Epitaxial Strain-Induced Growth of CuO at Cu 2 O/ZnO Interfaces

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1 pubs.acs.org/jpcc Epitaxial Strain-Induced Growth of CuO at Cu 2 O/ZnO Interfaces A. E. Gunnæs,*, S. Gorantla, O. M. Løvvik,, J. Gan, P. A. Carvalho, B. G. Svensson, E. V. Monakhov, K. Bergum, I. T. Jensen, and S. Diplas Department of Physics/Center for Materials Science and Nanotechnology, University of Oslo, P.O. Box 1048, Blindern, N-0316 Oslo, Norway SINTEF Materials and Chemistry, P.O. Box 124, Blindern, Forskningsveien 1, N-0314 Oslo, Norway ABSTRACT: Cu 2 O/ZnO has been envisaged as a potential material system for the next-generation thin film solar cells. Thus far, the experimental efforts to obtain conversion efficiencies close to the theoretically predicted value have failed. Combining aberration-corrected (scanning) transmission electron microscopy and density functional theory modeling, we studied the interfaces between single-crystal c- axis-oriented ZnO and high-quality magnetron-sputtered Cu 2 O films. Strikingly, our study shows that the first 5 nm of the Cu oxide films has the structure of the monoclinic CuO phase. The CuO layer is textured with the (111), (111 ), (1 11), (11 1 ) (1 1 1), (11 1), (1 11,) and (100) planes parallel to the (0001) and (0001 ) ZnO interfaces. The ionic arrangement on these planes resembles the hexagonal arrangement of the ZnO interface, and epitaxy exists across the interface. A continued epitaxial growth of [111]-oriented Cu 2 O follows resulting in epitaxial 180 rotation twins in the Cu 2 O layer. For the case with the (100) CuO interfacial plane we have (111)[11 0] Cu2O (100)[011] CuO (0001)[112 0] ZnO. Because of a closer lattice matching of CuO with ZnO and Cu 2 O, the total strain and energy is reduced compared to a pure (111) Cu2O (0001) ZnO interface. The existence of CuO is anticipated to be a contributing factor for the low conversion efficiencies obtained experimentally. INTRODUCTION Cuprous oxide or cuprite (Cu 2 O) is a low-cost and nontoxic intrinsic p-type semiconductor with a band gap of 2.1 ev. Combined with n-type ZnO, which has a band gap of 3.4 ev, a Cu 2 O/ZnO p n heterojunction would be obtained. Cu 2 O has a primitive cubic crystal structure (Pn3 m; a = nm) where the Cu + and O 2 ions are located on fcc and bcc sublattices, respectively. ZnO has a unit cell with a hexagonal wurzite crystal structure (P6 3 mc; a = nm, c = nm). A heteroepitaxial relationship between ZnO and Cu 2 O has been expected due to a relatively small lattice mismatch of 7.6% anticipated for the Cu 2 O(111)/ZnO(0001) interface. 1 A most interesting prospect for the Cu 2 O/ZnO materials system is related to its theoretically predicted conversion efficiency of 18%. 2 With its low cost, nontoxic constituents, and potentially high conversion efficiency it has been envisaged as a potential materials system for the next-generation thin-filmbased solar cells. 3 6 At the present, however, the highest reported conversion efficiencies are only 2% for Cu 2 O/ZnO heterojunction thin films 3 and 4.12% and 5.38% when an undoped Ga 2 O 3 or ZnO thin interfacial layer is present in AZO/n-semiconductor/p-Cu 2 O heterojunctions. 7,8 Hence, there is an increasing effort to find solutions to improve the conversion efficiency. The majority of the research efforts in this field are, however, focused on developing synthesis methods, and only few investigations are devoted to in-depth atomic-scale investigations. 1,9 11 In order to obtain a fundamental understanding about the reasons behind the discrepancy between the theoretical and the experimentally obtained conversion efficiencies of Cu 2 O/ZnO heterojunction solar cells, model systems of high-quality Cu 2 O films on O- and Zn-polar singlecrystal ZnO were used in the present study. A combination of atomically resolved scanning transmission electron microscopy (STEM), electron diffraction (ED), and density functional theory (DFT) based modeling was employed to investigate the material system and, in particular, the substrate film interface, as the performance of thin film-based devices is well known to be largely affected by the quality of its interfaces. EPERIMENTAL SECTION The Cu 2 O films investigated in the present study were grown by direct current/radio frequency (dc/rf) reactive magnetron sputtering on c-axis-oriented single-crystal ZnO wafers purchased from Tokyo Denpa Co., Ltd. Detailed description of the deposition parameters is reported by Gan et al. 12 In brief, the substrate wafers were double-side polished and pretreated to make sure they were free of any contamination before the sputtering deposition. One side of such a prepared wafer served as the Zn-polar substrate while the other side as the O-polar substrate. During the magnetron sputtering deposition, the substrate temperature was maintained at 400 C and the base pressure of the chamber was below Pa. Received: July 18, 2016 Revised: September 13, 2016 Published: September 19, American Chemical Society 23552

2 Cross-section TEM specimens were prepared by the conventional tripod wedge mechanical polishing protocol. The final thinning of the specimens was done by ion milling for 1 h in a Fischione 1010 instrument with beam energy and current of 5 kv and 5 ma, respectively, and milling angle of ±6 on both sides of the specimen. A JEOL 2100F microscope operated at 200 kv and an FEI Titan G microscope equipped with a DCOR probe Cs-aberration corrector, operated at 300 kv, were used in the present study. STEM high-angle annular dark field (HAADF) Z-contrast imaging was performed using the FEI microscope with a probe current of 100 pa and probe convergence and collection angles of 22 and mrad range, respectively. The resulting spatial resolution was 0.08 nm. HAADF images show primarily atomic number (Z) contrast, and in the present case, high-resolution (HR) HAADF images offer advantage over HRTEM ones owing to the direct correlation between the image contrast and the cation positions in the sample. Fast Fourier transform (FFT) analysis was used on HR images, and lattice strain measurements were carried out using the geometric phase analysis (GPA) plugin for the DigitalMicrograph software.13 The DFT calculations were performed with the plane-wavebased VASP code14,15 using the Perdew Burke Ernzerhof gradient approximation16 and the projector augmented wave method.17 The plane-wave cutoff was 400 ev, and the k-point density was at least 4 points per Å 1 in each unit cell direction. The force relaxation criterion was 0.05 ev/å, and the two outermost layers of one side of the slab structure were fixed. The interface structures were built using the ASE environment.18 The strain was divided equally between the two sides of the interface, and a 5 5 two-dimensional potential energy surface (PES) was generated by lateral translations of one of the sides relative to the other. Each PES point was obtained by relaxation of the whole interface structure, and the lowest energy was taken to represent the interface. A 1 nm thick layer of vacuum was inserted between the two phases on one side, so only one interface was formed. For reference energies, an additional layer of 1 nm vacuum was inserted between the two phases on the other side, creating two separate slabs divided by vacuum on both sides. The interface energy was defined as the energy difference between the interface model and the reference model. Figure 1. (a) HAADF STEM cross-section overview image of the Opolar Cu2O/ZnO, and (b) DF TEM image of the Zn-polar Cu2O/ ZnO films showing contrast variations in the Cu2O films consistent with polycrystalline films. Figure 2. SAD patterns from two specimen orientations where the ZnO substrate is oriented along the (a c) [11 00] and (d f) [1120 ]ZnO ZA. (a c) (111)[110 ]Cu2O (0001)[1120 ]ZnO orientation relationship; (e and f) Cu2O twin relationship. and 2/3 between {111} rows are due to multiple scattering, i.e., the two twin regions are partly overlapping and the electrons will, due to their short scattering mean free path, scatter from both twin orientations resulting in additional reflexions. From the indexed patterns in Figure 2, it can be concluded that the substrate and the Cu2O films has the two orientation variants: ( ) [ ] C u 2 O ( ) [ ] Z n O a n d ( ) [1 10]Cu2O (0001)[112 0]ZnO. The orientation relationship (111)[11 0 ] Cu2O (0001) [1120 ] ZnO is in agreement with a heteroepitaxial relationship between the two phases reported by others and attributed to energetically favorable periodic lattice coincidence between these two crystal lattices.10 Wang et al.22 reported that the atomic arrangement of the first atomic layer of Cu2O predicts the atomic arrangement throughout the film. This would imply that one get growth of textured Cu2O films but not necessarily an epitaxial relationship between Cu2O and ZnO. The observed twinned growth of Cu2O is, however, regarded as evidence of epitaxial growth. As shown in Figure 3, the Cu ions across the grain boundary have a changed stacking order of their {111} Cu planes (B, C, A, B, C... and C, B, A, B, C...). This is consistent with two different starting growth sites on a substrate representing an A layer. Such twins can be described as epitaxial twins and have been reported for other oxide-based epitaxial heterostructures.23 RESULTS AND DISCUSSION Figure 1 shows overview images of the cross sections of the (a) O-polar Cu2O/ZnO and (b) Zn-polar Cu2O/ZnO specimens. The Cu2O films are in both cases dense and of high quality. Prior RD analysis of the specimens showed strong (111)Cu2O and (0002)ZnO diffraction peaks12 consistent with previous reports where it has been interpreted as evidence for a heteroepitaxial relationship between Cu2O and ZnO due to the 7.6% mismatch between the two lattices.1 Representative selected area diffraction (SAD) patterns from the Cu2O films and the ZnO substrates are shown in Figure 2. The SAD pattern from the Cu2O films in Figure 2c can be indexed in agreement with a single-crystal zone axis pattern ([1 1 2]). However, the complex pattern seen in Figure 2f can be described as two single-crystal zone patterns ([1 10] and [110 ]) twin related with a 180 rotation around the common [111] axis. Such a twin relation has previously been reported for Cu2O nanoparticles and in Cu2O film upon oxidation of a single-crystal Cu The additional reflexions located 1/

3 to the much larger twin domains of the Cu2O film above (Figure 4a and 4b). The RD analysis and SAD from interfacial regions (Figure 2) did not show indications of secondary phases. In both cases this can be explained by the relative low volume fraction of the secondary phase and polycrystalline structure and in the case of RD small grain sizes. However, -ray energy-dispersive spectroscopy (EDS) and fast Fourier transform (FFT) image analysis of HR STEM images, as illustrated in Figure 4, performed at various positions along the interfacial regions showed that the interfacial layer was not consistent with Cu2O. However, the monoclinic CuO phase with space group C1c1 and lattice parameters a = nm, b = nm, c = nm, and β = did give a good fit to the FFT data and is consistent with the EDS data. This is unexpected as the CuO layer was grown under oxygen-poor deposition conditions optimized for Cu2O growth. On the basis of the FFT analysis of the HR STEM data, it is clear that the CuO grains grow on the ZnO substrate with preferred orientation relationships. Figure 4 illustrates the (111)[11 0]Cu2O (100)[011]CuO (0001)[112 0]ZnO orientation relationship which together with (0001)ZnO (111), (111 ), (1 11), (11 1 ) (1 1 1), (11 1), (1 11 )CuO (111)Cu2O represents the common interfaces found. All CuO planes ((111), (1 11), (11 1 ), (1 1 1) (11 1), (1 11 ), and (100)) represent dense planes of Cu ions with Cu arrangements similar, but distorted, to the one present in {111}Cu2O planes. In the {111}Cu2O planes the Cu Cu cation interionic bond length (CIBL) is nm. The Cu Cu CIBL values at the CuO interfaces are listed in Table 1, Figure 3. (a) STEM image illustrating the stacking order of A,C,B,A... and A,B,C,A... in the twinned Cu2O grains. (b) Illustration of the possible starting points with respect to orientation of the Cu2O lattice giving rise to 180 rotation twins where the green spheres represent Cu positions. The RD of the current specimens showed that the cell parameters of the Cu2O films were and nm for the O- and Zn-polar substrates, respectively.12 This shows that the lattice is only 1.1% and 0.7% strained relative to unstrained Cu2O (a = nm). The twin growth could be assumed to have relaxed the strain induced by epitaxy. However, structural defects located close to the heterointerface could also cause the strain in the grown film to be reduced. Figure 4 shows representative HAADF STEM images from Cu2O/ZnO interfacial regions where bright spots correspond to Zn- and Cu-atom positions in the substrate and film region, respectively. Independent of substrate polarity, the first 5 nm of the grown Cu O films show distinct different atomic arrangement and grain sizes (5 10 nm in diameter) compared Figure 4. STEM-HAADF images from the (a and b) O-polar and (c) Zn-polar Cu2O/ZnO interfaces where the corresponding FFT patterns from b show the (111)[11 0]Cu2O (100)[011]CuO (0001)[112 0]ZnO orientation relationship and d, e, and f the Cu, Zn, and O EDS elemental distributions in c

4 Table 1. Cu Cu Cation Interionic Bond Lengths (CIBL) on the Interfacial Planes of CuO Based on the CuO Lattice Parameters Cu Cu (nm) (100) (111), (111 ), (1 11), (11 1 ), (1 1 1) (11 1), (1 11) a = nm, b = nm, c = nm, and β = Corresponding CIBL for ZnO and Cu2O are and nm, respectively. a Figure 5. Projection of (100) planes in CuO, where blue and red spheres represent Cu and O ions, respectively. Crossed and the uncrossed Cu ions are on alternating planes. Arrows represent interionic Cu Cu distances: nm (black), nm (orange), and nm (red) based on ref 24. and the similarity in atomic configuration with {111} planes of the face-centered cubic Cu sub lattice of Cu2O (Figure 3) is illustrated for the (100)CuO plane in Figure 5. As such, either (100)CuO or any of the 111 type of planes of CuO ((111), (111 ), (1 11), (11 1 ) (1 1 1), (11 1), (1 11 )) can act as suitable surfaces for epitaxial growth of {111}Cu2O, and this behavior can justify the presence of Cu2O twin-related grains instead of randomly oriented (111) textured Cu2O grains one could expect with no epitaxial relations. To further rationalize the formation of a CuO interfacial layer under oxygen-poor deposition conditions, DFT modeling was performed. A series of stoichiometric CuO/ZnO, Cu2O/ ZnO, and CuO/Cu2O interface models were constructed after an automatic interface model search, based on the ASE environment,18 and relaxed to their low-energy configuration. Many hypothetical interface models with relatively small lattice mismatch were identified. However, in order to obtain a stable epitaxial interface, atomic columns have to match as much as possible at the interface plane and the local stoichiometry must be appropriate. This means that the most stable interface is not necessarily the one with the lowest lattice mismatch and that a hypothetical interface with low lattice mismatch may be impossible to realize in practice. This is demonstrated in Figure 6a and 6b, showing the Cu2O(111)/ZnO(0001) interface with the lowest volumetric strain (3.1%) found with our automatic routine. The low structural compatibility due to nonmatching atomic columns indicates that this model is not likely to form. The model shown in Figure 6c and 6d represents the potential compatibility between the two phases; however, in this case the interface exhibits significantly higher strain (7.8%). The interface models that are shown in Figure 6c and 6d (111) [11 0]Cu2O (0001)[112 0]ZnO, Figure 6e and 6f (111) [11 0]Cu2O (100)[011]CuO, and Figure 6g and 6h (100) [011]CuO (0001)[112 0]ZnO show excellent epitaxy in all Figure 6. Energetically relaxed interface model structures from DFT calculations showing (a d) (111) Cu2O (0001) ZnO, (e and f) (111)Cu2O (100)CuO, and (g and h) (100)CuO (0001)ZnO interfaces, seen in the (a, c, g) [11 00]ZnO, (b, d, h) [112 0]ZnO, (e) [1 1 2]Cu2O, and (f) [1 10]Cu2O projections. (a and b) Interface with the lowest strain obtained by automatic routine, showing very low structural compatibility. directions and the lowest strain. Supporting the validity of the selection scheme, they also represent two of the experimentally observed interfaces (Figure 4). A comparative evaluation of the stability of the competing interfaces has been performed based on the calculated interface energy Eint. Eint, defined as the energy required to pull apart the two constituting parts leaving two free surfaces, is largest for the 23555

5 The picture that emerges is the following: instead of a direct interface between ZnO and Cu2O as depicted in Figure 6a and 6b, CuO forms an intermediate layer to accommodate the strain between ZnO and Cu2O. The latter situation can be illustrated by combining Figure 6c and 6e or alternatively Figure 6d and 6f. This corresponds very closely to the experimental micrograph in Figure 4b. The strain at the (100)CuO/(0001)ZnO interface was further analyzed experimentally by geometric phase analysis (GPA). The GPA εxx map in Figure 7b shows the local strain at the CuO/ZnO interface. The color scheme clearly shows that the CuO lattice is mostly unstrained and that the CuO/ZnO lattice mismatch is relaxed by multiple misfit dislocations distributed along the interface and in its vicinity. Indeed, the measured CuO lattice is 4.8 ± 1.7% larger than the ZnO lattice along [1 100], which is in agreement with the expected unrelaxed theoretical 5.8% deviation for unstrained conditions. These results demonstrate that strain is minor in the CuO layer and mostly accumulated in the ZnO layer as there is much stronger variation of εxx contrast below the interface (Figure 7e). An analysis of the Burgers vectors in Figure 7a was performed using the right-hand start finish convention. All dislocations were found to be 90 edge dislocations with Burgers vector 1/2 [11 00]ZnO and parallel to the interface.25,26 The white and red arrows in Figure 7 correspond to the extra half plane in the ZnO lattice at (1 101 ) and (1 101) planes as seen in the Bragg-masked inverse FFT images (Figure 7f and 7g) where the two dislocation components of the 90 full edge dislocation can be clearly seen. The dislocation cores are spaced 4 nm apart along the interface (Zn dislocation core spaced with 14th, 16th, and 19th Zn Zn CIBL). The (100)CuO plane has in its unstrained state, as illustrated in Figure 5 and listed in Table 1, Cu Cu CIBL values of (black), (orange), and nm (red). In a [001]CuO projection, the Cu Cu CIBL alternates between and nm along the unstrained CuO/ZnO interface with an average Cu Cu CIBL of nm. An unstrained ZnO most stable interface. The resulting DFT-calculated interface energy Eint showed that both the (100)CuO/(0001)ZnO and the (111)Cu2O/(100)CuO display relatively high interface energies (2.5 and 3.1 ev), whereas the (111)Cu2O/(0001)ZnO interface energy is lower (1.5 ev). In other words, the hypothetical (111)Cu2O/(0001)ZnO interface is thermodynamically relatively unstable, favoring the presence of a CuO interfacial layer. As seen from Table 2, the main difference between the interface models is related to the strain between the Table 2. Strain and Calculated Interface Energy from the DFT Study where Vectors 1 and 2 Represent in-plane Atomic Bond Directions at the Interfaces strain (%) interface vector 1 vector 2 volume Eint (ev) (100)CuO/(0001)ZnO (111)Cu2O/(0001)ZnO (111)Cu2O/(100)CuO constituting phases. The strain is very high in the case of (111)Cu2O/(0001)ZnO; the linear strain is 7.8%, and the volume strain is 16.2%. This is associated with a large energy penalty, and it also means that if such an interface is formed, it would most likely abound with defects (such as dislocations) which may tend to relieve the build up of the stresses. With a calculated volumetric strain of 5.8%, the (100)CuO/ (0001)ZnO interface is found to be much more favorable and hence gives a strong impetus to form epitaxial CuO instead of Cu2O on top of ZnO. Nevertheless, the low oxygen content during synthesis means that there is a thermodynamic drive toward the formation of Cu2O. Thus, even if the epitaxial growth of CuO is preferred over Cu2O at the ZnO surface, it is reasonable to assume that Cu2O will start forming when strain is reduced to a given level during CuO growth. CuO is in this way acting as a buffer layer accommodating the strain between ZnO and Cu2O. Figure 7. (a and d) HRSTEM image showing burgers circuits around misfit dislocations at the CuO/ZnO interface, (b and e) corresponding GPA εxx map, (c) FFT of a indicating the spots used for GPA, and (f and g) inverse FFT moire images enhancing the (1 101 ) and (1 101) ZnO planes indicated by white and red arrows in d

6 lattice has, correspondingly, a CIBL of nm in the (001) plane. Hence, in an unstrained state, along a (100) CuO / (0001) ZnO interface, one has that 19* Zn Zn CIBL = 18* averaged Cu Cu CIBL. This is in good agreement with the observation of the dislocation spacing in Figure 7 pointing to total low strain in the CuO lattice. The presence of the misfit dislocations along the CuO/ZnO interface reduces the strain relative to the DFT-calculated value, where no dislocations were taken into account and increasing the driving force to grow CuO even more. It is thus clear that attempts to form an epitaxial layer of Cu 2 O directly on top of ZnO (0001) are likely to fail. This can be understood from the cubic crystal structure of Cu 2 O, which gives very few possible combinations of cell parameters (and atomic column arrangements). CuO has, on the other hand, a monoclinic crystal structure with three different lattice constants and can thus provide a large variety of interface unit cell sizes and shapes. This flexibility turns out to be more important than processing parameters when it comes to growing an epitaxial Cu O film on top of ZnO. Thus, a path toward epitaxial Cu 2 O on ZnO could be through an electronically benign buffer layer suppressing the formation CuO, as corroborated by an improved conversion efficiency of Cu 2 O/ZnO heterojunction solar cell having an intermediate interfacial layer. 7,8 CONCLUSIONS By atomically resolved STEM and FFT analysis it was revealed that single-crystalline (0001) and (0001 ) ZnO substrates have an unintentional 5 nm thin intermediate layer of CuO at the Cu 2 O/ZnO interfaces irrespective of ZnO surface polarity. The growth of CuO is textured on the ZnO substrates, showing evidence of a continuation of densely packed CuO planes of cations such as (111), (111 ), (1 11), (11 1 ), (1 1 1), (11 1), (1 11 ), and (100) parallel with the (0001) ZnO, (0001 ) ZnO, and (111) Cu2O interfaces. The densely packed CuO planes are found to act as a template, facilitating growth of epitaxial twinned Cu 2 O films. The DFT study of the theoretical (111)[11 0] Cu2O (0001)- [112 0] ZnO and the experimentally observed (100)[011] CuO (0001)[112 0] ZnO and (111)[11 0] Cu2O (100)[011] CuO interfaces showed that both the strain and the interface energy were in favor of epitaxial growth of CuO on the ZnO (0001) surface. GPA showed evidence of localized strain associated with 90 edge misfit dislocations with the Burgers vector 1/2 [11 00] ZnO along the (100) CuO /(0001) ZnO interface. The misfit dislocations reduce the strain in the CuO layer, making it even more energetically favorable to grow CuO relative to Cu 2 O. 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