r06_nandan_crystallization behavior of crystalline-amarphous diblock copolymers

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Crystallization Behavior of Crystalline-Amorphous

Diblock Copolymers Consisting of a Rubbery

Amorphous Block

BHANU NANDAN, JEN-YUNG HSU, AND

HSIN-LUNG CHEN

Department of Chemical Engineering, National Tsing Hua University,Hsin-Chu, Taiwan

Block copolymers show interesting phase behavior if at least one of the constituting

blocks is crystallizable. The morphology development becomes highly complex whensuch a block copolymer is crystallized above the glass transition temperature of theamorphous block due to competition between microphase separation and crystalliza-

tion during the structure evolution process. In this review, we focus on the morphologydevelopment in such crystalline-amorphous (C-A) diblock copolymers where theamorphous block remains rubbery during the crystallization process. Crystallization

behavior in bulk as well as in thin films is considered. The issue of crystal orientationand chain folding upon crystallization in these diblock copolymers has been discussed.

Moreover, the nucleation mechanism in these C-A diblock copolymers and its effect ontheir crystallization kinetics is also described. Some of the emerging areas of research

such as crystallization behavior in blend of C-A diblock copolymers has been brieflydiscussed and finally the future challenges, which holds promise for our further in-depth understanding of crystallization in block copolymers, has been identified.

Keywords block copolymer, microphase-separation, crystallization, chain-folding,homogenous nucleation, crystal orientation

1. Introduction

Block copolymers in which two or more chemically different sub-chains form a single

molecule are a fascinating class of soft materials with unique structural properties.

Their ability to self-assemble into a variety of ordered structures with domain sizes in

the nanometer range has attracted significant attention in recent years.1–5 The self-

assembly of block copolymers is governed by a delicate balance of the interaction

energy and the chain stretching. The repulsive interaction between the chemically

different blocks drives the system to phase separate, whereas the connectivity of

copolymer chains restricts phase separation to a molecular length scale. Hence, the

phase separation in block copolymers is generally known as “microphase separation.”

Received 8 October 2005; Accepted 11 January 2006.Address correspondence to Hsin-Lung Chen, Department of Chemical Engineering, National

Tsing Hua University, Hsin-Chu 30013, Taiwan. Fax: þ886-3-5715408; E-mail: [email protected]

Journal of Macromolecular Sciencew, Part C: Polymer Reviews, 46:143–172, 2006

Copyright # Taylor & Francis Group, LLC

ISSN 1558-3724 print/1558-3716 online

DOI: 10.1080/15321790600646802



The free energy minimization during microphase separation results in the formation of

various interesting and thermodynamically stable structures on nanometer length scale.

For diblock copolymers consisting of two amorphous block chains, these structures

include lamellae (lam), bi-continuous gyroid phase, hexagonally packed cylinders

(hex), body-centered cubic (bcc) packed spheres and can be controlled by varying the

volume fractions of the constituting blocks or the segregation strength between the

blocks. The phase behavior describing these self-assembled structures has been fairly

well understood theoretically as well as experimentally.

Incorporation of crystallizable blocks in the block copolymers introduces extra com-

plexity in their microphase separation behavior. The morphological development in crys-

talline block copolymers is controlled by two competing self-organizing mechanisms;

namely, microphase separation and crystallization. As a result these block copolymers

exhibit richer phase behavior which is more difficult to predict. Crystalline-amorphous

(C-A) diblock copolymers, where one of the blocks is crystalline and the other is

amorphous, are the most widely studied crystalline diblock copolymers. Depending on

the segregation strength and on the relative values of the glass transition temperature of

the A block (TgA), the crystallization temperature (Tc), and the order-disorder transition

temperature (TODT), many different morphologies can be generated in C-A diblock copo-

lymers. Figure 1 schematically illustrates the possible scenarios and its effect on the

structure formation in C-A diblock copolymers.

Crystallization behavior of C-A diblock copolymers has been widely studied. Some

excellent reviews covering these studies have recently been published.6–8 The first of

these reviews by Hamley6 was later updated with more recent studies by Loo and

Register.7 More recently, Muller et al.8 have further reviewed the crystallization

behavior in semicrystalline block copolymers where they emphasized more on aspects

such as thermal properties and their relationship to the block copolymer morphology.Their review also dealt with nucleation, crystallization, and morphology of more

complex materials like double crystalline AB diblock copolymers and ABC triblock

copolymers with one or two crystallizable blocks.

As can be noted from Fig. 1, the most interesting and widely studied scenario is for the

case when TODT. Tc . TgA for the C-A diblock copolymer. The melt structure generated by

microphase separation in this case will be perturbed by crystallization if C-blocks constitute

the matrix phase and this will result in a lamellar morphology of alternating crystalline and

amorphous layers. However, when C-blocks become the minor constituent, they become

confined in the ordered microdomains. In this case, the interplay between the segregation

strength and crystallization driving force has a strong impact on the extent of structural per-

turbations and the structural development during crystallization becomes more complex. Anumberof studies havebeen reported for this kind of C-A diblock copolymer which substanti-

ates our knowledge on the crystallization and microphase separation behavior. However, a

complete understanding of the morphological development and the crystallization kinetics

in C-A diblocks, with soft amorphous phase, still remains a scientific challenge. Since the

previous reviews on crystallization behavior of block copolymers were more broad in their

nature, we felt the need for a review which focuses primarily on C-A diblock copolymers

with soft amorphous phase. Such a review will not only highlight the interesting behavior

reported through a vast amount of literature but will also generate interest in solving future

challenges for this kind of system. The present review, therefore, deals only with C-A

diblock copolymer where TODT. Tc . TgA. We will start the review by first giving an

overview on the available theories of crystallization in diblock copolymers. Then we

will focus on the morphological evolution in C-A diblock copolymers where the

B. Nandan et al.144



crystallization in thin films will be considered as a special case. Subsequently the crystal

orientation and crystallization kinetics of the C-blocks in microphase-separated structures

will be discussed. We will further highlight a new area of research in this field which

consists of the blends of two C-A diblock copolymers. Finally we will discuss the chal-

lenges which are the focus of most of the current research in C-A diblock copolymers.

2. Theories of Block Copolymer Crystallization

The first theoretical treatment for crystallization in block copolymers was given by

Ashman and Booth.9 They extended the Flory-Vrij theory10 which describes the crystal-

lization from melt in homopolymers to the block copolymer case. Their model allows

for the determination of melting temperature, degree of crystallinity, and end interfacial

energies for crystalline block copolymers. A more detailed theoretical analysis for the

structure formation on crystallization in C-A diblock copolymers had been developed

Figure 1. Schematic illustration showing various possible structure development scenario after

crystallization in semicrystalline diblock copolymers.

Crystallization Behavior of Diblock Copolymers 145



by Dimarzio, Guttman, and Hoffman (DGH theory).11 According to their model, the

lamellar morphology consists of alternating layers of amorphous and crystalline blocks

with crystalline chains folded perpendicular to the interface. The most important predic-

tion by Dimarzio et al.11 was the existence of an equilibrium degree of chain folding in the

crystalline state governed by the balance of the thermodynamic forces between the crystal-

line and amorphous domains. As opposed to homopolymers where the chain folding is

metastable and annealing reduces the amount of chain folding so that in the limit of

infinite annealing time extended-chain crystals will result, diblock copolymers would

anneal to an equilibrium crystal thickness. It was suggested that reducing the number of

folds in the C block from the equilibrium degree of chain folding will result in an extra

stretching of the A blocks and hence the preference of the A block to gain entropy by

attaining a random coiled state opposes crystal thickening. The proposed model allowed

for the stretching of polymer chains, the change in packing entropy arising from

changes in orientation of bonds, and the space-filling properties of the chains. The theor-

etical prediction also indicated that the lamellae in the crystalline layers pack in a

monolayer since the bilayer system was found to have a higher free energy and was

hence unstable. Furthermore, expressions for the calculation of amorphous and crystalline

layer thickness were also presented. Compared to the Flory-Vrij theory10 which con-

sidered the enthalpic term also, this model was a purely entropic one. DGH theory 11

also provided a universal expression for the domain spacing, d, in C-A diblock copolymers

which was found to scale as

d NtNÀ1=3a ð1Þ

where Nt

is the total degree of polymerization and Na

is the degree of polymerization of the

amorphous block.

Whitemore and Noolandi12 presented a more rigorous thermodynamic analysis of

structure formation during crystallization in C-A diblock copolymers using self-consistent

mean-field theory. The diblock copolymer was assumed to have a lamellar structure of

alternating semicrystalline and amorphous layers with the chemical bonds which

connected the copolymer blocks lying in the interfacial regions between the layers. The

amorphous blocks were modeled as flexible chains and the crystalline blocks as folded

chains. The expression for free energy in the crystallized state was interpreted as a sum

of four main contributions: the free energy of the amorphous block, that of the crystalline

block, the interaction energy of the two blocks, and the reduced entropy due to the

localization of the copolymer joints to the interfaces. The first two were suggested to be

the significant factors that control the equilibrium lamellar thickness. The model con-

sidered both the monolayer and bilayer arrangement of lamellae in the crystalline layer.

In agreement with the model of Dimarzio et al.,11 this model also predicted higher free

energy for a bilayer arrangement in the crystalline lamellae in case of a fully crystallizable

C-block. However, the bilayer arrangement becomes probable when the C-blocks are

semicrystalline in nature. Furthermore, the model gave the lowest free energy for the

case when crystalline chains have integral number of foldings. Expressions for

thickness of the amorphous/crystalline regions and the number of folds were also

derived. The scaling prediction of the Whitemore and Noolandi theory for the lamellar

repeat distance, d, is

d NtNÀ5=12a ð2Þ

B. Nandan et al.146



The later experimental results have agreed with the prediction, though there have been a

few exceptions.

3. Morphological Development in C-A Diblock Copolymerswith a Rubbery A Block

Below the order-disorder transition temperature, block copolymers undergo microphase

separation resulting in the formation of various self-assembled ordered structures. Now

we consider the case of microphase separated C-A diblock copolymers; what will

happen once it is cooled below the melting temperature of the C block? Energetically,

there are two possible scenarios. The melt structure will be preserved on crystallization

if microphase separation dominates the structure development process. If, however, crys-

tallization is the stronger driving force for phase separation, it is expected that C-A block

copolymers will behave like semicrystalline homopolymers and cooling from the melt

would result in the formation of alternating crystalline-amorphous lamellae and spherulitic

superstructures. Since the energy associated with crystallization (0 100 J/g) is much

larger than that associated with microphase separation (0 1 J/g), it is expected that crys-

tallization will always dominate the structure development process. This means that the

structure of a C-A diblock copolymer in a crystallized state should be no different from

that in a semicrystalline homopolymer.

However, it is well known now that this is not always the case. If the amorphous block

is glassy during the crystallization process i.e. TgA. Tc

C (hard confinement), the crystal-

lization of C-block will be kinetically forced to occur under a physical confinement and

the melt structure in this case is preserved.7 But what will happen if the amorphous

block is liquid-like or rubbery during the crystallization process i.e. TgA, Tc

C. It has

been observed that the complex interplay between microphase separation and crystalliza-tion process in this case can lead to complex structure development scenarios. In this

section we will focus on our present knowledge regarding morphology development in

such systems during the crystallization process.

The early studies on C-A diblock copolymers with rubbery amorphous block were

mostly carried out on weakly segregated systems. Hence, as expected, these studies

revealed that the microphase separated structures of these diblock copolymers undergo

considerable structural rearrangement during crystallization and the final morphology

consists of alternating crystalline and amorphous lamellae irrespective of the initial

melt structure. Nojima et al.13 were the first to observe this behavior while studying

poly(1-caprolactone)-block -polybutadiene (PCL-b-PB), a weakly segregated diblock

copolymer. The SAXS analysis showed that the energetic gain on crystallization over-whelmed that on microphase separation, so that the melt microstructure was completely

destroyed by the subsequent crystallization of PCL block. The morphology observed on

crystallization in PCL-b-PB diblock copolymer with total molecular weight (Mw) in the

range of 9400–39400 was found to be the same and in each case morphological rearrange-

ment from the microdomain structure into the lamellar morphology was observed.14 As

compared to the case of PCL homopolymers, the lamellar thickness was significantly

reduced in the PCL-b-PB, indicating that the lamellar morphology was strongly affected

by PB blocks. In a later study on PCL-b-PB with a wider range of molecular weight

(8000 Mn 62000) and at different crystallization temperatures (220 Tc 458C),

Nojima et al. observed that the molecular weight drastically affected the final microdomain

morphology on crystallization.15 All the copolymers had cylindrical or spherical microdo-

main structure in the melt with PCL block forming the domains. For copolymers with




Mn 19000 (where Mn is the number average molecular weight) morphological transition

to lamellar morphology was observed. However, for PCL-b-PB copolymers with

Mn! 44000, the microdomain structure in the melt state was found to retain in the

experimental crystallization temperature range. Moreover, the PCL crystallinity was con-

siderably reduced or became almost zero even though the PCL block was highly crystalline

by nature. This showed that in the case of higher molecular weight copolymers, when the

effective segregation strength becomes relatively high, microphase separation may

dominate over crystallization in PCL-b-PB. For Mn ¼ 30000, though the morphological

transition into a lamellar morphology did occur, the principal repeating distance of the

morphology did not change, suggesting the possible epitaxial relationship between the

microdomain structure and lamellar morphology for all Tc’s investigated. In a separate

study Nojima et al. showed that if PB blocks in PCL- b-PB copolymer were crosslinked

then the PCL blocks crystallized within the structure and no morphological transformation

was observed.16 Lee et al. have also shown that in case of blends of a poly(ethylene oxide)-

block -poly(butadiene) (PEO-b-PB) with a PB homopolymer, if PB blocks were cross-

linked, the crystallization of PEO blocks occurred within the pre-existing domain

morphology.17

Rangarajan et al. reported the morphological transformation on crystallization in case

of weakly segregated polyethylene/head-to-head polypropylene (PE-b-hhPP) polyolefin

diblock copolymer using time-resolved small- and wide-angle X-ray scattering (SAXS

and WAXS).18 However, they noted that microphase separation presented a substantial

barrier to the large-scale structural reorganization which occurred on crystallization.

Ryan et al. observed that the lamellar and hexagonally-packed cylinder structures in

poly(ethylene)-block -poly(ethylethylene) (PE-b-PEE) and poly(ethylene)-block -poly

(ethylene-proyplene) (PE-b-PEP) diblock copolymers transformed into a lamellar

structure on crystallization where the melt-structure was destroyed.19,20 Again, in caseof poly(isoprene)-block -poly(ethylene oxide) (PI-b-PEO) diblock copolymer with PI

volume fraction in the range 0.25, f PI , 0.92, Floudas et al. observed a layered

structure on crystallization irrespective of the initial melt-structure.21,22 The transform-

ation from the hexagonally packed cylindrical structure to a layered structure on crystal-

lization proceeded via heterogeneous nucleation and growth process and resulted in the

formation of a spherulitic superstructure composed from stacks of lamellar crystal.

The results of these earlier studies led to a belief that crystallization may always

dominate over the microphase separation in C-A diblock copolymer when the A block

is rubbery during crystallization process. However, later studies showed that structure

formation in these systems is not as simple and the initial idea about crystallization as

the dominant structure directing process only stemmed from the fact that the C-Adiblock copolymers used in the earlier studies were weakly-segregated. The segregation

strength of a diblock copolymer can be increased either by using blocks with stronger

repulsive interactions or by increasing its molecular weight or both. The melt segregation

strength dependent structure formation on crystallization was first demonstrated by

Quiram et al.23–25 through their work on poly(ethylene)-block -poly(3-methyl-1-butene)

(PE-b-PMB) semicrystalline diblock copolymer. It was observed that melt segregation

strength and rate of crystallization could have a profound effect on the solid-state

structure. The melt segregation strength in this case was varied by selecting block copoly-

mers from a wide range of molecular weights. Also the composition of the diblocks was

such that they formed a hexagonally-packed PE cylindrical microdomain in the melt

state. It was observed that crystallization from the strongly segregated melt (high

molecular weight diblocks) was confined to the cylindrical microdomain and produced

B. Nandan et al.148



a morphology essentially independent of thermal history. However, the morphology

produced by crystallization from weakly-segregated melts (low molecular weight

diblocks) was found to be highly dependent upon thermal history. Faster cooling kineti-

cally confines the crystallization to cylinders, while slower cooling resulted in complete

disruption of the cylindrical melt mesophase upon crystallization, leading to a lamellar

morphology with a large domain spacing. Interestingly, despite the marked differences

in the final structure between the polymers where crystallization was constrained to

cylinders and those where it broke out to form lamellar microdomains, the crystallization

kinetics were not remarkably different.

Loo et al.26

studied the crystallization behavior in a sphere forming block copolymer

of poly(ethylene)-block -poly(styrene-r-ethylene-r-butene) (PE-b-PSEB) where a 70 wt%

styrene content in the PSEB block provided strong interblock repulsion with PE in the

melt, allowing segregation strengths more than triple of that at the ODT to be accessed

at reasonable molecular weights. The composition of the copolymer (PE/PSEB 9/55)

was such that the microdomain morphology in the melt state consisted of PE spheres of

25 nm diameter packed in a bcc lattice. The amorphous PSEB block remained in a

rubbery state during crystallization of PE block. It was observed that the melt structure

in this system was preserved in the solid state even for extremely slow crystallization.

Since the sphere diameter was only 25 nm and the crystallization process extended over

hours, it was unlikely that the morphology developed under these conditions reflected

any kinetic limitation imposed by hindered diffusion of block copolymer chains; more

likely, crystallization confined to spheres was the equilibrium morphology of PE/PSEB

9/55. It was also reported that confining crystallization to within block copolymer micro-

domains impacted its kinetics drastically and had significant effect on the nucleation

mechanism. This will be discussed in more detail later in this review.

The segregation strength of the PE-b-PSEB diblock copolymer was varied byreducing their molecular weight. The break-out behavior on crystallization was thus

observed with the lower-molecular-weight samples. The extent of break-out was

dependent on the molecular weight of the sample as well as crystallization conditions.

Motivated by the work of Quiram et al.23 where they observed that crystallization

kinetics in a cylinder-forming PE-b-PMB was the same in both confined and break out

crystallization; Loo et al.27 further investigated this system by taking the diblock with a

range of molecular weight covering weak to strong segregation regime. The composition

of the diblock was such that they formed hexagonally packed PE cylinders in the melt

state. From the studies carried out on this system it was shown for the first time that

C-A diblock with a rubbery amorphous block, apart from the usual confined and break

out crystallization, can also exhibit templated crystallization behavior. The templatedcrystallization was observed in diblocks with intermediate segregation strength. In the

templated crystallization, while the melt morphology was generally retained on cooling,

local distortions and connections between cylinders occurred due to crystallization. The

cylinders formed by microphase separation in the melt generally guided the growing

crystals but did not wholly confine them such that structural perturbation occurs on a

local scale. As will be discussed later, this makes the crystallization kinetics in

templated case very different from that observed in the confined crystallization.

Templated crystallization, however, was not observed in the sphere-forming samples. In

cylinder-forming samples, the occasional “rogue” crystals connecting cylinders did not

impact the overall structure significantly, since the cylinders were long. Two such connec-

tion events per cylinder, while sufficient to permit crystals to percolate throughout the

entire specimen, would have little effect on the average structure probed by SAXS. By




contrast, the isometric microdomains in sphere-forming samples would no longer be

spheres if each were connected to the positions of two neighboring microdomains.

From their work on PE-b-PSEB and PE-b-PMB diblock copolymers forming spheres

or cylinders of PE, Loo et al.27 were able to compile a “classification map” where the nor-

malized interblock segregation strength (the ratio of interblock segregation strength at the

crystallization temperature to that at its TODT) was plotted against the volume fraction of

the crystallizable component, as shown in Fig. 2. The map classified the regions where the

three modes of crystallization i.e. breakout, templated, and confined were found. In the

case of sphere-formers, crystallization was effectively confined within the microdomains

when the normalized interblock segregation strength was high. Below the segregation

strength of 3 (normalized), dramatic structural rearrangement was observed on crystalliza-

tion. For cylinder-formers, structural rearrangement was again observed at weak inter-

block segregation (,1.5, normalized) and confined crystallization was again observed

at strong interblock segregation (.4). However, templated crystallization occurred

between these two limits.

The three modes of crystallization i.e. confined, templated, and break out, since then

have been identified unambiguously in other C-A diblock systems also. Xu et al.28,29

reported a comprehensive study on poly(ethylene oxide)-block -poly(butylenes oxide)

(PEO-b-PBO) diblock copolymers blended with PBO homopolymer and having

different segregation strength and morphologies. They observed that crystallization can

be readily confined in spheres whereas in cylinders confined crystallization occurred

only in the case of most strongly segregated system. However, for lamellar morphology

crystallization always led to break out of the melt structure. In an analysis similar to

that made by Loo et al.,27 Xu et al.29 found that the relative segregation strength

x C/x ODT had an important influence on crystallization mechanism. The confined

crystallization was observed at x C/x ODT . 3 whereas breakout crystallization occurredwith x C/x ODT , 3. For the blends with x C/x ODT around 3 the mode of crystallization

Figure 2. Classification map of crystallization models in PE-based semicrystalline diblocks with

rubbery matrices. Segregation strength at the crystallization temperature, normalized to that at the

ODT, is indicated on the y axis. Volume fraction of ethylene block (vE) in each diblock is shown

on the x axis; polymers with vE , 0.19 form spheres of PE (circles), those with vE . 0.19 form

cylinders of PE (squares). Open symbols denote complete destruction of the melt mesophase

upon crystallization (breakout); filled symbols denote complete confinement; symbols with a vertical

hatch denote templated crystallization. (Reprinted with permission from.

27

Copyright 2002 by theAmerican Chemical Society.)

B. Nandan et al.150



was found to be dependent on the crystallization temperature (confined at lower tempera-

ture but breakout at higher temperature). Templated crystallization was considered by

them as a type of partially confined crystallization where only part of the morphology

in the melt was retained after crystallization and the other part is transformed into

lamellae. A further study done by Xu et al.30 on neat PEO-b-PBO block copolymer

with x C/x ODT , 3 showed that breakout crystallization always occurred irrespective of

the initial melt structure consistent with their previous work.

Chen et al.31 have further studied the structure formation on crystallization in the

intermediate segregation regime where the morphology after crystallization arises from

the balance between the crystallization and microphase separation driving forces. The

molten mesophase in this case was neither fully preserved nor completely transformed

into 1D stacked lamellae, but instead intermediate structures were generated through the

crystallization process. The studies were carried out on a wet-brush PEO-b-PB/PB

blend system. The PEO spherical microdomains in the melt state slightly deformed into

ellipsoid-like objects on crystallization as shown in Fig. 3. Recently, Ho et al.32

observed a unique undulated lamellar morphology in poly(styrene)-block -poly(L-lactide)

Figure 3. TEM micrographs of PEO-b-PB/h-PB blends ( f PEO ¼ 0.13) showing the shape of micro-

domains before and after crystallization

(a) in the melt the PEO microdomains have spherical shape;

(b) PEO microdomains becomes ellipsoid-like in shape after slowly crystallizing the sample at

58C/min;

(c) crystalline PEO microdomains with higher magnification illustrating the images of ellipsoid-like

domains more clearly. PB matrix was preferentially stained with OsO4. Spheres are drawn in the

corners of the micrographs to help distinguish between the actual shape of the microdomains and

that of a sphere;

(d) Schematic presentation proposing the structure of the ellipsoid-like crystalline microdomains.

The interface in the crystalline samples is thicker, and it may consist of the uncrystalline PEO

segments mixed with a minor portion of PB. (Reprinted with permission from.

31

Copyright 2002by the American Chemical Society.)




(PS-b-PLLA) diblock copolymer when PLLA block crystallizes in a soft PS amorphous

phase. Figure 4 represents the TEM micrographs showing the undulated lamellar

morphology observed by Ho et al.32 The amplitude and periodicity of the undulation

instability were dependent upon the orientation of microphase-separated lamellae.

Similarly, Lambreva et al.33 observed that the bulk melt morphology of hexagonally

packed cylinders in PEO-block -poly(ethylene/butylenes) (PEO-b-PBh) transforms into

hexagonal perforated lamellar phase upon crystallization. Chen et al.34 also observed

perforated lamellae microstructure in PEO-b-PB diblock copolymer on crystallization

as shown in Fig. 5. More strikingly, recently, Hsu et al.35 have observed a highly

twisted lamellar structure in a symmetric PCL-b-PB diblock copolymer on crystallization

at high undercoolings. This shows the kind of structural complexity which may result from

crystallization-induced deformation of microdomains.

Recently, there has been a tremendous interest in exploring the mechanism by which

templated or breakout crystallization occurs in C-A diblock copolymers. It is believed that

in the microphase-separated melt consisting of cylindrical or spherical microdomains,

coalescence or welding of individual domains must occur on the way to the formation

of large crystalline lamellae. Loo et al.27 showed that the spherical microdomains of PE

blocks in PE-b-PSEB diblock copolymer form disk- or rod-like domains through local

coalescence when crystallized slowly as shown in Fig. 6. It was suggested that sufficient

Figure 4. TEM micrographs of PS-b-PLLA diblock copolymer ( f PLLA ¼ 0.585) obtained after crys-

tallizing the samples at

(a) 858C;

(b) 1008C from ordered melt at 1008C.

The micrographs show that crystallization induces undulation in the initial lamellar microdomains

observed in melt. PS domains were preferentially stained with RuO4. (Reprinted in part withpermission from.32

Copyright 2004 by the American Chemical Society.)

B. Nandan et al.152



Figure 5. TEM micrographs of PEO-b-PB diblock copolymer ( f PEO ¼ 0.54) after crystallizing at

room temperature (ca. 278C). The micrographs show the formation of perforated lamellar structure

on crystallization. PB matrix was preferentially stained with OsO4. (Reprinted with permission

from.34 Copyright 2002 by the American Chemical Society.)

Figure 6. TEM micrographs of PE-b-PSEB diblock copolymer (wE ¼ 0.14) after undergoing

different crystallization histories:

(a) quenched from the melt to room temperature at 508C/min;

(b) isothermally crystallized at 648C for 20 min.

The fast crystallization results in a structure similar to that in the melt however on slow crystallization

the initial macrolattice completely destroys. The PE microdomains are elongated into rods and discs.

PSEB matrix was preferentially stained with RuO4. (Reprinted with permission from.

27

Copyright2002 by the American Chemical Society.




deformation of microdomain interface by the growing crystals could allow the interfaces of

neighboring domains to meet, thereby permitting the crystal growth fronts in the crystalliz-

ing domains to intrude into the nearby molten domains easily and consequently leads to

extended crystal growth. Another “dissociation/association” mechanism proposed by

Nojima et al.36 asserted that the driving force of crystal growth can pull C blocks out of

the molten domains; these block chains then diffuse across the matrix and eventually add

onto the growing crystal surface that appeared in the crystallizing domains, leading to

extended crystal growth from an originally nanoscaled domain. Huang et al.37 studied

the coalescence of PEO spherical microdomains in a wet-brush PEO-b-PB/h-PB blend

system near the onset of melting of the as-crystallized samples. They found that on crystal-

lization, two or three spherical microdomains merged into a highly elongated prolate or rod-

like domain through postannealing as shown in Fig. 7. Since nearly all microdomains

remained semicrystalline during the treatment, this coalescence process was distinguished

from that found in the crystallization process. It was suggested that the further development

of crystallinity and perturbation in crystal dimension through crystal thickening (or

thinning) during postannealing introduced enough interfacial deformation to destabilize

the microdomain interface and consequently induced domain coalescence to reduce the

interfacial tension. In a later study,38 they also compared the coalescence process during

Figure 7. TEM micrographs showing the morphology of the as-crystallized or annealed PEO-b-PB/h-PB samples ( f PEO ¼ 0.17) with following thermal treatments

(a) as-crystallized at2308C;

(b) annealed at 388C after crystallization at 2308C;

(c) annealed at 418C after crystallization at 2508C.

The slightly ellipsoidal PEO microdomains found in the as-crystallized sample are seen to transform

into highly elongated prolates or rodlike domains after annealing. (Reprinted with permissionfrom.37


B. Nandan et al.154



isothermal crystallization and postannealing in PEO-b-PB/h-PB blend with that in neat

PEO-b-PB diblock copolymer to reveal whether the presence of homopolymer in the

amorphous matrix hinder or facilitate microdomain coalescence driven by the crystalliza-

tion. Three samples with nearly the same volume fraction and length of PEO block were

adopted for the study; the first was an asymmetric PEO-b-PB that contains no h-PB, the

second was a blend containing 12 wt% of h-PB, and the third was a blend consisting of

63 wt% of h-PB. The representative TEM micrograph of these three samples after crystal-

lization and the post-annealing is shown in Fig. 8. It was found that the blend containing the

higher h-PB content exhibited the strongest resistance against microdomain coalescence.

Figure 8. TEM micrographs showing crystallization induced microdomain coalescence in PEO-b-

PB copolymer and PEO-b-PB/h-PB blends

(a) PEO-b-PB/PB blend (wh-PB ¼ 0.63, f PEO¼ 0.17) crystallized at 2238C;

(b) PEO-b-PB/h-PB blend (wh-PB ¼ 0.63, f PEO¼ 0.17) annealed at 388C after crystallizing it at

2238C;

(c) PEO-b-PB/h-PB blend (wh-PB ¼ 0.12, f PEO ¼ 0.17) crystallized at 2308C;

(d) PEO-b-PB/h-PB blend (wh-PB ¼ 0.12, f PEO¼ 0.17) annealed at 388C after crystallizing it at

2308C;

(e) PEO-b-PB diblock copolymer ( f PEO ¼ 0.17) crystallized at 2238C;

(f) PEO-b-PB diblock copolymer ( f PEO ¼ 0.17) annealed at 388C after crystallizing it at 2238C.

The micrographs show that the PB homopolymer hinder the coalescence of the PEO microdomains.(Reprinted with permission from.38





The resistance toward domain coalescence exerted by the homopolymer was proposed to

stem from the diffusion barrier associated with the rejection of a portion of homopolymer

out of the coronal regions of the micelles, while the corresponding gain in free energy of

mixing was minor compared with the reduction in interfacial free energy upon domain

coalescence.

Recently, Xu et al.39 reported structural changes of block copolymer micelles induced

by crystallization. They studied PEO-b-PBO/h-PBO blends which formed disordered

spherical microdomains of PEO in the melt state. It was shown that the crystallization

of the PEO block from the micelles can be divided into two steps: Firstly, the micelles

crystallized individually through homogenous nucleation. Secondly, crystallization

induced deformation of the micelles. The extent of deformation is strongly dependent

on microstructure of the block copolymers. For shorter block copolymer larger defor-

mation occurred and the deformed micelles could aggregate into macro-crystals, while

the micelles of the longer diblock copolymer experienced little deformation and the mor-

phology of micelle was retained after crystallization.

Hamley et al.40 provided further information on pathway of structure formation

during crystallization from their studies on oriented hex and gyroid (gyr) melt phases in

an asymmetric PEO-b-PI diblock copolymer. Crystallization resulted in the formation

of 1D stacked lamellae where the alignment of planes of cylinders in the initial shear

oriented hexagonal phase templates the orientation of lamellar planes. Similarly, an

aligned gyroid phase templated the crystallization in the crystalline lamellae phase, the

orientation of which matched that of the (220) planes of the gyroid structure. The

domain spacing increased by 40% on crystallization, indicating a transition that is

strictly not epitaxial. It was suggested that the stretching of chains that accompanies crys-

tallization presumably causes the increase in the length scale, while the pinning of block

junctions to interfaces ensured that crystallographic register between lattice planes ismaintained.

4. Crystal Orientation and Chain Folding in Crystalline State

The folding of crystallizable chain and its orientation after crystallization in the lamellar

crystal structure of diblock copolymers has received considerable attention. Depending on

the circumstances, the crystallites may be oriented either parallel or perpendicular to the

plane of the lamellae. Figure 9 illustrates the parallel and perpendicular orientation of

crystal stems in lamellar microdomains of a diblock copolymer. In case of C-A diblock

copolymer having a glassy amorphous phase, the folding of chains occur such that

crystal stems are always parallel to the interface in a lamellar microdomain. However,for the case of rubbery amorphous phase, it has been shown that crystal stems lie perpen-

dicular to the interface. Also in accordance with the theoretical predictions of Dimarzio

et al.11 it has been revealed that the crystalline block acquires an equilibrium degree of

chain folding in the crystallized microdomains.

In the early studies done by Hamley et al.41 on high molecular weight PE-b-PEP and

PE-b-PEE diblock copolymer, the PE crystal stem orientation was found to be parallel to

the lamellar interface. It was suggested that interfacial area per block junction was suffi-

ciently large for amorphous blocks of high molecular weight which allowed the PE stem to

fold in a parallel orientation. This then allowed crystallization to take place without an

overall change in length scale. However, later studies by Hong et al.42–44 on PEO-b-PB

diblock copolymers, in both bulk and thin film state, showed that the PEO crystalline

chains orient normal to the microphase separated lamellar domain interface. It was also

B. Nandan et al.156



noted that the crystallization of PEO in the diblocks resulted in nonintegral folded crystal-

lites since the increase of PEO lamellar thickness with decreasing undercooling was con-

tinuous and that the chain stretching energy in the amorphous blocks prevents the

formation of extended chain PEO crystallites. Recently Li et al.45 also observed a continu-

ous change of the long period with increasing crystallization temperature in PEO- b-PBdiblock copolymer, which indicated noninteger folded chain crystals. This is in sharp

contrast to low molecular mass PEO homopolymers for which noninteger-folded chain

crystals are metastable and are only observed at the beginning of the crystallization.

The relative stability of noninteger-folded chain crystals in block copolymer systems

was attributed to two reasons. First, due to the loss of entropy because of attendant stretch-

ing of the amorphous block, the Gibbs free-energy landscape between integer-folded chain

crystals (n, n þ 1) will be rather flat. This reduces the thermodynamic driving force toward

integer-folded chains in comparison with homopolymers. Second, kinetically, the thicken-

ing process must overcome not only internal friction within the PEO crystals but also that

within the amorphous part. The perpendicular orientation of crystal stems to the lamellar

interface has also been revealed by Ho et al.32

in case of PS-b-PLLA diblock copolymersusing combined 2D SAXS and WAXS experiments.

In a cylinder forming PE-b-PMB diblock copolymer, Quiram et al.25 observed

that when crystallization is confined in the cylindrical microdomain, the orientation of

the PE crystal was the same as was found in the glassy matrix in poly(ethylene)- block -

poly(vinylcyclohexane) (PE-b-PVCH) diblocks. The crystals aligned preferentially

within the semicrystalline cylinders, but the orientation depended on the ability of

chains to diffuse during the crystallization process. When chain diffusion was most

rapid, alignment was observed with the chain axis in the crystals perpendicular to the

cylinder axis and the b axis (fast growth axis) coincident with the cylinder axis.

However, when the chain mobility was limited, the crystal stems tilted with respect to a

plane which is normal to the cylinder axis, allowing better accommodation of the

amorphous material at the interphase.

Figure 9. Schematic illustration of perpendicular and parallel orientation of crystalline stems in the

self-assembled microdomain of block copolymers. (Reprinted with permission from.6 Copyright

1999 by the Springer Science and Business Media.)




It would also be worth mentioning here about recent dynamic Monte Carlo simulation

studies by Hu et al.46,47

on primary crystal nucleation under both hard and soft confine-

ment in lamellar phase of C-A diblock copolymers. In the case of crystallization under

hard confinement, Hu et al.46 found that local chain ordering near the block junction or,

alternatively, close to the microphase interface, facilitated primary crystal nucleation

and hence caused the emergence of perpendicular (or parallel) crystallite orientations.

The study showed that crystals in the lamellar geometry tend to nucleate with the chain

axis perpendicular to the lamellar plane. However, if in the same lamellar structure, the

junctions between the crystallizable and non-crystallizable blocks of the polymers are

broken then crystallites tend to align parallel to the lamellar plane. Furthermore, in the

case of soft confinement Hu et al.47 studied the temperature dependence of crystallite

orientation as well as its correlation with the occurrence of microdomain coalescence

during crystallization. The simulation was done under both limits of strong and weak

segregation. It was observed that from high to low temperatures the saturated isothermal

crystallization showed a transition from perpendicular to random in the preference of

crystallite orientation. Furthermore, they found that crystallite orientation can have an

important influence on the coalescence process during crystallization. It was found that

under high temperatures those crystallites which showed their orientational preferences

perpendicular to the lamellar microdomain were responsible for the occurrence of coalesc-

ence, whereas the segregation strength just played a subsidiary role. Moreover, during iso-

thermal crystallization, the primary crystallization only produces undulation of lamellar

domains, while the subsequent isothermal annealing, in particular, the thickening of

these perpendicularly oriented crystals, induced the coalescence process. Furthermore,

random orientations of crystallites were found to facilitate the stability of microdomains

at low temperatures.

The extensive work done by Ryan et al.48,49 on PEO-b-PBO diblock copolymers haveshown that in the crystallized state PEO blocks acquire an equilibrium degree of chain

folding. They used low molar mass block copolymers in order to quantify the chain

folding of PEO block. On crystallization, the copolymers that had disordered or

lamellar melt phases showed an increase in the characteristic length scale due to stretching

of the amorphous PBO block above that experienced in the melt. The copolymers formed

kinetically-determined, highly-folded structures on rapid crystallization. These metastable

structures were stable to annealing but could be melted and self-seeded to grow equili-

brium less-folded structures in which the extent of folding was determined by the

balance between the Gibbs energies of PEO-block folding and PBO-block stretching.

Recently, Lee et al.50 further experimentally verified the DGH theory11 by showing that

crystalline block in C-A diblock copolymers show an equilibrium degree of foldingwhich increased steadily with the length of the amorphous block (Figure 10). They

showed that the hydrogenated poly(norbornene) (hPN) block in diblock copolymers of

hPN and hydrogenated poly(ethylidene norbornene) (hPN-b-hPEN) could be induced to

fold up to four times as the amorphous block increased in length. Reproducible interdo-

main spacings and crystal thicknesses were achieved for a range of thermal histories,

strongly suggesting that these correspond closely to equilibrium values.

5. Crystallization in Block Copolymer Thin Films

Compared to crystallization studies in bulk that in thin films of semicrystalline diblock

copolymers has relatively been less studied. Most of these studies have been carried out

by Reiter and coworkers.51–56 In one of the earliest studies on crystallization in block

B. Nandan et al.158



copolymer thin films, Reiter et al.51,52 observed that crystalline lamellae in thin films of

low molecular weight PEO-b-PBh diblock copolymer can be oriented vertically if crystal-lized at some definite conditions. Alignment of these lamellae on large length scales was

found when crystallization occurred at boundaries created by a dewetting process. This

provided a simple technique for creating regularly patterned polymer surfaces on the

nanometer scale. By employing a different rate of crystallization they noted that the

vertical orientation of the lamellae was kinetically controlled but not thermodynamically

favored. Annealing at temperatures closer to the melting point or keeping the sample at

room temperature for several months allowed the formation of a lamellar structure

parallel to the substrate. It was suggested that the key parameters in controlling the crystal-

line morphology was the number of chain folds (selected by the kinetics of crystallization)

which ultimately determines the lamellar spacing and relaxation within the crystalline

state.It will be discussed in the next section that overall growth kinetics of highly

asymmetric diblock copolymers forming spherical or cylindrical mesophases differs

qualitatively from the kinetics of unconfined geometries. Crystallization is initiated by

homogeneous nucleation under strong confinement in these mesophases. Reiter et al.53

provided a direct proof of such a crystallization mechanism in real space by showing

using atomic force microscopy (AFM) that crystallization and melting of a 12 nm PEO

spherical microdomains in thin films of PEO-b-PBh diblock copolymers occurred indepen-

dently of other domains. The elasticity differences between amorphous PBh, amorphous

PEO, and crystalline PEO permitted the clear resolution and identification of

amorphous and crystallized PEO spheres within the PBh matrix as shown in Fig. 11.

Reiter et al. isothermally crystallized their samples at 238C and by directly imaging the

morphology with AFM they found that crystallization occurred in a random manner,

Figure 10. Interlamellar spacings d for hPN-b-hPEN diblocks, plotted in the scaling form suggested

by DGH theory. Nt is the total degree of polymerization and Na is the degree of polymerization of the

amorphous block. Each point represents a different block copolymer (all with hPN block Mn near

6 kg/mol). Dashed horizontal lines represent plateaus corresponding to a discrete number of folds

n of the hPN block, as indicated in the schematic structures. The results confirm the basic premise

underlying the DGH theory that an equilibrium degree of chain folding exists in C-A diblock

copolymers which increases with the length of the amorphous block. (Reprinted with permission

from.50





sphere by sphere, and that the fraction of crystalline cells increased with crystallization

time. The authors found that the crystallization followed first-order kinetics. Deviations

were, however, observed in the final stages of crystallization, where the rate slowed sub-

stantially. Crystallization studies under strong confinement in restricted geometries like

ultrathin films of PEO containing diblock copolymers also had been used to gain moreinformation on the fundamental mechanisms of crystal formation in polymers. Under

such conditions, the observed morphology and its temporal evolution can be directly

related to molecular processes and the kinetics of crystal growth.53–56

Opitz et al.57

studied confined crystallization of PEO-b-PBh diblock copolymers in

lamellar films using AFM, X-ray reflectivity, and optical microscopy. They observed

that crystallization of the PEO block leads to an increase in the lamellar thickness of

both blocks in order to accommodate an integer or half-integer number of folds in the ver-

tically oriented crystalline stems. As the density of PEO increases upon crystallization,

this effect is accompanied by a contraction in the lateral direction, which results in

cracking of the film. In the case of thin films of an asymmetric PEO- b-PBh diblock

copolymer having PEO cylinders oriented parallel to the substrate, they observed that

crystalline stems are oriented parallel to the cylinder axis.

Figure 11. AFM phase images of sphere-forming PEO-b-PBh diblock copolymer showing the

variation in the number and distribution of crystalline cells after crystallization at 2238C for

(a) 5 min;

(b) 15 min. White circles represent crystallized PEO spheres whereas dark circles are amorphous

PEO spheres.

(c) Shows, on semilogarithmic scales, the percentage of uncrystallized spheres as a function of time.

The straight line represents the fit to first-order kinetics. (Reprinted with permission from.53

Copyright 2001 by the American Physical Society.)

B. Nandan et al.160



Structural evolution in a PEO-b-PB diblock copolymer on crystallization in thin films

was studied by Hong et al.43,44

Apart from their observation that PEO crystalline chains

orient perpendicular to the lamellar layers of the microphase separated structure, they

noted some more interesting behavior. Hong et al.43,44 observed that when a crystallization

front moved through the PEO layers in the thin film, it did so in all three layers and with the

crystals in orientational registry, but with a time lag between the crystallization of succes-

sive layers, suggesting that PEO crystals nucleated in one of the lamellae can eventually

crystallize the materials in adjacent lamellae. This was remarkable since the crystals were

separated by approximately 10 nm thick amorphous PB layers. It was suggested that such

spreading of crystals from one layer to another could be facilitated by the presence of pre-

existing edge or screw dislocations in the molten film. When a propagating PEO crystallite

in a single layer encounters such a screw dislocation, it spreads to adjacent layers to

produce a multilayered structure which originated from a single nucleus and thus has

one crystallographic orientation.

6. Nucleation Mechanism and Kinetics of Crystallization

The primary nucleation mechanism and kinetics of crystallization in nanodomains of

microphase separated C-A block copolymers have shown striking differences compared

to that in homopolymers depending on the degree of confinement. The primary nucleation

process involves formation of a small amount of crystalline material due to fluctuation in

density or order in the supercooled melt and is the first step for polymer crystallization.

The primary nucleation can be either homogeneous or heterogeneous. If no second

surface or existing nuclei (i.e. any type of second phase) is present and the nuclei

formation takes place spontaneously only due to supercooling, the phenomenon is

referred to as homogenous nucleation. There is a large free energy barrier for the

formation of a critical nucleus in this case and hence homogeneous nucleation requires

large undercoolings. Heterogeneous nucleation mechanism, which is more common in

polymers, occurs due to the presence of a foreign particle. These particles serve as the

critical nucleus for further crystallization and hence heterogeneous nucleation occurs at

relatively lower undercoolings. As will be discussed below, it has been widely shown

that restricting crystallization on a nanometer length scale in C-A block copolymers sig-

nificantly affected the nucleation mechanism. In fact, it is generally believed that confined

crystallization follows a homogenous nucleation mechanism, since the number of impu-

rities, which act as heterogeneous nuclei, is very small when compared to the C

domains, and the fraction of C block initiated by heterogeneous nucleation accounts for

only a very minor part of the crystallizable materials. For example, the number of Cdomains/cm3 is of the order 1014, 1015, and 1017 in a typical diblock copolymer with

lam, hex, and bcc morphology, respectively. By contrast, the number of impurities was

estimated to be only of the order of 105 cm23.27

In the bulk melt of homopolymers where crystal growth can advance freely over a

macroscopic scale, both nucleation and crystal growth are operative simultaneously

during the crystallization. The temporal development of crystallinity at a given tempera-

ture is properly described by the Avrami equation,

xcðtÞ ¼ 1 À expðÀktnÞ ð3Þ

where xc(t) is the normalized degree of crystallinity that has formed at time t, k is the

overall crystallization rate constant containing contribution from both nucleation and




crystal growth, and n is the Avrami exponent relating to the mechanism of nucleation as well

as the growth geometry. Most homopolymers exhibit an Avrami exponent of n ¼ 2–3

which prescribes a sigmoidal shape in the plot of xc(t) vs. t. This sigmoidal shape implies

crystallization proceeding through heterogeneous nucleation and long-range crystal

growth. In the microphase-separated melt of block copolymers where the crystal growth

is strongly frustrated by the nanoscaled continuity, once a nucleus is formed in a given

microdomain, the subsequent growth pertaining to this nucleus is limited to a very short

range such that it completes instantaneously before a new nucleus can be created in this

microdomain. In this case, the rate of crystallization will simply be proportional to the

fraction of microdomains that have not yet nucleated, yielding an Avrami exponent n ¼ 1

and the crystallization kinetics in this case can be described by following equation,

xcðtÞ ¼ 1 À expðÀk NtÞ ð4Þ

where k N is the nucleation rate constant. Equation (4) prescribes xc(t) to follow a simple

exponential function instead of a sigmoidal curve. Confinement directed nucleation

mechanism and crystallization kinetics in block copolymers have been widely studied

and the advent of techniques like time resolved simultaneous SAXS/WAXS have

provided valuable insight into the nucleation modes and growth habits during the crystalliza-

tion of semicrystalline block copolymers, as will be discussed below.

The initial studies by Nojima et al.58 showed that the crystallization kinetics in C-A

diblock copolymers with a rubbery amorphous phase does not differ much from that in

homopolymers. They studied the crystallization behavior of microphase separated PCL-

b-PB diblock copolymer by SAXS employing synchrotron radiation. The Avrami

analysis at the early stage of crystallization showed an exponent n ranging from 2 to 3,

which was comparable to n evaluated for PCL homopolymer. However, in the latestage, the crystallization of PCL-b-PB was significantly retarded compared to the case

of PCL, and the rate was dependent on the microphase separated structure and/or

molecular characteristics of the copolymer. Similarly Ryan et al.19 also observed that in

case of a PE-b-PEE and PE-b-PEP diblock copolymer having lamellar and hexagonally

packed cylinder structures in the melt, the crystallization proceeded by a nucleation and

growth process since the Avrami exponent was close to 3. However, these early results

were understandable since they involved weakly segregated diblock copolymers. On crys-

tallization the melt structure was totally disrupted resulting in crystalline lamellar structure

and hence the kinetics was similar to that observed for homopolymers.

Quiram et al.,24 in the meantime, made some interesting observations during their

studies on PE-b-PMB diblock copolymer. They noted that in the case of a strongly segre-gated PE-b-PMB diblock copolymer, though the crystallization remains confined in the PE

cylindrical microdomains, the kinetics of crystallization as monitored by SAXS/WAXS

were not markedly different from those of homogenous or weakly segregated crystalliz-

able diblocks that do form spherulites. The diblock showed sigmoidal crystallization

kinetics with n ¼ 1.7 –3.4. This contrasted starkly with the case of C-A diblock

copolymer with glassy amorphous phase, where confined crystallization within

cylinders led to first-order crystallization kinetics. As was to be found later, this was a

case of templated crystallization where the cylinders formed by microphase separation

in the melt generally guide the growing crystals but do not wholly confine them. This inter-

esting observation was also made by Shiomi et al.59 in case of a cylinder forming PEO-b-

PB diblock copolymer. SAXS results showed that the melt structure was preserved upon

crystallization, but the Avrami exponents as evaluated from DSC results were found to be

B. Nandan et al.162



almost the same as those for PEO homopolymers. However, the overall rates of crystal-

lization and crystal growth were considerably suppressed and the apparent activation

energy was somewhat higher in the crystallization from cylindrical domains.

The first comprehensive study on crystallization kinetics when the crystallizable

block is strongly confined in the block copolymer microdomain was done by Loo

et al.26 Their studies were inspired by previous reports on crystallization in ultrathin

polymer films on a substrate where it was revealed that both the crystallinity and crystal-

lization kinetics can be strongly perturbed by such confinement. They used a relatively

higher molecular weight asymmetric PE-b-PSEB diblock copolymer where the resulting

strong interblock repulsion restricted crystallization to spheres of 25-nm diameter.

Using time-resolved SAXS and WAXS, Loo et al.26 found that the crystallization

kinetics of the strongly confined PE-b-PSEB diblock can be quantitatively well

described by a simple exponential decay, or n ¼ 1 in the Avrami equation. Such first-

order kinetics indicated that the rate of isothermal crystallization was simply proportional

to the fraction of spheres which had yet to crystallize, as anticipated if crystallization in the

diblock is confined within individual microdomains. The crystallization in the PE spheres

were homogenously nucleated since the number of microdomains (2 Â 1016 spheres/cm3) far exceeded the possible number of impurities in the sample (109 nuclei/cm3).

The nucleation process hence required deep undercooling and since the PE spheres in

the diblock were only 25 nm across, crystal growth from the nucleus to the microdomain

interface was essentially instantaneous. The temperature dependence of the crystallization

rate thus reflected the temperature dependence of the nucleation rate only.

These earlier studies showed that the relationship between crystallization kinetics and

morphology in this kind of block copolymers was highly complex especially when the

crystallization was confined in the microdomain. Motivated by this interesting behavior

and especially that of the earlier study by Quiram et al.,24 Loo et al.27 studied the crystal-lization behavior in a range of PE-b-PSEB and PE-b-PMB diblock copolymers differing in

their molecular weight and hence segregation strength. All the diblocks formed spherical

or cylindrical microdomains of PE in the melt. As mentioned in section 2 of this review,

depending on the morphological perturbations and resultant crystallization kinetics, Loo

et al.27 categorized the crystallization in the diblocks into breakout, templated, and

confined. In the breakout crystallization, morphological perturbations on a large-scale

disrupted the melt structure producing conventional sigmoidal crystallization kinetics.

However, in confined crystallization the melt structure was retained on crystallization.

The crystallization was initiated by homogenous nucleation and the crystallization

kinetics was non-sigmoidal with n ¼ 1. In the templated crystallization, though the

overall morphology of hexagonally packed cylinders in the melt was retained after crystal-lization, the crystallization kinetics was sigmoidal with n . 1. In this case Loo et al.27

observed using TEM that occasionally “rogue” crystals connecting different cylinders

develop which allow a large volume of material to be crystallized from a single

nucleus, producing conventional crystallization kinetics and an overall crystallization

rate dramatically faster than for confined crystallization.

Chen et al.60 also revealed that the crystallization kinetics in the nanoscaled microdo-

mains of a diblock system can be precisely controlled by its microdomain morphology.

They studied crystallization kinetics and crystalline morphology of PEO-b-PB/h-PB

blends. In the lamellar melt, crystallization of PEO blocks was analogous to the

common spherulitic crystallization in homopolymers where the process occurred

through a series of heterogeneous nucleations followed by the propagation of crystal

growth over a macroscopic scale. The Avrami plot as shown in Fig. 12(a) showed




Figure 12. Developments of crystallinity during isothermal crystallizations in

(a) lamellae-forming neat PEO-b-PB ( f PB ¼ 0.50) and PEO-b-PB/h-PB ( f PB ¼ 0.64) blend

crystallized at 408C;

(b) cylinder-forming PEO-b-PB/h-PB blend ( f PB ¼ 0.69) crystallized at different temperatures;

(c) sphere-forming PEO-b-PB/h-PB blend ( f PB ¼ 0.83) crystallized at different temperatures.

In the case of samples with lamellae morphology, the crystallization curves are sigmoidal whereas

the crystallization curves for samples exhibiting cylindrical and spherical morphology follow the

exponential function. (Reprinted with permission from.60 Copyright 2001 by the AmericanChemical Society.)

B. Nandan et al.164



sigmoidal shape properly fitted by Avrami equation with n % 2.5. The crystal growth was

of long range because the growth fronts could repeatedly thrust into the microdomains yet

to be crystallized, and such a repetitive intrusion generated a highly interconnected

lamellar morphology. However, the crystallinity developments in the blends containing

cylindrical and spherical microdomains followed first-order kinetics as anticipated when

the crystallization was controlled by homogenous nucleation. The Avrami plots, as

shown in Fig. 12(b) and (c) followed the exponential function prescribed by Eq. (4).

The corresponding melt structure was neither totally disrupted into a lamellar morphology

nor fully preserved upon crystallization.

Furthermore, Chen et al.61

made a very interesting observation that freezing

temperature (Tf ) of PEO, in a PEO-b-PB/h-PB diblock system, determined from DSC

cooling experiment display distinct transitions at the compositions corresponding to the

morphological transformation, and the undercoolings required to initiate crystallizations

in cylindrical and spherical morphologies were much larger than that associated with

lamellar melt due to homogenous nucleation-controlled mechanism. Figure 13 shows

the plot obtained by them when the samples were cooled from 808C at 58C/min. Tf for

the neat symmetric PEO-b-PB diblock was around 358C. As the volume fraction of PB

( f PB) was increased by blending the diblock with PB homopolymer, morphological trans-

formation occurred. Tf dropped almost discontinuously by as much as 558C at f PB ¼ 0.69,

where the melt morphology transformed into cylinders. Tf leveled off with further addition

of PB, but a depression of 98C was identified as the melt morphology transforms from

cylinders to spheres (i.e., at f PB ¼ 0.83). Figure 13 clearly depicts three regimes of crystal-

lization kinetics corresponding precisely to the three morphological patterns of PEO-b-

PB/PB blends. Xu et al.27 further showed that different modes of such Tf -composition

plots can be obtained in PEO-b-PBO/PBO blends depending upon the segregation

strengths prescribed by the block lengths. However, it must be noted that a substantialdecrease in Tf from its value of the homopolymer cannot always be taken as an

evidence for confined crystallization. Loo et al.7 have also recently suggested that

freezing points measured during dynamic cooling are only loosely related to the state of

confinement during isothermal crystallization, and that direct structural measurement

Figure 13. Plot showing the freezing temperature (Tf ) of the PEO blocks as a function of f PB in case

of PEO-b-PB/PB blends. Three regimes of crystallization kinetics corresponding precisely to the

three morphological patterns can be identified. (Reprinted with permission from.

61

Copyright2001 by the American Chemical Society.)




(e.g., by in-situ SAXS or TEM) are essential to properly confirm the state of confinement

imposed on the crystals.

Apart from crystallization behavior in the classical microdomain structures Chen

et al.34 have studied crystallization in some more special geometric confinement. Their

study involved dry-brush PEO-b-PB/h-PB blends where the global structure of the

lamellar unit formed cylindrical or spherical vesicles at high h-PB compositions with

the vesicle walls composed of PEO blocks. The kinetics study revealed that crystalliza-

tions in this dry-brush blend may be templated or confined depending upon the compo-

sition of h-PB. Long-range crystal growth was accessible to the neat PEO-b-PB and the

blends with weight fraction of h-PB (wh-PB) , 0.8 due to extensive connectivity of

PEO lamellae templated by the melt morphology. Templated crystallization persisted

until wh-PB reached 0.7 where most PEO microdomains formed walls of spherical

vesicles. The corresponding crystallization was effectively confined within the vesicle

walls, and the exceedingly large number of density of PEO domains led to a homogenous

nucleation-controlled crystallization. In general, the confinement effect exerted by dry-

brush blending was far less effective than the corresponding wet-brush blending in

which the confinement started to operate at wh-PB 0.48 (i.e., the composition where

the PEO domains transformed into cylinders).

Recently, Hsu et al.62 have shown that the phase behavior of the strongly-segregated

blend consisting of a C-A diblock copolymer and an amorphous homopolymers (h-A),

which depends on the degree of wetting of A blocks by h-A can be conveniently

probed by the crystallization kinetics of C block. They blended h-PB of different

molecular weight with a lamellae-forming PEO-b-PB diblock to yield blends exhibiting

wet-brush, partial dry-brush, and dry-brush phase behavior in the melt state. The crystal-

lization rate of the PEO blocks upon subsequent cooling, as manifested by the T f was

highly sensitive to the morphology and spatial connectivity of the microdomainsgoverned by the degree of wetting of PB blocks. As the weight fraction of h-PB

reached 0.48, for instance, Tf experienced an abrupt rise as the system entered from the

wet-brush to the dry-brush regime, because the crystallization in the PEO cylindrical

domains in the former required very large undercooling due to a homogenous nuclea-

tion-controlled mechanism while the process could occur at the normal undercooling in

the latter since PEO domains retained lamellar identity with extended spatial connectivity.

The studies, reviewed above, demonstrate that there is a strong correlation between

microdomain morphology and crystallization kinetics in C-A diblock copolymers.

Knowledge about one can facilitate the understanding of the other.

7. Crystalline-Amorphous Diblock Copolymer Blends

Although a number of studies have been carried out on crystallization behavior in neat

C-A diblocks and their blends with h-A, blends of two C-A diblocks have still not been

studied in detail. It is expected that the crystallization and the resulting phase behavior

in such systems may be highly complex.

Recently, Huang et al.63 reported a cocrystallization behavior in binary blends of C-A

diblock copolymers. Their studies were motivated by the well-known fact that homopoly-

mer mixtures from the same homologous series, differing sufficiently in length, undergo

fractionated crystallization where the long polymer chains segregate from the short

ones and crystallize separately. In some instances, however, the long and short chains crys-

tallize into the same crystalline lamellae leading to cocrystallization. However, the cocrys-

tallization observed in such homopolymer mixtures is a kinetically driven process which

B. Nandan et al.166



takes place under nonequilibrium conditions. Huang et al. investigated whether the crys-

tallization behavior will be the same in case of a binary blend of C a-b-Ab /Cg -b-Ad i.e.

whether Ca

and Cg

blocks originally mixed within the microdomains in the melt would

cocrystallize or phase segregate into their own crystalline lamellar structures. The

model system under study was the blends of a short nearly symmetric PEO-b-PB and a

long asymmetric PEO-b-PB. Using DSC and SAXS, Huang et al. found that the PEO

blocks of different lengths in the binary PEO-b-PB blends tended to cocrystallize

whereas the corresponding blends of PEO homopolymers showed phase-segregated crys-

tallization. Huang et al. presented a schematic model to explain their results which is

shown in Fig. 14. The melt structure (Fig. 14(a)) was formed by the intimate mixing of

the two diblocks, where each lamellar domain is constituted of two layers of brushes

lying on top of each other. The plausible structural scenario after phase-segregated crystal-

lization was presented (Fig. 14(b)) where the crystallites formed by the longer and shorter

Figure 14. Schematic illustrations of the structures in blends of a symmetric and asymmetric

PEO-b-PB diblock copolymers:

(a) the melt structure formed by the intimate mixing of the two diblocks, where each lamellar

domain is constituted of two layers of brushes lying on top of each other;

(b) a crystalline structure generated by the phase-segregated crystallization, where the crystallites

formed by the longer and shorter PEO blocks coexist within the lamellar domains upon fraction-

ation. The long PB blocks are highly stretched to maintain the normal density in the PB domain;

(c) the crystalline structure generated by cocrystallization.

This structure allows the lower interfacial energy and higher conformational entropy of the long

PB blocks in the melt state to be largely retained. (Reprinted with permission from.

63

Copyright2004 by the American Chemical Society.)




PEO blocks were shown to coexist within the lamellar domains upon fractionation.

However, this lamellar structure contained an entropic penalty due to excessive stretching

of the long PB blocks. It was also mentioned that phase segregated crystallization via

demixing of short and long diblocks was not possible since it will lead to an increase in

the interfacial energy. In light of the energetic and entropic penalties introduced by

phase-segregated crystallization, the cocystallized structure shown in Fig. 14(c), in

which the shorter and the longer PEO blocks cocrystallize uniformly in the microdomain

with similar lamellar thickness, was shown to become the thermodynamically favored

morphology of the system. The cocrystallization observed here represented a scenario

where a kinetics-dominated process found in homopolymer crystallization possibly

turned into a thermodynamically favored process in the C-A diblock systems.

Crystallization behavior in binary blends of PCL-b-PB copolymers having extremely

different crystallization rates has been studied by Tanimoto et al.64 using time-resolved

SAXS with synchrotron radiation and DSC. The binary blend formed a single micro-

domain structure in the melt over the whole composition range and the crystallization

proceeded with an intermediate rate between those of the constituent PCL-b-PB copoly-

mers to result in a single lamellar morphology. This indicated that the crystallization in

the blend was driven by a single crystallization mechanism as in the case of pure

PCL-b-PB copolymers. Also a steep change in the total crystallization rate with compo-

sition was observed which was ascribed to the difference in the stability of the preexisting

microdomain structure.

8. Future Challenges

Despite a plethora of studies, there is still a lot of room to pursue on our knowledge aboutthe phase behavior, crystallization kinetics, and mechanism in semicrystalline diblock

copolymers. Although crystallization in diblock copolymers with glassy amorphous

phase is relatively well understood, that with rubbery amorphous phase, which is dealt

with in this review, will need much more systematic study in the future.

In the weak-segregation regime extensive breakout of the pre-existing melt-structure

occurs during crystallization. The more interesting scenario is where the C-block is

confined in two or three dimensions viz. cylindrical and spherical microdomains in the

melt state. Crystallization in these cases extensively deforms the microdomains which

further undergo coalescence process resulting in the crystalline lamellar morphology.

However, how the coalescence process of different microdomains proceeds in breakout

crystallization has still not been well understood. Also we still do not know if the crystal-lization and coalescence process occurs simultaneously and that how large-scale crystal

growth occurs during this process. Recent theoretical predictions by Hu47 which

explains the coalescence process during crystallization may help in understanding the

breakout mechanism. Further, it is well known that order-order transitions (OOT) from

spherical to lamellar microdomains in case of amorphous diblocks involve some

transient phases. This raises the question if similar or more anomalous transient phases

can be observed during crystallization-induced OOT in semicrystalline block copolymers

as it also involves coalescence of microdomains in the initial stages. Also, in the inter-

mediate segregation regime crystallization is essentially confined, especially in

spherical or cylindrical microdomains, depending on the degree of undercooling. This

reveals that kinetic factors such as diffusion may also be playing a crucial role during crys-

tallization of block copolymers especially considering that the coalescence process

B. Nandan et al.168



involves microdomains which are well-separated. The role of such kinetic factors in

confined crystallization needs to be understood more properly.

Another problem which needs to be addressed in semicrystalline diblock copolymers

is the preference of crystallite orientations both in the initial stages of crystallization as

well as in the final crystallized state. Solving this problem is also crucial in order to

properly understand the breakout mechanism during crystallization. Recently, attempt

has been made to address this problem using theoretical means which we discussed in

the previous sections.47 However, experimental understanding of these theoretical predic-

tions will be a challenge for future studies. It will be especially important to understand the

effects of parameters like degree of undercooling, segregation strength of the diblock, and

dimensionality of confinement on crystallite orientations.

Crystallization induced deformation in semicrystalline block copolymers sometime

also leads to anomalous self-organized structures. For instance, spherical microdomains

becomes ellipsoid, perforated lamellar morphology has been observed after crystallization

in cylindrical microdomains, and classical lamellar structure becomes undulated. We

believe many more such anomalous structures driven by crystallization may appear in

future studies which may be highly complex. Recently, while studying the PCL- b-PB

diblock copolymer in our group, we observed such a complex structure which appears

to be helix in nature, though an unambiguous assignment of the correct morphology is

still in process. This just serves as an example for the kind of structural complexity

which we may expect in semicrystalline diblock copolymers.

The crystallization behavior in diblock blends may be an interesting area of research

for the future. It was reported that in contrast to kinetically trapped solid solutions formed

in mixtures of homologues homopolymers, the cocrystallization behavior observed in cor-

responding diblock mixtures may be a thermodynamically driven phenomena. Further

studies are required in order to arrive at a definite conclusion. Also, the cocrystallizationwas observed in the diblock blends when they were crystallized from a microphase

separated melt state. It will be interesting to investigate if the results will be similar

when the diblock blends are directly crystallized from homogenous melt state.

The cocrystallization was observed in a mixture of similar diblocks having different

molecular weight. It will be of interest to extend this concept to explore the possibility of

cocrystallizing two chemically different blocks (C and D) in the binary blends of C-b-A

and D-b-A. The diblock components chosen should preferably have identical amorphous

blocks to enhance the miscibility of the two copolymers. If solid solutions of C and D can

be formed, then this will open a novel route for preparing polymer crystalline alloys,

which is essentially implausible through blending of the corresponding homopolymers.

The fundamental aspects of crystallization behavior and its effect on the morphologi-cal features in semicrystalline block copolymers are just getting clear to us. However, a

number of problems still remain unresolved which demands more extensive and systema-

tic studies on these materials. This, together with the opening of some new directions of

research, further increases our curiosity regarding the future developments in this area.

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r06_nandan_crystallization behavior of crystalline-amarphous diblock copolymers

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