Tunability of transport properties in semi-exfoliated Ag_xCoO_2−δ (0.4 < x < 1) crystals†

Marc Kamel ^a, Hatem M. Titi ^b, Mohamad Ataya ^b, Antranik Jonderian ^b, Kirk H. Bevan *^ac and Eric McCalla *^bc
aDivision of Materials Engineering, Faculty of Engineering, McGill University, Montreal, Canada. E-mail: kirk.bevan@mcgill.ca
^bDepartment of Chemistry, McGill University, Montreal, Canada. E-mail: eric.mccalla@mcgill.ca
^cCentre for the Physics of Materials, McGill University, Montreal, Canada

First published on 30th July 2024

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1. Introduction

In the on-going search for new high-performing semiconductors for the electronics industry, it has become increasingly important to find highly conductive transparent semiconductors.¹ Having both n- and p-type transparent semiconductors with high mobilities would enable important new applications for technologies such as solar cells² and transistors.³ In particular, concerning the design of transistors pertaining to transparent electronics, n and p-type doping is crucial to addressing and implementing lower power operation.^4,5 While high mobility n-type transparent semiconductors such as BaSnO_3−δ have shown significant progress,^6,7 finding high mobility p-type transparent semiconductors has proven more elusive. Recently, delafossite materials have received increasing attention in this regard (e.g. ref. 8 demonstrated high hole conductivity in transparent PdCoO₂ thin films). Some delafossites also showcase high in-plane conductivity (e.g. a low resistivity of 3 μΩ cm for PtCoO₂⁹). In addition, delafossites have shown good transparency features¹⁰ which could be used for various applications such as solar cells,¹¹ gas sensors¹² or in high temperature electronics.¹³ Delafossites are layered structures with an ABO₂ stoichiometry where A and B are cations allowing a broad chemical flexibility (e.g. A cations include Pd, Ag, Cu, Pt, while B cations include Ga, In, Co, Al).¹⁴ Other transparent Delafossites include CuAl_0.5Fe_0.5O₂,¹⁴ CuAlO₂,¹⁵ both of which are examples of crucially needed p-type transparent conductive oxides.

Another delafossite that received considerable attention in the past is AgCoO₂.^1,16–19 Delafossite AgCoO₂ has been synthesized in both polycrystalline²⁰ and crystal forms¹⁶ by performing ion-exchange on the precursor NaCoO₂. Delmas et al. demonstrated that the stacking of the delafossite can be varied by using various NaCoO₂ polymorphs.¹⁷ Of highest interest for the current work, two different AgCoO₂ polymorphs (D2 and D3) can be obtained in this manner and their structures as well as the structure of their respective precursors are shown in Fig. 1. This shows that the ion-exchange process where Na is replaced by Ag has very little impact on the overall stacking of the CoO₂ slabs. Previous work on this delafossite also clearly showed that the materials synthesized by ion-exchange are p-type semiconductors, which is attributed to Ag vacancies acting as p-dopants.^18,19 Although this is qualitatively logical, metal vacancies do in fact act as p-dopants while oxygen vacancies act as n-dopants, it does not make quantitative sense as the stoichiometry identified is Ag_0.75CoO₂ and the charge carrier concentrations are only on the order of 10¹⁹ cm⁻³,^18,19,21,22 which certainly do not require 25% metal vacancies to be created, in fact 25% vacancies would result in an extremely high carrier concentration. There are therefore important questions to answer related to how the materials stoichiometry leads to the observed transport properties and this will be addressed in detail herein. To date, there has been no study that explored variations in the stoichiometry of the material and, as far as we know, only the x ∼ 0.75 material has been reported. It is also important to note that, to our knowledge, DFT calculations have only been performed on the stoichiometric AgCoO₂²³ which is shown to be an intrinsic semiconductor, again implying a gap with respect to experiments where clear p-type transport is observed. These earlier studies calculations were done using the GGA functional which is well known to under-estimate band gaps. Hence, we include DFT calculations using the more accurate hybrid functional herein on a variety of compositions to better understand how the values of x and δ affect the band structure and thus the transport properties in Ag_xCoO_2−δ.

Another result of note from prior work on this delafossite is that single-crystals have been reported by ion-exchange.¹⁶ Ref. 16 shows SEM images of crystalline AgCoO₂ with particle sizes of about 50 μm and showing only top view of the samples. As will be discussed thoroughly herein, when viewed from the side these crystals show significant fracturing during ion-exchange. Importantly, we obtain comparable images from the top using the same synthesis route, but the side views of the crystals that we report here for the first time show that the crystals are now semi-detached flakes and as such they are no longer single-crystals. Thus, in this work, we focus on updating our understanding of AgCoO₂ delafossites made from single-crystal precursors. The powder samples are well-known to take the D3 structure shown in Fig. 1, while the single-crystal samples take the D2 structure²⁰ using the Delmas notation. The structural differences between D2 and D3 types lie in the orientation of successive CoO₂ slabs as illustrated in Fig. 1. In the D2 structure, all CoO₂ cations align along the c-axis, and each CoO₂ slab is rotated by π/3 relative to the preceding one. In contrast, the D3 structure exhibits uniform orientation for all CoO₂ slabs.²⁴ It is certainly of note that one of the potential applications for AgCoO₂ is as a catalyst for oxidation of CO.²⁵ In such applications, high-surface areas are required and thus the D3 powder samples are ideal. However, here we focus on potential electronic applications and it is the samples made from single-crystal precursors (we refer to these as crystal samples) that show tunability in terms of electronic transport properties. By contrast, the powder samples show no variation in the current study and serve as a reference to better understand the changes that are only seen in the crystal samples.

2. Experimental methods

Polycrystalline AgCoO₂ was obtained by ion-exchange. First, NaCoO₂ was synthesised by mixing Na₂O and Co₃O₄ using mortar and pestle for 30 min and pressing the reagents into a pellet. Pellets were then heated in an oxygen tube furnace for 1.5 h at 650 °C with a heating/cooling rate of 2 °C min⁻¹. To make AgCoO₂, we mixed together the NaCoO₂ precursor with AgNO₃ (50% in excess to promote the reaction) and KNO₃ using a mortar and pestle for 30 min.²⁶ The mixture was then put in a quartz tube that was sealed under vacuum. The tube was then heated in a box furnace for 5 days at a temperature we denote as the ion-exchange temperature (T_IE), which ranged from 150 to 300 °C. It is worth noting that ion-exchange cannot be performed below 150 °C as the potassium nitrate flux does not melt at these temperatures. After the synthesis, the product was washed with distilled water then dried in a vacuum oven at approximately 80 °C.

For the synthesis of the larger crystals, the same method was utilized but with NaCoO₂ single crystals as the precursors. Here, to make NaCoO₂ single crystals, we followed the method of Fujita et al.²⁷ We mixed Na₂CO₃ with Co₃O₄ and used NaCl as a flux in a Co₃O₄:Na₂Co₃:NaCl mass ratio of 1:7.5:7. The mixture was placed in an Al₂O₃ crucible (ID = 1.1 cm, OD = 1.3 cm) and covered with a MgO plate and then placed in the furnace to be heated to 950 °C at a rate of 10 °C min⁻¹ with a dwell time of 12 hours followed by a slow cooling at 0.5 °C min⁻¹ down to 850 °C. NaCoO₂ crystals were then obtained by washing away the NaCl flux with distilled water. The same ion-exchange process was done for the crystal synthesis, except the crystals were simply covered above and below with pre-mixed AgNO₃ and KNO₃ powder. Fig. 1 details the experimental procedures, illustrating the workflow for both powders and crystals highlighting their differences throughout synthesis and characterization. This figure shows the contrast in the two precursors, illustrates the setup for the ion-exchange steps, demonstrates the differences in the final morphology and also shows how the contacts were made for the Hall measurements. The powders were pressed into pellets and then indium contacts were added while for the crystals we sputtered contacts onto the crystal using a mask.

Powder X-ray diffraction (PXRD) was performed on the powder samples in a Panalytical Empyrean diffractometer equipped with a Mo anode (60 kV, 40 mA, 10 min scan, 4–30° scattering angles) and a GalaPIX area detector. The crystals were characterized using single-crystal X-ray diffraction (SCXRD) data on a Bruker D8 Venture diffractometer equipped with a Photon 200 CMOS area detector, and IμS microfocus X-ray source (Bruker AXS, CuKα source).

The microstructure and elemental mapping of selected samples were observed by scanning electron microscopy (SEM) using a Hitachi SU3500 microscope at 10 kV equipped with an Al Kα micro-focused monochromator and an energy dispersive spectroscopy (EDS) detector. The X-ray photoelectron spectroscopy (XPS) measurements were done at a pressure of mbar with an X-ray spot size of 200 μm. The survey and high-resolution spectra were collected with a pass energy of 200 and 50 eV, respectively.

We also performed thermogravimetric analysis (TGA) with the TA Discovery 5500 on polycrystalline and crystal samples under oxygen atmosphere. The samples were heated to either 190 °C or 590 °C at a rate of 1 °C min⁻¹, dwelling there for an hour and finally cooling back down to room temperature at a rate of 10 °C min⁻¹. Finally, to measure electronic properties of our samples we used the Hall measurement technique in the Van der Pauw geometry using an Ecopia HMS-3000 device that allowed us to do measurements at room temperature and at 77 K by immersing the sample in liquid nitrogen. For the crystal samples, the contacts were made by first sputtering 10 nm of Ti, then 50 nm of Au, and finally attaching wires using silver epoxy on the corners as shown in Fig. 1f.

Finally, we performed DFT calculations on 4 different structures. The D2 and D3 AgCoO₂ structures were obtained from the materials project (mp-19178 for D2 and mp-753001for D3).²⁸ The silver and oxygen deficient structures (D2-Ag_0.5CoO₂ and D2-AgCoO_1.5, respectively) were generated using the program supercell²⁹ using the parent D2-AgCoO₂ structure. A 2 × 1 × 1 D2-Ag_0.5CoO₂ supercell was generated with 14 atoms to represent the silver deficient delafossite. Similarly, a 2 × 1 × 1 D2-AgCoO_1.5 oxygen deficient supercell with 14 atoms was generated and shown in Fig. S7 (ESI†) for the oxygen deficient AgCoO_1.5 delafossite. The use of the supercell was required in both cases to keep the vacancy concentration low enough to obtain a stable structure. Density functional theory (DFT) was carried out using the vienna ab initio simulation package (VASP) version 5.4.1. The generalized gradient approximation (GGA) in Perdew–Burke–Ernzerhof formulation was used for the structure relaxation. The density of states (DOS) and band diagram plots were generated using the hybrid functional, HSE06 with the standard 0.25 fraction of Hatree–Fock exchange. Based on the cutoff energy and K-point convergence test, the calculations were done with a cutoff energy of 550 eV in a 4 × 8 × 2 Γ-centered mesh. The DOS was plotted using the vasplab package in MATLAB.³⁰ We also used VASPKIT³¹ code to perform the band structure calculations. The values of the valence band maximum, conduction band minimum and bandgap were calculated using the pymatgen python library.³²

3. Results and discussion

The synthesis of both polycrystalline and crystal forms of semiconducting Ag_xCoO_2−δ delafossite was done using the same ion-exchange method. First, we discuss the structures as determined using XRD. For the NaCoO₂ polycrystalline precursor, we obtain the expected pattern (ICDD 99-008-8818) for the structure indexed in the Rm space group with lattice parameters a = b = 2.887(3) Å and c = 15.608(1) Å. After ion-exchange, the material indexed to D3-AgCoO₂ as shown in Fig. 2c for all ion-exchange temperatures used. The lattice parameters extracted by fitting the pattern for T_IE = 300 °C are a = b = 2.8703(2) Å and c = 18.3276(1) Å. Fig. 2c also shows that a small amount of silver metal (Ag₀) is seen in all powder samples, consistent with the literature.²³ Fig. S1 (ESI†) shows the SEM/EDX analysis of the polycrystalline sample, revealing that the particle sizes remain unchanged regardless of ion-exchange temperature and with a composition that also remains steady as shown in Table S1 (ESI†). The SEM images were analyzed to extract average particle sizes included in Table S2 (ESI†) which again show no variations across the powder samples. The XRD patterns were also analyzed to extract crystallite sizes after correcting the peak widths for machine-broadening. The results in Table S2 (ESI†) show that the crystallites are smaller than the SEM particles which is to be expected with the 850 °C synthesis used (battery cathode research has shown that over 1000 °C is required to convert particles to single crystallites³³). Again, crystallite size remains roughly constant across all powders. Similarly, the percentage of delafossite D3 phase only varies between 94 and 96% in these samples. The higher, the temperature the more our sample is made up of the delafossite phase. Finally, the particle sizes also show consistency across all ion-exchange temperatures.

By contrast to the powder samples, SCXRD shows significant variation in the crystal samples as we change T_IE. Fig. 3a shows that the single-crystal synthesis of the NaCoO₂ precursor was successful: we obtained a structure that yields sharp peaks and they index to the P1c space group with lattice parameters a = b = 2.8349(2) Å and c = 10.8916(8) Å consistent with ref. 34, the structure being found in the Crystallography Open Database as COD153298.³⁵ SEM images in Fig. S2 (ESI†) also show that the precursor is made up of large crystals that range in size with the largest being more than 2 mm across. Despite the single-crystal precursor, the SCXRD patterns after ion-exchange at both 180 and 300 °C show arcs in Fig. 3b instead of the individual points associated with scattering from single-crystals. However, the arcs are short in length, showing that the samples are by no means powders, but are no longer single-crystals either. By integrating over these arcs, we produce the simulated PXRD patterns shown in Fig. 2b for these crystals (we will refer to them as crystals throughout this article to distinguish them from the powder samples, but we remind the reader to not confuse them with single-crystals). The results index to the D2-AgCoO₂ structure with varying lattice parameters shown in Fig. 3d–f with larger cell volumes as the ion-exchange temperature, T_IE, increases. This shows that there is variation in the bulk structure as T_IE changes. It is also important to recognize the massive increase in c lattice parameter occurring during ion-exchange (10.89 Å to 12.35 Å represents a 12% increase), while the a-axis does not increase by more than 1.5%. The consequences of this will be discussed extensively below.

Fig. 4 shows SEM images of the D2-AgCoO₂ crystals synthesized with various T_IE as well as one made without any silver present during the ion-exchange. The results as viewed from above the particles are very consistent with previous literature: the samples look like the single-crystal precursors when viewed from above. However, the profile views show varying extent of exfoliation with layers separating noticeably during the ion-exchange process. This phenomenon was particularly pronounced in crystals subjected to ion-exchange at 300 °C and 180 °C, as evidenced by the SEM images (a, b, d, and e). However, Fig. 4f shows that the exfoliated sheets are first formed during the departure of sodium ions from the crystal, and this process was then further intensified when silver ions were introduced into the structure in Fig. 4d and e. Although this phenomenon has not been reported previously in AgCoO₂, in hindsight, it is not surprising considering the 12% increase in c-axis lattice parameter taking place during ion-exchange such that fracturing occurs in that direction only.

The exfoliation does not, however, explain why the lattice parameters increase with the ion-exchange temperature. For this, we turn to EDX measurements on the crystals shown in Fig. 2a. Fig. S3 and Table S1 (ESI†) show EDX results for the crystal samples, which indicate that all samples contain very little sodium after the ion-exchange, and the data also shows high tunability of silver content in the crystals. For the polycrystalline samples, the Ag:Co ratio (x) is constant at about 0.8 and in fact shows silver deficiency consistent with ref. 18 and 19. By contrast, in the crystals, the silver content is stoichiometric (x = 1) for high T_IE (300 °C) and decreases dramatically all the way down to x = 0.42 at the lowest T_IE. Fig. S4 (ESI†) shows that with an ion-exchange temperature of 150 °C, a noisy XRD signal is obtained, and for no silver present we are unable to identify the phases present. Although EDX is a surface measurement, the fact that the bulk lattice parameters increase with T_IE strongly supports the conclusion that this non-stoichiometry is systemic and not a surface effect. Whereas we expected lower T_IE to lead to slower ion-exchange, it instead leads to a lower equilibrium Ag content. This provides us with a unique opportunity to study a solid solution by controlling the synthesis conditions rather than the stoichiometry of the precursors utilized. It is also, to our knowledge, the first ever study of Ag_xCoO₂ with varied x values.

To further understand the chemistry of these new Ag-deficient structures, we perform XPS to determine the oxidation states of Ag and Co. Fig. 5 and Fig. S5 (ESI†) show XPS patterns of crystals and powders, respectively, synthesized at both low and high ion-exchange temperatures. The figures show the XPS fits and Table 1 and Table S3 (ESI†) show the extracted values for both silver (peaks identified using ref. 36 and 37) and cobalt (peaks identified according to ref. 38 and 39) at low and high T_IE. For the powder samples, the average oxidation states for Ag and Co are 1.16 and 2.41, respectively, at 300 °C, giving an overall formula Ag_0.8CoO_1.67 at the surface of the particles (this is based on SEM/XPS only). Although we are cautious to not fully trust average oxidation states extracted from XPS, this result for the powder samples certainly shows that both metal and oxygen vacancies will need to be taken into account in order to understand the transport in these materials. When doing the synthesis at T_IE = 150 °C, the crystals are even more silver deficient than the powder samples, therefore we expect such vacancies to be compensated by higher oxidation of cobalt and silver and/or lower oxygen content. At the lowest silver content of x = 0.42, Table 1 shows an average oxidation state of 2.54 for cobalt and this drops to 2.40 in the sample with x = 1 (T_IE = 300 °C) showing oxidation of cobalt does help compensate for the silver deficiency. Similarly, we could see the same trend for silver where the average oxidation state goes from 1.69 to 1.10 at low and high temperature of ion-exchange, respectively, again showing that silver oxidation also helps compensate for silver vacancies and it actually compensates more than the cobalt does. The data reported here therefore demonstrates that the material at the surface (per XPS sensitivity being limited to the surface region only) is both silver and oxygen deficient and the silver content can be tuned in the crystals over a large range (0.4 to 1.0). Below, we discuss consequences of these surface off-stoichiometries in terms of transport properties measured using the Hall technique. Since such measurements are heavily impacted by the surface of the materials, we feel this focus on the surface composition is highly justified. We leave the depth of these non-stoichiometries (i.e. do they extend fully into the bulk?) as an open question worthy of further study, with neutron diffraction being a particularly viable probe of both bulk oxygen and silver.

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From the structural study performed here we can therefore conclude that both the silver and oxygen content can be highly tuned in delafossite Ag_xCoO_2−δ, and both silver/oxygen non-stoichiometries may therefore play a role in determining charge carrier transport properties. To explore the potential consequences of these compositional variations, DFT was performed on various structures (Fig. 6). Table S4 (ESI†) shows the lattice parameters from DFT and from XRD, while Table S5 (ESI†) shows the relatively small changes in lattice parameters taking place during relaxation. The lattice parameters are in reasonable agreement between experiment and calculation. Band structure and density of states plots were generated for key structures. First, D2 and D3 AgCoO₂ show similar features in term of band gap which is around 2.7 eV, as well as the positions of the conduction and valence bands having similar values in both structures. Thus, the vacancy-free AgCoO₂ structures show features of intrinsic semiconductors. DFT was also performed on silver deficient D2-Ag_0.5CoO₂, which has a similar Ag content to the low-Ag crystals we obtained by ion-exchange (0.42). In this calculation, the Fermi level is clearly within the valence band demonstrating p-doping. This is not unexpected, oxygen vacancies are well-known n-dopants, so it follows that metal cation vacancies serve as p-dopants. The bandgap of this system is fairly close to that of the original structure, so we view the Ag-deficient delafossite as a p-doped semiconductor where the Ag deficiencies do not dramatically distort the band structure. Finally, we performed DFT calculation on oxygen vacant samples as we concluded from EDX/XPS that the samples are oxygen deficient. The band structure is dramatically restructured compared to the stoichiometric AgCoO₂ band diagrams, such that this level of oxygen deficiency should be viewed as an altered structure altogether. Of highest significance, the Fermi level again crosses the valence band as shown in Fig. 6d such that we expect an electrically metallic material based on this band diagram – with low energy indirect band offsets near the Fermi level that may allow for the promotion of hole carriers as well. This shows that the oxygen vacancies are not acting as simple dopants, there are so many of them that it results in a re-structuring of the bands entirely. The contribution of each element to the band structures can be seen in Fig. S8 (ESI†) showing the projected band diagrams. In the silver deficient material (Fig. S8c, ESI†), both Ag and Co show valence band maxima above the Fermi level. In the case of the low oxygen material, the band rearrangement results in cobalt being heavily present at the top of the valence band which again crosses the Fermi level giving a metallic like band structure.

To further understand/confirm the significance of the DFT results and the structural results above, we performed Hall measurements in the van der Pauw configuration in order to determine the transport properties in these delafossites. Fig. 7 and Table 2 show Hall measurement results on both crystals and polycrystalline samples at room temperature and at liquid nitrogen temperature. We extracted bulk charge carrier concentration, mobility, and conductivity for all samples. Firstly, all measurements show p-type conduction in the Hall measurements, largely in agreement with the DFT calculations. The results for the 300 °C crystal (silver content of 1, expected oxygen deficiency based on XPS) show p-type conduction therefore indicates that very low oxygen content can allow for the existence of hole carriers in such a phase as predicted by the DFT calculations. There are also clear trends in the charge carrier concentrations as a function of T_IE in Fig. 7. The bulk concentrations of the polycrystalline samples show a near-stable trend (ranging from 5.68 × 10¹⁹ to 3.3 × 10¹⁹ cm⁻³) which is consistent with the EDX results that showed a constant silver content in the polycrystalline samples (the small variation seen is attributed to the small difference in the amount of silver metal, Ag₀, that tends to decrease with a higher T_IE as seen in Table S2, ESI†). We observed the same trend for polycrystalline samples when measured at liquid nitrogen temperatures but with an overall lower bulk concentration. At room temperature there are more charge carriers (holes) which indicates a certain amount of free carrier freeze-out occurring at 77 K. Similar trends are seen in conductivity: a slight decrease in conductivity in the polycrystalline samples can be attributed to the difference in silver metal content depending on T_IE, though part of the effect (particularly at 77 K) is due to a small change in mobility. At room temperature, though, mobility does not change significantly (0.10 to 0.15 cm² V⁻¹ s⁻¹). Overall, the polycrystalline samples show very little tunability in the transport properties, consistent with the low tunability of silver content seen in the structure results above.

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By contrast, the crystals show a marked increase in tunability of transport properties. Fig. 7 shows that charge carrier concentrations can be varied by about an order of magnitude with the highest concentrations being obtained for the highest T_IE which results in the highest silver content (x ≈ 1). At low ion-exchange temperatures, crystals showed a bulk concentration close to that of the polycrystalline samples (which are also both silver and oxygen deficient at the surface). However, when increasing the temperature of ion-exchange, the bulk concentration of carriers increased gradually up to 3.2 × 10²⁰ cm⁻³. We observed the same trend in the conductivity which increased from 25 to 79 S cm⁻¹ as T_IE increased. There is no strong trend in the mobility (μ) of the crystals where it seems to be generally constant at room temperature, and increases slightly with T_IE at 77 K. Importantly, this indicates that band transport and not hopping transport is responsible for the carrier motion.⁴⁰ The trends in bulk concentration are similar at RT and 77 K, though the values are lower at 77 K (ranging between 1.74 × 10¹⁹ to 2.94 × 10¹⁹ cm⁻³). The conductivity of the crystals kept the same trend with an overall increase going from 25.74 S cm⁻¹ to 197.4 S cm⁻¹ at 77 K. In addition, to understand further the ion-exchange process, we measured the electronic properties of a sample that was synthesized without silver nitrate. Table S1 (ESI†) shows that during this process 92.2% of the sodium left the crystal such that this material is nearly Na free and of course has no silver. As one would expect, this sample showed the lowest conductivity with 1.80 S cm⁻¹ at room temperature and 2.81 S cm⁻¹ at 77 K, and a relatively low bulk concentration compared to the other silver-containing crystal samples with only 9.3 × 10¹⁹ cm⁻³ at room temperature and 8 × 10¹⁷ cm⁻³ at 77 K. This shows consistency with the rest of the samples (Fig. 7) where lower silver content in the crystal yields lower conductivities. This, however, is counterintuitive given that the DFT calculations show that Ag vacancies should yield holes based on the band structures (see Fig. 6c). The fact that the opposite happens, as is also seen for the carrier concentration as a function of the T_IE in Fig. 7, suggests that the silver vacancies may interfere with the oxygen content and prevent/limit the band reorganization seen in the oxygen deficient DFT calculation (see Fig. 6d). In this sense, we speculate that the two non-stoichiometries compete to a certain extent thereby setting a bound for how far the transport can be improved by varying Ag & O content. However, the fact that Hall measurements indicate holes are the dominant carrier indicates that Ag deficiency likely dominates the properties observed. From Fig. 6d we can see that strong oxygen deficiency would result in a metallic electronic structure reconstruction that would not be as amenable to generating a p-type Hall voltage (as for example, by contrast, the Ag deficiency is in Fig. 6c). This suggests that the band rearrangement predicted from DFT for the oxygen deficient material did not fully occur in our samples.

In addition, we performed TGA experiments in an oxygen atmosphere for both polycrystalline and crystal samples, shown in Fig. S6 (ESI†), in order to determine to what extent the oxygen content can be tuned in the materials after synthesis and the resulting impact on transport properties. Fig. S6 (ESI†) shows the results for two TGA runs: the first up to 190 °C shows a mass gain (uptake of oxygen) while the second up to 590 °C shows a mass decrease (release of oxygen). For both samples, it was confirmed that the XRD patterns did not vary (this was not the case for samples heated above 590 °C). Fig. 7 shows that the polycrystalline sample heated up to 190 °C gave an increase in the bulk concentration (light blue circle in Fig. 7) up to 8.1 × 10¹⁹ cm⁻³ consistent with an increase in holes by increasing the oxygen content (oxygen vacancies are n-dopants so oxygen excess acts as p-dopants). The sample had a negligible change in mobility (from 0.10 to 0.13 cm² V⁻¹ s⁻¹) and an increase in conductivity going from 0.75 to 1.45 S cm⁻¹ consistent with the charge carrier concentration increasing. On the other hand, the polycrystalline sample heated up to 590 °C (orange circle in Fig. 7) showed a decrease in the number of charge carriers down to 6.5 × 10¹⁸ cm⁻³ which we attribute to a decrease in oxygen content as seen in the TGA mass drop. The two TGA heated polycrystalline samples are therefore consistent with the transport properties being tunable by varying the oxygen content such that the hole concentration can either be reduced with less oxygen or increased with more oxygen.

Having established the temperatures at which oxygen is taken up and released on the powders, we repeated the experiments on two crystals (T_IE = 180 and 300 °C). It should be noted that the crystals do not weigh enough to obtain measurable mass changes in the TGA. Fig. 7 shows that the sample with T_IE = 300 °C gave the expected decrease in charge carrier density after heating to 590 °C, the bulk concentration decreased down to 1.6 × 10²⁰ cm⁻³ with a conductivity staying relatively stable with a value of 209.53 S cm⁻¹. These results are fully consistent with the results for the powders: decreased oxygen content yields a decrease in charge carrier densities. For the T_IE = 180 °C sample, however, the bulk concentration increased slightly, going up to 2.16 × 10²⁰ cm⁻³. This is the second counterintuitive transport trend seen in the samples with both Ag and O deficiency. It should however be mentioned that the increase in concentration is small and the values are within uncertainty of each other at RT. Nonetheless, the non-intuitive trends seen in the Ag/O deficient materials warrant further study, as do the impact of the competing cationic and anionic off-stoichiometries in the same complex oxide. This sample also showed an increase in conductivity from 25.55 S cm⁻¹ for the as synthesized sample to 137.38 S cm⁻¹ for a sample synthesized at the same temperature but that was then heated in the TGA. These results therefore show that the crystals’ carrier transport properties can indeed be further optimized by varying the oxygen content following post synthesis treatments. It is also important to re-emphasize that the cation oxidation states are also playing a critical role in tuning the transport properties. In particular, the reduction of Ag at the higher ion-exchange temperatures (as determined by XPS) results in a greater number of holes available for transport (accounting for some of the increase in the hole carrier concentration seen in Fig. 7). There is thus a delicate balance between cationic oxidation state and oxygen content that occurs in the Ag-deficient delafossites.

4. Conclusion

In this study, we focused on the electronic transport properties of Ag_xCoO_2−δ delafossite materials. The main objective of this study was to enhance the transport properties of delafossite AgCoO₂ by tuning the composition in large single-crystals. To this end, we demonstrated for the first time that crystals can be made over a remarkably wide silver content range (x = 0.4 to 1.0) and that all samples also showed oxygen deficiency at the surface. By contrast, polycrystalline materials showed an approximately constant silver content (x = 0.8) and also showed oxygen deficiency. The Ag deficiency can therefore only be achieved in very large crystals. DFT results show that the Ag vacancies provide p-type doping (at high doping concentrations holes are accessible at the Fermi level) and that excessively low oxygen content results in metallic transport and a full band reconstruction. The subsequent Hall measurements confirm that all synthesized samples are p-type materials and the stoichiometry variations result in about an order of magnitude variation in charge carrier densities in the crystals. TGA followed by Hall measurements indicate that further tuning of the transport properties is possible by careful manipulation of the oxygen content in the materials. The crystal made by ion-exchange at 300 °C gave the best transport properties with 3.20 × 10²⁰ cm⁻³ hole concentrations, a mobility of 1.51 cm² V⁻¹ s⁻¹ and a conductivity of 79 S cm⁻¹. Beyond making materials with promising transport properties, this study also demonstrated for the first time that the crystal materials made by ion-exchange on single-crystal precursors do not in fact yield single-crystals, but rather the products are semi-exfoliated layered materials due to the large amount of stress resulting from Ag insertion between the CoO₂ slabs. Thus, further optimizing this exfoliation process might provide 2D materials with novel properties – a topic for future study. Further exploration into the roles of both the exfoliation and the Ag/O non-stoichiometries is clearly warranted given the high degree of control both have on the transport properties.

Data availability

Some of the data supporting this artcle have been included in the ESI† section. The rest is available from the corresponding author upon reasonable request.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was funded by an FRQNT Teams grant and by the Natural Sciences and Engineering Research Council of Canada under the auspices of two Discovery grants. Thank you to QCAM/CQMF for supporting this project. Thanks to Petr Fiurasek for his help in TGA measurements, and Bobby Miglani for her assistance with DFT calculations. We also acknowledge both the Quebec Centre for Advanced Materials and the Centre for Green Chemistry and Catalysis, both funded by the FRQNT, who provided support in materials' characterization.

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† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4tc02403f

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