HA composites obtained in solution: Effect of the gamma radiation

NIM B Beam Interactions with Materials & Atoms

Nuclear Instruments and Methods in Physics Research B 247 (2006) 331–341 www.elsevier.com/locate/nimb

HDPE/HA composites obtained in solution: Effect of the gamma radiation Albano Carmen a,b,*, Karam Arquı´medes a, Perera Rosestela c, Gonza´lez Gema a, Domı´nguez Nohemy a, Gonza´lez Jeanette c, Sa´nchez Yanixia a a

Polymer Laboratory, Chemistry Center, and Materials Laboratory, Technology Center, Venezuelan Institute for Scientific Research (IVIC), P.O. Box 21827, Caracas 1020-A, Venezuela b Universidad Central de Venezuela, Facultad de Ingenierı´a, Escuela de Ingenierı´a Quı´mica, Caracas, Venezuela c Universidad Simo´n Bolı´var, Departamento de Meca´nica, Caracas, Venezuela Received 24 October 2005; received in revised form 23 February 2006 Available online 2 May 2006

Abstract Radiation is employed to sterilize composite materials used in the biomedical field. Due to the changes induced by radiation onto polymeric materials, it is important to study variations in their melt flow index (MFI), as well as in their mechanical and thermal properties. In this work, those previous parameters were determined in composites obtained via solution of a high-density polyethylene (HDPE) in decalin, with different amounts of hydroxyapatite (HA), varying from 10 to 30 parts per hundred, after being exposed to gamma radiation at absorbed doses between 25 and 100 kGy. After the irradiation, the MFI of HDPE dissolved in decalin and precipitated afterwards and without filler increased from 6 to 24 g/10 min at the highest absorbed doses. This behavior was also observed in composites with 10 pph of HA, being the increase less pronounced, specifically in the range between 50 and 100 kGy. Composites with 20 and 30 pph of HA showed a maximum MFI value at 50 kGy, which decreased at higher doses. This implies that the filler begin to exert an influence because it does not melt at the test temperature and consequently, it does not flow. It was observed that Young’s modulus increased with HA addition due to rigidity of the ceramic filler. Radiation did not significantly affect this tensile property. On the other hand, the tensile strength did not show significant variations at the different doses but the filler content did affect this property improving it. Finally, elongation at break showed a drastic decrease with filler addition. When the thermal behavior was studied it was noticed that crystallization and melting temperatures remained unchanged. Instead, crystallinity degree slightly increased in the composites, and a little decrease was obtained when they were irradiated. Ó 2006 Elsevier B.V. All rights reserved. PACS: 61.80.x; 81.05.Pj Keywords: HDPE/HA composites; Radiation; Mechanical behavior; Thermal behavior; Thermodegradative behavior

1. Introduction In the last years, composites made of polymers and ceramic fillers represent a new class of materials of high interest. * Corresponding author. Address: Polymer Laboratory, Chemistry Center, and Materials Laboratory, Technology Center, Venezuelan Institute for Scientific Research (IVIC), P.O. Box 21827, Caracas 1020-A, Venezuela. Tel.: +58 212 504 1636; fax: +58 212 504 1350. E-mail addresses: [email protected] (A. Carmen), akaram@quimica. ivic.ve (K. Arquı´medes).

0168-583X/$ - see front matter Ó 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.nimb.2006.03.004

High-density polyethylene (HDPE) usually shows low toxicity and has been recommended as a suitable material for bone tissue substitution [1,2]. On the other hand, hydroxyapatite is a biocompatible ceramic that is present in natural bones. Consequently, composite materials based on PE filled with HA have been studied by some authors. They concluded that these materials showed suitable hardness and rigidity, as well as high biocompatibility for tissue replacement [3,4]. The interfacial adhesion between the polymer and HA is one of the important factors governing the mechanical

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behavior of these composites. The control of the interfacial interactions is a difficult task. One of the methods employed is the use of coupling agents. Sousa et al. [5] studied HDPE–HA composites containing zirconate and titanate coupling agents. However, these additives may have detrimental effects on the biocompatibility properties of the material. Additionally, based on the potential that those materials have in the biomedical area, it is important to get an adequate dispersion of the filler (HA) into the polymer matrix because this drastically influences the mechanical behavior, which in turn will define the application of the final product [6]. Hence, the objective of this study was to obtain composites of HDPE–HA prepared via solution and to study their thermodegradative, mechanical, thermal and morphological behavior when subjected to irradiation. 2. Experimental 2.1. Materials In this research, a commercial high-density polyethylene (HDPE) with a melt flow index of 7.8 g/10 min supplied by Coramer C.A. was employed. Hydroxyapatite was synthesized via precipitation methods at room temperature [7]. The synthesis was carried out from an aqueous suspension of calcium hydroxide (Ca(OH)2) in a solution of di-ammonium hydrogen phosphate ((NH4)2HPO4). The products were thoroughly washed to neutral pH and dried at 75 °C for 48 h. The powdered HA was characterized using Xray diffraction (XRD) in a Siemens D5005 diffractometer using CuKa radiation and a Ni filter, at 40 kV and 20 mA. 2.2. Preparation and irradiation of the composites The HDPE/HA composites were prepared in a decalin (cis and trans decahydronaphtalene) solution. To do so, the polymer was dissolved in it at 120 °C under continuous stirring. Afterwards, the HA was added maintaining the agitation for 20 more minutes in order to get a good dispersion of the filler. After this time, ethanol was poured into the suspension to precipitate the composite, which was then filtered and washed with n-hexane to eliminate the remaining decalin. The same procedure was repeated without adding the filler into the polymeric solution, for comparison purposes. Different amounts of HA were used, i.e. 10, 20 and 30 parts per hundred (pph). Samples of each composite were compression molded and cut following the ASTM-638 standard. Specimens and powdered composites were irradiated with c-rays from a 60 cobalt source in air at a dose rate of 4.8 kGy/h. Integral doses were 25, 50 and 100 kGy. Non-irradiated samples were also tested for comparison purposes. 2.3. Characterization FTIR spectra were obtained using a Nicolet Magna-IR 750 spectrometer, to characterize the HA and to ascertain

the possibility of polymer degradation during the irradiation. Thus, powdered HA and compression molded films of irradiated and non-irradiated HDPE were analyzed in the range 4000–400 cm1 at a resolution of 4 cm1 after 32 scans. Melt flow index (MFI) measurements were taken from the irradiated and non-irradiated composites and from the powdered HDPE. A Ray Ran Advanced Melt Flow Systems with a weight of 2.16 kg was used at 190 °C. The mechanical properties of irradiated and non-irradiated HDPE and HDPE–HA composites were determined in an Instron Tensile Tester Machine (model 4204). Each experimental result reported represents an average of at least seven samples tested under identical conditions. Differential scanning calorimetry was also used to study the crystallization and melting behavior of the materials using a Metler Toledo DSC 821. Two samples of about 10 mg of each composite were tested under a pure nitrogen atmosphere. The sample was first heated up to 190 °C at 20 °C/min and kept at that temperature for 3 min. Then, it was cooled down to 25 °C and melted again afterwards. The first cooling and the second heating thermograms (after a common thermal history) were recorded at 10 °C/min. Moreover, the composites were analyzed using the successive self-nucleation and annealing technique (SSA) developed by Mu¨ller et al. [8]. That technique was designed to produce molecular segregation during self-nucleation, crystallization and annealing induced by a specific thermal treatment using DSC. In this technique, an initial standard thermal history is created melting the sample at 170 °C for 3 min and cooling it down to 25 °C at 10 °C/min. Afterwards, a heating scan at 10 °C/min is performed up to a selected self-seeding and annealing temperature denoted Ts, where the sample is isothermally kept for 5 min before cooling it down to 25 °C at 10 °C/min. Then, the sample is heated again up to a temperature which is 5 °C lower than the previous Ts, at 10 °C/min and held at this temperature for 5 more minutes, before cooling it down to 25 °C at 10 °C/min. This procedure is repeated while the temperature is being lowered at 5 °C intervals with respect to the previous step as many times as needed to cover the temperature range between the first chosen Ts and 25 °C. Finally, the melting behavior of the sample was recorded at 10 °C/ min when the conditioning was over. Thermogravimetric studies were carried out in a Thermogravimetric Analyzer (Mettler Toledo TGA851) at a heating rate of 10 °C/min, at temperatures up to a 700 °C under a pure nitrogen atmosphere. From the thermograms obtained in the TGA, the kinetic parameters such as the activation energy (Ea) and initial decomposition temperature (Tid) were determined. The values of Ea and Tid indicate the thermal stability of the materials, that is, if those values are high, the thermal stability will be high and the degradation process will be thus delayed. The activation energy was determined using the equation of McCallum– Tanner and the experimental data obtained in the TGA. The equation of McCallum–Tanner [9] is as follows:

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logðGðaÞÞ ¼ log 

AEa  0:483E0:435 a bR

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3. Results and discussion

ð0:449 þ 0:217Ea Þ103 ; T

ð1Þ

1n

where, GðaÞ ¼ 1ð1aÞ for all values of n 5 1 and 1n G(a) = ln(1  a) for n = 1. Where b is the rate of heating (degree/time), T is the temperature (K), a is the reacted fraction at the time t, Ea is the activation energy (kJ/mol), A is the pre-exponential factor and R is the gas constant. While, w the conversion of the sample is equal to a ¼ ww00w . Where w0 f is initial mass of the sample (g), w is the mass of the sample to a given temperature (g) and wf is the final mass of the sample (g). The initial decomposition temperatures were determined by the intersection of the tangent lines in the decomposition area of the thermogram. The molecular weights and molecular weight distributions of the HDPE before and after irradiation were determined through gel permeation chromatography (GPC) in 1,2,4-trichlorobenzene at 135 °C with 2,6 di-tert-butyl4-methylphenol added as antioxidant, using a Waters Alliance 2000 equipment with a differential refractive index detector. Three columns of Styragel HT 6E, HT5 and HT3 were employed. The calibration curve was obtained from polystyrene standards of narrow molecular weight distributions. Mw, Mn, and the polydispersity were recorded. Transmission electron microscopy (TEM) was carried out in a Phillips CM10 operating at 80 kV and used to analyze the microstructure of the composites. The samples were prepared by ultramicrotomy and observed without any further treatment.

The synthesized HA was characterized through XRD and FTIR techniques and the results confirmed that its structure coincides with that reported in the literature for HA [7]. Since the main objective of this work was to study the effects that gamma radiation has on the HDPE–HA composites, FTIR spectra were recorded from samples of non-irradiated and irradiated HA at different absorbed doses. Fig. 1 shows the spectrum of non-irradiated HA (HA-0). A band at 3427 cm1 attributed to the stretching of the hydroxyl group is noticed. Such band is slightly displaced from that of commercial HA, which appears at 3569 cm1 and could be a consequence of water absorption. The presence of carbonate ions is ascertained in the 1650 and 1300 cm1 region, whereas the phosphates are noticed in bands at 1046, 961, 660 and 520 cm1. A weak band at 472 cm1 is also detected. The spectrum is similar to that obtained by Rehman and Bonfield [10]. Fig. 1 also shows the spectra of irradiated HA at different absorbed doses. Those spectra are similar to that of non-irradiated HA, which leads to the conclusion that the irradiation process has no effect on the chemical structure of the HA. XRD results confirmed such fact. Since the blending of HDPE was carried out in solution using decalin, the effects that this solvent may produce in the HDPE were analyzed. Fig. 2(a) displays the MFI values of the HDPE and its composites after being irradiated. A sharp decrease in the MFI values is obtained when the polyethylene without decalin (HDPEwd) is irradiated, due to its crosslinking and/or long-chain branching. Nevertheless, an increase in the MFI values with the absorbed doses is obtained when HDPE is dissolved in decalin

HA-0 HA- 25

HA-50 HA-100

3500

3000

2500

2000

1500

1000

Wavenumbers (cm-1) Fig. 1. Infrared spectra of irradiated hydroxyapatite at different doses.

500

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Fig. 2. Melt flow index of irradiated samples: (a) effect of decalin in HDPE; (b) HDPE/HA composites.

(HDPEd), then precipitated and irradiated afterwards. Fig. 3 is used to help explain this behavior. Such figure shows the FTIR spectra of irradiated HDPEd. In the spectra, a band in the region of 1700–1750 cm1 appears after irradiation and is attributed to the carbonyl groups which are generated as a consequence of the oxidation of the polymer. This degradation process produces chain scissions that lower the polymer’s molecular weight thus decreasing its viscosity. The above mentioned band increases with the radiation dose. When decalin is used as a solvent, it may be kept occluded into the polymer and this fact produce a high radical concentration (determined by EPR [11]) when the sample is irradiated. If those radicals do not react among themselves, they produce chain scissions. Silva et al. [11] found that when decalin is present in the HDPE, high absorbed doses favors the formation of Dec+ radicals, which transfer the radical to the polymer, producing high radical concentrations. Then, chain scissions reactions take place extensively instead of crosslinking, lowering the molecular weight and viscosity of the polymer and consequently increasing its MFI.

To confirm this last fact, GPC measurements were taken. Figs. 4 and 5 show the molecular weight distribution curves of both HDPEwd and HDPEd. First of all, it can be seen in the figures that both samples in their non-irradiated state are similar, even when a slight decrease in Mw can be appreciated in HDPEd Mw value (Table 1). This indicates that decalin treatment is affecting the polymer by degrading it. In the same table it is also seen that, when the HDPEwd is irradiated at different absorbed doses (25 and 100 kGy), an increase in its average molecular weight (Mw) is produced, confirming that crosslinking or branching reactions are taken place as a consequence of the radiation process. At 100 kGy, a decrease in Mn was also detected, probably due to the fact that at higher doses chain scission reactions are also occurring, forming lower molecular weight fragments, thus increasing the polydispersity. This can also be appreciated in Fig. 4, in which a shoulder in the low molecular weights region appears. In the same figure, another shoulder at higher molecular weight is seen, indicating the presence of high molecular weight species.

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335

HDPEd - 0 HDPEd - 25 HDPEd - 50 HDPEd - 100

3500

3000

2500

a

2000

1500

1000

500

Wavenumber (cm-1)

HDPEd - 0 HDPEd - 25 HDPEd - 50 HDPEd - 100

1745

1735

1725

b

1715

1705

1695

Wavenumber (cm-1) Fig. 3. (a) Infrared spectra of irradiated HDPE; (b) zoom of 1700 cm1 area.

0 KGy 25 KGy 100 KGy dw/d(logMw)

dw/d(logMw)

1

0.5

0 3.0

1

0.5

0 3.0

4.0

5.0 (log Mw)

6.0

0 KGy 25 KGy 50 KGy 100 KGy

4.0

5.0

6.0

log Mw

Fig. 5. Effect of gamma radiation in decalin treated HDPE (HDPEd).

Fig. 4. Radiation effect on HDPEwd.

On the other hand, the molecular weights of HDPEd samples decrease when irradiated. As the absorbed dose is increased, the molecular weights decrease (Table 1). When their chromatograms are analyzed (Fig. 5), it is seen

that curves are displaced towards lower molecular weights when the absorbed doses increase, thus indicating a higher amount of low-molecular weight chains. In this case, traces of decalin make the chain scissions mechanism, especially

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Table 1 Average molecular weight of irradiated polyethylene Mn (g/mol) ± 10%

Radiation dose (kGy)

0 25 50 100

Mw (g/mol) ± 10%

HDPEwd

HDPEd

HDPEwd

HDPEd

HDPEwd

HDPEd

25 252 26 907 – 20 834

26 611 22 892 18 173 15 122

67 097 76 617 – 89 257

61 413 49 362 43 003 34 686

1.66 2.85 – 4.28

2.31 2.16 2.37 2.29

in longer molecules, the one dominating when the samples are irradiated. From the results above, it can be said that when traces of decalin are present, they favor the chain scission degradation mechanism in irradiated HDPE, whereas crosslinking and/or long-chain branching is produced by the same type of irradiation in HDPE when decalin is absent. Albano et al. [12] also reported crosslinking and/or chain branching in HDPE irradiated with c rays. Fig. 2(b) shows the melt flow index values of the composites made out of HDPEd and different amounts of hydroxyapatite (10–30 pph) irradiated at different absorbed doses (0–100 kGy). In non-irradiated composites, the melt flow index values tend to decrease as the amount of the HA in increased, due to the increase in their viscosities as a consequence of the flow resistance of the unmelted ceramic filler. On the contrary, the increase of the absorbed dose up to 50 kGy led to an increase of the composites MFI values. Nevertheless, at 100 kGy of absorbed dose, this behavior is not observed in the HDPEd/HA composites, fact that may be attributed to the presence of the filler. At the highest dose (100 kGy), the HA at concentrations equal or higher than 20 pph could be protecting the polyethylene against the radiation and hence, the above mentioned increase of the MFI was not as sharp as that in composites with the lowest amount of the filler. Table 2 displays the Young’s modulus values as a function of the amount of HA for each one of the absorbed doses. A slight increase in this tensile property at 0 kGy is noticed when HA content increase, evidencing a reinforcing effect exerted by the filler. When the radiation effect was studied, it was noticed that radiation did not modify in a significant way the composite’s Young’s modulus. This phenomena was also reported by Dole [13], who explained that radiation effect in doses up to 1000 kGy are relatively unimportant, because Young’s modulus and yield stress are

Table 2 Effect of radiation in composites Young’s modulus (E (MPa)) Sample

HDPEd PEHA 10 PEHA 20 PEHA 30

Polydispersity ± 10%

Radiation dose 0 kGy

25 kGy

50 kGy

100 kGy

1108 ± 72 1141 ± 56 1384 ± 108 1315 ± 20

1338 ± 130 1337 ± 149 1446 ± 104 1298 ± 189

1179 ± 81 1172 ± 90 1466 ± 86 1518 ± 328

1320 ± 98 1192 ± 135 1431 ± 161 1324 ± 220

strongly dependent on the crystallinity; so radiation induced crosslinks or chain scissions are minor perturbations. Only at very high levels of radiation (over 2000 kGy) these effects become important. Fig. 6 shows the composite’s morphology. In general, an excellent dispersion is achieved at low HA contents (Fig. 6(a)). However, at higher filler concentrations (30 pph, Fig. 6(b)) some agglomerates of the filler are detected, which in turn can determine the tensile properties hereby analyzed. Changes in the tensile strength of the composites with the absorbed dose and filler content are reported in Fig. 7. As it can be seen, the addition of the HA increases this mechanical parameter at all the absorbed doses. For each absorbed dose, the toughness of the polymer and the rigidity of the reinforcing filler (HA) are combined, thus rendering an improvement in the mechanical properties of the composites [14]. In general, the tensile strength increases with the addition of the HA in concentrations up to 20 pph. At 30 pph of the filler, the average tensile strength is lower than that of the composites with 10 and 20 pph of filler, probably due to the presence of agglomerates (Fig. 6(b)) which act as defects that may initiate cracks and hence the rupture of the material [3]. In the same figure, it is also seen that the effects of the absorbed doses are not significant on the strength of the different composites, which implies that in this case, the filler content and their crystallinity degree prevails determining this mechanical parameter. Table 3 shows the elongation at break of the composites subjected to irradiation. In unfilled HDPEd, this parameter sharply decreases as the absorbed dose increases, due to the decrease in the polymer’s molecular weight. When the filler is incorporated into the polymer, a sharp decrease in the elongation at break is also obtained. This behavior is attributed to the facts that the filler is a brittle material that has a low interfacial adhesion to the polymeric matrix where it is embedded in [15]. As there is not a good interaction between the filler particles and the polyethylene, their interface acts as a stress concentrator where the sample fractures and becomes fragile. For the composites, when the effects of different absorbed doses are compared (Table 3), it can be seen that there is not a significant change in the elongation at break with the absorbed dose. Some theoretical models were used to analyze the tensile data. The Young’s modulus was predicted using the Guth– Smallwood equation [16,17]:

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337

Fig. 6. TEM micrographs of HDPE/HA blends: (a) PEHA20 composite; (b) PEHA30 composite.

40 35

Tensile strength, MPa

30 25 HDPEd PEHA10 PEHA20 PEHA30

20 15 10 5 0

0

25

50

100

Integral dose, kGy Fig. 7. Tensile strength of irradiated composites.

Table 3 Elongation at break of composites exposed at different radiation doses Sample

HDPEd PEHA 10 PEHA 20 PEHA 30

eb (%) 0 kGy

25 kGy

50 kGy

100 kGy

249 ± 92 11.9 ± 1.5 6.9 ± 1.1 5.4 ± 0.8

83.0 ± 37.6 8.5 ± 1.2 7.1 ± 0.7 5.3 ± 0.8

85.6 ± 30.0 10.7 ± 1.1 5.9 ± 0.5 5.1 ± 0.4

27.3 ± 5.3 10.1 ± 1.6 6.3 ± 1.2 4.8 ± 0.6

Ec ¼ 1 þ 2:5Uf þ 14:1U2f ; Em

ð2Þ

where Ec is the Young’s modulus of the composite (MPa), Em is the Young’s modulus of the matrix (MPa) and Uf is the filler volume fraction. To calculate Uf, the values of 0.94 g/cm3 as the HDPE’s density [18] and 3.16 g/cm3 as the density of the HA were used [1,19]. Furthermore, to predict the theoretical values of the tensile strength, the equation of Nicolais–Narkis was employed [20]: rc 2=3 ¼ ð1  1:21Uf Þ; rm

ð3Þ

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where rc is the composite’s tensile strength (MPa) and rm is the matrix tensile strength (MPa). On the other hand, the equation of Nielsen [21] was used to predict the values of ec/em and to compare them to the experimental ones. Such equation indicates the existence of perfect adhesion, because the fracture tends to proceed from particle to particle. The equation is ec 1=3 ¼ ð1  Uf Þ; ð4Þ em where ec is the elongation at break of the composite (%) and em is the elongation at break of the matrix (%). Tables 4–6 display the dependence of the Young’s modulus, tensile strength and elongation at break values (Ec, rc, ec) of the composites normalized to the respective values of the Young’s modulus, tensile strength and elongation at break values of the matrix on the weight fraction of the filler and on the absorbed dose. When the results presented in Table 4 are analyzed, it can be concluded that the values of Ec/Em of the unirradiated composites obtained through Eq. (2) are reasonably well related to the experimental results, indicating that the incorporation of HA increases the rigidity of the HDPE, through the restriction in the mobility of the polymer molecules. On the other hand, the radiation of the composites has a random influence on these results, that is, they do not follow a certain trend. As it has already been mentioned, the radiation process produced a predominance of chain scission reactions in these composites obtained via solution, which, along with the presence of agglomerates at filler contents above 20 ppc and the HDPE high crystallinity degree, strongly determined the mechanical behavior of these materials. For this reason, the theoretically values predicted using the equation of Guth–Smallwood are not well correlated to the experimental data, owing to the fact

Table 4 Comparison of experimental data with the theoretical predictions from the Guth–Smallwood equation for different radiation doses Composites

Ec/Em theoretical

Ec/Em experimental for different doses (kGy) 0

25

50

100

PEHA 10 PEHA 20 PEHA 30

1.0840 1.1848 1.2995

1.0298 1.2491 1.1868

0.9993 1.0807 0.9701

0.9941 1.2434 1.2875

0.9030 1.0841 1.0030

Table 5 Comparison of experimental data with the theoretical predictions from the Nicolais–Narkis equation for different radiation doses Composites

PEHA 10 PEHA 20 PEHA 30

rc/rm theoretical

0.8861 0.8226 0.7717

rc/rm experimental for different doses (kGy) 0

25

50

100

2.2949 2.3294 1.9721

2.9725 3.2018 2.4217

2.3113 2.3889 2.1407

1.9608 2.0491 1.6601

Table 6 Comparison of experimental data with the theoretical predictions from the Nielsen’s simple model for different radiation doses Composites

PEHA 10 PEHA 20 PEHA 30

ec/em theoretical

0.6932 0.6171 0.5657

ec/em experimental for different doses (kGy) 0

25

50

100

0.0478 0.0277 0.0217

0.1024 0.0855 0.0639

0.1250 0.0689 0.0596

0.3700 0.2308 0.1758

that some assumptions were made by its authors before proposing such an equation which are not the case in this work. They are: the equation takes into account that the filler must be evenly distributed, there must not be filler–filler interactions and the particles must be rigid and spherical. These assumptions are not exactly true in this case as can be seen in the micrographs (Fig. 6), because HA particles are not spherical and rigid and tend to agglomerate themselves when in high proportions. It can be inferred that filler–filler interactions could be possible in the composites when the filler content increases. The tensile strength of the composites was analyzed using the theoretical model of Nicolais–Narkis, in order to understand the generation of discontinuities or weak points in the structure of theses two-phase systems. The predicted values of the ratio rc/rm are smaller than the experimental ones in all the range of HA contents and radiation doses, though there are a few exceptions. The Nicolais–Narkis equation describes structures where the adhesion is poor, because the weightage factor (1.21) is believed to be dependent on the adhesion quality between the matrix and the inclusion. A value of 1.21 of the weightage factor is stated to be valid for the extreme case of poor adhesion and spherical inclusions [22]. Eq. (3) was follow using a factor P instead of the value of 1.21, being the values of P negative as seen in Table 7, because the composite’s tensile strength is higher than that of the matrix. According to Maiti and Lopez [23], the polymer–filler adhesion improves as the ‘‘P’’ values decrease. Hence, it can be concluded that a good interfacial adhesion is present in HDPE–HA composites prepared via solution. This fact was confirmed when the Kunori and Geil equation was used [22]. This equation relates the tensile strength with a proportionality parameter ‘‘a’’, which is a stress concentration parameter. A higher value of ‘‘a’’ corresponds to a

Table 7 Values of the stress concentration parameter ‘‘P’’ (Eq. (3)) for different radiation doses Composites

PEHA 10 PEHA 20 PEHA 30

‘‘P’’ for different doses (kGy) 0

25

50

100

13.7544 9.0660 5.1534

20.9515 15.0152 7.5367

13.9286 9.4715 6.0469

10.2052 7.1541 3.4993

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stronger stress concentration. The Kunori and Geil equation is as follows: rc ¼ expðaUf Þ. ð5Þ rm

Table 10 Crystallization temperatures of irradiated HDPE/HA composites Sample

Tc (°C) ± 1 0 kGy

25 kGy

50 kGy

100 kGy

The values of ‘‘a’’ are shown in Table 8. As it can be seen, those values are negative, which leads to the same conclusion as before, that is, there seems to be a good interfacial adhesion in the composites. Tables 7 and 8, display the values of ‘‘P’’ and ‘‘a’’ increase when the composites are irradiated at 100 kGy of absorbed dose. This fact could be attributed to a slight hindrance of the polymer–filler interaction. Table 6 displays a sharper decrease in the ec/em values than what Nielsen’s equation predicts. This indicates that the filler produces discontinuities in the stress transfer and a weakening of the composite, bringing about the fracture at lower strains. These results contradict those obtained for the tensile strength. Some explanations could be the formation of HA agglomerates because of its nanometric size (demonstrated in Fig. 6) and the degradative process in the HDPE, which strongly decrease the values of elongation at break. Tables 9 and 10 exhibit the thermal properties of the composites. As seen there, the Tm (melting temperature) and Tc (crystallization temperature) remained unchanged after the radiation process. However, from Table 11 it is clear that the crystallinity (X) degree showed a slight increase with filler addition that could be attributed to a minor nucleating effect of nanometric HA particles. When samples were exposed to gamma radiation, a decrease in crystallinity was observed in all the materials, probably due to the interruption of the more linear sequences of the polyethylene resultant from the chain scission process and/or crosslinking. Fig. 8 displays the SSA curves of HDPEd irradiated at different absorbed doses. At 50 and 100 kGy of absorbed

HDPEd PEHA 10 PEHA 20 PEHA 30

116 116 116 117

115 116 116 116

115 115 116 116

115 116 116 115

Table 8 Values of the proportionality constant ‘‘a’’ (Eq. (5)) for different radiation doses Composites

PEHA 10 PEHA 20 PEHA 30

‘‘a’’ for different doses (kGy) 0

25

50

100

28.7566 15.0592 8.2890

37.7122 20.7240 10.7956

29.0030 15.5082 9.2900

23.3093 12.7755 6.1868

Table 9 Melting temperatures of irradiated HDPE/HA composites Sample

Tm (°C) ± 1 0 kGy

25 kGy

50 kGy

100 kGy

HDPEd PEHA 10 PEHA 20 PEHA 30

134 135 135 134

133 134 134 133

132 134 134 132

134 135 135 133

Table 11 Crystallinity degree of irradiated HDPE/HA composites Sample

X (%) ± 2 0 kGy

25 kGy

50 kGy

100 kGy

HDPEd PEHA 10 PEHA 20 PEHA 30

70 74 74 74

62 67 65 67

60 65 60 64

59 67 62 61

doses a broadening of the main melting peak and the appearance of a shoulder in the peak towards its lower temperature side are observed. It seems that the irradiation at 50 kGy, and even more at 100 kGy, could be producing an interruption and a slight decrease in the length of the crystallizing CH2 sequences, thus increasing the crystal populations that melt at lower temperatures (represented in that shoulder). Such a fact is consistent with the formation of some thinner lamellas at the expense of the disappearance of some of the thicker ones. Something similar happens in the HDPEwd when it is irradiated (Fig. 9). However, in the first case, chain scissions must be the responsible for such effect, as demonstrated by the results obtained from other property measurements, whereas in the last case, chain branching and/or crosslinking produced the slight increment in the thinner lamella populations (as demonstrated by the results obtained from the determination of other properties). The thermal stability of the composites was determined through thermogravimetric analyses. In polymers, some factors such as the polymer/filler compatibility and their structure could affect the decomposition process. The initial decomposition temperatures (Tid) allow determining the beginning of the decomposition process of the composites (Table 12). Tid tended to increase slightly with HA additions as a consequence of a stabilizing effect provided by filler. On the other hand, radiation caused a decrease slightly in Tid as a consequence of polymer degradation even when the absorbed dose did not caused significant variations in this parameter. Table 13 exhibits the activation energy of the composites at the different absorbed doses. The behavior is analogous to that of the degradation initial temperatures. In non-irradiated samples, the thermal stability of the composites is improved with HA inclusion, reaching a maximum at 20 pph and remaining constant at 30 pph. It was observed

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100 kGy

50 mW

50 kGy

25 kGy

0 kGy

40

60

80

100

120

140

160 °C

Fig. 8. Irradiated HDPEd DSC heating scans after applying the SSA technique.

100 kGy 50 mW

50 kGy 25 kGy

O kGy

40

60

80

100

120

140

160 °C

Fig. 9. DSC heating scans after applying the SSA to irradiated HDPEwd.

Table 12 Initial decomposition temperatures of HDPE/HA blends

Table 13 Activation energy of irradiated composites

Sample

Tid (°C) ± 1

Sample

Ea (J/mol) ± 5

0 kGy

25 kGy

50 kGy

100 kGy

0 kGy

25 kGy

50 kGy

100 kGy

HDPEd PEHA 10 PEHA 20 PEHA 30

452 458 460 461

448 456 455 455

447 458 455 454

449 457 454 453

HDPEd PEHA 10 PEHA 20 PEHA 30

325 338 385 380

295 296 352 371

280 265 268 318

294 297 322 337

4. Conclusions that HA is affecting the beginning of the degradation by retarding it. In general, all the composites irradiated at different absorbed doses are less stables and more prone to thermal degradation.

Composites of HDPE/HA prepared via solution showed an excellent filler dispersion, but decalin affected the polymer behavior when irradiated, causing a decrease in its

A. Carmen et al. / Nucl. Instr. and Meth. in Phys. Res. B 247 (2006) 331–341

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