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ROBERT A. SPURR, EDWARD H. ERATH, and HOWARD MYERS1 Research Laboratories, Hughes Aircraft Co., Culver City, Calif.
Curing Process in Phenolic Resin
X-Ray Diffraction Analysis Extent of cure, which, vitally affects physical properties, can now be measured by this new technique
erties of phenolic resins are influenced by the curing process, which is defined here as the heating of a resin in an oven subsequent to molding in a press. Elastic modulus, flexural strength. dielectric constant, and dielectric loss tangent are all affected by duration and temperature. To study these factors. nature of the curing process should be determined and an empirical measure of its progress should be developed. Therefore, curing in a phenolic resin was followed by x-ray diffraction, electron microscopy. and infrared and ultraviolet spectroscopy. The first two techniques were most useful. Many tests (2-6, 70, 72, 22, 24) have been developed to estimate degree of polymerization for phenolic resins and their applicability ( 7 ) was reviewed in 1955. Generally, no test sensitive to changes induced in thermosetting plastics by the curing process has been evolved; however, x-ray diffraction spectra can sensitively measure degree of cure in phenolic resins. Such spectra of phenolics have been studied by Megson ( 7 7 , ZO),Hunter (17), and their associates. But because of uncertainties inherent in determining with a densitometer maxima from photographic negatives, these investigators were unable to detect a reproducible change in the pattern with increasing cure. In this work, however. a highresolution diffractometer permitted measurement of shifts in the x-ray scattering maxima in the neighborhood of 18' (2 8. where 8 is the Bragg angle) as a function of degree of cure. These shifts are well outside the range of experimental error.
Experimental Procedure Phenolic samples were prepared by a reaction of 776 grams of formaldehyde, 1000 grams of phenol (molar ratio, 0.921), and 20 ml. of a 20yo solution of sodium hydroxide. This mixture was then heated under reflux and at the first sign of cloudiness (after about 5 hours), the reaction was halted by removing the heat source. The mixture was washed twice and the water was removed by de1 Present address, Douglas Aircraft Co., Santa Monica, Calif.
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cantation and evaporation under vacuum. The resinous reaction product, weighing 928 grams, was poured into a shallow mold preheated in a press at 149' C. After 30 seconds, when production of volatiles indicated a reaction: a pressure of 1000 pounds per square inch was immediately applied and maintained for 4 hours. The molded samples were removed from the press and heated in an oven at 176' C. for varying periods up to 480 hours. The cure samples were wafers approximately 25 mm. square and 0.5 mm. thick. Spectra were determined with a North American Philips x-ray diffractometer equipped with current and voltage stabilizers. The source was run at 35 kv. and 20 ma. Copper K, radiation was filtered by a nickel plate 0.0075 inch thick and collimated with '/Z-degree divergence and scatter slits, and a 0.003inch receiving slit. Each wafer was placed in the scanning goniometer and curves of intensity of scattered radiation us. the scattering angle, 2 8, were plotted at the rate of '/4 degree per minute with a Brown recording potentiometer. An alternative method of scanning the maxima manually and measuring intensity with a count-rate meter gave no greater precision. Orientation of the sample wafer in the goniometer affected the scattering spectrum considerably. When samples were ground to a rougelike powder, however, the diffraction maximum for each sample could be readily located within 0.05' (2 8 ) . These orientation effects can be minimized by using a rotating sample holder.
Results Periodicity d (Table I) is calculated from the Bragg equation for first-order reflection nX = 2d sin 0
where X is the wave length of the x-rays. The most frequently occurring distance in the resin is about 1.22d (78). Thus, distances associated with scattering maxima increase with increasing degree of cure. The data in Table I are fitted approximately by the empirical equation. o/eo = 1 - 0.0128 log (1 t )
INDUSTRIAL AND ENGINEERING CHEMISTRY
+
where 00 is the Bragg angle for no cure and tis cure time in hours at 176' C. A radial distribution function (F'g 11 ure 3), calculated from the scattering curve for an uncured sample, is represented by
where D is the radial distribution func-
15
0
20
25
3
DEGREES 26
Figure 1. X-ray diffraction spectrum and i t s shift after 480-hour cure 190
1 m
c4
I
II
180 I
IO0
0
I
I
I 2 00
II
II
300
400
500
C U R E TIME HOURS
Figure 2. Diffraction cure time at 176" C.
maximum
VS.
6
4
A
2
0 -2 -4
-6 0
2
4
6
8
IO
12
14
16
18
r,A
Figure 3.
Radial distribution function
2C
Table II, X-Ray Diffraction Analysis of the Curing Process in Phenolics Sample
Cure Hr., 176' C .
Maximum,
0 26.25 113.75 189.0 250.5 294.25 421.75 480.5
18.79 18.48 18.32 18.24 18.21 18.18 18.18 18.16
-2
e
Periodicity, A.
4.719 4.797 4.838 4.859 4.867 4.875 4.875 4.880
c
tion and r is the interatomic distance in the resin which gives rise to scattering. The parameter s is related to the Bragg angle by the equation s =
4n
- sin 0
x
The function i(s) is derived from the experimental intensity after correction for polarization, absorption, and incoherent scattering. The method of calculating this curve is described by Klug and Alexander (78).
model of the phenolic resin where aromatic rings joined by methylene groups are assumed coplanar. Lengths of the aromatic and aliphatic C-C bonds are assumed to be 1.38 and 1.54 A., respectively, and the methylene group is assumed to exhibit the tetrahedral bond angle of 109" 28'. Steric hindrance from hydrogens ortho to the methylene group actually prevents coplanarity, and some of the longer distances represented by bars, therefore, are slightly underestimated. Nevertheless, good correspondence generally exists between locations of the bars and maxima of the scattering curve. The peak in the radial distribution function at 4.8 A., however. is higher than expected from arrangement of the bars drawn from the model. Therefore, this peak probably represents chain separation in the polymer. The most prominent maximum in polymer radial distribution curves is usually related to interchain distances (74, 78). Moreover, models of phenolic polymers indicate important distances of amroximatelv 5 A. between atoms in neighboring nonbonded phenolic groups. That these intermolecular distances increase with increasing cure time (Table I) is surprising because
shrinkage generally occurs during cure of phenolic resins ( 9 ) . This may be explained by cross linking during the cure. In the early stages of polymerization, the phenolic resin remains liquid and gelation occurs rather abruptly. This suggests that formation of chains rather than three-dimensional networks is favored a t first. The same conclusion is reached from statistical and steric considerations. I n the resin studied, the phenol-formaldehyde ratio of approximately unity is suitable for synthesis of linear structures. Therefore, at the beginning of cure, the resin exhibits the characteristic intermolecular distance observed in models when linear chains are placed side by .side. The curing process proceeds with further cross linking and ramification of the polymer, so that packing becomes less efficient. Gross shrinkage of the resin is attributed to loss of vclatiles and eliminatian of comparatively large lacunae.
Acknowledgment
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Discussion The vertical bars of Figure 3 represent interatomic distances for a simplified
The authors are indebted to Donald S. Stang for preparing the resin and to Murray Bloom for preparing the cured samples.
ROBERT A. SPURR, EDWARD H. ERATH, and HOWARD MYERS Research Laboratories, Hughes Aircraft Co., Culver City, Calif. DANIEL C. PEASE Department of Anatomy, University of California, Los Angeles, Calif.
Curing Process in Phenolic Resin
Electron-Microscopic Analysis Existence of micelles is verified by observation, but rather than being units having definite boundaries, they result from differences in cohesiveness
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FUNDAMENTAL difficulty in the theory of high polymer structure is lack of agreement between theory and experiment in estimating tensile strength. De Boer, for example, has calculated ( 9 ) the theoretical tensile strength for phenol-formaldehyde polymers as 4300 kg. per sq. mm. on the assumption that primary valence bonds are broken in rupture; the value is greater than 39 kg. per sq. mm. if it is assumed that only van der Waals forces are involved. He quotes an experimental value of 7.8 kg. per sq. mm. Therefore, the resin is less than one fifth as strong as would be expected from a calculation based on weak surface forces. To account for this discrepancy, Houwink (75, 76) adapted to resins the theory of Smekal (23), according to which discrepancies between calculated and observed strengths of crystalline materials
can be explained by the presence of Lockerstellen (structural defects). Near these defects, concentration of stresses leads to the low strengths observed. These defects are thought to arise from the chemical process of condensation, which begins simultaneously at a number of reactive centers. As long as the molecular weight increases by adding small molecules-e.g., dimethylolphenols-a molecule will probably be added at each active site of the growing network. As material of low molecular weight in the vicinity of the network is depleted, the reaction is retarded, because the probability that a large molecule will be properly oriented for addition is slight. The resulting structure is pictured as a collection of dense regions held together by methylene links to form a spongelike skeleton (isogel) having
cavities filled with a viscous liquid condensate of lower molecular weight. Megson ( 7 7 , 79) and others have provided support for this picture by measuring molecular models. An alternative theory, proposed by Stager (25) and others, was developed by studying thin films swollen by acetone. According to Stager, resin molecules build u p into three-dimensional, intermeshed spherocolloidal particles which are embedded in a matrix of lower molecular weight. He believes there is no need to assume that molecules must be linked by primary (methylene) bonds to give rise to a rigid skeleton; gelation is assumed to result from association and interpenetration of molecules. Stager's model does not, however, adequately account for the abrupt change in physical properties and in solubility experienced VOL. 49, NO. 1 1
NOVEMBER 1957
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