Microelectronics Reliability

Microelectronics Reliability 49 (2009) 269 287 Contents lists available at ScienceDirect Microelectronics Reliability journal homepage: www.elsevier.com/locate/microrel Interfacial fracture toughness of Pb-free solders S.M. Hayes a,b, N. Chawla b, *, D.R. Frear a a Advanced Packaging and Systems Integration, Freescale Semiconductor, Tempe, AZ 85284, United States b School of Materials, Fulton School of Engineering, Arizona State University, Tempe, AZ 85287-8706, United States article info abstract Article history: Received 29 August 2008 Received in revised form 10 November 2008 Available online 19 December 2008 Increasing environmental concerns and pending government regulations have pressured microelectronic manufacturers to find suitable alternatives to Pb-bearing solders traditionally used in electronics packaging. Over recent years, Sn-rich solders have received significant attention as suitable replacements for Pbbearing solders. Understanding the behavior of intermetallics in Sn-rich solders is of particular concern as the microelectronics industry progresses towards Pb-free packaging. The formation of intermetallic compounds results from the reaction of the solder with the metallization on the substrate in the electronic package. While the presence of the intermetallic is an indication of good wetting, excessive growth of the intermetallic can have a dramatically adverse effect on the toughness and reliability of the solder joint. Understanding their fracture behavior will lend insight to their reliability under mechanical and thermomechanical strains. We investigated the intermetallic compound growth associated with Sn 0.7Cu and Sn 4.0Ag 0.5Cu solders on Ni Au, Ni Pd, and Cu substrates. (Ni,Cu) 3 Sn 4 was present at the Ni interface for both solders but was coarser in the case of Ni Pd. Cu 6 Sn 5 and Cu 3 Sn were observed for both solder types. The Cu 3 Sn layer was similar in thickness and appearance for both solders, but the Cu 6 Sn 5 was smoother and rounder in the case of Sn 0.7Cu. Additional time above liquidus resulted in growth of the Cu 6 Sn 5 layer and eventual spalling of the IMC grains. The effect of the intermetallic on the toughness (K Q ) of the solder joint was investigated using a modified compact tension specimen. Typical failure modes included bulk solder failure, intergranular separation, and intermetallic fracture, or cleavage. In some cases, additional time above solder liquidus was used to shift the dominant failure mode from that dominated by the bulk solder to interfacial delamination through the intermetallics. Solder joint fracture toughness was different between Ni Sn and Cu Sn interfacial intermetallics and was also affected by the relative intermetallic thickness. The relationship between solder and intermetallic microstructure and mechanical properties is discussed. Ó 2008 Elsevier Ltd. All rights reserved. 1. Introduction The microstructure and mechanical behavior of traditional Sn Pb solders has been well documented and is well understood [1 6]. Recently, environmental concerns, legislation, and even customer preference are driving the microelectronics industry towards implementation of Pb-free solders in packaging [7 10]. The electronics industry has settled on a few Sn-rich solders after careful consideration of mechanical behavior, cost, processibility, and integration. Sn 3.9Ag 0.5Cu, and Sn 0.7Cu are two common Pb-free solder candidates for replacing the traditional Sn Pb solder systems. Upon reflow, Sn-rich solders readily react with several metallizations to form interfacial intermetallics. Typical pad metallizations include Ni Au, Cu, or Ni Pd. When reflowing Sn-rich * Corresponding author. E-mail address: nchawla@asu.edu (N. Chawla). solders on such pad metallizations, several intermetallics may be formed, such as Cu 6 Sn 5,Ni 3 Sn 4,Cu 3 Sn, and Ag 3 Sn. While the presence of these intermetallic compounds is necessary to promote a proper bond and adhesion between the solder and metallization, at large thicknesses the intermetallic layer can significantly decrease the fracture toughness of the joint. Fig. 1 shows the relationship between apparent fracture toughness (measured using a compact tension geometry) and intermetallic thickness for several Pb-free and Pb-bearing solders reflowed on Cu undermetal (after Frear et al. [1]). Note that with increasing intermetallic layer thickness (particularly above 5 10 lm), the toughness of the joint decreases significantly. Thus, the brittle nature of most of these intermetallics, as well as the lower ductility of Pb-free solders compared to Pb Sn, have caused some concern regarding the reliability of Pb-free solders. In particular, the growth rate of the intermetallics is accelerated for Sn-rich solders due to larger amounts of Sn available for reaction with the pad metallizations. Furthermore, the reliability of Sn-rich solders is degraded in cases of high strain 0026-2714/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.microrel.2008.11.004

270 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 KQ (MPa*m 1/2 ) 12 10 8 6 4 2 0 0 5 10 15 20 25 30 Intermetallic Thickness (μm) 40In-40Sn-20Pb 50Sn-50In 50Pb-50In 97Sn-3Ag 95Sn-5Sb 60Sn-40Pb Fig. 1. Apparent fracture toughness, K Q, of the various solder alloys as a function of the thickness of the interfacial intermetallic compound layer (after Frear et al. [1]). W a Fig. 2. Modified compact tensile specimen for the measurement of interfacial fracture toughness where W = 2.54 cm, B = 0.953 cm and a = 0.8 cm. B rates (approximately >10 1 /s) that can be experienced primarily during pure mechanical shock rather than lower rates more common to thermal cycling. In the case of higher strain rate loading, failure analysis of the solder joint typically shows that fracture takes place at the intermetallics [1,7]. Several studies have been performed to evaluate the fracture behavior of Pb Sn solder joints. The majority of these studies have reported that Pb Sn solder joints have a tendency to fail through the bulk solder, although some studies suggested that the intermetallics at the solder-undermetal interface acted as crack initiation sites [1]. In the case of the bulk solder failures, fracture was attributed to a coarsening of the Pb Sn microstructure. Kang et al. [2] attempted to relate decreased fatigue life to the thickness of Ni 3 Sn 4 intermetallics formed in a Pb Sn solder joint on Ni metallization. They thermally cycled Si power transistors soldered to Ni using Pb Sn solders. They found that failure sometimes took place near the solder/intermetallic to undermetal interface, but within the solder rather than the intermetallic. Keller [3] found that tensile testing of previously thermal cycled joints to failure resulted in fracture through the intermetallic region while tests of non-cycled joints showed joint failure through the bulk solder. Thermal cycling caused further growth and coarsening of the intermetallics making them more susceptible to fracture. Hall et al. [5] have shown similar results from aged joints using 60Sn 40Pb. It appears that thermal aging induced intermetallic growth in Sn Pb on Cu solder joints increased interfacial failure and decreased failure through the solder itself. Frear and Vianco [7] showed that we can expect a similar trend in Pb-free solders, but at an accelerated rate, due to faster intermetallic growth rate associated with Sn-rich solders. They also indicated that shear tests performed at a strain rate of round 6.6 10 4 /s and with the intermetallic thickness less than 10% of the overall joint thickness, fracture took place primarily through the solder. In tests performed in tension, fracture in both the solder and within the interfacial intermetallic region was observed. More recently, Kurosaka et al. [11] related aging of Sn Pb solder joints on Ni pads to lower tensile strength. They also reported on the aging effects related to Sn 3Ag 0.5Cu on Ni joints. In both cases, the failure mode shifted from bulk solder failure to brittle intermetallic failure along with a corresponding decrease in tensile strength. Again, intermetallic growth and coarsening brought on by thermal aging decreased the strength of the solder joint. Understanding the kinetics of intermetallic formation in these solder and metallization systems is important to the prediction of solder joint reliability. For example, Kurosaka et al. [11] showed a significant change in the interfacial morphology of Sn 3.0Ag 0.5Cu on Ni when finishing the Ni with approximately 60 nm of Pd. They found the that the growth of (Cu,Ni) 6 Sn 5 was minimized by the small layer of Pd. Microstructure characterization of subsequent pull tests revealed ductile failures occurring within the solder while samples made without the Pd finish showed evidence of less desirable brittle failures occurring within the intermetallics. It has also been reported that thick Cu 6 Sn 5 tends to have deleterious effects on reliability while a thin, irregular layer of the same phase Fig. 3. Macro, optical microscope and SEM views of a polished fracture toughness sample surface before tensile testing.

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 271 Table 1 Thickness (lm) of metallizations used in this study. Cu Ni Au Ni Pd Sn 4.0Ag 0.5Cu 7.5 ± 2.5 6.5 ± 0.5/0.5 ± 0.1 4.0 ± 1.0/1.25 ± 0.25 Sn 0.7Cu 7.5 ± 2.5 6.5 ± 0.5/0.5 ± 0.1 4.0 ± 1.0/1.25 ± 0.25 Table 2 Liquidus and reflow temperatures of Pb-free solder alloys investigated. Liquidus ( C) Sn 4.0Ag 0.5Cu 217 220 260 ± 3 Sn 0.7Cu 227 267 ± 3 Reflow temperature ( C) still indicates good solder wetting but without the decreased reliability [12 14]. Furthermore, intermetallic thickness has been shown to affect the fracture behavior of solder joints under shear stress. Deng et al. [15] found that total intermetallic (Cu 6 Sn 5 and Cu 3 Sn) thickness greater than 20 lm tended to cause fracture along the intermetallic interfaces under shear failure. Deng et al. [15] and Chawla and Sidhu [16] used Finite-element analysis (FEA) to predict shear stress shear strain curves related to increased intermetallic thickness as well as the morphology. It was found that for both planar and nodular intermetallics, greater intermetallic thickness leads to higher shear stress at the same shear strain. They also showed that nodular morphology of the intermetallics resulted in higher shear strength than planar intermetallics of the same thickness. FEA modeling results of nodular intermetallics reported by Chawla and Sidhu [16] showed stress concentrations in the nodules suggesting that they serve as crack initiation sites. The conditions for mechanical testing will also have a significant influence on the failure modes for Sn-bearing solders. While very low strain rates will tend to cause ductile solder failures, higher deformation rates resulting from mechanical shock or drop tests can cause the failure modes to take place in the IMC region. Frear [17] studied the affect of deformation rate on Sn 40Pb on Cu fractures in a shear orientation. Frear [17] discovered that fracture occurred through the intermetallics for shear strain rates greater than 6.6 10 4 /s. Below that rate, failure tended to occur within the bulk solder. It should be noted that intermetallic thicknesses Fig. 4. Comparison of intermetallic morphology found on Ni Au and Ni Pd metallization for both Sn 0.7Cu and Sn 4.0Ag 0.5Cu solders after 45 s of solder reflow. Fig. 5. Evolution of intermetallic morphologies found on Cu metallization for various reflow times.

272 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 8 7 Sn0.7Cu Sn4.0Ag0.5Cu 2 Sample 1 Sample 2 Sample 3 A P max =P Q Total Intermetallic Thickness (μm) 6 5 4 3 2 1 0 0 50 100 150 200 Reflow Time (s) Fig. 6. Evolution of intermetallics on Cu metallization as a function of reflow time. Load (kn) 1.5 1 0.5 O 0 0 0.05 0.1 0.15 0.2 0.25 0.3 0.35 Displacement (mm) Fig. 8. Load-displacement curves collected for Sn 4.0Ag 0.5Cu samples after 180 s of reflow with P max and P Q indicated. will also have a significant effect on failure location with failures generally shifting from bulk solder to intermetallic as the intermetallic grows. While the intermetallics of Pb Sn and Pb-free solders on different undermetals have been studied, as summarized above, comprehensive studies characterizing the deformation behavior of the intermetallics and the resulting fracture toughness seem to be missing. In this study we have characterized the fracture behavior of several Sn-rich, Pb-free solder joints under fracture toughness testing. Sn 4.0Ag 0.5Cu and Sn 0.7Cu and three metallizations, Cu, Ni Au, and Ni Pd, were investigated to create six combinations. The growth of the intermetallics over extended reflow times of 45, 90, and 180 s was also studied on selected combinations. The relative fracture toughness was measured using a modified compact tension geometry. The fracture behavior was studied and related to the intermetallic morphology and microstructure, as well as solder microstructure. This study presents a thorough evaluation of the relationship between the intermetallics formed during reflow (described in the companion paper, part I), their fracture behavior, and the resulting solder joint fracture toughness in Cu and Ni-containing metallization. 2. Materials and experimental procedure A compact tension specimen, as specified in ASTM standard E399-90, was used to measure fracture toughness. The sample geometry is depicted in Fig. 2 and is modified from that described in the standard as it is comprised of two brass pieces joined by a solder joint. The brass pieces are machined according to the standard but were slightly longer in dimension B such that multiple samples could be created from a single piece. Geometric tolerances of the samples were taken into consideration and accommodated for by measurement of individual samples. While the actual geometry varied slightly from the ASTM standard, the K Q values obtained are still valid for comparison within this study. The bars were wire cut using electro-discharge machining (EDM) to an approximate width of 9.6 mm. Each side of the cut sample was mechanically ground and polished to a 1 lm surface finish. Fig. 3 shows views of a sample after polishing. Grinding Load, P A P max A P max A P max =P Q P 5 =P Q P 5 P Q P 5 Type I Type II Type III O O O Displacement, v Fig. 7. Method of determining P Q from the three possible types of load-displacement curves measured during mechanical testing.

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 273 Fig. 9a. Apparent fracture toughness (K Q ) of the solder joints investigated. The number following the cell description indicates seconds of reflow time. One sigma standard deviation error bars are included. and polishing resulted in sample widths of 9.53 mm ± 0.05 mm (B dimension). A suspension of 0.05 lm colloidal silica and a soft polishing cloth was then used to further polish and lightly etch the sample. The resulting surface finish allowed for measurement of intermetallic layers and observation of the crack opening and fracture analysis of the samples after the toughness test. The three different solderable finishes include 5 10 lm of electroplated Cu, 3 5 lm of electroplated Ni with 1.0 1.5 lm of electroless-pd and 6 7 lm of electroless Ni with 0.4 0.6 lm of electroplated Au. The difference in thickness of the electroless Ni and electroplated Ni layers was due to the inherent differences in the plating techniques and their specific process conditions. 10 9 The plated samples, still masked, were then prepared for soldering. Rosin-mild-activated flux was applied to the exposed plated faces of the samples and then pre-heated to 200 C. The sample bars were then briefly immersed in molten solder (Sn 0.7Cu or Sn 4.0Ag 0.5Cu) to form a thin solder layer of approximately 25 lm. The solder was held in the pot at approximately 40 C above the liquidus temperature of the solder type being processed. This initial thin layer of solder was necessary to prevent large voiding in the solder joint. The two bars were then recoated with RMA flux and clamped together with a 0.010 (254 lm) shim placed between the bars to hold the two pre-soldered faces at the distance of approximately 250 lm. Measurements from subsequent cross-sections showed the resulting joint thickness to be 254 lm ±10. The two clamped bars were then immersed as one into the molten solder for a prescribed time of 45, 90, or 180 s. The joined sample bars were removed and allowed to cool in air. A screw-driven mechanical testing system was used to perform fracture toughness experiments by the standard method 8 KQ (MPa*m 1/2 ) 7 6 5 Sn0.7Cu-Cu Sn4.0Ag0.5Cu-Cu Sn0.7Cu-NiPd Sn4.0Ag0.5Cu-NiPd Sn4.0Ag0.5Cu-NiAu 4 0 1 2 3 4 5 6 7 8 Intermetallic Thickness (μm) Fig. 9b. Apparent fracture toughness, K Q, of the two solder alloys and three undermetal types as a function of the interfacial intermetallic compound layer. Fig. 10. Profile view of a fractured Sn 0.7Cu on Ni Au joint after 45 s of reflow.

274 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 Fig. 11. SEM micrographs and corresponding EDS spectrums for opposing surfaces of the fracture surfaces, (a) and (b), of Sn 0.7Cu on Ni Au joint with 45 s of reflow. for plane strain fracture toughness of metallic materials, as described in ASTM E399. Polished samples were loaded into the fixtures and secured with pins inserted through holes in the sample and specially designed clevis grip. A 5 N preload was applied to remove any initial slack in the setup. The samples were then pulled to failure in tension at a rate of 2.1 10 3 mm/s or an apparent strain rate of 8.3 10 3 /s. This strain rate is in the range of mechanical shock [1]. The apparent displacement used to calculate the strain rate was measured by the movement of the cross head. Compliance of the testing hardware and the brass samples was measured and found to be linear over the load range experience by all the samples. Therefore, the compliance had a minimal impact on the determination of K Q. Typical failure loads were on the order of 2 kn and all tests were performed with a 5 kn load cell. At least three samples for each condition were tested. Fracture toughness values were calculated using ASTM E399. Details of this method are described in the results section. Once failure occurred, each of the fracture surfaces was carefully removed and collected for fracture surface analysis. Analysis was conducted to determine whether the fracture occurred in the bulk solder, through cleavage of an intermetallic layer, or delamination of the intermetallic solder interface. Samples from each cell were characterized using optical microscopy as well as Scanning Electron Microscopy (SEM) and Energy Dispersive Spectroscopy (EDS) (see Table 1). 3.2. Intermetallic formation on Ni metallization Sn 0.7Cu solder and Ni Au reacted to form a relatively thin, continuous layer of Ni 3 Sn 4 intermetallic when reflowed for 45 s. The intermetallic was found to be approximately 1.5 lm in thickness. There was no significant level of Au detected in the solder or near the interface so it is probable that the Au is predominantly in solution within the solder. Fig. 4 shows the microstructures formed after 45 s of reflow for Ni Au and Ni Pd on both Sn 0.7Cu and Sn 4.0Ag 0.5Cu solders. When reacting Sn 0.7Cu solder on Ni Pd for 45 s the (Ni,- Cu) 3 Sn 4 intermetallic was again formed. Although the average thickness was similar to that formed on Ni Au at 1.5 lm, the general appearance was much less regular and coarser. Some IMC nod- 3. Results and discussion 3.1. Microstructure characterization IMC thickness and morphology are very important in controlling the fracture toughness of the solder joint. Thus, the microstructure of the samples was characterized prior to fracture toughness testing in order to quantify the IMC thicknesses and morphology. Solder joint thickness measurements were performed on polished samples using an optical microscope. The solder joint thickness for all specimens was approximately 250 lm (see Table 2). Fig. 12. Sn 0.7Cu solder on Ni Pd with 45 s of reflow showing ductile failure through of the solder.

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 275 ules were around 4 5 lm in thickness. Upon 180 s of reflow the (Ni,Cu) 3 Sn 4 layer became thicker and more uniform with an average thickness of 2.5 lm. For both reflow times the Pd was completed consumed by the solder with no Pd detected by EDS near the interfacial region or in any significant formations within the solder. The Pd is likely in solution within the solder. For Sn 4.0Ag 0.5Cu on Ni Au after 45 s of reflow, the (Ni,- Cu) 3 Sn 4 layer was approximately 1 lm thick. Compared to the appearance of the Sn 0.7Cu solder on Ni Au, the (Ni,Cu) 3 Sn 4 layer in this case was somewhat finer but more irregular. Reactions of Sn and Ag in the bulk solder may have affected the interfacial reaction resulting in the difference in intermetallic appearance between Fig. 13. Sn 0.7Cu solder on Ni Pd with 180 s of reflow showing opposing fracture surfaces (I, II) of the (Ni,Cu) 3 Sn 4 phase. Fig. 14. Opposing fracture surfaces of opposing sides (I) and (II) showing (a and c) intermetallic cleavage very near the Ni undermetal and (b and d) intermetallic cleavage through the middle of the (Ni,Cu) 3 Sn 4 for Sn 4.0Ag 0.5Cu solder on Ni Au with 45 s of reflow.

276 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 in the thickness of the electroless-pd deposition could explain the variations in IMC thickness. Thinner Pd would be consumed sooner than thicker Pd regions resulting in greater Ni Sn reaction time and the associated thicker IMC. Again, there was no indication of Ag near the IMC region. In general, the (Ni,Cu) 3 Sn 4 formed on Ni Pd was found to be coarser and with greater variation in thickness than that found on Ni Au. Once again, the IMC thickness variation could be due to thickness uniformity of the electroless plated deposition of Pd. 3.3. Intermetallic formation on Cu metallization Fig. 15. Profile view of the crack opening through (Ni,Cu) 3 Sn 4 intermetallic for Sn 4.0Ag 0.5Cu solder on Ni Pd with a 45 s reflow. Sn 4.0Ag 0.5Cu and Sn 0.7Cu solders. There was no indication of Ag within the region of continuous interface intermetallic. Reacting Sn 4.0Ag 0.5Cu with Ni Pd resulted in a somewhat thicker (Ni,Cu) 3 Sn 4 layer, approximately 2 lm in thickness. However, there was a significant variation in the layer thickness with some continuous areas being as thick as 3 4 lm. Some variation The evolution of intermetallic growth on Cu metallization is shown in Fig. 5. Sn 0.7Cu solder formed the characteristic Cu 6 Sn 5 and Cu 3 Sn intermetallics when reflowed on Cu for 45 s. The Cu 6 Sn 5 layer was continuous and relatively regular in both morphology and thickness of 3.5 lm. The individual nodules appeared to be approximately 2.5 lm in diameter. Situated between the Cu 6 Sn 5 layer and the Cu undermetal was the Cu 3 Sn layer. This layer was approximately 0.75 lm thick with excellent uniformity. Increasing the reflow time to 90 s resulted in coarsening of the grains up to 4.5 lm in diameter and thickening of the Cu 6 Sn 5 layer to approximately 4.5 lm in thickness. After 180 s of reflow the Cu 6 Sn 5 continued to become even coarser with some grains approaching 7 lm in diameter. The thickness also continued to grow greater than 7 lm. There was some minor spalling observed with the larger grains as they grew larger and became unstable during reflow. Fig. 16. Opposing fracture surfaces of opposing sides (I) and (II) showing (a and c) intermetallic cleavage through the middle of the (Ni,Cu) 3 Sn 4 and (b and d) intermetallic cleavage very near the Ni undermetal for Sn 4.0Ag 0.5Cu on Ni Pd with 45 s of reflow.

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 277 Fig. 17. Typical ductile solder failure found with Sn 0.7Cu solder on Cu reflowed for 45 s. Formation of the characteristic Cu 6 Sn 5 and Cu 3 Sn intermetallics was also observed for Sn 4.0Ag 0.5Cu solder on Cu metallization. After 45 s the Cu 6 Sn 5 layer thickness averaged 2.5 lm with apparent grain diameters also around 2.5 lm. The underlying Cu 3 Sn layer was continuous and approximately 0.75 lm thick. The Cu 3 Sn was very similar in appearance to that of Sn 0.7Cu solder. The grains appear somewhat less round and smooth along their boundaries with the bulk solder. Even after 45 s of reflow they exhibited more faceted growth than what was found with Sn 0.7Cu solder. Plate-like formations of Ag 3 Sn were also seen in bulk solder. After 90 s of reflow the Cu 6 Sn 5 grains have elongated to as great as 5 lm while not showing significant signs of coarsening along the plane of the Cu undermetal. After 90 s of reflow there was evidence of Cu 6 Sn 5 spalling once the individual grains became long enough and they detached from their base during reflow. After 180 s of reflow the base of the Cu 6 Sn 5 layer became more continuous but the overall thickness of the grains did not increase significantly. This limiting effect was likely caused by the spalling that was first noted on the 90 s reflow samples. The spallation also created an IMC rich region within the solder but near the attached Cu 6 Sn 5 layer. This particle reinforced composite like structure may have an impact on the mechanical performance of this joint. Ag 3 Sn plate-like features formed during cooling were found in closer proximity to the Cu 6 Sn 5 layer than in samples with shorter reflow time. The Cu 3 Sn layer was unchanged from 45 to 90 to 180 s of reflow. Fig. 6 shows the measured total intermetallic thickness as a function of reflow time, for both alloys reflowed on copper. The main differences between the intermetallic formation created on Cu with Sn 0.7Cu and Sn 4.0Ag 0.5Cu solders was the faceted morphology and more elongated growth of the Cu 6 Sn 5 associated with the Ag-bearing solder. Fig. 18. (a and b) Right to left transition from ductile solder failure to interfacial delamination and (c and d) opposing surfaces of the interfacial failure for Sn 0.7Cu solder on Cu after 45 s of reflow.

278 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 Fig. 19. (a) Faceted Cu 6 Sn 5 grains and the corresponding (b) faceted depressions with smooth fracture and (c) rough fracture, all for Sn 0.7Cu on Cu with 45 s of reflow. 3.4. Fracture toughness of solder joints The pure mode I fracture toughness, K IC, can be calculated for monolithic materials with a known pre-crack length and a crack that propagates by mode I. Calculating K IC assumes linear elastic fracture mechanics with small scale yielding. While the crack in the solder joint may partially grow in mode I, some mixed mode crack growth and plasticity within the solder may be present. Therefore it is not reasonable to use the term K IC for these soldered samples. As a result, the fracture toughness of the solder joints, using a modified compact tension geometry tested under ASTM standard E399 (as mentioned in the experimental procedure), was quantified by the critical stress intensity factor K Q. The K Q values calculated in this study are useful for relative comparisons of the solder alloys, undermetallurgy, and corresponding intermetallics. K Q, or the relative fracture toughness, was calculated by the following equation: K Q ¼ P Q BW 1=2 f a W ; ð1þ where P Q is the load at failure, B is the specimen thickness, W is the specimen width, a is the crack length and f(a/w) is a geometric factor given by The load-displacement curves recorded for each sample were analyzed according to E399 to determine the failure load (P Q ). Fig. 7 shows the method of determining P Q from the load-displacement curves. Per ASTM E399, a line OP 5 is drawn with slope (P/ v) 5 = 0.95 (P/v) O where (P/v) O is the slope of line OA, as shown in Fig. 7. Line OP 5 represents a 5% deviation in from linearity for the load-displacement curve. If the load at every point preceding P 5 is lower than P 5, then P 5 equals P Q (Type I). However, if the maximum load precedes P 5 then the earlier maximum load is denoted as P Q (Types II and III). The ratio P max /P Q is then calculated. If this ratio exceeds 1.10 then the test is not valid and it is possible that P Q bears no relation to K IC. Fig. 8 shows typical load-displacement curves. The majority of data collected in this study was of Type III with a small portion exhibiting Type I. In all cases the ratio P max /P Q was found to be less than 1.10, making these valid K Q measurements. The evolution of the IMC with increasing reflow time was previously detailed in the paper in Part I, showing an increasing thickness with reflow time. The results for K Q as a function of solder alloy and metallization are shown in Fig. 9b. The relationship between increasing reflow time and decreasing K Q is clearly seen for all cases on Cu metallization as well as with the results of Sn 0.7Cu on Ni Pd. Fig. 9b shows the same relationship between K Q and the intermetallic thickness. a ð2 þ a=wþ 0:886 þ 4:64a=W 13:32a 2 =W 2 þ 14:72a 3 =W 3 5:6a 4 =W 4 f ¼ : ð2þ W ð1 a=wþ 3=2

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 279 Fig. 20. Various tilted views of the smooth fracture with Cu 6 Sn 5 exposed grains for Sn 0.7Cu on Cu with 45 s of reflow. 3.5. Analysis of the fracture surfaces on Ni metallizations Analysis of the fracture surfaces was performed to determine the nature of fracture in the joints. Optical microscopy, SEM, and EDS techniques were employed to identify the fracture surfaces. In all cases, multiple failure locations were studied. Inspection of the fracture surface of Sn 0.7Cu on Ni Au after 45 s of reflow showed failure closely associated with the intermetallics. Fig. 10 shows the profile of the failed sample indicating failure near the Ni 3 Sn 4 layer. SEM images in Fig. 11 show the opposing failure surfaces along with the corresponding EDS spectra. Fracture analysis of both surfaces is necessary to confirm whether fracture occurred within the intermetallic or not. Failure took place predominantly by delamination of the solder-intermetallic interface. The higher Cu content detected on the intermetallic side likely indicates the presence of the (Cu,Ni) 6 Sn 5 phase. The opposing side showed no signs of Ni thereby confirming no significant fracture of the IMC. Ni delamination from the NiP seed layer was also noted, resulting in relatively low values of K Q. This is a result of the poor Ni layer adhesion rather than that of the intermetallic layer itself. Sn 0.7Cu solder on Ni Pd with 45 s of reflow appeared to fail primarily through the bulk solder, as shown in Fig. 12, with a characteristic ductile fracture appearance. With 180 s of reflow and the corresponding thicker intermetallic layer, the failure shifted to the (Ni,- Cu) 3 Sn 4. SEM and EDS analysis indicated the presence of Ni and Cu on the opposing fracture sides shown in Fig. 13. The presence of Ni and Cu on both sides indicates cleavage of the (Ni,Cu) 3 Sn 4 phase. In Fig. 13-I, both the fracture through the mid-level of the intermetallic as well as fracture near the base of the intermetallic can be seen. No Pd was detectable close to the fracture surface. Analysis of the fracture of Sn 4.0Ag 0.5Cu solder on Ni Au metallization with 45 s of reflow showed failure occurring by intermetallic cleavage through the middle of the IMC and near the IMC Fig. 21. Profile view of Sn 0.7Cu on Cu after 90 s of reflow.

280 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 Fig. 22. Opposing fracture surfaces (a) and (b) for Sn 0.7Cu on Cu solder after 90 s of reflow and higher magnification images (c) and (d), respectively. interface with the Ni. Fig. 14 shows opposing fracture surfaces of both types of failures. The intermetallic cleavage shown in Fig. 14 is similar to that found on Sn 0.7Cu solder on Ni Pd after 180 s of reflow. EDS spectrum of the region depicted in Fig. 14a showed very high levels of Ni with some Sn and no significant levels of Cu. This suggests that failure occurred through the Ni 3 Sn 4 phase very near the Ni undermetal. The opposing surface for this failure is depicted in Fig. 14c and has a spectrum indicative of (Ni,- Cu) 3 Sn 4. Analysis of the region depicted in Fig. 14b indicated the presence of the (Ni,Cu) 3 Sn 4 phase while analysis of Fig. 14d also showed evidence of the (Ni,Cu) 3 Sn 4 phase but with a very high level of Sn suggesting that the solder was just beneath the exposed surface. For Sn 4.0Ag 0.5Cu solder on Ni Pd with 45 s of reflow the fracture occurred through the intermetallic as shown in the optical observation in Fig. 15. The fracture in this case was very similar to that of Sn 4.0Ag 0.5Cu on Ni Au with 45 s of reflow and of Sn 0.7Cu solder on Ni Pd after 180 s of reflow. Fig. 16 shows the opposing surfaces, I and II, of fractures through the (Ni,- Cu) 3 Sn 4 as shown in a and c and near the IMC interface with the Ni depicted in b and d. The fracture through the IMC resulted in rough and uneven fractured grain surfaces while the failure surface near the Ni and IMC interface exhibited somewhat smoother and rounded grain fracture. 3.6. Analysis of the fracture surfaces on Cu metallizations The Cu 6 Sn 5 intermetallic present on Cu metallization resulted in a unique fracture behavior. With the softer Sn 0.7Cu solder, lower reflow times resulted in predominantly bulk solder failures. Typical ductile fracture morphology was found as illustrated in Fig. 17 for this case after 45 s of reflow. Interfacial delamination between the solder and Cu 6 Sn 5 grains was found although to a much lesser degree than ductile solder failure. Fig. 18 shows a transition ridge (a and b) where the failure mode shifts from the solder to the interface as well as the opposing surfaces of the interfacial failure (c and d). Close inspection of Fig. 18b shows the detail of the transition ridge. The ductile failure in the solder is giving way to the interfacial failure and exposing the underlying Cu 6 Sn 5 grains. In Fig. 18c regular depressions in the solder can be seen which correspond to exposed Cu 6 Sn 5 grains. Fig. 19 shows higher magnification views of the nodules and corresponding holes on the opposite side of the fracture surface. At this magnification the faceted nature of the grains can be observed. Interestingly the depressions themselves mirror the faceted surface of the grains. This figure also shows smooth and rough surfaces of the solder for these regions with grain-caused depressions. The rough surface appears to be a transition between pure solder failure and the smoother failure closer to the IMC. Fig. 20 shows multiple tilted views of the smooth solder failure near the IMC. After 90 s of reflow the IMC coarsened and the layer thickness increased. With this growth, the prevalence of solder failure diminished and the dominant fracture location became delamination between the IMC and solder. Intermetallic cleavage as the failure mode also increased but was not as common as the delamination. Fig. 21 shows the profile of such a failure. Some Cu 6 Sn 5 grains can be seen intact but most of the grains have fractured and are

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 281 Fig. 23. Opposing fracture surfaces (a) and (b), for Sn 0.7Cu solder on Cu after 180 s of reflow and higher magnification images (c) and (d), respectively. embedded in the opposing fracture surface. Close inspection of the fracture surface itself shows much of the fracture of the Cu 6 Sn 5 grains occurring near their bases and close to the Cu 3 Sn layer. Fig. 22 shows opposing fracture surfaces illustrating this failure mode. In Fig. 22c, some intact grains are visible along with the cleavage surface of neighbor grains. Fig. 22d shows the opposing surface which includes the underside of the fractured grains still adhered to the solder and the depressions in the solder again at neighbor sites. Increasing the reflow time to 180 s produced significantly higher amounts of fracture through the intermetallic by cleavage. There was very little ductile solder failure but still some instances of delamination between the solder and IMC. The dominant intermetallic fracture surfaces were more continuous than was the case with 90 s of reflow. Fig. 23 shows opposing fracture surfaces with very few of the Cu 6 Sn 5 grains remaining intact. Additional, higher magnification micrographs of Fig. 23c and d are shown in Fig. 24. Many grain boundaries can be located on the fractured surfaces. In some cases the hexagonal Cu 6 Sn 5 can be clearly identified. EDS spectra of the fractured surfaces in Fig. 24a and b showed significantly higher Cu peaks suggesting that the EDS sampling depth was possibly including the Cu 3 Sn layer. Fig. 25 shows an EDS spectrum of the region contained with the clearly hexagonal feature of Fig. 24a. It cannot be conclusively stated that failure occurred at the interface of the Cu 6 Sn 5 and Cu 3 Sn phases but the failure was, at least, in very close proximity to the interface and likely within 1 lm. The fracture behavior of Sn 4.0Ag 0.5Cu on Cu metallization resulted in similar types of failure modes as Sn 0.7Cu did after fracture toughness testing. However, after 45 s of reflow there was very little evidence of ductile solder failure. Recall that with Sn 0.7Cu solder with 45 s of reflow the dominant failure mode was through the solder. For the case of Sn 4.0Ag 0.5Cu there were nearly equal amounts of failure through Cu 6 Sn 5 cleavage and delamination between the intermetallic and solder. These two types of failures can be seen in Fig. 26. Intact grains appear to be evenly dispersed amongst the fractured grains. Some solder remained between two larger, neighboring grains as the Cu 6 Sn 5 phase tended to become more elongated and less coarse as was found with Sn 0.7Cu solder. Another feature unique to this solder composition was the presence of the Ag 3 Sn phase. While these plate-like intermetallics were typically found within the bulk of the solder they were occasionally present at an interface. Fig. 27 shows a long plate-like Ag 3 Sn growth adhered to the top of the Cu 6 Sn 5 layer. Its presence at that location suggests that the adhesion of Ag 3 Sn to Cu 6 Sn 5 is greater than that to the solder. Fig. 28 shows opposing fracture surfaces which include a long Ag 3 Sn formation and its related void left in the solder. Again the Ag 3 Sn is seen adhered to the underlying Cu 6 Sn 5 layer. With 90 and 180 s of reflow the Cu 6 Sn 5 grains continued to elongate along with an increase in the continuity of the layer itself. With this progression of the IMC layer came an evident increase in cleavage fracture. Fig. 29 shows the increase in intermetallic fracture with additional reflow time. Fracture of the 45 s sample still

282 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 Fig. 24. Opposing fracture surfaces (a) and (b) as well as c and d for Sn 0.7Cu solder on Cu after 180 s of reflow. shows some Cu 6 Sn 5 grains intact. The sample that was reflowed for 90 s shows fracture entirely through the IMC layer but appears rough and not always through the base of the grains. After 180 s of reflow the fracture surface is smoother and can be found consistently closer to the base of the grains. Fig. 30 shows the detail in opposing fracture surfaces found after 90 s of reflow. Fractures close to the base of the grains as well as fracture through the middle of the grains can be seen. Also evident on these surfaces is the hexagonal nature of the Cu 6 Sn 5 with many 120 angles visible. Fig. 31 shows the fracture surfaces associated with Sn 4.0Ag 0.5Cu on Cu metallization after 180 s of reflow. In this case the fracture occurs almost entirely through the base of the Cu 6 Sn 5 intermetallic. The fracture surface is somewhat smoother than that after 90 s of reflow possibly related to the increase in the continuity of the layer near the interface with the Cu 3 Sn phase. Still evident is the hexagonal nature of the Cu 6 Sn 5 phase. 3.7. Relationship between fracture toughness and failure modes Fig. 25. EDS spectrum of (a) IMC cleavage found with Sn 0.7Cu solder on Cu metallization after 180 s of reflow along with a (b) reference spectrum of a the Cu 6 Sn 5 phase obtained under the same conditions. Relating the fracture toughness values to the observed fracture surfaces required quantification of the type and degree of failure modes, which were categorized as: (a) ductile failure in the bulk solder, (b) interfacial delamination at the solder/intermetallic interface and, (c) intermetallic cleavage. Upon studying the fracture surfaces all samples were found to contain a mixture of the failure modes rather than one distinct failure mode. Observations of each sample were made with an optical microscope to estimate

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 283 Fig. 26. Opposing fracture surfaces, (a) and (b), for Sn 4.0Ag 0.5Cu solder on Cu after 45 s of reflow and higher magnification images (c) and (d), respectively. Fig. 27. Cross section view of the Ag 3 Sn intermetallic adhered to the Cu 6 Sn 5 layer. the portion of each failure mode within the sample. The regions observed by optical microscope and assigned a failure mode were verified by SEM analysis. Fig. 32 shows the results of the fracture mode mixtures for each of the sample types. To properly interpret the relative fracture toughness values (K Q ) given in Fig. 3, the related failure modes must be considered. Beginning with the results on Ni metallization we see that the average fracture toughness value for Sn 0.7Cu solder on Ni Pd after 45 s of reflow was 9.13 MPa m 1/2 while after 180 s of reflow the relative toughness decreased approximately 10% 8.19 MPa m 1/2. Recall that after the additional reflow for this sample cell the intermetallic layer became thicker and more uniform. We also see from Fig. 32 that, with increasing reflow time, the main failure mode changed from ductile solder failure to cleavage of the Ni 3 Sn 4 layer. Although there is overlap of the standard deviation of the K Q values the downward trend is evident. When considering the shift in dominant failure mode, the decrease in fracture toughness values can be attributed to the increase thickness of the intermetallic layer which resulted in diminished strength of the solder joint. For Sn 4.0Ag 0.5Cu on Ni Pd reflowed for 45 s, the dominance of intermetallic cleavage is even greater. Although the (Ni,Cu) 3 Sn 4 thicknesses were very similar for both solders on Ni Pd after just 45 s of reflow, the fracture surfaces were quite different as were the toughness values. The lower ductility of the Sn 4.0Ag 0.5Cu solder, relative to Sn 0.7Cu, likely resulted in the shift from mostly ductile to mostly IMC cleavage failures. Thus, in addition to IMC thickness, the ductility of the solder also plays a role in the toughness of the joint. The decreased solder ductility may have transferred additional strain to the intermetallic resulting in the significantly lower average fracture toughness of 4.70 MPa m 1/2. Comparing the results of Sn 4.0Ag 0.5Cu on Ni Pd to those on Ni Au both after 45 s of reflow does not show a significant change in the toughness values. Although the average toughness value for this solder on Ni Au is over 20% higher than that on Ni Pd, the large standard deviation shows that there are no significant differences between these two cells. Studying the results for the cells with Cu metallization also yielded interesting trends. Recall that for Sn 0.7Cu solder on Cu metallization after 45 s of reflow the failure mode was predominantly ductile solder failure. Tracing the level of solder failure through 90 s and then 180 s of reflow for this solder and undermet-

284 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 Fig. 28. Opposing fracture surfaces of an (a) Ag 3 Sn formation and (b) the related void left in the solder for Sn 4.0Ag 0.5Cu solder on Cu metallization after 45 s of reflow. Fig. 29. Cross section views of fractures for Sn 4.0Ag 0.5Cu solder on Cu metallization after (a) 45, (b) 90 and (c) 180 s of reflow. al shows a steady decrease. At the same time there is an evident increase in the amount of cleavage of the Cu 6 Sn 5 layer. The increase in IMC cleavage was likely brought on by the progressive growth of the Cu 6 Sn5 layer. In the microstructure characterization section, we see the most significant change in the Cu 6 Sn 5 layer between 90 and 180 s of reflow. Interestingly we also see the most significant change in the amount of intermetallic cleavage between 90 and 180 s of reflow. Furthermore, the average relative fracture toughness decreased more than 22%, from 9.23 to 7.55 MPa m 1/2, respectively. These trends appear to be related and it can be concluded that the toughness of the Cu 6 Sn 5 layer is clearly lower than that of the Sn 0.7Cu solder. The same trends are easily observed for Sn 4.0Ag 0.5Cu solder on Cu metallization. In this case, ductile failure in the solder was less prevalent than with the Sn 0.7Cu solder. After 90 s of reflow the amount of ductile failure for Sn 4.0Ag 0.5Cu decreased significantly to a very low level similar to Sn 0.7Cu solder after 180 s of reflow. The effect of this drop in ductile failure between 45 and 90 s of reflow for the Sn 4.0Ag 0.5Cu solder on Cu can be seen in the average relative fracture toughness values. A significant drop from 8.97 to 6.81 MPa m 1/2, respectively, is noted. This drop is also similar to what was found for Sn 0.7Cu solder on Cu between 90 and 180 s of reflow. For the Sn 4.0Ag 0.5Cu solder between 90 and 180 s of reflow, the amount of cleavage in the Cu 6 Sn 5 layer increases to the highest level found in any of the cells investigated. Consequently, the relative fracture toughness also decreased between 90 to 180 s of reflow although only slightly. It is again clear that the decrease in relative fracture toughness is related to the increasing prevalence of solder joint failure though intermetallic cleavage. While the results of Sn 0.7Cu and Sn 4.0Ag 0.5Cu solders on Cu undermetal after 180 s

S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 285 Fig. 30. Opposing fracture surfaces of Sn 4.0Ag 0.5Cu solder on Cu metallization after 90 s of reflow. Fig. 31. Opposing fracture surfaces of Sn 4.0Ag 0.5Cu solder on Cu metallization after 180 s of reflow. Failure Mode Fractions 1.00 0.90 0.80 0.70 0.60 0.50 0.40 0.30 0.20 0.10 0.00 SnCu-NiPd-45 SnCu-NiPd180 Failure Modes v. Metallurgy IMC Cleavage IMC/Solder Delamination Bulk Solder SnCu-Cu-45 SnCu-Cu90 SnCu-Cu180 SnAgCu-NiAu-45 SnAgCu-NiPd-45 SnAgCu-Cu-45 Solder and Undermetallurgy Type SnAgCu-Cu90 SnAgCu-Cu180 Fig. 32. Estimated failure mode fractions for compact tensile fracture toughness samples.

286 S.M. Hayes et al. / Microelectronics Reliability 49 (2009) 269 287 of reflow both show the dominant failure mode to occur through the Cu 6 Sn 5, their average relative fracture toughness values are somewhat different. The values for Sn 0.7Cu were approximately 16% greater than found with Sn 4.0Ag 0.5Cu. This difference could be the result of the occasional presence of the Ag 3 Sn found at the interface for the samples with the Ag-bearing solder. The more directional growth observed with the Sn 4.0Ag 0.5Cu solder resulted in taller Cu 6 Sn 5 grains with a greater aspect ratio. This may have resulted in grains less capable of resisting the crack propagation. The lower ductility of the Sn 4.0Ag 0.5Cu solder and the additional transfer of strain to the IMC could also contribute to the difference in relative fracture toughness. Figs. 33 and 34 show the failure mode fraction as a function of intermetallic thicknesses. In Fig. 33, for Sn 0.7Cu solder on Cu metallization the relationship between the estimated failure mode fractions and the intermetallic thickness appears to be linear over the range of thicknesses investigated. Clearly an increase in Cu 6 Sn 5 thickness results in an increase in fracture through the IMC and a corresponding decrease in bulk solder failure. The intersection of the IMC fracture and bulk solder failure lines shows an apparent critical thickness of approximately 6 lm. Above this critical thickness the joint is likely to fail through the IMC rather than the bulk solder. A similar trend is noted for Sn 4.0Ag 0.5Cu solder on Cu metallization in Fig. 34, although the relationship is not linear. Recall the spallation effect that was previously noted for Sn 4.0Ag 0.5Cu on Cu that occurred when the Cu 6 Sn 5 reached approximately 5 lm in thickness. Despite the IMC not significantly increasing in thickness, fracturing within the IMC increased as spallation began. The spalling of the IMC can have a localized strengthening effect near the attached IMC layer. The spalled IMC may act as particle reinforcement which in turn will decrease the ductility of the region near the attached Cu 6 Sn 5 layer resulting in increasing failure within the intermetallic layer. Fig. 34 shows the apparent critical intermetallic thickness Sn 4.0Ag 0.5Cu on Cu to be 4 5 lm. The difference in critical thickness for the two solder types can be explained by ductility. The Sn 4.0Ag 0.5Cu solder is less ductile than the Sn 0.7Cu solder and therefore makes the joint less tolerant of increasing IMC thickness. Fraction of Failure 1 0.8 0.6 0.4 0.2 Fraction of Failure Through IMC Fraction of Failure in Solder 0 0 2 4 6 8 10 Intermetallic Thickness (μm) Fig. 33. Estimated failure mode fractions for Sn 0.7Cu solder on Cu undermetal and its dependence on intermetallic thickness. Fraction of Failure 1 0.8 0.6 0.4 0.2 0 0 2 4 6 8 10 Intermetallic Thickness (μm) Similar trends are noted for Ni metallization as were mentioned for Cu metallizations. Increasing Ni 3 Sn 4 thickness tends to cause more failure within the IMC and decreasing failure within the bulk of the solder. Critical thickness of the Ni 3 Sn 4 appears to be approximately 1.5 lm for Sn 4.0Ag 0.5Cu and 2.5 lm for Sn 0.7Cu. Greater ductility associated with the Sn 0.7Cu appears to make the joint more tolerant of the increasing IMC thickness. 4. Conclusions Microstructural characterization of Pb-free solders and interfacial intermetallic formation was quantified. Solder joints were fabricated using Sn 4.0Ag 0.5Cu and Sn 0.7Cu solders on Cu, Ni Au, and Ni Pd. (Ni,Cu) 3 Sn 4 was present at the Ni interface for both solders but was coarser in the case of Ni Pd. Cu 6 Sn 5 and Cu 3 Sn were observed for both solder types. The Cu 3 Sn layer was similar in thickness and appearance for both solders, but the Cu 6 Sn 5 was smoother and rounder in the case of Sn 0.7Cu. Additional time above liquidus resulted in growth of the Cu 6 Sn 5 layer and eventual spalling of the IMC grains. The fracture toughness behavior of Pb-free solder joints on Ni and Cu metalizations was quantified and fractographic analysis of the failure regions was performed. The morphology and thickness of the IMC was shown to affect the relative fracture toughness of the solder, K Q, as well as the failure mode. For the solders investigated, the critical thickness of Ni 3 Sn 4 that results in IMC failure instead of bulk solder failure is approximately 1.5 2.5 lm. The critical thickness for the same effect for Cu 6 Sn 5 is approximately 4 6 lm. Observation of the fracture toughness behavior showed that K Q can decrease by as much as 22% when failures shift from the bulk solder to the IMC. For a given thickness the fracture toughness of Cu 6 Sn 5 appears to be greater than that of Ni 3 Sn 4. However, Cu 6 Sn 5 experiences more rapid growth with extended exposure to reflow resulting in significant decreases in toughness. Acknowledgements Failure Fraction in IMC Failure Fraction in Solder Fig. 34. Estimated failure mode fractions for Sn 4.0Ag 0.5Cu solder on Cu undermetal and its dependence on intermetallic thickness. Research support for this study was provided by Freescale Semiconductor, and is gratefully acknowledged. The assistance of