Investigating the mechanical properties, creep and crack pattern of Cu6Sn5 and (Cu,Ni)6Sn5 on diverse crystal planes

Materials Science & Engineering A 566 (2013) 126–133

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Investigating the mechanical properties, creep and crack pattern of Cu6Sn5 and (Cu,Ni)6Sn5 on diverse crystal planes D. Mu a,b, H. Huang b,1, S.D. McDonald a,b, J. Read a,b, K. Nogita a,b,n a b

Nihon Superior Centre for the Manufacture of Electronic Materials, The University of Queensland, Brisbane, QLD 4072, Australia School of Mechanical and Mining Engineering, The University of Queensland, Brisbane, QLD 4072, Australia

a r t i c l e i n f o

a b s t r a c t

Article history: Received 5 September 2012 Received in revised form 19 December 2012 Accepted 22 December 2012 Available online 11 January 2013

Cu6Sn5 is an important intermetallic compound (IMC) commonly formed during lead-free soldering. It is known that Cu6Sn5 exhibits significantly different thermo-mechanical deformation behaviour compared to both bulk solder alloys and their substrates. In high-density 3-D electrical packages individual solder joints may contain only a few grains of Cu6Sn5. The knowledge of the mechanical properties, creep and crack behaviour of Cu6Sn5 on different crystal planes is therefore essential to understanding the deformation of lead-free solder joints in service. In this research, the mechanical properties, creep and crack patterns on diverse crystal planes of hexagonal Cu6Sn5 and (Cu,Ni)6Sn5 were investigated using electron back scattered diffraction (EBSD), scanning electron microscopy (SEM) and nanoindentation. It was found that the mechanical properties, creep and crack patterns of hexagonal Cu6Sn5 were strongly related to the crystal orientation. The addition of Ni was found to reduce the anisotropy in hardness and the creep of Cu6Sn5 and had a significant effect on the crack patterns of Cu6Sn5. & 2013 Elsevier B.V. All rights reserved.

Keywords: EBSD Nanoindentation Mechanical properties Orientation relationship Anisotropy Soldering

1. Introduction It is well known that the intermetallic compound (IMC) layers formed at the solder/substrate interface play an important role in the crack formation and deformation of a solder joint [1,2]. The recent progress of 3-D packaging technologies of modern electronic products has led to an increased volume fraction of IMCs in solder joints, often with only a few grains of the IMC in each individual joint [3,4]. Furthermore, in some 3D integrated circuits (ICs), the electronic joint consists entirely of IMCs; comprised mostly of Cu6Sn5 with some Cu3Sn. Cu6Sn5 is perhaps the most important IMC in electronic packaging and commonly forms during interface reactions between most Sn-based solders and Cu substrates [3,5]. As Cu6Sn5 exhibits a strong texture relationship to both singleand poly-crystalline Cu substrates [6–13], a knowledge of the relationship of the mechanical properties and the crystal orientation of Cu6Sn5 is essential in gaining an understanding of the deformation behaviour and failure mechanisms of lead-free solder joints. This is particularly true for the 3D ICs, where joint

n Corresponding author at: The University of Queensland, Nihon Superior Centre for the Manufacture of Electronic Materials, Brisbane, QLD 4072, Australia. Tel.: þ61 7 3365 3919; fax: þ 61 7 3365 3888. E-mail address: [email protected] (K. Nogita). 1 Tel.: þ61 7 3365 3583; fax: þ61 7 3365 4799.

0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.12.057

reliability is the most challenging problem for 3D packaging technology. Intensive research has been conducted to investigate the mechanical properties of Cu6Sn5 formed in diffusion couples [14], at the solder-substrate interface [15,16] and in bulk IMCs formed by either solidification or diffusion [17]; and these previous experimental results are in good agreement with theoretical computations [18]. Trace element additions can have significant effects on the properties of IMCs in lead-free solder alloys. For instance, there is a large variation in the mechanical properties of (Cu,Ni)6Sn5 formed between Sn–0.7Cu–0.05Ni solder and polycrystalline Cu substrates compared with Cu6Sn5 between Sn–0.7Cu and Sn–3Ag–0.5Cu solders and polycrystalline Cu substrates [15]. This variation is due to the anisotropic mechanical properties combined with differing crystal orientations of the Cu6Sn5 and (Cu,Ni)6Sn5 IMCs, rather than inhomogeneous chemical composition; as micro-XRF studies have confirmed the homogeneity of the Sn, Cu and Ni ratio in (Cu,Ni)6Sn5 [19]. At equilibrium, Cu6Sn5 has a monoclinic structure at temperatures below 186 1C and a hexagonal structure at temperatures above 186 1C [20]. A new long-ordered hexagonal-like monoclinic structure has also been identified for stoichiometric Cu6Sn5 produced under specific conditions [21]. Recent research has focused on the relationship between crystal orientation and the mechanical properties of Cu6Sn5. However, there is variation in the literature regarding which crystal planes have the highest

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elastic modulus and hardness. In a previous study [22], using nanoindentation and electron backscattered diffraction (EBSD) it was found that the higher elastic modulus and hardness occurred on the (0 0 1) plane of directionally solidified hexagonal Cu6Sn5. This result is in agreement with a study where the mechanical properties were measured using nanoindentation of Cu6Sn5 micropillars prepared by using an ion-lathe technique [23]; and first-principle theoretical computations of hexagonal-like monoclinic Cu6Sn5 [24,25]. This is in contrast to research which found that the higher elastic modulus and hardness of Cu6Sn5 grown directionally by liquid-electromigration technique was on the (1 1 0) plane [26]. The creep characteristics of Cu6Sn5 at room [27] and elevated temperatures [28] have been investigated using nanoindentation. It was found that an increase in temperature resulted in a remarkable decrease of the creep exponents of Cu6Sn5, due to its high homologous temperature at 125 1C and 150 1C. Ghosh investigated crack formation of Cu6Sn5 using nanoindentation testing with a Berkovich tip [29]; and the crack formation appeared to be influenced by the tip geometry. Zhang et al. studied the crack modes of Cu6Sn5 at the solder/substrate interface using insitu micro-hardness testing and found that the cracks in Cu6Sn5 formed underneath the indenter and at the interface between the Cu6Sn5 and the Cu substrate [30]. Jiang and Chawla also suggested that the crack formation was closely related to the crystal orientation by using nanoindentation on a Cu6Sn5 micro-pillar [31]. However, none of the existing literature gives direct observation on the creep and crack patterns on specific crystal planes of Cu6Sn5. The mechanical properties of Cu6Sn5 can be greatly influenced by the alloying additions to the base solder. Ni has proven to be an important alloying element in Sn-based lead-free solders and the addition of Ni modifies solidification microstructures, increases the volume fraction of the eutectic phase and reduces cracks in interfacial IMCs during service [32,33]. Nogita et al. also found that the addition of Ni stabilizes hexagonal Cu6Sn5 down to room temperature [34,35]. Furthermore, the solubility of Ni in Cu6Sn5 has been found to reduce its thermal expansion [36], modify the growth texture [12,13], increase the hardness and elastic modulus [15–17] and reduce the anisotropy in hardness of directionally solidified Cu6Sn5 [22]. Ni also has variable effects on the creep of Cu6Sn5 at room and elevated temperatures [28]; which is due to the increase of total energy in (Cu,Ni)6Sn5 [37] and the stabilization effect of Ni [34,35]. However, the effect of Ni addition on the anisotropy in mechanical properties, especially on the creep and crack patterns of Cu6Sn5, is not available in the literature yet. This paper addressed this knowledge gap by systematically investigating the mechanical properties, creep and crack patterns on different crystal planes of directionally solidified Cu6Sn5 and (Cu,Ni)6Sn5 using EBSD and nanoindentation techniques. The effect of these properties on the solder micro-joint reliability and the crack formation of Cu6Sn5 with and without Ni between the Cu substrate and lead-free Sn based solder alloys is discussed.

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Fig. 1. Schematic of sections taken through directionally solidified samples and the Cu6Sn5 crystal planes that are made available for analysis by the section.

The directionally solidified samples were sectioned at the angles of normal, 451 and parallel to the solidification direction, as shown in Fig. 1, and then polished using the traditional techniques for metallographic observation. Crystal orientation of Cu6Sn5 and (Cu,Ni)6Sn5 was detected using a JEOL 6460 (JEOL, Tokyo, Japan) SEM with the electron back scattered diffraction (EBSD) function. Subsequent to this, the mechanical properties on different crystal planes were measured using a Triboindenter (Hysitron, MN, USA). A Berkovich indenter of tip radius of 100 nm was used. The indentation load of 2000 mN was applied at a loading rate of 200 mN/s; the dwell time at peak load was 10 s and the unloading rate was 133.33 mN/s. For creep measurement, the loading and unloading rate remained the same as the indenting process described earlier, but the dwell time used was 60 s at the indentation load of 2000 mN. For detailed information on the determination of mechanical properties and creep, please refer to our previous study [28]. To investigate the cracking behaviour of Cu6Sn5, the nanoindentation tests were further performed using a conical tip of 20 mm in radius that generates an axisymmetric stress distribution. The maximum load for indenting was varied from 20 mN to 70 mN with a step increment of 5 mN. The indentation impressions were then examined using SEM.

3. Experimental results 3.1. Microscopic characterization 2. Experimental procedure Cu6Sn5 and (Cu,Ni)6Sn5 were prepared using Sn–4Cu–0/0.05/ 0.3Ni alloys supplied by Nihon Superior Co. Ltd. (Osaka, Japan). The solder alloys were melted at 470 1C, and drawn into a Pyrex tube by application of a mild vacuum. Directional solidification was conducted at 440 1C at a speed of 1 mm/min. Detailed information of the solidification conditions can be found elsewhere [38].

Parallel and transverse optical images of Sn–4Cu–0/0.05/0.3Ni alloys after directional solidification are shown in Fig. 2. The quenched liquid/eutectic interface was readily observed, as was the growth direction of the Cu6Sn5 and (Cu,Ni)6Sn5 IMCs. In Fig. 3, representative secondary electron SEM images associated with the EBSD identification of crystal planes are presented. On each grain, several spot analyses were conducted to ensure the accuracy of measurements.

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Fig. 2. Optical micrographs of Sn–4Cu–0/0.05/0.3Ni alloys after directional solidification.

From Figs. 2 and 3, it can be seen that both hexagonal Cu6Sn5 and (Cu,Ni)6Sn5 were directionally solidified with the normal to the (0 0 1) plane (the c axis) parallel to the thermal gradient. This could be due to the (0 0 1) plane of hexagonal Cu6Sn5 and (Cu,Ni)6Sn5 having a lower surface energy than other crystal planes. Moreover, because the hexagonal-like monoclinic phase was only found for stoichiometric Cu6Sn5 [21], the indexed diffraction pattern of EBSD also indicated that the Cu6Sn5 remained as the hexagonal crystal structure after directional solidification. Using energy-dispersive X-ray spectroscopy (EDS) spot analysis the, Ni contents in (Cu,Ni)6Sn5 from the directionally solidified Sn–4Cu–0.05Ni and Sn–4Cu–0.3Ni were found to be close to 1 at% and 5 at%, respectively. Previous EDS and micro-XRF mapping studies have shown the Ni was homogeneously distributed in Cu6Sn5 [19,28,39].

interesting to note that the increase of Ni content from 1 at% to 5 at% in the Cu6Sn5 directionally solidified from Sn–4Cu–0.3Ni alloy does not increase the indentation modulus on the (1 1 0) and (1 0 0) crystal planes, but preferentially increases the indentation modulus on the (0 0 1) plane. The highest hardness was also found to be on the (0 0 1) crystal plane of Cu6Sn5. The 1 at% Ni in the Cu6Sn5 directionally solidified from the Sn–4Cu–0.05Ni alloy was found to increase the hardness on all crystal planes and reduce the anisotropy of hardness of Cu6Sn5. After the 5 at% Ni addition in the Cu6Sn5 directionally solidified from the Sn–4Cu–0.3Ni alloy, the hardness was further increased and became almost isotropic. A comparison of the results shows that there is no significant difference in the mechanical properties on the (1 1 0) and (1 0 0) planes, as well as the (1 1 1) and (1 0 1) planes of Cu6Sn5 and (Cu,Ni)6Sn5.

3.2. Indentation modulus and hardness

3.3. Creep on different crystal planes

A representative load–displacement curve during the indentation is shown in Fig. 4. The creep phenomenon can be observed within the holding period. Vlassak and Nix have indicated that the indentation modulus was typically different from the elastic modulus in the direction of indentation and could be best approximated by a weighted average of the elastic constants [40]. Because the elastic constants of Cu6Sn5 are not available yet, the indentation modulus was used to represent the overall trends of elastic deformation behaviour of Cu6Sn5 on diverse crystal planes in this study. This treatment might introduce some uncertainty but would largely facilitate the analysis of elastic deformation in this hexagonal system, where the elastic constants are unknown. The indentation modulus and hardness measured on (0 0 1), (1 1 0), (1 0 0), (1 0 1) and (1 1 1) planes of Cu6Sn5 and (Cu,Ni)6Sn5 are presented in Table 1. These results are in agreement with the previous results showing that the (0 0 1) plane has a higher indentation modulus than other crystal planes [22,23] and Ni increases the indentation modulus [15–17]. In this study, the 1 at% Ni present in the Cu6Sn5 directionally solidified from the Sn–4Cu–0.05Ni alloy increases the indentation modulus on both the (0 0 1) and (1 1 0) crystal planes but does not significantly affect the anisotropy in indentation modulus of Cu6Sn5. It is

In Fig. 5, the representative creep displacements during dwell time are shown. It can be seen that the steady creep stage has been reached after around 10 s of dwell time. The rupture stage was not reached because of the relatively small indentation load of 2 mN and the short dwell time of 60 s. For the Ni-free Cu6Sn5, as shown in Fig. 5(a), the largest initial indentation depth and creep displacement during dwell time are on the (1 1 0) and (1 0 0) crystal planes; while the smallest initial indentation depth and creep displacement is on the (0 0 1) crystal plane of Cu6Sn5. On the (1 0 1) and (1 1 1) crystal planes, the initial indentation depth and creep displacements are slightly higher than on the (0 0 1) planes, but lower than that on the (1 1 0) and (1 0 0) planes. For the Cu6Sn5 solidified from Sn–4Cu–0.05Ni alloy, the creep displacements, as well as the indentation depth at the beginning of dwell time, were reduced as shown in Fig. 5(b). The greatest creep displacement occurred on the (1 1 0) and (1 0 0) planes; however, the difference between creep displacements on each crystal plane was clearly reduced. For the Cu6Sn5 solidified from Sn–4Cu–0.3Ni alloy, as shown in Fig. 5(c) the initial indentation depth and creep displacements on different crystal planes are nearly isotropic compared to the Ni-free Cu6Sn5.

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Fig. 3. SEM and EBSD Kikuchi pattern along with a schematic representation of the Cu6Sn5 orientation in Sn–4Cu–0/0.05/0.3Ni alloys after directional solidification.

The creep stress exponents of Cu6Sn5 and (Cu,Ni)6Sn5 on different crystal planes were determined using the method described in [26] and presented in Table 1. It was found that the lowest creep exponents were on the (1 1 0) and (1 0 0) crystal planes of Cu6Sn5, and the largest creep exponent was on the (0 0 1) plane. The addition of Ni increased the creep stress exponents on all crystal planes and significantly reduced the difference between the creep exponents on different crystal planes. 3.4. Crack patterns of Cu6Sn5 and (Cu,Ni)6Sn5

Fig. 4. A representative load–dsiplacement curve during indentation test.

The deformation and cracking patterns on the crystal planes parallel to the c-axis of Cu6Sn5 and (Cu,Ni)6Sn5 are shown in Fig. 6. For the Ni-free Cu6Sn5, it was found that the primary cracks initiated at a peak load close to 30 mN are perpendicular to the c-axis, which is approximately parallel to the (0 0 1) plane. At the indentation load of approximately 40 mN, the secondary cracks were observed parallel to the c-axis of Cu6Sn5. As the indentation

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Table 1 Elastic modulus, hardness and creep stress exponents measurements including standard deviations on crystal planes of Cu6Sn5 directionally solidified from Sn–4Cu–0/0.05/0.3Ni alloys. Crystal plane

Sn–4Cu

Elastic modulus (GPa) (0 0 1) 120.75 7 2.12 (1 1 1) 117.237 2.63 (1 0 1) 118.367 4.52 (1 1 0) 106.25 7 2.86 (1 0 0) 109.72 7 3.14 Hardness (GPa) (0 0 1) 6.22 7 0.26 (1 1 1) 6.14 7 0.18 (1 0 1) 6.07 7 0.12 (1 1 0) 5.74 7 0.15 (1 0 0) 5.67 7 0.23 Creep stress exponents (0 0 1) 40.2 7 3.4 (1 1 1) 38.3 7 2.7 (1 0 1) 37.9 7 4.2 (1 1 0) 29.7 7 3.2 (1 0 0) 30.1 7 1.9

Sn–4Cu–0.05Ni

Sn–4Cu–0.3Ni

127.25 74.12 122.85 73.98 121.14 74.71 116.8 72.37 114.63 72.49

136.3372.05 132.60 72.92 135.9772.70 114.08 73.88 118.7872.53

6.88 70.17 6.61 70.12 6.87 70.23 6.83 70.16 6.81 70.15

7.50 70.22 7.51 70.19 7.56 70.32 7.41 70.22 7.53 70.15

42.6 72.9 41.5 72.1 40.3 73.3 37.1 73.6 39.0 73.1

46.9 72.3 45.2 72.8 44.7 73.5 42.3 73.4 43.1 72.8

Fig. 5. Creep displacement during dwell time of Cu6Sn5 directionally solidified from (a) Sn–4Cu, (b) Sn–4Cu–0.05Ni and (c) Sn–4Cu–0.3Ni alloys.

load increased to 70 mN, the Cu6Sn5 completely ruptured and these rupture cracks are predominately perpendicular to the c-axis and close to the (0 0 1) plane of hexagonal Cu6Sn5. For the Cu6Sn5 solidified from Sn–4Cu–0.05Ni alloy, the primary cracks were not observed until the indentation load

increased to 70 mN as shown in Fig. 6. At an indentation load of 70 mN, the observed cracks were at approximately 451 to the caxis, approximately parallel to and close to the (1 1 1) or (1 0 1) crystal planes. For the Cu6Sn5 solidified from Sn–4Cu–0.3Ni alloy, cracks were not observed, even at the maximum indentation load of 70 mN. After Ni addition, the resistance of the intermetallic to cracking appeared to be increased substantially, as did the hardness and the creep stress exponent.

4. Discussion 4.1. Indentation modulus, hardness and creep By careful examining the load–displacement curve in Fig. 4 and the creep displacement curve in Fig. 5, it can be found that the creep reached a steady stage but some small reverse plasticity still remained at the beginning of the unloading process. As a result, the indentation moduli should be considered an upper bound estimate. The indentation moduli and hardness of different crystal planes were plotted against the misorientation angle, defined as the angle between the normal of the crystal plane and the c-axis. From Fig. 7(a), it can be seen that the measured indentation modulus of Cu6Sn5 decreased as the angle away from the normal of a particular crystal plane and the c-axis increased. This is in agreement with the previous experimental results obtained from directionally solidified Cu6Sn5 and Cu6Sn5 micropillars [22,23]. The previous work [22,23] also showed that the Cu6Sn5 exhibited the highest hardness on the (0 0 1) crystal plane, but the lowest on those planes perpendicular to the (0 0 1) crystal plane. In our study, similar trends were found and the hardness on the (0 0 1) plane was higher than on the crystal planes close to (1 1 0), (1 0 0), (1 1 1) and (1 0 1), as shown in Fig. 7(b). In Fig. 7(c), it can be seen that the creep of Cu6Sn5 is also anisotropic with the smallest creep stress exponent values on the planes parallel to the c-axis. Obviously, the mechanical properties and creep are dependent on the crystal orientation of hexagonal Cu6Sn5. In accordance with the ICDD (International Centre for Diffraction Data) number of 047–1575 (for Hexagonal, P63/mmc), the lattice structure of hexagonal Cu6Sn5, was drawn using the TOPAS software and is shown in Fig. 8. The Cu6Sn5 has a c-axis of 5.097 A˚ and a-axis of ˚ with the c/a ratio of 1.21. The inter-plane distances are 4.206 A, 2.546 A˚ between basal (0 0 1) plane and 3.643 A˚ between prism (1 0 0) and (1 1 0) planes. Therefore, when the indenter penetrated the surface normal to the close packed (0 0 1) plane, more atoms in the lattice would be displaced, for the same applied penetration depth, than normal to any other plane. As a result, an indentation on the (0 0 1) crystal plane would generate more inter-atomic stress than on other crystal planes. As the indentation modulus is supposed to be closely related to the elastic modulus, which is usually proportional to the inter-atomic forces in a solid, it is reasonable to attribute the anisotropy of indentation modulus to the difference in the inter-plane distance in hexagonal Cu6Sn5. The anisotropy in hardness could be attributed to the HCP structure of Cu6Sn5, on which slip occurs most readily on the (0 0 1) plane. When indenting on the (1 1 0) and (1 0 0) crystal planes, as the load direction was parallel to the (0 0 1) plane, the indentation thus directly generated shear stress along the (0 0 1) slip plane. As a result, the slip between (0 0 1) planes occurred more easily. As a result, the hardness values measured on the (1 1 0) and (1 0 0) crystal planes are lower than those on other crystal planes. The anisotropic creep of Cu6Sn5 can be explained as the creep is controlled by dislocation slip, because of the relatively low homologous temperature of Cu6Sn5 at room

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Fig. 6. Crack pattern on crystal planes parallel to c-axis of Cu6Sn5 directionally solidified from Sn–4Cu–0/0.05/0.3Ni alloys at 30 mN, 40 mN and 70 mN.

temperature and the high creep stress exponent obtained. During indentation on the (1 1 0) and (1 0 0) planes, dislocation slip will occur more easily on (0 0 1) planes as discussed above and result in higher creep displacement. As shown in Fig. 7(a), the Ni addition increased the indentation modulus but had no significant effect on the anisotropy of indentation modulus. The increase in Ni addition from 1 at% to 5 at% increased the indentation modulus on the (0 0 1), (1 1 1) and (1 0 1) crystal planes, but did not have a remarkable effect on the indentation modulus on the (1 0 0) or (1 1 0) crystal planes. The addition of Ni not only increased the hardness on each crystal plane, but also reduced the anisotropy in hardness on different crystal planes as shown in Fig. 7(b). In Fig. 7(c), the addition of Ni was found to increase the average value of the creep stress exponent and clearly reduce the anisotropy in creep stress exponents on different crystal planes. Hence, Ni demonstrated remarkable effects on the mechanical properties, especially the anisotropy in hardness and creep, of Cu6Sn5. The effect of Ni addition on mechanical properties and creep can be explained based on the Ni-induced distortions in lattice structure present in the (Cu,Ni)6Sn5. According to our previous first principle calculation [37], the Ni replaces the Cu atoms at the Cu2 position of Cu6Sn5, as indicated in Fig. 8. The Ni-induced substitutional defect resulted in lattice shrinkage [36,41] and generated inter-atomic stress around the substitutional Ni atoms in Cu6Sn5. Hence, the indentation modulus of Cu6Sn5 was increased by the Ni addition. Moreover, according to the previous studies using synchrotron X-ray diffraction [41], the incorporation of Ni into the lattice, preferentially reduced the a-axis but had only a minor effect on the c-axis, with additions of up to 10 at% Ni. For Cu6Sn5 with 5 at% Ni addition, the inter-plane distances are 2.545 A˚ between the basal (0 0 1) plane and 3.634 A˚ between prism (1 0 0) and (1 1 0) planes. As stated previously, an

indentation on the (0 0 1) crystal plane would still generate more internal stress compared with on other crystal planes. As a result, the Ni addition increased the indentation modulus on each crystal plane, but had no significant effect on the anisotropy of indentation modulus. The addition of Ni was found to increase the hardness on each crystal plane but also reduce the anisotropy in hardness on different crystal planes. This is likely to be due to Ni-induced solid solution strengthening. As mentioned above, Ni replaced Cu at the Cu2 position and resulted in the formation of an increased stress field around substitutional Ni atoms in Cu6Sn5 [37]. Because the Cu2 atom was at the centre of the Cu6Sn5 unit cell, the Ni-induced increased stress field impeded the movement of neighbouring Cu1 and Sn1 atoms. This in turn impeded the dislocation movements in different crystal directions, resulting in an increase in hardness and a reduction of the (hardness) anisotropy of Cu6Sn5. Because the creep of Cu6Sn5 was controlled by dislocation slip, the Ni-induced stress field in the centre of the Cu6Sn5 unit cell also impeded the dislocation slip in different crystal directions [37]; and hence increased the average value, but reduced the anisotropy in creep stress exponents on different crystal planes. 4.2. Crack patterns of Cu6Sn5 and (Cu,Ni)6Sn5 In Cu6Sn5, the cracks formed were of an open mode. As indicated in the current and previous works [30,31], there was no significant plastic deformation prior to cracking, and hence Cu6Sn5 should be considered as a brittle material. It is commonly accepted that brittle cracking is closely related to the cleavage in a crystal [42]. Also mentioned earlier, Cu6Sn5 has a HCP structure with the (0 0 1) as the slip plane. Thus, during indentation, the stress homogeneously distributed under a

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changed the slip plane from the (0 0 1) plane to close to the (1 0 1) or (1 1 1) planes of Cu6Sn5. 4.3. Implications for lead-free soldering

Fig. 7. Average (a) elastic modulus, (b) hardness and creep stress exponents against misorientation angle between normal of a crystal plane and c-axis of Cu6Sn5 directionally solidified from Sn–4Cu–0/0.05/0.3Ni alloys.

Fig. 8. Lattice structure of hexagonal Cu6Sn5.

conical tip, would place the (0 0 1) plane of Cu6Sn5 under a tensile strain at the edge of the tip. Once the stress increased to a critical threshold, cleavage commenced. This result is consistent with the observation that the Cu6Sn5 at the solder/substrate interface with a (0 0 1) growth texture was prone to crack in the direction parallel to the c-axis [32,33]. The presence of Ni was observed to modify the crack pattern of Cu6Sn5: first by changing the direction of the initial crack from perpendicular to 451 to the c-axis; and second by increasing the crack resistance. The Ni induced solid solution strengthening and the preferential reduction of the lattice parameter along the a-axis of Cu6Sn5 should be responsible for the change in cracking mechanics. In other words, the Ni reduced the deformation under the same indentation load and

For a lead free solder joint in a conventional ball grid array (BGA) arrangement, the intermetallic layer is commonly considered to the most probable location for crack initiation and formation because of its brittleness. However, the mechanical properties may also influence the overall deformation of a solder joint, depending on the applied strain rate and volume fraction of the intermetallic in the solder joint [10]. As mentioned in the introduction, current minimisation of modern electronic devices has led to an increase in the volume fraction of IMCs, including Cu6Sn5, in a lead-free solder joint. This places an increased emphasis on the importance of Cu6Sn5 on the overall deformation of a solder joint. According to the TTT curve [43], the majority of Cu6Sn5 at the solder/substrate interfaces remains in the metastable hexagonal structure immediately after soldering reactions. Although a new monoclinic phase has been identified for the stoichiometric Cu6Sn5 [21] it is still reasonable to discuss the implication of anisotropy in mechanical properties of Cu6Sn5 due to the similarity between the new monoclinic and hexagonal phases. Hence, the anisotropic deformation behaviour of Cu6Sn5, as shown in this research can be used in combination with the knowledge of the strong growth texture of the intermetallics [3,6,7] to design a solder joint with much greater reliability. First of all, the growth texture of Cu6Sn5 at the solder substrate interface should be carefully controlled. As shown in Fig. 9(a), the IMC layer provides the mechanical bonding between substrates and electronic components, and the shear stress resulting from the relative movement between electronic components and the substrate is expected to be the major deformation mode during service of electronic devices. According to our previous studies [32,33], cracks in Cu6Sn5 with a (0 0 1) growth orientation occurred mostly in a direction parallel to the poly-crystalline Cu substrate. Hence, it is desirable to avoid the (0 0 1) growth texture of a Ni-free Cu6Sn5 IMC layer; where an applied shear stress would occur on the easy-to-deform (1 1 0) and (1 0 0) planes and in the direction parallel to the (0 0 1) slip plane of hexagonal Cu6Sn5, as shown in Fig. 9(b). The modification of growth texture could be achieved either by controlling the crystal orientation of the Cu substrate [6–8], manipulating the soldering temperature [9] or adding alloying elements to the solder [12,13]. As the first two approaches would incur much greater processing costs the addition of alloying elements seems to be the most feasible method. As stated previously, the addition of Ni into solder alloys has many beneficial effects on the IMC layer, (Cu,Ni)6Sn5, including stabilisation of the hexagonal phase to room temperature; reduction of thermal expansion (through effectively reducing thermal stress of solder joints); reduced anisotropy of mechanical behaviour [27–30]; and as shown in this research, improved mechanical behaviour including increased creep stress exponents and a lower propensity for crack formation. Furthermore, the Ni induced the (Cu,Ni)6Sn5 to grow in a preferred (1 0 1) direction [12]. Such a growth direction indicated that any applied shear stress would occur on the (1 1 1) and (1 0 1) crystal planes, as shown in Fig. 9(c). As the (1 1 1) and (1 0 1) crystal planes of (Cu,Ni)6Sn5 had greater creep and exhibited more resistance to cracking than the (1 1 0) or (1 0 0) crystal planes of Ni-free Cu6Sn5, as shown in Fig. 9(b) and (c), crack initiation and propagation is more difficult. The addition of Ni has apparently enhanced the desired properties required for solder joint integrity. This benefit could

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Acknowledgements This research was conducted under an international cooperative research program between the University of Queensland (UQ), Australia and Nihon Superior Company, Japan, and was supported by an ARC-Linkage project (ID: LP100200250). The authors would like to thank Dr Y.Q. Wu, Mr R. Irwan, Ms Z.L. Lin and Ms M.Y. Lu for assistance with experimentation. The authors acknowledge the facilities, and the scientific and technical assistance, of the Australian Microscopy & Microanalysis Research Facility at the Centre for Microscopy and Microanalysis, UQ. DM acknowledges the financial support received in the form of an Australian Postgraduate Award (APA).

References

Fig. 9. Schematic diagram of stress distribution and crack pattern in Cu6Sn5 and (Cu,Ni)6Sn5 layers at solder and substrate interface [12,32].

be extrapolated to enhance the integrity of micro-bumps in 3D ICs, which consist entirely of IMCs. By making use of the preferred growth orientation it should be possible to maximise the desired service properties of such a micro-bump.

5. Conclusions In this study, the mechanical properties and creep of Cu6Sn5 and (Cu,Ni)6Sn5 on diverse crystal planes were investigated using EBSD and nanoindentation techniques. Strong anisotropy in indentation modulus, hardness and stress creep exponents of Cu6Sn5 was found. The addition of Ni increased the indentation modulus, hardness and creep stress exponents on different crystal planes. While the anisotropy in indentation modulus of hexagonal Cu6Sn5 was not significantly affected by the Ni addition, the anisotropy in hardness and stress creep exponents of Cu6Sn5 was significantly reduced. The primary crack on crystal planes parallel to c-axis of hexagonal Cu6Sn5 was found to be on the (0 0 1) crystal plane. The addition of small amounts of Ni in Cu6Sn5 resulted in cracks initiated in the directions 451 to the c-axis of Cu6Sn5. The implications of these effects associated with Ni addition were discussed with respect to applications for lead-free soldering.

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