Advances in Addition-cure Phenolic Resins

Advances in addition-cure phenolic resins

C.P. Reghunadhan Nair*

Propellant and Special Chemicals Group, Polymers and Special Chemicals Division, Vikram Sarabhai Space Centre,

Thiruvananthapuram 695 022, India

Received 15 January 2003; revised 12 December 2003; accepted 8 January 2004

Abstract

Recent developments in the area of addition curable phenolic resins are reviewed. The article highlights the chemistry of

addition-cure phenolic resins and discusses the different strategies involved in their molecular design. Structural modification

through incorporation of thermally stable, addition curable groups on the novolac backbone is one strategy. The transformation

of phenolic hydroxyl groups to addition curable functions forms an alternate approach. Cross-linking of novolac or its

derivatives with a suitable curative also leads to addition-curable phenolic resin systems. This article examines the synthesis,

characterization and curing of noted addition curable phenolic systems. Their thermal, physical and mechanical properties are

discussed and the structure–property correlations examined. In selected cases, the adhesive properties of the systems have been

examined. The review includes discussions on the properties of the composites in relevant cases. The systems discussed here

include mainly allyl- and maleimide-functional phenolics, epoxy–phenolic, polybenzoxazine, bisoxazoline–phenolic,

acetylene-functional and propargyl ether phenolics and phenolic-triazine. The relative advantages and demerits of these

systems are discussed and their application potentials are considered.

q 2004 Elsevier Ltd. All rights reserved.

Keywords: Addition curable polymers; Phenolic resins; Novolac resins; Epoxy–phenolic; Bismaleimides; Allyl phenolics; Xylok; Diallyl

bisphenol; Bismaleimides; Maleimide-functional phenolics; Bisoxazoline–phenolic; Polybenzoxazine; Propargyl ether phenolic; Acetylene-

terminated polymers; Phenyl ethynyl polymers; Phenolic-triazine; High char-yielding polymers; Thermally stable polymers; Film adhesives;

Polymer adhesives; Alder-ene reaction; Polymer matrix composites

Contents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 404

1.1. Strategies for designing addition-cure phenolics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 405

2. Allyl-functional phenolics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 405

2.1. Allyl phenolic–bismaleimide blend . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 405

2.1.1. Allyl phenol–maleimide reaction mechanism . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 406

2.2. High performance polymers based on allyl phenol–BMI systems. . . . . . . . . . . . . . . . . . . . . . . . 408

2.3. Adhesives based on allyl phenolics–BMI . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 413

3. Bisoxazoline–phenolics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 418

3.1. Commercial PBOX–phenolic systems. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 420

3.2. Blends and composites of BISOX/Phenolic . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 421

3.3. Structural modifications of bisoxazoline–phenolics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 422

0079-6700/03/$ - see front matter q 2004 Elsevier Ltd. All rights reserved.

doi:10.1016/j.progpolymsci.2004.01.004

Prog. Polym. Sci. 29 (2004) 401–498

www.elsevier.com/locate/ppolysci

* Fax: þ91-471-415236.

E-mail address: [email protected] (C.P. Reghunadhan Nair).

http://www.elsevier.com/locate/ppolysci

4. Polybenzoxazines (PBZ) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 422

4.1. Features of polybenzoxazines . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 422

4.2. Cure mechanism and cure kinetics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 424

4.3. Structure–property relations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 425

4.4. Reactive blending of polybenzoxazines . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 428

4.5. Non-reactive blends and composites of benzoxazine . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 431

4.6. Degradation of polybenzoxazine . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 435

4.6.1. Thermal stabilization and degradation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 435

4.6.2. Chemical degradation of PBZ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 437

4.6.3. UV stability of PBZ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 437

5. Phenol–epoxy systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 437

5.1. Epoxy–phenol cure kinetics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 437

5.2. Latent catalysis of epoxy–phenol reaction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 440

5.3. Structure–properties relations in epoxy–phenol . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 441

5.4. Flame resistant epoxy–phenolic systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 447

5.5. Miscellaneous curative for novolac . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 451

6. Phenolic resins with phenyl maleimide functions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 451

6.1. Maleimide–phenolic resin cured with allyl–phenolics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 455

6.2. Maleimide–phenolic (PMF)–epoxy blend. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 456

7. Pendant phenol functional linear polymers. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 460

7.1. Pendant phenol-functional thermoplastics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 460

7.2. Pendant phenol-functional addition-cure systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 462

8. Propargyl ether functional phenolics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 464

8.1. Curing of propargyl ether resins . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 465

8.2. Structure–property relation in propargyl phenolics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 466

8.3. High molar-mass PN resins . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 468

8.4. Thermal degradation behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 469

8.5. Propargyl ether resins based on oligomeric novolac . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 471

8.6. Propargyl novolac–epoxy blend . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 472

9. Phenolic resins with terminal acetylene groups . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 474

9.1. Curing of EPAN resins. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 475

9.2. Thermal characteristics of EPAN resins . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 475

10. Phenolic resins with phenyl ethynyl groups . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 477

10.1. Phenyl ethynyl functional addition-curable phenolic resins. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 477

10.1.1. Cure and thermal characteristics of PEPFN resins . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 478

10.2. Condensation–addition cure phenyl ethynyl phenolic resins. . . . . . . . . . . . . . . . . . . . . . . . . . . . 479

11. Comparative thermal property of PMF, PN, EPAN and PEPFN resins. . . . . . . . . . . . . . . . . . . . . . . . . 480

12. Phenolic–triazine resin (P–T resins) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 480

12.1. Features of P–T resins . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 481

12.2. Properties of P–T systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 484

12.3. Structurally modified P–T resins . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 487

12.4. P–T/epoxy blends . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 489

12.5. Thermal degradation of P–T resins. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 490

12.6. Applications of P–T resins. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 490

13. Outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 490

Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 491

References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 491

C.P. Reghunadhan Nair / Prog. Polym. Sci. 29 (2004) 401–498402

Nomenclature

ABPF allyl-functional novolac ofbisphenol-A

AE allyl–phenolic epoxy

AP allyl phenol

Ar-DOPO-N novolac from DOPO reacted with tere-

phthaldicarboxaldehyde and phenol

B-a bisphenol A-based benzoxazine

{bis(4-phenyl-3,4-dihydro-2H-1,3-

benzoxazinyl) isopropane}

BER bispropargyl ether resins

BHPP bis (3-hydroxyphenyl) phenyl phos-

phate

BisA-N bisphenol A-novolac

BME 4,40-bismaleimidodiphenyl ether

BMI bismaleimide

BMIP bisphenol A-bismaleimide

BMM 4,40-bismaleimido diphenyl methane

{4,40- methylene bis (maleimido ben-

zene)}

BMS 4,40-bismaleimidodiphenyl sulfone

BPA bisphenol A

BPBA bis propargyl ether bisphenol A

BPh bisphthalonitrile

BPK bis propargyl ether bisphenol ketone

BPS bis propargyl ether bisphenol sulfone

BZ benzoxazine

CAI compression after impact

CNE o-cresol novolac epoxy

CNH cresol novolac hardener

CTE coefficient of thermal expansion

DABA 2,20-diallyl bisphenol A

DCPDP dicyclopentadiene—phenolic resin

DDM diamino diphenylmethane

DGEBA diglycidyl ether of bisphenol A {2,20-

bis (4-glycidyloxy phenyl) propane}

Dk dielectric constant

DMA dynamic mechanical analysis

DMF dimethyl formamide

DMSO dimethyl sulfoxide

DOPO 9,10-dihydro-9-oxa-10-phosphaphe-

nanthrene-10-oxide

DOPO-MA melamine-modified Ar-DOPO-N

DOPO-PF Ar-DOPO-N,blendedwithPFnovolac

DOPO-PN novolac from DOPO and 4-hydroxy

benzaldehyde

DPn degree of polymerization (number

average)

DSC differential thermal analysis

EPAN ethynyl phenyl azo novolac

EPAP ethynyl phenyl azo phenol

EPN novolac epoxy resin

FTIR fourier transform infra red

FTMS fourier transform mechanical spec-

troscopy technique

GC/MS gas chromatography-mass spectrum

GIC fracture energy

GPC gel permeation chromatography

HDT heat distortion temperature

HPM 4-hydroxy phenyl maleimide

IDT initial decomposition temperature

ILSS inter laminar shear strength

KIC the fracture toughness (plain-strain

stress intensity factor)

LOI limiting oxygen index

LSS lap shear strength

MDI 4,40-diphenyl methane diisocyanate

MMT montmorillonite clay

MPN melamine – phenol formaldehyde

novolac

NBR nitrile rubber

ODOPB 2-(6-oxido-6H-dibenzo kc,eloxa-pho-

phorin-6-yl)1,4-benzene diol

ODOPM DOPO–formaldehyde reaction pro-

duct (2-(6-oxid-6H-dibenz kc,el k1,2loxaphosphorin-6-yl)-methanol)

ODOPM-MPN

melamine – phenol formaldehyde

novolac (MPN)-modified ODOPM

ODOPM-PN phenol formaldehyde novolac (PN)-

modified ODOPM

OMMT organically modified montmorillonite

clay

OPN oligomeric propargyl novolac

1,3-PBOX 1,3-phenylene bisoxazoline

PBOX poly(bisoxazoline)

PBZ poly(benzoxazine)

PC polycarbonate

PCL poly(1-caprolactone)

PCS poly(4-cyanato styrene)

PCS-BD copolymers of 4-cyanato styrene with

butadiene

PCS-MMA copolymers of 4-cyanato styrene with

MMA

PDT peak decomposition temperature

C.P. Reghunadhan Nair / Prog. Polym. Sci. 29 (2004) 401–498 403

1. Introduction

Despite the emergence of several new classes of

thermosets, high performance polymers and several

other new generation materials that are superior in

some respects, phenolic resins retain industrial and

commercial interest, a century after its introduction.

Phenolic resins are preferred in a wide range of

applications, from commodity and construction

materials to high technology aerospace industry.

This recognition emerges from the fact that these

resins have several desirable characteristics, such as

superior mechanical strength, heat resistance and

dimensional stability, as well as, high resistance

against various solvents, acids and water. They are

inherently flame resistant, and evolve low smoke upon

incineration. Although phenolics cannot be substitutes

for epoxies and polyimides in many engineering

areas, their composites still find a major market in

thermo-structural application in the aerospace indus-

try due to good heat and flame resistance, excellent

ablative properties and low cost. These key properties

add to their market growth, and as a result of

innovative research, new products and applications

continue to emerge, demonstrating the versatility

and the potential of phenol resins to cope with

the ever-changing requirements and challenges of

advanced technology [1–5].

PEAR poly ether amide resin

PEK polyether ketone

PEPFN phenyl ethynyl phenol–phenol for-

maldehyde novolac resin

PEPFR phenyl ethynyl phenol–formaldehyde

resole resin

PES polyether sulfone

PF phenol formaldehyde novolac resin

phr parts per hundred parts of resin

PHRR peak heat release rate

PMAF phenolic resin from HPM, allylphenol

and formaldehyde

PMF phenolic resin from HPM, phenol and

formaldehyde

PMI N-phenyl maleimide

PMR polymerizable monomeric reactants

PN propargyl novolac

PS polyethersulfone

P-T phenolic-triazine

Pth-PBZ phthalonitrile functional polyben-

zoxazines

PTMO poly(tetramethylene oxide)

PU polyurethane

PVP poly(N-vinyl-2-pyrrolidone)

RT room temperature

RTM resin transfer molding

SBSS short beam shear strength

SEM scanning electron microscopy

SIN simultaneous interpenetrating net-

works

T5 temperature at 5% weight-loss in

TGA

Te temperature of end of thermal

phenomenon

TEM transmission electron microscopy

Terp-Bz terpene diphenol-based benzoxazine

Tg glass transition temperature

TGA thermo gravimetric analysis

TGMDA tetra glycidlyl methylene dianiline

THF tetra hydro furan

Ti temperature of onset of thermal

phenomenon (curing or decompo-

sition)

Tm temperature of maximum of ther-

mal phenomenon (curing or

decomposition)

TPP triphenyl phosphine

TPS T-peel strength

TTT time–temperature-transformation

UD unidirectional

Va benzoxazine from poly( p-vinyl

phenol)

Vba benzoxazine from vinyl phenol–

butyl acrylate copolymer

VBP vinyl phenol–butyl acrylate copo-

lymer

VP poly( p-vinyl phenol)

VRTM vacuum resin transfer molding

WLF William–Landel–Ferry


Undisputedly, classical phenolic resins based on

resole and novolac dominate the resin market. How-

ever, their acceptance as a universal material in many

engineering areas is hampered by some of the inherent

qualities derived from their special chemical structures.

These resins cure at moderately high temperature by a

condensation mechanism with the evolution of vola-

tiles, which necessitates application of pressure during

molding to form void-free components. The need for

the use of catalyst for curing and the limited shelf life of

resin at ambient conditions are also major shortcomings

of these systems. When compared to many known

thermally stable polymers, their thermo-oxidative

stability is low. The rigid aromatic units tightly held

by the short methylene linkages make the matrix brittle.

In view of this, a new chemistry is needed to modify the

cure of phenolic resins, in particular, a new method is

needed to chain extend and/or to cross-link phenolic

resins without production of volatiles and allow for

extended shelf stability at ambient conditions for the

formulated thermosets. In doing so, it is imperative that

the modifications do not impair the thermo-mechanical

characteristics of the resultant system. The concept of

addition cure phenolics gains significance in this

context.

1.1. Strategies for designing addition-cure phenolics

Several approaches have been reported for modi-

fication of phenolic resins and their cure chemistry.

Structural modification to confer addition-cure char-

acter has been one thrust area of research [6,7].

Addition-curable phenolic resins with improved

thermal and pyrolysis characteristics will be the

desirable resins in composites for thermo-structural

applications [8]. Higher char-yield leads also to a

better heat shielding. Such high char phenolics could

be potential candidates as matrices in carbon/carbon

composites too with obvious advantages [9]. The

major strategies in designing addition-cure phenolics

are:

(i) Incorporation of thermally stable addition-cur-

able groups on to novolac backbone

(ii) Structural modification (transformation) invol-

ving phenolic hydroxyl groups

(iii) Curing of novolac by suitable curatives through

addition reactions of OH groups

(iv) Reactive blending of structurally modified

phenolic resin with a functional reactant

This article gives an account of recent research

efforts in these directions for realizing addition

curable phenolic resins.

2. Allyl-functional phenolics

Allyl phenol–formaldehyde novolac, synthesized

by the allylation of novolac can cure thermally at

180 8C without the evolution of volatiles. On heating,

the O-allyl derivative rearranges to the C-allyl

polymer prior to cross-linking. The thermal curing

of this resin takes place by polyaddition at allyl double

bonds. The curing rate and cross-link density depend

on the content of the reaction centers in the polymer

molecule [10]. The allyl derivatives of phenols have

been used for the manufacture of glass fiber-

reinforced plastics and moldings, casting or impre-

gnating compositions of high heat resistance,

mechanical strength and chemical resistance [11].

However, achieving complete allyl curing is difficult.

Moreover, the cured matrix is not thermally stable due

to the thermal fragility of the cross-links arising from

polymerization of allyl groups. It has been observed

that allyl phenyl groups generally require prolonged

heating (nearly 6 h) at 250 8C for meaningful extent of

curing, which risks degradation of other fragile groups

in the network [12].

2.1. Allyl phenolic–bismaleimide blend

While reducing the cure temperature, the thermal

stability of the allyl phenolic novolac resins can be

further improved by reactive blending with bismalei-

mide compounds (BMI). The reaction between allyl

phenol (AP) and phenyl maleimide groups has been

exploited to synthesize a variety of polymer systems

with useful properties. In such reactive blends, further

improvement in properties is possible by way of

structural modification of either the BMI or the

phenolic ring. The properties of the resultant matrix

depend on the relative ratio of the two reactants, and

the cure schedules.


2.1.1. Allyl phenol–maleimide reaction mechanism

Earlier studies by Enoki et al. [13] showed that the

reaction between the two components proceeds via

the Ene reaction. The unsaturated Ene adduct

intermediate undergoes a further Diels-Alder type

reaction with BMI to give the bis- and tris adducts.

The intermediate step (Diels-Alder) is sometimes

referred to as Wagner–Jauregg reaction too [14]. The

total reaction sequence is also referred to as Alder-ene

reaction. At very high temperatures, a retro-Diels-

Alder reaction is also suggested. Investigation of

Reyx et al. [15] through reaction of model com-

pounds, (i.e. phenylmaleimide with AP) and identifi-

cation of products confirmed the Enoki mechanism.

The reaction sequences are shown in Scheme 1. A

multitude of reactions occurs at different temperatures

and reactant stoichiometries. All the three structures

are possible, but their relative concentrations could

depend on the stoichiometry. Ideally, a ratio 1:3

(allyl:maleimide) gives a maximum cross linking and

enhanced thermal stability, but this could lead to

brittle matrices. Lower maleimide-content ensures the

Ene structure-dominated, tough matrix with good

flexural properties, but with reduced Tg and thermal

capability. In most of the cases, a compromise of

various properties is achieved at an allyl/BMI ratio of

1:2. The Ene reaction between maleimide and allyl

groups has been separately investigated by Cunning-

ham et al. [16]. Although an unambiguous mechanism

has not been evolved, the generally accepted mech-

anism is along the lines of Reyx et al. However, others

have proposed modified mechanisms, through inves-

tigation of the cure chemistry by various techniques.

Thus, curing of BMI/2,20-diallylbisphenol-A (DABA)

system under different heat schedules has been

investigated by fourier Transform Infra Red (FTIR),

fluorescence, and UV-reflectance spectroscopy [17].

FTIR was used to quantify the extent of succinimide

Scheme 1. Reaction sequences for 2-allylphenol–phenyl maleimide reaction (Alder-ene reaction) [6].


formation and to identify the cross-linking processes,

which occurred during high temperature curing

(250–260 8C). High-temperature curing processes

were also identified by reflection spectroscopy with

a tungsten source. Various reaction pathways were

discussed in terms of their consistency with the

spectroscopic data. Another investigation on the

mechanism and kinetics of cure of a two-component

BMI formulation, composed of 4,40-bis(maleimido)

diphenyl methane (BMM) and DABA suggested a

copolymerization of maleimide and allyl groups as the

major reaction [18]. In this study, an in-situ real time

study of the progress of reaction was conducted in the

temperature range from 140 to 250 8C, using remote

fiber optic near-infrared spectroscopy. The principal

reaction observed was an alternating copolymeriza-

tion involving maleimide and allyl double bonds.

Maleimide homopolymerization was detected only in

the initial stages of reaction at temperatures above

200 8C. The extent of self-condensation (or etherifica-

tion) of hydroxyl groups on the allyl component,

which leads to the cross-linking was observed to vary

with the reaction temperature. Yet another sequence is

predicted for this reaction when catalyzed by

triphenyl phosphine (TPP)[19]. Thus, the studies on

model compounds, N-phenylmaleimide (PMI) and

o-allylphenol (AP) showed that the products of the

PMI/TPP system are oligomers and polymers of PMI,

whereas the main product of the PMI/AP/TPP system

is the PMI trimer, which had the five-member ring

formed via the phosphonium ylide intermediate as

shown in Scheme 2.

In the BMI/TPP system, maleimides only poly-

merize above 175 8C. On the other hand, maleimides

trimerize above 120 8C in the presence of DABA and

TPP. The BMI polymerization was faster in the

presence of DABA. Since the network structure differs

in the presence of TPP, the polymer properties could

depend on the cure conditions, the presence of

catalysts, etc. In continued work, Shibahara et al.

[20] established that the product also depends on the

reactants’ molecular environment. Thus, while inves-

tigating the thermal reactions of PMI and AP (or

DABA) using 13C NMR and GPC, they found that

reactions of PMI and AP (1:1 and 3:1 adducts)

proceeded through an Ene-reaction, and that sequential

Diels-Alder reactions accompanied the polymer of

PMI and AP. On the other hand, the products from PMI

and DABA were the Ene-adduct and the polymer, but

the Diels-Alder adduct could not be detected, in

contrast to the PMI/AP system. This difference in

reactivity for PMI-AP and PMI-DABA was attributed

to steric repulsion of DABA. A slightly different

mechanism, including the condensation of OH-groups

at higher temperatures has been proposed by Morgan

et al. for the curing reaction in BMI/allylphenol

systems [21]. The glass-transition temperature was

investigated as a function of curing conversion for a

diaminodiphenylsulfone-modified DABA/BMI resin

system at different temperature regimes, and modeled

Scheme 2. Formation of maleimide trimer from phenyl maleimide in presence of triphenyl phosphine.


using a modified DiBenedetto equation [22]. Although

the relationship between the glass-transition tempera-

ture and conversion of the BMI system conforms to the

DiBenedetto equation for conversion ,0.6 and at

lower cure temperatures, the results deviated signifi-

cantly from the equation at higher cure temperatures;

thus, it was an inadequate model for the system. FTIR

analysis showed that the major cross linking reactions

did not occur during cure of the modified BMI at and

below 150 8C. However, as the cure temperature was

increased, the cross linking reactions responsible for

three-dimensional network structures became more

dominant. At 190 8C, formation of succinimide rings

occurred in the curing process. The higher cure

temperatures induced a significantly faster initial

cross-linking rate. However, further cross linking

was retarded because the increase in the cross-links

sterically slowed further cross-linking activity.

2.2. High performance polymers based on allyl

phenol–BMI systems

Commercial addition-cure formulations based on

co-reaction of diallylphenols and BMIs are available.

Thus, Matrimide-5292 of Ciba-Geigy typically con-

tains DABA and BMM (Scheme 3). This is one of the

leading matrix resins for carbon fiber composite for

advanced aerospace applications. The earlier formu-

lations contained a lower concentration of BMI. As

discussed earlier, the neat resin properties are

composition-dependent. But in a narrow concen-

tration band, the variation in properties are not

significant except for a moderately better thermo-

mechanical profile at higher BMI load, as evident

from the representative data given in Table 1 [23].

The thermal capability and mechanical perform-

ance of the cured resin also depends on the cure

schedule. Low temperature cure produces apparently

stronger and tougher systems, with a marginal penalty

in thermal capabilities. Data in Table 2 substantiate

this [24]. Under a given cure schedule, BMI enhances

the thermal capability at the cost of toughness and

flexural strength. A compromise in properties is

generally observed for an allyl:BMI ratio of 1:2.

Suitably formulated, the resin system made up

of 4,40-bismaleimidodiphenyl methane, DABA

and desirable catalysts can give a high Tg (,315 8C)

matrix, despite a post cure temperature of only 200 8C

[25]. It is stable up to 450 8C, with a char-yield of

,30% at 700 8C. T-300 carbon laminates retained

Scheme 3. Components of Matrimide 5292.

Table 1

Neat resin properties of BMI/DABA of different molar ratio (cure schedule: 180 8C/1 h þ 200 8C/2 h þ 250 8C/6 h) [23] (Reprinted from 29th

Int SAMPE Symp by permission, q1984 Society for the Advancement of Materials and Process Engineering)

Property (MPa) BMI/DABA-1/1 BMI/DABA-1.2/1 BMI/DABA 0.87/1 (RT)

RT 149 8C 204 8C RT 149 8C 204 8C

Tensile strength 85 53 41 97 72 74 80

Tensile modulus 4430 2529 2100 4030 2943 2418 4243

Elongation (%) 2.3 2.6 2.3 3 3.05 4.6 2.3

Flexural strength 172 – – 192 – – 160

Flexural modulus 4207 – – 4140 – – 4114

Compressive strength 21 – – 220 – – –

Compressive modulus 2485 – – 2570 – – –

HDT (8C) 273 – – 285 – – 295

Tg (8C, DMA) 295 – – 310 – – –

GIC (J/m2) 195 – – 217 – – –


51% of short beam shear strength (SBSS) at 230 8C.

The processibility of the blend can be enhanced by

reactive diluents. Thus, a modified BMI resin system

for resin transfer molding (RTM) was prepared using

diallyl bisphenol-A ether and 1,4-diallyl phenyl ether

as reactive diluents for BMI [26]. Further toughening

of Matrimide has been achieved by incorporation of

polyetherimide particulates and Kevlar-49-whisker

[27]. Such systems are also characterized by enhanced

glass transition temperatures. Silicone has been

incorporated into the maleimide-modified allyl novo-

lac by reaction with substituted siloxanes. These

compositions have better heat- and moisture resistance

and are useful for semiconductor applications [28].

The thermal stabilities of BMI–DABA blends

modified with high-performance amorphous thermo-

plastics such as bisphenol A (BPA) polysulfone (PS),

polyether ketone (PEK), and polyether sulfone (PES)

bearing a phthalidylidene group, etc. have been studied

by differential scanning calorimetry (DSC) and scan-

ning electron microscopy (SEM) [29]. The stability of

thermoplastic components were compared with the

area of the endothermic peak that appeared within the

glass transition region for thermoplastic components in

the cured blends aged at a temperature below Tg: The

stability of thermoplastic was improved by the

formation of semi-interpenetrating polymer networks.

The stability of thermoplastic with higher Tg was more

easily controlled. As for epoxies, the N-phenylmalei-

mide–styrene copolymers are useful toughening

agents for the BMI resin composed of BMM, DABA

and triallyl isocyanurate [30]. The matrix structure was

controlled by changing the equivalent ratio of the two

allyl components. Morphologies of the modified resins

changed from particulate to co-continuous and to

inverted phase structures, depending on the modifier-

content. The optimum matrix structure contained

about 20 wt% triallyl isocyanurate and 5 wt% of

polymer additive wherein, the fracture toughness

ðKICÞ for the modified resins increased 100% at a

moderate loss of flexural strength and with retention in

flexural modulus and the glass transition temperature.

The Alder-ene reaction has been used to derive

diverse high performance polymers. Thus, partially

allylated novolac, when mixed with BMM and molded

under pressure at high temperature gave products with

high decomposition temperature (465 8C) [31]. By a

similar strategy, a thermosetting resin system for resin-

transfer molding based on allyl novolac and BMI has

been developed [32]. In this case, the novolac resin was

allylated by conventional route and BMI was used as

the curing agent and allyl phenyl ether, as the diluent. It

was blended with BMM (in 1:1 ratio). On heating, the

O-allyl derivative rearranges to the C-allyl derivative

as shown in Scheme 4. This was followed by the Alder-

ene reaction.

Table 2

Properties of BMI/DABA system (cure schedule: 175 8C/3 h

þ 230 8C/4 h) [24] (Reprinted from Advances in Polymer Science

by permission, q1994 Springer-Verlag)

Property BMI/DABA molar ratio

1.2/1 1.5/1 02/1 03/1

Flexural strength (MPa) 186 188 174 131

Flexural modulus (GPa) 4.02 3.94 4.05 4.14

Deflection (%) 7.78 7.3 5.53 3.5

KIC (MN/m1.5) 0.97 0.86 0.8 0.64

GIC (J/m2) 197 158 133 83

Tg (8C) 279 282 288 288

Scheme 4. Synthesis protocol and curing for allyl novolac–

bismaleimide system.


The viscosity–temperature curve and the vis-

cosity–time curve were used to characterize the

processing property of the resin system. The resin

system had a long pot life at the injection temperature.

The cured resin showed outstanding heat resistance

and good flexural properties. The flexural strength of

122 MPa and corresponding modulus of 3.53 GPa

confirm their good strength characteristics. The Tg

was 288 8C. However, the system absorbed 4.5%

water. The glass cloth composite fabricated using

RTM technology showed a flexural strength of

413 MPa which was retained to 90% and 65 at 200

and 300 8C, respectively.

In a nearly identical strategy, aromatic hydro-

carbon-modified maleimides have been used with

allyl etherified-novolac to improve the flexural

modulus [33]. Condensation resin of aralkyl ether

and phenol, i.e. polyaralkyl phenol (Xylok) is used

commercially as a high heat resistant phenolic resin.

Xylok is a desirable matrix resin for advanced

composites [34]. This is primarily because of its

excellent insulation, heat resistance and good proces-

sibility. The major disadvantage of Xylok is its low

reactivity and the brittleness of the cured structure.

One of the methods used to modify xylok, to improve

its properties is by allylization to derive the allyl

Xylok as per Scheme 5. Allyl xylok can be thermally

polymerized through the unsaturated bonds in its

structure, but it yields only an oligo polymer. In order

to obtain highly cross linked structure, allyl xylok can

be co-cured with BMIs, resulting in a class of high

performance matrix resin for advanced composites.

Thus, BMI is a desirable comonomer for allyl xylok

[35]. The copolymer of the BMM and allyl xylok

exhibits excellent heat resistance (no mass-loss up to

490–500 8C) with good mechanical properties and

hot-wet resistance. The heat distortion temperature

(HDT) of the water-aged (for 10 h) copolymer is

280 8C and the polymer absorbs only 2.3% water.

Although the properties are not commendably

high, the excellent high temperature retention is

worth mention. The good water-ageing resistance is

also reflected in the properties. The flexural strength

drops from 114 to 78 MPa at 250 8C. For wet

conditions, the corresponding drop is from 92 to

55 MPa. For glass composites, the drop in flexural

strength is proportional, from 360 to 242 MPa. The

Tg of the system is very high, almost close to the

decomposition temperature. The Tg of 490 8C is the

highest for a thermoset and drops only to 478 8C

under wet conditions. The polymer starts to degrade

rapidly at 490 8C. The glass-reinforced, compression-

molded material showed good thermal and

mechanical properties and outstanding dielectric

characteristics. Thus, the material is projected as a

high performance insulator [36]. During curing, no

by-products of low molecular weight are generated.

Hence, the resulting cured structure is compact and

has few defects.

High performance blends of BMI–DABA resulted

on blending the BMM/DABA system with bismalei-

midodiphenyl ether (BME), BPA–BMI (BMIP),

Scheme 5. Synthesis of Allyl-xylok and curing with bismaleimide.


allyl–phenolic epoxy (AE) and a thermoplastic-

modified polyetherketone (PEK-C) [37]. The influ-

ences of the various modifying components on the

properties of BMI neat resin and BMI/T300 compo-

sites were investigated. Results indicated that the two

BMI neat resin systems modified with AE and PEK-C,

and BMIP, AE and PEK-C had outstanding impact

strength (.18 kJ/m2) and excellent heat resistance.

BMI/T300 carbon composites corresponding to these

modified BMI resin systems showed excellent tough-

ness (CAI .210 MPa), outstanding mechanical prop-

erties, good processibility and relatively high Tg: A

copolymer of BMI, DABA and diphenylsilanediol was

synthesized in a similar approach [38]. The copolymer

could be cured around 200 8C, and the cured resins had

good thermal stability. A higher curing temperature in

the range 170–210 8C was favorable to obtain a more

thermally stable resin. By reducing the content of

diphenylsilanediol cyclo-homopolymer in the resin,

thermal stability could be improved further. Allyl

naphthols can replace allylphenols in Alder-ene

adducts [39]. The simplest co-reactant prepared in

this work was 7-allyloxy-2-naphthol, and satisfactory

matrices were obtained with appropriate BMIs.

However, the laminate coupons made by using this

system had lower thermal stability in comparison to

Matrimid 5292. The performance of this new chem-

istry was further tested by incorporating the naphthol/

maleimide Diels-Alder addition structure into two

other co-reactants. The most successful of these

compounds (i.e. the Diels-Alder adduct made from

the diallyl ether 2,7-dihydroxynaphthalene and

4-hydroxyphenylmaleimide) produced a cured neat

resin having Tg ,30 8C higher than that of a com-

parable system cured with the standard BMI co-reac-

tant (BMM). The system also possessed marginally

higher tensile strength and modulus. The fact that these

laminate coupons had better thermal performance than

the system using commercial co-reactant showed that

the presence of the Diels-Alder adduct structure in the

resin backbone was not detrimental to the normal

performance of the Ene-cured BMIs.

The Alder-ene polymers can be conferred good

ablative properties by introducing boron into the

molecular backbone of allyl compounds. Such

systems manifest improved thermal resistance, with-

out altering the mechanical strength [40]. Linear

polymeric boron-allyl compounds were synthesized

from DABA, or its mixture with bisphenol-A (BPA)

or 4,40 dihydroxy diphenyl sulfone, and boracic acid

(Structures B to E). The condensation product of

boracic acid and 2-allyl phenol (structure A) or

boracic acid and DABA (structure B), reacted with

formaldehyde (structure E and F, respectively) also

served as a co-reactant for BMM. The different

compositions are shown in Scheme 6. These allyl

derivatives were blended with BMM in 2:3 weight

ratio and heat cured. The mechanical properties

showed marginal improvement for systems containing

compounds A–D, as did Tg and HDT. However,

incorporation of boron gave significant enhancement

of the thermal stability for all cases. The more cross-

linked systems based on compounds E and F showed

enhanced Tg and thermal stability at the cost of the

mechanical and impact strength. The Ti values

increased by more than 60 8C and the char-yield

shot up to about 56–60% (from 21% for the non-

boron system). The relevant property data are

compiled in Table 3. The tensile properties are

comparable to the earlier reported vales for similar

systems (refer Tables 1 and 2). However, the flexural

properties are inferior, including for the control

composition. This may be a result of the different

allyl/BMI ratio and the cure schedule. From the

thermal characteristics, the authors concluded that the

system might possess good ablative properties, but no

true ablative evaluation was performed.

Simultaneous interpenetrating networks (SIN),

based on polyurethane – allylnovolac have been

reported [41]. The PU components were prepared by

reacting 4,40-diphenyl methane diisocyanate with

poly(tetramethylene oxide) (PTMO) of different

molecular weight ranges (UT series). The phenolic

resin component was synthesized by substituting the

hydroxy groups of the phenolic resin with the allyl

groups. To prove that the alkene groups can be applied

as a binding element between the networks to improve

the network compatibility, trimethylol propane mono-

allyl ether was chosen as the PU chain extender in one

series of the PU/allyl novolac resin SINs (TUT series).

A detailed study of the thermal, mechanical, and

dynamic properties and morphology revealed that the

extent of phase mixing of the graft PU/allyl novolac

resin SINs (TUT series SINs) was significantly

improved over that of UT series SINs. This result

was consistent with the loss tangent shift in dynamic


Scheme 6. Synthesis strategy for Boron-containing allylphenol–bismaleimide system [40] (Reprinted from Journal of Applied Polymer Science

by permission, q1999 John Wiley and Sons).

Table 3

Mechanical and thermal properties of boron-containing alder-ene polymers. Control system contains DABA/BMM in the ratio 0.8/1 [40]

(Reprinted from Journal of Applied Polymer Science by permission, q1999 John Wiley and Sons)

Property A B C D E F Control

Tensile strength (MPa) 83.6 83.3 80.7 79.5 77.8 74.4 73

Tensile modulus (GPa) 3.72 3.77 3.7 3.78 3.78 3.79 3.61

Elongation (%) 2.3 2.32 2.45 2.25 2.25 2.17 2.2

Flexural strength (MPa) 124 139 129 122 116 98 112

Impact strength (kJ/m2) 12.6 10.9 9.6 17.7 11.8 8.9 13

Tg (8C) 277 282 278 283 285 330 274

Ti(8C) 489 490 487 495 495 499 426

Char-yield (%) 57 56 54 60 58 61 21


mechanical analysis (DMA) measurements and with

transmission electron microscope. The mechanical

properties of the graft SINs (TUT series) were lower

than those of the original SINs (UT series). In a

related work, SINs based on poly(urethane–epoxy)/

allyl novolac resins were prepared by cross-linking a

urethane–epoxy adduct with diamino diphenyl-

methane (DDM) and allyl novolac resin simul-

taneously [42]. A urethane – epoxy adduct was

synthesized by terminating the urethane prepolymer

composed of 4,40-diphenyl methane diisocyanate

(MDI) and PTMO with epoxy. The thermal, mechan-

ical and dynamical characteristics of these SINs

correlated well with morphological observations

from transmission electron microscopy (TEM). All

these SINs exhibited a two-phase structure and the

mechanical properties were correlated with the phase

behavior. Further modification in the structure was

effected by synthesizing polyurethane and phenolic

resin via a maleimide-terminated PU/allyl nonyl

novolac resin (ABCP) [43]. The polyurethane was

again a product of MDI and PTMO. The composition,

compatibility, mechanical and thermal properties of

these inhomogeneous network polymers were inves-

tigated. The miscibility was improved very little, but

the tensile strength had a clear improvement when the

PU’s soft segment, i.e. PTMO, was longer. This was

attributed to the entanglement between the com-

ponents. Attempts to improve the phase miscibility in

a maleimide-terminated polyurethane/allyl nonyl

novolac resin system by way of co-reaction with

BMM did not meet with success [44]. In fact, a higher

content of BMM promoted phase separation, although

a lower BMM-content enhanced the tensile strength.

2.3. Adhesives based on allyl phenolics–BMI

The Alder-ene chemistry was used to derive high

temperature phenolic adhesives, based on the reactive

blend of an allyl-functional bisphenol-A novolac

(ABPF) and BPA–BMI (BMIP) [45]. ABPF was

synthesized by reaction of DABA with formaldehyde

under acidic conditions. The reaction was followed by

GPC, and the synthesis conditions were optimized to

produce soluble polymer of desirable molecular

weight distribution. The polymers were characterized

by FTIR, NMR and GPC. ABPF was reactively

blended with BMIP and cured through the Alder-ene

reaction at high temperatures. The cure characteristics

of ABPF–BMIP blends were studied using DSC and

DMA, which evidenced the multi-step cure reactions

characteristic of the Alder-ene systems. The DSC

thermograms in Fig. 1 substantiate this.

The cure reaction showed two distinct exotherms

in the temperature ranges 100–170 and 180–270 8C,

respectively, and a less prominent exotherm initiated

at about 280 8C and extending beyond 300 8C. The

first broad exotherm encompasses the Ene reaction,

and the second unresolved one contains the Wagner–

Jauregg and final Diels-Alder steps. DSC of the resin

system already cured at 200 8C/3 h and 250 8C/2 h are

presented in the same figure. DSC of the 200 8C/3 h

cured system shows appreciable residual cure exo-

therm beyond 260 8C. Since this system can not

contain any unreacted maleimide, the exotherm

spreading from 260 8C can be assigned to the curing

of residual unsaturated groups. The exotherm onset at

260 8C is triggered by curing of the vinyl phenyl

groups (of probably the Ene-adduct), followed by the

residual allyl groups. The DSC of the 250 8C/2

h-cured one also shows the presence of less significant

residual cure exotherm around 300 8C, assignable to

the allylphenol polymerization. Except for these,

Fig. 1. DSC of bisphenol A bismaleimide(BMIP)-diallylbisphenol A

novolac(ABPF) blend; (A) before cure; (B) after cure at 200 8C/3 h

and (C) after cure at 250 8C/2 h [45] (Reprinted from Polymer

International by permission, q2001 Society of Chemical Industry).


the cure becomes practically completed at 250 8C, and

at this temperature, the network could retain part of

the unsaturated bonds derived from the initial reac-

tions. The non-isothermal DMA of the glass prepreg

(in Fig. 2) implied that the major cross linking process

is the Diels-Alder step occurring beyond 220 8C, since

no considerable modulus build-up occurs prior to this

temperature corresponding to the Ene-reaction.

The isothermal DMA of ABPF–BMIP system at

three different temperatures in Fig. 3 show that the

cure is accelerated at higher temperature, and the

ultimate modulus is dependent on the temperature of

cure. The cure is practically complete in about 90 min

at 250 8C. However, a 6 h curing was performed to

achieve maximum reaction at this temperature. This

cure schedule also ensures complete polymerization

of any residual allyl groups. The cure sequences are

identical to those shown Scheme 1, and extrapolating

the same chemistry, the cured structure of the resin

can be depicted as in Scheme 7.

The adhesive properties of the cured blend are

given in Table 4. Although the adhesive properties of

the system are only moderately good, it is remarkable

that the properties are retained to greater than 100% at

150 8C. Moderate cross linking achieved through a 1:1

maleimide–allyl phenol stoichiometry and a stepwise

cure up to a maximum of 250 8C for 2 h was found to

be very effective for the optimum LSS properties in

this series. The stoichiometry of BMIP/ABPF was

varied from 0.8:1.0 to 3.0:1.0 and the LSS at RT and

at 150 8C were determined. The dependence of

Fig. 2. Non-isothermal DMA of bisphenol A bismaleimide (BMIP)-

allylbisphenol A novolac (ABPF)/glass prepreg (heating rate: 5 8C/min,

frequency 1 Hz, N2 atmosphere [45] (Reprinted from Polymer

International by permission, q2001 Society of Chemical Industry).

Fig. 3. Isothermal DMA of bisphenol A bismaleimide (BMIP)-allylbisphenol A novolac (ABPF) blend, evaluated as glass prepreg, at different

temperatures: (A) 160 8C, (B) 200 8C, (C) 250 8C (frequency 1 Hz, N2 atmosphere, heating rate 5 8C/min for the dynamic part) [45] (Reprinted

from Polymer International by permission, q2001 Society of Chemical Industry).


adhesive properties on reactant stoichiometry is

shown in Fig. 4. It was observed that the LSS

properties optimized at a 1:1 ratio. The variation of

relative retention of LSS at 150 8C shown in the same

figure confirms the optimum value for this ratio. A

BMI concentration lower than this stoichiometry

leads to very low LSS at 150 8C, and also the lowest

high temperature retention (about 90%) of the LSS

due to the reduced number of cross-links in the

system. An excess of allyl groups in such case leads to

more chances of Ene homopolymerization resulting in

more linear structures than does the Diels-Alder

cyclo-adduct. A stoichiometric excess of BMIP could

lead to increased cross-linking reactions and the

homopolymerization could result in rigid, brittle

network structure, with decreased load-bearing capa-

bility. However, the higher cross linked system is

conducive to an excellent retention of properties at

150 8C. Interestingly, the 1:1 combination also

showed a good retention of the adhesive strength up

Scheme 7. Likely Alder-ene reaction products from 1:1 Dallylbisphenol–formaldehyde (ABPF)–bismaleimide (BMIP) blend [46] (Reprinted

from High Performance Polymers by permission, q2000 SAGE Publications).


to 250 8C. The adhesive properties of this system were

improved tremendously on matrix modification using

polysulfone (PS) and polycarbonate (PC) [46]. The

performance advantage was more in the case of PS,

showing an optimum improvement at a 20% loading,

as against PC exhibiting maximum properties at 10%

loading. The related adhesive data are compiled in

Table 5. The performance advantage of the additives

is clear in Fig. 5, showing a comparison of the

adhesive properties and thermo-adhesive profiles of

different systems.

The properties decrease beyond 200 8C in the case

of toughened systems, whereas the properties are

retained for the neat resin well above 250 8C. The

relatively better performance of a homogeneous blend

resulting from PS is manifested as a significant

improvement in the properties at ambient conditions,

but these decrease drastically at about 200 8C, in

contrast to the PC-modified system. An SEM analysis

of the modified formulations corroborated this relative

difference as resulting from different morphological

features in cases of PS and PC. A comparatively

uniform distribution of the thermoplastic component

leading to the co-existence of the resin-rich and

additive-rich phases was found to enhance the

toughness of the PS-system, whereas precipitation of

larger particles in PC-modified system was less

efficient for enhancing the adhesive properties.

DMA corroborated the observations made in SEM.

Existence of co-continuous phases of thermoplastic,

resin matrix and thermoplastic-dissolved matrix in

PS-modification and a clear phase separation in the

case of PC-modified system, manifested independent

glass transitions. The DMA for various systems are

shown in Fig. 6. Both the additives decreased

Table 4

Adhesive properties of cured BMIP–ABPF (1:1 stoichiometry)

under different cure conditions [45] (Reprinted from Polymer

International by permission, q2001 Society of Chemical Industry)

Sl. no. Cure conditions

(temperature/

time) (8C/h)

LSS at

RT (MPa)

LSS at

150 8C

(MPa)

Retention

of LSS

at 150 8C (%)

1 160/4 1.9 2.6 137

2 160/0.5

þ 200/3

2.8 3.4 121

3 160/0.5

þ 200/0.5

þ 250/2

4.1 4.8 117

4 160/0.5

þ 200/0.5

þ 250/6

3 4 133

Table 5

Adhesive properties of thermoplastic-modified BMIP-ABPF system (1:1 stoichiometry; cured at 160 8C/30 min þ 200 8C/30 min þ 250 8C/2 h)

[46] (Reprinted from High Performance Polymers by permission, q2000 SAGE Publications)

Properties Un-modified PES-content (phr) PC-content (phr)

10 20 30 10 20 30

LSS at RT (MPa) 4.1 13.9 19.3 14.8 11.3 8.8 8.3

LSS at 150 8C (MPa) 4.8 11 16.5 13.7 9.5 6.9 6.7

LSS at 200 8C (MPa) 5 7.1 11 10.7 8.4 6.2 5.5

LSS at 250 8C (MPa) 5.2 4.6 6.4 5.5 8 5.4 3.5

TPS at RT (kN/m) Poor 0.32 0.38 0.4 Poor Poor Poor

Fig. 4. Dependence of adhesive properties on reactant stoichiometry

for bisphenol A bismaleimide (BMIP)-diallylbisphenol A novolac

(ABPF)system (a) at ambient, (b) at 150 8C, (c) percent retention at

150 8C [46] (Reprinted from High Performance Polymers by

permission, q2000 SAGE Publications).


the modulus. The low temperature Tg is due to the

phase-separated additives. The co-existing phase

showed intermediate Tg: The transition zones of

different phases and corresponding Tg values are

given in Table 6. The unmodified systems manifested

high Tg; of the order of 350 8C, which increased to

390 8C on enhancing the maleimide ratio. All these

systems provided good adhesion up to 250 8C and can

be used for moderate load bearing applications at high

temperatures, as is required in certain aircraft and

defense applications.

The TGA of the blend of ABPF and BMIP with

varying maleimide to AP ratio indicated that the

thermal stability of the system was only marginally

improved by an increase in BMI stoichiometry [47].

The effect of the BMI structure on the adhesive

properties was also evaluated, using four different

BMIs, namely, BMIP, BMM, BME and 4,40-bismalei-

mido phenyl sulfone (BMS) [48]. The polar groups in

BMIS and BMIE contributed to better adherend wetting

and consequently better LSS at ambient. The high-

temperature LSS retention was comparatively better for

BMIP and BMS systems. The thermo-mechanical

properties of the blend, implied from DMA, showed

that BMM and BME systems with higher cross link

density manifested higher Tg: Although the absolute

values are not high, all systems exhibited remarkably

high retention of LSS (.100%) at 250 8C. The data for

different BMIs are given in Table 7.

DABA, the precursor diphenol for ABPF was also

reacted with BMIP stoichiometrically [48]. The LSS

values at temperatures up to 250 8C for this system in

comparison to those of its polymer analogue (i.e.

ABPF) are presented in Fig. 7. At low temperatures,

Fig. 5. Adhesive Performance advantage of polysulfone (PS)- and

polycarbonate (PC)- modified adduct of bisphenol A bismaleimi-

de(BMIP)-diallylbisphenol A novolac(ABPF) at 1:1 maleimide-

allyl phenol/stoichiometry.

Fig. 6. DMA of (—) bisphenol A bismaleimide (BMIP)-allylbi-

sphenol A novolac (ABPF); (- - -) bisphenol A bismaleimide

(BMIP)-diallylbisphenol A novolac (ABPF)-polysulfone-20(PS-

20); and (–--–--–) bisphenol A bismaleimide (BMIP)-diallylbi-

sphenol A novolac (ABPF)-polycarbonate-20 (PC-20), showing the

different Tg regimes [46] (Reprinted from High Performance

Polymers by permission, q2000 SAGE Publications).

Table 6

Tg data for different BMIP-ABPF systems from DMA [46]

(Reprinted from High Performance Polymers by permission,

q2000 SAGE Publications)

System BMIP/

ABPF

molar ratio

Tg (1)a

(8C)

Tg (2)b

(8C)

BMIP-

ABPF

01:01 335 355

BMIP-

ABPF

02:01 ,380 ,380

BMIP-

ABPF

03:01 ,380 ,390

BMIP-

ABPF-PS20

01:01 190,260,300 190,260,345

BMIP-

ABPF-PC20

01:01 150,210,320 150,210,350

a Based on the deflection point, tan d in curve.b Corresponding to tan d peak temperature.


the LSS values for the DABA-based systems are

better than those of the corresponding ABPF-based

ones. This is attributed to a better wetting of the

adherend by the low viscous, monomeric BMI-DABA

blend in contrast to the polymer-based one. However,

the thermo-adhesive profile is far superior for the

polymer version. Thus, the LSS dropped significantly

above 150 8C for DABA, whereas it continuously

increased for the ABPF-based system up to 250 8C

(Fig. 7). The better high temperature performance of

the polymeric (BMIP-ABPF) system is a consequence

of the comparatively higher cross-link density leading

to enhanced cohesive strength for the network. The

thermo-adhesion profiles of the two systems are in

tune with their relative DMA behavior included in

Fig. 6. The DABA-system showed a broader glass

transition initiated below 200 8C, with a Tg maximum

of ,300 8C ðtan dmaxÞ; in contrast to the ABPF system

showing a higher Tg of about 350 8C under identical

cure conditions.

3. Bisoxazoline–phenolics

The unusual addition co-reaction of novolac

phenolic resins with phenylene bisoxazoline has

been explored to derive a new class of non-

conventional phenolic thermosetting resin by Cul-

bertson et al. [49]. The polymerization involves a

tertiary phosphine-catalyzed reaction of bisoxazoline

with a phenol-free novolac resin leading to an ether–

amide copolymer as shown in Scheme 8.

The systems are suited for high performance

composite applications [50]. The key features,

which foretell the great usefulness of bisoxazoline–

phenolic resins in aerospace and other high perform-

ance application areas, include:

† No volatile bye products produced during curing

† Low cure exotherm, about 20% of epoxies and BMI

† Long term thermo-oxidative stability (10,000 h at

177 8C)

† Low cure shrinkage (,1%)

† High neat resin modulus (.500 MPa)

† High compression and shear strength for composites

† Excellent toughness (GIC is ,5 times greater than

those of epoxies and BMI)

Table 7

LSS properties (in MPa) of different BMI-ABPF systems (1:1

stoichiometry; cured at 160 8C/30 min þ 200 8C/30 min þ 250 8

C/2 h) [48] (Reprinted from Polymers and Polymer Composites by

permission, q2003 Rapra Technology)

System RT 150 8C 200 8C 250 8C

BMIP-ABPF 4.1 4.8 5 5.2

BMM-ABPF 4.2 4.3 4.3 4.2

BME-ABPF 5.4 5.5 5.5 5.6

BMS-ABPF 5.9 6.6 7.1 7.6

Fig. 7. Comparative adhesive properties of diallylbisphenol-A

(DABA)-bisphenol A bismaleimide (BMIP) and bisphenol A

bismaleimide (BMIP)-diallylbisphenol A novolac (ABPF) systems

at different temperatures. Scheme 8. Additionpolymerisationofbisoxazoline–phenolicsystem.


† Excellent adhesion to glass, reinforcing fibers and

particulates

† Long shelf life for resin and prepregs

† Low melt temperature and low viscosity for melt

† Low flammability and smoke release, meeting

aircraft use regulations

† Low CTE (,4 £ 1025/8C)

† Easy prepreg formation as claimed by the inventors

† High Tg (170–295 8C) and high service temperature

(275 8C)

The usual cure cycle is 175 8C with a post cure at

225 8C. Composites are processable by RTM tech-

nique. The fiber-reinforced copolymers possess the

low smoke and heat release requirements of materials

for aircraft interior applications [51]. Through selec-

tive use of catalysts, a very long shelf life for

unrefrigerated prepregs at typical aerospace autoclave

conditions (,176 8C) is achievable. Reactive, low

viscosity additives enhance the formation of prepregs

with the resin (as claimed by the inventors). Based on

this chemistry, several compositions with many

interesting properties have been patented. Electrical,

physical and mechanical properties of the neat resin

suggest that these new thermosets could be useful in a

variety of electrical applications. Further, their

chemical, physical and mechanical properties,

coupled with the ease of formulating tough, machine-

able materials (which can be highly filled with metals

and other fillers) confirm opportunities to use them in

a wide variety of plastic material and/or mould-

making applications. The glass transition temperature

of the system can be widely tuned through variation in

stoichiometry of the phenolic resin and the bisoxazo-

line. The Tg increases with the PBOX-content as

shown in Fig. 8.

The physical, thermal and mechanical properties

are also composition-dependent. The variation in

fracture toughness with composition is shown in

Fig. 9, which includes the toughness range for

aircraft grade commercial TGMDA/DDS systems

for comparison. Toughness becomes optimized at a

PBOX/Phenolic weight ratio of 40/60, and is much

higher than those of commercial epoxy systems.

Thus, successful bisoxazoline–phenolic compo-

sitions contain about 60% by weight of novolac.

Data in Table 8, on the comparative mechanical

performance of typical aerospace resins show

the superiority of PBOX–phenolic over the epoxies

[52]. The properties are even superior to those of

the toughened BMI. The 40/60 PBOX/PF compo-

sition has better mechanical characteristics than the

60/40 composition.

The mechanical data of composites of bisoxazo-

line–phenolic indicate that these resin systems have

excellent potential for use in a range of high

Fig. 8. Dependence of Tg on PBOX/phenolic composition [49].

Fig. 9. Composition dependence of fracture property of PBOX/

Phenolic system in comparison to epoxy system (TGDMA/DDS:

tetraglycidyl methylene diamine/diamiodiphenyl sulfone) [52].


performance applications, particularly when high

modulus, good compressive strength and high inter-

laminar shear strength are demanded. The glass

composite properties given in Table 9 show the

superiority of these systems over the conventional

resins [49]. The very small cure shrinkage (,0.1%)

and low CTE are believed to contribute greatly to the

good physical properties and excellent thermal shock

resistance of laminates. The shelf life of the prepregs

can be tuned by selective use of catalysts. The carbon

composite properties are either comparable or superior

to the best-improved epoxy systems, as evident from

the selected data given in Table 10.

3.1. Commercial PBOX–phenolic systems

Southwest Research Institute, USA has developed

and patented bisoxazoline–phenolic thermoset resins

that are tough and possessing low-flammability and

high service temperature (176 8C), (named as

PEARe, Poly Ether Amide Resin) and also their

carbon fiber- or fiberglass composites. The work is

aimed at developing materials for a wide range of

applications, including subsonic and supersonic air-

crafts, ground and marine transportation, lightweight

composite pipes, heat shields, and other high-strength,

non-conductive materials for the construction,

electrical, and oil industries [53]. PEARe is a

lightweight, strong, economical and versatile

Table 8

Comparative mechanical properties of neat PBOX–phenolic and high performance epoxy and BMI [52] (Reprinted from 20th Int SAMPE Tech

Conf by permission, q1988 Society for the Advancement of Materials and Process Engineering)

Property 1,3-PBOX /phenolic

(40/60)

1,3-PBOX /phenolic

(60/40)

TGMDA/DDS

(untoughened)

BMI, XU292/DDS

(improved BMI)

Flexural strength (MPa) 193.7 172.4 91.7 184.8

Flexural modulus (MPa) 5033 4909 3440 4000

Tensile strength (MPa ) 89.6 – 56.8 93.8

Elongation at break (%) 1.8 – 1.8 3

Tensile modulus (MPa) 5137 – 3737 3889

Compressive strength (MPa) 236.5 256 201.3 210

Compressive modulus (MPa) 4882 4789 1958 2482

Fracture energy (GIC; J/m2) 157–223 – 54 250

Table 9

Mechanical properties of glass-cloth reinforced laminates compared

to epoxy and phenolic systems [49] (Reprinted from 34th Int

SAMPE Symp by permission, q1989 Society for the Advancement

of Materials and Process Engineering)

Property 1,3-PBOX/

phenolic (40/60)

Epoxy

(Hexcel F-161)

Phenolic

Flexural

strength (MPa)

758.5 606.8 608.1

Flexural

modulus (GPa)

31 24.8 28.3

Tensile

strength (MPa)

451.6 489.5 403.4

Tensile

modulus (GPa)

28.3 24.1 –

SBSS (MPa) 76.5 – –

Table 10

Comparative mechanical properties of carbon composites of

PBOX–phenolic and high performance epoxy systems [52]

(Reprinted from 20th Int SAMPE Tech Conf by permission,

q1988 Society for the Advancement of Materials and Process

Engineering)

Property 1,3-PBOX/PF

(40/60)

and AS-4

DOW71788.00

epoxy (tough)/

DDS and AS-4

NAEMCO

RIGIDITE-

5225 (improved

EPOXY)

and CELION

Flexural

strength (MPa)

1593 2069 1793

Flexural

modulus (GPa)

113.1 115.1 124.1

SBSS (MPa) 108.3 89.6 111.7

Tensile

strength (MPa)

1924 1999 1724

Tensile

modulus (GPa)

140 122 129.6

Compressive

strength (MPa)

1524 1682 1448

Compressive

modulus (GPa)

115.8 132 126.2


composite matrix. This resin fuses a variety of

significant performance characteristics into a single

material. Composites made with the PEAR resin

system tends to experience longer wear resistance,

longer life expectancy, flame resistance, and the

ability to endure cyclic stress loading.

Some critical product qualities of PEARe are:

† Heat and flame resistance

† Low toxicity when exposed to high levels of heat

† Excellent mechanical properties

† Electrical insulating properties

† Dimensional stability

† Relative ease of processing

The manufacturer markets different varieties of

PEAR resins, depending on their composition,

solvent, etc. [54] (Refer to Table 11).

Some of the properties of PEAR resins are given in

Table 12 [54]. From the property data, it appears that

the resin formulation is close to 40/60 PBOX–

phenolic composition.

3.2. Blends and composites of BISOX/Phenolic

As the phenolic resin market for aircraft interiors

and mass transportation applications grows, fire

retardant standards in stringent applications must

be met by improved formulations. One method to

enhance flame resistance is to add siloxane to the

resin. Addition and exfoliation of montmorillonite

clay (MMT) is yet another approach. These two

techniques have been adapted in concert

for bisoxazoline–phenolic system [55]. Thus, the

co-reaction of an epoxy-terminated siloxane with

a novolac phenolic/bisoxazoline reduced the peak

heat release rate (PHRR) by 27% compared to a

control. At 8% siloxane, PHRR dropped from 225 to

164 kW/m2 at a 75 kW/m2 heat flux. DMA testing

showed a Tg of 220 8C for the 8% polysiloxane-

modified material, compared to the Tg of 248 8C for

the unmodified version. The clay (MMT) was then

incorporated into a novolac resin using solvent

fractionation techniques, and employing ethanol and

toluene. Clear plaques were produced, and X-ray

Table 11

Features of different polyether–amide from bisoxazoline [54]

PEARe

product

version

Manufacturing -

process

Typical

applications

Hot melt RTM resin

infusion

Aircraft and structural

components

Solvent Pre-preg

lamination

Sports equipment,

aircraft interiors

Low melt RTM Low stress applications,

consumer products

Table 12

Properties of PEAR resin [54]

Key properties

Volatiles during cure 0.0

Shrinkage during cure

Viscosity 100 cps

Fracture Toughness ðGICÞ 156–223 J/m2

Curing exotherm (J/g) 78.0

CTE 42 £ 1026 8/C

Tg Range 87–121 8C

Specific gravity

(neat resin)

1.25

Neat resin strength (MPa)

Tensile strength 90

Compressive strength 237

Flexural strength 194

Neat resin modulus (MPa)

Tensile modulus 5134

Compressive modulus 4879

Flexural modulus 5030

Composite strength glass 57% resin 43% (MPa)

Tensile strength 490

Flexural strength 758

Composite modulus (glass 57% resin 43% (GPa))

Tensile modulus 28.2

Flexural modulus 31

Composite short beam shear strength (glass 57% resin 43%) (MPa)

Room temperature—dry 108

93 8C Hot-wet 91

Thermal cycles

(glass cloth composite)

0

Cycles (MPa)

50

Cycles (MPa)

Flexural strength 784 779

93 8C Hot-wet 32.4 33.1

Total heat release (kW/m2) (FAA-OSU heat release calorimeter 35)

In two minutes 65.3

Maximum heat release 67.6

Flammability (kW/m2) (NIST cone calorimeter)

Maximum rate of heat

release

295.0


diffraction (XRD) showed exfoliation due to loss of

the clay peak at 19 A.

3.3. Structural modifications of bisoxazoline–

phenolics

Structurally modified bisoxazoline derivatives

have led to new poly(ether–ester–amide) multiblock

terpolymers and copolymers, which are hetero-phase

materials, endowed with elastomer properties [56].

Aliphatic and aromatic alternating poly(ether–amide)

copolymers represent potential engineering materials.

Binary and tertiary polyester/polyamide composites

reinforced in situ are prepared by reactive blending in

the melt. The chemical bonds formed between the

separate phases via diblock copolymers improve the

compatibility of blend components. Perfectly

branched and hyper branched poly(ether–amide)s

based on bisoxazoline have also been reported [57].

The nucleophilic ring-opening addition reaction of

phenol groups towards oxazoline units has been used

for the preparation of hyperbranched poly(ether–

amide)s [58]. For this, the AB2 monomer, viz. 2-(3,5-

bishydroxyphenyl)-1,3-oxazoline, was synthesized

and converted to a highly branched polymer in bulk

or in solution at temperatures above 190 8C. The

resulting hyperbranched polymers exhibit a degree of

branching of 50%, as verified by high-resolution

NMR spectroscopy, and are highly soluble in polar

organic solvents with low solution viscosity. Their

glass transition temperatures are in the region of

170 8C and degradation does not start below 300 8C.

Melt rheology measurements revealed a predomi-

nantly elastic behavior with a relatively high viscosity

at low frequency. A lower melt viscosity was

achieved by end-group modification. The hyper-

branched poly(ether–amide)s serve as effective

viscosity modifiers for polyamide-6 matrix. New

monomers for formulating thermosetting composites,

based on tetraphenyl-substituted bisoxazoline mono-

mers were synthesized via the direct reaction of

2-(diphenylmethyl)oxazoline with bromoalkyls, using

tert-butyllithium [59]. These bisoxazolines have

different melting points with varied molecular chain

flexibility. They functioned well as cross linkers when

heated with phenolic resins or poly(acrylic acid),

providing a path to new thermosetting materials

with controlled glass transition temperature. Six new

ether-linked bisoxazolines were synthesized via reac-

tion of p-hydroxyphenyl-2-oxazoline with dihalides

[60]. These bisoxazolines were used as chain

extenders or cross linkers for resins, monomers or

polymers containing various acidic groups, including

phenolics, via step-growth reactions. Thus, a novolac

and a bisphenol-A oligomer resin, as well as

poly( p-hydroxy styrene) were chain extended and

cross linked to produce thermosets with high glass

transition temperatures. The new bisoxazolines were

also polymerized with diphenols to generate linear or

branched oligomers and polymers. These new, ether-

linked bisoxazolines are claimed to be potential high

performance thermosets.

4. Polybenzoxazines (PBZ)

Another interesting addition-cure phenolic system

is based on oxazine-modified phenolic resin that

undergoes a ring-opening polymerization to give

polybenzoxazine, which is effectively a poly(amino-

phenol). The precursors are formed from phenol and

formaldehyde in the presence of amines. The choice

for phenol and amine permits design flexibility and

polymer property tailoring. The as-synthesized mix-

ture consists of monomer, and oligomers that contain

phenolic groups. For practical applications, the

mixture is good enough, but for controlled structure

and properties, the monomer is freed of the oligomers.

The ring-opening polymerization can be catalyzed by

acidic catalysts that permits a wide cure temperature.

In the presence of acidic catalysts (e.g. phenols), the

cure temperature window can be reduced from

160–220 8C to about 130–170 8C (i.e. a decrease of

30–50 8C). The synthesis and polymerization of BZ

are depicted in Scheme 9.

4.1. Features of polybenzoxazines

These new materials, belonging to the addition-

cure phenolics family were developed to combine the

thermal properties and flame retardance of phenolics

and the mechanical performance and molecular

design flexibility of advanced epoxy systems [61].

The polybenzoxazines overcome several short-

comings of conventional novolac and resole-type


phenolic resins, while retaining their benefits. PBZ

resins are expected to replace traditional phenolics,

polyesters, vinyl esters, epoxies, BMI, cyanate esters

and polyimides in many respects. The molecular

structure of PBZ offers superb design flexibility that

allows properties of the cured material to be

controlled for specific requirements of a wide variety

of individual requirements. The physical and mech-

anical properties of these new polybenzoxazines are

shown to compare very favorably with those of

conventional phenolic and epoxy resins. The resin

permits development of new applications by utilizing

some of their unique features such as [62,63]:

† Near zero volumetric change upon polymerization

† Low water absorption

† Tg much higher than cure temperature

† Fast mechanical property build-up as a function of

degree of polymerization

† High char-yield

† Low CTE

† Low viscosity

† Excellent electrical properties

Table 13 compares the properties of PBZ with

those of the state-of-the-art matrices. The relative

advantages of PBZ are obvious. They present a resin

system with the highest tensile properties, and Tg can

be boosted to as much as 340 8C through proper

choice of the precursor phenol. This new family of

phenolic resin features a wide range of mechanical

and physical properties that can be tailored to various

needs. DMA reveals that these candidate resins for

composite applications possess high moduli and glass

transition temperatures, at low cross-link densities.

Long-term immersion studies indicate that these

materials have a low rate of water absorption and

low saturation content. Impact, tensile, and flexural

properties are also good. Results of the dielectric

analysis on these polybenzoxazines demonstrate their

suitability for electrical applications [64]. BZs are

cured usually in the temperature window of

160–220 8C. The polymers exhibit Tg in the range

Scheme 9. General protocol for the synthesis and polymerisation of benzoxazine based on bisphenol A.

Table 13

Comparative properties of various high performance polymers

Property Epoxy Phenolics Toughened BMI Bisox–phen (40:60) Cyanate ester P–T resin PBZ

Density (g/cc) 1.2–1.25 1.24–1.32 1.2–1.3 1.3 1.1–1.35 1.25 1.19

Max use temperature (8C) 180 200 ,200 250 150–200 300–350 130–280

Tensile strength (MPa) 90–120 24–45 50–90 91 70–130 42 100–125

Tensile modulus (GPa) 3.1–3.8 03/05 3.5–4.5 4.6–5.1 3.1–3.4 4.1 3.8–4.5

Elongation (%) 3–4.3 0.3 3 1.8 02/04 2 2.3–2.9

Dielectric constant (1 MHz) 3.8–4.5 04/10 3.4–3.7 – 2.7–3.0 3.1 3–3.5

Cure temperature (8C) RT–180 150–190 220–300 175–225 180–250 177–316 160–220

Cure shrinkage (%) .3 0.002 0.007 ,1 ,3 ,3 ,0

TGA onset (8C) 260–340 300–360 360–400 370–390 400–420 410–450 380–400

Tg (8C) 150–220 170 230–380 160–295 250–270 300–400 170–340

GIC (J/m2) 54–100 – 160–250 157–223 – – 168

KIC (MPa m1/2) 0.6 – 0.85 – – – 0/94


160–340 8C depending on the structure, and have

higher thermal stability. The high TGA decompo-

sition onset temperature (for dihydroxy benzophe-

none–aniline system, it is ,400 8C) is attributed to

the very strong intramolecular H-bonding between

phenolic OH and the Mannich bridge. Char-yield as

high as 82% has been claimed. Their composites are

comparable to polyimides and other high performance

polymers, but are easily processable [65].

The ring-opening polymerization of these new

materials occurs with either near-zero shrinkage or

even with a slight expansion upon cure. It is proposed

that the volumetric expansion of the BZ resin is

mostly due to the consequence of molecular packing

influenced by inter- and intramolecular hydrogen

bonding. The role of hydrogen bonding on the

volumetric expansion has been studied by system-

atically changing the primary amine used in the BZ

monomer synthesis. In comparison to the other known

expanding monomers and spiro ortho compounds, this

resin has been shown to have a high potential for

structural/engineering applications [66]. Polybenzox-

azines have the lowest heat release during combustion

and is therefore, more flame resistant, surpassing that

of phenolics and polyetherimides, the current aero-

space matrices of choice. Table 14 gives the heat out

put for three low-heat-release, aerospace grade

polymers. PBZ possess the lowest heat release

characteristics.

4.2. Cure mechanism and cure kinetics

Polymerization of BZ with a free ortho position

can occur through a ring-opening reaction. No volatile

by-products are evolved and no strong catalyst is

required for the reaction [67]. The reaction site of the

ring-opening polymerization in monofunctional aro-

matic amine-based BZs has been investigated through

a systematic manipulation of the monomer chemistry

[68]. Thus, selective protection or activation of sites

on the arylamine ring towards electrophilic aromatic

substitution has allowed a series of materials to be

developed, which contain varying amounts of phe-

nolic Mannich base bridges, arylamine Mannich base

bridges, and methylene bridges. Electron-donating

alkyl substituent groups at one or both the meta posi-

tions on the arylamine ring facilitate ring-opening/

degradation at lower temperatures. This opening of

rings in a step other than the polymerization reaction

greatly increases the numbers of methylene linkages.

Confirmation of the reaction sites was obtained via 1H

and 13C NMR spectroscopy of the oligomeric species.

The reaction pathways on the curing reaction of 3-aryl

substituted benzoxazine was investigated for the

model reaction of 3,4-dihydro-6,8-dimethyl-3-phe-

nyl-2H-1,3-benzoxazine with 2,4-xylenol [69]. The

reaction was carried out at 140 8C for 5 h. N,N-Bis

(2-hydroxy-3,5-dimethylbenzyl) phenylamine and

N-(2-hydroxy-3,5-dimethylbenzyl) phenylamine

were isolated as the main products at an early stage.

The reaction of the first compound by the thermal

treatment was studied by proton-NMR spectroscopy.

This result indicates that this compound produces

several inter- and intramolecular rearrangement

products. Based on these data, some possible reaction

pathways were proposed. In the presence of p-toluene

sulfonic acid monohydrate, 3,30[4,40-methylene-

diphenyl]bis(3,4-dihydro-6,8-dimethyl-2H-1,3-benz-

oxazine) was isolated as one of the intermediates.

The polymerization generally manifests autocata-

lysis. The ring-opening polymerization of BZ was

monitored by rheological analysis utilizing both

conventional isochronic and a new multi frequency

approach that can observe the critical gel [70]. The

activation energies were calculated for the process.

The methylamine-based Bz has higher activation

energy for the gelation process than the aniline-based

one. The cure kinetics of BZ precursor has been

analyzed by torsional braid analysis [71], and gelation

and glass transition during curing were studied. Two

specific glass transitions were obtained at the curing

temperature of 220 8C. The time–temperature-trans-

formation spectra were also generated. BZ is also

cationically polymerizable [72]. Electrochemical

polymerization of PBZ in acetonitrile/alkaline

aqueous solution has also been reported [73].

Table 14

Comparative heat release for various polymer systems [66]

Polymer Peak heat

release rate

(W/g)

Total heat

release

(kJ/g)

Polyetherimide 200 10

Phenolic resin 100 7.5

PBZ 30 2.5


The resultant polymer showed good heat-resistance

properties.

4.3. Structure–property relations

The physical, mechanical and thermal properties of

polybenzoxazines are primarily decided by the nature

of the diphenol and the amine. The system derived

from 4,40-dihydroxy benzophenone and aniline pro-

vides a combination of high thermal stability (5%

weight loss at .400 8C) and high Tg (340 8C).

Table 15 typically illustrates the structure–property

relation for three different BZs [61]. It is seen that the

thermal capability for dihydroxy benzophenone-based

system is achieved at the cost of mechanical

performance. The impact property is good for PBZ

derived from aliphatic amine.

Similar to BZs, naphthoxazines were obtained

from hydroxy naphthalene with aniline and formal-

dehyde [74]. The polynaphthoxazine also showed a Tg

higher than the cure temperature. Thermal properties

in terms of the weight loss after isothermal ageing in

static air, the decomposition temperature from

thermogravimetric analysis, and the change of

dynamic storage moduli at high temperatures also

confirmed their superior thermal characteristics. The

dependence of thermal stability and mechanical

properties on the nature of the amine were also

examined. Thus, the thermal and mechanical proper-

ties of polybenzoxazine thermoset networks

containing varying amounts of phenolic Mannich

bridges, arylamine Mannich bridges, and methylene

bridges were investigated [75]. In materials based on,

m-toluidine and 3,5-xylidine, the onset of thermal

degradation is delayed until around 350 8C with no

significant effect on the final char-yield. Materials

with additional amounts of arylamine Mannich

bridges and methylene bridges show improved

mechanical properties, including higher cross link

densities and rubbery plateau moduli. Regulation of

the viscosity of a difunctional BZ resin is achieved by

addition of a monofunctional BZ monomer or a

difunctional epoxy monomer as reactive diluents to

further improve processibility [76]. The glassy state

properties, such as stiffness at room temperature, are

unaffected by the incorporation of the monofunctional

BZs. The thermal stability of the monofunctional-

modified polybenzoxazine is not significantly affected

below 200 8C. Properties sensitive to network struc-

ture, however, are affected. The incorporation of the

monofunctional BZ reduces cross link density and

produces a looser network structure, while the

difunctional epoxy increases cross link density and

leads to a more connected network structure.

The relatively low cross-link density of BZ has

been overcome by synthesizing pendant BZ func-

tional vinyl polymer by reacting poly(4-vinylphenol)

(VP) with formalin and aniline [77]. The pendant

phenol polymer was obtained by homopolymeriza-

tion of vinyl phenol (VP) or its copolymerization

with butyl acrylate (VBP). The corresponding BZs

(Va and VBa, respectively) were reacted with

bisphenol A epoxy (DGEBA) and o-cresol novolac

epoxy (CNE). The curing reaction proceeded rapidly

at higher temperatures without a curing accelerator.

The reaction induction time or cure time of the

molten mixture from VP-based BZ and epoxy resin

was found to decrease, in comparison to those of

conventional BPA-based BZ and epoxy resin. The

cured resins from VP-based BZ and epoxy resin

showed higher Tg; mechanical properties, electrical

insulation, and water resistance compared to the one

from VP and epoxy resin (phenol–epoxy reaction)

using imidazole as the catalyst. In the case of CNE,

the net gain in Tg was not substantial. The

improvement in flexural properties was also mar-

ginal. BZ–epoxy co-cure is described in detail in

Section 4.4. The properties of the cured resin are

complied in Table 16. The electrical requirements

Table 15

Structure-property relation for polybenzoxazines [61]

Properties Bis benzoxazine formed from

BPA þ methyl

amine

BPA

þ aniline

Dihydroxy

benzophenone

þ aniline

Tensile strength

(MPa)

103 126 6.2

Tensile modulus

(GPa)

3.8 4.5 .6

Strain at break (%) 2.6 2.9 2.3

Impact strength

(J/m)

31 18 –

Polymer density 1.122 1.195 1.250

Tg (8C) 180 170 340


of new printed wire boards are beyond the

capabilities of epoxies. The dielectric constant

ðDkÞ (3–3.5) and dissipation factor (0.006–0.012)

of selected and structurally modified PBZ systems

are inferior only to cyanate esters and they find

application in printed wire boards [78].

Whereas selection of suitable precursors can lead

to high char-yielding PBZ, this can also be achieved

Table 16

Properties of pendant phenol- and benzoxazine-functional linear polymers cured with different epoxies [77] (Reprinted from Journal of Applied

Polymer Science by permission, q2001 John Wiley and Sons)

Resin system* Tg (8C) KIC (MPa m1/2) Water

absorption (%)

Flexural

strength (MPa)

Flexural

modulus (GPa)

Va/DGEBA 174 0.63 0.32 – –

VBa/DGEBA 143 0.6 0.54 – –

VP/DGEBA 165 0.64 0.5 – –

VBP/DGEBA 118 0.6 1.39 – –

Va/CNE 209 – 0.11 77 7.8

VBa/CNE 177 – 0.16 70 6.9

VP/CNE 209 – 0.16 71 6.6

VBP/CNE 160 – 0.28 69 5.1

Table 17

Thermal properties of various acetylene functional PBZ [80] (Reprinted from 43rd Int SAMPE Symp by permission, q1988 Society for the

Advancement of Materials and Process Engineering)

Monomer structure T5 (8C) Char yield % (800 8C)

462

470

489

415

478

458

494

440

73

78

79

75

80

74

71

78

390 32

380 76

488 81


by the incorporation of additional curing sites such as

acetylene on the BZ backbone. The synthesis is

achieved using amino phenyl acetylene in place of

amine. They are polymerized in the range 190–

200 8C. DMA analyses showed a Tg of the order of

329–368 8C, much higher than those of the non-

acetylene BZs [79]. The high thermal stability of this

class of polybenzoxazines is a combined result of the

independent polymerization of both the terminal

acetylene groups and the BZ rings. The thermal

capability is a combined effect of the acetylene and

BZ as these properties are dependant also on the

backbone structure. The structure of various acetylene

functional polymer and their thermal properties are

compiled in Table 17 [80].

The high char-yield achieved for this class of

materials is in the range 71–81% by weight at 800 8C

in a nitrogen atmosphere, and 30% by weight at

700 8C in air, as determined by TGA. The increase in

char-yield is very significant. Temperature at 5%

weight loss ðT5Þ increases from 390 8C to as high as

490 8C for the acetylene-functional polymers. Tem-

perature at 10% weight loss ðT10Þ is in the range 520–

600 8C. Blending BZs with the acetylene-functional

BZ resulted in improved char-yield of the former [79,

81]. Mechanistic investigations revealed that the

acetylene group polymerizes to form short chains of

polyenes, almost simultaneously with the ring-open-

ing polymerization of BZ [81]. In fact the thermal

stability depends on the atmosphere of curing. Thus,

polymerization in air conferred better thermal stab-

ility and char-yield than the polymerization done in an

inert atmosphere. Both the ring-opening polymeriz-

ation and the acetylene polyene formation are faster in

air. The syntheses of two typical acetylene-containing

polymers (i) with terminal acetylene and (ii) with

phenyl acetylene groups are shown in Scheme 10 [81].

The phenyl ethynyl benzoxazine (structure ii)

Scheme 10. Synthesis protocol for benzoxazines containing (i) terminal acetylene (ii) phenyl ethynyl groups [81] (Reprinted from Polymer by

permission, q1999 Elsevier Science).


required very high cure temperature (i.e. 350 8C) for

cure completion, typical of phenyl ethynyl group

polymerization.

Another strategy to improve Tg and thermal

stability is by fluorination of BZ. Thus, a fluorinated

polybenzoxazine was synthesized by the ring-opening

polymerization of hexafluoroisopropylidene-contain-

ing BZ monomer. Substantial development of Tg

occurred at low degrees of conversion. The thermal

stability also improved upon fluorination [82].

Fluorinated PBZ with fluorine groups on the amine

(viz. 3,4-dihydro-3-pentafluorophenyl-2H-1,3-ben-

zoxazine) was obtained in high yield from penta-

fluoroaniline [83]. This monomer, synthesized by a

non-conventional route is a potential precursor for a

polybenzoxazine in electronic applications in view of

its low ðDkÞ, low flammability, low refractive index,

low coefficient of friction and high glass transition

temperature.

PBZ with phthalonitrile groups manifested good

thermal characteristics [84]. The phthalonitrile-func-

tional polybenzoxazines (Pth-PBZ, structure in

Scheme 11) showed good char retention up to 80%

at 800 8C. In air, the char-yield is up to 70% at 600 8C

with T5 in the range 380–420 8C. Tg is in the range

275–300 8C. Unlike conventional phthalonitrile poly-

mers, these systems require a relatively lower cure

temperature, of the order of 250 8C. The unreacted

nitrile groups react further during degradation

accounting for the high char residue. The polymers

are classified as flame resistant. The superior thermal

stability is evident from the comparative thermal

properties of different phthalonitrile-based polymer

given in Table 18. The gain in char-yield in air is

substantial.

Dendrictic macromolecules containing BZ moi-

eties have also been synthesized [85]. The dendritic

building blocks were made by coupling reactions via a

convergent approach. The BZs were synthesized by

various combinations of amines and phenols permit-

ting design flexibility by using derivatives of phenol,

primary amine and aldehyde. Dendrimers containing

several different BZ moieties were synthesized up to

generation 2. As the generation number increased, the

curing temperature decreased gradually.

4.4. Reactive blending of polybenzoxazines

Despite high modulus and Tg; polybenzoxazines

have surprisingly low cross-link density in compari-

son to other thermosets. Hydrogen-bonding is ade-

quate to induce rigidity and constrain the mobility in

the glassy state. The reactive blending with epoxy

resins allows the network to achieve higher cross-

linking. The co-reaction proceeds via the ring-open-

ing polymerization of BZ followed, by reaction of the

generated phenol with the epoxy, probably catalyzed

by the amine group. The proposed reaction sequences

can be found in Scheme 12.

Epoxy ring-opening by the generated amine is

totally not ruled out. Copolymerization leads to a

significant increase in the glass transition temperature

[86]. Reaction of DGEBA with BPA-based benzox-

azine (B-a) increased the Tg of the latter from about

143 to about 153 8C for an epoxy-content of 35%.

Further increase in epoxy decreases Tg to below that

of pure PBZ. This behavior is shown in Fig. 10. The

copolymer exhibits strain at breakage that is up toScheme 11. Structure of phthalonitrile-functional benzoxazine [84]

(Reprinted from Polymer by permission, q2000 Elsevier Science).

Table 18

Comparative thermal performance of various phthalonitrile poly-

mers [84]

Polymer N2 Air

T10 Char-yield

(%) 800 8C

T10 Char-yield

(%) 600 8C

Pth-PBZ 596 80 450 75

Pth-polyimidea 550 65 510 10

Pth-diamineb 500 70 500 15

a Phthalonitrile–polyimide.b Phthalonitrile cured by diamine (2%).


twice the strain for pure BZ and a flexural strength

significantly improved with minimal sacrifice in

modulus. The flexural strength increased from 125

to 170 MPa on enhancing the epoxy-content to 50%.

Interestingly, the increase is almost linear to the

epoxy-content as shown in Fig. 11.

The modulus decreased from 4.4 to 3.4 GPa,

translating to an increase in strain at break from 3 to

6.3%. A corresponding decrease occurred in the

storage modulus determined by DMA. By under-

standing the structural changes induced by variations

of epoxy-content and their effect on material proper-

ties, network can be tailored to specific performance

requirements. Thermally stable systems resulted on

using BPA based BZ as a hardener of the epoxy resin

[87]. The curing reaction proceeded without any

accelerator. The molding compound showed good

thermal stability under 150 8C, which corresponded to

the injection molding temperature. Above 150 8C, the

curing reaction proceeded rapidly. The cured epoxy

resin showed good heat resistance, water resistance,

electrical insulation, and mechanical properties com-

pared with the epoxy resin cured by the bisphenol-A

type novolac. PBZ–EPOXY system has been pro-

jected as a high performance matrix for several

applications [88]. These systems are processable by

traditional techniques such as prepreg, RTM and

VRTM etc. for composite fabrication. The system, not

requiring any refrigeration, is stable at ambient

conditions indefinitely. Suitably formulated, their Tg

Scheme 12. General expected reaction sequence for curing of benzoxazine and epoxy.

Fig. 10. Dependence of Tg on composition for diglycidyl ether of

bisphenol A (DGEBA)/ {bis(4-phenyl-3,4-dihydro-2H-1,3-benzox-

azinyl) isopropane} (B-a) blend [86] (Reprinted from Polymer with


Fig. 11. Relative increase in flexural strength with epoxy-content for

polybenzoxazine-epoxy blend [86].


can be boosted to ,200 8C, and they possess

excellent mechanical properties and low moisture

absorption. The mechanical properties of these

systems in comparison to pure BZ and a standard

epoxy can be found in Table 19. The blend shows

improved mechanical performance at ambient. The

trend is maintained in the wet and high temperature

conditions. However, in this case, the fracture

properties are adversely affected although Tg was

enhanced.

BZ–epoxy blend with superior heat- and water

resistance and electrical insulation resulted when a

terpene diphenol-based benzoxazine (Terp-Bz) was

cured by an epoxy (o-cresol novolac epoxy) quanti-

tatively above 180 8C [89]. The properties were

superior to bisphenol-A-based benzoxazine (B-a) or

bisphenol-novolac (BisA-N) system cured in the

presence of the same epoxy. The gain in properties,

is however only marginal (about 10% increase in

flexural strength). A ternary mixture of BZ, epoxy,

and phenolic novolac resins provided low melt

viscosity resins and void-free specimens with minimal

processing step [90,91]. The material properties were

highly dependent on the composition of the mono-

mers in the mixture. A Tg of 170 8C and considerable

thermal stability can be obtained from these systems.

Phenolic novolac resin acts mainly as an initiator for

these ternary systems, while low melt viscosity,

flexibility and improved cross link density of the

materials were attributed to the epoxy fraction.

Polybenzoxazine imparts thermal curability, mechan-

ical properties as well as low water uptake to the

ternary systems. The authors claim the materials as

promising candidates for application as underfilling

encapsulation and other highly filled systems. The

gelation behavior of this low melt viscosity ternary

blend was investigated by Fourier Transform Mech-

anical Spectroscopy technique (FTMS) in order to

study the effect of epoxy diluent on the rheological

property development before and after the gel points

[92]. The gel times range from 5 to 30 min at 140 8C

to less than 5 min at 180 8C for all tested ternary

system compositions. The gelation of the ternary

mixture showed Arrhenius-type behavior, permitting

prediction of the gel time.

BZ can be used for functionalization of other

polymeric systems as a means of cross linking. Thus,

polysiloxanes are fictionalized with these groups and

their adhesion properties in glass fiber-reinforced

composites are improved [93]. BPA-based BZ can be

cured by bisoxazoline in the presence of triphenylpho-

sphine as a catalyst at 170 8C [94]. The blend is

processable by melt impregnation, RTM etc. The

phenolic hydroxyl groups generated by the ring-

opening reaction of the BZ ring react with the

oxazoline ring at 200 8C as shown in Scheme 13.

The cure can be completed in less than 30 min at

this temperature as evident from DMA. The system

flows easily at 100–140 8C, and has good thermal

stability. The melt viscosity of the molding compound

Table 19

Mechanical properties of neat castings of PBZ, PBZ–epoxy and a control epoxy [88] (Reprinted from 46th Int SAMPE Symp by permission,

q2001 Society for the Advancement of Materials and Process Engineering)

Properties PBZa PBZZ–epoxy Epoxyb

RT Wet, RT 150 8C RT Wet, RT 150 8C RT Wet, RT 150 8C

Tensile strength (MPa) 31 31 21 52 45 52 52 30 37

Tensile modulus (MPa) 5334 5016 1290 4319 4513 2788 3995 3491 2401

Elongation (%) 1.2 1.5 3 1.6 1.1 2.4 1.4 1 2.2

Flexural strength (MPa) 132 103 44 115 84 89 142 61 77

Flexural modulus (MPa) 4602 4844 1538 4430 4610 3015 3712 3457 2518

Compress strength (MPa) 228 – – 232 – – 279 – –

Compress modulus (MPa) 3505 – – 3243 – – 1939 – –

KIC (MPa m1/2) 0.94 0.65 – 0.55

GIC (J/m2) 168 – – 83 – – 71 – –

Tg (8C) 171 – – 219 – 223 – –

a Undisclosed structure.b TGMDA þ DDS.


is around 0.1–1.0 Pa s at 140 8C, even after 1.5 h, but

increases rapidly at 180 8C. The cured resin showed

good heat resistance (Tg ,195 8C), water resistance

(0.28% absorption), electrical insulation (volume

resistivity of 3.5 £ 1016 ohm cm), and mechanical

properties (KIC ,0.93 MPa m1/2), that are only

marginally superior to cured resin from BPA-type

novolac and bisoxazoline. On the other hand, when

BPA-based benzoxazine (B-a) containing oligomers

(oligo-B-a) was cured with bisoxazoline, the cure time

and temperature could be lowered, compared with

those from pure B-a and bisoxazoline [95]. Above

160 8C, the curing reaction of oligo-B-a with

bisoxazoline (1,3-PBOX) proceeded more rapidly

than that of B-a with bisoxazoline. The cured resin

from oligo-B-a and bisoxazoline showed better heat

resistance (Tg- 232 8C) and water resistance (0.23%

absorption), than the cured resin from B-a and

bisoxazoline. However, the fracture toughness

showed a reverse trend (KIC-0.75 against

0.93 MPa m1/2 for the latter).

In another study, two poly functional BZ monomers,

viz. 8,80-bis (3,4-dihydro-3-phenyl-2H-1,3-ben-

zoxazine) and 6,60-bis(2,3-dihydro-3-phenyl-4H-1,

3-benzoxazinyl) ketone, were co-cured in an auto-

clave [96]. These two polybenzoxazines showed

mechanical and thermal properties similar to or better

than BMIs and some polyimides. They also showed

very high char-yield after carbonization in a nitrogen

atmosphere. Thermally stable blends result when

phenyl nitrile functional BZs are blended with nitrile

systems and cured [97]. Ortho-, meta-, and para-

phenylnitrile-functional BZs were polymerized at

different compositions with phthalonitrile-functional

monomers, providing copolybenzoxazines of high

thermal stability and easy processibility. The copoly-

mer char-yield increased from 59 to 77 wt% and Tg

from 180 to 294 8C with only 30 mol% of phthaloni-

trile-functional monomer.

4.5. Non-reactive blends and composites

of benzoxazine

PBZ is amenable for matrix modification by fillers,

reinforcements and polymer blends. The thermal

properties of physical blends containing BZ monomer

and polycarbonate (PC) were studied by non-isother-

mal DSC, TGA and FTIR [98]. The ring-opening

reaction and subsequent polymerization reaction of the

BZ were inhibited significantly by the presence of

polycarbonate. The glass-transition temperature of the

resulting blends decreased as the concentration of

polycarbonate increased and deviated markedly from

the Fox equation. An earlier degradation event

appeared in the blend with 11 and 33 wt% of PC.

Intermolecular hydrogen bonding between PC and

cured polybenzoxazine appeared after 1 h of isother-

mal curing at 180 8C, and continued throughout the

entire curing process. Subsequent studies confirmed a

possible cross-reaction between PC and the ring-

opened BZ that led to PC-grafted and PC-cross linked

PBZ [99]. The chain fragmentation in PC and

polybenzoxazine blend upon thermal polymerization

was investigated by size exclusion chromatography.

Molecular weight reduction of PC via trans esterifica-

tion between the hydroxyl groups of ring-opened BZ

and the carbonate groups from PC was observed. In

addition, excess heat of reaction compared to the

expected value was detected from DSC and was

assigned to the exotherm associated with the exchange

Scheme 13. Co-reaction between benzoxazine and bisoxazoline [95] (Reprinted from Journal of Applied Polymer Science by permission,

q2001 John Wiley and Sons).


reaction. The proposed PBZ-PC interaction leading to

grafts and cross links is shown in Scheme 14.

Improved systems are reported to result due to

synergism on melt blending BZ resins with poly(1-

caprolactone) (PCL) [100]. The PCL-content was

varied in the range 0–15 wt%. The Tgs of the BZ

blends were found to be slightly lower than that of

neat polybenzoxazine resin. The blends showed

improved mechanical properties, including higher

cross link densities, rubbery plateau moduli, and

flexural strengths compared to pure polybenzoxazine.

Unlike the previous case of PC-modification, the

thermal stability at the mid-temperature range was

enhanced, evident from the delayed onset of

decomposition temperature and the disappearance of

the first degradation event. However, at higher PCL-

loading, only one Tg resulted, and the Tg value of the

resulting blend appeared to be higher in the blend with

a greater amount of PCL [101]. Phase separation

occurred when a BZ monomer (B-a) was blended with

soluble poly(imide–siloxane)s [102]. The soluble

poly(imide–siloxane)s with and without pendent

phenolic groups (structures PISi–OH and PISi,

respectively, in Scheme 15) were prepared from the

reaction of 2,20-bis(3,4-dicarboxylphenyl)hexafluoro-

propane dianhydride with a,v-bis(aminopropyl)

dimethylsiloxane oligomer (PDMS; molecular

Scheme 14. Proposed reaction between polycarbonate and poly-

benzoxazine [99] (Reprinted from Journal of Applied Polymer

Science by permission, q2002 John Wiley and Sons).

Scheme 15. Poly (imide–siloxane) additives for polybenzoxazine matrix modification. Poly(imide–siloxane)s with pendent phenolic groups

(PISi–OH) and without pendent phenolic groups(PISi) [102], (Reprinted from Journal of Polymer Science Polymer Chemistry Edition by

permission, q2001 John Wiley and Sons).


weight ¼ 5000) and 3,30-dihydroxybenzidine or 4,40-

diaminodiphenyl ether.

In the presence of poly(imide–siloxane)s, the cure

onset shifted from 200–240 to 130–140 8C. Viscoe-

lastic measurements of the cured blends containing

poly(imide–siloxane) with OH functionality (i.e.

PISi–OH) showed two glass-transition temperatures,

one at 255 8C and another at around 250–300 8C,

indicating phase separation between PDMS and the

combined phase consisting of polyimide and poly-

benzoxazine components, due to the formation of AB-

cross linked polymer. For the blends containing PISi,

however, in addition to the Tg due to PDMS, two Tgs

were observed in high-temperature ranges, 230–260

and 300–350 8C, indicating further phase separation

between the polyimide and PBa components due to the

formation of semi-interpenetrating networks. In both

cases, Tg increased with increasing poly(imide–

siloxane)-content. Thus, the Tg of PBa increased

from 160 to about 300 8C on enhancing the polyimide

siloxane-content to 20% in the case of PiSi–OH. For

PISi, the Tg increase was to about 2258. As the

siloxane-content increased, tensile strength and mod-

ulus decreased implying an improved toughness

caused by the addition of poly(imide–siloxane).

The thermal stability of PBa also was enhanced by

the addition of poly(imide – siloxane)[102].

Urethane–benzoxazine copolymer films were pre-

pared by blending a monofunctional BZ monomer,

viz.3-phenyl-3,4-dihydro-2H-1,3-benzoxazine (Pa),

and polyurethane (PU) prepolymer that was syn-

thesized from 2,4-tolylene diisocyanate and poly-

ethylene adipate polyol (Mw , 1000) in 2:1 molar

ratio, followed by casting as films and thermal

curing at 190 8C [103]. Their synthesis is depicted in

Scheme 16.

The poly(urethane–benzoxazine) films exhibited

only one glass transition temperature, indicating

good miscibility between PU and polybenzoxazine

(Pba) due to the in-situ copolymerization of B-a

and PU prepolymers. The Tg of the polyurethane

was enhanced with increased Pba-content. Films

containing less than 50% of Pa had the character-

istics of an elastomer, with an elongation at break

of more than 200%. These elastic films exhibited

good resilience with excellent reinstating behavior.

Films containing more than 60% of Pa had the

characteristics of a plastic. Up to 15% Pba, the

films were elastomeric with tensile strength of

the order of 2–6 MPa. The value shot up to the

range 280 MPa when the Pa-content exceeded 15%.

The variations in tensile strength and elongation for

different copolymers illustrating their evolution

from elastomer to plastic nature, are shown in

Scheme 16. Synthesis of poly(urethane–benzoxazine)[103] (Reprinted from Journal of Polymer Science Polymer Chemistry Edition by


Fig. 12. Variation of mechanical properties with polybenzoxazine-

content for benzoxazine-modified PU film [103].


Fig. 12. Poly(urethane–benzoxazine) films showed

excellent resistance to solvents such as THF, DMF,

NMP and DMSO [104].

Amine-terminated butadiene–acrylonitrile copo-

lymer (ATBN) and CTBN were introduced to

polybenzoxazine by modification of the monomer

prior to curing. On a comparative scale, ATBN is

more effective than CTBN in improving the fracture

toughness of PBZ [105]. This was attributed to the

better distribution of rubber particles in an ATBN-

modified matrix than for the CTBN-modified one.

DMA confirmed the existence of two networks in the

ATBN-modified matrix [106]. The BZ resin permits

compounding with a variety of additives, including

nanoparticles such as montmorillonite, to make

nanocomposites [107,108]. Thus, polybenzoxazine–

clay hybrid composites were prepared from a BZ

precursor based on BPA and aniline or bis

(3-phenyl-3,4-dihydro-2H-1,3-benzoxazinyl) iso-

propane, (B-a, structure in Scheme 9) and surface-

treated montmorillonite. The dispersity of

organically modified montmorillonite (OMMT) at

the molecular level in the polybenzoxazine matrix

was confirmed by XRD. The Tg of the hybrid

material was enhanced vis-a-vis the base polymer.

The reinforcement effect of OMMT was reflected in

the enhanced high temperature storage modulus of

the composite. The additive also delayed the thermal

decomposition of the matrix in the high temperature

regime. The isothermal degradation studies indicated

good thermal stabilization. This was attributed to the

formation of OMMT nano-barriers minimizing the

permeability of volatile degradation products in

the material. Several types of polybenzoxazine/clay

hybrid nanocomposites have been prepared from

OMMT and mono- or bifunctional benzoxazine

{3-phenyl-3,4-dihydro-2H-1,3-benzoxazine (Pa) or

bis(3-phenyl-3,4-dihydro-2H-1,3-benzoxazinyl) iso-

propane (B-a), respectively} [109]. OMMT was

prepared by a cation exchange of MMT with

ammonium salts of amines such as tyramine,

phenylethylamine, aminolauric acid, and dodecyl

amine, with the structures given in Scheme 17.

Strong hydrogen-bonding existed between PBZ

and poly(N-vinyl-2-pyrrolidone) (PVP) on blending

the two polymers, via interaction between OH group

of phenol and the carbonyl groups of PVP. This

caused a positive deviation in the Tg-versus compo-

sition curve for the two systems. A Tg maximum was

observed at around 50/50 composition of the two

Scheme 17. Structures of benzoxazine monomers and organic modifiers for montmorillonite [109] (Reprinted from Polymer by permission,

q2002 Elsevier Science).


components as seen in Fig. 13. The Tg-composition

curve conformed to the Kiwi equation, accounting for

specific polymer–polymer interaction. These inter-

actions were corroborated by FTIR investigations

[110]. Polybenzoxazine/clay nanocomposites were

prepared by both melt method and solvent methods.

XRD measurements of the nano composites showed

that the blending method and the kind of solvent play

crucial roles in the dispersion of OMMT in the

polybenzoxazine matrix. The inclusion of any type of

OMMT significantly lowered the curing exotherm of

BZs. The hybrid nanocomposites exhibited higher Tg;

higher modulus and enhanced thermal stability. The

effect of clay was more significant at around 2–5 wt%

loading. The blending also reduced the exothermicity

of curing significantly. An immiscible polymer–clay

nanocomposite has been synthesized, consisting of

dispersed layers of OMMT in a polybenzoxazine

matrix, as shown by thermogravimetry, wide-angle

X-ray diffraction (WAXD), and TEM [111]. Proto-

nated amines showed high ion-exchangeability. The

spacing of the silicate layers was strongly dependent

on the size or molecular weight of the amine

derivative and the solvent type. Binary solvents (5%

methanol in toluene) exhibited superior ability to

swell OMMT. WAXD results revealed that the sili-

cate layer expansion of all polybenzoxazine-OMMT

nanocomposites prepared from either melt or solution

methods were similar, with about a 48 A increment.

This was correlated with TEM results that showed the

aggregation of silicate layers, indicating that all

prepared OMMT are immiscible with the polymer

matrix. The compatibility between amine modifying

agents and BZ dictated the characteristics of the

nanocomposites. The curing reaction in the synthesis

of polybenzoxazine–montmorillonite (MMT) nano-

composites shows autocatalytic characteristics for

fairly good conversion range [112]. Inclusion of boron

nitride, on the other hand, results in highly conducting

polybenzoxazine composites [113]. The reduced heat

of reaction and increased exothermic peak tempera-

ture indicated that the boron nitride surface is

inhibiting the BZ polymerization reaction [114].

Incorporation of Kenaf fiber in a polybenzoxazine

resin matrix to form a unidirectional reinforced

composite resulted in systems with enhanced flexural

property [115]. Compounding by calcium carbonate

leads to good particulate-filled PBZ composite [116].

4.6. Degradation of polybenzoxazine

4.6.1. Thermal stabilization and degradation

Polybenzoxazines are thermo-oxidatively stable.

This is exemplified by the superposition of the

thermograms of the cured resin in air and N2 [61]. It

has been shown that the thermal stability of poly-

benzoxazines is substantially improved further by

reactive amines. Various diphenols are found to

have some effect on the thermal stability of this

series of polybenzoxazines. Nitrogen-containing

phenolic resins are also non-flammable polymers

[117]. A study on their thermal decomposition

revealed that PBZ decompose by loss of amine

fragments [118]. Hence, one strategy for thermal

stabilization is to introduce cross-linking sites on

the amine moieties. Thus, propargylation was

an effective means for thermal stabilization.

Fig. 13. Dependence of Tg on blend composition for polybenzox-

azine/poly(N-vinyl-2-pyrrolidone)system [110] (Reprinted from

Polymer with permission, q2003 Elsevier Science).


The introduction of ethynyl [79–81,119] and nitrile

[84,97] groups in the backbone has been fruitful for

conferring thermal stability. It has been seen that

thermally stable phenol precursor such as naphthols

also ensures thermal stability for the resultant PBZ

[74]. The structural modifications and compounding

conducive to enhanced resistance to thermal degra-

dations of PBZ have been discussed at different

points in the preceding sections. Ishida et al

identified the decomposition product of aromatic

amine based polybenzoxazines through TGA and

GC-MS techniques [120]. Several degradation

products were identified, derived from the degra-

dation of the polymer and the recombination of the

degradation products. Benzene derivatives, amines,

phenolic compounds and Mannich base emerge

directly from the polymer. Benzofuran is derived

from further degradation of phenols. Biphenyl com-

pounds are obtained from recombination of phenyl

radicals after the loss of substituents form benzene,

amine and phenol derivatives. Isoquinoline and

phenathridine derivatives result from Mannich base

by loss of OH groups and dehydrogenation. The overall

degradation pattern is shown in Scheme 18. This team

also investigated the thermal decomposition processes

of a model compound containing a Mannich bridge and

a series of polybenzoxazine model dimer, with more-

or-less similar observations [121]. In this case, the 2,4-

dimethylphenol-based BZ dimers degraded into smal-

ler and highly volatile compounds, leaving no char at

Scheme 18. Degradation pattern for polybenzoxazine [120] (Reprinted from Polymer by permission, q2002 Elsevier Science).


the end of degradation. The p-cresol-based BZ dimers

also degraded into smaller and highly volatile

products. Some of these are able to undergo cross

linking and aromatization processes and form char.

The major decomposition products for modified,

p-cresol-based dimers are amines and ester compounds.

4.6.2. Chemical degradation of PBZ

As the network contains basic amine groups, the

stability of polybenzoxazines in acid medium is

doubtful. The chemical stability of typical polyben-

zoxazines based on bisphenol-A and primary amines

in a carboxylic acid solution has been studied [122]. It

was found that the Mannich base is stable. It is

proposed that the nature of the primary amine is

responsible for the interactions between the car-

boxylic acid and the Mannich-base model dimers. As

a result, the chemical stability of polybenzoxazines

may also be related to the nature of the amines, which,

in turn, influence the strength of the hydrogen-bonded

network structure that develops upon cure. While

aniline-based BZ was stable in acidic medium, that

based on methylamine disintegrated into small

fragments. The strong salt formation between the

more basic Mannich bases from the former disrupts

the hydrogen-bonding network and triggers the

degradation. The rapid degradation of certain BZs in

acid medium was attributed to macroscopic stress

cracking.

4.6.3. UV stability of PBZ

The UV stability of PBZ has also been investigated

[123]. Carbonyl-containing species were formed

when B-a resin was exposed to ultraviolet radiation

(l . 290 nm) in ambient air at room temperature.

The isopropylidene linkage was the reactive site of

cleavage and oxidation, resulting in the formation of a

2,6-disubstituted benzoquinone. Interestingly, the

Mannich bridge was neither cleaved nor oxidized. A

comparative UV exposure study has been performed

on polybenzoxazines containing the same amine

(methylamine), but with phenolic units possessing

various substituents on the para-positions of the

phenyl rings [124]. B-a is shown to have the highest

degree of substituted benzoquinone formation fol-

lowed by those polymers derived from hydroquinone,

4,40-(hexafluoroisopropylidene) diphenol, 4,40-thiodi-

phenol, 4,40-dihydroxybenzophenone, p-cresol

and phenol. The nature of the para-position in

phenolic substituents was found to have an impact

on the oxidation process affecting the degrees of

substituted benzoquinone formation. Some secondary

reactions were also found to occur as a result of photo-

oxidation.

5. Phenol–epoxy systems

Curing of epoxy with novolac type phenolic resin,

making use of the OH–epoxy reaction, appears to be

the simplest way to design addition-cure phenolic

system. Although less preferred, polyphenols are used

as curative for epoxies, since the addition-curing

results in void-free products which are comparatively

tougher due to the formation of flexible ether network

[125–128]. Phenol–epoxy thermosets are preferred

in void-free composite structures. The interest in these

systems has been revived further by the need for void-

free, low moisture absorbing matrices with low

dielectric properties for various electronic appli-

cations. The phenol–epoxy reaction is less facile

than amine-epoxy reaction, and is usually performed

in the presence of catalysts such as triphenyl

phosphine (TPP). The choice of phenolic resin and

the epoxy allows for a wide design flexibility and

property tailoring.

5.1. Epoxy–phenol cure kinetics

Since the cure kinetics control the morphology and

properties of the final polymer, they have been the

focus of research. Various researchers propose

different cure models and various hypotheses to

explain the observations. A few studies are presented

here. Thus, the cure kinetics and Tg of the products

have been investigated for the phenol–epoxy system,

where the phenols bear cyclopentadiene and xylok

moieties [129]. Investigation of the cure kinetics and

relationships between the glass transition temperature

and conversion of biphenyl epoxy resin (4,40-diglyci-

dyloxy-3,30,5,50-tetramethyl biphenyl) with these

phenolic hardeners were performed by differential

scanning calorimeter, using an isothermal approach

over the temperature range 120–150 8C. The results

indicated that the curing reaction of formulations

using xylok and dicyclopentadiene type phenolic


resins (DCPDP) as hardeners proceeds through a first-

order kinetic mechanism {Eq. (1)}, whereas the

curing reaction of formulations using novolac as a

hardener goes through an autocatalytic kinetic

mechanism {Eq. (2)},

da=dt ¼ k1ð1 2 aÞ ð1Þ

da=dt ¼ ðk1 þ k2amÞð1 2 aÞn ð2Þ

Differences of curing reaction with change of the

hardener in biphenyl epoxy resin systems were

explained with the relationship between Tg and the

reaction conversion, using the DiBenedetto equation

[130] as,

Tg ¼ Tgo þ ðTga 2 TgoÞlx=½1 2 ð1 2 lÞx� ð3Þ

where, Tgo and Tga are the glass transition tempera-

tures of the monomer and the fully cured network,

respectively, and l is an adjustable parameter.

In an analogous study, the cure kinetics of a

biphenyl epoxy (4,4-diglycidyloxy-3,30, 5,50-tetra-

methyl biphenyl)-xylok resin system with four

different catalysts was performed by a differential

scanning calorimeter using an isothermal approach

[131]. The curing reaction of the formulations using

TPP and 1-benzyl-2-methylimidazole as catalysts

proceeded through a first order kinetic mechanism,

whereas that of the formulations using diazabicy-

cloundecene and tetraphenyl phosphonium tetraphe-

nyl borate proceeded by an autocatalytic kinetic

mechanism. The effects of concentration of TPP as

catalyst on the curing of biphenyl epoxy and phenol

novolac resin system was also investigated. All

kinetic parameters were deduced from the isothermal

DSC analyses. A kinetic model with a diffusion term

was invoked to describe the reaction for the entire

conversion range [132]. Through the analyses of the

cure kinetics and the relationship between Tg and

conversion, it was concluded that TPP affects only the

kinetics and not the final cross link density.

The modified kinetic model takes the form:

da=dt ¼ ðk01 þ k02amÞð1 2 aÞn½B� ð4Þ

where k01 and k02 are the absolute rate constants for the

normal and auto catalytic reactions, and [B] is the

catalyst concentration. Although this equation

described the kinetics at any given concentration of

TPP, the authors introduced a diffusion term to

describe the kinetics over the entire conversion

range as in Eq. (5),

da=dt ¼ {ðk01 þ k02amÞð1 2 aÞn½B�}

{1=ð1 þ exp½Cða2 acÞ�} ð5Þ

where C is a material constant, ac is the critical

concentration beyond which the diffusion comes in to

effect and is temperature-dependent. The modified,

diffusion-corrected model clearly defined the kinetics

at all conversions and catalyst concentrations. The

cure kinetics of off-stoichiometric biphenyl epoxy

(4,4-diglycidyloxy-3,30,5,50-tetramethyl biphenyl)-

dicyclopentadiene phenolic resin system by DSC

using an isothermal approach was investigated [133].

To describe the curing reaction after the vitrification,

controlled by a diffusion factor, a semiempirical

relationship was used. A one-to-one relationship,

which is independent of the cure temperature between

the Tg and the fractional conversion, was interpreted

using the DiBenedetto equation. A Time–Tempera-

ture-Transformation (TTT) isothermal cure diagram

has been established at each equivalent weight ratio

using the kinetic model coupled with the diffusion

factor and the DiBenedetto equation. A study on

similar lines of the cure kinetics of the above biphenyl

epoxy reacted with a dicyclopentadiene type phenolic

resin was done with four different kinds of catalysts

[134]. The relative performance of the four catalysts

and their kinetics were identical to those in the

previous study [131].

The kinetics of reactions of certain phenol–epoxy

ternary systems has also interested the researchers.

Thus, the kinetics of cure of bisphenol-A diglycidyl

ether, bisphenol-A, and a sulphanilamide was inves-

tigated using HPLC [135]. The influence of tetrabutyl

phosphonium tetrafluoroborate as a catalyst for the

epoxy–phenol reaction was studied. For the same

initial composition, polymers differing by the cross-

link point distribution and the chain length between

cross-links were synthesized. Different solid-state

behavior is expected in these two categories of

networks. The phenol-epoxy reaction kinetics has

been described in detail by Biernath and Soanne [136].

The kinetics of triphenyl phosphine-catalyzed

thermal cure of the reactive blend of DABA and

epoxy novolac (EPN) was investigated using DSC

[137]. TPP catalyzed the phenol-epoxy reaction of


the dual cure system. The kinetic parameters for this

step were estimated by the variable heating rate

method of Kissinger. The activation energy ðEÞ and

pre-exponential factor ðAÞ depended on the catalyst

concentration. Although the variation in activation

energy with increase in catalyst concentration did not

follow a regular trend, the associated change in A

implied a systematic increase in the rate constant. The

apparent activation energies normalized to a fixed A

value ðEnÞ decreased with increasing catalyst concen-

tration. The catalysis effect was evident in the

computed rate constant ðk100Þ at 100 8C. The acti-

vation parameters were used to predict the cure profile

of the resin under given conditions of temperature and

catalyst concentration. Table 20 compiles the cure

parameters and the kinetic constants. Fig. 14 shows

the cure profiles at different catalyst concentration.

At these cure conditions, the allyl group remains

unreacted. This method can be used to synthesize allyl

functional phenoxy resins. A first order kinetics has

also been observed for esterified novolac-epoxy cure.

The resultant cured matrix is more moisture resistant

[138]. The structure–property correlation of phos-

phine-containing catalysts in phenol–epoxy reaction

has been separately investigated [139]. The relation-

ships of melting behavior and chemical structure were

discussed using the heat of formation calculated by a

semiempirical method. Two schemes are proposed for

the reactions with epoxy and phenol compounds: a

complex formation scheme, and an ionic scheme. The

reaction rate of epoxy and phenol compounds

decreased on introducing electron-withdrawing sub-

stituents such as chlorine in the basic catalyst (TPP).

The curing kinetics of epoxy–phenol novolac com-

pounds in the presence of a tertiary amine was studied

by isothermal calorimetry and rotation viscometry

[140]. The curing kinetics and mechanism, structural

organization on the chemical, topological, and

supramolecular levels, relaxation properties, rupture

mechanism, and physico-mechanical properties of the

resulting polymer were studied, as influenced by the

concentration of the tertiary amine. The reaction

kinetics with diffusion controlled mechanism, as well

as the volumetric change upon curing, of a cresol

novolac epoxy/o-cresol–formaldehyde novolac hard-

ener system were studied by Tai [141]. Simple

equations to model the change in linear coefficients

of thermal expansion with conversion were derived.

The true degree of conversion of this cross-linking

epoxy system was obtained on the basis of the heat of

reaction of monomeric monofunctional model com-

pounds. The reaction was then modeled as one of a

shifting order; autocatalytic followed by diffusion

controlled. The reaction in the diffusion-controlled

region was modeled by an nth order kinetic equation,

with its rate constant described by a WLF-type

equation. The experimental linear coefficients of

thermal expansion above and below the glass

transition temperature decreased linearly with the

degree of conversion, in agreement with the derived

equations. Yoon on the other hand, considered an

active complex-formation for the cure kinetic model

of biphenyl-type epoxy/phenol novolac resin system

[142]. Investigation of the cure kinetics of a biphenyl

epoxy–phenol resin system with different kinds of

Table 20

Cure temperatures and kinetic parameters for the TPP-catalyzed

cure of DABA-EPN [137]

TPP

(wt%)

E

(kJ/mol)

A (s21)

£ 1026

En (kJ/mol)

for A ¼ 1.36 £ 108

k100 (s21)

£ 103

0.5 68.3 2.7 80.4 0.75

1 79.7 145 79.5 1.01

2 79.4 306 76.9 2.31

3 74.5 89.8 75.7 3.36

Fig. 14. Time-conversion profiles for epoxy novolac/diallylbi-

sphenol A blend at varying concentration (wt%) of triphenyl

phosphine. Heating rate 10 8C/min, [137].


latent catalysts including encapsulated TPP was

performed by DSC using an isothermal approach

[143]. A combination of DSC and dynamic rheologi-

cal measurements has been employed to study the

chemo-rheological behavior of an epoxy–novolac

molding compound [144]. A procedure aiming at the

phenomenological description of cure kinetics was

developed. An empirical Arrhenius-type expression

was adopted for the description of the dependence of

the complex viscosity on temperature, frequency and

conversion by allowing the pre-exponential factor

and the flow activation energy to depend on frequency

and conversion. At low conversions (,5%), the

system behaved approximately as a thermoplastic

material. At higher conversions, the rheological

behavior of the system was dominated by the extent

of cure reaction. Similarly, the chemo-rheological

study of biphenyl-type epoxy/phenol novolac system

was performed at different isothermal curing tem-

peratures [145]. A modified WLF equation was used

to predict the chemo-viscosity during isothermal

curing reaction. The glass transition temperature

change due to curing reaction measured by DSC

could be expressed as a function of conversion. From

the measurements of isothermal viscosity, C1 and C2

terms in WLF equation were found to have an

Arrhenius-type dependence on temperature. Using

these relations in conjunction with the modified WLF

equation, a good agreement between predicted and

experimental viscosity has been claimed by the

authors. Like TPP, imidazole is another preferred

catalyst for the phenol-epoxy reaction. Thus, the cure

reaction between DGEBA and BPA in the presence of

the imidazole catalyst has been studied [146]. Two

initiation reactions were proposed. One is based on

the ionic complex formation between the epoxy and

the pyridine-type nitrogen of imidazole, and the other

on the ionic complex of BPA and imidazole. The

subsequent propagation steps were composed of three

steps, viz. the epoxy–phenol reaction, the acid–base

reaction and the epoxide and the alkoxide. A

generalized kinetic model was developed, which

satisfactorily defined the kinetics in a wide range of

imidazole and BPA concentrations. The kinetic model

was further validated through prediction of the

structural properties including epoxide conversion,

molecular weight, sol–gel fraction, branch points,

loop density of the gel part by Monte Carlo

simulation, and verification of the properties through

experimentation [147].

5.2. Latent catalysis of epoxy–phenol reaction

The latent catalysis in thermoset curing has many

implications in polymer processing. Latent catalysts

are inert under normal conditions, but show activity at

appropriate external stimulation, such as heating.

Development of efficient latent catalysis is important

for enhancement of both storage stability and

handling of thermosetting resins, because they can

simplify the curing operation to achieve one-pot

synthesis systems. This aspect of phenolic–epoxy

curing has also interested many researchers, as

evident from the large number of reports available

in this field. However, in the following discussion,

only a few notable works will be mentioned. Thus,

Park et al. [148] investigated the cure kinetics of the

DGEBA/novolac blend system of different phenolic

contents, initiated by a cationic latent thermal catalyst

viz. N-benzylpyrazinium hexafluoroantimonate

(BPH), by means of the analysis of isothermal

experiments using a DSC. Latent properties were

investigated by measuring the conversion as a

function of curing temperature. The results indicated

that the BPH has good latent thermal initiator

properties. The cure reaction of the blend system

using BPH as a curing agent was strongly dependent

on the cure temperature, and proceeded through an

autocatalytic kinetic mechanism that was accelerated

by the hydroxyl group produced through the reaction

between DGEBA and BPH. It was concluded that at a

specific conversion region, once vitrification took

place, the cure reaction of the epoxy/phenol novolac/

BPH blend system became a diffusion-controlled cure

reaction, rather an autocatalytic reaction. In a related

study, the team investigated the viscoelastic proper-

ties and gelation in epoxy/phenol-novolac blend

system initiated with BPH as a latent cationic thermal

initiator [149]. Latent behavior was investigated by

measuring the conversion as a function of curing

temperature using traditional curing agents, such as

ethylene diamine (EDA) and nadic methyl anhydride

(NMA) in comparison to BPH. The gelation of epoxy/

phenol–novolac blend system (as deduced from the

time of modulus crossover and hence apparent) was

dependent on frequency and cure temperature.


The activation energy for cross linking (Ec) from

rheological analysis increased within the composition

range 20–40 wt% novolac resin. The 40 wt% novolac

to epoxy resin showed the highest value in the blend

system, due to the three-dimensional cross-linking

that can take place between hydroxyl groups (within

the phenol resin or epoxides) and the epoxy resin.

Sulfonium salts are the most common latent catalysts,

with a latency that emerges from the thermal

dissociation of the carbon–sulfur bond to afford

alkyl cations [150]. Other onium salts such as

iodonium, ammonium, pyridinium and phosphonium

salts have also been developed with considerable

substituent and counter ions. Latent catalysis by

quaternary ammonium borates [151] possessing

good latency and storage stability, and by (triphenyl-

phosphinemethylene) boranes [152] has also been

disclosed. The latter study made use of the thermo

reversibility of the bond formation between the ylide

and boranes for the latent catalysis for polyaddition of

phenol to epoxy. The catalytic activity could be

controlled by introduction of substituents [153]. It was

found that among many substituents, an acyl group on

the ylide is more effective in enhancing catalysis.

Based on the substituent effect, the authors concluded

that the ease of B–C bond cleavage of the alkyl

borane and the Lewis acidity of the phosphonium ion

are critical in deciding the latent nature and the

catalytic activity. The dissociation and catalysis are

depicted in Scheme 19 [153].

An alternate mode for latent catalysis in phenolic–

epoxy matrix compositions is by embedding the

catalyst in the matrix [154]. Thus, tough, flame

retardant matrices for fiber-reinforced composites

have been realized. The strategy of embedding the

initiators for matrix cure provides a means for

ensuring stability (i.e. no reaction) of the phenolic–

epoxy matrix resins up to 140 8C while the matrix is

applied to the fiber preforms. Then, it is possible to

effect a rapid reaction of the composites at the cure

temperature of 180–200 8C. The cure times can be

significantly reduced since high initiator levels can be

employed with this approach. Reaction kinetics were

investigated by DSC to predict cure times of the

system. Initiators such as tris(2,4,6-trimethoxyphe-

nyl)phosphine encapsulated in thermoplastic poly-

imide fiber sizing yielded promising results. The

composite toughness and fatigue properties of these

flame-retardant composites were excellent, and com-

parable to systems without embedding the initiator in

the fiber sizing. The toughness was comparable to that

of toughened epoxies, and fatigue to that of vinyl ester

composites, unlike general flame-retardant epoxies.

The latent catalysis by encapsulated TPP in phenol–

epoxy has been discussed above [143].

5.3. Structure–properties relations in epoxy–phenol

The structure-property relationships of epoxy–

phenolic networks are quite well understood. Tysberg

et al. [128,155] developed a series of phenol–epoxy

networks and evaluated the structural and compo-

sition dependencies of the various properties. The

epoxy-phenol equivalent ratio was varied from 1:1 to

1:7.2. Generally, Tg decreased as the phenol-content

increased. A moderately high Tg (,150 8C) was

obtained for the bisphenol-A- and tetrabromo bisphe-

nol-A- based diepoxies. A siloxane diol-based epoxy

decreased the Tg tremendously (to 87–109 8C).

However, the phenol-content increased the

flame retardance significantly. Brominated-epoxy

Scheme 19. Mechanism of latent catalysis by (triphenyl phosphi-

nemethylene) boranes in phenol–epoxy reaction [153] (Reprinted

from Macromolecules by permission, q2002 American Chemical

Society).


and siloxane-epoxy exhibited better flame resistance.

The latter manifested the lowest heat release and

smoke toxicity. The fracture toughness in terms of the

plain-strain stress intensity factor ðKICÞ indicated an

increased resistance to crack propagation. A maxi-

mum KIC of 0.85 MPa m1/2 was observed for the ratio

1:3 for DGEBA. The properties of composites also

showed a remarkable improvement over the DDS-

cured control material. Network densities were

explored by measuring the moduli in the rubbery

regions, and these experimental values were com-

pared with those predicted from stoichiometry [156].

A phenolic hardener with an average functionality of

7.3 was used in compositions ranging from 50 to 80%.

The Tgs decreased, and toughness increased as the

phenolic novolac-content in the network was

increased. Both results were correlated to a decrease

in network densities along this series. An analysis of

the co-operativity of the networks suggested a

crossover in properties from two competing factors,

viz. the network density and intermolecular forces

(hydrogen bonding). Measured fracture toughness

values exceeded those of typical untoughened epoxy

networks, and far exceeded that for commercial

phenolic networks. In addition, an increase in the

novolac-content improved the flame retardance rather

dramatically. Thus, by controlling the novolac-con-

tent to reach an appropriate phenol to epoxy ratio, a

void-free system with both favorable mechanical

properties and flame retardance could be achieved.

The carbon composites-composites properties given

in Table 21 imply formation of strong composites.

Mesogenic epoxy–phenol systems could be rea-

lized by reaction of a biphenol with a variety of

catechol novolacs [156]. In the epoxy resin cured with

catechol novolac, which has a small substituent, such

as methyl group, the glass-rubber transition almost

disappeared and a characteristic pattern, such as a

Schlieren texture was clearly observed with a crossed

polarized optical microscope. On the other hand, the

epoxy resin cured with catechol novolacs that

contained large substituents, such as phenyl groups,

showed a well-defined glass–rubber transition. These

results showed that the motion of the network chains

is highly suppressed in the former system, because of

the orientation of mesogenic groups in the cured

resins. In the latter system, however, the orientation of

the mesogenic groups seems to be prevented by the

large substituents.

Networks containing both flexible segments and

rigid structures were synthesized on the basis of

bisphenol-A novolac and diglycidylether of butane-

diol using imidazole as an accelerator [157]. A

stoichiometric ratio between epoxy groups and

phenolic groups of the novolacs led to networks

with methylene bridges as network junctions. In

contrast to this, the same reaction with bisphenol-A

led to completely soluble products. The glass

transition temperature of this soluble material was

considerably lower than that of the networks.

Increasing the content of methylene bridges in the

novolacs led to an increased Tg of the networks, and to

a decrease of the specific heat ðCpÞ at the glass

transition. Furthermore, epoxy excess led to networks

with rubber-like properties. It was concluded that

intramolecular hydrogen bonding between phenolic

hydroxyl groups considerably influenced the reactiv-

ity of the novolac with the epoxy group. The

flexibility of the cured epoxy resin can be improved

by introducing specific moieties in the matrix.

Tricyclodecane is one such moiety [158]. This

group was introduced by way of curing epoxy resin

Table 21

Flexural strength of carbon fibre composite of epoxy–phenolic system [155] (Reprinted from Polymer by permission, q2000 Elsevier Science)

Novolac/epoxy

(wt/wt)

08 Flexural

strength

(MPa)

08 Flexural

modulus

(GPa)

908 Flexural

strength

(MPa)

908 Flexural

modulus

(GPa)

Warp flexural

strength

(MPa) 28 strands

Warp flexural

strength

(MPa) 16 strands

Epoxy/DDSa 1389 159 29 8.9 442 367

50/50 2051 156 63 12.1 436 351

70/30 2020 162 66 11.3 567 372

80/20 1808 174 39 11.1 379 261

a Control without phenol.


with nonylphenol(NP)/dicyclopentadiene(CPD)–for-

maldehyde adducts, along with other commercially

available curing agents such as Jeffamine D-400 and a

novolac. The effect of these tricyclodecane groups on

the properties of the cured epoxy resins is evident

from typical data in Table 22. The presence of the

flexible NP/CPD groups impaired the mechanical

performance of the PF-cured epoxy resin, whereas in

combination with Jeffamine D-400, the properties

were practically independent of the composition. The

low cross-linking caused by the NP/CPD moieties

could be the reason for this observation. In the

presence of Jeffamine D-400 (at the place of PF), all

mechanical properties improved. An associated

increase in elongation (almost two-fold) implies that

the impact properties are significantly improved in the

presence of Jeffamine D-400.

The proposed composition and structure of the

modified novolacs are shown in Scheme 20. Epoxy-

phenol systems with markedly improved fracture

toughness (KIC ¼ 1:32 MPa m1/2), higher Tg; lower

moisture absorption, and higher thermal decomposition

temperature result when the epoxy resin contains a 4,40-

biphenylene moiety (Bis-EBP) in the backbone and

cured with a novolac [159]. This is achieved without

sacrificing the mechanical properties significantly. The

improved thermal characteristics, despite a lower cross

link density are attributed to the restricted movement

Table 22

The properties of the epoxy resins cured by two types of nonylphenol/cyclopentadiene–formaldehyde adduct (NP/CPD) with PF resin

co-curative [158] (Reprinted from Journal of Applied Polymer Science by permission, q1999 John Wiley and Sons)

Composition Weight ratio Tensile

strength (MPa)

Elongation

(%)

Flexural

strength (MPa)

Flexural modulus (MPa)

PF/DGEBA 36/64 45.3 5.2 98.6 2903

A/PF/DGEBA 33/16.5/50.5 7.3 1.6 13.3 2826

A/PF/DGEBA 23/23/54 21.7 2.6 36.5 3085

A/PF/DGEBA 14/28/58 30.3 3.5 64.4 2928

B/PF/DGEBA 33/16.5/50.5 8.3 1.5 17.4 2726

B/PF/DGEBA 23/23/54 32.6 3.8 55.6 2859

B/PF/DGEBA 14/28/58 36.6 4.9 73.6 2932

A ¼ NP/CPD (2:1), B ¼ NP/CPD (3:2), PF ¼ Novolac (see Scheme 19).

Scheme 20. Different nonyl phenol/cyclopentadiene–epoxy compositions. A and B represent the cyclopentadiene–nonyl phenol copolymer in

different rations [158] (Reprinted from Journal of Applied Polymer Science by permission, q1999 John Wiley and Sons).


Scheme 21. General synthesis protocol and structure of epoxy resins containing pyrene, anthrylene and tetramethyl phenylene moieties [160]

(Reprinted from Journal of Applied Polymer Science by permission, q2000 John Wiley and Sons).


caused by the stiff 4,40- biphenylene unit. The moisture

absorption was reduced to 1.39% form about 2% for

the control material. The gain in Tg was substantial

(154 8C for Bis-EP and 172 8C for Bis-EBP in

comparison to 142 8C for the control).

The thermal stability of phenol–epoxy systems can

be enhanced by incorporation of certain specific

moieties in the backbone. Thus, incorporation of pyrene

in the backbone of an epoxy enhanced the thermal

properties of a novolac-cured system in comparison

with epoxy resins having an anthrylene or tetramethyl

phenylene moieties [160]. The pyrene-based system

showed a marginally higher glass transition tempera-

ture, lower coefficient of linear thermal expansion,

lower moisture absorption (1.28%), and higher anaero-

bic char-yield at 700 8C. It is not known how these

marginal gains in physical parameters affect the

mechanical performances, in view of the rigid back-

bone structures. Scheme 21 shows the structure of

various epoxy systems, and the common synthesis

protocol adopted for their syntheses.

Epoxy–phenol systems meant for electronic appli-

cations must meet certain stringent requirements in

electrical performance. One way to achieve this is by

siloxane modification of such systems. Thus, vinyl

siloxane (VS)-modified cresol novolac epoxy (CNE)

cured by cresol novolac hardener (CNH) resins results

in improved performance for electronic applications

[161]. The VS-modified CNE/CNH compound pos-

sessed a lower Young’s modulus, a lower linear

coefficient of thermal expansion (LCTE), and a higher

strain at break than its unmodified counterpart. The

combination of lower mechanical moduli and lower

LCTE resulted in a 33% reduction in thermal stress

caused by thermal mismatch. The incorporation of VS

incurred a 25% reduction in the equilibrium moisture

uptake and a 16% reduction in the coefficient of

diffusion for the system. The reaction kinetics were

studied for both components, to determine the

conditions required for simultaneous cross linking in

a designed synthesis procedure. The Tg of CNE/CNH

resins could be effectively controlled through careful

adjustment of a triphenylphosphine dosage [162].

Behavior characteristic of a diffusion-controlled

reaction were observed. The incorporation of VS

incurred a 35% reduction in the equilibrium moisture

uptake and a 20% reduction in the coefficient of

diffusion for the modified resin. This modified resin

could help alleviate the popcorn problems in inte-

grated circuit packages, which result from hygro-

thermal stresses. Another process has been described

to incorporate stable dispersed polysiloxane particles

into a PF-aralkyl novolac epoxy resin, used as an

ingredient in the encapsulant formulation to withstand

the thermal stress [163]. The siliconization was done

by hydrosilylation of allyl functional aralkyl epoxy

(as per Scheme 22), with curing by novolac (PF),

catalyzed by TPP. The Tg of the cured system was

unaltered by incorporation of siloxane (Tg ¼ 163 8C).

In this case, a ‘sea-island’ structure (‘islands’ of

silicone rubber dispersed in a ‘sea’ of an epoxy resin)

was observed in the cured rubber-modified epoxy

networks via SEM. The dispersed silicone rubber-

modified aralkyl novolac epoxy resin effectively

reduced the stress of cured epoxy molding compounds

by reducing the flexural modulus and the CTE.

Electronic devices encapsulated with the dispersed

silicone rubber-modified epoxy molding compounds

exhibited excellent resistance to the thermal shock

cycling test, resulting in an extended use life for the

devices. Whereas the control exhibited about 90%

failure at around 3000 cycles, the modified one

showed only 45% failure. The internal stress of the

encapsulant is significantly lowered by siliconization.

The molding composition containing about 17.5%

cresol novolac 9% phenolic hardener and about 68.5%

silica filler (control) was modified with 3% of

siloxane-modified aralkyl group. Their comparative

properties are given in Table 23.

Phenol–epoxy based adhesives with high strength

and high heat resistance have been developed for a

novel iron-core printed circuit board with high

mechanical strength and high heat radiation capability

[164]. A reticular pattern was formed on the surface of

an adhesive resin composed of nitrile rubber (NBR),

phenolic resin, and epoxy resin. This pattern was

formed as a result of phase separation of the epoxy

resin or phenolic resin from the NBR. Using a high-

molecular-weight epoxy resin led to high adhesive

strength and high heat resistance, without phase

separation.

Reactive blends of 2,20-diallyl bisphenol A

(DABA) and a novolac epoxy resin (EPN) were

investigated for their cure behavior, and rheological,

physical, mechanical and thermal properties [12].

Cure characterization done by DSC and DMA


Scheme 22. Synthesis protocol for dispersed silicone rubber modified aralkyl epoxy system [163] (Reprinted from European Polym Journal by



confirmed the dual curing through a sequential

phenol-epoxy reaction and allyl polymerization as

shown in Fig. 15. The cure sequences are shown in

Scheme 23. The former reaction was catalyzed by

triphenyl phosphine (TPP). By the regulation of cure

conditions, the phenol–epoxy reaction could be

completed while the allyl polymerization was limited

to 40%. The mechanical properties of the neat systems

with varying composition of EPN and DABA (Table 24)

showed that an increase of epoxy concentration in

the blend led to an improvement in the tensile strength

and flexural strength of the neat castings. The flexural

strength and interlaminar shear strength of the glass

laminate also showed an improvement with an

increase in EPN concentration. Although the cross-

link density of the neat casting was enhanced by

epoxy-concentration, this did not result in any

significant variations in Tg of the cured matrix,

which was in the range 78–82 8C. Complete polym-

erization of the allyl groups resulted in an increase in

Tg (from 76 to 92 8C) and thermal stability, with a

minor deterioration in the mechanical properties of

the neat system, but a considerable increase in the

mechanical properties of their glass composites. The

Tg of the composite was significantly higher than that

of the neat resin. The comparative data of the 40 and

100% cured systems are given in Table 25. The

increased ILSS at 100% allyl cure points to the

consolidation of the interphase by enhanced cross-

linking. A similar trend in ILSS and flexural strength

confirm the possible failure of the composite at the

interphase. This improved interphase strength

accounts for the enhanced composite properties,

despite the fact that the resin becomes more brittle

on effecting 100% allyl polymerization.

Distribution of the OH groups in the phenolic

curative could dictate the thermal mechanical charac-

teristics significantly. Thus, a weak glass transition

corresponding to a rubbery modulus was observed for

a biphenol type epoxy cured by catechol type novolac,

whereas the glass transition was clearly manifested

for phenol novolac. The suppression of the Brownian

movement in catechol-cured systems due to the

orientation of the biphenyl group (aided by

H-bonding) accounted also for a better high tempera-

ture tensile and bond strength in this case [165].

5.4. Flame resistant epoxy–phenolic systems

Although epoxy–phenolic systems are inherently

flame resistant at high phenol-content, for certain

stringent applications, this property requires further

amelioration. This can be achieved by way of

incorporation of specific elements such as phosphor-

ous, Si, etc. in the network. Different strategies are

adopted for their syntheses. Thus, a novel phosphorus-

containing novolac (DOPO-PN) from 9,10-dihydro-

9-oxa-10-phosphaphenanthrene-10-oxide (DOPO)

Table 23

Properties of siloxane-modified epoxy–phenol moulding com-

pound [163] (Reprinted from European Polymer Journal by

permission, q2001 Elsevier Science)

Property Control

epoxy

Siloxane-

modified

epoxy

Tg (8C) 164 163

CTE (,Tg) 2 £ 1025 1.8 £ 1025

CTE (.Tg) 7.3 £ 1025 7.5 £ 1025

Flexural

strength (MPa )

131.3 128.1

Flexural

modulus (GPa)

12.8 10.7

Thermal stress

parameter (kPa K21)

256 192.6

Moisture

absorption (%)

0.32 0.29

Fig. 15. Non-isothermal DMA and DSC of the epoxy novolac/dial-

lylbisphenol A blend, heating rate 10 8C/min [12] (Reprinted from

Polymers and Polymer Composites by permission, q2003 Rapra

Technology).


and 4-hydroxy benzaldehyde was obtained via a

simple addition reaction, per Scheme 24 [166]. The

DOPO-PN novolac with multi-phenol groups in the

molecular chain was used as a polyfunctional curing

agent for epoxies. Curing occurred over a broad

temperature range, from 160 to 275 8C. The activation

energies of the DOPO–PN curing reactions with

epoxies were in the range 76–85 kJ/mol from DSC

evaluation. High Tg (above 160 8C) and good thermal

stability (Ti above 300 8C) were observed for the

DOPO-PN/CNE200 (o-cresol novolac epoxy)-based

resins. However, increasing DOPO decreased Tg:

High char-yields and high limiting oxygen index

(LOI) values implied good flame retardance for

DOPO-PN-cured epoxy resins. The LOI values of

the epoxy resins increased from 21 to 36 on enhancing

the P-content from 0 to 5%. DOPO led to decreased

thermal stability although the char-content showed a

proportional increase from 20 to 40% in one case,

Scheme 23. Cure sequences for diallyl bisphenol A (DABA)-Epoxy novolac (EPN) system [12] (Reprinted from Polymers & Polymer

Composites by permission, q2003 Rapra Technology).

Table 24

Properties of DABA–EPN neat resin and composites (40% allyl curing) [12] (Reprinted from Polymers and Polymer Composites by

permission, q2003 Rapra Technology)

DABA/EPN equivalent ratio Neat moulding Glass laminate

Flexural strength

(MPa)

Tensile strength

(MPa)

Elongation

(%)

Tg

(8C)

Mc

(g/mol)

LSS

(MPa)

Flexural strength

(MPa)

01:00.8 77 81 3.5 77 58 30 361

01:00.9 91 105 3 81 41 – –

01:01.0 113 118 3.6 79 19 29 369

01:01.1 106 131 4 78 17 – –

01:01.2 116 130 3.6 82 26 34 438

Table 25

Effect of extent of allyl curing on properties of neat molding and

composite (1:1 composition) [12] (Reprinted from Polymers and

Polymer Composites by permission, q2003 Rapra Technology)

Property 40% Allyl cured 100% Allyl cured

Neat Composite Neat Composite

Tensile strength (MPa) 118 – 77 –

Elongation (%) 3.6 – 3.3 –

Flexural strength (MPa) 113 369 110 458

Tg (8C) 79 86 91 107

ILSS (MPa) 29 – 32


and from 35 to 45% in another. In a related work, Liu

[167] synthesized DOPO-based aralkyl novolac by

reaction of DOPO with terephthaldicarboxaldehyde

and phenol. The resultant product (Ar-DOPO-N, see

Scheme 25), blended with PF novolac (referred as

DOPO – PF) and melamine-modified novolac

(referred as DOPO–MA) were used as curative for

o-cresol formaldehyde novolac epoxy. The cured

system possessed moderately high Tg (159–177 8C)

and thermal stability (Ti . 320 8C). High char-yields

and good flame resistance (LOI ¼ 26–32.5) were also

observed. On replacing PF novolac with melamine-

modified PF novolac, the Tg enhanced further to 160–

186 8C. A phosphorous – nitrogen synergism is

believed to be the reason for the enhanced flame

resistance in this case. An almost linear relationship

was observed between LOI and phosphorous-content.

In this case, the LOI increased to 28–33.3, despite a

decreased char-yield in air caused by the melamine.

The reverse trends in LOI and char-yields imply a

possible vapor-phase action of the flame-retardant

elements. The linear relationship between LOI and P-

content and inverse relationship between LOI and

char-yield, substantiating the above hypothesis is

demonstrated in Fig. 16.

Another approach for flame retardance was based

on DOPO-formaldehyde reaction products. Thus,

DOPO was reacted with formaldehyde to produce a

reactive 2-(6-oxid-6H-dibenz kc,el k1,2l oxapho-

sphorin-6-yl)-methanol (ODOPM). Subsequently,

novel flame-retardant curing agents for epoxy resins,

Scheme 24. Synthesis of P-containing novolac (DOPO-PN) from

9,10-dihydro-9-oxa-10-phosphaphenanthrene-10-oxide (DOPO)

and 4-hydroxy benzaldehyde [166], (Reprinted from Polymer by


Scheme 25. Synthesis of P-containing aralkyl novolac (Ar-DOPO-N) from reaction of DOPO with terephthaldicarboxaldehyde and phenol

[167]. (Reprinted from Journal of Polymer Science Polymer Chemistry Edition by permission, q2002 John Wiley and Sons).


viz. ODOPM-PN and ODOPM-MPN were prepared

from phenol formaldehyde novolac (PN), melamine-

phenol formaldehyde novolac (MPN) and ODOPM,

respectively. The compounds ODOPM–PN and

ODOPM–MPN were used as flame-retardant hard-

ener for o-cresol formaldehyde novolac epoxy (CNE)

resin for electronic applications [168]. The phos-

phorus–nitrogen synergistic effect on flame retar-

dance, combined with the rigid structure of ODOPM

resulted in better flame retardance, higher Tg

and thermal stability for the phosphorus-nitrogen

containing epoxy resin system in comparison to the

regular phosphorus-containing flame-retardant epoxy

resin. The UL 94-VO rating could be achieved with a

phosphorus-content as low as 0.81 with 2.36%

nitrogen for the ODOPM-MPN cured epoxy resin

system. No fume and toxic gas emission was

observed. The same team investigated the physical

and flame retardant properties of o-cresol formal-

dehyde novolac resin cured in the presence of 2-(6-

Oxido-6H-dibenzokc,eloxa-phophorin-6-yl)1,4-ben-

zene diol (ODOPB) and bis (3-hydroxyphenyl) phenyl

phosphate (BHPP, structures in Scheme 26) [169].

The rigid cyclic structure of ODOPB caused an

enhanced Tg for the resultant cured material. The

phosphorous-containing epoxy exhibited higher Tg;

thermal stability and flame resistance when compared

to the state-of-the-art flame-retardant formulations

based on tetrabromobisphenol-A-based epoxies. In

this case, the UL-94VO rating was achieved with a

P-content as low as 1.1%. The ODOPB-epoxy system

gave Tg values in the range 183–187 8C, whereas the

BHPP only gave Tg in the range 125–155 8C. The

modulus, thermal stability and char-yield were also

higher for the ODOPB-cured resins. Flame-retardant

phenol–epoxy systems have been realized through

synergism of phosphorous/silicon and nitrogen/silicon

also. This was achieved by curing silicon-containing

epoxy resins with DOPO–PN-based and melamine-

based phenolic resins [170]. The silicon-containing

epoxies resulted on reacting diphenylsilandiol with

Fig. 16. Variation of LOI and char-content with P-content for flame

retardant epoxy-phenol polymer with two curatives. DOPO-MA:

Melamine-modified Ar-DOPO-N; DOPO-PF: Ar-DOPO-N blended

with PF novolac; Ar-DOPO-N: Novolac from DOPO reacted with

terephthaldicarboxaldehyde and phenol [167].

Scheme 26. Synthesis of P-containing phenolic curatives 2-(6-oxido-6H-dibenzokc,eloxa-phophorin-6-yl)1,4-benzene diol (ODOPB) and Bis

(3-hydroxyphenyl) phenyl phosphate (BHPP) for epoxy resin [169] (Reprinted from Journal of Polymer Science Polymer Chemistry Edition by



epoxy resins based on bisphenol-A and o-cresol

novolac in the presence of tin (II) chloride. Silicon,

in synergism with phosphorous, conferred significant

flame retardance to the systems. LOI values of the

order of 49 could be achieved at a phosphorous level of

4.71%. However, the Tg values were reduced when

compared to the DOPO–PF and DOPO–MA systems

described above.

5.5. Miscellaneous curative for novolac

Void-free thermosets have been prepared from

novolac, cured by bisphthalonitrile (BPh) compounds.

A BPh proportion of 15–20 wt% led to tough, high Tg

and extremely flame-resistant networks [171]. The

polymer properties depend on reactant ratios. The Tg

increased from 120 8C to about 210 8C on enhancing

the BPh-content from 5 to 25 wt%. TGA indicated a

slow degradation in air, beginning at about 500 8C and

extending to 750–800 8C. The char residue at 700 8C is

60–80%, which is significantly higher when compared

to the conventional thermosets including epoxy–

novolac. The degradation is non-thermo-oxidative in

nature. The system possessed good fracture toughness

(KIC , 0.8 MPa m1/2). Investigation of the cure mech-

anism using model compounds indicated formation of

diiminoisoindoline structure. Based on the study, a

product structure as depicted in Scheme 27 was

proposed, although the authors admit that this is

inconclusive.

6. Phenolic resins with phenyl maleimide functions

Novel phenolic novolac resins, bearing maleimide

groups (PMF resin) and capable of undergoing cure

principally through addition polymerization of these

groups were synthesized by polymerizing a mixture of

phenol and N-(4-hydroxy phenyl) maleimide (HPM)

with formaldehyde in the presence of an acid catalyst

Scheme 27. Possible reaction product of bisphthalonitrile and novolac [171] (Reprinted from Polymer by permission, q2002 Elsevier

Science).


[172]. The synthesis is shown in Scheme 28. The

maleimide-content was varied by regulation of the

stoichiometry in the feed. The resins were character-

ized by chemical, spectral and thermal analyses. DSC

and DMA revealed an unexpected two-stage curing

for these systems. The DSC of PMF resin of different

maleimide-content shows identical two-stage curing

as seen in Fig. 17. Whereas the cure at 225–275 8C

was attributable to the addition polymerization

reaction of maleimide groups, the exotherm at around

150–170 8C was ascribed to the condensation reac-

tion of methylol groups formed in minor quantities on

the phenyl ring of HPM. 13C NMR exhibited a signal

at 65 ppm due to CH2–OH groups. Polymerization

studies on non-hydroxy functional, N-phenyl malei-

mide (with formaldehyde) revealed that the phenyl

groups of these molecules are activated towards

electrophilic substitution reaction by the protonated

methylol intermediates (formed during the acid

catalyzed reaction of phenol and formaldehyde).

The methylol groups are formed by activation of the

alkyl substituted HPM molecule towards the proto-

nated formaldehyde molecule. The proposed reaction

pathways are depicted in Scheme 29.

It was also revealed that the presence of the

phenolic group on N-phenyl maleimides was not

imperative for its copolymerization with phenol and

formaldehyde. The cure characterization of the PMF

resin by DMA confirmed a two-stage cure. Although

DSC implied a prominence for the methylol conden-

sation, DMA showed that the contribution of this step

in the total cross-linking process is negligible,

although helpful in causing an early gelation for the

system. The methylol substitution amounted to about

6%. The isothermal DMA of the system (Fig. 18)

showed a minor increase in storage modulus at

Scheme 28. Synthesis and curing of phenol–maleimidophenol–formaldehyde (PMF) resin from hydroxy phenyl maleimide (HPM), phenol and

formaldehyde [172] (Reprinted from Journal of Applied Polymer Science by permission, q2000 John Wiley and Sons).

Fig. 17. DSC of various phenol–maleimidophenol–formaldehyde

(PMF) resins in N2. The number denotes the composition of the

polymer in terms of the maleimidophenol-content. Heating rate

10 8C/min [210].


the initial stage due to this reaction. The DMA also

confirmed the dominance of maleimide polymeriz-

ation over methylol condensation in the network

build-up process. The curing is practically complete in

about 2 h at 250 8C.

The kinetics of both the cure reactions substan-

tiated the proposed cure mechanism for each stage

[173]. Although the initial decomposition temperature

(IDT) of the cured resin was not significantly

improved, enhancing the cross-link density through

HPM improved the thermal stability of the material at

higher temperature regime. The anaerobic char-yield

also increased proportional to the maleimide-content.

Isothermal pyrolysis and analysis of the char

Scheme 29. Proposed reaction paths for formation of phenol–maleimidophenol–formaldehyde polymer [172] (Reprinted from Journal of

Polymer Science Polymer Chemistry Edition by permission, q2000 John Wiley and Sons).

Fig. 18. Isothermal DMA of phenol–maleimidophenol-formaldehyde (PMF-29) at 250 8C. Heating rate for dynamic part, 10 8C/min. Frequency

1 Hz [173]. (Reprinted from Journal of Applied Polymer Science by permission, q2001 John Wiley and Sons).


confirmed that the pyrolysis occurs by the loss of

hydrocarbon and nitrogenous products. The resin

served as effective matrices in silica- and glass fabric-

reinforced composites. Table 26 compiles the proper-

ties of the composites derived form PMF resins of

varying maleimide-content and cured at two different

temperatures. The resins cured at 200 8C naturally

possessed inferior mechanical properties due to

incomplete maleimide polymerization. On the con-

trary, higher cross-linking led to brittle failure. The

mechanical properties were optimum for moderately

cross-linked resins, wherein the composite failure was

found to occur through a combination of fiber

debonding and resin fracture. On a comparative

scale, the properties were better than those of resole

(resole cured at 150–180 8C). Changing the reinforce-

ment from silica to glass resulted in a composite with

improved strength.

In related work, aminophenols were allowed to

react with maleic and phthalic anhydrides, producing

high yields of the corresponding N-(hydroxyphenyl)

maleamic and phthalamic acids. Formaldehyde was

allowed to react with these products in the presence of

an acid catalyst. The resulting product was dehydrated

Table 26

Mechanical properties of laminate composites of PMF using silica reinforcement, effect of composition and cure temperature [173] (Reprinted

from Journal of Applied Polymer Science by permission, q2001 John Wiley and Sons)

Resina ILSS (MPa) 08 Compressive

strength (MPa)

Flexural strength

(MPa)

Resin-content (wt%)

200 8C 250 8C 200 8C 250 8C 200 8C 250 8C

PMF-19 11 17 39 104 84 178 30

PMF-23 16 23 74 202 147 190 30

PMF-29 18 24 55 160 166 170 30

PMF-35 – 22 – 150 – 200 31

PMF-29/glass – 34 – 290 – 190 30

Resole – 19 – 150 – – –

a Number signifies the extent of phenyl maleimide-content (wt%) in PMF resin.

Scheme 30. Synthesis of isoimide- and imide-containing phenolic resin [174] (Reprinted from Journal of Polymer Science Polymer Chemistry

Edition by permission, q2000 John Wiley and Sons).


to the corresponding maleisoimide- and phthalisoi-

mide-containing novel phenol formaldehyde resins

(Novolac-like). On treatment with a sodium carbonate

solution or on prolonged heating at 50 8C the resins

suffered rearrangement to the corresponding phenol-

formaldehyde resins with pendant maleimides [174].

This rearrangement was accompanied by an increase

in the softening point of the polymers. No other

polymer properties were reported. A typical synthesis

protocol is shown in Scheme 30.

6.1. Maleimide–phenolic resin cured with

allyl–phenolics

The thermal curing of the PMF system through

polymerization of maleimide group resulted in

comparatively brittle matrices. As a result, the

mechanical properties of the resultant composites

did not improve significantly over the conventional

phenolic resins. Hence, the concept of BMI toughen-

ing by way of reactive blending with allyl compounds

(discussed in Section 2.1) was extrapolated to the

PMF system. Thus, new addition-cure phenolic resin

systems were developed, based on the co-reaction of

maleimide-functional phenolic resin (PMF) with an

allyl-functional novolac (PAF) in varying proportions

[175]. The PAF resin was derived from 2-allyl phenol,

phenol and formaldehyde. The Alder-ene cure

sequences to form a cross-linked network system

were evidenced from the cure characterization studies

by DSC and DMA. Extrapolating the cure sequences,

the network structure of a 1:3 blend can be depicted as

in Scheme 31 for this system. Increasing allyl–phenol

content in the reactive blend decreased the cross

linking in the cured matrix, leading to enhanced

toughness and improved the resin-dominant

mechanical properties of the resultant silica laminates.

The mechanical properties of the composites of the

blend of different compositions are compiled in

Table 27. Under the cure conditions employed, the

properties are optimum at an allyl/maleimide ratio of

1:3. Changing the reinforcement from silica to glass

resulted in further amelioration of the resin-reinforce-

ment interaction, but the resin-dominant properties of

the composite remained unaltered. However, the

presence of AP diminished the thermal stability. The

cured polymers showed Tg in the range 170–190 8C.

In an analogous strategy, a one-component, self-cross

linkable polymer was synthesized by reacting a

mixture of phenol, HPM and 2-allyl phenol with

formaldehyde (PMAF) where, the allylphenol/malei-

mide ratio was maintained as 1:3. When compared to

the two-component system of similar composition,

this system showed inferior composite properties

[175] (data in Table 27). The structures of PMAF

resin and the likely structure of its cured product are

shown in Scheme 32.

Scheme 31. Cured structure of blend of phenol–maleimidophenol–formaldehyde (PMF) and Phenol–allylphenol–formaldehyde (PAF) [175]

(Reprinted from Journal of Applied Polymer Science by permission, q2001 John Wiley and Sons).

Table 27

Mechanical Properties of Silica Laminate of PMF-29, blended with

PAF resin [175] (Reprinted from Journal of Applied Polymer by

permission, q2001 John Wiley and Sons)

Blend Allyl/malei-

mide ratio

(equivalent)

ILSS

(MPa)

Flexural

strength

(MPa)

08 Compressive

strength (MPa)

MAP1 01:01 23.5 203 253

MAP2 01:02 25 209 222

MAP3 01:03 23 202 210

MAP4 01:04 29 204 251

MAP5 01:05 23.5 176 192

PMF 00:01 24 168 157

PMAF 01:03 20 165 171


6.2. Maleimide–phenolic (PMF)–epoxy blend

The majority of phenolic-based adhesive formu-

lations makes use of the condensation cure of resole

resins, which necessitates the application of high

pressure during the adhesive bonding to form void-free

glue lines. The limited shelf life of the resole-based

adhesives at ambient temperature is another short-

coming of phenolics. Addition curable phenolics-

based systems could be better suited for adhesive

applications than conventional phenolics. Thus, the

maleimide-functional novolac phenolic resin (PMF)

was cured with epoxy resin (EPN). It was evaluated for

the adhesive properties such as lap shear strength

(LSS) and T-peel strength (TPS) using aluminum

adherends, both in the thermally self-cured state as

well as co-cured with epoxy resins [176]. The adhesive

properties of the self-cured resin, although inferior at

ambient temperature, improved at high temperature

(150 8C) and the properties were found to depend on

the extent of maleimide curing. Table 28 lists the

adhesive properties of a typical PMF resin cured under

different conditions. The extent of maleimide curing

was estimated from studies on thermal polymerization

of the model compound, HPM. Although the adhesive

strength is not high, it is interesting that it increases at

high temperature due to softening of the rigid matrix.

The high temperature strength is also proportional to

the maleimide cross-linking (see Table 28). However,

on co-curing PMF with the EPN epoxy resin through

the phenol-epoxy reaction and subsequent maleimide

polymerization, the adhesive properties improved

significantly. The adhesive properties showed a strong

dependence on the nature of the epoxy resin used,

stoichiometry of the reactants, the concentration of

imide groups in the phenolic resin and the extent of

polymerization of the maleimide groups. The cure

sequences are shown in Scheme 33. The dependence of

the LSS at RT, 150 and 175 8C as well as the TPS at RT

on the epoxy–phenolic stoichiometry for the PMF-29/

EPN system is shown in Fig. 19.

The adhesive strength (both LSS and T-peel)

becomes optimized at a 1:1 equivalent of the two

components. The LSS at 150 and 175 8C as well as the

retention of the LSS at these temperatures also become

optimized at this stoichiometry. This shows the

dominance of the epoxy–phenol reaction in deciding

the adhesive characteristics of the system. The 1:1

stoichiometry ensures maximum cross-linking through

the phenol–epoxy reaction. A maximum extent of

phenol–epoxy cure is desirable to promote better

adhesion and better retention at higher temperature. An

excess of EPN leaves unreacted epoxy groups,

leading to matrix plasticization, and diminution of

properties. Although their absolute values were lower,

the phenol-rich systems exhibited better high-tempera-

ture retention of properties thanks to the presence of

thermally stable imido–phenol groups in the network.

Typical adhesive values of PMF resins with varying

Scheme 32. Structure of one-component phenol–allyl phenol–maleimidophenol–formaldehyde polymer (PMAF) and its cured product.

Table 28

Effects of cure temperature and maleimide polymerisation on LSS

of self-cured PMF-29. (initial cure: 170 8C/30 min; ultimate-cure:

30 min) [176] (Reprinted from Journal of Applied Polymer Science

by permission, q1999 John Wiley and Sons)

Post-cure

temperature

(8C)

Extent of

maleimide

polymerisation

(%)

LSS at

RT

(MPa)

LSS at

150 8C

(MPa)

Retention of

LSS at 150 8C

(%)

200 37 2.1 3.5 167

225 75 4.1 6.8 166

250 89 4.6 9.1 198


maleimide-content and measured at varying tempera-

tures are given in Table 29. The maleimide is

polymerized only to a minor extent (37%) under these

conditions, whereas the phenol-epoxy reaction is

complete. In comparison to the conventional novolac

(PF), the imide– novolac (PMF) contributed to

improved adhesion and better adhesive property

retention at higher temperature when cured with EPN.

Good bonding resulted for PMF-29 with moderate

maleimide-content.

In fact, for the PMF/EPN system, the high

temperature adhesion was directly proportional to

the extent of maleimide polymerization, which was

controlled by the cure temperature and time. Fig. 20

shows the direct relationship between high tempera-

ture (150 8C) retention of adhesion strength and the

extent of maleimide curing. The TPS (at RT) tends to

increase with maleimide-content, but the brittleness of

the system at higher maleimide-loading deteriorates

the strength. The comparative thermo-adhesive pro-

files of the PMF-epoxy and PF-epoxy, showing

Fig. 19. Effect of phenol/epoxy stoichiometry on Lap shear strength

and T-peel strength of phenol–maleimidophenol–formaldehyde

(PMF-29)/epoxy novolac (EPN) system (B) LSS at RT; (X) LSS at

150 8C; (O) LSS at 175 8C; (V) TPS at RT [176] (Reprinted from

Journal of Applied Polymer Science by permission, q1999 John

Wiley and Sons).

Scheme 33. Cure sequences of phenol–maleimidophenol–formaldehyde–epoxy novolac blend [176] (Reprinted from Journal of Applied

Polymer Science by permission, q1999 John Wiley and Sons).


the superiority of the former are shown in Fig. 21.

The absolute adhesion values and the thermal

resistance are good for the former.

Complimentary DMA studies led to the con-

clusion that the superior thermo-adhesive profile of

the PMF/EPN system is contributed by secondary

attraction forces induced by the polar imide groups.

The partial polymerization of the imide helps

consolidate these secondary forces of attraction

both within the resin system and at the bonding

interphase. The enhanced bonding characteristics of

the PMF–EPN system was manifested in a better

short beam shear strength (SBSS) of the carbon-UD

composite based on PMF/EPN (85.5 MPa) as

against that of PF/EPN (68.5 MPa). The study

also implied that the completion of epoxy–phenolic

reaction and moderate cross-linking through the

maleimide polymerization are conducive for

achieving optimum adhesive properties.

Although the adhesive properties of PMF/EPN

blends were good, attempts were made to improve it

further by blending with elastomeric modifiers, viz. a

liquid carboxyl-terminated butadiene acrylonitrile

(CTBN-L), a solid carboxyl-terminated butadiene

acrylonitrile (CTBN-S), epoxidized hydroxyl-termi-

nated polybutadiene (EHTPB) and epoxidized butyl

acrylate acrylonitrile polymer (EPOBAN) [177]. The

adhesive properties were found to depend on the nature

and concentration of the elastomers as well as on the

nature of the thermoset matrix being modified.

Fig. 21. Comparative thermo-adhesion profile for (B) phenol–

maleimidophenol–formaldehyde (PMF-29)/epoxy novolac (EPN)

and (X) phenol–formaldehyde (PF)/EPN systems [176] (Reprinted

from Journal of Applied Polymer Science by permission, q1999

John Wiley and Sons).

Fig. 20. Retention of lap shear strength (at 150 8C) and T-peel

strength at RT and their relationship to the extent of maleimide

polymerisation for phenol– maleimidophenol – formaldehyde

(PMF-29)/epoxy novolac (EPN) system. (B) Maleimide polym-

erisation; (X) LSS retention at 150 8C; (P) T-peel at RT.

Table 29

Adhesive properties of various PMF resins reacted with EPN (1:1 stoichiometry, cure: 170 8C/30 min and 200 8C/30 min.; phenol-epoxy

reaction is complete and extent of maleimide cure is 37%) [176] (Reprinted from Journal of Applied Polymer Science by permission, q1999

John Wiley and Sons)

System LSS at RT

(MPa)

LSS at 150 8C

(MPa)

Retention of LSS

at 150 8C (%)

LSS at 175 8C

(MPa)

Retention of LSS

at 175 8C

(%)

T-peel at ambient

(kN/m)

PF/EPN 15.5 7.2 47 3.1 20 0.4

PMF-19/EPN 15.3 5.6 37 2.8 18 0.43

PMF-29/EPN 17 14.3 84 8 47 0.56

PMF-42/EPN 11.3 8.2 73 3.9 35 0.08


The adhesive properties at ambient temperature of the

self-cured, brittle PMF resin were dramatically

improved by the inclusion of the elastomers, the

increase being substantial in the case of high molecular

weight CTBN. Typical results are shown in Fig. 22.

However, contrary to expectation, the adhesive

properties of the epoxy–PMF system were only

marginally improved by the high molecular weight

CTBN, while the other elastomers were nearly

ineffective. For both the self-cured and epoxy-cured

PMF systems, the inclusion of these elastomers

generally decreased the high temperature adhesive

properties, evidenced also from their dynamic mech-

anical spectra. This trend is evident from the LSS values

at 150 8C in case of CTBN-modified matrices shown in

Fig. 23. The decrease is sharper with CTBN-L.

The presence of phase-separated elastomer par-

ticles in the modified systems was evident in the SEM

analyses. The adhesive properties of the elastomer-

modified systems are given in Tables 30 and 31.

CTBN-L was not effective in enhancing the properties,

whereas high molecular weight CTBN-S was very

effective for both LSS and TPS. The relative increase

was more pronounced for the PMF self-cured system

rather than the PMF-EPN system. The epoxy-contain-

ing EHTPB and EPOBAN were capable of improving

the properties, but only to an insignificant extent. For a

given composition, the decrease in high temperature

property was sharper for CTBN-L than for CTBN-S,

due to the reduced compatibility of the former in

the matrix. Similarly, between EPOBAN and EHTPB,

Fig. 22. Effect of carboxyl terminated polybutadiene (CTBN) on the

adhesive characteristics of phenol–maleimidophenol–formaldehy-

de(PMF) and PMF/epoxy novolac blend; CTBN-S: solid, CTBN-L:

liquid [177] (Reprinted from Journal of Applied Polymer Science


Fig. 23. Effect of carboxyl terminated polybutadiene (CTBN) on the

high temperature adhesion of phenol–maleimidophenol–formal-

dehyde (PMF) and PMF/epoxy novolac systems. S ¼ CTBN-Solid,

L ¼ CTBN-Liquid [177] (Reprinted from Journal of Applied


Table 30

Effect of CTBN-S on the adhesive properties of PMF-29 and PMF-

29/EPN systems [177] (Reprinted from Journal of Applied Polymer

Science by permission, q1999 John Wiley and Sons)

Concentration

of the additive

(phr)

Self-cured PMF PMF/EPN

LSS (MPa) TPS

at RT

(kN/m)

LSS

(MPa)

TPS at

RT

(kN/m)

RT 150 8C RT 150 8C

0 2.1 3.5 0.01 17.0 14.3 0.56

10 5.3 5.2 0.43 18.5 10.6 1.20

15 8.0 4.5 0.64 20.0 9.4 1.70

20 11.3 4.3 0.70 20.0 7.8 2.00

30 14.3 3.9 1.00 20.0 6.5 2.00


the former was comparatively better in retaining the

LSS at high temperature; possibly due to its additional

cross linking through the cyclization reaction of the

nitrile groups under the cure conditions. A DMA study

of the elastomer-modified PMF/EPN system showed

that the absolute values of E0 are reduced with the

inclusion of these elastomers, whereas the temperature

for the maximum tanðdÞ remain almost unaffected

(Fig. 24). SEM analysis showed the single-phase

morphology for the unmodified PMF and PMF/EPN

systems and a phase-separated morphology for the

elastomer-modified materials. The size and state of

the precipitated secondary phase depend on the

molecular weight as well as the chemical reactivity

of the added elastomer. For the CTBN-S modified

systems, the dispersed phase was co-continuous and

this morphology led to the maximum improvement in

adhesive properties. For the CTBN-L-, EHTPB- and

EPOBAN-modified PMF/EPN systems, the dispersed

phase was non-uniform with a wide distribution of

particle sizes. These particles were too large to be

effective in improving the toughness of the system, and

hence reduced the adhesive properties. The elastomers

were, by and large, very effective in enhancing the peel

strength of the adhesives.

7. Pendant phenol functional linear polymers

7.1. Pendant phenol-functional thermoplastics

Linear vinyl polymers with pendant phenolic

groups were realized by free radical copolymeriza-

tion of N-(4-hydroxy phenyl) maleimide (HPM) with

Table 31

LSS (in MPa) of PMF-29 and PMF-29/EPN system modified by EHTPB and EPOBAN [177] (Reprinted from Journal of Applied Polymer

Science by permission, q1999 John Wiley and Sons)

Concentration of the additive (phr) EHTPB EPOBAN

Self-cured PMF PMF/EPN Self-cured PMF PMF/EPN

RT 150 8C RT 150 8C RT 150 8C RT 150 8C

0 2.1 3.5 17 14.3 2.1 3.5 17 14.3

10 3.5 2.2 16 6.3 2.2 2.6 15.4 4.9

15 4.6 1.8 14.7 5.3 – – – –

20 5.4 1.7 13.2 3.6 3.3 2.4 16.6 5.3

30 5.8 0.8 12.9 2.5 – – – –

Table 32

Composition, molecular and mechanical characteristics of BNM

polymers (BuA/AN weight ratio ¼ 2) [178] (Reprinted from

Journal of Adhesion Science and Technology by permission,

q2001 Koninklijke Brill NV)

Polymer

reference

HPM

(wt%)

Mn

£ 1024

Poly

disper-

sity

index

Tensile

strength

(MPa )

Elonga-

tion (%)

Softening

tempera-

ture (8C)

BNM-1 5.1 5.69 1.83 20 300 ,28

BNM-2 7.4 5.78 1.74 21.5 225 ,28

BNM-3 9.5 5.2 1.88 24.5 23 40

BNM-4 13.6 5.52 1.74 – – 45

Fig. 24. DMA spectra of the elastomer-modified phenol–maleimi-

dophenol–formaldehyde (PMF-29)/epoxy novolac blends(EPN)

cured in presence of epoxididsed hydroxy terminated polybutadiene

(EHTPB),Liquid carboxyl terminated butadiene–acrylonitrile poly-

mer (CTBN-L), Solid carboxyl terminated butadiene–acrylonitrile

polymer (CTBN-S) and epoxidised butylacrylate–acrylonitrile

polymer (EPOBAN)[177] (Reprinted from Journal of Applied



butyl acrylate (BuA) and acrylonitrile (AN) and were

characterized [178]. These thermoplastics (BNM)

could form good films, with mechanical and adhesive

properties dependent on the maleimido phenol-

content in the chain, as given in Table 32. The

structure of the terpolymers can be found in

Scheme 34. The polymer films could be directly

served as thermoplastic adhesives and their adhesive

properties were studied in detail. LSS were deter-

mined at ambient temperature, 50 8C and at 196 8C,

whereas T-peel strength (TPS) and flat wise-tensile

strength (FTS) were determined at ambient tempera-

ture. Enhancing the HPM-content increased Tg; the

tensile strength and modulus of the films, but

decreased their elongation. Nominal increase in

phenol-content was found conducive to improving

the adhesive properties of the films. At higher

concentrations of phenol, the film properties showed

a decreasing trend due to the embrittlement caused

by the rigid maleimide groups. Whereas the adhesive

property at 50 8C increased linearly with HPM-

content, due to increased Tg; a reverse trend was

observed for the property measured at 2196 8C due

to dominance of embrittlement effect. This trend is

demonstrated in Fig. 25. The reduced flow charac-

teristics of the high HPM-loaded systems led to

diminished tensile strength properties evaluated

using flat-wise bonded aluminium honeycomb adher-

ends. The T-peel was also optimum for BNM-3,

whereas FTS was good for BNM-2. The dependence

of TPS and FTS on maleimido-phenol-content of the

film adhesives are presented in Fig. 26.

On the contrary, enhanced HPM concentration

promoted the adhesive properties for vulcanization

bonding of NBR to aluminium. This is evident from

the data presented in Table 33. The failure mode is

also indicated. High maleimide-content promotes

bonding to the extent that failure is induced in the

rubber substrate for BNM-3 and BNM-4 in the peel

mode. Addition of silica filler marginally improved

the LSS for metal–metal system, but was detrimental

for metal–rubber bonding. A reverse trend was

observed for the carbon-filled BNM system, with the

diminished performance for metal – metal and

enhanced performance for metal–rubber systems

attributed to the weakening of the interphase in the

former, and to a possible reinforcement of the rubber

phase by carbon in the latter. The fillers generally

Scheme 34. Butyl acrylate–acrylonitrile–maleimidophenol (BNM) polymers and their thermosetting derivatives, A-triazine; B-propargyl

ether; C-epoxy.


improved the temperature capability of the adhesives.

This is shown in Fig. 27, which compares the adhesion

strength at 50 8C for BNM-3 filled with silica and

carbon. The retention is calculated based on the

strength at ambient temperature. The fillers impaired

the flow properties of the resin, and thereby adversely

affected the flat-wise bonded tensile strength in both

cases. A compromise between reinforcement and flow

property is reached at about 20 wt% of the filler

loading, where the adhesive property is maximum. In

general, the optimum LSS and T-peel properties were

obtained for BNM-3 and FTS properties for BNM-2.

7.2. Pendant phenol-functional addition-cure systems

The adhesives properties and their temperature

retention could be improved further on transforming

these film adhesives to addition-curable phenolic

thermoset films by reactive blending with an epoxy

Fig. 26. Dependence of T-peel strength (TPS) and flat-wise tensile

strength (FTS) on hydroxy phenyl maleimide-content for thermo-

plastic film adhesive [178] (Reprinted from Journal of Adhesion

Science and Technology by permission, q2001 Koninklijke Brill

NV).

Fig. 27. Effect of filler on adhesive (metal-metal) properties of butyl

acrylate–acrylonitrile–maleimidophenol (BNM-3) thermoplastic at

ambient and at 50 8C [178] (Reprinted from Journal of Adhesion

Science and Technology by permission, q2001 Koninklijke Brill NV).

Fig. 25. Evolution in lap shear strength at different temperatures of

butyl acrylate–acrylonitrile–maleimidophenol (BNM) polymers on

hydroxy phenyl maleimide (HPM)-content[178] (Reprinted from

Journal of Adhesion Science and Technology by permission, q2001

Koninklijke Brill NV).

Table 33

Aluminium-to-NBR vulcanization bonding properties of BNM

polymers (AF: cohesive failure in the adhesive; IF: adhesive–

rubber interface failure; RF: cohesive failure in the rubber) [178]

(Reprinted from Journal of Adhesion Science and Technology by

permission, q2001 Koninklijke, Brill NV)

Adhesive property BNM-1 BNM-2 BNM-3 BNM-4

LSS at ambient temperature

(MPa)

0.26 0.37 0.39 0.45

Nature of failure AF IF IF IF

TPS at ambient tempearture

(kN/m)

2.8 4.1 5.0 5.0

Nature of failure AF IF RF RF


resin. Alternatively transformation of the phenolic

OH groups to propargyl ether, cyanate etc and

subsequent curing also provided thermosets [179].

On cross-linking through reaction with a diepoxide,

the mechanical and adhesive properties of the film

adhesive (Ep-BNM) improved significantly. In this

case, the properties increased with the concentration

of imido–phenol and its cross-linking. However, high

imide-content and higher cross-linking were detri-

mental for the low temperature adhesive properties

(at 196 8C), although the retention at high temperature

was improved.

Values of the LSS at different temperatures

determined with Ep-BNM polymers (epoxy-cured

BNM-3) using aluminium substrates are shown in

Fig. 28, showing that the trend is the same for the LSS

at 25, 50 and 2196 8C. The initially observed

adhesive performance advantage is lost at higher

HPM-content. However, the effect is less pronounced

for the LSS at 2196 8C. As observed with thermo-

plastic BNM polymers [178], the increase in the LSS

with HPM-content for the Ep-BNM polymers is

attributed to the increased cohesive strength induced

by the polar hydroxy- and maleimide moieties.

Beyond 9.5% of HPM, the polymer becomes rather

brittle, leading to inferior strength. In comparison

with the corresponding thermoplastic BNM polymers,

Ep-BNM polymers exhibit higher LSS values at 25

and 50 8C, but inferior properties at 2196 8C. The

LSS and peel strength values of the different Ep-BNM

adhesives bonded aluminum-to-NBR joint are given

in Table 34. Both LSS and peel values increase with

increase in HPM-content and the failure mode

changes from cohesive failure in adhesive (AF) to

adhesive-rubber interface failure (IF), and finally to

cohesive failure in the rubber (RF) with polymers

having more than 7.4% HPM. Similar results were

obtained with the thermoplastic BNM polymers

where the LSS varied from 0.26 to 0.45 MPa and

peel strength varied from 2.8 to 5.0 kN/m. The effect

of introducing cross-links into the BNM polymer was

much less pronounced for the case of metal-to-rubber

bonding. Designing a one-component, self-curing

type thermoset (through phenol–epoxy reaction)

based on an acrylic polymer bearing both epoxy and

phenol groups (i.e. Gly–BNM) and was not very

effective in providing good adhesive properties.

Cross-linking the BNM polymers after chemical

reaction of the phenol group through propargylation

improved the adhesive properties at RT to a

comparable extent with that of epoxy cross-linked

resin (Ep-BNM). The propargyl–BNM polymer

exhibited good LSS values at 25 8C, only slightly

inferior to that of the Ep-BNM polymer; however, the

LSS properties at 50 8C were very poor and nearly the

same as with BNM. This could be due to incomplete

thermal curing of propargyl groups under these cure

conditions. Neither was chemical transformation of

Fig. 28. Lap shear strength (LSS) at 25, 50 and 2196 8C for

epoxy-cured butyl acrylate–acrylonitrile–maleimidophenol (Ep-

BNM) polymers as a function of Hydroxy phenyl maleimide

(HPM)-content in the polymer [179] (Reprinted from Journal of

Adhesion Science and Technology by permission, q2001

Koninklijke Brill NV).

Table 34

Aluminium-to-NBR vulcanisation bonding properties of Ep-BNM

polymers (AF: adhesive failure; IF: adhesive–rubber interface

failure; RF: rubber failure) [179] (Reprinted from Journal of

Adhesion Science and Technology by permission, q2001 Konink-

lijke, Brill NV)

Adhesive

property

Ep-BNM-1 Ep-BNM-2 Ep-BNM-3 Ep-

BNM-4

LSS at 25 8C

(MPa)

0.3 0.33 0.53 0.52

Type of failure AF AF RF RF

Peel strength

at 25 8C (kN/m)

2.8 4.2 5.1 5.0

Type of failure AF IF RF RF


the phenol to cyanate groups, and its subsequent

curing to polycyanurate very effective in improving

adhesive properties. The lowest LSS values were

exhibited by the cyanate–BNM polymer. All the

cross-linking reactions impaired the peel strength of

the film adhesive. Metal-to-NBR joint strength was

found to be nearly the same as that for BNM-3 for all

the modified polymers, indicating that the cross-links

introduced did not appreciably affect the strength.

Except Ep-BNM, all other polymers resulted in an

adhesive–rubber interfacial failure of the metal-to-

rubber joints, probably due to the weaker interactions

of the propargyl and cyanate ester groups with NBR,

in comparison to the epoxy. The various chemical

structures have been included in Scheme 34. The

adhesive properties of the thermosetting BNM poly-

mers are given in Table 35. The dominance of epoxy

curing in imparting the adhesive properties is clear.

8. Propargyl ether functional phenolics

Although less commercially exploited, propargyl

ether-functional phenolic resins were developed as a

potential hydrophobic substitute for epoxies in

Scheme 35. Synthesis and curing of propargyl novolac (PN) resins [197] (Reprinted from Polymer International by permission, q2001 Society

of Chemical Industry).

Table 35

Adhesive properties of modified BNM polymers compared to the unmodified thermoplastic polymer (IF: adhesive–rubber interface failure; RF:

rubber failure) [179] (Reprinted from Journal of Adhesion Science and Technology by permission, q2001 Koninklijke, Brill NV)

Adhesive property BNM-3 Ep-BNM-3 Gly-BNM Propargyl BNM Cyanate BNM

Metal-to-metal

LSS at 25 8C, (MPa) 17.5 20.2 8.5 18.0 8.0

LSS at 508C, (MPa) 6.9 14.5 6.0 6.0 2.8

Retention of LSS at 508C (%) 39 72 70 33 35

TPS at 25 8C (kN/m) 1.70 1.40 0.25 0.30 0.30

Aluminium-to-NBR

LSS at 25 8C( MPa) 0.39 0.53 0.39 0.42 0.42

Type of failure IF RF IF IF IF

TPS at 25 8C (kN/m) 5.0 5.1 4.0 4.3 3.8

Type of failure RF RF IF IF IF


advanced composites, electronics, adhesives and

coatings. The majority of thermosets, such as epoxy,

BMI, etc. absorb moisture up to 5%, resulting in low

hot/wet physico-chemical properties. For advanced

applications, the required hot/wet performances for

many composites are to exceed temperatures of about

230 8C. Hydrophilicity can lead to easy matrix

delamination too. Obviously, the high ðDkÞ of

absorbed water is detrimental for electronic appli-

cations of such polymers. Propargyl ether resins can

address some of these problems. The structural

similarity of propargyl ether to epoxy resins is useful

for their preparation, processing and development of

thermally stable polymers [180].

Derived from a phenolic backbone, resins based on

propargyl ether possess some of the basic features of

the parent polymer. Thus, propargyl phenolics can

offer a compromise matrix between epoxy and

phenolics. Propargyl ether resins are formed easily

from the phenolic precursor by the Williamson’s

reaction with the propargyl halide. The synthesis of

propargyl novolac (PN) is shown in Scheme 35. The

various cure mechanisms are also depicted in the

above scheme. More than PN resins, the synthesis and

curing of bispropargyl ether resins (BER) have been

quite well described in literature [181–183]. The

latter is formed from the corresponding bisphenol.

8.1. Curing of propargyl ether resins

The curing of propargyl ether resins proceeds by

Claisen rearrangement followed by addition polym-

erization of the resultant chromene. This has been

confirmed by studies on model propargyl ethers [184].

The structural dependence on cure kinetics of bis

propargyl ether resins (BER) has been investigated

[185]. Thus, bis propargyl ethers of bisphenol-A,

(BPBA), bisphenol ketone (BPK) and bisphenol

sulfone (BPS) were synthesized and characterized.

These monomers were thermally polymerized to the

corresponding poly(bischromenes). The cure beha-

vior, as monitored by DSC, depended on the structure

of the monomer. The non-isothermal kinetic analysis

of the cure reaction using four integral methods

revealed that the presence of an electron-withdrawing

group did not favor the cyclization reaction leading to

formation of chromene, which precedes the poly-

merization, and in agreement with the proposed

polymerization mechanism. Thus, the cure tempera-

ture and activation energy for the reaction increased in

the order BPBA , BPK , BPS. The cure profile

under isothermal and non-isothermal conditions

could be simulated from the kinetic parameters.

Typical isothermal cure predictions at varying tem-

peratures are shown in Fig. 29, along with an

experimental cure profile at one temperature for

comparison with the prediction.

The substituent group bridging the two phenyl

rings also influenced the thermal stability of the

resultant polymers. Thus, sulfone and ketone-contain-

ing polymers were more stable than the isopropyli-

dene-containing material. Kinetic analysis of thermal

decomposition of the major step involving the

chromene moieties revealed a two-stage degradation

mechanism. The computed activation energies

implied that the initiation of the degradation reaction

was disfavored by the electron-withdrawing nature of

the bridge unit, probably through destabilization of

the intermediate radical. This made such polymers

more thermally stable. The possible typical mechan-

ism for initiation of thermal degradation of the cross-

linked chromene is shown in Scheme 36, for polymer

of BPBA. The decomposition is proposed to initiate at

the crowded carbon, giving rise to benzyl type free

radicals, with a stability that is decreased by the

presence of electron-withdrawing substituents on the

phenyl ring. The same scheme depicts the structure of

BER, and its likely cured structure.

In the presence of a catalyst, the polymerization

mechanism is different, as are the properties of the

resultant polymers. Certain catalysts are found to favor

the linear polymerization of the acetylene groups. Thus,

polymerization of nitrophenyl propargyl ethers with

tungsten- and molybdenum-based metathesis catalyst

systems (MoOCl4/Me4Sn and MoCl5/EtAlCl2) gave

soluble linear high-molecular-weight poly(nitrophenyl

propargyl ether) (Mw ,4 £ 105) [186]. The expected

cyclotrimerization is favored when the system is heated

in the presence of the catalysts, such as cyclopentyl

cobalt dicarbonyl [181]. In other cases, a mixture of

cyclotrimerization, chromene formation and linear

polymerization occur. Consequently, the thermal

stability of polymer is strongly dependent on the

mode of polymerization The Claisen rearrangement

and the subsequent polymerization are so exothermic

(,1.200 kJ/g for PN resin) that curing has to be done


under controlled conditions, with a slow phase.

B-staging of the resin is possible at 185 8C. Processing

of the B-staged resins helps avoid the otherwise

excessive cure shrinkage (10–12%). The shrinkage of

the B-staged material to the final network is below 1%.

The B-staged BER systems have excellent tack and

drape and are easily processable. The isothermal

stability of propargyl ether resins is comparable to

that of acetylene terminated resins [187]. Final curing

of propargyl resin is done at high temperature (typical

case 208 8C/4 h). Post curing is carried out at 260 8C; Tg

is about 300 8C, and post curing boosts it to 360 8C. This

value is high when compared with other thermosets.

BER thermosets only absorb 0.3–0.4% moisture.

Typical tensile properties of BER resin are as follows:

tensile strength 103 MPa, modulus 4.55 GPa and

elongation 1.9%. Flexural strength and modulus are

105 MPa and 4.33 GPa, respectively. These are higher

than values for other thermosets.

8.2. Structure–property relation in propargyl

phenolics

Propargyl ether resins of cyclopentadiene–phenol

has been synthesized, but no property data are given

[180]. The mechanical strength of the bispropargyl

ether resin of diphenols are almost double those of

their corresponding acetylene-terminated analogues

[180], but with identical isothermal stability [187].

Propargyl etherified and glycidyl etherifed novolac

treated with a siloxane and modified with amino

silane, zinc stearate and carbon black provided

molding compositions with good mechanical

strength and moderately good thermal capability.

The flexural strength amounted to 173 MPa with a

flexural modulus of 12.5 GPa. About 35% strength

retention was observed at 260 8C. The polymer

possessed a Tg of 243 8C, excellent moisture

resistance and good solder crack resistance [188].

The adhesive strength of phenol–formaldehyde

resin increased on modification with propargyl

glycidyl ether. The adhesive strength increased to

45.8 kg/cm2, heat resistance by 50 8C, and Brinell

hardness to 33 kg/mm2 [189]. Thermosetting resins

with good dimensional stability and heat resistance

are obtained by mixing propargyl-etherified resin

with resole [190]. Resin compositions with good

workability and curability for laminates, heat- and

moisture resistances and low ðDkÞ are obtained using

propargyl-etherified phenolic resin-based formu-

lations [191]. A siloxane-modified, heat- and

moisture-resistant phenolic resin composition for

Fig. 29. Isothermal time-conversion profile prediction for bispro-

pargyl ether of bisphenolA (BPBA) at various temperatures. (B)

experimental data at 245 8C [185] (Reprinted from Polymer by


Scheme 36. Proposed mechanism for curing of BER and initiation

of thermal degradation of poly(bischromenes).


sealing semiconductor devices has been obtained by

blending propargyl etherified novolac with amine-

terminated polysiloxane [192].

Novel ester–imide prepolymers terminated by 3-

(4-aminophenoxy)-propane-1-yne as a thermally cur-

able group were synthesized [193]. The cured resins

showed excellent mechanical and thermal properties,

i.e. flexural strength of 269–370 MPa, Tg of 225–

269 8C, moisture absorption of 0.20–0.78%, and ðDkÞ

of 3.0–3.2. Glass-cloth-reinforced composites from

propargyl ether-terminated ester–imide prepolymers

demonstrated excellent mechanical, chemical, and

electrical properties [194]. The storage stability of the

varnish of the prepolymer was also good. The initial

flexural strength was well maintained even after

1000 h at 200, 220 and 240 8C. Glass-cloth-reinforced

composites prepared from prepolymers of blends of

propargylether-terminated ester–imide and a BMI

demonstrated that the Tg of the composite is directly

related to the weight ratio of BMI [195]. The

composites demonstrated excellent initial mechanical

properties. The chemical and electrical properties

under severe long-ageing conditions were also good.

Novel BZ monomers containing arylpropargyl ether

have also been reported [196]. On curing, thermally

stable polybenzoxazines resulted. One monomer is a

monofunctional BZ, i.e. 4-propargyloxyphenyl-3,4-

dihydro-2H-1,3-benzoxazine (P-appe), and the other

is a bifunctional BZ, bis(4-propargyloxyphenyl-3,4-

dihydro-2H-1,3-benzoxazinyl)isopropane (B-appe).

The synthesis protocol and structures of these

monomers and the resultant polymers are given in

Scheme 37.

The cure behavior of the P-appe and B-appe

monomers and the properties of the resulting

polymers were studied in comparison with

Scheme 37. Synthesis and curing of propargyl ether-functional benzoxazines [196]. (Reprinted from Macromolecules by permission, q2001

American Chemical Society).


4-phenyl-3,4-dihydro-2H-1,3-benzoxazine (P-a) and

bis(4-phenyl-3,4-dihydro-2H-1,3-benzoxazinyl)iso-

propane (B-a) as typical BZ monomers without

propargyl groups. The DSC cure of both P-appe and

B-appe showed a single exotherm corresponding to

the ring-opening polymerization of oxazine ring and

cross-linking of aryl propargyl ether group, at

almost the same temperature range as for P-a and

B-a. The Tg values of polybenzoxazines derived

from propargyl-containing monomers, PP-appe and

PB-appe, were higher by ,100 and 140 8C,

respectively, than for typical polybenzoxazines

without propargyl groups. The storage moduli of

Polymers of P-appe and B-appe were maintained

constant up to a higher temperature in comparison

to polymers of P-a and B-a. In other words,

propargylation led to enhanced thermal stability of

these structurally modified polybenzoxazines. The

T5% for poly(B-a) shot up from 342 to 362 8C for

poly(B-appe) and the increase in corresponding

char-yield was from 44 to 66%.

8.3. High molar-mass PN resins

PN resins, bearing varying extent of propargyl ether

groups were synthesized from high molecular weight

novolac and propargyl bromide [197]. The cure was

followed via DSC in the temperature range 165–

330 8C. The activation parameters for cure determined

by the integral method of Coats-Redfern are compiled

in Table 36. The activation energy for curing the

propargyl novolac was substantially higher than that

for model bispropargyl ether compounds and increased

marginally with the degree of functionalization.

The heat of curing increased proportional to the degree

of substitution. These were cured in the temperature

range 180–220 8C under isothermal condition.

The cure profile, extrapolated from non-isothermal

DSC kinetics studies was in league with the results

from DMA studies. The cure completion was

ascertained from the complete disappearance of the

xC–H group absorption at 3272 cm21 in FTIR. The

mechanical properties of the silica laminate of the

resins of varying propargyl-content revealed good

consolidation of the interphase, evident from the

initial gain in both interlaminar shear strength (ILSS)

and flexural strength with increase in the degree of

propargylation. However, the benefit of the better

resin-reinforcement interaction was not retained on

cross linking the resin further, whereon the composite

failed by a combination of fiber debonding and brittle

Table 36

Cure characteristics and related kinetic parameters of novolac-propargyl ethers [197] (Reprinted from Polymer International by permission,

q2001, Society of Chemical Industry)

Polymer

reference

Extent of

propargylation

(mol%)

DSC cure parameters DH Kinetic parameters

Ti (8C) Tm (8C) Te (8C) J/g kJ/mola E (kJ/mol) A £ 1026 (s21)

PN 18 18 165 243 330 244 27.5 85.4 3100

PN 45 45 173 245 335 630 77.5 103.9 2500

PN 54 54 174 249 320 683 86.5 109.1 8.8

PN 82 82 175 249 320 1086 148.9 118.6 10.7

Ti; cure onset temperature, Tm; maximum, Te; cure end temperature in DSC.a Per repeat unit.

Table 37

Mechanical properties of UD composites and laminates of PN resins

(glass reinforcement) [197] (Reprinted from Polymer International

by permission, q2001, Society of Chemical Industry)

Polymer

reference*

SBSS

of UD

composite

(MPa)

Mechanical properties of laminates

(MPa)

ILSS (MPa) Flexural

strength

(MPa)

Compressive

strength (08)

(MPa)

PN-18 35 20 169 248

PN-45 37 22 220 258

PN-54 51 23 167 262

PN-82 34 20 143 –

PN-82/silica – 18 136 –

PN-82/carbon – 38 562 –


fracture of the matrix. The properties of the

composites are given in Table 37. The mechanical

properties of the composites were not very good.

Changing the reinforcement to carbon enhanced the

properties, but were still lower than the industrial

standards. This was attributed to the fact that the PN

resins used possessed high molecular weight, pow-

dery in nature, whose fiber wetting capabilities might

not be very good.

8.4. Thermal degradation behavior

PN resins possess good thermal stability. The

thermograms of the resins with different degree of

propargylation along with that of cured resole are

shown in Fig. 30. TGA showed apparently single step

decomposition starting above 380 8C and ending at

around 640 8C for the cured PN resins. The decompo-

sition parameters obtained from TGA thermograms

are compiled in Table 38 [198]. The thermal stability

of the PN resins with respect to IDT ðTiÞ and char

residue at high temperature is significantly higher than

that of conventional resole systems. Whereas resole

starts to decompose below 300 8C, the decomposition

is triggered only at 380 8C for the PN resins. This

advantage in thermal stability must be a consequence

of the protection of the OH groups by etherification

that reduces the susceptibility of the methylene

protons for degradation. It was found that even

minor degree of propargylation was conducive for

boosting the Ti values by about 70 8C. Although

decomposition is initiated only at higher temperature,

the degradation is found to be quite rapid for all the

systems. Despite the increased cross-link density with

enhanced extent of propargylation, TGA manifested

nearly identical Ti values. Contrary to expectation, the

degradation becomes rapid and the char-yield

decreased with further increase in propargyl-content

and cross-link density. The rapidity is clearly

manifested in the systematic drift of the Tm to lower

temperature as the degree of propargylation increases

(Table 38). This unexpected behavior can be ascribed

to two reasons. As the propargyl-content increases,

the cured polymer possesses more aliphatic groups.

The enhanced cross linking achieved through the

aliphatic groups may not be conducive for increasing

the thermal stability, as these links are thermally

fragile. Moreover, the PN polymers synthesized at

different extent of propargyl etherification is likely to

possess varying substitution pattern, as shown in

Scheme 38.

At low degree of substitution, the propargyl

etherification should occur preferentially at the least

steric hindered terminal phenol groups and the ortho–

para substituted ones (Structure A). Then, the priority

of substitution is in the ortho–ortho methylene

substituted phenols (Structure B) and the least

preference should be for the 2,4,6-tris methylene

substituted phenyl molecules (Structure C, which acts

as branching position on the novolac backbone). The

thermal reactions of phenyl propargyl ether groups

Fig. 30. Thermograms of propargyl novolac polymers of different

propargyl-content in N2. Heating rate 10 8C/min. The number

denoted the extent of propargylation on the novolac [198]

(Reprinted from Journal of Macromolecular Science Pure and

Applied Chemistry by permission, q2003 Marcel Dekker).

Table 38

Thermal decomposition characteristics of PN resins (TGA) [198]

(Reprinted from Journal of Macrolecular Sciences Pure and Applied

Chemistry by permission, q2003 Marcel Dekker)

Polymer

reference

TI (8C) Tm (8C) Te (8C) Char

at 600 8C (%)

Resole 320 380 650 68

PN-18 390 510 640 74.3

PN-45 385 450 640 72.6

PN-54 380 360 650 70.3

PN-82 405 350 650 68.1


of different substitution pattern have been studied

[199–201]. The same mechanism can be extended to

propargyl novolac polymers. The type (A) propargyl

ether can rearrange to the chromene prior to

polymerization giving rise to comparatively thermally

stable cyclic structure. Type B giving the cyclic

ketone, and type C giving the polyene, are thermally

fragile entities. The curing reactions of various

propargyl groups are included in Scheme 35. The

cross-links generated by thermally fragile groups (by

structures resulting from B and C type substitutions)

undergo easy thermal degradation. Evidence for the

formation of the ketone structure (from Type B

structure) was obtained in the FTIR spectrum with a

broad absorption around 1740 cm21. This could

account for the rapid thermal degradation and lower

char-yield for the high-propargylated PN systems.

Thus, the thermal stability is found to be good only at

Scheme 38. Likely substitution pattern in propargyl novolac polymers at different degree of propargylation [198] (Reprinted from Journal of

Macromolecular Science, Pure and Applied Chemistry by permission, q2003 Marcel Dekker.


medium cross-linking. It may be remarked that these

addition curable PN re sins loose only ,2% mass on

curing, whereas resole looses 22–26% mass at the

initial stages of curing. In short, the overall char

residue of PN resins is superior to that of resole when

a comparison is done based on virgin resins.

8.5. Propargyl ether resins based on oligomeric

novolac

PN resins based on low molar mass, oligomeric

novolac (Mn , 300–600) provided tacky, flowing

resins (OPN resins) with viscosity in the range

2000–6000 mPa s [202]. The propargylation was

limited to about 85% as this was shown to result in

good thermal stability. A GPC analyses confirmed that

the resin possess oligomers other than monomeric

phenyl propargyl. The identical distribution pattern for

both the precursor novolac and the propargyl ether

indicated a uniform propargylation for all molar mass

species. The GPC pattern for the precursor and the PN

resin shown in Fig. 31 confirm this. The viscosity of the

resin decreased significantly on raising the temperature

as shown in Fig. 32. The temperature coefficient of

viscosity ðbÞ was obtained as 10.

These resins are suitable for solvent-free impreg-

nation. The resin cures at 200–220 8C, evident from

the DMA of the resin. The non-isothermal DMA

spectrum (shown in Fig. 33) was used to evaluate the

gel time and cure time. The resin properties are given

in Table 39. The time for stagnation of storage

modulus G0 or complex viscosity hp is indicative of

the cure time, and this decreases with increase in

temperature as shown in Fig. 34. The gel time is

obtained from the cross point of G0 and the loss

Fig. 31. GPC pattern for (– – –) novolac and (—) oligomeric

propargyl novolac (OPN) resins [202].

Fig. 32. Variation of viscosity with temperature for oligomeric

propargyl novolac polymer [202].

Fig. 33. Non-isothermal DMA of PN resin in air. Heating rate

5 8C/min. Parallel plate rheometry, controlled stress at 100 Pa. 4 Hz.

[202].


modulus G0; as typically shown for one case in Fig. 35.

Since the cross-point is frequency-dependent, the gel

points are only apparent. The gel time and cure times

at different temperatures, determined by rheometry

are given in Fig. 36. Both decrease with increase in

temperature. The glass laminate showed improved

mechanical performance. The resin has been success-

fully used for developing reaction-bonded SiC-based

ceramic components [203].

8.6. Propargyl novolac–epoxy blend

Partial propargylation of novolac permitted

co-reaction of the remaining OH groups with an

epoxy resin. Thus, partially propargylated oligomeric

novolac resin with an extent of propargylation around

70% was used to formulate a dual cure thermoset

when blended with a novolac epoxy resin [204]. The

dual cure through phenol–epoxy reaction and pro-

pargyl curing was evident both in DSC and DMA, as

shown in Figs. 37 and 38, respectively. In DSC,

Table 39

Typical thermal and physical properties of oligomeric PN resins

[202]

Sp gravity 1.16 g/cc

Extent of propargylation 85%

GPC molar mass Mn ¼ 600; Mp ¼ 2500;

Mw=Mn ¼ 5

Cure initiation 170 8C ( ex: DSC)

Cure completion 300 8C

DH of curing 1.2 kJ/g

h at 50 8C 3 Pa s

Temperature coefficient

of hðbÞ

10

Fig. 34. Cure profiles for propargyl novolac resin in air at different

temperatures by parallel plate rheometry in controlled strain mode.

[202].

Fig. 35. Determination of apparent gel time for oligomeric

propargyl novolac at 170 8C by rheometry [202].

Fig. 36. Dependence of gel time and cure time on temperature for

oligomeric propargyl novolac resin [202].


the epoxy–phenol cure peak occurs at ,135 8C and

that due to propargyl cure at ,235 8C. The phenol–

epoxy reaction could be catalyzed by TPP, but this did

not have any effect on the curing of propargyl ether

groups. DMA showed more or less the same cure

temperature regime, although the epoxy cure occurred

at a slightly higher temperature, due to the absence of

TPP. The DMA profile indicated the propargyl cure occurred in multi-steps. The cure sequences for the

resin system are shown in Scheme 39.

The isothermal DMA at 220 8C indicated that the

propargyl curing requires about 2 h at this tempera-

ture for reasonable network build-up. This is in

league with the observation for the pure PN resin

(Fig. 34). Table 40 lists the glass laminate properties

of the PN resin cured with different equivalent of

EPN resin [204]. The mechanical properties are not

significantly dependent on composition. As a whole,

the properties are better than those of the high molar

mass PN polymers. Presence of epoxy reduces the

gel time and makes the processing easier. The TGA

of the blend (in Fig. 39) showed a decreasing

thermal stability for the system on enhancing the

epoxy-content. Addition of epoxy diminishes both

the IDT ðTiÞ and the anaerobic char residue. The

char residue at 750 8C is inversely proportional to

the epoxy-content. DMA of the cured resin showed

Tg . 300 8C for the PN systems. Tg decreased in the

presence of epoxy. The adhesive characteristics of

the PPN–epoxy blend were also investigated.

Fig. 38. Non-isothermal DMA spectrum of epoxy/propargyl

novolac blend in presence of triphenyl phosphine. Heating rate

5 8C/min, (– – –) E00; (—) E0 [204] (Reprinted from Polymer and

Polymers Composites by permission, q2004 Rapra Technology).

Scheme 39. Dual cure sequences of partially propargylated

oligomeric propargyl novolac (OPN)-epoxy blend [204]. (Reprinted

from Polymers and Polymer Composites by permission, q2004

Rapra Technology).

Fig. 37. DSC of epoxy/propargyl novolac blends in presence of

triphenyl phosphine. Heating rate 10 8C/min [204] (Reprinted from

Polymer and Polymers Composites by permission, q2004 Rapra

Technology).


The variation of LSS with phenol–epoxy ratio is

shown in Fig. 40 for EPN and in Fig. 41 for

DGEBA. The graphs also show the percentage

retention of LSS at 100 8C. The adhesive properties

of the PPN resin were significantly enhanced by the

reaction with epoxy. The optimum adhesive strength

and retention of properties at high temperature were

observed with a phenol–epoxy equivalent ratio of

2:1 in both cases. A small proportion of epoxy

boosted the adhesion significantly. The strength was

higher for DGEBA due to its better flexibility and

lower cross-link density in this case, in contrast to

the rigid EPN system. However, the high tempera-

ture retention was nearly identical for these systems.

9. Phenolic resins with terminal acetylene groups

Addition curable phenolic resins, bearing terminal

ethynyl groups, anchored to a benzene ring through a

phenyl azo linkage (ethynyl phenyl azo novolac,

EPAN), were realized by a novel and simple synthesis

strategy involving the coupling reaction between

novolac and 3-ethynyl phenyl diazonium sulfate

[205]. The synthesis is shown in Scheme 40. The

diazo coupling was limited to the para position of the

novolac and occurred to a maximum of 50 mol%.

The molar mass, determined from GPC showed a down-

ward drift with increase in degree of functionalization.

Table 40

Glass laminate property of OPN (extent of propargylation 70%) and OPN-epoxy blend [204] (Reprinted from Polymers and Polymer

Composites by permission, q2004 Rapra Technology)

Polymer ref. Composition,

(OH/EPN

equivalent

ratio

ILSS

(MPa)

Flexural

strength

(MPa)

Compresive

strength

(MPa)

Tg (8C)

PNEX-100 100:0

(OPN)

30 390 165 .350

PNEX-8515 85:15 28 320 140 .350

PNEX-6535 65:35 30 375 140 300

PNEX-5050 50:50 25 340 174 270

Fig. 40. Variation of lap shear strength and its retention at 100 8C

with phenol–epoxy ratio for partially propargylated novolac

(PPN)/epoxy novolac system (EPN). [204] (Reprinted from

Polymers and Polymer Composites by permission, q2004 Rapra

Technology).

Fig. 39. TGA of propargyl novolac–epoxy blend of varying

composition(number denotes propargyl–epoxy ratio) in N2. (—)

PNEX-100, (– – –) PNEX-8515, (· · ·) PNEX-6535, (–·–·) PNEX-

5050, Heating rate 10 8C/min.[204] (Reprinted from Polymers and

Polymer Composites by permission, q2004 Rapra Technology).


This was attributed to the decreased hydrodynamic

volume of the polymer resulting from its comb-like

structure. This was confirmed form the trend in

intrinsic viscosity of the resins as a function of the

extent of diazo coupling. The molecular character-

istics of polymers with varying extent of functiona-

lization are given in Table 41.

9.1. Curing of EPAN resins

These resins showed a broad cure exotherm in

DSC in the range 140–240 8C due to the curing of

acetylene functions occurring by various reactions.

Like other acetylene-polymers these are also

hardened by a series of reactions, including

trimerization [206], Glaser coupling [207], Strauss

coupling [208], Diels-Alder reaction with partici-

pation of products of Glaser coupling and Strauss

coupling, Diels-Alder coupling with the aromatic

backbone, free radical polymerization with the

formation of linear and branched products [209],

etc.. The various mechanisms are depicted in

Scheme 41.

Spectral and GPC studies with model compound

(3-ethnyl phenyl azo phenol, EPAP) gave evidence

for Glaser coupling, trimerization and linear addition

reactions. Fig. 42 shows the high resolution GPC

profiles of the soluble part of the polymerization

product of EPAP, isolated in methanol and DMF.

Unreacted EPAP appeared at an elution time of

23.5 min. The component at 22 min is the dimer

formed by Glaser and Strauss coupling. The one at

20.1 min with approximately thrice the molar mass of

the monomer is attributed to the cyclic trimer. GPC

also showed the presence of a minor amount of higher

molar mass species in this fraction. The methanol

insoluble part (soluble in DMF) showed a multimodal

distribution and indicated the presence of EPAP. The

broad peak centered at 20 min encompasses the low

molar mass oligomers formed by various mechan-

isms. The high molar mass products appeared as well-

separated peaks at 13, 11 and 10.5 min. This part, in

all probability, contained the linearly polymerized

polyenes and branched polymers anticipated with the

proposed cure mechanism.

9.2. Thermal characteristics of EPAN resins

The polymers exhibited enhanced thermal stability

and anaerobic char in comparison to resole. The

thermal stability and anaerobic char-yield of the

polymers increased with enhanced cross-link density.

Against a char residue of 60–62% in resole, EPAP

gives 72–75% char at 700 8C. If the mass-loss during

curing is also considered, the net gain in char is

about 70% more than in resole. Isothermal pyrolysis

Fig. 41. Variation of LSS and LSS retention at 100 8C with

phenol–epoxy ratio for partially propargylated novolac (PPN)/

bisphenolA diglycidyl ether system (DGEBA). [204] (Reprinted

from Polymers and Polymer Composites by permission, q2004

Rapra Technology).

Scheme 40. Synthesis of ethynyl phenyl azo phenolic (EPAN) resin

[205] (Reprinted from Polymer by permission, q2002 Elsevier

Science).


at 700–900 8C showed that complete pyrolysis of the

EPAN system is not achievable under these con-

ditions. The TGA of resins with different acetylene-

content are shown in Fig. 43.

The evolution in elemental composition under

different pyrolysis conditions implied that the pyrol-

ysis occurs mainly by loss of nitrogen and hydro-

carbon. The higher proportion of char shows

the prospects for potential application of this resin

in ablative compositions and in carbon/carbon

composites. Non-isothermal kinetic analysis of the

degradation reaction confirmed that degradation

apparently occurred in a single kinetic step, with a

decomposition rate almost independent of cross-link

density [210]. These polymers yielded a considerably

higher proportion of char, whose XRD analysis

Scheme 41. Various cure possibilities for acetylene groups (uncatalysed) [205] (Reprinted from Polymer by permission, q2002 Elsevier

Science).

Table 41

Molecular characteristics of EPAN systems [205] (Reprinted from Polymer by permission, q2002 Elsevier Science)

Polymer ref. Extent of azo

coupling (mol%)

½h� in THF,

30 8C (dl g21)

Molecular weight by GPC,

(g mol21)

½h� £ Mn (dl mol21)

Mn Mw Mp

Novolac 0 0.155 700 1970 1840 108.5

EPAN-1 24 0.134 520 1570 2030 69.68

EPAN-2 35.8 0.085 430 1460 1950 36.55

EPAN-3 43 0.107 350 1240 1330 37.45

EPAN-4 49.7 0.074 330 1180 1110 24.42


confirmed the presence of partial crystalline character,

unlike the case for a conventional phenolic resin.

10. Phenolic resins with phenyl ethynyl groups

Replacement of terminal acetylene groups by

phenyl ethynyl function gives scope for improving

further the thermal stability of phenolic networks

due to the higher aromatic-content of the cross links.

Phenyl ethynyl groups have, of late, received a great

deal of attention as a means of thermally chain

extending and cross-linking polymers [211,212]. On

thermal curing, they provide a three-dimensional

network exhibiting an excellent combination of pro-

perties including high glass transition temperature,

good thermal stability, moisture- and solvent resis-

tance, good toughness and mechanical properties.

10.1. Phenyl ethynyl functional addition-curable

phenolic resins

Phenyl ethynyl functional phenol–formaldehyde

(novolac-type, PEPFN) addition curable resins were

synthesized by reacting a mixture of phenol and

3-(phenylethynyl)phenol (PEP) with formaldehyde in

Fig. 42. GPC traces of (—-—): ethynyl phenyl azo phenol polymer,

(—--—): Methanol soluble fraction, (—): methanol insoluble

fraction in DMF, UV (280 nm) detection [205] (Reprinted from

Polymer by permission, q2002 Elsevier Science).

Fig. 43. TGA thermograms of cured ethynyl phenyl azonovolac

resins in N2. Heating rate: 10 8C/min. (—): EPAN 1, (- - -): EPAN 2,

(— — —): EPAN 3, (-·-·-): EPAN 4, (—--—--—): Resole [205]

(Reprinted from Polymer by permission, q2002 Elsevier Science).

Fig. 44. Monomer conversion with time for the polymerisation

reaction of phenyl ethynyl phenol and phenol with formaldehyde

[213] (Reprinted from Journal of Applied Polymer Science by



the presence of an acid catalyst [213]. The poly-

merization reaction was performed at 75 8C, and the

reaction time was optimized by monitoring the

reaction mixture by GPC at different time intervals.

The molecular weight increased and the distribution

became broader as the reaction progressed. The

relative concentrations of the polymer and the starting

phenol reactants were determined by measuring the

area under the GPC curves corresponding to each

component. The product evolution with time is shown

in Fig. 44. The polymer was formed in about 74%

yield after 10 h. Relatively narrow molar-mass

distributed polymers were obtained in good yield.

The copolymer composition nearly matched the feed

composition. The presence of PEP led to reduced

molar-mass and narrow distribution of the copoly-

mers. The polymer properties are given in Table 42.

10.1.1. Cure and thermal characteristics of PEPFN

resins

The resin underwent thermal curing at around

250–275 8C, and the cure optimization was done by

isothermal DMA at 275 8C. The cure time of 1 h is

significantly lower than the cure time for conventional

phenyl ethynyl functional polymers. This points to an

altered cure mechanism in these resins. The cure

mechanism is proposed as a combination of acetylene

addition [214] and by addition of phenol to the triple

bond as implied in a model study [215]. Based on

these, the cure mechanism depicted in Scheme 42 was

proposed. The cure chemistry is not well understood.

In many cases, the polyene formation by linear

addition and cross linking is a widely accepted

Table 42

Composition and molecular weight characteristics of the PEPF

resins (conversion: 74–80%) [213] (Reprinted from Journal of

Applied Polymer Science by permission, q2003 John Wiley and

Sons)

Polymer

ref.

PEP in

feed

(wt%)

PEP in

copolymer

(wt%)

Molecular weight by GPC

Mn Mw Mp Mw=Mn DPn

PEPF 25 25 24 780 3220 1600 4.1 6.5

PEPF 50 50 49 820 2720 1400 3.3 5.9

PEPF 75 75 73.8 620 1510 860 2.5 4.4

PEPF100 100 100 650 1270 960 1.9 3.2

Scheme 42. Synthesis and likely cure mechanism for phenol–phenyl ethynyl phenol [213] (Reprinted from Journal of Applied Polymer Science



mechanism for the curing of phenyl ethynyl-contain-

ing polymers [216].

The thermal stability and anaerobic char residue of

the cured resins increased proportionate to the

phenylethynyl-content, and these properties were

improved over those of resol resin. These addition

cure phenolics provided an overall increase in char of

about 70% vis-a-vis resol resin when compared on the

basis of the uncured resins. The thermograms of resins

with varying degree of phenyl ethynyl functionaliza-

tion are given in Fig. 45. The variation of IDT and

anaerobic char with the phenyl ethynyl-content is

shown in Fig. 46.

10.2. Condensation–addition cure phenyl ethynyl

phenolic resins

The PEPFN resins have the disadvantage of the

need for a very high cure temperature. This causes a

processing difficulty in that the resin tend to bleed

away during molding at high temperatures. This can

be avoided by conferring a partial condensation

character to the resin, by synthesizing phenyl ethynyl

functional resoles (PEPFR). The curing of resole at

lower temperature facilitates the early gelation of the

resin. Thus, phenolic resins bearing methylol and

phenyl ethynyl functions and curing by a dual

mechanism of both condensation and addition were

synthesized by the reaction of 3-(phenyl ethynyl)

phenol (PEP) with formaldehyde under alkaline

conditions [217]. The synthesis and curing are

shown in Scheme 43.

Resins with varying relative concentration of the

two functional groups were synthesized and

characterized. GPC and NMR analyses confirmed

that the resin contained a mixture of multi-methylol

substituted phenols. It also contained some proportion

of partially chain extended molecules. A typical high

resolution GPC profile (of resoles derived from PEP)

showing the product distribution is shown in Fig. 47.

The molecular characteristics and the cure charac-

teristics of the resoles are given in Tables 43 and 44,

respectively. The methylol condensation occurred at

practically the same temperature range as for

conventional phenolic resins (,130–160 8C). The

ethynyl cure occurs at a lower temperature than the

model compound. A two-stage cure was confirmed in

both DSC and DMA analyses. The DMA spectrum is

shown in Fig. 48. The low temperature cure due to

methylol condensation was conducive to early gela-

tion of the system, at around 100–150 8C. The

ultimate curing through addition reaction of phenyl

Fig. 45. TGA traces of cured phenol–phenyl ethynylphenol–

formaldeheyde (PEPFN) polymers in N2. Heating rate 10 8C/min

[213] (Reprinted from Journal of Applied Polymer Science by


Fig. 46. Variation of Ti and char-yield (at 700 8C) for phenol–

phenyl ethynylphenol–formaldeheyde polymers with phenyl ethy-

nyl-content.


ethynyl group, however, required heating at 275 8C

for 2 h.

The cured resins exhibited better thermal stability

and anaerobic char residue in comparison to a

conventional resole. The thermal stability and char-

yielding property showed a diminishing trend with

enhanced methylol substitution. Resin with an F/P

ratio less than unity offered excellent thermal

stability, and anaerobic char-yield significantly higher

than that of the corresponding novolac (PEPFN).

Methylene groups favored the initial degradation,

whereas the higher temperature carbonization process

was independent of the network structure.

11. Comparative thermal property of PMF, PN,

EPAN and PEPFN resins

On comparing the thermal stability of the addition-

cure phenolic resins belonging to PMF, PN, EPAN

and PEPFN class as a function of composition, it was

observed that in many cases, thermal stability and

char-yield increase with increased cross linking via

enhanced functionalization. Exceptions were noted in

the case of the blend of PMF with allyl novolac and

for propargylated novolac (PN resins). In these two

cases, the thermal stability decreased with cross-

linking due to the increase in the aliphatic-content in

the cured polymer. The PMF resins exhibited least

thermal stability and the EPAN systems were the most

thermally stable. The maximum char-yield was

obtained for EPAN and PEPFR (with low methylol)

concentration. The comparative thermograms can be

found in Fig. 49 for resins with maximum thermal

stability in each series [218].

12. Phenolic–triazine resin (P–T resins)

Phenolic triazine (P–T) precursor resin is a

reaction product of novolac resin and cyanogen

Scheme 43. Synthesis and curing of phenol–phenyl ethynylphenol–formaldeheyde (PEPFR) resin (PEPFR, resole type) [217] (Reprinted from

Journal of Applied Polymer Science by permission, q2002 John Wiley and Sons).

Fig. 47. High resolution GPC of phenyl ethynylphenol–formalde-

heyde resole (PEPFR) resin.


halide. P–T network is formed by the thermal

cyclotrimerization of the cyanate ester of novolac as

shown in Scheme 44 [219,220]. It is an ideal matrix

system for composites, because it combines

the processibility convenience of epoxies and the

thermal capabilities of polyimides and fire resistance

of phenolics. The absence of volatile by-products

during cure renders them attractive matrices for void-

free moldings and composites.

12.1. Features of P–T resins

The key features of P–T resins are

† Low melt viscosity, epoxy like handling, resinous

consistency

† Long gel time and angle of flow, long-term thermal

stability

† High solubility in low boiling solvents

† Feasibility of mixing: dry, wet and impregnation,

ease of formulation

† Wide processing latitude, RTM processable

† Tg up to 399 8C depending on post cure

conditions

† Ultimate elongation and mechanical properties

equal to high performance polyimide systems

† Very low toxicity

† Low ðDkÞ and low moisture absorption and thermal

expansion

† Better thermo-oxidative stability than phenolics,

flame resistant and low smoke generation. Stability

comparable to polyimides

Table 44

Cure Characteristics of PEPFR resins from DSC. All temperatures are in 8C [217] (Reprinted from Journal of Applied Polymer Science by

permission, q2002, John Wiley and Sons)

Polymer reference Formaldehyde/

phenol ratio

Methylol condensation Ethynyl cure

DH (J/g) Ti Tm Te Ti Tm Te

PEPFR-1 1 91.3 170 196 230 260 300 320

PEPFR-1.5 1.5 100.9 170 198 230 270 320 340

PEPFR-2 2 142.3 160 202 220 255 280 325

PEP 0 – – – – 300 320 350

Fig. 48. DMA of phenyl ethynylphenol–formaldeheyde resole resin

showing two-stage curing [217]. (Reprinted from Journal of Applied


Table 43

Characteristics of PEPFR resins [217] (Reprinted from Journal of Applied Polymer Science by permission, q2002, John Wiley and Sons)

Polymer

reference

Formaldehyde/

phenol ratio

(F/P)

Relative ratio

–CH2–O–/–CH2-

from NMR

Different components (%) in PEPFR resin from GPC

Monomer Mono and

dimethylol

product

Dimer and

higher oligomers

PEPFR- 1 1 2.7 14.6 49 36.4

PEPFR- 1.5 1.5 3.5 18.7 50.3 31

PEPFR -2 2 5 14.9 54.7 30.4


P–T resin offers considerable processing flexibility

since their consistency ranges from low viscous liquids

to semi solids, with gel temperatures that can be tuned

by catalysis using a host of materials. However, the

ultimate cure temperature has to be high (.300 8) to

achieve optimum cure and higher Tg (above 300 8C). In

fact Tg is a function of the ultimate cure temperature as

shown in Fig. 50. Post curing at temperatures above

300 8C ensures Tg of about 400 8C [221]. PT resins

possess better thermo-oxidative stability and char-

yield than conventional phenolics, because they are

mostly cross-linked by triazine groups. The proximity

of the hydroxyl groups in phenolic resin renders these

methylene bridges thermo-oxidatively fragile, and the

degradation process is accelerated by the number of

dihydroxy phenyl methylene groups. The degradation

process of phenolic resins has been discussed in length

[222]. The widely accepted mechanism for oxidative

degradation is as shown in Scheme 45. The hydroxy

phenyl methylene group is the triggering point.

PT resins, on the other hand, are cross-linked mostly

by triazine phenyl ether linkages, which confer both

thermo-oxidative stability and toughness to the system.

The evidence for better thermo-oxidative stability is

obtained from the thermal behavior of the systems in

both air and inert atmosphere. The initial decompo-

sition pattern in air and in inert atmosphere essentially

superimpose, pointing to a non-oxidative mechanism

of degradation for PT systems (see the thermograms in

Fig. 51).

The decomposition starts at around 420 8C and the

char-yield is of the order of 65–70%. The char retains

about 5% nitrogen. This char-yielding quality implies

better prospects for application of this type of resin for

thermo-structural uses in aerospace, in place of

conventional phenolics. In fact, laser ablation studies

on a series of ablatives including PT resin have

confirmed their potentiality for such applications.

It was found that the ablation energy was highest for

the cyanate polymers, on exclusion of phthalocya-

nines, among the polymers tested [223]. Some

ablative formulations for rocket nozzle applications

contain P–T resins as one of the components [224].

Fig. 49. Comparative thermograms of various addition-curable

phenolic resin in N2, heating rate 10 8C/min. Phenol–phenyl

ethynylphenol–formaldeheyde novolac (PEPFN), phenyl ethynyl-

phenol–formaldeheyde resole (PEPFR); phenol–maleimidophe-

nol–formaldehyde (PMF); propargyl novolac (PN); Ethynyl phenyl

azo novolac (EPAN) and resole [218] (Reprinted from Polymer

Degradation and Stabilisation by permission, q2001 Elsevier

Science).

Scheme 44. General synthesis route for novolac cyanate and curing to phenolic–triazine.


PT resins possess the general characteristics of

cyanate esters with added thermal stability. The

relative strain capabilities and thermal performances

of some common thermosets are projected in Fig. 52.

The moderate strain capability and high temperature

capability conveys the comparative advantage of PT

system vis-a-vis the other thermosets.

The low moisture absorption, low dielectric

properties and better hygrothermal performance

project their potential for applications in many critical

aerospace structures. The low moisture absorption and

consequent reduced dimensional changes render it the

matrix of choice for composites for optical support

structures for satellite applications. Fig. 53 compares

the relative hygrothermal properties of common

thermosets, among which, the PT systems are seen

to occupy a comfortable position.

Scheme 45. Thermal degradation mechanism of phenol–formaldehyde resins.

Fig. 50. Dependence of Tg on cure temperature for phenolic–triazine

(P– T) resin [221] (Reprinted from Technical Brochure by

permission, q1996 Lonza Ltd).


12.2. Properties of P–T systems

P–T resin is commercially available under the

trade name Primaset P-T-15, P-T-30, P-T-60 and

P-T-90, which essentially differ in their molar

masses [221]. The characteristics of the resin are

listed in Table 45.

The gel time can be tuned by catalyst concen-

tration. The generally recommended catalysts are zinc

octoate/nonyl phenol, cobalt naphthenate, copper salts

etc. The neat resin properties of cured PT systems are

given in Table 46. Although PT possesses a very high

Tg; its neat resin mechanical properties are inferior to

those of other common thermosets. On the other hand,

the PT composites exhibit excellent mechanical

properties and thermo-mechanical profile, as is

evident from the typical data displayed in Table 47.

Although the properties diminish at high temperature,

Fig. 51. Thermograms of cured phenolic–triazine (PT-30) resin in

air and N2. Heating rate 10 8C/min [202].

Fig. 52. Comparative temperature- and strain capabilities of

common thermosets [238] (Reprinted from Advances in Polymer

Science by permission, q2001 Springer Verlag).

Fig. 53. Comparative hygrothermal performances of common

thermosets [238] (Reprinted from Advances in Polymer Science

by permission, q2001 Springer Verlag).

Table 45

Characteristics of different PT resins [221] (Compiled from

Technical Brochure by permission, q1996 Lonza Ltd)

Property P-T-15 P-T-30 P-T-60 P-T-90

Consistency Viscous

liquid

Viscous

liquid

Semi-

solid

Powder

Sp. gravity

(g/cc)

1.25 1.28 1.24

Viscosity

(mPa s),

93 8C

1.245 200 25000 –

Viscosity

(mPa s),

121 8C

3 80 1500 –

Gel-time

(min, 200 8C,

uncatalysed)

2 .20–30 20

Gel-time

(min,

catalyzed)

.30

(200 8C)

,5

(177 8C)

,5

(177 8C)

–


the values are still in the acceptable limits for many

structural applications. Their ageing characteristics

are also excellent (refer to data in Table 48). Das has

compared the thermal and mechanical properties of

phenolic and P–T resins [225]. Almost all of the data

in Table 49 speak about the superiority of the latter

system for both mechanical and thermal properties.

The high char-yielding property is conducive to better

flame resistance in actual fire situations. The LOI

comparison given in Fig. 54 shows its superiority even

to phenolics.

The P–T resin systems have been successfully

employed in filament winding of cylindrical structures

such as pressure bottles, which retain 83% of their

room temperature properties at 288 8C [219,226]. The

flexural strength of PT/T-300 rings is 339 MPa, which

only drops to 284 MPa at 288 8C. The apparent

modulus drops from 98.6 to 93.8 GPa. The mechanical

Table 46

Typical neat resin properties of PT resin [221] (Compiled from


Property Value

Tg (8C) Up to 400

Dk (1 MHz) 3.1

Dissipation factor 0.007

Ti (from TGA, 8C) 410–450

Char-yield (%) 65–70

CTE (8C21, 40–315 8C) 2.8 £ 1025

Tensile strength (MPa) 41.4

Tensile modulus (GPa) 4.07

Elongation (%) 2

Flexural strength (MPa) 110

Flexural modulus (GPa) 4.7

Compressive strength (MPa) 317

Table 47

Typical composite properties of PT resin [221] (Compiled from


Property E-Glass

laminate

Carbon UD composite

24 8C 260 8C 316 8C

SBSS

(MPa)

53.8 184 55.8 50.3

Flexural

strength (MPa)

558 1480 986 882

Flexural

modulus (GPa)

34.8 175 170 160

Table 48

Thermal ageing data on PT/carbon fibre UD composite [221]

(Compiled from Technical Brochure by permission, q1996 Lonza

Ltd)

Tempera-

ture (8C)

Ageing

time (h)

Flexural

strength

(MPa)

Flexural

modulus

(GPa)

SBSS

(MPa)

288 0 1650 124 151

100 1830 132 119

200 1520 123 112

500 1050 137 106

1000 1410 128 86

316 100 1750 120 93.8

200 1650 117 94.5

Table 49

Comparative mechanical and thermal properties of PT and phenolic

resins [225] (Reprinted from 30th Int SAMPE Tech Conf by

permission, q1998 Society for the Advancement of Materials and

Process Engineering)

Property PT resin PF resin (hexa-cured)

Ti (TGA) 8C 450 350

Tg; 8C (DMA) 400 121

Char-yield, % at 1100 8C 66–68 55

Ultimate elongation (%) 2 0.3

Flexural strength (MPa) 95.2 47.6

Flexural modulus (MPa) 4626 2517

Compressive strength (MPa) 306 102

CTE, 1025/8C 22 65

Rockwell hardness (M scale) 125 93

Fig. 54. Comparative limiting oxygen index values of common

thermosets [225] (Reprinted 30th Int SAMPE Tech Conf by

permission by permission, q1998 from Society for the Advance-

ment of Materials and Process Engineering).


properties of PT composites are comparable to those

of PMR-based composites, especially at high tem-

perature (refer to Table 50). Considering their ease of

processibility, PT has an edge over the PMR systems.

Their carbon composites have been experimented as

actuators in turbine engines, with significant advan-

tage over other high performance systems [227]. A

NASA report compared the mechanical properties of

composite and char residues of 27 modified phenolic

resins including PT resins, to those of conventional

phenolic resin [228]. Cyanate, epoxy, allyl, (meth)a-

crylate and ethynyl derivatives of phenolic oligomers

were reviewed. Novolac cyanate along with propar-

gyl–novolac resins provided an anaerobic char-yield

of 58% at 800 8C, whereas a modified epoxy novolac

provided 59% char. A phosphazene derivative was

effective in enhancing the char-yield. The novolac

cyanate, epoxy novolac and methacrylate–epoxy–

novolac were investigated for their composites.

The methacrylate–epoxy–novolac/graphite epoxy

Table 50

Comparative mechanical properties of P–T-carbon fibre composites (fiber: 58–64 vol%) [226]

Property P–T PMR-15

24 8C 330 8C 24 8C 315 8C

Flexural strength (Celion 6000) (MPa) 2482 1379 1930 965

Flexural modulus (GPa) 103–107 110 103–117 103

SBSS (CelionT-650/42) (MPa) 145–165 145–152 – –

SBSS (CelionT-650/42, polyimide sized) (MPa) 69 50 103 50

Scheme 46. Synthesis of cyanate ester and phenolic–triazine resin from cardanol novolac [231].


provided SBSS of 93.3 MPa at ambient conditions

whereas the novolac cyanate provided a value of

74 MPa. Novolac cyanate- and epoxy novolac-based

composites showed good property retention after

ageing at 204 8C for 12 weeks. A recent evaluation

of the composite materials of three high performance

matrices as a low-cost substitute for titanium

components for the control surfaces on hypersonic

aircraft showed the superiority of the PT systems. The

tensile, shear and compressive strength of the three

materials were tested at high temperature. The matrix

materials and test temperatures included (a) Poly-

etheramide (PEAR), 205–260 8C (b) BMI, 260–

427 8C and (c) Phenolic Triazine (PT), 370–538 8C.

While strength values for all materials decreased

substantially (25–100%) with increasing temperature,

the distinct advantage of PT systems were evidenced

in the evaluation [229]. The Lonza Group has claimed

some resin formulations based on P-T systems for

specific applications [230]. In a typical case, a patent

claims storage-stable prepolymer compositions that

are rapidly curable at elevated temperatures. This

comprises of a mixture of prepolymerized and non-

prepolymerized novolac cyanates of specified struc-

ture, useful as resin components for the manufacture

of printed circuit boards or in binders for abrasive

products. The mixture has a refractive index of 1.58,

and contains highly dispersible silicon dioxides and/or

particulate or fibrous fillers. The compositions are

solid at ambient temperature and can be ground to

powders that will flow.

12.3. Structurally modified P–T resins

Several structural modifications of novolac cyanate

have been attempted. Thus, introduction of a flexible

pentadecenyl group in a P–T network decreased the

shelf life of the precursor and the thermal stability of

the cured resin [231]. The resin was synthesized by

cyanation of cardanol–novolac or its mixed novolac

with phenol. The synthesis of the cyanate ester and the

resultant PT resin are shown in Scheme 46. The cross

linked cyanate ester of homo- and copolymers of

4-hydroxy phenyl maleimide (HPM) were syn-

thesized and thermally cured to imido–phenolic-

triazine [232]. The synthesis of a typical copolymer of

HPM, its cyanate ester and the resultant imido–

phenolic – triazine are depicted in Scheme 47.

However, the polymer showed inferior initial

decomposition properties, although the char-yield

was significantly higher. The cyanate ester of PMF

resins was also been synthesized [233]. The imide-

modified novolac (PMF) is described in Section 6.

The cyanate ester underwent a two-stage cure,

implying independent cure of both the cyanate and

maleimide groups. However, the cured imido–

phenolic–triazine exhibited poorer thermal stability

and anaerobic char residue, attributed to the inter-

ference of the rigid imide groups in the char forming

reactions of the triazine groups at higher temperature.

The structures of the cyanate ester and the imido–

phenolic–triazines are shown in Scheme 48. The

scheme has neglected the presence of a very minor

concentration of methylol groups present in the parent

phenolic resin (PMF). A patent claims preparation of

a low molar mass novolac cyanate ester prepolymer

(Mn ¼ 310 g/mol) from the corresponding novolac

resin. The prepolymers are claimed useful as coatings,

adhesives and as matrix in copper clad laminates for

printed circuit boards [234]. There are few reports on

cyanate esters of other polyhydric phenols. Different

grades of poly(4-cyanato styrene) (PCS, normal,

polymer and novolac grades) and copolymers of

4-cyanato styrene with butadiene (PCS-BD) or MMA

(PCS-MMA) as comonomers have been prepared by

Gilman et al. [235]. Flammability tests, performed

using micro scale combustion calorimeter, showed

significant differences in the flammability of the cured

polymers. The flammability decreased with increasing

branching of the cyanatophenyl styrene. The best

results were obtained for novolac grade polycyanato-

phenyl styrene. The structures of the various polymers

are shown in Scheme 49. Copolymer PCS-BD showed

properties similar to poly(cyanato styrene), probably

through cross linking of the unsaturated monomer at

high temperature. The PCS-MMA copolymer showed

the least flame resistance. On a comparative scale, the

PT resins exhibited the best flame resistance. The

thermal properties of these polymers were not

discussed. In a related work, copolymers of styrene

with 4-vinyl phenyl cyanate or 2,6-dimethyl-4-vinyl

phenyl cyanate were prepared via free radical

polymerization [236]. The copolymers were sensitive

to UV light and cross-linked on irradiation with

254 nm UV radiation. Interestingly, the cyanate groups

in the latter copolymer underwent rearrangement


Scheme 47. Cyanate ester of copolymer of hydroxy phenyl maleimide (HPM) and curing to imido–phenolic triazine [232] (Reprinted from

European Polym Journal by permission, q2000 Elsevier Science).

Scheme 48. Imido–phenolic–triazine derived from phenol–maleimidophenol–formaldehyde resin [233]. (Reprinted from European Polym

Journal by permission, q2001 Elsevier Science).


to the isocyanate during irradiation whereas, both the

copolymers yielded cyanurate networks on thermal

curing. Nano composites of PT and clay have also

been reported [237].

12.4. P–T/epoxy blends

For many applications, epoxy resins require

improvement in their properties. Their main drawbacks

in adhesive, coating, PCB and other industrial appli-

cations include, high moisture absorption, high ðDkÞ,

poor fire resistance, high smoke evolution, low Tg; etc.

Modification by way of a co-reaction with P–T resin

improves Tg; hot-wet performances, decreases ðDkÞ,

and moisture absorption and improves fire resistance.

Although many pathways have been proposed for

cyanate–epoxy co-reaction [238], the most accepted

mechanism involves formation of oxazoline groups in

the network as shown in Scheme 50.

The key features and advantages of P-T/epoxy

blends over the epoxy resin systems are [221]:

† Cost competitive

† Enhanced modulus, compressive strength and

hardness

† Improved electrical properties, low moisture

absorption

† Lower cure shrinkage

† Desirable fire, smoke and toxicity properties

† Enhanced Tg over epoxies (180–300 8C)

† No amine curative needed

† Stable materials for RTM hot-melt, and solution

prepregs

Scheme 49. Structures of various grades of cyanato phenyl styrene polymers [235] (Reprinted from 42nd Int SAMPE Symp by permission,

q1997 Society for the Advancement of Materials and Process Engineering).

Scheme 50. General co-reaction between cyanate ester and epoxy

group.


12.5. Thermal degradation of P–T resins

The thermal degradation phenolic triazine has been

investigated, along with that for eight other cyanate

esters [239]. The evolved gases were analyzed by

FTIR and GC/MS. A three-step degradation was

concluded. The steps are (1): random scission and

cross linking of the hydrocarbon backbone between

400 and 450 8C; (2) breakdown of the triazine ring

between 500 and 750 8C, with liberation of low

molecular weight volatile compounds and the for-

mation of a primary residue; and (3) decomposition of

the primary residue between 500 and 750 8C, with the

elimination of alkenes and hydrogen, leaving a

secondary carbonaceous char containing residual

oxygen and nitrogen. Substituted benzene and phenols

were identified in the decomposition products. A

strong peak in the gas chromatogram at m=z ¼ 44 was

attributed to a mixture of CO2 and HOCN. In PT resin,

this amounted to 38.52%. The benzene derivative and

phenol derivatives were 8.7 and 39.8%, respectively.

The rest 10.46% could be due to products such as aryl

cyanates, aryl cyanates, amines, etc.

12.6. Applications of P–T resins

The major areas where PT systems are preferred

are:

† Aero-space composite structures

† Low CME satellite structures and optical support

components

† Cylindrical structures and pressurized bottles

† Actuators in turbine engines

† Thermo-structural applications in nozzles

† Radoms and high signal speed printed circuit

boards

13. Outlook

The foregoing discussion has presented a con-

solidated view of the recent developments in

non-conventional, addition curable phenolic resins.

Phenolic resin still commands considerable research

and industrial interest. Innovative research is focused

on means to address the shortcoming of these

systems in terms of processibility and oxidative

resistance. The introduction of addition-cure pheno-

lics is a partial answer to these problems. The

allylphenol–BMI system is suited for void-free

composite systems for structural applications. In

this case, the advantages in mechanical performances

are usually achieved at the cost of the thermal

capabilities. Although the absolute adhesive strength

is not high, their high temperature retention is good;

rendering them suited for high temperature appli-

cations. Phenol-epoxy systems are excellent for high

strength structural and adhesive applications, but

their low Tg limits the application even at moderately

high temperatures. However, their amenability for

structural modification permits designing systems

with tailored properties. Polybenzoxazines, combin-

ing several salient features required of a high

performance matrix, possess good prospects for

application in several engineering areas. Their

flexibility for molecular design, amenability for

blending and compounding are added advantages.

Poly amide–ether resins derived from bisoxazoline–

phenolics also claim similar design flexibility and

performance profiles, but surprisingly, they are yet to

dominate the market.

Some of the resin systems described show good

prospects for immediate use in composites for aero-

space structural and thermo-structural applications.

P–T resin ranks top among them, with ease of

processibility, excellent thermo-oxidative stability

and thermo-mechanical profile. With good mechan-

ical performance, and high temperature capability

surpassing even the PMR resins, they appear to offer

an immediate solution for many challenging problems

in aerospace composite structural engineering.

Although their neat resin properties are not commend-

able, composites fabricated with them are very strong

and flame-and high temperature resistant. However,

the high cure temperature of P–T systems is an

impeding factor.

The comparative property of various common

thermosets in Table 13 is suggestive of the relative

merits of each system. The structural modifications of

novolacs by incorporation of groups such as ethynyl,

phenyl ethynyl, phenyl maleimide, propargyl etc.,

have been successful in conferring addition curable

nature to phenolics, resulted in enhanced thermal

stability and char-yield. Pending data on mechanical

properties, the industrial utilities of these systems are


yet to be judged. The prohibitive cost and the high cure

temperature of some of these new systems may impede

their easy acceptability in industrial applications.

Acknowledgements

At many phases of the work described here, the

author has associated with his colleagues, R.L. Bindu,

C. Gouri and Dona Mathew. The encouragement and

support received from V.C. Joseph, R. Ramaswamy,

K.N Ninan and K.S. Sastri are thankfully acknow-

ledged. Permissions granted by John Wiley and Sons,

Elsevier Sciences, Springer Verlag, M/s. Lonza Ltd,

SAMPE, SAGE Publications, Kluwer Academics,

Brill Academic Publishers, Rapra Technology and

American Chemical Society for reproduction of data

from their publications are gratefully acknowledged.

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Advances in Addition-cure Phenolic Resins