- Email: [email protected]
Effects of patterning on the interface toughness of wafer-level Cu–Cu bonds
Available online at www.sciencedirect.com
Acta Materialia 56 (2008) 438–447 www.elsevier.com/locate/actamat
Effects of patterning on the interface toughness of wafer-level Cu–Cu bonds Rajappa Tadepalli a, Kevin T. Turner b, Carl V. Thompson a,* a
Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA 02139, USA b Department of Mechanical Engineering, University of Wisconsin, Madison, WI 53706, USA Received 19 July 2007; received in revised form 1 October 2007; accepted 3 October 2007 Available online 11 December 2007
Abstract The chevron test has been employed to characterize the toughness of patterned wafer-level Cu–Cu thermocompression bonds created at 300 °C, for pattern sizes ranging from 2 to 500 lm. Features oriented perpendicular to the debond propagation direction (lines and pads) exhibited a significant increase in toughness (from 3 to 30 J m2 under mode I) with decreasing feature size (from 250 to 25 lm) for both mode I and mixed-mode loading, while no size-dependence was observed for debond propagation in the direction parallel to the bonded lines. The bond toughness was found to scale with the degree of discontinuity along the debond growth direction, due to greater energetic cost for multiple crack initiation events in the ductile bonded stack. Fractured surfaces of the discontinuous bonded interfaces exhibited ductile cohesive failure through the Cu film stack, signifying enhanced plastic energy dissipation leading to higher bond toughness. Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Bonding; Joining; Thin films; Toughness; Copper
1. Introduction Wafer bonding is an important process in the manufacture of semiconductor microelectronics components and microelectromechanical systems (MEMS). The principal applications for bonded wafers are in the fabrication of vertically integrated circuits [1] and MEMS devices [2], and the manufacture of advanced substrates (e.g. siliconon-insulator substrates) [3]. Common bonding mechanisms include direct bonding [4], anodic bonding [5] and thermocompression bonding [1,6]. With an increasing number of applications based on wafer bonding technology for the creation of three-dimensional structures, the need to bond wafers patterned by etching or thin film deposition has assumed greater importance in both direct bonding (e.g. fabrication of miniaturized gas turbine engine) [7,8] and thermocompression bonding (e.g. packaging for MEMS, three-dimensional *
Corresponding author. Tel.: +1 617 253 7652; fax: +1 617 258 6749. E-mail address: [email protected] (C.V. Thompson).
integrated circuits) [1,6]. Three-dimensional integrated circuit (3-D IC) technology based on the vertical integration of stacked device layers (with their associated metal interconnection layers) has emerged as a key area of application for patterned wafer bonding. 3-D ICs are expected to significantly reduce the composite lengths of interconnects, and thereby allow higher packing densities and smaller signal delays in wiring-limited ICs [9]. In particular, thermocompression bonding of Cu films is an attractive option for creating 3-D ICs, since Cu can serve as both the electrical (interconnect) and mechanical link between the individual device layers. Design of Cu interconnect structures for 3-D ICs requires an understanding of bond formation between patterned films. Qualitatively, it has been observed that Cu–Cu bond integrity is dependent on the size and area fraction (density) of patterns [10]. While quantitative characterization of Cu–Cu bonds between continuous films has been performed [11], the effect of patterning has not been examined in detail. Pattern effects in other thermocompression bond systems (Au–Au bonds) have been studied by
1359-6454/$30.00 Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2007.10.016
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
439
Tsau et al. [6]. In Au–Au studies, the bond quality was measured using four-point bend delamination fracture specimens. Toughness of discontinuous bond interfaces (bonded pads) was found to be significantly greater (250 vs. 50 J m2) than the toughness of continuous bonded interfaces (bonded lines). This difference was attributed to the high pressures achieved during bonding of pads (120 vs. 7 MPa), since the area density of pads was lower than that of the lines, and the applied force to the bond plate was kept constant. It is not clear from the above study, however, if there are fundamental differences between continuous and discontinuous bonded interfaces that are formed under similar bonding conditions. Effects of pattern size and orientation on adhesion of multi-layer thin film stacks have been investigated by Litteken et al. [12]. The contribution from plastic energy dissipation and crack-tip shielding to the relevant fracture energy was measured in patterned structures that contained arrays of Cu/PAE (polyaryl ether) or Cu/CDO (carbondoped oxide) lines. The bond quality was measured using four-point bend delamination fracture specimens. The fracture energy of Cu/CDO structures was found to be dependent on the orientation of the line width with respect to the direction of crack propagation. The width of the lines was on the order of the film thickness, thereby resulting in conditions under which dimensional constraints play a role in the plastic deformation of the ductile layers [12]. Common loading conditions for interface toughness characterization are tension (mode I), shear (mode II), and mixed-mode (mode I + mode II). The degree of mode-mixity is typically quantified by the phase angle w, which has a magnitude defined as rffiffiffiffiffiffiffi GII 1 w ¼ tan ; ð1Þ GI
Fig. 1. (a) Schematic diagram of a chevron test specimen for characterization of bonded interfaces. (b) Top view of a chevron specimen showing the important geometric characteristics.
where GI and GII are the mode I and mode II strain energy release rates, respectively. The phase angle varies from 0° (pure mode I) to 90° (pure mode II). Experimental values for fracture toughness, or the critical strain energy release rate, Gc, depend on material properties such as the atomiclevel bonding at the interface as well as the elastic–plastic constitutive behavior of adjacent materials and microstructures, and on mechanical parameters such as the loading mode-mixity near the crack tip. Generally, increasing mode-mixity leads to higher measured fracture toughness values, as a result of frictional effects in rough interfaces and the work of plastic deformation in ductile materials [13,14]. The bonded chevron specimen has been shown to be an effective means of characterizing the interface toughness of wafer bonds [15–18]. The specimen is simple to fabricate, amenable to miniaturization and easy to test. Fig. 1a shows a schematic of a chevron test specimen used for bonded interface characterization. The surface of one wafer, or of the interlayer, is patterned in the form of a V-shaped notch prior to bonding. The specimen is glued between two stubs
that allow the specimen to be gripped and loaded in tension during testing. The sharp chevron tip serves as a stressconcentration point and enables characterization of both weak and strong bonds. Fig. 1b shows the top view of the chevron specimen, and includes specification of important geometric parameters. The effective loading line is different from the specimen edge due to the finite area along which the stubs are glued. If the thickness or elastic properties of the two bonded layers of the specimen are not the same, a mixed-mode loading is produced at the interface. Recently, finite element (FE) modeling was used to analyze loading of a chevron specimen with layers of different thicknesses, in order to allow such specimens to be used for measurements of wafer bonds under mixed-mode conditions [17]. An experimental study of the mixed-mode interface toughness of Cu–Cu bonds has been completed using the chevron specimen [18]. It has been observed that the Cu–Cu bond toughness increases from 2.68 to 10.1 J m2 as the loading is changed from pure mode I to mixed-mode (w = 24°), for bonds created at 300 °C under a bonding pressure of
440
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
Fig. 2. Schematic diagrams of the four kinds of bonded interfaces studied. Blanket films and parallel lines involve crack propagation along continuously bonded interfaces, while pads and orthogonal lines require crack propagation across discontinuously bonded interfaces. ‘‘L” denotes the pitch of the patterns in the direction of crack propagation.
2.12 MPa. This is because the energy of plastic dissipation increases with increasing mode-mixity. In the present study, bonded chevron specimens were used to analyze the effects of patterning on Cu–Cu bond toughness under mode I and mixed-mode loading conditions. Two different kinds of bonded interfaces are considered: continuous and discontinuous (Fig. 2). Continuous interfaces present a homogeneous bond front along the direction of crack propagation, and are created by bonding blanket Cu films or by bonding patterned surfaces with lines aligned parallel to the crack propagation direction. Discontinuous interfaces have periodic bonded and unbonded regions along the direction of crack propagation. Two types of discontinuous interfaces are studied: bonded pads (discontinuous in two dimensions) and lines bonded orthogonal to the direction of crack propagation (orthogonal lines, discontinuous only along the propagation direction). ‘‘L” in Fig. 2 refers to the pitch of the patterns in the direction of crack propagation. The in-plane dimensions of the patterns are at least an order of magnitude greater than the film thickness, and therefore in-plane dimensional constraints are not expected to influence plasticity during debonding. 2. Experimental procedure 2.1. Specimen fabrication Thermally oxidized 100 mm diameter Si (1 0 0) wafers, with a nominal oxide thickness of 300 nm, were used as substrates for all thin film deposition and bonding studies. Two sets of wafers were fabricated for bonding studies. One set of wafers (set A) had patterned chevron structures, and the other set (set B) had patterned lines or pads. All bonding pairs comprised one wafer from each set. The chevron structures had the following dimensions: length
W (9 mm), width B (10 mm), notch angle h (=90°) and initial crack length a0 (=1 mm). The thickness ratio, g, is defined as the ratio between h2 and h1 (Fig. 1a). The thickness of the chevron wafers (set A) was kept constant (=475 lm). The thicknesses of wafers in set B were varied to induce mixed-mode loading conditions. Three different thicknesses were used: 475, 335 and 230 lm, corresponding to thickness ratios (g) of 1.0, 0.7 and 0.48 and phase angles (w) of 0°, 12.7° and 24°, respectively [17]. The chevron structures were patterned on oxidized wafers using standard photolithography. After patterning, the oxide film was etched using a buffered oxide etch (BOE) and deep reactive ion etching of the Si to a depth of 10 lm was carried out. Between adjacent chevron structures 0.5 mm wide alleys were patterned to enable dicing of individual specimens after bonding. After the creation of chevron structures, the wafers were cleaned in an O2 plasma, followed by a piranha clean (3:1 H2SO4:H2O2). The SiO2 film on the front side was then removed using BOE, and the chevron wafers were reoxidized thermally to an oxide thickness of 300 nm. The complementary set of wafers (set B) had patterned structures. Four different mask sets were used to create patterns. Pattern sizes/densities in the masks are outlined in Table 1. A Cu lift-off process was used to create patterned structures. After patterning image-reversal photoresist on wafer set B, 30 nm of Ta and 400 nm of Cu were successively deposited (without breaking vacuum) on both sets of wafers using e-beam evaporation. Ta served as both an adhesion promoter and a diffusion barrier between Cu and SiO2. The metal deposited on photoresist was lifted off the set B wafers by soaking them in NMP (N-methyl2-pyrrolidone) for at least 1 h, followed by light ultrasonic agitation in NMP and deionized water. The wafers from sets A and B were then prepared for thermocompression bonding. The detailed bonding procedure has been described elsewhere [18]. The patterned lines on the set A wafer were aligned to the chevron pattern in two different configurations: orthogonal and parallel. Schematic chevron structures with patterned lines and pads are shown in Fig. 3. The wafers were bonded at 300 °C for 1 h under a chamber vacuum of 103 Torr. The bonding pressure was maintained at a constant value of 2.12 MPa for all wafers. After the bonded pair was cooled to room temperature, the bonded wafers were taken out of the chamber and annealed in a tube furnace in a forming gas (95% Ar, 5% H2) ambient for 1 h. Fig. 4 shows a schematic cross-sectional view of the bonded stack. Individual
Table 1 Description of mask layouts used for the creation of patterns Pattern type
Pattern width (lm)
Spacing (lm)
Area fraction
Lines Pads Pads Pads
2–250 500 80 40
=width 100 40 20
0.5 0.69 0.44 0.44
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
441
Fig. 3. Schematic illustration of bonded chevron specimens with different pattern geometries. (a,b) Continuous interfaces (blanket films and parallel lines). (c,d) Discontinuous interfaces (orthogonal lines and pads). Note that the minimum and maximum ratios of the specimen width to the patterned dimensions were 20 and 5000, respectively.
Fig. 5. Typical load–displacement behavior from a chevron test. Fig. 4. Schematic cross-sectional view of a bonded chevron specimen. (The film thicknesses and substrate thicknesses are illustrated at different scales).
specimens for chevron testing were cut from the bonded wafers using a diamond saw. 2.2. Chevron testing Test specimens from bonded wafer-pairs were glued to aluminum stubs using a high-viscosity glue. Care was taken to prevent seepage of glue into the bonded interface. Approximately 15 specimens were tested for each geometric configuration. The specimens were selected to represent all regions on the bonded wafer. The glued specimens were tested under tension in an InstronTM 8848 electromechanical universal testing apparatus. The specimens were loaded at a constant displacement rate of 0.12 mm min1 and the load–displacement behavior was digitally recorded. A typical load–displacement curve for a chevronnotched specimen is shown in Fig. 5. After the initial elastic loading of the specimen, a crack initiates at the chevron
notch and propagates along the bonded interface. The triangle-shaped bonded area provides an increasing interfacial area with increasing crack length and load, and initially results in stable crack propagation. At some crack length the propagation becomes unstable and catastrophic failure of the specimen occurs; this coincides with the point of maximum load in Fig. 5. For a uniformly bonded interface, the bond toughness can be calculated from the maximum load in the test, Fmax, using [17]: GC ¼
F 2max Y 2min ; E B2 W
ð2Þ
where Ymin is the minimum value of the geometry function, E is the plane strain modulus, and B and W are the width and length of the specimen, respectively (Fig. 1b). For the anisotropic silicon specimens used in the current work, the effective plane strain modulus in the [0 0 1] direction in the (1 1 0) plane is used (E ¼ 144:7 GPa). FE analysis was used to determine the minimum value of the geometry function, which is dependent on the thickness of the
442
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
bonded layers and crack geometry [17]. Ymin is a function of the initial crack length ratio, a0 = a0/W, the total thickness of the specimen, htotal, and the thickness ratio, g: Y min ¼ ð25:71 þ 85:53a0 Þ sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ð2:91Þ ð2:91Þ htotal g htotal : þ 1þg 1þg
ð3Þ
The above equation is a fit to parametric FE results and estimates Ymin to within 1.8% of the FE calculations across a range of thicknesses (from 0.2 to 1.35 mm), thickness ratios (from 0.1 to 1), and initial crack length ratios (from 0.1 to 0.4) [17]. When the interface is patterned, the toughness value calculated from Eq. (2) must be corrected for the reduced bond area according to GC jpatterned ¼
1 GC junpatterned : Af
ð4Þ
The value of Af depends on the type and orientation of the pattern relative to the direction of crack propagation. When the pattern consists of lines oriented parallel to the direction of crack propagation, Af is the bonded area fraction. If the interface is patterned with lines that are oriented orthogonal to the direction of crack propagation, Af has a value of unity, as the bonded area is continuous along the crack front. For the bonded pad configuration, Af is the linear pad density along the crack front. From Eqs. (2)–(4), the interface toughness can be calculated from the maximum load measured in experiments and the geometry and elastic properties of the specimen. After testing, the specimens were dismounted from the studs using acetone and the studs were reused. Optical microscopy, scanning electron microscopy (SEM), atomic force microscopy and profilometry were then used to characterize the fracture surfaces.
Fig. 6. Mode I Fmax (failure load) vs. feature size for bonded blanket films, parallel and orthogonal lines, and pads. The failure load increases with decreasing feature size for discontinuous bonded interfaces, whereas the failure load for bonded parallel lines is independent of feature size. This trend is similar to that of bond toughness vs. feature size.
3. Results 3.1. Mode I bond toughness vs. pattern size Fig. 6 shows the variation in the failure load, Fmax (illustrated in Fig. 5), with respect to the pattern size. Since the specimen and chevron feature geometry parameters were the same across all tests, the failure load is related to the bond toughness by GC jpatterned /
1 2 F ; Af max
ð5Þ
based on Eqs. (2) and (4). Essentially, the failure load is a measure of the ‘‘global” toughness of the specimen, which, when normalized by the bonded area fraction, yields the ‘‘local” Cu–Cu bond toughness. Fig. 7 shows the variations in the mode I Cu–Cu bond toughness as a function of the feature size for all bonded configurations. For the parallel line configuration, the toughness was independent of the line width, and was mar-
Fig. 7. Mode I Cu–Cu bond toughness vs. feature size for bonded blanket films, parallel and orthogonal lines, and pads. For discontinuous bonded interfaces (orthogonal lines and pads), the toughness increases with decreasing feature size. For interfaces that are continuously bonded in the direction of crack propagation, the toughness is independent of feature size.
ginally higher than that of bonded blanket Cu films. On the other hand, the specimens bonded with orthogonal line configurations exhibited significant line width dependence, with the bond toughness increasing with decreasing line widths. Therefore, the degree of discontinuity along the crack propagation direction had a profound effect on the interfacial toughness. The other specimens with discontinuous interfaces, bonded pads, exhibited a similar trend in bond toughness with the bond toughness also increasing with decreasing pad width. The discontinuous bonded interfaces failed at higher loads as the pattern size decreased, while the bonded
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
443
parallel lines did not show any dependence of the failure load on pattern size. This trend shown in Fig. 6 is similar to that of the bond toughness (Fig. 7), even though the data in Fig. 6 are not normalized with respect to the bonded area fractions. The difference between Figs. 6 and 7 is that the bonded parallel lines, which failed at lower loads than bonded blanket films, exhibited higher toughness than the latter when bonded area fractions were invoked. 3.2. Mixed-mode bond toughness vs. pattern size The Cu–Cu bond toughness was found to be independent of line width for all mode-mixities under conditions for which the direction of crack propagation was parallel to the lines, as seen in Fig. 8. As seen in Fig. 9a, for the orthogonal line configuration, the toughness increased with line width for all mode-mixities and the magnitude of mode-mixity effect increased with decreasing line width. A common trend between the parallel and orthogonal line configurations is that the interfacial toughness increases with increasing phase angle, as discussed in Ref. [18]. This behavior is captured in Fig. 9b. 3.3. Failure loci The locus of failure was found to be dependent on the mode-mixity and the pattern orientation, with three different failure locations/mechanisms observed: at the Ta–SiO2 interface (adhesive failure), at the Cu–Cu bonded interface (cohesive failure) and in the bulk Cu (also cohesive). Roughly equal percentages of adhesive- and cohesivebonded interface failures were observed in mode I testing of parallel lines. The specimens tested under mixed-mode conditions had increasing incidence of adhesive failures, indicating crack deflection from the Cu–Cu bond interface
Fig. 8. Toughness of bonded Cu–Cu parallel lines as a function of line width for three different phase angles. Mode-mixity was achieved using asymmetric chevron specimens. Schematics of chevron specimens and cross-section of bonded stacks were shown in Figs. 3b and 4, respectively. For a given phase angle, the toughness is independent of line width for parallel lines.
Fig. 9. (a) Cu–Cu bond toughness as a function of line width for different mode-mixities. The orthogonal line configurations exhibit significant line width dependence, whereas the toughness of parallel lines is independent of line width. (b) Cu–Cu bond toughness as a function of phase angle for parallel and orthogonal lines. As reported in Ref. [18], the bond toughness is a strong function of the phase angle of loading.
to the Ta/SiO2 interface. Another difference between mode I and mixed-mode failure loci was that the mode I cohesive failures occurred along the bonded interface whereas the mixed-mode cohesive failures occurred within the bulk Cu, again confirming crack path deflection. Specimens with bonded orthogonal lines and pads exhibited predominantly adhesive failures, irrespective of the mode-mixity. Fig. 10 shows SEM images of the mode I fractured surfaces of the bonded lines. Failure at the bonded Cu–Cu interface for bonded parallel lines is shown in Fig. 10a and b, while bulk Cu failure (away from the bonded interface) in bonded orthogonal lines is shown in Fig. 10c and d. The fracture surfaces of orthogonal lines show rough features, indicative of enhanced plastic deformation. Fig. 11a and b show mixed mode cohesive failures of specimens tested under a phase angle of 12.7°, and Fig. 11c and d show
444
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
Fig. 10. Mode I fracture surfaces of bonded lines. (a,b) Parallel lines; (c,d) orthogonal lines (line width = 100 lm). Orthogonal lines show rough fracture surfaces, indicative of significant plastic deformation.
Fig. 11. Mixed-mode fracture surfaces of bonded lines. (a,b) w = 12.7°; (c) and (d) w = 24°. The roughness of the Cu fracture surfaces increases with increasing mode-mixity, indicating increased plasticity in ductile failure.
specimens tested under a phase angle of 24°. Evidence of ductile failure is particularly apparent in the samples loaded
with a 24° phase angle, for which a dimpled fracture surface morphology was observed (Fig. 11d).
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
445
Fig. 12. Mode I fracture surfaces of bonded pads. w = pad width. (a,b) w = 500 lm; (c,d) w = 80 lm; (e,f) w = 40 lm. Bright and dark pads in (e) represent Cu–Cu and Ta/SiO2 failure modes, respectively. (a,c,e) Mixed failure modes (Ta/SiO2 and Cu–Cu); (d,e,f) magnified images of the Cu–Cu fracture surfaces.
Fracture surfaces of bonded pads are shown in Fig. 12. Micrographs ‘a’ and ‘b’ show the surfaces of the widest pads (500 lm), ‘c’ and ‘d’ show pads of intermediate width (80 lm) and ‘e’ and ‘f’ show the smallest pads (40 lm). The roughness of the fracture surfaces increases with decreasing pad size, suggesting more plastic energy dissipation. A classic ductile cup-and-cone fracture surface can be seen in the micrographs of the smallest pads (Fig. 12f). 4. Discussion 4.1. Orthogonal lines and pads The orientation and size of features have a profound effect not only on the toughness of Cu–Cu bonds (Figs. 7
and 9), but also on the failure load (Fig. 6). The latter shows that patterning has a fundamental effect on the toughness of discontinuous bonded interfaces, since they fail at higher loads with decreasing pattern size, irrespective of the bonded area fraction. The pattern orientation determines the load required to fracture the interface, since the continuous patterns reduce the load compared with blanket interfaces, while the discontinuous patterns increase the load compared with blanket interfaces. Interestingly, pattern size dependence is observed only for bonded interfaces that are discontinuous in the direction of crack propagation, for which the bond toughness significantly increases as the pad or line size is decreased. Therefore, patterning effects can be traced to the nature of crack propagation along a heterogeneous
446
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
bonded interface. For the purpose of this discussion, only mode I toughness is considered, so that a direct comparison can be made between bonded lines and pads. However, the arguments below are expected to hold for mixed-mode toughness as well, the only difference being the mode-mixity contribution to enhanced plastic deformation, which has been discussed in [18]. Both orthogonal line and pad configurations involve crack propagation across periodic gaps in bonded regions. Since the bonded interface lies between ductile Cu layers on both sides, there is an energetic cost to crack initiation (or nucleation), involving void growth and coalescence ahead of the crack tip. The measured bond toughness is plotted as a function of the number of crack initiation events in Fig. 13. The number of crack initiation events is obtained by dividing the distance d traversed by the crack prior to complete specimen failure by the pitch of the patterns L (along the crack propagation direction), as shown schematically in Fig. 2. The value of d, estimated to be 3 mm from FE modeling [17], has been verified experimentally using a combination of load–displacement plots and fracture surface analyses. In Fig. 13 it can be seen that the bond toughness of orthogonal lines and pads increases monotonically with increasing number of crack initiation events, suggesting that more energy is absorbed in bond configurations for cracks that have to nucleate at multiple points. It was assumed that cracks nucleate simultaneously across all the pads along a row perpendicular to the crack propagation direction and therefore the row is considered as a single initiation event. However, there is still a discrepancy between the curves that relate toughness to the number of crack initiation events, as determined for pads versus lines, which is not clearly understood at this point.
Fracture surfaces of the orthogonal lines (Fig. 10c and d) and pads (Fig. 12) reveal rough features, indicative of ductile fracture. The roughness, and therefore, the plastic energy dissipation, are greater than that observed for parallel lines (Fig. 10a and b), providing further evidence for enhanced toughness due to a greater number of crack initiation events. Recently, Kubair and Spearing [19] used a spectral scheme with cohesive zone modeling to treat the effect of discontinuous patterning on the energy of interfacial fracture, using a square wave variation in the critical distance of separation of the interfaces. They predicted an increase in the apparent bond toughness associated with an increased (plastic) strain energy for decohesion of patterned bonded surfaces compared with unpatterned bonded surfaces, consistent with the results on bonded Au pads from [6]. However, they also predicted a weak dependence of the measured bond toughness on the repeat period of the pattern, which contrasts with the results found in the current study. The Kubair and Spearing treatment also does not apply to variations in bonded surfaces associated with two-dimensionally periodic patterns. 4.2. Parallel lines For the range of line widths considered (2–250 lm), the toughness of bonded parallel lines was independent of line width for all mode-mixities. Litteken and Dauskardt observed a line-width dependence of adhesion at interfaces of thin-film structures containing patterned polymer lines, even when the direction of crack propagation was parallel to the line axes [20]. In their work, decreasing the line width was found to significantly increase the interfacial fracture energy. The increase in adhesion was associated with increasing contributions from plastic energy dissipation in the patterned lines owing to the decrease in the lateral constraint of the lines with decreasing line width. The observed fracture resistance behavior was similar to the transition from plane strain to plane stress fracture commonly found in bulk metals and polymers. In the current work, the length scales (>2 lm) considered are still in the regime of plane strain for Cu films and, therefore, are not expected to have an effect on the interfacial adhesion. Based on the work of Litteken and Dauskardt, it is expected that narrower Cu lines (sub-micron widths) would show plane stress behavior and enhanced deformation and bond toughness, even under the parallel configuration. 5. Summary
Fig. 13. Toughness of bonded orthogonal lines and pads (from Fig. 6) plotted as a function of the number of crack initiation events. The number of crack initiation events is inversely proportional to the repeat-length of the features, L, along the direction of crack propagation. An increasing number of crack initiation events results in a greater energetic cost for crack propagation and, hence, higher toughness.
Mode I and mixed-mode chevron tests on patterned Cu– Cu bonded interfaces reveal interesting dependencies of bond toughness on pattern size, density, and orientation. For cracks running parallel to bonded lines, the interfacial adhesion was independent of line widths for widths as low as 2 lm. On the other hand, the bond toughness was a strong function of the line width/pad size when cracks propagated
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
across discontinuously bonded interfaces. In this case, for a given bonded area fraction, decreasing the feature size was found to significantly increase the decohesion energy. This increase was attributed to the greater number of crack initiation events for the smaller features. This semi-quantitative empirical model captures the effects of pattern size (repeat distance) and orientation on bond toughness. Mixed-mode bond toughness of orthogonal lines also followed similar trends as mode I toughness. As discussed in our earlier work [18], increased mode-mixity contributes to enhanced plastic deformation and, therefore, to greater bond toughness compared with specimens loaded in pure tension. References [1] Fan A, Rahman A, Reif R. Electrochem Solid State Lett 1999;2(10):534. [2] Schmidt MA. Proc IEEE 1998;86(8):1575. [3] Celler GK, Cristoloveanu S. J Appl Phys 2003;93(9):4955. [4] Plo¨ßl A, Kra¨uter G. Mater Sci Eng 1999;R25(1–2):1.
447
[5] Cozma A, Puers R. J Micromech Microeng 1995;5(2):98. [6] Tsau CH, Spearing SM, Schmidt MA. J Microelectromech Syst 2004;13(6):963. [7] Turner KT, Spearing SM. J Appl Phys 2002;92(12):7658. [8] Miki N, Zhang X, Khanna R, Ayon AA, Ward D, Spearing SM. Sensors Actuator A 2003;103(1–2):194. [9] Rahman A, Reif R. IEEE Trans VLSI Syst 2000;8(6):671. [10] Tan CS. Multi-layer three-dimensional silicon electronics enabled by wafer bonding, Ph.D. thesis, MIT, 2006. [11] Tadepalli R, Thompson CV. In: Proceedings of the IEEE 2003 international interconnect technical conference; 2003. p. 36. [12] Litteken CS, Dauskardt R, Scherban T, Xu G, Leu J, Gracias D, Sun B. In: Proceedings of the international interconnect technical conference; 2003. p. 168. [13] Evans AG, Hutchinson JW. Acta Metall Mater 1989;37(3):909. [14] Tvergaard V, Hutchinson JW. J Mech Phys Sol 1993;41(6):1119. [15] Hurd DS, Caretta R, Gerberich WW. J Mater Res 1995;10(2):387. [16] Bagdahn J, Petzold M, Plo¨ßl A, Wiemer M. Electrochem Soc Proc 2001;99(35):218. [17] Tadepalli R, Turner KT. Eng Fract Mech in press. [18] Tadepalli R, Turner KT, Thompson CV. J Mech Phys Solids, in press. [19] Kubair DV, Spearing SM. J Phys D 2006;39(6):1050. [20] Litteken CS, Dauskardt RH. Int J Fracture 2003;119/120:475.
Acta Materialia 56 (2008) 438–447 www.elsevier.com/locate/actamat
Effects of patterning on the interface toughness of wafer-level Cu–Cu bonds Rajappa Tadepalli a, Kevin T. Turner b, Carl V. Thompson a,* a
Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA 02139, USA b Department of Mechanical Engineering, University of Wisconsin, Madison, WI 53706, USA Received 19 July 2007; received in revised form 1 October 2007; accepted 3 October 2007 Available online 11 December 2007
Abstract The chevron test has been employed to characterize the toughness of patterned wafer-level Cu–Cu thermocompression bonds created at 300 °C, for pattern sizes ranging from 2 to 500 lm. Features oriented perpendicular to the debond propagation direction (lines and pads) exhibited a significant increase in toughness (from 3 to 30 J m2 under mode I) with decreasing feature size (from 250 to 25 lm) for both mode I and mixed-mode loading, while no size-dependence was observed for debond propagation in the direction parallel to the bonded lines. The bond toughness was found to scale with the degree of discontinuity along the debond growth direction, due to greater energetic cost for multiple crack initiation events in the ductile bonded stack. Fractured surfaces of the discontinuous bonded interfaces exhibited ductile cohesive failure through the Cu film stack, signifying enhanced plastic energy dissipation leading to higher bond toughness. Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Bonding; Joining; Thin films; Toughness; Copper
1. Introduction Wafer bonding is an important process in the manufacture of semiconductor microelectronics components and microelectromechanical systems (MEMS). The principal applications for bonded wafers are in the fabrication of vertically integrated circuits [1] and MEMS devices [2], and the manufacture of advanced substrates (e.g. siliconon-insulator substrates) [3]. Common bonding mechanisms include direct bonding [4], anodic bonding [5] and thermocompression bonding [1,6]. With an increasing number of applications based on wafer bonding technology for the creation of three-dimensional structures, the need to bond wafers patterned by etching or thin film deposition has assumed greater importance in both direct bonding (e.g. fabrication of miniaturized gas turbine engine) [7,8] and thermocompression bonding (e.g. packaging for MEMS, three-dimensional *
Corresponding author. Tel.: +1 617 253 7652; fax: +1 617 258 6749. E-mail address: [email protected] (C.V. Thompson).
integrated circuits) [1,6]. Three-dimensional integrated circuit (3-D IC) technology based on the vertical integration of stacked device layers (with their associated metal interconnection layers) has emerged as a key area of application for patterned wafer bonding. 3-D ICs are expected to significantly reduce the composite lengths of interconnects, and thereby allow higher packing densities and smaller signal delays in wiring-limited ICs [9]. In particular, thermocompression bonding of Cu films is an attractive option for creating 3-D ICs, since Cu can serve as both the electrical (interconnect) and mechanical link between the individual device layers. Design of Cu interconnect structures for 3-D ICs requires an understanding of bond formation between patterned films. Qualitatively, it has been observed that Cu–Cu bond integrity is dependent on the size and area fraction (density) of patterns [10]. While quantitative characterization of Cu–Cu bonds between continuous films has been performed [11], the effect of patterning has not been examined in detail. Pattern effects in other thermocompression bond systems (Au–Au bonds) have been studied by
1359-6454/$30.00 Ó 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2007.10.016
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
439
Tsau et al. [6]. In Au–Au studies, the bond quality was measured using four-point bend delamination fracture specimens. Toughness of discontinuous bond interfaces (bonded pads) was found to be significantly greater (250 vs. 50 J m2) than the toughness of continuous bonded interfaces (bonded lines). This difference was attributed to the high pressures achieved during bonding of pads (120 vs. 7 MPa), since the area density of pads was lower than that of the lines, and the applied force to the bond plate was kept constant. It is not clear from the above study, however, if there are fundamental differences between continuous and discontinuous bonded interfaces that are formed under similar bonding conditions. Effects of pattern size and orientation on adhesion of multi-layer thin film stacks have been investigated by Litteken et al. [12]. The contribution from plastic energy dissipation and crack-tip shielding to the relevant fracture energy was measured in patterned structures that contained arrays of Cu/PAE (polyaryl ether) or Cu/CDO (carbondoped oxide) lines. The bond quality was measured using four-point bend delamination fracture specimens. The fracture energy of Cu/CDO structures was found to be dependent on the orientation of the line width with respect to the direction of crack propagation. The width of the lines was on the order of the film thickness, thereby resulting in conditions under which dimensional constraints play a role in the plastic deformation of the ductile layers [12]. Common loading conditions for interface toughness characterization are tension (mode I), shear (mode II), and mixed-mode (mode I + mode II). The degree of mode-mixity is typically quantified by the phase angle w, which has a magnitude defined as rffiffiffiffiffiffiffi GII 1 w ¼ tan ; ð1Þ GI
Fig. 1. (a) Schematic diagram of a chevron test specimen for characterization of bonded interfaces. (b) Top view of a chevron specimen showing the important geometric characteristics.
where GI and GII are the mode I and mode II strain energy release rates, respectively. The phase angle varies from 0° (pure mode I) to 90° (pure mode II). Experimental values for fracture toughness, or the critical strain energy release rate, Gc, depend on material properties such as the atomiclevel bonding at the interface as well as the elastic–plastic constitutive behavior of adjacent materials and microstructures, and on mechanical parameters such as the loading mode-mixity near the crack tip. Generally, increasing mode-mixity leads to higher measured fracture toughness values, as a result of frictional effects in rough interfaces and the work of plastic deformation in ductile materials [13,14]. The bonded chevron specimen has been shown to be an effective means of characterizing the interface toughness of wafer bonds [15–18]. The specimen is simple to fabricate, amenable to miniaturization and easy to test. Fig. 1a shows a schematic of a chevron test specimen used for bonded interface characterization. The surface of one wafer, or of the interlayer, is patterned in the form of a V-shaped notch prior to bonding. The specimen is glued between two stubs
that allow the specimen to be gripped and loaded in tension during testing. The sharp chevron tip serves as a stressconcentration point and enables characterization of both weak and strong bonds. Fig. 1b shows the top view of the chevron specimen, and includes specification of important geometric parameters. The effective loading line is different from the specimen edge due to the finite area along which the stubs are glued. If the thickness or elastic properties of the two bonded layers of the specimen are not the same, a mixed-mode loading is produced at the interface. Recently, finite element (FE) modeling was used to analyze loading of a chevron specimen with layers of different thicknesses, in order to allow such specimens to be used for measurements of wafer bonds under mixed-mode conditions [17]. An experimental study of the mixed-mode interface toughness of Cu–Cu bonds has been completed using the chevron specimen [18]. It has been observed that the Cu–Cu bond toughness increases from 2.68 to 10.1 J m2 as the loading is changed from pure mode I to mixed-mode (w = 24°), for bonds created at 300 °C under a bonding pressure of
440
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
Fig. 2. Schematic diagrams of the four kinds of bonded interfaces studied. Blanket films and parallel lines involve crack propagation along continuously bonded interfaces, while pads and orthogonal lines require crack propagation across discontinuously bonded interfaces. ‘‘L” denotes the pitch of the patterns in the direction of crack propagation.
2.12 MPa. This is because the energy of plastic dissipation increases with increasing mode-mixity. In the present study, bonded chevron specimens were used to analyze the effects of patterning on Cu–Cu bond toughness under mode I and mixed-mode loading conditions. Two different kinds of bonded interfaces are considered: continuous and discontinuous (Fig. 2). Continuous interfaces present a homogeneous bond front along the direction of crack propagation, and are created by bonding blanket Cu films or by bonding patterned surfaces with lines aligned parallel to the crack propagation direction. Discontinuous interfaces have periodic bonded and unbonded regions along the direction of crack propagation. Two types of discontinuous interfaces are studied: bonded pads (discontinuous in two dimensions) and lines bonded orthogonal to the direction of crack propagation (orthogonal lines, discontinuous only along the propagation direction). ‘‘L” in Fig. 2 refers to the pitch of the patterns in the direction of crack propagation. The in-plane dimensions of the patterns are at least an order of magnitude greater than the film thickness, and therefore in-plane dimensional constraints are not expected to influence plasticity during debonding. 2. Experimental procedure 2.1. Specimen fabrication Thermally oxidized 100 mm diameter Si (1 0 0) wafers, with a nominal oxide thickness of 300 nm, were used as substrates for all thin film deposition and bonding studies. Two sets of wafers were fabricated for bonding studies. One set of wafers (set A) had patterned chevron structures, and the other set (set B) had patterned lines or pads. All bonding pairs comprised one wafer from each set. The chevron structures had the following dimensions: length
W (9 mm), width B (10 mm), notch angle h (=90°) and initial crack length a0 (=1 mm). The thickness ratio, g, is defined as the ratio between h2 and h1 (Fig. 1a). The thickness of the chevron wafers (set A) was kept constant (=475 lm). The thicknesses of wafers in set B were varied to induce mixed-mode loading conditions. Three different thicknesses were used: 475, 335 and 230 lm, corresponding to thickness ratios (g) of 1.0, 0.7 and 0.48 and phase angles (w) of 0°, 12.7° and 24°, respectively [17]. The chevron structures were patterned on oxidized wafers using standard photolithography. After patterning, the oxide film was etched using a buffered oxide etch (BOE) and deep reactive ion etching of the Si to a depth of 10 lm was carried out. Between adjacent chevron structures 0.5 mm wide alleys were patterned to enable dicing of individual specimens after bonding. After the creation of chevron structures, the wafers were cleaned in an O2 plasma, followed by a piranha clean (3:1 H2SO4:H2O2). The SiO2 film on the front side was then removed using BOE, and the chevron wafers were reoxidized thermally to an oxide thickness of 300 nm. The complementary set of wafers (set B) had patterned structures. Four different mask sets were used to create patterns. Pattern sizes/densities in the masks are outlined in Table 1. A Cu lift-off process was used to create patterned structures. After patterning image-reversal photoresist on wafer set B, 30 nm of Ta and 400 nm of Cu were successively deposited (without breaking vacuum) on both sets of wafers using e-beam evaporation. Ta served as both an adhesion promoter and a diffusion barrier between Cu and SiO2. The metal deposited on photoresist was lifted off the set B wafers by soaking them in NMP (N-methyl2-pyrrolidone) for at least 1 h, followed by light ultrasonic agitation in NMP and deionized water. The wafers from sets A and B were then prepared for thermocompression bonding. The detailed bonding procedure has been described elsewhere [18]. The patterned lines on the set A wafer were aligned to the chevron pattern in two different configurations: orthogonal and parallel. Schematic chevron structures with patterned lines and pads are shown in Fig. 3. The wafers were bonded at 300 °C for 1 h under a chamber vacuum of 103 Torr. The bonding pressure was maintained at a constant value of 2.12 MPa for all wafers. After the bonded pair was cooled to room temperature, the bonded wafers were taken out of the chamber and annealed in a tube furnace in a forming gas (95% Ar, 5% H2) ambient for 1 h. Fig. 4 shows a schematic cross-sectional view of the bonded stack. Individual
Table 1 Description of mask layouts used for the creation of patterns Pattern type
Pattern width (lm)
Spacing (lm)
Area fraction
Lines Pads Pads Pads
2–250 500 80 40
=width 100 40 20
0.5 0.69 0.44 0.44
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
441
Fig. 3. Schematic illustration of bonded chevron specimens with different pattern geometries. (a,b) Continuous interfaces (blanket films and parallel lines). (c,d) Discontinuous interfaces (orthogonal lines and pads). Note that the minimum and maximum ratios of the specimen width to the patterned dimensions were 20 and 5000, respectively.
Fig. 5. Typical load–displacement behavior from a chevron test. Fig. 4. Schematic cross-sectional view of a bonded chevron specimen. (The film thicknesses and substrate thicknesses are illustrated at different scales).
specimens for chevron testing were cut from the bonded wafers using a diamond saw. 2.2. Chevron testing Test specimens from bonded wafer-pairs were glued to aluminum stubs using a high-viscosity glue. Care was taken to prevent seepage of glue into the bonded interface. Approximately 15 specimens were tested for each geometric configuration. The specimens were selected to represent all regions on the bonded wafer. The glued specimens were tested under tension in an InstronTM 8848 electromechanical universal testing apparatus. The specimens were loaded at a constant displacement rate of 0.12 mm min1 and the load–displacement behavior was digitally recorded. A typical load–displacement curve for a chevronnotched specimen is shown in Fig. 5. After the initial elastic loading of the specimen, a crack initiates at the chevron
notch and propagates along the bonded interface. The triangle-shaped bonded area provides an increasing interfacial area with increasing crack length and load, and initially results in stable crack propagation. At some crack length the propagation becomes unstable and catastrophic failure of the specimen occurs; this coincides with the point of maximum load in Fig. 5. For a uniformly bonded interface, the bond toughness can be calculated from the maximum load in the test, Fmax, using [17]: GC ¼
F 2max Y 2min ; E B2 W
ð2Þ
where Ymin is the minimum value of the geometry function, E is the plane strain modulus, and B and W are the width and length of the specimen, respectively (Fig. 1b). For the anisotropic silicon specimens used in the current work, the effective plane strain modulus in the [0 0 1] direction in the (1 1 0) plane is used (E ¼ 144:7 GPa). FE analysis was used to determine the minimum value of the geometry function, which is dependent on the thickness of the
442
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
bonded layers and crack geometry [17]. Ymin is a function of the initial crack length ratio, a0 = a0/W, the total thickness of the specimen, htotal, and the thickness ratio, g: Y min ¼ ð25:71 þ 85:53a0 Þ sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ð2:91Þ ð2:91Þ htotal g htotal : þ 1þg 1þg
ð3Þ
The above equation is a fit to parametric FE results and estimates Ymin to within 1.8% of the FE calculations across a range of thicknesses (from 0.2 to 1.35 mm), thickness ratios (from 0.1 to 1), and initial crack length ratios (from 0.1 to 0.4) [17]. When the interface is patterned, the toughness value calculated from Eq. (2) must be corrected for the reduced bond area according to GC jpatterned ¼
1 GC junpatterned : Af
ð4Þ
The value of Af depends on the type and orientation of the pattern relative to the direction of crack propagation. When the pattern consists of lines oriented parallel to the direction of crack propagation, Af is the bonded area fraction. If the interface is patterned with lines that are oriented orthogonal to the direction of crack propagation, Af has a value of unity, as the bonded area is continuous along the crack front. For the bonded pad configuration, Af is the linear pad density along the crack front. From Eqs. (2)–(4), the interface toughness can be calculated from the maximum load measured in experiments and the geometry and elastic properties of the specimen. After testing, the specimens were dismounted from the studs using acetone and the studs were reused. Optical microscopy, scanning electron microscopy (SEM), atomic force microscopy and profilometry were then used to characterize the fracture surfaces.
Fig. 6. Mode I Fmax (failure load) vs. feature size for bonded blanket films, parallel and orthogonal lines, and pads. The failure load increases with decreasing feature size for discontinuous bonded interfaces, whereas the failure load for bonded parallel lines is independent of feature size. This trend is similar to that of bond toughness vs. feature size.
3. Results 3.1. Mode I bond toughness vs. pattern size Fig. 6 shows the variation in the failure load, Fmax (illustrated in Fig. 5), with respect to the pattern size. Since the specimen and chevron feature geometry parameters were the same across all tests, the failure load is related to the bond toughness by GC jpatterned /
1 2 F ; Af max
ð5Þ
based on Eqs. (2) and (4). Essentially, the failure load is a measure of the ‘‘global” toughness of the specimen, which, when normalized by the bonded area fraction, yields the ‘‘local” Cu–Cu bond toughness. Fig. 7 shows the variations in the mode I Cu–Cu bond toughness as a function of the feature size for all bonded configurations. For the parallel line configuration, the toughness was independent of the line width, and was mar-
Fig. 7. Mode I Cu–Cu bond toughness vs. feature size for bonded blanket films, parallel and orthogonal lines, and pads. For discontinuous bonded interfaces (orthogonal lines and pads), the toughness increases with decreasing feature size. For interfaces that are continuously bonded in the direction of crack propagation, the toughness is independent of feature size.
ginally higher than that of bonded blanket Cu films. On the other hand, the specimens bonded with orthogonal line configurations exhibited significant line width dependence, with the bond toughness increasing with decreasing line widths. Therefore, the degree of discontinuity along the crack propagation direction had a profound effect on the interfacial toughness. The other specimens with discontinuous interfaces, bonded pads, exhibited a similar trend in bond toughness with the bond toughness also increasing with decreasing pad width. The discontinuous bonded interfaces failed at higher loads as the pattern size decreased, while the bonded
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
443
parallel lines did not show any dependence of the failure load on pattern size. This trend shown in Fig. 6 is similar to that of the bond toughness (Fig. 7), even though the data in Fig. 6 are not normalized with respect to the bonded area fractions. The difference between Figs. 6 and 7 is that the bonded parallel lines, which failed at lower loads than bonded blanket films, exhibited higher toughness than the latter when bonded area fractions were invoked. 3.2. Mixed-mode bond toughness vs. pattern size The Cu–Cu bond toughness was found to be independent of line width for all mode-mixities under conditions for which the direction of crack propagation was parallel to the lines, as seen in Fig. 8. As seen in Fig. 9a, for the orthogonal line configuration, the toughness increased with line width for all mode-mixities and the magnitude of mode-mixity effect increased with decreasing line width. A common trend between the parallel and orthogonal line configurations is that the interfacial toughness increases with increasing phase angle, as discussed in Ref. [18]. This behavior is captured in Fig. 9b. 3.3. Failure loci The locus of failure was found to be dependent on the mode-mixity and the pattern orientation, with three different failure locations/mechanisms observed: at the Ta–SiO2 interface (adhesive failure), at the Cu–Cu bonded interface (cohesive failure) and in the bulk Cu (also cohesive). Roughly equal percentages of adhesive- and cohesivebonded interface failures were observed in mode I testing of parallel lines. The specimens tested under mixed-mode conditions had increasing incidence of adhesive failures, indicating crack deflection from the Cu–Cu bond interface
Fig. 8. Toughness of bonded Cu–Cu parallel lines as a function of line width for three different phase angles. Mode-mixity was achieved using asymmetric chevron specimens. Schematics of chevron specimens and cross-section of bonded stacks were shown in Figs. 3b and 4, respectively. For a given phase angle, the toughness is independent of line width for parallel lines.
Fig. 9. (a) Cu–Cu bond toughness as a function of line width for different mode-mixities. The orthogonal line configurations exhibit significant line width dependence, whereas the toughness of parallel lines is independent of line width. (b) Cu–Cu bond toughness as a function of phase angle for parallel and orthogonal lines. As reported in Ref. [18], the bond toughness is a strong function of the phase angle of loading.
to the Ta/SiO2 interface. Another difference between mode I and mixed-mode failure loci was that the mode I cohesive failures occurred along the bonded interface whereas the mixed-mode cohesive failures occurred within the bulk Cu, again confirming crack path deflection. Specimens with bonded orthogonal lines and pads exhibited predominantly adhesive failures, irrespective of the mode-mixity. Fig. 10 shows SEM images of the mode I fractured surfaces of the bonded lines. Failure at the bonded Cu–Cu interface for bonded parallel lines is shown in Fig. 10a and b, while bulk Cu failure (away from the bonded interface) in bonded orthogonal lines is shown in Fig. 10c and d. The fracture surfaces of orthogonal lines show rough features, indicative of enhanced plastic deformation. Fig. 11a and b show mixed mode cohesive failures of specimens tested under a phase angle of 12.7°, and Fig. 11c and d show
444
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
Fig. 10. Mode I fracture surfaces of bonded lines. (a,b) Parallel lines; (c,d) orthogonal lines (line width = 100 lm). Orthogonal lines show rough fracture surfaces, indicative of significant plastic deformation.
Fig. 11. Mixed-mode fracture surfaces of bonded lines. (a,b) w = 12.7°; (c) and (d) w = 24°. The roughness of the Cu fracture surfaces increases with increasing mode-mixity, indicating increased plasticity in ductile failure.
specimens tested under a phase angle of 24°. Evidence of ductile failure is particularly apparent in the samples loaded
with a 24° phase angle, for which a dimpled fracture surface morphology was observed (Fig. 11d).
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
445
Fig. 12. Mode I fracture surfaces of bonded pads. w = pad width. (a,b) w = 500 lm; (c,d) w = 80 lm; (e,f) w = 40 lm. Bright and dark pads in (e) represent Cu–Cu and Ta/SiO2 failure modes, respectively. (a,c,e) Mixed failure modes (Ta/SiO2 and Cu–Cu); (d,e,f) magnified images of the Cu–Cu fracture surfaces.
Fracture surfaces of bonded pads are shown in Fig. 12. Micrographs ‘a’ and ‘b’ show the surfaces of the widest pads (500 lm), ‘c’ and ‘d’ show pads of intermediate width (80 lm) and ‘e’ and ‘f’ show the smallest pads (40 lm). The roughness of the fracture surfaces increases with decreasing pad size, suggesting more plastic energy dissipation. A classic ductile cup-and-cone fracture surface can be seen in the micrographs of the smallest pads (Fig. 12f). 4. Discussion 4.1. Orthogonal lines and pads The orientation and size of features have a profound effect not only on the toughness of Cu–Cu bonds (Figs. 7
and 9), but also on the failure load (Fig. 6). The latter shows that patterning has a fundamental effect on the toughness of discontinuous bonded interfaces, since they fail at higher loads with decreasing pattern size, irrespective of the bonded area fraction. The pattern orientation determines the load required to fracture the interface, since the continuous patterns reduce the load compared with blanket interfaces, while the discontinuous patterns increase the load compared with blanket interfaces. Interestingly, pattern size dependence is observed only for bonded interfaces that are discontinuous in the direction of crack propagation, for which the bond toughness significantly increases as the pad or line size is decreased. Therefore, patterning effects can be traced to the nature of crack propagation along a heterogeneous
446
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
bonded interface. For the purpose of this discussion, only mode I toughness is considered, so that a direct comparison can be made between bonded lines and pads. However, the arguments below are expected to hold for mixed-mode toughness as well, the only difference being the mode-mixity contribution to enhanced plastic deformation, which has been discussed in [18]. Both orthogonal line and pad configurations involve crack propagation across periodic gaps in bonded regions. Since the bonded interface lies between ductile Cu layers on both sides, there is an energetic cost to crack initiation (or nucleation), involving void growth and coalescence ahead of the crack tip. The measured bond toughness is plotted as a function of the number of crack initiation events in Fig. 13. The number of crack initiation events is obtained by dividing the distance d traversed by the crack prior to complete specimen failure by the pitch of the patterns L (along the crack propagation direction), as shown schematically in Fig. 2. The value of d, estimated to be 3 mm from FE modeling [17], has been verified experimentally using a combination of load–displacement plots and fracture surface analyses. In Fig. 13 it can be seen that the bond toughness of orthogonal lines and pads increases monotonically with increasing number of crack initiation events, suggesting that more energy is absorbed in bond configurations for cracks that have to nucleate at multiple points. It was assumed that cracks nucleate simultaneously across all the pads along a row perpendicular to the crack propagation direction and therefore the row is considered as a single initiation event. However, there is still a discrepancy between the curves that relate toughness to the number of crack initiation events, as determined for pads versus lines, which is not clearly understood at this point.
Fracture surfaces of the orthogonal lines (Fig. 10c and d) and pads (Fig. 12) reveal rough features, indicative of ductile fracture. The roughness, and therefore, the plastic energy dissipation, are greater than that observed for parallel lines (Fig. 10a and b), providing further evidence for enhanced toughness due to a greater number of crack initiation events. Recently, Kubair and Spearing [19] used a spectral scheme with cohesive zone modeling to treat the effect of discontinuous patterning on the energy of interfacial fracture, using a square wave variation in the critical distance of separation of the interfaces. They predicted an increase in the apparent bond toughness associated with an increased (plastic) strain energy for decohesion of patterned bonded surfaces compared with unpatterned bonded surfaces, consistent with the results on bonded Au pads from [6]. However, they also predicted a weak dependence of the measured bond toughness on the repeat period of the pattern, which contrasts with the results found in the current study. The Kubair and Spearing treatment also does not apply to variations in bonded surfaces associated with two-dimensionally periodic patterns. 4.2. Parallel lines For the range of line widths considered (2–250 lm), the toughness of bonded parallel lines was independent of line width for all mode-mixities. Litteken and Dauskardt observed a line-width dependence of adhesion at interfaces of thin-film structures containing patterned polymer lines, even when the direction of crack propagation was parallel to the line axes [20]. In their work, decreasing the line width was found to significantly increase the interfacial fracture energy. The increase in adhesion was associated with increasing contributions from plastic energy dissipation in the patterned lines owing to the decrease in the lateral constraint of the lines with decreasing line width. The observed fracture resistance behavior was similar to the transition from plane strain to plane stress fracture commonly found in bulk metals and polymers. In the current work, the length scales (>2 lm) considered are still in the regime of plane strain for Cu films and, therefore, are not expected to have an effect on the interfacial adhesion. Based on the work of Litteken and Dauskardt, it is expected that narrower Cu lines (sub-micron widths) would show plane stress behavior and enhanced deformation and bond toughness, even under the parallel configuration. 5. Summary
Fig. 13. Toughness of bonded orthogonal lines and pads (from Fig. 6) plotted as a function of the number of crack initiation events. The number of crack initiation events is inversely proportional to the repeat-length of the features, L, along the direction of crack propagation. An increasing number of crack initiation events results in a greater energetic cost for crack propagation and, hence, higher toughness.
Mode I and mixed-mode chevron tests on patterned Cu– Cu bonded interfaces reveal interesting dependencies of bond toughness on pattern size, density, and orientation. For cracks running parallel to bonded lines, the interfacial adhesion was independent of line widths for widths as low as 2 lm. On the other hand, the bond toughness was a strong function of the line width/pad size when cracks propagated
R. Tadepalli et al. / Acta Materialia 56 (2008) 438–447
across discontinuously bonded interfaces. In this case, for a given bonded area fraction, decreasing the feature size was found to significantly increase the decohesion energy. This increase was attributed to the greater number of crack initiation events for the smaller features. This semi-quantitative empirical model captures the effects of pattern size (repeat distance) and orientation on bond toughness. Mixed-mode bond toughness of orthogonal lines also followed similar trends as mode I toughness. As discussed in our earlier work [18], increased mode-mixity contributes to enhanced plastic deformation and, therefore, to greater bond toughness compared with specimens loaded in pure tension. References [1] Fan A, Rahman A, Reif R. Electrochem Solid State Lett 1999;2(10):534. [2] Schmidt MA. Proc IEEE 1998;86(8):1575. [3] Celler GK, Cristoloveanu S. J Appl Phys 2003;93(9):4955. [4] Plo¨ßl A, Kra¨uter G. Mater Sci Eng 1999;R25(1–2):1.
447
[5] Cozma A, Puers R. J Micromech Microeng 1995;5(2):98. [6] Tsau CH, Spearing SM, Schmidt MA. J Microelectromech Syst 2004;13(6):963. [7] Turner KT, Spearing SM. J Appl Phys 2002;92(12):7658. [8] Miki N, Zhang X, Khanna R, Ayon AA, Ward D, Spearing SM. Sensors Actuator A 2003;103(1–2):194. [9] Rahman A, Reif R. IEEE Trans VLSI Syst 2000;8(6):671. [10] Tan CS. Multi-layer three-dimensional silicon electronics enabled by wafer bonding, Ph.D. thesis, MIT, 2006. [11] Tadepalli R, Thompson CV. In: Proceedings of the IEEE 2003 international interconnect technical conference; 2003. p. 36. [12] Litteken CS, Dauskardt R, Scherban T, Xu G, Leu J, Gracias D, Sun B. In: Proceedings of the international interconnect technical conference; 2003. p. 168. [13] Evans AG, Hutchinson JW. Acta Metall Mater 1989;37(3):909. [14] Tvergaard V, Hutchinson JW. J Mech Phys Sol 1993;41(6):1119. [15] Hurd DS, Caretta R, Gerberich WW. J Mater Res 1995;10(2):387. [16] Bagdahn J, Petzold M, Plo¨ßl A, Wiemer M. Electrochem Soc Proc 2001;99(35):218. [17] Tadepalli R, Turner KT. Eng Fract Mech in press. [18] Tadepalli R, Turner KT, Thompson CV. J Mech Phys Solids, in press. [19] Kubair DV, Spearing SM. J Phys D 2006;39(6):1050. [20] Litteken CS, Dauskardt RH. Int J Fracture 2003;119/120:475.