high performance epoxy-layered silicate nanocomposites

High Performance Epoxy-Layered Silicate Nanocomposites

X. KORNMAN”, R. THOMANN2, R. MJkHAUPT2, J. FINTER3, Qnd L. A. BERGLUNDl

‘Division ofPoZymer Engineering hleb University of Technology

S-97187 Luletc, Sweden

2Freiburg Materials Research Center and Institute for Macromolecular Chemistry

Albert-Ludwigs- Universitat Stefan-Meier-StraB 21fD-79104 Freibwg im Breisgau, Germany

3Vantic0 AG CH-4002 Basel, Switzerland

High performance epoxy-layered silicate nanocomposites based on tetra-glycidyl 4,4‘-diamino-diphenyl methane (TGDDM) resin cured with 4.4-diaminodiphemyl sulfone (DDS) have been successfully synthesized. Fluorohectorites modified by means of interlayer cation exchange of sodium cations for protonated dihydro-imidazolines and octadecylamine were used. Fluorohectorite exchanged with 1 -methyl- 2-norstea1yl-3-stearinoacid-amidoethyl-dihydro-imidazoli11ium ions was immiscible with the epoxy matrix. In contrast, fluorohectorites exchanged with hydroxyethyl- dihydro-imidazolinium (HEODI) and ricinyl-dihydro-imidazolinium ions (RDI) favored the formation of a nanocomposite structure. This is most likely due to the presence of -OH groups in their molecular structure, which has a catalytx effect on the polymerization occurring between the silicate layers. The diffusion of epoxy and curing agent molecules between the silicate layers is also promoted. Microscopy ob- servations revealed that the dispersion of the silicate aggregates on a microscale was proportional to the degree of separation of the silicate layers on a nanoscale. Decreased apparent glass transition temperature was observed in all the nanocornposites. Finally, mechanical property studies showed that epoxy-layered silicate nanocomposite formation could simultaneously improve fracture toughness and Young’s modulus, without adversely affecting tensile strength.

INTRODUCTION Nanocomposite technology using layered silicates as

igh performance epoxies based on tetra-glycidyl H 4,4’-diamino - dipheny 1 methane (TGDDM) resin and 4,4’-diamino-diphenyl sulfone (DDS) are exten- sively employed as matrices of carbon or aromatic fibers reinforced composites used in the aerospace in- dustry because they fulfill the requirements of high modulus and high temperature performance. How- ever, they are usually brittle as a result of their high crosslink density. Moreover, phenomena such as moisture absorption from high humidity environ- ments (1, 2) as well as residual stresses (3) can substantially lower the performance of the composites. Therefore, it is of interest to modify the epoxy matrix to improve the composite properties and extend the field of applications of polymeric composites.

nanoscale reinforcements offers an interesting alternative for the modification of polym.er matrix properties. Indeed, large improvements o:F mechanical and physical properties including modulus (4). barrier properties (3, flammability resistance (6), ablation performance (7) as well as reduction of residual stresses (8) have been reported for this type of materials at very low layered silicate content. The synthesis of this new kind of materials involves nanoscale dispersion of 1-nm-thick silicate layers of high aspect ratio (typi- cally 200) in a polymer matrix. The polymeric nanocomposites were first developed by researchers from Toyota Research Corp. to synthesize polymer-layered silicate nanocomposites based on polyamide-6 (9) and was extended later to in-situ polymerization of epoxies (10, 11).

POLYMER ENGINEERING AND SCIENCE, SEPTEMBER 2002, Vol. 42, No. 9 1815

X . Kornmann, R. Thomann, R. Mulhaupt, J. Finter, and L. A. Berglund

The typical in-situ polymerization procedure is the following. First, the layered silicate is rendered organophilic by exchanging the inorganic cations placed between the layers for organic surface modifiers. Then, the organosilicate is swollen in the mono- mer before adding the curing agent. If an appropriate balance between the polymerization occurring between the layers (intragallery polymerization) and the one occurring outside (extragallery polymerization) is established, separation of the silicate layers occurs ( 12). This balance is influenced by numerous factors, such as the curing temperature (12), the nature of the curing agent (13), the cation-exchange capacity of the layered silicate (14) as well as the nature the surface modifier (1 5). A critical parameter substantially favor- ing the separation of the silicate layers is the ability for the surface modifier to catalyze the intragallery polymerization. Thus, Ian et af. (16) demonstrated that the acidity of alkylammonium ions can account for catalytic epoxy homopolymerization between the silicate layers and that this has great influence on the extent of separation of the silicate layers. However, this type of catalysis may have a negative effect on the properties of the nanocomposites because the homopolymerization of epoxy can substantially change the structure of the polymer network. Therefore, it may be interesting to find alternative ways of catalyzing the intragallery polymerization.

Epoxy-layered silicate nanocomposites based on diglycidyl ether of bisphenol A (DGEBA] resin have been synthesized using a wide range of curing agents including aliphatic [ 13, 17) and aromatic diamines (12) but also anhydrides (18, 19). However, the synthesis of high performance epoxy-layered silicate nanocomposites based on highly crosslinked epoxies has not been reported.

In this work, nanocomposites based on high performance epoxy resin have been synthesized using fluoro-hectorite modified by means of interlayer cation exchange using cations such as protonated dihydro- imidazolines and octadecylamine. The influence of these surface modifiers on curing reactions and on the glass transition temperature of the epoxy matrix is discussed. The morphology and the mechanical properties of these high performance epoxy-layered silicate nanocomposites are also presented.

EXPERIMENTAL

Materials

The synthetic sodium fluorohectorite Somasif ME- 100, prepared by heating talcum in the presence of Na,SiF, for several hours in an electric furnace at high temperature, exhibits a cation-exchange capacity of 100 meq/ 1 OOg. Somasif ME- 100 was supplied by CO-OP Ltd., Japan.

Different surface modifiers were used in this study: octadecylamine (ODA) from Fluka Chemie GmbH, 1- methyl-2-norstearyl-3-stearinoacid-amidoethyl-dihydro-imidazolinium methosulfate (W75) (Rewoquat

W75) from Witco Surfactants GmbH (Germany), hydroxyethyl-dihydro-imidazolinium chloride (HEODI) (Servamine KO0 330) from Serve Delden (Holland), and ricinyl-dihydro-imidazolinium chloride (RDI) synthesized by the reaction of ricinoleic acid with ethylenedi- m i n e under nitrogen atmosphere at 250°C. Details about the synthesis are reported elsewhere (20). The molecular structures of the different surface modifiers are presented in Fig. 1.

The matrix was N,N,N,N-tetraglycidyl 4.4-diaminodiphenyl methane (TGDDM) resin, MY 9512, and was cured with 4,4'-diaminodiphenyl sulfone (DDS), HT 976. Both chemicals were supplied by Vantico AG (Germany).

Intercalation of Surface Modifiers

As reported previously for silicate modification with cyclic amidinium cations [20), 1.2 moles of the surface modifier were dispersed in deionised water at 80°C. Then, a stoichiometric amount of hydrochloric acid 37% was added to ensure protonation. After 30 rnin, 1 kg of Somasif ME-100 was added to the mixture. A white precipitate formed. I t was isolated and washed with deionised water by centrifugation until no chloride was detected in the centrifugate by one drop of 0.1 N AgNO, solution. For the synthesis of ME-W75, extensive washing was used since methosulfate ions could not be detected by the silver nitrate test. The organosilicate was dried 2 days at 80°C and then ground in a mill to produce powder with an aver- age particle size of - 80 pm.

Preparation of High Performance Epoxy-Layered Silicate Nanocomposites

The TGDDM resin was heated to 100°C under vacuum with the high shear mixer Labmax from Molteni AG. Then, a quantity of organosilicate corresponding to 5, 10 or 15 wt% of the high performance epoxy system (TGDDM + DDS) was added to the TGDDM resin and the mixture was stirred under vacuum for 3.5 h at 100°C to swell the organosilicate in the TGDDM resin. Then, the temperature of the mixture was raised to 140°C and a stoichiometric amount of the DDS curing agent was added. This mixture was stirred under vacuum for 30 min and then poured into an aluminum mold treated with release agent and previously preheated either to 80°C or 140°C. depending on the curing cycle used. It was then cured either 2h at 80°C + l h at 100°C + 4h at 150°C + 7h at 200°C (curing cycle No 1) or 2h at 140°C + 2h at 177°C + 7h at 200°C (curing cycle No 2). All the samples prepared for mechanical analyses were cured following curing cycle No 2.

Thermogravimetric Analysis Thermogravimetric analyses were performed with a

STA 409 from Netzsch. Thirty milligrams of sample were placed in a crucible and heated at lO"C/min to- gether with a reference under nitrogen atmosphere

1816 POLYMER ENGINEERING AND SCIENCE, SEPTEMBER 2002, Vol. 42, No. 9


ODA

w75

HEODI

RDI

NH;: CI-

, CH,OSO,- R = tallow alkyl 0

n

L N L / ' CI-

Flg. 1 . Surface modi@rs used in the synthesis of the nanocornposires: octadecyl-ammmonium chloride (ODA), 1 -rnethyl-2-norstearyl- 3-stearinoacid-athyl-dihydro-imidazolin~rn methosulfate W75), hydrawyethyl-dihydro-imidazoliniurn cl-lloride (HEODIJ, and ricinyl-dihydro-imidazolinium chloride lRDI). A tallow alkylfunction has the foUowing composition: 1% C,,, 4% C,,, 31% CJ6. and 64% Cl8'

(flow rate: 250 mL/min]. All samples were analyzed from 30°C to 700°C. I t was assumed that no organic species remained after the thermal analysis.

Differential Scanning Calorimetry

DSC measurements were performed with a DSC 200 from Netzsch under nitrogen atmosphere from 25°C to 350°C at a heating rate of 10"C/min. The different organosilicates were first swollen in the TGDDM resin under vacuum for 30 min at 140°C. A sample of the mixture was taken, quenched in liquid nitrogen, and then analyzed by DSC. Then, DDS curing agent was added to the mixture and it was mixed for another 30 min at 140°C. A second sample was taken, quenched in liquid nitrogen and then analyzed by DSC.

Wide Angle X-ray Scattering

Powder WAXS analyses were performed using a com- puter controlled Siemens D500 diffractometer with Cu radiation (50 kV, 40 mA). The scanning speed and the step size were 0.05"/5s and 0.05". The organosilicates were analyzed as produced.

Transmission Electron Microscopy

TEM samples were cut from nanocomposite blocks using an ultramicrotome Reichert & Jung Ultracut E equipped with a diamond knife at ambient conditions. Thin specimens of about 40 nm were cut from a mesa of about 0.5 X 0.5 mm2. They were collected in a through filled with water and placed on 400 mesh copper grids. Transmission electron micrographs were

taken with a LEO 912 apparatus at an acceleration voltage of 120 kV. The spacing between neighboring silicate layers was measured manually using the analysis software ESI-Vision (SIS) using images taken at a magmfication of 250,000. To receive valid results the mean value of about 200 measurements was cal- culated.

Scanning Electron Microscopy SEM samples were cut from nanocomposite and

micro-composite plates and their surfaces were pol- ished. After drying, the samples were coated with a thin layer of carbon and observed in a scanning electron microscope CAMSCAN S4-80CN at an acceleration voltage of 30 kV. The micrographs were obtained using backscattered imaging.

Dynamic Mechanical Thermal Analysis

Dynamic mechanical properties of the epoxy resin and the filled composites were measured with a Dy- namic Mechanid Analyzer RSA 11 from Rheometrics (dual cantilever beam test). Measurements covered temperatures from 30 up to 300°C at an heating rate of 2"C/min, a frequency of 1 Hz and 0.1% strain. The sample size was 4 X 2.5 X 50 mm3. Storage modulus E', loss modulus E" and damping factor tan 8 (= E"/E') were obtained from the experiment.

Tensile and Fracture Testa

Young's modulus, tensile strength and elongation at break were measured by means of tensile tests according to I S 0 527/95 using Zwick 1474 machine.


X . Kornrnunn, R. Thomann, R. Miilhaupt, J. Finter, and L. A. Berglund

The fracture toughness K,, was determined from bend notch tests using a Zwick Flexing Tester 1435 according to ISO/DIS 1358. The fracture energy G,, was computed from G,, = (Krc2/E).( l-u2) since fracture tests were performed under plane strain conditions. E is the Young's modulus and u is the Poisson's ratio of the material (u = 0.4).

The true silicate content, which is the content of silicate in the composites without surface modifiers, was determined by weighing the samples before and after pyrolysis. The conversion from w t o h to volYo was based on the densities of the polymer and the layered silicate, respectively, 1.2 and 2.74 g/mL. The true silicate content was used since surface modifiers con- tribute to the weight of the organosilicates. The error made in the determination of the true silicate content by pyrolysis was about 1 wtYo. The error originated from water molecules present between the layers as well as to the removal of OH groups in the layered silicate structure. In both tests, at least five samples of each composition were tested.

RESULTS AND DISCUSSION

Intercalation of Surface Modifiers

Sodium fluorohectorite (ME- 100) was modified by means of interlayer cation exchange of sodium cations for organic ions, thus rendering the silicate organophilic. Two protonated and one methylated dihydro-imidazolines were used because of their better thermal stability with respect to that of alkylammonium compounds: 1-methyl-2-norstearyl-3-stearinoacid-am- doethyl-dihydro-imidazolinium methosulfate (W75), hydroxyethyl-dihydro-imidazolinium chloride (HEODI) and ricinyl-dihydro-imidazolinium chloride (RDI), as well as a protonated monoamine, octadecyl-ammonium chloride [ODA). ODA is known to favor the formation of epoxy-layered silicate nanocomposites but has a lower thermal stability. The cyclic midinium cations were used successfully for modification of silicates and application in anhydride-cured epoxy systems and thermoplastic nanocomposites (20). Table 1 presents the results of thermogravimetric and WAXS analyses performed on the organosilicates.

The thermogravimetric analyzer determines (within ? 5 % ) the amount of organic cations exchanged between the silicate layers after the cation-exchange process. These values must be compared with the cation-exchange capacity of the fluorohectorite ME-100, which

is about 100 meq/ lOOg. Thus, one can observe that if ME-ODA, ME-HEODI and ME-RDI are fully exchanged, this is apparently not the case for ME-W75. ME-W75 retains only 67 meq/100g of surface modifier W75 between the silicate layers. However, W75 is a commercial product with tallow alkyl chains (1% C,,, 4% C,,, 31% C,,, and 64% CIS) so that it is diffi- cult to determine precisely the molecular weight of this surface modifier. Also it should be noted that control of the washing process could not be carried out by the test with silver nitrate. Therefore, the eval- uation of the quantity of W75 molecules retained between the silicate layers may not be accurate.

The secondzolumn in Table 1 shows that the surface modifiers present large differences in molecular weights. This means that the organosilicates have different organic content. Therefore, when mechanical properties are compared, it is important to evaluate the true silicate content of the composite. This is the content of silicate in the composites, not counting surface modifiers. The silicate content is evaluated by burning out the matrix and the organic ions and weighing the silicate residue.

The last column of Table I presents the interlamellar spacing of the organosilicates determined by WAXS. Since the interlameIlar spacing of dried fluorohectorite ME-100 is 0.94 nm, all the organosilicates show an increase in interlamellar spacing. This confirms that the organic cations are intercalated between the silicate layers.

Nanocomposite Synthesis The nanocomposites were cured using two different

curing cycles (No 1 or No 2) in order to look at their influence on the resulting nanocomposite structures. Table 2 presents the interlamellar spacing of the organosilicates in the epoxy matrix determined by TEM after each curing cycle.

Let us consider the results obtained with curing cycle No 1 (first column) and compare them with the interlamellar spacing of the pure organosilicates presented in Table 1. Once the organosilicates are swollen in the epoxy matrix, the interlamellar spacing of all the organosilicates increases systematically by 1.4 to 2.8 nm with the exception of ME-W75. This means that when ME-ODA, ME-HEODI and ME-RDI are used, the TGDDM and the DDS curing agent molecules diffuse and polymerize between the layers. This intragallery polymerization process separates the layers. Moreover,

Table 1. Organic Content and Interlamellar Spacing of the Different Organosilicates Determined, Respectively, by Thermogravimetry and WAXS.

~ ~~

Quantity of organic Molecular weight Interlamellar spacing of the surface modifier (@mole) cation exchanged (meqll 009) of the organosilicate (nm)

~~ ~

0.94 ME-100 - - ME-ODA 99 270 1.9 MEW75 67 637 3.2 ME-HEODI 90 352 2.2 ME-RDI 110 322 2.5



Table 2. Interlamellar Spacing of the Organosilicates in the Epoxy Matrix Determined by TEM After Curing Cycle No 1 (2h at 80°C + l h at 100°C + 4h at 150°C + 7h at 200°C) and Curing Cycle No 2 (2h at 140°C + 2h at 177°C + 7h at 200°C).

~ ~~~

Interlamellar spacing of the organosilicate (5 wt%) in the epoxy matrix determined by TEM (nm)

Curing cycle No 1 Curing cycle No 2

ME-ODA 4.0 4 .O ME-W75 3.0 2.5 ME-HEODI 3.6 4.13 ME-RDI 5.3 7.5b)

Ordered places. blrregular distribution.

the extent of separation of the silicate layers depends on the nature of the surface modifier. In the case of ME-W75, immiscibility between the organosilicate and the epoxy matrix was observed. It translates into a change of color of the resulting composite plates. In- deed, the composite plates produced with ME-W75 were brownish, whereas the ones produced with the other organosilicates and with the pure epoxy system were black.

If one now compares the effect of the two curing cycles, it is clear that for ME-HEODI and ME-RDI curing cycle No 2 has a larger influence on the separation of the silicate layers than curing cycle No 1. Thus, ME-RDI has an interlamellar spacing of 5.3 nm in the epoxy matrix after curing cycle No 1 whereas this spacing increases to 7.5 nm with curing cycle No 2. This effect was also observed previously in other epoxy systems (13, 21, 22). It is related to the speed of diffusion of the epoxy and curing agent molecules into the galleries during polymerization. Indeed, separation of the silicate layers occurs when there is a proper balance between the intragallery and the extragallery polymerization. In- creasing the initial curing temperature fkom 80°C (curing cycle No l) to 140°C (curing cycle No 2) favors the diffusion of epoxy and curing agent molecules into the galleries. This increases the cure kinetics of the intragallery polymerization and allows a larger separation of the silicate layers. However, it is surprising that chang- ing the curing cycle did not influence the interlamellar spacing of ME-ODA in the epoxy matrix since it remains constant (4 nm). When ME-W75 is used, a decrease of the interlamellar spacing is observed. I t means that ME-W75 does not swell in epoxy. The in- compatibility observed may find its origin in the fact that MEW75 was apparently not fully exchanged.

The degree of separation of the layers is intimately related to the curing reactions occurring in the vicinity of the silicate layers. For this reason, two series of DSC experiments were carried out to study the influence of the organosilicates on the homopolymerization of TGDDM and on the polymerization of the TGDDM/ DDS system.

Influence of Organosilicates on Curing Reactions The influence of the organosilicates on the homopoly-

merimtion of TGDDM and on the polymerization of the TGDDM/DDS system was studied by DSC. Figure 2

presents the exothermal peaks of the homopolymerization of TGDDM resin in the presence of the organosilicates.

The homopolymerizations of TGDDM and TGDDM with 10 wtYo of ME- 100 show similar behavior with an exothermal peak at 316°C. The difference in intensity of the exothermal peaks is caused by the presence of lower amount of polymer in the filled epoxy system. This suggests that the fluorohectorite ME-100 has no influence on the homopolymerization of TGDDM. On the other hand, all the organosilicates are shifting the exothermal peak to lower temperature. Thus, with 10 wtYo of ME-75, ME-HEODI or ME-RDI, the exothermal peak is shifted 6°C and with 10 wt?!) of ME-ODA, it is shifted 11°C. This illustrates that the surface modifiers have a significant influence on lbe homopolymerization of TGDDM resin. This was already reported by Lan et aL (16) with diglycidyl ether of bisphenol A resin and was attributed to the acidity of the surface modifier, which catalyses the epoxy ring opening polymerization. According to this interpretation, the protonated dihydro-imidazolines used in this study are not as acidic as the octadecyl-ammonium ions and therefore the homopolymerization of TG:DDM is catalyzed at a higher temperature. This is of great importance since the separation of the layers in epoxy-layered silicate nanocomposites has been attributed to this type of catalysis. Thus, these results sugest that ME-ODA should be the best surface modifiei- for the preparation of nanocomposites based on the TGDDM/DDS system. Moreover, it should be noted that adduct formation of deprotonated amines and amidines can occur and produce well-known catalysts for epoxy polymerization.

Figure 3 presents the exothermal peaks of the TGDDM/DDS system in the presence of the organosilicates. All the samples show an exotherm at about 270°C, which corresponds to the conventional polymerization of the epoxy matrix. However, at lower temperatures, the samples behave differently.

The fluorohectorite ME-100 has no apparent catalytic effect on the polymerization of the TGDDM/DDS (see top of Fig. 3). ME-W75 also appears to have only a small influence on the polymerization of the epoxy system. This means that the catalyzing effect of W75 on homopolymerization of epoxy does not significantly influence the polymerization.

POLYMER ENGINEERING AND SCIENCE, SEPTEMBER 2002, Yo/. 42, No. 9 1819

X. Kommann, R. Thomann, R. Mulhaupt , J. Finter, and L. A. Berglund

20 A A

A

A 0 TGDDM + 10 W% ME-100 + TGDDM + 10 wt% MEW75 O TGDDM + 10 W% ME-HEiODI X TGDDM + 10 wt% ME-RDI

n A 15 0% A

0 0 s f E '0 - TGDDM + 10 wt% MBODA

5

W

L, Q)

0

4 5

0 280 290 300 31 0 320 330

Temperature ("C) m. 2. Exotherms measured by DSC of the homopolymerization of Z D D M resin in the presence of the dtserent organosilicates.

In contrast, the exothermal peaks of epoxy mixed with ME-HEODI and ME-RDI present a large tail at low temperature (see bottom of Fig. 3). This suggests that these two protonated dihydro-imidazolines catalyze the polymerization of the epoxy. This catalysis may be attributed to the presence of -OH groups in their molecular structure. Indeed, it is well known that hydroxyl groups are able to catalyze epoxy-ring opening polymerization via hydrogen bonding and hydrogen transfer (23). Moreover, this catalysis effect of hydroxyl-substituted surface modifiers on the polymerization of epoxy resins has also been reported by Brown et aL (24).

ME-ODA also has a large effect on the polymerization (see bottom of Fig. 3). The presence of a shoulder around 180°C suggests that a secondary reaction occurs. It is apparently caused by the acidic catalysis of octadecylammonium ions on epoxy ring opening polymerization also observed in Fig. 2. However, increas- ing the initial curing temperature from 80°C (curing cycle No 1) to 140°C (curing cycle No 2) has no effect on the interlamellar spacing of ME-ODA since it remains constant at 4 nm (see Table 2) It means that ODA molecules do not provide a favorable environment for the diffusion of TGDDM and DDS molecules into the galleries even though ODA acts as a catalyst.

In contrast to the ME-ODA case, ME-HEODI and ME-RDI promote layer separation as the initial curing temperature is increased (see Table 2). Diffusion is apparently favored in this case. The presence of polar OH groups offers a more suitable environment for molecular diffusion into the galleries. Moreover, let us compare the interlamellar spacing of ME-HEODI and ME-RDI in epoxy after curing cycle No 2, 4.1 and 7.5 nm, respectively. It indicates that even the position of the OH groups in the molecular structure of the surface

modifier is of importance. Thus, RDI, with an OH group in the middle of its chain (see Fig. 1 ) promotes a larger layer separation than HEODI, which has an OH group on a side chain. It means that even the position of the OH group probably influences molecular diffusion. This is particularly interesting for future synthesis of surface modifiers promoting the formation of epoxy-layered silicate nanocomposites.

Morphology Figure 4 presents the transmission electron micro-

graphs of the composites based on ME-W75, ME- HEODI, ME-ODA and ME-RDI after curing cycle N" 2.

One can clearly see differences in the nanostruc- tures of these four materials. As previously discussed, the interlamellar spacing of the organosilicates in the epoxy matrix varies from 2.5 nm in the phase-sepa- rated composite based on ME-W75 (Fig. 4a) to 4 nm for the two nanocomposites based on ME-HEODI and ME-ODA (Figs. 4b and 4c). The largest interlamellar spacing, 7.5 nm, is observed for the nanocomposite based on ME-RDI (Fig. 4 4 . In all these materials, the silicate layers remain parallel. This long range ordered structure is typical for thermoset-layered silicate nm- ocomposites synthesized via in-situ polymerization. Indeed, it is the intragallery polymerization, which separates the layers. If the extent of this polymerization is uniform, the silicate layers remain parallel in the epoxy matrix during their separation. Contrary to the melt intercalation process, the lack of shear forces may explain why true exfoliation of layered silicate does not take place with in-situ polymerization of thermosets.

Figure 5 presents scanning electron micrographs of the composites based on ME-W75, ME-HEODI, ME-ODA and ME-RDI using backscattered imaging.

1820 POLYMER ENGlNEERlNG AND SCIENCE, SEPTEMBER 2002, V o f . 42, No. 9

High Performance Epoxy -Layered Silicate Nanocornposites

3 AA? A TGDDWDDS AA

0 EDDM/DDS + 10 wt% ME- 100 n M

+ TGDDWDDS + 10 wt% ME-W75 E

E E

1 2

3 1 8

W

L a,

0

0 1

140 180 220 260 300

Temperature ("C)

AbMA

AA A AA

A A TGDDM/DDS 0 TGDDM/DDS + 10 wt% ME-HEODI

bp

I

1 40 180 220 260 300

Temperature ("C) F ~ J . 3. Exothenns measured by DSC of the polymerization of TGDDM/DDS epoxy system in the presence of 10 wt% of ME-100, MEW75 [top), ME-HEODI, ME-R.DI, and ME-ODA (boftom).

The bright spots correspond to aggregates of layered silicate. It is apparent that the composites present different degrees of dispersion in the epoxy matrix. Thus, ME-W75 is poorly dispersed in the epoxy because of immiscibility (Q. 5 4 , whereas ME-HEODI and ME- ODA are more homogeneously dispersed in the epoxy matrix (Figs. 5b and 5c). The ME-RDI composite presents the best dispersion even though silicate aggregates can still be seen (Rg. 5d). Thus, there is a good correlation between the degree of dispersion of the silicate aggregates on a microscale and the degree of separation of the silicate layers on a nanoscale. In the

context of fracture properties, the degree of microscale dispersion is an important factor added to nanoscale structural effects.

Dynamic Mechanical Properties Rgre 6 presents the evolution of the maximum glass

transition temperature T&nax with true silicate content of the composites synthesized with fluorohectorite ME- 100 and the different organosilicates. The term T%, is used here because it does not represent the real glass transition temperature of the material. In thermosets, curing reactions become diffusion-controlled as the

POLYMER ENGINEERING AND SC/ENCf, SEPTEMBER 2002, Vof. 42, No . 9 1821

X . Kommann, R. Thomann, R. Mulhaupt, J. Finter, and L. A. BergIund

FYg. 4. 7Yansmisswn electron micrographs of the composites Synthesized with 5 wi% of ME-W75 [a). ME-HEODI (b), ME-ODA [c), and ME-RL)I fdj.

polymer reaches the glassy state so that for short curing times, the Tg of the final material cannot be substantially higher than the post-cure temperature. Since all the samples show a T& much higher than the post-cure temperature (200°C), it probably means that the samples were further cured during their dynamic mechanical thermal analysis. The systematic decrease of Tg,, observed with the nanocomposites may have two origins. It can either be due to a decrease of crosslink density of the epoxy network or result from the plasticizing effects of small free molecules present in the network. Usually, the height of the rubbery plateau is used to compare the crosslink density of different polymers. However, the nanocomposite structure provides large reinforcement of the rubbery modulus in epoxies so that such an evalua- tion cannot be performed. The presence of fluorohectorite ME-100 does not seem to substantially affect the epoxy network since Tg,, remains almost constant as the true silicate content increases. In contrast, the presence of organosilicates induces a grad- ual decrease of Tg,, as the true silicate content increases. Among all the organosilicates, ME-RDI and

ME-ODA have the highest impact on Tg-. Several explanations for this decrease may be suggested, including the thermal degradation of the surface modifier and network modification due to the homopolymerization of TGDDM or to the possible cation exchange of surface modifiers for curing agents during the mixing phase.

Tensile Properties Flgure 7 presents Young’s modulus a s a function of

true silicate content for the different composites. The Young’s modulus of the composite synthesized with ME-100 increases linearly with the true silicate content (discontinuous line).

This corresponds to the behavior of a conventional micro-composite. The composite based on ME-W75 presents a peculiar behavior. I ts modulus remains constant as the true silicate content increases. The immiscibility of ME-W75 in the epoxy matrix is responsible for this behavior. Poor interfacial adhesion is expected between the filler and the epoxy matrix. The Young’s moduli of the nanocomposites synthesized with ME-HEODI, ME-ODA and ME-RDI increase



3 260 - 0 v

X z 250 -

Fig. 5. Scanning electron micrographs of the composites synthesized with 5 wt% of ME-W75 (a), ME-HEODI (b), ML-ODA (c), and ME-RDI (d).

t

O O 0 0

A

0 0

0

A

0

0

A

240 ! I I I I I I I

0 1 2 3 4 5 6 7

True silicate content (~01%) Fig. 6. Evolution of the maximal glass transition temperature Tg- of the dicerent composites synthesized with ME-100 (.I, ML-W75 (01, ME-HEODI (O), ME-RDI (O), ME-ODA [A).


X . Kornmann, R. Thornann, R. Mulhaupt, J. Finter, and L. A. Berglund

5.5 h

6 i 5 9 m 1 1 II

3 4.5

9 4 G

E II)

M

3.5 0 1 2 3 4 5 6 7

True silicate content (~01%)

Rg. 7. Euolution of the Young’s modulus of the dizerent composites with true silicate content synthesized with ME-100 (W), ME-W75 (01, ME-HEODI (01, ME-RDl (O), ME-ODA (A). Disconiinmus and continuous lines are, respectively, fl to micro- and nanocomposites data

more steeply with the true silicate content [continuous lines). The stiffness improvements brought by the nanocomposite structure are substantially higher than that of the conventional micro-composites. The nanocomposite based on ME-RDI shows an increase in modulus of 32% for a true silicate content of 3.8 volVo. In comparison, the modulus of the conventional micro-composite based on ME-100 increases only by 14% for the same true silicate content. The difference between ME-RDI and ME- 100 is related to differences in effective volume fraction of reinforcement. The larger layer separation in ME-RDI causes a higher effective particle volume fraction, whereas particle stiffness is only marginally reduced by the intragallery polymer phase. This effect was also observed in previ- ous studies (15, 25) on nanocomposites prepared with DGEBA resin. In those cases, the Young’s modulus increased parabolically with true silicate content.

The ultimate properties of the material are only slightly affected by the nanocomposite structure. Thus, with 4 volVo of true silicate content, the tensile strength of the nanocomposites is lowered by 5%. This is ac- tually lower that the uncertainty of the measurement. The elongation at break of the nanocomposites is slightly lower than the pure epoxy.

Fracture Properties

FQwe 8 presents the evolution of the fracture toughness (bottom) and the fracture energy (top) of the different composites with true silicate content. The conventional composite based on ME-100 shows a lower increase in fracture toughness and fracture energy (discontinuous lines) with true silicate content than the nanocomposites (continuous lines).

The reason for this difference may be the apparent lower crosslink density of the epoxy matrix in the nanocomposites. Indeed, lowering the crosslink density of the epoxy may favor yielding in the matrix, which is known to improve toughness properties. However, it is also possible that the increase of effective particle volume fraction in the nanocomposites is responsible for part of the fracture property im- provement. Further experimental work is needed to iden* the fracture mechanisms of these materials. Previously, people have been reporting toughness im- provement in TGDDM/DDS resin using polyetherin- ide (PEI) as toughening agent (26). With 10 vol% of PEI, K,, was increased by 11 1%, Young’s modulus remained constant and the maximal glass transition temperature of the material was 264°C. If one consider the nanocomposite based on ME-HEODI with only 4.2 voWo of true silicate content, is increased by 112% and the nanocomposite has a Tg,, of 257°C but Young’s modulus is improved by 28%. Therefore, high performance epoxy-layered silicate nanocomposites seem good candidates to toughen TGDDM/DDS system and provide stiffness improvements at the same time. Moreover, the nanocomposite structure might provide additional advantages such as low water absorption and reduction of residual stresses.

CONCLUSIONS

High performance epoxy-layered silicate nanocomposites based on tetraglycidyl 4.4’-diaminodiphenyl methane WDDM) resin and cured with 4.4‘-diamh- odiphenyl sulfone (DDS] have been synthesized using fluorohectorite modifled with protonated dihydro-imidazolines and octadecylamine. The presence of OH


High Performunce Epoxy-Layered Silicate Nanocomposites

160 ______...- n .*.-- -E \ 120 t3 y 80

40

0 f 0 1 2 3 4 5 6 7

True silicate content (~01%)

1 n

0.8

0.6

8 0.4

-?

W

0.2 { o l I I I 1 I 8 f

0 1 2 3 4 5 6 7

True silicate content (~01%) Fig. 8. Evolution of thefracture toughness (bottom) and thefracture energy (top) of the di@ment composites with true silicate content synthesized with ME-100 (H), MEW75 (0). ME-HEODI (01, ME-RDI lo), ME-ODA (A). Discontinuous and continuous lines are, respectiuely, fit to micrw and nanocomposites data-

groups in the molecular structure of dihydro-imidazolines favored the formation of nanocomposites because it apparently had a catalytic effect on the polymerization occurring between the silicate layers and a positive influence on the diffusion of epoxy and curing agent molecules between the layers. Microscopy ob- servations revealed that there is a correlation between the separation of the layers on a nanoscale and the dispersion of the silicate aggregates on a microscale. However, an interlamellar spacing of 7.5 nm did not lead to a perfectly homogeneous structure of the nanocomposite on a microscale.

A systematic decrease of Tgmax was observed in all the nanocomposites. The reasons for this effect are not well understood but may include the thermal degradation of the surface modifier, the homopolymerization of the epoxy matrix as well as the possible cation exchange of surface modifiers for curing agent

molecules during the mixing phase. Nanocomposites based on TGDDM/DDS showed attractive and un- usual simultaneous improvements of fracture toughness and Young’s modulus without adversely affecting tensile strength. This offers new opportunities for impact modification via inorganic/organic hybrids without using conventional thermoplastics or rubbers as impact modifiers.

ACENOWLEDGMENTS

The authors thank the ‘Sonderforschungsbereich SFB 428” of the “Deutsche Forschungsgemeinschaft” (DFG) and the “Fonds der Chemischen Industrie” as well as the Swedish Engineering Research Council (TFR) for support of this project. Diplom-Chem. Jorg Frohlich from the Freiburg Materials Research Center is also acknowledged for fruitful discussions and for his help with experimental work.


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