NBSIR 75-760

Proceedings of Piezoelectric and

Pyroelectric Symposium-Worl(sliop

Martin G. Broadhurst, Coordinator

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

September 1 975

Final

Prepared for Office of Naval Research

Arlington, Va. 22044

U. S. Army Research Office

Durham, North Carolina 27706

NBSIR 76-760

PROCEEDINGS OF PIEZOELECTRIC AND

PYROELECTRIC SYMPOSIUM-WORKSHOP

Martin G. Broadhurst, Coordinator

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

September 1975

Final

Prepared for Office of Naval Research U. S. Army Research Office

Arlington, Va. 22044 Durham, North Carolina 27706

U.S. DEPARTMENT OF COMMERCE, Rogers C.B. Morton, Socretary

James A. Baker, III, Under Secretary

Dr. Betsy Ancker-Johnson, Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS. Ernest Ambler, Acting Director

I

\

PIEZOELECTRIC

AND

PYROELECTRIC

SYMPOSIUM - WORKSHOP

April 15 and 16, 1975

Sponsored by:

Office of Naval Research

Army Research Office

National Bureau of Standards

Lecture Room A

Administration Building

National Bureau of Standards

Gaithersburg, Maryland

PROGRAM

April 15. 1975

8:45 a.m. Welcome to NBS:

Dr. r.onald K. ?£by

Chief. Polymers Division

National Bureau of Standards

8:50 Introductory Remarks:

Dr. Kenneth Wynne,

Workshop Co-Chairman

Chemistry Program Office

Office of Naval Research

8:55 Guidelines for the Symposiun^-

Workshop:

Dr. Martin G. Broadhurst,

Workshop Co-Chairman.

Chief. Bulk Properties Section

Polymers Division

National Bureau of Standards

9:00 Technical Presentations

10:30 Coffee Break

10:45 Technical Presentations

1:00 p.m. Lunch

2:00 Technical Presentations

3:00 Coffee Break

3:15 Technical Presentations

5:00 End of Session

April 16. 1975

9:00 a.m. Technical Presentations

10:30 Coffee Break

10:45 "Specialty" Workshops

12:30 p.m. Lunch

1:30 Combined Workshops

3:00 Anticipated Conclusion of Conference

1
ORDER OF TECHNICAL

PRESENTATIONS

1. E. Fukada,

The Institute of Physical & Chemical Re- search, Japan

"Piezoelectricity in Polarized Polyvinylidene Fluoride"

2. J. Lando,

Case Western Reserve University,

Cleveland, Ohio

"Crystal Structures ol Polyvinylidene

Fluoride and its Copolymers"

3. R. Glen Kepler,

Sandia Laboratories, Albuquerque, New

Mexico

"Texture and Pyroelectricity in Polyvinylidene riuoride"

4. J. Ricca,

Army Materials & Mechanics Research Ctr., Watertown, Mass.

"Dependence of the Piezoelectric Activity

of Polyvinylidene Fluoride by High Speed

Uniaxial Stretching and Subsequent

Poling"

5. G. Pfister, Xerox Corporation, Rochester, New York

"Thermally Stimulated Currents and

Morphology in Polyvinylidene Fluoride"

6. G, T. Davis,

National Bureau of Standards,

Washington, D, C. "Piezoelectricity and Pyroelectricity in a Polyfvinylidene fluoride-tet/afluoroethylene) Copolymer"

7. J. Bergman,

Bftll Telephone Laboratories, Murray Hill,

New Jersey

"Polyvinylidene Fluoride as an Active

Device Element"

8. M. Litt. Case Western Reserve University,

Cleveland, Ohio

"Ferroelectric Model Compounds and

Polymers ( ? )"

9. S. Carr,

Northwestern University, Evanston,

Illinois "Origins of Persistent Electrical Polarizations of Polymer Solids"

10. R. Phelan,

Natidnal Bureau of Standards, Boulder

Colorado

"Pyroelectric Polymers Applied to Optical

Radiation Measurements"

11. H. Kawai,

Yokohama Municipal University,

Yokohama, Japan

12. M. Labes and R. Salomon,

Temple University, Philadelphia, Pa. "Pyroelectricity in PVF2"

13. J. Powers,

Naval Underwater Systems Center, New

London, Conn.

"Preliminary Investigations of Piezoelectric Polymers for Sonar Applications"

14. S. Edelman,

National Bureau of Standards

Washington, D. C.

"Piezoelectric Polymer Measuring Instru- ments"
TABLE OF CONTENTS

Page

Attendees List 1

Letter of Invitation 5

Introduction 6

"Piezoelectricity in Polarized Polyvinylidene Fluoride" 8

by E. Fukada and Hitsumasa Oshiki

"Piezoelectric Polymer Research in Japan" 34

by E. Fukada

"The Melting Temperature-Composition Curves of Poly (vinylidene fluoride) -Poly 35

(vinyl fluoride) Mixtures and Vinylidene Fluoride-Vinyl Fluoride Copol3miers"

by J. B. Lando

"Texture and Pyroelectricity in Polyvinylidene Fluoride" ...... 46

by R. G. Kepler

"Dependence of the Piezoelectric Activity of Polyvinylidene Fluoride Upon High 59

Speed Uniaxial Stretching and Subsequent Poling"

by R. J. Shuford, A. F. Wilde, J. J. Ricca and G. R. Thomas

"Thermally Stimulated Currents and Morphology in PVF2" 96

by M. Abkowitz, P. J. Luca, G. Pfister and W. M. Prest, Jr.

"Piezoelectricity and Pyroelectricity in a Poly (vinylidene f luoride-tetra- 120

f luoroethylene) Copol3nner"

by G. T. Davis

"Polyvinylidene Fluoride as an Active Device Element" 138

by J. G. Bergman

"Scanning for Ferroelectricity in Polycrystalline Materials" 165

by M. Litt, Che-hsiung Hsu, P. Basu and S. M. Aharoni

"Origins of Persistent Electrical Polarization in Poljnner Solids" 189

by S. H. Carr

"Pyroelectric Polymers Applied to Optical Radiation Measurements" 196

by R. J. Phelan, Jr.

"Pyroelectricity in PVF2" 199

by R. E. Salomon and M. M. Labes

"Preliminary Investigations of Piezoelectric Polymers" 209

by J. M. Powers

"Piezoelectric Polymer Measuring Instruments" 210

by S. Edelman

iii

- 1 -

ATTENDEES LIST

Piezoelectric and Pyroelectric Symposium-Workshop

April 15-16, 1975

Louis A. Abbagnaro

CBS Laboratories

227 High Ridge Road

Stamford, Conn. 06905

Stephen K. Carr

Northwestern University

2145 Sheridan Road

Evanston, Illinois 60201

William S. Barnhart

PennWalt Corporation

900 First Avenue

King of Prussia, Pa. 19406

Mario J. Cellarosi

National Bureau of Standards

Matls. Bldg. Rm. B322

Washington, D. C. 20234

John Bergman

Bell Telephone Laboratories

Holmdel, N. J. 07733

George L. Boyer

Office of Naval Research

800 N. Quincy Street

Arlington, Va. 22217

Phillip Bloomfield

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

Martin Broadhurst

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

Tony Clement

e

Bolt, Beranek, Newman

Cambridge, Mass. 02134

G. Thomas Davis

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

A. Dereggi

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

J. F. Ditter

Chemical Systems

1852 McGaw Avenue

Irvine, Calif. 92705

Robert B. Bunker

AFML

6500 Christy N-E

Albuquerque, New Mexico 87109

Anthony J. Bur

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

Ronald K. Eby

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

S. Edelman

National Bureau of Standards

Polymers Division

Washington, D. C. 20234
Barry Farmer

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

John Ferraris

National Bureau of Standards , Polymers Division

Washington, D. C. 20234

Richard Ferren

Pennwalt Corporation

900 First Avenue

King of Prussia, Pa. 19406

Robert B. Fox

Naval Research Laboratory

Washington, D. C. 20375

Freeman W. Fraim

Thermo Electron Corporation

101 First Avenue

Waltham, Mass. 02154

Eiichi Fukada

The Institute of Physical

and Chemical Research

Wako, Saitama, JAPAN

L. Garn

Night Vision Laboratory

Ft. Belvoir, Va. 22060

Daniel E. Gilbert

Research Triangle Institute

P.O. Box 12194

Research Triangle Park, N. C. 27709

John K. Gillham

Princeton University

Dept. of Chemical Engineering

Princeton, N. J. 08540

Thomas A. Giordano

CBS Laboratories

227 High Ridge Road

Stamford, Conn. 06905

V. S. Goel

Nuclear Regulatory Commission

Bethesda, Maryland

A. M. Glass

Bell Telephone Laboratories

Mountain Avenue

Murray Hill, N. J. 07974

James R. Griffith

U.S. Naval Research Laboratory

Washington, D. C. 20375

Jay Charles Hicks

Naval Undersea Center

San Diego, Calif. 92132

R. Glen Kepler

Sandia Laboratories

Organic Materials R&D Dept. 5810

Albuquerque, New Mexico 87115

F . Khoury

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

M. M. Labes

Temple University

13th & Norris Streets

Philadelphia, Pa. 19122

Jerome B. Lando

Case Western Reserve University

Dept. of Macromolecular Science

University Circle

Cleveland, Ohio 44106

Morton Litt

Case Western Reserve University

Dept. of Macromolecular Science

University Circle

Cleveland, Ohio 44106

Dick Madden

Bolt, Beranek, Newman

Cambridge, Mass. 02134

-2-
Rudolph J. Marcus

Office of Naval Research Branch Office

1030 East Green Street

Pasadena, Calf. 91106

G. Kirby Miller

GTE - Sylvania

Box 188

Mountain View, Calif. 94042

Jean A. Mo'ntemarano

Naval Ship R&D Center

Annapolis, Maryland 21402

Fredrick- Mopsik

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

Preston V. Murphy

Thermo Electron, S.A.

Infanta Carlota 32

Barcelona, SPAIN

Raymond Naar

Tufts University

Pearson Laboratory

Medford, Mass. 02155

George Neece

Office of Naval Research

800 N. Quincy Street

Arlington, Va. 22044

John H. Parker

903 Meadow Lane

Vienna, Va. 22180

Leighton H. Peebles, Jr.

Office of Naval Research

495 Summer Street

Boston, Mass. 02210

Robert J. Phelan, Jr.

National Bureau of Standards

Electromagnetics Division

325 Broadway

Boulder, Colorado 80302

Gustav R. Pfister

Xerox Corporation

Research Center

Webster, New York 14580

Ted 0. I'oehler

Johns Hopkins University

Applied Physics Laboratory

Baltimore, Maryland 21218

James M. Powers

Naval Underwater Systems Center

New London Lab

New London, Conn. 06320

William M. Prest

Xerox Corporation

800 Phillips Road

Webster, New York 14582

Peter F. Radice

Pennwalt Corporation

438 Springhouse Road

King of Prussia, Pa. 19406

Joseph P. Reardon

Naval Research Laboratory

Code 6170

Washington, D. C. 20375

Kennard Reynard

Horizons Inc.

23800 Mercantice Road

Cleveland, Ohio 44060

John Ricca

Army Materials & Mechanics Research Ctr.

Watertown, Mass. 02172

D. C. Robinson

National Bureau of Standards

EM 219

Washington, D. C. 20234

Steven Roth

National Bureau of Standards

Polymers Division

Washington, D. C. 20234

-3-
Charles F. Rowell

Office of Naval Research

Chicago Branch Office

536 S. Clark Street

Chicago, Illinois 60605

Barrie S. H. Royce

Princeton University

Materials Laboratory

D 416 Duf field Hall (E.Q.)

Princeton, N. J. 08540

Robert E. Salomon

Temple University

Chemistry Dept.

13th & Norris Streets

Philadelphia, Pa. 19122

E. Sharp

Night Vision Laboratory

Ft. Belvoir, Va. 22060

Richard J. Shuford

Army Materials & Mechanics Research Ctr.

Arsenal Street

Watertown, Mass. 02172

Joseph Simmons

Catholic University

Washington, D. C. 20017

Davis R. Squire

U.S. Army Research Office

Box CM, Duke Station

Durham, N. C. 27706

Charles Smyth

Princeton University

Princeton, N. J. 08540

E. M. Stanley

Naval Ship R&D Center

Annapolis, Maryland 21401

Alan Taylor

3M Company

St. Paul Minn.

C. C. Walker

Naval Sea Systems Command

Dept. of the Navy

Washington, D. C. 20362

Hugh Wright

Bolt, Beranek, Newman

Cambridge, Mass. 02134

Kenneth J. Wynne

Office of Naval Research

Chemistry Program

800 N. Quincy Street

Arlington, Va. 22044

-4-
UNITED STATES DEPARTMENT OF COMMERCE

National Bureau of Standards

Washington, D.C. 2Q234

LETTER OF INVITATION

We are in the process of organizing a symposium-workshop to be held at

the National Bureau of Standards on April 3 5-16, 1975. The meeting is

jointly sponsored by the Office of Naval Research, the Army Research

Office and NBS.

The purpose of the meeting is to bring together scientists and program

managers from DOD, university laboratories and industrial laboratories

involved in research on and manufacture of piezoelectric and pyroelectric

polymer materials and devices. He intend to exchange current information

on research and development activities in order to stimulate new efforts

in this field and promote the development and application of polymers for

transducer applications.

We look forv/ard to your reply to this invitation and hope you will be able

to accept. Included is a list of persons who have been invited and we welcome the names of other key people whom you believe should also be

included. General information about food, housing, transportation and

meeting rooms will be sent to you at a later date along with a meeting

schedule.

Sincerely yours,

MARTIN G. BROADHURST

Chief, Bulk Properties Section

Polymers Division

Enclosure

- 5
Introduction

This report contains manuscripts of talks that were presented at a

Symposium-Workshop on Piezoelectric Polymers which was held at the National

Bureau of Standards, Gaithersburg , Maryland on April 15 and 16, 1975. The

meeting was initiated by Dr. Kenneth Wynne, Chemistry Program Office of the

Office of Naval Research and the program vjas arranged by Dr. Martin Broadhurst,

Polymers Division, NBS. The meeting was cosponsored and supported by OMR,

Arlington, Virginia, ^^^^^ A::my lasaarcli Office, :our*iam„ lort-i Carolina, and

The purpose of this Symposium-Workshop was to bring together scientists

and program managers from federal (primarily DoD) , industrial, university, and

other laboratories who are involved in the research on and manufacture of

piezoelectric and pyroelectric polymer materials and devices, in order to

exchange information about and stimulate further work on the new and rapidly

developing measurement technology involving polymer transducers. This meeting

is particularly important at this time because highly active polymer films are

not yet available to device manufacturers in the U.S. or to potential users

such as DoD. Potential manufacturers are reluctant to produce polymer films

for this purpose because of uncertainty of their market value. Both manufacturers

and users are reluctant to invest heavily in polymer transducers because of a

lack of such important information as the nature of the underlying mechanism,

transducer reliability and life-time, ways of optimizing polymer sensitivity

and fabrication procedures.

This symposium is part of the ONR annual program review, the first session

of which was held at ONR in Arlington, Virginia on April 14. Only those papers

of most general interest in the area of piezoelectricity are included in this

Proceeding and other papers dealing v/ith potentially useful new polymers will

be compiled and distributed separately.

- 6 -
Offsetting these uncertainties is the growing awareness that simplicity

of fabrication, economy, flexibility, toughness and high sensitivity make

polymers attractive candidates for novel devices. In fact, piezoelectric

pcl^Tner headphones from Japan and pyroelectric pol>'^er radiation sensors in

the U.S. are already in connercial production anc potential new applications

ranging from patient pulse-monitoring transducers to ship hull coatings to make

lar~3 aperatura sonar ant annas ^-era discussed at t'r.e meeting.

Following a series of talks which are represented by the manuscripts

on the following pages, a timely workshop discussion confirmed that many

measurement and developzient problems are 3^"et to be overcome.

In order to speed up distribution of these Proceedings, the follo^ring

papers were reproduced as received from the authors with only minor modifications. Since these papers did not meet the uniform.ity requirements for

reproduction as an official NBS publication, this report is to be considered

as an informal account of the Symposium and x^rill be distributed to Symposium

attendees and other interested parties and to Do! as previously agreed.

Some of the work has been published previously elsewhere and some of

the papers are being submitted by the authors for publication elsewhere.

Preferences to information in this report should be r^de to:

Nat. Bur. Stand. (U.S.), Interagency Report,

jTBSIR 75-760, page nimber, (September 1975).

Copies in paper or microfiche form may be obtained from the National Technical

Information Service, (NTIS) , Springfield, Va . 22151.

Commercial materials are identified in this paper to adequately specify

the experimental procedure. Such identification does not imply recommendation

or endorsement by the National Bureau of Standards.

-7-
Piezoelectricity in Polarized Polyvinylidene Fluoride

Eiichi Fukada and Mitsumasa Oshiki*

The Institute of Physical and Chemical Research

Wako, Saitama 351, Japan

Abstract

The mathematical expressions for the complex piezoelectric stress-constant, e„, = e, ' - ie^,", the complex

31 31 31

electrostrictive constant, = <2i^

" ^^3l"

complex dielectric constant, £ = e

' - ie ", are derived

3 3 3 . by non-equilibrium thermodynamics. The basic assumption

is that the orientation of dipoles produces the residual

polarization. The determination of these quantities has

been made for polyvinylidene fluoride films over a

temperature range from -100°C to 100°C at a frequency

of 20 Hz. The residual polarization after heat, treatment

is found to be constant independent of temperature.

Introduction

The origin of the piezoelectricity in the stretched

and polarized films of polyvinylidene fluoride has been

discussed by many investigators A plausible explanation

is that the polarization originates in the stress-induced

reorientation of dipoles which have been preferentially

alligned in the direction normal to the film surface due

to stretching and subsequent poling procedures. Another

mechanism proposed is that the polarization is induced by

•^Department of Physics, Gakusimin University, Mejiro, Tokyo

-8-
the inhomogeneous strain coupled with the distribution of

space charge inside the film.

In this paper first we shall derive theoretical

expressions for the piezoelectric constant and electrostrictive constant in the electretized polymer and then

present some experimental data on the temperature variation

of some physical properties related to piezoelectricity forj

pdlyvinylidene fluoride electret film.s.

Theory

It has been demonstrated that the piezoelectric and

electrostrictive constants of polymers are complex quantities

and are represented as e = e' - ie" and < = <' - Ik".

The piezoelectric stress-constant, e, can be determined

either by measuring the polarization ,? ,under a given strain

f3Pl S , under short-circuited conditions,

measuring the stress, T, under a given electric f ield ,E , In the equilibrium state,

the equality of these two derivatives has been proved from

f 3T1

at the clamped condition, - ' —

12 the thermodynam.ic principles;

The electrostrictive constant is determined either by

measuring the stress proportional to the square of the

electric field, - -j j , or by measurinc the strain

dependence of the dielectric constant, -r-^! . The equality

of the equilibrium values of the electrostrictive constant

derived from, these two different observations has also

been proved by the thermodynamic principles!"^

-9-
Now, the piezoelectric constant for polymers is

defined as a complex quantity, i.e./ e = e' - ie". The

equality of this complex quantity measured by the direct

and converse piezoelectric effects, is not self-evident.

The equality of the complex electrostrictive constants

determined by the two different methods, is not proved,

either.

In order to derive the theoretical expressions for e

and K as complex quantities, we have to use non-equilibrium

thermodynamics^^ First we shall introduce the Helmholtz

free energy, F, as follows

F = (S-VP2)2 +la{P-PQC)2 +|3(p-p^^)2 +1x^2. (1)

The first term shows the elastic energy taking into account

the electrostrictive strain, vP2 , which is proportional to

the square of polarization p.. is the elastic constant

at P=0. The second and third terms show the dielectric

energy involving two internal variables ^ and z, . f; represents

an ordering parameter which changes slowly with time

during the poling process under a static electric field.

PqC is approximately the residual polarization, P^, after

poling. Pq is the polarization per unit of ^ or c» Since static polarization in an electret decays with time

very gradually, the state with P^ is not thermodynamically

stable. Thus we introduce the fourth term which represents

the increase of free energy due to f; . The third term in F

is the dielectric energy, which involves another ordering

parameter, l. , which varies in time scale of a rate comparable

-10-
with the period of the alternating excitation, under which

the dielectric, piezoelectric and electrostrictive measurements are made. When an external excitation is not

present, z, is given by P^/P^ . In eq. 1^ a and 6 are

coefficients related to the dielectric constant.

If we expand the first term of eq. 1 and neglect

4 the P term, then we have

F = |gPs2+ la(P-PoC)2 + |e(p-p^^)2 - ^SP2 + |xc2 (2)

where y=G'^v .

"

_ From eq. 2 we derive the stress, T, and the field, E, as

T =

E =

[8SJ

'9F

= GPS - YP2 (3)

S,C.?

= cx(P-P^C) + B(P-Po?) - 2ySP (4) 3P

The affinity, , conjugate to ^ is given by

'3F'

S,P,C = BP,(P-P,0. (5)

Solving for P from eq. 4, putting E=0 and T=0, and using the

approximation 2YS/(a+6) <<1/ we find that the residual

polarization, Pj. , and the residual strain, , are

Sr = YPr^GP . (7)

Now we shall consider the case when an external

excitation is applied to a sample which is specified by

Pj. / Sj. , Co ' • Since E, changes only during the poling

process, ^ will be considered a variable during the

measurement. The polarization, P, and the stress, T, are

expanded as functions of E, S, and ^ at P^ , S^ , Cq • The results are

2yPj. bp , 2y6P

AP =

-f^- + -TTp- AS + —I Ai; +^ EAS + ^ A^S (8)

-11-
4y^Pr^ 2YPr 2^6?

2

^ ^ Ac'

(a+3) (a+3)

Sy'Pr Sy^ePoPr 2y6Po

_ EAS ASAC r EAc (9)

(a+3) (a+3) (a+3)

In eqs 8 and 9, the expression for Ac is unknown.

The affinity, A^, is zero when C = Co* "^^^ temporal

»change of c is given by the following equation, introducing

a kinetic coefficient y,

AC = pA^ . (10)

Substituting eq 5 into eq. 10 we have

2yPj. 3Pq

AC = 3yPo{^ + AS +

^

AC

2y 2y3Po

+ — EAS + — ACAS - PqAc) dD

(a+3) (a+3)

For the sinusoidal excitations, we put

E = E^(e^^Et ^

^-ia)Etj

AS = S^(e^'^st ^

^-iiostj

Combining eqs. 10, 11 and 12, we have

E^e^'^Et 2YPrS,e^'^st

Ac = rT r +

(12)

aPo (l+iw^T ) aPo (l+icosT

)

aPo (a+3 ) 1 1+i (oje+^s ) x ) ^ a l+ia3ET

aPo (a+3) (l+i {iii^-i^s)T} a l+iojgX

+ AC (13)

-12-
where represents the complex conjugate of the previous

four terms. We have introduced a relaxation time,T/ which

is defined by

T = ., (14)

Introducing eq« 12 and 13 into eq. 8,

e' AP = (1 + ^ Yx^i )E.e a+3 a l+iojpT ^

(1 + 1- ^

)S,ei'^st

+

^

(1 + i - wi + ^ . - ^E s Q^(^E+^

(a+3) 2 a l+iojET^^-^ a 1+i (0)^+0)3) t^ a^a^

+ , (1 + ^ T-i^ ){1 + ^ Tir-r^- r-}EAS.ei('^E-^

(a+3) °^ I+10J2T a 1+1(03^-103)1 A A

+ AP (15)

1 Using P^'^e denote the amplitude of AP with the angular

frequency co^/the complex susceptibility is given by

and the complex dielectric constant is given by

e* = 1 + 4^ (1 + ^ YTT ) • (17) a+3 a I+iojeT

Using P ^ denote the amplitude of AP with angular , frequency / the complex piezoelectric stress-constant

is given by

Using P

"^^ P to denote the amplitudes of AP with angular

'

frequencies (uj^+Wg) and (to^-Ws^' complex electrostrictive

constant is given by

-13-
(k-)* = 47TPa~/EaS

= {1 + i 1 }{1 + i 1 , ^9)

Wjith st.itic strain (aig=0) , the electrostrictive constant

is expressed as

^* §£L (1 + B i_ )2 ,20)

* Next, from eq. 9, we obtain

a + 3 a l+i(jL)gT ^ A

2yP

a+3 ^ a l+ia)ET^^A®

(a+3) 2 °' I+io^et' a

+ T (21)

where T denotes the complex conjugates and higher terms.

Using T^^^ to denote the amplitude of stress with the angular

frequency cog, the complex elastic constant is given by

«5 )* = ^= ° - ^+I-<1 * I IT^^ '22)

Using to denote the amplitude of stress with the angular

frequency cog , the complex piezoelectric stress-constant is

given by

= - -eT" = ^ ^ I ra^' . . Using T^^'^^ to denote the amplitude of stress with the

angular frequency 2a)g , the complex electrostrictive

constant is

-14-
8ttT

Eq. 18 and eq. 2 3 show that the values of e* derived

from both the direct piezoelectric" effect and the converse

piezoelectric effect are equal complex quantities if the

exciting frequencies are the same. Eq. 2 0 and eq. 24 show

that K* derived from the strain-dependence of the

dielectric constant is equal as a complex quantity to that

derived from the component of stress proportional to

square of the electric field.

The following relations, which are known to hold for

equilibrium quantities, also apply to complex quantities.

(GE)* = - -^i- (25)

p. = ^^^

(26)

The latter equation is used to calculate from

experimental results.

Experiment

We used polyvinylidene fluoride films subjected to the

three different treatments shown in Table 1. The original

film was supplied by the Kureha Chemical Co. . The (n-n)

samples were prepared by annealing non-poled commercial

film at 135°C for 12 hours. The (s-n) samples were prepared

by stretching the (n-n) samples to 5 times their original

length and annealing them at IBCC for 12 hours. The (s-p)

samples were prepared by poling the (s-n) film at 85°C for

2 hours with a static electric field of 500 kV/cm. Infrared

absorptions observed at 510 cm~^ and 530 cm"^ shows

-15-
that the (n-n) samples contain both I and U crystalline

forms, but that the (s-n) and (s-p) Samples contain almost

entirely I crystalline form. The dimensions of the test

Scunple are about 15x10x0.05 mm. The length is the direction

of elongation. Silver was evaporated as electrodes on the

central area, 8x5 mm, on both surfaces of the sample.

A schematic diagram of experimental apparatus is shown

in Fig. 1. Both ends of the sample are clamped inside a

thermostat so that the uniaxial tension is applied in the

direction of length, which accords with the direction of

elongation. An a.c. voltage with the angular frequency o)

is applied between the electrodes of the sample, and the

resulting tensile stress of the sample in the length direction

is detected by a PZT ceramic element. The output of the

PZT detector is led to a charge amplifier and four phase

sensitive detectors, whose reference voltages have an

angular frequency of uj or 2a) and a phase angle of 0 or

tt/2 with respect to the output voltage of oscillator.

The outputs from these phase sensitive detectors give e',

e", k', and k". The real and imaginary components of the

dielectric constant, e', e", were also determined simultaneously

with the above quantities.

When the electric field is applied across the

thickness of the film, the tensile stress T^ produced

in the stretched direction is given by

were the third term on the right hand side represents the

Maxwell stress in the film and a is the Poisson's ratio.

16-
Thus, = 83^' - ie3^", K3, = K3i' - iK3i", and

£3 = £3' - ^£^3" ^s^s determined simultaneously as [' functions of temperature at a frequency of 20 Hz. |

• Fig. 2 shows the temperature dependence of £3' and

£3" for three kinds of samples. It is seen that the

stretching of the sample increases £3 ' and shifts the

dispersion to a higher temperature, and that the poling

process decreases £3 ' . This dispersion is caused by the

the

glass transition in^v^non-crystalline phase of polvinylidene

fluoride.

Fig. 3 shows the temperature dependence of ^3^^' ^^'^ ^or the three kinds of samples. As with [£3 ' / stretching increases ^3^^' and the poling decreases K3j^'«

According to eq. 19, k is the variation v/ith strain of

the dielectric constant. The density of the sample is

changed by j the strain. The variation of the

dielectric constant due to the variation of density can

be calculated from Clausius-Mossotti ' s equation.

This correction term was added to the observed value of

K in the subsequent calculations of the residual polarization.

Fig. 4 shows the temperature dependence of ^3^ ' and

63^" for the stretched and poled sample during the increase

of the temperature. The second run of the measurements gave

decreased values for e but the third run gave the

31

same results as the second run.

Using these experimental data, the residual polarization

Pj. was calculated according to eq. 26, The temperature

-1 7-
dependence of is given in Fig- 5. For the first run,

the decrease of was observed at the glass transtion

temperature. However, at the second and third runs, P^

was almost constant independent of temperature. The

decrease of P^ during the first run may indicate a weak

polarization in the non-crystalline phase. The persistent

observed at the second and third run may be caused by a

strong polarization associated with the crystalline phase.

We have undertaken the measurements of e , < and £ 31 31 3

during the poling process. If a d.c. bias field, E , is

D

applied to the film, the relation between the stress T^

and the field E^ is given by the following formula.

From eq. 27, we can see that there is a contribution to

the apparent piezoelectric constant due to the presence

of in addition to the ordinary piezoelectric constant

e . This piezoelectric constant observed with a d.c.

31

bias field, e ^, is given by

ob

Fig. 6 shows the time variation of e ' and e

" ob ob

during poling at 85 °C with a field of 200 kV/cm.

The lower figure shows the values of e ' and e

" 31 31

derived from eq. 28. It is seen that the piezoelectric

e-constant increases in two steps with relaxation times of

approximately 1 minute and 100 minutes.

Fig. 7 shows the temporal change of k^^ and

determined simultaneonsly with e . It is seen that k and

ob 31

-18-
decreases in two steps with relaxation times of

approximately 1 minute and 10 0 minutes.

Fig. 8 shows the temporal change of P^. calculated

from ^21' ^31' ^3 ^2^^*? 26. It is obvious that

Pj. increases in two steps with relaxation times of approximately

1 minute and 100 minutes . We suppose that the first

step with a relaxation time. of about 1 minute may

be due to the formation of polarization in the non-crystalline

phase either by dipole orientation or by space charge

injection. The following step may be due to the formation

of polarization in the crystalline phase or at the crystal

boundary surfaces, where the reorientation of dipoles or

the trapping of charge is induced.

Fig. 9 illustrates the decay of the piezoelectric

d-constant of electretized polyvinylidene fluoride films

left at room temperature. After an initial decay with a

relaxation time of approxim^ately 5 days , the d-constant

remains unchanged for about 10 months. The persisting and

constant piezoelectric effect could be caused by the

residual polarization associated with the crystalline phase^

I - of the polymer. 1

-19-
Summary

The mathematical expressions for the piezoelectric

stress-constant, ^3^/ the electrostrictictive constant,

dielectric constant, e^, all as complex

quantities, have been derived by means of non-equilibrium

thermodynamics on the assumption that the dipole orientation

produces a residual polarization. The temperature

dependence of these quantities for electretized

polyvinylidene fluoride films has been experimentally

determined at the same frequency of 20 Hz over a

temperature range from -100°C to lOCC. The residual

polarization calculated from these experimental data is

about lO^esu/cm^ independent of temperature. This polarization

probably originates from the alligned dipoles associated

the'

withy^crystalline phase. Measurements during poling

process show a gradual increase of the residual

polarization.

-20-
References

1. R. Hayakawa and Y. Wada, Adv.' Polym. Sci. 11, 1, (1973)

2c H. Kawai, Japan. J. Appl . Phys. S_, 975, (1969)

3. M. Oshiki and E. Fukada, J. Mat. Sci. 10, 1, (1975)

4. M, Tamura, K. Ogasawara, N. Ono, and S. Hagiwara,

J. Appl. Phys. 45, 3768, (1974)

5. N. Murayama and M. Segawa, J. Petroleum Soc . Japan

18, 77, (1975) (in Japanese)

6. M. C. Broadhurst, C. G. Malmburg, F. I. Mopsik and

W. P. Harris, Electrets, Charge Storage and Transport

"in Dielectrics (Ed. M. M. Perlman, The Electrochemical

Society) p. 492, (1973)

7. J. Cohen, S. Edelman and C. F. Vezzetti, ibid p. 505

(1973)

8. G. Dreyfus and J. Lewiner, ibid p. 517 (1973)

9. H. Sussner, D. Michas, A. Assfalg, S. Hunklinger and

K. Dransfeld, Physics Letters 45A, 475, (1973)

10. J. G. Bergman Jr. , J. H. McFee, and G. R. Crane,

Appl. Phys. Letters IQ^, 203, (1971)

11. G. Pfister and M. A. Abkonitz, J. Appl. Phys. 45 , 1001, (1974)

12. J. F. Nye, Physical Properties of Crystals,

Oxford Univ. Press (1957)

13. W. P. Mason, Piezoelectric Crystals and their

Application to Ulltasonics D. van Nostrand (1950)

14. S. R. DeGroot and P. Mazur, Non-equilibrium thermodynamics.

North Holland Co. (1962)

-21-
Table 1

Sample Treatment

n-n as-received film was

annealed at 130°C for 12 hrs

s-n n-n film was stretched by

5 times at 65°C and annealed

at ISO'C for 12 hrs

s-p s-n film was polarized

at 85*0 for 2 hrs with

a static electric field

500 kV/cm

-22-
Figure Captions

Fig. 1 The schematic diagram of the apparatus for

determining the dielectric constant e = e' - ie" , the

piezoelectric stress-constant, e = e' - ie" , and the

electrostrictive constant, k = k' - iK*'/ at a frequency of

20 Hz and over a temperature range from -100 "C to 100 **C.

Fig. 2 The temperature dependence of the dielectric

* ... constant, e = e ' - ie " , for differently treated films of

3 3 3

polyvinylidene fluoride (as to notations, see Table 1)

Fig. 3. The temperature dependence of the electrostrictive

constant, = k^i' " ^"^ai"'films of polyvinylidene

fluoride.

Fig. 4 The temperature dependence of the piezoelectric

stress-constant, e^^ = e^i' ~ ^®3i" ®lo^9^ted and

polarized film of polyvinylidene fluoride. After the

first run of measurement, during which the temperature is

increased from -lOO^C to 100 '^C, the magnitudes of e^i'

e^^" decreased, but the second and third runs of measurement

gave the same results.

Fig. 5 The temperature dependence of the residual

polarization, P^. = ®3]^(£3 ~ ^^^^^l' ^'^^ electretized

film of polyvinylidene fluoride. After the first run of

measurement^P^' decreased but the second and third runs gave

the same results, where P^ was independent of temperature.

-23-
Fig. 6 The time dependence during the poling process of

about 10 hours of the piezoelectric stress-constant,

%h = %b' " ^^ob" = ^31 ^'^31 ^^3)V^^' apparently

observed under a d.c. bias field E^^ . The values of

e^^ = " ^®3i" '^^^^ corrected using the values of

K^^ and shown in Fig. 7. The poling field E^^ is

200 KV/cm and the poling temperature is 85**C.

Fig. 7 The time dependence during the poling process of

about 10 hours of the electrostrictive constant,

=

<2i' " ^'^31 ^^'^ dielectric constant,

£3 = £3'- ie3".

Fig. 8 The time dependence during the poling process of

about 10 hours of the residual polarization calculated

from = 63^(^3 - l)/<^^.

Fig. 9 The aging of the piezoelectric strain-constant,

d3^ = ^3^^' - id3j^", for elongated and polarized films of

polyvinylidene fluoride for a period of 300 days.

-24-
L J o a

a, L X t J

Q Q a a

to tr m in

0< . 04 On

a

Am

-25-

' '

« -100 -50 0 50 100

Temperature ( )
-100 -50 0 50 100

Temperature ( )

-28-
lOr

Ql , , , 1_

-50 0 50 100

Temperature ( "C )

-29-


CO

(^uuD/nsa^Ol ) -'d

Piezoelectric Polymer Research in Japan

Unscheduled Presentation by E. Fukada

Following his scheduled presentation Dr. Fukada kindly showed slides

and discussed recent results and current theories of several other Japanese

scientists who have made substantial contributions to' the field of piezoelectric polymers. These scientists who were represented by Dr. Fukada were:

Professor H. Kawai

Department of Physics

Yokohama Municipal University

Yokohama, Japan

Dr. N. Ilurayama

' Nishiki Research Laboratories

Kureha Chemical Industry

Iwaki-shi, Fukushima-ken, Japan

Dr. H. Ohigashi

Basic Research Laboratories

Toray Industries, Inc.

Tebiro, Kamakura 248

Dr. M. Tamura

Acoustical Engineering Research Laboratory

Pioneer Electronic Corp.

4-2610, Hanazono

Tokorozawa, 359, Japan

Dr. Y. Wada

Department of Applied Physics

Faculty of Engineering

University of Tokyo

Bunkyo-ku

Tokyo , Japan

- 34 -
THE MELTING TEMPERATURE-COMPOSITION CURVES OF

POLY(VINYLIDENE FLUORIDE) -POLY (VINYL FLUORIDE)

MIXTURES AND VINYLIDENE FLUORIDE-VINYL

FLUORIDE COPOLYMERS

Jerome B. Lando - Case Western Reserve University

Cleveland, Ohio 44106

Recently Natta et. al. (1) have reported that mixtures of the homopolymers poly (vinylidene fluoride) (PVF2) and poly (vinyl fluoride) (PVF)

as well as vinylidene fluoride (VF2)-vinyl fluoride (VF) copolymers form

solid solutions over the entire composition range. However, no explanation

of the differences in the melting behavior of the two types of solid

solutions as a function of composition was presented. The data were

included in Figure 3 of that paper, which is reproduced herein as Figure 1.

These data can be explained using the equation for the freezing point

depression of dilute solutions in which the solid separates as a solid

solution, namely.

where AT^ is the freezing point depression, is the molal freezing point

lowering, m is the molality of the solution, and k is the ratio of

AT^ .= K^(l - k) m

2i0-

m180-

WO-

-35-
the mole fraction of solute in the solid solution to the mole fraction of

solute in the liquid. It can be seen that the value of k will determine

the initial slopes of the curves in Figure 1 from both composition extremes.

For k- less than unity there will be a freezing point depression (negative

slope) and for k greater than unity there will be a freezing point elevation

(positive slope)

.

It should be remembered that one of the polymorphic forms of PVF2 is

isdmorphous with the crystalline pihase of PVF and this form of PVF^ is

stabilized over the entire composition range studied by the presence of

either PVF homopolymer or VF comonomer (1)

.

Considering first the VF2 side of Figure 1, it can be seen that the

initial slope of curve (a), PVF-PVF2 mixture, is positive, indicating that

when the solid and liquid phases are in equilibrium the concentration of

PVF homopolymer in the solid state is greater than in the liquid. This is

not unreasonable since it means that the higher energy PVF^ chains are melting

preferentially. However, curve (b), VF-VF^ copolymer, has an almost zero

initial slope, indicating that the compositions of the melt and solid in

equilibrium are practically identical. Since the copolymerization of VF

with VF2 is ideal (1), there will be random placement of VF in the VF^ chains.

The chemical bonding between VF and VF2 will therefore preclude differences

in composition between the solid and the melt.

On the VF side of Figure 1 curve (a) has a slightly negative initial

slope indicating that there is more PVF2 in the melt than in the solid at

equilibrium. Here again the higher energy of the PVF_ chains causes them

-36-
to melt preferentially. The more negative initial slope of curve (b),

VF-VF2 copolymer indicating more VF^ in the' melt, can be explained

by considering the units as "defects" in the predominately VF chains.

Parts of the crystals containing a higher concentration of VF^ units will

melt preferentially. The initial slope is more negative in curve (b)

because the "defects" are chemically incorporated into the polymer chains

and do not exist as separate chains.

As would be expected, the differences in the two melting point-composition curves are related to the fact that in the copolymer there are

chemical bonds between the two components of the solid solution, whereas

such is not the case in the mixtures.

References

(1) C. Natta, C. Allegra, I. W. Bassi, D. Sianesi, G. Caporiccio and

E. Torti, J. Polymer Sci. A, 3, 4263 (1965).

-37-
Slides Presented at

Piezoelectric and Pyroelectrlc Symposium-Workshop

-38-
a-b projection of the unit cell determined by Lando,

Olf, and Peterlln for phase I.

Slide 3

-39-
(a)

MONOCUNIC

Slide A: (a) a-b projection of the unit

cell determined by Doll and Lando for

phase II.

Slide 5: (b) b-c projection of the crystal

structures proposed for phase II by Doll

and Lando

.

-40-
Initial Phase Treatment Final Phase

II None II Draw 50 I Draw 50 - .\nneal 180 II

II None II

Draw 50 I Draw 50 - .Anneal 190 II

I None I

I None I Draw 127 II

Draw 50 - ;\nneal 160 I

Variation of Crystalline Pliasc of Poly (A'inylidcne Fluoride)

and to Copolymers with Treati'.ient

.

-41-
•-I 1 1 1 r—

0 5 10 15 20

•

PERCENT DEFECTS

Slide 7: Potential energy of the

all-trans and TGTG' conformations

vs HHTT concentration.

._, 1 1 1 f- 0 5 10 15 20

PERCENT DEFECTS

Slide 8: Potential energy of the

all-trans and TGTG' conformations

vs VF, concentration.

-42-
5 10 15 % VF^ ADDED

Slide 9: Potential energy of the

all-trans and TGTG' conformations

vs the concentration of VF, added

to 5% HHTT defects.

-43-
Slide 10

Suini'.iary of the Lattice Constants of

Copolymers Determined from Power Photographs

Composition Phase

PVF^ II

91-9 VF2-VF3 II

PVF2 I

93-7 VT2-TFI: i;

83-17 Vr2-VF3 I

a b c

9.63 5.02 4.62

9.59 4.98 4.66

8.47 4.90 2.56

8.85 5,00 2.55

8.84 5.03 2.54

-44-
H H F y . H

-C(l)-C(4)-C(7)-C(10)-C(13)-

H H F F. H

Slide 11: Chemical structure of

head to head-tail to tail unit.

Slide 12: Melting temperatures of

(a) PVF^-PVF mixtures and of (b)

VF^-VF copolymers.

-45-
TEXTURE AND FYROELECTRICITY IN POLYVINYLIDENE FLUORIDE*

R. G. Kepler

Sandia Laboratories, Albuquerque, New Mexico 87II5

X-ray texture studies an-d mode locked laser experiments are

reported which show that the pyroelectricity in polyvinylidene

fluoride films results from a quite large temperature dependence

of the dipole moment of the unit cell. The mode locked laser

experiments also show that the pyroelectricity is predominantly

primary pyroelectricity.

*This work supported by the U.S. Energy Research and

Development Administration.

i

-46-
In recent years there have been a number of reports of piezoelectricity in polymer films'^ and in 1971 it was reported that films of

2 polyvinylidene fluoride ' could exhibit a large pyroelectric coefficient.

A number of studies have been carried out detailing some of the properties

3 of this pyroelectricity and one attempt has been made to develop a k theoretical model. In this paper_ we report on experiments on polyvinylidene fluoride which show that the poling process, which induces the pyroelectric effect in the films, changes the texture of the films and that

the pyroelectricity results from a change in the dipole moment of the

unit cell with temperature. It is also shown that the change in the

dipole moment with temperattire is not a result of the change in the xnnlt

cell dimensions with temperature. This is the first report of a macroscopic pyroelectric effect being induced by the application of an electric

field at elevated temperatures in which the texture of the material is

changed.

Polyvinylidene fluoride (PVP^) is a crystalline polymer which exists

in two stable polymorphs. In Phase I, the crystalline form studied

in this work, the molecules are in a planar zig-zag conformation where the

unit cell is orthorhombic, space group Amm2, with lattice constants

a = S.I+T, b = h.^O, and c (chain axis) = 2. 56 A, and two polymer chains

per cell.^ In this crystal structure the dipole moments of the monomer

units are all parallel, and since the dipole moment of the monomer unit is

2. ID, a single crystal of PVF^ which the dipoles were rigidly aligned

would have a spontaneous polarization of 13.2 |jC/cm . Experimentally it

-47-
has been shovm that a polarization of as much as 3-0 \iC/cm can be

achieved in polycrystalline films which have been poled by subjecting

them to high electric fields at elevated temperature. It has also been

shown that poled films can exhibit pyroelectric coefficients as large as

3 X 10'^ c/cm^.^

In order to determine the influence of the poling process on the

polymer films we have studied x-ray pole figiores of poled and unpoled

films and have found that poling changes the texture of the films. In

a poled sample the unit cells are preferentially oriented with their

dipole moments parallel to the applied field.

Sara.ples of PVF^ were prepared for this study by stretching, at 50°C,

7 mil thick films, available commercially from Pennwalt Corporation, to

five times their original length. After stretching, the films were annealed

at 130°C for one half hour. Raman light scattering studies have been

7 carried out on these films and it has been shown that they are essentially

all Phase I and that the molecular chain axes are highly aligned in the

draw direction. The stretched films were poled by applying 600 kv/cm at

100°C for 1/2 hour and then cooling the sample to room temperature with

the field applied.

g Reflection x-ray pole figures were determined for the most intense

x-ray diffraction peak for Cu Ka radiation. This peak has been indexed as

the sum of 2(110) and (200) reflections. The figures obtained for unpoled

films (lower half of figiire) and poled films (upper half of figure) are

shown in Fig. 1. The draw direction for each pole figure is vertical.

-48-
Each line in the pole figures is an intensity contour of the reflected

X rays in a stereographic projection.

The pole figures for- the unpoled samples show that the c-axis of

the tinit cells is highly aligned in the draw direction, in agreement

7 with the Raman experiments, but that the a and b axes are randomly

oriented in a plane perpendicular to the draw direction. In the poled

films it is clear that the reflected x-ray intensity has become strongly

peaked in the direction perpendicular to the surface of the film.

Calculations show that the structure factor contribution from the (llO)

planes is k.6 times more intense than the contribution from the (200)

planes. If the unit cells were perfectly aligned with their dipole

moment parallel to the applied field, the poles of the (110) planes would

appear in the pole figure at 0.27R on each side of the center, where R is

the radius of the projection. Therefore, with imperfect alignment a

strong broad peak will appear from these planes as is observed. It is

o 6 known that heating the samples to 120 depoles the PVF^ films and pole

figures of depoled samples are the same as unpoled films.

These results show that the poling and depoling process in PVP^

involves some premelting phenomena by which crystallites melt and

recrystallize in different orientations with no apparent change in crystal

structure. The melting point for PVF^ is near 170°C. We have examined

the Raman spectra of samples which have been heated to 150°C for l/2 hour

and have been unable to detect any evidence for conversion from Phase I to

-49-
Phase II at this temperatlire. Mechanical loss measurements have been

9 carried out on our samples by K. T. Gillen and he finds a very weak and

broad loss peak, around 50°C for low frequencies. These results appear

to imply that some backbone molecular motion becomes possible at kO or

50°C which permits the melting and recrystallization process.

• In order to distinguish between- possible mechanisms for the pyroelectric effect, the time dependence of the generation of the

pyroelectric charge was measured. Films of polyvinylidene fluoride

were dyed by adding the mode locking or Q switching dye bis (U-dimethylamlnodithiobenzil) nickel [Eastman 1^015] to a solution of polyvinylidene

fluoride in dimethylformamide . Films were cast by evaporating this

solution on glass. These films were poled by applying 3OO kV/cm at

100°C for 1/2 hour and then cooling them to room temperature with the

field applied.

The poled films were placed in the sample holder shown in Fig. 2,

The capacitance of the sample was typically less than 50 pf and the

signal was fed directly into a 50 CJ cable to a Tektronix k^h oscilloscope.

Light from a pulsed mode locked neodynium glass rod laser illuminated the

sample through the glass and gold mesh electrode. Figure 3 shows a

typical train of pulses observed when the sample is irradiated with a

train of mode locked pulses. Each voltage pulse results from a mode

locked light pulse. The sign of the voltage pulse depends on the orientation of the film as is expected for the pyroelectric effect.

-50-
It has not been possible to determine accurately the absolute magnitude of the fast response pyroelectric . effect . In general, a train of

10 to 20 mode locked pulses contained a total energy of approximate.ly

1 joule. In order to make the signal as large as possible it was desirable to absorb as much light as possible but for very short mode locked

pulses J optical bleaching can occur at high intensities. In all experiments reported here, even though all the available optical energy was

not incident on the sample, optical bleaching was an important effect.

15 2

A typical pyroelectric sample contained 5 x 10 dye molecules/cm .

Therefore, if the lifetime of the excited state of the dye molecule is

longer than the mode locked pulse, a pulse containing more than a few

15 15 times 10 photons could induce optical bleaching. Five x I9 photons

corresponds to about 10 joules.

In a typical experiment, if all the energy incident on the sample

had been absorbed, the total observed charge should have been 8 x 10 ''"^C,

a value which was about 30 times larger than that observed experimentally.

On the other hand, if it is assumed that half of the dye molecules are

excited once by each mode locked pulse, the observed charge was calculated

to be about 2 x lO'"'"'^ C. On this basis, the experimental results are in

reasonable agreement with expectations.

The observation that the major contribution to the pyroelectric charge

is induced in less than k nsec shows that the pyroelectricity of FVF^ is

predominantly primary pyroelectricity, the pyroelectricity which occurs

without a change in lattice dimensions .^^ The speed of sound in FVF- is

-51-
on the order of 2 x 10 cm/sec and, therefore, it would take from

25 nsec to 50 nsec for a. ^0 \m sample to relax mechanically after an

instantaneous uniformly absorbed heat pulse. Since the observed pyroelectric response exhibited at most a small contribution with a

25 to 30 nsec time constant, the primary pyroelectric coefficient is

substantially larger than the secondary pyroelectric coefficient.

In order to look for relatively slow thermal relaxation contributions to the pyroelectricity, the charge induced by a train of mode

locked pulses of laser radiation was integrated over times of the order

of 10 |isec. Contributions on the order of ICffo of the total induced

charge were observed with a time constant of the order of 1 ^sec. This

effect may arise from heat transfer from the electrodes to the sample.

The results presented in this paper show that the pyroelectricity

exhibited by PVT^ results from a preferential orientation of the unit

cells and a strong temperature dependence of the dipole moment of the

unit cell. Since the molecule consists of rigid dipoles the temperature

dependence of dipole moment of the unit cell apparently arises from

thermal motion (libration) of the rigid dipoles.

-52-
References

1. R. Hayakawa and Y. Wada, in Advances in Polymer Science , Vol. II,

(Springer Verlag, 1973) P. 1.

2. J. G. Bergman, J. H. McFee, and G. R. Crane, Appl. Phys. Lett. l8,

203 (1971).

3. A, M. Glass, J. H. McFee, and J. G. Bergman, Jr., J. Appl. Phys.

'if2, 5219 (1971); J. H. McFee, J. G. Bergman, Jr., and G. R. Crane,

Ferroelectrics 3, 305 (1972); H. Biirkard and G. Pfister, J. Appl.

Phys. U5, 3360 (197^); P. Buchman, Ferroelectrics 5, 39 (1973).

k. E. W. Aslaksen, J. Chem. Phys. 57, 2358 (1972).

5. J. B. Lando, H. G. Olf, and A. Peterlin, J. Poly. Sci. A-1, U,

$kl (1966).

6. R. G. Kepler, Photo. Sci. Eng., in press.

7. G. Cessac and J. G. Curro, J. Poly. Sci. 12, 695 (197^).

8. J. R. Holland, Handbook of X Rays , (McGraw-Hill, Nev York, I967).

9. K. T. Gillen, private communication.

10. J. F. Nye, Physical Properties of Crystals , (Oxford University

Press, 1969), p. 189.

-53-
Figure Captions

'

Fig. 1 - Pole figures for poled (top) and unpoled (bottom) samples

of polyvinyl!dene fluoride. The lowest intensity contour

line in all pole figures corresponds to 200 cps and the

contour interval is 100 cps.

Fig. 2 - The sample holder used for the fast response pyroelectric

effect measurements.

Fig. 3 - A train of current pulses observed from a poled and dyed

sample of polyvinylidene when it is irradiated with a

train of mode locked pulses from a neodynium laser. The

vertical sensitivity is 50 mV/cm across 50 0 and the sweep

speed is 20 nsec/cm.

-54-
-55-
-56-

Addendum - After I had finished writing the above it was pointed out

by L. A. Harrah that the argument about primary versus secondary

pyroelectricity requires that the sample's mechanical properties be

homogeneous. If crystalline regions of the polymer, which are probo ably less than 1000 A in diameter, can relax at the expense of

neighboring amorphous regions then the mechanical relaxation time

could be less than 0.1 usee. Therefore, before it can be concluded

that the pyroelectric effect in polyvinylidene fluoride is predominantly primary pyroelectricity, it has to be shown that the

inhomogeneity of the crystalline polymer is not responsible for the

observed fast response.

Dr. R. Glen Kepler

Sandia Laboratories

Albuquerque, New Mexico 87115

-58-
DEPENDENCE OF THE PIEZOELECTRIC ACTIVITY OF

POLYVINYLIDENE FLUORIDE UPON HIGH SPEED UNIAXIAL

STRETCHING AND SUBSEQUENT POLING

R. J. SHUFORD, A. F. WILDE,

O.J. RICCA AND G. R. THOMiAS

ORGANIC MATERIALS LABORATORY

ARMY MATERIALS & MECHANICS RESEARCH CENTER

WATERTOWN, MASSACHUSETTS 02172

Paper Presented at

Piezoelectric and Pyroelectric Symposi ui7,-Workshop

National Bureau of Standards

Gaithersburg, Maryland

April 15 and 16, 1975

-59-
INTRODUCTION

Much interest and research in piezoelect'ricty of synthetic polymers

1 2 was generated by the initial reports of Kawai and Fukada on the high

piezoelectric effect in polyvinyl idene fluoride (PVF^). This piezoelectric

(and pyroelectric) activity occurs as a result of stretching the polymer

film, poling in a static electric field at temperatures from 80 to 150°C,

3 4 5 6

and then cooling with the field applied. ' *

On a molecular level, polarization of polymer film in a static electric

field is thought to increase the true and/or polarization charges through

several mechanisms^: 1) alignment of C-F di poles in the crystalline

7 8 9 10 11 region ' ' ' ' and 2) injection of homocharges or space charges into the

12 13 film from the electrodes. ' The actual mechanism of polarization occuring

5 in PVF2 IS probably a combination of the two , the contribution of each

depending on the poling conditions.

Polyvinyl idene fluoride is a semi crystal line polymer that exists in

at least two different crystalline conformations: planar zig zag designated

phase I (ore) and trans-gauche-trans-gauche designated phase II (or a).

Polymer films of each molecular conformation can be obtained depending upon

5 14 15 16

the method of film preparation; » ' » i.e., casting from a solvent,

melt crystallization, annealing temperature, or extrusion. This polymorphism

has been determined by infrared spectroscopy^ confirmed by laserraman studies^^'^^ X-ray diffraction and NMR^^*^^"^''. The presence of

5 to 6% head-to-head defects in the polymer during polymerization limits

the degt^ee of crystal 1 inity that can be obtained and effects the thermodynamic

stabilities of the crystalline conformations^^'^^'^^.

-60-
Changes in the molecular conformation of PVF2 can occur by mechanical

stretching or polarization processes. Orientation of the polymer chains,

along with a conversion to crystalline phase I conformation, can be achieved

by uniaxial or biaxial stretching at low- tempera tures^^'^^'^'^. Increasing

the draw temperature increases the phase II content^^. The piezoelectric

activity is enhanced by stretching or rolling the film prior to poling .

Activation of the oriented films by- poling with heat also converts phase II 19

to phase I .

Other workers have investigated the dielectric properties of PVF2 as a 30-39

function of frequency, temperature and degree of crystal 1 inity . The

high temperature a relaxation peak occuring at 60 to 80°C at 300 Hz is due

to molecular motion in the crystalline regions of the polymer. The intensity

of this absorption peak increases as the degree of crystal 1 inity or rolling

increases. The B relaxation peak which occurs between -40 and -50°C at

300 Hz is due to the micro-Browian motion of the amorphous regions. A

lower temperature y relaxation peak has been assigned to molecular motion

in the amorphous regions. Anomalous dielectric behavior of PVF^ at low

frequencies and high temperatures is thought to arise from ionic conduction

35-39

or ionic impurities present in the polymer . Assignment of the molecular

modes to the dielectric relaxation has been confirmed by similar observations

of relaxation peaks from dynamic mechanical analysis of semi -crystal 1 ine

polymers.

The purpose of our research in piezoelectric polymers is to evaluate the

potential of making piezoelectric polymers transducers for application as

on-condition vibration sensors in aircraft. The large piezoelectric effect

-61-
reported for PVF^ suggested that it be a leading candidate for the above

application. Therefore our approach has been to use polymer processing

and characterization capabilities in our laboratory to impart and to

comprehensively study systematic changes in the structure and morphology

of PVF2 and to relate these to subsequent piezoelectric behavior in an

attempt to understand, optimize, and utilize the piezoelectric phenomenon.

Commercially available extruded P,VF2 film was uniaxially stretched to various

draw ratios up to 7/1 at 80°C. The degree of orientation and crystalline

chain conformation was determined by infrared spectroscopy, density,

birefringence, sonic velocity, and X-ray diffration. The stretched

films were poled in a high D.C. electric field under varying conditions.

The static peizoel ectric constant of polarized PVF^ films was determined

as a function of draw ratio, degree of orientation, poling temperatures,

polarization voltages, and poling time. The effects of stretching and

poling on the physical, mechanical, and electrical properties of this

film were also determined.

. EXPERIMENTAL

Film Orientation

Extruded homopolymer PVF2 film (Kynar, Pennwalt Corp.) was obtained

from Westlake Plastics. This film was 5.4 mils thick, had a density of

1.768 g/cc and intrinsic viscosity of 1.37, was slightly oriented (birefringence of 0.0045), and its crystalline portion was in the phase II

conformation as determined by infrared spectroscopy.

This PVF2 film was orineted to varying degrees (draw ratios from 2/1

to 7/1) by uniaxial stretching at high speed in a radiant oven with infrared

heaters, according to an AMMRC procedure previously described^^. The film

-62-
temperature during stretching was measured by an infrared sensing temperature

indicator (IRCON Model CH34LC) which was focused on the film just outside

the exit end of the oven. Typical runs- for the PVF2 film are given in Table

I.

TABLE I

Typical Run Conditions for PVF2 Film Stretching

Draw Ratio 6:1

Feed Rate 15 ft/mi

n

Drawing Drum Rate 90 ft/mi

n

Film Exit Temperature 80°C

Film Dimensions Inital Final

Width 5 inches 2.5 inches

Thickness 5.4 mils 2.6-2.7 mils

Infrared Spectroscopy

Infrared spectra of the stretched (and poled) films were measured

in the far infrared region from 200 to 1000 cm'^ on the Beckman IR 12

spectrometer. Polarized spectra were recorded in the far infrared region

using a silver wire grid polarizer.

Density and Intrinsic Viscosity

Densities of the polymer films were determined by the gradient density

column (carbon tetrachloride-ethylene bromide) at room temperature. Intrinsic

viscosity of the as received PVF2 film was determined in dimethyl formamide

at 25°C in an Ubbelohde viscometer.

-63-
Birefringence and Sonic Modulus

Total birefringence of the uniaxially stretched polymer films was

obtained by standard i nterferometric techniques from measurements of the

film thickness and retardation. Sonic modulus of the oriented films

(parallel to the draw direction) was determined by using a Dynamic Modulus

Tester fl.M. Morgan Co.)- To improve the accuracy of the results by this

method, films ranging in length from 10 to 20 cm. were used, and the

calculated sonic moduli were taken as the average of three separate

determinations.

X-Ray Diffraction

Wide angle x-ray diffraction photographs of the PVF2 films were

obtained with Cu Ka radiation on a flat plate camera having a sample-to-film

distance of 29.86 mm. X-ray photographs were taken in three different

directions, corresponding to each of the three mutually perpendicular axes

of the film.

Dynamic Mechanical Analysis

Dynamic mechanical properties were determined with the Rheovibron

DDV-I I B Dynamic Vi scoelastometer at a constant frequency of 110 Hz. Data

were obtained in the form of dynamic modulus and loss tangent as a function

of temperature, where the loss tangent (tan 5) is a measure of viscoelastic

energy absorption per cycle in a sample under oscillating strain. Low

temperature measurements of the stretched films were difficult to make

because of brittleness of the film and slippage in the grips.

-64-
Polarization

PVFg films (about 2 inches by 5 inches) were placed between two copper

plates (embedded in two electrically insulating blocks made of micarta),

clamped together and placed in the center of a controlled temperature oven

(Custom Scientific Instruments, Inc.) preheated to a specific temperature,

usually 110°C. The poling plates were connected to a high-voltage power

supply (Sorensen & Co., Model 9060) having a maximum output of 60 KV D.C.

Values of poling voltage were chosen to give poling field strengths generally

of 500 or 1000 kV/cm. The temperature of the polymer film during poling

which lagged the oven temperature (due to the mass of the insulating blocks)

was determined by using a thermocouple embedded in one of the insulating

blocks of the poling plates. Typically the poling voltage was not applied

until after a 5-minute equilibration in the preheated oven, at which time the

poling plate temperature was about 40°C. After 30 minutes poling time in a

preheated oven at 100 to no°C, the maximum film temperature was between 80

to 85°C. The sample was cooled (by opening oven door) with the field applied

for at least 15 minutes, at which time the final film temperature was about

65 to yo^'c.

Piezoelectric Activity

The first observations of piezoelectric activity were made by subjecting

the stretched poled PVF2 films to slow cyclic tensile loading in an Instron.

Two test modes were employed: sawtooth and sinusoidal. The resulting

piezoelectric signals were detected by an oscilloscope. These results served

only to demonstrate in a qualitative fashion the piezoelectric behavior of

PVF2.

-65-
Quantitative values of the static piezoelectric constant were obtained

by placing known compressive loads upon the stretch poled PVF2 film with an

Instron and measuring the resultant stress-induced output voltage from a

coulometer. The coulometer was constructed at AMMRC from a diagram furnished

by E. L. Church and H. Jenkinson of the Frankford Arsenal. The coulometer

consisted of an operational amplifier (Analog Devices, Inc., Model 310K)

with a feedback capacitor connected between the output and input circuits.

The coulometer output voltage was read with a Simpson voltmeter, and then

converted to values of piezoelectric charge. For loading, the polymer film

was placed between two metal plates serving as electrodes, thus avoiding

the need for electroding the film prior to measurement. From the incremental

values of static compressive load on the polymer film and the corresponding

coulometer output voltage, one could calculate the piezoelectric constant

(d^^) by the equation : 5

d33 = "^^^g^/film area ' •^^^/film area

in units of coulombs/newton. The piezoelectric constant was determined

at least one day after poling to allow dissipation of residual surface

charges on the film.

RESULTS AND DISCUSSION

It was found that systematic increases on the degree of PVF2 film

stretching produced systematic changes in the properties of the film. These

property changes were characterized by a number of methods as described in

the Experimental section.

-66-
Infrared Spectroscopy

IR spectroscopy in the far infrared region showed that the crystalline

component of the PVF2 film (as received) was predominantly in the Phase II 15-23

form . See Figure 1. This is evident' by the presence of absorption

bands at 615, 530, 410, 360, and 290 cm"''. As the degree of uniaxial

stretching was increased in steps up to 5/1, the IR spectra showed a

continuous decrease in the 530 cm""' and 410 cm""* bands and a concomitant

increase with stretching in the 445 cm"'' band. These results indicate a

gradual conversion of Phase II to Phase I due to the stretching process.

At further increases in the draw ratio up to 7/1, no further changes were

noted in the IR spectra. BDth the conventional IR spectra and the polarized

IR spectra exhibited no significant dependence upon poling of the PVF2 films

for films which had previously been highly stretched. This agrees with

22

Cessac and Curro , where according to their results and calculations

dipole orientations greater than 60 or 80% are not occuring during poling,

but where dipole orientations of less than 60% can not be ruled out.

Density

As the degree of film stretching was increased, the film density

also rose, but in an S-shaped fashion as shown in the curve of Figure 2.

The increase in density is attributed to the increased content of the

23-27

higher density phase I form and may also result from an increase in

the degree of crystal 1 inity.

Bi refri ngence

The total birefringence measured by these experiments is the sum of the

contributions due to the crystalline components, the amorphous components,

and the form birefringence. The latter contribution is usually assumed to

be small. The increase in birefringence upon stretching, noted in Figure 3,

-67-
is attributed primarily to the changes in degree of preferred orientation

in the crystalline phase. The leveling off seen at the higher draw ratio

suggests a saturation in the degree of crystalline orientation. For highly

stretched films no difference in birefringence was observed between poled

and unpoled specimens. This agrees with, the results of McFee, terrgman, and

Crane^.

Sonic Modulus

, The effect of stretching upon ,the sonic modulus also appears in Figure

3, where there is seen a continuous increase in modulus with draw ratio. In

a semi -crystal 1 i ne polymer the sonic modulus is the sum of the contributions

from the crystalline and amorphous regions. Previous work on oriented semicrystalline polymers has indicated that the sonic modulus is more sensitive

41

to orientation changes in the amorphous than in the crystalline regions

Vfe might therefore expect from Figure 3 that the orientation in the

amorphous region is continuing to increase at the higher draw ratios.

X-Ray Diffraction

The X-ray diffraction patterns seen in Figures 4-7 show that at low

degrees of stretching the individual crystallites are randomly oriented.

As the degree of stretching is increased, the crystallites assume an increasingly preferred orientation as indicated by the gradual breakup of

the diffraction rings into spots. Also the changes noticed in the diffraction

spacings (Table II) indicate a continuous conversion of Phase II to Phase I

with increasing degree of stretching. Wth the highly stretched PVF^ films

there was no difference produced in the X-ray diffraction patterns by poling.

22 This agrees with the results cited by Cessac and Curro .

-68-
Dynamic Mechanical Analyses

Dynamic mechanical measurements, including loss (tan 6) and

modulus, were determined for the unstretched and stretched PVF^

films as a function of temperature at a constant frequency of 110 Hz

(See Figures 8, 9, 10). A B transition peak occurs between -40 and

-20''C in both the unstretched and stretched PVF^ films. The ratio

of the height of the 3 peak compared to the a peak (50 to 70°C)

decreases in the stretched films'. With the stretched films there

were no significant differences in dynamic mechanical spectra

between upoled and poled specimens.

Piezoelectric Activity

The first observations of piezoelectric activity viere achieved

by the slow cyclic tensile loading of the stretched poled PVF^ films

in an Instron test machine. The piezoelectric signals were observed

by direct coupling to an oscilloscope. It was -found that the piezoelectric voltages measured in this way were proportional to the film

loading rates in the sawtooth test (see Figure 11), and were very

nearly proportional to the film cyclic test frequnecy in the

sinusoidal test (10 to 30 Hz). At the same time, these piezoelectric

voltages were lower than those to be expected from a simple calculation

based on a consideration of the PVF2 film as a paral 1 el -pi ate

capacitor. The observed rate (or frequency) dependence was similar to

42 43

that reported by Broadhurst et. al . ' who had measured the short-circuit

piezoelectric current for polyvinyl chloride films. The rate (or frequency)

.-£9-

/O
dependence of the AMMRC tests was explained by considering the low input

impedance of the oscilloscope as a charge leakage path which lowered the

instantaneous charge level on the PVF^ film. A mathematical treatment of

the circuit response to the charge generation and charge leakage processes

showed that the piezoelectric voltage should be nearly proportional to the

time derivative of the loading function at these low test rates. These

initial experiments served only as a qualitative demonstration of the

piezoelectric activity of PVF2.

The quantitative studies of the PVF2 films involved the determination

of the static piezoelectric constant as described in the Experimental section.

A typical result from this type of experiment is shown in Figure 12, where

the coulometer voltage outputs are plotted against the compressive loads

on the film specimens. The curves are highly linear during the loading

increments but become slightly non-linear during the unloading increments^

These results are quite reproducible after several loading and unloading

cycles. The first loading cycle generally produced larger voltages for

corresponding loading increments and, upon unloading, returned instead to

a voltage greater than zero at zero load. The polarity of the piezoelectric

charge was always the same, viz., the face of the PVF2 film which became

positive during the compression loading was the face of the film which was

adjacent to the negative poling plate during the poling process.

A few compression tests were performed in which two stretched poled

PVF2 films were stacked one above the other, i.e., mechanically in series.

Wien the films of equal piezoelectric activity were connected electrically

in parallel, the piezoelectric activity of the combination was double that

-70-
of either film. This agreed with the results of Church who has investigated

the piezoelectric activity of multi-layered PVF2 film assemblies. Wien we

connected our two PVF^ films electrically in series, we found that the

piezoelectric activity of the combination was equal to the average of that

of the two individual films. This latter result held for three different

cases: the two films were of equal activity, the two films were of unequal

activity, and one of the films had zero activity.

With PVF2 films poled under an effective field of 500 kV/cm, the piezoelectric constant was found to vary with the prior draw ratio in the manner

shown by Figure 13. The most rapid increase occurred between draw ratios

of 3/1 and 5/1, paralleling the conversion of Phase II crystalline conformation

to Phase I as a function of draw ratio as noted by IR spectroscopy. This

suggested that the Phase I crystalline material plays a considerable role

1 1 45

in the piezoelectric behavior ' in addition to any effect produced

by polymer orientation. The apparent leveling off of the piezoelectric

constant at the high draw ratios agrees with our results that there is

no further conversion of material to the Phase I form and there is no

further significant orientation of the crystalline regions at high draw

ratios under these experimental conditions. The further increase of

sonic modulus at the higher draw ratios indicates that the continuing

orientation of the amorphous material may contribute in only a minor way

to the piezoelectric constant.

Changes in density were produced by drawing, as previously discussed

for Figure 2. The piezoelectric constant, when plotted against density,

appears as in Figure 14. This monotonic relationship indicates that

the piezoelectric constant may be largely dependent upon the degree of

crystal! inity and/or the change in crystalline conformation with

-71-
stretching. The apparent leveling off of the piezoelectric activity

at high draw ratios (see Figure 13) and at high densities (suggested

by the cluster of points at high draw, ratios on Figure 14) may

indicate that the piezoelectric constant is limited by the degree of

n ?R ?Q crystallinity or the phase I content

A series of experiments was run to determine the effects of the

poling variables upon the subsequent piezoelectric activity of the

highly stretched PVF2 films. The effect of the poling voltage upon

the piezoelectric constant is shown in Figure 15, where one sees a

smoothly rising curve which begins to level off at the higher voltages.

This suggests that for these poling times and temperatures we are

approaching a limit to the piezoelectric activity that can be attained

by this method.

• The effect of poling times at final poling temperatures of 80°C

and above and fields of 500 kV/cm is seen in Figure 16. The times

plotted here refer to the length of time required for the film to

be heated through the specified temperature range. There was an

additional cooling time of at least 15 minutes to allow the film to

reach approximately 70°C. This graph shows that heating times of

5-10 minutes are sufficient to produce the maximum attainable piezoelectric activity under these poling conditions. Other results, not

shown here, indiate that longer poling times (60 to 90 minutes) have

little effect upon the piezoelectric activity.

The effect of poling temperature upon piezoelectric activity

at fixed poling fields (500 KV/cm) and poling times (total, 45 minutes)

appears in Figure 17. In these experiments the films were brought

-72-
to the specified temperature and annealed for 15 to 30 minutes prior

to applying the D.C. voltage. There is a marked effect of temperature,

where the piezoelectric constant ranges from 0.5 pcoul/nt at room

temperature to values of 6 to 8 pcouT/nf at 60 to 80°C where there is

an apparent leveling off of activity. • The dashed line in Figure 17

indicates the values to be expected from consideration of results in

Figures 15 and 16. Thus annealing of the film prior to poling seems

to have a deleterious effect on the piezoelectric activity.

The retention of the piezoelectric activity over a temperature

range for highly stretched PVF2 films poled at 85°C at 1000 KV/cm is

plotted in Figure 18. Initial film activity was about 19 pcoul/nt.

With no field applied, the films were then reheated between aluminum

plates for one hour at the specified temperature and then cooled to room

temperature before the piezoelectric activity was again determined. It

is seen that the piezoelectric activity was retained under these conditions

up to 85-90°C, beyond which there was a continuous decrease with temperature.

At 150°C about one-half of the original activity was still retained.

CONCLUSIONS

From the experiments and results described above, we can draw the

following conclusions:

1. Increases in the draw ratio of the PVF2 film produces continuous

conversions from crystalline Phase II to Phase I. At the same time the

increase in density may indicate a change in crystalline conformation

and/or simultaneous conversion of amorphorous polymer regions to a

-73-
crystalline form. The leveling off of the birefrigence at the high

draw ratios suggests that the crystalline phase is undergoing no

further orientation, whereas the continuing increase of sonic modulus

at high draw ratios indicates that the amorphous regions are still

undergoing further orientation.

.2. The piezoelectric activity increased with draw ratio, the

poling voltage, poling temperature, and poling time.

, 3. Significant amounts of oriented phase I crystalline material

are needed in order to achieve high degrees of piezoelectric activity

in PVF,,.

-74-
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^1 . Desper, C. R. "Technique for Measuring Orientation in Polymers".

CRC Critical Reviews in Macromolecular Science, v. 1, 1973, 503.

^2. Broadhurst, M. G. , Malmberg, C. G., Mopsik, F. I., and Harris, W. P.

"Piezo- and Pyroelectrici ty in Polymer Electrets". Presented at

Conference on Electrets, Charge Storage and Transport In Dielectrics,

Miami Beach, Fla., Oct. 1972.

^3- Broadhurst, M. G., Harris, W. P., Mopsik, F. I., and Malmberg, C. G.

"Piezoelectricity, Pyroelectrici ty and Electrostrict ion in Polymers".

Polymer Preprints, v. 1^*, 1973, P- 820.

kk. Church, E. L. Frankford Arsenal, Private communication, 1975.

^5. Tamura, M., Ogasawara, K. , Ono, N. , and Hagiwara, S. "Piezoelectricity

in Uniaxially Stretched Pol yv i nyl i dene Fluoride". J. Appl. Phys.,

V. ^5, 197^, p. 3768.

-79-
CVJ

LU

_)

QQ5C

U- Q_ O O o <

Q_

S5

UJ> • — X

f—I ' — O < <

ZDo ^ oI o

<: o

00 -

Ct:: ^ ro

00

<:

LU

Q_

u_O

J*. >-

htn

'C\J

CSJ oO

<:

on

eno

00 H00 o

00

CO +

00

onC

CDd

O)o

c:

ro o o

00 1-

-80-
Figure 1

-81-
Density (g/cc)

• • . • • • • *nJ -»j --sJ • —J O OO CX5 sO

VJl CD VJ1 O VJI CD

Figure 2

-82
Figure 3

-83-
mmmmmmmm

O I

UNORIENTED DRAW RATIO: 2/1

iiiKttiiiiiim iiririiniT^"-'''-' -•'-"^-'^'"iTiirfn-fi--" -'"^^

pi iwjwwi i. iip;.!,ij«|.. ,l 'AI'

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^F/6. 6

kMiitkitaliriitilllMi DRAW RATIO: 4/1

1

F/G. 7 DRAW RATIO: 6/1 WIDE ANGLE X-RAY DIFFRACTION OF STRETCHED PVF2 FILM

-84-
10.0^'

PVF2 ORIGINAL

lOll

1.0' lOlO

c

TO 0.1

10'

CM

c:>^

0.01 10^

0.001

-180 -120 0 120

Temperature ('O

10'

240

Figure 8
10.0^ PVFo DRAW RATIO 4/1

•

10

11

l.O* 10

10

CO

c

03 o.r

E

o

io9 I•o

o.or 10

8

0.001

-180 -120 0 120

Temperature {°C)

10'

240

Figure 9

-86-
10.0

1.0

0.1

0.01

0.001

PVF2 DRAW RATIO 6/1 tlO^l

4 io9

108

10^

-180 -120 0 120 240

. Temperature {°C)

Figure 10
Voltage O r\j ro v*J O —I ^ «—• oo \_n fX)

O

Figure 11

-88-
Voltage (dc) TO <JJ *k >J1

Figure 12

-89-
PZ Const (pcoul/nt)

Figure 13

-90-
-91-
PZ Const (pcoul/nt)

< <

3

Figure 15

-92-
PZ Const (pcoul/nt;

5' O

CO

3

CO

zn VjO

QJ

CD O

<—»- 13

O

o

Figure 15

-93-
PZ CONST ( pcoul/ nt )

Figure 17

-94-
?Z CO\STA'J ( pcoul/nt

)

Vjn 5 \ji

o

Figure 18

-95-
Thermally Stimulated Currents and Morphology in PVF

M. Abkowitz, P.J. Luca, G. Pfister and W.M. Prest, Jr.

Xerox Corporation, Webster, N.Y. 14580

(1) Introduction

This report will address three topics of interest to the study and understanding of tlie origin of pyro- and piezoelectric befiavior in polarized polymer

films, and in particular in PVFg. The first part v/ill show that thermally stimulated currents (TSC) are a viable technique to probe the morphological implications of the origin of the poling induced polarization. The a^-relaxation in

commercial films of PVF2 will serve as an example. The. TSC results will then be

related to ac current measurements which, in analogy to the TSC technique, are

run at constant frequency on a temperature axis. The relation between the heating

rate in the TSC experiment and the frequency in the ac experiment will be dem.onstrated. By then combining the TSC and ac techniques, we are able to track the

a^- and a^-relaxations in commercial PVF« films over 9 decades of frequency,

a c c Quite generally, with the method illustrated for the a^- and a^-relaxations in

PVFg, structure in the TSC spectrum can be related to structure observed in dielectric and mechanical relaxation studies. This way that structure in the TSC, which

is unique to molecular relaxations can be identified, thus overcoming what is often perceived to be a major limitation of an otherwise uniquely powerful technique

for studying low frequency nxDlecular motions in polymeric solids.

The second part of the presentation will discuss TSC measurements on samples

prepared from Pennwalt and Kureha powder materials. It will be shown that the

morphology of those samples, which was controlled by the crystallization tempera-

-96-
ture/time history strongly influences tiie TSC spectrum. In particular, the appearance of a third phcse, probably a modified form of the s-crystalline phase, drastically enhances the TSC amplitude.

Finally, the third part presents some fi el d/ter:perature history effects upon

the electrical properties of corr:5rcial PVF2 filr;£. Of particular interest are

those effects v/hicn v;ere induced at fields and temperatures typically used for

poling PVF2. Furthermiore , evidence for v/hat nay be field induced dipole pinning

in the amorphous phase of the polymer will be presented. The first end third part

of the presentation have been the subject of recent oublications , while the results

of the second part are in the stage of detailed analysis and sc~e data have to be

treated as preliminary.

(2) The g^ - an d a_-Rel axati ens

a C

In the TSC experiment, a sample which has previously been pcla^'ized by cooling

in a dc bias field, is heated at a constant rate under zero b^as. Tne electrodes

of the sample "are connected to a current mater and the current flov/ing in t-e external circuit as the polarized sample is heated, is continuously recorded versus

samiple tem.perature . The TSC often exhibits a structure with peaks occurring at

temperatures characteristic of m.olecular relaxations in the polymer. Peaks might

arise from reorientation of permanent or induced diocles or charge release from

traps associated with the relaxation process. In addition, of course, trap emptying unrelated to molecular relaxations might also occur.

In PVF2, the TSC exhibits a large peak at Tg '--Sl^'C, which is associated with

dipole relaxation, probably of the permanent (CH2CF2) dipoles (ci^-'r-el axati on) . The

evidence for dipole relaxation rather than charge release is illustrated in Fig. 1.

For our purpose, only the traces below room tem.perature are of interest. Trace a

shows the TSC following cooling of a 25um Kureha PVF^ film in 3V/ym. Trace b shows

the current measured if an unpoled film, i.e., a film which has been cooled to below Tg under zero bias, is heated with a 3V/-^m bias field applied. The traces a

-97-
and b are equal but ot opposite sign. These two experiments already suggest that

the TSC peak at -v-SCC is due to a dipole polarization. The observation of a peak

in trace b cannot be understood by trapped charge release since the film was cooled

under zero bias. In the experiment giving rise to trace b in Fig. 1, the orientation of the dipoles in tne polymers amorphous phase initially is random. On

heating the sample with the field applied, dipoles respond to the field as they

become unfrozen around Tg. This dipole alignment is observed by a corresponding

peak at Tg in the current flowing in the external circuit. Conversely, in the TSC

experiment (trace a) the dipoles initially are aligned and randomize at Tg when the

sample is heated under zero bias field. Additional experiments can be performed

to test this interpretation. For instance, heating an unpoled film sandwiched between blocking electrodes under appropriate bias reproduces trace b below 'vO°C

(see rig. 1, trace c). The position T^^ of the TSC peak can be analyzed as a function of the heating rate, applied field and poling temperature. On finds 'v^yiKcal/

mole as apparent activation energy for the relaxation process at Tg, T(Tg) '^500 sec

and T^(B^) = Tg 'v-STC. A value Ny^ ,>10^^ C^/cm is obtained for the average num,-

ber N of chain segments with average net dipole moment p responding to the dc field

at Tg.

Conventional ac dielectric experiments are performed under isothermal conditions and tneir relation to the nonisothermal TSC measurement is not straingt fomard.

The connection between TSC and ac experiments, however, is greatly facilitated if

the latter is performed at constant frequency as a continuous function of temperature. The problem of connecting the two techniques then reduces to relating the

heating rate 6 in the TSC experiment to the frequency w in the ac experiment. This

connection will be explored in the following paragraph.

Operationally, one can relate 6 and w via the well-known equation that defines

the temperature T^^ of the TSC maximum. In the usual framework for a exponentially

decaying polarization the condition for the TSC maximum is

^=J/6 (1)

-98-
where t is the relaxation time. Inserting for example, a exponential temperature

dependence for t Eq. (1; yields

1 . 6A

(2)

t2m

where A is the activation energy in °K. It is «;uggested then that a frequency

(jj = ^ can operationally be defined from Eq. (2). For typical values of 6 ^ 0.3 -

20°C/min and A 20-100 Kcal/mole, Eq. (2) would predict (^/2tt) ^ 10"^ - 10"''hz.

ac experiments are typically run in the 10 - 10 Hz frequency range, hence the

heating rate 6 and the lowest ac frequency to. Fig. 2 (lower part) shows the TSC

run at 9*'C/min. The upper part of Fig. 2 shows the real part of the ac current run

at 5 X 10~''hz. The similarity of the current peak around Tg is clearly indicated.

higher temperatures and becomes broader. This is shown in Fig. 3, upoor part,

-1 5

where the ac peak associated with Tg is shown for 5 x 10 and 2 x 10 Hz. The ac

current at Tg is proportional to the frequency of the ac driving field and thus

the units in Fig. 3 are relative. The lower part of Fig. 3 shows the linewidth

a(T^) of the ac peak as a function of the temperature T^^ at which the current

peaks. The linewidth a(T^) of the ac peak temperature T^ of the ac and TSC

a. - current peaks approach each other as the frequency of the ac experiment

a approaches the frequency which can operationally be defined from Eq. (2) for

the TSC experiment. Eq. (2) can be used to calculate Vt at the temperature T^^

from quantities determ.ined from tlie TSC experiment. For the glass transition,

the liLF relation is more approapriate to describe the temperature dependence of

T than a thermally, activated process, viz.

connection between TSC and ac experiment should be observable at the highest TSC

As the frequency of the ac experiment is increased, the a

~ peak rapidly shifts to

a

g ^^P^-C- (3)

-99-
in v^hich case Eq. (2) reduces to

^ 2 (4)

' (C2+TnrTg)2

In Eq. (3) C-j and C2 are constant v/hich is principle, can be obtained from a least

square fit. For our purpose, however, it is sufficient to use the "universal"

values C-j = 40, C2 = 51.6°K, which were found for a large number of materials. We

calculated Vt from Eq. (4) for various heating rates 6 and plotted log Vx versus Vt = Vt^ in Fig. (4) (a^-curve, open circles). For the analysis of the ac

peak we assumed a phenomenological distribution of relaxation times and so used

the Davidson-Cole relation

'

. WT = tan (J) (5)

where ({> = |- y is the associated distribution parameter. Although this relation nas been derived to analyze the peak of c'' versus w at constant temperature,

it can be used to derive Vx ^^'^^ the peak position of the ac current on the temperature axis if ^^p- << ^^7^' 3 condition which is often satisfied. Using Eq. (5)

we calculated Vt from the ac current peak measured at frequency The log Vx

values are plotted in Fig. 4 (a^-curve, full circles) as a function of Vj = "''^Tj^.

We note that the data derived from TSC and ac experiment join smoothly, demonstrating

self consistency of the analysis. Using both techniques, the relaxation behavior

of the -relaxation thus can be mapped out over ^ nine decades of frequency. The

a data clearly demonstrate the expected WLF temperature dependence (solid line calculated with the constants = 40, C2 = 51.6°K).

A similar analysis has been carried out for the a^-crystall ine relaxation in

commercial films of P\/F2. To study this relaxation the concentration of the a-phase

was enhanced by thermal treatment (see Sec. 3). Fig. 5 shows the temperature shift

of the a -relaxation with heating rote (TSC, upper part) and frequency (ac, lower

part). The corresponding logVx versus Vj = Vthi plot (Fig. 4, a^-curve) shows

again, a smooth connection of ac and TSC data and furthermore verifies the expected

-100-
thermal activation of the crystalline relaxation process (activation energy -vZlKcal/

mole).

(3) TSC on Samples v/ith Controlled Morphology

To gain further understanding of the origin of the significant pyro- and iezoelectric behavior it is necessary to separate- the effects of the different crystal

lographic phases and morphologies upon the poling induced polarization. Tn the

following, we examine these effects upon the TSC in unoriented samples with controlled

crystal phases and microstructure. The films are prepared from powder obtained

from Pennwalt Co. (powder number 881, 821 and 401) and from Kureha Chemical Co.

(KF-llOO). The powder was melted and then quenched to the crystallization temperature T . By varying T and the crystallization time t , various morphologies were

produced as verified by temperature dependent IR and DSC. IR absorption spectra

are used to distinguish three crystalline phases, the tg-tg' phase a, the planar

zig-zag 3 and a modification of the 6 phase called the y-phase. The a-phase, identified by IR bands at 530 , 615, 764 and 795cm" forms from the melt and has the

lowest melting temperature. The higher melting B-phase is also crystallized from

the melt and is identified by IR bands at 510 and 840cm"''. The highest melting yphase is formed by a crystal -crystal transformation from the a-phase. The y-phase

is identified by IR bands at 430 , 776 and 810cm which appear in addition to the

bands characteristic of the p-phase. Fig. 6 shows a Hof fman-l'eeks plot of the observed melting temperature Tj,^ versus crystallization temperature T^ for samples of

821 powder crystallized from the melt. Also indicated is the halftime of a-crystallization. It is noted that with increasing T the halftime decreases with progressive

rapidity. At higher crystallization temperatures B-crystal 1 i zati on occurs as a competitive process to a-crystallization. At T^ ^ 430°K formation of the y-phase is

observed in addition to the a- and B-phase. The concentrations of the three phases

for this particular 821 powder are given in the Table for different crystallization

-101-
temperatures. The relative concentrations and in particular that of the y-phase

are strongly dependent upon the crystallization time t . The films prepared this

way were typically 50ym thick. They were electroded with Au contacts and then

cooled in 'v3V/um to below Tg. The TSC was then run at 7.7°C/min. Fig. 7 shows TSC

traces for samples crystallized from 821 pov/.der. at the indicated temperatures T^.

It is evident that with the appearance y (T^ = 432. 5°K) the TSC increases significantly. For comparison we included a trace recorded on a biaxially oriented commercial film of Kureha PVF2, which contains the a- and 6-phases. We also note that

with increasing T the amplitude of the, a -peak decreases thus indicating an overall

increase in crystal linity of the sample. The TSC exhibits a well defined structure,

viz. at 'X'-SO^C the a^-peak, at 'v25°C the shoulder associated with the a^-relaxation

a c and at '\^15°C, shoulders which appear in B and y containing films. At still higher

temperatures (75-100°C) a peak and a shoulder is observed for the a and B rich samples, respectively. This latter peak would also be indicated for the y rich films

but would be off scale in Fig. 7. Assignments of the peaks which appear in addition to the well documented a^- and a^-relaxations is very preliminary. It is

a c suggested that the relaxation at %15°C is associated with the y-phase. That it

also appears in the T = 430°K sample is believed to indicate that this sample contains small traces of y, which are not observed with our IR spectrometer. The

higher temperature peak at 50-75°C is tentatively associated with the 6-phase on

the basis of the data on the commercial Kureha film. At this point, however, we

have not established the uniqueness of the relaxation modes related to the 6 and y

phases. They might well result from the same underlying mode occurring in differing

morphological enviroments. Fig. 8 shows similar TSC traces on samples prepared

from KF-1100 powder. The data in Fig. 7 and 8 thus unarit)i guously demonstrate that

the poling induced polarization is strongly dependent upon sample morphology. The

net effect upon the films pyroelectric and piezoelectric response when poled under

typical conditions ('vlOO°C, /?;300KV/cm) are currently being investigated.

-102- I
(4) Field/Temperature Effects

The temperature range above 100°C is of major technical interest because it

is here that poling is known to induce significant reversible pyro-and piezoelectric behavior. In the following, we describe results of ac and TSC measurements

performed in conjunction with applied dc bias amJ sequential in situ thermal cycling.

Our principal observations on commercial PVF2 .films are the following:

(i) Momentary application of electric fields at elevated temperatures significantly reduces the current measured at high temperatures. The field induced

current suppression by far exceeds the reduction achieved by merely thermal lizing

in zero electric field at the same temperature. The effect cannot be reversed by

reversing field polarity.

(ii) No enhancement of the absorption attributable to field exposure could

be observed.

(iii) Current levels suppressed by electric field exposure could in general,

be partially restored by heating under zero bias at temperatures about 2C°C below

the melting temperature. If the applied field was below the threshold field to induce reversible pyroelectricity, the field effect could be completely reversed in

a few thermal cycles but if it exceeded the threshold value, only marginal reversal

was attained after hours of persistent high temperature thermalizing.

(iv) Partial pinning of dipoles in the amorphous phase may occur as a consequence of field exposure.

In Fig. 9 some effects are illustrated for a lOHz ac trace recorded on a 2 m,i 1

Pennwalt film. The dashed line a is the first heating-cooling cycle of the asreceived film. At the amplifier gain setting chosen to record the traces shown in

Fig. 9, the Tg current peak was shut off and the current stabilized at 108°C to the

value indicated as a black dot on the figure without any indication of further drift.

A field of 16V/iim, which is well below the threshold field for reversible pyroelectricity was then switched on. The illustrated decrease in current occurred over a

span of 'v600ins. The sample was now cycled down in temperature under applied field

-103-
(horizontal arrow in Figure) to well below 'Tg, where the field vyas switched off.

Under zero bias, the sample was now recycled to 140°C and the broad current loop

b shown in the Figure was correspondingly recorded. The reduction in current level

at high temperature in the heating half of the cycle is a consequence of orevious

field exposure. The loop width observed on cooling is a consequence of heating

the' sample under zero field, a result of high temperature exposure in the thermal

cycle used to record the heating part of trace b. Vlith progressive cycling to 140°C

the effect of the field exposure could be substantially erased. Thermally reversing

the effeqt of field exposure became less effective as the field was increased to

poling levels, i.e., 30-40V/ym. This is illustrated in the lower part of the figure.

The film was exposed to a field of 37V/vm at 110°C for 3 hrs. and then cooled under

field. Curve c shows the initial current measured under zero dc bias on first heating

at an amplifier gain setting identical to that used in the upper figure. Curve d

is recorded after heating at HS^C under zero field for 17 hrs. Only partial restoration to that high temperature current level is indicated. The conductivity after

this treatment remains substantially below that measured on both an unpoled unannealed

film and a film given comparable thermal history but with no field applied.

Corraborati ve information from TSC measurements is shown in Fig. 10. An asreceived 25um Kureha film is cooled under zero bias to low temperature, where a

30V/ym film is switched on. The film is then thermally cycled under field giving

rise to loop a. The suppression of current level at high temperature observed in

comparing the heating and cooling currents at a given temperature in this range is

between two and three orders of magnitude as indicated on the logarithmic ordinate,

scale. Cooling the film in the presence of the strong dc bias field can freeze-in

a polarization and, as a result, the cooling current is suppressed relative to the

heating current. An estimate of the frozen-in polarization is obtained from the

depolarization current trace b shown in the lower part of Fig. 10. This trace was

recorded on heating the film under zero bias following completion of loop a and for

the purpose of illustration the negative of the recorded current was plotted. Up

-104-
to 'vO°C both heating trace a and trace b are- identical indicating that below 0°C

the difference in heating and cooling currents in loop a is due to a frozen-in

polarization. (See Sec. 2) At higher temperatures, however, trace b is orders of

magnitudes smaller than heating trace a. This clearly demonstrates that the width

of current loop a is not the result of a frozen-in polarization. If the film is

now cycled through four complete (a,b) sequences and then run under 30V/vim field,

again starting at low temperature, loop c is obtained. Loop c is similar to loop

a below 0°C but the effect of prolonged field exposure and subsequent thermal history is demonstrated when the high temperature loops c and a are compared. The

depolarization current following loop c is identical to that following loop a

(trace b., lower part, Fig. 10). While the evolution of the initial current trace

a into trace c revealed sensitivity to combined thermal and field exposure history,

the depolarization current density remained well below the total thermocurrent density measured under bias above 100°C. Furthermore, repeated thermal cycling at

30V/ym field tended to produce som.e additional current suppression at high temperatures and some additional loop narrowing relative to trace c. Heating currents,

therefore, always remained somewhat larger than cooling currents measured at the

same temperature.

Experiments presently unden-;ay, suggest lowering of high temperature conductivity by application of an electric field may involve more than one physical

mechanism. Whereas poling through blocking contacts permanently reduces ac conductivity, which suggest a field driven perturbation of film mi crostructure , recent

experiments appear to indicate that phenomena reminiscent of thermal switching are

also observed. These take the form of transient current instabilities, which appear when a film is cycled from low temperature under high dc bias into the poling

temperature range.

Pulsed field experiments indicate that reversible pyrcelectric behavior can

be induced on a .1-1 sec. time scale for fields and temperatures exceeding specified threshold values. The lowering of ac conductivity by application of a dc

-105-
field at poling temperatures (Fig. 9) evidently occurs on the same time scale.

The ac current peak in the temperature range of the glass transition was found

to be sensitive to applied dc bias. The effects of applied bias were thermally

reversible and not electrostri cti ve in origin. Typical results are illustrated in

Fig. 11 for a 12ym Kureha film measured at lOOHz. Trace a represents the ac current peak observed on a film cooled from 30°C to -180°C under zero dc bias. This

annealing procedure always reproduced peak a irrespective of any earlier treatment.

When a field of 40V/pm was applied at -180°C and the film then heated at 4°C/min.

under field to +50°C trace b was obtained. If the field was left on at +50°C and

the film cooled down to -180°C trace c resulted. The suppression of the peak amplitude saturated under additional cycling between these terminal temperatures at

the samie field intensity. The recovery of the current amplitude from that shown

on trace c to that shown on trace a after dc bias was removed, was very slow below

the glass transition temperature but occurred with increasing rapidity at temperature above Tg. The effect of applied dc bias field in depressing the ac current

peak amplitude was dependent on field intensity. It seems reasonable to tentatively

identify this effect with a field induced pinning of dipoles (realignment) in the

polymers amorphous phase. This realignment evidently occurs only in the neighborhood of the glass transition temperature and above. In the presence of a dc field,

this partial alignment remains stabilized at temperatures well in excess of Tg.

-106-
Trace a: Depolarization current (TSC), i.e., heating in zero field

of sample previously cooled under 3V/ym.

Trace b: Polarization current, i.e., heatina of unpoled sample in

3V/ym.

Trace c: Polarization current measured on sample sandv/i chod betv.'cen

Mylar blocking layers under appropriate bias. Heating rate: 9°C/niin.,

sample: Zo\ir^ Kureha PVF2.

ac current J^oC") '^'^'^ at 5 x 1.0" 'Hz (upper part) and TSC Jg(T) run at 9°C/min. to show convergence of TSC and ac lineshape and position

of the a^-relaxation.

a Upper part: Illustration of temperature shift and broadening of ac current a^-peak with increasing frequency.

Lower part: Temperature linewidth plotted versus teniperature T^ at

which ac current peaks. Convergence to the corresponding linewidth

of the TSC ag-peak is clearly indicated.

Combined plot of relaxation rate log 1/t versus temperature Tp^ calculated from peak position of ac (full points) and TSC (open points)

current peaks for a^- and ac-relaxati ens

.

Temperature shift of a^-rel axati on peek with heating rate (TSC experiment) and frequency (ac experim.ent) . Curves are shifted along ordinate for clarity, a^-peak is not shown for the ac current.

Hoffman-Weeks plot (melting temperature 1^ versus crystallization

temperature T^) fcr samples prepared from 821 Pennwalt powder. Also

shown is the helftime of a-crystallization. The a, 6 and y-phases

were identified by their IP. bands at 530 , 515 , 764 and 795cn-"' ; 510

and 840cm-T; and 430, 510, 776, 810 and 840cm-^ , respectively.

TSC of samples, crystal 1 ized at different temperature Ty.. Powder 821

Pennwalt. Also shown is the TSC of a com^mercial 25um Kureha PVF9.

Poling field 3V/um heating rate 7.7°C/min. The ordinate units are J/AsEp, where J, A, e and En are. the current (Amps), electrode area

(cm^), heating rate (°K/sec; and poling field (V/cm). The arrows indicate tentative assignment of TSC structure.

Same as Fig. 7, but samples prepared from KF-1100 (Kureha) powder.

ac current at lOHz measured on 2miil Pennwalt PVF2 subjected to various

field/temperature history.

(a) fresh film

(b) following application of 16V/ym at 108°C and cooling in field.

(c) following poling (110°C, 3 hrs., 37V/vim)

.

(d) following 17 hrs. anneal at 148°C.

TSC current m^easured on 25ym Kureha PVFp_ subjected to various field/

temperature history. Starting temperature for each cycle -100"C.

Heating rate 8.3°C/min.

-107-
FIGURE CAPTIONS (con't)

(a) fresh film under dc field

(b) depolarization current following (a).

(c) under field after four coPiplete not shov/n (a,b) sequences.

(d) depolarization current following (c).

Figure 11 ac current near Tg recorded at lOOHz for Kureha sample subjected

to various field/temperature history.

(a) heating trace fresh sample.

(b) heating trace for sample' (a) with AOV/um- applied,

(cj cooling trace for sample (bj with 40V/ym applied.

TABLE 1 - . . .

Concentrations (in percent of total weight) of a-, B- and y-phases calculated

from DSC traces for tne 821 samples discussed in the text. A comparable heat of

fusion (25 cal/gj was assumed for the three phases. The concentration of the y-phase

is strongly time (t ) dependent.

A

a 6 Y

420 . 46±4 -v^

430 44±4 2±1

432.5 40±4 4±1 2±1

435 37±4 4±1 5±1

-108-

I
-100 -50 0 50 100 150

-109-
emf(mV) -5-4-3-2-1 0 I 2 3 4 5

-150 -100 -50 0 25 50 75 100 125

T*»C

-110-
-111-
2.5 3.0 3.5 4.0 4.5

-112-

400 420 440 460

-114-
e 3

o

>

o<

ro

IO

^

2

Q.

UJ

<

Tx = 435^K

-100

Tx=432.5*»K

7

//////

Tx = 420*»K

a

KUREHA

25/im

-50 50 100

J_

150

-115-
1 —

F

-50 0 50 100 150

-116-
— 37V//tm JIO°C

I I I I U

60 80 100 120 140

-117-
-5n I I—I—f—1—I

\—

r

T(«C)

-118-
TEMPERATURE

-119-
Piezoelectricity and Pyrdelectricity in a Poly (vinylidenefluoride-tetrafluoroe thylene) Copolymer

G. Thomas Davis

National Bureau of Standards, Washington, D. C.

This paper ascribes the piezoelectric and pyroelectric activity of

polymers to frozen-in dipole alignment achieved during the poling procedure.

In such a model, electrical response to temperature or pressure change results

from a change in polarization arising from volume change and a change in

fluctuation amplitude of oscillating. dipoles . The development of the theory

will be reviewed briefly and then compared with results obtained for polyvinylchloride and a copolymer of vinylidenefluoride and tetraf luoroethylene.

The theory to be outlined here has been developed and presented in more

detail by Broadhurst et al^'''^ and by Mopsik and Broadhurst

Consider a thin film of polymer between two metal electrodes subjected

to an electric field of intensity E . The dielectric displacement D in a

P

direction normal to the surface of the film is given by:

D = e op. E + P (1)

where is the permittivity of vacuum and P is the polarization. Since D

at any point within the dielectric is defined as the product of E and the

permittivity ^' o'^s can express the resultant polarization as:

P = (e' -1) e E„ (2) o p

where is the relative permittivity. The relative permittivity of a

polar polymer which undergoes a glass transition varies with temperature as

shown schematically in Figure 1 where an increase in occurs in the

-120-
vicinity of the glass transition temperature. If an electric field is

applied at a temperature in the liquid range above T , where e' = the

polarization induced is expected to be:

If the film is then cooled below T before removing the electric field,

g

only that portion of the polarization that could have been induced at

tempetatures in the glassy region will' disappear. That polarization is

equal to (e' - 1) e E where is the relative Dermittivity at the ^ g o p g

temperature in the glassy region where the field is removed. The polarization remaining after the removal of the field is expected to be:

P = Ae'e E (4) r 0 p

•

where Ae'=£'-e'. I g

In the case of an amorphous polymer this phenomenological approach

leads to the result that the numerical value of is expected to increase

linearly with E^ . In order to predict how P is expected

to change with temperature or pressure, it is more helpful to consider P on

a molecular level.

In molecular terms, polarization is the net dipole moment per unit volume, i.e.

P = (5)

V

where <m> is the averaged moment considering both the vacuum dipole moment and

the induced moment due to the electric field in which the dipole finds itself and

'which is caused by the preferred orientation imposed on the system during the

poling process. N/V is the number of dipoles per unit volume. The problem of

-121- i

!
calculating <m> for frozen dipole alignment is similar to that solved by

Onsager for uniform liquids and has been worked out for this case by Mopsik

(2) and Broadhurst leading to the result that:

(eco + 2) Np . P = 3 <cos 9>. (6)

where is the dielectric constant for immobilized dipoles, is the

vacuum dipole moment of the molecule and <cos 9> is the averaged cosine of

the angle between the individual dipole moments and the total sample moment.

The polarization results in a compensating surface charge within the

electrodes on the surface of the polymer film. Real charges trapped within

the dielectric will also be compensated for by charges within the surface

electrodes but in this treatment we assume that trapped real charges remain

trapped and sample strains are uniform so that the trapped charges do not

contribute to any current flow between electrodes when the sample is strained.

The influence of any trapped charges on the induced dipole moment has also been

neglected in this model. A change in temperature or pressure causes a change

in volume polarization which results in a re-distribution of charge between the

surfaces under short-circuit conditions or generates a voltage between the

surfaces under open-circuit conditions. Since we treat the case where electrodes

are evaporated onto the surface of the polymer and consequently maintain the same

area (A) as the polymer, we can substitute the thickness (L) for V/A. An outline

of the change to be expected in P for a change in temperature (T) or pressure (P)

is presented in Figure 2 and may be found in more detail in Reference 2.

-122-
The piezoelectric coefficient d for a change in hydrostatic oressure

P as outlined in Figure 2 is

d .iifi = - P . « . (7)

p A dp " L

and the pyroelectric coefficient p is

P = ^ - ^ \ ^» [1 ^TnT ]. . (8)

00 L

Note that these coefficients are defined as the change in charge Q, the usual

measured quantity, as opposed to the classical definition which considers the

change in polarization P with strain. In the case of polyvinylchloride where

P is linear with the poling field, E^,

P = Ae'e opE (9)

so that: j a ' o r- d = Ae e £ 8, E (10)

p o "=° L p

Slid A ' T7 n J_ A^!^ , ,T - . p = Ae e £ E [1 + — j (11) ^ o " L p 2Ta £

which upon rearrangement yields:

d ^ . = A£'£^ (12)

p L o

- =,,-e (13)

E £ " 2Ta^

p L o L

The measurement of p and d^ have been described elsewhere ^''"^ but briefly,

one measures the amplified short-circuit current that flows from one electrode

of the polymer film to the other for a measured rate of temperature change or

a measured rate of pressure change supplied by a pressure fluid such as helium

gas. The thermal expansion coefficient and compressibility of the polymer films

have been measured by using three small pieces of the film to separate the plates

of a parallel plate capacitor. The fractional change in thickness as one changes

temperature or pressure is equal to the measured fractional change in capacitance

(corrected for e' of the helium in the chamber).

-123-
A summary of results obtained with poljrvinylchloride by Broadhurst

(1 2) et. al ' is presented in Table I. Since. the measured values of Ae'and

^00 ^^^^ polymer were 10 and 3 respectively, the product of 30 is in very

good agreement with the average value of 26.7 for the quantity d /E 6 e . p p L o At present, we have no direct measure of A<}» „ which is required for

comparison of the pyroelectric coefficients but a value of l6° at room

temperature would be required to account for the measurements reported in

the last column of Table I. Such a value seems very reasonable.

PoljrvinylidenefluorideCPVF^) and some of its copolymers [as well as

polyvinylf luoride (PVF)] can be made to exhibit piezoelectric and pyroelectric

responses much larger than those of PVC. These polymers differ from PVC in

that they are partially crystalline at room temperature and the glass transition

of the non-crystalline portion is far below room temperature, in the vicinity

of -AO^C. Furthermore, the polarization (as implied by the piezoelectric

response) induced during the poling procedure is not linear with applied field

(3 4) but appears to approach a saturation level .

For our initial investigations we have chosen a copolymer of vinylidenef luoride containing 27% tetraf luoroethylene. The advantage of using the copolymer

is that it crystallizes directly from the melt into a crystal form analogous

to that of form I (or B) of PVF2 homopolymer. This obviates the need to determine

crystal orientation (assumed random from the melt) and the relative amounts of

a and g crystal form in the sample. Furthermore, the copolymer is available

commerically from the Pennwalt Corporation as Kynar 7200. The copolymer is

soluble in a variety of polar solvents and also crystallizes from solution into

the potentially active 6- form crystal . Unfortunately the incorporation of so

-124-
much TFE reduces the maximum polarization to be expected and also lowers

the melting point to the vicinity of 130°C which may affect the stability

of the electret at room temperature.

Since is so far below room temperature, we assume that no dipole

orientation remains in the non-crystalline regions of the polymer at room

temperature after removal of the poling field. When dipoles within the

crystalline region of the polymer become aligned, it is unlikely that entire

crystallites can be moved under the influence of the poling field because

of the intercrystalline links which maintain the Integrity of the polymer

film. However, It is not unreasonable to assume that individual chain segments

within a crystallite can rotate about the chain axis—a mechanism advanced for

the so-called a relaxation in polyethylene A model depicting such a

rotation is sketched in Figure 3. The lefthand side of the figure depicts

portions of two different crystallites in the unpoled film. The upper crystallite

represents a chain orientation parallel to the plane of the film and

normal to the cross-section shown. The lower one depicts a crystallite

in which the molecular chains lie in a plane parallel to that of the paper

but Inclined to the plane of the film; furthermore, the b axis of the unit

cell is inclined 9 degrees away from the plane of the paper. In the poling

process, we envision rotation of the dipoles about the chain axis so as to

align as much as possible in the direction of the field, coupled with slight

translation relative to each other so that it "recrystallizes" with a reorientation of unit cells within the crystallites. On the right, the upper

-125-
crystallite exhibits maximum dipole alignment after poling. Chains within

the lower crystallite have been rotated so that the maximum dipole alignment

consistant with not changing the chain axis orientation is achieved. That

is, the b-axis of the unit cells (and therefore, the net dipole) now lies in

a plane normal to that of the film. However, only that component of the chain

lying in a plane parallel to the surface can contribute to the net polarization

upon poling since we have assumed that the chain axis orientation remains fixed.

.Since the polarization in the P,yE2-TFE copolymer is not linear with applied

field, one cannot infer the polarization from the expression used previously

for PVC. However, by using Equation 6 we can predict the maximum value of P

to be expected for a single crystal of the homopolymer and then modify it to

apply to the copolymer. When the dipole moment of vinylidenefluoride is taken

as an average of those reported for difluoroethane (2.27 D) and dif luoromethane

(1.97 T))^^\ and the unit cell volume is 53.9 A° ^ as reported by Hasegawa et al^^

2 the polarization of a single crystal of PVF^ would be 13.2 yC/cm which becomes

n 22 yC/cm^ when corrected for the interaction of the polarizable dipoles with the

surroundings according to (e^ + 2)/3 where = 3. The corrections to this

maximum value when applied to the copolymer are outlined in Figure 4.

In the ratio of volumes in Figure 4, V , is the volume occupied

copolymer

by an average copolymer unit calculated from the macroscopic density of the

sample and then expressed in per average repeat unit. In the model employed

here, only the crystalline portions of the polymer contribute to the polarization

so that the maximum polarization is reduced by the degree of crystallinity

.

-126-
We have assumed a uniforn distribution of TFE units between the crystalline

and non-crystalline regions of the copolymer and since only the vinylidene

fluoride portion contributes to the polarization, the maximum must be reduced

by the fraction of VT^ in the copolymer.- Even within the vinylidenef luoride

portion of the copolymer, only the usual head to tail additions of monomer

contribute to the net dipole. This has been assumed to be 95%. Finally, as

mentioned in reference to Figure 3, only that component of the chain axis

which ^ lies in the plane of the film is considered to contribute to the realizable

polarization. Since there is no "directionalty" required in the chain axis,

2 , we take the average of cos c where : is the angle between the chain axis

and the plane of the film. Assuming random orientation, this average is 2/3.

Measurements have been made on two types of filns; one cast from solution

and supplied to us by Pennwalt, the other pressed from molten polymer and

3 quenched into ice water. The density of 1.843 g/cm and a diffuse low angle

x-ray diffraction ring corresponding to spacings ranging from 120 to 155°A

were nearly identical in the two types of films. Film thickness ranged from

25 to 50 um. Metal electrodes were evaporated on the films and the poling

was done inside a can immersed in. an oil bath at constant temperature. After

the prescribed length of time at the poling temperature, the sample was cooled

by quenching the can into room temperature water with the electric field still

applied. At the time the experiments were started, it was assumed that all

of the "poling" took place at the elevated temperature and the time that the

field was applied at room temperature subsequent to that was not important.

More recently we have found that appreciable poling can be achieved at room

temperature. Polin? at room temperature has also been renorted for the

PVF horcopolymer

-127-
The results obtained using the cast film are summarized in Table II

and those for the melt-crystallized films are .summarized in Table III.

The poling tine was 35 minutes in all cases. Multiple entries indicate

results for replicated experiments and are entered in the same order in

the tables for both coefficients. (Note that the field strengths and the

poling temperatures are not exactly the same in Tables II and III.) Results

for the two types of sample preparation are essentially the same. At the

higher temperatures and higher field' strengths, the resulting electrical

response to strain reaches a nearly constant value—especially when the

pyroelectric coefficient is considered. There is appreciably more scatter

in the data for response to pressure perhaps due to the inherent temperature

changes which accompany the expansion or compression of the pressurizing

fluid. In the case of the melt-crystallized film where data was obtained

over a wider range of poling temperatures, the response is nearly independent

of poling temperature at the highest field strength, employed. There are

frequent break-downs at 500 kV/cm and elevated temperatures so higher field

strengths have not yet been tried.

To enable comparison of the results in Tables II and III with the expressions for d and p, (Equations 7 and 8) we have also measured a and 6, for the

P Li Li . two films indicated by the asterisk in the tables. For the cast film, 6^^ =

2.05 X

10~'^^ m^/N and a = 2.2 x 10~^ k""*" while for the melt-crystallized film we

Lt obtained 3. = 2.15 x lO""*"^ m^/N and = 1.76 x lO""^ k""*". (From volume dilatometry on

-4 -1 a sample of the copoljrmer which is about 44% crystalline, we find a = 1.72 x 10 K )

-128-
From the literature, the dielectric constant for immobilized dipoles has

been taken to be about 3. Unfortunately, we do not as yet have a direct

measure of P but we can now estimate what its value would have to be according

to this model to account for the maximum measured value of d . P

P = d /S^ £ = 1.92 X 10"^C/cm^

p L »

2 This value of P is well vrithin the maximum value of 3 yC/cm which one might

expect to achieve with this randomly -oriented copolymer.

From the ratio of d to p and the measured values of a and B, one can

P obtain an estimate of A?, the RMS amplitude of oscillation of the dipoles

in the copolymer. At 300 K, As would have to be 44° (<i>^ = 62°) (calculated

from tabulations of Bessel functions rather than the small angle approximation)

.

As opposed to the 14° found for PVC which seemed very reasonable, a value of

44° seems unusually large. Because of the seemingly large estimate of A^i,

the details of this model as applied to semicrystalline polymers may require

refinements.

Acknowledgement

Partial support of this work by the Office of Naval Research is

gratefully acknowledged.

-129-
References

1. M. G. Broadhurst, C. G. Malmberg, F. I. Mopsik and W. P. Harris, Electrets,

Charge Storage and Transport in Dielectrics, M. M. Perlman Ed., Proceedings

of an International Conference on Electrets sponsored by the Electrochemical

Society, October 1972, Miami Beach, Florida, p. 492.

2. F. I. Mopsik and M. G. Broadhurst, NBS Report No . 2 on ONR Task No. 74-622

("Molecular Dipole Electrets") and also submitted for publication in

J. Appl. Phys.

3. M,. Tamura, K. Ogasawara, N. Ono and S. Hagiwara, J. Appl. Phys. 45 , 3768 (1974).

4. M. Oshiki and E. Fukada, J. Materials Sci. 10, 1 (1975).

5. J. D. Hoffman, G. Williams, and E. Passaglia, J. Polymer Sci. C14, 173 (1966).

6. R. D. Nelson, Jr., D. R. Lide, Jr., and A. A. Maryott, "Selected Values of

Electric Dipole Moments for Molecules in the Gas Phase" NSRDS-NBIO; U. S. Government Printing Office, Washington, D. C. (1967).

^ 1. R. Hasegawa, Y. Takahashi, Y. Chatani, and H. Tadokoro, Polymer J. _3» 600 (1972)

8. H. Kakutani, J. Polymer Sci. A- 2 8^, 1177 (1970).

-130-
I

TEMPERATURE

Figure 1. Schematic representation of relative permittivity of an amorphous

polymer as a function of temperature at constant frequency.

-131-
AOWOOV

A

CD

i-J <

II >

0)oA

(D«ouVo +

to

(1) +

13 <A

COouV

3. +

CM A

(X>

COo

CJ +r

(0

OlH

-e- CM

De o

(-1 -e- 3C

O

CO -e-

-e?H

-i«ir 3

CM

M

A

CM

O

-e- V

u

<uuo

4J

CO

)-l •H

IDC

•H

A

05 ou

diUoo

l-l (UN

-e- <

to <o

CM

-e- <

CM

•e- < A

CD

COouVI

II I

CO

CM

-eCM

CM

A

CD I

13a

CO o

II

13

CO oV c u c 8 o CO o CO 3 CO OJ o 11 o CO .u U-l (0 »4 V 3. < CO > CO (J o 3 < z c <-i V iH V + CO PU o ' .H (1) CO o a o •H 1 CO (U > + CN U CO a V4 z II CO II 01 >> a + cn II . •O PQ >l A A c —s 8 a |fo o CD II II <D 1 •H CM —V

•H o- ^—•

II tH CM c CO /—s

CO II 0) + cn + 0 0 O o o H

II 1 1 O

H -e- -e- o 13 c 8 T3|'d V •a •

y V CO

w II 8 8 CO

O x: « w T3 >-) 13 rH|< u

4J 01H

T3 T3 s—'

II

O II H CO u o

cy < A.

CO a. 'H o 4J o

•H •H

V

U 4J

4J c u U CO CO

(U 3 0) iH o* c

0) •H O 01 N. (U H (U •H

O- • u

CO CO o 0) •rl iH C O o

•H •H

4J 13 13

CO c: > 13 CO •H <V

c:

u

0)

a

-a •H 0) rH •H

a> COO

•H

U-lo

M-l

U-l0) o e o

0) CJ

(U 4-1 c CO c

•H >^ o

iH CO •H

(J CO 3 CO c O CO u a. c o X <

CMM3

00
Q

LiJ _lO

60 QJ C -H

n] 4J U iH

rH

0) 4-1 (A n o >^

0-

u o o u

CU

u

<u

m

B

>^ c o

•H O- O 3

•Hw

CO

W -rt C -U

(U CO

CO 4J

(U

C

V4 (U

Cu. -rl 0) l-i Vj O O <U 3

•H iH

W

4J

O 0)

rt o- >J S -rl 0) "O rt 4=o fi w

Cfi -H rt
(e + 2) ^PVFP = :i

• — . X . X . f„ „ . <COs2 6> max 3 V v i H-T PVF^ copolymer

(^^ + 2)

U _ (3 + 2) 2.12 D 5 TO 1 Vic ^- Mc

Vp^p 3 53.9 3 cm^ cm^

VoTTx. /V 1 = 53.9/66.4 = 0.81 PVF2 copolymer

X = Fraction of sample which is crystalline = 0.36

X = Fraction of sample which is vinylidenef luoride = 0.73

f = Fraction of head to tail addition within VF. = 0.95

H-T 2 <cos^ 6> = Projection of chain axes into plane of film

= 2/3 for random orientation

P =3.0^^2

max cm

Figure 4. Maximum polarization of PVF^ single crystal corrected for

semicrystalline copolymer with random orientation.

I

-134-
Table I. Summary of piezoelectric and pyroelectric

data obtained with polyvinylchloride

.

Sample E , kV/cm

P d * pC/N

P

A4 40 0.079 "

A3. 80 0.173

A2 160 0.363

Al 320 0.619

PA 344 0.777

PB 350 0.735

PC 349 0.649

PD 351 0.604

a = 0.78 X 10 K

6^ = 0.86 X

10~-'-° m^/N

d

*+ , 2 2 E

D, nc/cm K E 6, e E e

' ' p L o p L o

0.0127 25.9 46.1

0.025 28.4 44.7

0.056 29.8 45.8

0.095 25.4 43.1

0.114 29.7 48.1

0.105 27.6 43.3

0.104 24.4 43.2

0.102 22.6 42.0

Avg. 26.7 44.5

*The values reported in reference 1 incorrectly measured a contribution to

the current from the region outside the area of the evaporated electrodes.

+ In addition to the area correction, the values in reference 1 were reported

p = C"^) whereas here p is defined as ^ *

-135-
Table II. Pyroelectrlc and piezoelectric coefficients for

PVF2-TFE copolymer film cast from solution.

2 p, nC/cm K

T, °C

E, kV/cm 50 60 70 • 80

200 l.A

300 1.0 , .2.6 2.7 2.5*

500 1.9 2.4

dp. pC/N

T, °C

E. kV/cm 50 60 70 80

200 5.1

300 3.9 10.4 11.2 8.5*

500 7.1 10.1

Thermal expansion and compressibility measured on this specimen.

-136-
Table III. Pyroelectric and piezoelectric coefficients for

PVF2-TFE copolymer film crystallized from the melt,

2 p, nC/cm K

T, °C

E. kV/cm 0 30 60 80

100 0.04 , .1.1 1.4 1.2

300 1.5 1.9 2. 6*, 2. 7 2.0,2.2

500 2.4,2.2,2.0 2.3 2.2

d^, pC/N

T, °C

E. kV/cm 0 30 60 80

100 0.5 2.3 5.4 5.7

300 5.4 7.2 9. 8*, 12.1 6.1,8.1

500 11.7,7.7,9.0 7.5 8.2

Thermal expansion and compressibility measured on this specimen.

-137-
POLYVINYLIDENE FLUORIDE AS AN ACTIVE DEVICE ELEMENT

by •

J. G. Bergman

- Bell Laboratopies

Holmdel, New Jersey 07733

ABSTRACT

» ... A general summary of some of the device applications

which have recently been found for po ly viny lidene fluoride,

such as infrared detector elements, nonlinear optical mediums,

as well as an active element in a copying process is given.

These applications represent some of the first instances where

the non-passive characteristics of polymers are successfully

employed. "

-138-
POLYVINYLIDENE FLUORIDE AS AN ACTIVE DEVICE ELEMENT

by

J. G. Bergman.

Bell Laboratories

Holmdel, New Jersey 07733

Over the last few decades polymer usage has shown

enormous growth, in most all of these applications however

one of the major rationales for usage has been the "passive"

character or chemical inertness of polymers. Now however,

the advent of polymers as "active" elements in state of the

art electronic and optical devices is upon us.' It is our

purpose here to review some of the recent pyroelectric and

nonlinear optical applications of one polymer in particular,

polyviny lidene fluoride (PVF2). We will first review some of

the structural and optical properties, then discuss some

pyroelectric applications and lastly a par"ticular device

application namely using PVF^ as the active element in a copying process will be discussed.

PVF2 is a highly crystalline polymer (CH2CF2)^

1 2 belonging to the acentric polar group mm2 , ' In Fig. 1

we see a model of PVF2 which shows that the C-F and C-H bond

dipoles have a resultant dipole moment which is perpendicular

to the polymer chain axis, i.e. parallel to z in Fig. 1,

(For later reference, a coordinate system is defined with Z parallel to the polar axis of the chain and Y parallel to the

-139-
chain as shown in Fig. 1). If we assume reasonable values-^

for the bond dlpoles I.e. [y(c-c) O', y(C-F) = 1.4D and

y(C-H) = O.^D] we find that the calculated spontaneous

polarization is between 8 and 16 yC/cm , depending on the

polarity of the C-H bond. The experimental value, obtained

ij 2 by Buchman is 8 yC/cm . It would be nice to be able to

conclude here that this now determines the sense of the C-H

bond polarity but this is not the case, since even in well

defined cases where good structural data and measured molecular

dipole moments exist (such as in the case of Thiourea ) the

calculated P is only half the size of the observed value,

s- The mechanical treatment which produces the texture

in the PVFp film does not, of course,, affect the sense of the

microscopic polarization (as many dlpoles point out one side

of the film as out the other side) . Electrical poling viz

6 6 (10 V/cm at '\>120°C) is required in order to make the films

have uniform polarity. A possible poling mechanism is shown

in Fig. 2. The model illustrates a l80° rotation of the

molecular dlpoles about the polymer chain axis.

The degree of ordering due to mechanical treatment

is shown in Fig. 3 where we compare the two birefringes as a function of the thickness of the film. One also sees that

the magnitude of the pyroelectric coefficient is also, as

-140-
expected, related to the degree of order. Induced in the

film by stretching. We see that the thinner biaxially

oriented films have the largest p.yroelectric coefficient

2

('^^2.4 nC/cm °C). In the case of the biaxjal films, there

are three different values of the principal indices n^, n^

and n . Here n is the film normal or polar axis and n

z z y

correspond to the preferred direction of the polymer chains.

A confirmation of the biaxial nature of the film is shown

In Pig. 4 which shows a conoscopic Interference figure. (A

stack of films was used in order to get a sufficiently long

light path in the material).

Once the films are poled they become activated for

second harmonic generation (SHG). A -schematic representation

of the experimental set up is shown in Fig. 5. A detail of

the sample configuration is shown in Pig. 6. A Nd:YAG laser

(X=1.06ym) with a peak power of '^100 watts was used to

generate harmonics at X = 0.53yni. Samples were cut from the

19um thick biaxially-or lent ed film in the form of small wedges

whose shape and orientation are shown in Pig. 6. The laser

beam propagates in the plane of the film ( | jk) and is focused

so that its waist is contained in the film. To prevent burnin

only "^15% of the full laser power was employed. The second

harmonic signal, separated from the fundamental by suitable

-141-
filter (Fig. 5), was then detected with a phot omult Ipl ler

.

According to the X-ray measurements, the polar phase of PVF^

has point symmetry mm2 (Z is the polar axis). The components

of the second harmonic polarization for mm2 symmetry are

given in terms of the fundamental electric field components by

P = 2d_,E E

X ^ 31 X z

P = 2d-„E E

y 32 y z

P = d_,E^ + d-^E^ + d^^E^.

z 31 X 32 y 33 z

Thus there are three independent second order nonlinear

coefficients, d^-]^, ^^'^ "^33* These equations predict

that when the laser is polarized normal to the film [|z), the

second harmonic will also be polarized in the z direction

2

(P = P = 0, P = d-,-,E ). This result is observed experiX y ' z 33 ^ mentally i.e. (d^^ coupling). The above equations also

predict that when the laser is polarized in the plane of the

film (||x), the second harmonic signal should be polarized

normal to the film i.e. orthogonal to the fundamental

2

(P = P = 0, P = d^,E ). This result (d^, coupling) is

X y ' z 31 '

V 1- <=> also observed experimentally. Finally, if a wedge is cut

having its length parallel to y and the laser beam is polarized

in y direction no SHG is observed. This result implies that

-142-
d^^ is very small compared to d^^ and d^^* The magnitudes

of the coefficients, relative d^^ (Quartz) were found to be

d^^(?YF^) ^ 2d^^(PVP2) d^^(Si02) >> d^^iFVF^)

We should note that the very small size of d^^ is

to be expected since from Fig. 1 we see that for E | | y we

"cannot couple to either the C^F or C-H bonds. Hence we must

get our signal from only the carbon bonds, which because of

their centro symmetric nature should give a very weak (if any)

signal

.

It 'is easy to use the poled PVP2 films as pyroelectric

detectors of electromagnetic radiation. The essentials of

such a detector are shown in Fig. 7. The poled film is cemented

with conductive epoxy to a metal heat sink. The conductive

front-face electrode is made to be either highly absorbing

(gold-black) or partially transmitting (very thin gold film)

for the incident radiation. If the load resistor, R, is small

compared to the resistance of the film and is also small

enough that the RC time constant of the load circuit is

negligible, the voltage response, , of the detector to a modulated beam is given by

dP ^dP\ dT

-143-
where A Is the electroded area of the film and P is the

polarization of the film. Since dP/dt is the pyroelectric

coefficient we see then that is proportional to dT/dt

,

the heating or cooling rate of the film. The response of

two detectors to a 90 Hz chopped beam is given in Pig. 8.

Detector A is constructed as in Pig. 7 with its entire back

sur^'ace in contact with the heat sink. In this case the

thermal response time is only a few milliseconds and thermal

equilibrium, as evidenced by the decreasing response, can

take place during the beam-on or beam-off periods. Detector B

has the PVP2 -film mounted in a drum head fashion with only

its periphery in contact with the heat sink. In this case

the thermal response time is approximately one second with

the result that the heating and cooling rates are constant

during the on- and off-periods, and a square-wave response

is obtained from detector B. The behavior exhibited in

Pig. 8 identifies the effect observed as true pyroelectriclty

i.e. (a reversible dependence of the polarization on temperature) and not simply depolarization i.e. (an irreversible

decay of the polarization with time when the sample is heated)

as is often studied in electrets. In Pig. 9 we show the

response of a PVP2 detector to a Q-switched laser pulse

of submicro second risetlme. Also shown is the same pulse as

-144-
detected by a Cu doped Ge photoconduc t Ive detector operating

at 4.2°K. In Fig. 10 we see the set up for pyroelectric

scanning experiment the results are _shown in Pig. 11. We

see from Fig. 11 that the pseudo 3-dimensional shows basically

no response outside the poled area and considerable nonuniformity in the poled region. One can thus conclude that

if domains exist in these samples they must be smaller than

'^'10 ym.

Common materials have pyroelectric coefficients

O p dP/dt 'V 10" C/cm °C. One can see that a localized 1°C

temperature rise can generate a localized electrostatic

—8 2 charge on the order of 10 C/cm . Since conventional

electrostatic copiers based on the photoconduc tlve effect

also involve local charges of this order, we have attempted

to employ the pyroelectric effect as a basis for a copier.

In Fig. 12 we see the simplified demonstration of the effect

in LlTaO^ . The bottom surface of the LiTaO^ crystal

C'^10 milxl cm ), coated with a conducting layer of aluminum,

is at ground potential. Initially, both blocks of copper

are at the same temperature, and the top surface of the

crystal is wiped free of stray charges with a Tesla coil

or salt solution. Then heat is applied to the left-hand

copper block until a temperature rise of about 6°C is obtained

on the left-hand side of the crystal, as shown in Fig. 13(a).

-145-
Since the pyroelectrlc coefficient of LiTaO^ is "^2 x 10 C/cm °C,

2

one generates a charge density of "^120 nC/cm . The approximate

charge distribution is shown in Fig." 13(b). A developer

consisting of negatively charged toner loosely bound to

positively charged steel shot is poured over a baffle and then

over the crystal surface. One then observes that the toner

(negative charge) is attracted preferentially to the region

where the electric field [Fig. 13(c)] is directed outward

from the surface. This effect is enhanced if the baffle is

at a negative potential (-1.2kV) in order to increase the

negative charge on the ink and is degraded if the baffle is at

a positive potential. The same effect was observed for a

0 . 020-in. -thick l-in.-diam ceramic disk of lanthanum-modified

lead zirconat e-titanate (PLZT). In the material used, 6% of

the lead was replaced by lanthanum and the Zr:Ti atom ratio

was 65:35. A more dramatic demonstration of the localized

heating and its concomitant charge distribution is shown in

Fig, 14, A 65O-W lamp passes light through the object slide,

projecting the image on the pyroelectrlc element, PVF2

(dP/dt%2xlO~^ C/cm^°C). The PVF2 is oriented so that the

heated region develops a negative charge and the cool region

has a (relative) positive charge. Hence, the toner (negative)

is attracted to the cool region. A photograph of the developed

image is shown in Fig. 15(b). Negative images were readily

-146-
obtained by neutralizing the pyroelectric element after the

image had been projected and then allowing the pyroelectric

elem.ent to cool dovm, thus reversing the charge distribution

in sign only. Images have been transferred from the PVF^

7 film to paper by standard techniques. A representative

printed copy is shov;n in Fig. 15(c).

If a crystalline (i-.e., ordered dipolar) homogeneous

medium is ass-oned, the expected resolution limit should be

of the order of the film thickness, since any projected line

whose width is thinner than the film thickness may be substantially broadened by thermal diffusion. The (0.8-mil)

PVF2 samples used in these experiments v/ere "^50% crystalline,

the crystalline phases being divided between the useful

orthorhombic phase and a second monoclinic phase. The distribution of the orthorhombic phase (i.e., homogeneity) in

the film is not kno^-m at this time. We are therefore hopeful

that improving ?" ^ or finding other materials (hence, with

even larger pyroelectric coefficients) will result in substantial

improvement s

.

We have demonstrated than an electrostatic copying

process based on the pyroelectric effect is indeed feasible.

Furthermore, since the process is inherently wavelength

independent (unlike a phot oconduct ive process), it is more

-147-
readily adaptable to color copying or to studying thermal

distributions in the Infrared. Further experiments on other

pyroelectrlc materials as well as" other developing processes

are In progress.

The author gratefully acknowledges the collaboration

of J. H. McFee, G. R. Crane and A. M. Glass without whom

these results would doubtless never be obtained.

-148-
REFERENCES

J. B. Lando, H. G. Olf, and A. Peterlln, J. Polymer Sci.

9^1 (1966).

H. Kawai, Japan J. Appl . Phys.. 8_, 975 (1969).

C, P. Smyth, "Dielectric Behavior and Structure",

McGraw Hill, N. Y. (1955) p. 244.

P. Buchman, Ferroelectric s ^, 39 (1973).

^G. J. Goldsmith and J, G./v^/hite, J. Chem. Phys. 31^,

1175 (1959); the Handbook of Chem. and Phys. gives

y (Thiourea) = 4 . 9D

.

More detailed descriptions of these processes can be

found in:. J. G. Bergman, J. H, McFee and G. R» Crane,

Appl. Phys. Lett. l8_, 203 (1971 ); J, H. McFee, J. G, Bergman

and G, R. Crane, Ferroelectrics 2> 305 (1972); A, M. Glass,

J. H. McFee and J. G. Bergman, J. Appl. Phys. 42_, 5219

(1971); J. G. Bergman, G. R. Crane, A. A. Ballman and

H. 0 'Bryan, Appl. Phys. Lett. 21, 497 (1972); and references

therein.


Birefringence and pyroelectric coefficient of

poled PVFg a function Cf the method of film

preparation.

Pyroelectric

. ' Coefficient

Description (n^-n^) (n^-n^) ^^^^^^2 o^)

0.4 mil

(blaxially stretched)

0.018 0.011 2.4 ± .7

0.8 mil

(blaxially stretched)

0.019 0.011 2.4 ± .7

2.0 mil

(uniaxially stretched

in the x direction)

0.000 0.013 0.7 ± .2

5.0 mil 0.000 0.000 .0.3 ± .1

Indices measured at 6328A using an Abbe refractometer;

n = 1.425; estimated errors in An are ±0.002: n

ave z corresponds to the film normal or polar axis while n

and n refer to the orthogonal directions in the plane

y

of the film having the smallest and largest indices,

respectively

.

All films obtained from Kureha Chemical Co., Tokyo, Japan.

-152-

LENS • FILTER

I

LASER

—>

V

2u>

CRYSTAL

2u;

POLARIZER

DETECTOR

ANALYZER

Fig. 5

-154-


-*| I*- 2xlO"^sec

Del. 3

r = 1 sec

Response of PVF, pwocleclric dcieciors lo 90 Hz

chopped beam from CO- laser. Deiector A: Tlierma! response

time%, = 2 X 10- - sec. Deiecior B : t, ^ 1 sec In boih cases the

load resistance R is small enough thai Eq. (11 is valid.
Fig. 9

-158-


50 WATT

HEATER

COOLING

WATER

L

^LiTa03

CRYSTAL

NEGATIVE TONER

COLLECTS HERE

UNDER CONDITIONS

SHOWN IN FIG. 2

Pig. 12

-161-
+5X10

(c) l^x 0

(V/cm)

-5x10" h

+6x10,-8

(b) 0

(COUL/cm^)

-6x10"® h

(a)

f

20

14

- ® ® • o o 0

^ HOT

REGION

•

1 1

COLD

REGION

9

'9

O • • O 9 9

1

i 2 3.4

POSITION IN 0.10" UNITS

Pig. 13

-162--
rG.E. DVY

LAMP

/PYROELECTRIC

\ ELEMENT

,12.7 cm f.l.

"^"^.^

OBJECT

25.4 cm—»4-^—25.4 cm —>-

6,35 cm

Fig. 14

I
-164-
Scanning for FerroelGCtric ity in

Polycrystalline Materials

by

M. Litt, Che-hsiung Hsu, P. Basu and S. M. Aharoni

Department of Macromolecular Science

Case Western Reserve University

Cleveland, Ohio 44106

; ABSTRACT

The dielectric behavior of polycrystalline pellets of 2-

alkyl ben: imidazoles are reported in this paper. The experimental technique consists of measuring the dielectric constant

of a polycrystalline pellet as a function of temperature. Dielectric anomalies were found for which the position of the

dielectric and loss tangent peaks were independent of frequency,

which showed a solid/solid first order transition. When the

sample was held under a voltage gradient while the temperature

was dropped through the transition temperature, reversible

charge storage showed poling of the sample, and therefore,

implies possible ferroelectric behavior.

The experimental procedure was validated on thiourea, a known ferroelectric. 2-Ethyl ben: imida-ole showed the same

phenomena at a transition temperature of about -90°C and is

probably a new ferroelectric material.
Scanning for Ferroelectr icity in

Polycrystalline Materials

by

M. Litt, Che-hsiung Hsu, P. Basu. and S. M. Aharoni

Department of Macromolecular Science

Case Western Reserve University

Cleveland, Ohio 44106

INTRODUCTIONWe have been interested in organic ferroelectric materials

for several years and have been endeavoring to find some new

classes in the group of materials which have infinite chains

of H-bonds. It has been shown theoretically and experimentally

that, when dipoles are associated to form a chain, the magnitude of the dielectric response increases with the number of

dipoles in the chain^>2. In addition to chains of K-bonds,

to be potentially ferroelectric, the bonding must be reversible.

Thus the compound should have two isomers of identical stability,

where the transformation from one isomer to the other is accomplishec

by proton switching. This also reverses the dipole involved.

Benzimidazoles , were reasonable candidates to study for the

following reasons. They are known to associate through 11-bonding

and have high dipole moments. NMR studies^ indicate that the

two nitrogen atoms in benzimidazole are electrostatically equivalent. This implies that the associated H-bonding dipoles could

be capable of reversal from one position of equilibrium to another

of equal energy. This is shown schematically in Fig. 1. Because of these factors, it was decided to investigate the dielectric properties of 2 -substituted bcnzimi dazolcs** . Since

-166"

! . I J
much time and strenuous effort is required to grow single crystals, and techniques for the study of dielectric constants and

pyroelectric effects in polycryst al 1 ine materials have already

been reported in the 1 i terature ^ ' °

, polycrys tall ine pellets

of benzimidazoles were used in this study. However, since in

ceramics the polarization directions are only 90° apart, there

is a much larger resultant polarization on poling than in our

case where the dipoles flip 180°. We feared initially that

the effect would be too small to observe. This report describes

the experimental method and its application to a control compound,

thiourea, a known ferroelectric^*®, and to benzimidazoles with

unusual dielectric properties. Furthermore, the pyroelectric

effect and its voltage dependence will also be discussed.

EXPERIMENTAL

A. Material Preparations

2-Ethylbenzimidazole (2-EBIA), obtained from Pfaltz § Bauer

Co., was twice treated with activated charcoal in boiling distilled

water and recrystallized from the filtered solution. Sample

was dried thoroughly by a vacuum oven and then zone refined

for 140 passes. The ground powder was pressed at 120°C while

being maintained under vacuum. The resultant 5^" diameter pellets

were completely translucent but not optically clear. Gold film

was deposited on both sides of pellets at a vacuum of lO''' torr.

2 - Isopropylbenzimidazole (2-IBIA) was synthesized according

to Philips' method^. Pellets were prepared from zone refined

-167-
powder at 200®C and were also gold shadowed.

Thiourea, obtained from Baker Co., was recrystallized

twice from ethanol and was dried thoroughly. The well ground

powder was pressed under vacuum at R.T. and gold shadowed.

B. Apparatus

• The dielectric constants were determined- by calculating

the ratio, e' = C/Cq, where C is the capacitance of the cell

containing the material, exclusive of lead and insulator capacitance, and Co is the corresponding vacuum capacitance.

Capacitance was measured with a General Radio model 1620A bridge

with a grounded shield using a three terminal configuration,

shown in Figure 2.

The pellet was sandwiched between two brass electrodes

of if" diameter. The sample chamber was placed in a grounded

copper can which could be cooled or heated between -150° and

150*'C. An electrically insulated thermocouple was mounted into

the low-voltage electrode so that the temperature of the sample

could be measured accurately. A Heath/Schlumberger Series EU 200

assembly consisting of a strip-chart recorder, a potent iometric

amplifier, an offset module, and a four-channel signal multiplexer was used. The chart recorder was interfaced with a Keithley 610C electrometer, thus enabling us to record temperature and current or charge on the same strip chart. The

electrometer leakage was about 10"^'* amps, and did not interfere with the measurements. A Eurotherm series consisting of

a temperature controller and a programmer was employed to program

-168-
the sample temperature. Thus the pyroelectric effect could

be examined.

RESULTS AND DISCUSSION

a) Dielectric Measurements : Figure 3 shows the dielectric

constant of the thiourea pellet as a function of temperature

at the frequency of 1 KHz. (Thiourea is a known ferroelectric).

The two transitions are in fairly good agreement with the positions of the two conspicuous peaks reported by other workers'''^.

A narrow transition reported for the single crystal at -96°C'^

was not seen even though the experiment was carried out continually recording the temperature and capacitance. This might

be attributed to the clamping of constituent microcrys tals

.

The magnitude of the dielectric constant peak at the lower

temperature was reduced 100 times when compared with that of

the single crystal, which is understandable because the sample

is randomly oriented. The field thus encounters regions of

both high and low permittivity, which is equivalent to incorporating a large amount of low dielectric constant material

in the pellet. Therefore, the dielectric constant will be

reduced"^ »

^ ° » ^

^

. The magnitudes of the dielectric constant peaks

observed in the polycrystalline material were roughly proportional to the logarithm of the magnitudes of the corresponding

peaks in the single crystal. Thus, dielectric constant studies

on polycrystalline materials can reveal the presence of a transition at a specific temperature, but cannot determine the absolute

magnitude of the peak.

-169-
Figure 4a shows the dielectric constant of the pelletized

2-EBIA as a function of temperature and frequency. A distinct

maximum in the dielectric constant was observed at about -90°C.

An additional peak also occurred at about -60°C, above which

the dielectric constant decreased till the ice temperature.

Above 0°, the dielectric constants first increased slowly and

then rapidly. The variations seemed, to be the same for three

different frequencies. The loss tangent curve is shown in

Figure 4b, It clearly indicates that a sharp transition occurred at about -90°C for all three frequencies. At this transition point, the loss tangent at a given temperature was higher

for the higher frequency. This means that the maximum dielectric

absorption frequency is greater than 10 KHz. The loss tangent

did not show a clear transition at -60°C" and it started to rise

very rapidly above room temperature. The rapid increase of

loss tangent and dielectric constant at higher temperature was

attributed primarily to the onset of conductivity.

Temperature dependence of loss tangent and dielectric constant of 2-isopropyl benzimidazole is shown in Figure 5. There

is a slight rise in dielectric constant and a very broad "peak"

centering at -100°C. There is also a transition at about 105°C.

However, this maximum is small and cannot be reproduced when

lowering the temperature. The loss tangent rose slightly below

room temperature and reached a plateau. Above room temperature,

it rose very steeply and showed a peak for 10 KHz at the temperature corresponding to the peak of the dielectric constant

-170-
curve. This indicates that there probably is a transition at

105°C.

b) Polar i zat ion : In order to understand the nature of

dielectric transition of 2-EBIA at -90°C, pyroelectric measurements were performed on a 16.6 mil thick pellet. The experiment was carried out in the following way. D.C. voltage of

300 volts was applied to the pellet from R.T. to a temperature

far below the transition point (about -140°C). The current

flow in the electric field as a function of temperature is

shown in Figure 6. It can be seen that at -90°C, while the

sample was being cooled, a large peak of current (current maximum ^

10'-'-'^ amps) was observed. The current flow was negligible

(^ 10"12 amps) below the transition temperature. On removing

the electric field, the sample displayed a pyroelectric current

and, on heating through the transition temperature, current

discharge was observed, indicating that charge was stored in

the sample. The sample was repoled from -60 to -140°C at 100

volts and 200 volts and the flow of charge was measured by

heating the sample. A steady increase of charge in coulombs

continued until the sample discharged at the transition point.

Figure 7. This curve indicated that the charge flow rose as

the temperature approached the transition point. For comparison,

the pellet was cooled down without being subjected to a D.C.

electric field. The flow of charge when heating the sample

was negligible.

Thiourea (33.0 mils thick) was tested in the same way for

-171-
pyroelectric effects at temperatures below its ferroelectric

transition. The result is shown in Figure 8. It shows a close

similarity with 2-EBIA except that the total charge stored for

a given gradient was one quarter that of 2-EBIA.

The voltage dependence of total charge stored for a 16.6

mi,l thick pellet of 2-EBIA is shown in Figure 9. It clearly

indicates that charge stored reached saturation at about 50

volts (1,200 V/cm) but started increasing linearly with voltage

after 100 volts. This might indicate that properly aligned

crystals in- the sample were poled at 50V, but higher voltages

either injected charge, poled crystallites which were not

oriented in the direction of the field, or poled clamped crystallites.

Figure 10 shows discharge loops measured at a heating/ •

cooling rate of 4°C/min. for a 16.6 mil thick 2-EBIA pellet

poled at 200 volts. Cycle 1 represents the first thermal cycle

between -144°C and -108°C after the pellet was poled. The

charge flow rose when the pellet was heated. Upon cooling,

the current reversed sign but only a fraction of charge was

recovered. The recovered charge kept decreasing after each

cycle. This was probably because stored charge was partly lost

when the temperature approached the transition point. The loss

is probably not due to migration of injected charge as Fig. 6

shows the charge flow is minimal at -70 and -110°C. The maximum in charge flow at -90° should be due to polarization as

tan 6 is largest at 10 KHz. Overall, the curve shows quali-

-172-
tatively that the transition at -90*0 is pyroelectric in nature.

The nature of the two transitions in 2-IBIA was also investigated by the foregoing techniques . No charge storage was

observed when the pellet (20.2 mils thick) was poled through

-100°C at 100 volts. The same pellet was poled at lOOV through

105°C; it showed a change of slope in the current/temperature

plot at the transition temperature. The transition might be

swamped due to the high protonic conductivity of the sample

at the high temperature.

The poled pellet was coDled to room temperature before

removing the field. ^sTien it was heated from room temperature

to 78°C the charge flow rose to 8 >^ 10"-'-'-' coulombs. Upon cooling,

the charge was not recovered at all. In order to definitely

establish the nature of high temperature transition, single,

crystal studies are required.

CONCLUSIONS

Dielectric measurements on polycrys tall ine pellets were

demonstrated to be a useful method for screening organic compounds for anomalous transitions. As a consequence of this

screening, 2-EBIA has been found to be an interesting material.

Polycrystalline 2-EBIA behaves py ro e 1 ec tr ical ly when poled

above -9Q°C in a D.C. electric field and cooled below the transition point before removing the field. It was shown by Aharoni

that, upon substitution of deuterium for hydrogen participating

in hydrogen bonding, the dielectric transition was shifted from

-90°C to -60°C^. From the above and the frequency independent

-173-
dielectric peak, it can be concluded that 2-EBIA shows a first

order transition at about -90°C, probably ferroelectric, which

results from the dipoles coupled through hydrogen bonding.

The nature of the transition in 2-EBIA at about -60°C has

not yet been investigated. It will be undertaken in the near

future. Single crystal growth' of 2-EBIA is in process. Dielectric studies on single crystals will determine the accurate

transition temperatures; these are not well defined in the pellet

studies due to the clamping of constituent microcrystals and

the random orientation of dipoles.

The combination of techniques, dielectric scanning at different frequencies and pyroelectric measurements, which had been

used for ceramics is a powerful tool for establishing ferroelectric

behavior in polycrystalline materials generally. The scanning establishes whether dielectric anomalies exist and whether they

are probable first order transitions. The associated pyroelectric

behavior, reversible energy storage at the transition, establishes

probable f erroelectricity . We have demonstrated it for 2-ethyl

benzimidazole after validating it on thiourea, a known ferroelectric .

We thank the Office of Naval Research for partial support

of this work. Vie also thank the Research Corporation for support .

-174-
REFERENCES -

1. R. A. Sack, Austral. J. Sci, Res., A3_, 1 35 (1932).

2. R. J. Meakins , Trans. Far. Soc . , 3_S, 1953 (1962).

3. G. S. Reddy, R. T. Hobgood, Jr. and J. M. Goldstein, J.

.Am. Chen. Soc, SJ_, 336 (1962).

4. S. M. Aharoni, Ph.D. Thesis, Case Western Reserve Universitv,

, (1972) . _

5. G. Shirane, R. Xewnhar, and R. Pepinskv, Phvs . Rev., 96 , 581 (1954).

6. W. R. Cook, Jr., D. A. Berlincourt and F. J. Scholz, J.

Appl. Phys., 34, 1392 (1363).

7. G. J. Goldsmith and J. G. IsTiite, J. Che-. Phys., _n, 11"5 (1959).

8. K. Gesi, J. Phys. Soc. Japan, 2^, 1Q7 (1969).

9. M. A. Philips, J. Cher.. Soc, 2^, 2393 :i92S).

10. R. C. Miller and C. P. Sr^.yth, J. .Am. Chen. Soc, 21> (1957).

11. W. G. Cadv, "Pie:oelectricitv" , McGra'-c-Hill , X. Y. (1946).

pp. 167-16S.

-175-
FIGURE CAPTIONS

Fig. 1. Structure of 2-alkyl benzimidazoles and the nature

of the H-bonding reversal of dipoles.

Fig. 2. Three terminal configuration for capacitance measurements.

Fig. 3. Dielectric constant at 1 KHz of polycrys tal 1 ine thiourea.

Fig. 4a. Dielectric constant of polycrys tal 1 ine 2-ethyl benzimidazole at 100 Hz, 1 KHz and 10 KHz.

4 Fig. 4b. Tan 6 for polycrys tal 1 ine 2-ethyl benzimidazole at

100 Hz, 1 KHz and 10 KHz.

Fig. 5a. Dielectric constant of polycrys tal 1 ine 2-isopropyl

benzimidazole at 100 Hz, 1 KHz and 10 KHz.

Fig. 5b. Tan 6 of po lycrys tal 1 ine 2-isopropyl benzimidazole

at 100 Hz, 1 KHz and 10 KHz.

Fig. 6. Current flow as a function of temperature (ramping

rate of -4°C/min) for polycrystalline 2-ethyl benzimidazole. Poling voltage gradient was 7.1 x 10^ V/cm.

Fig. 7. Current discharge of polycrystalline 2-ethyl benzimidazole after poling at indicated voltage gradient.

Fig. 8. Current discharge of poled polycrystalline thiourea.

Poling voltage gradient was 3.6 x 10^ V/cm.

Fig. 9. Current discharge in poled polycrystalline 2-ethyl

benzimidazole as a function of poling voltage.

Fig. 10. Pyroelectric current flow in poled polycrystalline

2-ethyl benzimidazole as a function of temperature

and cycle.

-176-
-177-
HIGH

SHIELD

GROUND o -I r SAMPLE

r 6

LOW

-178-
-179-
-180-
-181-
-182-
5L

\1

1

25

IS

N

I

S

IS

3Z'

05-

5L00 I-

-5^

IS iSl IS IS

• • • J n N

01130 NUi 00 1

s G

-

•

-183-

!
-184-
-185-
-186-
-187-

ORIGINS OF PERSISTENT ELECTRICAL POLARIZAT ION

IN POLYMER SOLIDS

by

Stephen H. Carr

Department of Materials Science and Engineering

Northwestern University

Evanston, Illinois 60201

Introduction :

Sensor-related properties of electrically polarized polymers, such as pyroelectricity and piezoelectricity, depend on a particular combination

of physical properties and persistent electrical polarization, P, that

exists in these solids. Efforts to control pyroelectric and piezoelectric

performance of such materials will necessarily involve a search for the

best chemical compositions and physical microstructures that can be obtained with polymeric substances. Such explorations will produce materials

having ranges of thermal expansion coefficients and compressibilities, but

the property of these solids that can be varied by many factors of ten (and

thus be the most effective parameter in controlling sensor-related activity)

is magnitude of P. Thus, the research results described on the following

pages is concerned specifically with evaluation of the total magnitude of

P and with its relationship to physical microstructure

.

Persistent electrical polarization can arise from a variety of sources

(Slide 1). These individual contributions can have widely differing magnitudes, and they can be of opposing sign, such that the net polarization of

the specimen could be very small. Examples of "electrets" created by each

of these origins may be cited, but it is often regarded that poling with

contacting electrodes simply results in permanent dipoles adopting a preferred orientation parallel to the applied poling field. Measurement of

P is somewhat difficult to do unambiguously, but (if one makes some simple

assumptions) evaluation of thermally stimulated discharge (TSD) currents

can be one of the best techniques for determining P. An example of such

currents is seen in Slide 2, where a family of curves represents depolarization of polyacrylonitrile (PAN) polarized at 130 C and 5x10^ Vcm~ for

successively longer intervals of time. The peak maximum at 90 C corresponds

to a mechanically active relaxation process, so it is suggested that the

peak represents the loss of a part of P which was due to preferrentially

oriented dipolar side groups in PAN. Additional information from TSD peaks

can be obtained by plotting current vs. reciprocal temperature for the lowtemperature limb of a TSD peak, as is shown in Slide 3. The slope of this

plot indicates an activation energy for the depolarization process of 28

kcal/mole, which is a value somewhat consistent with the mechanism mentioned

above. A more complete treatment of a TSD peak is represented by the analytical expression given in Slide 4. Here, the adjustable parameter is E,

the activation energy, and successful fit of experimental data by this

relationship can be regarded as evidence that 1) the polarization was uniformly distributed throughout the specimen thickness and 2) the distribution

of relaxation times is fairly narrow. Slides 5 and 6 show the match of

-189-
Cowell and Woods plots to TSD data obtained 6n PAN poled as in Slide 2.

The values of activation energy used to get successful fit were 28 kcal/mole

for the 90 C peak and 45 kcal/mole for the 185°C peak.

Of special interest is the relative strengths of these two TSD peaks.

As shCT^ in Slide 7, the 185 C peak is two orders of magnitude larger than

the 90 C peak. It is also one order of magnitude larger than the normalized

polarization one would expect if saturation polarization (~16 ^,coul/cm ) had

resulted due to perfect alignment of nitrile sdde groups parallel to the

polarizing field. Thus, one can infer that this peak is caused by relaxation

of some source of polarization other than oriented nitrile groups. Interestingly, changing molecular organization by the simple process of uniaxial

elongation prior to poling has a marked effect on the magnitude of both peaks,

indicating that both contributions to P are structure-sensitive (and therefore that changes in microstructure can be effective ways to control P in

PAN). Somewhat direct evidence for. dipolar orientation in poled PAN is seen

in birefringence data (Slide 8). It is observed that a higher value of

birefringence exists in stretched films after they have been polarized. This

is consistent with the concept that birefringence in oriented PAN is dominated by the nitrile groups and that they had undergone rotation into

directions parallel with the polarizing field during poling. Also of interest is the information obtained by X-ray scattering (Slide 9). These

data suggest that an initial "as-cast" film (a) is somewhat glasslike in

its scattering properties but that it will become (b) fairly ordered (onedimensional order that lacks crystallographic registry in directions normal

to chain backbones) as a result of annealing. However, the X-ray diffractometer scan shown in (c) , which is from a film that had been annealed and

polarized, differs very little (if at all) from that shown in (b) . Thus

it might be either 1) that the extent to which there is a preferred orientation of nitrile groups in polarized PAN is insufficient to create a change

in molecular organization or 2) that the same molecular organization exists

in films before and after being polarized. A plausible way to reconcile

both the birefringence data and the X-ray data is to postulate the existence

of ferroelectric-like domains.

Other evidence for interactions between internal field and molecular

organizations is seen in Slide 10. Here, the effects of poling at increasingly stronger electrical fields is seen not to lead simply to progressively

higher amounts of polarization (as judged by area under TSD peaks) . Specifically, it can be seen that itnpolarized PAN actually possesses some degree

of polarization. Also, fields above 5x10^ Vcm"''" produce a polarization

which has a component that is of opposite polarity to the strong peaks

whose maxima are in excess of 200 C. The exact explanation for these observations is not available at this point, but they suggest that poling may

be affecting (independently or dependently) several contributions to the

total value of P. It is possible that particular organizations of nitrile

groups, such as are listed in Slide 11, are playing an important role in

causing these effects and the final result may be nonuniformities of

polarization in thickness and in lateral directions of these "electret"

solids.

-190-
Slide 1. Possible Origins of Persistent Electrical Polarization

Homocharge

Injected species from surroundings

Chemical moities created from mobile ions.

Heterocharge

Oriented Dipoles

Composition gradient of ions

Slide 2.

DEPOLARIZATION TEMPERATURE, °C

-191--
Slide 4.

THERMALLY STIMULATED DISCHARGE

(TSD)

-2

i = A exp[-t - B exp[-t1 t

]

i = discharge current

A,B = constants

t = E/kT

E = activation energy for depolarization process

T.A.T. Cowell and J. Woods, Brit. J. Appl. Phys.

18, 1045 (1967).

-192-
. 40 60 80

TEMPERATURE

100

(»C) 120 140

Slide 6.

Slide 7.

120 140 160 180 TEMPERATURE Cc)

NORMALIZED POLARIZATIONS

OF POLYACRYLONITRILE FILMS

200 220

TSD Peak

90°C

Sample

Not stretched

Stretched

P/e E2_EL

120

210

185°C Not stretched

Stretched

13,000

51,000

-1°3~
Slide 8.

A BIREFRINGENCE ANALYSIS OF PAN

Stretched PAN:

Birefringence normal to draw: -0.005

Birefringence parallel to draw: -0.002

(Optic axis parallel to

transverse direction)

Stretched-polarized PAN:

Birefringence normal to draw: -0.009

Birefringence parallel to draw: -0.001

(Optic axis 25° from

transverse direction)

Slide 9.

15 20 25 30

2 6 (degrees:

-194-
90 10 130 150 170 190

TEMPERATURE,

Slide 11.

Associations of Nitrile Side Groups

Intramolecular Structure:

Causes chains to adopt a compact

(aperiodic) conformation

Intermolecular Structure:

Leads to fluctuations in local order.

May organize into ferroelectric-like

domains.

Nonuniformities of Polarization in Dielectric Foils

Local variations of physical microstructure

,

Local variations of chemical microstructure.

Lateral nonuniformity of poling field.

Variation of polarization structure

in thickness direction.

-195-
"Pyroelectric Polymers Applied to Optical Radiation Measurements"

Robert J. Phelan, Jr.

Electromagnetics Division, National Bureau of Standards

Boulder, Colorado 80302

Presented at the Symposium-Workshop on Piezoelectric Polymers, NBS,

Washington, D. C, April 15-16, 1975 .

We have demonstrated the potential of using pyroelectrics

in optical radiation measurement instruments that can be

electrically calibrated, fast, sensitive, accurate, and usable

over an extremely wide spectral range. In the development of

these instruments, we have made considerable use of polymer

pyroelectrics. To optimize and prove the systems has required

extensive characterizations of the pyroelectric polymers.

We have focused on structures compatible with optical measurements; properties and characteristics irrelevant to optimizing

detectors have been of secondary concern. In the presentation we

describe our m.easurements , indicate accuracies obtained, give the

parameters we have considered significant, and give what we believe

to be the state-of-the-art in applying pyroelectric polymers to

optical radiation measurements. The presentation is a summary of

a number of publications referenced at the end of this paper.

The principal points of the presentation are: (1) to compare

the polymers PVF and PVF2 with other pyroelectrics and show that

they have favorable properties; (2) to describe and demonstrate

the limits of spatial uniformities and their significances (extremely

uniform response detectors can be made) ; (3) the responsivities of

196
the detectors depend both on material properties and device fabrication details; (4) modulation frequency responses can be predicated in detail - except at high frequencies where we have

* some loose ends; (5) the D can be made equal to the highest

available from any pyroelectric ; (6) studies of long term

stabilities have shown that the drift in responsivity can be

less than 5% per year, but by adding electrical calibration

detector drift problems are eliminated; (7) the detectors are

useful at temperatures down to 77°K and probably even lower;

(8) spectral responses can be made flat to better than 2%

over the wavelength range of 1 to 12 ym; and (9) complete systems

based on a polymer pyroelectric are being evaluated as a new

standard for optical radiation measurements.

- 197 -
References to available published reports from the "Optical

Electronics Program" at NBS, Boulder - relevant to the application

of PVF and PVF^ pyroelectrics.

1. "High D Pyroelectric Polyvinylfluoride Detectors", Robert J.

Phelan, Jr., Robert J. Mahler, and Alan R. Cook, Applied Physics

Letters 1£, 337 (1971). .

.

2. "Absolute, Pyroelectric Radiometers and Two Dimensional Arrays",

" R. J. Phelan, Jr., R. L. Peterson, G. P. Klein, C. A. Hamilton,

and G. 17. Day, Proceedings of the Technical Program, ElectroOptical Systems Design Conference, Sept. 18-20, 1973, New York,

New York. . ^. . . v. • .

3. "Electrically Calibrated Pyroelectric Optica]..--Radiation Detector

Robert J. Phelan, Jr., and A. R. Cook, Applied Optics 12, 2-194

(1973) . . . . ;.; : .~"

4. . "The Polarization of PVF and PVF2 Pyroelectrics", R. ' J . Phe] an

,

R. L. Peterson, C. A. Hamilton, and G. W. Day, Ferroelectrics 7,

375 (1974). ... ;.v. ^/./^^^ "

.

5. - "Effects of poling conditions on responsivity and. uni f ormi ty . o

f

pol?.rization of PVF^ pyroelectric detectors". G. I'/, Day, C. A.

Hamilton, R. L. Peterson, R. J. Phelan, Jr., and L. 0. Mullen,

Applied Physics Letters 2£, 456 (1974). ... .

'

. •. v..

6. "Analysis of response of pyroelectric optical detectors", . R. L. Peterson, G. V/. Day, P. M. Gruzensky^ and R. J. Phelan, Jr

Journal of Applied Physics 4_5 , 3296 (1974).' . .1'-;

:

7. "A Pyroelectric Power Meter for the Measurement of Low Level

Laser Radiation", C. A. Hamilton and G. W.- Day, NBS Technical

Note #665, Feb. (1975). . : - ^^^^^

8. "Electrically Calibrated Detectors versus. Black Body Sources",

R. J. Phelan, Jr., Proceedings of the IRIS "Specialty Group on

Infrared Deteccors, Mar. (1975). ". : . ^...-^ [:

'

9. "A Flat Spectral Response Detector for the Visible to 12 yiii

Region", G. W . Day, C. A. Hamilton, and K. W. Pyatt, Proceedings

of the IRIS Specialty Group on Infrared Detectors, Mar. (1975).

IC. "An Absolute Pyroelectric Radiometer (A Reference for Detector

Responsivi-cics)" , C. A. Hamilton and G. W. Day, Proceedings of

the IRIS Specialty Group on Infrared Detectors, Mar. (1975).

198 -
Pyroelectricity in PVF^

R. E. Salomon and M. M. Labes

Department of Chemistry, Temple University, Philadelphia, Penna. 19122

We have underway a combined experimental and theoretical program

aimed at understanding the mechanism of pyroelectric phenomena in electrically

poled PVF2 in order to identify the key parameters which can be altered so

as to enhance the pyroelectric coefficient of this material.

Our experimental work has already led to the obser\'ation that a

copolymer of vinylidene fluoride (95%) and tetraf luoroethylene (5%) has a significantly higher pyroelectric coefficient than the poly (vinylidene fluoride)

homopolymer . ^ We furthermore find that the temperature dependence of the

pyroelectric coefficient of the copolymer is significantly greater than that

of the homopolymer. The effect of poling time and the nature of the first

cycle thermally stimulated currents are also compared. We find that in poling,

saturation times are significantly greater in the copolymer.

The theoretical modelling, which is the main subject to be discussed

here, has been primarily directed towards an understanding of the poling process.

The experimental findings to date have not really helped to distinguish between

the various poling mechanisms. It has seemed to us, from both the limited

structural information available as well as a general consideration of seraicrystalline polymers, that the rotation of polar monomers has two serious

restrictions which must be dealt with in any adequate theory. The first is

concerned with interactions within a chain. Except for small torsions, the

rotation of one monomer is strongly coupled to the next and so on. Hence, it

would seem that appreciable polarization would require that entire chains or

large segments of chains flip over. The second restriction concerns the inter-

-199-
actions between the chains, particularly within the polar crystalline phase.

It is clear that the chains can not reorient in a random manner.

Accordingly, we have developed four models which reflect these considerations. These models are referred to as the Rigid Rod Model, the Perimeter

of Separation Model, the Molecular Rearrangement Model and the Space Charge

Model. In the Rigid Rod Model, we envision polar cylinders whose axis correspond

to the chain axis. The cylinders are caused to rotate in the poling field.

The temperature is elevated so az to lower the viscosity sufficiently. Reversible

pyroelectric phenomena can be described by this, and the next three models, by

Introducing a Lorentz type of internal electric field. The Perimeter of Separation Model is a domain type of theory. We have recently been able to calculate

the domain shape and size for a simple assumed crystal morphology. The Molecular Rearrangement Model is based on a theory of field induced changes in

polymer conformation, namely, going from helical to planar zig-zag in the electric field. This model, if correct, would predict some changes in the diffuse

x-ray scattering which thus far have not been observed. The Space Charge Model

is based, to some extent, on the reported effects of blocking electrodes on

pyroelectricity . In this model, the poling process simply leads to a space

charge, while the pyroelectric current is a result of the temperature dependence

of the piezoelectric strain coefficient and the space charge distribution.

The first model has been used to predict the form of the kinetics

of poling and depoling, while the others are essentially equilibrium models.

The Space Charge Model treats the piezoelectric properties phenomenologically

and thereby suffers as a basic model. These treatments are currently being

refined so that they may be tested in a definitive manner against experimental

results.

-200-
Figure Captions

Figure 1. Measurements of pyroelectrlc current 1^ from the decay of the

total current after the heating cycle was stopped at t = 0.

Figure 2. Pyroelectrlc coefficients for PVF^ (• ) and VF^-TFE copolymer (O

)

at poling fields of A x 10^ ( ) and 6 x 10^ ( — ) V/cm.

T «= 80°C, t = 5 h.

P P

Figure 3. Rigid Rod Model. Space between rods considered as containing an

amorphous, viscous, apolar material.

Figure 4. Calculated polarization rate for Rigid Rod Model during poling

versus reduced time.

Figure 5. Perimeter of Separation Model.

Figure 6. Molecular Rearrangement Model.

Figure 7. Space Charge Model.

-201-
Current

-202-

RIGID ROD MODEL
-205-
Perimeter of Separation Model

"6=specific

perimeter

energy

Helmlioltz Free Energy ,

-206-
S = Helical

] and

J, = Planar

• T

> < s

s s s s s s

s s s s s s

s s s s s s

s s i s s s

s s i s- s s

t1 s T s s i

t1 t1 Iw s I I

t1 t1 i t1 Iw t1 1 iV t1 i iV

t t 1 t i i

t t 1 t i i

t 1 t

Zig- zag

s s s s s s

s s i s s s

s s i s s s

t s s s s

t s i s s

t s t s s

t s I t s i

I t

s i t t

t t t

t i t t

t t t

-207-

I
SPACE CHARGE MODEL

t

t4-

-J- :

-»• +

+4-4

+

t4-

t

t

(D

a^v PCX) ~ En

Xk= length of sample ander -field Ek

A.o=leTigth of sample itnder field E=o

For Pn=-QTiX v/ith Lf dX =

^-rot.

i = Pyroelec. Uurrent

:2

-208-
PRELIMINARY INVESTIGATIONS OF PIEZOELECTRIC POLYMERS

FOR SONAR APPLICATIONS

James M. Powers

Naval Underwater Systems Center

New London, Conn. 06320

ABSTRA.CT .

Two Navy sonar applications in which piezoelectric polymer might

be used as hydrophone elements are for small diameter towed arrays and

large aperture hull mounted arrays. In these applications polymer

hydrophones would offer the following advantages over piezoceramic;

availability of large dimensions allowing noise reduction by spatial

averaging, low density for increased buoyancy, ability to withstand

high pressure, and low cost. To date we have prepared various polymer

samples, set up measurement systems to evaluate piezoelectric g

coefficients, and have measured samples supplied by Dr. Edelman of

NBS which exhibit g, an order of magnitude higher than ceramic.

-209-
PIEZOELECTRIC POLYMER MEASURING INSTRUMENTS

S. Edelman

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

ABSTRACT

*Piezoelectric polymer sensors have advantages over conventional instruments for measuring dynamic mechanical quantities in special cases because

of their low density, flexibility, resistance to damage by mechanical shock,

thinness, and low cost. A number of examples are given and a typical polymer

sensor is described.

-210-
Piezoelectric Polymer Measuring Instruments

The polymer transducers group uses favorable properties of polymers to

advance the science of measurement. Considered as piezoelectric materials,

.polymers are less active than the ceramics which are used in conventional

measuring instruments. However, polymers have other properties which make

their use advantageous in particular kinds of measurements. There are some

cases of considerable importance where polymer instruments are uniquely

suited. Usually the advantages of polymer instruments show up in dynamic

measurements. A typical case is measuring the level of vibration at a point

on a thin metal sheet. Any point on the metal sheet to which an accelerometer is attached will have significantly greater surface density than the rest

of the sheet. The vibration pattern of the sheet will rearrange itself so that

the motion at that point is minimized and the measurement is unrepresentative.

On the other hand, cementing a small piece of polymer film to the metal sheet

need not change the surface density significantly and a meaningful measurement

can be made. The amplitude of vibration and the variation of vibration level

with frequency measured by the polymer are representative of what the level at

the point would be with no instrument attached. Conversely, both the vibration

amplitude and the spectrum of resonances measured by an accelerometer would be

affected by the presence of the instrument. Thus, it is practicable to distribute a number of polymer vibration gages over the surface of a panel and to

deduce from their reading the mode of vibration of the panel under various

conditions. This kind of study can not be performed effectively with conventional vibration measuring instruments. Polymers have somewhat similar advantages for studying the noise signatures of bearings, gears and transmission

-211-

T
systems. The mass of a conventional instrument Introduces a set of resonances

which confuse the noise spectrum being studied and which may hide the changes

of the spectrum which indicate the first signs of deterioration. Also, the size

of conventional instruments and their need for a threaded mounting hole may

dictate their location at some distance from the origin of the noise. A polymer

gage, almost always, can be mounted directly to the noise source with a much

better chance of picking up an uncluttered spectrum. Polymer gages are handy

for acoustic emission studies because, of the ease with which they can be bonded

to metal and because their thinness and high internal damping make their response

fairly flat with frequency up to several megahertz.

As strain gages, polymer gages have about two-thirds of the sensitivity of

semi-conductor gages but they can be used without the need for bias voltage or

bridge balancing. Also, they are not brittle and can be used where mechanical

shock occurs and on curved surfaces.

With the addition of lead foil as a seismic mass, polymer gages can be

used as flexible accelerometers with an unusually wide frequency range.

The fact that polymer g^ges can provide good coupling to soil can be used

as the basis for another group of important applications. At present, seismometers are used to study stress waves in the ground. Seismometers are large,

expensive instruments and usually they are made larger and more expensive by

being enclosed in a body of grout to provide better coupling to the soil. A

thin sheet of polymer can be installed much more easily and it provides good

coupling without any special effort. A figure of merit used by those who work

with stress waves in the ground is proportional to the ratio of the diameter

of the gage to its thickness. A polymer gage can be so thin that even a relatively small gage can have a good figure of merit. In some cases polymer gages

benefit from encapsulating to minimize unwanted signals due to bending or sliear

-212-
effects but, even so, the polymer gage is cheap, easily installed, and well

coupled compared with seismic instruments. We have been doing some work in

developing such gages in cooperation with the Air Force Weapons laboratory and

Bolt Beranek and Newman and also in connection with some work we are doing

for the Federal Highway Administration. We have discussed the use of similar

stress gages for earthquake studies and for monitoring the condition of earth

dams. 'Another possible application i§ in underground geophysical exploration.

The thinness and flexibility of polymer gages make them feel and act

mechanically very much like skin. This characteristic is used in automobile

crash studies. A pattern of polymer stress gages are to be put on the head

and chest of anthropomorphic dummies to detect the areas of contact and the time

history of the impact. Since the gages behave very much like skin, they do not

disturb the anthropomorphic behavior of the dummy and, since they are flexible,

they are not likely to be damaged in the crash. Conventional pressure gages

form concentrated masses whose behavior during a crash would be much different

from the behavior of the rest of the skin of the dummy and they are likely to

be damaged during a crash. Their impact is also likely to damage parts of the

automobile in ways that are not characteristic of real crashes.

Polymer gages are being considered to initiate air bag inflation and to

measure pressure and temperature pulses in tires.

The similarity of polymer gages to body tissues has other potential uses.

We are developing gages to be inserted into monkey brains to determine how much

of an impact to the outside of the head is felt as a pressure pulse in the brain.

The polymer gage is to be used here because its density is close to that of the

brain material. If a conventional pressure gage were used, Inertial effects

would produce more signal than the pressure change.

• -213-
Polymer gages can be applied like bandaids to monitor heart sounds and

pulse rates of patients during exercise. Conventional instruments can be

used. Their size and mass make the patient conscious of their presence and

his behavior is not entirely normal. He is much more likely to forget polymer

gages.

A different kind of application depends on the fact that poling a polymer

stores about one and a half times as much energy per unit volume as is stored

in piezoelectric ceramic under similar conditions. One example we have worked

on is development of detonators for ordnance where polymer sheet is preferable

to ceramic because it is lighter for a given energy storage, can be connected .

to provide a desired electrical output impedance more readily, and can be used

to line enough of a shell to provide detonation for any angle of incidence.

An obstacle in the use of polymers for this kind of application is the fact that

it is more difficult to release the energy stored in the polymer than that

stored in the ceramic because the ceramic has a greater electromechanical coupling

constant. The situation is analogous to comparing paper and wood as fuel. Paper

is easier to light but wood gives more energy per unit volume. In the case of

artillery there is plenty of energy available to release the stored energy.

We have considered a number of other applications where this v/ould be true; for

example, powering some prosthetic devices.

We have not done any work on devices to generate and detect sound in air

except for our own amusement. Dr. Fukada demonstrated a polymer microphone

several years ago and the Plessey Company, in England, has demonstrated a polymer

microphone. Recently, Pioneer Corp. has announced commercially available microphones, loudspeakers, and head sets.

-214-
We have worked on sound measurement under water. Pol>Tner hydrophones have

several obvious advantages over the conventional piezoelectric ceramics such

as a good acoustic impedance match to sea water, flexibility which makes them

capable of being unreeled behind a ship and rereeled after use, low cost, ease

'of fabrication and manipulation, and others. These advantages promise to make

polymer hydrophones useful for fish finding, fatho-etry, geophysical exploration

,

and military applications. Single sheets of PVF- are considerably less sensitive

than conventional hydrophones. T-Tnen we started working on hydrophones we were

advised that a sensitivity of -208 dB relative to lY/uPa would be needed before

polymer hydrophones could be considered seriously. Our attempts to reach this

level make up a kind of year-long detective story whose hero is Aime de Reggi.

When we started, it seemed clear that the way to increase sensitivity while

preserving the advantages of polymer hydrophones was to form a stack of sheets

connected in series electrically and in parallel -.echanlcally . IsTien we did this,

the sensitivity of different stacks varied greatly. Some stacks were no m.ore

active than single sheets. After much investigation, Dr. de Reggi found that the

output of the separate sheets in a stack differed in phase and that the outpuc of

a stack was the vectorial sum of the outputs of the separate sheets rather

than the arithmetic sum. Considerable further investigation showed that the

reason for the phase difference was the presence of microscopic bubbles at the

interface between sheets. The presence of only a few bubbles was enough to change

the output phase drastically, showing that the sensitivity of the small area of

material in the walls of the bubbles had output comparable with the output of the

much greater area where there were no bubbles. Dr. de Reggi deduced that the

greater activity of the material in the walls of bubbles arose because that

material was acting as a membrane and was bcinc strained by the sound pressure

against a resisting area consisting of the thickness of the pol>Tiier film multiplied by the diameter of the bubble while the portion of the surface which was

-215-
free of bubbles was subjected to a compressive stress which was resisted by the

whole surface area. This development suggested two courses for further progress.

One was to avoid bubbles in making stacks. That work is proceeding. The other

was to take advantage of the greater sensitivity supplied by using the membrane

inode of operation. A sensor was devised which was essentially a big bubble. It

consists of a pad of compressed foam rubber between two sheets of piezoelectric

polymei^. In this way, the polymer is 'held under tension by the foam rubber and

the sound pressure modulates the tension. The output of this type of gage is

about 30 dB greater than the output of single sheets of polymer subjected to

compressive stress so far as our preliminary measurements show, and the output'

Is reasonably flat with frequency to about 500 Hz. The sensitivity is about

10 dB better than our goal of -208 dB re IV/viPa. Much more work needs to be

done to allow us to understand the behavior of this gage completely and to

optimize the various design parameters. However, in its present preliminary

form it is within the range of conventional hydrophones in sensitivity while

providing the advantages of low density, flexibility, good impedance match with

water, ability to withstand shock, and low cost.

A typical sensor consists of a sandwich of two thin polymer sheets. Each

sheet has evaporated metal electrodes on both faces and the sheets are fused

together so that charges of the same polarity appear on the inner faces. The

center conductor of a coaxial cable is connected to the electrodes on these inner

faces and the shield of the cable is connected to the electrodes on the outer faces.

In- this way all exposed surfaces are at ground potential and the signal inside the

sensor is well shielded. Usually, the active area of the gage is connected to

the coaxial cable by a long thin strip of metal evaporated onto an inactive

-216-
portion of the same polymer sheet that contains the active part of the gage.

In this way, the masses of the connectors and cables are kept from affecting the

measurement. The active portion of the gage can be cut to any reasonable size

and shape needed for the measurement.

'217-
NBS-IMA (REV. 7-73)

U .S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO. NBSIR 75-760

2. Gov't Accession

No.

3. Recipient's Accession No.

4, TITLE AND SUBTITLE

Proceedings of Piezoelectric and Pyroelectric - Symposivnn-Workshop

5. Publication Date September 1975

6. Performing Organization Code

7.AUTH0R(S) Broadhurst 8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No.

11. Contract/Grant No.

12. Sponsoring Organization Name and Complete Address (Street, City, State, ZIP)

Office of Naval Research Army Research Office

Arlington, Va. 22044 Durham, North Carolina 27706

13. Type of Report & Period Covered

Final

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less factual summary o{ most si^ilicant information. If document includes a significant

bibliography or literature survey, mention it here.) The purpose of this Symposium-Workshop was to bring together scientists and

program managers from federal (primarily DoD) , industrial, university, and foreign

laboratories who are involved in the research on and manufacture of piezoelectric

and pyroelectric polymer materials and devices, in order to exchange information

about and stimulate further work on the new and rapidly developing measurement

technology involving polymer transducers. This meeting is particularly important

at this time because highly active polymer films are not yet available to device

manufacturers in the U.S. or to potential users such as DoD.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper

,,ame; separated by semicolons) Electrets; hydrophone; plezoelectr ic ; polymers;

polyvinylidene fluoride; pyroelectric; radiometer; transducers

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