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Epitaxial Cu3Ge Thin Film: Fabrication, Structure, and Property

Written By

Fan Wu and Nan Yao

Submitted: November 9th, 2015 Reviewed: May 4th, 2016 Published: July 27th, 2016

DOI: 10.5772/64060

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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 [57], Si [4, 810], Ge [10, 11], YBa2Cu3O7-x [12], SixGe1-x [8, 9, 13], Ta/TaN [14] and GaN [15]), but also various deposition methods have been used for Cu3Ge thin film fabrication. For example, multiple physical vapor deposition methods, such as electron beam deposition, sputtering [1, 2, 4, 11, 12, 16] and thermal evaporation [17], or chemical approaches such as vapor‐solid reaction [14], have been exploited. Despite the achievements in this field, the Cu3Ge films reported up to now were mostly polycrystalline with impurity phase [1, 2, 413]. Consequently, Cu3Ge films with better crystallinity are highly desired, in order to minimize diffusion paths (grain boundaries) and lower electrical resistivity.

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 + 0.02 eV, rendering it a desirable mid‐gap gate metal to be used for applications in CMOS devices. Since sapphire has been widely used as a substrate in high‐power and high‐frequency CMOS integrated circuits, especially for high‐power radio‐frequency applications, the integration of Cu3Ge as contacts or interconnections onto sapphire substrate will have potential applications in the semiconductor industry.

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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 [1826]), 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 θ–2θ scan analysis, with Cu‐Ka X‐ray source. Atomic force microscopy (AFM) is exploited to obtain height images for characterization of thin film morphology. All AFM scans are performed in tapping mode, using cantilevers with a resonance frequency close to 330 kHz. The AFM image resolution is 512 × 512 pixels, and scan rate is 1.5 × 1.5 µm2. Gwyddion is used to process the obtained AFM images.

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.

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3. Results and discussion

Figure 1.

XRD θ–2θ patterns for samples on c‐plane sapphire (Al2O3{0 0 1}). For all five samples, the laser shot frequency is 1 Hz, the deposition temperature is 400°C, and 90 repetitions of depositions of Cu/Ge were performed, while numbers of laser pulses for Cu and Ge depositions are different: sample 1: Cu:Ge=35:5; sample 2: Cu:Ge=25:5; sample 3: Cu:Ge=15:5; sample 4: Cu:Ge=14:2; and sample 5: Cu:Ge=7:1.

To characterize the microstructures of the deposited thin films, XRD θ–2θ patterns are obtained and shown in Figure 1. Since Cu/Ge layer thickness in each repetition is decreasing from sample 1 to sample 5, the crystallinities of Cu3Ge films are improved. The two peaks corresponding to c‐sapphire (0 0 6) and (0 0 1 2) planes exist in all XRD θ–2θ patterns, showing the substrate orientation. The film of sample 1 is polycrystalline Cu3Ge with impurity Ge phase, as evidenced by the peak at 27.32° belonging to Ge (1 1 1) plane, and multiple peaks belonging to Cu3Ge. Note that for sample 1, the deposited Cu/Ge layers in one deposition repetition were too thick to react with each other completely, such that extra Ge phase exists. The impurity Ge phase also exists for samples 2 and 3 (as evidenced by the peak at 27.32° for Ge (1 1 1)), due to the high Ge:Cu pulse ratio in each repetition. Furthermore, the crystallinities of Cu3Ge films are also improved for samples 2 and 3, as evidenced by the disappearance of peaks belonging to Cu3Ge (-1 3 1) and (0 0 4) planes. For samples 4 and 5, no impurity Ge phase exists, due to the reduced pulse numbers of Ge/Cu. More importantly, the crystallinities of Cu3Ge films are significantly improved with only two Cu3Ge peaks remaining, proving that bi‐epitaxial Cu3Ge thin films have been grown on c‐substrate.

Figure 2.

(a)–(e) The plan‐view images of samples (1)–(5) respectively, observed by optical microscope.

Figure 3.

Typical plan‐view SEM image of the Cu3Ge film in sample 5, showing a uniform Cu3Ge film grown via 3‐D growth mode, without the formation of chunks and excessive Ge grains.

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.

Figure 4.

Typical AFM image of Cu3Ge film on c‐sapphire in sample 5, showing the diameters of these islands to be ∼300–500 nm, with heights ranging from 180 to 240 nm.

Figure 5

(a) A typical bright‐field cross‐section TEM image showing the film morphology of sample 5 and (b) the [1 1 0] zone‐axis pattern (ZAP) of sample 5, clearly demonstrating an epitaxial relationship between the single‐crystal Cu3Ge island and c‐sapphire substrate [42].

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 c‐plane sapphire via ‘layer‐plus‐island growth’ mode following a two‐step process. Initially, the Cu3Ge grew in a layer‐by‐layer fashion on the c‐plane sapphire to form an ultrathin uniform film up to several monolayers thick; beyond a critical thickness, the Cu3Ge growth continued through the nucleation and coalescence of adsorbate islands. A sharp interface without secondary phases is observed between c‐sapphire and epitaxial Cu3Ge film, over a large area. Selected area diffraction pattern (SADP) analysis (Figure 5(b)) was performed at the Cu3Ge/sapphire interface for studying orientation relationships. Overlapping of the low‐order diffraction points of film and substrate corroborates their epitaxial relationship. A careful examination does show the splitting of higher order diffraction points of two phases.

Figure 6

(a) An HRTEM image showing the atomic structure at the interface region between the Cu3Ge island and c‐sapphire substrate. The viewing direction is [1 1 0] zone of c‐sapphire ([1120] zone in Miller‐Bravais Indices system), (b) the schematic illustration of the matching scenario at the interface between e1‐Cu3Ge (0 1 0) and c‐sapphire (0 0 1), and (c) the HRTEM image at a higher magnification showing the atomic structure at the interface in detail, at the region close to that of (a) [42].

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) [2832] 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).

Figure 7

A typical HRTEM image showing the planar defects in Cu3Ge thin film, in which twins and stacking faults coexist [42].

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 ϕCu3Ge was measured by KPFM, showing the fundamental electronic property which affects both electron emission through the surface and electronic trajectories near the surface [39]. KPFM measures the work function of solid surfaces at atomic or molecular scales, demonstrating information about composition and electronic state of the local structures on the surface. When the conducting tip and the sample are brought in contact, a net electric current would flow between them until the Fermi levels are aligned. During measurement, a voltage is applied between tip and sample, consisting of a DC‐bias and an AC‐voltage:

V=( V DC V CPD )+ V AC sin( ωt )E1

where VCPD is the contact potential difference between a conductive tip and a sample, VACsin(ωt) and VDC are the applied AC voltages of frequency ω and DC voltage on the tip. The electrostatic force in a capacitor can be written as:

F= 1 2 dC dz V 2 E2

where C is the tip‐sample capacitance, z is the tip‐sample distance, and V is the tip‐sample voltage. Combining the above two equations, the electrostatic force can be split up into three contributions:

F= F DC + F ω + F 2ω E3

The DC component (FDC) contributes to the topographical signal, the term Fω measures the contact potential, and the contribution F can be used for capacitance microscopy. They can be expressed as following:

F dc = dC dz [ 1 2 ( V DC V CPD ) 2 + 1 4 V AC 2 ]E4
F ω = dC dz ( V DC V CPD ) V AC sin( ωt )E5
F2ω=14dCdzVAC2cos(2ωt)E6

When VDC = VCPD, the electrostatic force component measured at frequency ω and the oscillating amplitude would be zero. Consequently, VDC can track VCPD at each point of the scan area using a feedback circuit. Once VCPD is obtained, the local work function of Cu3Ge thin film ϕCu3Ge can be calculated as:

eVCPD=φtipφCu3GeE7

where ϕtip is the work function of the conductive tip (∼4.90 + 0.02 eV) [40], calibrated with a freshly cleaved highly oriented pyrolytic graphite (HOPG, Grade 1 SPI, 4.65 eV) [40] before the experiment. Figure 8(a) and (b) shows the topographical image and the corresponding surface potential image obtained for epitaxial Cu3Ge thin film. Small grains of tens of nanometers can be seen clearly from the topographical image, while an almost uniform distribution of VCPD within each grain was observed from the contact potential image. The “band profiles” derived from the middle regions (between two white dashed lines) of Figure 8(a) and (b) are shown in Figure 8(c) and (d), demonstrating the height and VCPD variation. The average VCPD for epitaxial Cu3Ge thin film was measured to be ∼0.43 V. Therefore, the work function of Cu3Ge thin film is:

φCu3Ge=φtipeVCPD=(4.9±0.02)0.43=(4.47±0.02)eVE8

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].

Figure 8.

KPFM characterization of epitaxial Cu3Ge thin film. (a) The topographical image, (b) the corresponding surface potential image, (c) the band profile derived from the middle region (between two white dashed lines) of (a) and (d) the band profile derived from the middle region (between two white dashed lines) of (b).

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4. Conclusions

Epitaxial e1‐Cu3Ge thin films are fabricated on c‐plane sapphire substrate. The crystallinity of the as‐grown e1‐Cu3Ge thin film is improved, due to the modified deposition route. The e1‐Cu3Ge phase (orthorhombic) is grown by the paradigm of domain matching epitaxy on c‐plane sapphire, where two structures and their symmetries are completely different. In addition, twins are observed and studied in Cu3Ge thin films, and their formation mechanism is explained by the concept of generalized planar fault energies (GPFEs). Twinning will reduce the diffusion of Cu and Ge atoms into the substrate or adjacent layers, and enhance the mechanical stability of Cu3Ge thin films due to the improved strength and ductility and higher electromigration resistance. 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 to be used for applications in CMOS devices. Therefore, the present epitaxial Cu3Ge thin film with controlled crystallinity and defect structure is a promising candidate for the next‐generation metallization material in the semiconductor industry.

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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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Written By

Fan Wu and Nan Yao

Submitted: November 9th, 2015 Reviewed: May 4th, 2016 Published: July 27th, 2016