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Annealing effects on the mechanical properties of organic composite materials irradiated with γ-rays
Journal
146
ANNEALING MATERIALS S. EGUSA
EFFECTS ON THE MECHANICAL IRRADIATED WITH y-RAYS *
**, M.A.
KIRK
and
R.C.
of Nuclear
PROPERTIES
Materials 127 (1985) 146-152 North-Holland, Amsterdam
OF ORGANIC COMPOSITE
BIRTCHER
Materials Science and Technology Diuision, Argonne National Laboratory,
Argonne, Illinois
60439, USA
and M. HAGIWARA Takasaki Japan
Radiation
Chemistry
Research
Received
14 August
1984; accepted
Establishment,
17 September
Japan Atomic
Energy Research
Institute,
Takasaki,
Gunma 370- 12,
1984
Four kinds of cloth-filled organic composites (filler: glass or carbon fiber; matrix: epoxy or polyimide resin) and a unidirectional alumina fiber/epoxy composite were irradiated with 6oCo y-rays at room temperature, and then annealed in vacua 2 h at 180°C. These composites were examined with regard to the mechanical properties at room temperature before and after the annealing. Following irradiation the Young (tensile) modulus of these composites remains practically unchanged up to 2000 Mrad irrespective of the annealing. The shear modulus and the ultimate strength also remain unchanged up to this dose before annealing, but after annealing both of them begin to decrease significantly at doses below 2000 Mrad for all the composites except the glass/polyimide composite. This finding indicates that latent radiation damage is activated by the annealing, thus causing a decrease in the load transfer capacity at the fiber/matrix interface and, hence, losses in the shear modulus and the ultimate strength. As to the fracture behavior, on the other hand, the propagation energy increases only slightly with increasing absorbed dose before annealing, but after annealing it increases significantly, apparently accompanying a decrease in the ultimate strength. This result is also attributed to the annealing-activated damage at the fiber/matrix interface.
1. Introduction
degradation in the mechanical and electrical properties. Accordingly, the study of irradiation effects in organic composites will be of great importance. For this reason, several studies have been done recently on the irradiation effects, particularly from the viewpoint of organic insulators to be used in superconducting fusion magnets [7-111. In spite of these studies, however, questions still remain about the mechanism of radiation-induced degradation of organic composite materials. In a previous work on 2 MeV electron irradiations [12], we found that a radiation-induced decrease in the ultimate strength of composite materials is attended by a decreasing shear modulus and an unchanged Young modulus with the absorbed dose up to 15,000 Mrad. This finding suggests that the degradation in the mechanical properties is caused mainly by a decrease in the capacity of load transfer from the matrix to the fiber due to the radiation
Organic composite materials are leading candidates for mechanical supporters and electrical insulators in the construction of superconducting magnets for fusion reactors [l-3]. This is because organic materials are superior to inorganic materials in terms of cost and processing [4]. Organic composites are also suitable materials for space vehicles because they are light and strong [5,6]. If organic composites are employed for these purposes, however, they will be subjected to substantial quantities of high-energy radiation during several years or decades, thus leading to significant * Work supported by the US Department of Energy. ** Present address: Takasaki Radiation Chemistry Research Establishment, Japan Atomic Energy Research Institute, Takasaki, Gunma 370-12, Japan.
0022-3115/85/$03.30 0 Elsevier Science Publishers (North-Holland Physics Publishing Division)
B.V.
S. Egusa et al. / Organic composite materials damage at the interface, thus decreasing the ultimate strength and the shear modulus at the same time. The previous work demonstrated also that these mechanical properties exhibit an incubation dose of 2,OOt-10,000 Mrad before beginning to decrease, thus giving an appearance that no radiation damage is produced until the absorbed dose exceeds the incubation dose. In the present work, five kinds of organic composites were irradiated with 6oCo y-rays, and were examined with regard to the mechanical properties at room temperature. This study revealed that the mechanical properties are scarcely changed by the irradiation up to 2,000 Mrad, but are significantly degraded by the subsequent in-vacua annealing 2 h at 180°C depending on the kind of composites. The present paper mainly describes the dose dependence of the mechanical properties of the composites annealed after irradiation. Possible mechanisms for the annealing effects are also presented, together with an expression for the dose dependence of the ultimate strength.
2. Experimental 2.1. Materials Four kinds of cloth-filled organic composites were prepared by Sumitomo Bakelite Co., Ltd., using E-glass fiber cloth (Kanebo KS-1210) or carbon fiber cloth (Torayca No. 6142) as reinforcing filler, and epoxy (Sumiepoxy ELM-434 cured with p-aminophenyl sulfone, 1 : 1 by mol) or polyimide (Kerimid 601) as matrix resin. The 2.0 mm thick laminate sheets of these composites were cut into specimens of 6.4 x 2.0 x 70 mm3 dimensions. Another kind of composite was especially prepared by Sumitomo Kagaku Co., Ltd., using unidirectional alumina fibers [13] as reinforcing filler and epoxy as matrix resin. The epoxy was the same type as that described above. The 1.6 mm thick sheet was cut into specimens of 3.2 X 1.6 X 70 mm3 dimensions.
147
procedures as previously described [12]. All the tests were conducted at room temperature and at a crosshead speed of 0.5 mm/min, applying the load in the direction normal to the laminate planes of specimen. The Young (tensile) and shear moduli were determined usually for one specimen at each dose. For the ultimate strength and the fracture propagation energy, on the other hand, an average value of three failure tests was taken for each dose.
3. Results 3. I. Ultimate strength The ultimate strength calculated from [14]
of the composite,
u,, = 3Pf(I/h)/2bh,
a,,,
was
(1)
where Pr is the applied load at failure in the three-point bend test, I is the span length, b is the specimen width, and h is the specimen depth (thickness). The failure test was made at the span length of 12.7 mm. The ultimate strength thus determined is plotted as a function of absorbed dose in figs. l-5 separately for each composite studied here. Each data point indicates the average value of three failure tests, and the error bars show the standard deviation. Missing error bars mean that the deviation was too small to be shown. Following irradiation the ultimate strength is seen to be scarcely changed or slightly increased by the irradia-
2.2. Irradiations and mechanical tests 6oCo y-ray irradiations were made in air at ambient temperature, with a dose rate of 4.6 Mrad/h up to a maximum dose of 2,000 Mrad. The irradiated and control specimens were put into a quartz capsule, and then annealed in vacua 2 h at 180°C. The mechanical properties were examined by performing three-point bend tests using the same bend rig and
Absorbed dose (Mad) Fig. 1. Plot of the ultimate strength versus the absorbed dose in matrix for the glass/epoxy composite before (0) and after in-vacua annealing 2 h at 180°C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 24.5, k = 2.15 X 10W3 Mrad-‘, C = 31.6 kgf/mm’. and R = 45.7 kgf/mm2 (see text).
S. Egusa et al. / Organic composite materials
148 0 loo-
,
5
10
15
I
I
I
0
20 (MGy)
I
glasslpolyimide
“E _EfJo-
,
loo -
(MPa)
-8GO
5
10
I
I
Q-----Y
P
-
-
-0-
40-
annealed P---_
1000 1500 Absorbed dose Wad)
200
i
000 0
2000
Fig. 2. Plot of the ultimate strength versus the absorbed dose in matrix for the glass/polyimide composite before (0) and after in-vacua annealing 2 h at 180°C (0).
tion up to 2,000 Mrad for all the composites (filled circles in figs. l-5), thus indicating the existence of an incubation dose before the strength begins to decrease. After annealing, however, the ultimate strength comes to exhibit a clear dose dependence for all the composites except the glass/polyimide composite (open circles in figs. l-5). For the glass/epoxy composite (fig. l), for instance, the ultimate strength remains unchanged up to about 1,000 Mrad and then decreases with increasing absorbed dose. Such a characteristic of dose dependence is observed also for the carbon/epoxy,
600
400
200
,I0
Wa)
-
r60-
1
-
eo-~____.____.____.____e___--“o
F
5
20 (MGy)
‘carbonlpdtyimide
YE 5
15
500
1aoo 1500 Absorbed dose (Wad)
2ooo
Fig. 4. Plot of the ultimate strength versus the absorbed dose in matrix for the carbon/pctyimide composite before (0) and after in-vacua annealing 2 h at 180°C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 6.27, k = 3.11 X lo-’ Mrad-‘. C = 20.0 kgf/mm2, and R = 45.5 kgf/mm* (see text).
and alumina/ epoxy composites carbon/ polyimide, (figs. 3, 4, and 5). For these composites, the ultimate strength begins to decrease at a dose below 1,000 Mrad, and levels off at doses of 1,000 Mrad or above. For the glass/epoxy composite, the leveling off will be observed at a dose above 2,000 Mrad, as is supposed from a preliminary experiment of the present work using specimens irradiated with 2 MeV electrons up to 5,000 Mrad. It is worth noting that the ultimate strength of the control specimen is practically the same before and
200
0 ,
5
10 I
15
20 (MGy)
I alumina/epoxy
Wa)
T
L 010
40
‘alp-
200
20
0
1 400
anneal ed+-
loo0 Absorbed
1500 2000 dose (Mrad)
Fig. 3. Plot of the ultimate strength versus the absorbed dose in matrix for the carbon/epoxy composite before (0) and after in-vacua annealing 2 h at 18O’C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 3.72, k = 2.13 x 10e3 Mrad-‘, C = 35.9 kgf/mm2, and R = 29.7 kgf/mm* (see text).
o-o 0
500
loo0 Absorbed
1500 2ocO dose (t&ad)
Fig. 5. Plot of the ultimate strength versus the absorbed dose in matrix for the alumina/epoxy composite before (0) and after in-vacua annealing 2 h at 18O’C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 55.2, k = 6.80X 10m3 Mrad-‘, C = 46.2 kgf/mm2, and R = 82.2 kgf/mm2 (see text).
S. Eg-usaet al. / Organic composite materials
after the annealing for all the composites. This result indicates that the decrease in the ultimate strength after annealing is definitely attributable to latent radiation damage which can be activated by the annealing. 3.2. Young and shear mod& The Young (tensile) and shear moduli were determined by performing the three-point bend test at span lengths of 63.5, 50.8, 38.1, 25.4, 19.05, and 12.7 mm for each specimen, and by solving the simultaneous equations which were derived from [14]
where A is the deflection, P is the applied load, E is the Young modulus, and G is the shear modulus. Further details of the procedure are described elsewhere [12]. The Young modulus of the control specimen increased in the order of the glass/ polyimide, glass/epoxy, carbon/ polyimide, and carbon/ epoxy composites, reflecting a difference in the Young modulus of the constituent materials. The Young modulus of the alumina/epoxy composite was about 2 times higher than that of the carbon/epoxy composite, perhaps reflecting a difference in the reinforcing form of unidirectional fibers and fiber cloth. Following irradiation the Young modulus of these composites remained practically unchanged up to 2,000 Mrad (see fig. 2 in ref. [15]). This was also the case even for the specimens which were annealed in vacua 2 h at 180°C after the irradiation.
149
The shear modulus of the control specimen, on the other hand, was comparable to that for each of the other four kinds of cloth-filled composites, but was nearly 2 times higher for the unidirectional alumina fiber/epoxy composite. Following irradiation the shear modulus was scarcely changed or slightly increased by the irradiation up to 2,000 Mrad for all the composites studied here (see figs. 3-7 in ref. [15]). After annealing, however, the shear modulus came to exhibit a clear dose dependence for all the composites except the glass/ polyimide composite. The dose dependence was quite similar to that described above for the ultimate strength, thus suggesting that the ultimate strength and the shear modulus are correlated with each other. The correlation was, in fact, confirmed by the plot of the ultimate strength versus the shear modulus shown in fig. 6 for the composites annealed after irradiation. 3.3. Fracture propagation The fracture propagation energy was calculated by integrating the load-deflection curve over the deflection from the failure point to infinity [16]. After being normalized for the specimen cross-sectional area, the energy was plotted as a function of absorbed dose. A typical example of the plot is shown in fig. 7 for the glass/epoxy composite. Each data point indicates the average value of three tests made at the span length of 12.7 mm, and the error bars show the standard deviation. Missing error bars mean that the deviation was too small to be shown. Following irradiation the propagation energy is seen to increase only slightly with absorbed dose (filled circles
(GPa)
0
1
2
3
4
- 140“E 5 120-
5 I
10 I
15 I
20 (MGy)
I glass/epoxy
%0L b c 00E f
60-
$
40-
5
/
_/ /* /’
20 0-o 0
-
400
-
200
..
100 200 300 400 Shear modulus (kgf/mm2)
Fig. 6. Correlation between the shear modulus and the ultimate strength for the glass/epoxy (0). carbon/epoxy (0). carbon/ polyimide (O), and alumina/epoxy composites (M) annealed in vacua 2 h at 18OT after irradiation.
Absorbed
dose (Mud)
Fig. 7. Plot of the fracture propagatton energy versus the absorbed dose in matrix for the glass/epoxy composite before (0) and after in-vacua annealing 2 h at 180°C (0).
150
S. Egusa et al. / Organic composite materials
in fig. 7). After annealing (open circles), however, the propagation energy begins to increase at about 1,000 Mrad, apparently accompanying a decrease in the ultimate strength (see open circles in fig. 1). These characteristics of the propagation energy versus absorbed dose plot were observed also for the carbon/ polyimide and alumina/epoxy composites. For the glass/polyimide composite, no significant increase in the propagation energy after the annealing was observed within the dose range covered in the present work. For the carbon/epoxy composite, on the other hand, the exact evaluation of the propagation energy was impossible both before and after the annealing. This was because the load did not approach zero even at a large deflection after the failure [12].
creasing the number of polymer chains capable of load transfer at the fiber/matrix interface. In addition, internal shear stresses will be generated at the fiber/matrix interface during the cooldown to room temperature by a difference in the thermal expansion between the fiber and the matrix. Such stresses may also aid the entangled polymer chains to loosen. As suggested by Klabunde and Coltman, migration of gas evolved during irradiation and/or annealing may play an important role in activating latent radiation damage by accumulating preferentially at the fiber/matrix interface [8]. This possible mechanism of gas migration will be much more important during the annealing, because heating of irradiated polymers is known to increase the yield of gas evolution [ 181. 4.2. Dose dependence of ultimate strength
4. Discussion 4.1. Mechanism
of annealing-activated
degradation
The Young modulus of composites studied here was found to remain unchanged even after the irradiation up to 2,000 Mrad and the subsequent annealing in vacua 2 h at 18O’C. The shear modulus and the ultimate strength, however, were found to decrease simultaneously after the annealing (figs. l-6). It was also found that the fracture propagation energy increases after the annealing (fig. 7), apparently accompanying a decrease in the shear modulus or the ultimate strength. All of these characteristics of the degradation behavior after the annealing were observed even before the annealing for composites irradiated with 2 MeV electrons up to 15,000 Mrad [12]. These findings strongly suggest that the mechanism of annealing-activated degradation of irradiated composites will be essentially the same as that proposed earlier for electron irradiation [12]. Annealing of irradiated composites, therefore, is considered to activate latent radiation damage near the fiber/matrix interface, thus causing a decrease in the load transfer capacity at the interface and, consequently, losses in the shear modulus and the ultimate strength. When organic composites are irradiated at room temperature, rearrangement of polymer chains which have been severed and/or cross-linked is probably not completed because the glass transition temperature is well above 1OO’C for highly-cross-linked polymers like the present epoxy and polyimide resins [17]. The rearrangement, however, will be completed if the composites are annealed above the transition temperature. In the process of the rearrangement, polymer chains which have been entangled may be loosened, thus de-
The ultimate strength of composites studied here was found to have an incubation dose before beginning to decrease (figs. l-5). This was the case both before and after annealing, although the incubation dose was reduced considerably by the annealing. The existence of such an incubation dose may be associated with network structures of highly-cross-linked polymers in the fiber/matrix interface region. At low doses in the incubation range, each event such as main-chain scission of polymer will take place separately from one another, because as a first approximation the event will be regarded to occur at random in an irradiated system. Such an isolated event in the network structure is considered to cause no observable decrease in the load transfer capacity. At high doses above the incubation range, however, some accumulation of these events will destroy the network structure in certain localized regions, thus causing a decrease in the load transfer capacity and, consequently, losses in the ultimate strength and the shear modulus. Taking into consideration these points, we tried to formulate the dose dependence of ultimate strength of composite materials before and after annealing. As mentioned above, the mechanism proposed for the radiation-induced degradation of organic composites is considered to hold even after the annealing. From a simplistic standpoint, therefore, annealing of irradiated composites may be regarded merely to decrease the number of polymer chains capable of load transfer at the fiber/matrix interface, thus degrading the composites to such an extent that could be attained by further irradiation without annealing. Then the following expression for the dose dependence of the ultimate strength of the composite, a,,, may be valid after the
151
S. Egusa et al. / Organic composite materials
annealing
also:
UC”= aor,VrTJ + uG(l - Vt).
(3)
where a is the coefficient dependent on the form of fibers, (I~”is the ultimate tensile strength of fibers, 0: is the matrix stress at the composite failure, Vr is the volume fraction of fibers, and q is the load transfer capacity at the fiber/matrix interface. The details of this equation are described elsewhere [12]. In order to describe the dose dependence of 7. we tentatively adopted the target theory of radiation biology in a previous paper [12]. The theory was indeed able to describe the dose dependence of 9. but the identity of was not necessarily obvious in case of the “target” composite materials. In the present paper, therefore, we propose another expression for the dose dependence of 7~without using the idea of the target. From the above considerations on the incubation dose, it will be reasonably assumed that an observable decrease in n is determined not only by the number of main-chain scissions but also by some accumulation of the scissions. This assumption may be expressed by dn/dN
= -AN”,
(4)
where A and P are constants, and N is the concentration of main-chain scissions in the network structure of the fiber/matrix interface region. Integration of this equation and some rearrangement give s=(r~~.,>[l
-(N/N,)“]
+nm>
(5)
where n = Y + 1, v,, is the initial value of 7, and the subscript cc refers to the limiting case that the network structure has been destroyed completely. The kinetics of main-chain scissions, on the other hand, can be written as dN/dD=k(N,-N),
(6)
where D is the absorbed dose, and k is the parameter reflecting the efficiency of main-chain scissions due to irradiation. For composites annealed after irradiation, the efficiency of latent radiation damage being activated should also be included in the value of k. Integration of this equation and substitution of the resulting N/N, into eq. (5) give ?1=(90-~m)[l-(1-e-kD)n]+9m.
(7).
The constant n (n 2 1) may be termed the accumulation parameter, because this will reflect the degree of accumulation of main-chain scissions which is necessary to cause an observable decrease in n. At the same time, the parameter n greater than 1 can be regarded as implying some surplus of potential capacity of load
transfer over the minimum capacity essential for the reinforcing fibers. The value of n, therefore, is considered to depend not only on the network structure at the fiber/matrix interface but also on the strength of reinforcing fibers. For the parameters no and q,, on the other hand, the fiber/matrix interface is known to have at least two modes of load transfer, i.e., the chemical bond mode and the friction force mode [19]. If therefore the friction force mode is much less sensitive to radiation compared with the chemical bond mode, and is regarded to be practically dose-independent, then the values of (no - nm) and nm may reflect the initial contributions due to the chemical bond mode and the friction force mode, respectively. The fitting procedure of eqs. (3) and (7) was done for the four parameters n, k, C, and R, with the relationships given by C=~n,,Q(?a-nI,), R = (YCT~~I’~TJ~ + a;(1
(8) - Vr).
(9)
Simultaneous equations in n, k, C, and R were solved by using the non-linear least squares method referred to as the Marquardt Method. The solid curves in figs. 1 and 3-5 are the results of the fitting procedure made for the composites annealed after irradiation. The fit between the calculated and observed plots is seen to be excellent, and accordingly it is concluded that eqs. (3) and (7) can describe the dose dependence of the ultimate strength for the composites annealed after irradiation. For the composites before annealing (filled circles in figs. l-5), on the other hand, the fitting procedure is practically impossible at the present stage because the ultimate strength does not decrease at all at absorbed doses below 2,000 Mrad. The fitting procedure would be possible if further data points are obtained from a wider dose range than is covered in the present work. Then the annealing effects could be discussed quantitatively based on the values of n, k, C, and R.
5. Conclusions
The present work has shown that the mechanical properties of composite materials are scarcely changed by the irradiation up to 2,000 Mrad, but are significantly degraded by the subsequent annealing, thus indicating that latent radiation damage is activated by the annealing. The linear relationship found between the shear modulus and the ultimate strength for each composite strongly suggests that the decrease in the ultimate strength is attributable to a decrease in the load transfer capacity at the fiber/matrix interface due to the radia-
152
S. Egua
et al. / Organic composite materials
tion damage activated by the annealing. The increased propagation energy for composites annealed after irradiation is also explained by the annealing-activated damage at the interface. Based on these findings, possible mechanisms for the annealing effects of irradiated composites are presented, together with a new expression for the dose dependence of the ultimate strength. In order to confirm the presented mechanisms, however, further studies are required particularly on the influence of annealing temperature and time on the mechanical properties of irradiated composites.
Acknowledgements The authors wish to express their appreciation to the staff of Sumitomo Kagaku Co., Ltd. for providing specimens of the alumina/epoxy composite. Critical and valuable comments on this manuscript from Dr. C.E. Klabunde of Oak Ridge National Laboratory are gratefully acknowledged. One of the authors (S.E.) is especially grateful to Dr. B.S. Brown of Argonne National Laboratory for providing him an opportunity to work at Argonne. He also appreciates grant from Japan Atomic Energy Research Institute which enabled him to visit Argonne.
References [l] B.S. Brown, J. Nucl. Mater. 97 (1981) 1. [2] M.A. Abdou, J. Nucl. Mater. 72 (1978) 147. [3] F.W. Clinard, Jr. and G.F. Hurley, J. Nucl. Mater. 104 (1981) 705.
103 &
[4] R.D. Hay and E.J. Rapperport, Oak Ridge National Laboratory Report ORNL/TM-2643 (1976). [5] R. Yokota, Nihon Fukugo Zairyo Gakkai-shi 9 (1983) 92. [6] R.E. Fornes, J.D. Memory and N. Naranong, J. Appl. Polym. Sci. 26 (1981) 2061. [7] R.R. Coltman, Jr. and C.E. Klabunde, J. Nucl. Mater. 103 & 104 (1981) 717; 113 (1983) 268. [8] C.E. Klabunde and R.R. Coltman, Jr., J. Nucl. Mater. 117 (1983) 345. [9] S. Takamura and T. Kato, J. Nucl. Mater. 103 & 104 (1981) 729; Advances in Cryogenic Engineering Materials 30 (1984) 41. [lo] G.F. Hurley, J.D. Fowler and D.L. Rohr, Cryogenics 23 (1983) 415. [ll] H.W. Weber, E. Kubasta, W. Steiner, H. Benz and K. Nylund, J. Nucl. Mater. 115 (1983) 11. [12] S. Egusa, M.A. Kirk, R.C. Birtcher, M. Hagiwara and S. Kawanishi. J. Nucl. Mater. 119 (1983) 146; Nucl. Instruments Methods in Physics Research Bl (1984) 610. [13] E. Ichiki and Y. Abe, Kobunshi 30 (1981) 883. 1141 C. Zweben, W.S. Smith and M.W. Wardle. in: Composite Materials: Testing and Design (Fifth Conference), ASTM STP 674, Ed. S.W. Tsai (American Society for Testing and Materials, 1979) p. 228. [15] S. Egusa, M.A. Kirk and R.C. Birtcher, J. Nucl. Mater. 126 (1984) 152. [16] P.W.R. Beaumont, P.G. Riewald and C. Zweben, in: Foreign Object Impact Damage to Composites, ASTM STP 568 (American Society for Testing and Materials, 1974) p. 134. (171 T. Murayama, in: Dynamic Mechanical Analysis of Polymeric Material (Elsevier, Amsterdam, 1978) Chapter 3. [18] A. Chapiro, in: Radiation Chemistry of Polymeric Systems (Interscience, New York, 1962) p. 357. [19] C.C. Chamis, in: Composite Materials, Vol. 6, Ed. E.P. Plueddemann (Academic Press, New York, 1974) Chapter 2.
146
ANNEALING MATERIALS S. EGUSA
EFFECTS ON THE MECHANICAL IRRADIATED WITH y-RAYS *
**, M.A.
KIRK
and
R.C.
of Nuclear
PROPERTIES
Materials 127 (1985) 146-152 North-Holland, Amsterdam
OF ORGANIC COMPOSITE
BIRTCHER
Materials Science and Technology Diuision, Argonne National Laboratory,
Argonne, Illinois
60439, USA
and M. HAGIWARA Takasaki Japan
Radiation
Chemistry
Research
Received
14 August
1984; accepted
Establishment,
17 September
Japan Atomic
Energy Research
Institute,
Takasaki,
Gunma 370- 12,
1984
Four kinds of cloth-filled organic composites (filler: glass or carbon fiber; matrix: epoxy or polyimide resin) and a unidirectional alumina fiber/epoxy composite were irradiated with 6oCo y-rays at room temperature, and then annealed in vacua 2 h at 180°C. These composites were examined with regard to the mechanical properties at room temperature before and after the annealing. Following irradiation the Young (tensile) modulus of these composites remains practically unchanged up to 2000 Mrad irrespective of the annealing. The shear modulus and the ultimate strength also remain unchanged up to this dose before annealing, but after annealing both of them begin to decrease significantly at doses below 2000 Mrad for all the composites except the glass/polyimide composite. This finding indicates that latent radiation damage is activated by the annealing, thus causing a decrease in the load transfer capacity at the fiber/matrix interface and, hence, losses in the shear modulus and the ultimate strength. As to the fracture behavior, on the other hand, the propagation energy increases only slightly with increasing absorbed dose before annealing, but after annealing it increases significantly, apparently accompanying a decrease in the ultimate strength. This result is also attributed to the annealing-activated damage at the fiber/matrix interface.
1. Introduction
degradation in the mechanical and electrical properties. Accordingly, the study of irradiation effects in organic composites will be of great importance. For this reason, several studies have been done recently on the irradiation effects, particularly from the viewpoint of organic insulators to be used in superconducting fusion magnets [7-111. In spite of these studies, however, questions still remain about the mechanism of radiation-induced degradation of organic composite materials. In a previous work on 2 MeV electron irradiations [12], we found that a radiation-induced decrease in the ultimate strength of composite materials is attended by a decreasing shear modulus and an unchanged Young modulus with the absorbed dose up to 15,000 Mrad. This finding suggests that the degradation in the mechanical properties is caused mainly by a decrease in the capacity of load transfer from the matrix to the fiber due to the radiation
Organic composite materials are leading candidates for mechanical supporters and electrical insulators in the construction of superconducting magnets for fusion reactors [l-3]. This is because organic materials are superior to inorganic materials in terms of cost and processing [4]. Organic composites are also suitable materials for space vehicles because they are light and strong [5,6]. If organic composites are employed for these purposes, however, they will be subjected to substantial quantities of high-energy radiation during several years or decades, thus leading to significant * Work supported by the US Department of Energy. ** Present address: Takasaki Radiation Chemistry Research Establishment, Japan Atomic Energy Research Institute, Takasaki, Gunma 370-12, Japan.
0022-3115/85/$03.30 0 Elsevier Science Publishers (North-Holland Physics Publishing Division)
B.V.
S. Egusa et al. / Organic composite materials damage at the interface, thus decreasing the ultimate strength and the shear modulus at the same time. The previous work demonstrated also that these mechanical properties exhibit an incubation dose of 2,OOt-10,000 Mrad before beginning to decrease, thus giving an appearance that no radiation damage is produced until the absorbed dose exceeds the incubation dose. In the present work, five kinds of organic composites were irradiated with 6oCo y-rays, and were examined with regard to the mechanical properties at room temperature. This study revealed that the mechanical properties are scarcely changed by the irradiation up to 2,000 Mrad, but are significantly degraded by the subsequent in-vacua annealing 2 h at 180°C depending on the kind of composites. The present paper mainly describes the dose dependence of the mechanical properties of the composites annealed after irradiation. Possible mechanisms for the annealing effects are also presented, together with an expression for the dose dependence of the ultimate strength.
2. Experimental 2.1. Materials Four kinds of cloth-filled organic composites were prepared by Sumitomo Bakelite Co., Ltd., using E-glass fiber cloth (Kanebo KS-1210) or carbon fiber cloth (Torayca No. 6142) as reinforcing filler, and epoxy (Sumiepoxy ELM-434 cured with p-aminophenyl sulfone, 1 : 1 by mol) or polyimide (Kerimid 601) as matrix resin. The 2.0 mm thick laminate sheets of these composites were cut into specimens of 6.4 x 2.0 x 70 mm3 dimensions. Another kind of composite was especially prepared by Sumitomo Kagaku Co., Ltd., using unidirectional alumina fibers [13] as reinforcing filler and epoxy as matrix resin. The epoxy was the same type as that described above. The 1.6 mm thick sheet was cut into specimens of 3.2 X 1.6 X 70 mm3 dimensions.
147
procedures as previously described [12]. All the tests were conducted at room temperature and at a crosshead speed of 0.5 mm/min, applying the load in the direction normal to the laminate planes of specimen. The Young (tensile) and shear moduli were determined usually for one specimen at each dose. For the ultimate strength and the fracture propagation energy, on the other hand, an average value of three failure tests was taken for each dose.
3. Results 3. I. Ultimate strength The ultimate strength calculated from [14]
of the composite,
u,, = 3Pf(I/h)/2bh,
a,,,
was
(1)
where Pr is the applied load at failure in the three-point bend test, I is the span length, b is the specimen width, and h is the specimen depth (thickness). The failure test was made at the span length of 12.7 mm. The ultimate strength thus determined is plotted as a function of absorbed dose in figs. l-5 separately for each composite studied here. Each data point indicates the average value of three failure tests, and the error bars show the standard deviation. Missing error bars mean that the deviation was too small to be shown. Following irradiation the ultimate strength is seen to be scarcely changed or slightly increased by the irradia-
2.2. Irradiations and mechanical tests 6oCo y-ray irradiations were made in air at ambient temperature, with a dose rate of 4.6 Mrad/h up to a maximum dose of 2,000 Mrad. The irradiated and control specimens were put into a quartz capsule, and then annealed in vacua 2 h at 180°C. The mechanical properties were examined by performing three-point bend tests using the same bend rig and
Absorbed dose (Mad) Fig. 1. Plot of the ultimate strength versus the absorbed dose in matrix for the glass/epoxy composite before (0) and after in-vacua annealing 2 h at 180°C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 24.5, k = 2.15 X 10W3 Mrad-‘, C = 31.6 kgf/mm’. and R = 45.7 kgf/mm2 (see text).
S. Egusa et al. / Organic composite materials
148 0 loo-
,
5
10
15
I
I
I
0
20 (MGy)
I
glasslpolyimide
“E _EfJo-
,
loo -
(MPa)
-8GO
5
10
I
I
Q-----Y
P
-
-
-0-
40-
annealed P---_
1000 1500 Absorbed dose Wad)
200
i
000 0
2000
Fig. 2. Plot of the ultimate strength versus the absorbed dose in matrix for the glass/polyimide composite before (0) and after in-vacua annealing 2 h at 180°C (0).
tion up to 2,000 Mrad for all the composites (filled circles in figs. l-5), thus indicating the existence of an incubation dose before the strength begins to decrease. After annealing, however, the ultimate strength comes to exhibit a clear dose dependence for all the composites except the glass/polyimide composite (open circles in figs. l-5). For the glass/epoxy composite (fig. l), for instance, the ultimate strength remains unchanged up to about 1,000 Mrad and then decreases with increasing absorbed dose. Such a characteristic of dose dependence is observed also for the carbon/epoxy,
600
400
200
,I0
Wa)
-
r60-
1
-
eo-~____.____.____.____e___--“o
F
5
20 (MGy)
‘carbonlpdtyimide
YE 5
15
500
1aoo 1500 Absorbed dose (Wad)
2ooo
Fig. 4. Plot of the ultimate strength versus the absorbed dose in matrix for the carbon/pctyimide composite before (0) and after in-vacua annealing 2 h at 180°C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 6.27, k = 3.11 X lo-’ Mrad-‘. C = 20.0 kgf/mm2, and R = 45.5 kgf/mm* (see text).
and alumina/ epoxy composites carbon/ polyimide, (figs. 3, 4, and 5). For these composites, the ultimate strength begins to decrease at a dose below 1,000 Mrad, and levels off at doses of 1,000 Mrad or above. For the glass/epoxy composite, the leveling off will be observed at a dose above 2,000 Mrad, as is supposed from a preliminary experiment of the present work using specimens irradiated with 2 MeV electrons up to 5,000 Mrad. It is worth noting that the ultimate strength of the control specimen is practically the same before and
200
0 ,
5
10 I
15
20 (MGy)
I alumina/epoxy
Wa)
T
L 010
40
‘alp-
200
20
0
1 400
anneal ed+-
loo0 Absorbed
1500 2000 dose (Mrad)
Fig. 3. Plot of the ultimate strength versus the absorbed dose in matrix for the carbon/epoxy composite before (0) and after in-vacua annealing 2 h at 18O’C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 3.72, k = 2.13 x 10e3 Mrad-‘, C = 35.9 kgf/mm2, and R = 29.7 kgf/mm* (see text).
o-o 0
500
loo0 Absorbed
1500 2ocO dose (t&ad)
Fig. 5. Plot of the ultimate strength versus the absorbed dose in matrix for the alumina/epoxy composite before (0) and after in-vacua annealing 2 h at 18O’C (0). The solid curve is the result of fitting procedure made by using eqs. (3) and (7) with n = 55.2, k = 6.80X 10m3 Mrad-‘, C = 46.2 kgf/mm2, and R = 82.2 kgf/mm2 (see text).
S. Eg-usaet al. / Organic composite materials
after the annealing for all the composites. This result indicates that the decrease in the ultimate strength after annealing is definitely attributable to latent radiation damage which can be activated by the annealing. 3.2. Young and shear mod& The Young (tensile) and shear moduli were determined by performing the three-point bend test at span lengths of 63.5, 50.8, 38.1, 25.4, 19.05, and 12.7 mm for each specimen, and by solving the simultaneous equations which were derived from [14]
where A is the deflection, P is the applied load, E is the Young modulus, and G is the shear modulus. Further details of the procedure are described elsewhere [12]. The Young modulus of the control specimen increased in the order of the glass/ polyimide, glass/epoxy, carbon/ polyimide, and carbon/ epoxy composites, reflecting a difference in the Young modulus of the constituent materials. The Young modulus of the alumina/epoxy composite was about 2 times higher than that of the carbon/epoxy composite, perhaps reflecting a difference in the reinforcing form of unidirectional fibers and fiber cloth. Following irradiation the Young modulus of these composites remained practically unchanged up to 2,000 Mrad (see fig. 2 in ref. [15]). This was also the case even for the specimens which were annealed in vacua 2 h at 180°C after the irradiation.
149
The shear modulus of the control specimen, on the other hand, was comparable to that for each of the other four kinds of cloth-filled composites, but was nearly 2 times higher for the unidirectional alumina fiber/epoxy composite. Following irradiation the shear modulus was scarcely changed or slightly increased by the irradiation up to 2,000 Mrad for all the composites studied here (see figs. 3-7 in ref. [15]). After annealing, however, the shear modulus came to exhibit a clear dose dependence for all the composites except the glass/ polyimide composite. The dose dependence was quite similar to that described above for the ultimate strength, thus suggesting that the ultimate strength and the shear modulus are correlated with each other. The correlation was, in fact, confirmed by the plot of the ultimate strength versus the shear modulus shown in fig. 6 for the composites annealed after irradiation. 3.3. Fracture propagation The fracture propagation energy was calculated by integrating the load-deflection curve over the deflection from the failure point to infinity [16]. After being normalized for the specimen cross-sectional area, the energy was plotted as a function of absorbed dose. A typical example of the plot is shown in fig. 7 for the glass/epoxy composite. Each data point indicates the average value of three tests made at the span length of 12.7 mm, and the error bars show the standard deviation. Missing error bars mean that the deviation was too small to be shown. Following irradiation the propagation energy is seen to increase only slightly with absorbed dose (filled circles
(GPa)
0
1
2
3
4
- 140“E 5 120-
5 I
10 I
15 I
20 (MGy)
I glass/epoxy
%0L b c 00E f
60-
$
40-
5
/
_/ /* /’
20 0-o 0
-
400
-
200
..
100 200 300 400 Shear modulus (kgf/mm2)
Fig. 6. Correlation between the shear modulus and the ultimate strength for the glass/epoxy (0). carbon/epoxy (0). carbon/ polyimide (O), and alumina/epoxy composites (M) annealed in vacua 2 h at 18OT after irradiation.
Absorbed
dose (Mud)
Fig. 7. Plot of the fracture propagatton energy versus the absorbed dose in matrix for the glass/epoxy composite before (0) and after in-vacua annealing 2 h at 180°C (0).
150
S. Egusa et al. / Organic composite materials
in fig. 7). After annealing (open circles), however, the propagation energy begins to increase at about 1,000 Mrad, apparently accompanying a decrease in the ultimate strength (see open circles in fig. 1). These characteristics of the propagation energy versus absorbed dose plot were observed also for the carbon/ polyimide and alumina/epoxy composites. For the glass/polyimide composite, no significant increase in the propagation energy after the annealing was observed within the dose range covered in the present work. For the carbon/epoxy composite, on the other hand, the exact evaluation of the propagation energy was impossible both before and after the annealing. This was because the load did not approach zero even at a large deflection after the failure [12].
creasing the number of polymer chains capable of load transfer at the fiber/matrix interface. In addition, internal shear stresses will be generated at the fiber/matrix interface during the cooldown to room temperature by a difference in the thermal expansion between the fiber and the matrix. Such stresses may also aid the entangled polymer chains to loosen. As suggested by Klabunde and Coltman, migration of gas evolved during irradiation and/or annealing may play an important role in activating latent radiation damage by accumulating preferentially at the fiber/matrix interface [8]. This possible mechanism of gas migration will be much more important during the annealing, because heating of irradiated polymers is known to increase the yield of gas evolution [ 181. 4.2. Dose dependence of ultimate strength
4. Discussion 4.1. Mechanism
of annealing-activated
degradation
The Young modulus of composites studied here was found to remain unchanged even after the irradiation up to 2,000 Mrad and the subsequent annealing in vacua 2 h at 18O’C. The shear modulus and the ultimate strength, however, were found to decrease simultaneously after the annealing (figs. l-6). It was also found that the fracture propagation energy increases after the annealing (fig. 7), apparently accompanying a decrease in the shear modulus or the ultimate strength. All of these characteristics of the degradation behavior after the annealing were observed even before the annealing for composites irradiated with 2 MeV electrons up to 15,000 Mrad [12]. These findings strongly suggest that the mechanism of annealing-activated degradation of irradiated composites will be essentially the same as that proposed earlier for electron irradiation [12]. Annealing of irradiated composites, therefore, is considered to activate latent radiation damage near the fiber/matrix interface, thus causing a decrease in the load transfer capacity at the interface and, consequently, losses in the shear modulus and the ultimate strength. When organic composites are irradiated at room temperature, rearrangement of polymer chains which have been severed and/or cross-linked is probably not completed because the glass transition temperature is well above 1OO’C for highly-cross-linked polymers like the present epoxy and polyimide resins [17]. The rearrangement, however, will be completed if the composites are annealed above the transition temperature. In the process of the rearrangement, polymer chains which have been entangled may be loosened, thus de-
The ultimate strength of composites studied here was found to have an incubation dose before beginning to decrease (figs. l-5). This was the case both before and after annealing, although the incubation dose was reduced considerably by the annealing. The existence of such an incubation dose may be associated with network structures of highly-cross-linked polymers in the fiber/matrix interface region. At low doses in the incubation range, each event such as main-chain scission of polymer will take place separately from one another, because as a first approximation the event will be regarded to occur at random in an irradiated system. Such an isolated event in the network structure is considered to cause no observable decrease in the load transfer capacity. At high doses above the incubation range, however, some accumulation of these events will destroy the network structure in certain localized regions, thus causing a decrease in the load transfer capacity and, consequently, losses in the ultimate strength and the shear modulus. Taking into consideration these points, we tried to formulate the dose dependence of ultimate strength of composite materials before and after annealing. As mentioned above, the mechanism proposed for the radiation-induced degradation of organic composites is considered to hold even after the annealing. From a simplistic standpoint, therefore, annealing of irradiated composites may be regarded merely to decrease the number of polymer chains capable of load transfer at the fiber/matrix interface, thus degrading the composites to such an extent that could be attained by further irradiation without annealing. Then the following expression for the dose dependence of the ultimate strength of the composite, a,,, may be valid after the
151
S. Egusa et al. / Organic composite materials
annealing
also:
UC”= aor,VrTJ + uG(l - Vt).
(3)
where a is the coefficient dependent on the form of fibers, (I~”is the ultimate tensile strength of fibers, 0: is the matrix stress at the composite failure, Vr is the volume fraction of fibers, and q is the load transfer capacity at the fiber/matrix interface. The details of this equation are described elsewhere [12]. In order to describe the dose dependence of 7. we tentatively adopted the target theory of radiation biology in a previous paper [12]. The theory was indeed able to describe the dose dependence of 9. but the identity of was not necessarily obvious in case of the “target” composite materials. In the present paper, therefore, we propose another expression for the dose dependence of 7~without using the idea of the target. From the above considerations on the incubation dose, it will be reasonably assumed that an observable decrease in n is determined not only by the number of main-chain scissions but also by some accumulation of the scissions. This assumption may be expressed by dn/dN
= -AN”,
(4)
where A and P are constants, and N is the concentration of main-chain scissions in the network structure of the fiber/matrix interface region. Integration of this equation and some rearrangement give s=(r~~.,>[l
-(N/N,)“]
+nm>
(5)
where n = Y + 1, v,, is the initial value of 7, and the subscript cc refers to the limiting case that the network structure has been destroyed completely. The kinetics of main-chain scissions, on the other hand, can be written as dN/dD=k(N,-N),
(6)
where D is the absorbed dose, and k is the parameter reflecting the efficiency of main-chain scissions due to irradiation. For composites annealed after irradiation, the efficiency of latent radiation damage being activated should also be included in the value of k. Integration of this equation and substitution of the resulting N/N, into eq. (5) give ?1=(90-~m)[l-(1-e-kD)n]+9m.
(7).
The constant n (n 2 1) may be termed the accumulation parameter, because this will reflect the degree of accumulation of main-chain scissions which is necessary to cause an observable decrease in n. At the same time, the parameter n greater than 1 can be regarded as implying some surplus of potential capacity of load
transfer over the minimum capacity essential for the reinforcing fibers. The value of n, therefore, is considered to depend not only on the network structure at the fiber/matrix interface but also on the strength of reinforcing fibers. For the parameters no and q,, on the other hand, the fiber/matrix interface is known to have at least two modes of load transfer, i.e., the chemical bond mode and the friction force mode [19]. If therefore the friction force mode is much less sensitive to radiation compared with the chemical bond mode, and is regarded to be practically dose-independent, then the values of (no - nm) and nm may reflect the initial contributions due to the chemical bond mode and the friction force mode, respectively. The fitting procedure of eqs. (3) and (7) was done for the four parameters n, k, C, and R, with the relationships given by C=~n,,Q(?a-nI,), R = (YCT~~I’~TJ~ + a;(1
(8) - Vr).
(9)
Simultaneous equations in n, k, C, and R were solved by using the non-linear least squares method referred to as the Marquardt Method. The solid curves in figs. 1 and 3-5 are the results of the fitting procedure made for the composites annealed after irradiation. The fit between the calculated and observed plots is seen to be excellent, and accordingly it is concluded that eqs. (3) and (7) can describe the dose dependence of the ultimate strength for the composites annealed after irradiation. For the composites before annealing (filled circles in figs. l-5), on the other hand, the fitting procedure is practically impossible at the present stage because the ultimate strength does not decrease at all at absorbed doses below 2,000 Mrad. The fitting procedure would be possible if further data points are obtained from a wider dose range than is covered in the present work. Then the annealing effects could be discussed quantitatively based on the values of n, k, C, and R.
5. Conclusions
The present work has shown that the mechanical properties of composite materials are scarcely changed by the irradiation up to 2,000 Mrad, but are significantly degraded by the subsequent annealing, thus indicating that latent radiation damage is activated by the annealing. The linear relationship found between the shear modulus and the ultimate strength for each composite strongly suggests that the decrease in the ultimate strength is attributable to a decrease in the load transfer capacity at the fiber/matrix interface due to the radia-
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et al. / Organic composite materials
tion damage activated by the annealing. The increased propagation energy for composites annealed after irradiation is also explained by the annealing-activated damage at the interface. Based on these findings, possible mechanisms for the annealing effects of irradiated composites are presented, together with a new expression for the dose dependence of the ultimate strength. In order to confirm the presented mechanisms, however, further studies are required particularly on the influence of annealing temperature and time on the mechanical properties of irradiated composites.
Acknowledgements The authors wish to express their appreciation to the staff of Sumitomo Kagaku Co., Ltd. for providing specimens of the alumina/epoxy composite. Critical and valuable comments on this manuscript from Dr. C.E. Klabunde of Oak Ridge National Laboratory are gratefully acknowledged. One of the authors (S.E.) is especially grateful to Dr. B.S. Brown of Argonne National Laboratory for providing him an opportunity to work at Argonne. He also appreciates grant from Japan Atomic Energy Research Institute which enabled him to visit Argonne.
References [l] B.S. Brown, J. Nucl. Mater. 97 (1981) 1. [2] M.A. Abdou, J. Nucl. Mater. 72 (1978) 147. [3] F.W. Clinard, Jr. and G.F. Hurley, J. Nucl. Mater. 104 (1981) 705.
103 &
[4] R.D. Hay and E.J. Rapperport, Oak Ridge National Laboratory Report ORNL/TM-2643 (1976). [5] R. Yokota, Nihon Fukugo Zairyo Gakkai-shi 9 (1983) 92. [6] R.E. Fornes, J.D. Memory and N. Naranong, J. Appl. Polym. Sci. 26 (1981) 2061. [7] R.R. Coltman, Jr. and C.E. Klabunde, J. Nucl. Mater. 103 & 104 (1981) 717; 113 (1983) 268. [8] C.E. Klabunde and R.R. Coltman, Jr., J. Nucl. Mater. 117 (1983) 345. [9] S. Takamura and T. Kato, J. Nucl. Mater. 103 & 104 (1981) 729; Advances in Cryogenic Engineering Materials 30 (1984) 41. [lo] G.F. Hurley, J.D. Fowler and D.L. Rohr, Cryogenics 23 (1983) 415. [ll] H.W. Weber, E. Kubasta, W. Steiner, H. Benz and K. Nylund, J. Nucl. Mater. 115 (1983) 11. [12] S. Egusa, M.A. Kirk, R.C. Birtcher, M. Hagiwara and S. Kawanishi. J. Nucl. Mater. 119 (1983) 146; Nucl. Instruments Methods in Physics Research Bl (1984) 610. [13] E. Ichiki and Y. Abe, Kobunshi 30 (1981) 883. 1141 C. Zweben, W.S. Smith and M.W. Wardle. in: Composite Materials: Testing and Design (Fifth Conference), ASTM STP 674, Ed. S.W. Tsai (American Society for Testing and Materials, 1979) p. 228. [15] S. Egusa, M.A. Kirk and R.C. Birtcher, J. Nucl. Mater. 126 (1984) 152. [16] P.W.R. Beaumont, P.G. Riewald and C. Zweben, in: Foreign Object Impact Damage to Composites, ASTM STP 568 (American Society for Testing and Materials, 1974) p. 134. (171 T. Murayama, in: Dynamic Mechanical Analysis of Polymeric Material (Elsevier, Amsterdam, 1978) Chapter 3. [18] A. Chapiro, in: Radiation Chemistry of Polymeric Systems (Interscience, New York, 1962) p. 357. [19] C.C. Chamis, in: Composite Materials, Vol. 6, Ed. E.P. Plueddemann (Academic Press, New York, 1974) Chapter 2.