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Enhanced ductility and mechanical strength of Ni-doped Sn–3.0Ag–0.5Cu lead-free solders
Materials and Design 55 (2014) 309–318
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Enhanced ductility and mechanical strength of Ni-doped Sn–3.0Ag–0.5Cu lead-free solders A.A. El-Daly a, A.M. El-Taher a,⇑, T.R. Dalloul b a b
Physics Department, Faculty of Science, Zagazig Univ., Zagazig, Egypt Physics Department, Faculty of Science, Islamic Univ. of Gaza, Gaza, Palestine
a r t i c l e
i n f o
Article history: Received 25 July 2013 Accepted 2 October 2013 Available online 16 October 2013 Keywords: Lead free solder alloys Microstructure Mechanical properties
a b s t r a c t In this paper, the tensile tests were conducted to investigate the effect of adding a small amount of Ni on the microstructure, thermal and mechanical properties of 3.0Ag–0.5Cu(SAC 305) solder. The results indicated that addition of Ni can effectively decrease both the undercooling and the onset melting temperature of SAC(305) solder alloy. The strength and ductility of the SAC(305) solder depend significantly on Ni content. In general, the SAC(305)–0.5%Ni solder reveals superior mechanical properties in terms of maximum strength and ductility when compared to the high Ni-content or plain solders. Microstructure analysis revealed that a new g-(Cu,Ni)6Sn5 intermetallic compound (IMC) phase containing large amount of Ni was generated, while the initial Cu6Sn5 phase was converted into (Cu,Ni)6Sn5 phase after 0.5%Ni addition. Besides, the fine fiber-like Ag3Sn and finer dot-shaped precipitates rather than needle-like morphology have occurred at the surface of b-Sn matrix easily, which could provide more obstacles for dislocation pile up in the adjacent grains and enhanced the mechanical property. With increasing Ni addition, the Ni-doped SAC(305) solder showed a corresponding deterioration in their mechanical property due to the coarsening of (Cu,Ni)6Sn5 IMCs and increasing the inter-particle spacing of Ag3Sn IMCs in the eutectic colony. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Owing to considerations of environmental protection and human health, the adoption of lead-free solder has become an inevitable trend in the electronics industry. The lead-free solders based on Sn–Ag–Cu (SAC) system with Ni addition have become popular because of their superior soldering properties, as well as their comparatively low cost [1,2]. In view of that, the SAC solders are commonly used in surface mount technology (SMT) assembly for microelectronics into industrial production [3]. However, many issues with the SAC solders still remain unresolved, such as the best composition, the large undercooling in solidification, and the formation of large brittle IMC. Especially, alloys with high Ag content exhibit the formation of large Ag3Sn IMCs and short creeprupture life time in service [4]. For these reasons, efforts have been made to develop solders with low Ag and Cu contents. Reducing the Ag and Cu contents of the SAC alloys, such as Sn–3.0Ag– 0.5Cu (SAC305) alloy gives rise to more primary b-Sn phase (large b-Sn grains) and decreases the number of Ag3Sn IMC particles. It is therefore expected to result in lower elastic modulus and yield ⇑ Corresponding author. Tel.: +20 552327173/552303252 (O); fax: +20 552308213. E-mail addresses: [email protected], [email protected] (A.M. El-Taher). 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2013.10.009
strength than the high Ag content SAC alloys. It has been established also that the minor element additions such as: Ti, Zn, Al, Ni, Bi, In, Fe, Co, and rare-earth elements are a simple and effective way for modifying the properties of these lead-free solders [5–8]. One of the highly focused areas of research is the effect of Ni content on the microstructure and mechanical behavior of SAC solder alloys due to reliability concerns. The main concern is the existence of the (Cu,Ni)6Sn5 IMC particles, which possess high stability and inherently compatible geometry to the surrounding microstructure, i.e. bulk solder and substrate. The (Cu,Ni)6Sn5 appears at the bulk solder alloy when the Ni content is 0.05 mass% [9]. However, most of the experiments reported in the recent literatures focused on the reliability of Sn–Ag and SAC alloys with adding small amount of Ni. Except for that, attempts of adding different amounts of Ni (0.5–5.0 mass%) to develop solder joints with high reliability have been performed [10–12]. High Ni content is desired for thermal cycling performance because of the strengthening effect and stability of (Cu,Ni)6Sn5 at different temperatures. Previously, it was reported that the addition of 5 mass% Ni could completely prevent the formation of Cu3Sn in SAC alloys [13]. It has been suggested that the incorporation of Ni in Cu6Sn5 makes this phase more thermodynamically stable [2]. Gao et al. [14] found that the interaction parameter between Sn and Ni is much more negative than that between Sn and Cu. This implies that (Cu,Ni)6Sn5 is more stable than Cu6Sn5. Thus, in the presence of Ni atoms,
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Table 1 Chemical composition of the solders studied (wt.%). Alloy
Cu
Ag
Ni
Sn
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
0.5 0.5 0.5
3.0 3.0 3.0
– 0.5 1.0
Bal. Bal. Bal.
the growth of (Cu,Ni)6Sn5 is more favorable and hence faster. Consequently, the crack propagation in the (Cu,Ni)6Sn5 layer becomes more difficult [15]. Belyakov and Gourlay [16] studied the growth of NiSn4 during the solidification of Sn–Ni alloys containing 0– 0.45 wt.% Ni. They found that Primary NiSn4 crystal growth was promoted by high cooling rates and the Sn–NiSn4 eutectic could grew under all solidification conditions used. Although nickel has a strong influence on solderability and deformation behavior of solders, the detailed mechanism is not fully understood. Even for the ternary SAC system, there is very limited knowledge about the influence of Ni element on the properties of SAC(305) alloy. This paper focuses on the effect of Ni additions on the microstructure, thermal behavior and corresponding mechanical properties of SAC(305) alloy during high temperature and strain rate range. The key factors that affect the thermal and mechanical behavior of the solder alloy after Ni addition are discussed. 2. Experimental procedures The nominal compositions for both plain and Ni-doped solders are given in Table 1. Their compositions were Sn–3.0Ag–0.5 wt.% Cu, Sn–3.0Ag–0.5Cu–0.5 wt.% Ni and Sn–3.0Ag–0.5Cu–1.0 wt.% Ni (which are hereafter called SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni, respectively). The lead-free SAC(305)–xNi (x = 0– 1.0 wt.%) solders were prepared from pure Sn(99.99%), Ag(99.999%), Cu(99.98%) and Ni(99.99%) as raw materials. The process of melting was carried out in a vacuum arc furnace under high purity argon atmosphere to produce rod-like specimen with a diameter of approximately10 mm. The melt was held at 600 °C for 2 h to complete the dissolution of Sn, Cu, Ag and Ni and then poured in a steel mold to prepare the chill cast ingot. A cooling rate of 6–8 °C/s was achieved, so as to create the fine microstructure typically found in small solder joints in microelectronic packages. The microstructure was examined by scanning electron microscopy (SEM) JSM-5410, Japan. A solution of 3%HCl, 2%HNO3 and 95% (vol.%) Ethyl alcohol was prepared and used to etch the samples. Phase identification was based on Energy Dispersive X-ray Spectrometry (EDS). Because of different approaches in the
Table 2 Comparison of solidus temperature (Tonset) and liquidus temperature (Tend) for SAC305, SAC–0.5Ni and SAC–1.0Ni solder alloys from heating curve. Alloy
(Tonset) (°C)
Tend (°C)
Pasty range (Tend-Tonset) (°C)
Peak temperature (°C)
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
219.9 218.7 216.2
221.7 224.4 224.6
1.8 5.7 8.4
221.6 224.4 224.4
Table 3 Comparison of melting temperature (Tm) for heating, solidus temperature (Ts) and under cooling range for various solder alloys. Alloy
Tm (°C)
Ts (°C)
Undercooling (Tm–Ts) (°C)
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
219.9 218.7 216.2
195.7 203.0 201.9
24.2 15.7 14.3
specimen design, test setups and experimental methodology are necessary to investigate solder tensile behavior at different volumes. In the present work, the homogenized cast ingots were then mechanically machined into a wire samples with a gauge length marked 4 102 m for each samples and 2.5 mm diameter. Details are described in [17]. Before testing, the specimen was annealed at 120 °C for 60 min to reduce the residual stress induced in the sample preparation. Subsequently, tensile creep tests are conducted on a 3360 universal material testing system, and GWT 504 high temperature testing system respectively. The tests were carried out at various strain rates ranging from 105 to 103 s1 and constant temperature of 25 °C. Also, the tests were conducted at different temperatures ranging from 25 to 120 °C with a constant strain rate of 8.8 104 s1. The axial strain is measured in accordance with the ASTM:E83-10a, and ASTM D76/D76M-11 standard practices for force verification. Then, the mechanical properties were obtained by averaging testing data. The temperature variation inside the high temperature furnace is maintained within 1.5 °C. 3. Results and discussion 3.1. Thermal analysis For understanding the relationship between melt properties and Ni content, the thermal behavior of solders were analyzed using DSC analysis and the results are demonstrated in Fig. 1 and summarized in Tables 2 and 3. During the heating process, it is noteworthy that the Tonset decreases from 219.9 to 216.2 °C with
Fig. 1. DSC results during heating (endothermal) and cooling (exothermal) of SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solders.
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Eutectic
Ag3Sn
β-Sn
(Cu,Ni)6 Sn5
(d)
(Cu,Ni)6Sn5
Fracture
β-Sn Ag3Sn
β-Sn Fig. 2. Scanning electron micrographs (SEM) of as-cast: (a) SAC(305), (b) SAC(305)–0.5Ni, (c) SAC(305)–1.0Ni solders and (d) EDS analysis of (Cu,Ni)6Sn5 IMC particles.
(a)
(b)
Ag 3 Sn Cu6 Sn 5
(c)
Fig. 3. (a) High-magnification SEM microstructure of SAC (305) and EDS analyses of (b) Ag3Sn, and (c) Cu6Sn5 IMCs.
an increase in the Ni content. The prominent endothermic peak of Ni-free SAC(305) solder appeared at 221.6 °C, which corresponds to the eutectic temperature of the Sn–3.5Ag–0.25Cu ternary system as reported by Tsao and Chang [18]. However, the raising of
endothermic peak from 221.6 °C to 224.4 °C due to Ni addition can be attributed to the high melting temperature of Ni, which dissolved with a low solubility in the lead-free SAC(305) solder. Nevertheless, the pasty range of SAC(305), SAC(305)–0.5Ni and
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Fig. 4. Tensile properties of the solders SAC(305), SAC(305)–0.5Ni and SAC(305)– 1.0Ni solder alloys at T = 25 °C and e_ = 2 103 s1.
Table 4 Tensile properties of SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys. Alloy
UTS (MPa)
0.2%YS (MPa)
Elongation (%)
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
28.0 31.5 27.0
25.0 27.5 23.0
8.2 12.3 6.1
SAC(305)–1.0Ni solders, respectively, was 1.8, 5.7 and 8.4 °C, which is lower than 11.5 °C for Sn–Pb eutectic [17]. This narrow pasty range of SAC(305) solders may avoid manufacturing problems, such as increasing the sensitivity to vibration during wave soldering. Also, it can decrease the probability of fillet lifting phenomena, the tendency towards porosity and hot tearing due to the effect of alloy shrinkage. For these reasons, Ni-containing SAC(305) solders have displayed satisfactory reliability. Undercooling is defined as the difference between the melting point during heating and the solidification temperature during cooling, and relates to the difficulty of nucleating a solid phase in a liquid state [19]. By having a direct effect on nucleation during solidification, the Ni could affect the undercooling of the b-Sn phase, and this may cause significant repercussions on the microstructure of the solidified sample. In Fig. 1 and Table 3, the exothermal peaks upon cooling for both alloy samples appeared at lower temperature compared with their endothermic peaks due to the undercooling, but the reductions in temperature were not identical. It is worthy of notice that the degree of undercooling of SAC(305) solder was strongly decreased from 24.2 to 15.7 and 16.3 °C after Ni addition. This means that the solidification processes of Ni-containing alloys were initiated with the nucleation of new IMC precipitates. The literature showed also that a decrease in b-Sn undercooling resulted in reducing Ag3Sn formation in the SAC alloys, and it can be achieved by the addition of a fourth alloying elements such as Ni, Zn and Sb [17,19]. In this case, Ni could act as an additional nucleation sites, improving the rate of nucleation, leading to a reduced undercooling effect. As a result, one may expect that the microstructure of the Ni-solidified samples does not show any noticeable large Ag3Sn precipitates, and the Ag3Sn growth is probably only limited to small phases in the eutectic matrix due to the combined effects of streaming and enhanced nucleation. These results are quite consistent with the explanation proposed by Chen et al. [7] that primary IMCs might act as
Fig. 5. Effect of strain rate on: ultimate tensile strength (UTS), yield stress (0.2YS) and elongation (El.%) at T = 25 °C for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
heterogeneous nucleation sites for Sn dendrites upon solidification and are able to suppress the undercooling of Sn–Ag–Cu alloys. 3.2. Effect of Ni additions on the solidification microstructures Fig. 2 shows the microstructure of plain SAC(305) and Ni-containing solders in the as-cast condition. In case of Ni-free SAC(305) solder (Fig. 2a), the dark gray b-Sn primary grains with an average diameter of 40–70 lm are surrounded by the light gray area of fine eutectic network (Sn, Ag3Sn and Cu6Sn5) with the average spacing lamellae of 4 lm. These results are very similar to those of the SAC solders [20]. After 0.5 wt.% Ni addition to the
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Fig. 6. Effect of temperature on: ultimate tensile strength (UTS), yield stress (0.2YS) and elongation at strain rate of 8.8 104 s1 for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
SAC(305) solder, a different morphology was observed. The morphology tends to transform from needle-like to fiber-like Ag3Sn and fine dot-shaped precipitates at the surface of b-Sn matrix are observed in Fig. 2b. It can be seen that the average size and spacing of Ag3Sn particles were significantly reduced when compared with the SAC(305) solder. The refined IMC morphology also indicates the increase of nucleation rate and provides the justification for its better performance, which could restrict the dislocation motion. On the other hand, the microstructure could be clear evidence of the presence of new rod-like (Cu,Ni)6Sn5 IMCs induced at a network of eutectic area. All the particles investigated contain a significant amount of Ni with an average content of 17.2–35.7 at.% Ni, as
Fig. 7. Relationship between r and ln(e_) at (a) T = 25, (b) T = 70 and (c) T = 120 °C for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
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Fig. 8. Relationship between ln(r) and ln(e_) at (a) T = 25, (b) T = 70 and (c) T = 110 °C for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
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315
Fig. 9. Relationship between ln[sinh(ar)] and ln(e_) for determination stress exponent (n) values at (a) T = 25, (b) T = 70 and (c) T = 110 °C for SAC(305), SAC(305)0.5Ni and SAC(305)–1.0Ni solder alloys.
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Fig. 10. The activation energy (Q), values of SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
Table 5 Activation energy (Q) and stress exponent (n) values for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys. Alloy
Q (kJ/mol)
Temperature (°C)
a (MPa1)
n
SAC(305)
58.7
25 70 120
0.03739 0.05375 0.07231
7.0 6.0 5.3
SAC(105)–0.5Ni
62.8
25 70 120
0.04135 0.05543 0.07000
8.4 7.1 6.4
SAC(105)–1.0Ni
50.6
25 70 120
0.43243 0.05500 0.06667
9.3 8.0 7.2
confirmed by EDS analysis shown in Fig. 2d. It implied that the initial Cu6Sn5 phase was converted into g-(Cu,Ni)6Sn5 phase during cooling. The formation of rod-like g-(Cu,Ni)6Sn5 morphology may result from anisotropy of surface energies and compositional related to the presence of Ni, which influences the growth behavior of Cu6Sn5. However, the microstructure changes dramatically after 1.0 wt.% Ni addition. The volume fraction of the Ag3Sn IMC particles in the eutectic area tends to decrease and its interparticle distance tends to increase to the average value of 15 lm. Worthy of notice is that the b-Sn-rich phases became larger, and the average size of plate-like (Cu,Ni)6Sn5 IMCs was increased. Increasing the Ni content also can enhance the crack formation inside (Cu,Ni)6Sn5 IMCs during cooling, as seen in Fig. 2c. This will more dramatically degrade the mechanical properties and lifetimes of SAC(305)–1.0Ni solder joints than the other alloys. Fig. 3a shows the close-up view of plain SAC(305) with the corresponding EDS analysis, where the light needle-like morphology is the Ag3Sn (Fig. 3b) and scallop-like morphology phase is the Cu6Sn5 IMC, as shown in Fig. 3c. 3.3. Tensile tests The representative stress–strain plots of one temperature (RT) and one strain rate (2 103 s1) for the base alloy as well as the solders containing 0. 5 and 1.0 wt.% Ni are shown in Fig. 4. The average values of ultimate tensile strength (UTS), 0.2% yield strength (0.2YS), and percentage of elongation are given in Table 4. The SAC(305) solder had an UTS of 28 MPa, 0.2YS of 25 MPa, and elongation to failure of 8.2%. With addition of 0.5%Ni, a prom-
inent effect on both the strength and ductility of SAC(305) solder had occurred. The UTS and 0.2%YS of SAC(305)–0.5Ni alloy were 31.5 and 27.5 MPa, respectively, which were higher 1.1 times that of the baseline SAC(305). Besides, the total elongation of SAC(305)– 0.5Ni alloy is 12.3%, which was 52.3% higher than that of SAC(305) solder. The enhancement in both the strength and ductility of SAC(305)–0.5Ni solder is mainly attributed to the formation of (Cu,Ni)6Sn5 and the homogeneous distribution of fine and large eutectic mixture, as indicated in Fig. 2. A number of studies are available on the effect of nickel as an alloy addition to the Sn-based solders [21]. It has been suggested that the incorporation of Ni in the (Cu,Ni)6Sn5 IMC makes this phase more thermodynamically stable at room temperature, since Ni was found to retard the transformation of hexagonal structure to monoclinic structure at 186 °C, thereby stabilizing the hexagonal structure at room temperature. Accordingly, the stabilization of hexagonal (Cu,Ni)6Sn5 will suppress the volume change associated with phase transformation and hence will eventually prevent IMC cracking. This is expected to improve the solder performance. Nogita and Nishimura noticed that [15] the lattice structure of (Cu,Ni)6Sn5 is only thermally stable at room temperature as the hexagonal g-Cu6Sn5 when containing 9 at.% Ni, while that containing 2 at.% Ni is not stable and transforms to monoclinic during cooling. This result indicates that the high-temperature hexagonal g-(Cu,Ni)6Sn5 is only stable at room temperature when sufficient Ni is present in the (Cu,Ni)6Sn5 IMC. In view of these, the current results reveal hexagonal g(Cu,Ni)6Sn5 containing 17.2–35.7 at.% Ni, which support and agree well with the previous studies, and predicted that Ni would thermodynamically stabilize (Cu,Ni)6Sn5 down to room temperature. The stabilization of hexagonal (Cu,Ni)6Sn5 in Ni-containing solders could have significant industrial applications and points to future research directions. However, the extra 1.0 mass% Ni additions caused the UTS and 0.2YS to decrease to 27 and 23 MPa, respectively, and elongation was diminished to 6.1%. These values illustrate that under the same testing conditions, the UTS, 0.2YS and elongation of SAC(305) solder are significantly decreased with increasing Ni content. The reason may be that the volume fraction of the Ag3Sn IMC particles in the eutectic area tends to decrease and the interparticle distance tends to increase with the addition of 1 wt.% Ni. However, more dissolved Ni can combined with the bulk Cu6Sn5 particles and accumulate together in the solder matrix inhomogeneously to form coarse (Cu,Ni)6Sn5 IMCs during solidification. Another unusual finding is that the coarse (Cu,Ni)6Sn5 grain surface is fractured slightly due to inconsistent volume contraction at the solder/ (Cu,Ni)6Sn5 interface during cooling, as seen in Fig. 2c. Such structure of bulk (Cu,Ni)6Sn5 may degrade the mechanical properties and fracture strengths of solder joints, leading to poor reliability of electronic devices. 3.4. Effect of strain rate on the mechanical properties Fig. 5 presents the effect of strain rate on UTS, 0.2YS and percentage of elongation of SAC(305)–xNi solder specimens at room temperature. In general, the entire alloys displayed an increase in both UTS and 0.2% YS with increasing strain rate. The mechanical feature improvement was attributed to the interaction between IMCs or precipitates with the dislocation motion. These results are consistent with the previous findings [22] that the tensile behavior of SAC(205)–0.05Ni alloys often exhibits a strong dependence on the strain rate, or high strain rate sensitivity. At low strain rates, the precipitates hinder the fast moving dislocation. With increasing strain rate, the IMCs cannot capture the moving dislocation any more due to the faster velocity of the dislocation in the alloy. Therefore, when the stress is applied to the solder, the dislocations glide on the slip planes which can tolerate the
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deformation. Since the b-Sn-rich phase is the basic ductile phase of the SAC(305)-based solders, the IMCs formed inside the matrix make the solder exhibits high strength. Relevant investigations [22] also demonstrate that the (Cu,Ni)6Sn5 phases in Ni-containing samples have higher elastic modulus and hardness than Cu6Sn5. In consequence, dissolution of 0.5Ni in Cu6Sn5 phase could result in higher mechanical strength of Ni-containing samples as seen in Fig. 5. Compared with plain solder, 1.0 Ni-containing samples exhibited a lower mechanical strength and lower ability for plasticity, which can sustain forced deformation with crushing fracture due to the coarsening of (Cu,Ni)6Sn5 phase. However, examination of these curves reveals also that all the samples have experienced plastic deformation, and the plastic flow is directly related to the increase of strain rate, except at strain rate of 4 104 s1, where the total strain was decreased. One possible reason is the material structure is piecewise discontinuous and heterogeneous owing to the existence of micro-defects or damage. Based on the equivalence principle [23], the micro-defects can be ‘smeared out’ and the stress and strain state can be considered as homogeneous. It seems that the non-uniform microstructure excludes the possibility of microstructural homogeneities, which can result in consistent elongation values. However, no attempt is made to identify the physical nature of the damage parameters and to distinguish between different damage mechanisms. 3.5. Effect of temperature on the mechanical properties To achieve the temperature dependence of deformation behavior of the tested solders, their tensile properties were investigated in the temperatures range of 25–120 °C. Typical stress–strain curves of the alloys obtained at a strain rate of 8.8 104 s1 are shown in Fig. 6a–c. In contrast to the variation of total elongation, which increased with increasing temperature, the UTS and 0.2%YS of three alloys are dramatically decreased with increasing temperature. This is a consequence of dynamic recovery of lead-free solder, which are encouraged at high temperatures. Increasing the deformation temperature will initiate rearrangement of the dense dislocation networks formed by strain hardening into simple and more ordered ones [24]. This reduces the lattice energy, and therefore lowers the values of UTS and YS at higher temperatures. In addition, the dislocations have much more energy and can overcome tiny IMCs particles that existed in the b-Sn-matrix. However, Fig. 6 also reveals that the SAC(305)–0.5Ni solder possess the highest tensile strength values at all temperature range. In contrast, the SAC(305)–1.0Ni alloy exhibits the lowest one, which can be reflected by the microstructural change observed after 1.0Ni addition (Fig. 2c). The SAC(305)–1.0Ni alloy exhibited small amount of Ag3Sn and coarse (Cu,Ni)6Sn5 IMCs in the eutectic colony. In the meantime, the coarsening deteriorates the strengthening effect of the IMC particle, and thus the b-Sn structure is dominant in UTS at higher temperature. Those massive precipitation alloys, must be avoided because they will possible induce crack in the matrix under low stress or some cyclic stress under service conditions. Nevertheless, the ductility as measured by percent elongation has increased with inconsistent behaviors. The best ductility is usually obtained at the intermediate temperature region. This is recognized in many alloys and occurs at temperatures high enough for grain boundary sliding. It depends on many aspects such as; the matrix phase, IMC chemistries, the nature of interface formed between the IMCs and the matrix, the growth rate and the stability of hard IMCs in the alloy matrix [23]. 3.6. Kinetic analysis of SAC(305)–xNi solders during hot deformation The most comprehensive way of analyzing the deformation data incorporates both the stress exponent ‘n’ and activation
energy ‘Q’ of the operating deformation process. The data are analyzed based on the standard hyperbolic sine Eq. (1) [25]. The appeal of the hyperbolic sine analysis is that it is mathematically capable of incorporating both low- and high-stress data in a single model reflecting power-law and power-law breakdown behavior, respectively.
e_ ¼ A½sinhðarÞn expðQ =RTÞ
ð1Þ
where A (s1), a (MPa1), are material constants independent of temperature, a is stress level parameter, R is the universal gas constant, n is the stress exponent constant related to strain rate, T(K) is thermodynamic deformation temperature, r (MPa) is the steady state flow stress and Q (kJ/mol) is the activation energy of deformation. The logarithm was taken and Eq. (1) was rearranged:
ln e_ ¼ ln A þ n ln½sinhðarÞ Q =RT
ð2Þ
The value of a represents the stress reciprocal at which the material deformation changes from power to exponential stress dependence. An approximate value of a was calculated by a = b/ n1, where b and n1 are the average slopes of ln(e_) r and ln(e_) ln(r) lines at constant T, respectively. The plots of ln(e_) r and ln(e_) ln(r) lines at temperature of 25, 70 and 120 °C for all alloys are given in Figs. 7 and 8. The stress exponent values ‘n’ are calculated from the slope of the ln e_ against ln[sinh(ar)] plot for varying temperature (Fig. 9). The activation energy Q can be expressed as the slope of ln(e_) against 1/T (Fig. 10). The results summarized in Table 5 reveals a range of stress exponents of 5.3–9.3 and activation energies, with the predominance of activation energies falling in the range of 50.6–62.8 kJ/mol for dislocation core diffusion, in consistent with Q 40–65 kJ/mole for Sn self-diffusion through dislocation cores [25], and is almost half the value of that for lattice diffusion (approximately 100 kJ/mol) [26]. On the other hand, the decrease in n value with increasing temperature has been related to the instability of the microstructure, which occurs during high temperature deformation [24]. The anisotropy of the Sn unit cell and the differences in microstructural features and scale suggest that wide ranges of values for both n and Q are to be expected. Deformation mechanisms in the particle-strengthened or multiphase eutectic structures of the Pb-free alloys may thus be similar to those observed in pure tin, but operate at higher stress than in pure tin. According to a dispersion strengthening mechanism, the refinement of the Sn-rich phases, the refinement of eutectic area and uniformly dispersed IMC particles have positive effects on the enhanced mechanical performance of the SAC(205)–0.05Ni solder. They not only help inhibit the grain boundaries from sliding, but also impede the dislocations movement and raise dislocation densities. In contrast, the enhanced crack formation inside (Cu,Ni)6Sn5 IMCs during cooling and increased the inter-particle spacing inside the eutectic area have the negative effects on the mechanical performance of the SAC(205)–1.0Ni solder. However, by determining A, Q, n and a values for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni alloys, the creep strain rate, respectively, can be estimated from the constitutive equations represented by:
e0 ¼ 4:16 103 ½sinhð0:06088rÞ6:6 exp e0 ¼ 1:12 103 ½sinhð0:04967rÞ7:5 exp e0 ¼ 4:16 103 ½sinhð0:08078rÞ5:6 exp
58:7 RT
62:8 RT
50:6 RT
ð3Þ
ð4Þ
ð5Þ
This model fits the experimental tensile curves of SAC(305) alloys well. Hence, the steady-state strain rate and some temperatures can be accurately predicted for both solders by Eqs. (3)–(5).
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4. Conclusion The dependencies of UTS, 0.2YS and elongation to fracture on representative microstructural parameters were shown to be associated with the morphological changes of the Ag3Sn IMC developed along the SAC(305) alloys as well as the formation of new (Cu,Ni)6Sn5 particles in the SAC(305) solder due to Ni addition. Following conclusions can be made based on the results: (1) Addition of 0.5Ni not only refined the Ag3Sn IMC phase and decreased interparticle spacing, but also promoted the formation of new (Cu,Ni)6Sn5 particles in the SAC(305) solder, which could provide more obstacles for dislocation pile up in the adjacent grains. With addition of 1.0Ni, the volume fraction of Ag3Sn IMC particles in the eutectic area decreased and its interparticle distance increased. Besides, the average size of plate-like (Cu,Ni)6Sn5 IMCs was increased. (2) Additions of Ni into the lead-free SAC solder caused its undercooling and solidus temperature to decrease by about 7.9–8.5 °C and 1.2–3.7 °C, respectively. (3) Addition of 0.5Ni to SAC(305) alloy resulted in an increase in YS and UTS as well as the ductility at all temperatures and strain rates, while the benefits of mechanical behavior and ductility were reduced when the concentration of Ni exceeded 0.5 wt.%. (4) According to the obtained stress exponents and activation energies, it is proposed that the dominant deformation mechanism in SAC(305) solders is dislocation core diffusion over the whole temperature range investigated.
References [1] Cheng-Ying H, Jenq-Gong D. Wetting kinetics and wettability enhancement of Pd added electrolytic Ni surface with molten Sn–3.0Ag–0.5Cu solder. Mater Lett 2013;92:278–80. [2] Mu Dekui, Yasuda Hideyuki, Huang Han, Nogita Kazuhiro. Growth orientations and mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5 on poly-crystalline Cu. J Alloys Compd 2012;536:38–46. [3] Tsao LC. An investigation of microstructure and mechanical properties of novel Sn–3.5Ag–0.5Cu–xTiO2 composite solders as functions of alloy composition and cooling rate. Mater Sci Eng A 2011;529:41–8. [4] Chuang CL, Tsao LC, Lin HK, Feng LP. Effects of small amount of active Ti element additions on microstructure and property of Sn–3.5Ag–0.5Cu solder. Mater Sci Eng A 2012;558:478–84. [5] El-Daly AA, Hammad AE, Fawzy A, Nasrallh DA. Microstructure, mechanical properties, and deformation behavior of Sn–1.0Ag–0.5Cu solder after Ni and Sb additions. Mater Des 2013;43:40–9.
[6] Shnawah DA, Sabri MFM, Badruddin IA, Said S. A review on effect of minor alloying elements on thermal cycling and drop impact reliability of low-Ag Sn– Ag–Cu solder joints. Microelectron Int 2012;29:47–57. [7] Chen WM, Kang SK, Kao CR. Effects of Ti addition to Sn–Ag and Sn–Cu solders. J Alloys Compd 2012;520:244–9. [8] Shnawah DA, Said SBM, Sabri MFM, Badruddin IA, Che FX. Microstructure, mechanical, and thermal properties of the Sn–1Ag–0.5Cu solder alloy bearing Fe for electronics applications. Mater Sci Eng A 2012;551:160–8. [9] Hammad AE. Investigation of microstructure and mechanical properties of novel Sn–0.5Ag–0.7Cu solders containing small amount of Ni. Mater Des 2013;50:108–16. [10] Niranjani VL, Chandra Rao BSS, Sarkar R, Kamat SV. The influence of addition of nanosized molybdenum and nickel particles on creep behavior of Sn–Ag lead free solder alloy. J Alloy Compd 2012;542:136–41. [11] Gain AK, Chan YC. The influence of a small amount of Al and Ni nano-particles on the microstructure, kinetics and hardness of Sn–Ag–Cu solder on OSP-Cu pads. Intermetallics 2012;29:48–55. [12] Sona M, Prabhu KN. Review on microstructure evolution in Sn–Ag–Cu solders and its effect on mechanical integrity of solder joints. J Mater Sci: Mater Electron 2013;24:3149–69. [13] Mu D, Yasuda H, Huang H, Nogita K. Growth orientations and mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5. J Alloys Compd 2012;536:38–46. [14] Gao F, Takemoto T, Nishikawa H. Effects of Co and Ni addition on reactive diffusion between Sn–3.5Ag solder and Cu during soldering and annealing. Mater Sci Eng A 2006;420:39–46. [15] Nogita K, Nishimura T. Nickel-stabilized hexagonal (Cu,Ni)6Sn5 in Sn–Cu–Ni lead-free solder alloys. Scripta Mater 2008;59:191–4. [16] Belyakov SA, Gourlay CM. NiSn4 formation during the solidification of Sn–Ni alloys. Intermetallics 2012;25:48–59. [17] El-Daly AA, Hammad AE. Enhancement of creep resistance and thermal behavior of eutectic Sn–Cu lead-free solder alloy by Ag and In-additions. Matrer Des 2012;40:292–8. [18] Tsao LC, Chang SY. Effects of Nano-TiO2 additions on thermal analysis, microstructure and tensile properties of Sn–3.5Ag–0.25Cu solder. Mater Des 2010;31:990–3. [19] Chinnam RK, Fauteux C, Neuenschwander J, Janczak-Rusch J. Evolution of the microstructure of Sn–Ag–Cu solder joints exposed to ultrasonic waves during solidification. Acta Mater 2011;59:1474–81. [20] Wang M, Wang J, Feng H, Ke W. In-situ observation of fracture behavior of Sn– 3.0Ag–0.5Cu lead-free solder during three-point bending tests in ESEM. Mater Sci Eng A 2012;558:649–55. [21] Tay SL, Haseeb ASMA, Johan MR, Munroe PR, Quadir MZ. Influence of Ni nanoparticle on the morphology and growth of interfacial intermetallic compounds between Sn–3.8Ag–0.7Cu lead-free solder and copper substrate. Intermetallics 2013;33:8–15. [22] El-Daly AA, El-Taher AM. Improved strength of Ni and Zn-doped Sn–2.0Ag– 0.5Cu lead-free solder alloys under controlled processing parameters. Mater Des 2013;47:607–14. [23] El-Daly AA, Hammad AE. Development of high strength Sn–0.7Cu solders with the addition of small amount of Ag and In. J Alloys Compd 2011;509:8554–60. [24] El-Daly AA, Fawzy A, Mohamad AZ, El-Taher AM. Microstructural evolution and tensile properties of Sn–5Sb solder alloy containing small amount of Ag and Cu. J Alloys Compd 2011;509:4574–82. [25] Witkin D. Creep Behavior of Bi-Containing Lead-Free Solder Alloys. J Electron Mater 2012;41(2). pp. 190–03. [26] Geranmayeh AR, Nayyeri G, Mahmudi R. Microstructure and impression creep behavior of lead-free Sn–5Sb solder alloy containing Bi and Ag. Mater Sci Eng A 2012;547:110–9.
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Enhanced ductility and mechanical strength of Ni-doped Sn–3.0Ag–0.5Cu lead-free solders A.A. El-Daly a, A.M. El-Taher a,⇑, T.R. Dalloul b a b
Physics Department, Faculty of Science, Zagazig Univ., Zagazig, Egypt Physics Department, Faculty of Science, Islamic Univ. of Gaza, Gaza, Palestine
a r t i c l e
i n f o
Article history: Received 25 July 2013 Accepted 2 October 2013 Available online 16 October 2013 Keywords: Lead free solder alloys Microstructure Mechanical properties
a b s t r a c t In this paper, the tensile tests were conducted to investigate the effect of adding a small amount of Ni on the microstructure, thermal and mechanical properties of 3.0Ag–0.5Cu(SAC 305) solder. The results indicated that addition of Ni can effectively decrease both the undercooling and the onset melting temperature of SAC(305) solder alloy. The strength and ductility of the SAC(305) solder depend significantly on Ni content. In general, the SAC(305)–0.5%Ni solder reveals superior mechanical properties in terms of maximum strength and ductility when compared to the high Ni-content or plain solders. Microstructure analysis revealed that a new g-(Cu,Ni)6Sn5 intermetallic compound (IMC) phase containing large amount of Ni was generated, while the initial Cu6Sn5 phase was converted into (Cu,Ni)6Sn5 phase after 0.5%Ni addition. Besides, the fine fiber-like Ag3Sn and finer dot-shaped precipitates rather than needle-like morphology have occurred at the surface of b-Sn matrix easily, which could provide more obstacles for dislocation pile up in the adjacent grains and enhanced the mechanical property. With increasing Ni addition, the Ni-doped SAC(305) solder showed a corresponding deterioration in their mechanical property due to the coarsening of (Cu,Ni)6Sn5 IMCs and increasing the inter-particle spacing of Ag3Sn IMCs in the eutectic colony. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Owing to considerations of environmental protection and human health, the adoption of lead-free solder has become an inevitable trend in the electronics industry. The lead-free solders based on Sn–Ag–Cu (SAC) system with Ni addition have become popular because of their superior soldering properties, as well as their comparatively low cost [1,2]. In view of that, the SAC solders are commonly used in surface mount technology (SMT) assembly for microelectronics into industrial production [3]. However, many issues with the SAC solders still remain unresolved, such as the best composition, the large undercooling in solidification, and the formation of large brittle IMC. Especially, alloys with high Ag content exhibit the formation of large Ag3Sn IMCs and short creeprupture life time in service [4]. For these reasons, efforts have been made to develop solders with low Ag and Cu contents. Reducing the Ag and Cu contents of the SAC alloys, such as Sn–3.0Ag– 0.5Cu (SAC305) alloy gives rise to more primary b-Sn phase (large b-Sn grains) and decreases the number of Ag3Sn IMC particles. It is therefore expected to result in lower elastic modulus and yield ⇑ Corresponding author. Tel.: +20 552327173/552303252 (O); fax: +20 552308213. E-mail addresses: [email protected], [email protected] (A.M. El-Taher). 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2013.10.009
strength than the high Ag content SAC alloys. It has been established also that the minor element additions such as: Ti, Zn, Al, Ni, Bi, In, Fe, Co, and rare-earth elements are a simple and effective way for modifying the properties of these lead-free solders [5–8]. One of the highly focused areas of research is the effect of Ni content on the microstructure and mechanical behavior of SAC solder alloys due to reliability concerns. The main concern is the existence of the (Cu,Ni)6Sn5 IMC particles, which possess high stability and inherently compatible geometry to the surrounding microstructure, i.e. bulk solder and substrate. The (Cu,Ni)6Sn5 appears at the bulk solder alloy when the Ni content is 0.05 mass% [9]. However, most of the experiments reported in the recent literatures focused on the reliability of Sn–Ag and SAC alloys with adding small amount of Ni. Except for that, attempts of adding different amounts of Ni (0.5–5.0 mass%) to develop solder joints with high reliability have been performed [10–12]. High Ni content is desired for thermal cycling performance because of the strengthening effect and stability of (Cu,Ni)6Sn5 at different temperatures. Previously, it was reported that the addition of 5 mass% Ni could completely prevent the formation of Cu3Sn in SAC alloys [13]. It has been suggested that the incorporation of Ni in Cu6Sn5 makes this phase more thermodynamically stable [2]. Gao et al. [14] found that the interaction parameter between Sn and Ni is much more negative than that between Sn and Cu. This implies that (Cu,Ni)6Sn5 is more stable than Cu6Sn5. Thus, in the presence of Ni atoms,
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Table 1 Chemical composition of the solders studied (wt.%). Alloy
Cu
Ag
Ni
Sn
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
0.5 0.5 0.5
3.0 3.0 3.0
– 0.5 1.0
Bal. Bal. Bal.
the growth of (Cu,Ni)6Sn5 is more favorable and hence faster. Consequently, the crack propagation in the (Cu,Ni)6Sn5 layer becomes more difficult [15]. Belyakov and Gourlay [16] studied the growth of NiSn4 during the solidification of Sn–Ni alloys containing 0– 0.45 wt.% Ni. They found that Primary NiSn4 crystal growth was promoted by high cooling rates and the Sn–NiSn4 eutectic could grew under all solidification conditions used. Although nickel has a strong influence on solderability and deformation behavior of solders, the detailed mechanism is not fully understood. Even for the ternary SAC system, there is very limited knowledge about the influence of Ni element on the properties of SAC(305) alloy. This paper focuses on the effect of Ni additions on the microstructure, thermal behavior and corresponding mechanical properties of SAC(305) alloy during high temperature and strain rate range. The key factors that affect the thermal and mechanical behavior of the solder alloy after Ni addition are discussed. 2. Experimental procedures The nominal compositions for both plain and Ni-doped solders are given in Table 1. Their compositions were Sn–3.0Ag–0.5 wt.% Cu, Sn–3.0Ag–0.5Cu–0.5 wt.% Ni and Sn–3.0Ag–0.5Cu–1.0 wt.% Ni (which are hereafter called SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni, respectively). The lead-free SAC(305)–xNi (x = 0– 1.0 wt.%) solders were prepared from pure Sn(99.99%), Ag(99.999%), Cu(99.98%) and Ni(99.99%) as raw materials. The process of melting was carried out in a vacuum arc furnace under high purity argon atmosphere to produce rod-like specimen with a diameter of approximately10 mm. The melt was held at 600 °C for 2 h to complete the dissolution of Sn, Cu, Ag and Ni and then poured in a steel mold to prepare the chill cast ingot. A cooling rate of 6–8 °C/s was achieved, so as to create the fine microstructure typically found in small solder joints in microelectronic packages. The microstructure was examined by scanning electron microscopy (SEM) JSM-5410, Japan. A solution of 3%HCl, 2%HNO3 and 95% (vol.%) Ethyl alcohol was prepared and used to etch the samples. Phase identification was based on Energy Dispersive X-ray Spectrometry (EDS). Because of different approaches in the
Table 2 Comparison of solidus temperature (Tonset) and liquidus temperature (Tend) for SAC305, SAC–0.5Ni and SAC–1.0Ni solder alloys from heating curve. Alloy
(Tonset) (°C)
Tend (°C)
Pasty range (Tend-Tonset) (°C)
Peak temperature (°C)
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
219.9 218.7 216.2
221.7 224.4 224.6
1.8 5.7 8.4
221.6 224.4 224.4
Table 3 Comparison of melting temperature (Tm) for heating, solidus temperature (Ts) and under cooling range for various solder alloys. Alloy
Tm (°C)
Ts (°C)
Undercooling (Tm–Ts) (°C)
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
219.9 218.7 216.2
195.7 203.0 201.9
24.2 15.7 14.3
specimen design, test setups and experimental methodology are necessary to investigate solder tensile behavior at different volumes. In the present work, the homogenized cast ingots were then mechanically machined into a wire samples with a gauge length marked 4 102 m for each samples and 2.5 mm diameter. Details are described in [17]. Before testing, the specimen was annealed at 120 °C for 60 min to reduce the residual stress induced in the sample preparation. Subsequently, tensile creep tests are conducted on a 3360 universal material testing system, and GWT 504 high temperature testing system respectively. The tests were carried out at various strain rates ranging from 105 to 103 s1 and constant temperature of 25 °C. Also, the tests were conducted at different temperatures ranging from 25 to 120 °C with a constant strain rate of 8.8 104 s1. The axial strain is measured in accordance with the ASTM:E83-10a, and ASTM D76/D76M-11 standard practices for force verification. Then, the mechanical properties were obtained by averaging testing data. The temperature variation inside the high temperature furnace is maintained within 1.5 °C. 3. Results and discussion 3.1. Thermal analysis For understanding the relationship between melt properties and Ni content, the thermal behavior of solders were analyzed using DSC analysis and the results are demonstrated in Fig. 1 and summarized in Tables 2 and 3. During the heating process, it is noteworthy that the Tonset decreases from 219.9 to 216.2 °C with
Fig. 1. DSC results during heating (endothermal) and cooling (exothermal) of SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solders.
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Eutectic
Ag3Sn
β-Sn
(Cu,Ni)6 Sn5
(d)
(Cu,Ni)6Sn5
Fracture
β-Sn Ag3Sn
β-Sn Fig. 2. Scanning electron micrographs (SEM) of as-cast: (a) SAC(305), (b) SAC(305)–0.5Ni, (c) SAC(305)–1.0Ni solders and (d) EDS analysis of (Cu,Ni)6Sn5 IMC particles.
(a)
(b)
Ag 3 Sn Cu6 Sn 5
(c)
Fig. 3. (a) High-magnification SEM microstructure of SAC (305) and EDS analyses of (b) Ag3Sn, and (c) Cu6Sn5 IMCs.
an increase in the Ni content. The prominent endothermic peak of Ni-free SAC(305) solder appeared at 221.6 °C, which corresponds to the eutectic temperature of the Sn–3.5Ag–0.25Cu ternary system as reported by Tsao and Chang [18]. However, the raising of
endothermic peak from 221.6 °C to 224.4 °C due to Ni addition can be attributed to the high melting temperature of Ni, which dissolved with a low solubility in the lead-free SAC(305) solder. Nevertheless, the pasty range of SAC(305), SAC(305)–0.5Ni and
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Fig. 4. Tensile properties of the solders SAC(305), SAC(305)–0.5Ni and SAC(305)– 1.0Ni solder alloys at T = 25 °C and e_ = 2 103 s1.
Table 4 Tensile properties of SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys. Alloy
UTS (MPa)
0.2%YS (MPa)
Elongation (%)
SAC(305) SAC(305)–0.5Ni SAC(305)–1.0Ni
28.0 31.5 27.0
25.0 27.5 23.0
8.2 12.3 6.1
SAC(305)–1.0Ni solders, respectively, was 1.8, 5.7 and 8.4 °C, which is lower than 11.5 °C for Sn–Pb eutectic [17]. This narrow pasty range of SAC(305) solders may avoid manufacturing problems, such as increasing the sensitivity to vibration during wave soldering. Also, it can decrease the probability of fillet lifting phenomena, the tendency towards porosity and hot tearing due to the effect of alloy shrinkage. For these reasons, Ni-containing SAC(305) solders have displayed satisfactory reliability. Undercooling is defined as the difference between the melting point during heating and the solidification temperature during cooling, and relates to the difficulty of nucleating a solid phase in a liquid state [19]. By having a direct effect on nucleation during solidification, the Ni could affect the undercooling of the b-Sn phase, and this may cause significant repercussions on the microstructure of the solidified sample. In Fig. 1 and Table 3, the exothermal peaks upon cooling for both alloy samples appeared at lower temperature compared with their endothermic peaks due to the undercooling, but the reductions in temperature were not identical. It is worthy of notice that the degree of undercooling of SAC(305) solder was strongly decreased from 24.2 to 15.7 and 16.3 °C after Ni addition. This means that the solidification processes of Ni-containing alloys were initiated with the nucleation of new IMC precipitates. The literature showed also that a decrease in b-Sn undercooling resulted in reducing Ag3Sn formation in the SAC alloys, and it can be achieved by the addition of a fourth alloying elements such as Ni, Zn and Sb [17,19]. In this case, Ni could act as an additional nucleation sites, improving the rate of nucleation, leading to a reduced undercooling effect. As a result, one may expect that the microstructure of the Ni-solidified samples does not show any noticeable large Ag3Sn precipitates, and the Ag3Sn growth is probably only limited to small phases in the eutectic matrix due to the combined effects of streaming and enhanced nucleation. These results are quite consistent with the explanation proposed by Chen et al. [7] that primary IMCs might act as
Fig. 5. Effect of strain rate on: ultimate tensile strength (UTS), yield stress (0.2YS) and elongation (El.%) at T = 25 °C for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
heterogeneous nucleation sites for Sn dendrites upon solidification and are able to suppress the undercooling of Sn–Ag–Cu alloys. 3.2. Effect of Ni additions on the solidification microstructures Fig. 2 shows the microstructure of plain SAC(305) and Ni-containing solders in the as-cast condition. In case of Ni-free SAC(305) solder (Fig. 2a), the dark gray b-Sn primary grains with an average diameter of 40–70 lm are surrounded by the light gray area of fine eutectic network (Sn, Ag3Sn and Cu6Sn5) with the average spacing lamellae of 4 lm. These results are very similar to those of the SAC solders [20]. After 0.5 wt.% Ni addition to the
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Fig. 6. Effect of temperature on: ultimate tensile strength (UTS), yield stress (0.2YS) and elongation at strain rate of 8.8 104 s1 for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
SAC(305) solder, a different morphology was observed. The morphology tends to transform from needle-like to fiber-like Ag3Sn and fine dot-shaped precipitates at the surface of b-Sn matrix are observed in Fig. 2b. It can be seen that the average size and spacing of Ag3Sn particles were significantly reduced when compared with the SAC(305) solder. The refined IMC morphology also indicates the increase of nucleation rate and provides the justification for its better performance, which could restrict the dislocation motion. On the other hand, the microstructure could be clear evidence of the presence of new rod-like (Cu,Ni)6Sn5 IMCs induced at a network of eutectic area. All the particles investigated contain a significant amount of Ni with an average content of 17.2–35.7 at.% Ni, as
Fig. 7. Relationship between r and ln(e_) at (a) T = 25, (b) T = 70 and (c) T = 120 °C for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
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Fig. 8. Relationship between ln(r) and ln(e_) at (a) T = 25, (b) T = 70 and (c) T = 110 °C for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
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315
Fig. 9. Relationship between ln[sinh(ar)] and ln(e_) for determination stress exponent (n) values at (a) T = 25, (b) T = 70 and (c) T = 110 °C for SAC(305), SAC(305)0.5Ni and SAC(305)–1.0Ni solder alloys.
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Fig. 10. The activation energy (Q), values of SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys.
Table 5 Activation energy (Q) and stress exponent (n) values for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni solder alloys. Alloy
Q (kJ/mol)
Temperature (°C)
a (MPa1)
n
SAC(305)
58.7
25 70 120
0.03739 0.05375 0.07231
7.0 6.0 5.3
SAC(105)–0.5Ni
62.8
25 70 120
0.04135 0.05543 0.07000
8.4 7.1 6.4
SAC(105)–1.0Ni
50.6
25 70 120
0.43243 0.05500 0.06667
9.3 8.0 7.2
confirmed by EDS analysis shown in Fig. 2d. It implied that the initial Cu6Sn5 phase was converted into g-(Cu,Ni)6Sn5 phase during cooling. The formation of rod-like g-(Cu,Ni)6Sn5 morphology may result from anisotropy of surface energies and compositional related to the presence of Ni, which influences the growth behavior of Cu6Sn5. However, the microstructure changes dramatically after 1.0 wt.% Ni addition. The volume fraction of the Ag3Sn IMC particles in the eutectic area tends to decrease and its interparticle distance tends to increase to the average value of 15 lm. Worthy of notice is that the b-Sn-rich phases became larger, and the average size of plate-like (Cu,Ni)6Sn5 IMCs was increased. Increasing the Ni content also can enhance the crack formation inside (Cu,Ni)6Sn5 IMCs during cooling, as seen in Fig. 2c. This will more dramatically degrade the mechanical properties and lifetimes of SAC(305)–1.0Ni solder joints than the other alloys. Fig. 3a shows the close-up view of plain SAC(305) with the corresponding EDS analysis, where the light needle-like morphology is the Ag3Sn (Fig. 3b) and scallop-like morphology phase is the Cu6Sn5 IMC, as shown in Fig. 3c. 3.3. Tensile tests The representative stress–strain plots of one temperature (RT) and one strain rate (2 103 s1) for the base alloy as well as the solders containing 0. 5 and 1.0 wt.% Ni are shown in Fig. 4. The average values of ultimate tensile strength (UTS), 0.2% yield strength (0.2YS), and percentage of elongation are given in Table 4. The SAC(305) solder had an UTS of 28 MPa, 0.2YS of 25 MPa, and elongation to failure of 8.2%. With addition of 0.5%Ni, a prom-
inent effect on both the strength and ductility of SAC(305) solder had occurred. The UTS and 0.2%YS of SAC(305)–0.5Ni alloy were 31.5 and 27.5 MPa, respectively, which were higher 1.1 times that of the baseline SAC(305). Besides, the total elongation of SAC(305)– 0.5Ni alloy is 12.3%, which was 52.3% higher than that of SAC(305) solder. The enhancement in both the strength and ductility of SAC(305)–0.5Ni solder is mainly attributed to the formation of (Cu,Ni)6Sn5 and the homogeneous distribution of fine and large eutectic mixture, as indicated in Fig. 2. A number of studies are available on the effect of nickel as an alloy addition to the Sn-based solders [21]. It has been suggested that the incorporation of Ni in the (Cu,Ni)6Sn5 IMC makes this phase more thermodynamically stable at room temperature, since Ni was found to retard the transformation of hexagonal structure to monoclinic structure at 186 °C, thereby stabilizing the hexagonal structure at room temperature. Accordingly, the stabilization of hexagonal (Cu,Ni)6Sn5 will suppress the volume change associated with phase transformation and hence will eventually prevent IMC cracking. This is expected to improve the solder performance. Nogita and Nishimura noticed that [15] the lattice structure of (Cu,Ni)6Sn5 is only thermally stable at room temperature as the hexagonal g-Cu6Sn5 when containing 9 at.% Ni, while that containing 2 at.% Ni is not stable and transforms to monoclinic during cooling. This result indicates that the high-temperature hexagonal g-(Cu,Ni)6Sn5 is only stable at room temperature when sufficient Ni is present in the (Cu,Ni)6Sn5 IMC. In view of these, the current results reveal hexagonal g(Cu,Ni)6Sn5 containing 17.2–35.7 at.% Ni, which support and agree well with the previous studies, and predicted that Ni would thermodynamically stabilize (Cu,Ni)6Sn5 down to room temperature. The stabilization of hexagonal (Cu,Ni)6Sn5 in Ni-containing solders could have significant industrial applications and points to future research directions. However, the extra 1.0 mass% Ni additions caused the UTS and 0.2YS to decrease to 27 and 23 MPa, respectively, and elongation was diminished to 6.1%. These values illustrate that under the same testing conditions, the UTS, 0.2YS and elongation of SAC(305) solder are significantly decreased with increasing Ni content. The reason may be that the volume fraction of the Ag3Sn IMC particles in the eutectic area tends to decrease and the interparticle distance tends to increase with the addition of 1 wt.% Ni. However, more dissolved Ni can combined with the bulk Cu6Sn5 particles and accumulate together in the solder matrix inhomogeneously to form coarse (Cu,Ni)6Sn5 IMCs during solidification. Another unusual finding is that the coarse (Cu,Ni)6Sn5 grain surface is fractured slightly due to inconsistent volume contraction at the solder/ (Cu,Ni)6Sn5 interface during cooling, as seen in Fig. 2c. Such structure of bulk (Cu,Ni)6Sn5 may degrade the mechanical properties and fracture strengths of solder joints, leading to poor reliability of electronic devices. 3.4. Effect of strain rate on the mechanical properties Fig. 5 presents the effect of strain rate on UTS, 0.2YS and percentage of elongation of SAC(305)–xNi solder specimens at room temperature. In general, the entire alloys displayed an increase in both UTS and 0.2% YS with increasing strain rate. The mechanical feature improvement was attributed to the interaction between IMCs or precipitates with the dislocation motion. These results are consistent with the previous findings [22] that the tensile behavior of SAC(205)–0.05Ni alloys often exhibits a strong dependence on the strain rate, or high strain rate sensitivity. At low strain rates, the precipitates hinder the fast moving dislocation. With increasing strain rate, the IMCs cannot capture the moving dislocation any more due to the faster velocity of the dislocation in the alloy. Therefore, when the stress is applied to the solder, the dislocations glide on the slip planes which can tolerate the
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deformation. Since the b-Sn-rich phase is the basic ductile phase of the SAC(305)-based solders, the IMCs formed inside the matrix make the solder exhibits high strength. Relevant investigations [22] also demonstrate that the (Cu,Ni)6Sn5 phases in Ni-containing samples have higher elastic modulus and hardness than Cu6Sn5. In consequence, dissolution of 0.5Ni in Cu6Sn5 phase could result in higher mechanical strength of Ni-containing samples as seen in Fig. 5. Compared with plain solder, 1.0 Ni-containing samples exhibited a lower mechanical strength and lower ability for plasticity, which can sustain forced deformation with crushing fracture due to the coarsening of (Cu,Ni)6Sn5 phase. However, examination of these curves reveals also that all the samples have experienced plastic deformation, and the plastic flow is directly related to the increase of strain rate, except at strain rate of 4 104 s1, where the total strain was decreased. One possible reason is the material structure is piecewise discontinuous and heterogeneous owing to the existence of micro-defects or damage. Based on the equivalence principle [23], the micro-defects can be ‘smeared out’ and the stress and strain state can be considered as homogeneous. It seems that the non-uniform microstructure excludes the possibility of microstructural homogeneities, which can result in consistent elongation values. However, no attempt is made to identify the physical nature of the damage parameters and to distinguish between different damage mechanisms. 3.5. Effect of temperature on the mechanical properties To achieve the temperature dependence of deformation behavior of the tested solders, their tensile properties were investigated in the temperatures range of 25–120 °C. Typical stress–strain curves of the alloys obtained at a strain rate of 8.8 104 s1 are shown in Fig. 6a–c. In contrast to the variation of total elongation, which increased with increasing temperature, the UTS and 0.2%YS of three alloys are dramatically decreased with increasing temperature. This is a consequence of dynamic recovery of lead-free solder, which are encouraged at high temperatures. Increasing the deformation temperature will initiate rearrangement of the dense dislocation networks formed by strain hardening into simple and more ordered ones [24]. This reduces the lattice energy, and therefore lowers the values of UTS and YS at higher temperatures. In addition, the dislocations have much more energy and can overcome tiny IMCs particles that existed in the b-Sn-matrix. However, Fig. 6 also reveals that the SAC(305)–0.5Ni solder possess the highest tensile strength values at all temperature range. In contrast, the SAC(305)–1.0Ni alloy exhibits the lowest one, which can be reflected by the microstructural change observed after 1.0Ni addition (Fig. 2c). The SAC(305)–1.0Ni alloy exhibited small amount of Ag3Sn and coarse (Cu,Ni)6Sn5 IMCs in the eutectic colony. In the meantime, the coarsening deteriorates the strengthening effect of the IMC particle, and thus the b-Sn structure is dominant in UTS at higher temperature. Those massive precipitation alloys, must be avoided because they will possible induce crack in the matrix under low stress or some cyclic stress under service conditions. Nevertheless, the ductility as measured by percent elongation has increased with inconsistent behaviors. The best ductility is usually obtained at the intermediate temperature region. This is recognized in many alloys and occurs at temperatures high enough for grain boundary sliding. It depends on many aspects such as; the matrix phase, IMC chemistries, the nature of interface formed between the IMCs and the matrix, the growth rate and the stability of hard IMCs in the alloy matrix [23]. 3.6. Kinetic analysis of SAC(305)–xNi solders during hot deformation The most comprehensive way of analyzing the deformation data incorporates both the stress exponent ‘n’ and activation
energy ‘Q’ of the operating deformation process. The data are analyzed based on the standard hyperbolic sine Eq. (1) [25]. The appeal of the hyperbolic sine analysis is that it is mathematically capable of incorporating both low- and high-stress data in a single model reflecting power-law and power-law breakdown behavior, respectively.
e_ ¼ A½sinhðarÞn expðQ =RTÞ
ð1Þ
where A (s1), a (MPa1), are material constants independent of temperature, a is stress level parameter, R is the universal gas constant, n is the stress exponent constant related to strain rate, T(K) is thermodynamic deformation temperature, r (MPa) is the steady state flow stress and Q (kJ/mol) is the activation energy of deformation. The logarithm was taken and Eq. (1) was rearranged:
ln e_ ¼ ln A þ n ln½sinhðarÞ Q =RT
ð2Þ
The value of a represents the stress reciprocal at which the material deformation changes from power to exponential stress dependence. An approximate value of a was calculated by a = b/ n1, where b and n1 are the average slopes of ln(e_) r and ln(e_) ln(r) lines at constant T, respectively. The plots of ln(e_) r and ln(e_) ln(r) lines at temperature of 25, 70 and 120 °C for all alloys are given in Figs. 7 and 8. The stress exponent values ‘n’ are calculated from the slope of the ln e_ against ln[sinh(ar)] plot for varying temperature (Fig. 9). The activation energy Q can be expressed as the slope of ln(e_) against 1/T (Fig. 10). The results summarized in Table 5 reveals a range of stress exponents of 5.3–9.3 and activation energies, with the predominance of activation energies falling in the range of 50.6–62.8 kJ/mol for dislocation core diffusion, in consistent with Q 40–65 kJ/mole for Sn self-diffusion through dislocation cores [25], and is almost half the value of that for lattice diffusion (approximately 100 kJ/mol) [26]. On the other hand, the decrease in n value with increasing temperature has been related to the instability of the microstructure, which occurs during high temperature deformation [24]. The anisotropy of the Sn unit cell and the differences in microstructural features and scale suggest that wide ranges of values for both n and Q are to be expected. Deformation mechanisms in the particle-strengthened or multiphase eutectic structures of the Pb-free alloys may thus be similar to those observed in pure tin, but operate at higher stress than in pure tin. According to a dispersion strengthening mechanism, the refinement of the Sn-rich phases, the refinement of eutectic area and uniformly dispersed IMC particles have positive effects on the enhanced mechanical performance of the SAC(205)–0.05Ni solder. They not only help inhibit the grain boundaries from sliding, but also impede the dislocations movement and raise dislocation densities. In contrast, the enhanced crack formation inside (Cu,Ni)6Sn5 IMCs during cooling and increased the inter-particle spacing inside the eutectic area have the negative effects on the mechanical performance of the SAC(205)–1.0Ni solder. However, by determining A, Q, n and a values for SAC(305), SAC(305)–0.5Ni and SAC(305)–1.0Ni alloys, the creep strain rate, respectively, can be estimated from the constitutive equations represented by:
e0 ¼ 4:16 103 ½sinhð0:06088rÞ6:6 exp e0 ¼ 1:12 103 ½sinhð0:04967rÞ7:5 exp e0 ¼ 4:16 103 ½sinhð0:08078rÞ5:6 exp
58:7 RT
62:8 RT
50:6 RT
ð3Þ
ð4Þ
ð5Þ
This model fits the experimental tensile curves of SAC(305) alloys well. Hence, the steady-state strain rate and some temperatures can be accurately predicted for both solders by Eqs. (3)–(5).
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4. Conclusion The dependencies of UTS, 0.2YS and elongation to fracture on representative microstructural parameters were shown to be associated with the morphological changes of the Ag3Sn IMC developed along the SAC(305) alloys as well as the formation of new (Cu,Ni)6Sn5 particles in the SAC(305) solder due to Ni addition. Following conclusions can be made based on the results: (1) Addition of 0.5Ni not only refined the Ag3Sn IMC phase and decreased interparticle spacing, but also promoted the formation of new (Cu,Ni)6Sn5 particles in the SAC(305) solder, which could provide more obstacles for dislocation pile up in the adjacent grains. With addition of 1.0Ni, the volume fraction of Ag3Sn IMC particles in the eutectic area decreased and its interparticle distance increased. Besides, the average size of plate-like (Cu,Ni)6Sn5 IMCs was increased. (2) Additions of Ni into the lead-free SAC solder caused its undercooling and solidus temperature to decrease by about 7.9–8.5 °C and 1.2–3.7 °C, respectively. (3) Addition of 0.5Ni to SAC(305) alloy resulted in an increase in YS and UTS as well as the ductility at all temperatures and strain rates, while the benefits of mechanical behavior and ductility were reduced when the concentration of Ni exceeded 0.5 wt.%. (4) According to the obtained stress exponents and activation energies, it is proposed that the dominant deformation mechanism in SAC(305) solders is dislocation core diffusion over the whole temperature range investigated.
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