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Cu single shear lap solder joints
Materials Science and Engineering A 444 (2007) 75–83
Shear strength and interfacial microstructure of Sn–Ag–xNi/Cu single shear lap solder joints Yang-Hsien Lee a,1 , Hwa-Teng Lee b,∗ a
Department of Mechanical Engineering, Southern Taiwan University of Technology, Tainan 701, Taiwan b Department of Mechanical Engineering, National Cheng Kung University, Tainan 701, Taiwan Received 12 May 2006; accepted 11 August 2006
Abstract This study investigates composite lead-free solders fabricated by adding between 0.5 and 3 wt% of Ni particles in situ to Sn–3.5 wt%Ag lead-free solder. The single lap shear strength, fracture behavior and microstructural evolution characteristics of the as-reflowed specimens are examined and compared with those of specimens thermally aged at 150 ◦ C for various aging times. In general, it is found that the single lap shear strength of the joints increases with increasing Ni addition in the as-reflowed condition, but decreases with increasing storage time in the aged specimens. For Ni additions of 0.5 and 1 wt%, the specimens fracture in the solder near the intermetallic compound (IMC) layer/solder interface, which suggests that the solder matrix has a lower strength than the IMC layer. The presence of elongated dimple-like structures on the fracture surfaces of these specimens is indicative of a ductile failure mode. For Ni additions of more than 1 wt%, the specimens fracture with brittle characteristics at the solder/IMC interface, which indicates that an increased Ni addition increases the strength of the solder matrix beyond that of the interfacial layer. © 2006 Elsevier B.V. All rights reserved. Keywords: Shear strength; Microstructure; Single shear lap joint; Intermetallic compound (IMC)
1. Introduction Although traditional Sn–Pb solders have many advantageous properties, the toxicity of their Pb content presents a major health hazard. The need to protect worker safety and public health has led to a worldwide drive to develop new lead-free solder alloys for the packaging of common electronic assemblies [1]. The literature contains many studies of lead-free alloys designed for soldering applications. Of the various potential lead-free solders which have been proposed, eutectic Sn–3.5Ag and Sn–Ag–Cu solder has emerged as the leading candidate to replace Sn–Pb solder since it has no lead content, superior mechanical properties and a good wetting behavior on copper [2–5]. Solder joints provide both mechanical strength and electrical conductivity, and consequently, play an important role in the connection of electronic components to printed circuit boards. Fundamentally, a solder material should wet the substrate and
∗
Corresponding author. Tel.: +886 6 2757575x62154; fax: +886 6 2745698. E-mail addresses: [email protected] (Y.-H. Lee), [email protected] (H.-T. Lee). 1 Tel.: 886 6 2533131x3529; fax: 886 6 2537912. 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.08.065
provide good adhesion. Furthermore, solder joints must be capable of withstanding the high thermal and mechanical loadings which frequently occur during handling and system use at higher temperatures [6]. In solder joints between a Sn-based solder and a copper substrate, intermetallic compounds (IMCs) such as Cu6 Sn5 and Cu3 Sn form and grow both during the soldering operation itself and during the subsequent system use [7–10]. The IMC layer provides a metallurgical bonding between the solder and the substrate. Furthermore, the morphology of the IMC layer also has a fundamental effect on the solder joint reliability because crack initiation and propagation during thermal fatigue generally takes place at the solder/IMC interface [11]. Hence, a thick IMC layer containing coarse intermetallic compounds may adversely affect the mechanical properties and performance of the solder joint. It is often reasoned that the mechanical integrity of the solder joint is degraded by the inherent brittleness of the IMC layer since cracking can readily occur when the joint is mechanically or thermally stressed [12,13]. Therefore, the formation, growth, and fracture of the IMC layer and the microstructure of the bulk solder have a crucial influence on the solder joint reliability. The mechanical properties of the solder are important because they determine the fracture and thermal fatigue behavior
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Fig. 1. Single lap shear solder joint test specimen.
of the solder joint. Therefore, it is important that the mechanical properties of the solder, such as its shear strength, adhesive strength and ductility, during the soldering operation itself and during the subsequent system use are well understood [14–17]. Attempts have been made to modify the solder microstructure and to improve its mechanical properties by adding small quantities of alloying elements to Sn-based solders [3,4]. Producing small, constituent-phase IMC particles in the bulk solder alloy is often a key design objective when developing alloys designed to enhance the mechanical strength and deformation behavior of solder joints. Some researchers have investigated the use of alloying elements such as copper, nickel, indium, antimony, bismuth, and so on, as a means of reducing the melting point of eutectic Sn–3.5Ag solder while simultaneously improving its mechanical properties [18–26]. The results of these studies have confirmed that certain composite solders exhibit the nec-
essary combination of enhanced strength and other favorable mechanical properties required by the electronics industry. Nickel is commonly used to provide a diffusion barrier between Cu components and Sn or Sn-based solder alloys in order to prevent, or at least suppress, the formation of Cu3 Sn and Cu6 Sn5 intermetallic compounds. Ni is an effective additive since the stable phases, which it forms in the Ni–Sn binary system, grow more slowly than the Cu intermetallic compounds. Furthermore, Ni has good wetting characteristics with Sn. However, the choice of Ni as an alloying element is primarily due to its formation of additional intermetallic phases, which improve the mechanical properties of the solder. Accordingly, the present study investigates the effects of adding varying quantities of nickel particles to eutectic Sn–3.5Ag solder to form composite solders. The formation and growth of interfacial IMC compounds between the Nicontaining composite solders and a Cu substrate are studied. Additionally, the effects of the level of Ni addition and the thermal aging time on the single lap shear strength of the Sn–Ag–x wt%Ni composite solder/Cu substrate joints are investigated. 2. Experimental procedures The present composite solders were prepared by mixing 0.5, 1, 2, and 3 wt% of Ni particles (with an average particle size of 2–3 m) into a molten Sn–3.5 wt%Ag eutectic solder. Prior to mixing, a thin film of Rosin Mildly Activated (RMA) flux (Taiyo electric BS-10 flux, Japan) was coated on the surface of the Ni particles to remove the oxides, which naturally occur. The
Fig. 2. Optical micrographs of interfacial microstructure of: (a) 0.5Ni, (b) 1Ni, (c) 2Ni, and (d) 3Ni solder joints in as-reflowed condition.
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molten solder was maintained at a temperature of 300 ◦ C and stirred with a glass rod in an inert atmosphere for approximately 5 min to ensure a uniform particle distribution [3]. The solder was then poured into a metal box (15 mm × 10 mm × 10 mm) and allowed to cool in air. In this study, a small single shear lap solder joint is designed which better mimics the actual solder joints used in the automobile and microelectronics industries. The various composite solders were rolled into sheets with a thickness of approximately 200 m and were then punched to produce discs with a diameter of 1.5 mm. To ensure proper alignment of the solder joint, a fixture was constructed for sample preparation. The specimens were put into a furnace and hold 260 ◦ C for 60 s, and then removed from the furnace and cooled in air to room temperature. The heating rate was fixed at 1.5 ◦ C/s and the measured cooling rate between 260 and 150 ◦ C was approximately 0.85 ◦ C/s.
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Single shear lap joints with an area of approximately 1 mm2 were then fabricated by using the composite solders to join two copper substrates, as shown in Fig. 1. Following the soldering process, some of the specimens were retained in an as-reflowed condition, while the others were thermally aged at a temperature of 150 ◦ C for 100, 200, 500, or 1000 h, respectively. To obtain a top view of the intermetallic compounds, the majority of the solder on the specimens was ground away, and the specimens were then etched with a 20% HCl–80% C2 H5 OH solution for 30 min to selectively dissolve the remaining solder. A uniaxial micro-force test system (SHIMADZU AG-1, Japan) was used to carry out shear tests on the various solder joint specimens. The tests were performed at room temperature under a strain rate of 1 × 10−2 s−1 . To permit metallographic observations, the fractured samples were mounted in epoxy and then cross-sectioned in a direction perpendicular to the solder–copper
Fig. 3. Optical micrographs of interfacial microstructure of: (a) Sn–Ag, (b) 0.5Ni, (c) 1Ni, (d) 2Ni, and (e) 3Ni solder joints following thermal aging at 150 ◦ C for 1000 h.
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interface. The specimens were then successively ground down to 4000 grit on a silicon carbide paper under water cooling. Finally, the specimens were polished using 3 and 1 m diamond paste and 0.02 m aluminum oxide, and then etched with a 5% HCl solution. The microstructure and composition of the solder joints, and their fracture surfaces following the shear test, were analyzed using optical microscopy (OM), scanning electron microscopy (SEM), and energy dispersive X-ray spectrometer (EDS) techniques. 3. Results and discussion 3.1. Microstructure of composite solder joint Fig. 2 presents optical micrographs of the as-reflowed Sn–3.5Ag–x%Ni composite solder joints. It is observed that the microstructure of the solder consists of globular Sn-rich phase matrix with dispersed Ag3 Sn precipitates around the Sn-rich regions. A small number of Ni–Cu–Sn intermetallic particles are also observed in the solder matrix. The intermetallic compounds within the solder matrix are identified by EDS compositional analysis as (Ni,Cu)3 Sn4 . Comparing the various micrographs in Fig. 2, it is clear that the mean free path between the -Sn phases is reduced with increasing Ni content. Furthermore, a Ni–Cu–Sn intermetallic layer is evident at the interface of the solder and the Cu substrate. EDS analysis reveals that this compound contains Cu, Ni, and Sn and has an atomic compositional ratio close to that of Cu6 Sn5 intermetallic phase. These compounds are found from EDS analysis to be (Cu,Ni)6 Sn5 . This finding is consistent with the results reported by previous investigators [21–25]. Following prolonged aging at an elevated temperature, Sn in the area of the (Cu,Ni)6 Sn5 IMC closest to the copper substrate diffuses into the Cu side and prompts the formation of Cu3 Sn [11]. The optical micrographs presented in Fig. 3 reveal that after aging at 150 ◦ C for 1000 h, the (Cu,Ni)6 Sn5 intermetallic layer in the current specimens becomes noticeably thicker and a Cu3 Sn IMC layer is formed between the (Cu,Ni)6 Sn5 intermetallic layer and the Cu substrate. Fig. 4 shows a typical top view of the intermetallic compounds of an Sn–Ag–3 wt%Ni/Cu joint in the as-reflowed condition. The IMC contains two elongated rod-type Ni–Sn–Cu ternary IMCs. EDS elemental analysis reveals that the longer IMCs are (Cu,Ni)6 Sn5 while the shorter IMCs are Cu6 Sn5 . The nano-sized particles observed on the surface of the longer rodtype IMCs are found to be Ag3 Sn compounds. Fig. 5 shows the Sn–Ag–3 wt%Ni/Cu joint aged at 150 ◦ C for 200 and for 500 h the morphology of large Ag3 Sn IMC platelets located at the (Cu,Ni)6 Sn5 IMC layer. It is observed that the platelets partially penetrate the IMC layer. Since the interfacial layer is formed by the reaction of the molten solder and the substrate during the reflow treatment, it can be inferred that the large Ag3 Sn nucleate initially from the interfacial layer. Since Ag has a higher concentration compared to Cu; Ag3 Sn can nucleate and grow rapidly to a high volume. Large Ag3 Sn intermetallics were also observed in the bulk solder near the solder/Cu interface. These Ag3 Sn decreased in numbers but
Fig. 4. Top view morphology of long rod-like (Cu,Ni)6 Sn5 grains and short rodlike Cu6 Sn5 particles growing from interfacial layer of Sn–Ag–3Ni/Cu joint in the as-reflowed condition.
increased in size due to coarsening as the solder joints were aged. In the soldering process, the formation of intermetallic compounds between the metal substrate and the solder must be
Fig. 5. Morphologies of larger platelet Ag3 Sn and Cu–Ni–Sn IMC layer on Cu substrate. (a) Sn–Ag–3Ni composite solder/Cu substrate following aging at 150 ◦ C for 200 h and (b) 500 h.
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carefully controlled since the brittleness of the IMC layer may degrade the reliability of the solder joint [7,9,11,12,14]. The morphological change and increased thickness of the IMC layer observed in the aged specimens indicates that the interfacial reaction of the solder with the Cu substrate continues during the thermal aging process. The interface reactions between Sn–Ag solders and Cu can be categorized into two types. Cu6 Sn5 and Cu3 Sn are two of the main reaction products. In all cases Cu6 Sn5 is the first compound formed, and very little if any Cu3 Sn is present immediately after solidification. During aging treatment Cu3 Sn will form by a solid state reaction between Cu and Cu6 Sn5 , and after prolonged aging the Cu3 Sn layer becomes more evident. It has been reported [11,16] that the atomic diffusion of Cu and Sn through the interface IMC layer was the main controlling factor for the growth of the interface IMC during aging. The growth kinetic of interface IMC layer follows the square root time law expressed in the equation below: X = X0 + Dt1/2 . The Cu3 Sn layer facing a Cu substrate is formed thinner than the Cu6 Sn5 layer. Fig. 6 plots the variation of the average Cu6 Sn5 and Cu3 Sn IMC layer thickness with the square root of aging time at 150 ◦ C.
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In general, it can be seen that the thickness of both layers increases with an increasing Ni content. Guo et al. [25] have indicated significant IMC layers growth was observed in Ni composite solder joint under isothermal aging at 150 ◦ C. The addition of more than 1 wt% Ni appears to have a particularly profound influence on the thickness of the (Cu,Ni)6 Sn5 reaction layer. The IMC layer thickness of the Sn–Ag–3Ni joint specimen is approximately 1.5 times that of the Sn–Ag joint specimen in both the as-reflowed condition (i.e. Sn–Ag: 6.1 m, Sn–Ag–3Ni: 9.4 m) and following thermal storage for 1000 h (i.e. Sn–Ag: 13.6 m, Sn–Ag–3Ni: 19.7 m). Fig. 6 also shows that the thickness of the IMC layer increases with an increasing aging time. The thickness of the Cu6 Sn5 and (Cu,Ni)6 Sn5 intermetallic layer is greater than that of the Cu3 Sn layer in all cases. Significantly, the Cu3 Sn layer in the Sn–Ag solder joint is consistently thicker than that in any of the Sn–Ag–xNi solder joints following thermal aging. The formation of Cu3 Sn is affected by the phase stability of Cu6 Sn5 according to the reaction Cu6 Sn5 + Cu → Cu3 Sn. The thermodynamic affinity in Cu6 Sn5 is stronger between Ni and Sn than between Cu and Sn. Thus, the phase stability is greater for (Cu,Ni)6 Sn5 than for Cu6 Sn5 [26]. This finding suggests that
Fig. 6. Variation of interfacial intermetallic layer thickness with Ni content and aging period for storage at 150 ◦ C.
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Fig. 7. Variation of single lap shear strength with Ni addition as function of aging time for storage at 150 ◦ C.
the addition of Ni to the Sn–Ag solder effectively suppresses the growth of Cu3 Sn compounds. 3.2. Single shear lap strength Since solder joints are frequently subjected to mechanical loading during their service lives, their mechanical properties, e.g. their fatigue and shear strengths and their creep resistance, have a crucial effect on the solder joint reliability and hence
Fig. 8. The trend of tension strength and shear strength with Ni content.
the integrity of the overall electronic package. Therefore, a series of shear tests at room temperature under a strain rate of 1 × 10−2 s−1 evaluated the effect of the interfacial reactions on the strength of solder joints aged for 0, 100, 500 and 1000 h at 150 ◦ C. Fig. 7 presents the shear strength results obtained for the Sn–Ag, Sn–Ag–0.5Ni, Sn–Ag–1Ni, Sn–Ag–2Ni and Sn–Ag–3Ni solder alloys. In general, the as-reflowed joints have greater shear strength than the aged specimens. Fur-
Fig. 9. SEM fracture surface of Sn–Ag–1Ni/Cu joint specimens stored at 150 ◦ C for different aging times: (a) as-reflowed, (b) 100 h, and (c) 1000 h; (d) OM cross-section view of specimen aged for 1000 h.
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thermore, the shear strength of the as-reflowed solder joint improves with increasing Ni addition. The shear strength enhancement is most likely attributed to the additional formation of (Ni,Cu)3 Sn4 particles and a refinement of the solder microstructure as the level of Ni addition increases [3]. Fig. 7 demonstrates that thermal aging leads to a reduction in the shear strength of the solder specimens. Thermal aging has two important effects on the solder joint microstructure; it increases the IMC layer thickness and it leads to a coarsening of the solder microstructure [16,25–30]. Therefore, the reduction in shear strength observed in the current aged specimens can reasonably be attributed to an increased thickness of
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the brittle intermetallic layer, the agglomeration of (Ni,Cu)3 Sn4 compounds, or a coarsening of the Ag3 Sn particles. Fig. 7 shows that the shear strength drops most dramatically during the first 100 h of thermal aging. Comparing the tension strength and shear strength, Fig. 8 presents the tensile strength [4] increases with increasing Ni content, and variation of shear strength with Ni content also reveals a similar trend. In a microstructural sense, the reduction in strength after 100 h can be attributed to a significant coarsening of the microstructure of the joint. Fig. 3 (corresponding to 1000 h thermal storage) provides clear evidence of the degradation in the microstructure during aging, i.e. the images show a significant (Cu,Ni)6 Sn5 IMC layer growth at the interface, coarse (Ni,Cu)3 Sn4 , Ag3 Sn, and (Cu,Ni)6 Sn5
Fig. 10. (a) SEM fracture surface of Sn–Ag–2Ni/Cu joint (near Cu substrate); (b) OM cross-section view of Sn–Ag–2Ni/Cu joint specimens aged at 150 ◦ C for 1000 h, (c) high-magnification image of region A; (d) SEM fracture surface of Sn–Ag–3Ni/Cu joint; (e) high-magnification image of region B and (f) OM cross-section view of Sn–Ag–3Ni/Cu joint specimens aged at 150 ◦ C for 1000 h.
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phases in the bulk of the solder, and voids in the interfacial IMC layer. Figs. 6 and 7 show that the Sn–Ag–3%Ni solder joint has higher shear strength than any of the other solder joints following thermal aging for 1000 h despite having the thickest IMC layer. Therefore, it appears that the thickness of the IMC layer does not play a critical role in controlling the shear strength of single lap solder joints. Note that a similar finding was also reported in the studies of Deng et al. [16] and Yoon et al. [28], in which it was shown that the shear strength of solder joints is controlled primarily by the mechanical properties of the solder rather than the thickness of the intermetallic compound layer. 3.3. Fracture surface analysis The fracture surfaces of the solder/Cu joints provide a valuable insight into the deformation mechanisms, which lead to their failure. Therefore, the failed as-reflowed and aged joints were cross-sectioned in a perpendicular direction to the solder/Cu interface and their fracture surfaces inspected using SEM and OM techniques. Fig. 9 presents SEM images of the fracture surfaces of Sn–3.5Ag–1Ni/Cu joint specimens in the as-reflowed condition and following aging at 150 ◦ C for 100 and 1000 h, respectively. Fig. 9(d) presents an OM image of the fracture surface cross-section in an Sn–3.5Ag–1Ni/Cu joint aged for 1000 h. In all cases, it is observed that the specimens fracture within the Sn–3.5Ag–1Ni solder matrix. This suggests that the Sn–3.5Ag–1Ni solder matrix is weaker than the IMC layer between the solder and the Cu substrate. Furthermore, the fracture surfaces in Fig. 9(a–c) contain elongated dimplelike structures. Hence, it is apparent that the Sn–3.5Ag–1Ni/Cu joints all fail with a ductile fracture mode, irrespective of their thermal aging conditions. The trend in Fig. 9 is also representative for the cases of Sn–3.5Ag/Cu and Sn–3.5Ag–0.5Ni/Cu joint specimens fracture surfaces. Fig. 10 shows the fracture surfaces of the Sn–Ag–2Ni/Cu and Sn–Ag–3Ni/Cu joints aging at 150 ◦ C for 1000 h conditions. It can be seen that these specimens have a brittle fracture morphology rather than the ductile morphology observed in the Sn–3Ag and low Ni-containing composite solder specimens. Due to the combined effect of the brittle (Cu,Ni)6 Sn5 IMC layer and the stress concentration at the solder/intermetallic interface, (Cu,Ni)6 Sn5 nodular tips protrude into the fracture surface. Under shear conditions, these nodular tips cause a large shear strain localization effect, and hence the IMC interface becomes the preferred site for crack formation and propagation. Therefore, as shown in Fig. 10(b and f), partial debonding failures occur at the interface between the solder matrix and the (Cu,Ni)6 Sn5 IMC layer. The faceted appearance of the fracture surfaces in Fig. 10(a and d) are caused by cleavage fractures of the (Cu,Ni)6 Sn5 layer. When the cleavage planes in neighboring (Cu,Ni)6 Sn5 grains are approximately parallel, the fracture propagates easily through the inside of the IMC layer or along the interface of the (Cu,Ni)6 Sn5 layer and the solder matrix. Fig. 10(b and f) shows the debonding phenomenon in the Sn–Ag–2%Ni and
Sn–Ag–3%Ni joints at the interface between the solder matrix and the (Cu,Ni)6 Sn5 IMC layer. With 2 and 3% Ni addition, the solder matrix has a high (Ni,Cu)3 Sn4 and Ag3 Sn intermetallic phase content and hence has an enhanced strength. Therefore, the interface between the solder matrix and the IMC layer in the Sn–Ag–2%Ni and Sn–Ag–3%Ni joints is relatively weaker than that in the Sn–Ag and low Ni-containing composite solder joints. Consequently, as shown in Fig. 10(b and f), the shear fracture appears to follow a mixed path with an initial crack at the (Cu,Ni)6 Sn5 /solder interface region followed by propagation within the IMC layer or along the near interface ((Cu,Ni)6 Sn5 /solder) region. From the results presented in Figs. 9 and 10, it appears that the probability of debonding shear failure depends on the relative balance between the strength of the solder matrix and the strength of the (Cu,Ni)6 Sn5 /solder interface. In general, cracking occurs at the weakest point in the joint microstructure, i.e. in the solder matrix in the Sn–3.5Ag and low Ni-containing composite solder specimens and at the interface of the solder matrix and the IMC layer in the higher Ni-content specimens. 4. Conclusion This study has examined the microstructural evolution of the IMC layer formed between Sn–Ag–x%Ni composite solders and a Cu substrate and has investigated the single lap shear strength of specimens treated under various thermal aging conditions. The following conclusions can be drawn: 1. In the as-reflowed condition, the Sn–Ag–x%Ni composite solders contain only a (Cu,Ni)6 Sn5 layer at the interface between the solder matrix and the Cu substrate. However, when the joints are thermally aged, a Cu3 Sn layer is formed between the (Cu,Ni)6 Sn5 layer and the Cu substrate. In general, the thickness of the IMC layer is found to increase with an increasing Ni content and an increasing thermal storage time. 2. The shear strength of the single lap joints increases with increasing Ni addition in the as-reflowed specimens. However, the shear strength decreases as the joints are thermally aged regardless of their Ni content. The reduction in shear strength is caused primarily by a coarsening of the solder microstructure. The effect of the increased IMC layer thickness following thermal storage does not appear to have a significant effect of the shear strength of the joint. 3. For Ni additions of 0.5 and 1 wt%, the specimens fracture in the solder matrix near the IMC layer/solder interface. The fracture surfaces are characterized by elongated dimple-like structures, which suggest a ductile failure mode. However, when the level of Ni addition exceeds 1 wt%, the failure mode changes from ductile fracture to brittle fracture, and the specimens fail at the interface of the solder matrix and the IMC layer. Therefore, it appears that higher levels of Ni addition increase the strength of the solder matrix beyond that of the interface.
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Shear strength and interfacial microstructure of Sn–Ag–xNi/Cu single shear lap solder joints Yang-Hsien Lee a,1 , Hwa-Teng Lee b,∗ a
Department of Mechanical Engineering, Southern Taiwan University of Technology, Tainan 701, Taiwan b Department of Mechanical Engineering, National Cheng Kung University, Tainan 701, Taiwan Received 12 May 2006; accepted 11 August 2006
Abstract This study investigates composite lead-free solders fabricated by adding between 0.5 and 3 wt% of Ni particles in situ to Sn–3.5 wt%Ag lead-free solder. The single lap shear strength, fracture behavior and microstructural evolution characteristics of the as-reflowed specimens are examined and compared with those of specimens thermally aged at 150 ◦ C for various aging times. In general, it is found that the single lap shear strength of the joints increases with increasing Ni addition in the as-reflowed condition, but decreases with increasing storage time in the aged specimens. For Ni additions of 0.5 and 1 wt%, the specimens fracture in the solder near the intermetallic compound (IMC) layer/solder interface, which suggests that the solder matrix has a lower strength than the IMC layer. The presence of elongated dimple-like structures on the fracture surfaces of these specimens is indicative of a ductile failure mode. For Ni additions of more than 1 wt%, the specimens fracture with brittle characteristics at the solder/IMC interface, which indicates that an increased Ni addition increases the strength of the solder matrix beyond that of the interfacial layer. © 2006 Elsevier B.V. All rights reserved. Keywords: Shear strength; Microstructure; Single shear lap joint; Intermetallic compound (IMC)
1. Introduction Although traditional Sn–Pb solders have many advantageous properties, the toxicity of their Pb content presents a major health hazard. The need to protect worker safety and public health has led to a worldwide drive to develop new lead-free solder alloys for the packaging of common electronic assemblies [1]. The literature contains many studies of lead-free alloys designed for soldering applications. Of the various potential lead-free solders which have been proposed, eutectic Sn–3.5Ag and Sn–Ag–Cu solder has emerged as the leading candidate to replace Sn–Pb solder since it has no lead content, superior mechanical properties and a good wetting behavior on copper [2–5]. Solder joints provide both mechanical strength and electrical conductivity, and consequently, play an important role in the connection of electronic components to printed circuit boards. Fundamentally, a solder material should wet the substrate and
∗
Corresponding author. Tel.: +886 6 2757575x62154; fax: +886 6 2745698. E-mail addresses: [email protected] (Y.-H. Lee), [email protected] (H.-T. Lee). 1 Tel.: 886 6 2533131x3529; fax: 886 6 2537912. 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.08.065
provide good adhesion. Furthermore, solder joints must be capable of withstanding the high thermal and mechanical loadings which frequently occur during handling and system use at higher temperatures [6]. In solder joints between a Sn-based solder and a copper substrate, intermetallic compounds (IMCs) such as Cu6 Sn5 and Cu3 Sn form and grow both during the soldering operation itself and during the subsequent system use [7–10]. The IMC layer provides a metallurgical bonding between the solder and the substrate. Furthermore, the morphology of the IMC layer also has a fundamental effect on the solder joint reliability because crack initiation and propagation during thermal fatigue generally takes place at the solder/IMC interface [11]. Hence, a thick IMC layer containing coarse intermetallic compounds may adversely affect the mechanical properties and performance of the solder joint. It is often reasoned that the mechanical integrity of the solder joint is degraded by the inherent brittleness of the IMC layer since cracking can readily occur when the joint is mechanically or thermally stressed [12,13]. Therefore, the formation, growth, and fracture of the IMC layer and the microstructure of the bulk solder have a crucial influence on the solder joint reliability. The mechanical properties of the solder are important because they determine the fracture and thermal fatigue behavior
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Fig. 1. Single lap shear solder joint test specimen.
of the solder joint. Therefore, it is important that the mechanical properties of the solder, such as its shear strength, adhesive strength and ductility, during the soldering operation itself and during the subsequent system use are well understood [14–17]. Attempts have been made to modify the solder microstructure and to improve its mechanical properties by adding small quantities of alloying elements to Sn-based solders [3,4]. Producing small, constituent-phase IMC particles in the bulk solder alloy is often a key design objective when developing alloys designed to enhance the mechanical strength and deformation behavior of solder joints. Some researchers have investigated the use of alloying elements such as copper, nickel, indium, antimony, bismuth, and so on, as a means of reducing the melting point of eutectic Sn–3.5Ag solder while simultaneously improving its mechanical properties [18–26]. The results of these studies have confirmed that certain composite solders exhibit the nec-
essary combination of enhanced strength and other favorable mechanical properties required by the electronics industry. Nickel is commonly used to provide a diffusion barrier between Cu components and Sn or Sn-based solder alloys in order to prevent, or at least suppress, the formation of Cu3 Sn and Cu6 Sn5 intermetallic compounds. Ni is an effective additive since the stable phases, which it forms in the Ni–Sn binary system, grow more slowly than the Cu intermetallic compounds. Furthermore, Ni has good wetting characteristics with Sn. However, the choice of Ni as an alloying element is primarily due to its formation of additional intermetallic phases, which improve the mechanical properties of the solder. Accordingly, the present study investigates the effects of adding varying quantities of nickel particles to eutectic Sn–3.5Ag solder to form composite solders. The formation and growth of interfacial IMC compounds between the Nicontaining composite solders and a Cu substrate are studied. Additionally, the effects of the level of Ni addition and the thermal aging time on the single lap shear strength of the Sn–Ag–x wt%Ni composite solder/Cu substrate joints are investigated. 2. Experimental procedures The present composite solders were prepared by mixing 0.5, 1, 2, and 3 wt% of Ni particles (with an average particle size of 2–3 m) into a molten Sn–3.5 wt%Ag eutectic solder. Prior to mixing, a thin film of Rosin Mildly Activated (RMA) flux (Taiyo electric BS-10 flux, Japan) was coated on the surface of the Ni particles to remove the oxides, which naturally occur. The
Fig. 2. Optical micrographs of interfacial microstructure of: (a) 0.5Ni, (b) 1Ni, (c) 2Ni, and (d) 3Ni solder joints in as-reflowed condition.
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molten solder was maintained at a temperature of 300 ◦ C and stirred with a glass rod in an inert atmosphere for approximately 5 min to ensure a uniform particle distribution [3]. The solder was then poured into a metal box (15 mm × 10 mm × 10 mm) and allowed to cool in air. In this study, a small single shear lap solder joint is designed which better mimics the actual solder joints used in the automobile and microelectronics industries. The various composite solders were rolled into sheets with a thickness of approximately 200 m and were then punched to produce discs with a diameter of 1.5 mm. To ensure proper alignment of the solder joint, a fixture was constructed for sample preparation. The specimens were put into a furnace and hold 260 ◦ C for 60 s, and then removed from the furnace and cooled in air to room temperature. The heating rate was fixed at 1.5 ◦ C/s and the measured cooling rate between 260 and 150 ◦ C was approximately 0.85 ◦ C/s.
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Single shear lap joints with an area of approximately 1 mm2 were then fabricated by using the composite solders to join two copper substrates, as shown in Fig. 1. Following the soldering process, some of the specimens were retained in an as-reflowed condition, while the others were thermally aged at a temperature of 150 ◦ C for 100, 200, 500, or 1000 h, respectively. To obtain a top view of the intermetallic compounds, the majority of the solder on the specimens was ground away, and the specimens were then etched with a 20% HCl–80% C2 H5 OH solution for 30 min to selectively dissolve the remaining solder. A uniaxial micro-force test system (SHIMADZU AG-1, Japan) was used to carry out shear tests on the various solder joint specimens. The tests were performed at room temperature under a strain rate of 1 × 10−2 s−1 . To permit metallographic observations, the fractured samples were mounted in epoxy and then cross-sectioned in a direction perpendicular to the solder–copper
Fig. 3. Optical micrographs of interfacial microstructure of: (a) Sn–Ag, (b) 0.5Ni, (c) 1Ni, (d) 2Ni, and (e) 3Ni solder joints following thermal aging at 150 ◦ C for 1000 h.
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interface. The specimens were then successively ground down to 4000 grit on a silicon carbide paper under water cooling. Finally, the specimens were polished using 3 and 1 m diamond paste and 0.02 m aluminum oxide, and then etched with a 5% HCl solution. The microstructure and composition of the solder joints, and their fracture surfaces following the shear test, were analyzed using optical microscopy (OM), scanning electron microscopy (SEM), and energy dispersive X-ray spectrometer (EDS) techniques. 3. Results and discussion 3.1. Microstructure of composite solder joint Fig. 2 presents optical micrographs of the as-reflowed Sn–3.5Ag–x%Ni composite solder joints. It is observed that the microstructure of the solder consists of globular Sn-rich phase matrix with dispersed Ag3 Sn precipitates around the Sn-rich regions. A small number of Ni–Cu–Sn intermetallic particles are also observed in the solder matrix. The intermetallic compounds within the solder matrix are identified by EDS compositional analysis as (Ni,Cu)3 Sn4 . Comparing the various micrographs in Fig. 2, it is clear that the mean free path between the -Sn phases is reduced with increasing Ni content. Furthermore, a Ni–Cu–Sn intermetallic layer is evident at the interface of the solder and the Cu substrate. EDS analysis reveals that this compound contains Cu, Ni, and Sn and has an atomic compositional ratio close to that of Cu6 Sn5 intermetallic phase. These compounds are found from EDS analysis to be (Cu,Ni)6 Sn5 . This finding is consistent with the results reported by previous investigators [21–25]. Following prolonged aging at an elevated temperature, Sn in the area of the (Cu,Ni)6 Sn5 IMC closest to the copper substrate diffuses into the Cu side and prompts the formation of Cu3 Sn [11]. The optical micrographs presented in Fig. 3 reveal that after aging at 150 ◦ C for 1000 h, the (Cu,Ni)6 Sn5 intermetallic layer in the current specimens becomes noticeably thicker and a Cu3 Sn IMC layer is formed between the (Cu,Ni)6 Sn5 intermetallic layer and the Cu substrate. Fig. 4 shows a typical top view of the intermetallic compounds of an Sn–Ag–3 wt%Ni/Cu joint in the as-reflowed condition. The IMC contains two elongated rod-type Ni–Sn–Cu ternary IMCs. EDS elemental analysis reveals that the longer IMCs are (Cu,Ni)6 Sn5 while the shorter IMCs are Cu6 Sn5 . The nano-sized particles observed on the surface of the longer rodtype IMCs are found to be Ag3 Sn compounds. Fig. 5 shows the Sn–Ag–3 wt%Ni/Cu joint aged at 150 ◦ C for 200 and for 500 h the morphology of large Ag3 Sn IMC platelets located at the (Cu,Ni)6 Sn5 IMC layer. It is observed that the platelets partially penetrate the IMC layer. Since the interfacial layer is formed by the reaction of the molten solder and the substrate during the reflow treatment, it can be inferred that the large Ag3 Sn nucleate initially from the interfacial layer. Since Ag has a higher concentration compared to Cu; Ag3 Sn can nucleate and grow rapidly to a high volume. Large Ag3 Sn intermetallics were also observed in the bulk solder near the solder/Cu interface. These Ag3 Sn decreased in numbers but
Fig. 4. Top view morphology of long rod-like (Cu,Ni)6 Sn5 grains and short rodlike Cu6 Sn5 particles growing from interfacial layer of Sn–Ag–3Ni/Cu joint in the as-reflowed condition.
increased in size due to coarsening as the solder joints were aged. In the soldering process, the formation of intermetallic compounds between the metal substrate and the solder must be
Fig. 5. Morphologies of larger platelet Ag3 Sn and Cu–Ni–Sn IMC layer on Cu substrate. (a) Sn–Ag–3Ni composite solder/Cu substrate following aging at 150 ◦ C for 200 h and (b) 500 h.
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carefully controlled since the brittleness of the IMC layer may degrade the reliability of the solder joint [7,9,11,12,14]. The morphological change and increased thickness of the IMC layer observed in the aged specimens indicates that the interfacial reaction of the solder with the Cu substrate continues during the thermal aging process. The interface reactions between Sn–Ag solders and Cu can be categorized into two types. Cu6 Sn5 and Cu3 Sn are two of the main reaction products. In all cases Cu6 Sn5 is the first compound formed, and very little if any Cu3 Sn is present immediately after solidification. During aging treatment Cu3 Sn will form by a solid state reaction between Cu and Cu6 Sn5 , and after prolonged aging the Cu3 Sn layer becomes more evident. It has been reported [11,16] that the atomic diffusion of Cu and Sn through the interface IMC layer was the main controlling factor for the growth of the interface IMC during aging. The growth kinetic of interface IMC layer follows the square root time law expressed in the equation below: X = X0 + Dt1/2 . The Cu3 Sn layer facing a Cu substrate is formed thinner than the Cu6 Sn5 layer. Fig. 6 plots the variation of the average Cu6 Sn5 and Cu3 Sn IMC layer thickness with the square root of aging time at 150 ◦ C.
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In general, it can be seen that the thickness of both layers increases with an increasing Ni content. Guo et al. [25] have indicated significant IMC layers growth was observed in Ni composite solder joint under isothermal aging at 150 ◦ C. The addition of more than 1 wt% Ni appears to have a particularly profound influence on the thickness of the (Cu,Ni)6 Sn5 reaction layer. The IMC layer thickness of the Sn–Ag–3Ni joint specimen is approximately 1.5 times that of the Sn–Ag joint specimen in both the as-reflowed condition (i.e. Sn–Ag: 6.1 m, Sn–Ag–3Ni: 9.4 m) and following thermal storage for 1000 h (i.e. Sn–Ag: 13.6 m, Sn–Ag–3Ni: 19.7 m). Fig. 6 also shows that the thickness of the IMC layer increases with an increasing aging time. The thickness of the Cu6 Sn5 and (Cu,Ni)6 Sn5 intermetallic layer is greater than that of the Cu3 Sn layer in all cases. Significantly, the Cu3 Sn layer in the Sn–Ag solder joint is consistently thicker than that in any of the Sn–Ag–xNi solder joints following thermal aging. The formation of Cu3 Sn is affected by the phase stability of Cu6 Sn5 according to the reaction Cu6 Sn5 + Cu → Cu3 Sn. The thermodynamic affinity in Cu6 Sn5 is stronger between Ni and Sn than between Cu and Sn. Thus, the phase stability is greater for (Cu,Ni)6 Sn5 than for Cu6 Sn5 [26]. This finding suggests that
Fig. 6. Variation of interfacial intermetallic layer thickness with Ni content and aging period for storage at 150 ◦ C.
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Fig. 7. Variation of single lap shear strength with Ni addition as function of aging time for storage at 150 ◦ C.
the addition of Ni to the Sn–Ag solder effectively suppresses the growth of Cu3 Sn compounds. 3.2. Single shear lap strength Since solder joints are frequently subjected to mechanical loading during their service lives, their mechanical properties, e.g. their fatigue and shear strengths and their creep resistance, have a crucial effect on the solder joint reliability and hence
Fig. 8. The trend of tension strength and shear strength with Ni content.
the integrity of the overall electronic package. Therefore, a series of shear tests at room temperature under a strain rate of 1 × 10−2 s−1 evaluated the effect of the interfacial reactions on the strength of solder joints aged for 0, 100, 500 and 1000 h at 150 ◦ C. Fig. 7 presents the shear strength results obtained for the Sn–Ag, Sn–Ag–0.5Ni, Sn–Ag–1Ni, Sn–Ag–2Ni and Sn–Ag–3Ni solder alloys. In general, the as-reflowed joints have greater shear strength than the aged specimens. Fur-
Fig. 9. SEM fracture surface of Sn–Ag–1Ni/Cu joint specimens stored at 150 ◦ C for different aging times: (a) as-reflowed, (b) 100 h, and (c) 1000 h; (d) OM cross-section view of specimen aged for 1000 h.
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thermore, the shear strength of the as-reflowed solder joint improves with increasing Ni addition. The shear strength enhancement is most likely attributed to the additional formation of (Ni,Cu)3 Sn4 particles and a refinement of the solder microstructure as the level of Ni addition increases [3]. Fig. 7 demonstrates that thermal aging leads to a reduction in the shear strength of the solder specimens. Thermal aging has two important effects on the solder joint microstructure; it increases the IMC layer thickness and it leads to a coarsening of the solder microstructure [16,25–30]. Therefore, the reduction in shear strength observed in the current aged specimens can reasonably be attributed to an increased thickness of
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the brittle intermetallic layer, the agglomeration of (Ni,Cu)3 Sn4 compounds, or a coarsening of the Ag3 Sn particles. Fig. 7 shows that the shear strength drops most dramatically during the first 100 h of thermal aging. Comparing the tension strength and shear strength, Fig. 8 presents the tensile strength [4] increases with increasing Ni content, and variation of shear strength with Ni content also reveals a similar trend. In a microstructural sense, the reduction in strength after 100 h can be attributed to a significant coarsening of the microstructure of the joint. Fig. 3 (corresponding to 1000 h thermal storage) provides clear evidence of the degradation in the microstructure during aging, i.e. the images show a significant (Cu,Ni)6 Sn5 IMC layer growth at the interface, coarse (Ni,Cu)3 Sn4 , Ag3 Sn, and (Cu,Ni)6 Sn5
Fig. 10. (a) SEM fracture surface of Sn–Ag–2Ni/Cu joint (near Cu substrate); (b) OM cross-section view of Sn–Ag–2Ni/Cu joint specimens aged at 150 ◦ C for 1000 h, (c) high-magnification image of region A; (d) SEM fracture surface of Sn–Ag–3Ni/Cu joint; (e) high-magnification image of region B and (f) OM cross-section view of Sn–Ag–3Ni/Cu joint specimens aged at 150 ◦ C for 1000 h.
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phases in the bulk of the solder, and voids in the interfacial IMC layer. Figs. 6 and 7 show that the Sn–Ag–3%Ni solder joint has higher shear strength than any of the other solder joints following thermal aging for 1000 h despite having the thickest IMC layer. Therefore, it appears that the thickness of the IMC layer does not play a critical role in controlling the shear strength of single lap solder joints. Note that a similar finding was also reported in the studies of Deng et al. [16] and Yoon et al. [28], in which it was shown that the shear strength of solder joints is controlled primarily by the mechanical properties of the solder rather than the thickness of the intermetallic compound layer. 3.3. Fracture surface analysis The fracture surfaces of the solder/Cu joints provide a valuable insight into the deformation mechanisms, which lead to their failure. Therefore, the failed as-reflowed and aged joints were cross-sectioned in a perpendicular direction to the solder/Cu interface and their fracture surfaces inspected using SEM and OM techniques. Fig. 9 presents SEM images of the fracture surfaces of Sn–3.5Ag–1Ni/Cu joint specimens in the as-reflowed condition and following aging at 150 ◦ C for 100 and 1000 h, respectively. Fig. 9(d) presents an OM image of the fracture surface cross-section in an Sn–3.5Ag–1Ni/Cu joint aged for 1000 h. In all cases, it is observed that the specimens fracture within the Sn–3.5Ag–1Ni solder matrix. This suggests that the Sn–3.5Ag–1Ni solder matrix is weaker than the IMC layer between the solder and the Cu substrate. Furthermore, the fracture surfaces in Fig. 9(a–c) contain elongated dimplelike structures. Hence, it is apparent that the Sn–3.5Ag–1Ni/Cu joints all fail with a ductile fracture mode, irrespective of their thermal aging conditions. The trend in Fig. 9 is also representative for the cases of Sn–3.5Ag/Cu and Sn–3.5Ag–0.5Ni/Cu joint specimens fracture surfaces. Fig. 10 shows the fracture surfaces of the Sn–Ag–2Ni/Cu and Sn–Ag–3Ni/Cu joints aging at 150 ◦ C for 1000 h conditions. It can be seen that these specimens have a brittle fracture morphology rather than the ductile morphology observed in the Sn–3Ag and low Ni-containing composite solder specimens. Due to the combined effect of the brittle (Cu,Ni)6 Sn5 IMC layer and the stress concentration at the solder/intermetallic interface, (Cu,Ni)6 Sn5 nodular tips protrude into the fracture surface. Under shear conditions, these nodular tips cause a large shear strain localization effect, and hence the IMC interface becomes the preferred site for crack formation and propagation. Therefore, as shown in Fig. 10(b and f), partial debonding failures occur at the interface between the solder matrix and the (Cu,Ni)6 Sn5 IMC layer. The faceted appearance of the fracture surfaces in Fig. 10(a and d) are caused by cleavage fractures of the (Cu,Ni)6 Sn5 layer. When the cleavage planes in neighboring (Cu,Ni)6 Sn5 grains are approximately parallel, the fracture propagates easily through the inside of the IMC layer or along the interface of the (Cu,Ni)6 Sn5 layer and the solder matrix. Fig. 10(b and f) shows the debonding phenomenon in the Sn–Ag–2%Ni and
Sn–Ag–3%Ni joints at the interface between the solder matrix and the (Cu,Ni)6 Sn5 IMC layer. With 2 and 3% Ni addition, the solder matrix has a high (Ni,Cu)3 Sn4 and Ag3 Sn intermetallic phase content and hence has an enhanced strength. Therefore, the interface between the solder matrix and the IMC layer in the Sn–Ag–2%Ni and Sn–Ag–3%Ni joints is relatively weaker than that in the Sn–Ag and low Ni-containing composite solder joints. Consequently, as shown in Fig. 10(b and f), the shear fracture appears to follow a mixed path with an initial crack at the (Cu,Ni)6 Sn5 /solder interface region followed by propagation within the IMC layer or along the near interface ((Cu,Ni)6 Sn5 /solder) region. From the results presented in Figs. 9 and 10, it appears that the probability of debonding shear failure depends on the relative balance between the strength of the solder matrix and the strength of the (Cu,Ni)6 Sn5 /solder interface. In general, cracking occurs at the weakest point in the joint microstructure, i.e. in the solder matrix in the Sn–3.5Ag and low Ni-containing composite solder specimens and at the interface of the solder matrix and the IMC layer in the higher Ni-content specimens. 4. Conclusion This study has examined the microstructural evolution of the IMC layer formed between Sn–Ag–x%Ni composite solders and a Cu substrate and has investigated the single lap shear strength of specimens treated under various thermal aging conditions. The following conclusions can be drawn: 1. In the as-reflowed condition, the Sn–Ag–x%Ni composite solders contain only a (Cu,Ni)6 Sn5 layer at the interface between the solder matrix and the Cu substrate. However, when the joints are thermally aged, a Cu3 Sn layer is formed between the (Cu,Ni)6 Sn5 layer and the Cu substrate. In general, the thickness of the IMC layer is found to increase with an increasing Ni content and an increasing thermal storage time. 2. The shear strength of the single lap joints increases with increasing Ni addition in the as-reflowed specimens. However, the shear strength decreases as the joints are thermally aged regardless of their Ni content. The reduction in shear strength is caused primarily by a coarsening of the solder microstructure. The effect of the increased IMC layer thickness following thermal storage does not appear to have a significant effect of the shear strength of the joint. 3. For Ni additions of 0.5 and 1 wt%, the specimens fracture in the solder matrix near the IMC layer/solder interface. The fracture surfaces are characterized by elongated dimple-like structures, which suggest a ductile failure mode. However, when the level of Ni addition exceeds 1 wt%, the failure mode changes from ductile fracture to brittle fracture, and the specimens fail at the interface of the solder matrix and the IMC layer. Therefore, it appears that higher levels of Ni addition increase the strength of the solder matrix beyond that of the interface.
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