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Growth orientations and mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5 on poly-crystalline Cu
Journal of Alloys and Compounds 536 (2012) 38–46
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Growth orientations and mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5 on poly-crystalline Cu Dekui Mu a,⇑, Hideyuki Yasuda b, Han Huang a, Kazuhiro Nogita a,b a b
School of Mechanical and Mining Engineering, The University of Queensland, Brisbane, Qld 4072, Australia Department of Adaptive Machine Systems, Osaka University, Suita, Osaka 565-0871, Japan
a r t i c l e
i n f o
Article history: Received 12 March 2012 Received in revised form 25 April 2012 Accepted 25 April 2012 Available online 9 May 2012 Keywords: Growth orientation Pole figures Synchrotron radiation Mechanical properties Nanoindentation Intermetallics Crystal growth X-ray diffraction
a b s t r a c t Lead-free solders are important materials in current generation electrical packages, due to the increasingly stringent legislative requirement aimed at reducing the use of lead. The lead-free solders based on the Sn–Cu system with Ni addition have become popular because of their superior soldering properties, as well as their comparatively low cost. This research investigates the effect of Ni addition on the growth morphologies, crystal orientations and mechanical properties of Cu6Sn5 at the interface between hyper-eutectic Sn–Cu high-temperature lead-free solder alloys and Cu substrates, prior to and after aging, by the use of X-ray diffraction (XRD), scanning electron microscopy (SEM) and nanoindentation. The (Cu,Ni)6Sn5 in Sn–Cu–Ni/Cu solder joints showed a more strongly oriented (1 0 1) texture, compared to the Cu6Sn5 in Sn–Cu/Cu solder joints. The Ni-induced (1 0 1) texture contributes to higher and more scattered average values with larger standard deviations in both elastic modulus and hardness for (Cu,Ni)6Sn5. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Due to environmental concerns, the restriction of hazard substance (RoHS) directive has proposed the phase-out of lead in high-temperature soldering by 2014 in the European Union (EU). This has increased the demand for suitable high-temperature lead-free solders. One of the candidate alloy groups for high-temperature soldering is the hyper-eutectic Sn–Cu alloys, which have Cu contents ranging between the eutectic composition of 0.86 mass% [1] and an upper limit of around 7.6 mass%. These hyper-eutectic Sn–Cu alloys have been used for high temperature dip soldering and tinning of copper wire, polyurethane coated wire and component terminations and display minimal copper erosion at temperatures up to 400 °C [2]. During the interface reactions between the Sn–Cu alloy and Cu substrate, a layer of Cu6Sn5 forms. The Cu6Sn5 layer at the interface is commonly recognized to play a critical role for crack formation in a solder joint during service [3,4]. It has been observed that most cracks forming in the Cu6Sn5 layer are nearly parallel to the solder/substrate interface, while few cracks are present in alternate directions [5]. However, it is unclear if this is due to preferential crack paths existing in Cu6Sn5 combined with preferred growth textures, or the location of maximum
⇑ Corresponding author. Tel.: +61 7 3365 1387; fax: +61 7 3365 3888. E-mail addresses: [email protected], [email protected] (D. Mu). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.04.110
stress planes that occur during thermal cycling and associated dimensional changes [6]. To clarify this issue, detailed information on the morphology, crystal orientation and mechanical properties of Cu6Sn5 formed at solder/substrate interface is required. Cu6Sn5 is the most commonly formed intermetallic during leadfree soldering. Extensive research has been conducted to investigate the mechanical [7–10] and thermal properties [11], the formation [12,13], the evolution [14] and the crystal structure [15–17] of Cu6Sn5. Taking the growth orientations into account, several preferred crystal relationships between Cu6Sn5 and single-crystal Cu [18–20] or Ni [21] plates have been observed and these have been attributed to the minimization of distance between Cu atoms in Cu6Sn5 and the Cu or Ni atoms in the plates. For the Cu6Sn5 formed on a non-textured poly-crystalline Cu substrate, the intermetallic orientation has been studied in the systems of Sn–Pb [22], Sn– Ag–Cu [23], Sn–0.7Cu [24] and Sn–3.5Ag [25]. However, there is little information available on the growth morphology and crystal orientation of Cu6Sn5 formed during the interface reaction between Cu and hyper-eutectic high-temperature Sn–Cu solders. In particular, the crystal orientation and its effect on mechanical properties of Cu6Sn5 at the solder/substrate interface have never been discussed. In this research, the morphology, crystal orientation and mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5 intermetallics, formed between hyper-eutectic Sn–3/4/7Cu–0/0.05/0.1/0.3Ni (mass%) alloys and poly-crystalline Cu substrates, are systematically
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investigated using scanning electronic microscopy (SEM), X-ray diffraction (XRD) and nanoindentation. The crystal orientations are then discussed with respect to the pattern of crack formation and propagation, as well as mechanical properties, in Cu6Sn5 and (Cu,Ni)6Sn5 intermetallic layers. 2. Experiment The Cu6Sn5 and (Cu,Ni)6Sn5 formed during the interface reaction between polycrystalline Cu and a variety of lead-free Sn–3/4/7Cu–Ni solders, as shown in Table 1, were investigated. The samples were prepared by dipping the Cu plates (C1220P) of 10 mm 30 mm 0.3 mm with a common flux (Flux B) into the molten solder at a temperature of 340 °C for Sn–3Cu (mass%) alloys, 373 °C for Sn–4Cu alloys and 430 °C for Sn–7Cu alloys for 10 s. Some of the dipped samples were then aged at 120 °C for 1176 h. Selected as-dipped and aged samples were polished perpendicular to the solder interface using conventional techniques for sample preparation. Mechanical properties (elastic modulus and hardness) of the IMCs in cross-sectioned samples were measured using a Hysitron nanoindenter (Hysitron, MN, USA) [26,27]. A Berkovich diamond indenter of tip radius of 100 nm was used and the indentation load of 1000 lN was applied at a loading rate of 200 lN/s and with a dwell time at peak load of 5 s. The unloading rate was 200 lN/s. In order to investigate the morphology and growth orientation from the top view of the intermetallics (solder side), some of the samples, prior to and after aging, were etched in a solution of ortho-nitro-phenol (35 g) and NaOH (50 g) in 1 L of water at 80 °C to remove the solder alloys. The etched samples were then cut into a round plate of 9 mm diameter before XRD investigations. Synchrotron X-ray diffraction (XRD) was performed on the powder diffraction beamline at Australian Synchrotron using a Mythen-II detector. The synchrotron X-ray diffraction data were collected on the prepared samples in the 2h range 10–80° with the accelerate voltage of 20 keV and the wavelengths of 0.825 Å, respectively. The data collection period was 11 min, with 1 min for stabilization and 10 min for data collection. The X-ray peak data for Cu6Sn5 and (Cu,Ni)6Sn5 were analyzed using Search-Match software. As a reference crystallography and atomic coordination, ICDD (International Centre for Diffraction Data) number of 045-1488 (for monoclinic, C2/c) and 047-1575 (for hexagonal, P63/mmc) were used in association with Search-Match. In addition, (1 0 1) XRD pole figures were plotted for all 24 samples using a laboratory Rigaku machine (Rigaku, Tokyo, Japan). The 2h angle for the (1 0 1) plane of hexagonal Cu6Sn5 was calculated according to ICDD number of 047-1575 as 6.88° for Cu-Ka X-ray. For each pole figure, 64 scans were conducted with the rotation angle of 3°. SEM observation was conducted on both cross-sectioned and etched samples using a Philips XL30 machine (FEI, Hillsboro, USA) with an acceleration voltage of 20 kV.
3. Results and discussion 3.1. Growth morphologies During the interface reaction between near-eutectic Sn–Cu solders and Cu substrates, a scallop-type Cu6Sn5 layer firstly forms, followed by the formation of a two-phase Sn + Cu6Sn5 layer [4]. Park and Arróyave [12] simulated the early-stage nucleation of Cu6Sn5 and proposed that the precipitated Cu6Sn5 would grow parallel to the solder/substrate interface. The growth of the Cu6Sn5 layer in the direction perpendicular to the solder/substrate interface starts only after the precipitated Cu6Sn5 grains coalesce. According to the Sn–Cu phase diagram, the Cu3Sn phase may also Table 1 Composition of solders. Sample No.
Composition (mass%)
1 2 3 4 5 6 7 8 9 10 11 12
Sn–3Cu Sn–3Cu–0.05Ni Sn–3Cu–0.1Ni Sn–3Cu–0.3Ni Sn–4Cu Sn–4Cu–0.05Ni Sn–4Cu–0.1Ni Sn–4Cu–0.3Ni Sn–7Cu Sn–7Cu–0.05Ni Sn–7Cu–0.1Ni Sn–7Cu–0.3Ni
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form during the liquid–solid reaction depending on the solder compositions and reaction conditions. During aging, Cu3Sn grows and the total thickness of the intermetallics layer increases with coarsening of the intermetallic structures. Ni additions reportedly suppress the formation and growth of Cu3Sn [28]. As shown in Figs. 1(a) and 2(a), the Cu6Sn5 that forms in the Ni-free Sn–3/4Cu and Cu reaction couples shows a scallop-type morphology. This observation is in agreement with the previous studies on near-eutectic Sn–0.7Cu [3] and eutectic Sn–Pb solders [22], in which the growth of intermetallics was on poly-crystalline substrates. In Fig. 3(a), the Cu6Sn5 formed in the Ni-free Sn–7Cu alloy and Cu reaction couple has a faceted rod-like morphology. This is likely due to the enhanced intermetallic formation at higher dipping temperatures. For the Ni-containing Sn–3/4/7Cu and Cu reaction couples, the morphology of (Cu,Ni)6Sn5 was changed from scallop-type to hexagonal rod-like, as shown in Figs. 1–3(c), (e) and (g). The formation of rod-like (Cu,Ni)6Sn5 may result from anisotropy of surface energies and compositional supercooling related to the presence of Ni, which influences the growth behavior of Cu6Sn5. After aging, the morphologies of Cu6Sn5 and (Cu,Ni)6Sn5 become faceted scallop-type, as shown in Figs. 1–3(b), (d), (f) and (h). At temperatures below 350 °C, Cu3Sn can also be formed according to the Sn–Cu phase diagram. In the previous studies on eutectic Sn–Cu and Sn–Pb solders reacted with Cu substrates, the Cu3Sn phase cannot be observed until a sufficient aging treatment was given [3]. In this study, the formation of Cu3Sn was found to be dependent on both the Cu concentration and resultant reaction temperature. For the as-dipped Ni-free Sn–3Cu and Cu reaction couple, Cu3Sn cannot be detected by SEM. However, for the asdipped Ni-free Sn–4/7Cu and Cu reaction couples, Cu3Sn can be clearly observed as shown in Fig. 4(a). The addition of Ni can suppress the formation and growth of Cu3Sn. Wang et al. reported that the addition of 5 mass% Ni could completely prevent the formation of Cu3Sn in Sn–Ag–Cu alloys [28]. In our study, the formation of Cu3Sn was not observed for any as-dipped Ni-containing Sn–4/ 7Cu and Cu reaction couples. After aging, the Ni-containing samples also have less Cu3Sn compared to these Ni-free samples, as shown in Fig. 4(b) and (d). This can be explained from the kinetic viewpoint that Ni suppresses the formation of Cu3Sn by reducing the driving force, and hence the diffusion flux, of Cu in Cu3Sn [29]. Besides the effects observed in this study, the addition of a small amount of Ni has been found to reduce the thickness of the intermetallic layer [30] and stabilize the hexagonal Cu6Sn5 [16,17], which may effectively reduce the internal stress which results from phase transformations. 3.2. Growth orientations The synchrotron XRD patterns of Cu6Sn5 formed at the Sn–4Cu– (0.05Ni)/Cu interface, prior to and after aging, are shown in Fig. 5(a–c). It can be found that most Cu6Sn5 at the as-dipped Sn–4Cu/Cu interface remained in the hexagonal structure as shown in Fig. 5(a). From the Sn–Cu phase diagram, the Cu6Sn5 has two equilibrium phases: the monoclinic g0 phase stable at temperatures below 186 °C and the hexagonal g phase stable at temperatures above 186 °C [1]. According to the previous studies on the time–temperature–transformation (TTT) curve [31], the dipping time (10 s) and subsequent air cooling time in this study were both insufficiently long for the hexagonal Cu6Sn5 to completely transform into the low-temperature monoclinic phase. The Cu6Sn5 in as-dipped Ni-free Sn–Cu/Cu reaction couples is likely to be a mixture of monoclinic and meta-stable hexagonal phases, which can be confirmed from these small monoclinic peaks at high 2h angles as shown in Fig. 5(a). After aging at 120 °C for 1176 h, most hexagonal Cu6Sn5 transformed into the monoclinic phase as shown in
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Fig. 1. Top view SEM images of Cu6Sn5 and (Cu,Ni)6Sn5 in Sn–3Cu–xNi/Cu couples (x = 0.05, 0.1, 0.3).
Fig. 2. Top view SEM images of Cu6Sn5 and (Cu,Ni)6Sn5 in Sn–4Cu–xNi/Cu couples (x = 0.05, 0.1, 0.3).
Fig. 3. Top view SEM images of Cu6Sn5 and (Cu,Ni)6Sn5 in Sn–7Cu–xNi/Cu couples (x = 0.05, 0.1, 0.3).
Fig. 5(b). In both as-dipped and aged Sn–4Cu/Cu samples, small Cu3Sn peaks can be observed at low 2h angles, which is consistent with SEM observations. It has been reported that the Cu6Sn5 formed at the interface between Sn–Ag solder and Cu substrates has a (0 0 1) growth texture with the c axis perpendicular to the substrate surfaces [25]. In this study, the (0 0 1) texture of Cu6Sn5 at the Sn–4Cu/Cu interface can be confirmed by comparing the diffraction patterns of Sn–4Cu and Sn–4Cu–0.05Ni samples. In Fig. 5(a), Cu6Sn5 had higher X-ray diffraction intensity on (0 0 2) and (0 0 4) crystal planes, which means the c axis of Cu6Sn5 is also perpendicular to the substrate and there is a (0 0 1) growth texture. After Ni addition, the X-ray diffraction intensity on (1 0 1) crystal plane of (Cu,Ni)6Sn5 increased as shown in Fig. 5(c). From Fig. 5(a) and (c), it can be found that both Cu6Sn5 and (Cu,Ni)6Sn5 have the dominant growth on (1 0 2) plane. Therefore, the effect of Ni solubility on the crystal growth of Cu6Sn5 and (Cu,Ni)6Sn5 can be represented by the diffraction intensity ratios between
(1 0 1), (1 0 2) and (2 0 2) planes. The X-ray diffraction intensity ratios between (1 0 1), (1 0 2), and (2 0 2) crystal planes of Cu6Sn5 and (Cu,Ni)6Sn5 in as-dipped Sn–3Cu/Cu and Sn–4Cu(0, 0.05, 0.3Ni)/Cu reaction couples are shown in Fig. 5(d–f). Compared to the Cu6Sn5 formed in Ni-free Sn–3Cu/Cu and Sn–4Cu/Cu reaction couples, the addition of Ni increased the crystal growth on (1 0 1) plane of Cu6Sn5 by a factor of around 10 as shown in Fig. 5(d). In Fig. 5(e), the growth on (2 0 2) plane of Cu6Sn5 was also slightly increased by Ni addition. This observation is in agreement with previous study on near-eutectic Sn–0.7Cu solder and suggests the Ni addition resulted in the promotion of (1 0 1) growth orientation of Cu6Sn5. For a systematic analysis of the effect of Ni on Cu6Sn5 growth orientation, the (1 0 1) pole figures of hexagonal Cu6Sn5 and (Cu,Ni)6Sn5 in all 24 samples were measured by a laboratory XRD machine as shown in Figs. 6–9. In Fig. 6(a), the intensity distribution of Cu6Sn5 formed in Ni-free samples are almost random along the rolling direction (RD) and transverse direction (TD), which
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Fig. 4. Intermetallics at the solder/substrate interface: (a) Sn–4Cu as-dipped (b) Sn–4Cu aged, (c) Sn–4Cu–0.1Ni as-dipped and (d) Sn–47Cu–0.1Ni aged.
Fig. 5. X-ray diffraction results: (a) X-ray diffraction pattern for Sn–4Cu/Cu sample, (b) X-ray diffraction pattern for aged Sn–4Cu/Cu sample, (c) X-ray diffraction pattern for Sn–4Cu–0.05Ni/Cu sample, (d) X-ray diffraction intensity ratio of (1 0 1)/(1 0 2), (e) XRD intensity ratio of (2 0 2)/(1 0 2) and (f) X-ray diffraction intensity ratio of {(1 0 1) + (2 0 2)}/(1 0 2).
indicates that no preferred (1 0 1) crystal orientation exists. As the Cu content increases, the diffraction peaks in the (1 0 1) pole figures become stronger and more scattered, as shown in Fig. 6(a), (c) and (e), which is attributed to increased Cu6Sn5 formation and the presence of the Cu3Sn phase as observed by SEM and synchrotron XRD techniques. After aging, the diffraction intensity becomes stronger and more scattered compared to that of the as-dipped samples as shown in Fig. 6(b), (d) and (f). The greater degree of scatter of diffraction
peaks in the aged samples was due to not only the growth of Cu6Sn5, but also the increased volume of the monoclinic Cu6Sn5 phase, whose diffraction intensity naturally contains many small peaks. In addition, the growth of the Cu3Sn layer in the aged Ni-free Sn–4/7Cu/Cu reaction couples would also contribute to the scattered diffraction intensity. The (1 0 1) pole figures of Cu6Sn5 and (Cu,Ni)6Sn5 formed in Sn– 3/4/7Cu–Ni/Cu interface are shown in Figs. 7–9. While the Ni-free Cu6Sn5 has a (0 0 1) texture [25], the addition of Ni resulted in a
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Fig. 6. Pole figure of Ni-free samples: (a) Sn–3Cu, (b) aged Sn–3Cu, (c) Sn–4Cu, (d) aged Sn–4Cu, (e) Sn–7Cu and (f) aged Sn–7Cu.
Fig. 7. Pole figure of Sn–3Cu–xNi samples: (a) Sn–3Cu, (b) aged Sn–3Cu, (c) Sn–3Cu–0.05Ni, (d) aged Sn–3Cu–0.1Ni, (e) Sn–3Cu–0.1Ni, (f) aged Sn–3Cu–0.1Ni, (g) Sn–3Cu– 0.3Ni and (h) aged Sn–3Cu–0.3Ni.
Fig. 8. Pole figure of Sn–4Cu–xNi samples: (a) Sn–4Cu, (b) aged Sn–4Cu, (c) Sn–4Cu–0.05Ni, (d) aged Sn–4Cu–0.1Ni, (e) Sn–4Cu–0.1Ni, (f) aged Sn–4Cu–0.1Ni, (g) Sn–4Cu– 0.3Ni and (h) aged Sn–4Cu–0.3Ni.
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Fig. 9. Pole figures of Sn–7Cu–xNi samples: (a) Sn–7Cu, (b) aged Sn–7Cu, (c) Sn–7Cu–0.05Ni, (d) aged Sn–7Cu–0.1Ni, (e) Sn–7Cu–0.1Ni, (f) aged Sn–7Cu–0.1Ni, (g) Sn–7Cu– 0.3Ni and (h) aged Sn–7Cu–0.3Ni.
Fig. 10. Schematic diagram of crack propagation.
Fig. 11. Elastic modulus of Cu6Sn5 and (Cu,Ni)6Sn5 in as-dipped and aged samples.
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Fig. 12. Hardness of Cu6Sn5 and (Cu,Ni)6Sn5 in as-dipped and aged samples.
(1 0 1) texture of (Cu,Ni)6Sn5 in all the Ni-containing samples. As the Ni and/or Cu content increased, the (1 0 1) texture of (Cu,Ni)6Sn5 became stronger. The (1 0 1) pole figures showed that the (Cu,Ni)6Sn5 has the (1 0 1) plane parallel to the solder/substrate as shown in Figs. 7–9(c), (e) and (g). The hexagonal (Cu,Ni)6Sn5 therefore has the c-axis in a direction inclined to the normal of the solder/substrate interface. Using lattice parameters measured in the previous experiments [6] (a- and c- axis of hexagonal (Cu,Ni)6Sn5 of 0.42 and 0.51 nm, respectively) it is calculated there is an approximately 45° disorientation between the c-axis of (Cu,Ni)6Sn5 and the normal of the solder/substrate interface. Growth orientation normally results from preferred nucleation and/or growth mechanisms and the grain size of the substrate may influence these mechanisms [32]. In this study, the substrate is a common C122P polycrystalline Cu plate without heat treatments as given in previous studies [18–20]. Thus, the formation of the (1 0 1) texture of (Cu,Ni)6Sn5 in this study should be attributed to preferred growth resulting from the anisotropic surface energies. After random early-stage nucleation, the (Cu,Ni)6Sn5 grains with (1 0 1) planes parallel to the substrate surface have lower surface energies and consume other (Cu,Ni)6Sn5 grains during the ripening process, which results in the formation of a (1 0 1) texture. This explanation could be supported by the synchrotron XRD result that Ni promoted the (1 0 1) plane growth of Cu6Sn5 and our previous study on near-eutectic Sn–0.7Cu alloys [24]. Similar mechanism can also be used to explain the (0 0 1) texture in Cu6Sn5. Li et al. attributed the formation of (0 0 1) to (1 0 1) texture of Cu6Sn5 as temperature increasing from 240 to 280 °C to the difference in surface energy [25]. In this study, the Ni-free Cu6Sn5 has a (0 0 1) growth orientation at a temperature of 340 °C and above, which can be explained as the high Cu contents in solder alloys enhance the (0 0 1) growth of Cu6Sn5 [24]. After aging, the (1 0 1) pole figure of the aged samples displays a stronger (1 0 1) texture in the (Cu,Ni)6Sn5 layer, as shown in Figs. 7–9 (d), (f) and (h). As mentioned before, there is an allotropic transformation from high-temperature hexagonal g-Cu6Sn5 to low-temperature monoclinic g0 -Cu6Sn5 at an equilibrium temperature of 186 °C. According to the TTT curve of Cu6Sn5 [31], this allotropic transformation is complete after aging at 120 °C for 1176 h and the reported (0 0 1) texture of hexagonal Cu6Sn5 may not be observable. In this study, the (Cu,Ni)6Sn5 remains in the stabilized
hexagonal structure after aging and the stronger (1 0 1) texture is due to the continued growth of the stabilized (Cu,Ni)6Sn5. The observation that Ni affects the growth orientation of the interfacial Cu6Sn5 may have implications for failure mechanisms in lead-free solder joints. Firstly, it is known that the Cu6Sn5 formed at the Sn–Cu/Cu interface has a (0 0 1) growth orientation with the c-axis perpendicular to the solder/substrate interface. Moreover, the cracks formed in the Cu6Sn5 layer are prone to be parallel to the solder/substrate interface [5]. Using directional solidification samples, the elastic modulus and hardness on [0 0 1] crystal planes are found to be lower than [1 0 1] crystal planes [33]. Hence, it is reasonable to conclude the Cu6Sn5 at the dipped solder/substrate interface is easier to fracture on [0 0 1] crystal planes. As the crack propagates, the resultant stress is transferred to the adjacent grains because there is little barrier between adjacent grains. For the Cu6Sn5 with a (0 0 1) growth texture, as shown in Fig. 10(a), the stress will be directly applied on the adjacent [0 0 1] planes, on which cracking is easier to initiate. However, the (1 0 1) growth texture of (Cu,Ni)6Sn5 will result in an angle between [0 0 1] planes of adjacent grains, as shown in Fig. 10(b). As the crack forms and propagates from one grain to another, the associated stress will not be directly applied on the [0 0 1] planes, which is prone to fracture. As a result, the crack propagation in the (Cu,Ni)6Sn5 layer becomes more difficult. This hypothesis may be supported by the observation that the (1 0 1) texture in Cu6Sn5 leads to a higher shear strength of BGA solder balls [25] compared to the (0 0 1) Cu6Sn5 texture. 3.3. Mechanical properties The elastic modulus and hardness values of Cu6Sn5 and (Cu,Ni)6Sn5 measured in the cross-sections of the as-dipped and aged samples, along with the standard deviations, are shown in Figs. 11 and 12. For the Ni-free samples, the elastic modulus (GPa) and hardness (GPa) are 120.5 ± 5.1 and 7.0 ± 0.5 for Sn–3Cu, 123.9 ± 8.2 and 7.0 ± 0.6 for Sn–4Cu, and 121.3 ± 5.9 and 7.3 ± 0.5 for Sn–7Cu, respectively. The average values of Cu6Sn5 are in good agreement with the respective values reported in previous studies [7,34] and the computational result [35], and no apparent effects of Cu concentration can be observed. However, the standard deviations of elastic modulus and hardness of Cu6Sn5 are slightly higher
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than the results obtained in previous studies [8,33,34]. According to the (0 0 1) texture in Cu6Sn5, most indentations in this study were likely to be on a variety of crystal planes parallel to the c-axis of hexagonal Cu6Sn5 in the cross-sectioned samples. By nanoindentation on a directionally solidified sample, the Cu6Sn5 was found to have anisotropic mechanical properties with an elastic modulus (GPa) of 121.0 ± 2.5 on (0 0 1) plane and 105.0 ± 1.8 on (1 1 0) plane; hardness (GPa) of 6.0 ± 0.2 on (0 0 1) plane and 5.7 ± 0.1 on (1 1 0) plane [33]. In this study, the relatively greater standard deviation of mechanical properties is likely due to the anisotropic mechanical properties of the hexagonal Cu6Sn5 on the measured planes. After aging, the elastic modulus (GPa) and hardness (GPa) are 120.2 ± 7.3 and 7.0 ± 0.4 for Sn–3Cu, 119.9 ± 10.0 and 7.0 ± 0.5 for Sn–4Cu and 121.3 ± 9.9 and 7.3 ± 0.4 for Sn–7Cu. Hence, the aging process has no significant effect on mechanical properties, but slightly increased the standard deviations of both elastic modulus and hardness of Cu6Sn5 in the Ni-free samples, which may be due to the mechanical properties variation associated with the change of crystal structure during the hexagonal to monoclinic phase transformation. From the elastic modulus and hardness of (Cu,Ni)6Sn5 as shown in Figs. 11 and 12, two effects of Ni addition on the mechanical properties of Cu6Sn5 can be observed: firstly, the (Cu,Ni)6Sn5 phases in nickel-containing samples have higher elastic modulus and hardness; secondly the average values of elastic modulus and hardness of (Cu,Ni)6Sn5 are more scattered with greater standard deviations than Cu6Sn5. The increase in both elastic modulus and hardness of Cu6Sn5 with the Ni presence is in agreement with previous studies using etched [9] and multi-reflowed [8] BGA samples. Using the cast samples of stoichiometric compositions, a linear relationship has been reported between Ni content and elastic modulus/hardness of (Cu,Ni)6Sn5 [10]. For Ni-free Cu6Sn5, the elastic modulus and hardness were 111.2 ± 3.1 and 6.6 ± 0.4 (GPa). As the Ni content increased to 11.6 mass%, the elastic modulus and hardness were measured to be 140.2 ± 7.4 and 8.6 ± 0.7 (GPa) for the stoichiometric Cu4Ni2Sn5 samples. In this study, the elastic modulus and hardness of (Cu,Ni)6Sn5 are within the range of values measured on stoichiometric samples and the increase on mechanical properties can be explained by solid solution strengthening. The scattered mechanical properties of (Cu,Ni)6Sn5 have been observed on the cross-sectioned samples prior to aging [10]. Using the synchrotron X-ray fluorescence mapping, Ni was confirmed to be homogeneously distributed in the Cu6Sn5 intermetallic layer [36]. In this study, the scattered mechanical properties of (Cu,Ni)6Sn5 in Ni-contained samples can be attributed to the effect of crystal orientation. According to the (1 0 1) texture in (Cu,Ni)6Sn5, the c-axis of (Cu,Ni)6Sn5 is inclined to the substrate surface. After sample mounting and polishing, the c-axis of (Cu,Ni)6Sn5 will not be perfectly parallel to the cross-sectioned surfaces. As a result, the indentation on (Cu,Ni)6Sn5 are likely to be on a number of different crystal planes. Using directionally solidified samples, an anisotropic elastic modulus of (Cu,Ni)6Sn5 was found, but the hardness was more consistent on (0 0 1) and (1 1 0) crystal planes after Ni addition [33]. It is therefore reasonable to attribute the scattered elastic modulus of (Cu,Ni)6Sn5 to the (1 0 1) growth orientation and anisotropic mechanical properties of (Cu,Ni)6Sn5. Although no anisotropic hardness on (0 0 1) and (1 1 0) crystal planes of (Cu,Ni)6Sn5 has been found, the scattering in hardness of (Cu,Ni)6Sn5 in cross-sectioned samples could be explained by the combined effects of (1 0 1) growth orientation and the higher Ni contents used in current study.
4. Conclusions The morphologies, growth orientations and mechanical properties of the Cu6Sn5 and (Cu,Ni)6Sn5 formed at the interfaces of Sn–3/
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4/7Cu–0/0.05/0.1/0.3Ni alloys and poly-crystal Cu plates were systematically investigated. The morphology of Cu6Sn5 was found to be dependent on the Cu content. Scallop-like morphology was formed in the solders with Cu content of 3 and 4 mass%, while faceted rod-like morphology was formed in the solder with Cu content of 7 mass%. The addition of Ni changed the scallop-like morphology of the Cu6Sn5 into a hexagonal rod-like morphology. Aging also influenced the morphologies of the IMCs. Both Cu6Sn5 and (Cu,Ni)6Sn5 displayed faceted scallop-like morphology after aging. The formation of Cu3Sn was found to be dependent on the Cu concentration and the resultant dipping temperature of the solder alloys. The addition of Ni suppressed the formation of Cu3Sn. Our synchrotron XRD results showed that the Ni-free Cu6Sn5 had (0 0 1) texture, which was consistent with the previous study. The addition of a small amount of Ni resulted in a preferred (1 0 1) texture of (Cu,Ni)6Sn5, which might impede the crack propagation in intermetallic layers. As the Ni content increased, the (1 0 1) texture of (Cu,Ni)6Sn5 became stronger. The Cu concentration in solder alloys did not have a significant effect on the mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5, both prior to and after aging. For the Ni-free samples, the average values of mechanical properties of Cu6Sn5 were in good agreement with those reported previously. The (Cu,Ni)6Sn5 in the Ni-containing samples had greater, but more scattered, values of elastic modulus and hardness than the Ni-free Cu6Sn5. The greater standard deviation in modulus and hardness was attributed to the growth of the preferred (1 0 1) texture and the anisotropy in mechanical properties of (Cu,Ni)6Sn5. Acknowledgements Nihon Superior Co. Ltd. (NS) supplied samples under an international cooperative research program. Synchrotron XRD experiments were performed at the Australian Synchrotron (Project IDs: AS113/PDFI/4113 and AS112/PD/3712) under the Queensland Foundation Investor beamtime scheme. The authors would like to thank Dr. J. Kimpton in the Australian Synchrotron for XRD experiments, Dr. S.D. McDonald, Dr. Y.Q. Wu and Mr. J. Read at the University of Queensland (UQ) for stimulating discussions and suggestions. DM would like to thank the financial supports from the Australian Postgraduate Award (APA) Program, the UQ Graduate School International Travel Scholarship (GSITS) and the Frontier@lab Program at Osaka University. References [1] H. Okamoto, Phase Diagrams of Dilute Binary Alloys, vol. 157, ASM, International, 2002. [2] K. Nogita, M. Greaves, Guymer, B. Walsh, J. Kennedy, T. Nishimura, Trans. Jpn. Inst. Electron. Packag. 2 (2010) 104–109. [3] T. Laurila, V. Vuorinen, J.K. Kivilahti, Mater. Sci. Eng., R 49 (2005) 1–60. [4] Y.C. Chan, D. Yang, Prog. Mater Sci. 55 (2010) 428–475. [5] K. Nogita, C. Gourlay, T. Nishimura, JOM 61 (2009) 45–51. [6] D. Mu, J. Read, Y.F. Yang, K. Nogita, J. Mater. Res. 26 (2011) 2660–2664. [7] X. Deng, N. Chawla, K.K. Chawla, M. Koopman, Acta Mater. 52 (2004) 4291– 4303. [8] H. Tsukamoto, Z.G. Dong, H. Huang, T. Nishimura, Mater. Sci. Eng., B 164 (2009) 44–50. [9] L. Xu, J.H.L. Pang, Thin Solid Films 504 (2006) 362–366. [10] D. Mu, H. Tsukamoto, H. Huang, K. Nogita, Mater. Sci. Forum 654 (2010) 2450– 2454. [11] N. Jiang, J.A. Clum, R.R. Chromik, E.J. Cotts, Scr. Mater. 37 (1997) 1851–1854. [12] M. Park, R. Arróyave, Acta Mater. 58 (2010) 4900–4910. [13] R. Gagliano, G. Ghosh, M. Fine, J. Electron. Mater. 31 (2002) 1195–1202. [14] C.H. Wang, S.W. Chen, Acta Mater. 54 (2006) 247–253. [15] A.K. Larsson, L. Stenberg, S. Lidin, Acta Crystallogr. B 50 (1994) 636–643. [16] K. Nogita, T. Nishimura, Scr. Mater. 29 (2008) 191–194. [17] K. Nogita, Intermetallics 18 (2010) 145–149. [18] J.O. Suh, K.N. Tu, N. Tamura, J. App. Phys. 102 (2007) 063511-1–063511-7. [19] J.O. Suh, K.N. Tu, N. Tamura, App. Phys. Lett. 91 (2007) 051907-1–051907-3.
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[30] H. Nishikawa, J.Y. Piao, T. Takemoto, J. Electron. Mater. 35 (2006) 1127–1132. [31] K. Nogita, C.M. Gourlay, S.D. McDonald, Y.Q. Wu, J. Read, Q.F. Gu, Scr. Mater. 65 (2011) 922–925. [32] I. V. Markov, Crystal Growth for Beginners: Fundamentals of Nucleation, Crystal Growth and Epitaxy, World Scientific Pub. Co. Inc., 2003. [33] D. Mu, H. Huang, K. Nogita, Mater. Lett., submitted for publication. [34] P.F. Yang, Y.S. Lai, S.R. Jian, J. Chen, R.S. Chen, Mater. Sci. Eng., A 485 (2008) 305–310. [35] N. Lee, V. Tan, K. Lim, Appl. Phys. Lett. 88 (2006) 031913. [36] K. Nogita, H. Yasuda, C.M. Gourlay, S. Suenaga, H. Tsukamoto, S.D. Mcdonald, A. Takeuchi, K. Uesugi, Y. Suzuki, Trans. Jpn. Inst. Electron. Packag. 1 (2010) 40– 46.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Growth orientations and mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5 on poly-crystalline Cu Dekui Mu a,⇑, Hideyuki Yasuda b, Han Huang a, Kazuhiro Nogita a,b a b
School of Mechanical and Mining Engineering, The University of Queensland, Brisbane, Qld 4072, Australia Department of Adaptive Machine Systems, Osaka University, Suita, Osaka 565-0871, Japan
a r t i c l e
i n f o
Article history: Received 12 March 2012 Received in revised form 25 April 2012 Accepted 25 April 2012 Available online 9 May 2012 Keywords: Growth orientation Pole figures Synchrotron radiation Mechanical properties Nanoindentation Intermetallics Crystal growth X-ray diffraction
a b s t r a c t Lead-free solders are important materials in current generation electrical packages, due to the increasingly stringent legislative requirement aimed at reducing the use of lead. The lead-free solders based on the Sn–Cu system with Ni addition have become popular because of their superior soldering properties, as well as their comparatively low cost. This research investigates the effect of Ni addition on the growth morphologies, crystal orientations and mechanical properties of Cu6Sn5 at the interface between hyper-eutectic Sn–Cu high-temperature lead-free solder alloys and Cu substrates, prior to and after aging, by the use of X-ray diffraction (XRD), scanning electron microscopy (SEM) and nanoindentation. The (Cu,Ni)6Sn5 in Sn–Cu–Ni/Cu solder joints showed a more strongly oriented (1 0 1) texture, compared to the Cu6Sn5 in Sn–Cu/Cu solder joints. The Ni-induced (1 0 1) texture contributes to higher and more scattered average values with larger standard deviations in both elastic modulus and hardness for (Cu,Ni)6Sn5. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Due to environmental concerns, the restriction of hazard substance (RoHS) directive has proposed the phase-out of lead in high-temperature soldering by 2014 in the European Union (EU). This has increased the demand for suitable high-temperature lead-free solders. One of the candidate alloy groups for high-temperature soldering is the hyper-eutectic Sn–Cu alloys, which have Cu contents ranging between the eutectic composition of 0.86 mass% [1] and an upper limit of around 7.6 mass%. These hyper-eutectic Sn–Cu alloys have been used for high temperature dip soldering and tinning of copper wire, polyurethane coated wire and component terminations and display minimal copper erosion at temperatures up to 400 °C [2]. During the interface reactions between the Sn–Cu alloy and Cu substrate, a layer of Cu6Sn5 forms. The Cu6Sn5 layer at the interface is commonly recognized to play a critical role for crack formation in a solder joint during service [3,4]. It has been observed that most cracks forming in the Cu6Sn5 layer are nearly parallel to the solder/substrate interface, while few cracks are present in alternate directions [5]. However, it is unclear if this is due to preferential crack paths existing in Cu6Sn5 combined with preferred growth textures, or the location of maximum
⇑ Corresponding author. Tel.: +61 7 3365 1387; fax: +61 7 3365 3888. E-mail addresses: [email protected], [email protected] (D. Mu). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.04.110
stress planes that occur during thermal cycling and associated dimensional changes [6]. To clarify this issue, detailed information on the morphology, crystal orientation and mechanical properties of Cu6Sn5 formed at solder/substrate interface is required. Cu6Sn5 is the most commonly formed intermetallic during leadfree soldering. Extensive research has been conducted to investigate the mechanical [7–10] and thermal properties [11], the formation [12,13], the evolution [14] and the crystal structure [15–17] of Cu6Sn5. Taking the growth orientations into account, several preferred crystal relationships between Cu6Sn5 and single-crystal Cu [18–20] or Ni [21] plates have been observed and these have been attributed to the minimization of distance between Cu atoms in Cu6Sn5 and the Cu or Ni atoms in the plates. For the Cu6Sn5 formed on a non-textured poly-crystalline Cu substrate, the intermetallic orientation has been studied in the systems of Sn–Pb [22], Sn– Ag–Cu [23], Sn–0.7Cu [24] and Sn–3.5Ag [25]. However, there is little information available on the growth morphology and crystal orientation of Cu6Sn5 formed during the interface reaction between Cu and hyper-eutectic high-temperature Sn–Cu solders. In particular, the crystal orientation and its effect on mechanical properties of Cu6Sn5 at the solder/substrate interface have never been discussed. In this research, the morphology, crystal orientation and mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5 intermetallics, formed between hyper-eutectic Sn–3/4/7Cu–0/0.05/0.1/0.3Ni (mass%) alloys and poly-crystalline Cu substrates, are systematically
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investigated using scanning electronic microscopy (SEM), X-ray diffraction (XRD) and nanoindentation. The crystal orientations are then discussed with respect to the pattern of crack formation and propagation, as well as mechanical properties, in Cu6Sn5 and (Cu,Ni)6Sn5 intermetallic layers. 2. Experiment The Cu6Sn5 and (Cu,Ni)6Sn5 formed during the interface reaction between polycrystalline Cu and a variety of lead-free Sn–3/4/7Cu–Ni solders, as shown in Table 1, were investigated. The samples were prepared by dipping the Cu plates (C1220P) of 10 mm 30 mm 0.3 mm with a common flux (Flux B) into the molten solder at a temperature of 340 °C for Sn–3Cu (mass%) alloys, 373 °C for Sn–4Cu alloys and 430 °C for Sn–7Cu alloys for 10 s. Some of the dipped samples were then aged at 120 °C for 1176 h. Selected as-dipped and aged samples were polished perpendicular to the solder interface using conventional techniques for sample preparation. Mechanical properties (elastic modulus and hardness) of the IMCs in cross-sectioned samples were measured using a Hysitron nanoindenter (Hysitron, MN, USA) [26,27]. A Berkovich diamond indenter of tip radius of 100 nm was used and the indentation load of 1000 lN was applied at a loading rate of 200 lN/s and with a dwell time at peak load of 5 s. The unloading rate was 200 lN/s. In order to investigate the morphology and growth orientation from the top view of the intermetallics (solder side), some of the samples, prior to and after aging, were etched in a solution of ortho-nitro-phenol (35 g) and NaOH (50 g) in 1 L of water at 80 °C to remove the solder alloys. The etched samples were then cut into a round plate of 9 mm diameter before XRD investigations. Synchrotron X-ray diffraction (XRD) was performed on the powder diffraction beamline at Australian Synchrotron using a Mythen-II detector. The synchrotron X-ray diffraction data were collected on the prepared samples in the 2h range 10–80° with the accelerate voltage of 20 keV and the wavelengths of 0.825 Å, respectively. The data collection period was 11 min, with 1 min for stabilization and 10 min for data collection. The X-ray peak data for Cu6Sn5 and (Cu,Ni)6Sn5 were analyzed using Search-Match software. As a reference crystallography and atomic coordination, ICDD (International Centre for Diffraction Data) number of 045-1488 (for monoclinic, C2/c) and 047-1575 (for hexagonal, P63/mmc) were used in association with Search-Match. In addition, (1 0 1) XRD pole figures were plotted for all 24 samples using a laboratory Rigaku machine (Rigaku, Tokyo, Japan). The 2h angle for the (1 0 1) plane of hexagonal Cu6Sn5 was calculated according to ICDD number of 047-1575 as 6.88° for Cu-Ka X-ray. For each pole figure, 64 scans were conducted with the rotation angle of 3°. SEM observation was conducted on both cross-sectioned and etched samples using a Philips XL30 machine (FEI, Hillsboro, USA) with an acceleration voltage of 20 kV.
3. Results and discussion 3.1. Growth morphologies During the interface reaction between near-eutectic Sn–Cu solders and Cu substrates, a scallop-type Cu6Sn5 layer firstly forms, followed by the formation of a two-phase Sn + Cu6Sn5 layer [4]. Park and Arróyave [12] simulated the early-stage nucleation of Cu6Sn5 and proposed that the precipitated Cu6Sn5 would grow parallel to the solder/substrate interface. The growth of the Cu6Sn5 layer in the direction perpendicular to the solder/substrate interface starts only after the precipitated Cu6Sn5 grains coalesce. According to the Sn–Cu phase diagram, the Cu3Sn phase may also Table 1 Composition of solders. Sample No.
Composition (mass%)
1 2 3 4 5 6 7 8 9 10 11 12
Sn–3Cu Sn–3Cu–0.05Ni Sn–3Cu–0.1Ni Sn–3Cu–0.3Ni Sn–4Cu Sn–4Cu–0.05Ni Sn–4Cu–0.1Ni Sn–4Cu–0.3Ni Sn–7Cu Sn–7Cu–0.05Ni Sn–7Cu–0.1Ni Sn–7Cu–0.3Ni
39
form during the liquid–solid reaction depending on the solder compositions and reaction conditions. During aging, Cu3Sn grows and the total thickness of the intermetallics layer increases with coarsening of the intermetallic structures. Ni additions reportedly suppress the formation and growth of Cu3Sn [28]. As shown in Figs. 1(a) and 2(a), the Cu6Sn5 that forms in the Ni-free Sn–3/4Cu and Cu reaction couples shows a scallop-type morphology. This observation is in agreement with the previous studies on near-eutectic Sn–0.7Cu [3] and eutectic Sn–Pb solders [22], in which the growth of intermetallics was on poly-crystalline substrates. In Fig. 3(a), the Cu6Sn5 formed in the Ni-free Sn–7Cu alloy and Cu reaction couple has a faceted rod-like morphology. This is likely due to the enhanced intermetallic formation at higher dipping temperatures. For the Ni-containing Sn–3/4/7Cu and Cu reaction couples, the morphology of (Cu,Ni)6Sn5 was changed from scallop-type to hexagonal rod-like, as shown in Figs. 1–3(c), (e) and (g). The formation of rod-like (Cu,Ni)6Sn5 may result from anisotropy of surface energies and compositional supercooling related to the presence of Ni, which influences the growth behavior of Cu6Sn5. After aging, the morphologies of Cu6Sn5 and (Cu,Ni)6Sn5 become faceted scallop-type, as shown in Figs. 1–3(b), (d), (f) and (h). At temperatures below 350 °C, Cu3Sn can also be formed according to the Sn–Cu phase diagram. In the previous studies on eutectic Sn–Cu and Sn–Pb solders reacted with Cu substrates, the Cu3Sn phase cannot be observed until a sufficient aging treatment was given [3]. In this study, the formation of Cu3Sn was found to be dependent on both the Cu concentration and resultant reaction temperature. For the as-dipped Ni-free Sn–3Cu and Cu reaction couple, Cu3Sn cannot be detected by SEM. However, for the asdipped Ni-free Sn–4/7Cu and Cu reaction couples, Cu3Sn can be clearly observed as shown in Fig. 4(a). The addition of Ni can suppress the formation and growth of Cu3Sn. Wang et al. reported that the addition of 5 mass% Ni could completely prevent the formation of Cu3Sn in Sn–Ag–Cu alloys [28]. In our study, the formation of Cu3Sn was not observed for any as-dipped Ni-containing Sn–4/ 7Cu and Cu reaction couples. After aging, the Ni-containing samples also have less Cu3Sn compared to these Ni-free samples, as shown in Fig. 4(b) and (d). This can be explained from the kinetic viewpoint that Ni suppresses the formation of Cu3Sn by reducing the driving force, and hence the diffusion flux, of Cu in Cu3Sn [29]. Besides the effects observed in this study, the addition of a small amount of Ni has been found to reduce the thickness of the intermetallic layer [30] and stabilize the hexagonal Cu6Sn5 [16,17], which may effectively reduce the internal stress which results from phase transformations. 3.2. Growth orientations The synchrotron XRD patterns of Cu6Sn5 formed at the Sn–4Cu– (0.05Ni)/Cu interface, prior to and after aging, are shown in Fig. 5(a–c). It can be found that most Cu6Sn5 at the as-dipped Sn–4Cu/Cu interface remained in the hexagonal structure as shown in Fig. 5(a). From the Sn–Cu phase diagram, the Cu6Sn5 has two equilibrium phases: the monoclinic g0 phase stable at temperatures below 186 °C and the hexagonal g phase stable at temperatures above 186 °C [1]. According to the previous studies on the time–temperature–transformation (TTT) curve [31], the dipping time (10 s) and subsequent air cooling time in this study were both insufficiently long for the hexagonal Cu6Sn5 to completely transform into the low-temperature monoclinic phase. The Cu6Sn5 in as-dipped Ni-free Sn–Cu/Cu reaction couples is likely to be a mixture of monoclinic and meta-stable hexagonal phases, which can be confirmed from these small monoclinic peaks at high 2h angles as shown in Fig. 5(a). After aging at 120 °C for 1176 h, most hexagonal Cu6Sn5 transformed into the monoclinic phase as shown in
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Fig. 1. Top view SEM images of Cu6Sn5 and (Cu,Ni)6Sn5 in Sn–3Cu–xNi/Cu couples (x = 0.05, 0.1, 0.3).
Fig. 2. Top view SEM images of Cu6Sn5 and (Cu,Ni)6Sn5 in Sn–4Cu–xNi/Cu couples (x = 0.05, 0.1, 0.3).
Fig. 3. Top view SEM images of Cu6Sn5 and (Cu,Ni)6Sn5 in Sn–7Cu–xNi/Cu couples (x = 0.05, 0.1, 0.3).
Fig. 5(b). In both as-dipped and aged Sn–4Cu/Cu samples, small Cu3Sn peaks can be observed at low 2h angles, which is consistent with SEM observations. It has been reported that the Cu6Sn5 formed at the interface between Sn–Ag solder and Cu substrates has a (0 0 1) growth texture with the c axis perpendicular to the substrate surfaces [25]. In this study, the (0 0 1) texture of Cu6Sn5 at the Sn–4Cu/Cu interface can be confirmed by comparing the diffraction patterns of Sn–4Cu and Sn–4Cu–0.05Ni samples. In Fig. 5(a), Cu6Sn5 had higher X-ray diffraction intensity on (0 0 2) and (0 0 4) crystal planes, which means the c axis of Cu6Sn5 is also perpendicular to the substrate and there is a (0 0 1) growth texture. After Ni addition, the X-ray diffraction intensity on (1 0 1) crystal plane of (Cu,Ni)6Sn5 increased as shown in Fig. 5(c). From Fig. 5(a) and (c), it can be found that both Cu6Sn5 and (Cu,Ni)6Sn5 have the dominant growth on (1 0 2) plane. Therefore, the effect of Ni solubility on the crystal growth of Cu6Sn5 and (Cu,Ni)6Sn5 can be represented by the diffraction intensity ratios between
(1 0 1), (1 0 2) and (2 0 2) planes. The X-ray diffraction intensity ratios between (1 0 1), (1 0 2), and (2 0 2) crystal planes of Cu6Sn5 and (Cu,Ni)6Sn5 in as-dipped Sn–3Cu/Cu and Sn–4Cu(0, 0.05, 0.3Ni)/Cu reaction couples are shown in Fig. 5(d–f). Compared to the Cu6Sn5 formed in Ni-free Sn–3Cu/Cu and Sn–4Cu/Cu reaction couples, the addition of Ni increased the crystal growth on (1 0 1) plane of Cu6Sn5 by a factor of around 10 as shown in Fig. 5(d). In Fig. 5(e), the growth on (2 0 2) plane of Cu6Sn5 was also slightly increased by Ni addition. This observation is in agreement with previous study on near-eutectic Sn–0.7Cu solder and suggests the Ni addition resulted in the promotion of (1 0 1) growth orientation of Cu6Sn5. For a systematic analysis of the effect of Ni on Cu6Sn5 growth orientation, the (1 0 1) pole figures of hexagonal Cu6Sn5 and (Cu,Ni)6Sn5 in all 24 samples were measured by a laboratory XRD machine as shown in Figs. 6–9. In Fig. 6(a), the intensity distribution of Cu6Sn5 formed in Ni-free samples are almost random along the rolling direction (RD) and transverse direction (TD), which
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Fig. 4. Intermetallics at the solder/substrate interface: (a) Sn–4Cu as-dipped (b) Sn–4Cu aged, (c) Sn–4Cu–0.1Ni as-dipped and (d) Sn–47Cu–0.1Ni aged.
Fig. 5. X-ray diffraction results: (a) X-ray diffraction pattern for Sn–4Cu/Cu sample, (b) X-ray diffraction pattern for aged Sn–4Cu/Cu sample, (c) X-ray diffraction pattern for Sn–4Cu–0.05Ni/Cu sample, (d) X-ray diffraction intensity ratio of (1 0 1)/(1 0 2), (e) XRD intensity ratio of (2 0 2)/(1 0 2) and (f) X-ray diffraction intensity ratio of {(1 0 1) + (2 0 2)}/(1 0 2).
indicates that no preferred (1 0 1) crystal orientation exists. As the Cu content increases, the diffraction peaks in the (1 0 1) pole figures become stronger and more scattered, as shown in Fig. 6(a), (c) and (e), which is attributed to increased Cu6Sn5 formation and the presence of the Cu3Sn phase as observed by SEM and synchrotron XRD techniques. After aging, the diffraction intensity becomes stronger and more scattered compared to that of the as-dipped samples as shown in Fig. 6(b), (d) and (f). The greater degree of scatter of diffraction
peaks in the aged samples was due to not only the growth of Cu6Sn5, but also the increased volume of the monoclinic Cu6Sn5 phase, whose diffraction intensity naturally contains many small peaks. In addition, the growth of the Cu3Sn layer in the aged Ni-free Sn–4/7Cu/Cu reaction couples would also contribute to the scattered diffraction intensity. The (1 0 1) pole figures of Cu6Sn5 and (Cu,Ni)6Sn5 formed in Sn– 3/4/7Cu–Ni/Cu interface are shown in Figs. 7–9. While the Ni-free Cu6Sn5 has a (0 0 1) texture [25], the addition of Ni resulted in a
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Fig. 6. Pole figure of Ni-free samples: (a) Sn–3Cu, (b) aged Sn–3Cu, (c) Sn–4Cu, (d) aged Sn–4Cu, (e) Sn–7Cu and (f) aged Sn–7Cu.
Fig. 7. Pole figure of Sn–3Cu–xNi samples: (a) Sn–3Cu, (b) aged Sn–3Cu, (c) Sn–3Cu–0.05Ni, (d) aged Sn–3Cu–0.1Ni, (e) Sn–3Cu–0.1Ni, (f) aged Sn–3Cu–0.1Ni, (g) Sn–3Cu– 0.3Ni and (h) aged Sn–3Cu–0.3Ni.
Fig. 8. Pole figure of Sn–4Cu–xNi samples: (a) Sn–4Cu, (b) aged Sn–4Cu, (c) Sn–4Cu–0.05Ni, (d) aged Sn–4Cu–0.1Ni, (e) Sn–4Cu–0.1Ni, (f) aged Sn–4Cu–0.1Ni, (g) Sn–4Cu– 0.3Ni and (h) aged Sn–4Cu–0.3Ni.
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Fig. 9. Pole figures of Sn–7Cu–xNi samples: (a) Sn–7Cu, (b) aged Sn–7Cu, (c) Sn–7Cu–0.05Ni, (d) aged Sn–7Cu–0.1Ni, (e) Sn–7Cu–0.1Ni, (f) aged Sn–7Cu–0.1Ni, (g) Sn–7Cu– 0.3Ni and (h) aged Sn–7Cu–0.3Ni.
Fig. 10. Schematic diagram of crack propagation.
Fig. 11. Elastic modulus of Cu6Sn5 and (Cu,Ni)6Sn5 in as-dipped and aged samples.
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Fig. 12. Hardness of Cu6Sn5 and (Cu,Ni)6Sn5 in as-dipped and aged samples.
(1 0 1) texture of (Cu,Ni)6Sn5 in all the Ni-containing samples. As the Ni and/or Cu content increased, the (1 0 1) texture of (Cu,Ni)6Sn5 became stronger. The (1 0 1) pole figures showed that the (Cu,Ni)6Sn5 has the (1 0 1) plane parallel to the solder/substrate as shown in Figs. 7–9(c), (e) and (g). The hexagonal (Cu,Ni)6Sn5 therefore has the c-axis in a direction inclined to the normal of the solder/substrate interface. Using lattice parameters measured in the previous experiments [6] (a- and c- axis of hexagonal (Cu,Ni)6Sn5 of 0.42 and 0.51 nm, respectively) it is calculated there is an approximately 45° disorientation between the c-axis of (Cu,Ni)6Sn5 and the normal of the solder/substrate interface. Growth orientation normally results from preferred nucleation and/or growth mechanisms and the grain size of the substrate may influence these mechanisms [32]. In this study, the substrate is a common C122P polycrystalline Cu plate without heat treatments as given in previous studies [18–20]. Thus, the formation of the (1 0 1) texture of (Cu,Ni)6Sn5 in this study should be attributed to preferred growth resulting from the anisotropic surface energies. After random early-stage nucleation, the (Cu,Ni)6Sn5 grains with (1 0 1) planes parallel to the substrate surface have lower surface energies and consume other (Cu,Ni)6Sn5 grains during the ripening process, which results in the formation of a (1 0 1) texture. This explanation could be supported by the synchrotron XRD result that Ni promoted the (1 0 1) plane growth of Cu6Sn5 and our previous study on near-eutectic Sn–0.7Cu alloys [24]. Similar mechanism can also be used to explain the (0 0 1) texture in Cu6Sn5. Li et al. attributed the formation of (0 0 1) to (1 0 1) texture of Cu6Sn5 as temperature increasing from 240 to 280 °C to the difference in surface energy [25]. In this study, the Ni-free Cu6Sn5 has a (0 0 1) growth orientation at a temperature of 340 °C and above, which can be explained as the high Cu contents in solder alloys enhance the (0 0 1) growth of Cu6Sn5 [24]. After aging, the (1 0 1) pole figure of the aged samples displays a stronger (1 0 1) texture in the (Cu,Ni)6Sn5 layer, as shown in Figs. 7–9 (d), (f) and (h). As mentioned before, there is an allotropic transformation from high-temperature hexagonal g-Cu6Sn5 to low-temperature monoclinic g0 -Cu6Sn5 at an equilibrium temperature of 186 °C. According to the TTT curve of Cu6Sn5 [31], this allotropic transformation is complete after aging at 120 °C for 1176 h and the reported (0 0 1) texture of hexagonal Cu6Sn5 may not be observable. In this study, the (Cu,Ni)6Sn5 remains in the stabilized
hexagonal structure after aging and the stronger (1 0 1) texture is due to the continued growth of the stabilized (Cu,Ni)6Sn5. The observation that Ni affects the growth orientation of the interfacial Cu6Sn5 may have implications for failure mechanisms in lead-free solder joints. Firstly, it is known that the Cu6Sn5 formed at the Sn–Cu/Cu interface has a (0 0 1) growth orientation with the c-axis perpendicular to the solder/substrate interface. Moreover, the cracks formed in the Cu6Sn5 layer are prone to be parallel to the solder/substrate interface [5]. Using directional solidification samples, the elastic modulus and hardness on [0 0 1] crystal planes are found to be lower than [1 0 1] crystal planes [33]. Hence, it is reasonable to conclude the Cu6Sn5 at the dipped solder/substrate interface is easier to fracture on [0 0 1] crystal planes. As the crack propagates, the resultant stress is transferred to the adjacent grains because there is little barrier between adjacent grains. For the Cu6Sn5 with a (0 0 1) growth texture, as shown in Fig. 10(a), the stress will be directly applied on the adjacent [0 0 1] planes, on which cracking is easier to initiate. However, the (1 0 1) growth texture of (Cu,Ni)6Sn5 will result in an angle between [0 0 1] planes of adjacent grains, as shown in Fig. 10(b). As the crack forms and propagates from one grain to another, the associated stress will not be directly applied on the [0 0 1] planes, which is prone to fracture. As a result, the crack propagation in the (Cu,Ni)6Sn5 layer becomes more difficult. This hypothesis may be supported by the observation that the (1 0 1) texture in Cu6Sn5 leads to a higher shear strength of BGA solder balls [25] compared to the (0 0 1) Cu6Sn5 texture. 3.3. Mechanical properties The elastic modulus and hardness values of Cu6Sn5 and (Cu,Ni)6Sn5 measured in the cross-sections of the as-dipped and aged samples, along with the standard deviations, are shown in Figs. 11 and 12. For the Ni-free samples, the elastic modulus (GPa) and hardness (GPa) are 120.5 ± 5.1 and 7.0 ± 0.5 for Sn–3Cu, 123.9 ± 8.2 and 7.0 ± 0.6 for Sn–4Cu, and 121.3 ± 5.9 and 7.3 ± 0.5 for Sn–7Cu, respectively. The average values of Cu6Sn5 are in good agreement with the respective values reported in previous studies [7,34] and the computational result [35], and no apparent effects of Cu concentration can be observed. However, the standard deviations of elastic modulus and hardness of Cu6Sn5 are slightly higher
D. Mu et al. / Journal of Alloys and Compounds 536 (2012) 38–46
than the results obtained in previous studies [8,33,34]. According to the (0 0 1) texture in Cu6Sn5, most indentations in this study were likely to be on a variety of crystal planes parallel to the c-axis of hexagonal Cu6Sn5 in the cross-sectioned samples. By nanoindentation on a directionally solidified sample, the Cu6Sn5 was found to have anisotropic mechanical properties with an elastic modulus (GPa) of 121.0 ± 2.5 on (0 0 1) plane and 105.0 ± 1.8 on (1 1 0) plane; hardness (GPa) of 6.0 ± 0.2 on (0 0 1) plane and 5.7 ± 0.1 on (1 1 0) plane [33]. In this study, the relatively greater standard deviation of mechanical properties is likely due to the anisotropic mechanical properties of the hexagonal Cu6Sn5 on the measured planes. After aging, the elastic modulus (GPa) and hardness (GPa) are 120.2 ± 7.3 and 7.0 ± 0.4 for Sn–3Cu, 119.9 ± 10.0 and 7.0 ± 0.5 for Sn–4Cu and 121.3 ± 9.9 and 7.3 ± 0.4 for Sn–7Cu. Hence, the aging process has no significant effect on mechanical properties, but slightly increased the standard deviations of both elastic modulus and hardness of Cu6Sn5 in the Ni-free samples, which may be due to the mechanical properties variation associated with the change of crystal structure during the hexagonal to monoclinic phase transformation. From the elastic modulus and hardness of (Cu,Ni)6Sn5 as shown in Figs. 11 and 12, two effects of Ni addition on the mechanical properties of Cu6Sn5 can be observed: firstly, the (Cu,Ni)6Sn5 phases in nickel-containing samples have higher elastic modulus and hardness; secondly the average values of elastic modulus and hardness of (Cu,Ni)6Sn5 are more scattered with greater standard deviations than Cu6Sn5. The increase in both elastic modulus and hardness of Cu6Sn5 with the Ni presence is in agreement with previous studies using etched [9] and multi-reflowed [8] BGA samples. Using the cast samples of stoichiometric compositions, a linear relationship has been reported between Ni content and elastic modulus/hardness of (Cu,Ni)6Sn5 [10]. For Ni-free Cu6Sn5, the elastic modulus and hardness were 111.2 ± 3.1 and 6.6 ± 0.4 (GPa). As the Ni content increased to 11.6 mass%, the elastic modulus and hardness were measured to be 140.2 ± 7.4 and 8.6 ± 0.7 (GPa) for the stoichiometric Cu4Ni2Sn5 samples. In this study, the elastic modulus and hardness of (Cu,Ni)6Sn5 are within the range of values measured on stoichiometric samples and the increase on mechanical properties can be explained by solid solution strengthening. The scattered mechanical properties of (Cu,Ni)6Sn5 have been observed on the cross-sectioned samples prior to aging [10]. Using the synchrotron X-ray fluorescence mapping, Ni was confirmed to be homogeneously distributed in the Cu6Sn5 intermetallic layer [36]. In this study, the scattered mechanical properties of (Cu,Ni)6Sn5 in Ni-contained samples can be attributed to the effect of crystal orientation. According to the (1 0 1) texture in (Cu,Ni)6Sn5, the c-axis of (Cu,Ni)6Sn5 is inclined to the substrate surface. After sample mounting and polishing, the c-axis of (Cu,Ni)6Sn5 will not be perfectly parallel to the cross-sectioned surfaces. As a result, the indentation on (Cu,Ni)6Sn5 are likely to be on a number of different crystal planes. Using directionally solidified samples, an anisotropic elastic modulus of (Cu,Ni)6Sn5 was found, but the hardness was more consistent on (0 0 1) and (1 1 0) crystal planes after Ni addition [33]. It is therefore reasonable to attribute the scattered elastic modulus of (Cu,Ni)6Sn5 to the (1 0 1) growth orientation and anisotropic mechanical properties of (Cu,Ni)6Sn5. Although no anisotropic hardness on (0 0 1) and (1 1 0) crystal planes of (Cu,Ni)6Sn5 has been found, the scattering in hardness of (Cu,Ni)6Sn5 in cross-sectioned samples could be explained by the combined effects of (1 0 1) growth orientation and the higher Ni contents used in current study.
4. Conclusions The morphologies, growth orientations and mechanical properties of the Cu6Sn5 and (Cu,Ni)6Sn5 formed at the interfaces of Sn–3/
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4/7Cu–0/0.05/0.1/0.3Ni alloys and poly-crystal Cu plates were systematically investigated. The morphology of Cu6Sn5 was found to be dependent on the Cu content. Scallop-like morphology was formed in the solders with Cu content of 3 and 4 mass%, while faceted rod-like morphology was formed in the solder with Cu content of 7 mass%. The addition of Ni changed the scallop-like morphology of the Cu6Sn5 into a hexagonal rod-like morphology. Aging also influenced the morphologies of the IMCs. Both Cu6Sn5 and (Cu,Ni)6Sn5 displayed faceted scallop-like morphology after aging. The formation of Cu3Sn was found to be dependent on the Cu concentration and the resultant dipping temperature of the solder alloys. The addition of Ni suppressed the formation of Cu3Sn. Our synchrotron XRD results showed that the Ni-free Cu6Sn5 had (0 0 1) texture, which was consistent with the previous study. The addition of a small amount of Ni resulted in a preferred (1 0 1) texture of (Cu,Ni)6Sn5, which might impede the crack propagation in intermetallic layers. As the Ni content increased, the (1 0 1) texture of (Cu,Ni)6Sn5 became stronger. The Cu concentration in solder alloys did not have a significant effect on the mechanical properties of Cu6Sn5 and (Cu,Ni)6Sn5, both prior to and after aging. For the Ni-free samples, the average values of mechanical properties of Cu6Sn5 were in good agreement with those reported previously. The (Cu,Ni)6Sn5 in the Ni-containing samples had greater, but more scattered, values of elastic modulus and hardness than the Ni-free Cu6Sn5. The greater standard deviation in modulus and hardness was attributed to the growth of the preferred (1 0 1) texture and the anisotropy in mechanical properties of (Cu,Ni)6Sn5. Acknowledgements Nihon Superior Co. Ltd. (NS) supplied samples under an international cooperative research program. Synchrotron XRD experiments were performed at the Australian Synchrotron (Project IDs: AS113/PDFI/4113 and AS112/PD/3712) under the Queensland Foundation Investor beamtime scheme. The authors would like to thank Dr. J. Kimpton in the Australian Synchrotron for XRD experiments, Dr. S.D. McDonald, Dr. Y.Q. Wu and Mr. J. Read at the University of Queensland (UQ) for stimulating discussions and suggestions. DM would like to thank the financial supports from the Australian Postgraduate Award (APA) Program, the UQ Graduate School International Travel Scholarship (GSITS) and the Frontier@lab Program at Osaka University. References [1] H. Okamoto, Phase Diagrams of Dilute Binary Alloys, vol. 157, ASM, International, 2002. [2] K. Nogita, M. Greaves, Guymer, B. Walsh, J. Kennedy, T. Nishimura, Trans. Jpn. Inst. Electron. Packag. 2 (2010) 104–109. [3] T. Laurila, V. Vuorinen, J.K. Kivilahti, Mater. Sci. Eng., R 49 (2005) 1–60. [4] Y.C. Chan, D. Yang, Prog. Mater Sci. 55 (2010) 428–475. [5] K. Nogita, C. Gourlay, T. Nishimura, JOM 61 (2009) 45–51. [6] D. Mu, J. Read, Y.F. Yang, K. Nogita, J. Mater. Res. 26 (2011) 2660–2664. [7] X. Deng, N. Chawla, K.K. Chawla, M. Koopman, Acta Mater. 52 (2004) 4291– 4303. [8] H. Tsukamoto, Z.G. Dong, H. Huang, T. Nishimura, Mater. Sci. Eng., B 164 (2009) 44–50. [9] L. Xu, J.H.L. Pang, Thin Solid Films 504 (2006) 362–366. [10] D. Mu, H. Tsukamoto, H. Huang, K. Nogita, Mater. Sci. Forum 654 (2010) 2450– 2454. [11] N. Jiang, J.A. Clum, R.R. Chromik, E.J. Cotts, Scr. Mater. 37 (1997) 1851–1854. [12] M. Park, R. Arróyave, Acta Mater. 58 (2010) 4900–4910. [13] R. Gagliano, G. Ghosh, M. Fine, J. Electron. Mater. 31 (2002) 1195–1202. [14] C.H. Wang, S.W. Chen, Acta Mater. 54 (2006) 247–253. [15] A.K. Larsson, L. Stenberg, S. Lidin, Acta Crystallogr. B 50 (1994) 636–643. [16] K. Nogita, T. Nishimura, Scr. Mater. 29 (2008) 191–194. [17] K. Nogita, Intermetallics 18 (2010) 145–149. [18] J.O. Suh, K.N. Tu, N. Tamura, J. App. Phys. 102 (2007) 063511-1–063511-7. [19] J.O. Suh, K.N. Tu, N. Tamura, App. Phys. Lett. 91 (2007) 051907-1–051907-3.
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