Fire-Resistant, Strong and Green Polymer Nanocomposites Based on

Research Article pubs.acs.org/journal/ascecg

Fire-Resistant, Strong, and Green Polymer Nanocomposites Based on Poly(lactic acid) and Core−Shell Nanofibrous Flame Retardants Jiabin Feng,†,⊥ Yiqi Sun,‡,⊥ Pingan Song,*,‡,§ Weiwei Lei,∥ Qiang Wu,‡ Lina Liu,‡ Youming Yu,*,‡ and Hao Wang*,§ †

China-Australia Institute for Advanced Materials and Manufacture, Jiaxing University, 56 Yuexiu South Road, Jiaxing 314000, China Department of Materials, School of Engineering, Zhejiang A&F University, 88 Huancheng North Road, Hangzhou 311300, China § Center for Future Materials, University of Southern Queensland, West Street, Toowoomba, Queensland 4350, Australia ∥ Institute for Future Materials, Deakin University, Locked Bag 20000, Geelong, Victoria 3220, Australia ‡

S Supporting Information *

ABSTRACT: Despite extraordinary mechanical properties and excellent biodegradability, poly(lactic acid) (PLA) still suffers from a highly inherent flammability, restricting its wide applications in the electric and automobile fields. Although a wide range of flame retardants have been developed to reduce the flammability, so far, they normally compromise the mechanical strength of PLA. Herein, we have demonstrated the fabrication of a novel core−shell nanofibrous flameretardant system, PN-FR@CNF, through in situ chemically grafting the phosphorus−nitrogen-based polymer onto the cellulose nanofiber (CNF) surface. The results show that adding 10 wt % PN-FR@CNF enables PLA to achieve a V-0 flame resistance rating during vertical burning tests and to exhibit a dramatically reduced peak heat release rate in cone calorimetry measurements, indicating a significantly reduced flammability. In addition, the tensile strength of PLA also increases by around 24% (about 72 MPa). This work offers an innovative methodology for the design of the unique integration of extraordinary flame retardancy and mechanical reinforcement into one hierarchical nanostructured additive system for creating advanced green polymeric materials. KEYWORDS: Core−shell, Nanofibrous, Flame retardant, Poly(lactic acid) (PLA), Mechanical property

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the flammability of PLA.6−15 However, the resultant flameretardant PLA composites normally show compromised mechanical strength relative to the PLA bulk. In other words, these FRs normally improve the flame retardancy of PLA at the expense of the mechanical strength. For example, the combination of distiller’s dried grains and resorcinol bis(phenyl phosphate) could increase the flame retardancy, but simultaneously reduce the mechanical strength of PLA by around 15%.11 Song et al. reported that although the ammonium polyphosphate (APP)/poly(ethylene glycol) 6000 system enabled PLA to achieve a V-0 rating, the tensile strength was reduced by 27%.9 Meanwhile, they also found a slight decrease in the tensile strength relative to the PLA host, although adding only a 1.0 wt % concentration of a small molecular flame retardant (P-AA) could make PLA pass the V-0 rating.10 Apparently, these flame retardants usually lead to a certain decrease in the tensile strength of PLA in spite of exceptional flame-retardancy efficiency probably because of the poor

INTRODUCTION Nowadays, biobased polymers derived from renewable resources have increasingly replaced traditional petrochemical nonbiodegradable counterparts in many fields due to excellent biodegradability and mechanical properties.1−3 Among these green polymers, poly(lactic acid) (PLA) is regarded as one of most promising polymers due to its mechanical strength and stiffness being comparable to those of most conventional thermoplastics, including polypropylene and polycarbonate in addition to the high degree of transparency.2−5 These features enable PLA to gradually find extensive applications ranging from the electronic and electrical field as the casing of cell phones and computers to the automobile and aerospace sectors as the interior decoration materials.3−5 Unfortunately, PLA is inherently flammable with a low limited oxygen index (LOI) of only 19 vol % and tends to drip upon being ignited, extremely restricting its applications in the fields above where excellent flame-retardancy performance, typically a V-0 rating during UL94 tests, is required.3−5 Until now, a wide variety of flame retardants (FRs), particularly phosphorus-based FRs, have been successively developed, and some FRs are found highly efficient to reduce © 2017 American Chemical Society

Received: May 7, 2017 Revised: July 18, 2017 Published: July 25, 2017 7894

DOI: 10.1021/acssuschemeng.7b01430 ACS Sustainable Chem. Eng. 2017, 5, 7894−7904

Research Article

ACS Sustainable Chemistry & Engineering interfacial adhesion with the PLA matrix and the weak mechanical property itself. Therefore, much work has been recently performed to improve the mechanical property of the flame-retardant PLA composites by introducing the reinforcement components or preparing novel flame-retardant additives with one-dimensional (1D) nanostructure.16−22 A typical example is that Yu et al. added a 5 wt % concentration of functionalized carbon nanotubes into PLA composites filled with 10 wt % ramie; the reinforcement effect was however not observed due to weak interfacial interactions.15 Fox et al. have recently investigated the effect of nanofibrious cellulose (NFC) or modified NFC on APP−flame-retarded PLA composites, but the tensile strength of PLA composites dramatically decreased because of the strong polarity of both APP and NFC and their poor dispersion.16 Meanwhile, Qian et al. have recently reported that adding a 0.5 wt % concentration of rodlike aluminated mesoporous silica (Al-SBA) could impart a V-0 rating flame retardancy to PLA and increase the tensile strength to 60 MPa.21 Another recent work has reported that a 20 wt % concentration of a nanorod-shaped flame retardant based on aluminum hydroxide and benzenephosphinic acid could increase the tensile strength of PLA by around 23% in addition to giving it a V-0 rating when the loading level is above that.22 It seems that these 1D flame retardants however still have some key issues, such as the time-consuming fabrication and high loading levels (≥20 wt %). Therefore, the creation of 1D nanostructured flame retardants remains challenging despite their unique combination of exceptional flame retardancy and mechanical reinforcement. Recently, cellulose nanofiber (CNF) isolated from plants has been considered to be one of most potential reinforcing candidates for polymers owing to the outstanding mechanical strength and stiffness as well as high length/diameter ratio.23−25 Therefore, we herein attempt to adopt CNF as the reinforcing template for synthesizing a novel high-performance nanofibrous flame retardant by in situ chemically grafting the flameretardant shell from the CNF surface. The as-prepared flame retardant (PN-FR@CNF) exhibits exceptional flame retardancy and mechanical reinforcement effects in PLA. Especially, the addition of 10 wt % PN-FR@CNF enabled PLA to achieve a V0 flame retardancy rating and simultaneously to increase the tensile strength by 24%. This work offers a novel methodology to create hierarchical flame retardants combining exceptional flame retardancy and mechanical reinforcement functions by using the in situ template strategy.

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Scheme 1. Schematic Illustration for the Synthetic Route to Core−Shell Nanofibrous Flame Retardants, PN-FR@CNF

aqueous solution of NaOH. About 50 g of ECH was subsequently dropped into the solution, which was allowed to react for 12 h. When the reaction was over, the solution was centrifuged followed by washing with deionized water more than five times to remove all unreacted ECH until it turned neutral, and about 2.63 g of modified CNF, designated as E-CNF, was obtained. Then 2.6 g of E-CNF was dispersed into the deionized water to obtain a 1 wt % aqueous solution, into which 13.0 g of PEI was slowly added while the mixture was slowly heated to 80 °C. The reaction was stopped by cooling to room temperature after 8 h. The mixture was centrifuged and then washed with deionized water at least five times to remove all unreacted PEI. About 2.9 g of PEI-modified CNF (designated as NH2@CNF) was collected. As for the preparation of the target product, PN-FR@ CNF, typically 200 mL of an aqueous solution of NH2@CNF (2.0 g) was first adjusted to a pH value of ∼10.5, and both 0.0788 mol (∼2.36 g) of formaldehyde and 0.075 mol (∼10.35 g) of DEP were slowly dripped into this solution under violent stirring at a temperature of 85 °C. The reaction was terminated after 8 h, and the solution was treated by the same procedures as above. About 2.25 g of PN-FR@CNF was finally gained after the solution was calibrated. Nanocomposite Fabrication. PLA nanocomposites were fabricated by combining the solvent exchange and melt mixing approaches. To promote the dispersion of PN-FR@CNF in the polymer matrix, the aqueous suspension of PN-FR@CNF was solventexchanged with chloroform more than five times to completely remove the water. Meanwhile, a predesigned amount of PLA pellets was dissolved in chloroform and then mixed with the chloroform suspension of PN-FR@CNF under high-speed dispersion to produce a homogeneous mixture. The mixture was then allowed to dry by removing the solvent in the fume hood, leading to the preparation of the PLA/PN-FR@CNF master batch in which PN-FR@CNF accounted for 20 wt %. PLA nanocomposites were fabricated via melt compounding of the PLA resin and the PLA/PN-FR@CNF master batch using a ThermoHaake Torque rheometer at 170 °C for 10 min with a rotor speed of 60 rpm for each sample. In the designation of the samples, PLA/xPN-FR@CNF, x refers to the mass fraction of PN-FR@CNF in the nanocomposite; for instance, PLA/5.0PN-FR@CNF contains 95 wt % PLA and 5.0 wt % PN-FR@CNF. In addition, the PLA nanocomposite containing 5.0 wt % CNF (PLA/5.0CNF) was also prepared according to the same protocol for comparison. Characterization. FT-IR spectra were obtained on a Bruker Vector 22 FT-IR spectrometer with the KBr pellet pressing method. Xray photoelectron spectroscopy (XPS) spectra were obtained on a Thermo ESCALAB 250 spectrometer. X-ray diffraction (XRD) was carried out using a Rigaku X-ray generator (Cu Kα radiation with λ = 1.54 Å) at room temperature. Elemental analysis (EA) measurements were carried out on a Vario EL elemental analyzer (Elementar Analysensysteme GmbH, Germany). The phosphorus content was determined by inductively coupled plasma atomic emission specotro-

EXPERIMENTAL SECTION

Raw Materials. Poly(lactic acid) (PLA 4032D) was purchased from NatureWorks (Minnetonka, MN) with a density of 1.24 g/cm3. Cellulose nanofiber (CNF; KY100) solid was provided by Daicel Chemical Industries Co., Ltd. (Japan) and contained 25 wt % CNF and 75 wt % water. Other chemical agents, such as epoxy chloropropane (ECH), polyethylenimine (PEI) (Mw = 10000), diethyl phosphite (DEP), and formaldehyde (CH2O; 37%) were purchased from Aladdin Chemical Agent Co., Ltd. Sodium hydroxide (NaOH), hydrochloric acid (HCl; 37.5%), chloroform, and ethanol were analytical grade and were used as received without further purification. Synthesis of PN-FR@CNF. The schematic representation for the synthetic route to PN-FR@CNF is shown in Scheme 1. First, 2.5 g of CNF was dispersed into 247.5 mL of deionized water with the aid of high-speed dispersion (T18 IKA, Germany) to produce a 1.0 wt % transparent aqueous mixture of CNF. The mixture was then heated to 85 °C while the pH value was adjusted to ∼12 using a 1 mol/L 7895


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ACS Sustainable Chemistry & Engineering

Figure 1. (A) IR spectra, (B) XRD patterns, (C) TGA curves in nitrogen conditions, and (D) XPS surveys of PN-FR@CNF and its precursors. XPS surveys of the N 1s core of (E) NH2@CNF and (F) PN-FR@CNF. Ti refers to the initial degradation temperature. scopy (ICP-AES) on a PROFILE SPEC spectrometer (Leeman, United States). The morphologies of CNF and PN-FR@CNF as well as their dispersion in the PLA matrix were observed using transmission electron microscopy (TEM; JEM-1200EX). Scanning electron microscopy (SEM) images were recorded on an S4800 (FEI-SEM, Japan) at an accelerating voltage of 5 kV. The relative element content in the char residue was determined by energy-dispersive X-ray analysis (EDAX) using the mapping mode. Thermogrametric analysis (TGA) tests were performed on a TA SDTQ600 (TA Instruments) thermogravimetric analyzer. About 8.0 mg of sample was heated from room temperature to 700 °C at a heating rate of 20 °C under a N2 or an air atmosphere. Raman spectroscopy was performed on a Nicolet Almega dispersive Raman spectrometer (Thermo Scientific) at 514 nm. The flammability of samples with a size of 100 × 100 × 3.0 mm3 was evaluated using a cone calorimeter performed in an FTT UK device according to ISO 5660 with an incident flux of 35 kW/m2. Typical results from the cone calorimeter are reproducible to within 5%, and the data reported here were the means of triplicate experiments. The vertical burning tests (UL-94) were performed on a vertical burning test instrument (CZF-2-type) (Jiangning, China), and the samples were 130 × 13 × 3 mm3 according to measurement standard ASTM D3801. Limited oxygen index (LOI) values were determined using an HC-2 oxygen index instrument on 120 × 6 × 3 mm3 sheets according to the standard oxygen index test ASTM D2863. Tensile properties of the samples were determined on a WD-5 electronic universal tensile tester with a cross-head rate of 5.0 mm/ min, and all samples were measured in quintuplicate. The Charpy impact strength of the composites was measured by means of a Charpy impact test, according to the specification GB/T 1043-93 using unnotched rectangular specimen testing. The dimensions of the specimens were 80 × 10 × 4 mm3. At least 10 tests were conducted for each sample, and the results were averaged. Dynamic mechanical analysis (DMA) was conducted using a dynamic mechanical analyzer (DMA242C, TA) at a frequency of 1 Hz with a heating rate of 3 °C/ min. The temperature range was from −40 to +130 °C. The glass transition temperature (Tg) was determined from the peak of loss factor (tan δ)−temperature plots. To obtain the morphology information on both CNF and PN-FR@ CNF, a total of over 100 nanofibers or particles were chosen. These micrographs were analyzed with the aid of image analysis software, and the diameters of both CNF and PN-FR@CNF were calculated. The number-average diameter (Dn) of the nanofibers was calculated according to the following equation:

Dn =

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∑ ND i i / ∑ Ni

(1)

RESULTS AND DISCUSSION Characterization of PN-FR@CNF. Many analytical techniques are utilized to characterize the chemical structure, crystal changes, thermal property, chemical composition, and morphology of the target product, PN-FR@CNF. As shown in Figure 1, compared with the pristine CNF, the characteristic absorption peak of epoxy groups is clearly observed at 930 cm−1 in the FTIR spectrum of E-CNF, indicating that the epoxy groups are successfully introduced onto the CNF surface.24,27 As for NH2@CNF, two new weak stretching vibration peaks of amino groups (N−H) located at 3351 and 3285 cm−1 appear as compared with the stretching vibration of O−H groups (υO−H) at 3341 cm−1 in the IR spectra of both CNF and E-CNF,27 besides the disappearance of the absorption band of epoxy groups. In addition, the rocking vibration of C−N (δC−N) is also found at 1577 cm−1, and these changes suggest that PEI is chemically grafted onto the surface of NH2@CNF. In the case of PN-FR@CNF, several new absorption peaks, such as 906 cm−1 (υP−C), 1061 cm−1 (υP−O/C−O), and 1208 cm−1(υPO), are attributed to the Mannich reaction between DEP, formaldehyde, and NH2@ CNF, as also evidenced by the disappearance of the stretching vibration of N−H groups (see Figure 1A).29−32 Thus, the IR results strongly indicate the successful synthesis of the target product, PN-FR@CNF. X-ray diffraction patterns are used to investigate the crystal structure change of CNF after chemical grafting (Figure 1B). The pristine CNF shows four typical diffraction peaks at 2θ = 14.6°, 16.5°, 22.6°, and 34.1°, respectively, corresponding to (101), (101̅), (002), and (040) crystal planes of type I cellulose.33,34 It also shows another peak at 27.5° which is ascribed to the characteristic diffraction peak of hemicellulose, indicating the presence of a small amount of hemicellulose. As for E-CNF, the XRD pattern clearly shows that four peaks of cellulose remain and the relative intensity has no obvious change. However, the peak of hemicellulose disappears because 7896


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Table 1. Chemical Composition of PN-FR@CNF and Its Precursors Obtained from XPS, Element Analysis (EA), and ICP-AESa

a

run

[C] (wt %)

[O] (wt %)

[H] (wt %)

[N] (wt %)

[P] (wt %)

CNF E-CNF NH2@CNF PN-FR@CNF

45.6a/45.7b 46.1a/46.2b 46.6a/46.5b 44.4 a/43.8b

54.4a/49.1b 53.9a/48.5b 45.0a/45.1b 46.5a/43.0b

5.2 b 5.3b 5.5b 5.3b

4.2a/3.9b 3.6a/3.3b

5.5a/4.6c

Superscripts a, b, and c refer to the results obtained by the XPS, EA, and ICP-AES methods, respectively.

(−NH−), and a 25% content of tertiary amino groups in [email protected],30,31 In comparison, the relative contents of active amino groups, −NH2 and −NH−, respectively, decrease to 32% and 6% accompanied by a significant increase in the tertiary amino groups, about 62% in the N 1s survey of PNFR@CNF (Figure 1F).26,29,36 These changes strongly indicate the Mannich addition reaction between PEI, CH2O, and DEP and the successful synthesis of PN-FR@CNF. TEM images of CNF and PN-FR@CNF allow us to visually observe the morphology changes. As shown in Figure 2, the

of the strong alkali treatment. After long PEI chains are grafted onto the CNF surface, namely, NH2@CNF, besides the (002) crystal plane peak of the cellulose, the intensity of the other three characteristic peaks become too weak to be observed. Meanwhile, four new strong diffraction peaks at 2θ = 26.8°, 31.1°, 45.1°, and 56.1° appear, probably arising from the crystallinity peaks of PEI chains, suggesting the PEI chains are successfully introduced onto the CNF surface.35,36 As for the target product, PN-FR@CNF, its diffraction pattern displays four typical peaks of type I cellulose (similar to E-CNF) but accompanied by two new diffraction peaks at 2θ = 29.0° and 39.1°. These changes are likely due to the Mannich reaction between PEI and DEP, which considerably alters the crystallinity of PEI. Thus, the XRD results provide further evidence for the fabrication of PN-FR@CNF. TGA analysis is also employed to evaluate the thermal property of CNF before and after functionalization, as shown in Figure 1C. In nitrogen, pure CNF shows a very high thermal stability with an initial thermal degradation temperature (Ti) of 309 °C, but leaves a small amount of char residue of only 12.2 wt % at 700 °C, which is similar to that in the previous report.33 After introducing epoxy groups, the E-CNF displays a much lower Ti of only 183 °C than CNF, probably because of the lower thermal stability of epoxy groups on its surface. However, the hemicellulose removal enables the concentration of the char residue to be strikingly enhanced to 58.1 wt %. As for the NH2@CNF, Ti and the char residue concentration decrease to 118 °C and 24.7 wt %, respectively, due to the much lower thermal stability and char-forming capability of PEI. With regard to the PN-FR@CNF, Ti and the char residue concentration respectively rebound to 253 °C and 36.2 wt % because of the Mannich reaction and high char-forming capability of DEP. Interestingly, PN-FR@CNF also shows a high Ti of 267 °C in air (see Figure S1, Supporting Information). Such high thermal stability enables as-prepared PN-FR@CNF to basically bear the melt processing of most polymers, and the high char-forming capability can contribute to its flame retardancy in the condensed phase. XPS measurements allow us to determine the chemical composition and bond types of PN-FR@CNF and its precursors (see Figure 1D), and detailed results are listed in Table 1. The XPS results clearly show that E-CNF basically has a chemical composition similar to that of CNF (C, 46 wt %; O, 54 wt %; H, 5.2 wt %). As for NH2@CNF, 4.2 wt % N is determined besides C and O. By contrast, 3.6 wt % N and 5.5 wt % P are also found in the PN-FR@CNF, indicating the presence of phosphorus, agreeing well with the FTIR results. It is evident that the XPS results are consistent with the elemental analysis and ICP-AES results. The deviations of the relative element composition for PN-FR@CNF are attributed to different analysis methods. N 1s XPS patterns allow us to evaluate the relative content of different bond types of N elements. Figure 1E shows a 65% content of primary amino groups (−NH2), a 10% content of secondary amino groups

Figure 2. TEM images of (A, B) pristine CNF and (C, D) PN-FR@ CNF showing the morphology change before and after chemical grafting.

pure CNF clearly exhibits an entangled fibrous structure with dimensions of 2−20 μm in length and 20−100 nm in diameter (Dn = 60 ± 40 nm) (see Figure 2A,B). Despite the weblike structure, both the profile and boundary of individual nanofibers are readily distinguishable due to their high crystallinity.37 By contrast, for PN-FR@CNF, there is an amorphous PN-FR polymer phase coating the CNF surface, which makes it difficult to identify the boundary of the nanofibers (see Figure 2C). Upon observation at a higher magnification, the PN-FR@CNF displays a maximum thickness of around 280 nm in some regions of the fibers, and the nanofiber core is only about 100 nm thick, which means the thickness of the PN-FR polymer phase or the polymer layer (shell) can reach up to 180 nm (see Figure 2D). PN-FR@CNF shows a statistical number-average diameter (Dn) of 200 ± 80 nm, much higher than the Dn of pristine CNF, indicating the functionalization of CNF. In spite of that, it is obviously observed that there is great diameter variability for the as7897


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ACS Sustainable Chemistry & Engineering prepared core−shell nanofibrous flame retardant because the graft reaction occurs randomly on the surface of CNF (see Figure S2, Supporting Information). In other words, the PN polymer shell cannot uniformly coat the entire CNF fiber surface, thus leading to significant differences in the fiber diameter of the different zones of PN-FR@CNF. Despite a statistical diameter ranging from 120 to 280 nm, the shell thickness of PN-FR@CNF varies noticeably along the fiber. Therefore, there is some inhomogeneity and variability for the core−shell nanofibrous flame retardant. Similar morphology changes can also be observed in the atomic force microscopy (AFM) images (see Figure S3, Supporting Information). Thus, the microscopic morphology observations strongly indicate that we have successfully synthesized one novel core/shell-like nanofibrous flame retardant. Dispersion of PN-FR@CNF. It is widely recognized that the dispersion of nanomaterials in the polymer host directly determines the comprehensive performances of the resultant nanocomposites. As presented in Figure 3A,A1, the pristine

and PN-based polymer of weak polarity on the CNF surface.39 However, when the loading level of PN-FR@CNF increases to 10 wt %, some small agglomerates with a size of 2−5 μm appear, as presented in Figure 3C. Moreover, the statistical Dn values of CNF nanofibers are 200 ± 80 nm in the PLA/5PNFR@CNF and 400 ± 200 nm in the PLA/10PN-FR@CNF. At high magnification, the PN-based polymer phase (marked by the light blue arrow) and the individual CNF (marked by the blue arrow) can be identified at higher magnifications (Figure 3B1,C1). Moreover, as clearly seen in Figure 3B1, the fibrous dispersion phase displays an average diameter of around 260 nm, close to that of the PN-FR@CNF. This indicates that the PN-based polymer shell enables some individual PN-FR@CNF nanofibers to disperse within the PLA matrix by serving as the effective interfacial compatibilizer between PLA and CNF in addition to offering the flame-retardancy function. Flame Retardancy of PLA Nanocomposites. Cone calorimeter measurements are used to evaluate the flameretardancy efficiency of as-fabricated PN-FR@CNF in the PLA host since the cone calorimeter is able to simulate a real fire on a reduced scale.40−42 Figure 4A shows the heat release rate of

Figure 4. (A) Heat release rate and (B) normalized mass loss curves of PLA and its nanocomposites based on CNF and PN-FR@CNF at an incident heat flux of 35 kW/m2.

PLA and its nanocomposites during the cone test, with key parameters summarized in Table 2. Apparently, the PLA bulk shows a time to ignition (tign) of 69 s, burns very quickly after being ignited, reaches a peak heat release rate (pHRR) of 443 kW/m2, and finally generates a total heat release (THR) of 65.2 MJ/m2 (see the THR curves in Figure S4, Supporting Information). Meanwhile, the combustion process is accompanied by dripping, indicating the high flammability of PLA. After addition of 5 wt % CNF, although both tign and the heat release rates (pHRR and THR) of the resultant nanocomposites slightly decrease, the dripping phenomenon disappears due to the suppression action of the high length/ diameter ratio CNF on the melt flow of the PLA bulk during the cone tests. Similarly, adding PN-FR@CNF at a loading level of above 5 wt % also avoids the dripping. Interestingly, the tign sharply decreases by 12 s for PLA/5PN-FR@CNF and then steadily increases to 65 s as compared with that of the PLA host. Moreover, both pHRR and THR significantly decrease with increasing loading levels of PN-FR@CNF, for instance, a pHRR and THR of ∼304 kW/m2 and ∼60 MJ/m2 for PLA/ 10PN-FR@CNF and 270 kW/m2 and 57 MJ/m2 for PLA/ 15PN-FR@CNF, respectively. This means that incorporating 10 and 15 wt % PN-FR@CNF enables the pHRR value to decrease by 31% and 39%, respectively, relative to that of the PLA bulk, strongly suggesting the considerable reduction in the

Figure 3. Typical TEM images of (A, A1) PLA/5CNF, (B, B1) PLA/ 5PN-FR@CNF, and (C, C1) PLA/10PN-FR@CNF showing the CNF dispersion in the PLA matrix.

CNF basically exists in the form of large agglomerates (2−5 μm in length and 0.5−2.0 μm in diameter) in spite of the fact that CNF seems to disperse well in the PLA matrix. This poor dispersion is primarily due to the poor interfacial compatibility caused by the large polarity differences between the weak polar PLA and strong polar CNF with abundant hydroxyl groups on the surface.17 In comparison, for PLA/5PN-FR@CNF-based nanocomposites, it is apparent that the majority of PN-FR@ CNF fibers are able to homogeneously disperse within the PLA matrix, as shown in Figure 3B, which is because of the significantly improved interfacial compatibility between PLA 7898


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ACS Sustainable Chemistry & Engineering Table 2. Detailed Combustion Parameters Obtained from Cone Calorimeter and UL-94 Tests run PLA PLA/5CNF PLA/5PN-FR@CNF PLA/10PN-FR@CNF PLA/15PN-FR@CNF

tigna (s) 69 64 57 59 65

± ± ± ± ±

1 1 1 1 1

pHRRa (kW/m2) 443 381 339 304 270

± ± ± ± ±

25 22 20 18 15

THRa (MJ/m2) 65 62 61 60 57

± ± ± ± ±

1.2 1.3 1.1 1.2 1.0

AMLRa (g/s) 0.13 0.10 0.056 0.050 0.047

± ± ± ± ±

0.02 0.01 0.005 0.006 0.005

[char]a (wt %)

dripping (Y/N)

UL-94

± ± ± ± ±

Y N N N N

NR V-1 V-1 V-0 V-0

0.87 0.82 4.0 5.7 12

0.1 0.1 0.4 0.5 0.6

a

tign, pHRR, THR, AMLR, and [char] respectively represent the time to ignition, peak heat release rate, total heat release, and char residue concentration after the cone tests.

flammability of PLA. In comparison, adding a 20 wt % concentration of the mixture of phosphorus-containing compounds and micro/nanocrystalline cellulose (MCC/ NCC) only leads to a ∼20% reduction in the pHRR of PLA.43 Similar phenomena were observed in other polymer nanocompsites based on carbon nanotubes.3,39 In addition, the average mass loss rate (AMLR) also exhibits the same change trend as the heat release rate. In other words, the pure CNF basically does not affect the mass loss rate of PLA, whereas the AMLR monotonously decreases with increasing PN-FR@CNF loading levels, as shown in Figure 4B and Table 2. For instance, as compared with an AMLR of 0.13 g/s for the PLA matrix, the AMLR decreases to 0.056 and 0.047 g/s when the loading levels of PN-FR@CNF are 5 and 15 wt %, respectively. The remarkable reduction in AMLR indicates that the combustion of PLA strikingly turns slow in the presence of PN-FR@CNF. Likewise, the char residue concentration after the cone tests also increases monotonously with the increasing loading level of PN-FR@CNF and reaches 12 wt % for the PLA/15PN-FR@CNF system. Char residue in such high concentrations can act as the thermal barrier for both external heat flux and heat feedback, thus effectively protecting the underlying polymer bulk and improving the flame retardancy of PLA. This can be clearly evidenced by the positive relation between the pHRR reduction and the char concentration increase. As expected, UL-94 test results show that the PLA fails to pass the UL-94 test (no rating); adding 5 wt % CNF or PN-FR@CNF only enables PLA to achieve a V-1 rating. However, when the loading level of PN-FR@CNF is above 10 wt %, a V-0 rating is achieved, which means that the PLA nanocomposites can satisfy the industrial fire-retardancy requirements. In comparison, 20 wt % phosphorylated cellulose (MCC-P) was reported to allow PLA to pass a V-0 rating.43 In addition, LOI values also show a trend similar to tha tof the UL-94 results; 15 wt % PN-FR@CNF increases the LOI from 19.2 for the PLA matrix to 28.0 (see Table S1, Supporting Information). Thus, as-designed PN-FR@CNF is one highly effective flame retardant for polymers. Viscoelastic Behavior of PLA Nanocomposites. The viscoelastic behavior of polymer nanocomposites is extensively employed to examine the dispersion of nanomaterials within the polymer matrix and to understand the physical network that is closely correlated with the flame-retardancy property, specifically the peak heat release rate.39,42,43 As shown in Figure 5, the PLA host basically displays a typical viscoelastic behavior of a linear polymer; namely, both the storage modulus (G′) and loss modulus (G″) steadily increase whereas the complex viscosity (η*) decreases with increasing frequency. Unexpectedly, adding 5 wt % pure CNF leads to a noticeable decrease in the three viscoelastic parameters G′, G″, and η*. This interesting behavior is probably due to the strong repulsion interactions between large strong polar CNF

Figure 5. Frequency (ω) dependency of the (A) storage modulus (G′), (B) loss modulus (G″), and (C) complex viscosity (η*) and the linear relationship between the pHRR ratio and the G′ and η* ratios.

agglomerates (about 2−5 μm) and the PLA chains. This repulsion action can create a certain volume in the interface phase, which in turn provides some free space for PLA chains to easily move. By contrast, it is clear that, upon adding PNFR@CNF, all G′, G″, and η* parameters sharply increase, and nearly increase by 3 orders of magnitude especially at a frequency ω of below 0.1 rad/s relative to those of both the PLA host and PLA/5CNF. This evidently indicates a strong mechanical reinforcement effect of PN-FR@CNF on the PLA host (modulus, G′) and the mobility restriction of PLA chains during the melt state, thus leading to a dramatic increase in the melt viscosity (η*). More importantly, a so-called “second plateau” apparently appears in the low-frequency region once PN-FR@CNF is incorporated into the PLA host even if at a loading level of only 5 wt %. This plateau has been widely believed to be an indication for the formation of the physical networklike structure, which is extensively reported in many polymer nanocomposites based on nanomaterials such as clay and carbon nanotubes.39,42,43 The physical network not only fully reflects the uniform dispersion of nanofibrous PN-FR@CNF within the PLA matrix, agreeing well with the TEM observation 7899


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transformed into NH3 and completely released into the gas phase during combustion. In addition, the XRD results show that the PN-based polymer degrades before CNF to form a char layer which can provide thermal protection for the CNF core, whereas the CNF decomposes completely without the protection of the PN polymer, as evidenced by their Raman spectra (see Figure S5, Supporting Information). The above comprehensive analysis of char residues enables the establishment of possible char-forming models for PLA/ CNF and PLA/PN-FR@CNF systems. As presented in Figure 7, the PLA/CNF system (see Figure 7A) burns rapidly upon ignited and releases a lot of heat, and finally, only 0.8 wt % char is left because of the flammability of both PLA and CNF. Moreover, this char shows a discontinuous fragmented islandlike structure, as seen in the digital image. Therefore, CNF has a limited effect on the flammability of PLA. In comparison, for the PLA/PN-FR@CNF system (see Figure 7B), once it is ignited, the presence of uniformly dispersed PNFR@CNF dramatically increases the melt viscosity of the PLA nanocomposites, reduces the heat release rate, and slows the combustion process, thus remarkably reducing the flammability of PLA. Meanwhile, because of the lower thermal stability and high char-forming capability, the PN-FR@CNF shell degrades ahead of PLA and CNF and forms an intact char layer on the CNF surface. As a consequence, this enables part of the CNF core to leave and form a fibrous char, which can contribute to maintenance of the structural integrity of the char, as evidenced by the digital photo. Therefore, PN-FR@CNF can noticeably reduce the flammability of PLA and be an efficient flameretardant system for polymers. Mechanical Performances. Another key function of asdesigned PN-FR@CNF is the reinforcement effect on the PLA host. The mechanical performances of the resultant green polymer nanocomposites are evaluated by classic mechanical measurements. As shown in Figure 8 and Table 3, the PLA bulk shows an ultimate tensile strength (σt) of ∼58.0 MPa, an elastic modulus (Et) of ∼2.5 GPa, and a strain at break (ε) of only 8.0% without any yield behavior observed, indicating the typically brittle but strong characteristic of PLA. After addition of 5.0 wt % CNF, both σt and Et only marginally increase whereas ε slightly decreases. Such negligible mechanical changes are mainly due to the poor dispersion of CNF in the PLA matrix and the weak interfacial bonding. In comparison, as expected, incorporating the PN-FR@CNF significantly enhances the mechanical performances of PLA since the uniform dispersion and strong interfacial adhesion facilitate the effective transfer of external loads from the polymer matrix to the highstrength CNF.34,39,46 For instance, the presence of 5.0 wt % PN-FR@CNF enables the σt and Et to increase to 65.4 MPa and 2.68 GPa, respectively, in spite of the slight reduction in ε. When the loading level of PN-FR@CNF is 10 wt %, σt and Et are respectively enhanced by 24% (∼71.8 MPa) and 12% (∼2.75 GPa). Meanwhile, the tensile toughness (τ), calculated by eq 2,28,47 also increases from 2.73 MJ/m3 for the PLA bulk

in Figure 3B,C, but also strongly indicates that the high length/ diameter ratio PN-FR@CNF is able to create a threedimensional (3D) physical cross-link network within the polymer host. In addition, previous work has reported that such a physical cross-link network has a close relationship with the heat release rate during the cone tests. As expected, there nearly exists a linear relationship between η*c/η*m and pHRRc/ pHRRm, despite a relatively large deviation between G′c/G′m and pHRRc/pHRRm, as presented in Figure 5D. This means that the higher the complex viscosity, the lower the pHRR or the lower the flammability of the PLA nanocomposites, which is consistent with the results of previous reports.44,45 Analysis of the Char Residue. It is widely recognized that phosphorus−nitrogen-based intumescent flame retardants play the critical part in the condensed phase, and thus, examining the structure, morphology, and chemical composition of the char residue after the cone tests can contribute to the understanding how PN-FR@CNF works in the PLA host in fire. Although the char residue of PLA/CNF seems to still be of structural integrity, it shows a discontinuous structure with many voids or holes left (due to bubbling during combustion) and even large cracks at a higher magnification (see Figure 6A,A1). The EDX results indicate that 77 wt % carbon and 23

Figure 6. SEM images and corresponding chemical composition of the char residue for (A, A1, A2) PLA/5CNF and (B, B1, B2) PLA/5PNFR@CNF together with (C1, C2, C3) EDX mapping patterns for PLA/ 5PN-FR@CNF.

wt % oxygen are detected in the char. In comparison, the PLA/ 5PN-FR@CNF leaves an intact but fibrous-network-like char after the cone tests, and the fibrous char layer (marked by blue arrows) is probably partly inherited from the fibrous CNF, as can be clearly observed at a higher magnification (see Figure 6B,B1). This fibrous structure can strengthen the structural integrity and provide the mechanical strength for the char. Moreover, besides C and O, about 5.6 wt % P is also found in the char (see Figure 6B2), which can be visually confirmed by EDX mapping patterns (see Figure 6C1,C2,C3). In addition, the N element cannot be detected mainly because it is basically

εx = εb

τ=

∑ εx = 0

σεx (2)

to 3.11 MJ/m3, an increase of 14%, indicating an improved mechanical toughness. However, the further increasing loading level of PN-FR@CNF only leads to an improved elastic modulus, instead of simultaneously resulting in increased σt and 7900


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Figure 7. Schematic representations for the char-forming process for (A) PLA/5PN-FR@CNF and (B) PLA/5CNF during burning as well as their digital photos after cone tests at a heat flux of 35 kW/m2.

noticeably enhances the modulus in the whole temperature region (see Figure S6, Supporting Information). In brief, the tensile data fully demonstrate that as expected PN-FR@CNF is able to considerably improve the tensile properties of PLA. In addition, the moduli of the PLA nanocomposites are theoretically simulated by using the semiempirical Halpin− Kardos model modified by Van Es for high aspect ratio fiber composites (see eqs 3−7),48,49 and the predicted values are then compared with the experimental ones. E = Em

E⊥ = Em

1 + η ζϕf 1 − η ϕf

(3)

1 + 2η⊥ϕf 1 − η⊥ϕf

(4)

where

Figure 8. (A) Typical tensile stress-strain curves, (B) tensile strength, (C) elastic modulus, and (D) impact strength as a function of PNFR@CNF loading level for the PLA nanocomposites.

τ, which is likely due to the agglomeration of PN-FR@CNF at such a high loading level. The agglomeration is visually evidenced by the TEM image in Figure 3B. In addition, DMA results also show a trend similar to that of the tensile results. In other words, the addition of CNF only leads to a slight increase in the storage modulus, whereas adding PN-FR@CNF

E =

E f / Em − 1 E f / Em + ζ

(5)

E⊥ =

E f / Em − 1 E f / Em + 2

(6)

ζ = (0.5L /w)1.8

(7)

Ec = 0.184E + 0.816E⊥

(8)

E∥ and E⊥, respectively, refer to the longitudinal and transverse elastic moduli of the unidirectional composites. ϕf and ζ are the volume fraction of fibers and a shape factor dependent on the fiber geometry and orientation. Em, Ef, and Ec represent the elastic moduli of the PLA matrix, the CNF, and a randomly

Table 3. Detailed Mechanical Data of PLA Nanocomposites Obtained from Both Tensile and Impact Measurements run PLA PLA/5CNF PLA/5PN-FR@CNF PLA/10PN-FR@CNF PLA/15PN-FR@CNF a

σta (MPa) 58.0 58.5 65.4 71.8 70.3

± ± ± ± ±

1.0 0.7 0.6 1.0 0.9

Eta (GPa) 2.50 2.55 2.68 2.79 3.15

± ± ± ± ±

0.05 0.07 0.05 0.08 0.09

εa (%) 8.0 7.1 7.2 7.5 6.3

± ± ± ± ±

0.8 0.5 0.6 0.7 0.6

τa (MJ/m3) 2.73 2.54 2.75 3.11 2.46

± ± ± ± ±

0.09 0.04 0.08 0.07 0.06

σia (kJ/m2) 7.63 8.10 9.80 7.02 7.21

± ± ± ± ±

0.18 0.25 0.22 0.23 0.16

σt, Et, ε, τ, and σi refer to the ultimate tensile strength, Young modulus, strain at break, tensile toughness, and impact strength, respectively. 7901



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oriented composite based on the laminate theory. Figure 8C clearly shows that the theoretical values calculated by the Halpin−Kardos model basically fall between the elastic modulus obtained from tensile tests and the ones obtained from flexural tests in the whole range of loading levels in spite of some deviations. This deviation is reasonable because this model fails to consider the interactions between fibers since it is based on self-consistent theory by assuming a single fiber incased in a cylindrical shell of matrix resin. Other possible reasons include that the high entanglement makes it very difficult to accurately calculate the fiber shape factor (ζ), and in fact, the PN-FR@CNF is not individually dispersed within the polymer matrix. Therefore, the Halpin−Kardos model basically can predict the elastic modulus range of our system, which also indicates that PN-FR@CNF can be dispersed well in and effectively reinforce the PLA. Besides, impact measurement results show that the addition of 5 wt % PN-FR@CNF increases the impact strength (σi) from 7.63 kJ/m2 for the PLA host to 9.80 kJ/m2, about a 28% increase, whereas the same loading level of CNF only leads to a negligible improvement (about 8.10 kJ/m2). However, instead of continuing to increase, σi sharply decreases to ∼7.02 kJ/m2 when the loading level of PN-FR@CNF increases to 10 wt %, as also observed in previous studies on CNF-reinforced epoxy resin nanocomposites.38 This means that because of the strong interface and uniform dispersion, the fiber network at a low loading of less than 5 wt % can facilitate the dissipation of impact energy whereas high loading levels (above 10 wt %) of fibers are likely to negatively affect the energy absorption, thus leading to reduced impact toughness.

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Research Article

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acssuschemeng.7b01430. TGA curve of PN-FR@CNF in air, TEM images of PNFR@CNF, AFM images for CNF and PN-FR@CNF, total heat release of PLA and its composites, XRD and Raman data of char residues, and DMA and LOI values of PLA and its nanocomposites (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. *E-mail: [email protected]. ORCID

Pingan Song: 0000-0003-1082-652X Weiwei Lei: 0000-0003-2698-299X Author Contributions ⊥

J.F. and Y.S. contributed equally to this work and are co-first authors.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Scientific Research Foundation of Zhejiang A&F University (Grant No. 2055210012), the National Natural Science Foundation of China (Grant Nos. 51303162 and 51628302), the Zhejiang Provincial Natural Science Foundation of China (Grant No. Q15C160002), the Program for Key Science and Technology Team of Zhejiang Province (Grant No. 2013TD17), the Commonwealth Project of the Science and Technology Agency of Zhejiang Province of China (Grant No. 2017C37078), and the Industrial Transformation Training Centre (Grant No. IC170100032).

CONCLUSIONS

In this work, we have successfully created a novel core−shell nanofibrous flame-retardant system (PN-FR@CNF) via in situ chemical grafting of an ecofriendly phosphorus−nitrogen-based flame retardant onto the high-strength CNF surface. PN-FR@ CNF can be uniformly dispersed within the PLA matrix because the PN-based polymer shell can serve as an effective interfacial compatibilizer between the strongly polar CNF and weakly polar PLA. Adding 10 wt % PN-FR@CNF can significantly reduce the pHRR of PLA by 31% and simultaneously enables PLA to achieve a V-0 flame-retardant rating (the industrial flame-retardancy requirement) during vertical burning, while the pristine CNF only leads to a slight improvement of the flame retardancy of PLA. Moreover, as compared with the PLA host, incorporating 10 wt % PN-FR@ CNF increases the tensile strength and elastic modulus by 24% (∼71.8 MPa) and 12% (∼2.75 GPa), respectively, due to the homogeneous dispersion and strong interfacial adhesion between CNF and the PLA matrix. Thus, the as-synthesized PN-FR@CNF is a new-generation high-performance flameretardant system combining exceptional flame retardancy and mechanical reinforcement. The resultant flame-retardant and strong PLA nanocomposites are expected to find many wide applications in the electronic and electrical, automobile, and aerospace fields. More importantly, this work offers a unique methodology to create nanostructured multifunctional additives for polymers.

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Fire-Resistant, Strong and Green Polymer Nanocomposites Based on

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