ORI GIN AL PA PER

Radiation-induced modification of dielectric relaxationspectra of polyolefins: polyethylenes vs. polypropylene

Dejan Milicevic • Maja Micic • Edin Suljovrujic

Received: 25 February 2014 / Revised: 11 June 2014 / Accepted: 17 June 2014 /

Published online: 29 June 2014

� Springer-Verlag Berlin Heidelberg 2014

Abstract The molecular relaxation behaviour of polyolefins exposed to high-

energy radiation has been investigated by dielectric loss (tan d) analysis. Therefore,

low-density polyethylene, high-density polyethylene (HDPE), and isotactic poly-

propylene (iPP) were gamma-irradiated in air to various absorbed doses (up to

700 kGy). All relaxation zones (a, b, c, and d in the order of decreasing tempera-

ture), between 25 K and melting temperature, were studied. The radiation-induced

changes observed in the dielectric relaxation spectra were related to the modifica-

tions in the structural and morphological parameters attributed to exposure of the

polyolefins to radiation. Wide-angle X-ray diffraction, infrared spectroscopy, and

gel measurements were used to determine the radiation-induced changes in the

crystalline structure, oxidative degradation, and the degree of network formation,

respectively. The present study reveals high dielectric and/or relaxation sensitivity

of polyolefins to gamma radiation. Disappearance of some relaxations (such as brelaxation in HDPE and low temperature c and d relaxations in iPP) is clearly

observed with irradiation. For the other relaxations, besides the large changes in the

relaxation intensity, radiation also induces smaller/larger changes in the distribution

of relaxation times, peak position, and activation energy.

Introduction

Polyolefins have excellent mechanical and dielectric properties and, therefore, a

wide variety of industrial applications, including electrical ones. Due to their low

dielectric loss and good heat resistance they have been widely used as electrical

insulation, e.g. for cables and as a dielectric in power capacitors [1–3]. Considering

the molecular structure of apolar hydrocarbon polymers (such as polyethylene,

D. Milicevic � M. Micic � E. Suljovrujic (&)

Vinca Institute of Nuclear Sciences, University of Belgrade, Belgrade, Serbia

e-mail: edin@vinca.rs

123

Polym. Bull. (2014) 71:2317–2334

DOI 10.1007/s00289-014-1190-6

polypropylene, etc.), the dipole moments of the (C–H apolar) groups contained in

these polymers are very low (an order of 0.1 debye) and hardly detectable by the

usual dielectric techniques [3]. Despite this, apolar polymers exhibit measurable

dielectric spectra corresponding to the transitions measured by the mechanical

relaxation techniques. The measurable dielectric relaxations and losses are generally

ascribed to impurities and to the fact that these polymers are always slightly

oxidized and thus contain polar carbonyl, peroxy, or hydroperoxy groups. Among

impurities, residual catalysts and antioxidants have been reported to affect the

dielectric properties. However, for the electrical application of such polymers it is of

essential interest to understand the dielectric phenomena in them. Furthermore,

dielectric measurements can give valuable information about the structure and

dynamics of materials. It is well known that the dielectric response can be used as an

indicator of condition and ageing processes occurring in polymer insulation [4, 5].

In dielectric and mechanical relaxation studies, polyethylene and polypropylene

usually display three characteristic relaxation zones; these have conventionally been

designated as a, b, and c relaxations in order of decreasing temperature. Although

some detailed molecular assignments are still open for debate, the reality of the

basic relaxation processes is clear; these have been well summarized by Boyd [6]. In

addition, iPP may also exhibit a fourth relaxation, named the d process, which

occurs at lower temperatures compared to other relaxations [7].

In order to better investigate the molecular relaxations, different structural and

morphological modifications of polyolefins have been performed in the past, mainly

by drawing, thermal treatment, irradiation, ageing, artificial weathering [5], slight

oxidation of the chains, and by doping the polymer matrix with polar molecules and

particles as probes. In general, radiation-induced structure modifications can be

exploited not only from the standpoint of commercial applications but also as a

useful tool for highlighting some of the fundamental processes and properties of

polymers. Thus, exposition to radiation can be used to increase the dielectric

activity and to investigate low visible dielectric relaxations in apolar polyolefins [8].

Even though the overall radiation chemistry of polyolefins was investigated in detail

and several comprehensive reviews are available on this topic [9–12], the effects of

radiation on dielectric relaxation behaviour have not been investigated to an

appreciable extent [8, 13, 14]. A major application of high-energy radiation is

crosslinking of insulation; crosslinking to a gel content of 55 % was shown to be

beneficial for cable insulation [9]. By linking the macromolecules into a network,

the toughness, impact resistance, chemical resistance, and working temperatures are

improved [1]. A second major application of high-energy radiation is sterilization of

the medical disposables [15]. Furthermore, polyolefins are also used in many

applications (nuclear power plants, radiation equipment, sterilization systems,

space-based applications, etc.) where exposure to high-energy radiation can occur.

Radiation-induced changes greatly influence the dielectric properties (electric

strength, dielectric loss, permittivity, electric conductivity, charge state, etc.) and

dielectric relaxation spectra of apolar polymers such as polyolefins. The introduc-

tion of carbonyl, hydroperoxide, and other polar groups as a result of radiation-

induced oxidation intensifies dielectric losses. Furthermore, crosslinking and chain

scission affect the mobility of macromolecules, especially in the amorphous phase,

2318 Polym. Bull. (2014) 71:2317–2334

123

causing a shift of the relaxation maxima and a change in activation energy of the

dielectric relaxation to which the mentioned dipolar and molecular movements

contribute [16].

Our aim is to draw a complete dielectric relaxation map of virgin and gamma

irradiated LDPE, HDPE, and iPP, and to establish a connection between the

evolution of the dielectric relaxations and the radiation-induced changes in the

structure. Dielectric relaxation spectroscopy (DRS) and gamma radiation were used

as powerful methods for characterization and modification of polymer structure,

respectively. In the case of dielectric relaxation measurements, the polar groups that

were introduced by radiation were regarded as tracer groups whose motion reflected

the motion of polymer chains. A variety of supplementary measurements were made

to qualitatively determine the radiation-induced changes in the structure. The results

obtained by IR, WAXD and gel measurements were compared with the changes in

the intensity, position, and activation energy of dielectric relaxations.

Experimental

Three types of polyolefins were used in this study: LDPE (HIPTEN 22003A3,

q = 0.922 g/cm3, Mw = 110,000), HDPE (Hiplex HHM 5502, q = 0.955 g/cm3,

Mw = 300,000) and iPP (HIPOL MA2CR type C-7608 q = 0.906 g/cm3,

Mw = 136,000). Isotropic sheets were prepared by 20 min compression moulding

in a Carver laboratory press, at different temperatures (150 �C for LDPE, 160 �C for

HDPE, and 190 �C for iPP) and a gradual increase in pressure up to 3.28 MPa. The

moulded sheets were quenched in an ice-water mixture. The samples

(0.28 ± 0.02 mm thick) were wrapped in Al-foil and irradiated in a 60Co radiation

facility, in air, at room temperature, at a dose rate of 9 kGy/h, to various absorbed

doses up to 700 kGy.

Wide-angle X-ray diffractograms of the samples were obtained using a Bruker

D8 Advance Diffractometer (in normal mode, with Cu Ka emission). The parallel

beam optics was adjusted with a parabolic Gobel mirror (push plug Ni/C) with a

horizontal grazing incidence soller slit of 0.12� and an LiF monochromator;

diffractometer scans were taken in the angular range of 2h = 10�–45�, with a step

of 0.02�, and 10 s exposition per step. Furthermore, crystallinity was evaluated from

diffraction curves by resolving multiple peak data into individual crystalline peaks

and an amorphous halo. Quantitative analysis and fitting of multiple peaks in

experimental spectra were performed using standard software. For the gel

measurements, the samples were inserted into a 200 mesh stainless steel cloth

and immersed in xylene with 0.5 wt % antioxidant (Irganox 1010). The gel content

was determined by measuring the weight loss of the samples after solvent extraction

in the boiling xylene for 17 h, followed by drying the samples for 4 h in a vacuum

oven at 60 �C. The results are average values of five identically prepared samples.

A Carl-Zeiss Model 75IR Specord was used for recording infrared spectra. The

absorbencies at 1,720 and 1,715 cm-1 were determined from these spectra for PEs

and iPP, respectively. The oxidative (carbonyl) content was measured through

normalized absorbance (A/d values; A-absorbance; d-sample thickness).

Polym. Bull. (2014) 71:2317–2334 2319

123

The dielectric loss tangent (tan d) of the samples in the form of discs 1.3 cm in

diameter was measured on a Digital LCR Meter 4284A coupled with a 22C-

kriodin(R) cryosystem, as a function of temperature (25–405 K) and in the

frequency range 1 kHz–1 MHz. Dielectric measurements were taken at increments

of approximately 2 K during a heating run, with a heating rate of 1.7 K/min

between equilibrated temperatures. At each equilibrated temperature, measurements

of capacitance and tan d were taken at several frequencies from 1 kHz–1 MHz; data

acquisition over the frequency range required about 5 min.

Results and discussion

Figures 1, 2, 3 depict the dielectric relaxation spectra of virgin polyolefins (Fig. 1)

and those irradiated to absorbed doses of 100 (Fig. 2) and 700 kGy (Fig. 3).

Dielectric relaxation spectra for virgin PEs confirm the presence of three relaxation

zones (a, b, and c in the order of decreasing temperature), while iPP spectra show

additional, less visible d relaxation zone at low temperatures. At a frequency of

LDPE

tan

δδ[1

0-4]

βγ

(a)

α

T[K]100 200 300 400

0

2

4

100 kHz

T[K]100 200 300 400

tan

δ[1

0-4]

β

γ(b)

α

unirradiated

0

4

8

T[K]100 200 300 400

tan

δ[1

0-4]

(c)

2

4

1 MHz

0

2

0

αγ

δ

β

αγδ

β

100 kHz

100 kHz

HDPE

iPP

Fig. 1 a Dielectric loss tangentversus temperature for virgin(unirradiated) polyolefins:LDPE (a), HDPE (b),and iPP (c)

2320 Polym. Bull. (2014) 71:2317–2334

123

100 kHz, the high temperature a relaxation can be observed around 340, 370, and

330 K in LDPE, HDPE, and iPP, respectively. The dielectric a relaxation is

universally observed in all crystalline polymers and is usually attributed to the

motion of chain units within the crystalline portion but the amorphous phase in the

neighbourhood of the crystallites also contributes to this process. It has been

reported that the incorporation of structural and chemical factors into the chain, such

as chlorination, branching, or copolymerization with non-crystallisable units,

decreases the intensity of this transition and in some cases, with high chlorine or co-

unit concentration, the relaxation even disappears [17]. There is strong evidence,

from dielectric and NMR measurements, that the a process in PEs is dielectrically

active due to reorientation of carbonyl groups in the chains in the crystalline phase

[18]. Boyd [19] has proposed that the dielectric a process can be represented by

propagation process of a twisted defect along the chain within a crystal lattice,

leading to reorganization of the crystal surface. On the other hand, in the case of PP

the contribution of the amorphous phase in the neighbourhood of the crystallites to

the dielectric a relaxation cannot be excluded [14, 20]. Many studies indicate that

0

7

T[K]100 200 300 400

tan

δ[1

0-4]

14

γ α

β

0

7

14

T[K]100 200 300 400

β

γ αta

n δ

[10-4

]

0

35

tan

δ[1

0-4]

70

T[K]100 200 300 400

αβ

γδ

LDPE

D=100 kGy

HDPE

iPP

100 kHz

1 MHz

100 kHz

(a)

(b)

(c)

Fig. 2 a Dielectric loss tangent versus temperature for polyolefins irradiated in air to 100 kGy: LDPE(a), HDPE (b), and iPP (c)

Polym. Bull. (2014) 71:2317–2334 2321

123

the mechanical/dielectric a relaxation zone in polyolefins contains two or even more

relaxations with different origins, relaxation times, and activation energies. In

general, most authors admit that the dielectric a relaxation zone in PE is formed by a

single relaxation [18, 21], while for PP it is observed that this dielectric relaxation

zone has a complex nature consisting of two or even three independent relaxation

processes [14]. At higher frequencies (1 MHz), the multiple nature of the dielectric

a process is clearly evident for iPP (Fig. 1c). It was found by Pluta and Kryszewski

[22] that the morphology and structure differentiation significantly influence the

nature and the number of components of the mechanical a relaxation process in iPP.

The presence of a smectic phase as well as the decrease of both the spherulite size

and the structure perfection lead to enhancement of the mobility of crystallites, and

consequently to an increase of the contribution from the intralamellar regions to the

a relaxation process. A comparison between the dielectric relaxation spectra of

LDPE and HDPE shows a large difference in the dielectric a relaxation zone

(Fig. 1a, b). The increase in magnitude can clearly be connected with crystalline

content; for virgin LDPE v = 27 %, while for virgin HDPE v = 52 %. Besides the

D=700 kGy

100 kHzta

n δδ

[10-4

]

iPP1 MHz

0

35

70

tan

δ[1

0-4]

T[K]

100 200 300 400

βγ α

0

20

40

tan

δ[1

0-4]

T[K]

100 200 300 400

γ α

β100 kHz

LDPE

HDPE

0

50

100

T[K]

100 200 300 400

αβ

γδ

(a)

(b)

(c)

Fig. 3 a Dielectric loss tangent versus temperature for polyolefins irradiated in air to 700 kGy: LDPE(a), HDPE (b), and iPP (c)

2322 Polym. Bull. (2014) 71:2317–2334

123

changes in the relaxation intensity, the difference in the crystal phase also induces

changes in peak position. The dielectric a process occurs at much higher

temperatures in HDPE (370 K) than in LDPE (340 K). The position of the aprocesses seems to be governed by the mean thickness of the crystallites, and this is

the important feature of the a relaxation. It has been demonstrated by Popli et al.

[23] that the temperature of this transition increases with the crystallite thickness for

a series of branched, linear, and metallocene catalyzed PEs. Hereinafter, objective

values for the temperatures of the relaxation processes were obtained using curve

fitting from isochronal loss scans at several frequencies. Loss map for the dielectric

processes is presented in Fig. 4 for virgin LDPE, HDPE, and iPP. The a processes

show Arrhenius behaviour

fmax ¼ fmax;1 exp � Ea

kT

� �ð1Þ

where fmax,? is a dimensional parameter, Ea is the apparent activation energy and

k is Boltzmann’s constant. The calculated apparent activation energies for the

dielectric a relaxation in LDPE, HDPE, and iPP are 95.5, 108, and 115 kJ/mol,

respectively; these values are in good agreement with the literature data which

usually range from 90 to 170 kJ/mol [6, 8, 14, 24–26].

The intermediate dielectric b relaxation can be observed in LDPE, HDPE, and

iPP around 275, 270, and 300 K, respectively (Fig. 1). In general, this process has

its origins in the amorphous fraction, but only for the case of PP this relaxation is

undoubtedly attributed to the glass–rubber transition [25]. The reported apparent

activation energies are in the range of 300–700 kJ/mol for PP [24, 25, 27–33].

1000/T [K-1]

106

105

104

103

2.5 3.0 3.5 4 5 6

Los

s-pe

ak fr

eque

ncy

[Hz]

LDPE

HDPE

iPP

8 12 16

Fig. 4 Loss map for the a, b, c, and d processes in virgin polyolefins. The full curves are Arrhenius fits tothe dielectric a, b, c, and d relaxations as well as the VFTH fit to the dielectric b relaxation in iPP

Polym. Bull. (2014) 71:2317–2334 2323

123

For PEs, a connection between the b relaxation and glass transition is controversial.

According to many authors, the b relaxation is attributed to the cooperative

segmental mobility of disordered chains [6, 19, 23] and connected with the glass

transition [34], especially in the case of LPE [35]. On the other hand, 13C NMR

measurements have shown that there is no direct correlation between the

temperatures of glass transition and b relaxation [36]. Significant differences

between the reported activation energies of 50–115 kJ/mol [37, 38] and

180–500 kJ/mol [39, 40] suggest that in the former case the b relaxation in PE

should be treated as a motion in interfacial regions, and in the latter as a highly

cooperative process such as the glass transition. The comparison between the

dielectric relaxation spectra of LDPE and HDPE shows a large difference in the

dielectric b relaxation zone. The b relaxation in branched PE (LDPE) is quite

prominent (Fig. 1a), but in HDPE it is much less so (Fig. 1b). There are few reasons

for this. The first and the main are connected with the amorphous nature of this

relaxation—its magnitude decreases with increasing crystalline fraction. The second

concerns the effect of the semi-crystalline environment. The presence of the crystal

surfaces and the connections of the amorphous chains to them have an immobilizing

effect on the b relaxation. The constrained chains are unable to relax completely;

the relaxation strength is reduced. Through a careful examination of the crystallinity

dependence of the b relaxation process, Popli et al. [23] have demonstrated that this

relaxation results from the relaxation of chain units in the interfacial region. The

interlamellar content increases with increasing degree of branching, due to which

the b relaxation is more pronounced in branched PE, whereas in linear PE it may not

occur. The third reason is connected with the time temperature behaviour in

comparison with the other two relaxations. There is a relatively limited window of

frequency and temperature where the b process clearly can be observed in HDPE

[21]. However, objective values for the temperatures of the b process were obtained

using curve fitting. The loss map for the b relaxation in PEs is a little bent (Fig. 4),

indicating some cooperative behaviour. Despite this fact, the b relaxation can be

successfully fitted by the Arrhenius equation; the correlation coefficients of linear

regression are close to one (fc [ 0.98) and the calculated activation energies are

185 kJ/mol for LDPE and 280 kJ/mol for HDPE. It is apparent that the connections

to the crystal surface have a substantial constraining effect on the dynamics in the brelaxation zone, as it is suggested by Graff and Boyd [21]. On the other hand, the

Arrhenius plot of the b relaxation is significantly bent for iPP, which indicates

cooperative behaviour. This is a typical feature found with relaxations that are

related to glass transition. It confirms the hypothesis that the b relaxation in iPP is

related to the glass transition in the amorphous phase. The fmax(T) function of the brelaxation is fitted by the Vogel–Fulcher–Tammann–Hesse (VFTH) equation:

fmax ¼ fmax;1 exp � B

T � T1

� �ð2Þ

where fmax,?, B and the Vogel temperature T? are VTF parameters [41]. The Vogel

temperature is closely related to the dynamic glass-transition temperature Tg, which

is usually defined as the temperature where the relaxation time (s) is 100 s.

2324 Polym. Bull. (2014) 71:2317–2334

123

The VTF parameters are also used for determining the dynamic fragility m and the

apparent activation energy Ea at Tg. For virgin iPP, the values obtained for the glass-

transition temperature Tg = 281 K, dynamic fragility m = 103 and apparent acti-

vation energy at glass-transition Eg = Ea(Tg) = 560 kJ/mol are in good agreement

with the results published by Plazek and Ngai [42].

The low temperature dielectric c relaxation can be observed around 190 K in PEs

and 225 K in iPP, at a frequency of 100 kHz (Fig. 1). The c relaxation has its

origins in the amorphous fraction [20], although it has been proposed that this

relaxation takes place at least in part due to the motion of defects in the crystalline

regions [43] or/and the motion of disordered chain segments at surfaces of polymer

crystals [44]. In the case of PEs, this relaxation can be regarded as a sub-glass

transition attributed to the local motion of the central C–C bond of short chain

segments (by crank-shaft or flip-flop mechanism) and/or local motion of loose chain

ends in the amorphous phase [38]. Khanna et al. [35] have pointed out that this

relaxation involves the motion of a short segment (e.g. three to four CH2) belonging

to the amorphous phase but also the chain ends within the crystalline or amorphous

phases. However, the reported activation energies for the c relaxation are usually

low, i.e. 40–80 kJ/mol for PEs [16, 21] and 25 kJ/mol for iPP [45]. The Arrhenius

temperature dependence is observed for the c relaxation (Fig. 4), and calculated

apparent activation energies for this relaxation in LDPE, HDPE, and iPP are 48, 62,

and 35 kJ/mol, respectively. Dielectric relaxation spectra of virgin iPP also exhibit

the fourth relaxation observed at lowest temperatures (Fig. 1c). In general, the drelaxation is weak or absent and according to Sinnott it is attributed to the hindered

rotation of methyl groups [46]. This relaxation shows Arrhenius behaviour (Fig. 4)

and calculated apparent activation energy is 7 kJ/mol, which is close to the literature

data (5 kJ/mol) [45]. In addition, this relaxation is not visible in dielectric spectra of

irradiated iPP samples (Figs. 2, 3).

By comparing the dielectric loss scans for the virgin (Fig. 1) and irradiated

samples (Figs. 2, 3), it can be observed that radiation introduces significant

qualitative and/or quantitative changes in dielectric relaxation spectra. Different

origin/nature of the dielectric relaxations leads to different evolutions with gamma

radiation. Besides the most obvious changes in the intensity of the dielectric spectra,

radiation also induces disappearance of some dielectric relaxations. There are two or

even more reasons for the ‘‘vanishing’’ of a dielectric relaxation from the spectra.

The first is connected with the origin/nature of the molecular relaxation and the

structural changes introduced by radiation. Thus, the radiation-induced changes in

structure such as chemicrystallization and crosslinking can introduce a decrease in

the amount and/or restriction in the mobility of chain segments that contribute to a

specific relaxation. The second concerns the effect of uneven radiation-induced

oxidative degradation and participation of polar groups in each relaxation. Thus, the

intensity of the dielectric spectra of polyolefins depends on the number of polar

groups formed as a result of radiation, e.g. radiation-induced oxidation. On the other

hand, the participation of polar groups differs among different relaxations and is

closely related to the origin/nature of the relaxation. Namely, there is only a

contribution from the phase in which the relaxation occurs and polar groups that are

connected with the chain segments contributing to the relaxation. Radiation

Polym. Bull. (2014) 71:2317–2334 2325

123

introduces different changes in an amorphous/crystalline phase. In general, the

major effect of irradiation, either electron beam or gamma rays, on the crystalline

regions are some imperfections [47]. The macromolecules in these regions have

very small mobility and oxygen is almost unable to diffuse; diffusion constants for

crystalline regions are small, 8–9 orders of magnitude smaller than in the

amorphous region [10]. Because of that, radiation-induced oxidation takes place

mostly in the amorphous region and at the surface of crystallites.

The disappearance of the dielectric b relaxation in HDPE and low temperature

dielectric c and d relaxations in iPP with gamma radiation is evident from Figs. 2b,

3b, respectively. Complete disappearance of the already weak b relaxation with

irradiation is clearly evident for HDPE even for a low absorbed dose (Fig. 2b).

Since oxidation cannot induce a disappearance of this relaxation [48], an increase in

crystallinity and most probably gel formation are the main suspects (Fig. 6a, d).

This can be expected taking into account the fact that this relaxation is entirely

connected with interlamellar content. Restricted chain mobility in interlamellar

regions as a consequence of crosslinking, together with lower interlamellar content

in irradiated samples, will lead to the disappearance of the b relaxation with

radiation. According to Ratner et al. [49], this transition in HDPE also disappeared

upon the peroxide generated crosslinking in interlamellar regions. For the case of

iPP, the dielectric c relaxation has practically ‘‘vanished’’ with gamma irradiation in

air, as in the case of ultraviolet rays [50]. The dynamic mechanical investigation of

iPP thermo-oxidative degradation has indicated that the initiation of thermal

oxidation is concomitant with a partial vanishing of the c relaxation, too [51]. In

general, it looks that oxidative degradation of iPP structure plays a critical role in

the disappearance of the low temperature dielectric relaxations [7] but also puts a

significant accent on the high temperature a relaxation [14]. From Fig. 1c, it is

evident that the intensities of the dielectric a and b relaxations at a frequency of

1 MHz are similar in virgin iPP. On the other hand, a much larger intensity of the arelaxation can clearly be observed in dielectric spectra of irradiated samples at the

same frequency (Figs. 2c, 3c). Because of the nature of the dielectric a relaxation

and the fact that the radiation-induced oxidation takes place mostly in the

amorphous region and on boundary layers, it can be concluded that there is a

significant contribution of boundary layers between the amorphous phase and

crystallites to the dielectric a relaxation in iPP. It can also be concluded that the

radiation-induced oxidative degradation in iPP occurs greatly on boundary layers

between the amorphous and crystalline phase.

The intensity of the dielectric spectra of PEs is mainly determined by the

carbonyl groups [52], while in the case of PP additional contribution of

hydroperoxides to dielectric spectra cannot be neglected [7, 14, 53]. Dipolar

moment values and/or radiation-induced increase in concentration of other polar

groups are much lower. The quantification of various oxidation products of

polyolefins (as a result of gamma, photo, and thermal oxidation) was done by

Lacoste et al. [54]. The radiation-induced modifications that occur in the carbonyl

and hydroxyl region were investigated in our previous papers [7, 14, 16], too. The

carbonyl groups in PEs are mainly ketones (1,725–1,715 cm21) and aldehydes

(1,720–1,730 cm21); they are responsible for the absorption maxima at 1,720 cm21

2326 Polym. Bull. (2014) 71:2317–2334

123

(insert in Fig. 5b). Together with them, carboxylic acids (1,718–1,710 cm21) which

are responsible for the small shift of absorption maxima to 1,715 cm21 are the main

radiation-induced oxidation products observed in the carbonyl region of iPP (insert

in Fig. 5c). The evolution of carbonyl content (through the normalized absorbance

at 1,720 and 1,715 cm21) with absorbed dose is presented in Fig. 5. A linear

dependence of the carbonyl content is evident for lower doses, while for higher ones

([200 kGy) an intense deviation (saturation) from linear curve occurs for HDPE

and especially iPP. Probably the real reason for the saturation in carbonyl contents is

the fact that the oxygen present in the bulk and consumed due to radiation-induced

reactions is not supplied fast enough by diffusion at higher doses. Apparently, the

radiation-induced oxidation is limited by insufficient diffusion rate of oxygen and

0 250 5000

1

2

3

4A

1720

/d [1

03m

-1]

Absorbed dose [kGy]

LDPE0

1

2

3

HDPE

A17

20/d

[103

m-1

]

A17

15/d

[103

m-1

]

0

1

2

3

4

0 300 600

Absorbed dose [kGy]0 300 600

Absorbed dose [kGy]

iPP(a) (b) (c)ta

nδ

[10-4

]

αmax

max

max

0

10

20

30

40

50

LDPE

0 300 600

Absorbed dose [kGy]

0

10

20

30

40

αmax

max

HDPE

0 300 600

Absorbed dose [kGy]

0 300 600

Absorbed dose [kGy]

tan

δ[1

0-4]

tan

δ[1

0-4]

0

25

50

75

100

105 Hz106 Hz

αmax

iPP

Wavenum. [cm-1]

0 kGy

T [

%]

40

60

80

200 kGy

1900 1800 1700 1600

20

T [

%]

2000 1750 1500

30

60

90

200 kGy

Wavenum. [cm-1]

0 kGy

105 Hz 105 Hz(d) (e) (f)

Fig. 5 Normalized IR absorption intensity (A/d values; A = absorbance; d = sample thickness) incarbonyl region (at 1,720 and 1,715 cm-1 for PEs and iPP, respectively) versus absorbed dose for LDPE(a), HDPE (b), and iPP (c). Shown by the insert are the IR spectra in the carbonyl and vinylidene regionsfor virgin and irradiated (200 kGy) samples; the intensity of the dielectric relaxations versus absorbeddose for LDPE (a), HDPE (b), and iPP (c)

Polym. Bull. (2014) 71:2317–2334 2327

123

the accessibility of free radicals to atmospheric oxygen [55]. Post-radiation

annealing also plays a significant role and contributes to the saturation in carbonyl

and hydroperoxide contents. Annealing at elevated temperatures introduces a much

higher rate of thermal recombination of free radicals than the rate of oxygen

diffusion from the surface into the film; this effect is more pronounced for the higher

absorbed doses, i.e. for the higher concentrations of free radicals and in the samples

with higher crystallinity, since the free radicals trapped in the crystalline area are the

main cause of the post-radiation oxidation. Additionally, the saturation in the

carbonyl content for the iPP with absorbed dose coincides with the start of gelation

(Fig. 6a). Gavrila and Gosse [56] have also found for the iPP gamma-irradiated in

air that the amount of carbonyl groups declines sharply at the gel point. Domination

of chain scission reactions at low irradiation doses in iPP can be confirmed from

FTIR results (insert in Fig. 5c) where the increase in vinylidene unsaturated groups

(at 1,640 cm21), which are the products of chain scission reactions, is clearly

evident. In our previous study, we have discussed in detail the evolution of

vinylidene unsaturated groups as a function of absorbed dose. It has been revealed

that vinylidene concentration increases significantly at lower doses (B250 kGy).

With further increase in absorbed dose crosslinking reactions become dominant,

while saturation and decay in vinylidene content were observed [13]. In addition,

Veselovskii et al. [57] have reported that the net rate of vinylidene formation in

irradiated PP declines sharply at the gel point and suggested that a sudden increase

in vinylidene consumption is clearly associated with gelation. Intensive crosslinking

behaviour and the formation of net structure can have some but probably not a

decisive influence on the generation of oxidizing species. On the basis of a

comparison of IR spectroscopic and dielectric measurements, similar radiation-

induced evolution in the concentration of carbonyl groups (Fig. 5a–c) and in the

intensity (dielectric loss tangent maxima) of the relaxations is observed (Fig. 5d–f).

Despite good agreement, some differences between the intensities of the dielectric

relaxations and IR data are evident (Fig. 5). For the case of iPP, the deviation in the

intensity of the dielectric relaxations from linear dependence is even more

emphasized at higher absorbed doses than is the case with carbonyl content (Fig. 5c,

f). The larger deviation from the linear dependence in the intensity of the dielectric

relaxations than in the carbonyl content can probably be explained by a significant

additional contribution of hydroperoxides; the concentration of hydroperoxides in

iPP first increases for low absorbed doses (B100 kGy) and then levels off for

median doses and starts decaying with further increase in the radiation dose [14].

Besides the changes in the relaxation intensity, radiation also induces changes in

distribution of relaxation times, peak position, and activation energy. The variations

in the position and apparent activation energy of the dielectric b relaxation with

absorbed dose are shown in Fig. 6b, c, respectively. The increase in temperature and

the apparent activation energy for this relaxation in LDPE can be related to the

changed chain mobility in the amorphous phase induced by crosslinking (Fig. 6a).

On the other hand, the position of the dielectric b relaxation in iPP is slightly shifted

to lower temperatures at low irradiation doses (B200 kGy); this shift, together with

the decrease in the apparent activation energy and the dynamic fragility (insert in

Fig. 6b), can be attributed to the domination of chain scission reactions. For higher

2328 Polym. Bull. (2014) 71:2317–2334

123

doses at which the crosslinking reactions become dominant (Fig. 6a), a recovery of

temperature, dynamic fragility, and the apparent activation energy for this

relaxation in iPP is more than evident (Fig. 6b, c).

The variations in the position and apparent activation energy of the dielectric arelaxation with absorbed dose are shown in Fig. 6e, f, respectively. The shift in the

position and the increase in activation energy of the dielectric a relaxation in HDPE

at low doses (D B 200 kGy) are due to more intensive oxidative degradation and

radiation-induced changes connected with the crystalline phase. Namely, radiation-

induced breakage of macromolecules gives rise to a summation effect connected

with the crystalline phase: additional crystallization and increased perfection, as

well as the size of the crystallite [58]. At higher doses, lower crystallinity and

crystallite thickness were compensated by a high crosslinking efficiency that

influences the position and activation energy of the a relaxation in the same manner.

0 300 600

Gel

[%]

Absorbed dose [kGy]0 300 600

Absorbed dose [kGy]0 300 600

Absorbed dose [kGy]

0 300 600

Absorbed dose [kGy]0 300 600

Absorbed dose [kGy]

0 300 600

Absorbed dose [kGy]

0

25

50

75

LDPEHDPE

iPP

Τ β[K

]

260

280

300

320

Ε a[k

J/m

ol]

15

30

45

60

75

Cry

stal

linit

y [%

]

100

120

140

160

180

Ε a[k

J/m

ol]

HDPEiPP

340

360

380

αmax

Τ α[K

]

150

175

200

225400

500

600

LDPEiPP

(a) (b) (c)

LDPE

iPP

relaxation

max

105 Hz

m

0 300 60080

100

120

Absorbed dose [kGy]

LDPEHDPE

iPP

LDPEHDPE

iPP

(d) (e) (f)

Fig. 6 a Gel content versus absorbed dose; b dielectric relaxation loss maxima; and c apparent activationenergy (Ea) for the dielectric b processes versus absorbed dose. Shown by the insert in (b) is the dynamicfragility m as a function of absorbed dose; d Crystallinity versus absorbed dose; dielectric relaxation lossmaxima (e); and apparent activation energy (f) for the dielectric a processes versus absorbed dose

Polym. Bull. (2014) 71:2317–2334 2329

123

According to Danch and Osoba [59], different processes, i.e. annealing, drawing, or

irradiation, can restrict the mobility of the chains involved in the a relaxation,

introducing a shift in its position and increase in activation energy. The changes

observed in the position and apparent activation energy with irradiation are small

for LDPE (not presented) and correspond to relatively small changes in the

crystalline phase (Fig. 6a).

Contrary to this, the explanation for the large radiation-induced shift in the

position and increase in activation energy of the dielectric a relaxation is not so

simple for iPP. This shift is most intensive for lower doses (B100 kGy) at which

the gel content is zero and oxidative degradation dominates (Fig. 6a, d). Thus, it

is not possible to make the correlation between the shift in the position of the

dielectric a relaxation and the crosslinking. Furthermore, WAXD data have

indicated that the radiation-induced changes in crystallinity and crystal size can

have some but probably not a decisive influence on the position of the arelaxation. The crystallinity and crystal size for highly irradiated samples are

smaller than those for the unirradiated one (Fig. 6d) [14], but the dielectric arelaxation still occurs at much higher temperatures (Fig. 6e) and has a much

higher apparent activation energy (Fig. 6f). The most probable explanation for

the observed shift with radiation is connected with the complex and multiple

nature of this relaxation. A study of multiple dielectric a peaks in experimental

spectra of virgin, annealed and irradiated iPP samples is presented in our previous

paper [14]. The shift in the position and increase in apparent activation energy of

the dielectric a relaxation to higher temperatures can presumably be explained by

the prevalence of high temperature components (characterized by higher

activation energies) in this relaxation due to the radiation-induced oxidative

degradation [14]. Deviation from such behaviour is evident at higher doses at

which crosslinking dominates over the oxidative degradation and a significant

level of net formation is achieved. The changes in activation energies with

absorbed dose are relatively similar to those observed for the position of the

dielectric a relaxation (Fig. 6e, f).

Conclusions

Dielectric relaxation spectroscopy and gamma radiation were used as powerful

methods for characterization and modification of polyolefins, respectively. Radi-

ation introduces significant qualitative and quantitative changes in dielectric

relaxation spectra of LDPE, HDPE, and iPP. The radiation-induced oxidation of

apolar polyolefins causes a significant increase in the magnitude of the dielectric

spectra due to the increase in polymer polarity, e.g. amount of carbonyl,

hydroperoxide, and other polar groups. On the other hand, the participation of

polar groups differs among different relaxations and is closely related to the origin/

nature of the relaxation. The intensity of the dielectric spectra of PEs is mainly

determined by the carbonyl groups, while in the case of PP additional contribution

of hydroperoxides to dielectric spectra cannot be neglected. Besides the changes in

the relaxation intensity, radiation also induces disappearance of some relaxations.

2330 Polym. Bull. (2014) 71:2317–2334

123

For the case of HDPE, restricted chain mobility in interlamellar regions as a

consequence of crosslinking, together with lower interlamellar content in irradiated

samples, will lead to the disappearance of the already weak b relaxation with

gamma radiation. On the other hand, it looks that a large oxidative degradation of

iPP structure plays critical role in the disappearance of the low temperature c and ddielectric relaxations with gamma radiation.

Radiation also induces changes in the distribution of relaxation times, peak

position, and activation energy of some dielectric relaxations. Different origin/

nature of the dielectric relaxations leads to different evolutions with gamma

radiation. Radiation-induced changes in the amorphous phase can be related to the

evolution of the dielectric b relaxation. The increase in the temperature and

apparent activation energy for this relaxation in LDPE can be related to the

restricted chain mobility in the amorphous phase, induced by crosslinking and net

formation. The VFTH dependence observed for the b relaxation in iPP confirms the

hypothesis that this relaxation is clearly related to the glass transition in this

polyolefin. A decrease in temperature, dynamic fragility, and the apparent activation

energy at lower irradiation doses followed by their recovery at higher ones was

observed for this relaxation in iPP. This decrease can be attributed to the large

oxidative degradation and predominance of chain scission reactions. On the other

hand, the recovery at higher doses can be connected with the prevalence of

crosslinking and net formation.

Radiation-induced changes in the crystalline phase can be related to the

evolution of the dielectric a relaxation, but the contribution of the amorphous

phase in the neighbourhood of the crystallites to this relaxation should be taken

into account especially for iPP. Due to a large increase in magnitude of the

dielectric a relaxation, it can be concluded that the radiation-induced oxidative

degradation in iPP occurs greatly on boundary layers between the amorphous and

crystalline phase. The shift in the position and increase in activation energy of the

dielectric a relaxation in HDPE are due to more intensive oxidative degradation

and radiation-induced changes connected with the crystalline phase. At higher

doses, lower crystallinity and crystallite thickness were compensated by a high

crosslinking efficiency which influences the position and activation energy of the

a relaxation in the same manner. Contrary to this, the explanation for the large

radiation-induced shift in the position and increase in activation energy of the

dielectric a relaxation is much more complicated for iPP. Radiation-induced

changes in crystallinity and crystal size have some influence on the position and

activation energy of the dielectric a relaxation, but the observed behaviour of this

relaxation cannot be explained only by this. The complex and multiple nature of

this relaxation in iPP gives an additional explanation for the observed changes; it

can be found in the prevalence of high temperature relaxation components

(probably characterized by higher activation energies) in the dielectric a relaxation

zone of irradiated iPP.

Acknowledgments This work has been supported by the Ministry of Education, Science and

Technological development of the Republic of Serbia (Grant No. 172026).

Polym. Bull. (2014) 71:2317–2334 2331

123

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