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Inelastic deformation of polyimide-copper thin films
Materials Science and Engineering, A142 ( 199'1 ) 135-144
135
Inelastic deformation of polyimide-copper thin films D. M. Shinozaki and A. Klauzner Department of Materials Engineering, University of Western Ontario, London, Ontario N6A 5B9 (Canada)
P. C. Cheng Advanced Microscopy and Imaging Laboratory, Department of Electrical and Computer Engineering, State University of New York, Buffalo, NY (U.S.A.)
(Received January 22, 1991 )
Abstract
Thin-layered composites of polyimide-copper are tested in tension. The specimens are prepared using standard thin film preparation techniques. The stress relaxation behaviour of the two phases is measured and modelled in terms of a single activated rate process. The internal stress in the polyimide is measured and increases with tensile strain. The deformed samples are examined with laser-scanning confocal microscopy and standard optical microscopy. Microcracks in copper and shear bands in polyimide are observed and correlated qualitatively with the stress relaxation model.
1. Introduction
Recent developments in the fabricz.tion of microelectronic devices have led to very-largescale integration (VLSI). The interconnect structure in VLSI chips has become increasingly densely packed, with metal conductors separated by layers made of insulators such as polymers. For these applications, polyimides have a number of attractive properties, including low dielectric constant, stability at high temperatures (of the order of 400 °C), processability using standard thin film preparation methods (such as spinning or spraying) and patternability with lithographic methods. The interconnect wiring which is in contact with the insulating polyimide is often copper and the composite behaviour of polyimidecopper is therefore of some practical interest. A significant problem which affects the reliability of the device is the failure of these kinds of multiphase layered structures by deformation and fracture at or near the interface between the different materials [1-4]. The stresses which cause such failures can be generated during the processing steps, which involve changes in temperature [5]. The difference in coefficient of thermal expansion between the metal and the polymer can result in quite high stresses, which can cause 0921-5093/91/$3.50
slow crack growth or other forms of mechanical breakdown of the interface regions. The mechanical properties of polymers generally are known to be strongly time dependent, and stresses can relax rapidly, even at room temperature [6]. It is important, therefore, to characterize the stress relaxation behaviour of polyimide films quantitatively. It is also important to examine the mechanical response of thin metallic copper prepared by evaporation on the polyimide surface. Although there is a considerable body of literature on the plastic deformation of single crystal and bulk polycrystalline copper, there is much less detailed information on deformation of evaporated films, particul.arly those deposited on polyimide by evaporation.
2. Experimental methods
The material was prepared by spinning Dupont PI-2545 on a silicon wafer (silicon oxide surface) and baking to a temperature of 380 °C in nitrogen using carefully controlled heating and cooling steps. The steps used in the curing cycle were similar to those used in standard industrial processes. The careful control of these steps was necessary since it was known that the final © Elsevier Sequoia/Printed in The Netherlands
136
chemical structure depends on the exact curing process [7, 8]. For comparison, some samples were cured in air using a similar thermal process. The thickness of the polyimide after curing was measured (7.8/~m) using a profilometer at specific points. The uniformity of the thickness over large areas of the specimen was confirmed using optical interference microscopy. Copper was deposited at room temperature on the polyimide by vacuum evaporation using a standard oil diffusion-pumped chamber. The rate of deposition was approximately 10 nm min-1 and the final thickness confirmed using a profilometer (0.37/~m). Samples were cut into strips 5 x 35 mm 2 and carefully mounted in an Instron. The tensile tests were run at a strain rate of 4 x 10 - 3 S -1. The stress relaxation tests were made in tension at room temperature (24.5 °C). The force on the specimen was measured as a function of time for up to 105 s. Because the measured temperature was consistently stable only for approximately 104 s, most of the data were recorded over this shorter time. A large number of tests were made on different specimens and the consistency of the observations was confirmed. Laser-scanning confocal microscopy has been used to examine the morphology of the interface decohesion. The polyimide-copper specimen was viewed after tensile testing from the polyimide side of the composite. Standard optical microscopy has also been used, but three-dimensional information could not be obtained. Confocal optical microscopy was found to yield much higher quality images. The method has been used extensively in the biological sciences and has been explained in detail [9, 10]. A brief description of the principles is given as follows. A diffraction-limited spot is focused at the focal plane of the objective lens on to a specimen using a small source aperture. The light reflected from the focal plane is scattered back along the optical path and passes through the source aperture (or its conjugate). Light which is reflected from planes above or below the exact focal plane is focused at planes different from the plane occupied by the source aperture and is stopped from reaching the detector. The focused spot is rastered across the specimen and the detector signal recorded as a digitized twodimensional image. By changing the position of the focal plane of the objective lens, a series of images at different depths can be recorded. The
digitized image can be reconstructed in a computer to reveal the three-dimensional structure
[11]. The present work used a BioRad MRC500 confocal optical microscope with a 25 mW dualline argon ion laser as source. The detailed description of the modifications to eliminate instrumental artefacts has been presented elsewhere [11]. The images were recorded on an AT bus 80386 microcomputer and reconstructed and examined on a SUN Microsystems 4/260 workstation with a TAAC-1 accelerator. Stereo pairs were produced by taking the series of images (of the order of 20-30 single two-dimensional images) and digitally shifting the images to produce an artificial parallax. The relative threedimensional positions of the various features could thus be examined.
3. Experimental results The tensile stress-strain behaviour for polyimide films depended on the curing process (Fig. 1). The ultimate tensile strength and ductility were measurably higher for the material cured in a nitrogen atmosphere. However, it was interesting to note that the small strain secant modulus was greater for the air-cured sample. The force and elongation were measured and the force in the copper was calculated by assuming a parallel-coupled composite and subtracting the force acting on the polyimide from that acting on the composite (Fig. 2). The composite tensile behaviour was dominated by the copper at strains less than or equal to 0.025, while at larger strains 140 /A
120
,~xoo
.
~
8060_
r~
i I/
40200
0.0
0.~1
0.12 Strain
0.f3
Fig. I. Tensile o-e curves for polyimidc cured in nitrogen (A) and air (D).
137 140 /-
1205
~o/
A
0"
~ ' ~ B
f
A
/
.~-., 100
/-
~ ~I 00-
f
o'
m
I
f
/
60-
i, •
/
~3 40_ I /
c
I/
,
0.:
Strain Fig. 2. Force vs. strain in each phase in the composite.The forces in the polyimide(A) and copper (C) are calculated assuming parallel coupling. The overall force on the composite is B. 1000
800
000
~ 400-
200-
B
0 .~1
0 .t2
0 .'3
Strain Fig. 3. o - e curves for polyimide (A) and copper (C). The nominal stress for the composite (polyimide-copper) is shown for comparison (B).
the polyimide work hardened enough to carry the larger part of the load. The stress could then be plotted as in Fig. 3. The yield stress in the evaporated thin film copper was approximately 850 MPa. The tensile yield stress increases in metals as the grain size decreases, and this relatively high strength was expected since the grain size was known to be extremely small in evaporated metal films [12]. The stress-strain behaviour of the polyimide was consistent with other work on similar material [1, 13]. Compared to other polymers, the polyimide used had high strength and reasonable ductility and macroscopically deforms in a uniform fashion to fracture with little necking.
20- i/ ./ 0
0.0
0.~1
0
.'2
0
.~3
Strain Fig. 4. Starting stresses and strains for stress relaxation tests for polyimide (A) and polyimide-copper (B).
Stress relaxation tests from polyimide and from polyimide-copper specimens were started at different tensile strains (Fig. 4). Two different kinds of tests have been compared. In one set the tests involved a number of similar specimens, each tested to a different strain. In the second set the tests involved only one specimen, tested in tension to increasing strains, and stress relaxed in steps. The results in both tests were identical, with similar relaxation curves which depended on the starting stress. The relaxation curves for each of the phases were measured as shown in Fig. 5. As the strain increased, the rate of stress relaxation increased in the polyimide phase, while that in the copper decreased very slightly. For polyimide at large strains (0.18 or above) the relaxation rate increased and the curves intersected, which was not expected for normal relaxation behaviour and was not consistent with the behaviour of other polymers. In additional tests which are not shown here, the relaxation rates for air-cured polyimide were found to be much slower than those for nitrogencured polyimide when compared at similar levels of initial stress. The overall ductility was also reduced for air curing. It was clear that the curing process must be controlled carefully to control residual stresses in the polyimide. In the discussion to follow, the polyimide used was nitrogen cured unless otherwise described. At small strains there were no visible changes in the optically examined microstructures in both polyimide and polyimide-copper at small strains. Near e = 0.05 a general, fine-scale roughening of
138 120 ,
~..I00~
0
(a)
o
~
~
~
~. ~ Ln(t)
Ln(t)
~
10
900
~
700.
0
0.14
500.
3O0
(b) o
~
1o
Fig. 5. (a) Stress relaxation as a function of initial strain in polyimide. (b) Stress relaxation in copper.
the copper was detected. This was due to the plastic deformation of the fine-grained copper. After fracture of the composite specimen ( e ~ 0.24), some microcracks were visible in the copper phase. Since this was a manifestation of the interface failure which is important in these composites, these cracks were examined. The confocal optical micrographs of the microcracks and decohesion regions are shown in Fig. 6. The images shown are single two-dimensional images taken at one focal plane position in the polyimide. Each image is one of a series taken as successive focal plane positions through the polyimide. Each series has been reconstructed to produce a three-dimensional view. The stereo pair reconstructions are not shown here but the detailed interpretation of the images is given in Section 4. Micrographs using a standard optical microscope are shown for comparison in Fig. 7.
Fig. 6. Confocal optical micrographs of polyimide-copper after fracture. Microcracks are visible in the copper.
4. Analysis and discussion For the stress relaxation test the machine is assumed to have a stiffness much larger than the specimen. In this case the stress relaxes as the elastic strain in the specimen is converted to plastic strain. The plastic strain rate is propor-
139
tion in the reverse direction is generally much smaller than that in the forward direction. The forward plastic strain rate increases as the stress increases since the effective barrier is diminished: F - wa* /
]
gplastic = gC exp
(3)
where to is the shear strain activation volume [14] and a * = O'applied --O'internal is the effective stress [15]. The internal stress (O'internal) is the component of stress which opposes the forward operation of the plastic deformation mechanism. For a stress relaxation test the stress can be calculated as a function of time from eqns. (1) and (3) and is of the form [15]
o,=F+kTln( kT I_kTln(t+a) e) ~o ~Egcw] w
(4)
where t is the time and a is a constant of the form [16] kr Fig. 7. Standard optical micrographs of polyimide-copper taken under crossed polars. Faint shear bands in the polyimide are visible oriented at about 45 ° to the cracks.
tional to the measured rate of change of stress: 1 gplastic = -- E 0
(1)
The elastic strain is converted to plastic strain and the stress thus decays with time. The plastic strain rate in eqn. (1) depends on the exact mechanisms of deformation, which can be modelled as single activated rate processes. The plastic strain rate is related to the rate of transition of the entity involved in the deformation mechanism over an activation energy barrier: gplastic = gC exp - ~-~
(2)
where gplastic is the plastic strain rate, gc is the maximum attainable strain rate and F is the standard free energy of activation. The factor ~c involves the vibrational frequency, the strain associated with the operation of each mobile entity and the number of such mobile entities. For a finite applied stress the rate of plastic deforma-
a:gcEwexpk
-k7 ]
(5)
For a time-independent internal stress the applied stress can be used in eqn. (4) for most parts of the analysis, particularly those which deal with the activation volume, which is derived from the slope of the plotted version of eqn. (4). However, it should be noted that the driving force for the plastic relaxation mechanism is the effective stress and not the applied stress. The non-linearity in the o* vs. in(t) curve at short times is normally observed in stress relaxation of polymers and can be accounted for in this model through a. This constant is equal to the deviation between the relaxation data at short times in the test and the linear fit to the long-time curve (plotted as a vs. In(t)). At long times in the relaxation test the logarithmic time scale makes the deviation relatively insignificant [16]. The stress relaxation data obtained for both polyimide and polyimidecopper have been found in this work to be compatible with eqn. (5). The lineafity of the curves in Fig. 5 indicates that eqn. (4) is a reasonable function with which to describe the relaxation kinetics for both polyimide and polyimide-copper. At small strains the relaxation behaviour is as expected for both copper and polyimide phases. The rate of stress
140
decay increases as the initial stress increases. The copper relaxation rate decreases with increasing tensile strain because the copper phase shows strain softening (Fig. 3). At large strains (0.18 or more) the relaxation rate in the polyimide increases anomalously and the o vs. In(t) curves intersect the relaxation for lower initial stresses (Fig. 5(a)). The plastic deformation process which accommodates the stress relaxation is thus different at high and low strains. For a single kind of plastic mechanism it is expected that the stress relaxes according to eqn. (4) and that the slope of the curve is inversely proportional to the shear strain activation volume (to). In Fig. 8 to is plotted as a function of strain for polyimide. The activation volume is calculated in this figure only at strains less than or equal to 0.18, over which range the relaxation curves are well behaved. The activation volume changes smoothly with strain for a given polyimide, although for different cure cycles the values are different. This is expected since the molecular structure is different. The changes in activation volume with stress have been discussed by Kubat and Rigdahl in their review [17], but a rationale based on the dislocation analogue of plastic deformation is presented below.
4.1. Activation volume and plastic deformation mechanisms In metals the interpretation of the shear strain activation volume has been fully developed in terms of crystallographic dislocation motion [18, 19]. It is useful to describe this to show that the experimental observations in tests on polyimide
are consistent with a specific deformation mechanism. In Fig. 9 a dislocation segment moves up against an obstacle to position 2 under the influence of a stress. Enough energy must be supplied to the segment to move it to position 4, after which it passes the obstacle. The rate of change of the free energy with stress gives the shear stress activation volume
¢}F kTI t~ In(e)l
=-ao
(6)
/
This volume is equal to the area of the slip plane swept out by the dislocation in moving from 2 to 4 (the activation area) multiplied by the Burgers vector. The activation area so defined therefore depends on the stress and will decrease with increasing stress. This general relationship has been discussed by Saimoto and Sang in studies of the dislocation models of deformation in metals [20]. It has also been discussed earlier by Balasubramanian and Li [18] with reference to metals and by Li et al. with reference to polymers [14]. In this latter paper the mechanistic rationale to explain the applicability of this formalism to polymers is found in the possibility of describing plastic deformation as a shear band propagation process. Such amorphous dislocations had been proposed by Gilman [21]. In materials such as polystyrene, slip lines can be observed to terminate in the specimen. The boundary between slipped and unslipped regions is then similar elastically to a dislocation. It is shown in micrographs in a later section that such slip lines can be found in polyimide.
lOD
1234
v
0
Activation Area
84-
Obstacle Dislocat Movemen~
0
o.oo
0.65
o.lo
o.I~
o.zo
Strain Fig. 8. Shear strain activation volume (to) plotted as a function of strain for polyimide cured in nitrogen (A) and air (D).
/ 2
4
v
Position
Fig. 9. Schematic diagram to show that the shear strain activation volume is expected to decrease as the applied stress increases. (Following Balasubramanian and Li [18].)
141 The stress dependence of the activation volume is shown for polyimide and for evaporated copper in Figs. 10 and 11 respectively. The data were calculated only for relatively small strains ( e ~ 0 . 1 ) for which the observed mechanical response of the specimens were clearly well behaved. At larger strains, as mentioned earlier, anomalous stress relaxation was observed. The activation volumes for both the polyimide and the evaporated copper are inversely proportional to stress. This plastic response is consistent with the interpretation above and, if compared to the detailed models described by Li et al. [14], predicts a power-law relationship between the strain rate and stress.
6
45-
4.2. Internal stresses in polyimide The internal s t r e s s ( a i n t . . . . 1) in the polyimide can be calculated using White's approach, in which the intercept of the plot of ( 1/ t) ( d o / d In(t)) vs. a is the internal stress [22, 23] (Fig. 12). The non-linearity of the plots at small values of the dependent variable suggests that as time increases in the stress relaxation test, the stress does not saturate to a fixed internal stress but rather continues to decrease. This would be consistent with a time-dependent internal stress, which decays slowly, and which has been described earlier for both polyethylene [24] and polypropylene [25]. The internal stress is thus measured from the extrapolation of the straight line portion of the plot (Fig. 12) and is a measure of the internal stress component of the flow stress (Fig. 13). The work hardening in the polyimide is thus largely dominated by an increase in the internal stress in the material. The stress-strain curve can be partitioned into two components using a model described in detail for crystalline polymers [24,
25]: 0
~3-
O'applie d =
0
(7)
It is interesting to note, however, that while the internal stress increases with increasing strain, the other component (Oo) also increases. This is in contrast to the behaviour observed in polyethylene, in which Oo remains constant.
:,32t~
0 O. 1
0 .~2
1/0- (MPa-')
0.3
Fig. 10. In polyimide the shear strain activation volume is found to be inverselyproportional to the effectivestress.
g"
O o Jr- O'internal
4.3. Microcracks in polyimide-copper The stress relaxation curves for the polyimide-copper specimens did not show any sudden
0.40
30000-
"-"0.88 "~200000 ~ 0.36 0
0 .,=i
"--'~10000" I
.~ 0.34 .,o
0.32 0.001
i
i
,
I
i
i
i
i
0 J/ 0
i 0.002
1/0-
(MPa-')
Fig. 11. In copper the shear strain activation volume is relatively insensitiveto the stress.
2b
40
6'o
a'o
160
120
Stress (MPa) Fig. 12. Measurement of internal stress in polyimide as a function of strain (increasingtensile strain from left to right) using White's method [22, 25].
142
120 .w t
100
/~*"
I
*
0
I 0 I
flOg
0
I0
/g
-
/
m 60II1
/
-
/
,,o r~ 4-0-
/
20- /
/
,f 0.0
0.~1
0.~2
0.'3
Strain
stress in polyimide.
Fig. 14. Schematic diagram of microcracks opening in copper at A and generating large amounts of shear deformation along bands D.
change in inelastic processes even at strains up to 0.18 (Fig. 5(b)). The microcracks seen in Figs. 6 and 7 are generated at strains greater than 0.18 (Fig. 4) and are therefore immediate precursors to final failure. The micrographs were taken in reflected light, but if compared to images in transmitted light, showed clearly that the copper is cracked along directions perpendicular to the principal tensile stress. It is interesting to note that the thin microcrack has a large hole in the centre. Some of the copper is lost in the cracking process by crack branching and joining. The polyimide partially decoheres in the immediate vicinity of the crack in the copper (point A in Fig. 6(a)). The clear, closely spaced interference fringes in this region arise from the variation in spacing between the polyimide and the copper layers. The spacing can be calculated from the wavelength of the light and is of the order of 1 pm. These kinds of fringes are visible in normal optical microscopy. The decohesion spacing measured in such images is not simply related to the micromechanics of the failure process, since the specimen cracks and delaminates in the interface while under stress, and these images are recorded after unloading. The retractive strains clearly would vary in the region near the cracks. There is a second kind of interference fringe visible over the entire field of view in Fig. 6(b). Comparison with images examined in an interference microscope reveal that these are thickness fringes in the polyimide. The spun films are extremely uniform in thickness and normally produce one set of parallel fringes when
examined on a flat substrate such as a silicon wafer. The spacing between the fringes is a sensitive function of the thickness of the polyimide and the relative orientation of the specimen with respect to the optic axis of the microscope. The cracked and unloaded polyimide-copper composite produces local bending and distortion of the film, which results in the local changes in fringe spacing. In a monotonically loaded specimen this would be a convenient method to examine local strains during testing. This interpretation was confirmed by examining the three-dimensional images. The confocal micrographs in Fig. 6 are single two-dimensional images which are part of a series taken at different depths in the polyimide. The three-dimensional digital reconstruction of the entire series reveals the specimen strongly bulging in the immediate vicinity of each crack. This is due to the inhomogeneous retractive strains in the specimen on unloading. There is a third observation to be made at points such as B in Fig. 7(b). The crack in the copper is associated invariably with a triangular region of low contrast which represents a region of relatively large strain. This can be deduced by examining a number of such regions and observing that the regions are associated with the sides of the cracked copper. The mechanical response in these regions can be explained qualitatively as follows considering the schematic diagram in Fig. 14. The parallel coupled copper and polyimide carry identical total strains, although the elastic and plastic components are different and depend on the stress relaxation kinetics of each phase.
Fig. 13. Internal stress (open circles) compared to flow
143
The copper cracks and the polyimide locally stretches and work hardens to carry the additional load transferred from the copper. The polyimide is not fractured in these regions, so will be deformed locally more than the polyimide far away from the cracked copper. The difference in total strain between the copper and the polyimide varies from zero at a large distance from the crack to a large value at the crack (the point labelled A). In a direction along the applied stress direction, perpendicular to the copper crack, the polyimide is drawn into the crack region and undergoes a large amount of deformation. These regions are in many cases associated with shear bands which extend a considerable distance (several crack lengths) away from the crack tips and then stop. The contrast is extremely low and they can be seen as diffuse bands emanating from the crack tips at about 45 ° (Fig. 7(a)). At low magnification these bands are observed near each of the copper cracks. Micromechanically they are similar to the shear bands observed at the tip of a craze. The shear bands in polyimide-copper are observed with a polarizer inserted into the optical path, and the contrast can be reversed from darker than background (as shown in Fig. 7(b)) to lighter than background by rotating the orientation of the polarizer. In the deformed sample the overall brightness of the entire specimen can also be changed in this way, but not as strongly as observed in the shear band region. This implies a significant molecular orientation in these bands and in the triangular regions, consistent with the mechanical description above. It suggests that one mechanism of deformation in polyimide is shear band propagation, which would be consistent with a dislocation analogue suggested for polymers by Li et al. [14]. This in turn supports the single activated rate model suggested from the stress relaxation data.
could be modelled using a single activated rate process up to strains of the order of 0.18. This was consistent with the microstructural observation that shear bands were seen in the deformed samples. The relaxation rate in the polyimide became anomalously high at strains greater than this, indicating a change in deformation mechanism. The stress relaxation behaviour of the copper has been determined by assuming a simple parallel coupling of the metal and polymer phases. The composite sample failure was preceded by generalized microcracking in the copper. Each microcrack has associated with it a zone of shear and a region of delamination at the interface between the polyimide and copper. The microcracking and delamination were examined using laser-scanning confocal optical microscopy. The shear bands emanated from the ends of the cracks in the copper. These observations support the models of Ho and coworkers which relate the decohesion energies of the polyimide-copper interface to the tensile stress-strain curves [1, 2]. It is clear from the present work, however, that a considerable plastic work is generated by the polyimide shear band formation near each microcrack.
Acknowledgments The work has been supported by the Natural Sciences and Engineering Research Council of Canada. A.K. has been on sabbatical from R A F A E L in Israel and has been supported by the Ontario Center for Materials Research (A. E. Hamielec and J. D. Embury). Considerable assistance from Dr. Paul Ho of IBM (T. J. Watson Research Center) is gratefully acknowledged. The confocal optical microscopy has been supported by grants from the U.S. D O E (DE-AE-0888DP10782 and EE-FG03-89SF18012), the Whitaker Foundation, CAT-HIDI and NFWWTP.
5. Conclusions Composite specimens consisting of thin layers of polyimide and copper have been tested in tension and the inelastic deformation processes up to fracture have been studied. The composite films deformed uniformly on the macroscopic scale to quite large strains. The yield stress of the evaporated thin film copper was approximately 850 MPa while that of the polyimide was about 100 MPa. In polyimide the relaxation of stress
References 1 E S. Ho and E Faupel, Appl. Phys. Lett., 53(17) (1988) 1602. 2 S. T. Chen, C. H. Yang, F. Faupel and E S. Ho, J. AppL Phys., 64 (12) (1988) 6690. 3 Y. H. Kim, G. E Walker, J. Kim and J. Park, J. Adhesion Sci. Technol., 1 (1987) 331. 4 Y. H. Kim, J. Kim, G. F. Walker, C. Feger and S. P. Kowalczyk, J. Adhesion Sci. Technol., 2 (1988) 95. 5 C. Goldsmith, E Geldermans, E Bedetti and G. A.
144 Walker, J. Vac. Sci. Technol. A, 1 (2) (1983) 407. 6 I. M. Ward, Mechanical Properties of Solid Polymers, Wiley-lnterscience, New York, 1979, p. 142. 7 M. J. Brekner and C. Feger, J. Polym. Sci. A, 25 (1987) 2005. 8 M. J. Brekner and C. Feger, J. Polym. Sci. A, 25 (1987) 2479. 9 S. Inoue, in J. Pawley (ed.), The Handbook of Biological Confocal Microscopy, IMR Press, WI, 1989, p. 1. 10 J. G. White, W. B. Amos and E Fordham, J. Cell BioL,
105(1987)41. 11 D. M. Shinozaki, P. C. Cheng, S. Haridoss, R. Mitchell and A. Fenster, J. Mater. Sci., in the press. 12 R. Abermann and R. Koch, Thin Solid Films, 129(1985) 71. 13 E Maseeh and S. D. Senturia, in C. Feger, M. M. Khojzateh and J. E. McGrath (eds.), Polyimides: Materials, Chemistry and Characterization, Elsevier, Amsterdam, 1989, p. 575. 14 J. C. M. Li, C. A. Pampillo and L. A. Davis, in H. H. Kausch, J. A. Hassell and R. I. Jaffee (eds.), Deformation and Fracture of High Polymers, Plenum, New York,
1973, p. 239. 15 E Guiu and P. L. Pratt, Phys. Status Solidi, 6 (1964) 111. 16 D. M. Shinozaki, G. W. Groves and R. G. C. Arridge, Mater. Sci. Eng., 28(1977) 119. 17 J. Kubat and M. Rigdahl, in W. Brostow and R. D. Corneliussen (eds.), Failure of Plastics, Hanser, Munich, 1986, p. 60. 18 N. Balasumbramanian and J. C. M. Li, J. Mater. Sci., 5 (1970) 434. 19 D. M. Shinozaki and A. Klauzner, Proc. Materials
Research Soc., Electronic Packaging Materials Science, November 1990, Materials Research Society, Pittsburgh, PA, 1991. 20 S. Saimoto and H. Sang, Acta Metall., 31 (1983) 1873. 21 J.J. Gilman, in A. R. Rosenfield et al. (eds.), Dislocation Dynamics, McGraw-Hill, New York, 1968, p. 3. 22 J.R. White, Mater. Sci. Eng., 45 (1980) 35. 23 J. R. White, J. Mater. Sci., 16(1981) 3249. 24 D. M. Shinozaki and C. M. Sargent, Mater. Sci. Eng., 35 (1978) 213. 25 D. M. Shinozaki, C. M. Sargent and S. Fehr, Mater. Sci. Eng., 51 (1981) 93.
135
Inelastic deformation of polyimide-copper thin films D. M. Shinozaki and A. Klauzner Department of Materials Engineering, University of Western Ontario, London, Ontario N6A 5B9 (Canada)
P. C. Cheng Advanced Microscopy and Imaging Laboratory, Department of Electrical and Computer Engineering, State University of New York, Buffalo, NY (U.S.A.)
(Received January 22, 1991 )
Abstract
Thin-layered composites of polyimide-copper are tested in tension. The specimens are prepared using standard thin film preparation techniques. The stress relaxation behaviour of the two phases is measured and modelled in terms of a single activated rate process. The internal stress in the polyimide is measured and increases with tensile strain. The deformed samples are examined with laser-scanning confocal microscopy and standard optical microscopy. Microcracks in copper and shear bands in polyimide are observed and correlated qualitatively with the stress relaxation model.
1. Introduction
Recent developments in the fabricz.tion of microelectronic devices have led to very-largescale integration (VLSI). The interconnect structure in VLSI chips has become increasingly densely packed, with metal conductors separated by layers made of insulators such as polymers. For these applications, polyimides have a number of attractive properties, including low dielectric constant, stability at high temperatures (of the order of 400 °C), processability using standard thin film preparation methods (such as spinning or spraying) and patternability with lithographic methods. The interconnect wiring which is in contact with the insulating polyimide is often copper and the composite behaviour of polyimidecopper is therefore of some practical interest. A significant problem which affects the reliability of the device is the failure of these kinds of multiphase layered structures by deformation and fracture at or near the interface between the different materials [1-4]. The stresses which cause such failures can be generated during the processing steps, which involve changes in temperature [5]. The difference in coefficient of thermal expansion between the metal and the polymer can result in quite high stresses, which can cause 0921-5093/91/$3.50
slow crack growth or other forms of mechanical breakdown of the interface regions. The mechanical properties of polymers generally are known to be strongly time dependent, and stresses can relax rapidly, even at room temperature [6]. It is important, therefore, to characterize the stress relaxation behaviour of polyimide films quantitatively. It is also important to examine the mechanical response of thin metallic copper prepared by evaporation on the polyimide surface. Although there is a considerable body of literature on the plastic deformation of single crystal and bulk polycrystalline copper, there is much less detailed information on deformation of evaporated films, particul.arly those deposited on polyimide by evaporation.
2. Experimental methods
The material was prepared by spinning Dupont PI-2545 on a silicon wafer (silicon oxide surface) and baking to a temperature of 380 °C in nitrogen using carefully controlled heating and cooling steps. The steps used in the curing cycle were similar to those used in standard industrial processes. The careful control of these steps was necessary since it was known that the final © Elsevier Sequoia/Printed in The Netherlands
136
chemical structure depends on the exact curing process [7, 8]. For comparison, some samples were cured in air using a similar thermal process. The thickness of the polyimide after curing was measured (7.8/~m) using a profilometer at specific points. The uniformity of the thickness over large areas of the specimen was confirmed using optical interference microscopy. Copper was deposited at room temperature on the polyimide by vacuum evaporation using a standard oil diffusion-pumped chamber. The rate of deposition was approximately 10 nm min-1 and the final thickness confirmed using a profilometer (0.37/~m). Samples were cut into strips 5 x 35 mm 2 and carefully mounted in an Instron. The tensile tests were run at a strain rate of 4 x 10 - 3 S -1. The stress relaxation tests were made in tension at room temperature (24.5 °C). The force on the specimen was measured as a function of time for up to 105 s. Because the measured temperature was consistently stable only for approximately 104 s, most of the data were recorded over this shorter time. A large number of tests were made on different specimens and the consistency of the observations was confirmed. Laser-scanning confocal microscopy has been used to examine the morphology of the interface decohesion. The polyimide-copper specimen was viewed after tensile testing from the polyimide side of the composite. Standard optical microscopy has also been used, but three-dimensional information could not be obtained. Confocal optical microscopy was found to yield much higher quality images. The method has been used extensively in the biological sciences and has been explained in detail [9, 10]. A brief description of the principles is given as follows. A diffraction-limited spot is focused at the focal plane of the objective lens on to a specimen using a small source aperture. The light reflected from the focal plane is scattered back along the optical path and passes through the source aperture (or its conjugate). Light which is reflected from planes above or below the exact focal plane is focused at planes different from the plane occupied by the source aperture and is stopped from reaching the detector. The focused spot is rastered across the specimen and the detector signal recorded as a digitized twodimensional image. By changing the position of the focal plane of the objective lens, a series of images at different depths can be recorded. The
digitized image can be reconstructed in a computer to reveal the three-dimensional structure
[11]. The present work used a BioRad MRC500 confocal optical microscope with a 25 mW dualline argon ion laser as source. The detailed description of the modifications to eliminate instrumental artefacts has been presented elsewhere [11]. The images were recorded on an AT bus 80386 microcomputer and reconstructed and examined on a SUN Microsystems 4/260 workstation with a TAAC-1 accelerator. Stereo pairs were produced by taking the series of images (of the order of 20-30 single two-dimensional images) and digitally shifting the images to produce an artificial parallax. The relative threedimensional positions of the various features could thus be examined.
3. Experimental results The tensile stress-strain behaviour for polyimide films depended on the curing process (Fig. 1). The ultimate tensile strength and ductility were measurably higher for the material cured in a nitrogen atmosphere. However, it was interesting to note that the small strain secant modulus was greater for the air-cured sample. The force and elongation were measured and the force in the copper was calculated by assuming a parallel-coupled composite and subtracting the force acting on the polyimide from that acting on the composite (Fig. 2). The composite tensile behaviour was dominated by the copper at strains less than or equal to 0.025, while at larger strains 140 /A
120
,~xoo
.
~
8060_
r~
i I/
40200
0.0
0.~1
0.12 Strain
0.f3
Fig. I. Tensile o-e curves for polyimidc cured in nitrogen (A) and air (D).
137 140 /-
1205
~o/
A
0"
~ ' ~ B
f
A
/
.~-., 100
/-
~ ~I 00-
f
o'
m
I
f
/
60-
i, •
/
~3 40_ I /
c
I/
,
0.:
Strain Fig. 2. Force vs. strain in each phase in the composite.The forces in the polyimide(A) and copper (C) are calculated assuming parallel coupling. The overall force on the composite is B. 1000
800
000
~ 400-
200-
B
0 .~1
0 .t2
0 .'3
Strain Fig. 3. o - e curves for polyimide (A) and copper (C). The nominal stress for the composite (polyimide-copper) is shown for comparison (B).
the polyimide work hardened enough to carry the larger part of the load. The stress could then be plotted as in Fig. 3. The yield stress in the evaporated thin film copper was approximately 850 MPa. The tensile yield stress increases in metals as the grain size decreases, and this relatively high strength was expected since the grain size was known to be extremely small in evaporated metal films [12]. The stress-strain behaviour of the polyimide was consistent with other work on similar material [1, 13]. Compared to other polymers, the polyimide used had high strength and reasonable ductility and macroscopically deforms in a uniform fashion to fracture with little necking.
20- i/ ./ 0
0.0
0.~1
0
.'2
0
.~3
Strain Fig. 4. Starting stresses and strains for stress relaxation tests for polyimide (A) and polyimide-copper (B).
Stress relaxation tests from polyimide and from polyimide-copper specimens were started at different tensile strains (Fig. 4). Two different kinds of tests have been compared. In one set the tests involved a number of similar specimens, each tested to a different strain. In the second set the tests involved only one specimen, tested in tension to increasing strains, and stress relaxed in steps. The results in both tests were identical, with similar relaxation curves which depended on the starting stress. The relaxation curves for each of the phases were measured as shown in Fig. 5. As the strain increased, the rate of stress relaxation increased in the polyimide phase, while that in the copper decreased very slightly. For polyimide at large strains (0.18 or above) the relaxation rate increased and the curves intersected, which was not expected for normal relaxation behaviour and was not consistent with the behaviour of other polymers. In additional tests which are not shown here, the relaxation rates for air-cured polyimide were found to be much slower than those for nitrogencured polyimide when compared at similar levels of initial stress. The overall ductility was also reduced for air curing. It was clear that the curing process must be controlled carefully to control residual stresses in the polyimide. In the discussion to follow, the polyimide used was nitrogen cured unless otherwise described. At small strains there were no visible changes in the optically examined microstructures in both polyimide and polyimide-copper at small strains. Near e = 0.05 a general, fine-scale roughening of
138 120 ,
~..I00~
0
(a)
o
~
~
~
~. ~ Ln(t)
Ln(t)
~
10
900
~
700.
0
0.14
500.
3O0
(b) o
~
1o
Fig. 5. (a) Stress relaxation as a function of initial strain in polyimide. (b) Stress relaxation in copper.
the copper was detected. This was due to the plastic deformation of the fine-grained copper. After fracture of the composite specimen ( e ~ 0.24), some microcracks were visible in the copper phase. Since this was a manifestation of the interface failure which is important in these composites, these cracks were examined. The confocal optical micrographs of the microcracks and decohesion regions are shown in Fig. 6. The images shown are single two-dimensional images taken at one focal plane position in the polyimide. Each image is one of a series taken as successive focal plane positions through the polyimide. Each series has been reconstructed to produce a three-dimensional view. The stereo pair reconstructions are not shown here but the detailed interpretation of the images is given in Section 4. Micrographs using a standard optical microscope are shown for comparison in Fig. 7.
Fig. 6. Confocal optical micrographs of polyimide-copper after fracture. Microcracks are visible in the copper.
4. Analysis and discussion For the stress relaxation test the machine is assumed to have a stiffness much larger than the specimen. In this case the stress relaxes as the elastic strain in the specimen is converted to plastic strain. The plastic strain rate is propor-
139
tion in the reverse direction is generally much smaller than that in the forward direction. The forward plastic strain rate increases as the stress increases since the effective barrier is diminished: F - wa* /
]
gplastic = gC exp
(3)
where to is the shear strain activation volume [14] and a * = O'applied --O'internal is the effective stress [15]. The internal stress (O'internal) is the component of stress which opposes the forward operation of the plastic deformation mechanism. For a stress relaxation test the stress can be calculated as a function of time from eqns. (1) and (3) and is of the form [15]
o,=F+kTln( kT I_kTln(t+a) e) ~o ~Egcw] w
(4)
where t is the time and a is a constant of the form [16] kr Fig. 7. Standard optical micrographs of polyimide-copper taken under crossed polars. Faint shear bands in the polyimide are visible oriented at about 45 ° to the cracks.
tional to the measured rate of change of stress: 1 gplastic = -- E 0
(1)
The elastic strain is converted to plastic strain and the stress thus decays with time. The plastic strain rate in eqn. (1) depends on the exact mechanisms of deformation, which can be modelled as single activated rate processes. The plastic strain rate is related to the rate of transition of the entity involved in the deformation mechanism over an activation energy barrier: gplastic = gC exp - ~-~
(2)
where gplastic is the plastic strain rate, gc is the maximum attainable strain rate and F is the standard free energy of activation. The factor ~c involves the vibrational frequency, the strain associated with the operation of each mobile entity and the number of such mobile entities. For a finite applied stress the rate of plastic deforma-
a:gcEwexpk
-k7 ]
(5)
For a time-independent internal stress the applied stress can be used in eqn. (4) for most parts of the analysis, particularly those which deal with the activation volume, which is derived from the slope of the plotted version of eqn. (4). However, it should be noted that the driving force for the plastic relaxation mechanism is the effective stress and not the applied stress. The non-linearity in the o* vs. in(t) curve at short times is normally observed in stress relaxation of polymers and can be accounted for in this model through a. This constant is equal to the deviation between the relaxation data at short times in the test and the linear fit to the long-time curve (plotted as a vs. In(t)). At long times in the relaxation test the logarithmic time scale makes the deviation relatively insignificant [16]. The stress relaxation data obtained for both polyimide and polyimidecopper have been found in this work to be compatible with eqn. (5). The lineafity of the curves in Fig. 5 indicates that eqn. (4) is a reasonable function with which to describe the relaxation kinetics for both polyimide and polyimide-copper. At small strains the relaxation behaviour is as expected for both copper and polyimide phases. The rate of stress
140
decay increases as the initial stress increases. The copper relaxation rate decreases with increasing tensile strain because the copper phase shows strain softening (Fig. 3). At large strains (0.18 or more) the relaxation rate in the polyimide increases anomalously and the o vs. In(t) curves intersect the relaxation for lower initial stresses (Fig. 5(a)). The plastic deformation process which accommodates the stress relaxation is thus different at high and low strains. For a single kind of plastic mechanism it is expected that the stress relaxes according to eqn. (4) and that the slope of the curve is inversely proportional to the shear strain activation volume (to). In Fig. 8 to is plotted as a function of strain for polyimide. The activation volume is calculated in this figure only at strains less than or equal to 0.18, over which range the relaxation curves are well behaved. The activation volume changes smoothly with strain for a given polyimide, although for different cure cycles the values are different. This is expected since the molecular structure is different. The changes in activation volume with stress have been discussed by Kubat and Rigdahl in their review [17], but a rationale based on the dislocation analogue of plastic deformation is presented below.
4.1. Activation volume and plastic deformation mechanisms In metals the interpretation of the shear strain activation volume has been fully developed in terms of crystallographic dislocation motion [18, 19]. It is useful to describe this to show that the experimental observations in tests on polyimide
are consistent with a specific deformation mechanism. In Fig. 9 a dislocation segment moves up against an obstacle to position 2 under the influence of a stress. Enough energy must be supplied to the segment to move it to position 4, after which it passes the obstacle. The rate of change of the free energy with stress gives the shear stress activation volume
¢}F kTI t~ In(e)l
=-ao
(6)
/
This volume is equal to the area of the slip plane swept out by the dislocation in moving from 2 to 4 (the activation area) multiplied by the Burgers vector. The activation area so defined therefore depends on the stress and will decrease with increasing stress. This general relationship has been discussed by Saimoto and Sang in studies of the dislocation models of deformation in metals [20]. It has also been discussed earlier by Balasubramanian and Li [18] with reference to metals and by Li et al. with reference to polymers [14]. In this latter paper the mechanistic rationale to explain the applicability of this formalism to polymers is found in the possibility of describing plastic deformation as a shear band propagation process. Such amorphous dislocations had been proposed by Gilman [21]. In materials such as polystyrene, slip lines can be observed to terminate in the specimen. The boundary between slipped and unslipped regions is then similar elastically to a dislocation. It is shown in micrographs in a later section that such slip lines can be found in polyimide.
lOD
1234
v
0
Activation Area
84-
Obstacle Dislocat Movemen~
0
o.oo
0.65
o.lo
o.I~
o.zo
Strain Fig. 8. Shear strain activation volume (to) plotted as a function of strain for polyimide cured in nitrogen (A) and air (D).
/ 2
4
v
Position
Fig. 9. Schematic diagram to show that the shear strain activation volume is expected to decrease as the applied stress increases. (Following Balasubramanian and Li [18].)
141 The stress dependence of the activation volume is shown for polyimide and for evaporated copper in Figs. 10 and 11 respectively. The data were calculated only for relatively small strains ( e ~ 0 . 1 ) for which the observed mechanical response of the specimens were clearly well behaved. At larger strains, as mentioned earlier, anomalous stress relaxation was observed. The activation volumes for both the polyimide and the evaporated copper are inversely proportional to stress. This plastic response is consistent with the interpretation above and, if compared to the detailed models described by Li et al. [14], predicts a power-law relationship between the strain rate and stress.
6
45-
4.2. Internal stresses in polyimide The internal s t r e s s ( a i n t . . . . 1) in the polyimide can be calculated using White's approach, in which the intercept of the plot of ( 1/ t) ( d o / d In(t)) vs. a is the internal stress [22, 23] (Fig. 12). The non-linearity of the plots at small values of the dependent variable suggests that as time increases in the stress relaxation test, the stress does not saturate to a fixed internal stress but rather continues to decrease. This would be consistent with a time-dependent internal stress, which decays slowly, and which has been described earlier for both polyethylene [24] and polypropylene [25]. The internal stress is thus measured from the extrapolation of the straight line portion of the plot (Fig. 12) and is a measure of the internal stress component of the flow stress (Fig. 13). The work hardening in the polyimide is thus largely dominated by an increase in the internal stress in the material. The stress-strain curve can be partitioned into two components using a model described in detail for crystalline polymers [24,
25]: 0
~3-
O'applie d =
0
(7)
It is interesting to note, however, that while the internal stress increases with increasing strain, the other component (Oo) also increases. This is in contrast to the behaviour observed in polyethylene, in which Oo remains constant.
:,32t~
0 O. 1
0 .~2
1/0- (MPa-')
0.3
Fig. 10. In polyimide the shear strain activation volume is found to be inverselyproportional to the effectivestress.
g"
O o Jr- O'internal
4.3. Microcracks in polyimide-copper The stress relaxation curves for the polyimide-copper specimens did not show any sudden
0.40
30000-
"-"0.88 "~200000 ~ 0.36 0
0 .,=i
"--'~10000" I
.~ 0.34 .,o
0.32 0.001
i
i
,
I
i
i
i
i
0 J/ 0
i 0.002
1/0-
(MPa-')
Fig. 11. In copper the shear strain activation volume is relatively insensitiveto the stress.
2b
40
6'o
a'o
160
120
Stress (MPa) Fig. 12. Measurement of internal stress in polyimide as a function of strain (increasingtensile strain from left to right) using White's method [22, 25].
142
120 .w t
100
/~*"
I
*
0
I 0 I
flOg
0
I0
/g
-
/
m 60II1
/
-
/
,,o r~ 4-0-
/
20- /
/
,f 0.0
0.~1
0.~2
0.'3
Strain
stress in polyimide.
Fig. 14. Schematic diagram of microcracks opening in copper at A and generating large amounts of shear deformation along bands D.
change in inelastic processes even at strains up to 0.18 (Fig. 5(b)). The microcracks seen in Figs. 6 and 7 are generated at strains greater than 0.18 (Fig. 4) and are therefore immediate precursors to final failure. The micrographs were taken in reflected light, but if compared to images in transmitted light, showed clearly that the copper is cracked along directions perpendicular to the principal tensile stress. It is interesting to note that the thin microcrack has a large hole in the centre. Some of the copper is lost in the cracking process by crack branching and joining. The polyimide partially decoheres in the immediate vicinity of the crack in the copper (point A in Fig. 6(a)). The clear, closely spaced interference fringes in this region arise from the variation in spacing between the polyimide and the copper layers. The spacing can be calculated from the wavelength of the light and is of the order of 1 pm. These kinds of fringes are visible in normal optical microscopy. The decohesion spacing measured in such images is not simply related to the micromechanics of the failure process, since the specimen cracks and delaminates in the interface while under stress, and these images are recorded after unloading. The retractive strains clearly would vary in the region near the cracks. There is a second kind of interference fringe visible over the entire field of view in Fig. 6(b). Comparison with images examined in an interference microscope reveal that these are thickness fringes in the polyimide. The spun films are extremely uniform in thickness and normally produce one set of parallel fringes when
examined on a flat substrate such as a silicon wafer. The spacing between the fringes is a sensitive function of the thickness of the polyimide and the relative orientation of the specimen with respect to the optic axis of the microscope. The cracked and unloaded polyimide-copper composite produces local bending and distortion of the film, which results in the local changes in fringe spacing. In a monotonically loaded specimen this would be a convenient method to examine local strains during testing. This interpretation was confirmed by examining the three-dimensional images. The confocal micrographs in Fig. 6 are single two-dimensional images which are part of a series taken at different depths in the polyimide. The three-dimensional digital reconstruction of the entire series reveals the specimen strongly bulging in the immediate vicinity of each crack. This is due to the inhomogeneous retractive strains in the specimen on unloading. There is a third observation to be made at points such as B in Fig. 7(b). The crack in the copper is associated invariably with a triangular region of low contrast which represents a region of relatively large strain. This can be deduced by examining a number of such regions and observing that the regions are associated with the sides of the cracked copper. The mechanical response in these regions can be explained qualitatively as follows considering the schematic diagram in Fig. 14. The parallel coupled copper and polyimide carry identical total strains, although the elastic and plastic components are different and depend on the stress relaxation kinetics of each phase.
Fig. 13. Internal stress (open circles) compared to flow
143
The copper cracks and the polyimide locally stretches and work hardens to carry the additional load transferred from the copper. The polyimide is not fractured in these regions, so will be deformed locally more than the polyimide far away from the cracked copper. The difference in total strain between the copper and the polyimide varies from zero at a large distance from the crack to a large value at the crack (the point labelled A). In a direction along the applied stress direction, perpendicular to the copper crack, the polyimide is drawn into the crack region and undergoes a large amount of deformation. These regions are in many cases associated with shear bands which extend a considerable distance (several crack lengths) away from the crack tips and then stop. The contrast is extremely low and they can be seen as diffuse bands emanating from the crack tips at about 45 ° (Fig. 7(a)). At low magnification these bands are observed near each of the copper cracks. Micromechanically they are similar to the shear bands observed at the tip of a craze. The shear bands in polyimide-copper are observed with a polarizer inserted into the optical path, and the contrast can be reversed from darker than background (as shown in Fig. 7(b)) to lighter than background by rotating the orientation of the polarizer. In the deformed sample the overall brightness of the entire specimen can also be changed in this way, but not as strongly as observed in the shear band region. This implies a significant molecular orientation in these bands and in the triangular regions, consistent with the mechanical description above. It suggests that one mechanism of deformation in polyimide is shear band propagation, which would be consistent with a dislocation analogue suggested for polymers by Li et al. [14]. This in turn supports the single activated rate model suggested from the stress relaxation data.
could be modelled using a single activated rate process up to strains of the order of 0.18. This was consistent with the microstructural observation that shear bands were seen in the deformed samples. The relaxation rate in the polyimide became anomalously high at strains greater than this, indicating a change in deformation mechanism. The stress relaxation behaviour of the copper has been determined by assuming a simple parallel coupling of the metal and polymer phases. The composite sample failure was preceded by generalized microcracking in the copper. Each microcrack has associated with it a zone of shear and a region of delamination at the interface between the polyimide and copper. The microcracking and delamination were examined using laser-scanning confocal optical microscopy. The shear bands emanated from the ends of the cracks in the copper. These observations support the models of Ho and coworkers which relate the decohesion energies of the polyimide-copper interface to the tensile stress-strain curves [1, 2]. It is clear from the present work, however, that a considerable plastic work is generated by the polyimide shear band formation near each microcrack.
Acknowledgments The work has been supported by the Natural Sciences and Engineering Research Council of Canada. A.K. has been on sabbatical from R A F A E L in Israel and has been supported by the Ontario Center for Materials Research (A. E. Hamielec and J. D. Embury). Considerable assistance from Dr. Paul Ho of IBM (T. J. Watson Research Center) is gratefully acknowledged. The confocal optical microscopy has been supported by grants from the U.S. D O E (DE-AE-0888DP10782 and EE-FG03-89SF18012), the Whitaker Foundation, CAT-HIDI and NFWWTP.
5. Conclusions Composite specimens consisting of thin layers of polyimide and copper have been tested in tension and the inelastic deformation processes up to fracture have been studied. The composite films deformed uniformly on the macroscopic scale to quite large strains. The yield stress of the evaporated thin film copper was approximately 850 MPa while that of the polyimide was about 100 MPa. In polyimide the relaxation of stress
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