imposed change in the temperature. As mentioned, the thermal stability of the natural fibers apparently bears a correlation with their chemical constituents [45].

The thermal behavior of the CS/PE composite samples was investigated at a heating rate of 5℃/min from ambient temperature to 800℃ under nitrogen atmosphere at a flow rate of 60 ml/min. Figure 7(a) shows the thermal degradation pattern of the untreated, alkali-treated and silane–treated CS fiber-reinforced composites. The initial degradation temperature was higher for the untreated CS fiber. It is evident from the figure that the initial degradation (10 wt%) was started at 254.4 ℃ and final degradation was noticed at 490 ℃ for UTCs. At the same time, the initial degradation temperature of the alkali-treated composite was decreased to 120 ℃, whereas for silane-treated composites it rose to 238.7 ℃. This may be due to the changes in the rate of dehydration of the fibers after the chemical treatment. Figure 7(a) also reveals that the untreated CS covered by the waxes, greasy layer and the surface impurities restricted the rate of heat flow along the fiber orientation and caused the slow decomposition process which was observed from the larger slope in TGA curve. From the curve, the first range of weight loss which began from 30 to 110 ℃ was due to the loss of water content by the fiber. At about 200–350 ℃, the degradation occurred which was associated with the thermal decomposition of the lignin and the hemicelluloses. The final stage of weight loss occurred in the temperature range of about 350–400 ℃ indicating the degradation of α-cellulose and other non-cellulosic compounds [46]. During alkalization, it was observed that the CS fiber started degrading at 200 ℃ and displayed a lower thermal stability and it might be due to the less dehydration by the appearance of the amorphous cellulose and the lignin content [23]. It could have more hydroxyl groups from the hydrogen bonding of the lignin and the semisolid nature of amorphous cellulose. However, as far as the major weight loss (100 %) is concerned, there was a substantial enhancement in the thermal stability noticed for both alkali- and silane-treated composites. The alkali-treated composites could raise the final degradation temperature to 608.2 ℃, giving a charred residue of 4.8%, whereas for the silane-treated, it rose to 713.4 ℃, giving a charred residue of 3.2%. This could happen due to the better interfacial bonding and the breakage of decomposition products (cellulose, hemicelluloses, lignin and waxes) at the second stage leading to the formation of charred residue. OPEN IN VIEWER

The comparative thermal degradation behavior of the CS/PE composites (untreated, alkali and silane) between without and with 5 wt% of clay loading is shown in Figure 7(b–d). The TGA of all the composites followed two-stage decomposition behavior. Figure 7(b) shows the TGA for the unfilled and clay-filled UTC/PE hybrid nanocomposites. Initial thermal decomposition (10% of mass loss) of the UTC/PE composites without the addition of clay started at 254.4 ℃. Further, the addition of 5 wt% clay in the UTC composites did not show any improvement in decomposition and, in fact, it decreased the initial degradation temperature to 236 ℃. The formation of the agglomeration with 5 wt% of clay content could not penetrate into the surface of UTC due to the coverage of the inorganic contaminates (waxes and dusts). It could make a weak interface in the UTC composite and led to the increase in weight loss. However, for the ATC and the STC samples, the enhanced thermal stability was observed even for the addition of 5 wt% clay as shown in Figure 7(c and d). This could be due to the improved fiber/matrix adhesion resulting from the rough fiber surface after the removal of the waxy layers. A significant shift in the degradation temperature for the initial and the final decomposition was noticed for the alkali- and silane-treated composites with the addition of nanoclay. The enhancement in the thermal stability of the hybrid nanocomposites was predominant at higher temperature (>400 ℃) and it could happen due to the interaction between the disordered clay and PE matrix. In order to further confirm the increase in thermal stability of the composites due to the addition of nanoclay, the derivative thermogram of the PE/coupling agent-treated CS/nanoclay hybrid composite in the 200℃ to 430℃ was recorded and is presented in Figure 8. From Figure 8, it can be observed that two inflection temperatures (temperature at which the degradation rate is maximum) existed, the first one corresponding to the degradation of hemicellulose and lignin and the second one to the degradation of α-cellulose. Further, the second inflection temperature of the hybrid nanocomposite was higher than that of the PE/coupling agent-treated CS composite indicating the improvement of the thermal stability by nanoclay. Among all the composites, the final residue in the CS/PE clay-filled hybrid nanocomposites was comparatively higher than the clay-unfilled composites. This result had good agreement with the mass loss obtained from the cone calorimetry data results presented in Table 1. OPEN IN VIEWER