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Microstructure analysis of IrO2 thin films
Author’s Accepted Manuscript Microstructure analysis of IrO2 thin films Xiuyi Hou, Ryota Takahashi, Takahisa Yamamoto, Mikk Lippmaa
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S0022-0248(16)30974-5 http://dx.doi.org/10.1016/j.jcrysgro.2016.12.104 CRYS23963
To appear in: Journal of Crystal Growth Cite this article as: Xiuyi Hou, Ryota Takahashi, Takahisa Yamamoto and Mikk Lippmaa, Microstructure analysis of IrO2 thin films, Journal of Crystal Growth, http://dx.doi.org/10.1016/j.jcrysgro.2016.12.104 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructure analysis of IrO2 thin films Xiuyi Houa,∗, Ryota Takahashia,b , Takahisa Yamamotoc , Mikk Lippmaaa a Institute
for Solid State Physics, University of Tokyo, Chiba 277-8581, Japan b JST, PRESTO, Saitama 332-0012, Japan c Graduate School of Engineering, Nagoya University, Aichi 464-8603, Japan
Abstract We have grown IrO2 thin films on TiO2 (110) substrates to determine the pulsed laser deposition growth window for iridates. Relaxed IrO2 films were obtained at a growth temperature of 500◦ C and background oxygen pressure of 100 mTorr; otherwise, either pure Ir metal films or evaporative Ir loss were observed. Although x-ray Φ-scan measurement indicated that the films were epitaxial, a distinct grain structure was seen by atomic force microscopy and transmission electron microscopy. The grain boundaries were found to limit the conductivity of films at low temperature. It appeared that strain relaxation leads to stacking faults at grain boundaries. Keywords: Laser epitaxy, Defects, Oxides, X-ray diffraction
1. Introduction Recent interest in Ir oxides has been stoked by the prediction of topological states appearing in pyrochlore-type iridate films strained along the [111] crystal direction.[1, 2] Among the iridate pyrochlores, Pr2 Ir2 O7 has been shown by 5
angle-resolved photoemission spectroscopy (ARPES) to possess a single-point Fermi node at the Γ-point[3], making it a good candidate for epitaxial strain experiments if stoichiometric and atomically flat surfaces can be prepared. However, a common problem for the growth of iridate thin films by physical vapor
∗ Corresponding
author URL: [email protected] (Xiuyi Hou)
Preprint submitted to Journal of Crystal Growth
December 24, 2016
deposition techniques such as pulsed laser deposition (PLD), is the volatility of 10
nonstoichiometric iridium oxides. This can lead to severe stoichiometry deviations in thin films and it is the main reason why pyrochlore iridate thin films have so far been fabricated mostly by solid-phase epitaxy, where the film precursor material is deposited at room temperature and crytallization occurs during post-annealing.[4] While this process produces thin films with excellent struc-
15
tural properties, it unavoidably leads to a relaxed lattice with a compositionally degraded surface, making the films unsuitable for ARPES studies. We therefore explore in this work the thin film growth process of the simplest iridate, IrO2 , to determine the PLD process parameter window where the Ir4+ state can be stabilized. Besides the process optimization, we recognize that IrO2 is
20
an interesting material for a number of reasons, and we analyze the microstructure, transport behavior, and the strain relaxation mechanism in IrO2 (110) films grown on TiO2 (110) substrates. Iridium dioxide belongs to a small family of oxides that are good metals, together with RuO2 , ReO2 , OsO2 , MoO2 , and WO2 .[5] The Fermi level in IrO2
25
is located close to the top of the Ir 5d t2g band, leading to nonmagnetic metallic behavior.[6, 7, 8] IrO2 thin films have been used as a spin-to-current detector in spin valves [9] and may be useful in devices with magnetically switchable majority carrier type.[10] Other well-known applications of bulk or polycrystalline IrO2 are biocompatible electrodes,[11] oxygen barriers in semiconductor
30
devices,[12] electrodes in ferroelectric devices,[13] and electrocatalytic oxygen evolution eletrodes for water splitting.[14] The main difficulty in growing IrO2 films is the difficulty of oxidizing Ir metal and the volatility of nonstoichiometric Ir oxides.[15, 16] IrO2 films have been grown by various techniques, including molecular beam epitaxy (MBE),[10]
35
sputtering,[17] and PLD.[18, 19] Although it is common to use a metallic Ir target for film deposition and allow the ambient oxygen to oxidize the deposited metal at the film surface, it is known that in PLD growth, oxygen is also transported to the film from an oxide target.[20, 21] In an attempt to obtain stoichiometric films over a wide range of process parameters and for the results to 2
40
be relevant for oxide ablation, we chose to use an IrO2 target instead of Ir metal to grow IrO2 films by PLD. Unlike bulk iridate synthesis, which is often done in an enclosed quartz tube, thin film growth is a non-equilibrium process where the evaporation of volatile species and resputtering from the film surface by the ablation plume cannot be
45
prevented. A particularly severe problem for iridate growth is the formation of a metastable IrO3 phase, which is volatile even at low temperatures and therefore severely restricts the useful parameter space where a stoichiometric IrO2 phase can form. Although the Ir oxide volatility is not unique to IrO2 and applies to pyrochlore iridate thin film growth as well, the loss of Ir from the film surface
50
appears to be dependent on the particular crystal structure and the surface layer composition. For example, the growth of perovskite iridates such as SrIrO3 and related Ruddlesden-Popper phases appears to be less affected by Ir loss. We show here that it is possible to grow IrO2 films that are stoichiometric and possess an atomically well-defined surface morphology. The grain structure
55
of the films is associated with strain relaxation, which occurs by the nucleation of misfit dislocation boundaries. Some of the misfit dislocations lead to extended defects that can be observed on the film surface as distinct grain boundaries and affect the low-temperature transport characteristics of the films.
2. Experiments 60
The IrO2 films were grown by pulsed laser deposition[22] on TiO2 (110) single crystal substrates that were annealed in air for one hour at 900◦ C before film growth to form a regular step-and-terrace surface.[23] The ablation target was a nominally stoichiometric IrO2 pellet, which was ablatied with a KrF (Thin Film Star, λ = 248 nm) laser operating at 5 Hz. Although an oxide target
65
was used instead of Ir metal, the deposition rate was relatively low. The film growth experiments were therefore done at a laser fluence of 0.9 J/cm2 , which was just below the damage threshold of the target surface. This laser fluence yielded a growth rate of 70 pulses per unit cell (3.18 ˚ A) along the (110) growth
3
100 106
(b)
103 100 106
IrO2(330)
IrO2(220)
IrO2(110) Ir(200)
10
(a)
3
Ir(111)
Intensity (arb. units)
106
(c)
103 100 20
40
60 80 100 120 2 (deg)
Figure 1: XRD θ/2θ scans of typical Ir oxide films grown on TiO2 (110) at (a) 500◦ C and 100 mTorr, (b) 900◦ C and 100 mTorr, and (c) 500◦ C and 500 mTorr. Black filled circles mark the TiO2 (110) substrate peak positions.
direction. Visual inspection of the target surface showed that the normally black 70
IrO2 surface became metallic gray after the growth of a single sample, indicating that the target surface tends to be reduced to Ir metal, regardless of the ambient oxygen pressure. The target surface was therefore re-polished after each film growth experiment. The film surface morphology was measured by atomic force microscopy (AFM). Structural analysis was done by symmetric X-ray diffraction
75
(XRD), Φ-scan measurements, and reciprocal space mapping. Resistivity of the films was measured in a four-point configuration with Ag electrodes grown on the film surface by thermal evaporation.
3. Results and discussions The oxidation rate of an IrOx film depends on the growth temperature and 80
the ambient oxygen pressure. The structure and composition of films was therefore mapped over a range of growth temperatures, from 500◦ C to 900◦ C and oxygen pressures, from 4 mTorr to 500 mTorr. Fig. 1 shows a comparison of 4
several characteristic Cu Kα θ/2θ XRD patterns of IrOx films deposited at different growth conditions. Clear IrO2 diffraction peaks were obtained at 500◦ C 85
and 100 mTorr. All film peaks were sharp and corresponded to a (110)-oriented film with no other observable crystal orientations. The full width at half maximum of the rocking curve of the (220) reflection was 1.2◦ . The film peak position of IrO2 (220) was 57.88◦ , which is close to the expected bulk position of 57.94◦ (d = 1.59 ˚ A), which means that the films grown at optimal conditions were
90
fully relaxed on the TiO2 (110) substrate. As illustrated by the XRD pattern in Fig. 1(b), no IrO2 diffraction peaks were observed in a film grown at 900◦ C and 100 mTorr, with Ir metal peaks appearing instead. This indicates that the deposited IrO2 was reduced to Ir metal at high temperature and a polycrystalline mirror-like Ir metal coating was obtained instead of an epitaxial oxide
95
film. Besides the maximum growth temperature limit, the IrO2 growth process window also has oxygen pressure limits. A film grown at 500◦ C and 500 mTorr showed only substrate peaks with no observable film growth. It is known that Ir can form a volatile IrO3 phase for which no stable solid phase exists, even at room temperature.[15] If the Ir or IrOx precursors on the substrate surface are
100
rapidly oxidized to IrO3 before crystal formation can occur, the re-evaporation rate from the substrate exceeds the deposition rate and no film grows on the substrate. The effects of growth temperature and background oxygen pressure are summarized in a phase stability diagram for IrO2 films in Fig. 2. Ir metal films with
105
a metallic mirror surface were obtained at low oxygen pressures and high growth temperatures. A mixed Ir/IrO2 region appeared in the low-temperature end of the Ir metal region. Increasing the growth temperature at a fixed pressure of 100 mTorr yielded films that did not exhibit clear IrO2 peaks in the XRD patterns, although a broadened base was observed around the substrate diffraction
110
peaks. The volatile IrO3 phase may form and evaporate at high oxygen pressures. No film growth was observed at ambient oxygen pressures of 200 mTorr and above. AFM images of samples grown at 500 mTorr and 500◦ C or 700◦ C still 5
O2 Pressure (mTorr)
4 2 100 4 2 10
No films
IrO2 Ir metal
4 400 500 600 700 800 900 Temperature (°C)
Figure 2: A phase stability diagram for IrO2 films. Each filled circle represents a sample. The region marked with a yellow dashed line corresponds to the optimal growth conditions. The best films in terms of IrO2 crystallinity were obtained at the point marked with an arrow.
showed the original step-and-terrace surface of the TiO2 substrate with no de115
tectable IrO2 film formation. The surface morphology of the IrO2 films was analyzed by AFM. Even under optimal growth conditions of 100 mTorr and 500◦ C, multigrain surface morphologies were obtained, as shown in Fig. 3. The maximum height difference between grains was 25 nm for an average film thickness of 60 nm. Although
120
large grains were present in the films, the top surface of each individual grain appeared to be very flat, which is consistent with IrO2 films grown on TiO2 (110) maintaining the tetragonal rutile structure. The grains appeared to be separated by deep trenches, which suggests that the film growth proceeded in a columnar fashion, where each grain is essentially a single crystal.
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Strain relaxation is the most common reason for the appearance of lattice defects and grain boundaries in heteroepitaxial thin films. If strain relaxation occurs gradually as the film thickness increases beyond a critical value, the surface morphology should depend on the film thickness. However, the AFM images in Fig. 3 indicate that the grain size did not vary significantly between
130
1 nm and 60 nm-thick films, which suggests that the strain relaxation occurs
6
(a)
(b)
B
A
Height (nm)
(c)
(d)
20
B
A
15 10 5
0
50
100
150 200 Distance (nm)
250
300
Figure 3: Surface morphologies of IrO2 films with different film thicknesses: (a) 60 nm, (b) 10 nm, (c) 1 nm, and (d) 0.25 nm. A cross-section height profile is shown for (a). Image size is 1µm×1µm.
7
q110 (1/Å)
0.64
IrO2(310)
0.62 TiO2(310)
0.60 0.30
0.32
0.34
q110 (1/Å) Figure 4: Reciprocal space map of a 60 nm thick IrO2 (110) film in the vicinity of the (310) diffractin peak. The IrO2 film peak is strongly broadened in the [110] direction. The red cross marks the bulk IrO2 (310) reflection.
during the initial growth stage at the substrate interface. Nucleation seeds can be seen in Fig. 3(d) for a nominal film thickness of 0.25 nm, i.e., the height of the nucleated islands is no larger than two atomic layers. The areal density of the seed nuclei is similar to the grain density in thicker films, which suggests 135
that the grain boundary defects form during the initial IrO2 nucleation on the substrate surface. This result was unexpected, considering the epitaxial nature of the film, which would appear to rule out the formation of low-angle grain boundaries or extended stacking-fault type defects in thin strained layers. The strain state of the films was further analyzed by reciprocal space map-
140
ping and pole figure scans. Reciprocal space maps were measured near the rutile (310) and (332) reflections for a 60-nm-thick film grown at 100 mTorr and 500◦ C. The bulk lattice parameters for IrO2 [24] are a = 4.4983 ˚ A, c = 3.1544 ˚ A and for TiO2 ,[25] a = 4.5941 ˚ A, c = 2.9589 ˚ A. The in-plane lattice mismatch for a (110) interface is thus +6.2% in the (001) direction and −2.1% in the (110)
145
direction. A reciprocal space map for the (310) reflection in Fig. 4 shows the
8
105 Intensity (arb. units)
(a)
10
4
49°
131°
52°
128°
TiO2
103
(c) IrO2 (110)
52°
128°
IrO2
102
Ir O
3.18 Å
101 100
[110] 0
(b)
60
120
180 240 (deg)
300
360
(d) TiO2 (110)
101
Plane B 49°
131°
IrO2(330)
TiO2(330)
IrO2(220)
TiO2(220)
IrO2(110)
TiO2(110)
Intensity (arb. units)
103
102
Plane A
Ti O
3.25 Å 10
0
20
40 2
60 80 χ(deg)
100 Plane A
Plane B
Figure 5: (a) XRD Φ-scans for a 60-nm-thick IrO2 film and the TiO2 substrate (101) peaks. (b) In-plane grazing incidence diffraction pattern of a 2-nm-thick (110)-oriented IrO2 film. The peak spacings in (a) are determined by the angle between planes A and B, shown in (c) for IrO2 (110) and in (d) for TiO2 (110).
strain relaxation behavior in the (110) direction. The center-of-mass position of the film peak coincides with the expected bulk value, showing that the film is fully relaxed. The same result was obtained for the (001) direction map. Φ-scans of the same film sample were used to verify that there are no grains 150
in the film with unexpected in-plane orientations. The Φ-scans of the substrate and film (101) peaks are compared in Fig. 5(a). The film peak positions in the reciprocal space map (Fig. 4) matched the bulk values, showing that the films were fully relaxed. The scans in Fig. 5(a) show that the peak angle distances match exactly the expected angles between the A (112) and B (112) crystal
155
planes for bulk IrO2 and TiO2 structures [5, 26] as illustrated in Figs. 5(c) and 5(d). We can therefore conclude from the Φ-scans that the films were indeed epitaxial with the correct in-plane orientation, but fully relaxed on the TiO2
9
[001]
[110]
(a)
50 nm IrO2 TiO2 5 nm
(b)
[110] [110]
IrO2
[001]
Figure 6: (a) Wide-area HRTEM view of a nominally 15 nm thick IrO2 film, (b) a zoomed STEM view of the topmost layer of the IrO2 film.
substrates. Since the thickness of the film used in the reciprocal space and Φscan measurements was 60 nm, it was not possible to determine if a strained 160
layer exists at the substrate interface or not. Since the thickness-dependent film morphology analysis suggested that the grain structure forms at the initial film growth stage, the strain state of ultrathin layers was further studied by grazing-incidence in-plane x-ray diffraction (grazing angele = 0.15◦ ). The inplane diffraction pattern for a 2 nm IrO2 film (Rigaku SmartLab) is shown in
165
Fig. 5(b). Although the nominal film thickness was only 6 unit cells, clear film diffraction peaks were obtained. The measured in-plane lattice parameter in the (110) direction was 3.183 ˚ A, nearly equal to the bulk value of 3.181 ˚ A. We can thus conclude that the grain formation and strain relaxation occurs during the growth of the first few unit cells of an IrO2 film.
170
High-resolution cross-section transmission electron microscopy (HRTEM) and scanning electron microscopy (STEM) were used for closer analysis of the grain formation in IrO2 films. Fig. 6(a) shows a wide-scale HRTEM image of a nominally 15-nm-thick film. A clear grain structure is visible in the image, with
10
neighboring grains having height differences of up to 5 nm. As was already seen 175
in the AFM images, despite the large height differences, all grains appear to have atomically flat surfaces. This can be seen better in a narrow-scale STEM image of the surface region of a single grain in Fig. 6(b). This result indicates that despite the volatility problems associated with iridates, atomically flat and well ordered surfaces can be grown in a physical vapor deposition process.
180
The XRD analysis showed that the grain boundaries form within the first few unit cell layers of the IrO2 film. A high-resolution STEM image in Fig. 7(a) shows the details of the strain relaxation at the film - substrate interface along the [110] direction. The first atomic row of the IrO2 film can be easily identified due to the increased brightness in the high-angle annular dark-field (HAADF)
185
STEM image. An added crystal row can be seen in the middle of the imaging area. Similar-sized islands can be seen in the AFM image in Fig. 3d. Vertical misfit dislocations nucleate at either end of the island. An additional dislocation core can be seen in the upper left part of the image. The position of the added rows and columns can be more easily seen in the schematic lattice diagram in
190
Fig. 7(b). It is clear that an adatom [110] direction column and a [110] row always nucleate at the same point in the lattice. The reason for this particular dislocation structure can be understood by considering the rutile structure, illustrated in Figs. 7(c,d). If a single vertical adatom column is added in a crystal along the [110] direction, an adatom (shown in red in Fig. 7(c)) cannot be added
195
into the rutile structure. However, if the number of lattice rows differs in the [110] direction by one on either side of the dislocation, a continuous lattice can form, as shown in Fig. 7(d). The -2.1% in-plane misfit relaxation along the [110] of the IrO2 film on the TiO2 substrate surface thus leads to large out-of-plane disorder, as was seen by the large spread of the film diffraction peak in the
200
reciprocal space map in Fig. 4. The grain boundary formed by a stacking-faultlike structure along a misfit column can extend over a long distance in the film, as shown in Fig. 7(e), where a grain boundary can be traced over a distance of about 15 nm. Although most of the grain boundaries relax within the first 10 nm or less, some extended boundaries remain, giving rise to the distinct grain 11
(a)
(b)
10 nm [110]
10 nm [110]
[110]
[110] [001]
[001]
IrO2 TiO2
(c)
(d) Ir
[110]
[110]
[110]
[001]
[110] [001]
(e)
[110]
IrO2
[110] [001]
5 nm
TiO2
Figure 7: (a) STEM image of the dislocation structure of an IrO2 film. Three dislocation cores are marked in the image. (b) A schematic diagram of the same STEM image area, showing that each [110] direction dislocation column is accompanied by an added row along the [110] direction. (c) A dislocation column model. Ir atoms are shown in yellow, the position of the first atom of an adatom column is shown in red. The red atom does not fit the rutile structure. (d) The misfit column can be accommodated if an additional lattice row is added on one side of the dislocation. (e) An extended grain boundary in the IrO2 film. A vertical half-atomic-row shift can be seen between the IrO2 lattice on the left and right sides of the dislocation.
12
Ω
10
(b)
2 -4
8 6
800°C 6 hrs
10
4
Ω
(a)
2
0
8 6 4
15 nm 60 nm
2
400°C 6 hrs
10-5
8 6
2 -4
100
200 T (K)
10-5
300
8 6
0
30 nm 100 200 T (K)
300
Figure 8: Temperature dependence of resistivity of IrO2 films. (a) The effect of air annealing on the film resistivity, comparing an as-deposited film grown at 500◦ C and 100 mTorr (black) with those post-annealed in air at 400◦ C (red) or 800◦ C (blue) for 6 hours. (b) The effect of film thickness on low-temperature resistivity. The film thickness was 15 nm (blue), 30 nm (red), and 60 nm (black).
205
structure seen in the AFM images (Fig. 3) The low-temperature resistivity behavior of the films is shown in Fig. 8. While the room-temperature resistivities of the films are close to the expected bulk single-crystal values,[27] the saturation at low temperatures does not follow the bulk behavior. The residual resistance ratio (RRR) of an as-deposited film
210
grown at 500◦ C and 100 mTorr, shown in black in Fig. 8(a), was about 3.2, which is similar to other published reports on IrO2 films grown by PLD [19] and by molecular beam epitaxy.[10] Several possibilities can be considered as the cause for the unexpectedly high resistivity at low temperatures. Since the growth window was limited in terms
215
of temperature and oxygen pressure, oxygen nonstoichiometry is one possible reason. As-deposited films were therefore post-annealed in air at either 400◦ C or 800◦ C for 6 hours and the resistivity measurement was repeated, as shown in Fig. 8(a). No significant differences were found in the temperature dependence, although a resistivity increase was seen after the 800◦ C annealing treatment.
220
Ir volatility may have caused this increase due to the formation of IrO3 at the surface. Another possibility that was considered is the variation of oxygen stoichiometry as a function of film thickness, i.e., the film growth time. A series
13
of measurements were therefore done for films with thicknesses varying from 15 to 60 nm, but no significant differences were seen either, as shown in Fig. 8(b). 225
The resistivity analysis appears to show that the increased low-temperature resistivity is caused purely by structural effects, i.e., grain-boundary scattering.
4. Conclusion We have determined the growth window for the formation of IrO2 thin films by PLD on TiO2 substrates. Optimal films were obtained at 500◦ C at an oxygen 230
pressure of 100 mTorr. The resistivity behavior of the films is dominated by the presence of stacking-fault grain boundaries in the film. The formation of the grain structure is driven by strain relaxation. This strain relaxation leads to the formation of columnar grains with relatively high grain boundary resistance. However, despite the grain structure, the individual grain crystal quality was
235
excellent close to the surface and the grains had atomically flat surfaces. The main motivation for this study was to fabricate iridate films that would suitable for angle-resolved photoelectron spectroscopy studies, and for this particular purpose, the well-defined grain surface structure is more important than the effect of the grain boundaries on the low-temperature resistivity.
240
Acknowledgments The authors thank Shintaro Kobayashi of Rigaku Co. for grazing-incidence in-plane XRD measurements. This work was supported by a Grant-in-Aid for Scientific Research (Grant Nos. 25706022, and 26105002) from the Japan Society for the Promotion of Science. X. H. was supported by the Japan Society
245
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PII: DOI: Reference:
S0022-0248(16)30974-5 http://dx.doi.org/10.1016/j.jcrysgro.2016.12.104 CRYS23963
To appear in: Journal of Crystal Growth Cite this article as: Xiuyi Hou, Ryota Takahashi, Takahisa Yamamoto and Mikk Lippmaa, Microstructure analysis of IrO2 thin films, Journal of Crystal Growth, http://dx.doi.org/10.1016/j.jcrysgro.2016.12.104 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Microstructure analysis of IrO2 thin films Xiuyi Houa,∗, Ryota Takahashia,b , Takahisa Yamamotoc , Mikk Lippmaaa a Institute
for Solid State Physics, University of Tokyo, Chiba 277-8581, Japan b JST, PRESTO, Saitama 332-0012, Japan c Graduate School of Engineering, Nagoya University, Aichi 464-8603, Japan
Abstract We have grown IrO2 thin films on TiO2 (110) substrates to determine the pulsed laser deposition growth window for iridates. Relaxed IrO2 films were obtained at a growth temperature of 500◦ C and background oxygen pressure of 100 mTorr; otherwise, either pure Ir metal films or evaporative Ir loss were observed. Although x-ray Φ-scan measurement indicated that the films were epitaxial, a distinct grain structure was seen by atomic force microscopy and transmission electron microscopy. The grain boundaries were found to limit the conductivity of films at low temperature. It appeared that strain relaxation leads to stacking faults at grain boundaries. Keywords: Laser epitaxy, Defects, Oxides, X-ray diffraction
1. Introduction Recent interest in Ir oxides has been stoked by the prediction of topological states appearing in pyrochlore-type iridate films strained along the [111] crystal direction.[1, 2] Among the iridate pyrochlores, Pr2 Ir2 O7 has been shown by 5
angle-resolved photoemission spectroscopy (ARPES) to possess a single-point Fermi node at the Γ-point[3], making it a good candidate for epitaxial strain experiments if stoichiometric and atomically flat surfaces can be prepared. However, a common problem for the growth of iridate thin films by physical vapor
∗ Corresponding
author URL: [email protected] (Xiuyi Hou)
Preprint submitted to Journal of Crystal Growth
December 24, 2016
deposition techniques such as pulsed laser deposition (PLD), is the volatility of 10
nonstoichiometric iridium oxides. This can lead to severe stoichiometry deviations in thin films and it is the main reason why pyrochlore iridate thin films have so far been fabricated mostly by solid-phase epitaxy, where the film precursor material is deposited at room temperature and crytallization occurs during post-annealing.[4] While this process produces thin films with excellent struc-
15
tural properties, it unavoidably leads to a relaxed lattice with a compositionally degraded surface, making the films unsuitable for ARPES studies. We therefore explore in this work the thin film growth process of the simplest iridate, IrO2 , to determine the PLD process parameter window where the Ir4+ state can be stabilized. Besides the process optimization, we recognize that IrO2 is
20
an interesting material for a number of reasons, and we analyze the microstructure, transport behavior, and the strain relaxation mechanism in IrO2 (110) films grown on TiO2 (110) substrates. Iridium dioxide belongs to a small family of oxides that are good metals, together with RuO2 , ReO2 , OsO2 , MoO2 , and WO2 .[5] The Fermi level in IrO2
25
is located close to the top of the Ir 5d t2g band, leading to nonmagnetic metallic behavior.[6, 7, 8] IrO2 thin films have been used as a spin-to-current detector in spin valves [9] and may be useful in devices with magnetically switchable majority carrier type.[10] Other well-known applications of bulk or polycrystalline IrO2 are biocompatible electrodes,[11] oxygen barriers in semiconductor
30
devices,[12] electrodes in ferroelectric devices,[13] and electrocatalytic oxygen evolution eletrodes for water splitting.[14] The main difficulty in growing IrO2 films is the difficulty of oxidizing Ir metal and the volatility of nonstoichiometric Ir oxides.[15, 16] IrO2 films have been grown by various techniques, including molecular beam epitaxy (MBE),[10]
35
sputtering,[17] and PLD.[18, 19] Although it is common to use a metallic Ir target for film deposition and allow the ambient oxygen to oxidize the deposited metal at the film surface, it is known that in PLD growth, oxygen is also transported to the film from an oxide target.[20, 21] In an attempt to obtain stoichiometric films over a wide range of process parameters and for the results to 2
40
be relevant for oxide ablation, we chose to use an IrO2 target instead of Ir metal to grow IrO2 films by PLD. Unlike bulk iridate synthesis, which is often done in an enclosed quartz tube, thin film growth is a non-equilibrium process where the evaporation of volatile species and resputtering from the film surface by the ablation plume cannot be
45
prevented. A particularly severe problem for iridate growth is the formation of a metastable IrO3 phase, which is volatile even at low temperatures and therefore severely restricts the useful parameter space where a stoichiometric IrO2 phase can form. Although the Ir oxide volatility is not unique to IrO2 and applies to pyrochlore iridate thin film growth as well, the loss of Ir from the film surface
50
appears to be dependent on the particular crystal structure and the surface layer composition. For example, the growth of perovskite iridates such as SrIrO3 and related Ruddlesden-Popper phases appears to be less affected by Ir loss. We show here that it is possible to grow IrO2 films that are stoichiometric and possess an atomically well-defined surface morphology. The grain structure
55
of the films is associated with strain relaxation, which occurs by the nucleation of misfit dislocation boundaries. Some of the misfit dislocations lead to extended defects that can be observed on the film surface as distinct grain boundaries and affect the low-temperature transport characteristics of the films.
2. Experiments 60
The IrO2 films were grown by pulsed laser deposition[22] on TiO2 (110) single crystal substrates that were annealed in air for one hour at 900◦ C before film growth to form a regular step-and-terrace surface.[23] The ablation target was a nominally stoichiometric IrO2 pellet, which was ablatied with a KrF (Thin Film Star, λ = 248 nm) laser operating at 5 Hz. Although an oxide target
65
was used instead of Ir metal, the deposition rate was relatively low. The film growth experiments were therefore done at a laser fluence of 0.9 J/cm2 , which was just below the damage threshold of the target surface. This laser fluence yielded a growth rate of 70 pulses per unit cell (3.18 ˚ A) along the (110) growth
3
100 106
(b)
103 100 106
IrO2(330)
IrO2(220)
IrO2(110) Ir(200)
10
(a)
3
Ir(111)
Intensity (arb. units)
106
(c)
103 100 20
40
60 80 100 120 2 (deg)
Figure 1: XRD θ/2θ scans of typical Ir oxide films grown on TiO2 (110) at (a) 500◦ C and 100 mTorr, (b) 900◦ C and 100 mTorr, and (c) 500◦ C and 500 mTorr. Black filled circles mark the TiO2 (110) substrate peak positions.
direction. Visual inspection of the target surface showed that the normally black 70
IrO2 surface became metallic gray after the growth of a single sample, indicating that the target surface tends to be reduced to Ir metal, regardless of the ambient oxygen pressure. The target surface was therefore re-polished after each film growth experiment. The film surface morphology was measured by atomic force microscopy (AFM). Structural analysis was done by symmetric X-ray diffraction
75
(XRD), Φ-scan measurements, and reciprocal space mapping. Resistivity of the films was measured in a four-point configuration with Ag electrodes grown on the film surface by thermal evaporation.
3. Results and discussions The oxidation rate of an IrOx film depends on the growth temperature and 80
the ambient oxygen pressure. The structure and composition of films was therefore mapped over a range of growth temperatures, from 500◦ C to 900◦ C and oxygen pressures, from 4 mTorr to 500 mTorr. Fig. 1 shows a comparison of 4
several characteristic Cu Kα θ/2θ XRD patterns of IrOx films deposited at different growth conditions. Clear IrO2 diffraction peaks were obtained at 500◦ C 85
and 100 mTorr. All film peaks were sharp and corresponded to a (110)-oriented film with no other observable crystal orientations. The full width at half maximum of the rocking curve of the (220) reflection was 1.2◦ . The film peak position of IrO2 (220) was 57.88◦ , which is close to the expected bulk position of 57.94◦ (d = 1.59 ˚ A), which means that the films grown at optimal conditions were
90
fully relaxed on the TiO2 (110) substrate. As illustrated by the XRD pattern in Fig. 1(b), no IrO2 diffraction peaks were observed in a film grown at 900◦ C and 100 mTorr, with Ir metal peaks appearing instead. This indicates that the deposited IrO2 was reduced to Ir metal at high temperature and a polycrystalline mirror-like Ir metal coating was obtained instead of an epitaxial oxide
95
film. Besides the maximum growth temperature limit, the IrO2 growth process window also has oxygen pressure limits. A film grown at 500◦ C and 500 mTorr showed only substrate peaks with no observable film growth. It is known that Ir can form a volatile IrO3 phase for which no stable solid phase exists, even at room temperature.[15] If the Ir or IrOx precursors on the substrate surface are
100
rapidly oxidized to IrO3 before crystal formation can occur, the re-evaporation rate from the substrate exceeds the deposition rate and no film grows on the substrate. The effects of growth temperature and background oxygen pressure are summarized in a phase stability diagram for IrO2 films in Fig. 2. Ir metal films with
105
a metallic mirror surface were obtained at low oxygen pressures and high growth temperatures. A mixed Ir/IrO2 region appeared in the low-temperature end of the Ir metal region. Increasing the growth temperature at a fixed pressure of 100 mTorr yielded films that did not exhibit clear IrO2 peaks in the XRD patterns, although a broadened base was observed around the substrate diffraction
110
peaks. The volatile IrO3 phase may form and evaporate at high oxygen pressures. No film growth was observed at ambient oxygen pressures of 200 mTorr and above. AFM images of samples grown at 500 mTorr and 500◦ C or 700◦ C still 5
O2 Pressure (mTorr)
4 2 100 4 2 10
No films
IrO2 Ir metal
4 400 500 600 700 800 900 Temperature (°C)
Figure 2: A phase stability diagram for IrO2 films. Each filled circle represents a sample. The region marked with a yellow dashed line corresponds to the optimal growth conditions. The best films in terms of IrO2 crystallinity were obtained at the point marked with an arrow.
showed the original step-and-terrace surface of the TiO2 substrate with no de115
tectable IrO2 film formation. The surface morphology of the IrO2 films was analyzed by AFM. Even under optimal growth conditions of 100 mTorr and 500◦ C, multigrain surface morphologies were obtained, as shown in Fig. 3. The maximum height difference between grains was 25 nm for an average film thickness of 60 nm. Although
120
large grains were present in the films, the top surface of each individual grain appeared to be very flat, which is consistent with IrO2 films grown on TiO2 (110) maintaining the tetragonal rutile structure. The grains appeared to be separated by deep trenches, which suggests that the film growth proceeded in a columnar fashion, where each grain is essentially a single crystal.
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Strain relaxation is the most common reason for the appearance of lattice defects and grain boundaries in heteroepitaxial thin films. If strain relaxation occurs gradually as the film thickness increases beyond a critical value, the surface morphology should depend on the film thickness. However, the AFM images in Fig. 3 indicate that the grain size did not vary significantly between
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1 nm and 60 nm-thick films, which suggests that the strain relaxation occurs
6
(a)
(b)
B
A
Height (nm)
(c)
(d)
20
B
A
15 10 5
0
50
100
150 200 Distance (nm)
250
300
Figure 3: Surface morphologies of IrO2 films with different film thicknesses: (a) 60 nm, (b) 10 nm, (c) 1 nm, and (d) 0.25 nm. A cross-section height profile is shown for (a). Image size is 1µm×1µm.
7
q110 (1/Å)
0.64
IrO2(310)
0.62 TiO2(310)
0.60 0.30
0.32
0.34
q110 (1/Å) Figure 4: Reciprocal space map of a 60 nm thick IrO2 (110) film in the vicinity of the (310) diffractin peak. The IrO2 film peak is strongly broadened in the [110] direction. The red cross marks the bulk IrO2 (310) reflection.
during the initial growth stage at the substrate interface. Nucleation seeds can be seen in Fig. 3(d) for a nominal film thickness of 0.25 nm, i.e., the height of the nucleated islands is no larger than two atomic layers. The areal density of the seed nuclei is similar to the grain density in thicker films, which suggests 135
that the grain boundary defects form during the initial IrO2 nucleation on the substrate surface. This result was unexpected, considering the epitaxial nature of the film, which would appear to rule out the formation of low-angle grain boundaries or extended stacking-fault type defects in thin strained layers. The strain state of the films was further analyzed by reciprocal space map-
140
ping and pole figure scans. Reciprocal space maps were measured near the rutile (310) and (332) reflections for a 60-nm-thick film grown at 100 mTorr and 500◦ C. The bulk lattice parameters for IrO2 [24] are a = 4.4983 ˚ A, c = 3.1544 ˚ A and for TiO2 ,[25] a = 4.5941 ˚ A, c = 2.9589 ˚ A. The in-plane lattice mismatch for a (110) interface is thus +6.2% in the (001) direction and −2.1% in the (110)
145
direction. A reciprocal space map for the (310) reflection in Fig. 4 shows the
8
105 Intensity (arb. units)
(a)
10
4
49°
131°
52°
128°
TiO2
103
(c) IrO2 (110)
52°
128°
IrO2
102
Ir O
3.18 Å
101 100
[110] 0
(b)
60
120
180 240 (deg)
300
360
(d) TiO2 (110)
101
Plane B 49°
131°
IrO2(330)
TiO2(330)
IrO2(220)
TiO2(220)
IrO2(110)
TiO2(110)
Intensity (arb. units)
103
102
Plane A
Ti O
3.25 Å 10
0
20
40 2
60 80 χ(deg)
100 Plane A
Plane B
Figure 5: (a) XRD Φ-scans for a 60-nm-thick IrO2 film and the TiO2 substrate (101) peaks. (b) In-plane grazing incidence diffraction pattern of a 2-nm-thick (110)-oriented IrO2 film. The peak spacings in (a) are determined by the angle between planes A and B, shown in (c) for IrO2 (110) and in (d) for TiO2 (110).
strain relaxation behavior in the (110) direction. The center-of-mass position of the film peak coincides with the expected bulk value, showing that the film is fully relaxed. The same result was obtained for the (001) direction map. Φ-scans of the same film sample were used to verify that there are no grains 150
in the film with unexpected in-plane orientations. The Φ-scans of the substrate and film (101) peaks are compared in Fig. 5(a). The film peak positions in the reciprocal space map (Fig. 4) matched the bulk values, showing that the films were fully relaxed. The scans in Fig. 5(a) show that the peak angle distances match exactly the expected angles between the A (112) and B (112) crystal
155
planes for bulk IrO2 and TiO2 structures [5, 26] as illustrated in Figs. 5(c) and 5(d). We can therefore conclude from the Φ-scans that the films were indeed epitaxial with the correct in-plane orientation, but fully relaxed on the TiO2
9
[001]
[110]
(a)
50 nm IrO2 TiO2 5 nm
(b)
[110] [110]
IrO2
[001]
Figure 6: (a) Wide-area HRTEM view of a nominally 15 nm thick IrO2 film, (b) a zoomed STEM view of the topmost layer of the IrO2 film.
substrates. Since the thickness of the film used in the reciprocal space and Φscan measurements was 60 nm, it was not possible to determine if a strained 160
layer exists at the substrate interface or not. Since the thickness-dependent film morphology analysis suggested that the grain structure forms at the initial film growth stage, the strain state of ultrathin layers was further studied by grazing-incidence in-plane x-ray diffraction (grazing angele = 0.15◦ ). The inplane diffraction pattern for a 2 nm IrO2 film (Rigaku SmartLab) is shown in
165
Fig. 5(b). Although the nominal film thickness was only 6 unit cells, clear film diffraction peaks were obtained. The measured in-plane lattice parameter in the (110) direction was 3.183 ˚ A, nearly equal to the bulk value of 3.181 ˚ A. We can thus conclude that the grain formation and strain relaxation occurs during the growth of the first few unit cells of an IrO2 film.
170
High-resolution cross-section transmission electron microscopy (HRTEM) and scanning electron microscopy (STEM) were used for closer analysis of the grain formation in IrO2 films. Fig. 6(a) shows a wide-scale HRTEM image of a nominally 15-nm-thick film. A clear grain structure is visible in the image, with
10
neighboring grains having height differences of up to 5 nm. As was already seen 175
in the AFM images, despite the large height differences, all grains appear to have atomically flat surfaces. This can be seen better in a narrow-scale STEM image of the surface region of a single grain in Fig. 6(b). This result indicates that despite the volatility problems associated with iridates, atomically flat and well ordered surfaces can be grown in a physical vapor deposition process.
180
The XRD analysis showed that the grain boundaries form within the first few unit cell layers of the IrO2 film. A high-resolution STEM image in Fig. 7(a) shows the details of the strain relaxation at the film - substrate interface along the [110] direction. The first atomic row of the IrO2 film can be easily identified due to the increased brightness in the high-angle annular dark-field (HAADF)
185
STEM image. An added crystal row can be seen in the middle of the imaging area. Similar-sized islands can be seen in the AFM image in Fig. 3d. Vertical misfit dislocations nucleate at either end of the island. An additional dislocation core can be seen in the upper left part of the image. The position of the added rows and columns can be more easily seen in the schematic lattice diagram in
190
Fig. 7(b). It is clear that an adatom [110] direction column and a [110] row always nucleate at the same point in the lattice. The reason for this particular dislocation structure can be understood by considering the rutile structure, illustrated in Figs. 7(c,d). If a single vertical adatom column is added in a crystal along the [110] direction, an adatom (shown in red in Fig. 7(c)) cannot be added
195
into the rutile structure. However, if the number of lattice rows differs in the [110] direction by one on either side of the dislocation, a continuous lattice can form, as shown in Fig. 7(d). The -2.1% in-plane misfit relaxation along the [110] of the IrO2 film on the TiO2 substrate surface thus leads to large out-of-plane disorder, as was seen by the large spread of the film diffraction peak in the
200
reciprocal space map in Fig. 4. The grain boundary formed by a stacking-faultlike structure along a misfit column can extend over a long distance in the film, as shown in Fig. 7(e), where a grain boundary can be traced over a distance of about 15 nm. Although most of the grain boundaries relax within the first 10 nm or less, some extended boundaries remain, giving rise to the distinct grain 11
(a)
(b)
10 nm [110]
10 nm [110]
[110]
[110] [001]
[001]
IrO2 TiO2
(c)
(d) Ir
[110]
[110]
[110]
[001]
[110] [001]
(e)
[110]
IrO2
[110] [001]
5 nm
TiO2
Figure 7: (a) STEM image of the dislocation structure of an IrO2 film. Three dislocation cores are marked in the image. (b) A schematic diagram of the same STEM image area, showing that each [110] direction dislocation column is accompanied by an added row along the [110] direction. (c) A dislocation column model. Ir atoms are shown in yellow, the position of the first atom of an adatom column is shown in red. The red atom does not fit the rutile structure. (d) The misfit column can be accommodated if an additional lattice row is added on one side of the dislocation. (e) An extended grain boundary in the IrO2 film. A vertical half-atomic-row shift can be seen between the IrO2 lattice on the left and right sides of the dislocation.
12
Ω
10
(b)
2 -4
8 6
800°C 6 hrs
10
4
Ω
(a)
2
0
8 6 4
15 nm 60 nm
2
400°C 6 hrs
10-5
8 6
2 -4
100
200 T (K)
10-5
300
8 6
0
30 nm 100 200 T (K)
300
Figure 8: Temperature dependence of resistivity of IrO2 films. (a) The effect of air annealing on the film resistivity, comparing an as-deposited film grown at 500◦ C and 100 mTorr (black) with those post-annealed in air at 400◦ C (red) or 800◦ C (blue) for 6 hours. (b) The effect of film thickness on low-temperature resistivity. The film thickness was 15 nm (blue), 30 nm (red), and 60 nm (black).
205
structure seen in the AFM images (Fig. 3) The low-temperature resistivity behavior of the films is shown in Fig. 8. While the room-temperature resistivities of the films are close to the expected bulk single-crystal values,[27] the saturation at low temperatures does not follow the bulk behavior. The residual resistance ratio (RRR) of an as-deposited film
210
grown at 500◦ C and 100 mTorr, shown in black in Fig. 8(a), was about 3.2, which is similar to other published reports on IrO2 films grown by PLD [19] and by molecular beam epitaxy.[10] Several possibilities can be considered as the cause for the unexpectedly high resistivity at low temperatures. Since the growth window was limited in terms
215
of temperature and oxygen pressure, oxygen nonstoichiometry is one possible reason. As-deposited films were therefore post-annealed in air at either 400◦ C or 800◦ C for 6 hours and the resistivity measurement was repeated, as shown in Fig. 8(a). No significant differences were found in the temperature dependence, although a resistivity increase was seen after the 800◦ C annealing treatment.
220
Ir volatility may have caused this increase due to the formation of IrO3 at the surface. Another possibility that was considered is the variation of oxygen stoichiometry as a function of film thickness, i.e., the film growth time. A series
13
of measurements were therefore done for films with thicknesses varying from 15 to 60 nm, but no significant differences were seen either, as shown in Fig. 8(b). 225
The resistivity analysis appears to show that the increased low-temperature resistivity is caused purely by structural effects, i.e., grain-boundary scattering.
4. Conclusion We have determined the growth window for the formation of IrO2 thin films by PLD on TiO2 substrates. Optimal films were obtained at 500◦ C at an oxygen 230
pressure of 100 mTorr. The resistivity behavior of the films is dominated by the presence of stacking-fault grain boundaries in the film. The formation of the grain structure is driven by strain relaxation. This strain relaxation leads to the formation of columnar grains with relatively high grain boundary resistance. However, despite the grain structure, the individual grain crystal quality was
235
excellent close to the surface and the grains had atomically flat surfaces. The main motivation for this study was to fabricate iridate films that would suitable for angle-resolved photoelectron spectroscopy studies, and for this particular purpose, the well-defined grain surface structure is more important than the effect of the grain boundaries on the low-temperature resistivity.
240
Acknowledgments The authors thank Shintaro Kobayashi of Rigaku Co. for grazing-incidence in-plane XRD measurements. This work was supported by a Grant-in-Aid for Scientific Research (Grant Nos. 25706022, and 26105002) from the Japan Society for the Promotion of Science. X. H. was supported by the Japan Society
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