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A new phase in stoichiometric Cu6Sn5
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Acta Materialia 60 (2012) 6581–6591 www.elsevier.com/locate/actamat
A new phase in stoichiometric Cu6Sn5 Y.Q. Wu a,b, J.C. Barry c, T. Yamamoto d, Q.F. Gu e, S.D. McDonald a,b, S. Matsumura d, H. Huang b, K. Nogita a,b,⇑ a
Nihon Superior Centre for the Manufacture of Electronic Materials (NS CMEM), The University of Queensland, QLD 4072, Australia b School of Mechanical and Mining Engineering, The University of Queensland, QLD 4072, Australia c Science and Engineering Faculty, Queensland University of Technology, QLD 4000, Australia d Department of Applied Quantum Physics and Nuclear Engineering, Kyushu University, Motooka 744, Nishi-ku, Fukuoka 819-0395, Japan e Powder Diffraction Beamline, The Australian Synchrotron, Clayton, VIC 3168, Australia Received 19 June 2012; accepted 12 August 2012 Available online 17 September 2012
Abstract The intermetallic compound Cu6Sn5 is a significant microstructural feature of many electronic devices where it is present at the solder–substrate interfaces. The time- and temperature-dependent thermomechanical properties of Cu6Sn5 are dependent on the nature and stability of its crystal structure, which has been shown to exist in at least four variants (g, g0 , g6 and g8). This research details an additional newly identified monoclinic-based structure in directly alloyed stoichiometric Cu6Sn5 using variable-temperature synchrotron X-ray diffraction (XRD) and transmission electron microscopy. The phase is associated with a departure from the equilibrium temperature of the polymorphic monoclinic–hexagonal transformation temperature. The new monoclinic phase can be treated as a modulation of four ˚ , b = 7.311 A ˚, g8-Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell. It has been labeled as g4+1 and has cell parameters of a = 92.241 A ˚ and b = 118.95° determined from electron diffraction patterns. The XRD results could be fitted well to a Rietveld refinement c = 9.880 A using the new crystal parameters. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Cu6Sn5; Phase transformation; X-ray diffraction (XRD); Electron diffraction pattern; Crystallography
1. Introduction Electronic devices containing soldered components are ubiquitous in modern life. Due to environmental concerns, the majority of solder alloys have changed composition to eliminate the presence of lead [1–4]. Numerous Pb-free solder alloys are currently in use or being developed, including a range of Sn–Cu-based solder alloys [5–9]. In these alloys and many other Sn-based solders significant quantities of intermetallic compound (IMC) of Cu6Sn5 form during processing and cooling. As the scale of devices is minimized, the IMC occupies an increasing volume fraction of the solder joint and, in an extreme situation, all of the Sn might be consumed to form the IMC [10]. The IMC layer is integral ⇑ Corresponding author.
E-mail address: [email protected] (K. Nogita).
to microelectronic solder reliability and a fundamental understanding of Cu6Sn5 crystallography, which influences chemical, electrical and mechanical properties [11], is essential to the reliable manufacture and service of a large number of electronic devices. In addition, Cu6Sn5 has been proposed as a promising anode material in Li-ion batteries [12–14] and large quantities of Cu6Sn5 material would be required for this application. Cu6Sn5 (54.5 at.% Cu) exists in two different crystal structures according to the binary Cu–Sn phase diagram [15], as partly shown in Fig. 1. At equilibrium, monoclinic ˚, g0 -Cu6Sn5 (C2/c) with cell parameters of a = 11.022 A ˚ ˚ b = 7.282 A, c = 9.827 A, b = 98.84° [16] is the lower-temperature phase, while hexagonal g-Cu6Sn5 (P63/mmc) with ˚ and c = 5.086 A ˚ [17] is stacell parameters of a = 4.190 A 0 ble at higher temperatures. A g–g allotropic phase transformation occurs as the temperature falls below 186 °C
1359-6454/$36.00 Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.actamat.2012.08.024
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electronic materials, including solder alloys and Li-ion battery technology. 2. Experimental details 2.1. Preparation of samples
Fig. 1. Part of the binary phase diagram of the Cu–Sn system (from Ref. [15]).
[5,11,18]. It can be clearly seen from the phase diagram, as indicated by the red1 solid arrow in Fig. 1, that the g-phase region is a narrow band rather than a line, and thus composition variation does occur in the g-phase region. In addition, it is noted that the phase field is skewed towards Cu-rich compositions as the temperature increases. The composition range for these phases may result in the potential for crystal structure variations. The high-temperature modulation of g phase above 186 °C has been systematically studied by Larsson et al. [19] and two superstructures were proposed, namely g6 and g8, which are slightly more Cu rich (55.56 at.% Cu) compared to Cu6Sn5 (54.5 at.% Cu). The g6 structure was found above 350 °C and has ˚, the space group C2 and cell parameters of a = 12.60 A 8 ˚ ˚ b = 7.27 A, c = 10.20 A and b = 90.0°, while g is stable above 186 °C and has space group P21/c and cell parame˚ , b = 7.27 A ˚ and b = 62.5°. Similarly, ters of a = c = 9.83 A 0 the g -phase region is again a narrow band instead of a line, as indicated by the green dotted arrow in Fig. 1; thus there is a possibility that different crystal structures may exist in the g0 -phase. In fact, the X-ray diffraction (XRD) pattern of the low-temperature phase containing small peaks could not be indexed as the existing g0 in the previous work [20], providing the motivation for a more thorough understanding of the modulation of the g0 -phase. This paper investigates the crystal structure of directly alloyed stoichiometric Cu6Sn5 using variable-temperature synchrotron XRD and transmission electron microscopy (TEM). A new monoclinic-based phase was identified in the stoichiometric Cu6Sn5. The results provide fundamental insights into the properties of directly alloyed stoichiometric Cu6Sn5, which are of significance in the field of 1 For interpretation of color in Fig. 1, the reader is referred to the web version of this article.
High-purity stoichiometric Cu6Sn5 intermetallic phases were obtained by direct alloying. Stoichiometric ratios of high-purity copper and tin wires (Alfa Aesar) were stored in evacuated and sealed quartz ampules, which were then heated to 700 °C in an electric resistance furnace and held at this temperature overnight. The ampules were then quenched in water and heat treated by annealing at 400 °C for 528 h before cooling to room temperature in air together with the crucible, which allowed a slow cooling. As a comparison, Cu6Sn5 samples were also prepared by dissolving the Sn from Sn–4Cu ingot bars supplied by Nihon Superior Co. using selective etching, as detailed elsewhere [5]. Another sample of stoichiometric Cu6Sn5 with 2 at.% Ni addition was prepared under the same procedure by adding 2 at.% high-purity Ni. 2.2. Synchrotron XRD The alloyed samples were crushed in an agate mortar and loaded into a quartz capillary sample cell (0.3 or 0.5 mm in diameter) for the XRD experiments. Highresolution X-ray powder diffraction data were collected in the 2h range of 10–80° on the Powder Diffraction Beamline at the Australian Synchrotron using the Mythen-II detector with a 15 keV beam. For the experiments at elevated temperature, samples were heated using a hot-air stream from 30 to 170, 200, 220 and 250 °C sequentially, at a rate of 6 °C min1. The samples were maintained at each temperature of interest for 1 min for stabilization purposes followed by 10 min for data collection. For calibration, a Si standard (NIST640C) cell was measured for 3 min at room temperature. The wavelength calibrated by the Si standard is 0.07728 nm. The Rietveld refinement was performed using TOPAS v. 4.2. 2.3. TEM characterization Samples for TEM characterization were prepared using the crushed sample to ensure that the samples were exactly the same as that used for the synchrontron XRD experiment. The specimens were prepared by further crushing under ethanol in a mortar and pestle and then depositing the suspension of crushed crystal onto a holey carbon film and air drying. The specimens were also prepared using a focused ion beam (FIB) technique on an FEI Quanta 3D 200 system. Specimens approximately 5 5 lm2 in size having a thin region with a thickness of <100 nm were obtained. It should be noted that the grain size of the directly alloyed sample is in the range of hundreds of micrometers, and therefore the specimens fabricated for
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TEM observation were most likely to be from a single grain. High-resolution TEM images of the samples were obtained using a JEOL 2100 with a HR pole-piece (Cs = 1.0 mm, structure resolution limit = 0.23 nm, information limit = 0.19 nm) and JEOL JEM3200F with an X filter, operating at 300 kV. The TEM images were captured on a Gatan Orius 2k 4k SC1000 bottom-mount CCD camera (JEOL 2100) or imaging plates (JEM3200F). High-resolution lattice images were obtained for crystals tilted to a number of zone-axis orientations.
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Table 1 Relationship of indexing between hexagonal NiAs and monoclinic g8Cu5Sn4 structures. NiAs structure (hexagonal, P63/mmc, No. 194) ½1 1 0 [0 1 0] = [1 1 0] ½2 2 1 ½2 2 1 ½6 2 1 ½2 6 1
Cu5Sn4-g8 structure (monoclinic, P2i/c, No. 14) [0 1 0] ½1 0 1 ½1 4 1 ½1 4 1 ½0 4 1 ½1 4 0
3. Results and discussion
Fig. 2 shows the XRD patterns for stoichiometric Cu6Sn5 as the temperature varies from 30 to 250 °C. In order to magnify the relatively weak superstructure peaks, the heights for the dominating peaks have been truncated and only the 2h range from 14° to 25° is shown. A significant number of small peaks can be observed in the XRD pattern acquired at room temperature (30 °C), indicating a possible monoclinic phase. This monoclinic phase was thermally stable up to 200 °C. The small peaks of the monoclinic phase disappeared when the sample was heated to 220 °C or above, indicating that a phase transformation occurred at a temperature in the range of 200–220 °C. The main peaks of the XRD pattern acquired at 220 °C or above were identified as the hexagonal g-Cu6Sn5, indicating the occurrence of a phase transformation from a monoclinic phase to a hexagonal phase. It is of interest to note that the polymorphic transformation temperature is significantly higher than the 186 °C at which the well-documented g0 –g phase transformation occurs [5,11,18]. The different phase transformation temperature might indicate a different phase transformation route, suggesting that this monoclinic phase could differ from the monoclinic g0 phase. This observation is not without precedent and reactions in
this vicinity of 210 °C in the e-Cu3Sn and Cu6Sn5 phase field were debated in an early publication [21] but seem to have been neglected since. It should be noted that the slight shift of the whole XRD pattern to the lower 2h angle at higher temperatures was attributed to the thermal expansion of the lattice at the elevated temperature. The difference between the two monoclinic phases (i.e. the newly discovered phase in this work compared to the g0 -Cu6Sn5) is verified by comparing the room-temperature XRD patterns for stoichiometric Cu6Sn5 and Cu6Sn5 prepared by the alternative route of selectively dissolving Sn, as shown in Fig. 3a and b, respectively. It is generally accepted that the g0 phase exists in this latter Cu6Sn5 sample at room temperature [16] and the corresponding XRD pattern can be indexed matching the g0 phase [7], as shown in Fig. 3b; however, the XRD spectrum in Fig. 3a cannot be indexed to g0 [16,22], suggesting that a new monoclinic phase was formed in stoichiometric Cu6Sn5. Two extra peaks at around 18.51° and 20.57° started to emerge at a temperature of 200 °C and became obvious at 220 °C, as shown in Fig. 2. These small peaks can be attributed to the diffractions from the e-Cu3Sn phase which is also frequently observed at the interface between the soldered joints and Cu substrates [23,24]. In this work, the occurrence of this phase indicates the composition of the IMC is likely to be on the Cu-rich side of the phase-field.
Fig. 2. XRD peak profile for stoichiometric Cu6Sn5 at temperatures ranging from 30 to 250 °C. The background was subtracted.
Fig. 3. XRD peak profile with the background subtracted for (a) stoichiometric Cu6Sn5 and (b) dissolved Cu6Sn5 at room temperature.
3.1. XRD analysis
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3.2. Crystal structure determination by TEM characterization The lack of existing XRD data makes the indexing of this new phase difficult. Thus, further investigations using selected-area electron diffraction (SAED) and highresolution transmission electron microscopy (HRTEM) were performed. The images and diffraction patterns were indexed based upon the hexagonal NiAs-type parent structure of g phase, as well with reference to the g8 phase. The relationship of directions indexed relative to hexagonal and monoclinic phases is given in Table 1. Fig. 4a shows a diffraction pattern in the [0 1 0]H zone ð¼ ½1 0 1g8 Þ acquired using the finely crushed sample. The diffraction patterns are indexed using the parent NiAs-type
hexagonal structure. The structure has an 8 d114 superlattice (associated with the 1=8ð114ÞH diffraction spot), which is characteristic of g8-Cu5Sn4. The diffraction spots of Fig. 4a are quite sharp, which indicates that this region of crystal is reasonably strain free. The strain-free appearance of diffraction patterns is not typical of the sample. Fig. 4b shows a more typical example of diffraction in the [0 1 0]H zone. In this region of the crystal the parent structure is in the same orientation as Fig. 4a but the superlattice is in a twin-related orientation. That is, the superlattice associated with the 1=8ð1 1 4ÞH diffraction spot is twinrelated with respect to the 1=8ð1 1 4ÞH superlattice of Fig. 4a. The diffraction spots are not sharp, which indicates that there is a large amount of strain in the lattice. The least strain occurs on the (1 1 2)H spots; however, the (1 1 0)H and
zone (=[1 0 0]g8). (a) Diffraction pattern showing 1=8ð1 1 4Þ superlattice. (b) Diffraction with twin-related Fig. 4. Diffraction and images in the ½1 10 H H superlattice = 1=8ð1 1 4ÞH . (c) High-resolution lattice image of area associated with the diffraction pattern in (a). (d) The image of (c) Fourier processed to include only the (monoclinic) superlattice while excluding the hexagonal sublattice. (e) A schematic drawing of the g8-Cu5Sn4 structure.
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Fig. 5. Diffraction patterns in [0 1 0]H zone ð¼ ½1 0 1g8 Þ. A comparison of (a) Cu6Sn5–0 at.% Ni (monoclinic) and (b) (Cu,Ni)6Sn5–2 at.% Ni (hexagonal). The monoclinic phase has a high level of strain and the hexagonal phase does not.
superlattice spots are misshapen, which indicates that the lattice has strain. The degree of lattice strain in the monoclinic lattice of our sample (with nominal composition Cu6Sn5–0 at.% Ni) can be best appreciated by comparing it with another sample (with nominal composition (Cu,Ni)6Sn5–2 at.% Ni) which maintains the hexagonal structure at room temperature (as will be discussed elsewhere [25]). Fig. 5a shows diffraction in the [0 1 0]H zone for the monoclinic sample. There is a great deal of strain in the lattice as evidenced by the misshapen diffraction spots, and there is latticepoint splitting on some of the spots. Fig. 5b shows diffraction from the [0 1 0]H zone for the hexagonal sample. The diffraction pattern is much cleaner, which indicates that there is no lattice strain. An image associated with the Fig. 4a diffraction pattern is show in Fig. 4c. In the image the superlattice fades under continued exposure to the beam. It is clear that the extra copper atoms associated with the superlattice are moving about and are being randomized by the intense electron beam (the beam intensity is far greater for imaging as compared with diffraction). In Fig. 4c the 8 ð1 1 4ÞH spacing is hardly visible, although the 8 (3 3 4)H superlattice remains strong. In Fig. 4d the image of Fig. 4c has been processed to remove the sublattice periodicity and to enhance the superlattice. In Fig. 4d the 8 ð1 1 4ÞH is visible but is very much weaker than the 8 (3 3 4)H. It is worth noting that the 8 (3 3 4)H superlattice is common to both g8-Cu5Sn4 and g0 -Cu6Sn5, but the g0 -Cu6Sn5 structure has a 5 (1 1 4)H superlattice instead of 8 (1 1 4)H. Fig. 4e shows a schematic drawing of the g8-Cu5Sn4 structure. The interstitial Cu atoms are labeled as i-Cu. From this figure it can be seen that the 8 ð1 1 4ÞH (hexagonal cell) is equivalent to (1 0 0)g8 (monoclinic cell); and 4 (1 1 4)H is equivalent to (0 0 2)g8. In g8-Cu5Sn4 there is a double layer of the interstitial Cu associated with the 8 ð1 1 4ÞH spacing, and this double-layer is easily disrupted under the influence of a high-intensity electron beam whereas the 8 (3 3 4)H is more stable under the beam. The structures of g8-Cu5Sn4 and g0 -Cu6Sn5 are closely related. Fig. 6a and b shows schematic visualizations of
g8-Cu5Sn4 and g0 -Cu6Sn5, respectively. In these schematics each non-interstitial Cu atom is surrounded by six Sn atoms in an octahedral arrangement. The unit cell dimensions c and b are the same in both structures and a is different. The repeat distance of the g8-Cu5Sn4 Cu interstitial double layer along the a direction is 8 (1 1 4)H, whereas the repeat distance of the g0 -Cu6Sn5 Cu interstitial single layer along the a direction is 5 (1 1 4)H. In order to assist in analyzing the structures we use the scheme of Larsson et al. [16,19], where the structure is represented in terms of the stacking of the Cu–Sn octahedra. The g8-Cu5Sn4 schematic of Fig. 6a is drawn again using the scheme of Larsson et al. [19] in Fig. 6c. The bright and dark diamonds represent the areas where a pair of octahedra are adjacent and stacked face-to-face (light) and areas in between octahedra (dark). Based on the data discussed so far it has been possible to conclude that the structure of the previously unreported low-temperature phase in the stoichiometric Cu6Sn5 is partly disordered, which is reflected by the amount of strain in the lattice, and that the structure is similar to g8-Cu5Sn4. However, it has been possible to go further than this via a careful analysis of a set of diffraction data from a particularly well-ordered crystal. The set of data in Fig. 7 acquired from the sample made by FIB shows a diffraction tilt series which has SAED of the same crystal in seven different zone axis orientations. It is interesting to see that in the set of diffraction patterns in Fig. 7 the superlattice is incommensurate with the sublattice. Normally an incommensurate superlattice cannot be related to a fixed atomic arrangement, but in this case it can be. In fact, the incommensurate superlattice is due to a regular intergrowth of g8-Cu5Sn4 and g0 -Cu6Sn5 as will be described further below. In order to derive a structure based upon the diffraction patterns of Fig. 7 we need to look closely at the incommensurate spacing. Fig. 8a is a ½1 10H zone diffraction pattern and is a close-up of the diffraction pattern of Fig. 7b. It should be noted that the diffraction patterns shown in Fig. 8 were rotated to make the systematic row of (1 1 4)H reflections horizontal. The incommensurate nature of the superlattice is quite obvious when sighting along a line
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Fig. 6. Structural drawings and comparison of (a) g8-Cu5Sn4 structure and (b) g0 -Cu6Sn5 structure. (c) The g8-Cu5Sn4 schematic of (a) drawn again using the scheme of Larsson et al. [19].
between the center of the diffraction pattern and the (1 1 4)H spot. Starting from the center, spots 1, 2, 4, 5 and 7 are shifted slightly to the left or right of the line and spots 3 and 6 are on the line. That is, the superlattice planes are not parallel with (1 1 4)H planes but rather are tilted by approximately 1.5° relative to the (1 1 4)H. The expected superlattice d-spacing for the 8 (1 1 4)H is 0.872 nm, but
the measured d-spacing of the actual superlattice is smaller at 0.805 ± 0.003 nm. The ratio of the measured to expected is 1.08 ± 0.005. For comparison with Fig. 8a and b shows a simulated 1 0H zone diffraction pattern for the g8-Cu5Sn4 in a ½1 (=[0 1 0]g8). The g8 diffraction pattern has the exact 8 (1 1 4)H superlattice pattern as expected. For g8 the superlattice spacing is different, 0.872 nm for g8 compared with 0.805 nm in Fig. 8a, but the superlattice spot intensities are similar. It is clear that the crystal structure is close to g8 diffraction in spite of the incommensurate superlattice. If we sight along the line towards the ð1 1 4ÞH spot on Fig. 8a (the other set of 1 1 4 planes in the ½1 1 0H zone) it is seen that the 1/4ð1 1 4ÞH superlattice spot does not lie on the line. However, the 1/4ð1 1 4ÞH distance is exactly 1/4 of the distance between the central spot and ð1 1 4ÞH . The crystal tilted to the ½6 2 1H zone (Fig. 8c) contains the systematic diffraction line from the (1 1 4)H that is common to (1 1 4)H in the ½1 10H diffraction pattern (Fig. 8a). In the ½6 2 1H the conclusions about the incommensurabity of the 8 (1 1 4)H are the same. The superlattice planes are tilted by approximately 1.5° relative to the (1 1 4)H and the ratio of measured to expected 8 (1 1 4)H is 1.08 ± 0.005. A simulation of g8 in ½6 2 1H ð¼ ½0 4 1g8 Þ can be compared with the diffraction pattern, as shown in Fig. 8d. The superlattice spot intensities of g8 are similar to the experiment. The crystal tilted to the ½2 6 1H zone (Fig. 8e) contains the systematic diffraction line from the ð1 1 4ÞH that is the same as ð1 1 4ÞH in the ½1 1 0H diffraction pattern (Fig. 8a). In this diffraction pattern there is a superlattice of 1=8ð1 1 4ÞH , whereas in the ½1 1 0H the superlattice was 1=4ð1 1 4ÞH . The superlattice is exactly 1/8 of the ð1 1 4ÞH spacing but the superlattice spots are not on the line between the center of the diffraction pattern and ð1 1 4ÞH but lie slightly to the left or right of that line. How can we interpret the incommensurate superlattice? The ratio of the measured to the expected superlattice spacing is an important clue. The ratio 1.08 is very close to the ratio 40/37 (40/37 = 1.081). It was noted earlier (Fig. 6 and discussion thereof) that g8-Cu5Sn4 unit cell has a period of 8 (1 1 4)H between Cu interstitial double layers, and that g0 -Cu6Sn5 has a period of 5 (1 1 4)H between layers. The supercell (Fig. 9) is incommensurate by a factor of 37/40 which can be obtained by a periodic stacking of four g8Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell. We call this a 4 + 1 block (4 by g8 = 32, plus 1 by g0 = 5; 32 + 5 = 37), thus the new monoclinic phase will be labeled as g4+1. In Fig. 9 the 4 + 1 block is drawn using the scheme of Larsson et al. [16,19]. At first pass the diagram may look a bit complicated; however, it is based upon a simple system. In drawing out the 4 + 1 block structure it is clear that we need to shift the second 4 + 1 block by 1/8[3 0 1]g8. In addition, the second 4 + 1 block is inverted along the [0 1 0]g8 direction (i.e. [0 1 0]g8 becomes ½0 10g8 Þ, so there is doubling of the unit cell along [1 0 0]S. The unit cell of the 2 (4 + 1) superlattice is [1 0 0]S = 9.25[1 0 0]g8; [0 1 0]S =
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Fig. 7. Series of SAED patterns at different zone axes. In the sequence of (a) ! (d) ! (f), the diffraction spots pair (marked by the red arrows and A) were kept visible when rotating the sample. Similarly, in the sequences of (a) ! (b) ! (c) and (a) ! (e) ! (g), diffraction spots pairs (marked by green arrows and B, yellow arrows and C) were kept visible during rotating, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
[0 1 0]g8; and [0 01]S = [0 01]g8. The detailed coordinates of atoms in the unit cell can be found in the online Supplementary information. In Fig. 9 the unit cell outlines of the g8 and one g0 unit cells are shown and the legend on the figure shows the height along the b-axis (y-direction) of the interstitial Cu atoms. In the figure the average spacing between the (circled) g8 Cu-pairs is (37/40) 8 d114. Moreover, the plane upon which the circled Cu-pairs lie is, on average, rotated by 1° relative to the ð3 3 4ÞH planes. It is the average spacing of these g8 Cu-pairs that gives rise to the incommensurate superlattice. The 4 + 1 block system used in Fig. 9 gives a structure which can explain the incommensurability of the superlattice, i.e. the structure explains the 40/37 ratio and the approximate 1° rotation relative the expected (1 1 4)H alignment. However, the structure of Fig. 9 is an idealization of the actual structure. If the ideal structure did occur exactly as postulated in Fig. 9, then there would be a series of very fine superlattice spots in the diffraction patterns with a spacing of 1/37(1 1 4)H, and these are not observed. It seems that the
4 + 1 blocks are the average structure: the actual structure has a degree of randomization, where the average 2 (4 + 1) may be as a result of 5 + 1 followed by 3 + 1, or 6 + 1 followed by 2 + 1, for example. A simulated diffraction pattern of the g4+1 monoclinic structure is shown schematically in Fig. 10. The simulation is for the [0 1 0] orientation, for comparison with Fig. 8a. The simulation shows splitting of diffraction spots (arrowed) which is similar to the splitting in the experimental diffraction pattern. The diffraction spots are not evenly spaced along the (1 1 4)H direction as shown by a comparison of the distances marked in Fig. 10. In the simulation the ratio of distance second relative to first is 1.82 (i.e. <2.0). In the experimental diffraction pattern (Fig. 8a) the ratio is 1.80. The crystal with the incommensurate superlattice that we found provides us with an explanation for the strain that is observed in the majority of diffraction patterns from our samples with the nominal Cu6Sn5 composition. It appears likely that the strain arises from the mixing of one g0 with several g8 layers where it is the g8 that provides the predominant contribution to the superlattice.
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Fig. 8. Close-ups for analysis of some of the diffraction patterns from Fig. 7. (a) ½1 1 0H on zone diffraction pattern. (b) Simulated diffraction pattern for 1 0H (=[0 1 0]g8). (c) ½6 2 1H on zone diffraction pattern. (d) Simulated diffraction pattern for the g8 in ½6 2 1H ð¼ ½0 4 1g8 Þ. (e) ½2 6 1H on zone the g8 in ½1 diffraction pattern.
3.3. Reanalysis of the XRD pattern using the proposed atomic model After developing the atomic model of the new monoclinic phase (see online Supplementary information), Rietveld refinement was performed on the XRD profile for stoichiometric Cu6Sn5 at a temperature of 30 °C. In Fig. 11a, the black profile is the experimental result; the red profile shows the refinement result. The XRD pattern at 2h ranging from 14° to 23° is enlarged in the inset and it is apparent that even the small peaks are well described by the model. The Rietveld refinement gave Rp = 3.55 ˚, and Rwp = 5.35 with cell parameters a = 92.2411 A ˚ ˚ b = 7.3108 A, c = 9.8800 A and b = 118.9521°. The detailed atomic coordinates in the unit cell after Rietveld refinement can be found in the online Supplementary information, along with those from the model.
Rietveld refinements for the XRD profile of stoichiometric Cu6Sn5 at room temperature were also performed using other phases (g, g6, g8, g0 ) for comparison and the refinement results are shown in Fig. 11b along with the refinement result using g4+1. The refinements gave Rwp values of 7.11, 7.06, 6.29, 6.35 and 5.35 for g, g6, g8, g0 and g4+1 phases, respectively. It is clear that the g4+1 shows the best refinement in terms of Rwp and curve fitting, suggesting that the proposed crystal structure for the new phase g4+1 is sound. It is interesting to note that g8 also gives a reasonably good refinement, which is expected. The g8 phase is believe to be stable at elevated temperature (higher than 186 °C) [19], while in contrast, the new phase transforms to g phase at a temperature between 200 and 220 °C. The XRD profiles for stoichiometric Cu6Sn5 at temperatures ranging from 170 to 250 °C were also refined using Rietveld analysis. The proposed new monoclinic phase
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Fig. 9. The proposed superlattice structure based upon the 4 + 1 blocks (four g8-Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell). The g8 and g0 unit cells are drawn in outline, and the relative orientation and unit cell axes of the hexagonal cell, the g8 cell and the g4+1 cell are given.
Fig. 10. Simulation of diffraction pattern for the [0 1 0]H orientation of g4+1 phase. The double-ended arrows indicate the relative spacing of the first and second superlattice spots. Superlattice splitting (single-ended arrows) is not present in g8 (Fig. 8b) but is present in the experimental diffraction pattern (Fig. 8a).
was used for XRD profiles at temperatures of 170 and 200 °C, while hexagonal g phase was used for temperatures of 220 and 250 °C. The refinement results are shown in Table 2. Clearly, the dominant phase at temperatures below 200 °C was the monoclinic g4+1 phase, while the main phase at temperatures higher than 220 °C is the hexagonal g phase. 3.4. Discussion on the formation mechanism of g4+1 phase The importance of the finding and structure determination of this new monoclinic phase should be emphasized. It
Fig. 11. (a) Rietveld refinement of the XRD peak profile for stoichiometric Cu6Sn5 at a temperature of 30 °C: the blue profile is the experimental result, while the red one is the refinement result (the differential is also displayed). The XRD pattern at the 2h ranging from 14° to 23° is enlarged in the inset, showing the fit for the smaller peaks. (b) Rietveld refinements for XRD profile of stoichiometric Cu6Sn5 at a temperature of 30 °C using different phases g, g6, g8, g0 , g4+1 from top to bottom. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Table 2 Unit cell parameters of the XRD profile for stoichiometric Cu6Sn5 at temperatures ranging from 30 to 250 °C after Rietveld refinement. ˚) ˚) ˚) Temperature (°C) Crystal structure Space group a (A b (A c (A a (°) b (°) c (°) Rp 4þ1
30 170 200 220 250
g g4þ1 g4þ1 g g
P1 P1 P1 P63/mmc P63/mmc
92.2411 92.5664 92.6589 4.2372 4.2345
is widely assumed that the Cu6Sn5 phase in the near-eutectic solder alloys is identical to that at the interface of the solder–substrate joint [18], although with identification of a new monoclinic phase in the directly alloyed stoichiometric Cu6Sn5 samples, this assumption must be questioned. Most previous studies have assumed the intermetallic Cu6Sn5 was of the g0 phase, although this may only be the case for Cu6Sn5 forming as part of a eutectic reaction [7,18,26], and may not hold true for that formed at the solder–substrate joint [27,28]. Numerous processing routes can be used to obtain Cu6Sn5 including primary, eutectic and peritectic solidification reactions, solid-state diffusion and other synthetic routes such as spray pyrolysis [13] and ball-milling [29]. Differences in crystal structure may result from compositional variations or other influences related to these processing routes. It is still unclear why the new monoclinic phase only exists in the directly alloyed stoichiometric Cu6Sn5 samples while it is absent in the samples of Cu6Sn5 that have been collected by selectively dissolving the Sn phase from near-eutectic solder alloys. Further work is required to clarify the nature of the new monoclinic phase and the formation mechanisms. 4. Conclusions The following conclusions can be drawn from this work: A new monoclinic-based phase with cell parameters of ˚ , b = 7.311 A ˚ , c = 9.880 A ˚ and b = 118.95° a = 92.241 A was formed in directly alloyed stoichiometric Cu6Sn5 samples. The new monoclinic phase has structurally similar features to both g0 -Cu6Sn5 and g8-Cu5Sn4, and thus can be treated as a modulation of these two phases. It has been labeled as g4+1 because, on average, it is represented by the periodic stacking of four g8-Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell (i.e. a 4 + 1 block). g4+1 phase transformed to a hexagonal g-Cu6Sn5 phase at a temperature between 200 and 220 °C, which is higher than the 186 °C at which the well-documented g0 –g allotropic phase transformation occurs. These findings contribute to a growing body of knowledge revealing the complex behavior of the Cu6Sn5 intermetallic. In particular, they show that the composition, the thermal history and the processing or fabrication route strongly influence the structure adopted by Cu6Sn5. A more thorough understanding of the underlying
7.3108 7.3314 7.3347 a=b a=b
9.8800 9.9062 9.9104 5.1349 5.1339
90 90 90 90 90
118.9521 118.9399 118.9453 90 90
90 90 90 120 120
3.55 3.76 4.03 2.92 2.96
Rwp 5.35 6.07 6.55 5.19 5.35
mechanisms could have significant implications for leadfree solder alloy development. 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 program (project ID: LP100200250) and JSPS Fellowship for Research in Japan (project ID: S11730). XRD experiments were performed at the Australian Synchrotron (AS) Powder Diffraction Beamline (project ID: AS101/PD2249 and AS111/PD3364). TEM experiments were performed at the Centre for Microscopy and Microanalysis (CMM) at UQ and the Research Laboratory for High Voltage Electron Microscopy (HVEM) at Kyushu University (KU) with Nanotechnology Network program Japan (project ID: T23-019). J.C.B. acknowledges the support of and use of facilities at QUT. The authors would like to thank Mr. J. Read (UQ) and Mr. D. Mu (UQ) for help on XRD experiments and Mr. T. Daio (KU) and Dr. E. Tanaka (KU) on TEM experiments. Y.W. is funded by the ARC-Linkage project. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/ j.actamat.2012.08.024. References [1] Hsiao HY, Liu CM, Lin HW, Liu TC, Lu CL, Huang YS, et al. Science 2012;336:1007. [2] Zeng K, Tu KN. Mater Sci Eng R 2002;38:55. [3] Li Y, Wong CP. Mater Sci Eng R 2006;51:1. [4] Laurila T, Vuorinen V, Paulasto-Krockel M. Mater Sci Eng R 2010;68:1. [5] Nogita K. Intermetallics 2010;18:145. [6] Nogita K, Nishimura T. Scripta Mater 2008;59:191. [7] Schwingenschlogl U, Di Paola C, Nogita K, Gourlay CM. Appl Phys Lett 2010;96:061908. [8] Suh JO, Tu KN, Lutsenko GV, Gusak AM. Acta Mater 2008;56:1075. [9] Suh JO, Tu KN, Tamura N. Appl Phys Lett 2007;91:051907. [10] Li JF, Agyakwa PA, Johnson CM. Acta Mater 2011;59:1198. [11] Ghosh G, Asta M. J Mater Res 2005;20:3102. [12] Sarakonsri T, Apirattanawan T, Tungprasurt S, Tunkasiri T. J Mater Sci 2006;41:4749. [13] Ju SH, Jang HC, Kang YC. J Power Sources 2009;189:163.
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Hu RZ, Zeng MQ, Zhu M. Electrochim Acta 2009;54:2843. Saunders N, Miodownik AP. Bull Alloy Phase Diagrams 1990;11:278. Larsson AK, Stenberg L, Lidin S. Acta Crystallogr B 1994;50:636. Peplinski B, Schulz G, Schultze D, Schierhorn E. Mater Sci Forum 1996;228:577. Li MY, Zhang ZH, Kim JM. Appl Phys Lett 2011;98:201901. Larsson AK, Stenberg L, Lidin S. Z Kristallogr 1995;210:832. Nogita K, Mu D, McDonald SD, Read J, Wu YQ. Intermetallics 2012;26:78. Haughton JL. J Inst Metal 1921;25:309 [Annual General Meeting, London].
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[22] Lidin S, Larsson A-K. J Solid State Chem 1995;118:313. [23] Shang PJ, Liu ZQ, Li DX, Shang JK. Scripta Mater 2008;59: 317. [24] Shang PJ, Liu ZQ, Pang XY, Li DX, Shang JK. Acta Mater 2009;57:4697. [25] Wu YQ, McDonald SD, Huang H, Read J, Nogita K, in preparation. [26] Mardare CC, Hassel AW. Phys Status Solidi A 2012;209:825. [27] Zou HF, Yang HJ, Zhang ZF. Acta Mater 2008;56:2649. [28] Tu KN. Acta Metall 1973;21:347. [29] Zhang J, Zhang Y, Zhang X, Xia Y. J Power Sources 2007;167:171.
Acta Materialia 60 (2012) 6581–6591 www.elsevier.com/locate/actamat
A new phase in stoichiometric Cu6Sn5 Y.Q. Wu a,b, J.C. Barry c, T. Yamamoto d, Q.F. Gu e, S.D. McDonald a,b, S. Matsumura d, H. Huang b, K. Nogita a,b,⇑ a
Nihon Superior Centre for the Manufacture of Electronic Materials (NS CMEM), The University of Queensland, QLD 4072, Australia b School of Mechanical and Mining Engineering, The University of Queensland, QLD 4072, Australia c Science and Engineering Faculty, Queensland University of Technology, QLD 4000, Australia d Department of Applied Quantum Physics and Nuclear Engineering, Kyushu University, Motooka 744, Nishi-ku, Fukuoka 819-0395, Japan e Powder Diffraction Beamline, The Australian Synchrotron, Clayton, VIC 3168, Australia Received 19 June 2012; accepted 12 August 2012 Available online 17 September 2012
Abstract The intermetallic compound Cu6Sn5 is a significant microstructural feature of many electronic devices where it is present at the solder–substrate interfaces. The time- and temperature-dependent thermomechanical properties of Cu6Sn5 are dependent on the nature and stability of its crystal structure, which has been shown to exist in at least four variants (g, g0 , g6 and g8). This research details an additional newly identified monoclinic-based structure in directly alloyed stoichiometric Cu6Sn5 using variable-temperature synchrotron X-ray diffraction (XRD) and transmission electron microscopy. The phase is associated with a departure from the equilibrium temperature of the polymorphic monoclinic–hexagonal transformation temperature. The new monoclinic phase can be treated as a modulation of four ˚ , b = 7.311 A ˚, g8-Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell. It has been labeled as g4+1 and has cell parameters of a = 92.241 A ˚ and b = 118.95° determined from electron diffraction patterns. The XRD results could be fitted well to a Rietveld refinement c = 9.880 A using the new crystal parameters. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Cu6Sn5; Phase transformation; X-ray diffraction (XRD); Electron diffraction pattern; Crystallography
1. Introduction Electronic devices containing soldered components are ubiquitous in modern life. Due to environmental concerns, the majority of solder alloys have changed composition to eliminate the presence of lead [1–4]. Numerous Pb-free solder alloys are currently in use or being developed, including a range of Sn–Cu-based solder alloys [5–9]. In these alloys and many other Sn-based solders significant quantities of intermetallic compound (IMC) of Cu6Sn5 form during processing and cooling. As the scale of devices is minimized, the IMC occupies an increasing volume fraction of the solder joint and, in an extreme situation, all of the Sn might be consumed to form the IMC [10]. The IMC layer is integral ⇑ Corresponding author.
E-mail address: [email protected] (K. Nogita).
to microelectronic solder reliability and a fundamental understanding of Cu6Sn5 crystallography, which influences chemical, electrical and mechanical properties [11], is essential to the reliable manufacture and service of a large number of electronic devices. In addition, Cu6Sn5 has been proposed as a promising anode material in Li-ion batteries [12–14] and large quantities of Cu6Sn5 material would be required for this application. Cu6Sn5 (54.5 at.% Cu) exists in two different crystal structures according to the binary Cu–Sn phase diagram [15], as partly shown in Fig. 1. At equilibrium, monoclinic ˚, g0 -Cu6Sn5 (C2/c) with cell parameters of a = 11.022 A ˚ ˚ b = 7.282 A, c = 9.827 A, b = 98.84° [16] is the lower-temperature phase, while hexagonal g-Cu6Sn5 (P63/mmc) with ˚ and c = 5.086 A ˚ [17] is stacell parameters of a = 4.190 A 0 ble at higher temperatures. A g–g allotropic phase transformation occurs as the temperature falls below 186 °C
1359-6454/$36.00 Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.actamat.2012.08.024
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electronic materials, including solder alloys and Li-ion battery technology. 2. Experimental details 2.1. Preparation of samples
Fig. 1. Part of the binary phase diagram of the Cu–Sn system (from Ref. [15]).
[5,11,18]. It can be clearly seen from the phase diagram, as indicated by the red1 solid arrow in Fig. 1, that the g-phase region is a narrow band rather than a line, and thus composition variation does occur in the g-phase region. In addition, it is noted that the phase field is skewed towards Cu-rich compositions as the temperature increases. The composition range for these phases may result in the potential for crystal structure variations. The high-temperature modulation of g phase above 186 °C has been systematically studied by Larsson et al. [19] and two superstructures were proposed, namely g6 and g8, which are slightly more Cu rich (55.56 at.% Cu) compared to Cu6Sn5 (54.5 at.% Cu). The g6 structure was found above 350 °C and has ˚, the space group C2 and cell parameters of a = 12.60 A 8 ˚ ˚ b = 7.27 A, c = 10.20 A and b = 90.0°, while g is stable above 186 °C and has space group P21/c and cell parame˚ , b = 7.27 A ˚ and b = 62.5°. Similarly, ters of a = c = 9.83 A 0 the g -phase region is again a narrow band instead of a line, as indicated by the green dotted arrow in Fig. 1; thus there is a possibility that different crystal structures may exist in the g0 -phase. In fact, the X-ray diffraction (XRD) pattern of the low-temperature phase containing small peaks could not be indexed as the existing g0 in the previous work [20], providing the motivation for a more thorough understanding of the modulation of the g0 -phase. This paper investigates the crystal structure of directly alloyed stoichiometric Cu6Sn5 using variable-temperature synchrotron XRD and transmission electron microscopy (TEM). A new monoclinic-based phase was identified in the stoichiometric Cu6Sn5. The results provide fundamental insights into the properties of directly alloyed stoichiometric Cu6Sn5, which are of significance in the field of 1 For interpretation of color in Fig. 1, the reader is referred to the web version of this article.
High-purity stoichiometric Cu6Sn5 intermetallic phases were obtained by direct alloying. Stoichiometric ratios of high-purity copper and tin wires (Alfa Aesar) were stored in evacuated and sealed quartz ampules, which were then heated to 700 °C in an electric resistance furnace and held at this temperature overnight. The ampules were then quenched in water and heat treated by annealing at 400 °C for 528 h before cooling to room temperature in air together with the crucible, which allowed a slow cooling. As a comparison, Cu6Sn5 samples were also prepared by dissolving the Sn from Sn–4Cu ingot bars supplied by Nihon Superior Co. using selective etching, as detailed elsewhere [5]. Another sample of stoichiometric Cu6Sn5 with 2 at.% Ni addition was prepared under the same procedure by adding 2 at.% high-purity Ni. 2.2. Synchrotron XRD The alloyed samples were crushed in an agate mortar and loaded into a quartz capillary sample cell (0.3 or 0.5 mm in diameter) for the XRD experiments. Highresolution X-ray powder diffraction data were collected in the 2h range of 10–80° on the Powder Diffraction Beamline at the Australian Synchrotron using the Mythen-II detector with a 15 keV beam. For the experiments at elevated temperature, samples were heated using a hot-air stream from 30 to 170, 200, 220 and 250 °C sequentially, at a rate of 6 °C min1. The samples were maintained at each temperature of interest for 1 min for stabilization purposes followed by 10 min for data collection. For calibration, a Si standard (NIST640C) cell was measured for 3 min at room temperature. The wavelength calibrated by the Si standard is 0.07728 nm. The Rietveld refinement was performed using TOPAS v. 4.2. 2.3. TEM characterization Samples for TEM characterization were prepared using the crushed sample to ensure that the samples were exactly the same as that used for the synchrontron XRD experiment. The specimens were prepared by further crushing under ethanol in a mortar and pestle and then depositing the suspension of crushed crystal onto a holey carbon film and air drying. The specimens were also prepared using a focused ion beam (FIB) technique on an FEI Quanta 3D 200 system. Specimens approximately 5 5 lm2 in size having a thin region with a thickness of <100 nm were obtained. It should be noted that the grain size of the directly alloyed sample is in the range of hundreds of micrometers, and therefore the specimens fabricated for
Y.Q. Wu et al. / Acta Materialia 60 (2012) 6581–6591
TEM observation were most likely to be from a single grain. High-resolution TEM images of the samples were obtained using a JEOL 2100 with a HR pole-piece (Cs = 1.0 mm, structure resolution limit = 0.23 nm, information limit = 0.19 nm) and JEOL JEM3200F with an X filter, operating at 300 kV. The TEM images were captured on a Gatan Orius 2k 4k SC1000 bottom-mount CCD camera (JEOL 2100) or imaging plates (JEM3200F). High-resolution lattice images were obtained for crystals tilted to a number of zone-axis orientations.
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Table 1 Relationship of indexing between hexagonal NiAs and monoclinic g8Cu5Sn4 structures. NiAs structure (hexagonal, P63/mmc, No. 194) ½1 1 0 [0 1 0] = [1 1 0] ½2 2 1 ½2 2 1 ½6 2 1 ½2 6 1
Cu5Sn4-g8 structure (monoclinic, P2i/c, No. 14) [0 1 0] ½1 0 1 ½1 4 1 ½1 4 1 ½0 4 1 ½1 4 0
3. Results and discussion
Fig. 2 shows the XRD patterns for stoichiometric Cu6Sn5 as the temperature varies from 30 to 250 °C. In order to magnify the relatively weak superstructure peaks, the heights for the dominating peaks have been truncated and only the 2h range from 14° to 25° is shown. A significant number of small peaks can be observed in the XRD pattern acquired at room temperature (30 °C), indicating a possible monoclinic phase. This monoclinic phase was thermally stable up to 200 °C. The small peaks of the monoclinic phase disappeared when the sample was heated to 220 °C or above, indicating that a phase transformation occurred at a temperature in the range of 200–220 °C. The main peaks of the XRD pattern acquired at 220 °C or above were identified as the hexagonal g-Cu6Sn5, indicating the occurrence of a phase transformation from a monoclinic phase to a hexagonal phase. It is of interest to note that the polymorphic transformation temperature is significantly higher than the 186 °C at which the well-documented g0 –g phase transformation occurs [5,11,18]. The different phase transformation temperature might indicate a different phase transformation route, suggesting that this monoclinic phase could differ from the monoclinic g0 phase. This observation is not without precedent and reactions in
this vicinity of 210 °C in the e-Cu3Sn and Cu6Sn5 phase field were debated in an early publication [21] but seem to have been neglected since. It should be noted that the slight shift of the whole XRD pattern to the lower 2h angle at higher temperatures was attributed to the thermal expansion of the lattice at the elevated temperature. The difference between the two monoclinic phases (i.e. the newly discovered phase in this work compared to the g0 -Cu6Sn5) is verified by comparing the room-temperature XRD patterns for stoichiometric Cu6Sn5 and Cu6Sn5 prepared by the alternative route of selectively dissolving Sn, as shown in Fig. 3a and b, respectively. It is generally accepted that the g0 phase exists in this latter Cu6Sn5 sample at room temperature [16] and the corresponding XRD pattern can be indexed matching the g0 phase [7], as shown in Fig. 3b; however, the XRD spectrum in Fig. 3a cannot be indexed to g0 [16,22], suggesting that a new monoclinic phase was formed in stoichiometric Cu6Sn5. Two extra peaks at around 18.51° and 20.57° started to emerge at a temperature of 200 °C and became obvious at 220 °C, as shown in Fig. 2. These small peaks can be attributed to the diffractions from the e-Cu3Sn phase which is also frequently observed at the interface between the soldered joints and Cu substrates [23,24]. In this work, the occurrence of this phase indicates the composition of the IMC is likely to be on the Cu-rich side of the phase-field.
Fig. 2. XRD peak profile for stoichiometric Cu6Sn5 at temperatures ranging from 30 to 250 °C. The background was subtracted.
Fig. 3. XRD peak profile with the background subtracted for (a) stoichiometric Cu6Sn5 and (b) dissolved Cu6Sn5 at room temperature.
3.1. XRD analysis
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3.2. Crystal structure determination by TEM characterization The lack of existing XRD data makes the indexing of this new phase difficult. Thus, further investigations using selected-area electron diffraction (SAED) and highresolution transmission electron microscopy (HRTEM) were performed. The images and diffraction patterns were indexed based upon the hexagonal NiAs-type parent structure of g phase, as well with reference to the g8 phase. The relationship of directions indexed relative to hexagonal and monoclinic phases is given in Table 1. Fig. 4a shows a diffraction pattern in the [0 1 0]H zone ð¼ ½1 0 1g8 Þ acquired using the finely crushed sample. The diffraction patterns are indexed using the parent NiAs-type
hexagonal structure. The structure has an 8 d114 superlattice (associated with the 1=8ð114ÞH diffraction spot), which is characteristic of g8-Cu5Sn4. The diffraction spots of Fig. 4a are quite sharp, which indicates that this region of crystal is reasonably strain free. The strain-free appearance of diffraction patterns is not typical of the sample. Fig. 4b shows a more typical example of diffraction in the [0 1 0]H zone. In this region of the crystal the parent structure is in the same orientation as Fig. 4a but the superlattice is in a twin-related orientation. That is, the superlattice associated with the 1=8ð1 1 4ÞH diffraction spot is twinrelated with respect to the 1=8ð1 1 4ÞH superlattice of Fig. 4a. The diffraction spots are not sharp, which indicates that there is a large amount of strain in the lattice. The least strain occurs on the (1 1 2)H spots; however, the (1 1 0)H and
zone (=[1 0 0]g8). (a) Diffraction pattern showing 1=8ð1 1 4Þ superlattice. (b) Diffraction with twin-related Fig. 4. Diffraction and images in the ½1 10 H H superlattice = 1=8ð1 1 4ÞH . (c) High-resolution lattice image of area associated with the diffraction pattern in (a). (d) The image of (c) Fourier processed to include only the (monoclinic) superlattice while excluding the hexagonal sublattice. (e) A schematic drawing of the g8-Cu5Sn4 structure.
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Fig. 5. Diffraction patterns in [0 1 0]H zone ð¼ ½1 0 1g8 Þ. A comparison of (a) Cu6Sn5–0 at.% Ni (monoclinic) and (b) (Cu,Ni)6Sn5–2 at.% Ni (hexagonal). The monoclinic phase has a high level of strain and the hexagonal phase does not.
superlattice spots are misshapen, which indicates that the lattice has strain. The degree of lattice strain in the monoclinic lattice of our sample (with nominal composition Cu6Sn5–0 at.% Ni) can be best appreciated by comparing it with another sample (with nominal composition (Cu,Ni)6Sn5–2 at.% Ni) which maintains the hexagonal structure at room temperature (as will be discussed elsewhere [25]). Fig. 5a shows diffraction in the [0 1 0]H zone for the monoclinic sample. There is a great deal of strain in the lattice as evidenced by the misshapen diffraction spots, and there is latticepoint splitting on some of the spots. Fig. 5b shows diffraction from the [0 1 0]H zone for the hexagonal sample. The diffraction pattern is much cleaner, which indicates that there is no lattice strain. An image associated with the Fig. 4a diffraction pattern is show in Fig. 4c. In the image the superlattice fades under continued exposure to the beam. It is clear that the extra copper atoms associated with the superlattice are moving about and are being randomized by the intense electron beam (the beam intensity is far greater for imaging as compared with diffraction). In Fig. 4c the 8 ð1 1 4ÞH spacing is hardly visible, although the 8 (3 3 4)H superlattice remains strong. In Fig. 4d the image of Fig. 4c has been processed to remove the sublattice periodicity and to enhance the superlattice. In Fig. 4d the 8 ð1 1 4ÞH is visible but is very much weaker than the 8 (3 3 4)H. It is worth noting that the 8 (3 3 4)H superlattice is common to both g8-Cu5Sn4 and g0 -Cu6Sn5, but the g0 -Cu6Sn5 structure has a 5 (1 1 4)H superlattice instead of 8 (1 1 4)H. Fig. 4e shows a schematic drawing of the g8-Cu5Sn4 structure. The interstitial Cu atoms are labeled as i-Cu. From this figure it can be seen that the 8 ð1 1 4ÞH (hexagonal cell) is equivalent to (1 0 0)g8 (monoclinic cell); and 4 (1 1 4)H is equivalent to (0 0 2)g8. In g8-Cu5Sn4 there is a double layer of the interstitial Cu associated with the 8 ð1 1 4ÞH spacing, and this double-layer is easily disrupted under the influence of a high-intensity electron beam whereas the 8 (3 3 4)H is more stable under the beam. The structures of g8-Cu5Sn4 and g0 -Cu6Sn5 are closely related. Fig. 6a and b shows schematic visualizations of
g8-Cu5Sn4 and g0 -Cu6Sn5, respectively. In these schematics each non-interstitial Cu atom is surrounded by six Sn atoms in an octahedral arrangement. The unit cell dimensions c and b are the same in both structures and a is different. The repeat distance of the g8-Cu5Sn4 Cu interstitial double layer along the a direction is 8 (1 1 4)H, whereas the repeat distance of the g0 -Cu6Sn5 Cu interstitial single layer along the a direction is 5 (1 1 4)H. In order to assist in analyzing the structures we use the scheme of Larsson et al. [16,19], where the structure is represented in terms of the stacking of the Cu–Sn octahedra. The g8-Cu5Sn4 schematic of Fig. 6a is drawn again using the scheme of Larsson et al. [19] in Fig. 6c. The bright and dark diamonds represent the areas where a pair of octahedra are adjacent and stacked face-to-face (light) and areas in between octahedra (dark). Based on the data discussed so far it has been possible to conclude that the structure of the previously unreported low-temperature phase in the stoichiometric Cu6Sn5 is partly disordered, which is reflected by the amount of strain in the lattice, and that the structure is similar to g8-Cu5Sn4. However, it has been possible to go further than this via a careful analysis of a set of diffraction data from a particularly well-ordered crystal. The set of data in Fig. 7 acquired from the sample made by FIB shows a diffraction tilt series which has SAED of the same crystal in seven different zone axis orientations. It is interesting to see that in the set of diffraction patterns in Fig. 7 the superlattice is incommensurate with the sublattice. Normally an incommensurate superlattice cannot be related to a fixed atomic arrangement, but in this case it can be. In fact, the incommensurate superlattice is due to a regular intergrowth of g8-Cu5Sn4 and g0 -Cu6Sn5 as will be described further below. In order to derive a structure based upon the diffraction patterns of Fig. 7 we need to look closely at the incommensurate spacing. Fig. 8a is a ½1 10H zone diffraction pattern and is a close-up of the diffraction pattern of Fig. 7b. It should be noted that the diffraction patterns shown in Fig. 8 were rotated to make the systematic row of (1 1 4)H reflections horizontal. The incommensurate nature of the superlattice is quite obvious when sighting along a line
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Fig. 6. Structural drawings and comparison of (a) g8-Cu5Sn4 structure and (b) g0 -Cu6Sn5 structure. (c) The g8-Cu5Sn4 schematic of (a) drawn again using the scheme of Larsson et al. [19].
between the center of the diffraction pattern and the (1 1 4)H spot. Starting from the center, spots 1, 2, 4, 5 and 7 are shifted slightly to the left or right of the line and spots 3 and 6 are on the line. That is, the superlattice planes are not parallel with (1 1 4)H planes but rather are tilted by approximately 1.5° relative to the (1 1 4)H. The expected superlattice d-spacing for the 8 (1 1 4)H is 0.872 nm, but
the measured d-spacing of the actual superlattice is smaller at 0.805 ± 0.003 nm. The ratio of the measured to expected is 1.08 ± 0.005. For comparison with Fig. 8a and b shows a simulated 1 0H zone diffraction pattern for the g8-Cu5Sn4 in a ½1 (=[0 1 0]g8). The g8 diffraction pattern has the exact 8 (1 1 4)H superlattice pattern as expected. For g8 the superlattice spacing is different, 0.872 nm for g8 compared with 0.805 nm in Fig. 8a, but the superlattice spot intensities are similar. It is clear that the crystal structure is close to g8 diffraction in spite of the incommensurate superlattice. If we sight along the line towards the ð1 1 4ÞH spot on Fig. 8a (the other set of 1 1 4 planes in the ½1 1 0H zone) it is seen that the 1/4ð1 1 4ÞH superlattice spot does not lie on the line. However, the 1/4ð1 1 4ÞH distance is exactly 1/4 of the distance between the central spot and ð1 1 4ÞH . The crystal tilted to the ½6 2 1H zone (Fig. 8c) contains the systematic diffraction line from the (1 1 4)H that is common to (1 1 4)H in the ½1 10H diffraction pattern (Fig. 8a). In the ½6 2 1H the conclusions about the incommensurabity of the 8 (1 1 4)H are the same. The superlattice planes are tilted by approximately 1.5° relative to the (1 1 4)H and the ratio of measured to expected 8 (1 1 4)H is 1.08 ± 0.005. A simulation of g8 in ½6 2 1H ð¼ ½0 4 1g8 Þ can be compared with the diffraction pattern, as shown in Fig. 8d. The superlattice spot intensities of g8 are similar to the experiment. The crystal tilted to the ½2 6 1H zone (Fig. 8e) contains the systematic diffraction line from the ð1 1 4ÞH that is the same as ð1 1 4ÞH in the ½1 1 0H diffraction pattern (Fig. 8a). In this diffraction pattern there is a superlattice of 1=8ð1 1 4ÞH , whereas in the ½1 1 0H the superlattice was 1=4ð1 1 4ÞH . The superlattice is exactly 1/8 of the ð1 1 4ÞH spacing but the superlattice spots are not on the line between the center of the diffraction pattern and ð1 1 4ÞH but lie slightly to the left or right of that line. How can we interpret the incommensurate superlattice? The ratio of the measured to the expected superlattice spacing is an important clue. The ratio 1.08 is very close to the ratio 40/37 (40/37 = 1.081). It was noted earlier (Fig. 6 and discussion thereof) that g8-Cu5Sn4 unit cell has a period of 8 (1 1 4)H between Cu interstitial double layers, and that g0 -Cu6Sn5 has a period of 5 (1 1 4)H between layers. The supercell (Fig. 9) is incommensurate by a factor of 37/40 which can be obtained by a periodic stacking of four g8Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell. We call this a 4 + 1 block (4 by g8 = 32, plus 1 by g0 = 5; 32 + 5 = 37), thus the new monoclinic phase will be labeled as g4+1. In Fig. 9 the 4 + 1 block is drawn using the scheme of Larsson et al. [16,19]. At first pass the diagram may look a bit complicated; however, it is based upon a simple system. In drawing out the 4 + 1 block structure it is clear that we need to shift the second 4 + 1 block by 1/8[3 0 1]g8. In addition, the second 4 + 1 block is inverted along the [0 1 0]g8 direction (i.e. [0 1 0]g8 becomes ½0 10g8 Þ, so there is doubling of the unit cell along [1 0 0]S. The unit cell of the 2 (4 + 1) superlattice is [1 0 0]S = 9.25[1 0 0]g8; [0 1 0]S =
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Fig. 7. Series of SAED patterns at different zone axes. In the sequence of (a) ! (d) ! (f), the diffraction spots pair (marked by the red arrows and A) were kept visible when rotating the sample. Similarly, in the sequences of (a) ! (b) ! (c) and (a) ! (e) ! (g), diffraction spots pairs (marked by green arrows and B, yellow arrows and C) were kept visible during rotating, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
[0 1 0]g8; and [0 01]S = [0 01]g8. The detailed coordinates of atoms in the unit cell can be found in the online Supplementary information. In Fig. 9 the unit cell outlines of the g8 and one g0 unit cells are shown and the legend on the figure shows the height along the b-axis (y-direction) of the interstitial Cu atoms. In the figure the average spacing between the (circled) g8 Cu-pairs is (37/40) 8 d114. Moreover, the plane upon which the circled Cu-pairs lie is, on average, rotated by 1° relative to the ð3 3 4ÞH planes. It is the average spacing of these g8 Cu-pairs that gives rise to the incommensurate superlattice. The 4 + 1 block system used in Fig. 9 gives a structure which can explain the incommensurability of the superlattice, i.e. the structure explains the 40/37 ratio and the approximate 1° rotation relative the expected (1 1 4)H alignment. However, the structure of Fig. 9 is an idealization of the actual structure. If the ideal structure did occur exactly as postulated in Fig. 9, then there would be a series of very fine superlattice spots in the diffraction patterns with a spacing of 1/37(1 1 4)H, and these are not observed. It seems that the
4 + 1 blocks are the average structure: the actual structure has a degree of randomization, where the average 2 (4 + 1) may be as a result of 5 + 1 followed by 3 + 1, or 6 + 1 followed by 2 + 1, for example. A simulated diffraction pattern of the g4+1 monoclinic structure is shown schematically in Fig. 10. The simulation is for the [0 1 0] orientation, for comparison with Fig. 8a. The simulation shows splitting of diffraction spots (arrowed) which is similar to the splitting in the experimental diffraction pattern. The diffraction spots are not evenly spaced along the (1 1 4)H direction as shown by a comparison of the distances marked in Fig. 10. In the simulation the ratio of distance second relative to first is 1.82 (i.e. <2.0). In the experimental diffraction pattern (Fig. 8a) the ratio is 1.80. The crystal with the incommensurate superlattice that we found provides us with an explanation for the strain that is observed in the majority of diffraction patterns from our samples with the nominal Cu6Sn5 composition. It appears likely that the strain arises from the mixing of one g0 with several g8 layers where it is the g8 that provides the predominant contribution to the superlattice.
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Fig. 8. Close-ups for analysis of some of the diffraction patterns from Fig. 7. (a) ½1 1 0H on zone diffraction pattern. (b) Simulated diffraction pattern for 1 0H (=[0 1 0]g8). (c) ½6 2 1H on zone diffraction pattern. (d) Simulated diffraction pattern for the g8 in ½6 2 1H ð¼ ½0 4 1g8 Þ. (e) ½2 6 1H on zone the g8 in ½1 diffraction pattern.
3.3. Reanalysis of the XRD pattern using the proposed atomic model After developing the atomic model of the new monoclinic phase (see online Supplementary information), Rietveld refinement was performed on the XRD profile for stoichiometric Cu6Sn5 at a temperature of 30 °C. In Fig. 11a, the black profile is the experimental result; the red profile shows the refinement result. The XRD pattern at 2h ranging from 14° to 23° is enlarged in the inset and it is apparent that even the small peaks are well described by the model. The Rietveld refinement gave Rp = 3.55 ˚, and Rwp = 5.35 with cell parameters a = 92.2411 A ˚ ˚ b = 7.3108 A, c = 9.8800 A and b = 118.9521°. The detailed atomic coordinates in the unit cell after Rietveld refinement can be found in the online Supplementary information, along with those from the model.
Rietveld refinements for the XRD profile of stoichiometric Cu6Sn5 at room temperature were also performed using other phases (g, g6, g8, g0 ) for comparison and the refinement results are shown in Fig. 11b along with the refinement result using g4+1. The refinements gave Rwp values of 7.11, 7.06, 6.29, 6.35 and 5.35 for g, g6, g8, g0 and g4+1 phases, respectively. It is clear that the g4+1 shows the best refinement in terms of Rwp and curve fitting, suggesting that the proposed crystal structure for the new phase g4+1 is sound. It is interesting to note that g8 also gives a reasonably good refinement, which is expected. The g8 phase is believe to be stable at elevated temperature (higher than 186 °C) [19], while in contrast, the new phase transforms to g phase at a temperature between 200 and 220 °C. The XRD profiles for stoichiometric Cu6Sn5 at temperatures ranging from 170 to 250 °C were also refined using Rietveld analysis. The proposed new monoclinic phase
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Fig. 9. The proposed superlattice structure based upon the 4 + 1 blocks (four g8-Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell). The g8 and g0 unit cells are drawn in outline, and the relative orientation and unit cell axes of the hexagonal cell, the g8 cell and the g4+1 cell are given.
Fig. 10. Simulation of diffraction pattern for the [0 1 0]H orientation of g4+1 phase. The double-ended arrows indicate the relative spacing of the first and second superlattice spots. Superlattice splitting (single-ended arrows) is not present in g8 (Fig. 8b) but is present in the experimental diffraction pattern (Fig. 8a).
was used for XRD profiles at temperatures of 170 and 200 °C, while hexagonal g phase was used for temperatures of 220 and 250 °C. The refinement results are shown in Table 2. Clearly, the dominant phase at temperatures below 200 °C was the monoclinic g4+1 phase, while the main phase at temperatures higher than 220 °C is the hexagonal g phase. 3.4. Discussion on the formation mechanism of g4+1 phase The importance of the finding and structure determination of this new monoclinic phase should be emphasized. It
Fig. 11. (a) Rietveld refinement of the XRD peak profile for stoichiometric Cu6Sn5 at a temperature of 30 °C: the blue profile is the experimental result, while the red one is the refinement result (the differential is also displayed). The XRD pattern at the 2h ranging from 14° to 23° is enlarged in the inset, showing the fit for the smaller peaks. (b) Rietveld refinements for XRD profile of stoichiometric Cu6Sn5 at a temperature of 30 °C using different phases g, g6, g8, g0 , g4+1 from top to bottom. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Table 2 Unit cell parameters of the XRD profile for stoichiometric Cu6Sn5 at temperatures ranging from 30 to 250 °C after Rietveld refinement. ˚) ˚) ˚) Temperature (°C) Crystal structure Space group a (A b (A c (A a (°) b (°) c (°) Rp 4þ1
30 170 200 220 250
g g4þ1 g4þ1 g g
P1 P1 P1 P63/mmc P63/mmc
92.2411 92.5664 92.6589 4.2372 4.2345
is widely assumed that the Cu6Sn5 phase in the near-eutectic solder alloys is identical to that at the interface of the solder–substrate joint [18], although with identification of a new monoclinic phase in the directly alloyed stoichiometric Cu6Sn5 samples, this assumption must be questioned. Most previous studies have assumed the intermetallic Cu6Sn5 was of the g0 phase, although this may only be the case for Cu6Sn5 forming as part of a eutectic reaction [7,18,26], and may not hold true for that formed at the solder–substrate joint [27,28]. Numerous processing routes can be used to obtain Cu6Sn5 including primary, eutectic and peritectic solidification reactions, solid-state diffusion and other synthetic routes such as spray pyrolysis [13] and ball-milling [29]. Differences in crystal structure may result from compositional variations or other influences related to these processing routes. It is still unclear why the new monoclinic phase only exists in the directly alloyed stoichiometric Cu6Sn5 samples while it is absent in the samples of Cu6Sn5 that have been collected by selectively dissolving the Sn phase from near-eutectic solder alloys. Further work is required to clarify the nature of the new monoclinic phase and the formation mechanisms. 4. Conclusions The following conclusions can be drawn from this work: A new monoclinic-based phase with cell parameters of ˚ , b = 7.311 A ˚ , c = 9.880 A ˚ and b = 118.95° a = 92.241 A was formed in directly alloyed stoichiometric Cu6Sn5 samples. The new monoclinic phase has structurally similar features to both g0 -Cu6Sn5 and g8-Cu5Sn4, and thus can be treated as a modulation of these two phases. It has been labeled as g4+1 because, on average, it is represented by the periodic stacking of four g8-Cu5Sn4 unit cells plus one g0 -Cu6Sn5 unit cell (i.e. a 4 + 1 block). g4+1 phase transformed to a hexagonal g-Cu6Sn5 phase at a temperature between 200 and 220 °C, which is higher than the 186 °C at which the well-documented g0 –g allotropic phase transformation occurs. These findings contribute to a growing body of knowledge revealing the complex behavior of the Cu6Sn5 intermetallic. In particular, they show that the composition, the thermal history and the processing or fabrication route strongly influence the structure adopted by Cu6Sn5. A more thorough understanding of the underlying
7.3108 7.3314 7.3347 a=b a=b
9.8800 9.9062 9.9104 5.1349 5.1339
90 90 90 90 90
118.9521 118.9399 118.9453 90 90
90 90 90 120 120
3.55 3.76 4.03 2.92 2.96
Rwp 5.35 6.07 6.55 5.19 5.35
mechanisms could have significant implications for leadfree solder alloy development. 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 program (project ID: LP100200250) and JSPS Fellowship for Research in Japan (project ID: S11730). XRD experiments were performed at the Australian Synchrotron (AS) Powder Diffraction Beamline (project ID: AS101/PD2249 and AS111/PD3364). TEM experiments were performed at the Centre for Microscopy and Microanalysis (CMM) at UQ and the Research Laboratory for High Voltage Electron Microscopy (HVEM) at Kyushu University (KU) with Nanotechnology Network program Japan (project ID: T23-019). J.C.B. acknowledges the support of and use of facilities at QUT. The authors would like to thank Mr. J. Read (UQ) and Mr. D. Mu (UQ) for help on XRD experiments and Mr. T. Daio (KU) and Dr. E. Tanaka (KU) on TEM experiments. Y.W. is funded by the ARC-Linkage project. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/ j.actamat.2012.08.024. References [1] Hsiao HY, Liu CM, Lin HW, Liu TC, Lu CL, Huang YS, et al. Science 2012;336:1007. [2] Zeng K, Tu KN. Mater Sci Eng R 2002;38:55. [3] Li Y, Wong CP. Mater Sci Eng R 2006;51:1. [4] Laurila T, Vuorinen V, Paulasto-Krockel M. Mater Sci Eng R 2010;68:1. [5] Nogita K. Intermetallics 2010;18:145. [6] Nogita K, Nishimura T. Scripta Mater 2008;59:191. [7] Schwingenschlogl U, Di Paola C, Nogita K, Gourlay CM. Appl Phys Lett 2010;96:061908. [8] Suh JO, Tu KN, Lutsenko GV, Gusak AM. Acta Mater 2008;56:1075. [9] Suh JO, Tu KN, Tamura N. Appl Phys Lett 2007;91:051907. [10] Li JF, Agyakwa PA, Johnson CM. Acta Mater 2011;59:1198. [11] Ghosh G, Asta M. J Mater Res 2005;20:3102. [12] Sarakonsri T, Apirattanawan T, Tungprasurt S, Tunkasiri T. J Mater Sci 2006;41:4749. [13] Ju SH, Jang HC, Kang YC. J Power Sources 2009;189:163.
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