Abstract
In this paper, the fabrication and electrical property characterization of epitaxial Cu3Ge thin film are performed. By adjusting deposition parameters, the crystallinity of the as‐grown Cu3Ge thin films is improved, with the formation of twins within it. The average work function of epitaxial Cu3Ge thin film is measured to be ∼4.47 + 0.02 eV, rendering it a desirable mid‐gap gate metal for applications in complementary metal‐oxide semiconductor (CMOS) devices. The present study therefore shows an epitaxial Cu3Ge thin film that is promising for applications.
Keywords
- Cu3Ge thin film
- sapphire
- twin
- pulsed laser deposition
- semiconductor metallization
1. Introduction
Cu3Ge is a promising candidate and an alternative to Cu for contacts and interconnections in advanced integrated circuit devices. It has a relatively low bulk resistivity throughout the compositional window of 25–35 at% Ge [1, 2]. Its thermal stability against oxidation is also excellent [3, 4]. Furthermore, the service life of Cu3Ge is considerably longer than Cu because the out‐diffusion of Cu [4] is reduced. Not only different substrates have been used for growth of polycrystalline Cu3Ge films (including GaAs [5–7], Si [4, 8–10], Ge [10, 11], YBa2Cu3O7-
In this chapter, we show the fabrication of epitaxial Cu3Ge thin films with significantly improved crystallinity, due to a modified deposition route of Cu3Ge thin films. The average work function of epitaxial Cu3Ge thin film is measured to be ∼4.47
2. Experimental procedure
Pulsed laser deposition (PLD) was used to deposit Ge and Cu thin films repetitively on sapphire substrates. For Cu atoms to mix and bond with Ge atoms intimately, ultrathin Ge and Cu layers were deposited in each repetition. Cu3Ge film is thus expected to have improved crystallinity as laser ablation provides excess kinetic energies of Cu/Ge atoms. A series of five Cu3Ge films are fabricated at 400 ± 10°C, with systematically changed deposition parameters for investigating their correlation with the crystallinity of Cu3Ge films. Specifically, 90 repetitions of Ge and Cu layers are deposited for all five samples, with a changing pulse number of Ge and Cu deposition for sample 1–5. For example, 35 pulses of Cu and five pulses of Ge are deposited in one repetition for sample 1, and a total of 90 repetitions are performed. Similarly, pulses of Cu in one repetition for sample 2–5 are 25, 15, 14, and 7, while pulses of Ge in one repetition for sample 2–5 are 5, 5, 2, and 1. Before deposition, multi‐step cleaning is performed on the substrates. They are first cleaned with boiled acetone for 5 min, and then ultrasonically cleaned in acetone and methanol for 5 min each. Nitrogen gun is used to dry the cleaned substrates, which are then mounted into the chamber 4 cm away from the target. The pulsed KrF excimer laser used has a wavelength of 248 nm and pulse duration of 25 ns. The pre‐deposition vacuum is ∼10-7 Torr, while the actual deposition vacuum is ∼3.0 × 10-4 Torr due to heating. The laser source provides a laser with constant exciting voltage (23.8 keV [18–26]), and hits the pure Cu and Ge targets (from ESPI) at an incidence angle of 45°. Since the spot size of the landing laser is ∼2 mm × 3 mm and the landing laser beam energy is ∼0.29–0.30 J, the laser beam energy density is ∼4.8–5 J cm-2.
High‐resolution transmission electron microscope (HRTEM) JEOL‐2010F with a point‐to‐point resolution of 0.18 nm is used to characterize the microstructure of the deposited films. Focus ion beam (FIB) is used for preparing TEM samples. Rigaku X‐ray diffractometer (XRD) is used to perform XRD
The local work function of epitaxial Cu3Ge thin film was measured by Kelvin probe force microscopy (KPFM). In KPFM measurements, the conductive tip (NSG03/Pt) was working in tapping mode and lift mode scan. Reliable topographic images were obtained before the KPFM measurements. In KPFM tests, an AC and a DC bias were applied on the cantilever while the sample remained grounded. The frequency of the AC signal was set at ∼2kHz lower than the resonance frequency of the cantilever and the amplitude AC voltage was 560 mV. The lift scan height was 50 nm and the scan rate was 0.75 Hz.
3. Results and discussion
To characterize the microstructures of the deposited thin films, XRD
The quality and crystallinity of Cu3Ge films in sample 1–5 are also demonstrated by the plan‐view images of optical microscope images, shown in Figure 2. Figure 2(a)–(e) correspond to the plan‐view images of samples 1–5, respectively. Chunks appear in the plan‐view images of samples 1, 2, and 3, possibly due to the existence of excessive Ge grains in addition to the Cu3Ge phase. In contrast, uniform films are observed in Figure 2(e) and (f), proving the high quality of the Cu3Ge films with good crystallinity for samples 4 and 5. SEM and AFM analyses are performed to show more detailed surface morphology of sample 5. The plan‐view SEM image (Figure 3) shows the surface morphology of Cu3Ge film, with an average diameter of Cu3Ge islands to be ∼300–500 nm. The complementary AFM height map (Figure 4) demonstrates the average heights of Cu3Ge islands to be ∼180–240 nm.
TEM is performed to further study the epitaxial film morphology, as demonstrated in Figure 5(a). Note that the Au/Pt layers on top of Cu3Ge are deposited during ample preparation for protection. Based on the morphology of the Cu3Ge film, it may grow on the
HRTEM investigation was performed to study the atomic structure at the Cu3Ge/c‐sapphire interface, as shown in Figure 6(a). The Cu3Ge/c‐sapphire interface is semi‐coherent [27], due to visible periodic contrast along the interface induced by misfit dislocations, which accommodates the misfit strain between the two phases. For c‐sapphire and Cu3Ge, the interface planes are (0 0 1) and (0 1 0) planes respectively, and the viewing direction is [1 1 0] zone for both phases. Figure 6(b) shows the schematic illustration of the matching scenario at the interface. The atoms of sapphire within each lattice on (0 0 1) plane are located at the four corners of a rectangle (width = 4.759 Å, length = 8.243 Å), and the lattice of (0 1 0) plane Cu3Ge is also in rectangular shape (width = 4.54 Å, length = 4.22 Å). Therefore, two lattices of Cu3Ge (0 1 0) plane match with one lattice of the c‐sapphire (0 0 1) plane, as illustrated in Figure 6(b). The lattice misfits of the width and length between two systems are ∼4.60% and ∼-2.33%, respectively. The interface region close to that of Figure 6(a) is enlarged for observation using HRTEM and shown in Figure 6(c), to study the atomic structure in detail. The crystal planes of the film and the substrate connect with each other, but slightly bend around the misfit dislocations. Domain matching epitaxy (DME) [28–32] can be used to explain the matching of two phases, since no pseudomorphic Cu3Ge has grown across the interface and the misfit strain is accommodated within several atomic layers at the interface. 9/8 and 8/7 domains alternate with a relative frequency of 0.5 to accommodate the misfit strain perpendicular to the interface, as shown in Figure 6(c).
Planar defects, including twins and stacking faults (SF), are observed to exist in the as‐grown epitaxial Cu3Ge film and shown by HRTEM image (Figure 7). Extrinsic stacking fault is observed to exist close to the lower twin boundary (TB). Individual SF is observed within one grain. The densities of twins and SFs in the as‐deposited Cu3Ge film are close to that of pure Cu thin films with internal strain [33]. Previously in PLD‐deposited films, deformation twins have been shown to be dominant [34]. Therefore, the concept of “generalized planar fault energies (GPFEs)” can explain the formation mechanism of the twins and SFs in the as‐deposited Cu3Ge film. In this concept, multiple intrinsic material properties, including stacking fault energy (SFE), unstable stacking fault energy (USFE), and unstable twin fault energy (UTE), play together to affect the twinning possibility [35]. Alloying elements have been reported to lower the SFE of metals [36], and Cu alloy has been chosen to study twin interaction phenomena in previous studies [33] for the same reason. Thus, a leading partial Cu3Ge is more likely to nucleate and slip due to a small energy barrier. Further, the combined effect of SFE, USFE, UTE, as well as the magnitude and orientation of the local shear stress determines whether the SF is annihilated, maintained, or transformed into a twin. If the nucleation and gliding barrier for the trailing partial is low, it will be easy to annihilate an SF. Due to the existence of SF in our films, the difference between USFE and SFE should be high for Cu3Ge since it determines the energy barrier for the trailing partial to move. Therefore, stacking faults are generated by the movement of leading partials without trailing partials. Twins are then easily generated once a leading partial is emitted and an SF is formed, since the UTE is not quite higher than the USFE.
Grain size is another parameter to affect deformation twinning in nanocrystalline metals and alloys, in addition to GPFE curves. Deformation twinning is easiest at an optimum grain size [37]. Twins in metallic thin films have been reported to lead to excellent properties, including high electromigration resistance [35], high strength and ductility, good mechanical stability, and low electrical resistivity. They can also accommodate residual strains [38]. Furthermore, Cu3Ge with twins will be a better metallization material as the twin boundaries will deter the diffusion.
The local work function of epitaxial Cu3Ge thin film
where
where
The DC component (
When
where
which is between the work functions of n+ and p+‐polysilicon [41]. This value is desirable for epitaxial Cu3Ge thin film to be used as a mid‐gap gate metal even at very low temperatures for applications in CMOS devices, because it would require minimal and symmetric channel implants even at linewidths below 0.5 µm [41].
4. Conclusions
Epitaxial e1‐Cu3Ge thin films are fabricated on
Acknowledgments
FW and NY acknowledge the partial support by the National Science Foundation‐MRSEC program through the Princeton Center for Complex Materials (DMR‐0819860).
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