Surface segregation in model symmetric polyolefin diblock copolymer melts

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Surface segregation in model symmetric polyolefin diblock copolymer melts

Mohan Sikka, Navjot Singh, Frank Bates, Alamgir Karim, Sushil Satija, Charles Majkrzak

To cite this version:

Mohan Sikka, Navjot Singh, Frank Bates, Alamgir Karim, Sushil Satija, et al.. Surface segregation in

model symmetric polyolefin diblock copolymer melts. Journal de Physique II, EDP Sciences, 1994, 4

(12), pp.2231-2248. �10.1051/jp2:1994258�. �jpa-00248128�

Classification Physics Abstracts

61.12E 61.40 82.65D 68.90

Surface segregation in model symmetric polyolefin ^diblock copolymer melts

Mohan Sikka

(I), Navjot Singh (I),

Frank S. Bates

(I), Alamgir

Karim

(2),

Sushil

Satija (2)

and Charles F.

Majkrzak (2)

(')

_Department of Chemical

Engineering

and Materials Science,

University

of Minnesota,

Minneapolis~

Minnesota 55455. U-S-A-

(2) ^{National Institute of} Standards and Technology~

Gaithersburg,

Maryland, U-S-A- (Receii>ed 5 March 1994, received in final form 5 July J99J, accepted 29 July J994)

Rdsumd. Nous _avons dtud16 le comportement de s6gr6gation _en surface de couches minces de

polyoldfines

(copolymbres modbles h blocs) contenant _une

phase

lamellaire. Les films _sont pr6pards _{par «}

spin coating

_{» sur} diff6rents substrats, la microstructure _est

analysde

_par r6flexion de neutrons, r6flexion de rayon X,

ellipsom6trie, microscopie optique

et

spectroscopie

de _masse d'ions secondaires. Chacun des

copolymbres

h blocs _est suffisamment

proche

^{de la} transition ordre-d6sordre (TOD) _{pour que} le

profit

de

composition

dans la

phase

lamellaire puisse dtre

approxim6

par une fonction sinusoidale. L'interface _avec l'air _ou le solide est

toujours

enrichie par le mdme bloc, mdme lorsque I'£nergie de surface du substrat est

plus

61ev£e que celles des deux

blocs du copolymbre. Nous _{ne sommes} pas capables de relier _ce rdsultat _avec l'enrichissement pr6f6rentiel bass seulement _sur la diff£rence de tension de surface

dispersive

entre [es diff6rents blocs. Toutefois, _{nous remarquons que} dans tous les _cas, le bloc ayant la plus

petite

conformation est isold _en surface. Ceci suggbre qu'un facteur entropique, attribud h

l'asymdtrie

conformation- nelle~

joue

_un r61e dans le comportement en

proche

surface des

copolymdres

_en milieu fondu.

Abstract. Surface segregation in thin films of model symmetric

polyolefin

diblock

copolymers exhibiting

a lamellar

morphology

has been

investigated.

Films _were

spin

coated _{on a}

variety

of

substrates and the

resulting

microstructure

analyzed using

neutron reflection, X-ray reflection,

ellipsometry~

light

microscopy

and secondary ^ion mass spectrometry. Each of the block copolymers is sufficiently close to the order-disorder transition (CDT) _so that the

composition profile

within the lamellar domains _can be

represented by

_a nearly sinusoidal function. In all _cases, the _same block _was found to enrich both the

polymer/solid

and the

polymerlair

interfaces, _even

when the surface energy of the solid substrate

was

higher

than that of either block of the copolymer. These results _cannot be reconciled with preferential enrichment based

solely

_on the difference in the dispersive surface tension between blocks. In all instances, _we find that the

conforrrationally

smaller block is the surface-active _one. We propose that _an entropic

driving

force, attributable _to conformational asymmetry, plays an

important

role in the near-surface

behavior of block

copolymer

melts.

1. Introduction.

Block

copolymers

_are useful _as

large

surfactant molecules because of their

ability

to straddle

phase

boundaries and interfaces.

Applications

that _can be visualized _to

exploit

this interfacial

@Les

Editions de

Physique

¹⁹⁹⁴

activity

include

stabilizing

colloidal

dispersions Ii

and

compatibilizing

the interface between immiscible

homopolymers [2].

The surface

properties

of

polyolefin (saturated hydrocarbon) copolymers

are of

particular

interest _{to us} for several _reasons. In

1992, poly(ethylene)

and

poly(propylene) together represented

about 40§b of the total volumetric

production

of

thermoplastic

resins in the United States

[3]. Polyolefins

_are

widely

used in

products

where interfacial

properties

_are

important,

such _as

adhesives, packaging,

lubricants and

coatings.

Also, polyolefins

have been

recognized

_as

superior

materials for fundamental studies of

polymer properties,

and _a

variety

^of

compounds

with controlled molecular architecture have

recently

^been

synthesized

and characterized

[4, 5].

An

ability

to control intersurface interactions

using

block

copolymers

is dictated

by

_our

understanding

of the

physical

and

thermodynamic

behavior of these materials in the bulk state.

Considerable theoretical and

experimental

efforts have been made in the last twenty years ^to

categorize

the distinctive microstructure and

ordering

transitions

displayed by

block

copolym-

ers in the melt, as reviewed

recently by

Bates and Fredrickson

[6].

Of

particular

interest in this class of materials _are diblock

copolymers, composed

^of sequences of distinct monomeric units A and B

covalently

linked at _one

point.

The overall

degree

of

polymerization N,

the fraction

f

of block A in the chain, and the A-B segment-segment interaction parameter x, are three

parameters

^{that control the}

phase

behavior of bulk diblock

copolymers.

The

product xN

and

f

_are

generally

used _to

classify

the bulk

phase

state.

Recently,

Bates and coworkers have used model

polyolefin

diblocks and

homopolymer

blends _to

experimentally

demonstrate two facets of

phase

behavior not

anticipated by

conventional mean-field treatment. First, the

existence of

multiple

ordered

phases

in the weak

segregation

limit

[7]

for

polyolefin

diblocks that _are

compositionally off-symmetry,

I.e. when

f

is

significantly

different from 0.5

[8, 9].

Second, the correlation of

binary phase

behavior with _a parameter e

(see below)

that

provides

_a

measure of the conformational asymmetry ^{between the} two blocks _{or two}

homopolymers [10-

l2].

In _a thin diblock

copolymer

film

(with

film thickness

comparable

to several times the bulk

domain

spacing),

_{one may expect} bulk

thermodynamics

_to be modified

by

the presence of

symmetry breaking

^interfaces that _can influence the bulk

equilibrium configuration.

Most

experimental

work _to date in this _area has been done _on

nearly symmetric ~f

m 0.5 diblocks that form the lamellar

morphology

in the bulk. The first

study,

_on

poly(styrene)-poly(methyl-

methacrylate) (PS-PMMA)

diblocks _{cast on} silicon dioxide substrates and annealed above the

glass transition,

indicated _a lateral

alignment

of lamellar microdomains

parallel

to the film surfaces for

xN

_>

(xN)o~~ [13-16].

Here the order-disorder transition

(ODT)

refers _to the

point

in

xN

_space above which the

copolymer

would form the lamellar

morphology

from _a disordered melt. This lateral orientation has since been observed in other systems

[17-19].

A second effect that has been

reported [17-20]

is the presence of surface induced orientational order at values of

xN

_~

~XN )o~~,

^where the bulk _state

(I.e., essentially infinitely thick)

is

disordered. Another is

that,

for

XN

_>

(xN )o~~, affinity

^{for the} same or different blocks at the air and substrate interfaces leads to the

symmetric [17, 18,

2

Ii

_or the

asymmetric [13-16, 19]

thin-film

geometries depicted

in

figure

I. In

addition,

the

deposition

of

non-integral

and _non-

half-integral

amounts of material

(t)

_# dn _or

(t)

_#

d(n -1/2)

for the

symmetric

and

asymmetric

_cases,

respectively,

where

(t)

is the average film thickness and _n is _a

positive integer]

leads to a surface

topology composed

^of islands _or holes of thickness d. These defects

were first observed for PS-PMMA diblock films

using light

interference measurements

[22].

Currently,

_{we are}

studying

the effect of reduced

dimensionality

_on

phase

transitions and domain size for

off-symmetry ~f

_#

0.5) polyolefin

diblocks that exhibit non-lamellar and

multiple

ordered

phases

^{in the bulk}

[8, 9,

23,

24].

Related issues that emerge are whether surface

topology (holes

^and

islands)

_can be observed for non-lamellar

morphologies

cast in the

a.

Asymmetric

Film b.

Symmetric

Film

I ~

~

^~

~t3 _~~z

~~~~~

~

^~

~,

' ~~3

..~ ~+.. ~..~i.'~'"?~. iii? :J~~S.

~

~~z ^~' Z ~.

~jj~jm(()j#jjJ%~ ~

,,,,,,,,,,,,,,,,,,,,,

Substrate Substrate

(Il 1/2) d~ 'l£~fi

,

~, ,''~ ,§~

^'

^l'

^"

l'~ l~

/ /

',

/

a ^' ~

~ i

j

~

i i '

I j

' ' _'

I ' I i I ^' ' f

,' ',

_{,I ,}

,_,

,~'

,,

z z

Fig.

I. The two possible thin-film

geometries

for symmetric

~f=

0.5) diblock

copolymers.

d is the lamellar

period

and _n is

an integer. (a) Enrichment of different blocks _at

polymerlair

and

polymer/substrate

interfaces~ and (b) enrichment of the _same block at both interfaces. Each type of film

can have two different

composition profiles,

phase shifted by 180°,

depending

on which block is surface- active.

thin-film

configuration,

and whether the formation of surface _structures

depends

_on the

majority

component

being

surface-active. We _are also interested in

examining

how confine- ment of _a structured diblock

copolymer

^{melt between}

parallel plates

influences its microstruc- ture

[25],

since the film will be unable _to form holes and islands in such _a

configuration.

From the discussion above, three interrelated aspects emerge in

characterizing

the behavior of block

copolymers

in thin films and _near interfaces

I)

internal microstructure for _a system that may no

longer

be three-dimensional in nature,

it)

lateral

topology

at the free

(air)

surface of _a

film,

and

iii)

enrichment _{at a} free surface _{or a} solid interface. The

coupling

between two or

more of these

phenomena

introduces _a

degree

of

complexity

that makes thin-film research

challenging,

_as will be discussed in _a

forthcoming publication [26].

For

example,

_we have found that internal microstructure and surface

topology change

with temperature ^for at least three _{reasons :} thermal

expansivity changes

the relative dimensions of the overall film thickness and the domain

periodicity,

coil dimensions

(hence,

the bulk

periodicity)

themselves

change

with temperature

[27],

and

phase

transitions between ordered

morphologies

can occur

[8, 9].

We also observe that microstructure and

topology

for non-lamellar ordered _mor-

phologies,

cast as thin films, _are

susceptible

to

dynamic

barriers to true

equilibrium [26].

Examples

of

trapped complex morphologies

in solution cast films _are

reported by Hasegawa

and coworkers

[28]

for the

poly(styrene)-poly(isoprene)

system. ^The observation of

quasi-

ordered structures and the

suppression

of island formation in

asymmetric ~f

m 0.77 PEP-PEE

thin films have

already

been

reported by

_us

recently [24].

In this paper, the [east

complicated

of the three issues discussed above, surface

segregation

in ordered diblock

copolymer melts,

is examined in detail. A _recent letter summarized _our initial

findings

_on this

subject [18]. Compositionally symmetric J

m 0.5

polyolefin

diblocks

were used with molecular

weights

that

placed

them in the ordered _state~

forming

the lamellar

morphology

in the bulk, _as confirmed

by small-angle

neutron

scattering (SANS)

and

rheology experiments.

Thin films

supported

_on

just

_one side

by

a solid substrate _were studied. Molecular

weights

were chosen to make all the diblocks

relatively

close to the ODT _at the measurement

temperature ^the

composition profiles closely

resemble those

reported

^for ordered diblock

specimens

near TODT

l17],

and

proximity

to the ODT _ensures that the films achieve _an

equilibrium

state. Also, the

samples

_were annealed

sufficiently

above the

glass

transition temperature

(T~)

^{and, where}

appropriate

the

crystalline-amorphous

transition

(T~),

of the blocks _to eliminate kinetic barriers _to

equilibration

associated with the

glassy

and

crystalline

states. It is

pertinent

to stress that _an

equilibrated

lamellar

morphology,

with

alignment

of

microdomains in the lateral film

direction,

is

important

in

studying

surface

segregation

in _a direct way. This

configuration

_ensures

wetting

^of both film interfaces

by

the associated

surface-active block [8~

9]

without

distorting

^the bulk

morphology

_as observed in _more

complex

systems, for

example,

where _a

wetting

^{block is} a

minority

_component

[6].

Also,

limiting

the

study

to model

polyolefins

has enabled _{us to} examine the effect of molecular

architecture and coil size _on

polymer

surface

activity

in _a controlled way.

2.

Experimental.

2.I MATERIALS. Four model

polyolefins

_were chosen for this

study polyethylene (PE),

poly(ethylene-propylene) (PEP), poly(vinylcyclohexane) (PVCH),

and

poly(ethylethylene)

(PEE). ^The microstructure and relevant

physical properties characterizing

these

polymers

are

listed in table I. The statistical segment

length

b is of relevance _to this paper. It is defined

by

the

unperturbed

Gaussian coil radius of

gyration,

R~

=

b(N/6)~'~, _(l

where N is the

degree

of

polymerization

defined _as the molecular

weight

^divided

by

the

monomer mass. This parameter ^has been determined for PE, PEP, PVCH and PEE in the bulk

state

by small-angle

neutron

scattering [4, 10].

For each of these

polymers,

we

recognize

that

conformational asymmetry ^{effects should be defined in} an

N-independent

_way. We _can

combine the volume of the molecule

(V)

and space

filling (R~)

parameter ^into a

single

parameter Ii,

~

R( b2

~

V 6 vo ^' ~~~

where _vo is the volume

occupied by

a statistical segment. Conformational asymmetry ^between

two

species

is then defined

by

:

p

2

E=

~)) ₍₃₎

Table I also lists the value of the parameter

p~,

_{for each} of _our

polyolefins

_at 25 °C.

The number average molecular

weights, _M~,

for the diblock

copolymers

used in this

study

are listed in table II. Also listed _are the

isotope

content, the block

copolymer composition

Table I.

Polyolefin

molecular characteristics.

~°~~°~~~~ ^{~~~~~ ~'~~~°~~~}

g/cm3

~~l~~~

_K-1

PE

0.855~.b 7.5~ 8.9b>h

-0.58k

12.I

PEP

(Pclyefllylene-

0.854~ 7.0f

8.l~

-0.58~ 8.I

propylene)

~PC'Yetl1yl

~ ($

o_869~ 6.0f _4.9l 0.2i 3.8

e~~e)

~~~

~2

cH~

PUGH

~CH~-CH~

~Pc'Yvinyl N o_947b.4t 4.0~t 7.ld =0° 4.4

cycloltexwe)

° PE and PEP flso conwin small percen&ges ofrandomly dbtribttted branched repeat ttni~ as reportedin Reference 10. a Reference 29.

bEmapo~tttedfrom above the melt tempenture. CDentity gra&ent column measuremem. ^d Reference 12.e Reference 30. fReference 31.

i Based _{on a} chemical repeat unit of the homopolymer melt as determined by BANS. h Reference 32. i Reference 33.I Reference 34.

k Reference 35.

Table II.

Polyolefin

dib/ock

copolymers.

A-B 103Mn,

IA

TOD~°C Bulklamellar Ref

Diblock gmol"~

period,

d,

I

PVCH(d8)-PEE

52 0.48 212 248 (25°C) 12, 36

PEP-PEE(d6)

55 0.55 125 341 (25°C) 27, 37

PE(d6)-PEE

26 0.47 182 246 (130°C) 38

PE-PVCH(d8)

15 0.48 238 180 (25°C) 12, 36

PE(d~)-PEP

100 0.50 139 534 (130°C) 39

f, order-disor4er

transition

To~~,

and the bulk lamellar

spacing

d. These materials _were

synthesized

and characterized _as described in separate

publications [4, 5, 12].

Methods for

obtaining To~~

^and d in the bulk state _are also discussed elsewhere

[27, 40].

2.2 SAMPLE PREPARATION. Thin film

samples

_were

prepared by spin coating

from solution

onto either 2- _or 4-inch flat substrates. Silicon surfaces _were

prepared

in _a clean _room in the

following

_way. Polished silicon wafers were cleaned either with

methylene

chloride and methanol _or with _a 3 _: 7

hydrogen peroxide/sulfuric

acid solution at 120 °C. The wafers _were then rinsed

thoroughly

with distilled water for several minutes and then etched with _a buffered 1: 9

HF/NH~OH

solution to _remove the native oxide. After

thorough rinsing

with distilled

water once

again

and

drying

with

high purity nitrogen,

the wafers _were

immediately

coated.

Smooth

polystyrene

substrates _were

generated by spin coating

this

polymer

from toluene solutions _onto silicon wafers followed

by drying

ⁱⁿ a vacuum oven ; film

uniformity

and flatness _were confirmed

by X-ray

^reflection measurements. Silver surfaces _were obtained

by evaporating

500

A

_{of metal onto silicon wafers}

using

an E-beam metal evaporator. ^Polished

optical

quartz ^disks were cleaned under _an ultra-violet

lamp

before

being

coated.

PE-PEE and PE-PEP diblock

copolymers containing

a

crystallizable

block

(PE

melts _at

108

°C)

_were

spin

coated from hot

o-xylene

solutions onto heated substrates. PVCH-PEE and PE-PVCH diblocks

(PVCH

has _a

glass

transition _at 140

°C) [4]

were

spin

coated from

o-xylene

solutions at room temperature. ^PEP-PEE

films,

which _are characterized

by glass

transitions

near 56 °C and 20

°C, [5]

_were

spin

coated from

low-molecular-weight

alkane _or aromatic solutions at room temperature. ^{Film thickness could be controlled with} a

precision

of better than 50

A _by

a careful choice of

solvent,

solution concentration and

spinning speed.

All films _were dried in _vacuum for several hours after

coating

to remove residual solvent.

Films made from the

crystallizable polymers

that _were used for SIMS measurements were

heated above T~~j~ for PE, but below the ODT for the

copolymer,

and then

quenched

in

liquid nitrogen.

This

methodology

ensured that the

morphology

in the film _was

sufficiently

undistorted

by

the

crystallization

_process for

depth profiling

at room temperature

using

SIMS.

Neutron reflection measurements on films made from PE-PEE and PE-PEP diblocks _were

made in situ at 130 °C in a vacuum

chamber,

to ensure that both blocks _were in the

amorphous

state. Measurements _on PEP-PEE films _were made _at 25±2°C since the material is

amorphous

and well

annealed

_at

room temperature. ^{Films made from PVCH block}

polymers

were annealed for up to 24h _at 170 °C and then

quenched

_to the

glassy

_{state at room}

temperature ^{where all} measurements were made. The

profiles

in these

samples correspond

to those associated with the diblock _at the

glass

transition of PVCH.

2.3 INSTRUMENTATION. Neutron reflection

(NR)

measurements were

performed

at the National Institute of Standards and

Technology (NIST) using

the reflectometers

[41]

_on beam tube BT7

(with

fixed

wavelength,

A

=

2.35

A)

and beam tube NG7

(with

fixed

wavelength,

A

= 4.I

A). _{Background scattering}

was measured for several values of k

=

2 gr sin 0/A,

A~

~, the incident _momentum

perpendicular

to the surface where 0 is the

angle

of

incidence,

and _was subtracted from the data when it exceeded

m 10 9b of the total

specular reflectivity.

Measurement of film _{structure at} elevated temperature

using

NR _were made in situ

using

_a

vacuum

sample

chamber with controlled

heating

surfaces both behind and in front of the film

in order _to

optimize temperature uniformity.

Nominal film temperatures

reported

here _are

subject

to an error not

exceeding

10 °C _at

higher

temperatures. ^{For films that} were heated _to

higher temperatures

^for measurement on the neutron

beam,

a substantial

portion

of the

reflectivity

_{curve was}

repeatedly

monitored until it _no

longer changed

with time before

recording

the _curves used for

analysis.

Film thickness _was measured after

casting using

_a

Sopra~

ES4G

spectroscopic ellipsometer

at the Center for Interfacial

Engineering (CIE), University

of Minnesota _{or a}

Scintagg~ _X-ray

reflectometer _at

NIST,

with _an estimated _{error not}

exceeding

_± 10

A. Analysis techniques

to

extract film thickness _can be found in the literature for

ellipsometry [42]

and

X-ray reflectivity

[41].

For films with _an

incomplete

top

layer,

^the measured thickness

(t) corresponds

to an

average

single layer

as seen

by

the

ellipsometer.

Interference

micrographs

of the surface

topology

^{of films} were

acquired using

_an

Olympus~ light microscope

in _a reflection

configuration. Secondary

ions _mass spectrometry

(SIMS)

_was

performed

in the static mode

[43]

on some films _to characterize the

polymerlair

surface For

this,

_a Perkin

Elmer~

_PHI

6300 SIMS machine _was used with 7 kev Xe+

primary

ions rastered _{over a} 1.5 _x 1.5

mm~

area at a beam current of 0,15 nA. The _same machine _was used in the

dynamic

mode for

depth

profiling

measurements

[14]

on other films with 3 kev

O( primary

ions rastered _{over a}

700 _x 700 ~Lm~ area at a beam _current of 70 nA.

3. Results.

Figure

2 shows the surface

topology

of thin films of PEP-PEE _{cast on} silicon _at 23 °C. The

films _are close _to 3.35 and 4 bulk lamellar domains

(d~)

in thickness _as measured

by

fi§

O

.

.

~

~

~

10~m

Fig.

2. Interference

micrographs

of the surface of films of M~ 55,000 PEP-PEE

spin-coated

_on

hydrogen

passivated

^silicon. (a)

Sample

with measured

it)

=

385 _± 5

1

₄

db.

A

mainly

featureless surface is observed. (b) sample with measured

(t)

₌

1421

=

3.35 d~.

Light

areas

corresponding

to

/

it)

=

4d~ ^show up as islands on a darker

background

that

corresponds

to

it )

= 3d~. d~ ^{is the}

bull

lamellar period, estimated _as

3411

_{at 23 °C} from small angle neutron scattering [27, 40].

spectroscopic ellipsometry immediately

after

casting.

The

light

interference

micrographs

were

taken after the films _were allowed _to anneal at room temperature

(23 °C)

for two weeks. The film close _{to an}

integral

^number of

bilayers

in thickness shows _a

relatively

featureless surface

whereas the film with measured thickness close _{to a} half

integral

number of

bilayers

shows islands

(light

structures on dark

background)

that _{are one} lamellar domain in

height.

The contrast that makes the surface

topology

visible under _a reflection

microscope

arises from the constructive interference of

light

at two different colors for the top ^{and the base of the islands}

(lighter

at the top, ^where

it)

=

4

d~,

^and darker at the

base,

where

(t)

=

3

d~).

This result

clearly

demonstrates that the PEP-PEE diblock _assumes the

symmetric

thin-film geometry

depicted

in

figure

I with the _same block

(A

or

B)

situated _at both air and substrate interfaces.

For this geometry~ as discussed in the introduction

section,

the

deposition

of

non-integral

(a)

i

PEP

@

=

U ' ,

4J

qQ _~, '

4J

f~ ,

"

' ',

",

"-,

10~k

,

i~' (b)

_Air

si i

~PEE o

I

, ,

', /, i_{, ,} ,

, ,

,/,

/, i, i,

i ,

,

, ^' , , , 1 ^,

, , _, ^,

'

, , , , ,

, , , ^,

' , ^, , , ,

, , ,

, '

, , , , ,

~ ^{i , ,}

° 900 1800

distance from air surface,

A

Fig. 3. Example of fitting neutron reflectivity data to determine the surface-active block of _a diblock

copolymer

_system. (a) NR data

(open symbols)

for _a M~ ₌ ^{55, 000 PEP-PEE film}

spin-coated

_on

stripped

silicon. The solid line _is the fit corresponding to the optimized profile shown in (b) _{a~ a} solid _curve. The dashed line in (al

corresponds

to the

composition profile phase

shifted

by

180°, shown _as the dashed

curve in (b). The best fit to the data clearly shows that both film surfaces _are enriched by PEE.

amounts of material

((t)

_#

nd~,

where

(t)

is the average film thickness measured

by ellipsometry

and _n is _an

integer)

would lead _{to a} surface

topology composed

of holes and islands of

height d~,

as observed.

The _same result is _seen for all _our

polyolefin

diblocks cast on etched silicon each of _our

copolymers

forms the

symmetric

lamellar geometry ^where

(t)

=

nd~,

with the _same block

enriching

both air and substrate interfaces. This result is further substantiated

by

NR and SIMS

measurements as described below. For these

experiments

we have

prepared,

as far _as

possible,

integral

thickness films.

Determination of the surface active block in the diblocks _was

accomplished using

_neutron

reflectivity (NR)

measurements.

Techniques

for

analysis

of NR data, _as well _as the

sensitivity

of this

technique

to surface

composition,

have been discussed

by

us and other authors in separate

publications [17,

18, 41,

44].

To

recapitulate,

the NR data _can be

satisfactorily

fit

only

when the _correct

phase (I.e.,

surface-active

block)

is included in the

composition profile

used to simulate the

reflectivity.

As _an

example, figure

3b shows two

composition profiles

that

are used _to fit the data for PEP-PEE in

figure

3a _{; one}

profile

shows _an enrichment of PEE

(solid curve)

and the other shows _a

depletion

of PEE

(dashed curve)

at the films surfaces. The

corresponding

reflectivities calculated for these

composition profiles

are drawn _as the solid and

PVCH-@

b PE

-@

fl

C

ij it

I PE

@

PE

-fl

0 2 3 4 5 6

io~k

,

A-i

Fig.

4.

Experimental

neutron reflectivity curves (open symbols) ^for representative

polyolefin

diblock

copolymer

films _{cast on} etched silicon, listed in table III. Progressive vertical shifts of 4 orders of

magnitude

have been

applied

to the

plots.

Solid lines _are the calculated reflectivities for the optimized concentration

profile,

The surface-active block is outlined above each data _set. The temperature at which

each measurement was done is listed in table III.

Table III, Thin

film

characteristics,

Diblock Subsmte

Temp,°C

_102fl

(

^d

, n Surface #s

copolymer

block

A-B

PVCH(d8)-PEE tsj

^{25 4.4/3,8 230 4 PEE 0,96}

PEP-PEE(d~) tsi

25 8.1/3.8 348±5 5

PEE(d~)

0.90

'~ 25 ^'~ 330 3

PEE(d~)

0.87

"

Ag

25 ^'~ 325 5

PEE(d~)

^0.83

"

Quartz

25 '~ 330 5

PEE(d6)

^0.87

tsi 130 9.9/3.7 255 12 PEE 0.87

tsj ^{25 12,I/4A 168 8}

PVCH(d8)

0.94

tsi

130 9.9/6.7 535 6 PEP 0.72

tNafive odde removed usinga1:9HF:NII~OH ^e~h.

dashed _curves in

figure

3a. It is clear that

only

the

profile

with PEE

enriching

both air and Si lamellar

period

and the number of

periods

for each film. Also included in table III _are results for PEP-PEE films _{cast on a} number of substrates besides etched silicon. These

experiments

were

performed

to see if surface

segregation depends

_on the surface energy of the solid substrate. Neutron reflection data for these is not

presented

since

composition profiles

for these films matched the

optimized profile

in

figure

3b and the surface

segregation

is identical _to that

observed for etched silicon.

For diblocks where resolution _was

adequate,

direct

imaging

^{of film}

profiles

_was done

using dynamic

SIMS. In this

technique, primary O(

ions _are used _to sputter

through

_a

specimen

and

secondary

ion intensities recorded _{as a} function of time

(depth)

from the surface. The

interfaces

gives good

_agreement between calculated and measured reflectivities

(open

symbols).

The

plots

in

figure

4 show

reflectivity

data

(open symbols)

^{for films of the other} diblock

copolymers

listed in table II cast on etched

silicon, along

with calculated fits for

optimized composition profiles (solid

curves).

Table III summarizes the thin film characteristics extracted from the NR data

analysis

the value of p ^~ for each of the blocks at the temperature at which the NR data was taken, the type and volume fraction

#~

^of the surface-active block for all the diblock

films,

and the measured

secondary

H, D~ and Si~ ion

yields

for _two of the films _{as a} function of

sputtering

time is shown in

figures

5 and 6. Since _one of the blocks in each of _our

copolymers

is

partially deuterated,

the H~ and D~

signals provide

information about the

composition profile

of both chemical

species

in the thin films. The Si~

signal gives

_an estimate of the

point

in time that the substrate is reached after

sputtering through

the full thickness of the

sample. Figure

5 shows the SIMS

profile

for _a three

bilayer

PEP-PEE

(d~)

film _{cast on} etched silicon. The presence of

three oscillations in both the D~ and H~ spectra ^is clear, as is the enrichment of the deuterated

species (PEE)

at the silicon interface. Instrumental factors

beyond

_our control make the

etching

rate non-uniform in the first 100

A

_{of the film.}

Consequently,

the H~ and D~ oscillations _are not distinct in this

region

and it is not

possible

_to tell

directly

if there is _an H~ _{or a}

D~ maximum at the

polymerlair

interface. However, _a

comparison

of the

wavelength

of the

oscillations and the thickness of the

film,

_as measured

by ellipsometry,

makes it clear that there should be enrichment of D~

,

and therefore

PEE,

_at the air interface _as well.

Figure

^{6 shows} the SIMS

profile

for _{a two}

bilayer PE(d~)-PEP

film cast on etched silicon. For this

sample, following

the

example

above, ^{H~ maxima} are seen at both the air and substrate interfaces

which indicates that PEP is the surface-active

species

for this diblock. Both these results _are consistent with _our conclusions from NR for the _same diblocks.

For each of _our diblock films,

qualitative

confirmation of the surface-active

species

at the

polymerlair

interface _was also

accomplished using

static SIMS. This

technique

_{uses a}

primary

Xe+ beam _{at a} energy level less than that used for

dynamic profiling.

The

secondary

ions _are

collected

only

from about the top ¹⁵

A

_{of the}

_film,

_{and the}

presence of

mainly

deuterated _or

mainly hydrogenated

ionic

species

in the SIMS spectrum ^{indicates which of the blocks} enriches the air surface.

~ Si

e

I

B I

°

~

z

g

I

uttering

Fig. 5. SIMS

profile

at 23 °C of _{a ca.}

0001thick

_PEP-PEE

(d~)

diblock sample annealed 24 h _at

room temperature ⁱⁿ vacuum, shown _as secondary ion counts versus time. The vertical line _at

higher

etching times marks the

position

^of the

copolymer/silicon

substrate. The C~

signal

(not shown) remains

constant throughout the whole film, while H~ and D~ signals exhibit pronounced oscillations that

correspond _to lamellar domain~ oriented parallel to the film surface.

H

Si U I

U E

~ O

°

# fl E

~

D

Sputtering Time (sec)

Fig.

6. sIMs

profile

at 23 °C of _{a ca.}

10001thick

PE(d~ )-PEP ^diblock sample annealed 24 h _at

120 °C in _vacuum, shown _as secondary ion _{counts i>ersus} time. The vertical line at higher etching ^times

marks the position of the

copolymer/silicon

substrate.

4. Discussion.

Conventionally

surface enrichment in _a system ^with two chemical

species

has been understood in _terms of _a minimization of interfacial tension _{at a} system

boundary.

For

example~

in work

by

Russell and coworkers

[13-16],

_on

symmetric

PS-PMMA

diblocks,

_an enrichment of PS _at the air surface is

explained

in _terms of its lower surface

tension,

and the

segregation

of PMMA _at the silicon dioxide substrate is understood in _terms of its chemical

affinity

for this material. If

specific

chemical effects _are absent, and interfacial tension is associated

only

with

enthalpic (van

der

Waal's) effects,

the block with the

higher

surface energy is

expected

to segregate to the

(high energy)

substrate and the block with the lower surface tension is

expected

to

prefer

the free surface of the film. This follows from _a conventional definition of interfacial tension

between _a melt and _a solid

phase

that relies _on the difference between the surface tensions of the _two

phases.

An

example

of such _an effect _can be _seen in _recent work

by

Krausch and coworkers

[45]

_on mixtures of

poly(ethylene-propylene) (hPEP)

and the deuterated

homologue (dPEP)

cast as films _on etched silicon surfaces dPEP is found _to enrich the

polymerlair

interface, while hPEP is found to be surface-active _at the

stripped

Si surface. This behavior _can be attributed to the

slightly

lower cohesive energy

density

that characterizes the heavier

isotope [46]

and the fact that the solid surface has _a

higher

surface energy than either

polymeric species.

Table IV lists known values of the

dispersive

component of the surface energy

(y~)

of the chemical

species

present ⁱⁿ our

diblocks, along

with the substrates used for

spin-

Table IV.

Surface energies.

Yd,

PEP

30.6b

PEE (20-34)~

PE 35d

PUGH 35t

Hydrogen passivated

^44.7b

silicon~

> PEPf

> PEB

Silicon w1tll native

36.5b

oxide

Poly(styrene)

38d

Silver 1500'

a Native oxide removed with _a llydrofluoric acid etch.

b Reference 47. CAn experimentally determined value of the stirtace tension for PEE is unavailable in the literature. The lower estimate corresponds to the critical surface tension for closely packed methyl groups (Reference 53). The higher value is estimated using group contributions flom Reference 48, d Reference 49. 'Measured using contact angleswith water and methylene iodide, fReference 47, i Deduced flom isotopic segregation (See text).^h Reference 50.

coating.

The

dispersive portion

^{of the} surface energy is the relevant

quantity

in our

analysis

since

polar

and ionic forces _are

expected

to be

unimportant

in

simple

saturated

hydrocarbons.

For

hydrogen passivated silicon,

the substrate used in most of _our

experiments,

we are

interested in _a

quantitative

estimate of y~ as well _{as a} definitive _{answer to} the

question

does etched silicon have _a

higher dispersive

surface _energy than _our

polyolefin

melts ?

Comparison

of measured surface

energies,

_we

feel,

is _an

inadequate

_{way to answer} this

question

since

techniques

to measure the surface energy of solids and the surface tension of

polymer

melts are~ in

general,

different and

rely

_on different definitions of these

quantities [49].

^Part of the

answer is

provided by

^{Krausch and coworkers PEP} has _a lower surface energy than

stripped

silicon since dPEP and hPEP segregate to the

polymerlair

^and

polymer/silicon

interfaces of thin films,

respectively.

We

performed

_a similar

experiment using (h~)

PE with _a molecular

weight

of 1.8 _x

10~ mol/g,

and the deuterated

homologue (d~)

PE. A _ca. 4 000

ji

_{thick film}

was

spin-coated

at

high

temperature ^from a

o-xylene

solution that contained

equal weight

percent ^{of both}

polymers.

The film _was annealed _at 160 °C for 10 h under _vacuum and then

analyzed

at 23 °C with

dynamic

SIMS. The

resulting

H~ and D~

profiles,

shown in

figure 7,

H H

o~

fl

f

~ o

°

Si

# fl

3

D

i

Spunering Time (sec)

Fig.

7. SIMS

profile

at 23 °C of _{a ca.} 4

5001thick

_{film of PE and the} deuterated homologue, dPE cast _on

stripped

silicon. The film _was annealed at 160 °C for 10 h in _vacuum and then

quenched

in

liquid nitrogen

before measurement. The

asymmetric segregation

of dPE to the air surface and hPE to the silicon

substrate is

clearly

seen.

show

clearly

that hPE is

preferred

at the

polymer/silicon

interface and dPE _at the

polymerlair

interface. This establishes that

stripped

silicon has _a

higher

surface energy than PE _as well.

Segregation

based _{on an}

enthalpic

rationale alone _cannot

explain

the enrichment of both low energy

(polymerlair)

and

high

_energy

(polymer/substrate)

interfaces

by

the _same block. A

possible explanation

for _our

experimental

observations is that blocks that have different space

filling

characteristics will

experience unequal

conformational

perturbation

in the presence of _a

symmetry breaking

interface. This statement is illustrated

graphically

in

figure

8. Block A in _a

symmetric

A-B diblock

copolymer

has the _same molecular volume _as block

B,

_a condition that di1be cintrolled

by synthesis stochiometry.

This necessitates that the lamellar domains of size d will be

composed

of A- and B -rich subdomains that _are

equal

in lateral size

(d~

=

dB ).

If the

overall

density

remains constant

(which

it does to a close

approximation)

then the I-

dimensional division of space dictates this condition.

However,

the blocks _{are con-}

formationally _asymmetric

_as

captured

in the different

p

_parameter for each chain type

(Eq. (4)).

A

large

value of this parameter

signifies

_an

unperturbed

chain that is _more

conformationally expansive (I.e. occupies

less space per unit conformational

volume)

than a

chain that has a smaller p ^value. In this

example, p(

~

p(.

Thus B chains _are thicker and shorter than A chains. To preserve

incompressibility,

and the

equality

in subdomain

size,

chains from different sides of _an A-A interface will be _more intermixed than chains _{across a} B- B interface

(dotted

^lines in

Fig. 8).

^The creation of _a free

interface, therefore,

will _occur at

higher

conformational

penalty

at an A-A interface than at a B-B

interface,

and this

entropic

cost should be factored into the calculation of the overall surface free energy.

Here,

_we have

d~

~

d~

~

degree

of

interpenetration

between blocks from

neighboring

molecules. Such a

compositional

state would be _more

susceptible

to conformational

perturbation

than _a

strongly segregated

system ^with

sharp

interfaces. This raises an

interesting

issue. As _{X or} N

(I.e. XN) increase,

and the

degree

of

overlap

between

opposing

interfaces within microdomains decreases

(Fig. 9) [53],

the non-local

entropic driving

force for surface

segregation

should decrease. If _a

competing enthalpic

or local

entropic

factor exists, this could lead to an inversion in the

wetting tendency,

which would

provide

_a definitive test of _our concept.

~

~ _'l# ~~j$

~

Weak segregation Limit (WSL) strong segregation Limit (ssL)

zN

= 10

XN

» 10

Fig.

9. Schematic illu~tration showing decreasing overlap between like chains from

opposing

microdomains in _an A-B diblock copolymer as the product XN increases~ and interfacial widths become

narrower [53]. The influence of conformational asymmetry or preferential wetting should therefore

decrease _as the SSL is

approached.

~~ is the volume fraction of

siecies

A.

Segregation

based _on conformational asymmetry ^is

supported by

results from all five diblock systems ^for each _case the block with the smaller p value is surface-active _near the ODT

(see

Tab.

III). Particularly compelling

_are the results from the PVCH diblocks. In _one

instance, PE-PVCH,

the PVCH block is surface active~ while in _a second,

PVCH-PEE,

it is

not. In both _cases, the PVCH block is deuterium labeled while the other

polyolefin

block is

hydrogenous (Here

_we note that there is _no correlation between

isotope

content and surface

activity

in the results listed in Tab.

II).

A

comparison

of the relative

magnitude

of

enthalpic

and

entropic

components ^{of the surface}

free energy can be made

by considering

the PE-PEP system, ^{since the}

dispersive

and

conformational characteristics for both blocks _are well established

experimentally (Tab.

IV).

Enthalpically

driven

surface-activity

based _{on a} y~ value of _ca.

31dynes/cm

for PEP and 35

dynes/cm

for PE should lead _to the

(asymmetric) segregation

of PEP _to the air and PE _to the substrate interface.

Entropy,

on the other

hand,

favors the

(symmetric)

enrichment of the

conformationally

smaller PEP block _at both interfaces. Since the latter situation is observed

experimentally~

_we conclude that conformational

activity

is dominant for _an

~ value of 1.5

(given by

_~

=

p (~/p(~p)

and _a surface tension difference of 4-5

dynes/cm

between the blocks.

In

closing,

we make a brief remark about local

packing

effects

[52].

While the correlation between conformational asymmetry ^{and surface}

segregation

is rather

compelling,

we do not