ZnO Interfaces

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Epitaxial Strain Induced Growth of CuO at CuO/ZnO Interfaces Anette Eleonora Gunnæs, Sandeep Gorantla, Ole Martin Lovvik, Jiantuo Gan, Patricia Almeida Carvalho, Bengt Gunnar Svensson, Edouard V. Monakhov, Kristin Bergum, Ingvild T. Jensen, and Spyros Diplas J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b07197 • Publication Date (Web): 19 Sep 2016 Downloaded from http://pubs.acs.org on September 30, 2016

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The Journal of Physical Chemistry

Epitaxial Strain Induced Growth of CuO at Cu2O/ZnO Interfaces.

A.E. Gunnæsǂ*, S. Gorantlaǂ, O.M. Løvvikǂ,§, J. Ganǂ, P.A. Carvalhob, B.G. Svenssonǂ, E.V. Monakhovǂ, K. Bergumǂ, I. T. Jensen§ and S. Diplas§

ǂ

University of Oslo, Department of Physics/Center for Materials Science and

Nanotechnology, P.O. Box 1048 Blindern, N-0316 Oslo, Norway §

SINTEF Materials and Chemistry, P.O. Box 124 Blindern, Forskningsveien 1, N-0314

Oslo, Norway

Corresponding author: Anette Eleonora Gunnæs: [email protected], phone: +47 91514080 Mail list: [email protected], [email protected], [email protected], [email protected] [email protected], [email protected], [email protected], [email protected], [email protected]

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Abstract

Cu2O/ZnO has been envisaged as a potential material system for the next generation thin film solar cells. So 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 modelling, we have studied the interfaces between single crystal c-axis oriented ZnO and high-quality magnetron sputtered Cu2O films. Strikingly, our study shows that the first ~5 nm of the Cu-oxide films have the structure of the monoclinic CuO phase. The CuO layer is textured with the (111), (111), (111), (111), (111), (111), (111) 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 Cu2O follows resulting in epitaxial 180o rotation twins in the Cu2O layer. For the case with the (100)CuO interfacial plane we have: (111)[110]Cu2O II (100)[011]CuO II (0001)[1120]ZnO. Because of a closer lattice matching of CuO with ZnO and Cu2O, the total strain and energy is reduced compared to a pure (111)Cu2O II (0001)ZnO interface. The existence of CuO is anticipated to be a contributing factor for the low conversion efficiencies obtained experimentally.


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Introduction Cuprous oxide or cuprite (Cu2O) is a low-cost and nontoxic ‘intrinsic’ p-type semiconductor with a bandgap of ~ 2.1 eV. Combined with n-type ZnO, which has a bandgap of 3.4 eV, a Cu2O/ZnO p-n heterojunction would be obtained. Cu2O has a primitive cubic crystal structure (Pn3m; a = 0.427 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 (P63mc; a = 0.325 nm, c = 0.521 nm). A heteroepitaxial relationship between ZnO and Cu2O has been expected due to a relatively small lattice mismatch of 7.6% anticipated for the Cu2O(111)/ZnO(0001) interface.1

A most interesting prospect for the Cu2O/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-film based solar cells.3 – 6 At the present, however, the highest reported conversion efficiencies are only 2% for Cu2O/ZnO heterojunction thin films,3 and 4.12% and 5.38% when an un-doped Ga2O3 or ZnO thin interfacial layer is present in AZO/n-semiconductor/p-Cu2O 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, focussed 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 3 ACS Paragon Plus Environment


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behind the discrepancy between the theoretical and experimentally obtained conversion efficiencies of Cu2O/ZnO heterojunction solar cells, model systems of high quality Cu2O films on O- and Zn- polar single crystal 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 modelling 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.

Experimental The Cu2O 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 [0001] wafers purchased from Tokyo Denpa Co., Ltd. Detailed description of the deposition parameters is reported by Gan, J. et al.12 In brief, the substrate wafers were double-side polished and pre-treated to make sure they were free of any contamination before the sputtering deposition. One side of such a prepared wafer served as the Znpolar 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 4 x 10-4 Pa.


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Cross-section TEM specimens were prepared by the conventional tripod wedge mechanical polishing protocol. The final thinning of the specimens was done by ionmilling for 1 hr 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 G2 60-300 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 mrad and 76-200 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-wave based VASP code 14, 15 using the Perdew-Burke-Ernzerhof gradient approximation16 and the projector augmented wave method.17 The plane-wave cut-off 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, 5 ACS Paragon Plus Environment


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and the two outermost layers of one side 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 5x5 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.

Results and Discussion Figure 1 show 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 XRD 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


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Figure 1. (a) HAADF STEM cross section overview image of the O-polar Cu2O/ZnO and (b) DF TEM image of the Zn-polar Cu2O/ZnO films showing contrast variations in the Cu2O films consistent with polycrystaline films.

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 (c) can be indexed in agreement with a single crystal zone axis pattern ([1 1 2]). However, the complex pattern seen (f) can be described as two single crystal zone patterns ([1 10] and [11 0]) twin related with a 180o 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.19 – 21 The additional reflexions located 1/3 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



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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 have two orientation variants: (111)[110]Cu2O II (0001)[1120]ZnO and 111110 Cu

O 2

II 00011120

ZnO.

Figure 2. SAD patterns from two specimen orientations where the ZnO substrate is oriented along the (a-c) [1100] and (d-e) [1120]ZnO ZA respectively. (a)-(c) show the (111)[110]Cu2O || (0001)[1120]ZnO orientation relationship whereas (d)-(e) the Cu2O twin relationship.


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The orientation relationship 111110 Cu

O 2

|| 00011120 ZnO is in agreement with a

heteroepitaxial relationship between the two phases reported by others and attributed to energetically favourable periodic lattice coincidence between these two crystal lattices.10 Wang et al.22 have 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

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 180o rotation twins where the green spheres represents Cu positions.



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The XRD of the current specimens showed that the cell parameters of the Cu2O films were 0.4318 and 0.4301 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 = 0.427 nm). The twin growth could be assumed to have relaxed the strain induced by epitaxy. However, structural defects located close to the hetero interface could also cause the strain in the grown film to be reduced.

Figure 4 show 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 to the much larger twin domains of the Cu2O film above (Figure 4 (b)).

The XRD 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, polycrystalline structure and in the case of XRD small grain sizes. However, x-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


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layer was not consistent with Cu2O. However, the monoclinic CuO phase with space group C1c1 and lattice parameters: a = 0.46927 nm, b = 0.34283 nm, c = 0.51370 nm and β = 99.546o 24 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 the Cu2O growth.

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)[110]Cu2O II (100)[011]CuO II (0001)[1120]ZnO orientation relationship and (d), (e) and (f) the Cu, Zn and O EDS elemental distributions in (c).



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Based on 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)[110]Cu2O II (100)[011]CuO II (0001)[1120]ZnO orientation relationship which together with (0001)ZnO II (111), (111), (111), (111), (111), (111), (111)CuO II (111)Cu2O represents the common interfaces found. All the CuO planes ((111), (111), (111), (111), (111), (111), (111) 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 0.30066 nm. The Cu-Cu CIBL values at the CuO interfaces are listed in Tab. 1 and the similarity in atomic configuration with {111} planes of the face centred 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), (111), (111), (111), (111), (111)) can act as suitable surfaces for epitaxial growth of {111}Cu2O and this behaviour can justify the presence of Cu2O twin-related grains instead of randomly oriented (111) textured Cu2O grains one could expect with no epitaxial relations.


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Figure 5. A projection of (100) planes in CuO, where the blue and red spheres represent Cu- and O- ions, respectively. The crossed and the uncrossed Cu- ions are on alternating planes. The arrows represents inter ionic Cu-Cu distances: 0.307546 nm (black), 0.310058 nm (orange), 0.34283 nm (red) based on 24.

Table 1. Cu-Cu cation interionic bond lengths (CIBL) on the interfacial planes of CuO based on the CuO lattice parameters: a = 0.46927 nm, b = 0.34283 nm, c = 0.51370 nm and β =99.546o 24. Corresponding CIBL for ZnO and Cu2O are 0.32494 and 0.30066 nm, respectively.

Cu-Cu (nm)

0.375526 0.34283 0.317871 0.310058 0.307546 0.29058

(100)

X

X

X

X

X

X

X

X

X

(111), (111), X

(111), (111), (111) (111), (111)

X

To further rationalize the formation of a CuO interfacial layer under oxygen-poor deposition conditions, DFT modelling was performed. A series of stoichiometric 13 ACS Paragon Plus Environment


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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 lowest lattice mismatch and that a hypothetical interface with low lattice mismatch may be impossible to realize in practice. This is demonstrated in Figure 6 (a) and (b) showing the Cu2O(111)/ZnO(0001) interface with the lowest 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 6 (c) and (d) represents the potential compatibility between the two phases; however, in this case the interface exhibits significantly higher strain (7.8%).

The interface models shown in Figure 6 (c) and (d) (111) [110]Cu2O II (0001)[1120]ZnO, (e) and (f) (111)[110]Cu2O II (100)[011]CuO and (g) and (h) (100)[011]CuO II (0001)[1120]ZnO all show excellent epitaxy in all 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


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apart the two constituting parts leaving two free surfaces, is largest for the 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 constituting phases. The strain is very high in the case of (111)Cu2O / (0001)ZnO; the directional 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 built 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 towards the formation of Cu2O. So 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 the CuO growth. CuO is in this way acting as a buffer layer accommodating the strain between ZnO and Cu2O. 15 ACS Paragon Plus Environment


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The picture that emerges is the following: instead of a direct interface between ZnO and Cu2O as depicted in Figure 6(a) and (b), CuO forms an intermediate layer to accommodate the strain between ZnO and Cu2O. The latter situation can be illustrated by combining Figure 6(c) and (e) or alternatively Figure 6(d) and (f). This corresponds very closely to the experimental micrograph in Figure 4(b).


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Figure 6. Energetically relaxed interface model structures from DFT calculations showing (a), (b) ,(c), (d) (111)Cu2O || (0001)ZnO, (e), (f) (111)Cu2O || (100)CuO and (g), (h) (100)CuO || (0001)ZnO interfaces, seen in the (a), (c), (g) [1100]ZnO (b), (d), (h) [1120]ZnO, (e) [112]Cu2O and (f) [110]Cu2O projections. (a), (b) Interface with the lowest strain obtained by automatic routine, showing very low structural compatibility.

Table 2. Strain and calculated interface energy from the DFT study where vector 1 and 2 represents in plane atomic bond directions at the surface interfaces.

Interface

Eint (eV)

Strain (%)

Vector 1 Vector 2

Volume

(100)CuO / (0001)ZnO

0.47

5.34

5.84

2.50

(111)Cu2O / (0001)ZnO

7.80

7.80

16.22

1.47

(111)Cu2O / (100)CuO

5.01

1.64

6.74

3.14

The strain at the (100)CuO / (0001)ZnO interface was further analyzed experimental by Geometric Phase Analysis (GPA). The GPA ᵋxx 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



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lattice is 4.8 ± 1.7 % larger than the ZnO lattice along Vector 2, which is in agreement with the expected 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 ᵋxx contrast below the interface.

Figure 7. (a),(d) HRSTEM image showing burgers circuits around misfit dislocations at the CuO/ZnO interface, (b), (e) shows corresponding GPA ᵋxx map, (c) FFT of (a) indicating the spots used for GPA and (f), (g) inverse FFT moire images enhancing the (1101) and (1101) ZnO planes indicated by white and red arrows in (d).

An analysis of the Burgers vectors in Figure 7(a) was performed using the right-hand start-finish convention. All dislocations were found to be 90° edge dislocations with the Burgers vector ½ [1100]ZnO and parallel to the interface.25, 26 The white and red arrows in


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Figure 7 correspond to the extra half plane in the ZnO lattice at (1101) and (1101) planes as seen in the Bragg-masked inverse FFT images ((f) and (g)) 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 Tab. 1, Cu-Cu CIBL values of 0.307546 nm (black), 0.310058 nm (orange) and 0.34283 nm (red). In a [001]CuO projection, the Cu-Cu CIBL alternates between 0.307546 nm and 0.310058 nm along the unstrained CuO/ZnO interface with an average Cu-Cu CIBL of 0.308802 nm. An unstrained ZnO lattice has, correspondingly, a CIBL equal to 0.32494 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 Cu2O directly on top of ZnO (0001) are likely to fail. This can be understood from the cubic crystal structure of Cu2O, 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 19 ACS Paragon Plus Environment


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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 towards epitaxial Cu2O on ZnO could be through an electronically benign buffer layer suppressing the formation CuO, as corroborated by an improved conversion efficiency pf Cu2O/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 Cu2O / 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), (111), (111), (111), (111), (111) 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 Cu2O films.

The DFT study of the theoretical (111) [110]Cu2O II (0001)[1120]ZnO and the experimentally observed (100)[011]CuO II (0001)[1120]ZnO and (111)[110]Cu2O II


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(100)[011]CuO interfaces showed that both the strain and interface energy were in favor of epitaxial growth of CuO on the ZnO (0001) surface.

GPA showed evidence of strain in CuO film across CuO/ZnO interfaces is negligible and presence of 90° edge misfit dislocations with the Burgers vector ½ [1100]ZnO parallel to the (100)CuO/(0001)ZnO interface. The misfit dislocations reduce the strain in the CuO layer, making it even more energetically favourable to grow CuO relative to Cu2O.

The CuO interlayer, not resolved by XRD, is anticipated to be inherent to the interface between crystalline Cu2O and ZnO films in heterojunction structures. The CuO layer is highly detrimental for the current rectification and open circuit voltage of the Cu2O/ZnO heterojunction due to the differences in the band alignment and the band gap (2.1 eV vs 1.2 eV) and can be one of the causes for the poor conversion efficiency obtained experimentally for Cu2O/ZnO solar cells relative to that predicted theoretically.

Acknowledgement This work was conducted under the research project ES483391 Development of a Hetero-Junction Oxide-Based Solar Cell Device (HeteroSolar), financially supported by the Research Council of Norway (RCN) through the RENERGI program.



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ZnO Interfaces

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