NBSIR 83-2682

Molecular Processes Affecting

Mechanical Failure of Polymeric

Materials

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Center for Materials Science

Polymer Science and Standards Division

Washington, DC 20234

Annual Report Covering the Period

August 1, 1982 - December 31, 1932

April 1 983

Prepared for Materials Science Division

Defense Sciences Office

Defense Advanced Research Project Agency

1400 Wilson Blvd.

Arlington, Virginia 22209

NBSIR 83-2682

MOLECULAR PROCESSES AFFECTING

MECHANICAL FAILURE OF POLYMERIC

MATERIALS

Bruno M. Fanconi

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Center for Materials Science

Polymer Science and Standards Division

Washington, DC 20234

Annual Report Covering the Period

August 1, 1982 - December 31, 1982

April 1 983

Prepared for Materials Science Division

Defense Sciences Office

Defense Advanced Research Project Agency

1 400 Wilson Blvd.

Arlington, Virginia 22209

U.S. DEPARTMENT OF COMMERCE, Malcolm Baldrige, Secretary

NATIONAL BUREAU OF STANDARDS, Ernest Ambler. Director

Molecular Processes Affecting Mechanical Failure of Polymeric Materials

SUMMARY

Preliminary observations have established correlations between the

buildup in the concentration of molecular end groups and macroscopic failure

in polymers. Current work has examined the effects of stress history,

molecular weight and the dependence of the growth in end group concentrations

with time during the course of a mechanical degradation experiment. In cyclic

fatigue under uniaxial loading conditions the concentration of end groups was

found to exhibit a sigmoidal dependence on the number of cycles. After an

initial rapid rise to about 60 % of its value at failure the number of chain

scissions was observed to increase linearly until sample failure. Such a

relationship holds promise as the basis of methods to predict mechanical

failure. The effects of frequency, peak stress, and molecular weight have

also been examined and are reported.

INTRODUCTION

Polymeric materials are attractive for use in military applications due

to their low density and cost as well as unique properties such as resistance

to corrosive environments. However, the use of these materials has grown so

rapidly in new applications that frequently their full exploitation, particularly

in load bearing applications, is limited by inadequate knowledge and test methods

relevant to long term performance in the use environment. Furthermore,

techniques for nondestructive evaluation of polymeric materials have not been

developed to the extent found for metals and ceramics because the processes

affecting durability and performance are poorly understood. An improved
2

understanding of molecular processes controlling the deterioration of

mechanical properties would provide the basis for development of predictive

models of long term performance, of methods to assess the remaining useful

service life (nondestructive evaluation) , and the development of new polymeric

materials with improved long term performance.

The deterioration of mechanical properties has as its origin molecular

processes affecting the chemical bonding arrangements. These processes may

include disruption of hydrogen bonds and the van der Waals-type forces

between polymer molecules, and scission of valence bonds that link the atoms

together to form the polymer molecules. Catastrophic failure involves

scission of the carbon-carbon bonds forming the polymer backbone. This process

is thought to occur at the earlier stages of mechanical degradation but its

importance has not been fully evaluated due to conflicting experimental data.

In the model advanced by Zhurkov and co-workers (1) the backbone carbon-carbon

bond scission (chain scission) is regarded as the precursor to microvoid

formation from which cracks initiated leading to macroscopic failure. A critical

aspect of this model, and one that needs verification, is that the free radicals

produced as a result of the primary chain scission in turn initiate many

additional chain scissions before recombinations form stable chemical species.

The extent to which additional chain scission results from attack by free

radicals is important to the design of polymeric materials with improved stability

The origins of molecular processes associated with mechanical degradation

and fracture derive from the structures exhibited by polymers in the solid state.

Two basic morphologies exhibited by crystallizable polymers are the spherulitic

structure obtained by crystallization from the melt and the fiber structure associ

ated with high degrees of chain orientation. The first is prevelant in crystallizable polymers prepared in such a way that there is little or no preferred chain
- 3 -

orientation, and the second morphological type is formed by mechanical

deformation either during or subsequent to solidification from the liquid

state, or from solution. In Fig. la is shown an example of the underlying lamellar

texture in which a single polymer molecule, represented by the meandering

solid line, has both crystalline and disordered regions. The crystalline

regions are confined to thin lamellae of the order of tens of nanometers

in the chain direction, and since the molecular length is much larger, of

3 the order of 10 nm, it is apparent that one molecule must fold back and forth

in one lamella or span several lamellae.

Small angle neutron scattering experiments have confirmed the essential

features of the model shown in Fig. la for semicrystalline polymers such

as polyethylene. The lamellae are further organized into spherulitic

structures such as shown in Fig. lb in which the lamellae radiate outward

with the chain axis in the lamellae oriented tangentially to the radial

direction. The loops and entanglements formed in the disordered interlaminar

region are illustrated in Fig. lb. An important morphological feature is

the chain segments that bridge lamellae. The interlaminar region offers the

path of least resistance to fracture, and as the spherulitic structure is

deformed the predominant deformation on the local scale occurs in the

interlaminar regions. In this manner the molecular segments in the

disordered regions bear disproportionately large stresses that deform the

molecular geometry and weaken some of the valence bonds between the backbone

carbon atoms of the polymer molecule. This forms the incipient point for bond

scission.

The technological importance of chain orientation is that materials

with high tensile properties. Young's moduli and tensile strengths, exhibit

high chain orientation and fibrillar texture. The polyaramids used as fiber

reinforcements in advanced composites are an example of a highly oriented polymer.
Figure la - Lamellar texture of semicrystalline polymers.
5

Entangled

interSamellar

links

Branch

points

Figure lb - Spherulitic structure illustrating interlaminar links and entanglements.
- 6 -

The fiber structure of a highly oriented polymer is shown in Fig. 2.

In this case, lamellar texture is still evident for crystallizable polymers

and an essential feature is again the molecular segments that bridge the

lamellae. The molecules comprising the interlamellar links are nearly extended

so that subsequent deformation distorts the carbon-carbon backbond bonds

leading to chain scissions.

At the molecular level the effect of the applied mechanical stress is

to distort the equilibrium configuration, and in the case of valence bonds

change the electronic distributions. The effect on the potential energy

surface of the carbon-carbon bond is shown schematically in Fig. 3 where the

potential energy is plotted against the interatomic distance. The states

A and B refer to bonded and nonbonded states, respectively. In the absence of

applied stress, top diagram, the barrier energy G*Ag

which must be overcome

to cause chain scission is much larger than thermal fluctuations and the

probability of a chain scission event is negligible. In the bottom diagram

of Fig. 3 is depicted the situation when macroscopic stress is applied. In

the presence of an applied stress the potential energy surface is distorted,

moving the equilibrium bond distance to larger values and greatly reducing

the barrier energy G*^. The result is that G*^g is much smaller which

facilitates thermally activated bond scission.

As the potential energy surfaces of the vibrating atoms are determined

by the electronic distributions the vibrational frequencies of modes involving

carbon-carbon bond stretch should be affected by the applied mechanical stress.

For some polymers investigated, most notably polypropylene, it has been observed

that vibrational frequencies associated with motions of the backbone carbon

atoms vary linearly with applied mechanical stress (2,3). In the kinetic theory

of fracture the rate of bond rupture is given by

v = NkkT/h) exp (-G*^g/kT) exp (sgc/kT) ( 1 )
7

I

\

j

I

I

*k

1

)i

If

»

I

i

i

I

I

t

I

I

i

I

1

I

J

J

1

Figure 2 - Fiber morphology of uniaxially oriented polymers.
8

Kinetic Theory of Damage

Bond fracture as a consequence of thermal fluctuations,

0 stress i

Bond loading

a stress

T

'AB

i

t— Gab

Figure 3 - Potential energ3r surfaces of the carbon-carbon bond: top curve,

no applied stress; bottom curve, applied stress, c.
- 9

where N is the number of bonds, < is the transmission constant, k is the

Boltzmann constant, T is the temperature, h is Plank’s constant, G*^-g is

the energy for activation of bond scission in the absence of applied stress,

o is the applied stress, and s$ is the activation volume. If a critical

number of chain ruptures, NQ

, leads to macroscopic failure then the time to

failure, t ^ , can be identified as

According to equations (2) and (3) there is a logarithmic relationship

between time-to-failure and applied stress.

The logarithmic relationship between time-to-failure and applied

stress has been observed to hold over limited ranges of applied stress.

For example, time-to-failure data as a function of applied tensile load

for a linear polyethylene is shown in Fig. 4. The two linear regions

are indicative of two modes of macroscopic failure: at low stresses

brittle fracture dominates, whereas at higher stress failure is due to

inhomogeneous deformation associated with neck formation and propagation.

The downturn of the curves at the highest loadings can be attributed to a

heating effect such that the average temperature of the specimens is higher

than the ambient value. Increasing the temperature lowers the lifetime

and hence the curves shift towards shorter lifetimes.

A useful representation of the stress-strain-time behavior of polymers

is illustrated in Fig. 5 where the data are plotted for a linear

polyethylene. The boundary of the fracture envelop gives the stress-straintime at failure (physical separation). Below the a-line, brittle fracture

dominates, between a and 3, failure is by inhomogeneous deformation, and

( 2 )

where

K

AB

= K ^ exp ( “G*AB + s ^ a )/kT (3)
TIME

TO

FAIL

(s)

5 10 15 20 25 30 3

STRESS (MPa)

'Figure 4 - Time-to-failure data for a linear polyethylene in uniaxial load

condition as a function of applied stress and temperature.

CJl
ELONGATION

- 11

STRESS (MPa)

^Figure 5 - Fracture envelope of a linear polyethylene. a region, failure by

brittle fracture; 3 region, inhomogeneous deformation; y region,

brittle fracture of elongated specimens.
-12 -

between 6 and y, the elongated specimens again fail in a brittle-like

fashion. Fracture envelopes serve as a useful means of characterizing

macroscopic failure, and are useful in comparing results from analysis

of molecular level events. The amount of chain scission associated with

mechanical degradation may differ depending on the type of macroscopic

failure mode operative. Furthermore, the manner by which the number of

chain scissions accumulate to failure could vary.

Applied stresses that exhibit a periodic time dependence, such as cyclic

fatigue, are another type of loading history relevant to in-use conditions,

and for which laboratory studies have been conducted. Time-to-f ailure data on polyethylene (4) and poly (methymethacrylate) (5) subjected to stresses varying

sinusoidally in time have indicated that neither cycle dependent nor

accumulative damage failure are operative. In Fig. 6ab are compared

the failure data for ultra high molecular weight polyethylene subjected to

sinusoidal tensile loading. The dashed curve in Fig. 6a gives the

predicted failure behavior with frequency if failure occurred at a critical

number of cycles. Clearly, this does not happen. In Fig. 6b are compared

the observed failure times for fatigue loading with failure times predicted

from times-to-f ailure data under dead load conditions. This comnarison indicates

that damage leading to failure is not additive in terms of the fractional

failure times determined by static loading experiments. We have examined

how the number of chain scissions accumulate in cyclic loading in an

attempt to understand damage accumulation in fatigue experiments. We have

also determined the dependence of the number of chain scissions at failure

with test frequency, peak stress level, and molecular weight.
13

IQ"4 10-3 10-2 10

Test Frequency, Hz

Figure 6a - Time- to- failure data against test frequency as function of peak

load. The dashed curve gives the predicted failure behavior if

failure occurs at critical number of cycles.
- 14 -

Figure 6d - Times-to- failure compared with predictions based on constant

load data assuming the additivity of damage concept.
-15-

MOLECULAR DAMAGE CHARACTERIZATION

The evidence for chain scission associated with mechanical degradation

comes from three rather different types of experiments. The chemistry of

chain scission is depicted in Fig. 7 for polyethylene. The initial rupture

of the carbon-carbon backbone bond produces two unpaired electrons, or

primary free radicals, one on either end of the ruptured chain. These

molecular entities are highly reactive, but can be trapped by conducting the

mechanical degradation at cryogenic temperatures. Electron spin resonance

(ESR) has been used to determine the type and concentrations of free radicals

produced during mechanical degradation. Fig. 8a shows how the concentration of

free radicals builds up during straining of nylon-6 fibers. Up to about 7%

there is no detectible ESR signal, but the signal then rises slowly at first

and more abruptly at higher strains until the specimen failure at about 14%

strain. For most polymers ESR measurements must be carried out at cryogenic

temperatures, therefore this technique is not applicable to polymers in most

service environments. ESR spectra of polyethylene degraded under liquid nitrogen

and held at low temperatures have been interpreted as arising from the

primary radicals (6). As the PE specimens were slowly heated the ESR spectra

changed and spectral analysis indicated the presence of the alkyl radical,

-CH2 CHCH

2 -, shown in the second reaction illustrated in Fig. 7 (6). In the

absence of atmospheric oxygen this free radical decomposes into a primary

radical and a vinyl terminated chain end. In the presence of O

2 the reaction

proceeds through a peroxy radical that can decompose into an aldehyde group

and another oxygen containing free radical which may undergo further reactions

to produce alkyl radicals. The important aspects of the free radical

reactions shown in Fig. 7 are that the initial free radicals produced by

chain scission of over stressed chains may initiate many additional chain
Chain

Scission

Reactions

CM CM X 1 X 1 1 o 1 CM O CM

CM 1 CMX 1 X

X 1CM X

X o CM o 01 o1 o CM

-CH 1CM 1CM o CM X X X 1 X j— o CM o o 1CM o 1 zc 1 1CM CH 1CM

CM X o•

CM

X

X o o 01 1CM x 01 1CM • 1CM 1CN CM X

CN

X

X 01 X o o o 01 01 1CM CM 1 1CM

"7" 1CM CM CM X

X

mJmm

X

X o 01 o o o o 1 1 iCM X0

<N 10 • 1 o X I 0-0 ICM X01CM X01 I I CM CM X X 0 o1 I CM CM X X 0 o1 I cm n X X 0 o1 X CM O XI o

-r 1 X CM O X I oI o4i

CM

X CM

X

O

x 0 i o 1 £ I CM X CM X O X 0 I o 1 4? . CM X X o x 0 «o 1CM X CM _L. 0X0 I o

-rX CM1 —X O X o I O I I 11CMICM

CM I I CM CM . = =

CM

OO

X . , o TM ' s x CM O O x ' o o £ .

CM

O

X01CM X oX01CM CM O x X

CMI o o I I X CM CM o x X I o o I I I I . CM CM ImI i x o o O ' • •CM — X o 01 co X£ 1m X o X o I! O x ' o x % ^ Ox 1 . 5 o CM w I T I CM n cm x V x o

Figure 7-

Chain

scission

reactions

in

polyethylene.
- 17 -

Electron Spin Resonance

coE

-H

J2

03

OMB

T3

03

cn

o

0)w

LL

Figure 8a - Generation of free radicals in nylon-6 fibers by straining
-18 -

scissions, and in the process decrease the average length of the polymer

chains, as well as generate new molecular end groups. Evidence for decrease

in the average chain length comes from measures of molecular weight. Increases

in the concentration of molecular end groups are determined by infrared

spectroscopy. In Fig. 8b is shown how molecular weight changes with strain

in nylon-6 fibers. The behavior is similar to the ESR data.

The initial free radicals produced by mechanical excitation of CC bonds

may recombined with no net changes in the structure. This event is rather

unlikely since the overstressed chain will relax upon chain rupture further

increasing the distance between the two free radicals and decreasing the

probability of recombination.

Another mechanism is transfer of a proton from one side of the

ruptured chain to the other which has the effect of generating two new

chain ends, one of which is a vinyl end group (R-CH=CH

2

) and the other one

a methyl group. In this case, each pair of free radicals results in a

pair of new end groups, one of which will involve a carbon double bond.

In the presence of atmospheric oxygen the two new end groups may involve

oxygen, as an aldehyde group, for example.

A major unresolved question is the extent to which free radical

propagation reactions, such as those depicted in Fig. 7, produce additional

chain scissions from the initial stress induced chain rupture. If a

significant number of such reactions occurs a high local concentration of

chain rupture will exist and these may serve as sites for microvoid

formation. This situation is depicted schematically in Fig. 9 where

an initial stress- induced chain rupture has initiated additional chain

scissions. Even if the molecular segments cleaved by free radical attack

are not load bearing at the time of the stress-induced rupture they are

removed from any future load bearing role that would result from subsequent

local deformation.
19

Viscosity —Averaged

Molecular Weight

«

3

O

c -c — CD

a © w

C3 £^

o

l.

o

0

Q

Figure 8b - Degradation of the molecular weight of nvlon-5 fiber

by straining
20

Figure 9

- CH2 -

H H H

I I

- Schematic of microvoid formation as a result of free radical

propagation reactions.

i

;,l
- 21 -

The extent of free radical propagation reactions may be evaluated

through comparison of the concentration of free radicals produced during

the fracture process and the resulting number of molecular chain scissions.

The latter quantity may be determined in one of two ways. The type and

concentrat.ion of chemical moities at the chain ends can be determined by

infrared spectroscopy. In this technique, infrared absorption bands

identifiable with molecular vibrations of end groups are monitored before

and after application of mechanical stress. The difference in absorption

can then be related to changes in end group concentrations, and from this

information the concentration of chain scission found. The infrared data

may be susceptible to errors from a variety of sources. A discussion of these

in light of recently reported values in the literature will be given.

A second method for estimating the concentration of chain ruptures is

molecular weight degradation studies. Each chain scission lowers the average

length of the chain involved. The cummulative effect is to lower the

molecular weight, or more accurately, the number average molecular weight.

Molecular weight and molecular weight distribution measurements are carried

out on polymer solutions. Hence, these measurements are destructive as far

as solid polymers are concerned. Furthermore, viscometry, the usual

method of determining molecular weight, yields a measure of molecular weight

that is not readily interpretable in terms of number average molecular

weights, and hence concentration of chain scissions.

Zhurkov and colleagues were the first to compare ESR measurements of

the number of free radicals produced by fracturing polymers with the

number of chain scissions as determined by infrared spectroscopy (1,7). In these

studies it was reported that each free radical might represent a large

number of chain scissions. In fact, the number was reported to be several
22

thousand for polyethylene indicating substantially more molecular damage

during fracture than generally thought (1). On the other hand, determinations

of the number of chain scissions during fracture of polyethylene by viscosity

average molecular weight changes yielded values three orders of magnitude

lower (8). These results were reported by a different group who were using

polyethylene from another source (8).

The extremely large discrepancies between data from the two measurement

methods cannot be explained in terms of inherent approximations used to

interpret numbers of chain scissions from viscosity average molecular weight

changes. In addition, infrared and molecular weight degradation studies had

not been performed on the same polymer samples so that the effects of initial

molecular weight, morphology, etc. could not be adequately ascertained. Our

work reported herein has included ESR, infrared spectroscopic, and molecular

weight degradation measurements on the same polymer specimens. In this way,

for the first time, direct comparisons can be made to evaluate the extent of

free radical propagation reactions in polymer fracture and failure.

A major problem encountered in the infrared spectroscopic determination

of new end group concentrations resulting from fracture has been the

sensitivity of the IR measurements to factors other than concentrations of

particular groups. Deforming a polymer to fracture affects sample thickness,

surface conditions, transparency, degree of crystallinity, and chain

orientation. All of these factors can affect the ir measurements thereby

obscuring or exaggerating changes in end group concentrations. Although

in principle it should be possible to separate the spectral changes that are

a consequence of alterations in end group concentrations this has proven to

be a difficult task. In our investigations we used sample preparation

techniques that minimized effects from extraneous factors (9).
- 23

RESULTS

In our research we have used the highly sensitive technique of

Fourier transform infrared spectroscopy (FT-IR) to measure small changes

in concentrations of end groups due to polymer fracture. The FT-IR

technique is several order of magnitude more sensitive than conventional

dispersive infrared techniques, and hence is ideally suited for these

studies (10). Furthermore, the spectroscopic data is obtained in digital form

which facilitates spectral subtraction and data analyses required to determine

changes in end group concentrations. The signal-to-noise advantage of FT-IR

can be utilized to extend investigations of chain scission to less severe

conditions of mechanical degradation.

Polyethylene was chosen for this study because it had been used in

previous studies that used conventional dispersive ir instruments, and the IR

spectrum of polyethylene is relatively free of absorption bands. Commercial

polyethyl enes contain end groups of the kind generated during polymer

fracture as a product of the polymerization process. New end group

concentrations for fractured polyethylene are in the range of one per 10^ to

10° carbon-carbon bonds. Our studies using FT-IR have indicated that

concentrations of at least an order of magnitude lower than this can be

observed. Ideally, one would like the initial end group concentration to be

£ smaller than one per 10 carbon-carbon bonds.

We have characterized the amount of chain scission in polyethylenes

that have been mechanically degraded by (1) grinding under liquid nitrogen,

(2) straining to fracture in tensile loading, and (3) fatiguing under cyclic

loading up to and including failure. The last mentioned series of tests were

carried out at 4 different peak stress levels and 3 different mechanical

test frequencies. Since initial concentrations of end groups of the

same order of magnitude or larger as that produced by mechanical degradation

can greatly complicate the infrared spectroscopic analysis we have examined
- 24

polyethylenes having a range of molecular weights and hence, end group

6

concentrations. The ultra high molecular weight polyethylenes, Mw > 10

i were best suited for the FT-IR studies.

Initial FT-IR and ESR studies were carried out on commercial

polyethylenes in the form of rods that were mechanically degraded by

grinding under liquid and oriented filaments that were strained to

failure in uniaxial tension (9). IR spectra for these two polyethylenes

are shown in Figs 10 and 11. In each of these figures the top trace

• - '

is that of the sample, the middle trace is from the reference material,

and the bottom trace is obtained by digitally subtracting the middle trace

I from the top trace while using the band at 1309 cm as an internal

standard. It is apparant from visual comparison of the top two traces of

either rig. 10 or 11 that the changes in end group concentrations are

small. In the difference spectrum of Fig. 10 we have indicated those

bands attributable to new end groups. These are due to the aldehyde group

at 1735 cm ^

, the methyl group at 1370 cm \ and the vinyl group at 909 cm ^

A complete list of spectral changes are given in Table where comparisons

are made with two previous reports (7,11). The comparisons point to

the principal cause of the unusually large concentrations of end groups

previously reported (7). Note that Zhurkov et al (7) reported a significant incre.,

in the concentration of vinylene groups (RCH=CHR' ) whereas neither our work

nor that of Tabb (11) yielded any changes in this band. Unlike the polyethylene

' specimens used by us and Tabb those used by Zhurkov and co-workers contained

a significant concentration of vinylene groups as evidence by an intense

ir band at 965 cm * (12). We attribute the spectral changes reported by Zhurkov

for the vinylene band to artifacts such as sample thickness, orientation,

void formation that were accentuated by the significant absorption at 965 cm

in the reference material.
poiyetnyiene

tensile

fracture

25

2>DUDqjosqv

Figure

10

-

Fourier

transform

Infrared

spectra

of

tensile

fractured

polyethylene,

the

reference

spectrum

material,

and

their

difference

spectrum.
1/1

JQ

E

c

>

o

ao

c

a;-

(U

1—

0)>-O

13

u

r*5

0!>>

o

Sj • cO S

3

u '-i O JJo

3 03 4J C- O M

033- 03

33 O

-o a

>3 iJ

s- S3

3 3

'- 3

I

a

•—M
27

TABLE 1

IR Spectral Changes in

Tensile Fractured Polyethylene

Spectral Changes*

Frequency

(cm~l) Assignment This Work

Zhurkov

Ref. 1

Tabb

Ref .11

909 RCH = CH

2 + + +• + + +

965 trans R CH = CHR' - + -

990 RCH = ch

2 - - +

1370 ch

3

def. RCH

3 + + + + +

1642 c=c str. RCH = ch

2 - NR + +

1717 c=o str. RCOR' + + -

1735 c=o str. RCHO + + + + +

1898 Combination mode, crystalline + NR +

2020 Combination mode, crystalline + NR + + •

* Magnitude of change indicated by number of + signs, the larger the

number, the greater the change. No change is showned as - and NR if

not reported.

To circumvent artifacts due to chain orientation, etc. we prepared

films for ir measurements by dissolving the material to be investigated

in a suitable solvent from which films were cast for infrared

spectroscopic analysis. Alternately, we melt pressed films for IR

spectroscopic study directly from the mechanically degraded materials. Our

findings were the same with either method of specimen preparation.

A third specimen examined in our preliminary work was an ultra high

molecular weight polyethylene which was also ground under liquid N . The

spectra of this polyethylene are shown in Fig. 12 where again the top

trace is from the fractured material, the middle trace is from the reference,

and the bottom trace is their difference. Note that there is virtually no
. Figure 12 - Fourier transform infrared spectra of ground ultra high

molecular weight polyethylene, the reference material and

their difference spectrum.
29

absorption in the reference spectrum at either the aldehyde group frequency,

1735 cm

1

, or the vinyl moiety frequency at 909 cm ^ . Compared to the

lower molecular weight polyethylenes , Figs. 10 & 11, ultra high molecular

weight polyethylene is an excellent candidate material for spectroscopic study

since small changes in end group concentrations can be more readily detected.

Concentrations of new end groups attributable to mechanical degradation

and determined from the difference spectrum of Fig. 10 are presented in

Table 2. Integrated band intensities were determined by digital integration

of the baseline corrected absorption bands. Values for absolute intensities

were taken from the literature (13,14), and the Beer's law relationship was assumed.

The number of chain scissions was estimated from end group concentrations by

summation of the number of aldehyde and vinyl groups. Each chain rupture

that leads to two end groups produces one carbon double bond, as shown in

Fig. 7.

TABLE 2

New End Groups Formed During Fracture

of Polyethylene

Group

IR Band Freq.

(cm-1

) Concentrations

(No . /cc)

RCH=CH

2 909

RCH

3 1370

1735

1 . 3± . 4 x 1017

7.31.4 x 1017

RCHO 1.91.6 x 1017
30

The number of chain scissions per gram are listed in Table 3 where

comparison is made to the number of free radicals produced during

the degradation process. Ratios of these two concentrations provides a

measure of the extent of free radical reactions. Values of the ratio of

chain scissions to free radicals range from 11-125 depending on molecular

weight and type of mechanical degradation. These numbers are much smaller,

by at least a factor of 20, than those previously reported from comparison

of IR and ESR data (7). Our values are larger by 1 to 2 orders of magnitude

than that reported by Stoeckel et al from viscosity-average molecular weight

and ESR data on tensile fractured specimens (8).

TABLE 3

Chain Scissions per Free Radical

Ground PE

Tensile Zhurkov,

Commercial UHMW Fractured PE et . al

.

No. Chain Scissions

Per Gram 5.1 ± lxlO17 1.2 ± . lxlO18 3.4 ± lxlO17 9.9 x 10

18

No. Free Radicals

Per Gram 1.2 x 10 16 1 x 1017 3.5 x 10

15

5 x 10

15

Chain Scissions

Per Radical 34 - 50 11 - 13 68 - 125 1,980

Viscosity average molecular weight measurements do not provide a

direct measure of the number of chain scissions because viscosity depends on

weight average rather than number average molecular weight. The analysis

of changes in viscosity average molecular weight to determine number of chain

scissions requires assumptions of the dependence of relative probabilities of

rupture on molecular length, as well as the molecular weight distribution.
31 -

We have used gel permeation chromatography, GPC, to determine the change

in number average molecular weight with mechanical degradation. Plots of

the molecular weight distributions of the tensile fractured polyethylene

and the reference material are shown in Fig. 13, and the number of chain

scissions deduced from these distributions are given in Table 4 along with

the ratio to the number of free radicals. Unlike previous determinations the

chain scission concentration determined from molecular weight changes is in

good agreement with the results of the infrared data. Thus, the discrepancy

between previous measurements of chain scission concentrations may have been

caused by uncertainties in viscosity average molecular weight measurements,

method of mechanical deformation, or, molecular weight distribution of the

original material.

TABLE 4

Chain Scissions Per Free Radical

IR GPC

No. Chain Scissions

Per Gram 3.4 ± lx 1017 5.4 x 1017

No. Free Radicals

Per Gram 3.5 x 10 15 3.5 x 10

15

Chain Scissions

Per Radical 68 ~ 125 154

Zhurkov and co-workers (1) examined the mechanisms leading to the development of microvoids in stressed polymers by comparing the number of free radicals
Fractional

Amount

32

j Figure 13 - Molecular weight distribution of tensile fractured

polyethylene and its reference material.
33 -

produced during mechanical degradation with the concentration of chain

scissions and the size and concentration of microvoids generated. The

microvoid concentration and size were deduced from small angle x-ray

scattering data using a two-phase model - polymer and microvoid. It was

reported that the number of microvoids was approximately equal to the number

of free radicals, and that the average cross-sectional area of the microvoids

was consistent with the product of the number of chains ruptured per free

radical and the cross-sectional area per chain. Thus, their data suggested

a model in which one microvoid is generated per mechanically-induced free

radical and the size of the microvoid is controlled by the number of chain

scissions produced per mechanically generated free radical.

In light of our finding that the number of chain scissions produced

in tensile fracture polyethylene per radical is lower by at least an order

of magnitude we are planning to carry out x-ray scattering studies on our

samples. In analyzing the small angle x-ray data we will assume a three

phase system - crystalline, amorphous, and microvoid, instead of the two

phase system used by Zhurkov. The mathematical equations for a three

component system has been worked out and applied to a particulate composite

assumed to be composed of polymer, particulate reinforcement, and voids (15).

FATIGUE OF ULTRA HIGH MOLECULAR WEIGHT POLYETHYLENE

The second aspect of our work in this contract period has been concerned

with investigating the relationship between growth in the chain scission

concentration with the state of mechanical damage and whether failure could

be associated with a critical concentration of chain scissions. This work

was carried out on ultra high molecular weight polyethylene in order to take

advantage of the initial low concentration of end groups in these materials.
34 -

Mechanical degradation was controlled in uniaxial fatigue tests in which

both frequency of the applied load and its peak value were varied to

obtain failure times over several decades. The applied stress was of the

form

a(t) = A (1 + sin wt)

so that the applied stress varied from zero to 2k. The mechanical testing

results are summarized in Figs. 6a & b in which it is seen that

time-to-failure of a particular test frequency can vary from one to two

orders of magnitude depending on peak stress.

In equation 2 a relationship is derived between the time to failure and

a critical concentration of chain ruptures. In the additivity of damage

concept the amount of damage in a material increases in time at

a rate determined by the applied stress and temperature. At a particular

stress level, the fractional damage is given by the ratio of the time the

specimen is held at that stress to the time it would take the specimen to

fail at the same stress level. The latter times-to-failure are determined

in dead load tests. From the dead load failure times the times-to-failure

under cyclic loading conditions can be found and these are compared to

experimental fatigue data in Fig. 6. It is seen in Fig. 6 that the

experimental failure times are significantly higher than the predicted

times which raises questions concerning the notion of a critical concentration

of chain ruptures determining failure, the additivity of damage concept, or both.

We have examined the concept of failure associated with a critical

concentration of chain scissions by determining the amount of chain scission

at failure for polyethylene specimens fatigued to failure at different

peak stresses and that exhibited quite different failure times.
35 -

The spectrum of an ultra high molecular weight polyethylene fatigued

to failure at 0.002 Hz is shown in Fig. 14 along with that of the reference

material and their difference. Both aldehyde and methyl end groups

have increased in concentration, and consistent with the findings on

ground ultra high molecular weight polyethylene, virtually no increase in

vinyl group concentration has been detected. The chain scission

concentrations are found from the integrated intensities of the aldehyde

group band at 1735 cm ^ . The difference spectra from specimens fatigued

to failure at peak stresses of 35 to 45 MPa are shown in Fig. 15. The

areas of these peaks are virtually the same even though times to failure

varied by almost a factor of 20. The concentrations found from these peaks

are given in Table 5, and their equality is indicative of failure at a

critical concentration of chain scissions.

TABLE 5

Fatigue of Ultrahigh MW Polyethylene

Stress (MPA) Concentration (No./CC)

C=0 rch=ch2 30 1 . 95xl018 -

35 1 . 82xl018 3xl016

40 2 . 03xl018 -

45 1. 97xl018 Ixl01&

Spectroscopic investigations have been carried out on a second ultra

high molecular weight polyethylene with a manufacturer's reported molecular

weight that was one-half the value of the material used in the aforementioned

study. For this material it was found that failure occur at a mechanically

induced chain rupture concentration of about 1 * 10-^/gram instead of

2 * 10 /gram. This finding further suggests that molecular weight is a

determining factor in the level on molecular damage that can be tolerated

before failure occurs.
ABSORBANCE

.000

.030

.060

.090
37

STRESS LEVEL

45 MPA

40 MPA

35 MPA

30 MPA

1800 1750 1700 1650

WAVENUMBERS

igure 15 - Difference spectra due to aldehyde groups obtained from

ultra high molecular weight polyethylene fatigued to

failure at 0.002 Hz and at the peak stresses indicated.
38

Our finding that macroscopic failure occurs at a critical concentration

of chain ruptures suggests that monitoring the buildup in end group

concentrations may serve as the basis for lifetime prediction. In order

to fully assess the potential of this approach we have examined the increase

in chain scission concentration during the course of a f atigue-to-failure

test.

In this experiment the test specimen was subjected to a predetermine

number of stress cycles, removed from the mechanical tester, and placed in

the interferometer for infrared spectroscopic analysis. The specimen was

then returned to the mechanical tester for another series of stress cycles,

and the entire process repeated until the specimen failed. The increase in

aldehyde group concentration as a function of the number of stress cycles

is shown in Figs. 16ab for a specimen fatigued at 0.002 Hz and at a teak stress

level of 3_5 MPa. During the initial stress cycles the aldehyde concentration

increased only slightly. This was followed by a rapid rise to about 60%

of the aldehyde concentration at failure after which it rose linearly with

the number of stress cycles until macroscopic failure. In Figs. 16ab the

aldehyde concentration is plotted as its ratio to that at failure. A total

of three specimens were examined under identical test conditions, and all

displayed similar dependence of aldehyde concentration on number of fatigue

cycles.

We have also investigated the effect of fatigue test frequency on the

concentration of chain scissions at failure. These results are summarized

in Fig. 17 for the lower molecular weight ultra high molecular weight

polyethylene. The weak dependence on test frequency may be caused by a

slight rise in sample temperature as the fatigue frequency increases.

Although according to equation 1 the probability of chain rupture increases

with increasing temperature the critical concentration would also be expected
Ratio

To

Final

Carbonyl

Concentration

39

i Figure 16a - Initial rise in aldehyde group concentration.
uojiejiuaouoo jAuoqjgQ

!SU!J 01 ouey
Number

of

Carbonyls/cc

x

10

- 41 -

(Figure 17 - Aldehyde group concentration at failure as a function

of fatigue test frequency. Peak stress level was 35 MPa.
42

to rise due to a greater probability for free radical termination reactions

as a result of more molecular mobility. The extent of free radical propagation

reactions would decrease at higher temperatures forming damage zones that

would not be of sufficient dimensions to coalesce into the microvoids that

are the precursors to macroscopic failure.

The growth in the aldehyde concentration with the number of stress

cycles has also been measured for specimens fatigued at 0.01 Hz. These

experiments yielded a similar behavior with number of stress cycles as that

shown in Fig. 16b. Thus, the general features of the rise in chain scission

concentrations appear to be independent of the fatigue test frequency.

The fatigue experiments that have been carried out are constant peak

stress tests. During the fatigue expriment the initially isotropic test

specimen is elongated and preferential chain axis orientation in the

direction of the applied stress occurs. The sample morphology is transformed

from spherulitic texture. Fig. lb to the fiber morphology of Fig. 2. The

number of highly stressed molecules would be expected to change as these

morphological changes occur. Thus, the rise in chain scission concentrations

may correlate more directly with the imposed strain than the number of stress

cycles in a fatigue experiment. Furthermore, the behavior of chain scission

concentration with the first few cycles in fatigue (Fig .16a) is very

similar to the strain induced buildup in free radical concentration or

degradation of molecular weight.

Future work will address the relationship between chain scission

concentration and strain in a stress relaxation experiment.

CONCLUSION

Three important aspects of molecular processes associated with polymer

fracture have been investigated. Detailed experiments combining ESR

measurements of free radicals concentrations , infrared spectroscopic

determinations of new end group formation, and GPC measurements of molecular
- 43

weight degradation have been carried out on tensile fractured polyethylene.

Excellent agreement has been found between the number of chain scissions

associated with mechanical degradation as determined by increases in end group

concentrations and by the change in molecular weight . When the

concentration of chain scissions is ratioed to the concentration of free

radicals it is found that each free radical produces an additional 100-150

chain scissions. This high local concentration of chain scissions is

thought to be the precursor to microvoid formation from which macroscopic

cracks that lead to failure originate. Future experiments will use small angle

x-ray scattering to measure the size and concentration of these microvoids.

Previous work had suggested that the number of microvoids should approximately

equal the number of free radicals and the size of the microvoids should

relate to the number of chain scissions produced per free radicals.

The second aspect of our present work has examined the effect of

molecular weight and test frequency on the critical concentration of chain

ruptures at macroscopic failure in a fatigue test. It was observed that

decreasing the molecular weight by a factor of 2 decreased the critical

concentration by a factor of 2 also. The concentration of chain scissions

at failure increased slightly with increasing frequency.

The third aspect concerned an investigation of how the end group

concentration, or chain scission concentration, builds up during fatigue.

It was found that after a slow rise during the initial few stress cycles,

the concentration grew rapidly to about 60% of its value at failure.

After the abrupt rise the concentration increased linearly with the number

of fatigue cycles until failure. It is concluded that monitoring the number

of chain scissions through infrared spectroscopy may be a useful procedure

for estimating the remaining lifetime.

References

1. S. N. Zhurkov, V. A. Zakrevskii, V. E. Korsukov, and V. S. Kuksenko,

Sov. Phys. Solid State, 13, 1680 (1973).

2. "Developments in Polymer Fracture-1", E. H. Andrews, Editor, Applied

Science Publishers, London 1979.

3. V. M. Voroboyev, I. V. Razumovskaya, and V. I. Vettegren, Polymer, 19 , 1267 (1978).

4. J. M. Crissman, F. A. Khoury, and G. B. McKenna, NBSIR 82-2493, U.S.

Department of Commerce, National Bureau of Standards, 1981.

5. G. B. McKenna and R. W. Penn, Polymer 21, 213 (1980).

6. T. Kawashuna, S. Shimada, H. Kashiwabara, and J. Sohma, Poly. J. , _5,

135 (1973)

7. S. N. Zhurkov and V. E. Korsukov, J. Polym. Sci. Polym. Phys. Ed 12 , 385 (1974).

8. T. M. Stoeckel, J. Blasius, and B. Crist, J. Polym. Sci. Polym. Phys. Ed.

16, 485 (1978).

9. B. M. Fanconi, K. L. DeVries, and R. H. Smith, Polymer _23_, 1027 (1982).

10. R. J. Bell, Introductory Fourier Transform Spectroscopy, Academic Press,

New York, 1972.

11. D. L. Tabb, PhD Thesis, 1974, Case Western Reserve University.

12. V. E. Korsukov, V. I. Vettegren, and I. I. Novak, J. Polym. Sci.

Symposium No. 42, 1299 (1973).

13. H. S. Wexler, Spectrochim. Acta 21 , 1725 (1965).

14. V. I. Vakhlyueva, A. G. Finkel, L. M. Sverdlow and A. I. Andreeva,

Opt. Spectrosc., 25, 234 (1968).

15.

W-L. Wu, Polymer, 23, 1907 (1982).

NBS-114A (REV. 2-80

U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET (See instructions)

1. PUBLICATION OR

REPORT NO. NBSIR 83-2682

2. Performing Organ. Report No 3. Publication Date April 1983

4. TITLE and subtitle Annual Report to Defense Advanced Research Project Agency, Covering

Period 8/1/82-12/31/82, Molecular Processes Affecting Mechanical Failure of Polymeric

Materials.

5. AUTHOR(S)

Dr. Bruno M. Fanconi

6. PERFORMING ORGANIZATION (If joint or other than NBS, see in struction s)

national bureau of standards

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

7. Contract/Grant No. ARPA 4609

8. Type of Report & Period Covered

Annual

8/1/82-12/31/82

9.

SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City, State. ZIP)

10.

SUPPLEMENTARY NOTES

| | Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11.

ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bi bl iography or literature survey, mention it here) Preliminary observations have established correlations between the buildup in the

concentration of molecular end groups and macroscopic failure in polymers. Current

work has examined the effects of stress history, molecular weight and the dependence

of the growth in end group concentrations with time during the course of a mechanical

degradation experiment. In cyclic fatigue under uniaxial loading conditions the

concentration of end groups was found to exhibit a sigmoidal dependence on the number

of cycles. After an initial rapid use to about 60% of its value at failure the number

of chain scissions was observed to increase linearly until sample failure. Such a relationship holds promise as the basis of methods to predict mechanical failure. The

effects of frequency, peak stress, and molecular weight have also been examined and are reported.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolon s) infrared spectroscopy, chain scission, polyethylene, mechanical degradation, molecular

weight.

13. AVAILABILITY 14. NO. OF

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