Tin (Sn) plating is popularly used as a lead-finish technique for discrete components like chip resistors and chip capacitors [1]. The requirements for the surface finish are excellent oxidation protection of Cu substrates over the leads, consistent solderability, protection against abrasion and good appearance of the leads [2, 3, 4]. One of the common problems in the Sn finish encountered in the field is the discoloration of the surface. The discolored Sn finishes are classified as visual defects and related to the solderability loss in many cases. The discoloration of the Sn finishes is caused mostly by the surface contamination such as organic substances. In the present study, we have investigated a case of the discoloration occurred in the field that was not related to the surface contamination. A layer of Sn was electroplated using the rack plating technique on a Cu alloy substrate (Cu 93.9 wt. %, Sn 6 wt. %, P 0.1 wt. %). The components with the electroplated Sn surface finish were hermetically sealed using by vacuum packaging and subsequent N2 filling. We have selected two Sn surface finished samples which were produced at the same month and kept for about five months at room temperature. The Sn surface of one sample was discolored while the other was not. We employed a focused ion beam (FIB), X-ray energy dispersive spectroscopy (XEDS), transmission electron microscopy (TEM) and X-ray photoelectron spectroscopy (XPS) to analyze the microstructure and morphologies of the Sn layer as well as the interface between Sn and Cu substrate. Fig. 1 shows the scanning TEM (STEM) micrographs obtained from two Sn finishes, acquired by using a high-angle annular dark-field detector (HAADF). Fig. 1(a) was obtained from a sample without any discoloration and Fig. 1(b) was obtained from a discolored sample. We found a significant diffusion of Cu into the Sn layer along the grain boundaries, which reached almost the top surface and confirmed the existence of the Sn-Cu intermetallic compounds (IMCs) and Kirkendall voids from both samples. Fig. 2 (b) indicates the existence of the Sn oxide on the surface of the discolored sample. These experimental results indicate that the primary cause for the discoloration was the oxidation of the Sn surface. The presence of the thin oxide film resulted in the color on the surface due to the constructive and destructive interference of light. The color can vary as the thickness of the oxide layer. In addition, the excessive Cu diffusion along the grain boundary observed in the present study may accelerate the oxidation of the surface Sn layer through galvanic corrosion if it is exposed to humidity. This is due to the dissimilar galvanic potential between Sn and the exposed Cu on the leads. The diffusion of Cu should, therefore, be suppressed for long-term reliability.

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Base metal electrode (BME) multilayer ceramic capacitors (MLCCs) are mass-produced worldwide due to their low cost and shrinking form factor compared to noble metal electrodes (Pd/Ag, Pt) [1-3]. In these BME MLCCs, it is very important to understand the interfacial reaction including the diffusion of oxygen vacancies at the Ni/BaTiO 3 interfaces since they must be fired in a reducing atmosphere to protect the Ni from oxidation where BaTiO 3 is slightly reduced to result in formation of doubly ionized oxygen vacancies. Recently, Yang et al. have reported the existence of a discrete metallic alloy layer containing Ni, Ti and Ba between Ni and BaTiO 3 co-fired in reducing atmospheres along with an oxygen-depleted zone in BaTiO 3 adjacent to the metallic layer [4,5]. The metallic alloy layer was identified by highresolution transmission electron microscopy (HRTEM) observations and electron energy-loss spectroscopy (EELS) analysis. In the present paper, we present the similar observation of the interfacial layer to Yang et al.'s metallic alloy layer but different findings. We have confirmed the existence of the interfacial layer in the Ni electrode along the Ni/BaTiO 3 interfaces as shown in Fig. 1. Both Ba and Ni were detected by EELS in the interfacial layer as shown in Fig. 2 (a). Fig. 1 also shows the similar fringes in the Ni electrode region which Yang et al. suggested as Moiré fringes due to surface oxidation of Ni. However, we found that the interfacial layer grew as the electron-beam irradiation time increased as shown in Fig. 1 (a)-(c). At the beginning, the thickness of the interfacial layer was about 1 nm and grows to 10nm after about 7 min under the electron beam. Our experimental results suggest that the interfacial layer was not formed during the fabrication of the MLCCs but formed during the TEM experiments under an electron-beam irradiation. We have concluded that the interfacial layer in the HRTEM images can be formed by an electron-beam induced sputtering of surface oxide layer and the resulting removal of the Moiré fringes. The sputtering of the surface oxide can be evidenced by little oxygen concentration observed in the interfacial layer compared to that of the Ni electrode region as shown in Fig. 2. The presence of Ba in the interfacial layer may be explained by the diffusion of Ba across the Ni/BaTiO 3 interface driven by the electron-beam induced heating of the TEM specimen.