The Role of Ultrafine Crystalline Behavior and Trace Impurities in

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The Role of Ultrafine Crystalline Behavior and Trace Impurities in Copper on Intermetallic Void Formation Glenn Ross, Per Malmberg, Vesa Vuorinen, and Mervi Paulasto-Kröckel ACS Appl. Electron. Mater., Just Accepted Manuscript • DOI: 10.1021/acsaelm.8b00029 • Publication Date (Web): 10 Dec 2018 Downloaded from http://pubs.acs.org on December 14, 2018

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ACS Applied Electronic Materials

The Role of Ultrafine Crystalline Behavior and Trace Impurities in Copper on Intermetallic Void Formation Glenn Ross,*† Per Malmberg, ‡ Vesa Vuorinen,† and Mervi Paulasto-Kröckel†

† Department

of Electrical Engineering and Automation, Aalto University, P.O.Box

13500, FIN-00076 Aalto, Finland

‡

Chemistry and Chemical Engineering, Chalmers University of Technology, Kemivägen

10, SE-412 96 Gothenburg, Sweden

*Corresponding author: [email protected]

ABSTRACT

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In the microelectronic component industry, due to the miniaturization of functional units, the interaction of materials and interfaces play a much more significant role in their performance. The result of this is that ultrafine crystalline and trace impurity behavior impact on a devices operation to a much greater extent. This is the case with microconnects for 3D integration, such as micro-bumps. Any unwanted crystalline behavior or interfacial segregated impurities can drastically alter a micro-connects performance. A particular issue being intermetallic void formation, often known as Kirkendall voiding. Currently it is unclear under what conditions voids form and how to prevent them. This work studies the microstructural and compositional differences between samples with different voiding densities. Results show that samples that exhibit an ultrafine crystalline have a higher propensity to exhibit voiding. Also, there is a high concentration of trace impurities located in the electrochemically deposited Cu layer. After isothermal annealing, high concentrations of impurities are located at the interface between Cu and the Cu-Sn intermetallic of Cu3Sn. An alternative explanation to the traditional Kirkendall void formation theory is presented. The explanation is based on the interaction of trace impurities from the electroplating process, and the microstructural evolution.


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KEYWORDS: Intermetallic Voids, Kirkendall Voids, Microstructure, Impurities, Residual Stress, Scanning Transmission Electron Microscopy, X-Ray Diffraction, Time-of-Flight Secondary Ion Mass Spectroscopy.

1. Introduction

As the semiconductor industry scales its materials and interfaces to match the rate of transistor miniaturization, ultrafine crystalline behavior and trace impurities begin to play a major role in materials and interfacial behavior.1 Texturing and impurities, often in the parts-per-million ppm or parts-per-billion ppb range, can lead to a range of undesirable effects on the performance of high density micro-connects. For example with microbumps and solid-liquid interdiffusion (SLID) bonding of micromechanical systems (MEMS) for 3D integration. Such undesirable effects can include interfacial weakening, cracking and delamination. Impurities can exacerbate this behavior by causing segregated precipitates and voiding that changes the microstructural performance. One


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effect, and the focus of this paper, is intermetallic void formation often referred to as Kirkendall voiding. Intermetallic voids form at the interface between Cu and the Cu-Sn intermetallic compound (IMC) of Cu3Sn (ε) and have been cited as the root-cause of reduced structural performance of interconnects.2–5 Traditionally intermetallic void formation is believed to be a result of the Kirkendall effect, which is the diffusion imbalance between two adjacent metals.6 The Cu-Sn system studied in this work is an example of this. Numerous studies7–9 have shown that at Cu-Sn solidus temperatures, Cu in the ε-phase is the dominant diffusing species. As a consequence of the diffusion imbalance and the vacancy mechanism for substitutional diffusion, in order to maintain a flux-balance, Sn and vacancies must diffuse towards Cu (JCu-JSn-JVa=0). This results in the supersaturation of vacancies at the Cu/ε interface that eventually aggregate and result in the formation of voids. This is the basis of Kirkendall void formation theory.10 However, the flaw in this hypothesis is that frequently the propensity for void formation is not consistent between Cu-Sn samples. A number of studies11,12 have shown that when oxygen-free high conductive (OFHC) Cu is combined with an electroplated, or reflowed


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Sn layer, the interface experiences little to no propensity to void. Whereas electrochemically deposited Cu has, depending on the processing parameters, exhibited varying voiding densities.12 More recently research has shown a strong dependence on the electroplating parameters, such as the Cu electroplating additives and current density, to name a few.12–17 These results suggest that even though the Kirkendall effect may contribute to void formation, there must also be other contributing factors. If the Cu electroplating parameters affect the voiding propensity of a Cu-Sn interface, they must also impact the underlying microstructure and chemical composition. These properties need to be understood in order to better understand the void forming mechanisms. Developed in this paper is an understanding of the microstructure and chemical composition of the Cu-Sn interface for different voiding conditions. Three different electroplating chemistries are examined which produce different voiding propensities. Microstructural properties are analyzed using scanning transmission electron microscopy (STEM), electron diffraction and x-ray diffraction (XRD). The material chemistry is probed with time-of-flight secondary ion mass spectroscopy (ToF-SIMS). From these results an alternative intermetallic void formation hypothesis is proposed.


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2. Materials and Methods

2.1 Sample preparation

All samples were prepared on 100 mm p-Type Si wafers. Two sample sets were fabricated: (i) samples for SEM, TEM/STEM cross-sectional analysis and XRD and (ii) samples for the ToF-SIMS analysis. The fabrication steps for the two sample sets are indicated in Figure 1. Sample set (i) was prepared by sputtering a thin 20 nm Cr barrier/adhesion followed by a 100 nm of Cu seed-layer. Following magnetron sputtering, a 3 µm layer of Cu and 2 µm layer of Sn was electrochemically deposited. To influence the voiding propensity, a selection of additives commonly used in electroplating chemistries was selected for the electrochemical Cu bath: Bis-(3-sulfopropyl)-disulfide or SPS, Poly(ethylene glycol) or PEG and a synergistic combination of SPS and PEG, SPS+PEG. An overview of the bath compositions are given in Table 1. Electrochemical Sn was deposited from a common commercial solution (NB Semiplate Sn 100).


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Sample set (ii) was fabricated by first depositing a 120 nm thick plasma-enhanced chemical vapor deposed (PECVD) SiO2 layer. Subsequently, a magnetron sputter layer of 100 nm Pt, 20 nm Cr and 100 nm Cu was deposited. A 1 µm thin layer of electrochemical Cu and 5 µm layer of electrochemical Sn was deposited. The electrochemical Cu and Sn were deposited under the same conditions and using the same baths as in sample set (i). To probe the interface of interest with ToF-SIMS with the maximum depth resolution, the interface must be located close to the surface, so the weak adhesion interface (SiO2/Pt) was delaminated by exfoliation. The delaminated layer was remounted on a brass substrate using low-outgassing epoxy adhesive Loctite EA 9492. As the Pt layer is now on the surface and there is a thin electroplated Cu layer (1 µm), depth resolution is enhanced. Both sample sets included non-annealed and isothermally annealed samples. Isothermal annealing was done at 423K (150°C) for a duration of 4 hours.


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Figure 1. Illustration of the structure of (i) cross-section, (S)TEM and XRD samples and (ii) ToF-SIMS samples.

Table 1. Overview of samples and the sample electrochemical deposition chemistries.


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Sample Concentration

Chemical

0.25 M Copper sulphate SPS

1.8 M Sulphuric acid 1.92 mM Hydrochloric acid 0.01 mM Bis-(3-sulfopropyl)-disulfide (SPS) 0.25 M Copper sulphate

PEG

1.8 M Sulphuric acid 1.92 mM Hydrochloric acid 0.059 mM Poly(ethylene glycol) (PEG) 0.25 M Copper sulphate

SPS+PEG

1.8 M 1.92 mM 0.01 mM 0.059 mM

Formula CuSO4·5H2O

Aldrich 209198

H2SO4 HCl C6H12Na2O6S4

Aldrich 40306 Aldrich 40304 Raschig RALU®PLATE SPS

CuSO4·5H2O

Aldrich 209198

H2SO4 HCl H(OCH2CH2)nOH

Aldrich 40306 Aldrich 40304 Aldrich 202444

CuSO4·5H2O

Aldrich 209198

H2SO4 Sulphuric acid Hydrochloric acid HCl C6H12Na2O6S4 Bis-(3-sulfopropyl)-disulfide (SPS) H(OCH2CH2)nOH Poly(ethylene glycol) (PEG)

Supplier

Aldrich 40306 Aldrich 40304 Raschig RALU®PLATE SPS Aldrich 202444

2.2 Characterization methods

Scanning electron microscopy (SEM) cross-sections from the sample set (i) was done to determine the microstructure and voiding density in the non-annealed and annealed samples. Samples were cross-sectioned and ion-polished using the Gatan Ilion +. A JEOL JSM-6330F field emission SEM was used for the cross-sectional imaging with a back scattering electron detector. TEM lamellas were prepared using a focused ion beam (FIB), employing a in-situ lamella lift-out process. A dual-beam (SEM-FIB) Helios NanoLab 600 was used for the lamella preparation. The transmission electron microscopy (TEM) and scanning


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transmission electron microscopy (STEM) was conducted using the JEOL JEM-2800 operating at 200 kV.

The x-ray diffraction (XRD) measurements were performed with a Rigaku SmartLab operating with a 9 kW rotating Cu anode source. The XRD was performed in two measurement configurations for both a wide-2D diffraction map and residual stress sin2ψ. The wide-2D diffraction maps were taken with a focused incident beam and a 2D hybrid pixel array detector operated in 2D with an active area of 3000 mm2. The wide2D diffraction measurements were performed at incremental χ angles, the measurements stitched together and the results are presented in 2θ-χ space. The specimen and laboratory reference frames used in the XRD wide-2D diffraction maps can be seen in Figure 2 subfigure (a). The diffraction vector in this measurement moves in the S3-S1 plane at increasing values of χ. The residual stress measurements were carried out in parallel beam mode to avoid instrumental alignment errors and aberrations. The diffracted beam was passed through a 0.114° parallel slit analyzer (PSA) to the 2D hybrid pixel array detector operating in 0D mode. The diffraction vector


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is moved in the S3-S2 plane at increasing values of ψ, which can be seen in Figure 2 subfigure (b). The range of ψ was limited by the arrangement of the systems’ goniometer. Multiple higher-order reflections were acquired in order to achieve: (i) best least-squares data fit and (ii) larger d-spacing change as a function of ψ.


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Figure 2. XRD specimen (Sx) and laboratory (Lx) reference frames used in the (a) wide2D diffraction scan, where the diffraction vector moves in the S3-S1 plane at increasing values of χ and (b) residual stress sin2ψ measurement, where the diffraction vector moves in the S3-S2 plane at increasing values of ψ.

TOF-SIMS analysis was performed using a ToF-SIMS V instrument (ION-ToF, GmbH, Münster, Germany) equipped with a 25 keV Bismuth LMIG (Liquid metal ion gun) and a Cs sputter source. Depth profiling was performed in the interlaced mode with an analysis area of 100 µm x 100 µm (128x128 pixels) and a sputter area of 300 µm x 300 µm. Analysis was performed in the negative ion mode using the high current bunched mode, with pulsed primary ion beam (Bi3+, 0.4 pA) with a focus of approximately 2 μm and with a mass resolution of at least M/ΔM = 6000 fwhm at m/z 500. The Cs source was used at 3keV at 16 nA. The total Cs ion dose density of ranged between 1e18 and 2e18 for the different depth profiling experiments. All depth profiles were acquired and processed with the Surface Lab software (version 6.7, ION-TOF GmbH, Münster, Germany).


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3. Results and Discussion

To show the impact of the different electropating chemistires on the voiding propensity of the Cu-Sn interfaces, SEM micrographs of the six samples are presented in Figure 3. Sub-figures (a) to (c) are the non-annealed samples, followed by (d) to (f) that have been isothermally annealed at 423K (150°C) for 4 hours. Clearly illustrated is an electrochemical Cu chemisty dependancy on void formation. The SPS samples exhibit moderate void formation, PEG undergoes dense void formation, particularly after isothermal annealing, that forms a partial crack like void at the interface. In contrast, the SPS+PEG experiences little to no voiding. In the Cu-Sn system at the tempeatures that are used, there are two stoichiometric compounds, Cu3Sn (ε) and Cu6Sn5 (η’).18 It can be seen that in the non-annealed samples macroscopic voids have already began to form. Also, there is rapid diffusion of Cu into Sn, predominantly at the grain-boundaries of Sn. The grain boundary diffusion component is significant at room temperautre,19,20 and substantial formation of η’-phase can be seen from the micrographs (and the ε-


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phase as will be seen in the subsquent STEM micrographs). Void formation always occurs within the ε-phase.

Figure 3. SEM-BSE micrographs of the non-annealed samples (a) SPS, (b) PEG and (c) SPS+PEG, followed by the isothermally annealed samples at 423K (150°C) for 4 hours of (d) SPS, (e) PEG and (f) SPS+PEG. The structure of the samples is presented on the left side of the figure.


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The SEM cross-sections reveal that there is a voiding propensity difference between the samples and that there is a change during isothermal annealing. STEM cross-sections show the microstructural differences. Figure 4 shows STEM-BF micrographs for the six samples. Subfigure (a) to (c) are the non-annealed samples and (d) to (f) are the isothermally annealed samples. Once again, there is a clear voiding difference between the samples, and notable are the crystal size differences. SPS and PEG samples exhibit an ultrafine crystalline microstructure, whereas the SPS+PEG sample has a significantly larger crystal size. Interestingly there is little change in the crystal structure of all three samples as a function of isothermal annealing, particularly for the SPS and PEG samples where one would expect crystal growth to occur. In all samples it appears that the crystal growth is being inhibited. One additional feature of the microstructure of the SPS and the PEG samples is the significant number of crystal defects within the Cu grains themselves, these appear to increase after isothermal annealing, as twinning and stacking faults. This suggests a certain level of residual stress induced plastic


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deformation within the Cu layer, which will be determined by the XRD residual stress analysis.

Figure 4. STEM-BF micrographs of non-annealed samples (a) SPS, (b) PEG and (c) SPS+PEG, followed by the isothermally annealed samples 423K (150°C) for 4 hours (d) SPS, (e) PEG and (f) SPS+PEG. The structure of the samples is presented on the left of the figure.


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TEM select area diffraction patterns (SADP) of the Cu were taken, which are presented in Figure 5. Si was included in the area to reference the texture direction of the Cu crystals. The SADP confirms the ultrafine crystalline structure of the SPS and PEG samples. The discrete diffraction spots of the SPS+PEG SADP indicates a comparatively large crystal size. The SADP shows there are slight texture differences between the SPS and PEG samples, as the Cu(111) lattice planes in the SPS sample appear to have a narrow distribution around the substrate perpendicular direction, whereas the PEG sample appears to have a wider distribution.


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Figure 5. (a) HRTEM micrograph of the non-annealed SPS samples indicating the SADP region. The SADP of the non-annealed samples (b) SPS, (c) PEG and (d) SPS+PEG.

To better determine the crystral behavior over a larger sample area, XRD was used to map the diffraction patterns in 2θ-χ space. The results of the wide-2D diffraction maps can be seen in Figure 6. As noted in the SADP, there are crystallographic texture differences between the samples. However there does not appear to be a specific


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texturing behavior that could be associated with void formation, suggesting that texturing does not play a role. This is in contrast two previous papers21,22 that suggest voiding favors Cu samples with preferred orientations. SPS has a preferential orientation of the Cu(111) peak, whereas the PEG exhibits a more random orientation. The SPS+PEG sample exhibits, to a certain degree, Cu(111) texturing and also the Cu(200). Interestingly, in all samples after isothermal annealing, there are two common phenomena occurring. Firstly the full width at half maximum (FWHM) appears to increase significantly. The increased FWHM indicates there is a wider distribution of lattice d-spacings. The major contributor to this wider distribution is an increase in the average grain size. The second phenomenon is a shift in the 2θ value with increasing χ values. This indicates that as the diffraction vector moves parallel to the sample surface the d-spacing increases, suggesting there is a residual stress component.


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Figure 6. XRD wide-2D diffraction scan, (a) SPS samples, (b) PEG samples and (c) SPS+PEG samples of the non-annealed samples (upper) and the annealed samples (lower). The indicated peaks are from Cu (111), (200) and (220).

Residual sin2ψ stress measurements were made from the six samples to assess the impact residual stress has on the voiding propensity. For the residual stress determination the methods presented by Leeuwen et al.23 and extended by Welzel et

al.24 were utilized. A comprehensive overview of the methods can be found in these


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articles, and also in a review article by Welzel et al.25. The effective grain-interaction model, that includes the Vook-Witt, Inverse Vook-Witt, Reuss and Voigt models, has been applied using the non-linear least squares method to the sin2ψ plots of the Cu(220), Cu(400) and Cu(331) reflections. The result of the data fitting can be seen in Figure 7 and the residual stress values are listed in Table 2. Results show that all samples have a rotationally symmetric tensile residual stress component, with the SPS+PEG sample having the smallest value. After isothermal annealing all tensile stress values increased. This is somewhat counterintuitive, as one would expect the film to relax during the thermal process.


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Figure 7. The Sin2ψ results for the Cu(220), Cu(400) and Cu(331) reflections of the nonannealed and annealed (a) SPS, (b) PEG and (c) SPS+PEG samples.


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Table 2. Results of the least squares fit of the rotationally symmetric average residual stress for the six samples.

Sample SPS PEG SPS+PEG

Annealing Hours 0 4 0 4 0 4

||S MPa 40.0 ±16.5 128.7 ±15.0 51.4 ±15.8 151.5 ±9.1 12.1 ±4.5 72.4 ±7.1

Δ||S MPa +88.6 +100.1 +60.3

In the non-annealed samples, the Cu film is in a relatively “stress-free” state. The SPS and PEG samples exhibit a slightly higher tensile stress value, that appears to be due to the inclusion of impurities and a more complex microstructural state. During isothermal annealing the situation becomes more complicated. Due to the coefficient of thermal expansion (CTE) difference between the Si substrate and the metallic layers, after heating the Cu-Sn layers will enter a more compressive stress state. Strain hardening occurs to counter this compressive stress by grain size reduction, that hinders


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dislocation motion. In addition to this, solid-state interdiffusion and reaction of Cu and Sn results in the formation of the intermetallic compounds that produces a volumetric change. The volumetric change occurring during the early formation of the intermetallic compounds can be estimated by (Eq1 and Eq2).

𝑆𝑛 ∆𝑉 = 𝑉𝜀𝑚 ― 3𝑉𝐶𝑢 𝑚 ― 𝑉𝑚

(1)

𝑆𝑛 ∆𝑉 = 𝑉𝜂𝑚′ ― 6𝑉𝐶𝑢 𝑚 ― 5𝑉𝑚

(2)

′) Where ∆Vis the volumetric change due to the formation of the ε and η’-phases, 𝑉(𝜀,𝜂 is 𝑚

the molar volume of the ε (𝑉𝜀𝑚 = 27.30) and η’ (𝑉𝜂𝑚′ = 118.01)26 and 𝑉(𝐶𝑢,𝑆𝑛) is the molar 𝑚 𝑆𝑛 3 volume of Cu (𝑉𝐶𝑢 𝑚 = 7.11 ) and Sn (𝑉𝑚 = 16.29 ) in cm /mol. When determining the

volumetric change during the formation of the IMCs, there is a net decrease of -10.32 cm3/mol or -27.4% and -6.1 cm3/mol or -4.9% due to the formation of the ε and η’phases respectively. This is a significant volume shrinkage and would undoubtedly result in a significant local tensile stresses at the Cu/ε interface. Finally the samples are cooled to room temperature. Due to strain hardening of the Cu layer at the elevated temperature, when cooled the CTE mismatch between Si and Cu results in the


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increased average tensile stresses seen in Table 2. Certainly, if significant tensile stresses are present, this can provide the additional energy required to form new surfaces and contribute to void growth after the nucleation phase.27

To understand trace impurities in the Cu-Sn system, more specifically at the Cu/ε interface, ToF-SIMS has been used in the dynamic mode, to study the impurities as a function of depth from the exfoliated Pt surface. ToF-SIMS was chosen for its detection sensitivity in the ppm and ppb range.28 The trace impurities that are of interest are the components of the additives that breakdown during electroplating, such as C, Cl, O and S. It must be noted that as the sputtering depth increases the depth resolution decreases, due to differing sputtering rates in the dissimilar materials within the polycrystalline material itself and at the non-uniform interfaces (i.e. Cu/ε interface). The results of the measurements can be seen in Figure 8.

A striking feature in the ToF-SIMS results is the difference in the trace impurities between the SPS+PEG, and the SPS and PEG samples. Both SPS and PEG samples show large trace impurity content in the Cu layer that include C, Cl, O and S.


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Additionally, after isothermal annealing the impurities occupy primarily the Cu and εphase with the Cu/ε interface showing, in both the SPS and PEG the greatest concentration.


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Figure 8. ToF-SIMS results from SPS (a) non-annealed and (b) annealed, PEG (c) nonannealed and (d) annealed, and SPS+PEG (e) non-annealed and (f) annealed. The estimated phase boundaries are indicated by dashed lines.


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Results have shown that a Cu-Sn system that exhibits a high propensity to void has an ultrafine crystalline Cu microstructure that appears to experience grain size reduction during thermal annealing due to thermally induced stresses. In addition to the grain structure, the number of defects within the Cu microstructure appears to be higher in SPS and PEG samples.

The XRD results revealed a narrow distribution of Cu(111) orientation for SPS around the substrate in the direction perpendicular to the substrate, for PEG the distribution was much larger. The SPS+PEG appeared to have a competing preferred orientation between Cu(111) and Cu (200). Voided samples showed relatively higher tensile residual stress components, that corresponds to the higher defect density and impurities in the Cu microstructure. The ToF-SIMS results clearly showed that there is a high trace impurity concentration of C, Cl, O and S in the Cu electrochemically deposited Cu and in the Cu/ε after isothermal annealing. From these results an alternative intermetallic


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void formation hypothesis is proposed. A graphical illustration of this is presented in Figure 9.


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Figure 9. Proposed void formation mechanisms, (i) is the co-deposition of impurities that occurs during Cu electroplating, (ii) is the segregation of impurities to defects and grainboundaries, (iii) is the formation of voids and (iv) is the proposed localized void formation mechanisms as a function of time (t1 to t4).

From (i) to (iii) are the proposed progressive phases of intermetallic void formation, with (vi) being the localized illustration of void formation as a function of time (t1 to t4). Stage (i) is the deposition phase where trace electroplating impurities have been trapped into the Cu matrix due to the non-synergistic behavior of the Cu electroplating chemistry. These trace impurities are distributed throughout the Cu layer. Stage (ii) is the early formation of the IMCs through Cu dominated diffusion, in addition segregation of impurities to the Cu grain boundaries and other lattice defects occurs. Previous studies29 have shown that the formation energies of dilute solutions (containing Cl) are less energetically favorable to form within the Cu matrix, causing trace impurities to migrate to more thermodynamically favorable sites such a lattice defects. Stage (iii) is


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the intermetallic void formation stage. As Cu further diffuses into Sn, the moving Cu/ε interface reaches a lattice defect, such as a grain boundary. The lattice defects containing impurities have multiple effects on the impinging interface. Firstly, the segregated impurity disrupts the grain boundary component of Cu diffusion, and secondly it causes pinning of the moving Cu/ε interface to occur. This pinning produces a dragging force on the impinging interface. Behind the lattice defect is an inconsistency that produces a coalescence of vacancies. After the initial nucleation, tensile stresses cause new surfaces to form enabling further growth and eventually resulting in a void.

4. Conclusions

In concussion this work has presented the microstructure and trace impurity content of three samples that exhibit different intermetallic voiding propensities. The results of a high voiding propensity sample revealed a microstructure that had an ultrafine crystalline structure with a large number of crystal defects saturated by trace impurities. Voided samples experienced varying levels of preferential Cu (111) texturing when compared to the non-voided sample that had competing preferential texturing of


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Cu(111) and Cu(200) perpendicular to the substrate surface. From these results it appears that texturing does not play a major role in void formation. Average tensile residual stresses were slightly higher in the voided samples, which increased during thermal annealing and IMC formation, due to the volumetric change in the ε- and η’phases and CTE mismatch between the metal layers and the Si substrate. A revised model not solely based on the Kirkendall void theory was presented. The model proposed that trace impurities trapped during the electroplating process segregate to lattice defects such as grain boundaries. When the moving Cu/ε interface reaches one of these lattice defects, pinning of the interface occurs, causing interfacial drag. In addition, the contribution of grain boundary diffusion is inhibited due to the saturation of trace impurities. This leaves behind an inconsistency that allows the coalescence of saturated vacancies and subsequent tensile stress induced void growth.

ACKNOWLEDGMENT


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The authors would like to thank Okemtic for the supply of Si wafers. We acknowledge the provision of facilities and technical support by Aalto University at OtaNano – Microscopy Center (Aalto-NMC).

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TOC graphic


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Segregated impurities

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η’ 1 2ε 3 4 Cu 5 6

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The Role of Ultrafine Crystalline Behavior and Trace Impurities in

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