Reinforcement of Adhesion and Development ofMorphology at Polymer–Polymer Interface via ReactiveCompatibilization: A Review

Genjie Jiang, Hong Wu, Shaoyun GuoThe State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute of Sichuan University,Chengdu 610065, China

Polymer blending is a common and effective way todevelop new materials with desirable physical and me-chanical properties. Since most polymer pairs are im-miscible, reactive compatibilization has been exten-sively studied to stabilize morphology of polymerblends and improve their mechanical properties. In thepast several years, considerable interest has beenexpressed in understanding the fundamental kineticsand mechanisms of the interfacial reaction, investigat-ing the reinforcement of the interfacial adhesion andthe development of morphological structure at poly-mer–polymer interface induced by the interfacial reac-tion. The present review focused on some theoreticaland experimental results that include the formationand growth of copolymers at the interface, and alsothe major factors such as reaction conditions, the con-centration and bulk properties of the functionalizedpolymer, the thermodynamic interaction between thefunctionalized polymer and the matrix, which can influ-ence the interfacial adhesion and morphological devel-opment. POLYM. ENG. SCI., 50:2273–2286, 2010. ª 2010Society of Plastics Engineers

INTRODUCTION

To prepare polymer materials with desirable properties,

blending two or more polymers have been used in many

aspects. The most important point in designing such mate-

rials is that most polymer pairs are thermodynamically

immiscible, which results in poor interfacial adhesion.

Commonly, one way to improve interfacial adhesion

between two immiscible polymers is the addition of a

third component (often a diblock copolymer), which can

connect immiscible polymer molecules through enhanced

entanglement on both side of the interface [1–10]. How-

ever, it is almost impossible for all the added copolymer

molecules to reach the interface during melt blending.

Therefore, in most cases, a number of the copolymer mol-

ecules prefer to form micelles [11–13] in the blends rather

than locate at the interface, consequently, the compatibili-

zation efficiency is low. Another way to improve the

interfacial adhesion is in situ reactive compatibilization

[14–30]. When two polymer chains with reactive func-

tional groups meet at the interface during melt blending,

they can react to form block or graft copolymers, which

can reduce interfacial tension and stabilize the morphol-

ogy of the blend. The compatibilization efficiency of this

strategy is relatively high [31] because the in situ formed

copolymers are pinned at the interface. Given the impor-

tance of reactive compatibilization in theoretical fields

and industrial applications, it is desirable to understand

the fundamental factors influencing the formation of the

copolymers at the interface, the relationship between the

interfacial adhesion and the in situ formed copolymers.

Since Ide and Hasegawa [32] found that the mechanical

properties of nylon 6/polypropylene (PA6/PP) blends were

remarkably improved with increasing the concentration of

maleic anhydride added to the blend in 1974, there have

been many articles published to study the reactive compa-

tibilization in different polymer blend systems theoreti-

cally [33–38] and experimentally [39–50] during melt

mixing, where there exists complex flow such as a com-

bined shear/elongational flow. To exclude the effect of

complex flow on the reactive compatibilization, a planar

interface in layered sandwich [51–55] was used as a

model for theoretical studies on the coupling reaction

mechanism of the functionalized polymers in the past dec-

ade. The correlative research was mainly concentrated on

the following two aspects. One is how the reaction

kinetics depends on functional group reactivity, the diffu-

sion of the functionalized polymer chains, and the

increased amount of in situ formed copolymers at the

interface. The other is how the interfacial tension, mor-

phology, and adhesion are influenced by the in situ

formed copolymers. Macosko et al. [56] reviewed recent

research progress in their group on the major factors

Correspondence to: Shaoyun Guo; e-mail: nic7702@scu.edu.cn

Contract grant sponsor: National Natural Science Foundation of China;

contract grant number: 50703027; contract grant sponsor: Special Funds

for Major State Basic Research Projects of China; contract grant number:

2005CB623800.

DOI 10.1002/pen.21686

Published online in Wiley Online Library (wileyonlinelibrary.com).

VVC 2010 Society of Plastics Engineers

POLYMER ENGINEERING AND SCIENCE—-2010

influencing the interfacial reaction such as the inherent

reactivity of functional polymers [57, 58], thermodynamic

interaction between polymers [59], functional groups loca-

tion along a chain [60], and the effect of processing flows

[61], however, some interesting experimental results

reported by other groups were not involved in this review.

Creton et al. [62] discussed recent theoretical develop-

ment of the relationship between macroscopic fracture

toughness and interfacial stress transfer in immiscible

polymer blends, whose interfaces are compatibilized by

the addition of block or random copolymers.

In this article, we review recent theoretical and experi-

mental developments in model polymer interfaces rein-

forced through in situ reactive compatibilization in the ab-

sence of flow. The simple schematic of the reaction is

given in Fig. 1. For simplicity, the functional groups are

all located at the end of polymer chains. When the model

interfaces are annealed under certain conditions, the cou-

pling reaction will occur between two reactive functional

groups of polymer A and B once they meet each other at

the interface. As a result of that, the copolymers are

formed in the interfacial region. The article mainly high-

lights the interfacial reaction kinetics, the mechanisms for

the formation of the copolymers at the interface, as well

as the relationship between interfacial reaction, macro-

scopic fracture toughness, and microscopic interfacial

structure.

REACTION KINETICS AND MECHANISMS

It is well known that the chemical reactions between

small molecules are usually under diffusion control due to

the reduced mobility of reactive sites. When both the re-

active agent and the reaction media are highly entangled

polymeric chains, the reaction kinetics and the mechanism

become even more complicated. Theoretical analyses of

the polymer–polymer interfacial reaction kinetics have

been explored by the simple model system, which is

designed as the irreversible reaction between end-func-

tionalized polymers at a planar interface.

In the past several years, there are two main theories

which were used to depict the reaction kinetics and the

mechanism of the coupling reaction at polymer–polymer

interface: one is diffusion-controlled (DC) models sug-

gested by Fredrickson and coworkers [53, 63, 64], and the

other is reaction-controlled (RC) models established by

O’shaughnessy and coworkers [54, 55, 65–67]. For sim-

plicity, it was assumed in both models that the degree of

the polymerization (N) of two polymers is the same, a flat

interface was formed between them under quiescent con-

ditions, and reactive groups were located at the chain end

of two polymers. The increase of the number of copoly-

mer chains per unit area (S) with the increasing reaction

time (t) can be expressed as dS/dt ¼ k(t)nAnB, where k(t)is the reaction rate, nA and nB are the content of the reac-

tive chains with function groups A and B, respectively.

Diffusion-Controlled Model

For DC reaction, the formation of the copolymers at

the interface is controlled by the mass transfer of the

functionalized polymer chains. The reaction rate constant

(k) was described by [53]:

k � DRgð2RgSV= lnNÞ (1)

where D, Rg, and N are the diffusion coefficient, radius of

gyration, and degree of polymerization of a reactive

chain, respectively. Sv is the ration of the interfacial area

to total volume fraction. The factor 2RgSV/ln N in Eq. 1can be interpreted as the volume fraction of interfacial

region of an inhomogeneous system that is accessible to

A–B coupling. Reduction of this accessible volume frac-

tion by the factor of 1/ln N � 1/ln t1 is associated with

the subdiffusive Rouse dynamics on the time scale that

local equilibrium is achieved inside the interface. It can

be obtained that k scales with molecular weight in the

present Rouse regime as k � 1/ln N for unentangled

chains (NhNei). When the polymer chains exceed the

entanglement threshold (N ‡ Ne), the reputation time of

chain t � N3 according to tube model of Gennes, resul-

tantly, k � 1 (N ln N).On the basis of the previous theoretical framework,

Fredrickson and Milner [63] proposed that there are three

time regimes to characterize the growth of the copoly-

FIG. 1. Schematic of reactive compatibilization between end-function-

alized polymer chains A and B to form copolymers at a flat interface

between two immiscible polymers.

2274 POLYMER ENGINEERING AND SCIENCE—-2010 DOI 10.1002/pen

mers. The initial growth rate of the copolymers is

described by:

ds=dt ¼ K0r20 (2)

K0 ¼ 40:7R4g=t lnðt=t0Þ (3)

where s(t) is the number of copolymer chains per area of

interface that have been formed in the time interval t afterinitiating the reaction. r0 is the concentration of the reactivechains with function groups A and B. t0 is a reference time

whose definition depends on the state of entanglement.

In the early stages of the reaction (s � t � sq), whenthe density of reactive chains in the interfacial region can

be considered to be the same as the value in the bulk,

coupling reaction is controlled by the quasilocal reaction

rate coefficient.

riðtÞ � r0 (4)

sðtÞ � K0r02t (5)

In the intermediate regime (sq � t � sr), a depletion

hole of reactants in the interfacial region is formed with

the progression of reaction,

riðtÞ � r0ðpt=trÞ�1=4(6)

sðtÞ � ð2=p�1=2Þs�ðt=tsÞ1=2 ¼ ð2=p�1=2Þr0ðD0tÞ1=2 (7)

Therefore, the reaction is controlled by center-of-mass

diffusion of the reactive chains to fill the depletion hole.

Finally, the reaction time corresponds to s ‡ sr, the

potential barrier arising from previously formed copoly-

mers at the interface suppressed the reaction in the last re-

gime, therefore, the diffusion of the functionalize chains

across the barrier plays a very important role in the fur-

ther coupling reaction at the interface. The copolymer

grows slowly as

sðtÞ � s�ln1=2½N1=2t=ðtslnNÞ� (8)

Therefore, the predicted time dependence is a linear

growth of r(t) in the initial time regime, a t1/2 growth in

the intermediate regime, and a (ln t)1/2 growth in the last

regime, respectively. In most circumstances, only the

crossover from center-of-mass diffusion-limited behavior

(t1/2) to saturation behavior [(ln t)1/2] can be observed as

expected. Furthermore, Muller [68] investigated the first

two regimes using Mote Carlo simulations and found that

the simulant reaction rate agreed with the theoretical pre-

diction for the intermediate regime of the DC reaction,

however, the simulation indicated a higher reaction rate

than predicted by the theory.

According to the Fredrickson’s model, if the interfacial

coupling reaction was diffusion-controlled, the functional-

ized chains located in the vicinity of the interface would

first participate into the reaction. As the reaction going on,

more functionalized chains need to diffuse from the bulk to

the interface, then a concentration gradient of the function-

alized chains between the interface region and the bulk

would be established. Clarke et al. [69, 70] monitored the

physical adsorption progress of deuterated polystyrene end-

functionalized with carboxylic acid (dPS-COOH) from PS

matrix with various molecular weight to a special oxide

layer of a silicon substrate, which has immobile reactive

sites. It was shown that depletion zones in the concentration

profile of dPS-COOH were up to four times longer than

necessary for the chains to diffuse from surface to interface,

indicating the progress could not be depicted simply in the

term of diffusion model. Norton et al. [71] investigated the

similar interface composed of dPS-COOH chains and func-

tionalized epoxy network, and also found that the extent of

the reaction could be explained more than DC model.

However, Harton et al. [72, 73] provided direct evi-

dence of depletion hole of the functionalized chains at an

interface by studying the coupling reaction between deu-

terated hydrozy terminated PS (dPS-OH) and methyl

methacrylate and methacrylic acide copolymer (PMMA-

MAA) at a relatively low temperature of 1208C. In their

experiment, dynamic secondary ion mass spectrometry

(DSIMS) was used to examine the dPS volume fraction

(udPS) at the interface. The real-space DSIMS depth pro-

files of the concentration of dPS-OH at the interface from

0 to 100 h were shown in Fig. 2.

From Fig. 2a–e, it could be observed that a depletion

hole formed and grew to a width with 40 nm apparently

from 0 to 24 h, which was a direct evidence of DC cou-

pling reaction at the interface. When the reaction time

was up to 100 h, shown in Fig. 2f, the depletion hole

disappeared completely and the dPS-OH/PS layer ap-

proached diffusion equilibrium. On the basis of the previ-

ous experimental results, they concluded that the coupling

reaction mechanism at an immiscible polymer–polymer

interface was DC, at least in the early stage.

Therefore, the in situ chemical reaction at the interfaces

is controlled by DC mechanisms or not depending on the

diffusion step of the polymers put in contact and interac-

tions between them. It has been experimentally verified

that the diffusion step of polymer chains at interface of

two polymers is clearly dependent on the structure of the

polymers [74, 75], the relaxation process of chains [76],

the molecular weight and distribution of polymers [77–

81], and the composition [82, 83] and compatibility of the

interfaces [84] besides the experimental conditions such as

time [85], temperature [86, 87], and shearing flow [88,

89]. Bousmina et al. [90] also studied the effect of surface

topology and chemistry on the diffusion step by a novel

rheometry technique and theoretical analysis that allow

determination of diffusion coefficients of both polymer–

polymer interfaces and bulk polymers. It was revealed that

the diffusion step is strongly dependent on the initial chain

ends distribution at the surface before contact.

DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE—-2010 2275

Reaction-Controlled Model

The RC mechanisms for interfacial reactive compatibi-

lization were theoretically depicted in detail by O’shaugh-

nessy and Sawhney [65, 66]. They considered that the

reaction depends on the reactivity of the functional groups

and the stage of the reaction. They predicted that the

reaction kinetics obeys mean field theory in the reaction

between functional groups with weak reactivity (Q « Q*):

kðtÞ ¼ Qa3h (9)

where Q, a, and h were the local functional group reactivity,the group size, and the interfacial thickness, respectively.

When the functional groups are sufficiently reactive

(Q » Q*), k(t) would be controlled by the diffusion of the

reactive chains.

For unentangled melts (NhNe)

kðtÞ ¼ ða4=taÞ½1=lnðR=hÞ� � 1=lnN (10)

and for entangled melts (N ‡ Ne)

kðtÞ ¼ ða4=taÞð1=½ðN=NeÞðlnðN=NeÞÞ�Þ � 1=N lnN (11)

where R was the coil size, ta was the relaxation time of a

single chain unit, and Ne was the critical degree of poly-

merization for entanglements.

Finally, the crowded copolymer chains formed at the

interfaces suppressed k(t) to exponentially small values.

Furthermore, O’shaughnessy and Vavylonis [66] system-

atically investigated four types of reaction kinetics,

through all time regimes, as a function of the reactivity

and bulk reactive group densities. Three distinct kinetic

sequences which gave rise to different types of reaction

kinetics as a function of reaction time were established.

Each of these three regions involved a different sequence

of reaction kinetics regimes with increasing reaction time.

At short time scale, simple mean field (MF) kinetics was

applied, the second order rate constant was then inde-

pendent of time. After this MF phase, there were two pos-

sibilities, a direct transition to first-order kinetics may

occur. However, if the functional groups were very reac-

tive, then a second-order diffusion regime would onset. In

this time regime, it was predicted that S � t/(ln t) for

unentangled chains and S � t/(ln t) during the two t1/4

regimes, whereas S � t1/2 for the t1/8 regime for

entangled chains. In long time regimes, first-order DC

kinetics would be observed, time dependencies during

these regimes are S � t1/4 for unentangled chains,

whereas for entangled systems, in chronological order S� t1/4, S � t1/8, and S � t1/4. When the accumulating co-

polymer ultimately saturated the interface, first-order DC

kinetics with S � t1/2 were predicted.

Oyama et al. [91, 92] proposed a new reaction model

and kinetic equation for interfacial reaction at polymer–

polymer interface by discussing the reaction kinetics and

mechanisms of interfacial reaction between amorphous

PA and polysulfon end-functionalized with phthalic anhy-

dride (PAH), and triazine. They considered that the cou-

pling reaction at polymer–polymer interface was the pro-

gression of the formation of a two-dimensional monolayer

at the interface between the polymers. The vacant sites

available for reaction at interface was defined as the pa-

rameter (S* 2 S), where S corresponds to the areal den-

sity of copolymers formed at time t and S* is a critical

value of the areal density of copolymers which can form

the monolayer at the interface. It was demonstrated that

the coupling reaction at polymer–polymer interface was

similar to the reaction at the gas/solid interface in surface

science, and the reaction followed pseudo-first-order

kinetics in the parameter (S* 2 S) over the whole time

scale based on the assumption that the coupling reaction

would be terminated at S ¼ S*. Except for the systems in

their experiment, the reaction kinetics in published articles

for reactive polymer interfaces such as amorphous styrene-

maleic anhydride copolymer/amine-terminated butadiene-

acrylonitrile copolymer (SMA/ATBA), SMA/PA11 [93],

and carboxyl end-functional polystyrene/precured epoxy

containing excess epoxide groups [94] were all analyzed by

using their model. The results showed that all the experi-

mental data agreed quite well with presudo-first-order

FIG. 2. Real-space DSIMS depth profiles of the dPS volume fraction

(udPS), including both unreacted dPS-OH within the PS layer and

PMMA-g-dPS located at the interface, are shown as a function of sample

depth for (a) 0 h, (b) 3 h, (c) 6 h, (d) 12 h, (e) 24 h, and (f) 100 h at

1208C. (Reproduced from Ref. 73).

2276 POLYMER ENGINEERING AND SCIENCE—-2010 DOI 10.1002/pen

kinetics. So, it was suggested that the coupling reaction pro-

cess between polymers was probably reaction-controlled,

and the spatial restriction at the interface played an impor-

tant role during the coupling reaction.

Schulze et al. [95, 96] measured the interfacial excess

(Z*/Rg) at three molecular weights of dPS end-functional-

ized with an amine group (dPS-NH2-22, 237, and 292)

and PMMA end-functionalized with anhydride (PMMA-

anh) interface using forward recoil spectrometry (FRES).

The measured interface (shown in Fig. 3a) was prepared

by floating films containing 8.4 wt% dPS-NH2 and PS

onto PMMA-anh layers that had been spun from the solu-

tion onto Si wafer. In addition, a special designed trilayer

interface geometry shown in Fig. 3b was used to observe

independently the diffusion of dPS-NH2 chains through

the PS matrix and the reaction with PMMA-anh eventu-

ally. The results of the measured growth of Z*/Rg with

time for all three molecular weights could be described

with polymer chain reactivity, rather than the diffusion of

the functionalized chains. Compared with the concentra-

tion profiles obtained from these two different designed

interface geometries indicated that the dPS-NH2 could dif-

fuse through the PS layer many times before significant

reaction could be observed under experimental condition.

These results proved that the coupling reaction rate at PS-

NH2/PMMA-anh interface was not limited by the diffu-

sion of the functionalized polymer chains.

Later, Yin et al. [97–99] also probed the interfacial

reaction between PMMA-anh and PS-NH2 with high and

low molecular weight by size exclusion chromatography

with a UV detector (SEC-UV). Reactive polymers were

labeled with anthracene to increase the sensitivity. It was

found that the conversion of PS-NH2 was faster when the

reactive chains were shorter. It likely showed that the

reaction kinetics was DC controlled because the diffusion

coefficient depends on the chain length. However, when

the dependence of the interfacial roughness on the molec-

ular weight of the reactive chains was observed by AFM

and TEM, it was found that the interfacial modification

was slower than the diffusion of the reactive chains. The

same conclusion was obtained by them as well as Schulze

et al. that the interfacial reaction between PS-NH2 and

PMMA-anh was not controlled by diffusion but rather by

the reaction kinetics of the anhydride and amine func-

tional groups.

The activation energy of the reaction is another clue,

which can be used to judge whether the coupling reaction

kinetics at polymer interface are controlled by diffusion

or reactivity. It is believed that the DC reactions usually

have low activation energy (\30 kJ/mol) because it is a

physical process, whereas the RC mechanisms have

higher activation energy attributed to a chemical process.

Oyama and Inoue [91] estimated the activation energy for

the reaction at SMA/ATBA interface was 120 kJ/mol,

indicating that the coupling reaction at this interface was

reaction-controlled. Kim et al. [100] investigated the

effect of surface functionalization on the interfacial adhe-

sion between PP and PA6. The functional groups of PP

surface modified by low-energy ion-beam irradiation

method are concentrated on a very shallow surface layer

(less than 70 A), which makes the diffusion of reactive

chains, can be neglected significantly. The interfacial

reaction kinetics in this system can be considered as reac-

tion-controlled. The calculated activation energy was 97.8

kJ/mol, which was also in favor of reaction-controlled

mechanisms. Kim et al. [101, 102] studied the effect of

molecular weight of dPS-NH2 and reaction temperature

on the reaction kinetics at PS/poly(2-vinylpyridine)

(P2VP) interface. It was experimentally showed that the

activation energy of the reaction was 156 6 7 kJ/mol,

and the value was not affected by the molecular weight

of dPS-NH2. Jiao [103] measured the activation energy of

the reaction between dPS-NH2 and PSMA interface, and

the value was 207 kJ/mol. For the reaction between the

same functional groups, the experimental values of the

activation energy were not the same, but all the experi-

mental results suggested that the interfacial reaction

kinetics was controlled by the reactivity of the functional

groups.

Interfacial Adhesion (Gc)

When two immiscible polymer chains are connected

by the covalent bond of copolymers formed at the inter-

face, the interfacial adhesion between two immiscible

phases will be enhanced. In the case of planar interfaces

enhanced via reactive compatibilization, the interfacial ad-

hesion is characterized by the fracture toughness (Gc) of

that. It has been shown that an asymmetric double cantile-

ver beam (ADCB) test is a reliable test for measuring Gc

of polymer–polymer interfaces [104–106]. It involves

propagating a crack along the interface between two poly-

mer beams using a wedge. The specific configuration of

the test is shown in Fig. 4. The wedge, typically a razor

blade, is inserted at the interface to induce a crack. The

elasticity of the two polymer beams enhances the stress at

the crack tip, causing the crack to be advanced some dis-

tance ahead of the wedge. Adhesion acts to suppress

crack propagation, counteracting the elasticity. Therefore,

the distance of the crack advances ahead of the wedge is

an indication of the adhesion. Generally, constant rate

mode is introduced, in the method, the razor blade of

FIG. 3. Schematic of two sample geometries used in Ref. 95.

DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE—-2010 2277

width D is driven at a constant rate of 3 lm/s, and the

crack length (a) at the interface is continuously moni-

tored. Gc is then calculated according to the following

equation [107, 108]

Gc ¼ 3D2E1h13E2h2

3

8a4E1h1

3C22 þ E2h2

3C12

ðE1h13C23 þ E2h23C1

3Þ2" #

(12)

where Ei and hi denote the Young’s modulus and the

thickness of material i, D is the thickness of the razor

blade, a is the crack length ahead of the blade, Ci is the

correctional factor for material i, which is expressed as

follows:

Ci ¼ ½1þ 1:92ðhi=aÞ þ 1:22ðhi=aÞþ 0:39ðhi=aÞ�=½1þ 0:64ðhi=aÞ� ð13Þ

On the basis of the above measured method and other

analysis techniques, the effect of reaction conditions and

the bulk properties of polymers on the Gc of model reac-

tive compatibilization interfaces were extensively studied,

including reaction time and temperature, the crystalline

structure near the interface, molecular weight, and archi-

tecture of the formed copolymers, as well as the relation-

ship between Gc and the areal density of copolymers at

the interface (S).

Effect of Reaction Time and Temperature on Gc

Theoretically, when the reaction temperature is higher

than the glass translation temperature of two polymers,

interfacial reaction can occur since the polymer chains

near the interface can diffuse each other. But very long

time is needed to form effective adhesion at low reaction

temperature. When the reaction temperature is above the

melt temperature of both polymers, Gc will increase rap-

idly within first several minutes. To control the variation

in Gc over wide order of magnitude as a function of suita-

ble reaction time, the reaction temperature is usually cho-

sen between the melt temperatures of two polymers near

the interface.

Boucher et al. [109, 110] investigated Gc of PP/PA6

interface annealed at the temperature between 1858C and

2238C as a function of the annealing time, and 5 wt% an-

hydride grafted PP was blended with PP for reactive com-

patibilization. As shown in Fig. 5, Gc increased with the

annealing time and then reached a level of saturation for

all annealing temperature. Many other researchers

obtained similar experimental results in other experimen-

tal systems [111–113]. When the annealing temperature

was a litter higher than the melt temperature of PA6, i.e.,

2238C, the rate and the level of saturation of Gc increased

dramatically. Laurens et al. [114] also studied the effect

of annealing temperature on Gc of PP/PA6 by using func-

tionalized PP (PPf) with different molecular weight as the

compatibilizer. It was found that annealing the samples

above the PA6 melting temperature increased strongly the

adhesion in samples containing high molecular weight

PPf, while such an increase was not observed in other

samples. Therefore, they believed that a sufficient molec-

ular weight of functionalized polymer chains should be a

necessary condition to obtain a dramatic increase in inter-

facial adhesion at high reaction temperature.

However, some other researchers observed a maximum

in the saturation Gc at a certain reaction temperature when

they studied the similar experimental systems. Seo and

Ninh [115] found Gc of PP/PA6 compatibilized with

1 wt% maleic anhydride grafted PP showed a maximum

around 2208C and then decreased at higher annealing

temperature. Cho and Li [116] also observed an unex-

pected maximum in the saturation Gc of PP and amor-

phous PA6 interface at around 1708C, which was ascribed

to a different interfacial morphology around this tempera-

ture. Kim et al. [100] obtained a maximum saturation

value of Gc at 2008C when they studied the effect of sur-

face functionalization by low-energy ion-beam irradiation

on the enhancement of interfacial adhesion between PP

and PA6. Since the PP molecules were all in melt state,

FIG. 4. The schematic drawing of the geometry used in the ADCB

test.

FIG. 5. Gc versus annealing time for PP/PA6 interface at four anneal-

ing temperatures. (Reproduced from Ref. 109).

2278 POLYMER ENGINEERING AND SCIENCE—-2010 DOI 10.1002/pen

the influence of the PP crystallization on the Gc could be

eliminated. They argued that the appearance of a maxi-

mum at 2008C was attributed to the change of fracture

mechanisms at the interface.

Effect of Functionalized Chains Content on Gc

In general, 5 wt% functional polymer is enough for re-

active compatibilization in immiscible polymer systems

prepared by melt blending, because new reactive interface

can be formed continually due to shear effect. But the re-

active interface with model planar geometry is constant,

whether more functionalized polymer or not need to be

added for reactive compatibilization in this system. Cho

and Li [116] investigated the effect of maleic anhydride

grafted PP (mPP) content on the interfacial fracture

toughness of amorphous PA6/PP at the same bonding

temperature. It was shown that Gc could not be enhanced

when the amount of mPP was lower than 0.5 wt%, even

at a bonding time of 2 h. As the mPP content increased,

Gc increased obviously which was ascribed to the in situ

formation of copolymers at the interface. When the mPP

concentration reached 3 wt%, Gc became slightly leveled

off. Gc of the samples as a function of bonding time for 3

and 5 wt% mPP were nearly overlapped, indicating 3

wt% mPP were enough to compatibilize this system effec-

tively.

Koriyama et al. [117] examined the interfacial thick-

ness and adhesion of the reactive polysulfon (PSU) and

amorphous PA6 interfaces as a function of MAH group

content in PSU. It was observed that the interfacial thick-

ness was almost independent of the MAH group content,

especially when the MAH group content was more than

0.56 wt%. According to their results, the interfacial adhe-

sion increased with increasing interfacial thickness fol-

lowing the relationship, Gc � k2. So, they believed that

MAH group content of 0.2 wt% was sufficient to form a

thick interface, thus, efficiency interfacial reinforcement.

Effect of Molecular Architecture of FunctionalizedPolymer on Gc

When the miscibility between functionalized polymer

and the matrix is really poor, it will induce a lack of

entanglements or cocrystallization between the in situ

formed copolymers and the matrix. Therefore, it is not ef-

ficient to transfer stress across the interface during frac-

ture test. On the contrary, when the functionalized poly-

mer has a good compatibility with the matrix, the block

of the copolymers can entangle or cocrystallize with the

matrix chains well, leading to an effective reinforcement

of interface by developing a plastic deformation zones at

the craze tip. The molecular weight of the functionalized

polymer also influences the entanglement between the

copolymers and the matrix near the interfaces. Laurens

[118] synthesized a series of succinic anhydride function-

alized PP with different molecular architectures (isotatic-

PP, isotatic-PP with 5 wt% PE, syndiotactic-PP, metallo-

cene-PP) and discussed the role of the architecture of PP

block of the in situ formed copolymers on the interfacial

fracture toughness enhancement for isotactic-PP/PA6

interface. It was found that the kinetics of fracture tough-

ness enhancement of samples annealing at the same con-

ditions was strongly influenced by the miscibility of func-

tionalized PP (PPf) with the matrix PP. When the PPf was

isotatic-PP, Gc of i-PP/PA6 interface increased rapidly

and then reached saturation value with annealing time due

to its good compatibility with the matrix. When the par-

tially compatible PPf (PE-PPf) was incorporated, the

extent of the increase in Gc became less obvious and a

plateau was not always observed. In addition, the molecu-

lar weight of PPf also influenced Gc, which was shown

that the maximum Gc was larger for the high molecular

weight PPf. They proposed that long copolymer chains

were more efficient than short copolymer ones to rein-

force the interface, which was attributed to the fact that

long copolymer chains can link several lamellae in the

vicinity of the interface. Eastwood and Dadmun [119]

compared the ability of a series of styrene and methyl

methacrylate copolymers with varying architectures to

compatibilize PS/PMMA interface. It was demonstrated

that the ability of reinforcement of the interface decreased

in the order pentablock i triblock i diblock i heptablocks irandom copolymers, and the multiblock copolymers

except heptablocks provided good interfacial adhesion,

which was attributed to multiple interface crossings

between blocks of monomer and homopolymers. To form

adequate entanglements with homopolymer, a critical mo-

lecular weight was required for block length. This was

exactly the reason that why the heptablock copolymers

give the relatively weak interface.

Besides the aforementioned factors, The enhanced ad-

hesion between immiscible polymer interfaces through in

situ reactive compatibilization depends on many other pa-

rameters, such as cooling rate [116], chains orientation

[120], and crystalline structure at the interface [121],

especially for semicrystalline polymer interfaces.

Relationship Between Gc and R

It is evident that the covalent bond connection at the

interface will build up with the increasing S, thus, Gc

increases as the progress of interfacial reaction. The

changes of Gc are strongly depended on the fracture

mechanisms at the interface. To understand which mecha-

nisms are responsible for the interface reinforcement, it is

useful to examine the relationship between Gc and S. Inthe case of glassy polymer interfaces [122, 123], two dif-

ferent relationships between Gc and S have been

observed: when the fracture mechanisms of the interface

are simple chain scission or chain pull-out without any

plastic deformation, Gc is found to vary as S linearly;

when crazing and plastics deformation are initiated at the

interface, then Gc is observed to be scaled with S2. For

DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE—-2010 2279

semicrystalline polymer interfaces, the relationship

between Gc and S is more complicated because the crys-

talline structure at the interface has profound effects

[124–127] on the plastic deformation properties of the

polymers. Boucher et al. [109] first established a theoreti-

cal model to correlate Gc with S at semicrystalline PA6/

PP interfaces. By the combination of the results of ADCB

test and XPS analysis, they were surprised to obtain

that all date plotted on log–log scale in agreement Gc !S2.03 6 0.18 scaling, which was similar to that observed

for glassy polymer interface in the crazing regime pre-

dicted by Brown [104].

But in further work by Boucher [110], the relationship

between Gc and S was not fit to the above model when

PP/PA6 interface was annealed a litter above the melt

temperature of PA6, especially when high molecular

weight PPf was incorporated as compatibilizer. As shown

in Fig. 6, when the annealing temperature was below

2208C, it was likely that the relation between Gc and Swas not influenced by the conditions for preparing the

samples, and the interfacial fracture mechanisms were

similar for all values of S. However, it was not right in

the case of the samples prepared at a temperature of

2238C. According to Boucher’s further investigation, the

change in fracture mechanisms for sample with high mo-

lecular weight made at 2238C was related to the fact that

the b-form PP crystalline structure was developed at the

interface. It implied that the interfacial fracture mecha-

nisms were not only influenced by the copolymers formed

but also by the crystalline structure near the interfaces.

Lauren [114, 118] reported that the relationships between

Gc and S were not consistent with Boucher’s model in

some experimental systems. He proposed that the relation-

ship between Gc and S depends not only on annealing

conditions but also on the molecular weight of functional-

ized polymers and the miscibility between the functional-

ized polymer and the matrix. The structure of the in situ

formed copolymers and the crystalline and orientation

behavior near the interfaces should also be considered.

Because of the complicated influencing factors, a consist-

ent and universal model has not been established to

describe the relationship between Gc and S for reactive

compatibilized semicrystalline polymer interface.

Interfacial Morphological Development by Reaction

Some investigation has shown that the copolymers

generated at the interface can form polymeric surfactants,

thus, decrease the interfacial tension between the incom-

patible polymer phases and result in a significant diminu-

tion in the particle size of dispersed phase [128–131].

Many studies have found that the final morphology of

polymer blends with an in situ reactive compatibilizer

depends on the bulk properties of the components, blend

composition, the amount and the molecular weight of the

formed copolymer, viscosity ratio between different

phases, and the processing parameters during melt blend.

Among them, the shear force is a very important factor

because it can induce droplet breakup and bring continual

new reactive interfaces. For flat polymer interface system,

discussed in this article, there is no shear effect during

reaction compatibilization. How does the morphology de-

velop in this kind of experimental condition is another

interesting topic.

From the theoretical studies [63, 65] concerning the

kinetics of coupling reaction at a melt polymer interface,

a free energy barrier (l*/kBT) to the coupling reaction

will develop due to the entropy loss involved in stretching

the ‘‘brush’’ of grafted chains as the ratio of interface

excess of the grafting chain to its radius of gyration (Z*/Rg) increases. The buildup of such a barrier during the

reaction will suppress the coupling reaction and ultimately

limit the Z*/Rg achievable within experimental time.

O’shaughnessy and Sawhney [65] asserted that the reac-

tion rates will slow down to near zero before the reaction

can form sufficient copolymers at the melt polymer inter-

face to effectively diminish the interfacial tension. How-

ever, Jiao et al. [132] demonstrated that it is possible to

make interfacial tension at polymer interface decrease sig-

nificantly by reactive compatibilization for short function-

alized chains, and the decreases of the interfacial tension

is large enough to cause polymer interface instability,

which leads to interfacial roughness and eventually the

formation of microemulsions. A theoretical model corre-

lated to 2Dc/c0 and Z*/Rg was developed to predict the

transition of the interface from flat to corrugated, where

the interfacial tension becomes to zero.

On the basis of the self consistent mean field calcula-

tions of Shull [133], the decrease in interfacial tension

(2Dc) due to the increased interfacial excess of copoly-

mer chains can be written in the following equation:

FIG. 6. Gc of PP/PA6 interface versus S in a log–log plot at four

annealing temperatures. (Reproduced from Ref. 110).

2280 POLYMER ENGINEERING AND SCIENCE—-2010 DOI 10.1002/pen

�DgkBT

¼ r0affiffiffiffiffiffi6N

p aZ�

Rg

� �(14)

where r0 is the monomer number density, a is the statisti-

cal segment length or polymer, a is a function of Z*/Rg

tabulated by Shull for the limit where the unreactive ma-

trix chains are much longer than the end-functional

chains.

For the high molecular weight polymers without graft

copolymers, the initial interfacial tension (c0) is given by

the Helfand and Tagami results [134],

g0kBT

¼ r0a

ffiffiffiw6

r(15)

where w is the Flory–Huggins interaction parameter.

So, the ratio of the decrease in interfacial tension to

the initial interfacial tension can be expressed as

�Dgg0

¼ 1ffiffiffiffiffiffiwN

p aZ�

Rg

� �(16)

if w and a are known, the model depicted by the equation

can be used to predict the �Dg/g0 versus Z*/Rg for func-

tionalized chains with different molecular weight. In

Jiao’s work [132], the critical value of Z*/Rg (Z*c/Rg) for

vanishing interfacial tension of PS-NH2/PSMA were

investigated based on a self-consistent mean field theory.

For reactive PS-NH2/PSMA interface, it was found that

Z*/Rg decreased with decreasing N of PS-NH2, and the

theoretical predication for two dPS-NH2 chains were Z*c/Rg ¼ 1.5, 2.4 for N ¼ 55 and 270 chains, respectively.

On the other hand, the onset of the interface instability

occurred at Z*c/Rg of about 1.8 for N ¼ 55 and 2.5 for N¼ 270 PS-NH2 chains by determining the root-mean-

square (RMS) surface roughness, as shown in Fig. 7. The

experimental results of transition value of Z*c/Rg are in

qualitatively agreement with theoretical predication [135].

Some recent experimental results [98, 102, 136] also

proved that the decrease of interfacial tension due to reac-

tive compatibilization can result in the interfacial rough-

ening or interfacial emulsification at a planar polymer

interfaces in the absence of flow. The development of

morphology is closely related to the reaction conditions,

the concentration, and molecular weight of reactive func-

tionalized polymer.

Lyu et al. [137, 138] examined the morphological de-

velopment at layered reactive PS-NH2/PMMA-anh inter-

face as a function of annealing time by AFM and TEM

observation. According to TEM results, the interface

between PS and PMMA was flat before annealing, shown

in Fig. 8a. The interface became quite rough after anneal-

ing the sample for only 20 min, and some parts of the PS

domain appeared to have pinched off at the interface and

moved into PMMA phase, as shown in Fig. 8b. When the

annealing time increased to 1 h, the interfacial roughness

increased further and the magnitude of the width of the

roughening zone was about 0.5 lm shown in Fig. 8c. The

interfacial roughening during annealing was also observed

by AFM. It was observed that the magnitude of the width

of the roughening zone increased from 0 to 0.2 lm when

the sample was annealed for 1 h.

Zhang et al. [139] studied the relationship between the

interfacial roughening process and the extent of coupling

reaction at PS/PMMA interfaces. The coupling reaction

was controlled by varying the concentration of PS-NH2 in

an unreactive PS matrix, while maintaining the same pure

PMMA-anh layer. It was obtained that interfacial rough-

ness was strongly influenced by the concentration of PS-

NH2 in the PS layer. For 10 wt% PS-NH2 sample, the

RMS roughness increased to about 1.4 nm in the first

5 min and then remained the same value within 120 min.

When the concentration of PS-NH2 was high, take 75

wt% for example, the RMS roughness increased to above

10 nm dramatically in about 10 min and then increased to

above 15 nm very slowly over 120 min. There was an ab-

rupt increase in RMS roughness when the PS-NH2 con-

centration changed from 25 to 30 wt% within the same

annealing time of 1 h at 1758C, just shown in Fig. 9. To

explain the development of interfacial roughness, the

maximum interfacial coverage (S*) was defined. They got

a conclusion that when the concentrations of PS-NH2 was

low, S \ S*, the interfacial roughness increased with

reaction conversion but remained low. When the concen-

tration was increased to 30 wt%, S [ S*, the interfacial

roughness increased dramatically, and interfacial emulsifi-

cation phenomena happened, which was proved by TEM

observation of the interfacial morphology.

The length of the functionalized chains, i.e., the molec-

ular weight also influence the interfacial roughness [98, 102].

FIG. 7. RMS roughness of the PS/PSMA interface after washing with

cyclohexane measured by SFM as a function of Z�c /Rg. N ¼ 55 (&) and

N ¼ 270 (*). (Reproduced from Refs. 132 and 135).

DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE—-2010 2281

On the one hand, the short reactive chains can diffuse to

interface faster than long ones, on the other hand, when

long functionalized chains are introduced, the high molec-

ular weight copolymers formed at the interface do not

leave the interface, and suppress the reaction rate to zero,

whereas, the copolymers formed by short chains can dif-

fuse across the interface, which is favorable to the reac-

tion progress and interfacial roughness.

When reactive compatibilized polymer blends are pre-

pared in the presence of shear (or elongational) effect,

micelles are usually formed in the matrix phase depending

strongly on the amount of in situ formed copolymers as

well as the molecular structure of those [140–143]. In the

Jiao’s and Zhang’s research, the interfacial roughness

both remain nearly constant after the transition point

though the extent of reaction increases. They attributed this

to the fact that the entire emulsified region (copolymers

coated droplets) formed after the transition point was

removed by the selective solvent during the preparation of

sample for AFM observation. Direct TEM cross-sectional

imaging of the interface supported their assumption. Kim

and coworkers [144, 145] first reported the formation of

microemulsions at reactive interface of PS-mCOOH/

PMMA-GMA in the absence of shear when the samples

were annealed at 1808C for 17 h. A simple illustration

given in Fig. 10 was used to explain the interfacial morpho-

logical change with reaction stages and the formation

FIG. 8. Representative morphologies of PS-NH2/PMMA-ah interface

after static reaction of (a) 0, (b) 20, and (c) 60 min. (Reproduced from

Ref. 137).

FIG. 9. RMS roughness versus PS-NH2 concentrations after 1 h anneal-

ing of PS-NH2/PMMA-anh-pyr. (Reproduced from Ref. 139).

FIG. 10. Schematic describing variations of interfacial morphology de-

velopment for (PS-mCOOH)/(PMMA-GMA) bilayer at 1808C. (Repro-duced from Ref. 145).

2282 POLYMER ENGINEERING AND SCIENCE—-2010 DOI 10.1002/pen

progress of microemulsions in the studied system. Three

distinct stages were considered: in Stage I, the functional-

ized chains reacted each other mainly near the interface and

the interface was not roughened. When the in situ formed

copolymers were covered at least a single layer at the inter-

face (Stage II), the functionalized chains diffuse into the

brush like copolymer layer, the interface became corru-

gated. At the beginning of the Stage III, the reactant chains

diffuse through the brush like copolymer layer, the inter-

face became more corrugated with the process of reaction.

When t[ tmicroemulsion, the copolymers began to pinch off

and then PS chains were encapsulated by PMMA chains,

which finally became microemulsions. They verified the

hypothetic process by TEM observation [144].

They suggested that the microemulsions structure in

the planar interface depended on the reaction time and the

molecular weight of the functionalized chains. Longer

reaction time and lower molecular weight more likely to

form microemulsions structure. Since the above experi-

ment was performed in plate rheometer, a very small os-

cillatory shear force was applied on reactive bilayer sam-

ple during reaction, it might be argued that whether the

microemulsion formation was induced by this very small

shear effect. According to the experimental results of

Kim, the morphological development near the interface of

the (PS-mCOOH)/(PMMA-GMA) bilayer sample anneal-

ing in the presence of oscillatory shearing used in their

rheological measurements was essentially the same as that

in the absence of oscillatory shearing. Microemulsions

were also observed in the bilayer sample annealing at

1708C under quiescent conditions.

In a word, the copolymers formed at the planar inter-

face is expected to decrease the interfacial tension of this

polymer interface. Instability and roughening of the inter-

face is observed when the interfacial tension decreases to

a critical value. The extent of interfacial roughness

depends on the concentration and molecular weight of

functionalized chains and reaction conditions. There exists

a critical values for Z*/Rg and the concentration of the

functionalized chains, where the interfacial tension van-

ishes and the interfacial morphology changes dramatically

from flat to corrugated one.

CONCLUSIONS

This article has intended to provide a summary of the

recent developments in the field of interfacial reactive

compatibilization for simple model systems, which are

assembled by bilayer or trilayer polymer films or sheets

with well-defined interfaces. Such systems enable one to

determine the kinetics parameters and mechanisms of the

reaction, to study the effect of the in situ formed copoly-

mers on the enhanced interfacial adhesion, to observe the

interfacial morphological development induced by reac-

tive compatibilization, etc. Since the interfacial reactive

compatibilization is performed at a circumstance exclud-

ing the complex effect of shear and elongational flows, it

is possible to correlate the progress of interfacial reactive

compatibilization directly with some factors such as proc-

essing conditions, the bulk properties of polymer matrix,

and functionalized chains. The following achievements

related to different aspects are worth mentioning.

Though the question related to the reaction kinetics at

the interfaces are controlled by the reactivity or diffusion

of functionalized chains is still in discussion theoretically

and experimentally, the reactivity and diffusion of the

functional chains at the interfaces play an important role

in the reaction kinetics.

The measured fracture toughness of reinforced interfa-

ces is crucially dependent on the details of annealing con-

ditions, the content and molecular weight of the function-

alized chains, and the miscibility between the formed

copolymers and matrix. The reinforcing mechanisms can

be depicted by the Brown’s model in some systems. How-

ever, in the case of some semicrystalline polymer interfa-

ces prepared under some conditions, the relationship

between Gc and S is deviated from the theoretical model,

which is attributed to the morphology and the crystallinity

near the enhanced interface.

With the progress of interfacial reaction, the interfacial

tension at the interfaces will decrease to negative value at

a certain critical grafting density, which is related to the

content and molecular weight of the functionalized chains,

thus, leads to the transition from flat interface to the cor-

rugated one. Possibility of the formation of the microe-

mulsions at reactive interface in the absence of flow is

found in some cases.

Important theoretical and experimental developments

in reactive compatibilization have been achieved by

studying the coupling reaction at simple polymer inter-

face, but some important problems related to that are still

worth paying more attention. More accurate theoretical

model should be established to describe the kinetics and

mechanisms of coupling reaction at polymer–polymer

interfaces, and to predict the interfacial adhesion due to

the formation of copolymers; Reliable analysis techniques

are necessary to be developed to characterize the crystal-

line structure and morphology, and to understand the rein-

forced effect between semicrystalline polymer interfaces

clearly. Given the importance of shear effect in commer-

cial reactive processing, another important aspect is the

problem concerning the effect of shear force applied on

the reactive polymer systems. It was reported that Kim

[145] found that high strain and high frequencies inhibit

the interdiffusion process. Other studies [146, 147]

showed that large strains may generate high density of

chain ends, which in turn enhance the interdiffusion pro-

cess. Until now, the links between the shear flow and the

interdiffusion process or interfacial reaction are still a

controversial point. The interfacial morphology in coex-

truded multilayer is well defined by number of layers and

layer thickness. The large number of layers offers large

amount of specific reactive area, and the shear effect can

be controlled by the rate of extrusion, which makes it a

DOI 10.1002/pen POLYMER ENGINEERING AND SCIENCE—-2010 2283

suitable model system to study interfacial reaction [148,

149]. According to one of our recent work [150], as the

layer number of PE/PE-MAH/PA6 microlayer samples

increased, the areal density of the formation of copoly-

mers at the interface increased obviously because of the

stronger and more durable shearing and elongational

forces [151] on the melts flowing through the laminating-

multiplying elements. This caused that the interfacial ad-

hesion between PA6 and HDPE layers was improved.

Further deep study was in progress. The ultimate aim of

the work is a creation of the theoretical and experimental

results that would enable one to control the reactive proc-

essing as a powerful way to produce polymer composites

with desire properties.

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