Magnesium Doping for the Promotion of Rutile Phase Formation in the Pulsed Laser Deposition of TiO₂ Thin Films

Abstract

The preparation of a transparent and smooth rutile-type TiO₂ thin film without the use of the crystallographical effect of the substrate is a challenge for the advanced utilization of TiO₂ in the fields of optics and solid state ionics. Because acceptor doping leads to the formation of oxygen vacancies, this method has promise as a new approach to promote the formation of rutile-type TiO₂. Mg²⁺-doped TiO₂ thin films were prepared by pulsed laser deposition, and the effects of Mg²⁺ doping on the phases present, the microstructure, the optical properties, and the surface roughness of the films were investigated. Particular attention was paid to the Mg²⁺ distribution in the prepared films. The formation of the rutile phase was promoted by 2.7 mol% and 5.5 mol%Mg²⁺ doping. The negligible segregation of Mg²⁺ and absence of change in the extinction coefficient by Mg²⁺ doping indicate that Mg²⁺ worked as the acceptor and induced oxygen vacancies for charge compensation, which promoted the formation of the rutile phase. Given that Mg²⁺ is a doubly charged acceptor, Mg²⁺ doping is a more effective method for promoting the formation of the rutile phase than trivalence doping. Besides the excellent optical properties (n ≈ 3.03 and k < 0.02 at λ = 400 nm) of the 2.7%Mg²⁺-doped rutile-type TiO₂ thin film deposited at 350℃, the films were smooth, with a roughness index of only approximately 0.8 nm. This method of preparing smooth rutile-type TiO₂ thin films has potential for the further development of TiO₂-based resistive memory devices.

1. Introduction

TiO₂ has long been studied since it can be used in a wide variety of applications: in white pigments¹⁾, high-refractive-index coatings²⁾ and resistive switches^3,4), in the anodes of Li-ion batteries⁵⁾, in the electrolytes in the proton exchange membrane fuel cell⁶⁾ and in photocatalytic processes⁷⁾. The functionality of TiO₂ improves when its phases are controlled. A typical example is the higher photocatalytic activity of metastable anatase-type TiO₂^8,9). The anode performance of this particular form of TiO₂ in Li-ion batteries has been shown to be superior to that of a thermodynamically stable rutile phase^5,10). While the determination of the optimum phase should be done for all applications of TiO₂, particularly for resistive memory elements or surface proton conductors, little attention has been paid to rutile-type TiO₂. This is presumably due to a combination of the necessity of thin TiO₂ films for reducing electrical resistance and the rarity of rutile-type TiO₂ crystals in TiO₂ thin films^11–14). In order to investigate the optimum phase in applications to become standard practice, and for the rutile phase to be considered, an evolution in the preparation technique of rutile-type TiO₂ thin films is essential.

In contrast to the field of the solid state ionics, rutile-type TiO₂ has been the focus of significant attention in the field of optics due to its transparency and a higher refractive index (n) than that of the anatase phase (n_Rutile ≈ 2.75, n_Anatase ≈ 2.54 at λ = 550 nm¹⁵⁾). This higher-n of transparent rutile-type TiO₂ is expected to improve the controllability of the reflection spectrum for optical coatings¹⁶⁾. In the optics field, therefore, there is a clear demand for better methods for the preparation of rutile-type TiO₂ thin films. As mentioned above, TiO₂ thin films are generally employ anatase-type TiO₂. This is largely because the rutile phase has higher surface energy than the anatase phase¹⁷⁾. Three different preparation techniques have been reported for preparing rutile-type TiO₂ thin films: 1) deposition at high temperature around 700℃^11,18), 2) deposition on single crystal substrates^19,20), and 3) the introduction of oxygen vacancies, which promotes the formation of the rutile phase^21,22) based on the following equation^23–25):   

\[ {\rm O}_{\rm O}^{\rm X} \to {\rm V}^{\bullet\bullet}_{\rm O} + 2{\rm e}^{-} + \frac{1}{2}{\rm O}_{2} \](1)

However, each of these three techniques is associated with serious flaws. The high temperature required for the first technique gives a rough surface, which results in significant optical loss^11,16), the cost of single crystal substrates is impractically high, and the introduction of the oxygen vacancies through eq. (1) results in a decrease in the transparency of the TiO₂ thin film due to the excess of conduction electrons^11,26). Preparation of transparent rutile-type TiO₂ thin films at low temperature without the use of the crystallographical information of the substrate has proved to be challenging.

In a new approach to the low-temperature preparation of transparent rutile-type TiO₂ thin films, acceptor doping in the TiO₂ thin films has been investigated by our group. This approach was adopted because the acceptor M^x+ (1 ≤ x ≤ 3) can lead to the formation of oxygen vacancies without the introduction of excess electrons according to the following²⁷⁾:   

\[ {\rm M_{2}O}_{3} \to {\rm 2M}^{'}_{\rm Ti} + {\rm V}^{\bullet\bullet}_{\rm O} + {\rm 3O}_{\rm O} \ ({\rm when}\ x = 3\ {\rm in\ M}^{x + }) \](2)

It has been reported that doping with 10 mol% Al³⁺, a single-phase rutile-type TiO₂ thin film with a high n ≈ 3.05 (at λ = 400 nm) and a low extinction coefficient (k, ≈0.01) was successfully prepared on glass by pulsed laser deposition at 350℃. This single-phase rutile-type TiO₂ thin film showed negligible surface roughness (≈0.8 nm) and optical loss was insignificant¹⁶⁾. This suggests that acceptor doping in the TiO₂ thin films makes it possible to produce optical-coating-ready high-n rutile-type TiO₂ thin film at low temperature without the use of the crystallographical effect of the substrate. In the case of rutile-type TiO₂ thin film prepared by the 10 mol%Al³⁺ doping, this trivalent substitution would have resulted in a decline in the n value due to the decrease in the Ti⁴⁺ content ($n_{\rm Al_{2}O_{3}} \ll n_{\rm TiO_{2}}$)¹⁶⁾. By decreasing the concentration of the doping element, it is expected that the decline in the n value will be minimized.

In this investigation, the focus is on Mg²⁺ doping as an alternative to Al³⁺ doping in the production of rutile-type TiO₂ thin films. Because ${\rm Mg}^{''}_{\rm Ti}$ is a doubly charged acceptor, the appropriate dopant concentration of Mg²⁺ for the preparation of the rutile-type TiO₂ thin film can likely be reduced by half that required for Al³⁺ doping. Even though the TiO₂–MgO phase diagram indicates the solubility of Mg²⁺ into TiO₂ is very limited²⁸⁾, solubility can be maximized using the pulsed laser deposition (PLD) method for the preparation of TiO₂ thin films containing Mg²⁺ as the acceptor. The PLD method involves a rapid cooling process, and therefore the thermal equilibration would be insufficient for the segregation of Mg²⁺ to take place. This study of the TiO₂ thin film prepared by PLD investigates the effects of Mg²⁺ doping on the phases, the microstructure, and in particular of the distribution of Mg²⁺. The optical properties and the surface roughness were also investigated. The phase control technique for preparing TiO₂ thin films revealed by this study opens a new direction in the development of rutile-type TiO₂-based applications in the field of solid state ionics.

2. Experimental Procedure

Undoped, 2.7 and 5.5 mol%Mg²⁺-doped TiO₂ targets were prepared by a solid-state reaction of rutile-type TiO₂ powder (Kojundo Chemical Laboratory Co., Ltd: purity 99.9%) and MgO powder (Kojundo Chemical Laboratory Co., Ltd: purity 99.9%). The sintering condition for the undoped-TiO₂ target was at 1300℃ for 10 h in air, and those for the Mg²⁺-doped TiO₂ targets were at 1500℃ for 36 h in air. The PLD process was conducted using a KrF excimer laser (COMPexPro205, COHERENT, Inc: λ = 248 nm) and a highly vacuumed deposition chamber (≈10⁻⁵ Pa) (PLAD-242, AOV Co., Ltd). The laser was emitted at a frequency of 5 Hz at a fluence of 5.8 J·cm⁻² for 30 min. The substrate was an alkali-free glass (OA-10G, Nippon Electric Glass Co., Ltd: 15 mm × 15 mm × 0.7 mm) cleaned by using a sonicator with acetone. The deposition temperature was controlled by heating the substrates with an infrared lamp heater, and a Si plate was mounted on the backside of the substrates to transmit the heat. The distance between the substrate and the targets was kept at 50 mm in all cases. The deposition atmosphere was controlled using oxygen gas (purity 6N) at 0.5 Pa, which has been shown to be suitable for the preparation of the transparent rutile-type TiO₂ thin films¹⁶⁾. The thickness of the TiO₂-based thin films was controlled at approximately 100 to 150 nm under these conditions.

The phases present in the TiO₂-based targets was clarified by θ–2θ X-ray diffraction method (XRD, D8 Advance, Bruker AXS) and that for the thin films was clarified by grazing-incidence (α = 2°) XRD. To examine the phases present of the TiO₂-based thin films in detail, in-situ high-temperature XRD (HT-XRD) was also carried out in air. The X-ray source was Cu-Kα radiation (λ = 1.5418 Å). Micro-Raman spectroscopy (HR-800, Horiba Jobin Yvon S.A.S.) was also conducted to support the phase identification. The Raman spectra were generated using a He-Ne laser (λ = 632.8 nm). The microstructure was observed by field emission scanning electron microscope (FE-SEM, JSM-7800F, JEOL Ltd) and transmission electron microscope (TEM)–aberration-corrected annular dark-field scanning TEM (ADF-STEM) combined apparatus (JEM-ARM200F, JEOL Ltd) with an energy dispersive spectrometer (EDS). The surface morphology was observed by atomic force microscope (AFM, JSPM- 5200, JEOL Ltd).

The optical properties were clarified by spectroscopic ellipsometry (M-2000, J. A. Woollam Co., Ltd). This measurement included obtaining an ellipsometric parameter (Φ, $\Delta$) with incident angles from 50° to 70° and obtaining transmittance spectrum for a wavelength range of 250–1000 nm. Using these two data and by taking the back-side reflection of the substrate into account, the wavelength dispersion of n and k values was determined. In this study, the values of n and k at λ = 400 nm are discussed as typical values to ensure the transparency of the film.

3. Results and Discussions

To clarify the effect of Mg²⁺ doping on the phases present in the TiO₂ thin films, the Mg²⁺ concentration dependence on the XRD patterns was investigated. Figure 1 shows the XRD patterns of the undoped, 2.7 mol% and 5.5 mol% Mg²⁺-doped TiO₂ targets (Fig. 1(a)) and thin films deposited at 350℃ (Fig. 1(b)). In Fig. 1(a), XRD peaks attributed to the rutile-type TiO₂ were observed in the undoped TiO₂ target. In addition, MgTi₂O₅ was observed in the Mg²⁺-doped TiO₂ target. These phases present of the targets were consistent with the TiO₂-MgO phase diagram²⁸⁾. Meanwhile, as shown in Fig. 1(b), the XRD patterns of the Mg²⁺-doped TiO₂ thin films showed the peaks only from the rutile phase. It should be also noted that a single anatase phase was clearly confirmed in the undoped-TiO₂ thin film. This suggests that the dissolution of Mg²⁺ into TiO₂ matrix occurs when using the PLD process and this dissolution promotes the formation of the rutile phase, as is the case with Al³⁺ doping¹⁶⁾. For the XRD peaks of the Mg²⁺-doped TiO₂ thin films, i.e. (110) at 27.4° and (200) at 39.2°, no clear shift is observed in comparison with those of the undoped rutile-type TiO₂ thin film¹¹⁾. This is attributed to the low Mg²⁺ concentration and the similarity in the ionic radii between Ti⁴⁺ and Mg²⁺ ($r_{{\rm Ti}^{4+}} = 68\,{\rm pm}$ and $r_{{\rm Mg}^{2+}} = 65\,{\rm pm}$²⁹⁾).

Fig. 1

XRD patterns of the undoped, 2.7 mol% and 5.5 mol% Mg²⁺-doped TiO₂ (a) targets and (b) thin films deposited at 350℃.

To confirm the formation of the rutile phase in the Mg²⁺-doped TiO₂ thin films, HT-XRD and Raman analyses were carried out. Figures 2(a) and (b) show the HT-XRD patterns of the 2.7 mol% and 5.5 mol% Mg²⁺-doped TiO₂ thin films, respectively. The growth of the rutile phase can be recognized at around 550℃. Meanwhile, Ishii et al. have reported that the rutile phase does not emerge up to 600℃ in the case of amorphous and the anatase-type transparent TiO₂ thin films¹¹⁾. Figure 2(c) shows the Raman spectra of the undoped, 2.7 mol% and 5.5 mol% Mg²⁺-doped TiO₂ thin films. For the undoped one, a sharp peak at 144 cm⁻¹, which is attributed to an E_g mode of the anatase crystal³⁰⁾, is observed with weak peaks attributed to a brookite phase. Meanwhile, for the Mg²⁺-doped TiO₂ thin films, two peaks at 445 cm⁻¹ and 608 cm⁻¹ were emerged, which are attributed to the E_g and A_1g modes of the rutile crystal³¹⁾, respectively. This result also supports the formation of single rutile phase in the Mg²⁺-doped TiO₂ thin films.

Fig. 2

HT-XRD patterns of (a) 2.7 mol% and (b) 5.5 mol% Mg²⁺-doped TiO₂ thin films deposited at 350℃. (c) Raman spectra of the undoped, 2.7 mol% and 5.5 mol% Mg²⁺-doped TiO₂ thin films.

To clarify the phases present and to determine the exact mechanism of the formation of the rutile phase in the Mg²⁺-doped TiO₂ thin films, microstructural observation using an electron microscope was carried out. Figure 3 shows the microstructure of the 2.7 mol%Mg²⁺-doped TiO₂ thin film deposited at 350℃. While the weak contrast in the in-plane ADF-STEM image in Fig. 3 (a) may be attributed to differences in composition, no apparent segregation of the Mg element was observed in the EDS analysis, as can be seen in Fig. 3 (b). This indicates that most of the Mg²⁺ was incorporated into the TiO₂ matrix as the acceptor dopant, which then led to the formation of the rutile phase. Though a phase separation was observed in the case of the 10 mol%Al³⁺ doping¹⁶⁾, almost all of the Mg²⁺ apparently dissolved presumably because the dopant concentration was smaller in the case of Mg²⁺ doping, and the ionic radius of Mg²⁺ is closer to that of Ti⁴⁺ $(r_{{\rm Mg}^{2+}} = 65\,{\rm pm}$, $r_{{\rm Ti}^{4+}} = 68\,{\rm pm}$, $r_{{\rm Al}^{3+}} = 50\,{\rm pm}$²⁹⁾). In Fig. 3 (c), it can be seen that the crystal grains of the Mg²⁺-doped TiO₂ thin film are finer than those of undoped-TiO₂ thin films at ranges of 30–140 nm and 70–200 nm, respectively. This is most likely due to the pinning effect at the grain boundaries by the nanoparticles indicated by arrows in Fig. 3 (d). These nanoparticles are possibly composed of the Mg-rich phase (e.g. MgO, MgTi₂O₅); however, no reliable compositional difference between the nanoparticles and the matrix was observed. It appears that most of the Mg²⁺ was incorporated into the TiO₂ matrix as the acceptor dopant, resulting in the promotion of the rutile-phase formation in the TiO₂ thin film.

Fig. 3

Microstructure of the 2.7 mol% Mg²⁺-doped TiO₂ thin films deposited at 350℃ observed by (a) ADF-STEM and (b) its EDS mapping, (c) FE-SEM compared with the undoped-TiO₂ thin films deposited at 350℃, and (d) TEM.

To further discuss the origin and role of oxygen vacancy, the extinction coefficient of the 2.7 mol% Mg²⁺-doped TiO₂ thin film deposited under P(O₂) of 0.5 Pa was measured. Figure 4(a) shows the extinction coefficient of the 2.7 mol% Mg²⁺-doped TiO₂ thin film as deposited and after annealing at 600℃ in air. The as-deposited 2.7 mol% Mg²⁺-doped TiO₂ thin film shows a small extinction coefficient due to oxygen vacancies compensated by electrons, while the film after annealing shows a negligibly small extinction coefficient in the same wavelength range. This indicates that 1) excess oxygen vacancies are introduced through eq. (1) during deposition under P(O₂) of 0.5 Pa, and 2) the excess oxygen vacancies can be oxidized and removed by annealing under air. The excess oxygen vacancies, however, do not seem to affect the rutile phase formation. Figure 4(b) shows XRD patterns of undoped TiO₂ thin films deposited at 400℃ under various oxygen partial pressure reported by Ishii et al.¹⁶⁾ Without the acceptor dopants, the anatase phase was formed at P(O₂) of 0.5 Pa. In addition, given that 1) most Mg²⁺ is dissolved into the TiO₂ matrix (Fig. 3), 2) no clear XRD peak shift suggesting interstitial-type defects is observed (Fig. 1(b)), and 3) Mg²⁺ is not compensated by electron hole (as shown later in Fig. 5), the defect equilibrium of $[{\rm Mg}^{''}_{\rm Ti}] = [V^{\bullet\bullet}_{\rm O}]$ is most likely as expected, even though further study is required for clarifying the defect chemistry in detail.

Fig. 4

(a) Extinction coefficient of 2.7 mol% Mg²⁺-doped TiO₂ thin film deposited at 350℃ before/after the annealing up to 600℃. The inset is enlarged view. (b) XRD patterns of undoped-TiO₂ thin films deposited at 400℃ under vacuum and 0.5–7 Pa of oxygen partial pressure. Adapted from Ref. 16) with permission.

Fig. 5

Optical properties (at λ = 400 nm) of the undoped, 2.7 mol% Mg²⁺-doped and 10 mol% Al³⁺-doped TiO₂ thin films¹⁶⁾ as a function of temperature.

The refractive index of the TiO₂ thin films was expected to rise in TiO₂ thin films with a rutile phase formed by Mg²⁺ doping. Figure 5 shows the optical properties of the undoped, 2.7 mol% Mg²⁺-doped and previously reported 10 mol% Al³⁺-doped TiO₂ thin films¹⁶⁾ as a function of temperature. The 2.7 mol% Mg²⁺-doped TiO₂ thin films showed higher n values than the undoped-TiO₂ thin films, and at 350℃, its high n value of 3.03 is comparable to that of the 10 mol% Al³⁺-doped TiO₂ thin films. This is because the formation of the rutile phase was promoted by Mg²⁺ doping in the same manner as that which occurs with Al³⁺ doping. The k values of the Mg²⁺-doped samples are close to zero (less than 0.02) and independent of temperature, which indicates that the charge compensation of ${\rm Mg}^{''}_{\rm Ti}$ is not carried out by ${\rm h}^{\bullet}$ but by ${\rm V}^{\bullet\bullet}_{\rm O}$. Note that at 300℃, the Mg²⁺-doped sample has a higher n value than the Al³⁺-doped sample. Given that the phases in both samples are identical and the difference of atomic polarizability between Mg²⁺ and Al³⁺ is quite small ($n_{\rm MgO} \approx 1.76$³²⁾, $n_{{\rm Al_{2}O}_{3}} \approx 1.79$³³⁾ at $\lambda = 400\,{\rm nm}$), the higher n value of the Mg²⁺-doped sample can likely be attributed to the smaller dopant concentration. This indicates that rutile-type TiO₂ thin films prepared by Mg²⁺ doping can be expected to superior to those prepared by Al³⁺ doping in that they will have a higher n value.

The surface of the rutile-type TiO₂ thin film should be smooth to avoid diffuse reflection. Figure 6 shows an AFM image of the 2.7 mol% Mg²⁺-doped TiO₂ thin films deposited at 350℃. The morphology is smooth, which implies homogeneous formation of the rutile phase¹⁶⁾. The arithmetic mean roughness calculated using Fig. 6 is approximately 0.8 nm. This roughness value is much lower than that of the rutile-type TiO₂ thin film prepared without acceptor doping (≈7 nm¹¹⁾), and is the same as 10 mol% Al³⁺-doped rutile-type TiO₂ thin films which suffer negligible optical loss due to diffuse reflection¹⁶⁾.

Fig. 6

AFM image of the 2.7 mol% Mg²⁺-doped TiO₂ thin films deposited at 350℃.

Based on these results, it can be said that Mg²⁺ doping has a similar effect to that of Al³⁺ doping on TiO₂ thin films prepared by PLD. The formation of the rutile phase is promoted, a higher n value is achieved even at 350℃, no change in the k value occurs, and the surface is smooth. The 2.7%Mg²⁺-doped rutile-type TiO₂ thin film deposited at 350℃ showed not only excellent optical properties (n ≈ 3.03 and k < 0.02 at λ = 400 nm) but also a smooth surface at approximately 0.8 nm, which is comparable to that of the 10 mol%Al³⁺-doped rutile-type TiO₂ thin film deposited at 350℃¹⁶⁾. The advantage of Mg²⁺ doping over Al³⁺ doping is its effectiveness: the doubly charged acceptor means that the required concentration of the dopant Mg²⁺ is half that of the dopant Al³⁺, which minimizes the reduction in the refractive index.

These results suggest that smooth rutile-type TiO₂ thin films prepared by acceptor doping will contribute to the development of TiO₂-based resistive memory devices. It is well-known that the resistive switching of TiO₂ is based on a reversible phase transformation between amorphous-/anatase- type TiO₂ and the magnéli phase Ti_nO_2n−1^34–36), and that the broad dispersion between the resistive switching parameters is problematic, particularly with regard to the set voltage for each set/reset cycles. However, if rutile-type TiO₂ thin films are applied into the resistive switching element, a significant reduction in dispersion can be expected since the rutile phase is more thermodynamically stable than the anatase phase and is structurally similar to the magnéli phase³⁷⁾. Furthermore, as reported by Liu et al., the accepter in the TiO₂ modifies the distribution of the oxygen vacancies and enhances the formation and rupture effect of the conducting filaments, resulting in uniform set voltage for each of the set/reset cycles³⁸⁾. The resistance switching behavior of acceptor-doped rutile-type TiO₂ thin films is a topic for further investigation.

4. Conclusion

The effects of Mg²⁺ doping in the preparation of TiO₂ thin films were investigated. The phases present, and the microstructure, with a focus on the Mg²⁺ distribution were investigated, and the optical properties and surface roughness of these films were also determined. The formation of the rutile phase was promoted by the 2.7 mol% and 5.5 mol%Mg²⁺ doping. Negligible segregation of Mg²⁺ occurred and no notable change in the extinction coefficient due to Mg²⁺ doping was noted, suggesting that Mg²⁺ worked as the acceptor and induced oxygen vacancies as charge compensation, which resulted in the promotion of the rutile phase formation. The 2.7 mol%Mg²⁺-doped rutile-type TiO₂ thin film deposited at 350℃ was shown to meet all the critical requirements: high n ≈ 3.03, low k < 0.02 (at λ = 400 nm) and smooth roughness ≈0.8 nm. These values are comparable to those of the optical-coating-ready 10 mol%Al³⁺-doped rutile-type TiO₂ thin film¹⁶⁾. Mg²⁺ doping more effectively promotes the formation of the rutile phase than Al³⁺ doping since Mg²⁺ is a doubly charged acceptor. This acceptor-doped rutile-type TiO₂ thin film is expected to contribute to developments in superior optical coatings and resistive memory.

Acknowledgments

We would like to thank Dr. K. Kobayashi for taking the SEM, TEM and STEM images. AI would appreciate the financial support from the Grant-in-Aid for JSPS Research Fellow and the Interdepartmental Doctoral Degree Program for Multi-dimensional Materials Science Leaders in Tohoku University. HT would also like to acknowledge the financial support provided by JSPS (26249103).

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