Polymers for Microelectronics - American Chemical Society

Chapter 33

Intrinsic and Thermal Stress in Polyimide Thin Films

Downloaded by EAST CAROLINA UNIV on December 16, 2016 | http://pubs.acs.org Publication Date: November 23, 1993 | doi: 10.1021/bk-1994-0537.ch033

M.

Ree and D. P. Kirby

1

IBM Technology Products Division, 74 Creamery Road, Hopewell Junction, NY 12533

In-situ measurements of intrinsic stress, as well as overall internal stress, were performed for three soluble preimidized polyimides as a function of temperature over the range of 25-400 °C using a wafer bending technique: Sixef-44 (6FDA-4,4'-6F), Sixef-33 (6FDA-3,3'-6F), and Probimide 412. For polyimide films with ca. 12 µm thick, intrinsic stress is 29-31 MPa at room temperature, at which the films have been prepared, indicating that at the temperature of the film preparation the intrinsic stress is not sensitive to the type of backbone chemistry among these polyimides. The intrinsic stress varied with temperature in the first heating run, where its variation with temperature was strongly dependent upon the properties of the polyimide, including mechanical properties, polymer chain stiffness, molecular order, and glass transition temperature. The measured intrinsic stress is not small enough to be neglected, as one usually does. In addition, the thermal stress was estimated from the measured intrinsic and overall stresses as a function of temperature. Its variation with temperature was dependent upon the temperature regime: regime I (above T ),regimeII(115°C-T orT ),and regime III (below 115 °C). As is well known, thermal stress is developed below T and, based on our results, apparently increases linearly with descending temperature only in regime II. However, the slope of the thermal stress profiles in regime II is not same with that of the overall stress profiles. In addition, T 's of the polyimides were estimated from the overall stress-temperature profiles. g

g

f

g

g

Aromatic polyimides have found wide application in the microelectronics industry as alpha particle protection, passivation, and intermetallic dielectric layers, owing to their excellent thermal stability, mechanical properties and dielectric properties (2-3). Many microelectronic devices, such as VLSI semiconductor chips and advanced multi-chip modules (5), are composed of multilayer structures. In multilayered structures, one of the serious concerns related to reliability is residual stress caused by thermal and loading histories generated through processing and use, since polyimides have different properties (i.e., mechanical properties, thermal expansion coefficient, and phase transition temperature) from the metal conductors and substrates (ceramic, silicon, and plastic) com1

Current address: Barnett Institute, Northeastern University, 360 Huntington Avenue, Boston, MA 02115 0097-6156/94/0537-0482$06.00/0 © 1994 American Chemical Society Thompson et al.; Polymers for Microelectronics ACS Symposium Series; American Chemical Society: Washington, DC, 1993.


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Intrinsic and Thermal Stress in Polyimide Thin Films

483

monly employed. In general, residual stress in a polyimide film consists of two major components (4). One is the so-called thermal stress due to the mismatch of thermal expansion coefficients between film and substrate or metal layer, as well as the thermal history and mechanical properties. The other is intrinsic stress resulting from volume change due to solvent evaporation and shrinkage, from molecular structural ordering during the film formation process, and perhaps from the physical properties of the formed film. The overall internal stress of a polymer film is usually assumed to be the thermal stress, neglecting the contribution of intrinsic stress. Here, it should be noted that the intrinsic stress in a polymer film is not understood in detail yet. In the present study we have chosen several soluble preimidized polyimides in order to understand intrinsic stress behavior and its contribution to overall stress: Sixef-44, Sixef-33, and Probimide 4 1 2 (see Figure 1). These soluble preimidized polymers are good candidate polymer systems for studying the behaviors of stress components, because completely dried films can be made without any solvent complexation and imidization which occur in polyimide precursors. For these preimidized polymers, fully dried films were prepared on silicon wafers through room temperature drying under a nitrogen flow and then followed by vacuum drying to avoid the involvement of any thermal history. For these dried polyimide films, intrinsic stress was measured in-situ during heating as a function of temperature over the range of 2 5 - 4 0 0 °C using a wafer bending technique ( 5 , 6). During subsequent cooling after baking at a certain temperature, overall stress was also measured dynamically as a function of temperature. The thermal stress component was estimated from the measured intrinsic and overall stresses. INTRINSIC AND THERMAL STRESS Residual stress (i.e., interfacial stress) in a polymer thin film is commonly determined by measuring the curvature (deflection) of the bilayer composite structure of the film and a substrate after film deposition. The stress of the film in equilibrium with the resultant strain can be calculated from the curvature and mechanical parameters of the substrate using the following simple plate equation (7):

6(1-v )t \R s

F

F

*J

( 1 )

Here, the subscripts F and S denote polyimide film and substrate, respectively. The symbols cr, E, v, and t are stress, Young's modulus, Poisson's ratio, and thickness of each layer of material. R and are radii of a substrate with and without a polyimide film, respectively. For Si(100) wafers, biaxial modulus, E /(1 - v ), is 1.805 X 1 0 MPa (8). Eq (1) has been driven under the assumption that the stress is isotropic and uniform in the film plane. The application of this equation is limited to bending displacements smaller than the thickness of F

5

s

s

Thompson et al.; Polymers for Microelectronics ACS Symposium Series; American Chemical Society: Washington, DC, 1993.

484

POLYMERS FOR MICROELECTRONICS

the substrate. In other words, the thickness of a polymer film should be much smaller than the thickness of substrate. The residual stress (cr ) is known to consist of two major components: thermal stress and intrinsic stress (4). The thermal stress (a ) results from the mismatched thermal expansion coefficients (TECs) of the film and substrate, as well as the mechanical properties and thermal history of the film and can be estimated from the following equation (9): F


t

with TECs (a and a ), final heat-treat temperature (7}), stress measurement temperature ( D , modulus (E ) and Poisson's ratio (v ). In contrast to the thermal stress due to thermal history, the intrinsic stress arises mainly from constraints on molecular movement during film formation. A polymer film is commonly fabricated by applying the polymer in solution on to a substrate and subsequently drying it. During drying, the wet polymer film concentrates and its viscosity drastically increases due to solvent evaporation. The wet film starts to solidify when its viscosity reaches a gel point. Below the gel point, the molecules in the film are mobile enough to flow and thus residual stress can not be generated. However, above the gel point the film is extremely viscous and its glass transition temperature (T ) increases. The increase of viscosity and T in the film restricts the molecular motion and results in stress. The shrinkage due to solvent evaporation takes place in the direction of film thickness but is constrained in the direction of the film plane, becuase of the interfacial adhesion between the film and the substrate. Thus, interfacial stress develops in the film plane. This stress (o* ) generated by the polymer film deposition can be expressed by the following equation (10): s

F

F

F

g

g

d

E

(s ~ r)

F

a

"

(1 -

V) f

U

3(1 -

where is the volume fraction of solvent at which the film solidifies and is the volume fraction of solvent retained in the film. As is expressed in Eq (3), the stress (, overall stress; cr,, thermal stress; a intrinsic stress. The heating and cooling rates were 2.0 °C/min and 1.0 °C/min, respectively. iy



488


baked at 250 °C for 30 min. On heating, intrinsic stress decreased monotonically with temperature from 31 MPa at room temperature to 5.2 MPa at 250 °C. During aging at 250 °C for 30 min, intrinsic stress increased slightly from 5.2 MPa to 6.1 MPa. This stress increase is much smaller than that of the film aged at 150 °C. On cooling after the bake, stress increased linearly with decreasing temperature from 6.1 MPa at 250 °C to 41.4 MPa at room temperature. The thermal stress component was estimated in the same way as described above. The thermal stress exhibited two different behaviors, depending on the temperature region (see Figure 3). The thermal stress initially increased linearly with decreasing temperature from zero to 14.9 MPa in the regime of 115-250 °C and then decreased very slowly from 14.9 MPa to 12.5 MPa in the regime of 25-115 °C. The shape of the stress-temperature profile over 25-150 °C resembles that of the film baked at 150 °C. However, the stress level is 2 times higher than that of the 150 °C baked film. Figure 4 shows the stress behavior of the film baked at 400 °C. On heating, the intrinsic stress of 31 MPa at room temperature decreased monotonically with temperature to 2.2 MPa at 400 °C. During aging at 400 °C for 30 min, stress increased slightly from 2.2 MPa to 3.0 MPa. The increase in stress on aging is only 0.8 MPa. Similar aging behavior was previously observed at 250 °C. That is, the aging effect is very small above 250 °C. In contrast, this effect was more significant at 150 °C (a highly supercooled state). Therefore, the aging effect on the stress of Sixef-44 films becomes significant at temperatures below 250 °C. On cooling after baking at 400 °C, overall stress remained at ca. 3 MPa (intrinsic stress only) until 300 °C, and thereafter increased rapidly with decreasing temperature to 50 MPa at room temperature. The final stress is much higher than that of the films baked at 150 °C or 250 °C, indicating that the thermal stress component becomes more significant in the film baked at a higher temperature. The estimated thermal stress is compared with the intrinsic stress as well as the overall stress in Figure 4 and shows three different temperature regime behaviors. The thermal stress, which was developing during cooling, was zero or less until ca. 285 °C (regime I). On further cooling the stress increased rapidly with temperature down to ca. 115 °C (regime II) and thereafter decreased very slowly (regime III). In regime I (above T ), thermal stress could not develop because of its rapid relaxation due to the high mobility of polymer chains causing easy deformation of the film. In regime II (below T ), thermal stress was generated with decreasing temperature as predicted by Eq(2). In this regime, the Poisson's ratio (v ) term contributes negatively to the thermal stress, because it decreases slightly on cooling from the liquid state into the frozen glassy state (11). Also, the TEC (a ) term contributes negatively to the thermal stress, due to the nature of its temperature dependency. However, despite the negative contribution of those two terms (v and a ), the thermal stress has increased overall with decreasing temperature in regime II, indicating that the degree of supercooling (AT = 7} - T or T - 7") and Young's modulus (E ) are predominant contributors to the thermal stress here. The increase of thermal stress in regime II was not continued into regime III where stress varied only slightly with temperature. In this regime, the modulus (E ) and Poisson's ratio (v ) might be much less sensitive to temperature compared with regime II. The stress behavior g

g

F

F

F

F

g

F

F

F


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Intrinsic and Thermal Stress in Polyimide Thin Films 489


Sixef-44

250°C bake 0

50

100

150

200

250

Figure 3. Stress versus temperature behavior of the Sixef-44 film dried at room temperature measured on a Si(100) wafer during baking up to 250 °C and subsequent cooling: o>, overall stress; cr„ thermal stress; a intrinsic stress. The heating and cooling rates were 2.0 °C/min and 1.0 °C/min, respectively. i9

400°C bake 0

100

200 T(°C)

300

400

Figure 4. Stress versus temperature behavior of the Sixef-44 film dried at room temperature measured on a Si(100) wafer during baking up to 400 °C and subsequent cooling: o>, overall stress; o , thermal stress; a intrinsic stress. The heating and cooling rates were 2.0 °C/min and 1.0 °C/min, respectively. t

i9


490


in regime III might result from either the negative contribution of a just cancelling the positive contribution of the AT or its contribution being slightly more dominant than that of AT. The in-situ stress measurement was extended to Sixef-33 and Probimide 412 films. Representative results are shown in Figs. 5 and 6. The intrinsic stress of Sixef-33 films was 31 MPa at room temperature. This stress level is comparable to that of the Sixef-44 film. However, its variation with temperature in the first heating runs is different from that of the Sixef-44 films. As shown in Figure 5, on heating the intrinsic stress rapidly relaxed out with increasing temperature and then leveled off at 1.5 MPa above 150 °C. The fast stress relaxation might result from a relatively high mobility of the polymer chains. In comparison with the Sixef-44 film, the Sixef-33 film is expected to have a lower 7^. On cooling, the overall stress started to increase from 230 °C, due to the contribution of thermal stress generated in the supercooled state below T , and finally reached 45 MPa at room temperature. From these stress results, the thermal stress component was estimated as shown in Figure 5. The thermal stress remained at zero until 230 °C (regime I), increased with decreasing temperature down to ca. 115 °C (regime II), and then turned to decrease slowly to 16 MPa at room temperature (regime III). That is, the Sixef-33 film thermal stress variation with temperature showed a similar temperature regime behavior as that observed for the Sixef-44 film. Similar stress-temperature behaviors were observed for Probimide 412 films. The intrinsic stress was 29 MPa at room temperature. In the heating run, the stress of 29 MPa decreased to ca. 4 MPa at 400 °C (see Figure 6). On cooling after the bake, the overall stress increased continuously with temperature and reached 56 MPa at room temperature. From this stress profile, one expects that the T of the Probimide 412 is > 400 °C. The thermal stress of the film was also estimated from the intrinsic and overall stress profiles as a function of temperature. The thermal stress was not generated above 370 °C (regime I), however increased with deacreasing temperature over 115-370 °C (regime II) and then leveled off or decreased slightly below 115 °C (regime III). As described above, the three polyimides exhibited almost the same intrinsic stress (29-31 MPa) at the drying temperature, regardless of the different chemical backbone. This indicates that the mechanical properties (particularly Young's modulus and Poisson's ratio) of those polyimides are nearly the same. On heating, the intrinsic stress relaxed out with temperature. The stress relaxation was strongly dependent on T , that is, polymer chain flexibility. The T of the Sixef-33 is relatively low so that its stress relaxation is faster than that of the others (see Figs. 4-6). The stress-temperature profiles indicate that T is in the decreasing order of Sixef-33 < Sixef-44 < Probimide 412. On cooling after thermal treatments, overall stress was built up again in the polyimide films by recovery of intrinsic stress and generation of thermal stress (see Figs. 4-6). The overall stress of the polyimide films was sensitive to thermal history, due to the contribution of the themal stress component. In contrast to the intrinsic stress, the overall stress of baked films varied nearly linearly with temperature below T . For the films baked at 400 °C, overall stress at room temperature was 45 MPa for the Sixef-33, 50 MPa for the Sixef-44, and 56 MPa


F

g

g

g

g

g

g



33.

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400

Figure 5. Stress versus temperature behavior of the Sixef-33 film dried at room temperature measured on a Si(100) wafer during baking up to 400 °C and subsequent cooling: o>, overall stress; o~„ thermal stress; oj, intrinsic stress. The heating and cooling rates were 2.0 °C/min and 1.0 °C/min, respectively. 60

Probimide 412

40 CO

a.,

s i

820 3510 400°C bake 100

200 T(°C)

300

400

Figure 6. Stress versus temperature behavior of the Probimide 412 film dried at room temperature measured on a Si(100) wafer during baking up to 400 °C and subsequent cooling: o>, overall stress; o*„ thermal stress; cr intrinsic stress. The heating and cooling rates were 2.0 °C/min and 1.0 °C/min, respectively. i9



492


for the Probimide 412. Their intrinsic stress components were almost the same, 29-31 MPa. Therefore, the thermal stress component was 14 MPa for the Sixef-33, 19 MPa for the Sixef-44, and 27 MPa for the Probimide 412. In both the Sixef-33 and the Sixef-44 films, the intrinsic stress component was higher than the thermal stress component. In the Probimide 412, the intrinsic stress was comparable with the thermal stress component. The contribution of intrinsic stress to the overall stress was more significant in the polyimide films baked at low temperatures (see Figs. 2-4). This is evidence that in polymer films, particularly the polyimides studied here, the intrinsic stress component is not small enough to be neglected. Furthermore, the slope in the thermal stress variation with temperature generally is not same with that of the overall stress variation as shown in Figs. 4-6. The thermal stress variation with temperature is strongly dependent upon the temperature regime. The thermal stress varies nearly linearly only in regime II. However, even in this regime the slope of the thermal stress variation is not the same as that of the overall stress variation. Therefore, it may not be a good approach in the estimation of the biaxial modulus and thermal expansion coefficient of a polymer film that the overall stress measured for a thermally treated polymer film assumes to be the thermal stress itself. The three polyimide films baked at 400 °C exhibited ca. 3.0 GPa Young's modulus at room temperature (72), regardless of the polymer backbone. Equivalent Poisson's ratios are expected for these polyimides. Therefore, the difference in the overall stress of the polyimide films may result from the differences in the T s as well as in the TECs. Considering polymer chain flexibility and T T E C may be relatively high in the Sixef-33, intermediate in the Sixef-44, and low in the Probimide 412 film. However, the thermal stress was in the increasing order of Probimide 412 > Sixef-44 > Sixef-33. Consequently, the difference in the overall stress was predominantly driven by the difference in AT's due to different T s. In addition, the T of the films can be estimated from the stress-temperature profiles. In general, stress is not built up above T because of high polymer chain mobility. However, the chain mobility is restricted in the glassy state below T allowing the generation of stress. In the overall stress-temperature profile measured on cooling, the temperature at which the stress started to increase was chosen as T . For the films baked at 400 °C, T was 235 °C for the Sixef-33, 300 °C for the Sixef-44, and 400 °C for the Probimide 412. 9

g

g>

g

g

g

g

g

g

CONCLUSIONS The intrinsic stress, as well as overall internal stress, in three different polyimide films was dynamically measured during heating and subsequent cooling as a function of temperature over the range of 25-400 °C. The thermal stress, which was generated during cooling, was estimated from the intrinsic and overall internal stresses as a function of temperature. For all the polyimide films studied here, the intrinsic stress was 29-31 MPa at room temperature. The initial level of intrinsic stress, which had been generated mainly by the solidification of cast polyimide solution through solvent evaporation, is apparently not sensitive to the backbone chemistry. However, the variation of intrinsic stress with temperature


33.

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is strongly dependent upon the backbone chemistry, indicating that stress is correlated with the nature of the polyimide chain, including chain stiffness, molecular order and glass transition temperature. For these polyimides, the intrinsic stress is not small enough to be neglected. On the other hand, the thermal stress in the polyimide films was always developed below T on cooling after thermal treatment. Its variation with temperature was dependent upon the temperature regime: regime I (above T ), regime II: (from 115 °C to T or 7}), and regime III (below 115 °C). That is, below T or the final baking temperature (7}), the thermal stress does not vary linearly with temperature as one usually assumes. Only in regime II does the thermal stress apparently change linearly with temperature. However, even in this temperature regime the slope in the thermal stress profile is not same with that of the overall stress profile. Therefore, one should be carefully when estimating the biaxial modulus and thermal expansion coefficient of a film directly from the overall stress profile. In addition, the T s of the polyimides were estimated from the overall stress-temperature profiles. g

g

g


g

g

REFERENCES 1. Sroog, C. E. J. Polym. Sci.: Macromol. Rev. 1976, 11, 161. 2. Mittal, K. L., Ed. Polyimides:; Plenum Press: New York, NY, 1984. 3. Tummala, R. R.; Rymaszewski, E. J., Eds. Microelectronics Packaging Handbook; van Nostrand Reinhold: New York, 1989. 4. Hoffman, W.R. in Physics of Thin Film; Hass, G.; Thun, R. E., Eds.; Academic: New York, 1966; Vol.3, p 211, . 5. Ree, M.; Swanson, S.; Volksen, W ACS Polym. Preprints 1991, 32(3), 308. 6. Ree, M.; Nunes, T. L.; Volksen, W.; Czornyj, G. Polymer, 1992, 33, 1228. 7. Jaccodine, R. J.; Schlegel, W. A . J. Appl. Phys., 1966, 37, 2429. 8. Wortman, J. J.; Evans, R. A. J. Appl. Phys., 1965, 36, 153. 9. Timoshenko, S. J. Opt. Soc. Am., 1925, 11, 233. 10. Croll, S. G. J. Appl. Polym. Sci, 1979, 23, 847. 11. Brandrup, J.; Immergut, E. H., Eds. Polymer Handbook; Wiley-Interscience: New York, 1975; Chapter V. 12. Ree, M.; Chen, K. J. (unplublished results). Received December 30, 1992


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