Ultralow-Carbon Nanotube-Toughened Epoxy - ACS Publications

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Ultra-low carbon nanotube toughened epoxy: The critical role of a double-layer interface Jingwei Liu, Chao Chen, Yuezhan Feng, Yonggui Liao, Yun-Sheng Ye, Xiaolin Xie, and Yiu-Wing Mai ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b14767 • Publication Date (Web): 13 Dec 2017 Downloaded from http://pubs.acs.org on December 13, 2017

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ACS Applied Materials & Interfaces

Ultra-low carbon nanotube toughened epoxy: The critical role of a double-layer interface Jingwei Liu,† Chao Chen,§ Yuezhan Feng,† Yonggui Liao,† Yunsheng Ye,*† Xiaolin Xie,*† Yiu-Wing Mai‡ †

State Key Laboratory of Material Processing and Die&Mould Technology, School of Chemistry and

Chemical Engineering, Huazhong University of Science and Technology, Wuhan 430074, China. c

Ministry-of-Education Key Laboratory for the Green Preparation and Application of Functional Mate-

rials, Faculty of Materials Science and Engineering, Hubei University, Wuhan 430062, China ‡

Centre for Advanced Materials Technology (CAMT), School of Aerospace, Mechanical and Mecha-

tronic Engineering J07, The University of Sydney, Sydney, NSW 2006, Australia. ABSTRACT: Understanding the chemistry and structure of interfaces within epoxy resins is important for studying the mechanical properties of nanofiller-filled nanocomposites as well as for developing high performance polymer nanocomposites. Despite intensive efforts to construct nanofillers/matrix interfaces, few studies have demonstrated both an enhanced stress transferring efficiency while avoiding unfavorable deformation, due to undesirable interface fractures. Here, we report an optimized method to prepare epoxy-based nanocomposites whose interfaces are chemically modulated by poly(glycidyl methacrylate)-block-poly(hexyl methacrylate) (PGMA-b-PHMA) functionalized multi-walled carbon nanotubes (bc@fMWNT) and also offer a fundamental explanation of crack growth behavior and the toughening mechanism of the resulting nanocomposites. The presence of block copolymers on the surface of MWNT results in a promising double-layered interface, in which: (1) the outer-layered PGMA

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segment provides good dispersion and strong interface bonding with the epoxy matrix, which enhances load transfer efficiency and debonding stress, and (2) the inter-layered rubbery PHMA segment around MWNT provides the maximum removable space for nanotubes, as well as triggering cavitation, while promoting local plastic matrix deformation e.g. shear banding to dissipate fracture energy. An outstanding toughening effect is achieved with only a 0.05 wt% carbon nanotube (CNT) loading with bc@fMWNT, i.e. needing only a 20-times lower loading to obtain improvements in fracture toughness comparable to epoxy-based nanocomposite. The enhancement of their corresponding ultimate mode I fracture toughnesses and fracture energies are 4-times higher than pristine MWNT-filled epoxy. These results demonstrate that a MWNT/epoxy interface could be optimized by changing the component structure of grafted modifiers, thereby facilitating the transfer of both mechanical load and energy dissipation across the nanofiller/matrix interface. This work provides a new route for the rational design and development of polymer nanocomposites with exceptional mechanical performance. KEYWORDS: epoxy nanocomposite, carbon nanotubes, double layer interface, block copolymer, toughness INTRODUCTION Due to their outstanding performance, processability and low cost, epoxy resins have become a key material for many engineering applications, ranging from adhesives to the highly complex structural parts of aeroplanes.1,2 Nonetheless, their high crosslinking density contributes to lower fracture toughness, there still is a huge contrast between the promise and reality of using epoxy resins in diverse, practical applications. Rubber and thermoplastic soft particles are two traditional “tougheners” used to enhance the epoxies by forming a ductile inter-phase between the soft particle and epoxy matrix; however, a high particle loading is required to achieve maximum toughness, resulting in significantly reduced modulus and glass transition temperatures (Tg), due to its intrinsic low Tg.3 Different from soft particles; the use of rigid nanofillers, e.g. silica nanoparticles, nanoclays and carbon-based nanomaterials with


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high stiffness and high interface areas, in the matrix has led to improvements in fracture toughness as well as the modulus and Tg for epoxy-based nanocomposites.3,4 For example, an obvious increase in fracture toughness (KIC) (74-140 %) can be obtained by possible proposed toughening mechanisms such as crack pinning, shear yielding and void growth, in silica nanoparticle-filled epoxies,5,6 while the incorporation of nanoclay into epoxies improves the KIC by 15-125 % - due to the dominance of either microcracks or plastic deformation leading to energy dissipation during crack propagation.7,8 However, the enhancement in KIC obtained with these inorganic nanofillers is mostly less than that seen with soft particles and a high nanofiller loading (2.5-20 wt %) is required, which also results in a high viscosity and a reduction of processibility. Carbon nanotubes (CNT) and graphene nanosheets are two of the toughest cabon-based nanomaterials commonly introduced into epoxies to improve fracture toughness at relatively small loadings (98%, diameters: 10-15 nm, length: 0.5-2 µm. Glycidyl methacrylate (GMA, 98%) and hexyl methacrylate (HMA 198.5 %) both from Aldrich were distilled before use. 3-azidopropyl 4-cyano-4-((phenylcarbonothioyl)thio) butanoate (AzidoCTAs) was prepared according to a similar method that we previously reported.22 2,2’-azobis-(2methylpropionitrile) (AIBN, 98 % from Aldrich) was purified by recrystallization in ethanol and dried at room temperature under vacuum before use. N, N’-dimethylformamide (DMF, 94 %), thionyl chloride (SOCl2), and tetrahydrofuran (THF, 199 %) supplied by Sinopharm Chemical Reagent Co., Ltd were dried and distilled prior to use. All other chemical regents were used as-received. The 0.22 µm polycarbonate membrane and 0.22 µm Teflon membrane were purchased from Shanghai Xingya Purification Material Factory, China. In this work, typical diglycidyl ether of bisphenol-A based E-51 epoxy with an epoxide equivalent weight of 192 (Shanghai Resin Factory Co., Ltd., China) was used as the polymer matrix. The hardener was consisted of the methylhexahydrophthalic anhydride (MHHPA) and benzyldimethylamine catalyst (BDA) with a weight ratio of 40: 1. The resin to hardner ratio equaled 1: 0.82 for all nanocomposites. Synthesis of Alkyne@fMWNT. A mixture of MWNT (1.2 g), HNO3 (30 mL, 60%) and H2SO4 (90 mL, 98%) was refluxed at 80 °C for 3 h. After cooling to room temperature, the mixture was diluted with distilled water and vacuum-filtered through a 0.22 µm polycarbonate membrane and washed with distilled water until the aqueous layer reached neutral. The filtered solid was dried under vacuum overnight at 60 °C to give carboxylic acid-functionalized MWNT (COOH@fMWNT). COOH@fMWNT (1.0 g) was reacted with a large excess of SOCl2 in anhydrous THF under reflux for 24 h. After remov-


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ing the liquid under reduced pressure, the remaining solid was washed with anhydrous THF to remove the residual thionyl chloride. Finally, Acyl chloride functionalized MWNT (COCl@fMWNT) was mixed with 100 mL anhydrous DMF. After sonicating at 40 °C for 30 min, 100 mL propargyl alcohol was added dropwise to the mixture, which was then heated and maintained at 80 °C overnight. After cooling to room temperature, the suspension was vacuum-filtered through a 0.22 µm Teflon membrane and washed with DMF until the filtrate was clear. The resulting black powder (Alkyne@fMWNT) was dried at 60 °C for 24h. RAFT Polymerization of HMA or GMA with Azido Chain Transfer Agent (CTA-N3). Polymerizations were conducted in dioxane using AIBN as the primary radical source and CTA-N3 as the functional chain transfer agent. In a typical polymerization, HMA (2.55 g, 15.0 mmol), CTA-N3 (0.108 g, 0.3 mmol), and AIBN (9.9 mg, 0.06 mmol) were weighed in a 25 mL Schlenk flask. 2.2 mL of 1,4-dioxane was added. The mixture was degassed by three freeze-pump-thaw cycles and then heated at 70 °C for 24 h. The polymerization reaction was quenched by removal the reaction from heat and exposure to air. Then the reaction mixture was diluted with THF and precipitated into a large volume of ice ethanol, then precipitate was isolated by filtration, dried under vacuum for 24 h to give the PHMA as a red solid. (see Scheme S1 in Supporting information) The molecular characteristics of PHMA were determined by gel permeation chromatography (GPC). PGMA was prepared using a similar process-see Table 1. RAFT Polymerization of GMA using Azide-terminated PHMA as Macro-CTA. Macro-CTA (3.0 g, 0.67 mmol) in 15.3 mL of dioxane, monomer GMA (3.8g, 26.7 mmol), and AIBN (22 mg, 0.13 mmol) were charged into a Schlenk flask. After degassing by three freeze-pump-thaw cycles, the flask was immersed into an oil bath preheated to 70 oC for 8 h. The reaction was quenched by placing the flask into an ice- water bath and exposure to air. The mixture was diluted in 15 mL of THF followed by precipitating in ice ethanol (twice), and then dried under vacuum for 24 h to give the block copolymer PGMA-b-PHMA -N3 as a pink solid.



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Click Reaction of Azide-terminated polymer onto Alkyne Bering MWNT (Alkyne@fMWNT) Using the “Grafting to” Approach. Typically, Alkyne@fMWNT (200 mg) and DMF (60 mL) were added into a 150 mL flask, and sonicated at 40 oC for 30 min. Then, azide-functionalized PGMA-bPHMA [PGMA-b-PHMA-N3; Mn

(GPC)

= 10500 g mol-1] (1.7 g, 0.019 mmol), CuI (14 mg, 0.2 mmol),

and DBU (1.4 g, 0.2 mmol) were added. The solution was deoxygenated by nitrogen bubbling for at least 30 min. The flask was immersed into an oil bath preheated to 70 oC for 24 h. The solid was separated from the mixture by centrifugation. The collected solid was redispersed in THF (50 mL) and separated by centrifugation three times. The solid mass (PGMA-b-PHMA@fMWNT) was dried at 60 oC overnight in a vacuum oven. The specific conditions utilized for praration of PHMA@fMWNT and PGMA@fMWNT and the corresponding results are detailed in Table 2. Fabrication of Epoxy Nanocomposites. A desired amount of polymer@fMWNT nanofillers was first weighed and dispersed in THF by ultra-sonication for 2 h at 40 oC. The neat epoxy resin was added into the polymer@fMWNT dispersion and mixed by stirring and sonicating for 60 min at 50 oC. After completely removal of solvent with rotary evaporator and vacuum pump at 70 oC, the epoxy resin and curing agent were mechanically mixed and degassed at 60 oC for 30 min. Then the mixture was cast in preheated steel molds (60 oC) and cured at 80 oC for 120 min, 120 oC for 60 min, 140 oC for 60 min, and finally 160 oC for 120 min. Characterization. 1H NMR spectra were recorded using tetramethylsilane (TMS) as an internal standard in DMSO-d6 or CD3COCD3 on a Burker AV400 NMR spectrometer. The number-average molecular weight (Mn) and Mw/Mn of the polymers were measured by GPC equipped with an Agilent 1100 GPC using a PL gel 79911GP-104 (7.5300 mm, 10 mm beads' size) column. GPC measurements were carried out at room temperature using THF as the eluent at a flow rate of 1 mL min-1. Linear polystyrene standards were used for calibration. Fourier transform infrared (FTIR) spectra were recorded using a Nicolet Avatar 320 FTIR spectrometer. Thermo gravimetric analysis (TGA) of the fillers was conducted using a Perking Elmer Instruments TGA4000 from 35 to 800 oC at 20 oC min-1 in flowing N2. Raman


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spectra were investigated using a SENTERRA Micro Raman Spectrometer (Bruker Instruments, Germany) with an excitation wavelength of 532 nm. High-resolution X-ray photoelectron spectra (XPS) measurements were made [ESCA 2000 (VG Microtech)] using a monochromatized Al Ka anode. Viscosity measurements were taken in parallel plates with a diameter of 25 mm (CP50 and space 0.5 mm) under oscillatory shear on a MCR302 (Anton Paar, Austria) rheometer equipped with a temperaturecontrolled oven (Anton Paar P-PTD200). Viscosity data were collected as the shear rate was increased in logarithmic increments from 1 to1000 s−1. Transmission electron microscopy (TEM Tecnai G220 electron microscope (FEI Co., Netherlands)) was used to characterize the morphological structures of polymer@fMWNT and epoxy/polymer@fMWNT nanocomposites. For the epoxy/p-MWNT and epoxy/polymer@fMWNT nanocomposites, thin sections with a thickness about 90 nm were prepared by ultra-microtomy and gathered on a copper grid coated with carbon film. Glass transition temperatures (Tg) were measured using a TA Instrument Q2000 differential scanning calorimeter (DSC) from 20 to 200 oC at a heating rate of 10 oC min-1 in a nitrogen atmosphere. According to GB/T 1040-2006 (Chinese Standard), dumbbell-shaped specimens (75.0 mm long h 5.0 mm wide h 2.0 mm thick) were made for tensile tests. The tensile test for epoxy/p-MWNT and epoxy/polymer@fMWNT nanocomposites was conducted on CMT4104 (SANS, China) at a tensile rate of 1 mm min-1. The Young’s modulus was calculated from the slope between the stress and the strain. At least five samples of each blend were tested, from which the mean values and error bars were calculated. Fracture toughness was determined on single-edge-notch bending (SENB) specimens according to ASTM D5045-99. A precrack with an average length of 5 ± 0.5 mm was introduced in the SENB specimens (70.0 mm long h 10.0 mm wide h 4.0 mm thick) by tapping a fresh razor blade into the notch. Three-point bending test was conducted on the SENB specimens at a cross-head speed of 10 mm/min and a span of 40 mm. The morphologies of the fracture surfaces of samples were sputtered with platinum and observed through scanning electronic microscopy (SEM) (Sirion 200, FEI Company, Netherland) with an acceleration voltage of 10 kV. The



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mode-I critical-stress intensity factor (KIC) and critical strain energy release rate (GIC) were calculated using the following equations.  P max  K IC =  f ( x) 1/ 2   BW  f ( x ) = 6 x1/ 2

GIC

(1)

1.99 − x (1 − x)(2.15 − 3.93 x + 2.7 x 2 ) (1 + 2 x )(1 − x )3/ 2

(2)

(1 − ν 2 ) K IC 2 = E

(3)

Where x = a/W, and a is the length of initial crack. Pmax is the maximum load at specimen’s fracture. E is young’s modulus determined from the tensile test. For the given specimen geometry, Poisson’s ratio (v) of neat epoxy, (i.e., 0.3523), is used.

RESULTS AND DISCUSSION Synthesis of Azide-terminated PHMA-b-PGMA (PGMA-b-PHMA-N3) and PGMA-b-PHMA Functionalized MWNT (PGMA-b-PHMA@fMWNT). In this study, we used RAFT polymerization to synthesize well-defined PGMA-b-PHMA-N3 because the process is known to be tolerant to a wide range of monomers, functional groups, and reaction conditions. Additionally, it enables the preparation of high purity functional block copolymers with a targeted molar mass and a narrow molar mass distribution without the need to perform complicated post-polymerization transformations.24,25The polymerizations of HMA was synthesized according to our previously published report,

the

synthesized homopolymer of azide-terminated PHMA (PHMA-N3) with a well-controlled molecular weight (Mn = 4480 g mol-1) and a low polydispersity (PDI = 1.09) (Table 1, code 1) at a desired [M]/[chain transfer agent (CTA)] ratio can be used as macro-CTA for further chain extension reactions. To ensure MWNT with good dispersion and interphase in the epoxy nanocomposites, chain extension with controlled chain length from macro-CTA of PHMA-N3 to GMA monomer was attempted to synthesize a block copolymer with adequate chain length for PGMA segment (Mn > 5000 g mol-1)26. GPC results showed that the obtained PGMA-b-PHMA has a Mn of 10700 g mol-1 and low ACS Paragon Plus Environment

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polydispersity of 1.12 indicating the PGMA segment in block copolymer is well-controlled to the target chain length (Mn = 6220 g mol-1). To understand the influence of individual segments from block copolymer on the mechanical properties by epoxy/polymer-functionalized MWNT (polymer@fMWNT) nanocomposites, the individual azide-terminated PHMA or PGMA homopolymers were also prepared and grafted onto MWNT surface for comparison. However, their chain length must be well controlled as this is crucial for MWNT dispersion and interface properties. Thus, the homopolymers in this study were made to have a similar chain length compared to the block copolymer (Table 1, code 2 and 3). 1H NMR was used to confirm the formation of the homopolymer and block copolymer (shown in Figure S2), the results showing characteristic proton resonances. More detailed information is given in the supporting information.

Table 1. Preparation of various polymer-N3 by RAFT polymerization mediated by CTAs-N3 Code

Sample

Mono-

[M]: [CTA]: [I]

Con(%)

Mn, tha)

mer/[M]

Mn,GPCb)

MW/Mnb

(g/mol)

)

1

PHMA -N3(24)

HMA/3

25: 1: 0.2

94

3995

4480

1.09

2

PHMA-N3 (63)

HMA/3

70: 1: 0.2

93

11060

11050

1.10

3

PGMA-N3 (57)

GMA/2

80: 1: 0.2

85

9650

8500

1.08

4

PHMA(24)-b-

GMA/1

35: 1: 0.2

88

8850

10700

1.12

PGMA(44)-N3 a)

Mn,th = ([M]0: [CTA]0) × Mmonomer × Con % + MCTA, where MHMA=170.2, MGMA=142.1,

MCTA(for 1,2,3) = 361, MCTA(for 4) = 4480 respectively;

b)

Determined by gel permeation

chromatography (GPC) in THF relative to monodispersed polystyrene standards. Click coupling reactions are ideal “modular methodology” for grafting polymer chains on the surface of nanomaterials, such as polyhedral oligomeric silsesquioxane (POSS), CNTs and graphene nanosheets, because the relatively mild reaction conditions, high yields and selectivity allow the incorporation of various functional groups on the polymer without the risk of side reactions.27,28 To ensure effective grafting reaction for surface modification of MWNTs, we also employed click coupling reactions to a



prepare

MWNT-based

molecular

brush.

Scheme

S1

shows

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an

alkyne-bearing

MWNT

(Alkyne@fMWNT) being employed to modify MWNT with polymers by the grafting to strategy. A small amount of co-catalyst [CuI/DBU] and excess polymer was used in an attempt to drive the click reactions to high conversion, with repeated washing with an aqueous ammonium hydroxide solution, DMF, and THF, ensured that all catalysts and ungrafted polymers were completely removed. The successful polymer functionalization of Alkyne@fMWNT by click chemistry was confirmed by TGA, FTIR, XPS, and TEM measurements.

Characterization of polymer@fMWNT. To gain a quantitative picture of the extent of MWNT functionalization, TGA was performed on the reaction product (Figure 2). It can be observed that all samples, except for pristine MWNT (p-MWNT), have two weight loss stages in the range 200-800 °C. The first weight loss from 250 to 600 °C can be attributed to the thermal degradation of unstable groups or polymers on the surface of MWNT and the second weight loss from 600 to 800 °C arises from the thermal degradation of MWNT. The pure block copolymer can be completely decomposed at a temperature of 600 °C as indicated in Figure 2 (curve F); therefore, the mass loss of polymer@fMWNT samples at 600 °C is used to estimate the amount of polymer that covalently attached to the MWNT surfaces. As listed in Table 1, the grafting density of modified MWNT (Āmg) can be determined from equation 426.


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100

A: p-MWNT B: Alkyne@fMWNT

80

Weight (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


60

C: PHMA@fMWNT E: PGMA-b-PHMA@fMWNT

D: PGMA@fMWNT

40 20

F: PGMA-b-PHMA

0 100 200 300 400 500 600 700 800

Temperature (oC) Figure 2. TGA traces of (A) p-MWNT (B) Alkyne@fMWNT, (C) PHMA@fMWNT, (D) PGMA@fMWNT, (E) PGMA-b-PHMA@fMWNT and (F) PGMA-b-PHMA-N3.

Āmg = 10000 MCWP ⁄MPWC (chains per 104 carbons)

(4)

where MC is the relative molar mass of carbon (MC = 12 g mol-1), MP is the molecular weight of functionalized polymers (MPHMA = 11050 g mol-1; MPGMA = 8500 g mol-1; MPGMA-b-PHMA = 10700), and WP/WC is the weight fractions of the polymers and MWNT backbone (not including unstable groups of MWNT), respectively. WC/WP can be readily obtained from the TGA curves. According to our previous study26,29, the use of the “grafting to” approach and click chemistry strategy allows full control over the various types of grafted polymer chains within limited chain length (Mn = 2415~9600 g mol-1), while permitting a high grafting density (1.6~2.1 chains per 104 carbons) to a single rGO surface. Table 2 shows that Alkyne@fMWNT grafted with PGMA exhibits a high grafting density (12.9 chains per 104 carbons), indicating the method is equally, or more, effective than polymer grafting to rGO surfaces. However, the grafting density is decreased to 2.3 and 7.6 chains per 104 carbons with MWNTs grafted with PHMA or PHMA-b-PGMA, respectively. Although these two types of polymers have a similar chain length to PGMA homopolymer, their solubility is strongly dependent on the nature of their repeated units; for example, DMF is a good solvent for PGMA homopolymer, but not for PHMA ACS Paragon Plus Environment


Page 14 of 37

homopolymer. Thus, the reactivity between the azide groups at the polymer chain end and Alkyne@fMWNT decreased as the PHMA segment increases in grafted polymers, due to the increase of the coil structure caused by poor solubility of the PHMA segment in DMF. To enhance the chain-end azide accessibility, the grafting reactions were also carried out in DMF/THF or THF systems; however, the grafting density of the resulting polymer@fMWNT was not increased as we expected, which may be due to the poorer dispersibility of Alkyne@fMWNT in these systems when compared to DMF. Although differences in solubility for grafted polymers affect the grafting reaction, the polymer brush chain grafting density remains >2.3 chains per 104 carbons, which meets our requirement for polymerfMWNT use in epoxy-based nanocomposites.

Table 2. Grafting density of MWNT functionalized with different polymer-N3 brushes. Code

Name

Polymer-N3 Mn (g

Grafting ratio

Āmg (chains per 104

mol-1)a)

(wt%)b)

carbons)c)

1

PHMA63@fMWNT

11050

18.0

2.3

2

PGMA57@fMWNT

8500

48.0

12.9

3

PGMA44-b-

10600

40.3

7.6

PHMA24@fMWNT a)

Determined by gel permeation chromatography (GPC) in THF relative to monodispersed polystyrene

standards; b) Determined from the weight loss at 600 °C of TGA curves; c) Caculated by Equation 4. The covalent modification of MWNT by polymers can be verified by their FTIR spectra, see Figure S3. For the Alkyne@fMWNT (curve A), the absorption bands of 1725 cm-1 and 1147 cm-1 corresponding to the C=O stretching and C-O-C stretching respectively, which confirm the existance of carboxylic groups on the surface of MWNT. Additionally, the signals at 1074, 1150, 1725, 2940 and 2097 cm-1 present in PHMA24-b-PGMA44 (curve B), assign to C=S stretching,25 C-O stretching, C=O in carboxylic groups, CH2 stretching and azide bonding stretching, respectively. After polymer grafting,


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the disappearance of the peak at 2097 cm-1 and the significant expansion of the peaks at 1150, 1725, 2923 cm-1 suggest that the block copolymer was successfully grafted onto the surface of MWNT. Raman spectroscopic and XPS analyses were employed to provide an indication of the level of structural perturbation of the MWNTs after modification. The obvious two peaks can be seen in all the samples, related to the D (defect-induced mode) and G bands (E2g2 graphite mode) at 1334 and 1581 cm1

, respectively.30 The R value of MWNT (defined as ID/IG, the integrated intensity of the D band divided

by the integrated intensity of the G band) increased from 0.79 to ~1.10 (Figure S4) after the polymer functionalization, indicating that a certain number of structural defects have been introduced to the MWNT by the formation of covalent bonding on the sidewall of the MWNT.31 At the same time, oxygen-related functional groups and polymers can be expected to be introduced on the outer surface of the tubes during surface modification.30 XPS spectra in wide scan mode (Figure 3a) shows that pMWNT has a major carbon signal and a weak peak, due to low concentrations of oxygen-containing impurities.32 After oxidation, the increase in the relative intensities of O 1s signal and the decrease of the C/O atomic ratio value (calculated from the peak intensity of C 1s to O 1s) in the spectrum of COOH@fMWNT are consistent with the successful introduction of COOH groups on the MWNT surface. In contrast, the C/O atomic ratio in the spectrum of Alkyne@fMWNT increases due to the introduced alkyne groups. After polymer grafting, two notable features, a new peak appearing at ~401 eV and a decreasing of the C/O atomic ratio value in polymer@fMWNT (from 9.4 in Alkyne@fMWNT to 9.1, 5.4, 6.5 in PHMA63@fMWNT, PGMA57@fMWNT and PGMA44-b-PHMA24@fMWNT, respectively), can be clearly observed in the XPS spectrum. A new peak may result from the presence of 1,2,3-triazole linking units between MWNT and grafted polymers after the click reaction; a decrease in the C/O atomic ratio value is attributed to the fact that the grafted polymer chains possess a low C/O atomic ratio (2.0~4.5). Deconvolution of C 1s high resolution spectra (Figure 3b-e) for functionalized MWNTs was used to obtain information about the surface chemical components. Figure 3b shows six typical chemically shifted components,26 sp2 C=C (285.0 eV), sp3 C-C (285.5 eV), C-O (286.6 eV),



C=O (287.7 eV), O-C=O (289.3 eV), ±-±* (291.9 eV), observed in Alkyne@fMWNT. After polymer functionalization, all polymer@fMWNT samples show a new peak at 285.9 eV corresponding to C-N and C=N peaks, and significant increases in C-C, C-O, C=O and O-C=O, indicating the presence of grafted polymer molecules in the resulting materials. These XPS results show that MWNT was successfully functionalized by different polymer molecules, which is in a good agreement with previous FTIR and TGA results.

PGMA -b -PHMA @f MWNT

C/O=9.1

Alkyne@ fMWNT

C/O=9.4

COOH@ fMWNT

C/O=7.8

p-MWNT

C/O=28.1

800

(c)

C/O=5.4

PHMA@f MWNT

600 400 200 Binding energy (eV)

285.5 eV 285.0 eV C-C C=C 285.9 eV C-N and C=N 286.5 eV 287.6 eV C-O C=O 291.9 eV 289.5 eV ±-±* O-C=O

(d) Intensity

Counts

PGMA@f MWNT

(b) C/O=6.5

0

285.5 eV C-C

285.0 eV C=C

286.6 eV C-O 287.7 eV C=O 291.9 eV 289.3 eV ±-±* O-C=O

294 292 290 288 286 284 Binding Energy (eV)

285.5 eV 285.9 eV C-C 285.0 eV C-N and C=N C=C 286.5 eV 287.6 eV C-O C=O 289.1 eV O-C=O

(e) Intensity

N1s C1s

Intensity

O1s

(a)

Intensity

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 16 of 37

285.7 eV284.8 eV 286 eV C-C C=C C-N and C=N 286.6 eV 287.5 eV C-O C=O 289 eV O-C=O

291.4 eV ±-±*




Figure 3. (a) XPS results in a wide scan; carbon 1s XPS profile of: (b) Alkyne@fMWNT, (c) PHMA@fMWNT, (d) PGMA@fMWNT and (e) PGMA-b-PHMA@fMWNT. The morphologies of the polymer@fMWNTs were characterized by TEM. Figure 4a shows that the morphology of pristine MWNT shows a tube-like one-dimensional structure with a diameter of ~15-20 nanometers. After functionalization with PGMA44-b-PHMA24, a dense and uniform polymer layer ( 6 nm) was generated on the surface of the MWNTs (Figure 4b). It can be seen that the obtained block copolymer grafted MWNT (bc@fMWNT) maintained its structural fidelity, implying that surface modification did not induce a significant destruction of the MWNT’s sp2 structure, which is match with a little increase in R value (ID/IG ratio), as discussed in the above Raman analysis. ACS Paragon Plus Environment

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Figure 4. TEM images of (a) p-MWNT and (b) bc@fMWNT. Dispersion of Polymer@fMWNT in the Epoxy-based Nanocomposites. The dispersion level of MWNT in epoxy matrices is a key factor in delivering the high performance of composites. Rheological property measurements are useful for examining changes in the dispersion state and interfacial bonding resulting from surface modification. The effects of functionalization on the viscosity of epoxy/p-MWNT or epoxy/polymer@fMWNT suspensions were investigated at 25 oC with different shear rates. To understand differences in surface functionalities in epoxy resins a 2 wt% MWNT loading was chosen as an intermediate quantity that allows qualitative observations of viscosity changes under fluid conditions.33 Generally, an increased viscosity with improved dispersion and interfacial bonding by surface modification of nanofillers is observed – this is due to the increased spatial confinement of the epoxy resin thorough strong resin-nanofiller interactions.34,35As we expected, the resin containing PHMA63@fMWNT has a lower viscosity when compared to p-MWNT at low shear rates, suggesting either a poorly dispersed system or weaker resin-nanotube interactions (Figure 5). It is interesting to note that the viscosity of the suspensions decreased at low shear rates as the grafted polymer from MWNT, containing PGMA segments, reached lower values for the PGMA57@fMWNT. We speculated that the better MWNT dispersion seen by grafting highly compatible PGMA segments promotes the MWNT alignment, which allows the resin to flow more easily, thus reducing internal friction between epoxy molecules.36,31,37 After curing with hardener, the degree of dispersion of p-MWNT and bc@fMWNT in the epoxy matrix ACS Paragon Plus Environment


was examined by TEM (Figure 6). For the incorporated 0.05 wt% MWNT composite, two dispersion levels, i.e. good dispersion of a few tubes, and the agglomerates of most MWNTs larger than a few microns, can be clearly seen in Figure 6a. When a representative bunch is magnified (see Figure 6b), the MWNTs are seen to be inter-tangled. For the nanocomposites containing the bc@fMWNT, Figure 6c shows that MWNT grafted with block copolymer is homogeneously dispersed in the epoxy matrix due to the outer block of PGMA segment, which is in a good agreement with our previous rheological observations. With further additions of MWNTs (to 0.2 wt%), the dispersion of bc@fMWNT became denser in the epoxy matrix, with most being unwrapped making the boundaries between the adjacent MWNTs easy to distinguish. Neat EP EP/p-MWNT EP/PHMA@fMWNT EP/PGMA@fMWNT EP/PGMA-b-PHMA@fMWNT

140

Viscosity(Pa.s)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 18 of 37

120 100 80 60 40 20 0

1

10 100 Shear rates(s-1)

1000

Figure 5. Steady shear viscosity of epoxy resin suspensions with 2 wt% MWNT loading.


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Figure 6. TEM images of: (a-b) epoxy/p-MWNT, (c-e) epoxy/bc@fMWNT 0.05 wt% and (f-h) epoxy/bc@fMWNT 0.2 wt%.

Thermal and Mechanical Properties of Epoxy-based Nanocomposites. Figure 7 shows DSC curves for the epoxy-based nanocomposites with different contents of p-MWNT and bc@fMWNT, respectively. For the epoxy/p-MWNT nanocomposites (Figure 7a), the Tg decreased by 1.5 °C (133.6 °C) after 0.025 wt% MWNT was added into the epoxy. However, upon addition of more p-MWNT, the Tg values increased to,

137.7 °C and

138.5 °C for 0.1 and 0.2 wt % MWNT, respectively. Adding a

small amount of p-MWNT might result in an increase in the interface volume between the nanotubes and the epoxy, leading to enhanced molecular mobility, thereby lowering Tg of the nanocomposites. With further additions of MWNT, an increase in Tg may result from the restriction of molecular motion by MWNT. For the epoxy/bc@fMWNT nanocomposites (Figure 7b), their Tg first increased markedly from 135.1 to 138.0 °C with a 0.025 wt% MWNT loading, then a slow growth took place with further loading (from 138.0 to 139.2 oC). These results show that the epoxy-based nanocomposites incorporating bc@fMWNT have a higher Tg value compared to incorporated p-MWNT nanocomposites with similar filler loadings, which can be explained by the aggregation of p-MWNT in epoxy. However, the grafted PGMA segment from bc@fMWNT containing abundant epoxy groups provides a large number



of covalent bonding sites between the MWNT and the epoxy matrix.38 Therefore, the mobility of the polymer chain can be additionally restricted by a stronger adhesive interface and a better dispersion of bc@fMWNT in the epoxy matrix as compared with epoxy/p-MWNT nanocomposites.38-39

Heat Flow (W/g)

(a) EP/p-MWNT Neat EP O

135.1 C 0.025 wt% O

133.6 C 0.05 wt%

133.8 C O

0.1 wt% O

137.7 C 0.2 wt%

60

O

138.5 C

80

100 120 140 160 180 Temperature ( C) O

(b) EP/bc@fMWNT

Heat Flow (W/g)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 20 of 37

Neat EP O

0.025 wt%

135.1 C O

138.0 C

0.05 wt%

O

138.4 C

0.1 wt%

O

139.1 C

0.2 wt%

139.2 C O

60

80

100 120 140 160 180 Temperature ( C) O

Figure 7. DSC curves of: (a) epoxy/p-MWNT and (b) epoxy/bc@fMWNT nanocomposites. To maximize the MWNT strengthening and toughening effect, we designed a double-layered interface by surface grafting a block copolymer of PGMA44-b-PHMA24. An outer block of PGMA containing abundant epoxy groups would give a good dispersion and a strong covalent interfacial bonding between the nanotubes and the resin. This interfacial bonding could constrain the mobility of the segments and configure low deformation around the nanotubes, thus producing good load transfer at the interface. In


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comparison, an inter block of PHMA segment with a Tg between -10 oC and 0 oC,40 offers the possibility of significantly higher deformation around the MWNT. This is beneficial for enhancing the mobility of MWNTs in the host polymer and for energy dissipation during deformation. The influence of bc@fMWNT on the mechanical properties of the resulting epoxy-based nanocomposites were determined by tensile tests and crack-opening tests on SENB specimens. Figure 8 shows the tensile properties and toughness of epoxy-based nanocomposites with different nanofiller loading contents. Comparison of p-MWNT and bc@fMWNT shows that the later exhibits a dramatic improvement in tensile strength and fracture toughness, as well as an increase in modulus. The tensile strength of epoxy/pMWNT nanocomposites was increased slightly by adding 0.025 wt% MWNT, and then decreased with further loading. The negative effect of p-MWNT on tensile strength is due to the poor dispersion of MWNT and weak interface between the inert surface of p-MWNT and epoxy. An obvious increase in tensile strength was observed for epoxy/bc@fMWNT nanocomposites; especially, its tensile strength was nearly 15% higher than that of neat epoxy to 95 MPa with the addition of only 0.05 wt% MWNT, resulting from the well dispersed nanotubes and the covalently bonded epoxy/MWNT interface. Such improved dispersion and an improved interface for nanocomposites is also reflected in the mechanical modulus of the resulting nanocomposites, thus a more pronounced increase in modulus was observed for epoxy/bc@fMWNT nanocomposite as compared to the incorporated p-MWNT nanocomposites. In this study, we investigated the fracture toughness and fracture energy of the incorporated p-MWNT and bc@fMWNT nanocomposites, see Figure 8c, d. Crack-opening tests on single-edge-notch bending samples were performed to measure the mode I fracture toughness (KIC) strain energy release rate (GIC) of the neat epoxy, epoxy/p-MWNT and epoxy/bc@fMWNT nanocomposites with various nanofiller contents. The KIC of incorporated p-MWNT nanocomposites increased by ~10% for all samples compared to the base value of neat epoxy. In comparison, the presence of bc@fMWNT in nanocomposites results in a remarkable increase in fracture toughness. A sharp increase of KIC was observed with only a 0.05 wt% MWNT loading, followed by a tendency of slight increase when the loading was further



raised. A maximum improvement in fracture toughness was obtained with a loading of 0.05 wt%, at

(a)

EP/p-MWNT EP/bc@fMWNT

Strength (MPa)

100 90 80 70

Young's modulus (GPa)

which point KIC and GIC increased to 45% and 84%, respectively.

1.4

0.9

(c)

(b)


1.3 1.2 1.1

0.00 0.05 0.10 0.15 0.20 Loading (%)

0.00 0.05 0.10 0.15 0.20 Loading (wt%)

(d)



500

0.8

GIC (J m-2)

KIC (MPa m1/2)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 22 of 37

0.7 0.6 0.5

400 300 200

0.00

0.05 0.10 0.15 Loading (wt%)

0.20

0.00 0.05 0.10 0.15 0.20 Loading (wt%)

Figure 8. (a) Tensile strength, (b) Young’s modulus, (c) fracture toughness and (d) fracture energy of epoxy-based nanocomposites. How the double-layered interfaces, constructed by grafted block copolymer, influences the interfacial features and the state of the surrounding matrix network remains unknown; however, the surface functionalization of MWNT with PHMA and PGMA seems to be a good approch. For comparison, two homopolymers, i.e. functionalized MWNT, PHMA63@fMWNT and PGMA57@fMWNT, were prepared and filled into epoxy at a loading of 0.05 wt%, allowing both the higher tensile strength and better fracture toughness of epoxy/bc@fMWNT nanocomposites to be obtained. The tensile properties and fracture toughness of the nanocomposites with different epoxy/polymer@fMWNT nanocomposites are listed in Figure 9. All loadings of 0.05 wt%. Epoxy/PHMA63@fMWNT gave a slightly lower tensile strength and modulus, but slightly higher KIC and GIC values compared to epoxy/p-MWNT. The insignificant effect of PHMA63@fMWNT on tensile properties and fracture toughness is due to the weak epoxy/MWNT interface and low grafting density of PHMA on the surface of MWNT. In contrast, the


Page 23 of 37

filled PGMA57@fMWNT with a strong epoxy/MWNT interface turns out the dramatically enhancement in tensile strength and Young’s modulus compared to the neat epoxy, i.e. 25% (103.1 MPa), and 22% (1.36 GPa), respectively. However, the fracture toughness of epoxy/PGMA57@fMWNT cannot reinforce monotonically by simply improving interfacial adhesion and only displayed a slight increase in fracture toughness compared to the epoxy/p-MWNT. As a result, the double-layered interfaces produced by the PGMA-b-PHMA segments induce special interfacial failure modes. The above mechanical results of various composites show that that epoxy/MWNT composites with high mechanical performance can be achieved by constructing an interphase by the covalent functionalization of double-layered interfaces onto the surface of MWNT.

110

Strength

(a)

1.5

Young's modulus

1.4

90

1.3

80

1.2

70

1.1

Strength (MPa)

100

60

T EP NT WNT -MW A@fMWN WNT @fM EP/p M c@fM H GMA EP/b EP/P EP/P

0.9

KIC

1.0

(b)

GIC

500

0.8 400

GIC (J.m-2)

KIC (MPa.m1/2)

Young's modulus (GPa)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.7

300

0.6 0.5

EP

EP/p

NT NT WN T WN T -MW fMW @fM @fM /bc@ GMA HMA P /P /P E P P E E

200

Figure 9. Mechanical properties of the neat epoxy and epoxy-based nanocomposites filled with different polymer@fMWNT: (a) Tensile properties; (b) fracture toughness.



Page 24 of 37

Toughening Mechanism of MWNT Constructed with Double-layered Interfaces. It is well known that the lack of energy dissipation events during the crack propagation makes highly crosslinked epoxy possess an intrinsic brittleness and leads to catastrophic fracture under stress. Enhanced toughening of epoxy-based nanocomposites has been widely studied by introducing soft or/and rigid nanofillers into the epoxy matrix for either binary or ternary composites.4,41 Incorporation of soft rubber particles in crosslinked epoxy has shown great effectiveness in overcoming the inherent brittleness of these materials, because the plastic void growth during debonding with the matrix leads to massive plastic deformation of the matrix that can dissipate a significant amount of energy, thereby effectively increasing its toughness.42 CNTs are considered as a potential rigid nanofiller able to improve the toughness of crosslinked epoxy as well as its Young’s modulus and tensile strength. For epoxies filled with CNTs, the mechanism involves crack-bridging and the pull-out of CNTs (see schematic in Figure 10), both of which have the ability to dissipate energy during crack propagation, resulting in improved composite toughness.43,44 The bridging can suppress a further crack-tip opening displacement when the crack grows further, thereby restricting crack advance, compared to the situation without crack-bridging. The crack propagation energy can be further dissipated by applying the pull-out process to the CNTs, which is dependent on the interfacial adhesion between the CNTs and the matrix; therefore, a maximum fracture toughness may obtained by constructing optimized CNTs/matrix interface.45


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Figure 10. The schematics for the possible failure modes of p-MWNT and bc@fMWNT in epoxy matrix.

Figure 11. SEM fractographs of fracture surfaces of SENB epoxy-based nanocomposites with 0.2 wt% (a, b) p-MWNT and (c,d) bc@fMWNT. White arrows indicate the pulled-out MWNT or holes.



Page 26 of 37

Figure 12. SEM and TEM fractographs of bc@fMWNT nanocomposites. In this study, we employed SEM fractography to investigate the fracture mechanism of incorporated bc@fMWNT epoxy-based nanocomposites, while epoxy/p-MWNT nanocomposites were used for comparison. It is well known that the neat epoxy displays a featureless and smooth fracture surface after crack propagation, because of the intrinsic brittleness of the epoxy network.46 After the introduction of p-MWNT or bc@fMWNT in epoxy, it is clear that obvious geometric markings appear on the fracture surface of the resulting nanocomposites after the tensile test (Figure S5), moreover, they become rougher by increasing the MWNT content. The epoxy/bc@fMWNT nanocomposites deliver a rougher fracture surface (Figure S5e-h) when compared to epoxy/p-MWNT nanocomposites with the same MWNT loading level. From the high magnification image of the fracture surfaces of the epoxy-based nanocomposites in Figure 11a and 11b, the facile de-bonding between MWNT and epoxy matrix was revealed by most of the MWNT showing pull-out characteristics at the surface with long nanotube segments. By comparison, good homogeneity and dispersion on the fracture surface were shown by the epoxy/bc@fMWNT system (Figure 11c). The high magnification image (Figure 11d) revealed strong interfacial bonding between the MWNT and the epoxy matrix by the relatively short MWNT segments left on the fracture surface when compared to the epoxy/p-MWNT system (Figure 11b) and the relative long MWNT segments left on the fracture surface compared to the epoxy/PGMA@fMWNT system (Figure S6), suggesting that most of nanotube was embedded within the matrix. Moreover, Figure 11d and Figure 12 show that the diameters of bare MWNT and left holes in the epoxy/bc@fMWNT/ system


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(30~45 nm) appear to be much larger than those of MWNT in the epoxy/p-MWNT system (15~20 nm), which are seemingly caused by attachment of matrix to the MWNT when it is pulled-out.47 These results reflect the fact that MWNT in epoxy/bc@fMWNT may experience a longer period of MWNT/matrix interfacial de-bonding and frictional sliding processes before the pull out process of MWNT took place, as well as a reinforcement effect in a larger area away from the direct interface in comparison to the epoxy/p-MWNT.21 Additionally, it is worth noting that obvious microcracks are observed in the high magnification image of epoxy/bc@fMWNT/, see Figure 11, which cannot be found on the fracture surface of epoxy/p-MWNT nanocomposite. Such microcrack formation in the resulting nanocomposite also reflects the existence of impactful interfacial bonding between the MWNT and epoxy, which contributes to the maximum toughening at low MWNT loadings.20 Although crack-bridging and pulling-out as well as microcracks, crack-pinning and crack-deflection are still acting to toughen the epoxy; however, the propagation of the primary crack by the coalescing of microcracks may effectively weaken the toughening effect of bc@fMWNT leading to a limited increase in KIC and GIC with further loading. Based on our previous study, the significant improvement in toughness of epoxy-based nanocomposites with bc@fMWNT can be explained as follows (see Figure 10 for proposed toughening mechanisms). (i) The soft inter-layered PHMA around MWNT not only provides a removable space for nanotubes, but can also absorb elastic energy through self-deformation, while the de-bonding of the rubbery segment PHMA with epoxy was able to trigger cavitation and promote the local plastic deformation of matrix-e.g. shear banding to dissipate fracture energy.42 (ii) The outer-layered PGMA in bc@fMWNT contributes to the good dispersion and strong interfacial bonding leading to a significant improvement in interfacial adhesion thereby maximizing the pull-out energy.48 (iii) Combing both advantages of PHMA and PGMA to construct double-layered interfaces, endows the MWNTs with high de-bonding stress and high interfacial friction as well as increasing the pull-out displacement and bridging length in the nanocomposites, thus the negative influences caused by a strong (low pull-out displacement) or a weak (low pull-out energy) interfacial bond strength can be avoided.45



Page 28 of 37

It is notable that the tensile strength and toughness of incorporated CNTs epoxy-based nanocomposites depends on such variable as: (i) the size and type of CNTs, (ii) CNTs dispersibility during processing, and (iii) interfacial characteristics between the CNTs and the matrix. In order to fully understand the interfacial effects on the resulting mechanical properties, MWNTs with a short length (0.5-2 µm) were selected as starting material for nanocomposite preparation. To put our work in the perspective of published studies of tensile strength and fracture toughness in epoxy/CNT nanocomposites, Figure 13 shows that introducing CNTs improved either tensile strength or fracture toughness to a certain extent (blue dashed circle) at 0.1-1 wt% loadings. Surface treatments of CNTs with organic molecules or polymers resulted in further improved fracture toughness at 0.2-1 wt% (red and green dashed circle) - the toughness enhancements can be obtained with relatively small loadings (0.2 wt%) in the case of epoxy/TGAP@fMWNT system (Table S1, code 7). Among the epoxy/CNT nanocomposites, the best reported enhancement in KIC was 52.2 % made by incorporating 1 wt % ozone@fMWNT in epoxy resin (Table S1, code 8). In contrast, the bc@fMWNT only required a 20times lower loading of CNTs to obtain the improvement in KIC or GIC achieved, indicating our materials have optimized surface functionalization and interfacial design. This promising double-layered interface to polymer nanocomposites can be extended to the preparation of many other distinctive types of polymer-functionalized nanomaterials, paving the way for the preparation of polymer nanocomposites with high mechanical properties.


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Figure 13. Comparison of maximum increase of KIC and increase of tensile strength as a function of CNT content between epoxy/CNTs nanocomposites taken from literature (our results are shown as orange sphere).

CONCLUSIONS Well-defined

azide-terminated

PGMA-b-PHMA

copolymers

and

their

individual

homopolymer were synthesized by RAFT polymerization before grafting with MWNT. Then, various polymer@fMWNTs were prepared by click coupling reactions, followed by mixing with epoxy to prepare epoxy/MWNTs nanocomposites. A variety of characterization methods, including TGA, FTIR, Raman spectra, XPS, and TEM, proved that block copolymers and their individual homopolymers were successfully grafted onto the surface of the MWNTs without significant destruction of the nanotube sp2 structures. The well-dispersed bc@fMWNT and improved interfaces between MWNT and epoxy matrix in the resulting epoxy-based nanocomposites shown by rheology, TEM and SEM, resulted from the high dispersibility and covalent bonds formed between the outer block segments of PGMA from the MWNT with the epoxy matrix. The epoxy/bc@fMWNTs nanocomposites showed higher Tg and stronger tensile strength than epoxy/p-MWNT and neat epoxy. A remarkable toughening effect was obtained for



Page 30 of 37

nanocomposites filled with less than 0.1 wt% of MWNT, i.e. when filled with bc@fMWNT a 45% improvement in KIC and a corresponding 84% improvement in GIC with only 0.05 wt% of MWNT loading was found. The significant improvement in the toughness of bc@fMWNT-filled nanocomposite can be attributed to the construction of a double-layered interface in the nanocomposites, which not only enhanced load transfer efficiency and de-bonding stress but also absorbed elastic energy by self-deformation. Due to the nature of mutual exclusion, strength and fracture toughness of nanocomposites are always sacrificed for the sake of the other in most cases, thus it is challenging to prepare materials with simultaneous reinforcement and toughening effects. In this work, we illustrated that substantial enhancements in both tensile properties and fracture toughness were realized by introducing only a small amount of bc@fMWNT into an epoxy matrix. ASSOCIATED CONTENT

Supporting Information. Detailed chemical structures, schematics of the synthetic strategy of PGMAb-PHMA@fMWNT, GPC traces of polymers, 1H NMR spectra of polymers, FTIR spectra and Raman spectra of MWNTs, SEM fractographs of composites, comparison table of maximum increase of KIC and increase of tensile strength between epoxy/CNTs nanocomposites taken from literature. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]; [email protected]. Tel.: +86 27-87793241. Fax:+86 27-87543632.

Present Addresses †

State Key Laboratory of Material Processing and Die&Mould Technology, School of Chemistry and

Chemical Engineering, Huazhong University of Science and Technology, Wuhan 430074, China.


Page 31 of 37 §

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Ministry-of-Education Key Laboratory for the Green Preparation and Application of Functional Mate-

rials, Faculty of Materials Science and Engineering, Hubei University, Wuhan 430062, China ‡

Centre for Advanced Materials Technology (CAMT), School of Aerospace, Mechanical and Mecha-

tronic Engineering J07, The University of Sydney, Sydney, NSW 2006, Australia.

ACKNOWLEDGMENT The authors acknowledge financial support from National Natural Science Foundation of China (51503071 and 51673076). We also acknowledge access to XPS, SEM, TGA and TEM facilities of the Analytical and Testing Center of Huazhong University of Science and Technology.

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Ultralow-Carbon Nanotube-Toughened Epoxy - ACS Publications

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