Functionalized graphene oxide/phosphoramide oligomer hybrids flame retardant prepared via in situ polymerization for improving the fire safety of polypropylene

Article abstract of DOI:10.1039/c3ra45945d

A novel strategy based on functionalized graphene oxide (FGO)/phosphoramide oligomer flame retardant was developed to overcome the challenges of the dispersion of graphene sheets in polymer matrix and the ease of the burn-out of graphene under air atmosphere. Graphene oxide (GO) was modified by 4,4-diaminodiphenyl methane (DDM) and then in situ incorporated into phosphoramide oligomer, resulting in a nanocomposite flame retardant (FRs-FGO) containing exfoliated graphene. Subsequently, the flame retardant (FRs-FGO) was incorporated into polypropylene (PP) and simultaneously compatilized with PP-grafted maleic anhydride. TEM results showed that the FGO was dispersed more uniformly in PP than the bare GO because of the strong interfacial interaction and previous exfoliation of FGO in FRs before blending. The thermal properties investigated by thermogravimetric analysis (TGA) indicated that the addition of FRs-FGO into PP resulted in a significant improvement of thermal stability at elevated temperature with higher char yields. Moreover, the crystallization and fire safety properties of PP composites were also improved by the incorporation of FRs-FGO, including increased crystallization temperature (11.4 °C increase), reduced peak heat release rate (66.9% reduction) and decreased total heat release (24.4% decrease), and decreased fire growth rate index (73.0% decrease). The cone results indicated the simple blending of GO with FRs and exhibited less improvement in fire safety properties than FRs-FGO, which resulted from the improved dispersion and thermal stability of FGO sheets. The flame retardant mechanism was because of the shielding effect of FGO and char layers, which could reduce the release of combustible gases and inhibit the mass and heat transfer between the gas phase and condensed phase.

Full text of DOI:10.1039/c3ra45945d

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PAPER
Functionalized graphene oxide/phosphoramide
oligomer hybrids flame retardant prepared via in
situ polymerization for improving the fire safety of
polypropylene†
Cite this: RSC Adv., 2014, 4, 31782
ab
a
ab
a
ab
a
Bin Yu, Xin Wang, Xiaodong Qian, Weiyi Xing, Hongyu Yang, Liyan Ma,
a
ab
a
ab
bc
Ying Lin, Saihua Jiang, Lei Song, Yuan Hu* and Siuming Lo
A novel strategy based on functionalized graphene oxide (FGO)/phosphoramide oligomer flame retardant
was developed to overcome the challenges of the dispersion of graphene sheets in polymer matrix and the
ease of the burn-out of graphene under air atmosphere. Graphene oxide (GO) was modified by
4
,4-diaminodiphenyl methane (DDM) and then in situ incorporated into phosphoramide oligomer,
resulting in a nanocomposite flame retardant (FRs-FGO) containing exfoliated graphene. Subsequently,
the flame retardant (FRs-FGO) was incorporated into polypropylene (PP) and simultaneously
compatilized with PP-grafted maleic anhydride. TEM results showed that the FGO was dispersed more
uniformly in PP than the bare GO because of the strong interfacial interaction and previous exfoliation of
FGO in FRs before blending. The thermal properties investigated by thermogravimetric analysis (TGA)
indicated that the addition of FRs-FGO into PP resulted in a significant improvement of thermal stability
at elevated temperature with higher char yields. Moreover, the crystallization and fire safety properties of
PP composites were also improved by the incorporation of FRs-FGO, including increased crystallization
ꢀ
temperature (11.4 C increase), reduced peak heat release rate (66.9% reduction) and decreased total
heat release (24.4% decrease), and decreased fire growth rate index (73.0% decrease). The cone results
indicated the simple blending of GO with FRs and exhibited less improvement in fire safety properties
than FRs-FGO, which resulted from the improved dispersion and thermal stability of FGO sheets. The
flame retardant mechanism was because of the shielding effect of FGO and char layers, which could
reduce the release of combustible gases and inhibit the mass and heat transfer between the gas phase
and condensed phase.
Received 19th October 2013
Accepted 9th June 2014
DOI: 10.1039/c3ra45945d
www.rsc.org/advances
non-ammable or ame-retardant. Among various halogen-free
FRs, intumescent ame retardants (IFRs) have been proven to
1
. Introduction
In today's society, organic polymers are extensively used in be effective ame retardants for thermoplastic resins such as PP
different elds, including electronics, coatings, cable, etc., and poly-ethylene (PE), and are widely used in the industrial
3
–5
because of their excellent properties such as low density, applications. In general, a typical intumescent ame retar-
1,2
resistance to erosion, ease of processing, etc. However, the dant (IFR) system consists of three main components, namely
6
intrinsic ammability of polymers, which has resulted in huge an acid source, a char-forming agent and a blowing agent.
loss of human lives and property, limits their use in some areas. However, the IFR systems containing an APP and two other
Thus, it is necessary and urgent to make commercial polymers basic components are hardly evenly dispersed in the matrix
according to the addition ratio, thus the ame retardant effi-
ciency of these systems is reduced. Thus, some researchers have
a
State Key Laboratory of Fire Science, University of Science and Technology of China, been exploring novel IFRs with these three basic components
9
6 Jinzhai Road, Hefei, Anhui 230026, P.R. China. E-mail: yuanhu@ustc.edu.cn; Fax:
86-551-3601664; Tel: +86-551-3601664
7–9
on a small molecular compound or oligomer. However, a
+
fairly high loading is usually required to achieve good re
resistance that will deteriorate the physical and mechanical
properties of the polymeric materials. Therefore, to improve the
b
USTC-CityU Joint Advanced Research Centre, Suzhou Key Laboratory of Urban Public
Safety, Suzhou Institute for Advanced Study, University of Science and Technology of
China, 166 Ren'ai Road, Suzhou, Jiangsu 215123, P.R. China
c
Department of Building and Construction, City University of Hong Kong, Tat Chee ame retardant efficiency and reduce the amounts of IFRs used
Avenue Kowloon, Hong Kong
in polymers, the combination of nano technology and the IFR
strategy has been adopted.
†
Electronic supplementary information (ESI) available. See DOI:
0.1039/c3ra45945d
1
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Graphene, a one-atom-thick, two-dimensional (2-D) sheet
RSC Advances
In the present work, a kind of phosphoramide oligomer
2
composed of sp carbon atoms in a hexagonal lattice, has nanocomposite FR containing exfoliated and reduced graphene
attracted wide attention from both scientic and industrial oxide (FRs-FGO) was successfully prepared via in situ polymer-

elds because of its exceptional electronic, thermal and ization. Then, FRs-FGO was incorporated into PP to obtain the
10–13
mechanical properties.
With these fascinating properties, ame retardant PP nanocomposites. To further improve the
graphene has motivated tremendous interest in preparing gra- dispersion and compatibility of both graphene oxide and FR in
phene-based polymer composites. Recently, graphene with a the PP matrix, PP-graed maleic anhydride (PP-g-MA) was used
unique sheet structure, similar to the layered silicate, has been as a compatilizer. Graphene oxide was rstly treated in the
used as
composites.
a ame-retardant additive in polymer nano- presence of 4,4-diaminodiphenyl methane. The structure and
14–16
The huge challenge for the applications of morphology of FGO was characterized by FTIR, XPS, XRD and
graphene/polymer nanocomposites at present is the dispersion TEM. The crystallization, thermal and combustion properties of
of graphene in the polymer matrix due to the intrinsic van der PP nanocomposites were investigated by DSC, TGA, cone tests,
Waals force between graphene sheets, which results in the re- and the ame retardant mechanism was discussed.
17–19
aggregating and restacking of graphene.
Moreover, the bare
graphene is easily burnt out upon exposure to a heat ow under
air atmosphere, which could not act as an effective barrier to
2
. Experimental
prevent the escape of organic volatiles. The chemical modi- 2.1 Raw materials
cation of graphene with ame retardants is a better choice
PP (homopolymer, melt-ow rate ¼ 2.5 g/10 min) was supplied
because of the existing functional groups (hydroxyl and
carboxyl) on the basal plane and edges of graphene, which can
by Yangzi Petrochemical Co., Ltd. (China). PP-graed maleic
anhydride (PP-g-MA, 0.6–1.0 MA%) was purchased from Ningbo
Nengzhiguang New Materials Technology Co., Ltd. 4,4-Dia-
mino-diphenyl methane (DDM) was purchased from Sino-
pharm Chemical Reagent Co. Ltd. (Shanghai, China). Reagent
grade phosphorus oxychloride was provided by Shanghai HuaYi
Group HuaYuan Chemical Industry Co. Ltd. (Shanghai, China).
Pentaerythritol, acetonitrile, trichloromethane and diethyl
ether, N,N-dimethylformamide (DMF, AP), graphite powder,
20–22
provide reactive sites for chemical functionalization.
The

ame-retardant elements on the surface of graphene can not
only facilitate the dispersion of graphene but also generate
some char layers during combustion, as suggested by our
23
previous work. Unfortunately, the previous studies have
demonstrated that the amount of oxygen-containing groups on
the surface of graphene was limited (oxygen atom percentage of
approximately 30% for graphene oxide), and thus only limited
potassium permanganate (KMnO
CP), sulfuric acid (H SO , 98%), sodium nitrate (NaNO
hydrogen peroxide(H , 30% aq.) and hydrazine hydrate (85%,
4
, AP), hydrochloric acid (HCl,
24,25

ame retardants could be graed onto graphene.
To address
2
4
3
, AP),
this problem, a novel strategy of graphene wrapped with large
amounts of ame retardants is emerging, which can broaden
the applications of polymer/graphene nanocomposites for re
retardancy.
2 2
O
AP) were all reagent grade and purchased from Sinopharm
Chemical Reagent Co. Ltd. (Shanghai, China).
Presently, the dispersion of graphene in polymers is carried
out by three methods, namely in situ polymerization, solvent
blending and melt blending. Melt blending lead to poor
2
.2 Preparation of FRs and FRs-FGO nanocomposite ame
retardants
dispersion, whereas in situ polymerization and solvent blending Spirocyclic pentaerythritol bisphosphorate disphosphoryl
30
usually result in good dispersion. The combination of nano- chloride (SPDPC) was prepared in our group and poly-
composite technology and the IFR strategy is generally achieved (diaminodiphenyl methane spirocyclic pentaerythritol
by the melt blending method, but rarely via in situ polymeri- bisphosphonate) (FRs) was prepared according to the reported
31
zation method. As a result, new ideas regarding phosphoramide method, as shown in Scheme 1. Typically, SPDPC (30 g, 0.101
oligomer nanocomposite FR containing exfoliated and reduced mol) and DDM (20.8 g, 0.105 mol) were added in a glass ask and
graphene oxide sheets via in situ polymerization, aiming at dispersed by acetonitrile (150 ml). The mixture was stirred and
improving the dispersion and interfacial interactions in poly- gradually heated under nitrogen protection. The reaction was
ꢀ
mer matrix, is suggested.
completed aer 6 h at 80 C, when no HCl could be detected. The
Polypropylene (PP), one of the most widely used thermo- obtained product was ltered and puried with acetonitrile and
ꢀ
plastic materials, has been used in several elds, such as then dried to a constant weight at 80 C in a vacuum oven. The
furniture, electric casings, electronics, interior decorations, pale yellow solid powder was obtained (yield: 92%).
insulation, because of its low density, easy process property,
Graphite oxide was prepared from graphite powder using a
26–28
32
chemical resistance, etc.
However, PP is highly ammable modied Hummers' method. GO was modied and reduced
and drips during combustion. Moreover, it emits smoke and by DDM and the preparation process was as follows. Briey,
poisonous gases while burning, which limits the range of its 1.0 g of graphite oxide was dispersed in 300 ml DMF by ultra-
applications, especially in elds where high ame retardancy is sonication for 60 min at room temperature. Then, the obtained
4,29
required such as electrical industries. Therefore, the reduc- suspension was introduced into a 500 ml three-necked ask
tion of the re hazards of PP and expansion of its application in equipped with a nitrogen inlet, mechanical stirrer and reux
some limited elds are rather imperative and signicant.
condenser. Subsequently, 3.0 g of DDM were added to the ask
and the reaction was conducted under nitrogen at 90 C for
ꢀ
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Scheme 1 Synthetic route of poly(4,4-diaminodiphenyl methane spirocyclicpentaerythritol bisphosphonate) (FRs).
ꢀ
2
4 h. Aer the completion of this reaction, the slurry mixture a vacuum oven at 80 C overnight. PP-g-MA was maintained at
was ltered through
a 0.2 mm polytetrauoroethylene the same loading of 3 wt% during the preparation of all the
membrane and thoroughly washed with DMF and trichloro- composites. Briey, the PP nanocomposite containing 20 wt%
methane to remove the residual reagent. The ltercake was loading of FRs-FGO was prepared as follows: PP, PP-g-MA and
ꢀ
dried under vacuum at 80 C overnight. The preparation FRs-FGO were melt-mixed in a twin-roller mill for 10 min; the
ꢀ
procedure of FRs-FGO nanocomposite ame retardants is temperature of the mill was maintained at 180 C, and the roller
illustrated in Scheme 2. FGO (1.0 g) was dispersed in anhydrous speed was 100 rpm. Aer mixing, the sample was hot-pressed at
ꢀ
acetonitrile (100 ml) with sonication for 1 h, leading to the about 190 C under 10 MPa for 10 min into sheets with a
formation of a homogeneous suspension. Then, SPDPC (7.05 g) thickness of 3.0 mm for characterization. Similar to the proce-
and DDM (4.88 g) were charged into the abovementioned dure described above, the other two samples (PP-1, PP-2) were
suspension and the reaction mixture was reuxed and stirred at obtained.
ꢀ
8
0 C under a nitrogen atmosphere until no HCl was emitted.
The FRs-FGO nanocomposite was ltered and washed with
acetonitrile for three times. Finally, the lter cake was collected Table 1 Composition of PP and its flame retardant composites
ꢀ
by overnight drying at 80 C in a vacuum oven (yield: 87%).
Sample
PP (g)
PP-g-MA (g)
FRs (g)
GO (g)
FRs-FGO
2
.3 Preparation of FRs/PP composites and FRs-FGO/PP
PP-0
PP-1
PP-2
PP-3
100
77
77
—
3
3
—
20
18
—
—
—
2
—
—
—
20
nanocomposites
The composition of PP and its ame retardant composites are
presented in Table 1. Prior to use, PP and PP-g-MA were dried in
77
3
—
Scheme 2 Schematic illustration of the preparation of PP nanocomposite.
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.4 Characterizations
RSC Advances
2
The melting and crystallization behavior of the samples were
investigated by differential scanning calorimetry (DSC Q2000,
TA Instruments Inc., USA). Isothermal crystallization curves
Fourier transform infrared (FTIR) spectra were obtained using a
Nicolet 6700 (Nicolet Instrument Company, USA) spectropho-
tometer with the wavenumber ranging from 500–4000 cm
Samples were mixed with KBr powders and pressed into tablets
for characterization. The concentration for the preparation of
the pellets was maintained in the range of 0.1–1 wt%.
ꢀ
ꢁ1
were recorded at 135 C in nitrogen, while non-isothermal
.
ꢀ
crystallization traces were collected at a cooling rate of 10 C
ꢁ
1
min in nitrogen. Samples (2–5 mg) were heated from room
temperature to 180 C and then cooled to room temperature at a
linear rate of 10 C min . This heating–cooling cycle was
ꢀ
ꢀ
ꢁ1
Wide-angle X-ray diffraction patterns of the samples were
recorded on an X-ray diffractometer (Rigaku Dmax/rA), using a
Cu Ka tube and an Ni lter (l ¼ 0.1542 nm) at 40 kV and 20 mA.
repeated twice.
X-ray photoelectron spectroscopy (XPS) was conducted on a 3. Results and discussion
VG EscalabMark II spectrometer (Thermo-VG Scientic Ltd.
3
.1 Structural and morphological characterization
UK), using Al Ka excitation radiation (hn ¼ 1253.6 eV). All the
Fig. 1 shows the XRD patterns of GO, FGO and FRs-FGO. The
XPS spectra were calibrated using the C1s line at 284.8 eV. Curve
ꢀ
characteristic diffraction peak of GO is located at 2q ¼ 10.3 ,

tting and background subtraction were achieved using XPS
33
corresponding to the (002) reection of GO. A weak peak at
peak t soware.
ꢀ
2
theta ꢃ42 is attributed to the (100) reection of GO. Upon the
The morphology and structure of GO, FGO, FRs-FGO and PP
nanocomposites were studied by transmission electron
microscopy (TEM, JEM-2100F, Japan Electron Optics Laboratory
Co., Ltd, Japan) at an acceleration voltage of 100 kV. GO and
FGO were dispersed in deionized water under 0.5 h of ultra-
sonication. FRs-FGO was puried by extracting over DMF, and
then dispersed in deionized water. All the obtained dispersions
were dripped onto copper grids for observation. These ultrathin
sections of PP nanocomposites with a thickness of 50–70 nm
were obtained using a CM1900 microtome (Leica, Germany)
with water. Then, the ultrathin sections were placed on copper
grids before observation.
chemical modication of GO with DDM, FGO exhibits a
ꢀ
diffraction peak at a smaller diffraction angle of 8.64 , sug-
gesting a signicant enlargement in interlayer spacing. The
enlarged intra-gallery space of GO is attributed to the interca-
lation of DDM via a ring-opening reaction. In the case of FRs-
FGO, no visible diffraction peak is observed in XRD patterns,
indicating that the FGO sheets are exfoliated or dispersed
disorderly in the FRs via an in situ reaction prior to blending
with the polymer matrix. The previously dispersed FGO sheets
in FRs can simultaneously facilitate uniform dispersion and
improve the interfacial adhesion of FGO in PP.
To further conrm the structure of GO and FGO, X-ray
photoelectron spectroscopy (XPS) was performed to investi-
gate the surface composition of GO before and aer the
modication. The XPS C1s spectra of (a) GO, (b) FGO, (c) FRs-
FGO, (d) XPS survey scans of GO, FGO and FRs-FGO are
presented in Fig. 2. The relative area percentages, full-width at
half-maximum (FWHM) of the different C1s peaks and the
atomic percentages of C, O, N and P for GO, FGO and
FRs-FGO are summarized in Tables S1 and S2,† respectively
(see the ESI). As shown in Fig. 2a, GO has a considerable
Thermogravimetric analysis (TGA) was carried out on a TGA-
Q5000IR (TA Instruments, USA) thermo-analyzer instrument.
About 3.0 mg of the sample was measured from room tempera-
ꢀ
ꢀ
ꢁ1
ture to 700 C at a linear heating rate of 20 C min under both
ꢁ
1
air and nitrogen atmosphere (gas ow rate of 150 ml min ).
The ammability of the samples was investigated by cone
calorimeter (Fire Testing Technology, UK) according to ISO
5
660. Each specimen, of dimensions 100 ꢂ 100 ꢂ 3 mm, was
wrapped in an aluminum foil and exposed horizontally to
ꢁ
2
35 kW m external heat ux.
Thermogravimetric analysis/infrared spectrometry (TG-IR) of
the samples was performed using the TGA Q5000 IR thermog-
ravimetric analyzer that was interfaced to the Nicolet 6700 FTIR
spectrophotometer. About 5.0 mg of the sample was put in an
ꢀ
alumina crucible and heated from 30 to 600 C at a heating rate
ꢀ
ꢁ1
ꢁ1
of 20 C min (nitrogen atmosphere, ow rate of 45 ml min ).
The stainless steel transfer pipe and gas cell were heated at
ꢀ
2
30 C to avoid the condensation of volatile compounds.
Real-time Fourier transform infrared spectra (RT-IR) were
recorded using the Nicolet 6700 FT-IR spectrophotometer
equipped with a ventilated oven equipped with a heating device.
The cured samples were mixed with KBr powders, and the
mixture was pressed into a disc with a concentration in the range
of 0.1–1.0 wt%, and then placed into the oven. The temperature
ꢀ
ꢁ1
of the oven was increased at a heating rate of about 10 C min
.
Raman spectra were carried out on a LabRAM-HR Confocal
Raman Microprobe (JobinYvon Instruments, France) in the
ꢁ
1
range of 500–2000 cm using a 514.5 nm argon ion laser.
Fig. 1 XRD patterns of GO, FGO and FRs-FGO.
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Fig. 2 XPS C1s spectra of (a) GO, (b) FGO, (c) FRs-FGO, (d) XPS survey scans of GO and FGO and (e) digital photos of GO and FGO dispersed in
water and DMF, respectively.
degree of oxidation with a C/O atomic ratio of 63.2%/36.8%, the ESI†). To directly observe the reduction of GO, the color
and the C1s XPS spectrum of GO can be deconvoluted into changes before and aer the reaction are shown in the digital
four peaks, corresponding to the carbon atoms in different photos of GO and FGO dispersed in water and DMF, respec-
functional groups. Characteristic peaks at 284.6 (FWHM 1.45 tively (Fig. 2e). Because of the presence of hydrophilic oxygen
eV), 286.6 (FWHM 1.45 eV) and 288.4 eV (FWHM 1.45 eV) are functional groups on its surface and edge (carboxyl, hydroxyl),
attributed to un-oxidized graphite carbon (C–C), C atoms in GO is easily exfoliated in aqueous media, exhibiting a
hydroxyl and epoxy/ether groups (C–O), and carboxyl groups light yellow color. In the reduction process, the graing of
34
(
O–C]O), respectively. Aer the modication of GO with DDM on the GO surface makes the hydrophilic GO hydro-
DDM, the relative peak areas for the oxygen containing groups phobic and there are color changes from light yellow GO in
(
64.2%, Table S1†) are reduced signicantly, especially the water to black, which is considered to be a sign of GO
36
peak of C–O for the epoxy group and C]O for the carboxyl reduction.
group. The reduced relative peak areas for oxygen-containing
The FTIR spectra of (a) GO, (b) FGO, (c) FRs and (d) FRs-FGO
groups indicate that DDM can not only be graed on the GO are shown in Fig. 3. It is obvious that some typical absorption
surface but can also partially reduce GO, which was also peaks of oxygen functional groups appear in the spectrum of
3
5
ꢁ1
observed in the previous work. Moreover, a new peak at GO:O–H stretching vibration (3379 cm ), C]O stretching
ꢁ
1
2
85.6 eV (C–N) is observed in the C1s spectra of FGO, indi- vibration in carboxylic acid and carbonyl moieties (1724 cm ),
ꢁ
1
cating the successful attachment of DDM to the surface of GO C]C stretching vibration (1623 cm ), O–H deformation
ꢁ1
sheets. In the C1s spectrum of FRs-FGO, three types of carbon vibration (1420 cm ), C–O stretching vibration in carboxylic
with different chemical states are observed, which appear at (1226 cm ) and C–O–C stretching vibration in epoxy groups
around 284.2 (C–C), 285.6 (C–N) and 286.3 eV (C–O). From the (1043 cm ). Aer the reaction with DDM, FGO displays some
XPS survey scans of GO, FGO and FRs-FGO (Fig. 2d), a new new peaks, as shown in Fig. 3b: C–H stretching vibration
peak (N1s) at 400 eV is observed for FGO compared to pure (2920 cm , 2850 cm ) and aromatic vibration (1580 cm
ꢁ
1
ꢁ
1
37
ꢁ
1
ꢁ1
ꢁ1
,
ꢁ
1
GO, which further proves the reaction between GO and DDM. 1510 cm ) in DDM. The appearance of these new peaks origi-
In addition to the N1s peak, two new peaks at 191.2 and 134.1 nating from DDM conrms successful modication. Further-
ꢁ1
eV attributable to the P2s and P2p, respectively, are present in more, the peak at 1043 cm
attributed to the stretching
21
the XPS spectrum of FRs-FGO. Based on the XPS C1s spec- vibration of epoxy groups could not be detected in the spectrum
trum and the survey scans of FRs-FGO, the structure of of FGO because of the nucleophilic substitution reaction
phosphamide nanocomposites can be veried. To further between the primary amine group of DDM with the epoxy groups
investigate the structure of GO, FGO, FRs-FGO, the high- of GO, which further provides evidence for successful func-
resolution O1s spectra of (a) GO, (b) FGO and (c) FRs-FGO tionalization. For the FTIR spectrum of FRs (Fig. 1c), the
ꢁ1
and N1s spectra of (d) FGO and (e) FRs-FGO and P2p spectra absorption peaks at 3030 cm (unsaturated hydrocarbon), 2850
ꢁ
1
ꢁ1
ꢁ1
ꢁ1
ꢁ1
of (f) FRs-FGO are shown in Fig. S1 (the detailed description cm (–CH
2
–), 1620 cm (–NH
2
), 1510 cm (C]C), 1245 cm
ꢁ1
and discussion of assignment of XPS peaks are available in (P]O), 1027 cm (P–O–C) and 1085 cm (P–N) are observed,
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phosphoramide oligomers, the surface and edge of FRs-FGO
sheets are fuzzier than that of FGO. Moreover, a few stacking
graphene sheets are clearly observed. The reason for the
morphological change in FRs-FGO relative to FGO is that the
phosphoramide is successfully graed and wrapped on FGO.
The dispersed states of GO and FGO in PP are shown in Fig. 4c
and d, respectively. GO sheets are dispersed in PP mainly in the
form of multilayered stacked structures, and only minor exfo-
liated structures are observed. The obscure structures are the
exfoliated structures and the vivid ones are the stacked struc-
tures. In comparison, FGO dispersed in PP sample (Fig. 4d) is
much better than those of GO in PP, and the exfoliated struc-
tures of FGO are prominent in PP. The improved dispersion of
FGO can be attributed to two aspects: the surface modication
of GO can improve the dispersion of FGO in common organic
solvents and the interfacial interaction with FRs, resulting in
the exfoliation of FGO in FRs; in situ polymerization is regarded
as one of the most effective methods to prepare polymer
nanocomposites, which is benecial in improving the disper-
sion of FGO in FRs. Therefore, the combination of chemical
modication and in situ polymerization nally produces PP
nanocomposites with homogeneously dispersed FGO.
Fig. 3 FTIR spectra of (a) GO, (b) FGO, (c) FRs and (d) FRs-FGO.
31
which is consistent with the FTIR data previously reported. In
Fig. 3d, the FRs-FGO presents the same absorption peaks as
those of FRs because of the overlapping of the peaks of FGO and
FRs and the weak absorption intensities of FGO.
The crucial factors affecting the reinforcement of the prop-
erties of polymer nanocomposites are dispersion state and the
formation of strong interfacial interaction of nano-llers in the
3.2 Thermal and combustion behaviors
polymer matrix. To observe the morphology and evaluate the Thermogravimetric analysis (TGA) was employed to evaluate the
dispersion state of GO and FGO in PP matrix, the TEM images of thermal stability and degradation behaviors of nanollers and
(
a) GO, (b) FGO, (c) PP-2 and (d) PP-3 are presented in Fig. 4. As polymer composites. The TGA curves of GO, FGO, FRs, FRs-
can be seen from Fig. 4a, the morphology of GO is wrinkled and FGO, PP and its composites under air and nitrogen atmosphere
almost transparent, and the surface of GO sheets is fairly are shown in Fig. 5 and 6, respectively. The initial degradation
smooth. The ultrathin nature of GO sheets makes them nearly temperature of samples is the temperature at which 10 wt%
invisible, whereas the multilayered stacks of the GO make them mass loss occurs (T0.1). The temperature of maximum weight-
relatively clear. In the case of FGO (Fig. 4b), the surface of FGO loss rate (Tmax) is obtained from the peak maxima of DTG
appears to be rough and slightly fuzzy, which is attributed to the curves. The detailed data obtained from the TGA curves are
graing of DDM along with the stacking of GO aer reduction. listed in Table 2.
ꢀ
In Fig. 4c, upon the in situ incorporation of FGO into
Under air atmosphere, GO starts to lose weight below 100 C
36
because of the release of adsorbed water. The main mass loss
Fig. 4 TEM images of (a) GO, (b) FGO, (c) FRs-FGO, (d) PP-2 and (e) PP-3.
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Fig.
5 TGA curves of GO, FGO, FRs, FRs-FGO, PP and their
composites under air atmosphere.
Fig.
6
TGA curves of GO, FGO, FRs, FRs-FGO, PP and their
composites under nitrogen atmosphere.
ꢀ
of GO has two stages in the temperature ranges of 200–300 C
ꢀ
and 550–650 C. The rst stage is attributed to the decompo-
sition of oxygen-containing groups and the other to the oxida-
tion of char residue. Reduced by DDM, FGO is more stable than
Table 2 TGA data of PP and the resulting PP composites
ꢀ
ꢀ
ꢀ
Char (700 C, wt%)
T
0.1 ( C)
T
max ( C)
GO and exhibits higher initial degradation temperature (T0.1
)
and char residues. In the thermal degradation of FRs, the minor Sample
N
2
Air
N
2
Air
N
2
Air
ꢀ
mass loss occurring in the range of 150–250 C is attributed to
PP-0
the release of water or unreacted monomer. With an increase in
PP-1
436
365
370
364
283
327
318
302
464
476
477
477
339
381
412
398
1.53
8.29
8.56
8.62
1.64
4.35
5.34
6.55
temperature, the main decomposition process occurs above 280
C, where some complex chemical reactions occur, including PP-3
PP-2
ꢀ
the scission of phosphate ester bonds and the formation of
31
ꢀ
intumescent char. The char yield of FRs at 700 C is as high as
8.1%. The FRs-FGO exhibits similar degradation behaviors as
3
FRs during blending, resulting in higher thermal stability. The
higher char residues are attributed to the degradation of
phosphorus–nitrogen containing FRs generating more heat-
resistant residues. With respect to the char yields of various
composites, it is clear that the char residues of PP-3 containing
FRs but higher thermal stability with more char residues at 700
ꢀ
C (39.8 wt%). Pure PP undergoes thermal degradation in the
ꢀ
temperature range of around 250–380 C, which corresponds to
the random scission of PP chains. The T0.1, Tmax and char yield
ꢀ
ꢀ
of PP is 280 C, 339 C and 1.64 wt%, respectively. For the PP
composites, the incorporation of ame retardant additives
increases T0.1, Tmax and char yields, indicating the improvement
of thermo-oxidation resistance. For example, at the FRs-FGO
loading of 20 wt%, the T0.1, Tmax and char yield of PP-3 are 302
20 wt% FRs-FGO are higher than those of PP-1 and PP-2. The
results indicate that the presence of graphene sheets increases
the ratio of graphitic carbon and amorphous carbon among
char residues, thus improving the thermal resistance of char
residues; signicantly better dispersion of FGO in PP than GO
leads to a better barrier effect by preventing the escape of
pyrolysis products.
ꢀ
ꢀ
C, 398 C and 6.55 wt%, respectively. The improvement in T
0
.1
is attributed to the release of water or unreacted monomer for
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As shown in Fig. 6, nanollers (GO, FGO) and ame Table 3 Cone results of PP and its flame retardant composites
retardant additives in nitrogen show similar degradation
behaviors because they are heated in air atmosphere.
However, nanollers leave more char residues than those in
TTI
(s)
Peak HRR
THR
(MJ m
TSP
(m )
FIGRA
(kW (m s ))
ꢁ2
ꢁ2
2
2 ꢁ1
Sample
(kW m
)
)
air atmosphere because of the high thermal stability of char. PP-0
The virgin PP exhibited an initial decomposition temperature PP-1
(T0.1) of 436 C, which is higher than PP in air atmosphere.
54
69
64
66
1199
620
473
397
97.8
78.5
79.0
73.9
7.28
17.05
19.90
15.90
22.2
9.0
7.4
6.0
ꢀ
PP-2
PP-3
The reason for the lower thermal stability of PP under air
atmosphere is that the oxygen promotes the scission of the
side chain of PP at low temperature. With the incorporation of
additives, the T0.1 of PP composites shis toward lower
temperature, which is attributed to the degradation of
unstable P–O–C bonds in FRs at low temperature. However,
the char yield and Tmax both increase notably, suggesting
improved thermal stability at high temperature. High char
residues formed during combustion will decrease the release
of combustible gas, inhibit heat diffusion and protect the
polymer against re.
Cone calorimetry was used to evaluate the re performance
of these composites. Fig. 7 presents the heat release rate (HRR)
and total heat release (THR) versus time curves of PP and its
composites. Various parameters obtained from cone tests,
including time-to-ignition (TTI), peak heat release rate (PHRR),
total heat release (THR), total smoke production (TSP) and re
growth rate index (FIGRA) are also summarized in Table 3. It
can be observed from Fig. 7a that the pure PP burns very rapidly
aer ignition and the peak heat rate (PHRR) is as high as 1199
38
ꢁ2
kW m . With the incorporation of 20 wt% FRs, the PHRR of PP-
ꢁ2
1
decreases signicantly to 620 kW m . When the FRs is
partially replaced by 2 wt% GO, the PHRR of PP-2 is further
reduced to 473 kW m . Compared to PP-1 and PP-2, PP-3
ꢁ2
exhibits the lowest PHRR on adding 20 wt% FRs-FGO nano-
composite ame retardant. The PHRR values for PP-1, PP-2 and
PP-3 are reduced by 48.3%, 60.6% and 66.9% relative to the pure
PP, respectively. Similar to the HRR curves, PP-3 exhibits
considerably reduced total heat release (THR) (24.4% reduc-
tion), which is signicantly lower than those of PP-1 and PP-2.
The signicant reduction of PHRR and THR for PP-3 can be
explained by the fact that FRs, forming a cohesive and compact
char layer during combustion, acts as an insulating barrier
between re and PP matrix; FGO sheets act as barriers to
prevent the permeation of oxygen and the escape of volatile
degradation products. The presence of FGO promotes the
formation of graphitic carbon, which reinforces the barrier
effect of char layer. As compared with the simple blending of
FRs with GO, FRs-FGO performs much better on improving the

re safety properties of PP, which can be attributed to more
uniform dispersion and stronger interfacial interaction of FGO
in PP than bare GO.
Cone results (Table 3) show that all the composites con-
taining ame retardant additives present prolonged time-to-
ignition (TTI) and increased total smoke production (TSP)
compared with pure PP. Prolonged TTI suggests that PP
composites were harder to ignite than PP, which is probably
because of the increased thermal stability of PP composites.
The increase in TSP is probably attributed to the incomplete
combustion of PP matrix. The re growth rate index (FIGRA) is
used to evaluate the burning propensity of a material, which is
calculated as follows: FIGRA ¼ PHRR/time to PHRR. The lower
the FIGRA values, the higher the re safety of materials. Virgin
ꢁ
2
ꢁ1
PP exhibits the highest FIGRA value of 22.2 kW m
s . As
expected, adding ame retardant additives noticeably reduces
the FIGRA values of composites and the values follow the
sequence: PP-3 (6.0) < PP-2 (7.4) < PP-1 (9.0) < PP-0 (22.2). The
lowest FIGRA value for PP-3 containing 20 wt% FRs-FGO
among all the composites suggests the lowest burning
propensity.
Because the FRs-FGO nanocomposites exhibit superior
ame retardant efficiency and good dispersion of graphene, it is
Fig. 7 Heat release rate (HRR) and total heat release (THR) versus time
curves of PP and its composites.

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important to highlight the advantages of this process. The combustion is attributed to the shielding effect of FGO and char
comparison of this composite with other PP-based composites layers and the formation of more char residues.
reported in literature in terms of performance and preparation
39–42
method is summarized in Table S3† (see the ESI).
Because of
3.3 Condensed phase analysis
the high heat conductivity of graphene sheets, graphene
incorporation led to decreased TTI and slightly decreased THR. Real time FTIR was employed to investigate the solid pyrolysis
However, the FRs-FGO nanocomposites not only increase TTI, products of PP and its composites at various temperatures. The
but also signicantly decrease THR by 24.4%, demonstrating FTIR spectra of PP-0 and PP-3 at different temperatures are
the remarkably enhanced re safety of PP described in this presented in Fig. 10 and 11, respectively. PP shows character-
work. Moreover, the PHRR is also considerably decreased. In istic absorption peaks at 2960, 2920, 2880, 2840, 1460, 1380,
ꢁ
1
ꢀ
general, the dispersion of graphene in polymer matrix via the in 1305, 1255, 1166, 998 and 840 cm . At 200 C, absorptions at
ꢁ1
situ polymerization and solvent blending method is better than 1305, 998 and 840 cm , associated with the crystallization of
that of melt blending-based methods. Herein, the graphene PP, disappear, suggesting the melting of PP crystal. Mean-
sheets were previously exfoliated in ame retardants via in situ while, the appearance of new peaks at 1720 cm resulting from
polymerization, and subsequently incorporated into PP matrix the C]O absorption indicate that PP was oxidized. When the
using PP-g-MA as a compatilizer by melt blending. As expected, temperature increases up to 300 C, the relative intensities of all
45
ꢁ
1
ꢀ
PP/FRs-FGO exhibits much better dispersion than PP-2 by the absorption peaks decrease dramatically, indicating the
directly melt blending GO in PP.
occurrence of main decomposition. It is worth noting that no
ꢀ
To investigate the effect of graphene-based nanocomposites obvious absorption is observed at 450 C, implying that PP

ame retardant on the evolved gas products during pyrolysis, decomposes completely. It is clear from Fig. 11 that some new
ꢁ1
ꢁ1
TG-FTIR was used to analyze the pyrolysis products. The 3D TG- absorption peaks at 1602 cm (–NH ), 1510 cm (C]C), 1257
FTIR spectra of gas phase in the thermal degradation of (a) PP- cm (P]O), 1025 cm (P–O–C) and 1085 cm (P–N) appear
0
and PP-3 at the maximum decomposition rate are shown in temperature is over 280 C, the intensity of the peak at 1025
Fig. 8. As observed from Fig. 8a, the characteristic peaks in the cm is sharply decreased because of the decomposition of
regions of around 2800–3100 cm , 1600–1500 cm and 1550– unstable P–O–C bonds, which is consistent with the TGA results
2
ꢁ
1
ꢁ1
ꢁ1
31
and (b) PP-3 and (c) FTIR spectra of pyrolysis products for PP in the spectra of PP-3 as compared to PP-0. When the pyrolysis
ꢀ
ꢁ
1
ꢁ1
ꢁ1
ꢁ1
1300 cm are noted. According to a previous report, the vola- obtained in air. Previous work has reported that phosphate
tilized products are unambiguously identied by characteristic groups will degrade at low temperature to form some phos-
ꢁ
1
FTIR signals such as C–H compounds (2800–3100 cm ). For phoric acid derivatives, which will catalyze polymer degradation
46
ꢀ
the spectrum of PP-3 (Fig. 8b), the absorptions have no to produce more char residues. Above 330 C, the rapid
noticeable difference from those of PP-0, which are probably decrease in all the peak intensities indicates that the main
ꢀ
because of the overlapping of the absorptions of the volatile decomposition occurs. At 450 C, different from control PP,
components from FRs. For example, the peaks at around 970 some specic absorption is observed in the condensed phase of
ꢁ
1
ꢁ1
ꢁ1
cm and 885 cm are probably attributed to N–H deformation PP-3. The peak at 1290 cm assigned to the P]O vibration in
43,44
vibration in NH and P–O stretching vibration, respectively.
P–O–benzene structure indicates the formation of P–O–C
3
Similar peaks in the FTIR spectra of the samples indicate a complexes with higher thermal stability. Moreover, the
ꢁ1
similar composition of the volatile products. In Fig. 8c, the FTIR appearance of the peaks at 1085 and 880 cm resulting from
38
spectra of the pyrolysis products at maximum decomposition the P–O–P groups conrms the presence of pyrophosphate.
show that the main products of the thermal decomposition of Hence, the char layer composed of phosphorus–carbon
PP-0 and PP-3 are compounds containing hydrocarbons and complexes forms during combustion, which acts as an effective
ꢁ
1
unsaturated hydrocarbons. The peaks at 2958 and 2920 cm
,
barrier for protecting the matrix from decomposing at high
temperature.
ꢁ
1
ꢁ1
1
645 cm , 1460 and 1380 cm correspond to the CH /CH
2
3
stretching, C]C stretching and CH /CH bending vibrations,
respectively.
Raman spectroscopy is a powerful tool for carbonaceous
materials to characterize their structures. Fig. 12 shows the
2
3
43,44
The absorbance of the pyrolysis products for PP-0 and PP-3 Raman spectra of the char residues of PP-1, PP-2 and PP-3. As
vs. time: (a) total, (b) CH /CH (c) unsaturated hydrocarbons are can be seen from the gure, the Raman spectra of all the
presented in Fig. 9. It can be seen that the intensity of pyrolysis samples exhibit two prominent peaks at about 1360 and 1600
2
3
ꢁ
1
products for PP-0 reaches the maximum at about 21.6 min, cm , which correspond to D and G band, respectively. The
whereas that for PP-3 reaches the maximum at about 22.2 min. former is attributed to a breathing mode of K-point photons of
This phenomenon is attributed to the fact that FRs-FGO can A1g symmetry, whereas the latter is attributed to the rst order
2
47
improve the thermal stability of PP nanocomposites at high scattering of the E2g phonon of sp C atoms. In general, the
temperature, which is consistent with the TGA results. More- ratio of D to G band intensity (I /I ) is employed to estimate the
D
G
over, the absorbance intensity of the pyrolysis products for PP graphitization degree of the char, where ID and I are the inte-
G
composites is signicantly lower than that for PP. Conse- grated intensities of the D and G bands, respectively. A lower
quently, the addition of FRs-FGO can reduce the release of ratio of ID/IG implies the higher graphitization degree and
combustible volatiles, and thus leads to reduced heat release thermal stability of char structure. For PP-1, the intensity ratio
rate. The reduced release of organic volatiles during
I
D
/I
G
of the char is about 3.04. The incorporation of GO has a
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Fig. 8 3D TG-FTIR spectra of gas phase in the thermal degradation of (a) PP-0 and (b) PP-3 and (c) FTIR spectra of pyrolysis products for PP and
PP-3 at the maximum decomposition rate.
2 3
Fig. 9 Absorbance of pyrolysis products for PP-0 and PP-3 vs. time: (a) total, (b) CH /CH , (c) unsaturated hydrocarbons.
slight impact on the values of I
FRs-FGO markedly reduces the I
D
/I
/I
G
, whereas the presence of
values of the char from PP- temperature compared with that of the neat PP (T
T
cp of the PP composites (PP-1, PP-2 and PP-3) shi to higher
ꢀ
D
G
c
¼ 115.7 C,
ꢀ
3
to 1.85, indicating the increase of graphitized carbons in the Tcp ¼ 112.4 C). For example, the T and T values of PP-1 (T ¼
c
cp
c
ꢀ
ꢀ
ꢀ
ꢀ
residual char. The reduction in ID/IG ratio arises from the 123.6 C, T ¼ 120.0 C); PP-2 (T ¼ 125.4 C, T ¼ 121.3 C);
cp
c
cp
ꢀ
ꢀ
presence of graphene sheets. The I
D
/I
G
ratio of PP-2 is much PP-3 (T
c
¼ 123.9 C, Tcp ¼ 119.7 C). The higher T
c
and Tcp values
higher than that of PP-3 probably because GO is prone to of PP-1 relative to those of PP-0 can be explained by the fact that
oxidation under air atmosphere; however, FRs wrapped FGO is FRs act as nucleating agents accelerating the crystallization of
c
more thermally stable because the formation of char from FRs PP. The T and Tcp values of PP-2 are lower than those of PP-3,
during degradation process prevents the FGO from oxidation. indicating that GO particles are more efficient in accelerating
The effect of various additives on the crystallization behavior the crystallization of PP matrix than FGO. GO is an effective
of the PP composites was investigated by DSC. Fig. 13 shows (a) nucleating agent compared with FGO because FGO sheets are
ꢀ
non-isothermal crystallization curves (at a cooling rate of 10 C wrapped by FRs and exfoliated in FRs, which limits their
ꢁ
1
ꢀ
min ) and (b) isothermal crystallization (at 135 C) of PP and nuclear capability. It is clear from the isothermal crystallization
its composites. Crystallization temperature (T ) is the temper- curves (Fig. 9b) that all the additives accelerate the crystalliza-
c
ature at the crossing point of the tangents of the baseline and tion of the matrix: with the incorporation of 20 wt% additives,
the high temperature side of the exotherm. The crystallization the time for complete crystallization is signicantly reduced
peak temperature (Tcp) represents the peak temperature of the and the time follows the sequence: PP-0 > PP-1 > PP-3 > PP-2.
exotherm, which can be obtained from the non-isothermal The results are consistent with those obtained from the non-
crystallization curves (Fig. 13a). As shown in Fig. 13a, the T and isothermal crystallization curves.
c
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Fig. 10 RT-FTIR spectra of sample PP-0 at different temperatures.
Fig. 13 (a) Non-isothermal crystallization curves (at a cooling rate of
ꢀ
ꢁ1
ꢀ
1
0 C min ) and (b) isothermal crystallization (at 135 C) of PP and its
composites.
4. Conclusion
In conclusion, a novel strategy was demonstrated herein to
prepare novel nanocomposite ame retardants (FRs-FGO) con-
taining exfoliated graphene via in situ polymerization. Prior
exfoliation and the formation of strong interfacial adhesion for
FGO in the FRs are the main reasons for the homodispersion of
FGO in the PP matrix, as suggested by the TEM images. The
FRs-FGO was incorporated into the PP matrix compatilized with
Fig. 11 RT-FTIR spectra of sample PP-3 at different temperatures.
Fig. 12 Raman spectra of char residue of PP-1, PP-2 and PP-3.
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PP-graed maleic anhydride to improve its re safety. With the 15 S.-H. Liao, P.-L. Liu, M.-C. Hsiao, C.-C. Teng, C.-A. Wang,
addition of 20 wt% FRs-FGO in PP, a signicant improvement in
thermal stability at elevated temperature with higher char yields
M.-D. Ger and C.-L. Chiang, Ind. Eng. Chem. Res., 2012, 51,
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was achieved under both air and nitrogen atmosphere. The 16 X. Qian, L. Song, B. Yu, B. Wang, B. Yuan, Y. Shi, Y. Hu and
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ꢀ
ꢀ
signicantly reduced, including reduced PHRR, THR and 18 J. Zhu, Nat. Nanotechnol., 2008, 3, 528–529.
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results indicated that FRs-FGO could act as efficient nucleating
agents, which can accelerate the crystallization of PP, thus
improving the crystallization temperature and decreasing the
time for complete crystallization. The approach presented here
M. Herrera-Alonso, R. D. Piner, D. H. Adamson,
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2
2
2 H. J. Salavagione, M. n. A. G ´o mez and G. Mart ´ı nez,
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Acknowledgements
This work was nancially supported by National Basic Research
Program of China (973 Program) (2012CB719701), the joint 24 P. Khanra, T. Kuila, N. H. Kim, S. H. Bae, D.-s. Yu and
fund of NSFC and CAAC (no. 61079015), National Natural J. H. Lee, Chem. Eng. J., 2012, 183, 526–533.
Science Foundation of China (no. 51303167) and National 25 Z. Lei, L. Lu and X. S. Zhao, Energy Environ. Sci., 2012, 5,
Natural Science Foundation of China (no. 51203146).
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2
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